Selected Systems from Co-Fe-Si to Cu-Fe-Pt Landolt-Bornstein

Landolt-Börnstein Numerical Data and Functional Relationships in Science and Technology New Series / Editor in Chief: W. Martienssen Group IV: Physical Chemistry Volume 11

Ternary Alloy Systems Phase Diagrams, Crystallographic and Thermodynamic Data critically evaluated by MSIT® Subvolume D Iron Systems Part 3 Selected Systems from Co-Fe-Si to Cu-Fe-Pt Editors G. Effenberg and S. Ilyenko Authors Materials Science and International Team, MSIT®

ISSN 1615-2018 (Physical Chemistry) ISBN 978-3-540-74197-8 Springer Berlin Heidelberg New York Library of Congress Cataloging in Publication Data Zahlenwerte und Funktionen aus Naturwissenschaften und Technik, Neue Serie Editor in Chief: W. Martienssen Vol. IV/11D3: Editors: G. Effenberg, S. Ilyenko At head of title: Landolt-Börnstein. Added t.p.: Numerical data and functional relationships in science and technology. Tables chiefly in English. Intended to supersede the Physikalisch-chemische Tabellen by H. Landolt and R. Börnstein of which the 6th ed. began publication in 1950 under title: Zahlenwerte und Funktionen aus Physik, Chemie, Astronomie, Geophysik und Technik. Vols. published after v. 1 of group I have imprint: Berlin, New York, Springer-Verlag Includes bibliographies. 1. Physics--Tables. 2. Chemistry--Tables. 3. Engineering--Tables. I. Börnstein, R. (Richard), 1852-1913. II. Landolt, H. (Hans), 1831-1910. III. Physikalisch-chemische Tabellen. IV. Title: Numerical data and functional relationships in science and technology. QC61.23 502'.12 62-53136 This work is subject to copyright. All rights are reserved, whether the whole or part of the material is concerned, specifically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilm or in other ways, and storage in data banks. Duplication of this publication or parts thereof is permitted only under the provisions of the German Copyright Law of September 9, 1965, in its current version, and permission for use must always be obtained from Springer-Verlag. Violations are liable for prosecution act under German Copyright Law. Springer is a part of Springer Science+Business Media springeronline.com © Springer-Verlag Berlin Heidelberg 2008 Printed in Germany The use of general descriptive names, registered names, trademarks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. Product Liability: The data and other information in this handbook have been carefully extracted and evaluated by experts from the original literature. Furthermore, they have been checked for correctness by authors and the editorial staff before printing. Nevertheless, the publisher can give no guarantee for the correctness of the data and information provided. In any individual case of application, the respective user must check the correctness by consulting other relevant sources of information. Cover layout: Erich Kirchner, Heidelberg Typesetting: Materials Science International Services GmbH, Stuttgart Printing and Binding: AZ Druck, Kempten/Allgäu SPIN: 1185 9086 63/3020 - 5 4 3 2 1 0 – Printed on acid-free paper

Editors: Günter Effenberg Svitlana Ilyenko Associate Editor: Oleksandr Dovbenko MSI, Materials Science International Services GmbH Postfach 800749, D-70507, Stuttgart, Germany http://www.matport.com Authors: Materials Science International Team, MSIT®

The present series of books results from collaborative evaluation programs performed by MSI and authored by MSIT®. In this program data and knowledge are contributed by many individuals and accumulated over almost twenty years, now. The content of this volume is a subset of the ongoing MSIT® Evaluation Programs. Authors of this volume are:

Nataliya Bochvar, Moscow, Russia

Anatoliy Bondar, Kyiv, Ukraine

Lesley Cornish, Randburg, South Africa

Simona Delsante, Genova, Italy

Tatyana Dobatkina, Moscow, Russia

Gautam Ghosh, Evanston, USA

Joachim Gröbner, Clausthal-Zellerfeld, Germany

K.C. Hari Kumar, Chennai, India

Volodymyr Ivanchenko, Kyiv, Ukraine

Kostyantyn Korniyenko, Kyiv, Ukraine

Artem Kozlov, Clausthal-Zellerfeld, Germany

Viktor Kuznetsov, Moscow, Russia

Nathalie Lebrun, Lille, France

Pierre Perrot, Lille, France

Tatiana Pryadko, Kyiv, Ukraine

Rainer Schmid-Fetzer, Clausthal-Zellerfeld, Germany

Elena Semenova, Kyiv, Ukraine

Elena Sheftel, Moscow, Russia

Nuri Solak, Stuttgart, Germany

Jean-Claude Tedenac, Montpellier, France

Vasyl Tomashik, Kyiv, Ukraine

Michail Turchanin, Kramatorsk, Ukraine

Tamara Velikanova, Kyiv, Ukraine

Tatyana Velikanova, Kyiv, Ukraine

Andy Watson, Leeds, U.K.

Institutions The content of this volume is produced by MSI, Materials Science International Services GmbH and the international team of materials scientists, MSIT®. Contributions to this volume have been made from the following institutions: The Baikov Institute of Metallurgy, Academy of Sciences, Moscow, Russia Donbass State Mechanical Engineering Academy, Kramatorsk, Ukraine I.M. Frantsevich Institute for Problems of Materials Science, National Academy of Sciences, Kyiv, Ukraine Indian Institute of Technology Madras, Department of Metallurgical Engineering, Chennai, India Institute for Semiconductor Physics, National Academy of Sciences, Kyiv, Ukraine G.V. Kurdyumov Institute for Metal Physics, National Academy of Sciences, Kyiv, Ukraine Max-Planck-Institut für Metallforschung, Institut für Werkstoffwissenschaft, Pulvermetallurgisches Laboratorium, Stuttgart, Germany Moscow State University, Department of General Chemistry, Moscow, Russia

School of Chemical and Metallurgical Engineering, The University of the Witwatersrand, DST/NRF Centre of Excellence for Strong Material, South Afrika Northwestern University, Department of Materials Science and Engineering, Evanston, USA Technische Universität Clausthal, Metallurgisches Zentrum, Clausthal-Zellerfeld, Germany Universite de Montpellier II, Laboratoire de Physico-chimie de la Matiere, Montpellier, France Universita di Genova, Dipartimento di Chimica, Genova, Italy Universite de Lille I, Laboratoire de Metallurgie Physique, Villeneuve d’ASCQ, France University of Leeds, Department of Materials, School of Process, Environmental and Materials Engineering, Leeds, UK

Preface The sub-series Ternary Alloy Systems of the Landolt-Börnstein New Series provides reliable and

comprehensive descriptions of the materials constitution, based on critical intellectual evaluations of all data available at the time and it critically weights the different findings, also with respect to their compatibility with today’s edge binary phase diagrams. Selected are ternary systems of importance to alloy development and systems which gained in the recent years otherwise scientific interest. In one ternary materials system, however, one may find alloys for various applications, depending on the chosen composition.

Reliable phase diagrams provide scientists and engineers with basic information of eminent importance for fundamental research and for the development and optimization of materials. So collections of such diagrams are extremely useful, if the data on which they are based have been subjected to critical evaluation, like in these volumes. Critical evaluation means: there where contradictory information is published data and conclusions are being analyzed, broken down to the firm facts and re-interpreted in the light of all present knowledge. Depending on the information available this can be a very difficult task to achieve. Critical evaluations establish descriptions of reliably known phase configurations and related data.

The evaluations are performed by MSIT®, Materials Science International Team, a group of scientists working together since 1984. Within this team skilled expertise is available for a broad range of methods, materials and applications. This joint competence is employed in the critical evaluation of the often conflicting literature data. Particularly helpful in this are targeted thermodynamic and atomistic calculations for individual equilibria, driving forces or complete phase diagram sections.

Conclusions on phase equilibria may be drawn from direct observations e.g. by microscope, from monitoring caloric or thermal effects or measuring properties such as electric resistivity, electro-magnetic or mechanical properties. Other examples of useful methods in materials chemistry are mass-spectrometry, thermo-gravimetry, measurement of electro-motive forces, X-ray and microprobe analyses. In each published case the applicability of the chosen method has to be validated, the way of actually performing the experiment or computer modeling has to be validated as well and the interpretation of the results with regard to the material’s chemistry has to be verified. Therefore insight in materials constitution and phase reactions is gained from many distinctly different types of experiments, calculation and observations. Intellectual evaluations which interpret all data simultaneously reveal the chemistry of the materials system best.

An additional degree of complexity is introduced by the material itself, as the state of the material under test depends heavily on its history, in particular on the way of homogenization, thermal and mechanical treatments. All this is taken into account in an MSIT® expert evaluation.

To include binary data in the ternary evaluation is mandatory. Each of the three-dimensional ternary phase diagrams has edge binary systems as boundary planes; their data have to match the ternary data smoothly. At the same time each of the edge binary systems A-B is a boundary plane for many other ternary A-B-X systems. Therefore combining systematically binary and ternary evaluations increases confidence and reliability in both ternary and binary phase diagrams. This has started systematically for the first time here, by the MSIT® Evaluation Programs applied to the Landolt-Börnstein New Series. The degree of success, however, depends on both the nature of materials and scientists!

The multitude of correlated or inter-dependant data requires special care. Within MSIT® an evaluation routine has been established that proceeds knowledge driven and applies both, human based expertise and electronically formatted data and software tools. MSIT® internal discussions take place in almost all evaluation works and on many different specific questions the competence of a team is added to the work of individual authors. In some cases the authors of earlier published work contributed to the knowledge

base by making their original data records available for re-interpretation. All evaluation reports published here have undergone a thorough review process in which the reviewers had access to all the original data.

In publishing we have adopted a standard format that presents the reader with the data for each ternary system in a concise and consistent manner, as applied in the “MSIT® Workplace Phase Diagrams Online”. The standard format and special features of the Landolt-Börnstein compendium are explained in the Introduction to the volume.

In spite of the skill and labor that have been put into this volume, it will not be faultless. All criticisms

and suggestions that can help us to improve our work are very welcome. Please contact us via [email protected]. We hope that this volume will prove to be as useful for the materials scientist and engineer as the other volumes of Landolt-Börnstein New Series and the previous works of MSIT® have been. We hope that the Landolt Börnstein Sub-series, Ternary Alloy Systems will be well received by our colleagues in research and industry.

On behalf of the participating authors we want to thank all those who contributed their comments and

insight during the evaluation process. In particular we thank the reviewers - Hans Leo Lukas, Marina Bulanova, Paola Riani, Lazar Rokhlin, Anatolii Bondar, Yong Du, Olga Fabrichnaya, Artem Kozlov, K.C. Hari Kumar, Viktor Kuznetsov, Ludmila Tretyachenko and Tamara Velikanova.

We all gratefully acknowledge the dedicated scientific desk editing by Oleksandra Berezhnytska and Oleksandr Rogovtsov.

Günter Effenberg, Svitlana Ilyenko and Oleksandr Dovbenko Stuttgart, July 2007

Foreword

Can you imagine a world without iron and steel? No? I can’t either. The story of mankind is intimately linked to the discovery and successful use of metals and their

alloys. Amongst them iron and steel - we could define steel as ‘a generally hard, strong, durable, malleable alloy of iron and carbon, usually containing between 0.2 and 1.5 percent carbon, often with other constituents such as manganese, Chromium, nickel, molybdenum, copper, tungsten, Cobalt, or silicon, depending on the desired alloy properties, and widely used as a structural material’, have shaped our material world.

The story of iron takes us back to the period of the Hittite Empire around 1300 BC, when iron started

to replace bronze as the chief metal used for weapons and tools. Until today the story remains uncompleted and the social and economic impact of the iron and steel industry is now beyond imagination. In the year 2005 1.13 billion tons of crude steel were produced. Compared to 2004 this is an increase of 6.8%. That same year the steel production in China increased from 280.5 to almost 350 million tons. Concerning stainless steel: according to the International Stainless Steel Forum (ISSF), the global production forecast for 2006 now stands at 27.8 million metric tons of stainless crude steel, up 14.3% compared to 2005.

An English poem from the 19th century tells us

Gold is for the mistress Silver for the maid Copper for the craftsman Cunning at his trade Good said the baron Sitting in his hall But iron, cold iron Is master of them all

It is still actual and true. The list of different steel grades and related applications is impressive and still growing: low carbon

strip steels for automotive applications, low carbon structural steels, engineering steels, stainless steels, cast irons, and, more recently: dual phase steels, TRIP-steels, TWIP-steels, maraging steels, …

The list of applications seems endless: a wide range of properties from corrosion resistance to high tensile strength is covered. These properties depend on the percentage of carbon, the alloying elements, and increasingly on the thermo-mechanical treatments that aim at optimizing the microstructure.

Yet many potential improvements remain unexplored, also due to the increasing complexity of the

new steel grades. For instance, a recently patent protected new die steel for hot deformation has the following composition specifications: C 0.46 – 0.58; Si 0.18 – 0.40; Mn 0.45 – 0.75, Cr 0.80 – 1.20; Ni 1.30 – 1.70; Mo 0.35 – 0.65; V 0.18 – 0.25; Al 0.01 – 0.04; Ti 0.002 – 0.04; B 0.001 – 0.003; Zr 0.02 – 0.04; Fe remaining.

Although many properties of steel are directly related to non-equilibrium states, it remains a fact that

the equilibrium state creates the reference frame for all changes that might occur in any material - and consequently would effect its properties in use - that is actually not in its thermodynamic equilibrium state. This is what these volumes in the Landolt-Börnstein series stand for: they have collected the most reliable data on the possible phase equilibria in ternary iron based alloys. Therefore this first volume of data, as well as the other ones in a series of four to appear, is of immeasurable value for metallurgists and materials engineers that improve the properties of existing steels and develop new and more complex steel grades. It is about materials, it is about quality of life.

The well-recognized quality label of MSIT®, the Materials Science International Team, also applies to the present volume of the Landolt-Börnstein series. It should be available for every materials engineer, scientist and student.

Prof. Dr. ir. Patrick Wollants Chairman - Department of Metallurgy and Materials Engineering Katholieke Universiteit Leuven Belgium

Contents IV/11D3 Ternary Alloy Systems Phase Diagrams, Crystallographic and Thermodynamic Data Subvolume D Iron Systems Part 3 Selected Systems from Co-Fe-Si to Cu-Fe-Pt Introduction

Data Covered ................................................................................................................................... XIII General............................................................................................................................................. XIII Structure of a System Report ........................................................................................................... XIII

Introduction.............................................................................................................................. XIII Binary Systems ........................................................................................................................ XIII Solid Phases .............................................................................................................................XIV Quasibinary Systems................................................................................................................. XV Invariant Equilibria ................................................................................................................... XV Liquidus, Solidus, Solvus Surfaces........................................................................................... XV Isothermal Sections................................................................................................................... XV Temperature – Composition Sections ....................................................................................... XV Thermodynamics....................................................................................................................... XV Notes on Materials Properties and Applications....................................................................... XV Miscellaneous ........................................................................................................................... XV References............................................................................................................................. XVIII

General References ..........................................................................................................................XIX Ternary Systems

Co – Fe – Si (Cobalt – Iron – Silicon)..................................................................................................1 Co – Fe – V (Cobalt – Iron – Vanadium)...........................................................................................20 Co – Fe – W (Cobalt – Iron – Tungsten)............................................................................................42 Cr – Cu – Fe (Chromium – Copper – Iron) ........................................................................................57 Cr – Fe – H (Chromium – Iron – Hydrogen)......................................................................................84 Cr – Fe – Mn (Chromium – Iron – Manganese).................................................................................91 Cr – Fe – Mo (Chromium – Iron – Molybdenum) ...........................................................................106 Cr – Fe – N (Chromium – Iron – Nitrogen) .....................................................................................127 Cr – Fe – Nb (Chromium – Iron – Niobium) ...................................................................................145 Cr – Fe – Ni (Chromium – Iron – Nickel)........................................................................................154 Cr – Fe – O (Chromium – Iron – Oxygen).......................................................................................179 Cr – Fe – P (Chromium – Iron – Phosphorus)..................................................................................200 Cr – Fe – S (Chromium – Iron – Sulfur) ..........................................................................................215 Cr – Fe – Si (Chromium – Iron – Silicon)........................................................................................242 Cr – Fe – Ti (Chromium – Iron – Titanium) ....................................................................................269 Cr – Fe – V (Chromium – Iron – Vanadium) ...................................................................................283

Cr – Fe – Zr (Chromium – Iron – Zirconium)..................................................................................298 Cs – Fe – O (Cesium – Iron – Oxygen)............................................................................................308 Cu – Fe – H (Copper – Iron – Hydrogen) ........................................................................................315 Cu – Fe – Mn (Copper – Iron – Manganese) ...................................................................................320 Cu – Fe – Mo (Copper – Iron – Molybdenum) ................................................................................333 Cu – Fe – Nb (Copper – Iron – Niobium) ........................................................................................343 Cu – Fe – Ni (Copper – Iron – Nickel).............................................................................................352 Cu – Fe – O (Copper – Iron – Oxygen)............................................................................................379 Cu – Fe – P (Copper – Iron – Phosphorus) ......................................................................................403 Cu – Fe – Pt (Copper – Iron – Platinum)..........................................................................................422

Introduction

Data Covered

The series focuses on light metal ternary systems and includes phase equilibria of importance for alloydevelopment, processing or application, reporting on selected ternary systems of importance to industriallight alloy development and systems which gained otherwise scientific interest in the recent years.

General

The series provides consistent phase diagram descriptions for individual ternary systems. The representationof the equilibria of ternary systems as a function of temperature results in spacial diagrams whose sectionsand projections are generally published in the literature. Phase equilibria are described in terms of liquidus,solidus and solvus projections, isothermal and quasibinary sections; data on invariant equilibria are gener-ally given in the form of tables.The world literature is thoroughly and systematically searched back to the year 1900. Then, the publisheddata are critically evaluated by experts in materials science and reviewed. Conflicting information is com-mented upon and errors and inconsistencies removed wherever possible. It considers those, and only thosedata, which are firmly established, comments on questionable findings and justifies re-interpretations madeby the authors of the evaluation reports.In general, the approach used to discuss the phase relationships is to consider changes in state and phasereactions which occur with decreasing temperature. This has influenced the terminology employed and isreflected in the tables and the reaction schemes presented.The system reports present concise descriptions and hence do not repeat in the text facts which can clearlybe read from the diagrams. For most purposes the use of the compendium is expected to be self-sufficient.However, a detailed bibliography of all cited references is given to enable original sources of information tobe studied if required.

Structure of a System Report

The constitutional description of an alloy system consists of text and a table/diagram section which are sepa-rated by the bibliography referring to the original literature (see Fig. 1). The tables and diagrams carry theessential constitutional information and are commented on in the text if necessary.Where published data allow, the following sections are provided in each report:

Introduction

The opening text reviews briefly the status of knowledge published on the system and outlines the experi-mental methods that have been applied. Furthermore, attention may be drawn to questions which are stillopen or to cases where conclusions from the evaluation work modified the published phase diagram.

Binary Systems

Where binary systems are accepted from standard compilations reference is made to these compilations. Inother cases the accepted binary phase diagrams are reproduced for the convenience of the reader. The selec-tion of the binary systems used as a basis for the evaluation of the ternary system was at the discretion of theassessor.

Solid Phases

The tabular listing of solid phases incorporates knowledge of the phases which is necessary or helpful forunderstanding the text and diagrams. Throughout a system report a unique phase name and abbreviationis allocated to each phase.

Introduction 1

Landolt-BörnsteinNew Series IV/11D3

DOI: 10.1007/978-3-540-74199-2_1# Springer 2008MSIT®

Phases with the same formulae but different space lattices (e.g. allotropic transformation) are distinguishedby:

– small letters (h), high temperature modification (h2 > h1)(r), room temperature modification(1), low temperature modification (l1 > l2)

– Greek letters, e.g., ε, ε′– Roman numerals, e.g., (I) and (II) for different pressure modifications.

In the table “Solid Phases” ternary phases are denoted by * and different phases are separated by horizontallines.

Quasibinary Systems

Quasibinary (pseudobinary) sections describe equilibria and can be read in the same way as binary dia-grams. The notation used in quasibinary systems is the same as that of vertical sections, which are reportedunder “Temperature – Composition Sections”.

Fig. 1. Structure of a system report

2 Introduction

DOI: 10.1007/978-3-540-74199-2_1 Landolt-Börnstein# Springer 2008 New Series IV/11D3MSIT®

Fig.2.

Typicalreactio

nscheme

Introduction 3



Invariant Equilibria

The invariant equilibria of a system are listed in the table “Invariant Equilibria” and, where possible, aredescribed by a constitutional “Reaction Scheme” (Fig. 2).The sequential numbering of invariant equilibria increases with decreasing temperature, one numbering forall binaries together and one for the ternary system.Equilibria notations are used to indicate the reactions by which phases will be

– decomposed (e- and E-type reactions)– formed (p- and P-type reactions)– transformed (U-type reactions)

For transition reactions the letter U (Übergangsreaktion) is used in order to reserve the letter T to denotetemperature. The letters d and D indicate degenerate equilibria which do not allow a distinction accordingto the above classes.

Liquidus, Solidus, Solvus Surfaces

The phase equilibria are commonly shown in triangular coordinates which allow a reading of the concentra-tion of the constituents in at.%. In some cases mass% scaling is used for better data readability (see Figs. 3and 4).In the polythermal projection of the liquidus surface, monovariant liquidus grooves separate phase regionsof primary crystallization and, where available, isothermal lines contour the liquidus surface (see Fig. 3).

Fig. 3. Hypothetical liqudus surface showing notation employed

4 Introduction


Isothermal Sections

Phase equilibria at constant temperatures are plotted in the form of isothermal sections (see Fig. 4).

Temperature – Composition Sections

Non-quasibinary T-x sections (or vertical sections, isopleths, polythermal sections) show the phase fieldswhere generally the tie lines are not in the same plane as the section. The notation employed for the latter(see Fig. 5) is the same as that used for binary and quasibinary phase diagrams.

Thermodynamics

Experimental ternary data are reported in some system reports and reference to thermodynamic modeling ismade.

Notes on Materials Properties and Applications

Noteworthy physical and chemical materials properties and application areas are briefly reported if theywere given in the original constitutional and phase diagram literature.

Fig. 4. Hypotheticcal isothermal section showing notation employed

Introduction 5



Miscellaneous

In this section noteworthy features are reported which are not described in preceding paragraphs. Theseinclude graphical data not covered by the general report format, such as lattice spacing – composition data,p-T-x diagrams, etc.

References

The publications which form the bases of the assessments are listed in the following manner:[1974Hay] Hayashi, M., Azakami, T., Kamed, M., “Effects of Third Elements on the Activity of Lead inLiquid Copper Base Alloys” (in Japanese), Nippon Kogyo Kaishi, 90, 51–56 (1974) (Experimental, Ther-modyn., 16)This paper, for example, whose title is given in English, is actually written in Japanese. It was published in1974 on pages 51- 56, volume 90 of Nippon Kogyo Kaishi, the Journal of the Mining and MetallurgicalInstitute of Japan. It reports on experimental work that leads to thermodynamic data and it refers to 16cross-references.Additional conventions used in citing are:# to indicate the source of accepted phase diagrams* to indicate key papers that significantly contributed to the understanding of the system.Standard reference works given in the list “General References” are cited using their abbreviations and arenot included in the reference list of each individual system.

Fig. 5. Hypothetical vertical section showing notation employed

6 Introduction


General References[C.A.] Chemical Abstracts - pathways to published research in the world's journal and patent litera-

ture - http://www.cas.org/[Curr.Cont.]

Current Contents - bibliographic multidisciplinary current awareness Web resource - http://www.isinet.com/products/cap/ccc/

[E] Elliott, R.P., Constitution of Binary Alloys, First Supplement, McGraw-Hill, New York (1965)[G] Gmelin Handbook of Inorganic Chemistry, 8th ed., Springer-Verlag, Berlin[H] Hansen, M. and Anderko, K., Constitution of Binary Alloys, McGraw-Hill, New York (1958)[L-B] Landolt-Boernstein, Numerical Data and Functional Relationships in Science and Technology

(New Series). Group 3 (Crystal and Solid State Physics), Vol. 6, Eckerlin, P., Kandler, H. andStegherr, A., Structure Data of Elements and Intermetallic Phases (1971); Vol. 7, Pies, W. andWeiss, A., Crystal Structure of Inorganic Compounds, Part c, Key Elements: N, P, As, Sb, Bi,C (1979); Group 4: Macroscopic and Technical Properties of Matter, Vol. 5, Predel, B., PhaseEquilibria, Crystallographic and Thermodynamic Data of Binary Alloys, Subvol. a: Ac-Au ...Au-Zr (1991); Springer-Verlag, Berlin.

[Mas] Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, ASM, Metals Park, Ohio (1986)[Mas2] Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals

Park, Ohio (1990)[P] Pearson, W.B., A Handbook of Lattice Spacings and Structures of Metals and Alloys, Perga-

mon Press, New York, Vol. 1 (1958), Vol. 2 (1967)[S] Shunk, F.A., Constitution of Binary Alloys, Second Supplement, McGraw-Hill, New York

(1969)[V-C] Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic

Phases, ASM, Metals Park, Ohio (1985)[V-C2] Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic

Phases, 2nd edition, ASM, Metals Park, Ohio (1991)

Introduction 7



http://www.cas.org/

http://www.isinet.com/products/cap/ccc/

http://www.isinet.com/products/cap/ccc/

Index of alloy systems

Index of Ternary Iron Alloy Systems Co-Fe-Si to Cu-Fe-Pt

Co – Fe – Si (Cobalt – Iron – Silicon)Co – Fe – V (Cobalt – Iron – Vanadium)Co – Fe – W (Cobalt – Iron – Tungsten)Cr – Cu – Fe (Chromium – Copper – Iron)Cr – Fe – H (Chromium – Iron – Hydrogen)Cr – Fe – Mn (Chromium – Iron – Manganese)Cr – Fe – Mo (Chromium – Iron – Molybdenum)Cr – Fe – N (Chromium – Iron – Nitrogen)Cr – Fe – Nb (Chromium – Iron – Niobium)Cr – Fe – Ni (Chromium – Iron – Nickel)Cr – Fe – O (Chromium – Iron – Oxygen)Cr – Fe – P (Chromium – Iron – Phosphorus)Cr – Fe – S (Chromium – Iron – Sulfur)Cr – Fe – Si (Chromium – Iron – Silicon)Cr – Fe – Ti (Chromium – Iron – Titanium)Cr – Fe – V (Chromium – Iron – Vanadium)Cr – Fe – Zr (Chromium – Iron – Zirconium)Cs – Fe – O (Cesium – Iron – Oxygen)Cu – Fe – H (Copper – Iron – Hydrogen)Cu – Fe – Mn (Copper – Iron – Manganese)Cu – Fe – Mo (Copper – Iron – Molybdenum)Cu – Fe – Nb (Copper – Iron – Niobium)Cu – Fe – Ni (Copper – Iron – Nickel)Cu – Fe – O (Copper – Iron – Oxygen)Cu – Fe – P (Copper – Iron – Phosphorus)Cu – Fe – Pt (Copper – Iron – Platinum)

Index of alloy systems 1



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Cobalt – Iron – Silicon

Lazar Rokhlin

Introduction

This system is of interest for a number of industrial applications, principally commercial alloyed steels(Invar alloys), which lie in the Co-Fe rich part of the system, and thermoelectric materials, compositionsof which lie in the silicon rich part.In the oldest and quite detailed experimental investigation [1935Vog], the Co-CoSi-FeSi-Fe region of thesystem (up to ~35 mass% (~50 at.%) Si) was studied. Phase equilibria in this part of the system were deter-mined and partial liquidus and solidus surfaces along with nine polythermal sections (three of which beingquasibinary) were constructed. Further significant additions to the Co-Fe-Si phase diagram were made fol-lowing the study of solid state phase equilibria across the whole concentration range, carried out by[1975Fed], the constitution of the CoSi-FeSi and CoSi2-FeSi2 sections [1961Wit, 1964Asa, 1965Zel,1970Hes, 1971Uga] and the limits of the extension of the ordered phases based on the bcc-Fe solid solution(αδ phase) [1955Gri, 1989Koz, 1990Koz, 1990Fuk, 1991Fuk].[1949Jae] presented a short review of the Co-Fe-Si phase diagram based on [1935Vog]. Later, [1988Ray]gave a detailed review of the Co-Fe-Si phase diagram, which included the work of [1935Vog] and other ear-lier studies. This was later updated by [1994Rag], who included details of the ordering of the Fe rich bccsolid solution (αδ phase).Details of the experimental studies are reported in Table 1.First principles computational method was used by [1991Mot] in the study of the phase diagram of theFe1–xCoxSi2 section. The focus of the study was the phase equilibria involving phases with the CaF2 typestructure.

Binary Systems

The three binary systems Co-Fe, Co-Si, Fe-Si are accepted from [Mas2].

Solid Phases

Details of the solid phases in the system are reported in Table 2. The system is characterized by extendedsolid solutions, some of them being continuous. Only binary phases with ternary extensions are presentin this system. Among them is the Co2Si based solid solution where Fe can replace Co up to the composi-tion FeCoSi, which was initially assumed to be a ternary compound [1935Vog, 1998Lan]. The foundationof this assumption came from the solidification of the FeCoSi alloy and its polymorphous transformation atconstant temperature [1935Vog]. Also, it belongs to the TiNiSi series of compounds [1998Lan]. The FeCoSi‘compound’ is named in this assessment as τ. CsCl type ordering takes place in the α bcc (αFe base) solidsolution. This ordering observed in both the Co-Fe and Fe-Si systems over a large range of Co/Si concen-trations. The CsCl ordered phases in both binary systems (α’ in Co-Fe and α2 in Fe-Si) join across the tern-ary system. The second ordered phase, α1 of the BiF3 type, is observed only in the Fe-Si binary system butextends significantly into the ternary system.

Quasibinary Systems

The FeSi-CoSi vertical section of the ternary Co-Fe-Si is quasibinary. It is shown in Fig. 1. The system ischaracterized by a continuous solid solution between the FeSi and CoSi phases (θ) and an absence of anyinvariant reaction. The liquidus and solidus lines in Fig. 1 are shown after [1935Vog] and assuming themelting points for FeSi and CoSi according to the accepted Fe-Si and CoSi binaries [Mas2]. In effect, theFeSi-CoSi section divides the Co-Fe-Si system into two subsystems: Co-CoSi-FeSi-Fe and CoSi-Si-FeSi.

Co–Fe–Si 1



A second quasibinary system exists; the FeSi-FeCoSi(τ) vertical section. It is presented in Fig. 2 after[1935Vog] with the same minor amendments to take into account the accepted Co-Si and Fe-Si binary sys-tems [Mas2].The FeSi2-CoSi2 vertical section was indicted to be quasibinary by [1965Zel, 1988Ray]. However, accord-ing to the accepted Fe-Si and Co-Si binary systems [Mas2], there is only one allotropic form of the com-pound CoSi2, but two allotropic forms of the compound FeSi2. The high temperature βFeSi2 form meltscongruently and decomposes by a eutectoid reaction at 937°C. The composition of this high temperatureform (72 at.% Si at the melting maximum) differs slightly from the composition of the low temperatureαFeSi2 form corresponding to a strict stoichiometry of 66.7 at.% Si. Therefore, at temperatures below937°C, the FeSi2-CoSi2 section presented in [1965Zel, 1988Ray] will cross a two-phase region with a phaseother than αFeSi2 on the Fe-Si side. Consequently, the FeSi2-CoSi2 section presented in Fig. 3 after[1988Ray] is only partially quasibinary, meaning at temperatures higher than 937°C only.In [1935Vog], one more vertical section was indicated as quasibinary; Co2Si-FeCoSi(τ). As in the case ofthe previous section, the Co2Si-FeCoSi section can be considered to be only partially quasibinary as a con-sequence of the Co-Si binary phase diagram [Mas2]. According to [1935Vog], continuous solid solutionsexist between the high and low temperature allotropic forms of Co2Si and FeCoSi. However, accordingto the accepted Co-Si phase diagram [Mas2], the low temperature αCo2Si is formed during cooling fromthe melt by the peritectic reaction l + βCo2Si ⇌αCo2Si at ~1320°C; the high temperature βCo2Si phasehaving a congruent melting point of 1334°C. The binary composition of the βCo2Si phase is shifted slightly(by ~1 at.% Si) with respect to that of αCo2Si. Therefore, in the narrow region at ~1320-1334°C adjoiningthe Co2Si side, the Co2Si-FeCoSi section cannot be quasibinary. The Co2Si-FeCoSi section is presented inFig. 4 with the deviation from quasibinary nature being shown by dashed lines. However, [1988Ray] wasdubious over the solubility of Fe in βCo2Si and αCo2Si reaching the ‘ternary compound’ FeCoSi, referringto the work of [1975Fed] who determined the solubility of Fe in αCo2Si at 800°C to be only ~9 mass%(8 at.%). [1988Ray] accepted arbitrarily the solubility of Fe in βCo2Si and αCo2Si after solidification tobe about 17 mass% (15 at.%) and 10 mass% (8.8 at.%), respectively (as compared with 39 mass% Fe inFeCoSi (33.3 at.%)). According to [1988Ray], the ‘ternary compound’ FeCoSi (τ) does not exist, and there-fore, the sections FeSi-FeCoSi and Co2Si-FeCoSi cannot be quasibinary. Meanwhile in [1975Fed], only asingle isothermal section (at 800°C) was studied, using X-ray diffraction and microstructure investigationand no thermal analysis was conducted. The results of [1975Fed] cannot be considered as definitive owingto the contradiction between them and the conclusions of [1935Vog] and therefore additional study isrequired. Electron constitution and the sizes of the Fe and Co atoms suggest the possibility of significantsolubility of Fe in αCo2Si and βCo2Si. Therefore, the formation of ‘FeCoSi’ compound is quite probable.


The invariant equilibria were established reliably in the subsystem Co-CoSi-FeSi-Fe [1935Vog, 1988Ray].They are listed in Table 3 with amendments to take into account the accepted binary Co-Si phase diagram[Mas2] with the peritectic reaction l + (αCo)⇌ (εCo). Table 3 contains additionally the four-phase invariantequilibrium U3 involving (εCo), which must be present taking into account the above peritectic reaction.The three-phase equilibria given in Table 3 correspond to the quasibinary system FeSi-FeCoSi establishedin [1935Vog]. In Figs. 5a and 5b, the reaction scheme for the Co-CoSi-Fe-Si-Fe subsystem is presented after[1988Ray] with the same amendments to take into account the peritectic equilibrium l + (αCo) ⇌ (εCo) inthe accepted binary Co-Si phase diagram [Mas2]. Following [1988Ray] the reactions with the participationof the ternary compound τ are not included. In the CoSi2-Si-FeSi2 subsystem, the three-phase invariantequilibrium L ⇌ βFeSi2 + CoSi2 was established experimentally [1965Zel]. This invariant equilibrium isalso included in Table 3.

Liquidus, Solidus and Solvus Surfaces

The projection of the liquidus surface of the subsystem Co-CoSi-FeSi-Fe is presented in Fig. 6. It is con-structed after [1935Vog] with the addition of the corrections of [1988Ray] and amendments to maintain con-sistency with the binary systems [Mas2]. The proposed monovariant line L + (αCo) ⇌ (εCo) is shown as

2 Co–Fe–Si


dashed on the liquidus surface projection. This line runs from the liquid point p3 of the binary invariantequilibrium l + (αCo) ⇌ (εCo) [Mas2] to the proposed invariant four-phase point U3. Following[1988Ray] the reactions with the participation of the ternary compound τ are not included.Although the liquidus surface of the CoSi-Si-FeSi subsystem was not determined experimentally,[1988Ray] presented a ‘hypothetical’ version, shown in Fig. 7.Fig. 8 presents the projection of the composition of the solid phases which separate on cooling in the Co-Co2Si-Fe2Si-Fe subsystem. It shows the monovariant lines that bound parts of the solidus surface and solidphase regions. The projection was suggested by [1988Ray] based on the work of [1935Vog, 1975Fed].Some minor corrections have been made to ensure consistency with the binary systems [Mas2]. Followingthe conclusions of [1988Ray], the ‘FeCoSi phase’ has been omitted from the projection.

Isothermal Sections

Two isothermal sections were proposed by [1988Ray], for temperatures of 1160 and 800°C. These sectionswere based on the published works of [1935Vog, 1975Fed] with some amendments in order to reconcile thecontradictions between these works concerning the solubility of Fe in αCo2Si and βCo2Si and existence ofthe FeCoSi ‘compound’ following the work of [1975Fed]. The sections are presented in Figs. 9 and 10 withsome corrections to give consistency with the accepted binary systems [Mas2]. Following [1935Vog], the1160°C section contains the Co3Si region although in the Co-Si binary [Mas2] this compound exists onlyin the temperature range 1214-1204°C. Therefore, [1988Ray] suggested that the Co3Si based solid solutionexists in the ternary system at the lower temperature of 1160°C, but in a region of the diagram away fromand not including the Co-Si binary edge. Also in this section, tentative regions for (εCo) have been includedbecause of the existence of this phase in the binary Co-Si [Mas2]. The section at 800°C shows probablephase boundaries for regions where the ordered phases α2 and α1 can be formed in accordance with theCo-Fe and Fe-Si binary systems [Mas2]. For both sections, at 1160 and 800°C, [1988Ray] assumed thesolubility of Fe in αCo2Si to be less than in FeCoSi in contradiction with [1935Vog]. The partial isothermalsection at 550°C for the Fe corner of the phase diagram is presented in Fig. 11. The section shows the fieldswhere the ordered phases α2(α’) and α1 are stable. It is drawn following the results of experimental worksby [1989Koz, 1990Fuk, 1990Koz, 1994Koz, 1994Rag]. The results of the experiments were confirmed bycalculations.


Two vertical sections of the Co-CoSi-FeSi-Fe subsystem are presented in Figs. 12 and 13. They are drawnafter [1935Vog] with minor corrections to maintain consistency with the Co-Fe, Co-Si and Fe-Si binary sys-tems [Mas2].

Thermodynamics

The free energy of Fe base Co-Fe-Si ordering alloys was estimated by [1990Koz] based on the statisticalapproach of the Bragg-Williams-Gorsky approximation. The calculated isothermal section at 550°C showedgood agreement with the experimental one constructed by [1990Fuk].[2003Bol] investigated thermodynamic properties of the phases in the CoSi-FeSi and CoSi2-FeSi2 sectionsby chemical vapor transport methods. For the first section the composition dependences of enthalpy,entropy and heat capacity were given.


It is well known that βFeSi2 is a good candidate for thermoelectric applications and can be used at high tem-peratures, in the range 427-727°C [2006Ito, 2003Kim, 2003Ito, 2003Zhu, 2002Aru, 2002Tan, 2002Ur,2000Bel]. [1964Asa] was the first to study the physical properties of this system. βFexCo1–xSi2 (with x =0.03-0.05) is an n-type semiconductor, and it has been prepared using a powder metallurgy technique by[2003Kim]. Mechanochemical synthesis was used to prepare this material by [2000Bel]. Semi-metallicproperties of Co1–xFexSi solid solutions were studied by resistivity and thermoelectric measurements.

Co–Fe–Si 3



Owing to the presence of two magnetic elements in the ternary materials, special magnetic behavior isnoted. The magnetic properties of the FeSi compound and solid solutions based on CoSi were studied by[1977Che, 1977Nic, 1983Bus, 1986Mat, 1998Sch]. The magnetic properties of the Co-doped n-typeβ-FeSi2.5 single crystals were studied by [2002Aru].Thermal expansion and weak itinerant magnetism in Fe1–xCoxSi solid solutions were investigated by[1987Gel, 1986Gel]. The magnetization and magnetoresistance of Fe1–xCoxSi alloys are presented in[2002Cha1, 2002Cha2]. Later the L12 Heusler phase was studied by [2006Wur].FeCo-Si multilayers were studied by [2003Cho] using neutron diffraction, for applications in supermirrors.Mössbauer spectroscopy has been used widely in the study of these compounds. [1979Mey, 1982Gel] deter-mined the atomic configurations in Fe0.5–xCoxSi0.5 and in Fe1–xCoxSi by quantitative Mössbauer spectro-scopy. The phase separation of Co-Fe-Si alloys was studied using Mössbauer spectroscopy by[1996Mor]. The site occupation of dilute Co impurities in Fe3Si determined using the Mössbauer techniqueis presented in [1976Bla].Ordering in the bcc solid solutions (αδ phase) studied by neutron diffraction and Mössbauer is described in[1979Mey, 1979Ind].Anomalous regions in the magnetic phase diagram of (Fe,Co)Si was established and investigated in[1990Ish].

Miscellaneous

[1984Lan] suggested equations for the determination of the thermodynamic variables in the Co-Fe-Si sys-tem on the basis of its gas (hydrogen) absorbing capacity. [1991Nis] presents an experimental study of theinteraction parameter for carbon in Co-FeSi system in iron rich alloys.In [1974Bur, 1975Pic], the preferential lattice sites for occupation by dotted Co atoms in the Fe3Si structurewere presented.

Table 1. Experimental Investigations of the Co-Fe-Si Phase Relations, Structures, Thermodynamics

Reference Method / Experimental Technique Temperature / Composition /Phase Range Studied

[1955Gri] Thermal analysis <1000°C/ up to 5 mass% Si atequal Fe and Co contents

[1961Wit] X-ray diffraction Room temperature/ sectionsFeSi-CoSi and FeSi2-CoSi2

[1964Asa] Electrical resistivity, Hall coefficient, thermocouple power 4.2-800 K/ section FeSi2-CoSi2

[1965Zel] Optical microscopy, X-ray diffraction, thermal analysis 1000-1300°C/section FeSi2-CoSi2

[1970Hes] X-ray diffraction, thermal conductivity, Hall effect 750°C/Fe1–xCoxSi2,0 ≤ x ≤ 0.2

[1971Uga] Thermo-emf, electrical conductivity 25-1000°C/ Fe1–xCoxSi2, 0.005≤ x ≤ 0.1

[1975Fed] X-ray diffraction 800°C/whole compositionrange in the ternary

[1975Ver] X-ray diffraction, optical microscopy 1000°C/ (Fe1–xCox)3Si with0 ≤ x ≤ 0.6

(continued)

4 Co–Fe–Si


Reference Method / Experimental Technique Temperature / Composition /Phase Range Studied

[1989Koz] Transmission electron microscopy 550°C/ up to ~30 at.% Co, ~15at.% Si



[1998Lan] X-ray diffraction Room temperature/ CoFeSi

[2000Bel] X-ray diffraction, DTA, thermoelectric properties,electrical conductivity, electron microscopy

Room temperature/Fe1–xCoxSi2, 0.02 ≤ x ≤ 0.1

[2002Ers] X-ray diffraction anomalous fine structure (DAFS)spectra/High- resolution transmission electron microscopy(HRTEM)

Room temperature, 657°C/FexCo1–xSi2

[2002Fet] Mössbauer spectrometry Room temperature, 1000°C/Fe1–xCoxSi2

[2003Zhu] X-ray diffraction, thermoelectric properties, electricalresistance

50-500°C/ Fe1.86Co0.14Si5

[2003Kim] X-ray diffraction, SEM-EDX, physical properties ~100-700°C/ Fe1–xCoxSi2,0.01 ≤ x ≤ 0.03

[2003Ito] X-ray diffraction, SEM, EDX analysis, physical properties From RT to ~900°C/Fe0.98Co0.02Si2

[2006Wur] Magnetization/SQUID (superconducting quantuminterference device) magnetometry + VSM (vibrating-sample magnetometry)

5 K, 700 - 1150°C/ Co2FeSi/L12 phase

Table 2. Crystallographic Data of Solid Phases

Phase/TemperatureRange [°C]

PearsonSymbol/SpaceGroup/Prototype

LatticeParameters[pm]

Comments/References

γ, (γFe,αCo)

(αCo)1495 - 422(γFe)1394 - 912

cF4Fm�3mCu

a = 356.88

a = 364.67

continuous solid solution between γFe andαCo, dissolves up to 16.5 at.%Si [Mas2]

pure Co at 520°C [V-C2, Mas2]

pure Fe at 915°C [V-C2, Mas2]

(continued)

Co–Fe–Si 5






Comments/References

(εCo)< 1250

hP2P63/mmcMg

a = 250.71c = 406.86

dissolves up to 18.4 at.% Si at 1214°C [Mas2]

pure Co at 25°C [Mas2]

αδ, (αFe,δFe)< 1538

(δFe)1538 - 1394(αFe)< 912

cI2Im�3mW

a = 293.78

a = 286.65

continuous solid solution between δFe andαFe, dissolves up to ~78 at.% Co at 500°Cand 19.5 at.% Si [Mas2]pure Fe at 1480°C [V-C, Mas2]

pure Fe at 25°C [Mas2]

(εFe) hP2P63/mmcMg

a = 246.8c = 396.0

pure Fe at 25°C, 13 GPa [Mas2]

(αSi) cF8Fd�3mC(diamond)

a = 543.6 pure Si at 25°C [Mas2]

(βSi) tI4I41/amdβSn

a = 468.6c = 258.5

pure Si, at 25°C, p>9.5 GPa [Mas2]

(γSi) cI16Im�3mγSi

a = 663.6 pure Si at 25°C, p>16 GPa [Mas2]

(δSi) hP4P63/mmcαLa

a = 380c = 628

pure Si at 25°C, 16 GPa → 1 atm [Mas2]

α2, FeCo< 1302

cP2Pm�3mCsCl

a = 285.27 to284.34

ordered αδ,at 25-75 at.% Fe in Co-Fe [Mas2, V-C2],dissolves up to ~10-22 at.% Si in Fe-Si[Mas2, 1982Kub, 1989Koz, 1990Koz]

Co3Si1214 - 1193

t** - [Mas2]

αFe1–xCo1+xSi≲ 1320

αCo2Si≲ 1320

oP12PnmaCo2Si(τ, TiNiSi)

a = 491.8b = 373.7c = 710.9

0 ≤ x ≤ 1at ~32-34 at.% Si [1935Vog, Mas2]

x = 1 [P]at ~32-34 at.% Si

(continued)

6 Co–Fe–Si





Comments/References

τ(l), FeCoSi(l)< 970

a = 494.2 ± 0.2b = 377.6 ± 0.1c = 717.2 ± 0.2

x = 0 [1935Vog, 1998Lan]

βFe1–xCo1+xSi1334 - 970βCo2Si1334 - 1238τ(h), FeCoSi(h)1245 - 970

- -

-

-

0 ≤ x ≤ 1 [1935Vog, Mas2]

x = 1at ~32-35.8 at.% Si, [Mas2]x = 0, dissolves up to 28.4 mass% Feand 2.9 mass% Si [1935Vog, Mas2]

θ, Fe1–xCoxSi< 1460

ε, FeSi< 1410CoSi< 1460

cP8P213FeSi

a = 448.3

a = 444.5

0 ≤ x ≤ 1 [1935Vog, 1988Ray, Mas2]

x = 0, [Mas2, V-C2]

x = 1,at 49 - ~52 at.% Si [V-C2, Mas2]

CoSi2< 1326

cF12Fm�3mCaF2

a = 535 dissolves up to 13 at.% Fe[V-C2, Mas2, 2002Ers]

Co2Si3 tP20P4c2Ru2Sn3

a = 491.8 ± 0.3b = 373.8 ± 0.3c = 710.9 ± 0.3

metastable,prepared at 4 GPa[Mas2, V-C2]

Co3Si hP8P63/mmcNi3Sn

a = 497.6 ± 0.2c = 406.96 ± 0.03

metastable, HT, 1193-1214°C[Mas2, V-C2]

γCo2Si o** - metastable [Mas2]

Co4Si - - metastable [Mas2]

α1(l), Fe3Si< 1156

cF16Fm�3mBiF3

a = 560.8 ordered αδ,at ~10-30 at.% Si in Fe-Si[Mas2, V-C2, 1982Kub]dissolves up to ~27 at.% Co[1989Koz, 1990Koz]or ~45 at.% Co [1975Ver]

α1(h)1250 - 965

- - ordered α, at ~23-31 at.% Si[Mas2, 1982Kub]

αFeSi2< 982

oC48CmcaαFeSi2

a = 986.3 ± 0.7b = 779.1 ± 0.6c = 783.3 ± 0.6

dissolves up to 4 at.% Co at 750°C[Mas2, V-C2, 1970Hes]

(continued)

Co–Fe–Si 7






Comments/References

βFeSi21220 - 937

tP3P4/mmmβFeSi2

a = 269.37 to268.72c = 513.9 to 512.8

at 69.5-73.5 at.% Si in Fe-Si, dissolvesup to ~5 at.% Co [Mas2, V-C2, 1982Kub,1965Zel]

β, Fe2Si1212 - 1040

hP6P�3m1Fe2Si

a = 405.2 ± 0.2c = 5088. ± 0.2

homogeneity range ~1% Si [Mas2, V-C2,1982Kub]

η, Fe5Si31060 - 825

hP16P63/mcmMn5Si3

a = 675.9 ± 0.5c = 472.0 ± 0.5

[Mas2, V-C2, 1982Kub]

Table 3. Invariant Equilibria

Reaction T [°C] Type Phase Composition (mass%)

Co Fe Si

L ⇌ FeSi + βFeCoSi 1215 e6 LFeSiβFeCoSi

32.50~32.5

45.266.5~45.2

22.333.5~22.3

L ⇌ βFeSi2 + CoSi2 1190 e7 LβFeSi2CoSi2

23.84.750.1?

26.745.31.1?

49.550.048.8?

L + Fe2Si ⇌ αδ + Fe1–xCoxSi ? U1 - - - -

L + αδ ⇌ γ + Fe1–xCoxSi 1185 U2 L 27.7 52.1 20.2

L + (εCo) ⇌ γ + Co3Si ? U3 - - - -

L ⇌ γ + βCo2Si + Fe1–xCoxSi 1170 E1 L 33 46.8 20.2

L + Co3Si ⇌ (εCo) + αCo2Si ~1170 U4 L 75 13 12

L ⇌ γ + βCo2Si + αCo2Si 1165 E2 L 73 15 12

βFeCoSi ⇌ FeSi + αFeCoSi 880 e8 βFeCoSiFeSiαFeCoSi

37041.3

4166.539.1

2233.519.6

8 Co–Fe–Si


Fig. 1. Co-Fe-Si. Quasibinary system FeSi - CoSi

Co–Fe–Si 9



Fig. 2. Co-Fe-Si. Quasibinary system FeSi - FeCoSi (τ)

10 Co–Fe–Si


Fig. 3. Co-Fe-Si. Partially quasibinary system FeSi2-CoSi2

Co–Fe–Si 11



Fig. 4. Co-Fe-Si. Partially quasibinary system FeCoSi - Co2Si

12 Co–Fe–Si


Fig.5a

.Co-Fe-Si.Reactionschemeof

theCo-CoS

i-FeSi-Fesubsystem,part1

Co–Fe–Si 13



Fig.5b

.Co-Fe-Si.Reactionschemeof

theCo-CoS

i-FeSi-Fesubsystem,part2

14 Co–Fe–Si


Fig. 6. Co-Fe-Si. Liquidus surface projection of the Co-CoSi-FeSi-Fe subsystem

Co–Fe–Si 15



Fig. 7. Co-Fe-Si. Hypothetical liquidus surface projection of the CoSi-FeSi-Si subsystem

16 Co–Fe–Si


Fig. 8. Co-Fe-Si. Solidus surface projection of the Co-Co2Si-Fe2Si-Fe subsystem

Co–Fe–Si 17



Fig. 9. Co-Fe-Si. Isothermal section of the Co-CoSi-FeSi-Fe subsystem at 1160°C

18 Co–Fe–Si


Fig. 10. Co-Fe-Si. Isothermal section at 800°C

Co–Fe–Si 19



Fig. 11. Co-Fe-Si. Partial isothermal section at 550°C showing fields of the ordered phases α1 and α2

20 Co–Fe–Si


Fig. 12. Co-Fe-Si. Vertical section from Fe to (Fe1–xCoxSi) for the ratio (mass%) Co : Fe = 88.3 : 11.7, plotted in at.%

Co–Fe–Si 21



Fig. 13. Co-Fe-Si. Vertical section from Fe to (Fe1–xCoxSi) for the ratio (mass%) Co : Fe = 55.4 : 44.6, plotted in at.%

22 Co–Fe–Si


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[1975Ver] Vereshchagin, Yu.A., Babanova, E.N., Sidorenko, F.A., “Procedure for Synthesis, Homoge-neity Regions and Lattice Spacings of (Fe1–xMx)3Si (M = Mn, Co, Ni) Solid Solutions with aD03 Structure” (in Russian), Fiz. Metody Issled. Tverdogo Tela, (1), 40–42 (1975) (Crys.Structure, Experimental, 9)

[1976Bla] Blaauw, C., MacKay, G.R., Leiper, W., “Confirmation of Selective Site Occupation of DiluteCo Impurities in Fe3Si Using the Moessbauer Effect Technique”, Solid State Commun., 18,729–730 (1976) (Experimental, 9)

[1977Che] Chechernikova, O.I., Panteleimonov, L.A., Badtiev, E.B., Petrushkova, O.S., “The MagneticProperties of the FeSi Compound and Solid Solutions on the Basis of FeSi, CoSi and NiSi”(in Russian), Vestn.Mosk. Univ., Ser. 2: Khim., 18(4), 433–435 (1977) (Experimental, Magn.Prop., 3)

[1977Nic] Niculescu, V., Budnick, J.I., “Limits of Solubility, Magnetic Properties and Electron Concen-tration in Fe3–xTxSi System”, Solid State Commun., 24(9), 631–634 (1977) (Crys. Structure,Experimental, Magn. Prop., 17)

[1979Ind] Inden, G., “Determination of the Interchange Energies wCoSi(1), wCoSi(2) for bcc SolidSolutions from High Temperature Neutron Diffraction on Ternary Fe-Co-Si Crystals”, Phys.Status Solidi A, 56(1), 177–186 (1979) (Calculation, Crys. Structure, Experimental, Magn.Prop., 19)

Co–Fe–Si 23



[1979Mey] Meyer, W.O., Inden, G., “Quantitative Determination of Ordered Configurations in bccFe-Co-Si Solid Solutions by Mossbauer Spectroscopy”, Phys. Status Solidi A, 56(2),481–486 (1979) (Crys. Structure, Experimental, 6)

[1982Gel] Gel’d, P.V., Povzner, A.A., Romasheva, L.F., “Magnetic Susceptibility of the Fe1-xCoxSiSolid Solutions” (in Russian), Dokl. Akad. Nauk SSSR, 265(6), 1379–1381 (1982) (Experi-mental, Magn. Prop., 8)

[1982Kub] Kubaschewski, O., “Iron-Silicon” in “Iron-binary Phase Diagrams”, Springer Verlag, Berlin,136–139 (1982) (Phase Diagram, Review, 23)

[1983Bus] Buschow, K.H.J., van Engen, P.G., Jongebreur, R., “Magneto-Optical Properties of MetallisFerromagnetic Materials”, J. Magn. Magn. Mater., 38, 1–22 (1983) (Magn. Prop., OpticalProp., 23)

[1984Lan] Lange, K.W., Mitra M., “Determination of the Thermodynamic Variables (δG’E,M, δH’M,δS’E,M) of a Ternary Metallic System on the Basis of its Gas Absorbing Capacity”(in German), Arch. Eisenhuettenwes., 55(8), 359–364 (1984) (Experimental, Thermo-dyn., 21)

[1986Gel] Gel’d, P.V., Povzner, A.A., Kortov, S.V., Safonov, V.N., “Band Magnetism of the Fe1–xCoxSiSolid Solutions” (in Russian), Dokl. Akad. Nauk SSSR, 289(2), 351–354 (1986) (Experimen-tal, 8)

[1986Mat] Matysina, Z.A., Milyan, M.I., “Ordering in Ternary Magntic Alloys”, Phys. Met. Metallogr.(Engl. Transl.), 62(4), 44–50 (1986) (Calculation, Magn. Prop., 13)

[1987Gel] Gel’d, P.V., Povzner, A.A., Kortov, S.V., Krentsis, R.P., “Thermal Expansion and Weak Itin-erant Magnetism in Fe1–yMnySi and Fe1–xCoxSi Solid Solutions”, Sov. Phys. -Dokl., 32(12),1006–1008 (1987), translated from Dokl. Akad. Nauk SSSR, 297, 1359–1363 (1987)(Experimental, Thermodyn., 10)

[1988Ray] Raynor, G.V., Rivlin, V.G., “Co-Fe-Si” in “Phase Equilibria in Iron Ternary Alloys”, Inst.Metals, London, 406, 256–267 (1988) (Phase Diagram, Phase Relations, Review, 15)

[1989Koz] Kozakai, T., Zhao, P.Z., Miyazaki, T., “Phase Separations in Fe Rich Fe Base Ternary Order-ing Alloy Systems”, Met. Abstr. Light Metals and Alloys, 23, 32–33 (1989/1990) (Abstract,Crys. Structure, Experimental, Phase Diagram, 0)

[1990Fuk] Fukaya, M., Miyazaki, T., Pi Zhi Zhao, Kozakai, T., “A Statistical Evaluation of the FreeEnergy of Fe Base Ternary Ordering Alloys”, J. Mater. Sci., 25(1B), 522–528 (1990) (Calcu-lation, Magn. Prop., Phase Diagram, Phase Relations, Thermodyn., 42)

[1990Ish] Ishimoto, K., Yamauchi, H., Yamaguchi, Y., Suzuki, J., Arai, M., Furusaka, M., Endoh, Y.,“Anomalous Region in the Magnetic Phase Diagram of (Fe,Co)Si”, J. Magn. Magn. Mater.,90–91, 163–165 (1990) (Abstract, Magn. Prop., 10)

[1990Koz] Kozakai, T., Zhao, P.Z., Miyazaki, T., “Phase Separations in Fe Rich Fe Base Ternary Order-ing Alloy Systems”, Met. Abstr. Light Met. Alloys, 23 (1989-1990), Osaka, 23, 32–33 (1990)(Assessment, Experimental, Phase Diagram, 0)

[1991Fuk] Fukaya, M., Miyazaki, T., Kozakai, T., “Phase Diagrams Calculated for Fe rich Fe-Si-Co andFe-Si-Al Ordering Systems”, J. Mater. Sci., 26(2), 5420–5426 (1991) (Calculation, PhaseDiagram, 42)

[1991Mot] Motta, N., Christensen, N.E., “Phase Diagram of the Fe1–xCoxSi2 Alloy in the FluoriteForm”, Phys. Rev. B, 43(6), 4902–4907 (1991) (Crys. Structure, Experimental, Phase Dia-gram, 17)

[1991Nis] Nishizawa, T., Ishida, K., Ohtani, H., Kami, C., Suwa, M., “Experimental Study on Interac-tion Parameter for Carbon and Alloying Elements in Austenite and Ferrite”, Scand. J. Metall.,20, 62–71 (1991) (Calculation, Experimental, Phase Diagram, Phase Relations, Thermo-dyn., 34)

[1994Koz] Kozakai, T., Miyazaki, T., “Experimental and Theoretical Investigations on Phase Diagramsof Fe Base Ternary Ordering Alloys”, ISIJ Int., 34(5), 373–383 (1994) (Calculation, Magn.Prop., Phase Diagram, 18)

24 Co–Fe–Si


[1994Rag] Raghavan, V., “Co-Fe-Si (Cobalt-Iron-Silicon)”, J. Phase Equilib., 15(5), 527–528 (1994)(Phase Diagram, Review, 7)

[1996Mor] Moriya, T., Ukai, K., Kato, K., Kimura, M., Yazaki, K., Isokane, Y., Miyazaki, T., Kozakai,T., Koyama, T., “Moessbauer Study on the Phase Separation of Fe-Co-Si Alloys”, Mater.Trans. JIM, 37(5), 965–969 (1996) (Experimental, Electronic Structure, Phase Relations, 11)

[1998Lan] Landrum, G.A., Hoffmann, R., Evers, J., Boysen, H., “The TiNiSi Family of Compounds:Structure and Bonding”, Inorg. Chem., 37(22), 5754–5763 (1998) (Crys. Structure, Experi-mental, 34)

[1998Sch] Schneeweiss, O., Zak, T., Jiraskova, Y., Havlicek, S., Solyom, A., Marko, P., “Spin Textureand Magnetic Properties of Fe-Si-Co Alloys”, Acta Phys. Slovaca, 48(6), 707–710 (1998)(Crys. Structure, Experimental, Magn. Prop., 7)

[2000Bel] Belyaev, E., Mamylov, S., Lomovsky, O., “Mechanochemical Synthesis and Properties ofThermoelectric Material β-FeSi2”, J. Mater. Sci., 35, 2029–2035 (2000) (Crys. Structure,Electr. Prop., Experimental, Phase Relations, 18)

[2002Aru] Arushanov, E., Ivanenko, L., “Magnetic Properties of Undoped and Co-Doped n-Typeβ-FeSi2.5 Single Crystals”, J. Mater. Res., 17(11), 2960–2965 (2002) (Experimental, Magn.Prop., 28)

[2002Cha1] Chattopadhyay, M.K., Roy, S.B., Chaudhary, S., “Magnetic Properties of Fe1–xCoxSiAlloys”, Phys. Rev. B: Condens. Matter, 65(13), 132409-1–4 (2002) (Experimental, Magn.Prop., Semiconduct., 17)

[2002Cha2] Chattopadhyay, M. K., Roy, S. B., Chaudhary, S., Singh, K. J., Nigam, A. K., “MagneticResponse of Fe(1–x)CoxSi Alloys: A Detailed Study of Magnetization and Magnetoresis-tance”, Phys. Rev. B, 66(17), 174421_1-174421_7 (2002) (Experimental, Magn. Prop., Ther-modyn., 57)

[2002Ers] Ersen, O., Ulhaq-Bouillet, C., Pierron-Bohnes, V., Tuilier, M.H., Berling, D., Bertoncini, D.,Pirri, C., Gailhanou, M., Thiaudiere, D., “Evidence of a Ternary Co(1–x)FexSi2 Phase with aCaF2-Type Structure: High-Resolution Transmission Electron Microscopy and DiffractionAnomalous Fine Structure Study”, Appl. Phys. Lett., 81(13), 2346–2348 (2002) (Crys. Struc-ture, Experimental, 16)

[2002Fet] Fetzer, C., Dezsi, I., Vantomme, A., Wu, M.F., Jin, S., Bender, H., “Ternary CoxFe(1–x)Si2 andNixFe(1–x)Si2 Formed by Ion Implantation in Silicon”, J. Appl. Phys., 92(7), 3688–3693(2002) (Crys. Structure, Electronic Structure, Experimental, 27)

[2002Tan] Tani, J.-I., Kido, H., “First Princple Calculation of the Geometrical and Electronic Structureof Impurity-Doped β-FeSi2 Semiconductors”, J. Solid State Chem., 163, 248–252 (2002)(Calculation, Crys. Structure, Phys. Prop., 23)

[2002Ur] Ur, S.-C., Kim, I.-H., “Phase Transformation and Thermoelectric Properties of n-TypeFe0.98Co0.02Si2 Processed by Mechanical Alloying”, Mater. Lett., 57(3), 543–551 (2002)(Crys. Structure, Experimental, Morphology, Phase Relations, Transport Phenomena, 22)

[2003Bol] Boldt, R., Reichelt, W., Bosholm, O., Oppermann, H., “Investigation on Chemical VapourTransport of Intermetallic Phases in the System Co/Fe/Si”, Z. Anorg. Allg. Chem., 629(10),1839–1845 (2003) (Experimental, Thermodyn., 12)

[2003Cho] Cho, S.J., Krist, T., Mezei, F., “Determination of Interface Growth With Atomic Resolutionin FeCo-Si Multilayers”, Thin Solid Films, 434(1-2), 136–144 (2003) (Experimental, Inter-face Phenomena, 8)

[2003Ito] Ito, M., Tada, T., Katsuyama, S., “Thermoelectric Properties of Fe0.98Co0.02Si2 with ZrO2 andRare-Earth Oxide Dispersion by Mechanical Alloying”, J. Alloys Compd., 350, 296–302(2003) (Crys. Structure, Electr. Prop., Experimental, Transport Phenomena, 11)

[2003Kim] Kim, S.W., Cho, M.K., Mishima, y., Choi, D.C., “High Temperature Thermoelectric Proper-ties of p- and n-Type β-FeSi2 with Some Dopants”, Intermetallics, 11(5), 399–405 (2003)(Electr. Prop., Experimental, Transport Phenomena, 15)

Co–Fe–Si 25



[2003Zhu] Zhu, T.-J., Zhao, X.B., Lu, L., “Preparation and Thermoelectric Properties of Melt-SpunFe2Si5 Based Alloys”, Mater. Sci. Forum, 437–438, 471–474 (2003) (Electr. Prop., Experi-mental, Phase Relations, 7)

[2006Ito] Ito, M., Tada, T., Hara, S., “Thermoelectric Properties of Hot-pressed β-FeSi2 with YttriaDispersion by Mechanical Alloying”, J. Alloys Compd., 408–412, 363–367 (2006) (Crys.Structure, Electr. Prop., Experimental, Transport Phenomena, 9)

[2006Wur] Wurmehl, S., Fecher, G.H., Ksenofontov, V., Casper, F., Stumm, U., Felser, C., Lin, H.-J.,Hwu, Y., “Half-metallic Ferromagnetism with High Magnetic Moment and High Curie Tem-perature in Co2FeSi”, J. Appl. Phys., 99(8J), 08J103 (2006) (Experimental, Magn. Prop.,Phase Relations, 15)

[Mas2] Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, MetalsPark, Ohio (1990)

[V-C2] Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for IntermetallicPhases, 2nd edition, ASM, Metals Park, Ohio (1991)

26 Co–Fe–Si


Cobalt – Iron – Vanadium

Andy Watson, Lesley Cornish

Introduction

Apart from interest relating to steels, the main impetus with regard to research in this system comes from themagnetic properties of Co-Fe alloys and the influence of V additions. A composition of Fe-30 at.% Co hashighest saturated magnetization (Bs) for any known binary alloy, but the equiatomic composition has higherpermeability and lower coercivity whilst Bs is not significantly reduced [1991Rez]. This material is used inmany soft magnetic applications. However, alloys of this composition become brittle at room temperatureowing to ordering at T < 730°C. Ternary additions such as V are made in an attempt to slow ordering reac-tion and hence improve the mechanical properties.The phase relationships in this system have been reviewed by [1949Jae, 1981Ray, 1983Ray, 1988Ray,2002Rag], and are based mainly on early studies by [1938Koe, 1952Mar, 1967Jos, 1955Koe1,1955Koe2]. [2005Sou] gives an excellent review of Fe-Co-X systems, not only from a phase equilibriumpoint of view, but also magnetic and mechanical properties.Details of studies of the system are given in Table 1.

Binary Systems

The Co-Fe system is accepted from a Calphad assessment given by [2002Ohn]. The phase diagram is essen-tially the same as given in [Mas2], but it extends to lower temperatures following phase equilibrium studiesusing thin films. The Co-V and Fe-V binary phase diagrams are taken from [Mas2]. However, [2006Oka]reported on recent work by [2005Ust] who studied the extent of the σ phase in the binary Fe-V system usingXRD and electron microscopy. From their results they suggest that the σ phase decomposes below about650°C with phase separation of the bcc phase at temperatures below this. However, in the phase diagramshown by [2006Oka], there would seem to be a narrow strip of single phase bcc between the σ phaseand the region of phase separation. As pointed out by [2006Oka], this would seem to be unlikely, and mostprobably, the σ phase would decompose eutectoidally to give Fe rich and V rich bcc phases, much like inthe Cr-Fe system. As this remains uncertain, the feature has been ignored in the present work.

Solid Phases

No ternary phases have been reported for this system. However, the solid solution phases of the binary sys-tems penetrate deep into the ternary, and at appropriate temperatures, form complete series of solid solutionsacross the system (γ-fcc, αFe-bcc and σ). The (αFe) phase in the Co-Fe orders at temperatures below 730°Cabout the equiatomic composition. According to combined neutron and X-ray powder diffraction studiesundertaken by [1990Wil], V will substitute for Co on the B2 lattice, as predicted by [2002Boz] usingBozzolo-Ferrante-Smith first principles analysis. The extension of the ordering reaction into the three com-ponent system was subject to a theoretical study by [1986Mat], which was recorded in the review of[2002Rag]. The addition of 3 at.% V to an equiatomic Fe-Co alloy is predicted to reduce the peak orderingtemperature by about 100°C, with a shift slightly towards the Fe-side. However, earlier work by [1976Hag]suggested a reduced effect on the critical temperature with V addition. Using DTA studies, they calculatedthe location of the order-disorder transition lines in the ternary system for temperatures between 720and 640°C (Fig. 1). In contrast to [1986Mat], they found a transition temperature of around 704°C for a3 at.% Vaddition to the equiatomic Fe-Co alloy. [1973Cle] studied the kinetics of the ordering process usingspecific heat, lattice parameter and magnetic saturation measurements. They found that at additions ofless than 2.5 at.% V to the equiatomic alloy, little or no change in the ordering kinetics was observed. Addi-tions in excess of 2.5 at.% V resulted in a retardation of the process. They also discovered that the criticaltemperature decreased 11°C per at.% V added, which is more in agreement with the later work of[1976Hag]. This is somewhat in contrast to [1977Ale] who used electron diffraction and nuclear gamma

Co–Fe–V 1



resonance methods to find that Vaddition decreases the rate of ordering. [1985Nes] found the order-disorderkinetics to be extremely rapid, following neutron diffraction studies of a Fe-Co- 2 mass% V alloy. Theorder-disorder transformation is accompanied by a complex topological structure of interpenetrat-ing domains.The crystal structure and site occupancy of the VCo3 compound has been of much interest. [1963Zeg] con-cluded that virtually no Fe could substitute for Co in the compound. [1969Sin] also studied the crystal struc-ture of VCo3 and how this changes with the addition of Fe to the lattice, following the work of [1966Vuc].Arc-melted samples were homogenized in sealed argon filled ampoules at temperatures between 750 and1200°C for between 3 and 10 days. XRD patterns were produced using a Guinier-de Wolff or a Debye-Scherrer camera. Lattice parameters and stacking sequence were measured for two alloys, V(Fe0.3Co0.7)3and V(Fe0.1Co0.9)3. It is not clear from the article what the annealing temperature was for these two samples,but the results suggest different crystal structures; AuCu3 type and VCo3 type, respectively. This reflectsfeatures in the Co-V binary system where the VCo3 compound undergoes an allotropic transformation froma cubic AuCu3 type to hexagonal at around 1025°C [Mas2], with decreasing temperature. The work of[1969Sin] would tend to suggest that the AuCu3 structure is stabilized by the addition of Fe. Long rangeorder parameters of FeCo as a function of V content were determined by [1975Mal] using neutron scatter-ing. Their conclusion was that V interacts more strongly with Co than Fe. More recently, [1991Yao1,1991Yao2] studied the microalloying of this compound and found that the ordered structure of the alloycan be controlled by adjusting the electron density (e/a) of the compound. They found that the L12 structurewas stable when the e/a ratio is in the region of 7.625-7.850.Crystallographic details of the phases in the system are given in Table 2.


The addition of V to the Co-Fe system lowers the peritectic temperature for the reaction L + (δFe)⇌ (γFe).The liquidus line extends into the ternary system meeting liquidus lines from the Co-V system emanatingfrom the L⇌ (αCo) + σ and L + (V)⇌ σ reactions in a ternary transition reaction, L + α⇌ σ + γ; denotedas U1 in Figs. 2 and 3. However, the nature of the monovariant changes as it moves into the ternary system;from peritectic to eutectic, the switch occurring at about 20 mass% V, 42 mass% Co and a temperature ofabout 1400°C. Compositions of the phases taking part in the ternary invariant reaction are given in Table 3,and Fig. 2 gives the reaction scheme for the ternary system; based on the review of the system by[1988Ray]. Some amendments have been made to ensure compatibility with the accepted binary phase dia-grams.


The only study of the liquidus was carried out by [1955Koe1]. Samples were prepared from Armco Fe,Würfel Co (99.5% Co) and ferrovanadium alloys with compositions 91.7V-7.1Fe-1.14Si-0.014Aland 79.8V-18.7Fe-1.37Si-0.1Al, by sintering in corundum crucibles under either argon or hydrogen in aTamman furnace. Thermal analysis was used to determine the liquidus surface for alloys containing lessthan 60% V. These data were reviewed by [1988Ray] and a liquidus projection was postulated. Figure 3shows the liquidus surface presented by [1988Ray], comprising liquidus lines (denoted by double arrows),solidus lines (denoted by single arrows) and isotherms (denoted by dashed lines). Slight amendments havebeen made to ensure that the liquidus and solidus lines are compatible with the accepted phase diagrams ofthe binary systems. By plotting the liquidus lines and the solidus lines on the same diagram, it is possible tosee how the nature of the L + (δFe) ⇌ γ peritectic reaction changes to eutectic as it extends into the ternarysystem, where the liquidus line crosses over the γ solidus line. For clarity, Fig. 4 shows a magnified view ofthe transition reaction, U1, taken from [1988Ray]. The four phase plane representing the ternary invariantreaction L + α ⇌ σ + γ is shown as a trapezoid at the centre of the figure.

Isothermal Sections

[1955Koe1, 1955Koe2] also studied solid state equilibria in the ternary system. Alloys, prepared in a similarmanner to those for investigation of the liquidus surface were firstly homogenized for 2 d at 1200°C. Optical

2 Co–Fe–V


microscopy, dilatometry and XRD were used enabling partial isothermal sections to be drawn for asequence of temperatures of 1300, 1200, 1000, and 600°C (Figs. 5-7 and 11). The main feature of the partialsection at 1300°C (Fig. 5) is in the region around the transition reaction composition. At this temperature,which is below that of U1, there can be seen the two three-phase regions, L+γ+σ and α+γ+σ, which atthe transition reaction temperature, meet to form the four phase plane. Between these three-phase regionsis the two-phase region, γ+σ. With decreasing temperature, the two three-phase regions move apart untilthe liquid phase is no longer stable, leaving just one of the three-phase regions remaining; Fig. 6, 1200°C. Below 1025°C, the γ’ becomes stable resulting in a new three-phase equilibrium, between γ’+α+σ,Fig. 7 1000°C. γf and γp refer to ferromagnetic and paramagnetic γ, respectively. At low temperatures,the section becomes more complex. Figure 11 shows the equilibria for 600°C. This section is the resultof annealing specimens for very long times (~1 year) and is characterized by a large three-phase region,α+γ+γ(r), the γ’ having undergone a transformation to the hexagonal form. This three-phase region, alongwith that corresponding γ(r)+α+σ is the product of a solid state transition reaction that must take placebetween 600 and 1000°C. This reaction, γ+ σ⇌ α+ γ(r) was given at a temperature of 860°C by[1955Koe2, 1988Ray]. Also shown in this section is the ordered α’ phase, which begins to extend intothe ternary system from the Fe-V binary. Equilibria for V rich alloys have been added as the V3Co phasebecomes stable below 1025°C, but the phase boundaries are shown as dashed lines owing to their tentativenature. It should be noted that the γ’ phase is not stable at this temperature, according to [Mas2]. Thisphase, given in the original work, has been amended to the low temperature hexagonal phase, γ(r) here.[1974Ben1, 1974Ben2] studied samples with compositions of less than 30 at.% V at temperatures of 950,925 and 900°C. Samples were annealed at appropriate temperatures for 6 h before quenching into iced-brine. Metallography and microprobe analysis were used for phase analysis. Figures 8 to 10 show partialisothermal sections for these temperatures taken from the review of [1988Ray], which are based on theseresults along with data taken from isopleths produced by [1955Koe1, 1955Koe2]. However, amendmentshave been made to the Fe rich region of the diagram, where two three-phase equilibria have been added,shown by dashed lines. In [1988Ray], this region was shown as one two-phase region of α + γ. At this tem-perature, there must be two three-phase equilibria with an α + α region in between. The two three-phasefields will appear at 985°C, the temperature of the congruent transformation of α ⇌ γ in the Co-Fe binarysystem. In the partial isothermal section for 925°C shown in Fig. 9, as the Fe rich γ region has shrunk quitedramatically, the three-phase triangle representing the equilibrium between the α and γ phases richer in Fewill have moved right to the binary edge of the section. At 900°C, this three-phase field has now disap-peared, the temperature being below the α ⇌ γ transition temperature for pure Fe.The extent of the σ phase in this system at 1200°C was investigated by [1957Dar] using XRD and micro-structural examination of samples prepared by vacuum induction melting and long time annealing. Only thehomogeneity region of the σ phase was studied in any detail, but the results were in very good agreementwith the work of [1955Koe1, 1955Koe2].Figure 12 shows a metastable isothermal section for 600°C taken from the work of [1955Koe2], whichrepresents the phase constitution for alloys produced from furnace cooling conditions. The first thing to noteis that the equilibrium hexagonal variant of VCo3 (γ(r)) is not seen, rather the ordered fcc variant, γ’. Theγ + γ’ field extends appreciably into the three component system; much further than the γ + γ(r) field doesin the equilibrium diagram (Fig. 11) owing to the persistence of non-equilibrium γ phase to much higher Fecontents, the limit given by the composition γi. Therefore, V poor alloys have a high proportion of non-equilibrium γ in the structure. Alloys less rich in V than the dotted line γi - γh undergo the γ-α transforma-tion on heating and cooling, and are therefore subject to considerable hysteresis. Alloys with the γ structureat high temperatures would tend to retain the phase on cooling. The dashed lines correspond to the Curietemperature contours.

Temperature - Composition Sections

Vertical sections for the system were produced by [1955Koe1, 1955Koe2], at 35, 40, 52 and 60 mass% Co,and at 5, 10, 15 and 20 mass% V. A section was also prepared for V40Fe60 - Co. These are reproduced inFigs. 13 to 22. The f and p subscripts refer to ferromagnetic and paramagnetic variants of the phases, respec-tively. Amendments have been made where necessary to ensure agreement with the accepted binary phase

Co–Fe–V 3



diagrams. There is considerable uncertainty over the nature of the γ phase in the sections. In the originalwork, only the disordered and the ordered phases were considered (γ and γ’), whereas, in considerationof the accepted binary phase diagram, it is likely that many of the γ’ phase labels actually refer to the hex-agonal γ(r). Although some of these have been amended in the figures they should still be treated with anelement of caution as it is unclear where the transition between the two variants will take place within theternary system.

Thermodynamics

Surprisingly, there is no complete Calphad assessment of this system available in the literature. [1989Kau]calculated 6 isotherms for the ternary system by combining available thermodynamic descriptions for thebinary systems. Some of the modeling was quite primitive by present standards, particularly with respectto the σ phase, which was modeled as a substitutional solid solution as opposed to the compound energymodel that is used currently. Nevertheless, the calculated and experimental isothermal section for 900°Cwere in reasonable qualitative agreement. Sadly, the same could not be said for the sections at 600°C.


Co-Fe-V alloys, particularly those based on equiatomic Co-Fe, are used extensively because of their mag-netic properties. Low V additions result in a material with high saturation and high permeability character-istics [1952Mar]. Permendur is based on 50/50 FeCo with a 2 mass% V addition. A material with a similarcomposition is Remendur, which has been used for some time as a remanent reed sealed contact for thecommunications industry. The fabrication process for this material, also often referred to as FeCo-2V hasa direct effect on its properties, which depend on the precipitation of ferromagnetic phases. Commercial pro-duction involves hot rolling at a temperature between 1100 and 1200°C followed by air cooling. The mate-rial is then reheated to 900 - 950°C for around 0.5 h before quenching into iced brine. Subsequent aginggives the material the appropriate phase constitution. [1998Ust] found that the ferromagnetic B2 phase(α’) is formed only during the aging phase, and only in air (not vacuum) and at the surface of the material.Vicalloy I and II are alloys with a higher proportion of V. Vicalloy I has 9.5 mass% V giving a material sui-table for permanent magnet applications, with HC of 2388 A/m, Br of 0.9 T and BHmax of 8·10

3 Jm–3 fol-lowing an aging treatment of 2 h at 600°C [1952Mar]. Vicalloy II has a higher V content (13 mass%) andhas HC of 3.9·104 A/m, Br of 1 T and BHmax of 2.8·10

4 Jm–3 following a similar aging treatment. As goodas these materials are, their application is restricted owing to temperature sensitivity and poor mechanicalproperties. In order to address these problems, research has been conducted into producing a material withthe excellent magnetic properties of FeCo-2V but with improved mechanical properties and operating tem-perature range, particularly for aerospace applications. An alternative approach in the development of softmagnetic materials is through microstructural control. Ultra fine grained material has been studied by[2003Duc]. Grain sizes studied were 100, 150 and 290 nm, which was produced from cold rolled material(93%) that was then annealed at 438°C for 5 h (100 nm), 600°C for 1 h (150 nm) and 650°C for 1 h(290 nm). Tensile tests were performed on samples of the materials at temperatures from room temperatureto 500°C. Yield strengths of up to 2.1 GPa were measured at room temperature with ductility between 3 and13%. Strength was found to decline gradually with increasing temperature, with ductility increasing up to22%. [2004Sun1] reviews recent work on alloying Co-Fe-V alloys with a 4th component such as Nior Mo in order to improve the physical properties of the material without significant loss in magnetic qual-ity. The age hardening behavior of a Fe-40Co-5V-0.005B-0.015C-0.5Mo-0.5Nb (at.%) alloy was investi-gated by [2004Sun2, 2005Sun2]. Maximum hardening and coercivity was observed when the materialwas aged at 600°C, giving an electrical resistivity of 70-75 μΩcm. The improvement in the mechanicaland magnetic properties was attributed to the precipitation of a fine dispersion of ordered paramagneticγ’ phase.The use of amorphous and nanocrystalline materials and also thin films is becoming increasingly importantand recent developments have been reviewed by [1999Mch, 2005Sun1].The recovery and recrystallization of on ordered V(Fe22Co78)3 alloy was investigated by [1991Cah]. Thematerial was cold rolled to 25-50% reduction and annealed at various temperatures above and below the

4 Co–Fe–V


critical temperature for ordering. The recovery and recrystallization process was monitored by hardness andtensile testing, and by optical microscopy and TEM. It was found that annealing at temperatures below thecritical temperature the recovery process was severely retarded. Annealing at temperatures around the cri-tical temperature, the recrystallization rate was 300 times slower in the ordered region than in the disorderedregion. Above the critical temperature the recrystallization was very rapid. The retarded recrystallization inthe ordered region was attributed to recovery softening of the unrecrystallized material. At temperatures clo-ser to the critical temperature, some of this recovery softening was replaced by recovery hardening.The oxidation behavior of a FeCo-2V (at.%) alloy was studied by [2004Sun3]. Alloys were prepared by arcmelting, followed by a homogenization treatment for 1.5 h at 1100°C. The material was then forged beforesamples were cut for isothermal oxidation experiments at 500, 550 and 600°C. It was found that the oxida-tion kinetics followed a parabolic rate law. Microstructural and XRD analysis revealed the formation of anFe rich outer oxide layer together with an inner solute rich layer containing oxides rich in Vand Co. The Ferich layer is semiconducting and so may not have sufficient insulating properties to minimize eddy currentlosses during AC applications.Details of materials property studies are given in Table 4.

Miscellaneous

[1971Sht] studied the effect of pressure on phase transformation. By applying high isostatic pressure, theyexpected that transformation involving volume change to be suppressed. They used a Vicalloy specimen(Fe-52Co-12V-0.1C mass%). Samples were annealed at 1000°C before being drawn in a high pressureliquid. By adjusting the degree of deformation and the pressure in the liquid, it was found that applying apressure of 2 GPa reduces the amount of ferromagnetic α phase precipitated following tempering, as deter-mined by coercive force measurement, by 3-4 times.[1977Ash] studied the aging characteristics of an equiatomic Fe-Co alloy with 2 mass% V. Specimens fromcold-rolled sheet were heat treated at 850°C for 10ks before quenching in iced brine. Neutron diffractionstudies revealed a microstructure of metastable disordered α phase. Samples were then encapsulated inargon filled capsules and aged at temperatures between 477 and 627°C. On aging, the disordered α phaseorders to the α’ phase, with an ordered γ’ fcc phase (L12) precipitating preferentially on antiphase grainboundaries. It was found that material deformed by 25-50% before aging exhibited more rapid precipitationof the γ’ phase.Permendur and lead zirconate titanate (PZT) have been combined in a bilayer to produce a structure that hasmagnetoelectric properties [2004Lal]. These materials were chosen because of the low resistivity, high mag-netization (2.34 T), Curie temperature (940°C), permeability and magnetostriction (7·10–5) of the Permen-dur combined with the high ferroelectric Curie temperature and piezoelectric coupling constant of the PZT.Giant magnetoelectric interactions were observed at electromechanical resonance suggesting very high fieldconversion efficiency.[1977Bra] successfully predicted the occurrence and extent of the σ phase in Cr-Fe-V using a cluster model.

Table 1. Investigations of the Co-Fe-V Phase Relations, Structures and Thermodynamics

Reference Method/Experimental Technique Temperature/Composition/Phase RangeStudied

[1938Koe,1955Koe1,1955Koe2]

Thermal analysis, dilatometry, magneticproperty, microscopy and XRD

Temperatures up to 1600°C, V contents< 60 mass%.

[1952Mar] XRD, thermal analysis, hardnessmeasurement, magnetic propertymeasurement

Equilibria at V < 20 mass%

[1954Bae] XRD, dilatometry, electrical resistance Phase equilibria up to 4 mass% V

(continued)

Co–Fe–V 5




[1957Dar] XRD, metallography Homogeneity range of the σ phase at 1200°C

[1963Zeg] XRD Dissolution of Fe into VCo3

[1967Jos] Dilatometry Phase transformations in Fe-52-Co-8V(mass%)

[1969Sin] XRD - Guinier-de Wolff or Debye-Scherrer camera

VCo3 compound with the addition of Fe

[1973Cle] Specific heat, lattice parameter andmagnetic saturation

Kinetics and critical temperature for orderingin Co-Fe equiatomic alloys + V

[1974Ben1,1974Ben2]

Metallography, microprobe Samples with compositions of less than 30 at.% V at temperatures of 950, 925 and 900°C

[1975Mal] Neutron scattering Long range order parameter of FeCo as afunction of added V

[1976Hag] DTA Order-disorder transition of Co-Fe equiatomicalloys + V

[1977Ale] Electron diffraction, nuclear gammaresonance

Order-disorder kinetics in Co-Fe equiatomicalloys + V

[1977Ash] TEM, XRD Aging of Co-Fe equiatomic alloys + 2 mass%V

[1985Nes] Neutron diffraction Order-disorder kinetics of Fe-Co-2 mass% V

[1986Mat] Theoretical study of the mutual influenceof atomic ordering and magnetic moment

Order-disorder transition of Co-Fe equiatomicalloys + V

[1989Kau] Calphad technique - combination ofthermodynamic descriptions of binarysystems

Calculated isothermal sections for 1527, 1427,1327, 1200, 900 and 600°C

[1990Wil] Combined neutron/X-ray powderdiffraction studies

Site occupancy of V in Fe0.5Co0.48V0.02

[1991Rez] XRD, Mössbauer spectroscopy Dependence of lattice parameter with Vcontent in FeCo

[2002Boz] Bozzolo-Ferrante-Smith first principlesanalysis

Site occupancy in ordered Co-Fe equiatomicalloys + V

6 Co–Fe–V



Phase/Temperature Range[°C]

Pearson Symbol/Space Group/Prototype

Lattice Parameters[pm]

Comments/References

(εCo) hP2P63/mmcMg

a = 250.71c = 406.86

at 25°C [Mas2].Dissolves ~2 at.% V at 238°C [2002Ohn].

γ,(γFe, αCo)

(αCo)1495 - 422(γFe)1394 - 912

cF4Fm�3mCu a = 354.47

a = 364.67

[V-C2, Mas2]

[V-C2, Mas2] at 915°C

α,(δαFe,V)

(δFe)1538 - 1394(αFe)< 912(V)< 1910

cI2Im�3mW a = 293.15

a = 286.65

a = 302.40

1394°C [Mas2]Dissolves up to 17 at.% Co at 1499°C.at 25°C [Mas2]

at 25°C [Mas2]. Dissolves up to 22 at.%Co at 1422°C.

α‘< 730

cP2Pm�3mCsCl

28 to 75 at.% Fe in the Co-Fe system at500°C [Mas2]. Critical temperature of730°C in the Co-Fe binary, decreasing11°C for every 1 at.% V added [1973Cle].

σσVCo< 1422

σVFe< 1252

tP30P42/mnmσCrFe a = 883.4

c = 458.65

a = 896.5c = 463.3

at V0.62Co0.38 [V-C2].Contains 84 at.% V at 1025°C and 45 at.%V at 1248°C [Mas2].at V0.5Fe0.5 [V-C2].29.6 - 60.1 at.% V [Mas2].

γ’,VCo31070 - ~1025

V(Fe0.3Co0.7)3

cP4Pm�3mAuCu3

a = 356

a = 357.9

[V-C2], [1977Ash]

14 - 31 at.% V [Mas2].[1969Sin]

γr,VCo3 (hex)< 1025

V(Fe0.1Co0.9)3

hP24P63/mmcPuAl3

a = 503.8c = 1229

a = 503.4c = 1228

[V-C2]

14 - 31 at.% V [Mas2]. [1969Sin]

V3Co< 1025

cP8Pm�3nCr3Si

a = 467.6 [V-C2]

Co–Fe–V 7




Reaction T [°C] Type Phase Composition (at.%/mass%)

Co Fe V

L + α ⇌ σ + γ 1330 U1 Lασγ

46.2/48.944.0/46.643.9/46.547.4/50

22.1/22.124.3/24.423.3/23.422.4/22.4

31.7/2931.7/2932.8/30.130.2/27.6

σ + γ ⇌ α + γ(r) 860 U2 - - -

Table 4. Investigations of the Co-Fe-V Materials Properties

Reference Method / Experimental Technique Type of Property

[1971Sht] Coercive force measurement Amount of ferromagnetic phase (α) precipitated onaging as a function of applied pressure.

[1991Cah] Tensile and hardness testing.Optical microscopy and TEMstudies

Recrystallization behavior of deformed V(Fe22Co78)3.

[1991Yao1,1991Yao2]

Tensile test V(Fe,Co)3, ductility as a function of ordering.

[2003Duc] Tensile test Strength and ductility of ultra fine grain sized V(FeCo)3 with temperature.

[2004Sun2,2005Sun2]

Tensile test, hardness test, VSM Age hardening behavior and magnetic properties ofFe-40Co-5V-0.005B-0.015C-0.5Mo-0.5Nb (at.%)alloy

[2004Sun3] Isothermal heat treatment, XRD,SEM, EDX

Oxidation kinetics of FeCo-2V (at.%)

8 Co–Fe–V


Fig. 1. Co-Fe-V. Calculated order-disorder transition temperatures

Co–Fe–V 9



Fig.2.

Co-Fe-V.Reactionscheme

10 Co–Fe–V


Fig. 3. Co-Fe-V. Liquidus surface projection showing solidus lines

Co–Fe–V 11



Fig. 4. Co-Fe-V. Liquidus and solidus lines in the region of the invariant reaction

12 Co–Fe–V


Fig. 5. Co-Fe-V. Isothermal section at 1300°C

Co–Fe–V 13




14 Co–Fe–V



Co–Fe–V 15




16 Co–Fe–V



Co–Fe–V 17




18 Co–Fe–V



Co–Fe–V 19



Fig. 12. Co-Fe-V. Metastable isothermal section at 600°C

20 Co–Fe–V


Fig. 13. Co-Fe-V. Vertical section at 35 mass% Co

Co–Fe–V 21




22 Co–Fe–V



Co–Fe–V 23




24 Co–Fe–V


Fig. 17. Co-Fe-V. Vertical section at 5 mass% V

Co–Fe–V 25




26 Co–Fe–V



Co–Fe–V 27




28 Co–Fe–V


Fig. 21. Co-Fe-V. Vertical section VFe-Co

Co–Fe–V 29



Fig. 22. Co-Fe-V. Vertical section V40Fe60-Co

30 Co–Fe–V


References[1938Koe] Koester, W., Lang, K., “Cobalt Corner of the Iron-Cobalt-Vanadium System” (in German),

Z. Metallkd., 30, 350–352 (1938) (Phase Diagram, Phase Relations, Experimental, 4)[1949Jae] Jaenecke, E., “Co–Fe–V” (in German) in “Kurzgefasstes Handbuch aller Legierungen”,

Winter Verlag, Heidelberg, 634 (1949) (Phase Diagram, Phase Relations, Review, 1)[1952Mar] Martin, D.L., Geisler, A.H., “Constitution and Properties of Cobalt-Iron-Vanadium Alloys”,

Trans. Amer. Soc. Metals, 44, 461–483 (1952) (Crys. Structure, Experimental, Magn. Prop.,Morphology, Phase Diagram, Phase Relations, 7)

[1954Bae] Baer, G., Thomas, H., “Properties and Structure of the Iron-Cobalt Alloys with 50% Co andsmall Vanadium Content” (in German), Z. Metallkd., 45(11), 651–655 (1954) (Electr. Prop.,Experimental, Morphology, Phase Diagram, Phase Relations, 11)

[1955Koe1] Koester, W., Schmid, H., “Iron-Cobalt-Vanadium Ternary System. Part I” (in German), Arch.Eisenhuettenwes., 26(6), 345–353 (1955) (Phase Diagram, Experimental, Phase Relations, #,*, 26)

[1955Koe2] Koester, W., Schmid, H., “Iron-Cobalt-Vanadium Ternary System. Part II” (in German),Arch. Eisenhuettenwes., 26(7), 421–425 (1955) (Phase Diagram, Experimental, Phase Rela-tions, #, *, 7)

[1957Dar] Darby, J.B., Beck, P.A., “Sigma Phase in Certain Ternary Systems with Vanadium”, Trans.AIME, 209, 69–72 (1957) (Experimental, Phase Diagram, Phase Relations, 8)

[1963Zeg] Zegler, S.T., Downey, J.W., “Ternary Cr3O-Type Phases with Vanadium”, Trans. Met. Soc.AIME, 227, 1407–1411 (1963) (Phase Diagram, Phase Relations, 16)

[1966Vuc] van Vucht, J.H.N., “Influence of Radius Ratio on the Structure of Intermetallic Compoundsof the AB3 Type”, J. Less-Common Met., 11, 308–322 (1966) (Crys. Structure, Experimen-tal, 21)

[1967Jos] Josso, E., “Analysis of the Polymorphic Transformation in Iron Alloys with 52 % Co and 8 %V” (in French), Mem. Sci. Rev. Metall., 64(12), 1045–1051 (1967) (Phase Relations, Experi-mental, 7)

[1969Sin] Sinha, A.K., “Close-Packed Ordered AB3 Structures in Ternary Alloys of Certain TransitionMetals”, Trans. Met. Soc. AIME, 245, 911–917 (1969) (Crys. Structure, Experimental, 16)

[1971Sht] Shtolts, Ye.V., Yeshchenko, R.N., “Influence of Hydrostatic Compression on Phase Transfor-mation of an Iron-Cobalt-Vanadium Alloy During Deformation”, Phys. Met. Metallogr.(Engl. Transl.), 32(4), 207–209 (1971), translated from Fiz. Met. Metalloved., 32(4),876–878 (1971) (Experimental, Thermodyn., 6)

[1973Cle] Clegg, D.W., Buckley, R.A., “The Disorder-Order Transformation in Iron-Cobalt-BasedAlloys”, Met. Sci. Journ., 7, 48–54 (1973) (Crys. Structure, Experimental, 28)

[1974Ben1] Bennett, J.E., Pinnel, M.R., “Aspects of Phase Equilibria in Fe/Co/2.5 to 3.0% V Alloys”,J. Mater. Sci., 9(7), 1083–1090 (1974) (Experimental, Morphology, Phase Relations, 13)

[1974Ben2] Bennett, J.E., Pinnel, M.R., “Equilibrium Phases and Transformations in Fe/Co/2-3% VAlloys.”,Microstruct. Sci., 2, 29–33 (1974) (Experimental, Morphology, Phase Relations, 3)

[1975Mal] Mal’tsev, Ye.I., Goman’kov, V.I., Mokhov, B.H., Puzey, I.M., Nogin, N.I., “Influence ofAlloying Elements on the Fe-Co Superlattice”, Phys. Met. Metallogr., 40(2), 190–193(1975), translated from Fiz. Metal. Metalloved., 40(2), 443–445 (1975) (Crys. Structure,Experimental, 11)

[1976Hag] Hagiwara, M., Suzuki, T., “The Effect of the Addition of a Third Element (Cr, Mn, V) on theOrder-Disorder Transition Temperature in FeCo” (in Japanese), J. Japan. Inst. Metals, 40(7),738–743 (1976) (Crys. Structure, Phase Relations, Calculation, Experimental, 34)

[1977Ale] Alekseev, L.A., Dzhavadov, D.M., Tyapkin, Yu.D., Levi, R.B., “Investigation of the Influ-ence of Vanadium on the Structure of Fe-Co Alloys by Electron Microscopy and NGR Meth-ods” (in Russian), Fiz. Met. Metalloved., 43(6), 1235–1244 (1977) (Crys. Structure,Experimental, Morphology, 7)

[1977Ash] Ashby, J.A., Flower, H.M., Rawlings, R.D., “Gamma Phase in an Fe-Co-2% VAlloy”,MetalSci., 11, 91–96 (1977) (Crys. Structure, Experimental, Morphology, Phase Diagram, PhaseRelations, 20)

Co–Fe–V 31



[1977Bra] Brauwers, M., “Occurrence of the Sigma Phase Computed from a Cluster Model”, J. Phys. F:Met. Phys., 7(6), 921–927 (1977) (Calculation, Phase Diagram, Phase Relations, Theory, 17)

[1981Ray] Raynor, G.V., Rivlin, V.G., “Critical Evaluation of Constitutions of Certain Ternary AlloysContaining Iron, Tungsten and a Third Metal”, Int. Met. Rev., 26, 213–249 (1981) (Crys.Structure, Phase Diagram, Phase Relations, Review, 43)

[1983Ray] Raynor, G.V., Rivlin, V.G., “Phase Equilibria in Iron Ternary Alloys. 10. Critical Evaluationof Constitution of Cobalt-Iron-Vanadium System”, Int. Met. Rev., 28(4), 211–227 (1983)(Phase Diagram, Phase Relations, Review, 29)

[1985Nes] Nesterenko, E.G., Zoteyev, O.E., Perekos, A.E., Chumachenko, V.K., “Study of Atom Order-ing Mechanism in Fe-Co-2%V Alloy by Neutronographic Methods” (in Russian), Akad.Nauk Ukr. SSR, Metallofizika, 7(4), 45–49 (1985) (Crys. Structure, Experimental, 13)

[1986Mat] Matysina, Z.A., Milyan, M.I., “Ordering in Ternary Magnetic Alloys”, Phys. Met. Metallogr.(Engl. Transl.), 62(4), 44–50 (1986) (Crys. Structure, Calculation, Magn. Prop., 13)

[1988Ray] Raynor, G.V., Rivlin, V.G., “Co-Fe-V” in “Phase Equilibria in Iron Ternary Alloys”,Inst. Metals, London, 406, 268–283 (1988) (Crys. Structure, Phase Diagram, Phase Rela-tions, Review, Thermodyn., 10)

[1989Kau] Kaufman, L., “Computer Bases Thermochemical Modeling of Multicomponent Phase Dia-grams” in “Alloys Phase Stability”, Stocks, G.M., Gonis, A. (Eds.), Kluwer Acad. Publ.,145–175 (1989) (Calculation, Phase Diagram, Phase Relations, Theory, 43)

[1990Wil] Williams, A., Kwei, G.H., Ortiz, A.T., Karnowski, M., Warburton, W.K., “Combined Neu-tron and X-ray Powder Diffraction Study of Fe0.50Co0.48V0.02”, J. Mater. Res., 5(6),1197–1200 (1990) (Crys. Structure, Experimental, 11)

[1991Cah] Cahn, R.W., Takeyama, M., Horton, J.A., Liu, C.T., “Recovery and Recrystallization of theDeformed, Orderable Alloy (Co78Fe22)3V”, J. Mater. Res., 6(1), 57–70 (1991) (Crys. Struc-ture, Experimental, 30)

[1991Rez] Rezende, M.F.S., Mansur, R.A., Pfannes, H.-D., Persiano, A.I.C., “Phase Characterization inthe Fe-Co-V and Fe-Co-Nb Systems”, Hyperfine Interact., 66(1-4), 319–324 (1991) (Crys.Structure, Experimental, Phase Relations, 14)

[1991Yao1] Yao, X., Chen, N., “Microstructures and Mechanical Properties of (Fe,Co,Ni)3V Alloys”,Mater. Mechan. Eng., 15(2), 27–31 (1991) (Crys. Structure, Magn. Prop., 10)

[1991Yao2] Yao, X., Kang, F., Chen, N., “The Ordering Process of (Co80Fe20)3V Intermetallic Com-pound”, Trans. Met. Heat Treatm., 12(4), 17–22 (1991) (Crys. Structure, 6)

[1998Ust] Ustinovshikov, Y., Tresheva, S., “Character of Transformations in Fe-Co System”, Mater.Sci. Eng. A, A248, 238–244 (1998) (Experimental, Phase Relations, 12)

[1999Mch] McHenry, M.E., Willard, M.A., Laughlin, D.E., “Amorphous and Nanocrystalline Materialsfor Applications as Soft Magnets”, Prog. Mater. Sci., 44(4), 291–433 (1999) (Crys. Struc-ture, Electr. Prop., Experimental, Magn. Prop., Phase Relations, Review, 302)

[2002Boz] Bozzolo, G.H., Noebe, R.D., Amador, C., “Site Occupancy of Ternary Additions to B2Alloys”, Intermetallics, 10, 149–159 (2002) (Crys. Structure, Calculation, Electronic Struc-ture, 27)

[2002Ohn] Ohnuma, I., Enoki, H., Ikeda, O., Kainuma, R., Ohtani, H., Sundman, B., Ishida, K., “PhaseEquilibria in the Fe-Co Binary System”, Acta Mater., 50, 379–393 (2002) (Assessment, Cal-culation, Experimental, Phase Relations, Thermodyn., 50)

[2002Rag] Raghavan, V., “Co-Fe-V (Cobalt-Iron-Vanadium)”, J. Phase Equilib., 23(5), 442 (2002)(Phase Relations, Review, 4)

[2003Duc] Duckham, A., Zhang, D.Z., Liang, D., Luzin, V., Cammarata, R.C., Leheny, R.L., Chien, C.L.,Weihs, T.P., “Temperature DependentMechanical Properties of Ultra-Fine Grained FeCo-2V”,Acta Mater., 51(14), 4083–4093 (2003) (Experimental, Mechan. Prop., Morphology, 20)

[2004Lal] Laletsin, U., Padubnaya, N., Srinivasan, G., Devreugd, C.P., “Frequency Dependence ofMagnetoelectric Interactions in Layered Structures of Ferromagnetic Alloys and Piezoelec-tric Oxides”, Appl. Phys. A, 78(1), 33–36 (2004) (Experimental, Magn. Prop., 14)

32 Co–Fe–V


[2004Sun1] Sundar, R.S., Deevi, S.C., “Influence of Alloying Elements on the Mechanical Properties ofFeCo-V Alloys”, Intermetallics, 12(7-9), 921–927 (2004) (Magn. Prop., Mechan. Prop.,Morphology, Review, 32)

[2004Sun2] Sundar, R.S., Deevi, S.C., “Effect of Heat-Treatment on the Room Temperature Ductility ofan Ordered Intermetallic Fe-Co-V Alloy”, Mater. Sci. Eng. A, 369(1-2), 164–169 (2004)(Experimental, Mechan. Prop., Morphology, Phase Relations, 22)

[2004Sun3] Sundar, R.S., Deevi, S.C., “Isothermal Oxidation Behavior of FeCo-2V Intermetallic Alloy”,Intermetallics, 12(12), 1311–1316 (2004) (Crys. Structure, Experimental, Interface Phenom-ena, Kinetics, Morphology, Phase Relations, 12)

[2005Sou] Sourmail, T., “Near Equiatomic FeCo Alloys: Constitution, Mechanical and Magnetic Prop-erties”, Prog. Mater. Sci., 50(7), 816–880 (2005) (Crys. Structure, Electr. Prop., Kinetics,Magn. Prop., Mechan. Prop., Morphology, Phase Diagram, Phase Relations, Review, 122)

[2005Sun1] Sundar, R.S., Deevi, S.C., “Soft Magnetic Fe-Co Alloys: Alloy Development, Processing,and Properties”, Int. Mater. Rev., 50(3), 157–192 (2005) (Crys. Structure, Electr. Prop.,Magn. Prop., Mechan. Prop., Phase Diagram, Phase Relations, Review, 258)

[2005Sun2] Sundar, R.S., Deevi, S.C., Reddy, B.V., “High Strength FeCo-V Intermetallic Alloy: Electri-cal and Magnetic Properties”, J. Mater. Res., 20(6), 1515–1522 (2005) (Crys. Structure,Electr. Prop., Experimental, Magn. Prop., Mechan. Prop., Morphology, Phase Relations, 34)

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Co–Fe–V 33



Cobalt – Iron – Tungsten

Lesley Cornish, Andy Watson

Introduction

The system is of interest because of its magnetic properties (it is the basis for magnetic steels), which wereinitially studied by [1932Koe1, 1932Koe2, 1933Rog]. Also of interest is the precipitation behavior[1932Koe2, 1933Rog] of the material, which facilitates hardening.Most of the phase diagram was initially derived by [1932Koe1, 1937Syk], and some of the boundaries areslightly different from the currently accepted binaries (see below). Although [1988Ray] used the verticalsections of [1932Koe1, 1937Syk] to derive the liquidus surface and underlying solidification reactions,the former were considered too small-scale to be used to obtain the temperatures sufficiently accurately.However, both works [1932Koe1, 1937Syk] were extensive, with at least 78 and 28 compositions studied,respectively. The samples of [1932Koe1] covered most of the system (although they were below 70 mass% W, 85 mass% Fe and ~82 mass% Co), whereas those of [1937Syk] were in the Fe rich corner only.The system was reviewed by [1981Ray, 1988Ray, 1994Rag], with the two early works using the data of[1932Koe1, 1937Syk]. The later review [1994Rag] contained more recent data for the Fe2Wand Co3W bin-ary phases, including these phases in the diagrams; and the diagrams were taken from the calculations of[1988Gui].The calculation by [1988Gui] compared relatively well with experimental results [1932Koe1, 1937Syk,1986Jin, 1986Sel], although there were some discrepancies in the extent of the (δαFe) and (αCo,γFe)phases and the shape and position of the (δαFe) + (αCo,γFe) + μ phases. The Fe rich (δαFe) phase of[1932Koe1] penetrated furthest into the system, followed by that of [1937Syk], and lastly the phase fieldof [1988Gui]. Unfortunately, there were no recent data for that phase field. However, [1988Gui] admittedthat the calculated results were compromised by lack of data; especially thermodynamic data.The present work agrees mainly with that of [1994Rag] using the work of [1988Gui] as a basis, although thediagrams have been altered slightly to agree with the selected binary systems and the available experimentalwork [1932Koe1, 1937Syk, 1986Jin, 1986Sel]. However, the calculations at and just below the liquiduswere deemed to be less reliable since there was less agreement with the ternary experimental work. Itwas felt reasonable to use the calculation as the basis for this evaluation however, with necessary alterations,since some of the vertical sections [1932Koe1, 1937Syk] were not always consistent with each other, (albeitto a minor extent), even in the same work. Additionally, the earlier workers were using quite different ver-sions of the Co-Fe and Fe-W phase diagrams, which could have led to different interpretations than if theywere using those currently accepted. Some of the temperature-composition sections of [1937Syk] were verycomplex and were studied without using sufficient samples. Also, the carbon content (allegedly “carbon-free”) was high enough to promote pearlite in at least one composition.Details of studies of phase equilibria and thermodynamics are listed in Table 1.

Binary Systems

The Co-Fe binary phase diagram is accepted from [Mas2]. There are no true intermetallic phases in this sys-tem, although there is ordering and a magnetic transformation. The Co-W system is more complex and hastwo intermetallic phases. A Calphad type assessment of the Co-W system has been produced by [1989Gui],which is essentially the same as given in [Mas2]. However, there is one major difference in that the calcu-lated Co-W diagram exhibits phase separation in the (αCo) phase caused by the magnetic transition in Co(Fig. 1). This has the appearance of a ‘horn’ and is absent in the diagram of [Mas2]. Also, the homogeneityrange of the μ phase is slightly different (41.8-45.5 at.% W [1989Gui], 43.1–48.5 at.% W [Mas2]). Studiesin the ternary system by [1986Sel] would seem to confirm the homogeneity range given in [Mas2]. Also, theCo3W phase was modeled as a stoichiometric phase in [1989Gui], despite it having a homogeneity range.For the purposes of this assessment, the diagram of [Mas2] is preferred, despite the missing miscibility gap

Co–Fe–W 1



in the (αCo). In the Fe-W system according to [Mas2], there are three intermetallic compounds, μ(Fe7W6),δ(FeW) and λ(Fe2W), which is represented as a line compound (as well as being considered metastable).The presence of the δ phase is attributable to the work of [1986Nag], but earlier work [1932Koe1,1937Syk, 1986Jin, 1985Kos, 1986Sel] did not observe this phase in their studies. The system was assessedby [1987Gus], omitting the δ phase. There was also some disagreement in the location of the γ loop andαFe solvus between the assessed phase diagram and [Mas2]. The assessed curves seem to fit the experimen-tal data for the binary system very well, and so the [1987Gus] version of the phase diagram is accepted.

Solid Phases

The solid phases are given in Table 2. The Fe3W2 and CoW phases of [1932Koe1] defined then as the θsolid solution are now designated as Fe7W6 and Co7W6 and the μ solid solution respectively. The γ ⇌ αtransformation was studied further by [1943Syk]. [1981Uru] studied the metastable precipitates and misci-bility gap.The λ and δ phases were not observed by [1932Koe1]. Although [1943Syk] was aware of the extra phases,he did not formally record them because they were finely dispersed and he could not coarsen them suffi-ciently to obtain satisfactory XRD results.The structure of martensite was studied by [1974Edn, 1993Nik].


The μ (formerly designated θ) phase (W6(Fe,Co)7) taking part in the L + (δαFe)⇌ μ + (αCo,γFe) invarianttransition reaction was reported by [1932Koe1, 1988Ray]. The composition of the liquid phase was given as22 mass% Co, 27 mass% W (Table 3). The composition of the liquid resulting from the assessment of[1988Gui] was somewhat different (9.75 mass% Co, 25 mass% W). A reaction scheme was preparedby [1988Ray] and is shown in Fig. 2. However, amendments were made to include the minimum in theL+(W)+μ monovariant line as calculated by [1988Gui].


The liquidus (together with the reaction projection) and some solidus and solvus surfaces were first studiedby [1932Koe1], but is now inconsistent with the currently accepted binaries because the μ phase lies at aslightly different W content than was designated for θ, and the fact that the peritectic reaction in Co-Feoccurs at higher Fe content. Within these changes, [1986Zak] confirmed the liquidus surfaceof [1932Koe1]. [1949Jae] redrew the 20°C isothermal section slightly, and combined it with the liquidusprojection.A liquidus surface projection was calculated by [1988Gui], and compared well with that of [1932Koe1](especially the minimum in the L+μ+(γFe) monovariant line), there was less good agreement with theextension of the bcc liquidus surface. It was the author’s conclusion that the calculation was compromisedby the lack of data, especially thermodynamic data.The liquidus projection presented in Fig. 3 was drawn from [1988Ray]. The solid lines represent the liqui-dus lines and the dashed lines are solidus lines. Also included is the four phase invariant plane. Modifica-tions have been made to the positions of the lines in order to be consistent with the accepted binary systemsand to accommodate a narrower four phase invariant plane that is more consistent with most of the experi-mental data. The minimum in the L + (W) ⇌ μ monovariant taken from [1988Gui] is drawn, even though itwas not shown in the diagrams of [1932Koe1] or [1988Ray]. The inclusion of the minimum is reasonablebecause there is one in the nearby L+μ+γ monovariant line, and also the liquidus of the Co-Fe system[Mas2] is very flat, also showing a minimum.

Isothermal Sections

[1932Koe1] gave complete isothermal sections at room temperature and 1400°C, whereas [1932Koe2] gaveadditional data at 1300°C. The 20°C isothermal section was redrawn slightly by [1949Jae], and combinedwith the liquidus projection.

2 Co–Fe–W


Isothermal sections were calculated by [1988Gui] at: 800, 1000 (to compare with [1986Sel]), 1107 (to com-pare with [1986Jin]), 1200, 1400 (to compare with [1986Sel]) and 1477°C. The agreement with the experi-mental results was good, but not exact. In most cases, the calculated μ limits were less W rich than theexperimental results. The main reason for this is the composition limits of the μ phase in the binary descrip-tions used in the calculations.The isothermal sections presented here were redrawn slightly from [1988Gui] to be as compatible as possi-ble with the available experimental work [1932Koe1, 1937Syk, 1986Jin, 1986Sel]. Figures 4 and 5 showthe isothermal sections at 1000°C and 1107°C respectively. The phase boundaries of the isothermal sectionat 1400°C (Fig. 6) have not been changed.


Temperature-composition sections were produced by [1932Koe1] at 10, 15, 19, 28, 37 and 48 mass%Wandalso at Fe:Co ratios of 8:2 and 7:3. More sections were produced at 20 mass%W, together with sections at 5,8, 10, 15, 20, 25 and 30 mass% Co [1937Syk]. [1986Zak] drew sections across the W rich corner at 90 and80 mass% W. [1988Gui] presented calculated vertical sections to compare with the results of [1932Koe1](apart from the 19 mass% W section) and found good agreement with the experimental data. These sectionsare reproduced in Figs. 7 to 13. The ‘p’ and ‘f’ subscripts refer to the paramagnetic and ferromagnetic formsof the phases.

Thermodynamics

The system was assessed by [1988Gui] by extrapolating from the constituent binary systems and adding aternary excess term for the liquid phase.


Magnetic properties were first studied by [1932Koe2] and the change in coercivity was attributed to the pre-cipitation of another phase. The effect of W on the magnetic transition in Co-Fe bcc alloys was studied by[1979Ko, 1981Uru]. Although near equiatomic FeCo alloys have exceptional magnetic properties, they arevery brittle, and adding W conferred no advantages according to [2005Sou], but improved the workability(up to 2 mass% W) according to [2005Sun]. Magnetic properties for a wide range of alloys were examinedby [1980Kum, 1981Kum, 1982Kum1, 1982Kum2, 1982Kum3] and showed that high remanence semihardCo-Fe-W alloys had potential for relay or switch applications.Hardness was also studied by [1932Koe2]. As well as phase diagram studies, the effect of W content andheat treatment on the hardness was studied by [1937Syk]. [1951Tut] examined the hardness of ninewidely-dispersed alloys. Dispersion hardening was studied by [1971Ers]. [1979Ko, 1981Uru] studied theeffect of W on the formation of metastable precipitates. Sintered powder metallurgy Fe-Co-W alloys werecompared with the calculated phase diagram, and it was demonstrated that hardening was achievable byheat treatment for alloys between 20-30 mass% W and 20 mass% Co [2004Gal]. [1979Dem] studied theeffect of welding on the structures and phase transformations.[2003Cap] studied Co-Fe and Co-Fe-W alloys to find a more wear resistant plating than hard chrome; theattempt was not successful, although [2004He] found some success.Fe-Co-W cathodes have been shown to be promising candidates for fuel cell applications [1998Ram].Investigations of materials properties are listed in Table 4.

Miscellaneous

Mössbauer studies of electrodeposited Co-Fe-W coatings indicated a complex structure comprising at leasttwo magnetic phases and two non-magnetic phases [1994Vas].Due to the potential application of ultra-thin ferromagnetic films for spintronic devices, [2005Pra] studiedthe growth characteristics of thin layers of Co1–xFex on a W (1 1 0) substrate. Although a homogeneous ran-dom monolayer was produced, annealing at intermediate temperatures (237°C) gave SRO Fe3Co and FeOsuperstructures.

Co–Fe–W 3



Table 1. Investigations of the Co-Fe-W Phase Relations, Structures and Thermodynamics

Reference Method / Experimental Technique Temperature / Composition / Phase Range Studied

[1932Koe1] Metallography, XRD, thermalanalysis, dilatometry.

Up to melting points, entire system

[1932Koe2] Metallography, thermal analysis,Brinell hardness and magneticcoercivity.

Up to 1300°C, entire system

[1937Syk] Metallography, dilatometry, hardness Fe rich corner

[1943Syk] Metallography, hardness Fe rich corner

[1937Cor] Metallography, dilatometry, thermalanalysis, hardness and tensile tests

69.4 Co: 10 Fe: 20.4 W (mass%) 55.83 Co: 19 Fe:25.24 W (mass%) 38.9 Co: 20 Fe: 40.7 W (mass%)

[1981Uru] Electron microscopy, DTA, XRD andhardness

5 - 30 mass% W and 0 - 50 mass% Co

[1986Zak] Metallography, XRD Fe2W-Co7W6 μ region

[1988Gui] Calculation, using sublattice modelfor intermetallic phases

Complete


Phase/Temperature Range [°C]



Comments/References

(δαFe)(δFe)1538 - 1394(αFe)< 912

cI2Im�3mW

a = 293.15

a = 286.65

1394°C [Mas2]

at 25°C [Mas2]

γ, (αCo,γFe)(αCo)1495 - 422

(γFe)1394 - 912

cF4Fm�3mCu

a = 354.47a = 356.88

a = 364.67

[V-C2, Mas2]at 520°C [V-C2]

at 915°C [V-C2, Mas2]

(εCo)< 422

hP2P63/mmcMg

a = 250.71c = 406.86

at 25°C [Mas2]

(W)< 3422

cI2Im�3mW

a = 316.52 at 25°C [Mas2]

α´, FeCo cP2 a = 285.04 [V-C2]

(continued)

4 Co–Fe–W





Comments/References

< 730 Pm�3mCsCl

25 to 75 at.% Fe

WCo3< 1093

hP8P63/mmcNi3Sn

a = 512.0c = 412.0

at 25°C [Mas2, V-C2]

μ, W6(Fe,Co)7

W6Co7< 1689W6Fe7< 1637

hR39R�3mFe7W6 a = 475.5

c = 2568.17a = 475.63 ± 0.05c = 2572.8 ± 0.3

designated θ[1932Koe1]

[Mas2, V-C2]

[Mas2, V-C2]

δ, WFe≲ 1215

oP56P212121MoNi

- [Mas2], orthorhombic

Metastable / high pressure phases

(εFe) hP2P63/mmcMg

a = 246.8c = 396.0

at 25°C, 13 GPa [Mas2]

λ, Fe2W≲ 1060

hP2P63/mmcMgZn2

a = 473.7c = 769.4

[V-C2].First recognized by [1937Syk]

η hP4P63/mmLa

- < 35 mass% Co [1981Uru]


Reaction T [°C] Type Phase Composition (at.%)

Co Fe W

L + (δαFe) ⇌μ + (αCo,γFe) 1465 U L(δαFe)μ(αCo,γFe)

2622.111.723.2

63.767.147.367.5

10.310.8419.3

Co–Fe–W 5



Table 4. Investigations of the Co-Fe-W Materials Properties


[1932Koe2] Magnetic field, Brinell hardness Magnetic properties, hardness

[1933Rog] Fischer comparator, immersion in water,Rockwell tester.

Coercive force, residual magnetism, fluxdensity, electrical conductivity, density,hardness, Young’s modulus

[1979Ko] XRD, microprobe analysis Magnetic transition and precipitateformation

[1981Uru] Vickers hardness Hardness and precipitate formation

[2005Pra] Scanning tunnelling microscopy, scanningtunnelling spectroscopy and low-levelenergy diffraction (LEED).

Growth of Co1–xFex films on W (1 1 0)

Fig. 1. Co-Fe-W. The Co-W binary diagram

6 Co–Fe–W


Fig.2.

Co-Fe-W.Reactionscheme

Co–Fe–W 7



Fig. 3. Co-Fe-W. Liquidus and solidus surface projections

8 Co–Fe–W


Fig. 4. Co-Fe-W. Isothermal section at 1000°C

Co–Fe–W 9



Fig. 5. Co-Fe-W. Isothermal section 1107°C

10 Co–Fe–W


Fig. 6. Co-Fe-W. Isothermal section at 1400°C

Co–Fe–W 11



Fig. 7. Co-Fe-W. Vertical section at 10 mass% W

12 Co–Fe–W



Co–Fe–W 13




14 Co–Fe–W



Co–Fe–W 15




16 Co–Fe–W


Fig. 12. Co-Fe-W. Vertical section W-Fe80Co20

Co–Fe–W 17



Fig. 13. Co-Fe-W. Vertical section W-Fe70Co30

18 Co–Fe–W


References[1932Koe1] Koester, W., Tonn, W., “The Fe-Co-W System” (in German), Arch. Eisenhuettenwes., 5(8),

431–440 (1932) (Phase Diagram, Phase Relations, Experimental, *, 7)[1932Koe2] Koester, W., “Mechanical and Magnetic Age-Hardening of the Iron-Cobalt-Tungsten and

Iron-Cobalt-Molybdenum Alloys” (in German), Arch. Eisenhuettenwes., 6(1), 17–23(1932) (Phase Relations, Mechan. Prop., Experimental, 10)

[1933Rog] Rogers, B.A., “Magnetic Properties of Iron-Cobalt-Tungsten Alloys”, Met. Alloys, 4, 69–74(1933) (Experimental, Electr. Prop., Mechan. Prop., 15)

[1937Cor] Cornelius, H., Osswald, E., Bollenrath, F., “About the Processes During the Temper Hard-ening of Some Co-W-Fe Alloys” (in German), Metallwirtschaft, 16(17), 393–399 (1937)(Phase Relations, Mechan. Prop., Phys. Prop., 21)

[1937Syk] Sykes, W.P., “Structural and Hardening Characteristics of Some Iron-Cobalt-TungstenAlloys”, Trans. Am. Soc. Met., 25, 953–1012 (1937) (Phase Diagram, Phase Relations,Experimental, Mechan. Prop., Phys. Prop., *, 16)

[1943Syk] Sykes, W.P., “The Ar´´ Range in Some Iron-Cobalt-Tungsten Alloys”, Trans. Am. Soc. Met.,31, 284–302 (1943) (Phase Relations, Experimental, Mechan. Prop., Phys. Prop., 4)

[1949Jae] Jaenecke, E., “Co-Fe-W” (in German) in “Kurzgefasstes Handbuch aller Legierungen”,Winter Verlag, Heidelberg, 637–638 (1949) (Phase Diagram, Phase Relations, Review, 3)

[1951Tut] Tuteva, N.D., Klementev, A.D., “To the Question About Precipitation Hardening Alloy”(in Russian), Izv. Tomsk. Politekh. Inst., 68(1), 57–64 (1951) (Phase Relations, Experimen-tal, Mechan. Prop., 0)

[1971Ers] Ershova, L.S., “Dispersion Hardening of a Fe-Co-W Alloy”, Russ. Metall., 1, 103–110(1971) (Phase Relations, Experimental, Mechan. Prop., 8)

[1974Edn] Edneral, A.F., Zhukov, O.P., Perkas, M.D., “Effect of Cobalt on Ageing of Martensite andFerrite in Fe-Co-W and Fe-Co-Mo Alloys”, Met. Sci. Heat Treat., 18(9-10), 840–843(1974) (Phase Relations, Experimental, 6)

[1979Dem] Demyantsevich, V.P., Kryzhanovskii, A.S., “Structural and Phase-transformations in theDeposited Metal of the Fe-Co-W and Fe-Co-Mo Systems”, Welding Production, 26(5)3–5 (1979) (Experimental, Mechan. Prop., 5)

[1979Ko] Ko, M., Nishizawa, T., “Effect of Cobalt on the Solubility Anomaly due to Magnetic Tran-sition in αFe” (in Japanese), Nippon Kinzoku Gak-kai-Si, 43(2), 126–135 (1979) (Experi-mental, Phys. Prop., 5)

[1980Kum] Kumasaka, K., Ono, K., “Semihard Magnetic Properties of Fe-Co-WAlloys” (in Japanese),J. Jpn. Inst. Met., 44(11), 1267–1273 (1980) (Experimental, Electr. Prop., 19)

[1981Kum] Kumasaka, K., Ono, K., “Magnetic Anisotropy of an Fe-Co-W Semihard Alloy”, J. Jpn.Inst. Met., 45(1), 82–86 (1981) (Experimental, Electr. Prop., 10)

[1981Ray] Raynor, G.V., Rivlin, V.G., “Phase Equilibria in Iron Ternary Alloys. Critical Evaluation ofConstitutions of Certain Ternary Alloys Containing Iron, Tungsten, and a Third Metal”,Int. Met. Rev., 26(4) 213–249 (1981) (Phase Diagram, Review, #, *, 43)

[1981Uru] Urushihara, F., Sato, S., “The Miscibility Gap and the Metastable Precipitates of Fe-W-CoTernary Alloys” (in Japanese), Nippon Kinzoku Gak-kai-Si, 45(7), 723–731 (1981) (Experi-mental, Phys. Prop., 21)

[1982Kum1] Kumasaka, K., Ono, K., “High Remanence Fe-Co-W Semihard Magnetic Alloys for Relayor Switch Applications”, Review of the Electrical Communication Laboratories, Japan,30(4), 712–723 (1982) (Experimental, Electr. Prop., 15)

[1982Kum2] Kumasaka, K., Ono, K., “High Remanence Fe-Co-W Semihard Magnetic Alloys for Relayor Switch Application”, Electrical Communication Laboratories Technical Journal, Japan,31(5), 1015–1028 (1982) (Experimental, Electr. Prop., 17)

[1982Kum3] Kumasaka, K., Ono, K., “Low Cobalt Fe-Co-W Semihard Magnetic Alloys for RemanentReed Switch Applications”, IEEE Trans. Magn., MAG-18(4), 941–944 (1982) (Experi-mental, Electr. Prop., 7)

[1985Kos] Kostakis, G., “Intermetallic Phases of the System Fe-W”, Z. Metallkd, 76(1), 34–36 (1985)(Crys. Structure, Phase Diagram, Experimental, 13)

Co–Fe–W 19



[1986Jin] Jin, Z.P., Dept. Materials Science, Central-South University of Technology, Cjangsha,Hunan, China, Unpublished research, private communication to Guillermet, A.F., as quotedin [1988Gui]

[1986Nag] Nagender Naidu, S.V., Sriramyrthy, A.M., Rama Rao, P.R., “The Iron-Tungsten System”,J. Alloy Phase Diagrams, 2(3), 176–188 (1986) (Phase Diagram, Assessment, 80)

[1986Sel] Selleby, M., Bachelor Thesis, Division of Physical Metallurgy, Royal Institute of Technol-ogy, Stockholm, Sweden (1986), as quoted in [1988Gui]

[1986Zak] Zakharov, A.M., Parshikov, V.G., Godovannaya, E.B., “The Tungsten Corner of the W-Fe-Co system at the Temperatures higher 1200°C” (in Russian), Izv. Vyss. Uchebn. Zaved.,Tsvetn. Metall., 6, 76–78 (1986) (Phase Diagram, Experimental, *, 9)

[1987Gus] Gustafson, P., “AThermodynamic Evaluation of the C-Fe-W System”, Met. Trans. A, 18A,175–188 (1987) (Phase Diagram, Thermodyn., Assessment, Calculation, 53)

[1988Gui] Guillermet, A.F., “Thermodynamic Calculation of the Fe-Co-W Phase Diagram”,Z. Metallkd., 79(10), 633–642 (1988) (Phase Diagram, Calculation, Thermodyn., *, #, 34)

[1988Ray] Raynor, G.V., Rivlin, V.G., “Co-Fe-W” in “Phase Equilibria in Iron Ternary Alloys”, Inst.Metals, London, 283–288 (1988) (Phase Diagram, #, *, 2)

[1989Gui] Guillermet, A.F., “Thermodynamic Properties of the Co-W-C System”, Met. Trans. A, 20A(5), 935-956, (1989) (Phase Diagram, Thermodyn., Assessment, Calculation, 69)

[1993Nik] Nikolin, B.I., Babkevich, A.Yu., Izdkovskaya, V., Petrova, S.N., “Effect of Heat Treatmenton the Crystalline Structure of Martensite in Iron-, Nickel-, Manganese- and Silicon-dopedCo-W and Co-Mo Alloys”, Acta Met. Mater., 41(2) 513–515 (1993) (Crys. Structure,Experimental, 7)

[1994Rag] Raghavan, V., “Co-Fe-W (Cobalt-Iron-Tungsten)”, J. Phase Equilib., 15(5), 528–529(1994) (Phase Diagram, Review, 5)

[1994Vas] Vasilev, E.A., Tkachenko, T.M., Fedosyuk, V.M., Kasyutich, O.I., Dmitrieva, E.A., “Möss-bauer Study of Amorphous Co-Fe-W Films”, Russ. Metall., 2, 88–91 (1994) (Experimental,Phys. Prop., 7)

[1998Ram] Ramesh, L., Sheshadri, B.S., Mayanna, S.M., “Development of Fe-Co-WAlloys as CathodeMaterials for Fuel Cell Application”, Trans. Inst. Metal Finishing, 76(3) 101–104 (1998)(Experimental, Electr. Prop., 16)

[2003Cap] Capel, H., Shipway, P.H., Harris, S.J., “Sliding Wear Behaviour of Electrodepositied Cobalt-Tungsten and Cobalt-Tungsten-Iron Alloys”, Wear, 255(2), 917–923 (2003) (Morphology,Experimental, Mechan. Prop., 22)

[2004Gal] Galimberti, P., Antoni-Zdziobek, A., “Mechanical Properties and Microstructure of Fe-Co-WSintered Alloys after Heat Treatment” in “Powder Metallurgy World Congress and Exhibition(PM2004)”, European Powder Metallurgy Assoc., Shrewsbury, UK, 6 (2004) (Phase Rela-tions, Experimental, Mechan. Prop., 13)

[2004He] He, F.-J., Lei, J.-T., Lu, X., Huang, Y.-N., “Friction and Wear Behaviour of Electrodeposi-tied Fe-Co-W Deposits”, Trans. Nonferrous Mat. Soc. China, 14(5) 901–906 (2004) (PhaseRelations, Experimental, Mechan. Prop., 16)

[2005Pra] Pratzer, M., Elmers, H.J., “Heteroeptiaxial Growth of Co1–xFex Alloy Monolayers on W(1 1 0)”, J. Cryst. Growth, 275(1-2), 150–156 (2005) (Phase Relations, Experimental, 20)

[2005Sou] Sourmail, T., “Near Equiatomic FeCo Alloys: Constitution, Mechanical and Magnetic Prop-erties”, Prog. Mater. Sci., 50(7), 816–880 (2005) (Crys. Structure, Electr. Prop., Kinetics,Magn. Prop., Mechan. Prop., Morphology, Phase Diagram, Phase Relations, Review, 122)

[2005Sun] Sundar, R.S., Deevi, S.C., “Soft Magnetic Co-Fe Alloys: Alloy Development, Processing,and Properties”, Int. Mater. Rev., 50(3), 157–192 (2005) (Crys. Structure, Electr. Prop.,Magn. Prop., Mechan. Prop., Phase Diagram, Phase Relations, Review, 258)

[Mas2] Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International,Metals Park, Ohio (1990)

[V-C2] Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetal-lic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)

20 Co–Fe–W


Chromium – Copper – Iron

Tamara Velikanova, Mikhail Turchanin

Introduction

Phase equilibria in the Cr-Cu-Fe alloys are of importance both for the development of Cr-Cu-Fe alloys andfor their possible use in recycling processes in metallurgy. The Cr-Cu-Fe system is of interest for ferrous andnonferrous metallurgy and for the development of Cu based composites with favorable mechanical and elec-trical properties. The Cu rich alloys of the system are promising as new cost effective, high strength, highconductivity copper alloys. Copper has widely been used as an alloying element in ferrous metallurgy toimprove strength of steels under aging, their corrosion resistance and antifriction properties. On the otherhand, the presence of copper in ferrous alloys has negative aspects as well. It is known that a small amountof residual Cu in steel scrap, which is difficult to remove by conventional steelmaking process, is a reason of“surface fissures” during hot rolling process. The solubility of Cu in solid Fe, and therefore the influence ofCr on the solubility, and liquid-solid wettability are considered as important factors affecting surface fissuresformation. Features of phase relations in the Cr-Cu-Fe alloys appeared to help solving the problem of Cuseparation from steel scrap, which is important for both metallurgical processing difficulties and environ-mental problems. All these aspects stimulated intensive investigation of the Cr-Cu-Fe alloy phase diagram.The main works are listed in Table 1.Seven temperature-composition sections in the Fe rich corner of the diagram, three at constant chromiumand four at constant copper content, were reported by [1939Mor]. It is worth noting that the alloys investi-gated below the solidus were contaminated with carbon up to 0.1 mass%. Additional information on the ver-tical sections at 2 and 4 mass% Cu in a high Fe content composition range was reported by [1970Ahm,1974Zap] on alloys obtained by powder-metallurgical process and later the sections were calculated by[1998Mie]. The calculated temperature-composition sections at 1 and 4 mass% Cu in the whole concentra-tion range were reported by [2004Wan] on the basis of own experimental data. The phase equilibria werestudied in a wide composition range, and a number of isothermal sections were constructed in the tempera-ture range from 800 to 1300°C by [1967Sal, 1993Hao, 1997Oht, 2001Fer1, 2002Wan]. The effect of Cr onthe solubility of Cu in iron was investigated by [1967Sal] using alloys based on mild steel. All the availableexperimental data on the phase equilibria in the ternary Cr-Cu-Fe system generally well agree. Miscibilitygap in liquid of the system was studied by [2004Wan].The results of investigations of physical, chemical and mechanical properties of Cr-Cu-Fe alloys obtainedby both conventional and powder metallurgical methods were presented by [1967Sal, 1970Ahm,1974Zap, 1971Yam, 2001Les, 2001Fer1, 2001Fer2, 2001Fer3, 2001Son, 2002Kim].The Cr-Cu-Fe system was first critically assessed by [1979Cha]. This assessment was mainly based on theresults of [1939Mor, 1970Ahm, 1974Zap]. The same data set was used for modeling thermodynamic prop-erties of (Cu), α and γ phases of the system in the framework of development of approximate thermody-namic solution phase data for steels [1998Mie]. The thermodynamic parameters of the ternary systemwere evaluated by [1997Oht] on the basis of own experimental results for the solubility ranges of phasesin the temperature interval 1100-1300°C, and an isothermal section of the system at 1200°C was presented.The thermodynamic assessment of the phase equilibria in the Cr-Cu-Fe system was carried out by[2002Wan] on the base of own experimental results and data of [1997Oht] taking into account data of pre-vious works, excluding [1967Sal] who investigated the Cr-Cu-mild steel alloys instead of Cr-Cu-Fe as men-tioned above. The results of calculation agree well enough with the experimental data available. Theoptimized and consistent thermodynamic description of the Cr-Cu-Fe system of [2002Wan] was used forcalculation of the majority of figures on phase relations and reaction scheme presented in this assessment.

Cr–Cu–Fe 1



Binary Systems

The assessments of the Cr-Fe system by [1987And], of the Cr-Cu system by [1990Ham] and of the Cu-Fesystem by [1995Che] are accepted. These works are in good consent with [2002Ans, 2007Tur, Mas2]. Thethermodynamic data sets of [1987And, 1990Ham, 1995Che] were used by [2002Wan] for thermodynamicassessment of the ternary Cr-Cu-Fe system.

Solid Phases

Table 2 summarizes the crystallographic data on the Cr-Cu-Fe phases and their temperature and concentrationranges of stability. The (Cu), α and γ phases have marked homogeneity ranges in the ternary system.Cr and Fe decrease the solubility of each other in (Cu). The solubility of Cr in the (Cu) phase is not signifi-cantly decreased by addition up to 0.3 mass% Fe after [2001Fer]. The solubility of Fe in the (Cu) phase issignificantly reduced due to the presence of Cr. At 750°C and below this tendency becomes less brightlyexpressed. This indicates that at lower aging temperatures less precipitation occurs during sample quenching.An increase in the Cr solubility in the γ phase with increasing Cu concentration and a decrease in the Cusolubility in the γ phase with increasing Cr concentration was found by [1939Mor, 1967Sal, 1970Ahm,1974Zap, 1993Hao, 1997Oht, 2002Wan] to be general tendencies. The minimum of Cu solubility in auste-nite of mild steel with 5-7 mass% Cr reached in equilibrium with liquid at 1250°C was reported by[1967Sal]. It was not confirmed by the experimental investigations of [1993Hao, 1997Oht, 2002Wan]and the thermodynamic calculations of [1997Oht, 2002Wan] for the ternary Cr-Cu-Fe alloys. Accordingto our calculation based on the thermodynamic models of [2002Wan], the γ phase homogeneity range pene-trates down to ~80 at.% Fe at 1100 to 1300°C in the ternary system.The calculated maximal saturation of the α phase by copper is about 8 at.% Cu at 1380°C. This value cor-relates well with the experimental results of [1993Hao, 1997Oht].The (Cu), α and γ phases were treated as ternary solid solution phases in the thermodynamic assessment of[2002Wan]. For the σ phase the homogeneity range only in the binary system was taken into account in thecalculation.


The data on the invariant equilibria given in Table 3 follow from the thermodynamic calculation carried outin the present assessment. The reaction scheme after the calculation is given in Fig. 1. The existence of themiscibility gap of the liquid phase in the ternary system despite its absence in the boundary binary systemsis an interesting peculiarity of the ternary. The existence of the point corresponding to the cupola top of themiscibility gap of liquid, c1, and of two critical invariant points, c2 and c3, corresponding to the origin ofthe three-phase monovariant equilibria of the α and γ phases with two liquid phases rich in Fe or Cu givenin Table 3, reflects such a specific interaction of the components in the ternary. Invariant four-phase mono-tectic equilibrium of the transition type, U1, and invariant equilibrium of the transition type, U2, are shown.The calculated temperature of the U2 equilibrium, 1088°C, agrees perfectly with that measured experimen-tally by [1939Mor], 1084°C. The calculated and experimental compositions of the α and γ phases are alsoclose.

Liquidus and Solidus Surfaces

A liquidus projection is given in Fig. 2a (in the whole composition range) and in Fig. 2b (enlarged part nearthe Cu corner) according to the thermodynamic calculation performed in this assessment. Awide two-phaseL′ + L″ phase as well as addition of Fe to the Cr-Cu alloys decreases the liquidus of Cr-Cu alloys (liquidusof α-phase) down to the monotectic valleys L′ ⇌ L″ + γ and L′ ⇌ L″ + α, respectively. The above men-tioned feature gives rise to the minimum fold on the liquidus surface of the α phase in the ternary system.The fold extends from the point of the temperature minimum in the solid-liquid equilibrium region of thebinary Cr-Fe up to the invariant e1(min) point at the monovariant monotectic L′⇌ L″ + α line in the ternarysystem. Accordingly, the tie lines L′L″ in the equilibria L′ + L″ + α, which radiate from the Cu corner tor-wards the Cr-Fe side, go from the c2 and c3 critical points to meet one another at the minimal temperature on

2 Cr–Cu–Fe


the ruled surface L′L″ at 1362°C. The composition range of the primary crystallisation surface of (Cu) isshown to be very restricted in agreement with the data for the binary Cr-Cu and Cu-Fe phase diagrams.The calculated solidus surface projection is given in Fig. 3. The narrow γ + α + (Cu) three-phase regionshifted to the Cu-Fe edge separates the two-phase γ + (Cu) and α + (Cu) regions. A very slight decreasein solidus temperature from the Cu-Fe to Cr-Cu boundary systems is observed. According to the calculation,the solubility of Cu in the (αFe) and (γFe) phases at solidus decreases from 5.8 and 6.9 at.% in the binaryCu-Fe after [2007Tur] down to 3.9 at 18.3 at.% Cr and 5.7 at.% at 14.7 at.% Cr at 1088°C, respectively. Thecontent of Cr in (Cu) solid solution is rather small being of 0.3 at.% at 3.4 at.% Fe.

Miscibility Gap Surface

The calculated miscibility gap of the liquid phase in the ternary system in the stable and metastable ranges isgiven in Fig. 4. The isotherms at 1127 to 1477°C are taken from [2004Wan]. The set of the vertical sectionsof the miscibility gap parallel to the Cu-Fe edge is given in Fig. 5. The curve at 4 mass% Cr is taken from[2004Wan]. The others are shown after the present calculation. Figures 4 and 5 show that additions of Cr toCu-Fe alloys and Fe to Cr-Cu increase the critical temperature of the miscibility gap (metastable in the bin-aries Cu-Fe and Cr-Cu). Above 1.7 at.% Cr addition to Cu-Fe (critical point c3 at 1431°C) the stable mis-cibility gap appears. The minimal addition of Fe to Cr-Cu alloys to stabilize the miscibility gap is 9.7 at.%(critical point c2 at 1516°C). One can see that the stable miscibility gap exists in a wide composition regionin the ternary system, forming the rather flat cupola with maximum at 1523°C.

Isothermal Sections

Isothermal sections at 1500, 1250, 1050, and 800°C are given in Figs. 6 to 9 according to the thermody-namic calculations. The calculated isothermal sections at 1000, 1100 and 1300°C were reported by[2002Wan]. Additionally the isothermal sections at 1600, 1500, 1400, 1250, 1200, 1050, 900 and 800°Cwere calculated in the present assessment using the thermodynamic models of [2002Wan]. The results ofthe calculations correlate well with the experimental investigations of [1993Hao, 1997Oht, 2001Fer1,2002Wan].No experimental data on phase equilibria at temperatures above 1300°C are published. Strong bend of theL / L + α boundary towards the Cr-Fe edge giving a minimum Cu content at about 20 at.% Fe is seen in thecalculated section at 1500°C, Fig. 6. This feature agrees with the location of the invariant minimum at 1513°C and 21 at.% Cr in the boundary Cr-Fe system.The composition of the equilibrium phases (Cu) / (αFe) at 800°C, (Cu) / (γFe) at 900, 1000°C and L /(γFe)at 1200°C were reported by [2002Wan] for the Fe rich alloys of the Cr-Cu-Fe system. The alloys were pre-pared by melting and subsequently hot-rolled at 900°C, solution-treated at 900°C for 24 h, heat-treated at800 to 1000°C for 336-1680 h and then quenched in iced water. It was tested and confirmed that the speci-mens had reached an equilibrium state. The equilibrium compositions were determined up to a Cr content of10.1, 10.79, 10.85 mass% in the Fe based phases and 0.22, 0.21, 0.49 mass% in the (Cu) phase at 800, 900and 1000°C, respectively. The (γFe) solid solution of the composition 9.20Cr-3.38Cu (mass%) was foundto be in equilibrium with liquid of the composition of 0.12Cr-97.77Cu (mass%). The equilibrium composi-tions of the solid and liquid phases at 1100, 1200 and 1300°C were measured by [1997Oht] using solid-liquid diffusion couples held at the above mentioned temperatures for 24-48 h and subsequently quenchedin iced brine. The average composition of the frozen solid and liquid at the interphase boundaries of the dif-fusion couples were obtained up to a Cr content of 28.81, 15.33, 16.40 mass% in solid and of 0.52, 0.47,and 0.80 mass% in liquid at 1100, 1200 and 1300°C, respectively. It was shown that the addition of Crdecreases the solubility of Cu in the Fe based phases. This tendency becomes remarkable with increasingtemperature. Being obtained using diffusion couples technique the experimental data of [1997Oht] mightbe less reliable than those of [2002Wan] who used conventional metallurgical methods and well equilibratedalloys. However, the data of both [1997Oht] and [2002Wan] perfectly agree, and good agreement wasachieved between the calculated locations of the phase boundaries and the experimental results.The isotermal sections at 1200°C constructed by [1997Oht] is cited in the review by [2002Rag].

Cr–Cu–Fe 3



The Cr effect on the solubility of Cu in austenite of mild steel at 900, 1100 and 1250°C was studied by[1967Sal]. The measured solubility of Cu in Fe based phases is slightly lower compared with the resultsof [1997Oht, 2002Wan], and a slight minimum of the solubility upon addition of about 5-6 mass% Crwas found. The minimum point of Cu solubility in Fe based solid solution, reported by [1967Sal], wasnot confirmed by the calculation for the Cr-Cu-Fe system.The calculated sections at 1300, 1250,1200 and 1100°C demonstrate the equilibria of the α, γ and liquid(very rich in Cu) phases and differ from each other in the equilibrium compositions of phases only. Thephase equilibria at 1250°C are shown in Fig. 7. The equilibrium composition of phases at this and the abovementioned temperatures are given in Table 4. The calculated isothermal section at 1050°C is given in Fig. 8.The same feature of the phase equilibria formed by α, γ and (Cu) phases is shown in the calculated isother-mal sections at 1000 and 900°C. Only the compositions of equilibrium phases differ. They are given inTable 5. The isothermal section at 1000°C, reported by [1993Hao], shows slightly elevated Cu contentin the α and γ phases in equilibrium with (Cu) phase comparing with the data of [2002Wan]. [1993Hao]based on own study of three ternary diffusion couples annealed for 150 h with subsequent quenching inwater. The (Cu) / (Cu) + α boundary at 1050°C after the calculation is shown in Fig. 8. It agrees well withthe experimental results of [2001Fer1] obtained using WDS method. One can see in Fig. 8, that Fedecreases the solubility of Cr in (Cu) at 1050°C, and the solubility of Cr reduces more markedly withincreasing Fe content. Similar tendency was found by [2001Fer1] for lower temperatures (down to500°C) using resistivity measurements on aged and subsequently quenched in water alloys. However, thevalues of the solubility measured by this method at 1050°C were found to be noticeably lower than thoseafter WDS method. Consequently, the isotherms at 1025, 1000, 950, 850, 750, 650 and 500°C reportedby [2001Fer1] cannot be considered for description of the equilibria at the indicated temperatures.The calculated isothermal section at 800°C is shown in Fig. 9. The (Cu) + (αFe) / (αFe) phase boundaryshown in Fig. 9 agrees with that of [2002Wan] obtained up to about 10 mass% Cr. There are no experimen-tal data on a possible Cu solubility in the σ phase. It is assumed to be small taking into account the sizeof the Cu atomic radii on one hand, and Cr and Fe radii, on the other hand. Isothermal sections in the tem-perature interval 831 to 510°C, where the σ phase is stable in the Cr-Fe binary system, have to differ in theσ + (Cu) and neighboring α + (Cu) + σ regions only.


The calculated vertical sections are shown in Figs. 10 to 24. The section at 1 mass% Cr is taken from[2004Wan], and the others are given after the calculation performed in the present assessment correspondingto [2002Wan, 2004Wan]. Satisfactory agreement of the published calculations [1997Oht, 2002Wan,2004Wan] in a wide temperature range including the region of the equilibria with liquid phase and the avail-able experimental data of [1939Mor, 1970Ahm, 1974Zap, 1997Oht, 2002Wan, 2004Wan] was achieved asmentioned above. Thus one should believe that the thermodynamic models evaluated by [1997Oht,2002Wan] fit well to the thermodynamic description of the vertical sections in the whole composition rangeat the temperatures under consideration, and the calculated vertical sections are reliable.The first detailed experimental investigation of the temperature-composition sections of the Cr-Cu-Fe sys-tem was carried by [1939Mor]. Contamination with 0.04 to 0.11 mass% carbon was reported for the alloysused in the study of phase equilibria in the solid state. Seven partial vertical sections for the Fe rich partof the system were presented: at 2 mass% Cr (up to about 20 mass% Cu), at 5 and 14 mass% Cr (up to40 mass% Cu), and at 1, 2, 4 and 5 mass% Cu (up to 20 mass% Cr). The results of the later works of[1970Ahm, 1974Zap, 1993Hao, 1997Oht, 2002Wan] concerning the mutual solubility of Cr and Cu inthe γ phase and the solidus temperatures generally agree with the presented vertical sections. [1939Mor]found that the γ phase homogeneity range is enlarged by addition of copper so that the γ/γ + α boundarylies between 14 mass% Cr at 5 mass% Cu and 19 mass% Cr at 4 mass% Cu at about 1084°C. With increas-ing Cr content, the solubility of Cu in the γ phase and the Cu content at eutectoid composition at firstdecrease and then increase, while the eutectoid temperature changes slighly. [1970Ahm, 1974Zap] studiedthe influence of Cr on phase relations in the Fe rich alloys along the sections at 2 and 4 mass% Cu up to 18and 15 mass% Cr in the temperature range from the solidus down to 700 and 1100°C, respectively.

4 Cr–Cu–Fe


The calculated vertical sections at 2 and 5 mass% Cr given in Figs. 11 and 13, at 1 and 4 mass% Cr[2004Wan] given in Figs. 10 and 12, as well as at 1, 2, 4 and 6 mass% Cu, Figs. 18 to 21, confirm the abovementioned a peculiarity of the ternary system, namely, the widening of the γ phase homogeneity range inthe equilibria with the α phase when Cu is added. The miscibility gap of the liquid phase is shown inthe sections parallel to the Cu-Fe edge at 2 mass% Cr and more (including the highest content of Cr upto 40 at.%), Figs. 11 to 17. The same is seen in the sections parallel to the Cr-Fe edge at 2 mass% Cuand more given in Figs. 19 to 22. The isothermal line at 1395.5°C, corresponding to the monotectic invar-iant equilibrium U1, appears in the sections at 4 and 5 mass% Cr as well as at 20 at.% Cu additionally to theL' + L'' region, as shown in Figs. 12, 13 and 22a, 22b. The isothermal plane of the invariant four-phaseequilibria corresponding to the α + γ + (Cu) three-phase alloys solidification is intersected by the verticalsections at 1, 2, 4, 5 and 14 mass% Cr, Figs. 10 to 14.The minimum on the liquidus surface in the sections at 1, 2, 4 and 6 mass% Cu, Figs. 18, 19a, 20 and 21,near 20 at.% Cr at the temperature close to 1500°C is associated with the invariant minimum on the liqui-dus/solidus of the Cr-Fe binary system at 21 at.% Cu and 1513°C.The very narrow L' + L'' + α and L' + L'' + γ ranges spearing the L + α and L + γ fields, respectively, asshown in the sections at 2, 4 and 6 mass% Cu, Figs. 19a, 20 and 21, seem to be very unusual. At higher Cucontent, a minimum at 1362°C within the curves separating the L' + L'', L' + L'' + α and L + α phase regionsappears, Figs. 22a, b. It corresponds to the intersection of the degenerated into the line L' + L'' + α tie trianglewith the plane of this vertical section.

Thermodynamics

There are no experimental data about thermodynamic properties of ternary solution phases in the system. In[1974Sig] and [1988Uen] the thermodynamic properties of liquid alloys in the ternary system were modeledon the basis of theoretical ideas and equations for interaction parameters in ternary solution were developed.In [1997Oht, 1998Mie, 2002Wan] thermodynamic properties of the ternary solution phases, L, (Cu), α andγ, were modeled on the base of data on phase equilibria between these phases. All the results of the calcula-tion, both theoretical modeling and thermodynamic optimization of phase equilibria, demonstrate strongpositive deviations of the thermodynamic properties of the solution phases from the ideality. In all caseswhen ternary interaction parameters were taken into account they have highly positive values. Thus, posi-tive deviations from the ideality, inherent in phases of the boundary Cr-Cu and Cu-Fe binary systems,become more significant in the ternary. As a result, the miscibility gap in the liquid phase, which is meta-stable in the binary Cr-Cu and Cu-Fe melts, appears in the ternary system.The thermodynamic assessment of the Cr-Cu-Fe system was carried out by [2002Wan] using theCALPHAD approach. Binary interaction parameters for the Cr-Cu, Cr-Fe and Cu-Fe systems were takenfrom previous works of [1987And, 1990Ham, 1995Che]. Thermodynamic descriptions for the liquid,(Cu) and γ phases were taken from [1997Oht]. Thermodynamic description for the α phase was obtainedby [2002Wan] on the basis of the experimental data of [1997Oht, 2002Wan]. The ternary parameter forliquid phase was equal to zero. The excess thermodynamic properties of solid solution phases took intoaccount the mixing enthalpy and the mixing entropy of components.Due to the lack of experimental information, the solubility of σ phase in the ternary system is not consideredin [2002Wan]. The thermodynamic assessment was carried out for temperatures above 900°C.


The experimental works devoted to study of materials properties in the Cr-Cu-Fe system are listed inTable 6.New cost effective, high-strength, high-conductivity and resistant to softening copper rich Cr-Cu-Fe alloyswere developed and their properties were investigated by [2001Fer1, 2001Fer2, 2001Fer3]. Determinationof the temperature dependence of the solid solubility of Fe and Cr in copper at 500 to 1050°C for alloyscontaining 0 to 1 mass% Cr and 0 to 1 mass% Fe was undertaken by [2001Fer1] to help interpreting theobserved mechanical and electrical properties of alloys. The linear relationship between resistivity and chro-mium and iron concentrations in the Cu based solid solution was found by [2001Fer1]. Similarly, the

Cr–Cu–Fe 5



relationship between the resistivity and composition of alloys was also approximately linear after the solidsolubility limit of (Cu) phase has been exceeded. Using these data the limit composition of (Cu) solid solu-tion was determined in aged and subsequently quenched alloys.The evident difference in the compositions of the (Cu) solid solution coexisting with α (or γ) phase wasobserved depending on experimental technique, namely WDS method or electrical resistivity measure-ments. The lower solubility observed by the resistivity measurements was explained by the partial (Cu)solid solution decomposition accompanied by small secondary precipitations under quenching experiments.The WDS method gives an average signal including concentration of the above mentioned precipitations.The method based on resistivity measurements gives only composition of (Cu) solid solution remainingunder quenching. Thus the curves shown in Fig. 23 by thin lines are not isotherms of solubility at given tem-peratures. They correspond to unknown lower temperatures.The mechanical properties of the Cu -0.7 mass% Cr - 0.3 mass% Fe alloy which indicated large precipita-tion hardening response combined with the ability to stabilize cold worked microstructures to high tempera-tures with a high electrical conductivity remaining were reported by [2001Fer2]. The age hardeningresponse of the Cu - 0.7 mass% Cr - 2.0 mass% Fe alloy was minimal, but the resistance to softeningwas superior to that reported for any commercial high-strength, high-conductivity copper alloy with com-parable mechanical and electrical properties. For example, an excess of 85% of the original hardness ofthe 40 % cold worked alloy is retained after holding at 700°C for 1 h, whereas commercial high-strength,high-conductivity Cu-Fe-P alloys have been reported to soften significantly after 1 h exposure below500°C. The Cu - 0.7 mass% Cr - 2.0 mass% Fe alloy was expected to be more suitable for applications witha significant risk of exposure to elevated temperatures. Optical microscope examination of cold worked andaged microstructures confirmed the high resistance to recrystallization for Cu - 0.7 mass% Cr - 2.0 mass%Fe. The Zener-Smith drag term, predicting the pinning effect of second phase particles on dislocations incold worked microstructures, was calculated using the precipitate characteristics obtained from TEM,WDS and resistivity measurements [2001Fer2, 2001Fer3]. The pinning effect of the precipitate dispersionsin the peak-aged condition was determined to be essentially equivalent for the Cu - 0.7 mass% Cr - 0.3 mass% Fe and Cu - 0.7 mass% Cr - 2.0 mass% Fe alloys. A lower recrystallization temperature in the Cu - 0.7mass% Cr - 0.3 mass% Fe alloy was therefore attributed to faster coarsening kinetics of the secondary pre-cipitates resulting from a higher Cr concentration in the precipitates at a lower iron content [2001Fer3].The microstructure and mechanical properties of the deformation processed Cr-Cu-Fe microcompositewires combined with intermediate heat treatments have been investigated by [2001Son, 2002Kim]. Theinvestigations were performed using samples of the Cu-1.2Cr-9Fe (mass%) composition. The primaryand secondary dendrite arms were aligned along the deformation axis and elongated into filaments duringcold drawing. It was shown that Cr atoms were mostly located in the Fe based filaments and the coppermatrix was almost free of Cr atoms in agreement with the partition coefficient of Cr between the (Cu)and α or (Cu) and γ phases according to the alloy phase diagram. The filaments in Cr-Cu-Fe microcompo-sites were strengthened by the addition of Cr atoms and their refinement is relatively difficult due to thestrengthening of filaments by Cr. Thermo-mechanical treatments have been employed by [2002Kim] tooptimize the strength and conductivity of Cr-Cu-Fe microcomposites. The ultimate tensile strength andthe conductivity of the Cu-1.2Cr-9Fe (mass%) wires drawn to the cold drawing strain η = 4.8 without inter-mediate heat treatments were observed to be 920 MPa and 33.8% IACS, respectively, and those with heattreatments were 891 MPa and 41% IACS. Further drawing wires to the cold drawing strain η = 6.3 after anadditional heat treatment increased the conductivity from 43.1 to 53.3% IACS with a slight increase in hard-ness. The precipitation of impurities and alloying elements during intermediate heat treatments is thought toincrease the conductivity due to reduced impurity scattering. Fig. 24 shows the variations of the strength andthe ductility as a function of heat treatment temperature. The activation volumes for deformation increasedfrom 138b3 in the as-drawn wire to 230b3 in the wire annealed at 500°C. Numerous particles were observedin Cu matrix, and the spacing between these particles was found to be slightly smaller than the activationlength (138b = 35 nm). The most probable rate controlling mechanism of Cr-Cu-Fe microcomposites is sug-gested to be the interaction between dislocations and precipitates in Cu matrix.The processes of sintering of Cr-Cu-Fe alloys with 2 to 18 mass% Cr and 2 and 4 mass% Cu using iron,copper and chromium or ferrochromium polvers as starting materials were studied by [1970Ahm,1974Zap]. A notable homogenization degree of the alloys during the sintering process was reached at

6 Cr–Cu–Fe


1300°C after 12 h holding only. Examination of the diffusion phenomena and corrosion tests in nitric acidand magnesium chloride solution were carried out. While diffusion of Cr in Fe was found to slightlyincrease with Cu addition, microhardness and corrosion resistance increased with Cr addition markedly.The obtained diffusion coefficients of chromium in the α and γ phases as well as the corresponding activa-tion energies are listed in Table 7.The diffusion coefficients of copper and chromium in liquid Cr-Cu-Fe alloys at 1550°C were determi-ned by [1975Wan], and the following values were obtained: DCu,Cu = (5.5 ± 0.1)·10–5 cm2·s–1,DCu,Cr = (–0.1 ± 0.03)·10–5 cm2·s–1, DCr,Cr = (3.1 ± 0.2)·10–5 cm2·s–1, DCr,Cu = (2.5 ± 0.8)·10–5 cm2·s–1.Composite materials constituted of a Cu rich phase with a high electric conductivity and of a Fe rich phasewith a high strength have an obvious advantage in terms of various properties. The preparation of such com-posite materials, however, is not easy by the powder metallurgical process because of diffusion problems,poor wetting between Cu and Fe [2001Les] and the high cost for fabrication. The formation of the core-typemacroscopic morphology of as-cast Cr-Cu-Fe alloys due to the stable miscibility gap of the liquid phase isconsidered as useful peculiarity of the system for development of easy process of manufacture of the naturalcomposites on the base of the conventional casting process [2004Wan].

Miscellaneous

The problem of hot shortness (cracking) of steels, induced by subscale enrichment of alloying elements dur-ing reheating before hot working, stimulated the investigation of the Cr effect on the Cu solubility in aus-tenite and on the penetration of liquid copper rich phase down austenite grain boundaries. A slight decreasein the Cu solubility in austenite of mild steel when Cr content increased with a slight solubility minimum at5 to 6 mass% Cr was found by [1967Sal] as mentioned above. Although the Cu solubility increased athigher Cr additions, but, with 9.38 mass% Cr, at 1250°C, it was still lower than that with no additions. Thus,the effect of Cr appears to be nearly neutral. However, because of a slight negative effect observed particu-larly over the critical temperature range at about 1100°C, Cr may be classified as a detrimental rather thanbeneficial addition.The dihedral angle of the copper rich phase at the γ phase grain boundaries was measured by [1967Sal] toassess the effect of temperature on the penetration of the liquid copper rich phase into grain boundaries. Thealloys containing more than 3 mass% chromium display a sharp decrease in the dihedral angle at about1100°C. Minimum dihedral angle values, giving rise to maximum grain boundary penetration and probablymaximum susceptibility of the steel to hot shortness are between 1075°C and 1175°C.The multilayers Fe-X/Cu-8/Fe-Y/Cr/Fe-Y/Cu-8 show an inverse giant magnetoresistance [2001Mil]. Theelectrical conductivity of such multilayers decreases with the applied magnetic field. The electronic bandcontribution to the giant magnetoresistance for Fe-3/Cu-4/Fe/Cr/Fe/Cu-4 and Fe-3/Cu-4 multilayers wascalculated within the semiclassical approximation. The results show a large change in the giant magnetore-sistance behavior when one layer of Cr is introduced within the Fe layers. The dependence of impurity vsband effects in the appearance of inverse giant magnetoresistance in Cu/Fe superlattices with Cr was studiedby [2002Mil]. The calculated giant magnetoresistance ratios have been compared with the experimentalresults, and it was concluded that the experimental data can only be explained by taking into account Crbands.Laser surface remelting/resolidifying treatment on a powder metallurgically manufactured Cr-Cu-Fe contactmaterial was studied by [2000Gen]. A compact remelting/resolidifying layer was obtained with appropriatelaser treatment conditions and a suitable surface absorption coating. After the treatment, the Cr-Cu-Femicrostructure of alloy on surface was greatly refined and the α phase was uniformly dispersed in the Curich matrix with a fine spherical or near spherical form. Improved compactness and microstructure of thelaser remelted Cr-Cu-Fe material yielded increased hardness (up to 80%), wear resistance, and a reducedfriction coefficient compared with the base material. The mechanism of laser strengthening was concernedwith the microstructural features of the Cr-Cu-Fe material.

Cr–Cu–Fe 7



Table 1. Investigations of the Cr-Cu-Fe Phase Relations and Structures


[1939Mor] DTA, dilatometric measurements, opticalmicroscopy

Temperature-composition sections:at 2 mass% Cr and 0 to 20 mass% Cu;at 5 mass% Cr and 0 to 40 mass% Cu;at 14 mass% Cr and 0 to 40 mass% Cu;at 1 mass% Cu and 4 to 20 mass% Cr;at 2 mass% Cu and 4 to 20 mass% Cr;at 4 mass% Cu and 4 to 20 mass% Cr;at 6 mass% Cu and 4 to 20 mass% Cr

[1967Sal] Optical microscopy, EPMA, Cr-Cu-mildsteel alloys

Partial isothermal sections:at 900°C ((Cu)/γ phase boundaries);at 1100 and 1250°C ((L)/γ phase boundaries)

[1970Ahm,1974Zap]

Optical microscopy, dilatometricmeasurements, DTA, X-ray analysis

Temperature-composition sections:at 2 mass% Cu and 2 to 18 mass% Cr;at 4 mass% Cu and 2 to 18 mass% Crα/γ phase boundaries

[1971Yam] Optical microscopy, DTA Influence of Cu additions on the miscibility gapin the α-phase

[1993Hao] Diffusion couples technique, opticalmicroscopy, EPMA

Isothermal section at 1000°C, (Cu)/γ, (Cu)/αand α/γ phase boundaries

[1997Oht] Solid-liquid diffusion couples technique,SEM-Electron Dispersive X-ray analysis

Partial isothermal sections at 1100, 1200, 1300°C,(L)/γ and (L)/α phase boundaries

[2001Fer1] Wavelength dispersive spectroscopy,resistivity measurements

Partial isothermal section at 1050°C, (Cu)-phase

[2002Wan] SEM-Electron Dispersive X-ray analysis Partial isothermal sections:at 800°C, (Cu)/α phase boundaries;at 900°C, (Cu)/γ phase boundaries;at 1000°C, (Cu)/γ phase boundaries;at 1200°C, (L)/γ phase boundaries

[2004Wan] Optical microscopy 49 to 65 mass% Cu, 34 to 49.5 mass% Fe, 2 to7.2 mass% Cr, miscibility gap in liquid phase

8 Cr–Cu–Fe




Pearson Symbol/ SpaceGroup/ Prototype


Comments/References

(Cu), CrxCuyFe1–x–y< 1084.62

cF4Fm�3mCu

a = 361.46

at x = 0 and 1096°C, 0.955 ≤ y ≤ 1[1995Che]at x + y = 1 and 1075°C, 0.9992 ≤ y ≤1 [1990Ham]at x = 0, y = 1 and 25°C[V-C2, Mas2]

γ, CrxCuyFe1–x–y1486 - 849

(γFe)1394 - 912

cF4Fm�3mCu

a = 364.68

at x = 0, 0 ≤ y ≤ 0.13 [1995Che]at y = 0 and 977°C, 0 ≤ x ≤ 0.12[1987And]at x = 0, y = 0 and 912°C [V-C2, Mas2]

α, CrxCuyFe1–x–y< 1863

(Cr)< 1863(δFe)1538 - 1394(αFe)< 912

cI2Im�3mW

a = 288.48

a = 293.15

a = 293.22

at x = 0 and 1486°C, 0 ≤ y ≤ 0.063at x = 0 and 849°C, 0 ≤ y ≤ 0.017[1995Che]at y = 0, 0 ≤ x ≤ 1 [1987And]at x + y = 1 and 1557°C,0 ≤ y ≤ 0.0034 [1990Ham]at x = 1, y = 0 and 25°C [V-C2, Mas2]

at x = 0, y = 0 and 1394°C [V-C2,Mas2]at x = 0, y = 0 and 25°C [V-C2, Mas2]

σ, CrxFe1–x831 - 510

tP30P42/mnmCrFe

a = 879.66c = 455.82

at 831°C, x = 0.45at 510°C, x = 0.51at 700°C, 0.44 ≤ x ≤ 0.5 [1987And]at 650 - 790°C and x = 0.495 [V-C2]



Cu Cr Fe

L, L', L'' 1523 c1(critical) L 48.0 33.0 19.0

L', L'', α 1516 c2(critical) L', L''α

48.81.0

41.584.4

9.714.6

L', L'', γ 1431 c3(critical) L', L''γ

53.414.0

1.72.1

44.983.9

(continued)

Cr–Cu–Fe 9




Cu Cr Fe

L' + γ ⇌ L'' + α 1396 U1 L'γL''α

25.111.177.08.6

9.78.62.9

10.2

65.280.320.181.2

L' ⇌ L'' + α 1362 e1(min) L'L''α

18.884.26.3

25.75.7

29.7

55.510.164.0

L + γ ⇌ α + (Cu) 1088 U2 Lγα(Cu)

96.75.73.9

96.3

0.714.718.30.3

2.679.677.83.4

Table 4. Equilibrium Compositions of the α, γ and Liquid Phases in the Three-Phase Region

T [°C] Phase Composition (at.%)

Cu Cr Fe

1300 Lαγ

89.77.19.5

1.813.411.3

8.579.579.3

1250 Lαγ

92.26.38.5

1.614.912.3

6.278.879.2

1200 Lαγ

94.15.57.6

1.216.113.2

4.778.479.2

1100 Lαγ

96.54.05.9

0.818.214.5

2.777.879.6

Table 5. Equilibrium Compositions of the α, γ and (Cu) Phases in the Three-Phase Region


Cu Cr Fe

1050 (Cu)αγ

96.93.34.9

0.218.614.8

2.978.180.3

1000 (Cu)αγ

97.52.63.9

0.218.815.0

2.378.681.1

(continued)

10 Cr–Cu–Fe



Cu Cr Fe

900 (Cu)αγ

98.51.62.5

0.117.814.4

1.480.683.1

850 (Cu)αγ

98.81.22.0

0.115.813.3

1.183.084.7

Table 6. Investigations of the Cr-Cu-Fe Materials Properties

Reference Method/Experimental Technique Type of Property

[1967Sal] Optical microscopy Dihedral angle measurements at 1 to 15mass% Cu, 0 to 9.38 mass% Cr and at 900 to1250°C

[1970Ahm,1974Zap]

Diffusion couple method

Corrosion tests, micro-hardness tests

Homogenization degree of sintered of Cr-Cu-Fe alloys with 2 to 18 mass% Cr and 2 to4 mass% Cu at 1050 to 1410°C, diffusioncoefficient and activation energy ofdiffusion of Cr in α and γ phases.Microhardness and corrosion resistance

[1975Wan] Diffusion couple method Diffusion coefficients of copper andchromium in liquid Cr-Cu-Fe alloys at1550°C

[2000Gen] Optical microscopy, friction tests, resistancetests, hardness tests

Mechanical and electrical properties of lasersurface remelting/resolidifying processedCr-Cu-Fe contact material manufactured bypowder metallurgy

[2001Fer1] Resistivity measurements Temperature dependence of resistivity of Curich alloys at < 1 mass% Fe, < 0.6 mass% Crand at 500 to 1050°C

[2001Fer2] Electrical conductivity measurements,optical microscopy, TEM, WDS, hardnesstests

Microstructure, electrical and mechanicalproperties of Cu - 0.7 mass% Cr - 0.3 mass%Fe, Cu - 0.7 mass% Cr - 0.8 mass% Fe andCu - 0.7 mass% Cr - 2.0 mass% Fe alloys

[2001Fer3] TEM, WDS, hardness tests, resistivitymeasurements

Microstructure, electrical and mechanicalproperties of Cu - 0.7 mass% Cr - 2.0 mass%Fe cold-worked alloy

[2001Les] Stationary-drop method tests of wettability,optical microscopy, electron microscopy,EPMA

Adhesive characteristics and formation ofphase boundaries in Cu rich compositematerials manufactured by powdermetallurgical process at 1100-1300°C

(continued)

Cr–Cu–Fe 11




[2001Son] Mechanical properties tests usingextensometer, optical microscopy, SEM,resistivity tests using four-probe technique

Microstructure, mechanical and electricalproperties of deformation processed Cu - 1.2mass% Cr - 9 mass% Fe microcompositewires

[2002Kim] Mechanical strength tests using machineequipped with extensometer, Vickers micro-hardness tests, image analysis, TEM

Ultimate tensile strength and theconductivity of deformation processedCu - 1.2 mass% Cr - 9 mass% Femicrocomposite wires

[2004Wan] Optical microscopy Core type macroscopic morphologies

Table 7. Diffusion Coefficient and Activation Energy of Diffusion of Cr in α and γ Phases

DiffusionCouple

Diffusion Coefficient [cm2· s–1] at Sintering Temperature Qα

[kJ·g-atom–1]

Qγ

[kJ·g-atom–1]1410°C 1250°C 1050°C

Dα Dα Dγ Dα Dγ

(Fe-Cr)/Fe 1.7·10–7 3.2·10–8 4.6·10–10 1.9·10–9 4.2·10–11 239.3248 212.1288

(Fe-Cr + 2%Cu)/Fe

2.2·10–7 2.8·10–8 1.7·10–9 8.2·10–10 1.7·10–10 298.7376 194.9744

(Fe-Cr)/(Fe + 2%Cu)

1.2·10–7 1.3·10–8 4.2·10–10 3.0·10–10 3.0·10–11 308.7792 233.0488

(Fe-Cr + 2%Cu)/(Fe + 2%Cu)

1.5·10–7 2.8·10–8 5.0·10–10 1.4·10–9 6.5·10–11 251.8768 176.9832

(Fe-Cr + 4%Cu)/Fe

2.2·10–7 4.2·10–8 3.6·10–9 3.0·10–9 3.9·10–10 227.6096 195.3928

(Fe-Cr)/(Fe + 4%Cu)

1.5·10–7 2.7·10–8 8.5·10–10 1.4·10–9 9.4·10–11 256.0608 187.4432

(Fe-Cr + 4%Cu)/(Fe + 4%Cu)

1.5·10–7 1.6·10–8 8.0·10–10 3.1·10–10 6.5·10–11 329.6992 216.7312

(Fe-Cr)/Cu/Fe - 1.5·10–8 3.8·10–9 - - 239.3248 212.1288

12 Cr–Cu–Fe


Fig. 1. Cr-Cu-Fe. Reaction scheme

Cr–Cu–Fe 13



Fig. 2a. Cr-Cu-Fe. Calculated liquidus projection

14 Cr–Cu–Fe


Fig. 2b. Cr-Cu-Fe. Enlarged part of the liquidus projection

Cr–Cu–Fe 15



Fig. 3. Cr-Cu-Fe. Solidus surface projection

16 Cr–Cu–Fe


Fig. 4. Cr-Cu-Fe. Calculated isotherms of the cupola of stable and metastable miscibility gap of the liquid phase.Dashed lines are the tie lines at 1127°C.

Cr–Cu–Fe 17



Fig. 5. Cr-Cu-Fe. Calculated vertical sections of the miscibility gap. Solid lines correspond to the stable range anddashed lines correspond to the metastable range, section at 4 mass% Cr is taken from [2004Wan]

18 Cr–Cu–Fe


Fig. 6. Cr-Cu-Fe. Calculated isothermal section at 1500°C

Cr–Cu–Fe 19




20 Cr–Cu–Fe



Cr–Cu–Fe 21




22 Cr–Cu–Fe


Fig. 10. Cr-Cu-Fe. Calculated temperature-composition section at 1 mass% Cr, plotted in at.%

Cr–Cu–Fe 23




24 Cr–Cu–Fe



Cr–Cu–Fe 25




26 Cr–Cu–Fe



Cr–Cu–Fe 27



Fig. 15. Cr-Cu-Fe. Calculated temperature-composition section at 20 at.% Cr

28 Cr–Cu–Fe



Cr–Cu–Fe 29




30 Cr–Cu–Fe


Fig. 18. Cr-Cu-Fe. Calculated temperature-composition section at 1 mass% Cu, plotted in at.%

Cr–Cu–Fe 31



Fig. 19a. Cr-Cu-Fe. Calculated temperature-composition section at 2 mass% Cu, plotted in at.%

32 Cr–Cu–Fe


Fig. 19b. Cr-Cu-Fe. Enlarged part of the calculated temperature-composition section at 2 mass% Cu, plotted in at.%

Cr–Cu–Fe 33




34 Cr–Cu–Fe



Cr–Cu–Fe 35



Fig. 22a. Cr-Cu-Fe. Calculated temperature-composition section at 20 at.% Cu

36 Cr–Cu–Fe


Fig. 22b. Cr-Cu-Fe. Enlarged part of the calculated temperature-composition section at 20 at.% Cu

Cr–Cu–Fe 37



Fig. 23. Cr-Cu-Fe. Composition of (Cu) solid solution in alloys quenched from aging temperature after resistivitymeasurements

38 Cr–Cu–Fe


Fig. 24. Cr-Cu-Fe. Variations of ultimate tensile strength (U.T.S.) and ductility of Cu 1.2 mass% Cr 9 mass% Fe wiresas a function of heat treatment temperature

Cr–Cu–Fe 39



References[1939Mor] Moriwaki, K., “The Equilibrium Diagram of the Ternary System, Iron-Chromium-Copper”,

Tetsu to Hagane, 25, 396–403 (1939) (Phase Diagram, Experimental, 3)[1967Sal] Salter, W.J.M., “Effect of Chromium on Solubility of Copper in Meld Steel”, J. Iron Steel

Inst. Jpn., 205, 1156–1160 (1967) (Phase Diagram, Experimental, 30)[1970Ahm] Ahmed, M., Thuemmler, F., Zapf, G., “Metallographic Investigations of the Fe-Cr-Cu Alloys

Prepared by Powder Metallurgical Methods”, Arch. Eisenhuettenwes., 41, 797–803 (1970)(Phase Diagram, Mechan. Prop., Experimental, 25)

[1971Yam] Yamaguchi, M., Ymakoshi, Y., Mima G., “Miscibility Gap in the Iron-Chromium-Metal(X=Copper, Manganese, Molybdenum, Nickel, Vanadium, Silicon and Aluminium) System”,Proc. Int. Conf. Sci. Technol. Iron Steel, Tokyo, 2, 1015–1019 (1971) (Phase Relations,Experimental, 35)

[1974Sig] Sigworth, G.K., Elliott, J.F., “The Thermodynamics of Liquid Dilute Iron Alloys”, Met. Sci.,8, 298–310 (1974) (Thermodyn., Review, 249)

[1974Zap] Zapf, G., Ahmed, M., “Research Report of the State of Noth Rhine-Westphalia, No. 2430:Investigations on the Sinter Alloying of the Binary Iron-Chromium System and the TernaryIron-Chromium-Copper System (Forschungsbericht Des Landes Nordrheim)”, WestdeutscherVerlag, Opladen, Germany, 149 pp (1974) (Phase Diagram, Mechan. Properties, Experimen-tal, 88)

[1975Wan] Wanibe, Y., Takagi, T., Sakao, H., “Coupling Phenomenon in the Ternary Isothermal Diffu-sion of Liquid Iron-Chromium-Copper Alloys”, Arch. Eisenhuettenwes., 46(9), 561–565(1975) (Transport Phenomena, Experimental, 7)

[1979Cha] Chang, Y.A., Neumann, J.P., Mikula, A., Goldberg, D., “Cr-Cu-Fe” in “INCRA MonographSeries 6. Phase Diagrams and Thermodynamic Properties of Ternary Copper-Metall Sys-tems”, Uni. Wisconsin-Milwaukee, USA, 439–446 (1979) (Phase Diagram, Review, 6)

[1987And] Anderson, J.O., Sundman, B., “Thermodynamic Properties of the Cr-Fe System”, Calphad,11, 83–92 (1987) (Calculation, Phase Diagram, Thermodyn., Assessment) as cited in[2002Wan]

[1988Uen] Ueno, S., Waseda, Y., Jacob, K.T., Tamaki, S., “Theoretical Treatment of Interaction Para-meters in Multicomponent Metallic Solutions”, Steel Res., 59(11), 474–483 (1988) (Thermo-dyn., Theory, Calculation, 44)

[1990Ham] Haemaelainen, M., Jaaskelainen, K., Luoma, R., Nuotio, M., Taskinen, P., Teppo, O.A.,“Thermodynamic Analysis of the Binary Alloy Systems Cu-Cr, Cu-Nb and Cu-V”, Calphad,14(2), 125–137 (1990) (Calculation, Phase Diagram, Thermodyn., Assessment, 52)

[1993Hao] Hao, S.M., Jiang, M., “Cr-Cu-Fe”, Proc. 7thNat. Symp. Phase Diagrams, Chinese Phys. Soc.,Shanghai, 11–13 (1993) (Phase Diagram, Experimental, 0)

[1995Che] Chen, Q., Jin, Z., “The Fe-Cu System: a Thermodynamic Evaluation”, Metal. Mater. Trans.A, 26A(2), 417–426 (1995) (Calculation, Phase Diagram, Thermodyn., Assessment, 55)

[1997Oht] Ohtani, H., Suda, H., Ishida, K., “Solid/Liquid Equilibria in Fe-Cu Based Ternary Systems”,ISIJ Int., 37(3), 207–216 (1997) (Experimental, Calculation, Phase Relations, Review, Ther-modyn., 47)

[1998Mie] Miettinen, J., “Approximate Thermodynamic Solution Phase Data for Steels”, Calphad, 22(2), 275–300 (1998) (Review, Calculation, 83)

[2000Gen] Geng, H.R., Liu, Y., Chen, C.Z., Sun, M.H., Gao, Y.Q., “Laser Surface Remelting of Cu-Cr-Fe Contact Material”, Mater. Sci. Technol., 16(5), 564–567 (2000) (Experimental, Morphol-ogy, Mechan. Prop., Electr. Prop.) cited from abstract

[2001Fer1] Fernee, H., Nairn, J., Atrens, A., “Cu-Rich Corner of the Cu-Fe-Cr Phase Diagram”, J. Mater.Sci. Lett., 20, 2213–2215 (2001) (Experimental, Electr. Prop., Phase Relations, 6)

[2001Fer2] Fernee, H., Nairn, J., Atrens, A., “Cold Worked Cu-Fe-Cr Alloys”, J. Mater. Sci., 36(22),5497–5510 (2001) (Electr. Prop., Mechan. Prop., Experimental, 19)

[2001Fer3] Fernee, H., Nairn, J., Atrens, A., “Precepitation Hardening of Cu-Fe-Cr Alloys - Part 1 -Mechanical and Electrical Properties”, J. Mater. Sci., 36(11), 2711–2719 (2001) (Experimen-tal, Electr. Prop., Mechan. Prop., 19)

40 Cr–Cu–Fe


[2001Les] Lesnik, N.D., Minakova, R.V., Khomenko, E.V., “Chromium-Copper System: AdhesionCharacteristics, Doping, the Structure of Phase Boundary and Composites”, Powder Metall.Met. Ceram., 40(7–8), 432–440 (2001) (Experimental, Morphology, Phys. Prop., InterfacePhenomena, 12)

[2001Mil] Milano, J. Llois, A.M., “From Direct to Inverse GMR: Introduction of Cr in Fe/Cu Superlat-tices”, J. Mag. Mag. Mater., 226, 1755–1757 (2001) (Calculation, Magn. Prop.) cited fromabstract

[2001Son] Song, J.S., Hong, S.I., Kim, H.S., “Heavily Drawn Cu-Fe-Ag and Cu-Fe-Cr Microcompo-sites”, J. Mat. Proc. Tech., 113(1–3), 610–616 (2001) (Experimental, Mechan. Prop., Mor-phology, 21)

[2002Ans] Ansara, I., Ivanchenko, V., “Cr - Cu (Chromium - Copper)”, MSIT Binary Evaluation Pro-gram, in MSIT Workplace, Effenberg, G. (Ed.) MSI, Materials Science International ServicesGmbH, Stuttgart; Document ID: 20.19588.1.20, (2002) (Crys. Structure, Phase Diagram,Assessment, 29)

[2002Kim] Kim, Y.S., Song, J.S., Hong, S.I., “Thermo-Mechanical Properties of Cu-Fe-Cr Microcompo-sites”, J. Mat. Proc. Tech., 130, 278–282 (2002) (Experimental, Mechan. Prop., Morphology,23)

[2002Mil] Milano, J., Llois, A.M., Steren, L.B., “Combined Impurity and Band Effects on the Appear-ance of Inverse Giant Magnetoresistance in Cu/Fe Multilayers with Cr”, Phys. Rev. B, 66(13),Art. No. 134405 (2002) (Calculation, Magn. Prop.) cited from abstract

[2002Rag] Raghavan, V., “Cr-Cu-Fe (Chromium-Copper-Iron)”, J. Phase Equilib., 23(3), 257–258(2002) (Review, Phase Relations, 5)

[2002Wan] Wang, C.P., Liu, X.J., Ohnuma, I., Kainuma, R., Ishida, K., “Phase Equilibria in Fe-Cu-X(X: Co, Cr, Si, V) Ternary Systems”, J. Phase Equilib., 23(3), 236–245 (2002) (Experimen-tal, Calculation, Phase Diagram, Thermodyn., #, 38)

[2004Wan] Wang, C.P., Liu, X.J., Takaku, Y., Ohnuma, I., Kainuma, R., Ishida, K., “Formation of Core-Type Macroscopic Morphologies in Cu-Fe Base Alloys with Liquid Miscibility Gap”,Metall.Mater. Trans. A, 35A(4), 1243–1253 (2004) (Experimental, Calculation, Morphology, PhaseDiagram, Thermodyn., 31)

[2007Tur] Turchanin, M., Agraval P., “Cu - Fe (Copper - Iron)”, MSIT Binary Evaluation Program, inMSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH,Stuttgart; Document ID: 20.11107.1 (2007) (Crys. Structure, Phase Diagram, Thermodyn.,Assessment, 36)

[Mas2] Massalski, T.B. (Ed.) Binary Alloy Phase Diagrams, 2nd edition. ASM International, MetalsPark, Ohio (1990)


Cr–Cu–Fe 41



Chromium – Iron – Hydrogen

Pierre Perrot

Introduction

The harmful effect of hydrogen on mechanical properties of Cr-Fe alloys has long been known. Thesealloys, in the presence of H do not form hydrides, but interstitial solid solutions which play an importantrole in steelmaking because of embrittlement of steel products. Most of the investigations have been direc-ted towards the solubility of H in Cr-Fe melts, solid alloys and on the effect of hydrogen in mechanicalproperties of Cr-Fe alloys. The main experimental results are gathered in Table 1.

Binary Systems

The well known Cr-Fe system is accepted from [Mas2]. The Fe-H system under 0.1 MPa of hydrogen pres-sure is accepted from [1990San]. The Fe-H system has been carefully investigated by [2003Fuk] up to 10GPa of hydrogen pressure and 1500°C. The hydrogen pressure temperature diagram, presents a drastic low-ering of the melting point down to 800°C at 3 GPa. The triple point α-γ-ε of iron is shifted from 8.4 GPaand 430°C for pure iron to 5 GPa and 260°C for the Fe-H system. The hydrogen solubility under high pres-sure has been investigated by [2005Hir] in the γ region. Under 2.5 GPa, the hydrogen solubility increaseswith the temperature (from H/Fe = 0.12 at 700°C to H/Fe = 0.2 at 1000°C; under 4 GPa, H/Fe = 0.4 bet-ween 400 and 1000°C; under 6 GPa, the hydrogen solubility decreases when the temperature increases(from H/Fe = 0.7 at 700°C to H/Fe = 0.6 at 1000°C). The Cr-H system has been assessed by [1991Ven].The hydrogen solubility in liquid and solid chromium are given by:log10(at.% H) = 1.97 – 5530 K / T + 0.5 log10 (pH2 / bar) (in liquid Cr, T > 1800 K, pH2 < 100 bar)log10(at.% H) = 0.13 – 2620 K / T + 0.5 log10 (pH2 / bar) (in solid Cr, T > 1000 K, pH2 < 100 bar).Below 1000 K, the measured solubilities are higher than that predicted by the above equation, probablyowing the trapping of H by lattice defects. Chromium hydrides CrH and CrH2 are stable under hydrogenpressures higher than 20 kbar (2 MPa). The existence of CrH3 is doubtful. The temperature-pressure phasediagram proposed for the Cr-H system [2005Fuk] in given in Fig. 1.

Solid Phases

No ternary compounds are known at room temperature. However hydrides (Fe1–xCrx)Hy (x = 0.05, 0.25 and0.50; y < 1) have been prepared at 325°C under hydrogen pressure up to 7 GPa. They are metastable at roomtemperature and present a hexagonal lattice structure like CrH [2002Ant]. The solid phases are presented inTable 2. The H solubility in solid alloys was measured by [1965Sch] between 400 and 1000°C and evalu-ated at 1200 and 1400°C by [1966Bur].


A number of investigations on the solubility of H in liquid Cr-Fe alloys are reported, as shown in Table 1.Investigations on the solubility of H in solid Cr-Fe alloys are scarcer [1965Sch, 1975Col, 1977Arc,1977Bes] but indicate also that Cr increases its solubility at a given H2 pressure as shown in Fig. 2 whichis taken from [1965Sch]. The solubility of hydrogen is higher in the γ phase (lower chromium content at900 and 1000°C) than in the α phase (higher chromium content at 900 and 1000°C). The solubility of Hin liquid Cr-Fe alloys between 1550 and 1750°C under 1 bar of hydrogen pressure is given in Fig. 3. Itrepresents the best curves proposed by [1981Sch] from the compilation of former measurements. Up to100 bar (10 MPa) of hydrogen pressure, the Sievert’s law may be applied and the hydrogen solubilitymay be expressed by: xH2 (under p bar) = xH2 (under 1 bar) · (pH2 / bar)

1/2.

Cr–Fe–H 1



Thermodynamics

The entropy and enthalpy of dissolution of H gas into α, γ and σ(Fe,Cr) alloys, evaluated from the hydro-gen solubility measurements carried out by [1975Col, 1977Bes] are presented in Table 3. It must be pointedout that the enthalpy of dissolution of ½ H2 in solid alloy may be obtained from the values given by[1975Col] in Table 3 by adding 218 kJ·mol–1, which represents half of the dissociation enthalpy of H2.The dissolution of H2 in Cr-Fe alloys is an endothermal phenomenon: the hydrogen in alloys increases withthe temperature [1977Kar]. Measurements of [1977Arc] under higher hydrogen pressures (5 to 80 MPa)confirms the observation of an enthalpy of dissolution weakly dependent on the crystal structure of thealloy.The Wagner first order interaction parameter deduced from the measured solubility of hydrogen in the liquidalloy, defined by eH

(Cr) = ∂ log10 fH(Cr) / ∂ (mass% Cr), with fH

(Cr) = (mass% H in pure Fe) / (mass% H inalloy), was found to be –0.031 at 1600°C [1961Mae], value probably too negative, not confirmed by latermeasurements [1964Gun]. The interaction parameter is more probably eH

(Cr) = – 0.0024 ± 0.0001 between1550 and 1650°C [1963Wei, 1964Gun, 1972Ngi, 1974Boo], whereas [1967Ban] proposes –0.0056.[1970Fuk, 1970Kat] pointed out that the result depends on the measurement method used. For instance,at 1600°C, the values given by the sampling method and by the Sievert’s method are –0.0038 and–0.0022, respectively. All authors observe that Chromium decreases the activity coefficient of hydrogenand thus increases its solubility in liquid iron [1965Bur, 1968Bez, 1972Ngi]. Such a behavior is explainedby the positive departure from ideal behavior shown by the Cr-Fe alloys [1975Col, 1977Arc]. [1999Din]calculates, at 1600°C εH

(Cr) = – 0.6, with εH(Cr) = ∂ log10 γH

(Cr) / ∂ xCr, and γH(Cr) = xH(in pure Fe) / xH(in

alloy), which is more negative than the experimental value εH(Cr) = – 0.3.


High chromium in martensitic steels have strengths comparable to austenitic steels up to 500°C and presenta greater resistance to the effect of radiation damage. For this reason, it is important to obtain parameterssuch as the solubility and permeability of hydrogen and its isotopes.The diffusivity of hydrogen in (αFe) is very slightly lowered by alloying with Cr up to 30 mass% Cr[1965Sch]. The diffusivity was measured at 2·10–4 and 5·10–5 cm2·s–1 at 850 and 350°C respectivelybetween 0 and 20 mass% Cr in the alloy. At room temperature, the diffusivity decreases sharply with theCr content of the alloy, from 4·10–7 to 3·10–9 cm2·s–1 for pure Fe and Fe-30 mass% Cr alloy, respectively.Such a trend is experimentally confirmed by [1986Ven]. However, the diffusivity at room temperature aremeasured at 6.2 · 10–5 and 2.8 · 10–5 cm2·s–1 for pure Fe and Fe-5 mass% Cr. Hydrogen diffusivity in alloysseems strongly dependent on the sample preparation and of the input fugacity of hydrogen. [1997For] useda gas phase permeation technique to measure the deuterium diffusion between 194 and 465°C under 50 kPaof D pressure. The diffusion cannot be described by an Arrhenius law; below 330°C, the diffusion is greatlyreduced by trapping effects. The following expressions are proposed:Below 330°C: D / cm2·s–1 = 5.07 · 10–4 exp (– 13 500 / R T)Above 330°C: D / cm2·s–1 = 4.30 · 10–4 exp (– 13 000 / R T)The diffusion of H is obtained by multiplying the preexponential factor by 1.4 to take into account the iso-topic effect.

Table 1. Investigations of the Cr-Fe-H Phase Relations, Structures and Thermodynamics


[1961Mae] Hydrogen solubility measurements by thesampling method

1600°C, < 5 mass% Cr, 0.1 MPa ofhydrogen pressure

[1963Wei] Hydrogen solubility measurements by theSievert’s method


(continued)

2 Cr–Fe–H



[1964Gun] Hydrogen solubility measurements by the hotvolume method

1550-1650°C, < 30 mass% Cr, 0.1 MPaof hydrogen pressure

[1965Bag] Hydrogen solubility measurements by adifferential solubility method


[1965Sch] Solubility and mobility measured by aneffusion method

25-1000°C, < 30 mass% Cr

[1967Ban] Hydrogen solubility measurements by theSievert’s method

1550-1670°C, 9-50 mass% Cr, 0.1 MPaof hydrogen pressure

[1968Bez] Hydrogen solubility measurements by theSievert’s method

1600-1900°C, 40-70 mass% Cr,< 0.1 MPa of hydrogen pressure

[1970Fuk,1970Kat]

Hydrogen solubility measured by Sievert’sand sampling methods


[1972Ngi] Hydrogen solubility measured by Sievert’sand sampling methods

1550-1600°C, < 15 mass% Cr, < 0.1 MPaof hydrogen pressure

[1974Boo] Hydrogen solubility measurements by aconstant volume method


[1975Col] Hydrogen solubility measurements by thesampling method

500-1200°C, 5 to 85 mass% Cr, 0.1 MPaof hydrogen pressure

[1977Arc] Hydrogen solubility measurements by thesampling method

400-700°C, 3.5 to 25 mass% Cr, 5 to80 MPa of hydrogen pressure

[1977Bes] Hydrogen solubility measurements by thesampling method


[1986Ven] Hydrogen diffusion in solid Cr-Fe alloys Room temperature, < 5 mass% Cr

[1997For] Deuterium diffusion and solubilitymeasurements

194-465°C, 10 mass% Cr, 50 kPa ofdeuterium pressure

[2002Ant] Mössbauer study of ternary (Cr,Fe,H) alloys 325°C, 5, 25 and 50 at.% Cr, < 7 GPa ofhydrogen pressure





Comments/References

α, (αFe,Cr)(αFe)< 1538(Cr)< 1983

cI2Im�3mW

a = 286.65

a = 288.48

pure Fe at 20°C [Mas2] (A2 structure)dissolves from 1 ppm H at 25°C

to 0.1 mass% at 1538°C [2002Ant]pure Cr at 20°C [Mas2]

(continued)

Cr–Fe–H 3






Comments/References

γ, (γFe)1394 - 590

cF4Fm�3mCu

a = 293.16 at 915°C [Mas2, V-C2]

σ, CrFe830 - 440

tP30P42/mnmσCrFe

a = 879.4c = 455.2

44.5-50 at.% Cr[Mas2]

CrH hP4P63/mmcAnti NiAs

a = 271.7 ± 0.3c = 442.3 ± 0.5

47 to 50 at.% H [1991Ven]

CrH2 cF12Fm�3mCaF2

a = 385.9 ± 0.2 55 to 67 at.% H[1991Ven]

Table 3. Thermodynamic Data of Reaction or Transformation

Reaction or Transformation Temperature[°C]

Quantity, per mole of atoms[kJ, mol, K]

Comments

H (gas) ⇌ H (dissolved in α alloy) 1000 ΔdissH = – 170 ± 5ΔdissS

xs = 0.030 ± 0.005[1975Col]

H (gas) ⇌ H (dissolved in γ alloy) 1000 ΔdissH = – 175 ± 2ΔdissS

xs = 0.019 ± 0.005[1975Col]

H (gas) ⇌ H (dissolved in σ alloy) 750 ΔdissH = – 155 ± 2ΔdissS

xs = 0.033 ± 0.005[1975Col]

½ H2 (gas) ⇌ H (dissolved in γ alloy) 1092 ΔdissH = 30.4 ± 1.5ΔdissS = 0.038 ± 0.001

[1977Bes]

4 Cr–Fe–H


Fig. 1. Cr-Fe-H. The Cr-H binary system under high hydrogen pressures

Cr–Fe–H 5



Fig. 2. Cr-Fe-H. The hydrogen solubility in solid (Fe,Cr) alloys from 400 to 1000°C under 0.1 MPa of hydrogenpressure

6 Cr–Fe–H


Fig. 3. Cr-Fe-H. The hydrogen solubility in liquid (Fe,Cr) alloys at 1550, 1650 and 1750°C under 0.1 MPa ofhydrogen pressure

Cr–Fe–H 7



References[1961Mae] Maekawa, S., Nakagawa, Y., “The Effect of some Alloying Elements on the Solubility of

Hydrogen in Liquid Iron” (in Japanese), Nippon Kinzoku Gakkai-shi, 25(9), 577–580(1961) (Experimental, Phase Relations, Thermodyn., 6)

[1963Wei] Weinstein, M., Elliott, J.F., “Solubility of Hydrogen in Liquid Iron Alloys”, Trans. Met. Soc.AIME, 227, 382–393 (1963) (Experimental, Phase Relations, Thermodyn., 27)

[1964Gun] Gunji, K., Ono, K., Aoki, Y., “The Effect of Various Elements on the Solubility of Hydrogenin Liquid Pure Iron”, Trans. Nat. Res. Inst. Met. (Jpn.), 6(5), 209–213 (1964), translated fromJ. Jpn. Inst. Met., 28, 64-68, (1964) (Experimental, Phase Relations, Thermodyn., 7)

[1965Bag] Bagshaw, T., Engledow, D., Mitchell, A., “Solubility of Hydrogen in Some Liquid Iron-basedAlloys”, J. Iron Steel Inst., 203, 160–165 (1965) (Experimental, Phase Relations, Thermodyn,28)

[1965Bur] Burylev, B.P., “Solubility of Hydrogen in Liquid Iron Alloys” (in Russian), Izv. Vyss. Uchebn.Zaved., Chern. Metall., 8(2), 17–22 (1965) (Review, Phase Relations, Thermodyn., 13)

[1965Sch] Schwarz, W., Zittter H., “Solubility and Diffusion of Hydrogen in Iron Alloys” (in German),Arch. Eisenhuettenwes., 36(5), 343–349 (1965) (Experimental, Phase Relations, TransportPhenomena, 16)

[1966Bur] Burylev, B.P., “Solubility of Hydrogen in Solid Iron Alloys”, Russ. J. Phys. Chem., 40(4),442–445 (1966) (Phase Relations, Calculation, 12)

[1967Ban] Ban-ya, S., Fuwa, T., Ono, K., “Solubility of Hydrogen in Liquid Iron Alloys” (in Japanese),Tetsu to Hagane, 53(2), 101–116 (1967) (Experimental, Phase Relations, Thermodyn., 18)

[1968Bez] Bezobrazov, S.V., Danilovich, Yu.A., Charushnikova, G.V., Morozov, A.N., “Solubility of Hin Molten Fe-Cr” (in Russian), Izv. Akad. Nauk SSSR, Met., (3), 65–70 (1968) (Experimental,Kinetics, 12)

[1970Fuk] Fukuda, S., Sugiyama, T., Furukawa, T., Kato, E., “Solubility of Hydrogen in Liquid IronAlloys”, Rep. Casting Research Lab.,Waseda Univ., (21), 35–46 (1970) (Calculation, Experi-mental, Phase Relations, Thermodyn., 22)

[1970Kat] Kato, E., Fukuda, S., Sugiyama, T., Furukawa, T., “Solubility of Hydrogen in LiquidIron Alloys” (in Japanese), Tetsu to Hagane, 56, 521–535 (1970) (Experimental, Phase Rela-tions, 28)

[1972Ngi] Ngia, N., Yavoyskiy, V.I., Kosterev, L.B., Afanas’yev, M.I., “Hydrogen Solubility in BinaryIron-Base Alloys”, Russ. Metall., (4), 11–15 (1972), translated from Izv. Akad. Nauk SSSR,Met., (4),18-22, (1972) (Experimental, Phase Relations, Thermodyn., 19)

[1974Boo] Boorstein, W.M., Pehlke, R.D., “Measurement of Hydrogen Solubility in Liquid Iron AlloysEmploying a Constant Volume Technique”, Metall. Trans., 5, 399–405 (1974) (Calculation,Experimental, Phase Relations, 29)

[1975Col] Coldwell, D.M., McLellan, R.B., “Thermodynamic Properties of Fe-Cr-H Ternary Solid Solu-tions”, Acta Metall., 23, 57–61 (1975) (Experimental, Calculations, Thermodyn., 28)

[1977Arc] Archakov, Yu.I., Vanina, T.N., “Effect of Chromium on the Solubility in of Hydrogen in Iron atHigh Temperatures and Pressures”, Inorg. Mater., 50(6), 1166–1169 (1977), translated from Zh.Priklad. Khim., 50(6), 1209–1212, (1977) (Experimental, Phase Relations, Thermodyn., 8)

[1977Bes] Bester, H., Lange, K.W., “H Solubility in Fe and in Liquid Fe-Mn, Fe-Cr, and Fe-Si Alloys”(in German), Stahl Eisen, 97, 1037–1039 (1977) (Experimental, Phase Relations, Thermo-dyn., 53)

[1977Kar] Karamysheva, G.A., Men, A.N., “Calculation of the Concentration-Dependence of HydrogenSolubility in Binary Alloys Using the Cluster Component Method”, Russ. Metall., (2), 82–83(1977), translated from Izv. Akad. Nauk SSSR, Met., (2), 95–96 (1977) (Calculation, PhaseRelations, 5)

[1981Sch] Schuermann, E., Kaettlitz, W., “Equivalent Effect of the Alloying Elements on the Concentration-and Temperature-Dependent Hydrogen Solubility in Iron-rich Ternary and Multicompon-ent Melts” (in German), Arch. Eisenhuettenwes., 52(8), 295–301 (1981) (Calculation, PhaseRelations, 20)

8 Cr–Fe–H


[1986Ven] Veniali, F., Szklarska-Smialowska, Z., “A Study of the Diffusion and Trapping of Hydrogen inFe-3Cr and Fe-5Cr Alloys”,Mater. Chem. Phys., 15(6), 545–557 (1986) (Experimental, Inter-face Phenomena, Kinetics, 19)

[1990San] San Martin, A., Manchester, F.D., “The Fe-H (Iron-Hydrogen) System”, Bull. Alloys PhaseDiagrams, 11(2), 173–184 (1990) (Phase Diagram, Phase Relations, Review, Thermodyn., 86)

[1991Ven] Venkatraman, M., Neumann, J.P., “The Cr-H (Chromium-Hydrogen System)”, J. Phase Equi-lib., 12(6), 672–677 (1991) (Phase Relation, Crys. Structure, Thermodyn., Review, 58)

[1997For] Forcey, K.S., Iordanova, I., Yaneva, M., “The Diffusivity and Solubility of Deuterium in aHigh Chromium Martensitic Steel”, J. Nucl. Mater., 240, 118–123 (1997) (Experimental,Phase Relations, Transport Phenomena, 12)

[1999Din] Ding, X., Fan, P., Wang, W., “Thermodynamic Calculation for Alloy Systems”, Metall. Mat.Trans. B, 30B(2), 271–277 (1989) (Phase Relations, Thermodyn., Calculation, 18)

[2002Ant] Antonov, V.E., Baier, M., Dorner, B., Fedotov, V.K., Grosse, G., Kolesnikov, A.I., Ponya-tovsky, E.G., Schneider, G., Wagner, F.E., “High-Pressure Hydrides of Iron and its Alloys”,J. Phys.: Condens. Matter, 14, 6427–6445 (2002) (Crys. Structure, Experimental, Phase Rela-tions, Review, 46)

[2003Fuk] Fukai, Y., Mori, K., Shinomiya, H., “The Phase Diagram and Superabundant Vacancy Forma-tion in Fe-H Alloys under High Hydrogen Pressures”, J. Alloys Compd., 348, 105–109 (2003)(Phase Diagram, Phase Relations, Experimental, 42)

[2005Fuk] Fukai, Y., “The Structure and Phase Diagram of M-H Systems at High Chemical PotentialHigh Pressure and Electrochemical Synthesis”, J. Alloys Compd., 404/406, 7–15 (2005)(Phase Diagram, Phase Relations, Thermodyn., Review, 40)

[2005Hir] Hiroi, T., Fukai, Y., Mori, K., “The Phase Diagram and Superabundant Vacancy Formation inthe Fe-H Alloy Revisited”, J. Alloys Compd., 404/406, 252–255 (2005) (Phase Diagram,Experimental, 42)



Cr–Fe–H 9



Chromium – Iron – Manganese

Jan Vreštál

Introduction

The first extensive investigation of the Cr-Fe-Mn system was the determination of the polytherms at Cr con-tent up to 20 mass% and Mn up to 30 mass% by [1934Koe] which made possible to determine the solidusunivariant reaction lines and to draw a hypothetical liquidus projection relied upon interpolation betweenthe binary axes. [1938Bur] deduced the isothermal section at 650°C from microscopic examination of alloysannealed for very long period of time at temperatures between 600°C and 1000°C. In this section the σphase region was denoted. [1939Sch] confirmed the σ phase region in their isothermal section at 700°C.Unfortunately, their work was performed using alloys of the low-purity. Further experimental results werereported by [1949Gri], working with materials of medium purity (99.9 mass%). Their data enable interpola-tion of the σ phase boundaries over a wide range of temperatures. In spite of the X-ray studies performed,the structure of the σ phase was not exactly determined.Later studies were performed with materials of higher purity: [1968Tav] investigated 700°C and 1100°Cisothermal sections using materials with the total amount of carbon and nitrogen below 0.08 mass%.[1975Shv1] studied γ-σ equilibria. [1975Shv2] studied equilibria γ-α with materials containing0.05 mass% C.[1971Kir] and [1974Kir] determined precisely the γ-α equilibria using metals of the purity 99.99 mass% at750-950°C. Isothermal sections at 800-1100°C in the Fe rich part of the diagram (up to 50 mass% Fe) werepresented also in [1973Kra], [1989Oka]. They agree reasonably with the results of [1971Kir] and [1974Kir].So, the phase boundaries determined in these works are considered very accurate and serve as a guide to thereliability of other data.The γ-(γ+σ) equilibrium which is of technological interest, was studied in [1988Abe, 1989Oka, 1990Yuk,1990Mur] using metals of 99.9 mass% purity. The isothermal section experimentally determined by[1988Abe] at 650°C is different from the old data by [1939Sch]. [1989Oka] investigated this equilibriumand found the γ-(γ+σ) phase boundary between 10 and 12 mass% Cr at 15 mass% Mn, while [1990Yuk]reported rather different results for the γ-(γ+σ) phase boundary: between 4 and 6 mass% Cr at the sameconditions. The only apparent difference between these works was the materials used by [1990Yuk](99.9 mass%) and that of [1989Oka] (99.98 mass%) which was, therefore, preferred.The addition of 10 mass% Cr in the Fe-14, 18, 22 Mn (mass%) alloys decreases the As(ε-γ) temperature byabout 45 to 60°C [1980Geo].The liquid-α equilibrium was determined in [1986Kun] using alloys of 99.9 mass% purity. This first experi-mental determination of the liquidus in the system was confined to the liquid-α equilibrium, although theparameters from the model used by [1986Kun] could be used to calculate the liquid-γ equilibrium, too.The only thermodynamic data for ternary alloys are presented in [1978Muk] for the liquid phase at 1570°Cwhere a variation of the activity coefficient of manganese with the manganese content was determined. Cri-tical assessment of the system may be found in [1958Pot], later in [1985Riv] and [1988Ray]. In theseassessments, the development of the knowledge of the binary phase diagrams caused that some results ofthe early works, though correct, had to be incorporated into isotherms very different from that envisagedby the original investigators.The first Calphad type assessment was made by [1974Kir] who was able to reproduce the α-γ equilibriumdata between 750-950°C based on their analysis of the previous evaluations of the Fe-Mn and Cr-Fe sys-tems [1973Kir]. An improvement of the parameters published in [1974Kir] was presented by [1977Hil].The first calculation of the whole phase diagram by the Calphad method has been carried out by[1993Lee]. [1994Rag, 2003Rag] based his critical evaluation on the data of [1986Kun, 1990Mur, 1993Lee].A summary of investigations on phase relations, structure and thermodynamics is given in Table 1.

Cr–Fe–Mn 1



Binary Systems

Phase diagram of the binary boundary system Cr-Fe is accepted from [Mas2]. The Cr-Mn binary phase dia-gram (Fig. 1) is taken from [1993Lee] relied on the solubility data of Cr in (γMn) from [1957Hel] and γ-αequilibrium data from [1974Kir]. However it should be mentioned that several simplifications were done by[1993Lee] in comparison with the evaluation of [Mas2]. Phases σ´ and σ´´ were considered as one phaseσ´´. The α´ and α´´ also were considered as one phase α´ neglecting their homogeneity ranges.Using new experimental thermodynamic data, the Fe-Mn system has been recently reassessed and the Cal-phad description was updated by [2004Wit]. This new reevaluation of the phase equilibria leads to consis-tently better fits to the available experimental data. There is however, a typographical error in [2004Wit] inthat the Mn rich invariant reaction involving the liquid phase is given as a peritectic type reaction in thetable of invariants. This reaction should be denoted as eutectic, as confirmed by [2007Wit], and have beenwritten as L ⇌ (δMn) + (γMn,γFe). Consequently, the Fe-Mn system is accepted from [2004Wit].

Solid Phases

The crystallographic data of the Cr-Fe-Mn phases and their stability ranges are listed in Table 2.It should be mentioned that since the phase diagrams from [1993Lee] are accepted for the Cr-Mn andCr-Fe-Mn, in the present evaluation, the α´ and α´´ as well as σ´ and σ´´ will be indicated as α´ or σ´´phase, respectively, in the phase diagrams and in the text. However, crystallographic data for α´´ as wellas for σ´ are also presented in Table 2 to take into account information from [Mas2].The (Cr) and (αδFe) form a continuous series of solid solutions (α phase) in the temperature rangeof ~ 830-846°C. The α solid solutions widely extend into the ternary system: 21Cr-49Fe-30Mn at 1000°C. It should be mentioned that in the Cr-Mn system, in spite of having the same crystal structure and highmutual solubility, the (αCr) and (δMn) solid solutions never form a continuous solution because they areseparated by a field of the stability of the σ phase. However, in the Cr-Fe-Mn ternary system, they aremerged together forming a continuous solid solution.The (γFe) and (γMn) form a continuous series of solid solutions at the temperatures above ~1100°C. How-ever, solubility of Cr in the γ phase at 1000°C does not exceed 15 at.%.The polymorphic modifications of Mn dissolve Cr and Fe. In (αMn), the maximum solubility is ~31 at.% Feand 6 at.% Cr at 650°C. Maximum solubility of Fe and Cr in (βMn) is ~20 at.% Fe and 7 at.% Cr at 1000°Cand ~27 at.% Fe and 3 at.% Cr at 800°C.An important question is whether the σ´´ phases in the Cr-Fe and Cr-Mn systems form a continuous solidsolution in the temperature range 650-800°C. Considering the results of [1949Gri] and [1975Shv1] it can beconcluded that a formation of a continuous solid solution is possible as it was assumed also by [1958Pot,1985Riv, 1988Ray] and it was confirmed also in [1988Abe] and [1989Oka]. The solubility of Fe in the σphase is about 40 at.% at 1200°C and about 47 at.% at 1000°C.


The only invariant equilibrium given in Table 3 is calculated using the dataset of [1993Lee].


The liquidus surface projection is accepted from [1993Lee]. Thermodynamic parameters of the liquid phasewere assessed by [1993Lee] using the experimental phase equilibria data of [1986Kun] and thermodynamicdata of [1978Muk]. Agreement between the calculated liquidus temperature and the experimental results of[1986Kun] is very good near the Fe-Mn binary side and reasonably good near the Cr-Fe binary side.Liquidus isotherms are calculated in the present work using the data set of [1993Lee]. The liquidus surfacealong with the liquidus isotherms is presented in Fig. 2. The monovariant line limiting the σ phase primarycrystallization field exhibits a maximum at a temperature above 1350°C.

2 Cr–Fe–Mn


Isothermal Sections

Isothermal sections are calculated using the dataset of [1993Lee], because this thermodynamic descriptionreproduces reasonably well most of the available experimental data.For the description of the α-γ equilibrium, the [1974Kir] and [1989Oka] data were taken. The phase bound-ary information was given more weight than the tie-lines direction for their extended scatter. The abovementioned phase equilibria data are represented by the calculations of [1993Lee] well.Considering the results of [1949Gri] and [1975Shv1] it can be concluded that a continuous series of solidsolutions between the σ´´ phases is possible as it was assumed also by [1958Pot, 1985Riv, 1988Ray]. Alsoexperimental data of [1982Lem] and [1990Mur] confirm this assumption.Two modifications, high-temperature (σ) and low-temperature (σ´´), were considered as independent phasesand corresponding parameters were evaluated. Unidentified ternary phase proposed by [1984Fri] was notconfirmed in any further study and it was not considered in the present evaluation. The experimental databy [1988Abe] and [1989Oka] show good agreement with the present calculation. But a discrepancy existsbetween the calculation and the experimental data of [1990Mur]. It partly comes from a disagreement in theFe-Mn binary system between [1990Mur] and [1993Lee]. Another cause of the disagreement may be a pos-sible presence of the metastable ε (hcp) phase after samples annealing in [1990Mur]. Agreement with thedata of [1949Gri] is poor but their annealing time at 600°C was only 170 h, which seems to be too short.New experimental data of [1992Sch] for the isothermal section at 600°C agree reasonably with the calcula-tion by [1993Lee]. [1997Sop] confirmed calculated diagram at 800°C, but indicated that the calculated γphase field is slightly smaller than the one experimentally detected at 1100°C. [2000Pir] presented the iso-thermal section at 1100°C as a scheme with no experimental points, but the declared tendency of thedecreasing Cr content in the γ phase with increasing Mn content confirms the calculation results of[1993Lee]. The isothermal sections are presented in Figs. 3 - 6.


The temperature - composition sections at 6, 16 and 28 at.% Mn calculated using the dataset of [1993Lee]agree well with the experimental data by [1989Oka]. They are presented in Figs. 7 - 9.

Thermodynamics

The only thermodynamic data available for the system are those from [1978Muk], gained by the specialvapor pressure method (closed chamber method) for the liquid phase at 1570°C. The value of (dlnγMn /dxMn)= 0.9 for xCr < 0.034.The first attempt to calculate phase equilibria in this system on the thermodynamic basis (Calphad method)was made by [1974Kir] who was able to reproduce the α-γ equilibrium data between 750°C and 950°Cbased on their analysis of the previous evaluations of the Fe-Mn and Cr-Fe systems [1973Kir]. They intro-duced the regular solution parameters L for α and γ phases of the Cr-Mn system and parameters L for the αand γ phases of the Cr-Fe-Mn system, but they did not attempt to evaluate the binary parameters for the Cr-Mn system from any information on this binary system. They used all four parameters to describe theirexperimental information which was confined to the Fe rich corner of the ternary system. For reconcilingthe calculation results with experiment they had to introduce the temperature dependence of the obtainedparameters which is so strong that seems to be non-physical.[1977Hil] removed the strong temperature dependence of the above mentioned parameters and arbitrarilyassumed both ternary L parameters to be equal to each other. Binary L parameters for the α and γ phaseswere assumed to be temperature independent and the calculation results reproduced experimental phaseequilibrium data in the Fe rich corner reasonably.The first calculation of the whole phase diagram by the CALPHAD method has been carried outby [1993Lee]. The results were presented as isothermal sections at 1000, 950, 900, 850, 800, 750 and650°;C and vertical section at 6, 16 and 28 mass% Mn. [1993Lee] has treated α´ phase as stoichiometricphase due to the lack of information. The thermodynamic parameters L for the Cr-Fe system and L forthe Cr-Fe-Mn system in (αMn), (βMn) and L for Cr-Fe-Mn in hcp phases were given the value zero.

Cr–Fe–Mn 3



For the liquidus description, the experimental phase equilibria data of [1986Kun] and thermodynamic dataof [1978Muk] were used. The liquidus projection shows no invariant reaction. Similar results could beobtained also by using Calphad type parameters given by [1991Har].


Cr-Fe-Mn steels are considered as materials for structural components for fusion reactors because of theirlow induced radioactivity compared with other stainless steels. The development of non-magnetic steel withhigh stability of austenitic phase and strong resistance to irradiation was described in [1998Tak].Cr-Fe-Mn steels are considered to be suitable materials for cryogenic work because of their high plasticityand strength at low temperatures. Phase diagrams are plotted for the Cr-Fe-Mn alloys for the content of Crup to 14 at.% and Mn up to 30 at.% after cooling and deformation at 20 K and the diagrams of their mechan-ical properties are constructed in [2004Gri].Cr-Fe-Mn steels have high damping capacity, as was found in [2000Miy], where the effect of microstructureon the damping capacity has been clarified. [2003Iga] has investigated martensitic transformation throughγ-α’ and γ-ε-α’ reactions and obtained maximum damping capacity for the ε phase.The Cr-Fe-Mn system was studied also for its magnetic properties. [2000Som] has found that Mn enhancesantiferromagnetism of Cr and has determined temperature and concentration dependence of magnetoresis-tance in Cr0.9–xFexMn1–x and Cr0.55FexMn0.45–x alloys. In the former works [1938Bur] and [1992Sch], themagnetic properties determination was used for phase characterization.Mechanical properties of the Cr-Fe-Mn alloys were studied in [1987Now], where M7C3 carbide (M=Cr, Fe,Mn) in low cost Cr-Fe-Mn-C alloys was investigated with regard to the thermal stability, thermal expansionand its orientation in the γ matrix. Highly anisotropic behavior of the M7C3 carbide was found. [1975Shv2]has studied structure of Cr-Fe-Mn alloys for Cr 4 and 8 mass% and Mn in the range 0 to 40 mass% (C about0.05 mass%) and found maximum hardness at 10 mass% Mn. It was shown that hardness of ε martensite islower than that of α’ martensite. In the former works [1938Bur], [1939Sch], [1949Gri], [1968Tav], and[1990Mur], the mechanical testing was used for the phase characterization.Physical property investigations are listed in Table 4.

Miscellaneous

[1995Rag] successfully used datasets of [1993Lee] to calculate phase equilibria in the Cr-Fe-Mn-Ni system.Knowledge on the solubility of N in Cr-Fe-Mn alloys is important for various applications, where combina-tion of high strength, ductility and corrosion resistance is needed. It was investigated in [1994Fri], wherethermodynamic description of the Cr-Fe-Mn-N system was presented, in [1994Kun], where equilibriumsolubility of nitrogen in Cr-Fe, Fe-Mn and Cr-Fe-Mn alloys in 1000-1200°C range was measured experi-mentally, in [2003Tsu], where adsorption process of N in a solid solution of the Cr-Fe-Mn system was ana-lyzed to fabricate ultra-high nitrogen austenitic steels (>1 mass% N) and in [2005Sal], where mechanicalproperties (strength, ductility, corrosion resistance) were measured.Influence of some additive elements on the properties of Cr-Fe-Mn system is presented in [1991Tak] (effectof Al) and in [2000Pir] (effect of Si). Mössbauer spectra and X-ray diffraction were measured in [2000Sat].In spite of the many published experimental data on the Cr-Fe-Mn system, it may be recommended to deter-mine solubility of the components in the α´ phase and any thermodynamic data of the Cr-Fe-Mn system.Any information on the phase relations in the Cr and Mn rich corners and on the σ phase structure isdesired.

4 Cr–Fe–Mn


Table 1. Investigations of the Cr-Fe-Mn Phase Relations, Structures and Thermodynamics

Reference Method/Experimental Technique Temperature/Composition/Phase Range Studied

[1934Koe] Microscopy, dilatometry 800-1200°C/ 0 to 30 mass% Cr, 0 to 40 mass%Mn / γ-α equilibrium

[1938Bur] Microscopy, X-ray diffraction 600-1000°C/ 0 to 59 mass% Cr, 0 to 49 mass% Mn,0.1 mass% C, 0.3 mass% Si/ γ-α, γ-σ, α-σequilibrium

[1939Sch] Microscopy, phase analysis 700-800°C/ 0 to 60 mass% Cr, 0 to 64 mass% Mn,1 mass% Si/ γ-α, α-σ, γ-σ equilibrium

[1949Gri] Microscopy, TA, dilatometry 600°C, 1200°C/ Vertical sections 6 to 28 mass%Mn / γ-α, α-σ, γ-σ equilibrium

[1968Tav] Microscopy 700°C,1100°C/ 8 to 19 mass% Cr, 10 to 20 mass%Mn/ γ-α, α-σ, γ-σ equilibrium

[1971Kir] Isothermal annealing, electronmicroprobe

750°C/ 0 to 25 at.% Cr, 0 to 8 at.% Mn/ γ-αequilibrium

[1973Kra] Isothermal annealing, electronmicroprobe

1100°C/ 12 to 17 at.% Cr, 7 to 26 at.% Mn/ γ-αequilibrium

[1974Kir] Isothermal annealing, electronmicroprobe, Calphad typecalculations

750-950°C/ 5 to 19 at.% Cr, 0 to 7 at.% Mn/ γ-αequilibrium

[1975Shv1] Microscopy, electron microscopy,X-ray diffraction

650°C, 750°C, 1100°C/ 10 to 20 mass% Cr, 18 to48 mass% Mn/ γ-σ equilibrium

[1975Shv2] Microscopy, X-ray diffraction,dilatometry

650-1100°C/ 4 and 8 mass% Cr, 0 to 15 mass% Mn/γ-α equilibrium

[1977Hil] Calphad type calculations γ-α equilibrium

[1978Muk] Closed chamber method (vaporpressure)

1570°C/ 0 to 0.034 at.% Mn/ liquid thermodynamics

[1986Kun] DTA, directional solidification 1477-1537°C/ 0 to 25 mass% Cr, 0 to 12 mass% Mn/liquid-α, liquid-γ equilibrium

[1988Abe] Microscopy 10 at.% Cr, 30 at.%Mn/ σ phase

[1989Oka] Microscopy, X-ray diffraction 500-700°C/ 8 to 12 at.% Cr, 5 to 30 at.% Mn/γ-σequilibrium

[1990Mur] Isothermal annealing, electronmicroprobe, X-ray analysis

650°C / 0 to 20 mass% Cr, 5 to 63 mass% Mn/ γ - α,γ-αMn, α - σ equilibrium

[1991Har] Calphad type calculations 600-1800°C /Complete Cr-Fe-Mn system

[1992Sch] Isothermal annealing/DTA, X-rayanalysis

600°C/ 0 to 40 mass% Mn, 0 to 18 mass% Cr/ γ-αequilibrium

[1993Lee] Calphad type calculations 600-1800°C / Complete Cr-Fe-Mn system

(continued)

Cr–Fe–Mn 5




[1997Sop] Isothermal annealing/electronmicroprobe

800-1100°C/11 to 20 mass% Cr, 5 to 16 mass% Mn/γ-α, α-σ, γ-σ equilibrium

[2000Pir] Metallography, X-ray analysis 1100°C/0 to 30 at.% Cr, 0 to 50 at.% Mn/γ-α, α-σ,γ-σ equilibrium





Comments/References

α, (Cr,αδFe,δMn)

(Cr)≤ 1863(αFe)(r)≤ 912(δFe)(h2)1538 - 1394(δMn)1246 - 1138

cI2Im�3mW

a = 288.48

a = 286.65

a = 293.15

a = 308.0

continuous solid solution.maximum solubility at 1000°Cis ~21Cr-49Fe-30Mnpure Cr at 25°C, [Mas2]

pure Fe at 25°C, [Mas2]


pure Mn [Mas2]

γ, (γFe,γMn)

(γFe)1394 - 846(γMn)1138 - 1100

cF4Fm�3mCu a = 364.67

a = 386.0

continuous solid solution


pure Mn [Mas2]

(βMn)1100 - 727

cP20P4132βMn

a = 631.52 pure Mn [Mas2]

(αMn)< 727

cI58I�43mαMn

a = 891.26 pure Mn at 25°C [Mas2]

σ (h3)1)

Cr1–x–yMnxFey1323.02 - 991.291312 - 999 [Mas2]

tP30P42/mnmσCrFe

- [1993Lee]

at 1200°C, 0.75 < x < 0.82 at y = 00 < y < 0.4 at y = 0.4 x = 0.35

σ´ (h2)1)

Cr1–xMnx1006 - 800 [Mas2]

tP30P42/mnmσCrFe

- [Mas2]0.75 < x < 0.79

(continued)

6 Cr–Fe–Mn





Comments/References

σ´´ (h1)1)

Cr1–x–yMnxFey

σ´´, Cr1–xMnx< 994.99 [1993Lee]< 800 [Mas2]σ´´, Cr1–yFey830 - ~440

tP30P42/mnmσCrFe

a = 886.2c = 459.5a = 880.0c = 454.4

continuous solid solution below~800°Cat 1200°C Fe solubility is ~40 at.%

[1993Lee] 0.71 < x < 0.81at x = 0.77 [1986Ven]x = 0, 0.5 < y < 0.555 [Mas2]at y = 0.535 [1954Ber]

α´ (h2)Cr1–xMnx924.66 - 600

cI58I�43mαMn

- [1993Lee]x = 0.625

α´´ (h1)Cr1–xMnx< 600

cI58I�43mαMn

- [Mas2]0.60 < x < 0.65

ε hP2P63/mmcMg

a = 253.0c = 407.9

metastable[Mas2, V-C2]

1)Three modifications σ, σ´ and σ´´ exhibit different states of order



Cr Fe Mn

(δMn) ⇌ (γFe) + σ + (βMn) < 1019 E1 σ(γFe)(βMn)(δMn)

1876

12

13201817

69737671

Table 4. Investigations of the Cr-Fe-Mn Materials Properties


[1938Bur] Mechanical testing, magnetometry Hardness, magnetism

[1939Sch] Mechanical testing Hardness

[1949Gri] Mechanical testing Hardness

[1968Tav] Mechanical testing, resistometry Hardness, electrical resistivity

[1975Shv2] Mechanical testing Hardness

[1990Mur] Mechanical testing Hardness

[1992Sch] Resistometry, magnetometry Electrical resistivity, magnetism

(continued)

Cr–Fe–Mn 7




[1998Tak] Radioactivity measurement Induced radioactivity

[2000Miy] Mechanical testing Damping capacity

[2000Som] Magnetometry Magnetism

[2004Gri] Mechanical testing Mechanical properties at low temperatures

Fig. 1. Cr-Fe-Mn. Phase diagram of the Cr-Mn binary system

8 Cr–Fe–Mn


Fig. 2. Cr-Fe-Mn. Liquidus surface projection. Dashed lines are Calphad predictions

Cr–Fe–Mn 9



Fig. 3. Cr-Fe-Mn. Isothermal section at 1200°C. Dashed lines are Calphad predictions

10 Cr–Fe–Mn



Cr–Fe–Mn 11




12 Cr–Fe–Mn



Cr–Fe–Mn 13



Fig. 7. Cr-Fe-Mn. Vertical section at 28 at.% Mn. Dashed lines are Calphad predictions

14 Cr–Fe–Mn



Cr–Fe–Mn 15




16 Cr–Fe–Mn


References[1934Koe] Köster, W., “The Iron Corner of the Fe-Mn-Cr System” (in German), Arch. Eisenhuettenwes.,

7(12), 687–688 (1934) (Phase Diagrams, Phase Relations, Experimental, 3)[1938Bur] Burgess, C.O., Forgeng, W.D., “Constitution of Iron-Chromium-Manganese Alloys”, Am.

Inst. Min. Metall. Eng. Techn., Publ. Nr. 911, 5(3), 1–22 (1938) (Phase Diagram, Experimen-tal, 14)

[1939Sch] Schafmeister, P., Ergang, R., “Brittle Sigma Phase in Three Component Fe-Cr-Mn System”(in German), Arch. Eisenhuettenwes., 12(10), 507–510 (1939) (Phase Diagrams, Experimen-tal, 6)

[1949Gri] Grigorev, A.T., Gruzdeva, N.M., “Investigation of the Irons Alloys with Manganese andChromium” (in Russian), Izv. Sekt. Fiz.-Khim. Anal., 18, 92–116 (1949) (Phase Diagram,Experimental, Magn. Prop., Mechan. Prop., Morphology, Phase Relations, 29)

[1954Ber] Bergman, G., Shoemaker, D.P., “The Determination of the Crystal Structure of the σ Phase inthe Iron-Chromium and Iron-Molybdenum Systems”, Acta Cryst., 7, 857–865 (1954) (Crys.Structure, Experimental, 46)

[1957Hel] Hellawell, A., Hume-Rothery, W., “The Constitution of Alloys of Iron and Manganese withTransition Elements of the First Long Period”, Phil. Trans. R. Soc., 9A, 417–459 (1957)(Phase Diagram, Experimental, 27)

[1958Pot] Potucek, B., “The Constitution of Fe-Cr-Mn Alloys” (in Czech), Hutn. Listy, 13, 1070–1076(1958) (Phase Diagrams, Review, 44)

[1968Tav] Tavadze, F.N., Pirtskhalaishvili, V.A., Nabichvrishvili, M.A., “A Study of the Structure ofAlloys of the Fe Corner of the Fe-Cr-Mn System” (in Russian), Soobsh. Akad. Nauk Gruz.SSR, 49(3), 641–646 (1968) (Phase Diagram, Experimental, Phase Relations, 7)

[1971Kir] Kirchner, G., Larbo, G., Uhrenius, B., “An Investigation into the Distribution of Cr and MnBetween Ferrite and Austenite Using Experimental Measurements and Thermodynamic Cal-culations”, (in German), Prakt. Metallogr., 8, 641–654 (1971) (Phase Relations, Experimen-tal, 4)

[1973Kra] Kralik, F., Kovacova, K., “Determination of the α+γ /γ Phase Region Boundary in Fe-Cr-Ni,Fe-Cr-Co and Fe-Cr-Mn System” (in Czech), Kovove Mat., 11, 6–15 (1973) (Phase Diagram,Phase Relations, Calculation, Experimental, 7)

[1973Kir] Kirchner, G., Nishizawa, T., Uhrenius, B., “Distribution of Chromium Between Ferrite andAustenite and Thermodynamics of α-γ Equilibrium in Fe-Cr and Fe-Mn Systems”, Met.Trans., 4(1), 167–174 (1973) as quoted in [1977Hil]

[1974Kir] Kirchner, G., Uhrenius, B., “Experimental Study of the Ferrite / Austenite Equilibrium in theFe-Cr-Mn System and the Optimization of Thermodynamic Parameters by Means of a Gen-eral Mathematical Method”, Acta Metall., 22, 523–532 (1974) (Phase Diagram, Experimen-tal, #, *, 24)

[1975Shv1] Shvedov, L.I., Pavlenko, Z.D., “σ-Phase in Fe-Cr-Mn Alloys” (in Russian), Izv. Akad. NaukBeloruss SSR, Fiz-Tekhn., (2), 14–17 (1975) (Crys. Structure, Experimental, Mechan. Prop.,Morphology, 6)

[1975Shv2] Shvedov, L.I., Pavlenko, Z.D., “Structure and Phase Composition of Fe-Cr-Mn Alloys” (inRussian), Izv. Akad. Nauk Beloruss. SSR, (Fiz -Tekhn.), (2), 22 (1975) (Phase Diagram,Experimental, Magn. Prop., Mechan. Prop., Morphology, Phase Relations, 1)

[1977Hil] Hillert, M., Waldenstroem, M., “Gibbs Energy of Solid Solutions of C, Cr, Mn, Mo, and Ni inFe”, Scand. J. Met., 6, 211–218 (1977) (Calculation, Phase Diagram, Phase Relations, Ther-modyn., 33)

[1978Muk] Mukai, K., Uchida, A.,Tagani, T., Wasai, Y., Proc. 3rd Int. Iron and Steel Congress, ASM,Metals Park, OH, 266–276 (1978) (Thermodyn., Experimental, 17)

[1980Geo] Georgieva, I.Ya., Sorokina, N.A., Galtsova, V.I., Fiz. Met. Metalloved., 49, 206–209 (1980)as quoted by [1993Lee]

[1982Lem] Lemkey, F.D., Thompson, E.R., Schuster, J.C., Nowotny, H., Mater. Res. Soc. Symp. Proc.,31–50 (1982) (Phase Diagram, Experimental) as quoted by [1993Lee]

Cr–Fe–Mn 17



[1984Fri] Fritscher, K., Hammelrath, H., “New Ternary Phase in Cr-Fe-Mn Alloys”, Naturwissenschaf-ten, 71(11), 583 (1984) as quoted by [1993Lee]

[1985Riv] Rivlin, V.G., “Assessment of Phase Equilibria in Ternary Alloys of Iron”, J. Less-CommonMet., 114(1), 111–121 (1985) (Phase Diagram, Phase Relations, Assessment, 4)

[1986Kun] Kundrat, D.M., “Phase Realtionships in the Fe-Cr-Mn-Ni-C System at Solidification Tem-peratures”, Metall. Trans. A, 17A, 1825–1835 (1986) (Phase Diagram, Phase Relations,Experimental, Calculation, #, *, 22)

[1986Ven] Venkatraman, M., Neumann, J.P., “The Cr-Mn System”, Bull. Alloy Phase Diagrams, 7(5),457–462 (1986) (Crys. Structure, Phase Diagram, Review, 32)

[1987Now] Nowotny, H., Lemkey, F.D., Wayne, S.F., Pearson, D.D., Gupta, H., “ThermomechanicalProperties of Low-Cost Aligned Fe-Cr-Mn-C Alloys”, High Temp.-High Pressures, 19(5),501–508 (1987) (Experimental, Thermodyn., 12)

[1988Ray] Raynor, G.V., Rivlin, V.G., “Cr-Fe-Mn” in “Phase Equilibria in Iron Ternary Alloys”, TheInst. of Metals, London, 288–299 (1988) (Phase Diagram, Phase Relations, Rewiev, #, *, 11)

[1988Abe] Abe, F., Araki, H., Noda, T., “Discontinuous Precipitation of σ Phase During Recrystalliza-tion in Cold-rolled Fe-10Cr-30Mn Austenite”, Mater. Sci. Technol., 4(10), 885–893 (1988)(Phase Relations, Experimental, 38) as quoted by [1993Lee]

[1989Oka] Okazaki, Y., Miyaharaa, K., Hosoi, Y., Tanino, M., Komatsu, H., “Effect of Alloying Ele-ments of σ Phase Formation in Fe-Cr-Mn Alloys” (in Japanese), J. Jpn. Inst.Met., 53(5),512–521 (1989) (Phase Diagram, Experimental, Crys. Structure, 15) as quoted by [1993Lee]

[1990Mur] Murata, Y., Koyama, K., Matsumoto, Y., Morinaga, M., Yukawa, N., “Equilibrium PhaseDiagram of Fe-Cr-Mn Ternary System”, Trans. Iron Steel Inst. Jpn., 30(11), 927–936(1990) (Crys. Structure, Experimental, Phase Diagram, #, *, 19)

[1990Yuk] Yukawa, N., Morinaga, M., Nishiyama, K., Matsumoto, Y., Murata, Y., Ezaki, H., ASTM STP1047, Klueh, R.E., Gelles, D.S., Okada, M., Packen, N.H., (Eds.), ASTM, Philadelphia, PA,1990, pp.30–46, (Phase Diagram, Experimental) as quoted by [1993Lee]

[1991Har] Hari Kumar, K.C., PhD Thesis, Indian Institute of Technology, Delhi, India, 1991, (PhaseDiagram, Calculation) as quoted by [1995Rag]

[1991Tak] Takahashi, H., Ohnuki, S., Kinoshita, H., Nakahigashi, S., “Effect of Alloying Elements onPhase Stability in Neutron-Irradiated Fe-Cr-Mn Model Alloys”, J. Nucl. Mater., 629–632,179–181 (1991) (Experimental, Phase Diagram, 13)

[1992Sch] Schule, W., Lang, E., “A Contribution to the Phase Diagram of Iron-Manganese-ChromiumAlloys”, ASTM Spec. Techn. Publ. (Eff. Radiat. Mater.), (Stp. 1125), 945–957 (1992) (PhaseDiagram, Phase Relations, Experimental, Electr. Prop., Magn. Prop., 28)

[1993Lee] Lee, B.-J., “A Thermodynamic Evaluation of the Cr-Mn and Fe-Cr-Mn Systems”, Metall.Trans. A, 24A, 1919–1933 (1993) (Phase Diagram, Calculation, #, *, 63)

[1994Kun] Kunze, J., Rothe, I., “Solubility of Nitrogen in Austenitic FeCrMn Alloys”, Steel Res., 65(8),331–337 (1994) (Calculation, Experimental, Thermodyn., 25)

[1994Fri] Frisk, K., Caian, Q., “AThermodynamic Evaluation of the Solubility of N in Solid and LiquidCr-Fe-Mn Alloys”, Z. Metallkd., 85(1), 60–69 (1994) (Phase Diagram, Phase Relations, Cal-culation, Thermodyn., 26)

[1994Rag] Raghavan, V., “Cr-Fe-Mn (Chromium-Iron-Manganese)”, J. Phase Equilib., 15(5), 530–531(1994) (Phase Diagram, Review, 15)

[1995Rag] Raghavan, V., “Effect of Manganese on the Stability of Austenite in Fe-Cr-Ni Alloys”,Metall. Mater. Trans. A, 26A(2), 237–242 (1995) (Experimental, Phase Relations, Thermo-dyn., 18)

[1997Sop] Sopousek, J., Vrestal, J., Kunze, J., “Experimental Study of Phase Equilibria in the Fe-Cr-MnSystem in the Temperature Range 1073 to 1373 K”, Z. Metallkd., 88(3), 246–249 (1997)(Phase Relations, Experimental, 17)

[1998Tak] Takahashi, H., Shindo, Y., Kinoshita, H., Shibayama, T., Ishiyama, S., Fukaya, K., Eto, M.,Kusuhashi, M., Hatakeyama T., Sato, I., “Mechanical Properties and Damage Behavior of

18 Cr–Fe–Mn


Non-Magnetic High Manganese Austenitic Steels”, J. Nucl. Mater., 263(Part B), 1644–1650(1998) (Mechan. Prop., Experimental, 14) cited from abstract

[2000Miy] Miyahara, K., Wu, K., Okada, K., Kang, Ch.-Y., Sasaki M., Igata, N, “Microstructural Effectson the Damping Capacity of Fe-Cr-Mn Alloys”, Proc. 2nd Int. Conf. Proces. Mater. for Prop-erties, TMS - Miner. Metals & Materials Soc., Warrendale, PA, USA, 195–198 (2000)(Mechan. Prop., Experimental, 4) cited from abstract

[2000Sat] Satula, D., Szymanski, K., Waliszewski, J., Dobrzynski, L., Prus, B., “Moessbauer and X-rayDiffraction Study of Cr-Fe-Mn Alloys”, Molecular Physics Reports, Osrodek WydawnictwNaukowych, 30, 151–158 (2000) cited from abstract

[2000Pir] Pirtskhalayshvily, B.A., “Restoring Austenite-Forming Properties of Mn in Fe-Cr-Mn andFe-Cr-Mn-Ni Alloys by Dopping a Silicium” (in Russian), Metally, (2), 65–67 (2000)(Assessment, Mechan. Prop., Phase Diagram, 2)

[2000Som] Somsen, Ch., Acet, M., Nepecks, G., Wassermann, E.F., “The Effect of Magnetic Orderingon the Giant Magnetoresistance of Cr-Fe-V and Cr-Fe-Mn”, J. Magn. Magn. Mater., 208,191–206 (2000) (Experimental, Magn. Prop., Phase Relations, 26)

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[2003Rag] Raghavan, V., “Cr-Fe-Mn (Chromium-Iron-Manganese)”, J. Phase Equilib., 24(3), 259–260(2003) (Assessment, Crys. Structure, Phase Diagram, Phase Relations, 7)

[2003Tsu] Tsuchiyama, T., Takaki, S., “Thermodynamics of Nitrogen Absorption into Solid Solution inFe-Cr-Mn Ternary Alloys”, Mater. Sci. Forum, 426–432, 957–962 (2003) (Experimental,Mechan. Prop., Phase Diagram, Phase Relations, Thermodyn., 8)

[2004Gri] Grikurov, G., Antropov, N., Baratashvili, I., Skibina, L., Chernik, M, Yushchenko, K., “PhaseStability of the Fe-Cr-Mn System and the Problem of Development of Stainless Steels on itsBasis”, AIP Conf. Proc., no.711, USA, pt.1, (2004) 93–97 (Phase Diagrams, Mechan. Prop.,Experimental, 5) cited from abstract

[2004Wit] Witusiewicz, V.T., Sommer, F., Mittemeijer, E.J., “Reevaluation of the Fe-Mn Phase Dia-gram”, J. Phase Equilib. Diff., 25(4), 346–354 (2004) (Experimental, Phase Diagram, PhaseRelations, Calculation, Thermodyn., #, 34)

[2005Sal] Saller, G., Bernauer, J., Leitner, H., Clemens, H., “On Restricting Aspects in the Productionof Nonmagnetic Cr-Mn-N-Alloyed Steels”, Steel Res. Int., 76(11), 769–774 (2005) (Experi-mental, Phase Relations, 18)

[2007Wit] Witusiewicz, V.T., private communication to MSI, (2007)[Mas2] Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals

Park, Ohio (1990)[V-C2] Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic


Cr–Fe–Mn 19



Chromium – Iron – Molybdenum

Ales Kroupa

Introduction

The Cr-Fe-Mo system has been studied extensively. Both Cr and Mo are important elements used in steels,in the past and in the present days. Knowledge of the phase diagram of such a system is therefore of signifi-cant interest, especially now, when the number and amount of alloying elements in steels and other alloys isincreasing, and intermetallic phases are playing a more important role.One of the earliest experimental studies was carried out by [1931Wev], who examined the α/γ region in thisternary system. This information was reviewed by [1949Jae], who published the phase diagram showing theα/γ loop in the Cr-Fe-Mo system. The first information about the ternary phase, τ1, was published by[1949And, 1951And], and the structure was correctly identified. [1950Put] studied the σ phase region inthe ternary system. The first study of the complete ternary system was carried out by [1951Put]. They stu-died about 30 ternary alloys by thermal analysis in order to determine the melting temperatures in this sys-tem and established partial isothermal sections of the phase diagram. A second experimental study of theliquidus surface of the system was published by [1957Tak], who employed a similar experimental techniqueto that used by [1951Put]. Further experimental studies were published by [1950Duw, 1952Bue, 1953Koh,1954Kas, 1954McM, 1957Age, 1957Bec, 1957Gol, 1962Gri, 1964Alf, 1966Kim, 1971Yam, 1976Kie,1981Zha]. They studied mainly phase relations and properties of solid phases in the Cr-Fe-Mo system.The only study of thermodynamic properties of this system, namely the specific heat at low temperatures,was published by [1971Bau].Nevertheless, great inconsistencies were found in the results of various works published before 1980. Thefirst review of this system was presented by [1971Wes], who pointed out many discrepancies in the existingdata, which were discussed in extensive reviews by [1984Ray, 1988Ray]. One problem was that differentversions of the Fe–Mo binary phase diagram were available at different times, resulting in misinterpretationof intermediate phases found in the ternary system. Also, the identification of the ternary phases existing inthis system was complicated because of uncertainties concerning the number of such phases, their stabilityregions and crystal structure.The lack of information available in the reviews [1984Ray, 1988Ray] prompted more recent experimentaland theoretical studies to be carried out by several research groups. A comprehensive work was publishedby [1988And], who experimentally and theoretically studied the complete Cr-Fe-Mo system in the tempera-ture range 950–1200°C. Further experimental works were published by [1986Liu, 1988Liu, 1989Liu,1990Liu], who studied partial or complete isothermal sections at 850, 1050, 1100, 1180, 1200 and1250°C. A short review of these studies has been given by [1994Rag]. The lattice parameter of the τ1 phasewas reported by [1990Ere]. Since then, no other systematic experimental study of this system has beenpublished.Several authors have also modeled the phase equilibria in the Cr-Fe-Mo system theoretically. [1975Kau]applied the CALPHAD approach to this system. All intermetallic phases were described as stoichiometric.The most detailed theoretical modeling was carried out by [1988And], later augmented by [1992Qiu]. Theyalso used the CALPHAD method.Currently, more attention is being paid to properties of (Cr,Mo) steels and other commercial multicompo-nent alloys containing Cr and Mo, however, these works do not contain new information about the phasediagram of the Cr-Fe-Mo system. Some studies concerning intermetallic phases existing in the ternary sys-tem will be discussed briefly in the Miscellaneous section.The basic studies of phase equilibria and crystal structure of the phases are listed in Table 1.

Binary Systems

The binary Cr-Fe, Cr-Mo and Fe-Mo phase diagrams are accepted from [Mas2].

Cr–Fe–Mo 1



Solid Phases

Two solid solutions, α (bcc (A2) and γ (fcc (A1), and four intermediate binary phases σ, μ (Fe7Mo6), λ(Laves phase, Fe2Mo) and the high temperature R phase, all with significant solubility of the third element,exist in this system. One ternary phase, τ1, was found. Miscibility gaps are found in the α phase in both theCr-Fe and Cr-Mo systems. The crystallographic information available is shown in Table 2.The solubility of elements in the Cr rich α phase was studied by [1952Bue].The exact identification of the intermetallic phases in the system took a relatively long time and was com-plicated. An incorrect version of the Fe-Mo binary diagram was used in the analysis of most of the earliestexperimental studies [1951Bae, 1951Put, 1952Gol, 1953Koh, 1954McM, 1957Bec, 1957Tak, 1962Gri,1964Alf]. This binary phase diagram did not take into account the existence of the R phase and two inter-mediate phases, λ and μ, were considered as one phase, designated ε (Fe3Mo2). Its crystallographic struc-ture was given as trigonal [1951Put] or hexagonal [1957Bec]. The first study covering the whole phasediagram was published by [1951Put] and as the incorrect version of the Fe-Mo phase diagram was used,erroneous results were obtained. No ternary phases were found. Therefore, these results were not incorpo-rated into later assessments [1971Wes, 1984Ray, 1988Ray]. A complete and correct identification of allintermetallic phases existing in the system was given there.The ternary phase τ1 was first reported by [1949And], who correctly established its crystal structure as cor-responding to the αMn type and also measured its lattice parameter (Table 2). They also determined its che-mical composition and concluded that it is very close to the composition of the σ phase. This phase is nowusually denoted as χ, but it will be designated as τ1 here. The crystal structure, composition and lattice para-meters of the ternary phase were also studied by [1953Koh, 1954Kas, 1954McM, 1957Bec, 1964Alf,1966Kim, 1976Kie, 1990Ere]. The most detailed study of the τ1 phase was carried out by [1954Kas] and[1966Kim] and its lattice parameter and site occupancy were established. [1954Kas] also proposed its com-position to be Cr12Fe36Mo10, but the phase is considerably nonstoichiometric. The temperature range of thisphase is still uncertain. The lower temperature limit was determined experimentally by [1976Kie], whofound this phase to be stable up to approximately 550°C for compositions 20-28 mass% Cr and 2-5 mass%Mo. The information relating to the upper temperature limit is ambiguous. Some early studies found theternary phase to be stable up to approximately 1450°C [1957Tak]. They studied the liquidus surface of thissystem and found a τ1 primary solidification region below 1455°C. [1957Gol] and [1964Alf] also found thephase to be stable at temperatures around 1450°C. Nevertheless, they did not find the primary solidificationregion of the ternary phase.The latest systematic studies of the Cr-Fe-Mo diagram established the upper temperature limit of τ1 to bebetween 1180 and 1200°C [1986Liu, 1988Liu, 1989Liu] and also, there is no primary solidification regionfor this phase. They also confirmed the existence of the τ1 phase to be down to at least 850°C. This wasconfirmed by [1988And], who did not find the τ1 phase at 1200°C. [1986Liu, 1988Liu, 1989Liu,1988And] did not study this system at lower temperatures.All of the old studies [1957Gol, 1957Tak, 1964Alf] used an incorrect version of the Fe-Mo binary phasediagram, and it is quite possible that the results were misinterpreted. For instance, [1964Alf] identified onlythe τ1 and σ intermetallic phases and the α solid solution in their work, despite studying the concentrationand temperature regions where other intermetallic phases (especially λ and R) should also appear. Also, thetemperature stability of the σ phase does not agree with more recent results. [1957Gol] identified anotherternary phase, denoted N in their paper, which was not found elsewhere.For these reasons, the newer experimental measurements were accepted in the present evaluation and thetemperature range of the τ1 phase is given as approximately 550-1200°C.The σ phase was found to exist in the Cr-Fe-Mo system, creating a complete region of miscibility betweenthe terminal high temperature σFe,Mo phase and the low temperature σCr,Fe [1952Bue]. The crystal structureand lattice parameters were measured by [1950Duw], who defined the tentative tetragonal structure for theσ phase and measured the composition dependence of lattice parameters on Mo content at 650°C. Unfortu-nately, their measurements do not agree with later results that were obtained after the correct crystal struc-ture had been found, e.g. [1976Kie]. The phase equilibria between the terminal solid solutions and theσ phase were studied by [1957Age, 1964Alf], but their results are influenced by the above mentioned dis-crepancies in phase identification. Consistent results concerning the temperature and concentration range

2 Cr–Fe–Mo


of the σ phase were obtained by [1988And, 1986Liu, 1988Liu, 1989Liu] and they are accepted in thisreview. The σ phase is stable across the whole experimentally studied region and the nonstoichiometryincreases significantly with the addition of the third element in comparison with the Cr-Fe and the Fe-Mobinary systems.The presence of the λ phase in the system was first reported by [1957Bec], who correctly defined the struc-ture of the phase and determined its upper temperature limit to be 954°C; slightly higher than the currentlyaccepted value for it in the Fe-Mo binary system. Nevertheless, the results were generally not taken intoaccount in later reviews, e.g. [1961Eng, 1964Alf]. The present Fe-Mo binary phase diagram first appearedin the review of [1971Wes]. It is generally accepted that the λ phase exists in this system, but nevertheless,no systematic study of the low temperature regions where it exists has been carried out. The behavior of thisphase in the ternary system and the probable influence of Cr on its temperature and concentration range wasestimated by [1984Ray, 1988Ray], who corrected the results of previous studies [1954McM, 1957Tak].A similar situation exists also for the μ and R phases, which were misinterpreted in the earlier works. The μphase was mistaken for the ε phase in the past. Its composition was later corrected to Fe7Mo6. Nevertheless,the independent high temperature R phase was still not taken into account. The existence of a possible tern-ary phase, denoted ρ, was reported in [1957Bec]. This phase could not be identified from X-ray diffractionpatterns. Similarities between its structure and that of the binary R phase were found in the patterns, andtherefore, there is a strong reason to identify this phase as the Cr stabilized binary Fe-Mo R phase. The sta-bility range of both phases was studied in detail in the latest experimental work [1986Liu, 1988And,1988Liu, 1989Liu]. These results are reasonably consistent as to the temperature stability of the intermetal-lic phases, especially for the μ phase. There is a slight uncertainty concerning the lower temperature limit ofthe R phase. [1986Liu, 1988Liu, 1989Liu] reported the phase to be stable from the solidus to 930°C. Theresults of [1988And] agree with the upper limit (they identified this phase at 1200°C), but they did not findthe R phase at 950°C. On the other hand, [1990Liu] reported the precipitation of metastable R phase, even atlower temperatures. This phase transformed to other intermetallic phases during ageing. Therefore, it is pos-sible that the R phase observed at 930°C is also metastable and would also transform after a much longerannealing time, and being very close to the true lower temperature limit, the kinetics are very slow. Thelower temperature limit of R phase remains therefore uncertain.Goldschmidt reported in [1952Gol, 1957Gol] the existence of the ternary N phase (CrFe4Mo2) mentionedabove. However, this phase was not confirmed by any other study, even if the composition of some experi-mental alloys, e.g. [1957Bec] and [1954McM], were in the same region. Particularly, [1957Bec] excludedthe possibility of N phase being consistent with their ρ phase. Therefore, this phase was not taken intoaccount in earlier reviews [1984Ray, 1988Ray] as well as in the present evaluation.

Liquidus Surface

Melting temperatures in the Cr-Fe-Mo system were measured by [1951Put]. The existence of a region oflow freezing points of around 1460°C was found close to the composition 20 mass% Cr, 60 mass% Feand 20 mass% Mo. As mentioned above, the existence of the ternary phase was not taken into account inthis work and the information given does not allow the identification of the primary solidification phases.[1957Tak] also published a study of the liquidus surface, but in this case, the τ1 ternary phase was correctlyidentified in the system. Both studies used high purity materials for alloy preparation, and the melting tem-peratures were measured by thermal analysis. Because of the incorrect interpretation of the binary inter-mediate phases in the Fe-Mo system (confusion involving the λ, μ, and R phases), the liquidus surfaceof [1957Tak] was redrawn by [1984Ray] with respect to the new information. The regions of primary pre-cipitation also included the ternary τ1 phase. However, this was not confirmed by the later experimentalwork of [1986Liu, 1988And, 1988Liu, 1989Liu], who did not find the ternary phase to be stable above1200°C. In particular, [1986Liu] studied the composition region where [1957Tak] observed the ternaryphase but did not find it at 1200°C. As the recent papers are mutually consistent and there are still significantdoubts about the identification of the intermetallic phases in the older papers, the liquidus projection from[1988And] (Fig. 1) was accepted in the present evaluation. This liquidus projection was calculated by[1988And] taking into account their own experimental results and those of [1981Zha]. The authors reached

Cr–Fe–Mo 3



very good agreement between the calculated and the experimental phase diagrams for the temperature range950-1200°C, and therefore, the prediction of the liquidus projection should be reliable.


Details of invariant equilibria presented by [1957Tak] were assessed by [1988Ray]. Nevertheless, the dis-crepancies in the identification of the primary solidification regions remained. Solidification of the μ andτ1 phases was assumed by [1957Tak] and it was later changed to R and τ1 by [1984Ray], as the existenceof the high temperature R phase was confirmed later. The binary Fe-Mo edge of the liquidus surface did notcorrespond to the accepted binary Fe-Mo diagram. [1984Ray] proposed a peritectic reaction L + σ + R⇌ τ1at 1455°C and two other invariant reactions involving the liquid: the transition (U type) reactionL + σ ⇌ α + τ1 at 1385°C and the eutectic reaction L ⇌ α + τ1 + R at 1345°C. Considering the newerexperimental data, these reactions cannot be correct, as the ternary τ1 phase is not stable above 1200°C.There are no new experimental data on the liquidus surface, but it was predicted by [1988And]. Accordingto theoretical modeling, there is only one invariant reaction involving the liquid, namely the U type reactionL + σ ⇌ α + R at 1471°C.Two invariant reactions were also found in the solid phase region by [1957Bec], who published isopleths for70 and 80 mass% Fe, and [1986Liu], who presented the section for 16 mass% Mo. The exact reaction tem-peratures were not established in these studies, and there are discrepancies in the positioning of the reactionsand in the reaction sequence in comparison with the work of [1988And]. Therefore, the theoretical assess-ment based on [1988And] was used in combination with their experimental results to establish the probabletype of invariant reactions. The invariant reactions are given in Table 3. The probable reaction sequence isshown in Fig. 2.

Isothermal Sections

The phase relations between the α and γ solid solutions in the Cr-Fe-Mo system were studied by[1931Wev], who analyzed the influence of Cr and Mo on the shape of closed γ loop in the Fe rich corner.He found the expected temperature decrease of the lower α/α + γ phase boundary with increasing Cr com-pared with the binary Fe-Mo system. The transformation temperature was observed to be at about 870°Cand at 4-5 mass% Cr for the Cr:Mo ratio 5:1. The results of this study are shown in Fig. 3. Another experi-mental study of the Fe rich part of the diagram was carried out by [1967Bun], who studied mainly the influ-ence of Mo (up to 2 mass%) on the C-Cr-Fe system at 1050°C, but some alloys did not contain C.The isothermal sections between 1900-700°C (Figs. 4-13) were constructed and calculated using the experi-mental data from [1931Wev, 1954McM, 1976Kie, 1984Ray, 1986Liu, 1988Liu, 1988And, 1989Liu]. Theisothermal sections for lower temperatures - down to 850°C - are limited by early measurements. Partial iso-thermal sections of the Cr-Fe-Mo system were published by [1984Ray, 1988Ray]. They used the data avail-able at that time [1950Duw, 1951Bae, 1952Bue, 1953Koh, 1954Kas, 1954McM, 1957Bec, 1957Tak,1957Gol, 1981Zha] and some other information relating to commercial Cr-Mo steels and alloys. Theyhad to redraw all the phase diagrams published in the above mentioned papers to adjust them in relationto the new Fe-Mo phase diagram. However, the result was not completely satisfactory, as no systematicstudy of this system was available, which would have solved inconsistencies in the earlier works. Someinformation on the medium Cr region for 700 and 800°C is also available in [1976Kie] (Fig. 13 showsthe 700°C section).Experimental and theoretical isothermal sections of the Cr-Fe-Mo system in the temperature range 950-1200°C were published later by [1988And] in their comprehensive work. Partial and complete experimentalisothermal sections were also published by [1986Liu, 1988Liu, 1989Liu] for 850, 1050, 1100, 1180, 1200and 1250°C. The results of these works are consistent, confirming the temperature regions of existence forthe intermetallic phases τ1, μ and σ. The insignificant discrepancy mentioned above was found only for theR phase. The phase diagrams published by [1988And] and [1986Liu, 1988Liu, 1989Liu] differ slightly inthe extention of the composition regions, where particular intermetallic phases are stable.The isothermal sections selected for publication in this evaluation for temperatures between 900°C and1900°C mainly are based on [1988And], as their results are more consistent with the accepted binary

4 Cr–Fe–Mo


diagrams from [Mas2]. The projections of phase boundaries of the intermetallic phases (especially μ) intothe Fe-Mo binary system in [1986Liu, 1988Liu, 1989Liu] do not quite correspond to the generally acceptedvalues published by [Mas2]. Despite the discrepancies mentioned above, the isothermal sections at 1250and 850°C published in [1989Liu] are shown in Figs. 6 and 12. The agreement with the results of[1988And] is good and the positions of phase boundaries for the R and μ phases on the Fe-Mo axis wereadjusted to be consistent with the accepted binary diagram.The Cr-Fe-Mo system was also theoretically assessed by [1988And] using the CALPHAD method. Theyused their own experimental data and those from [1981Zha]. The γ loop in the Fe rich corner of the isother-mal sections published in this paper was omitted. The experimental results [1931Wev] and [1967Bun] wereused by [1992Qiu] for the theoretical assessment. The experimental and modeled phase diagrams are in verygood agreement. As the calculated isothermal sections at given temperatures were also used in [1988And]as a background for the presentation of their experimental tie lines, the results of the theoretical modelingare shown in Figs. 7-9, 10 and 11. The prediction of the isothermal sections at 1600°C and 1900°C are alsoshown in Figs. 4 and 5. No experimental information is available, but the tendency of Cr to stabilize all bin-ary intermediate phases is quite clear. The calculated isothermal section at 1000°C from [1988And] (Fig. 7)was also published by [1990Hil].


Only very limited information is available for temperature-composition sections in the literature. Two iso-pleths for 70 and 80 mass% Fe were presented by [1957Bec]. Various temperature-composition sections forseveral Mo contents and Cr/Mo ratios were published by [1964Alf] and the isopleths for three different Crcontents (20, 24, 28 mass%) were constructed by [1976Kie]. It was pointed out by [1988Ray] that theresults of [1976Kie] disagree with the accepted Cr-Fe binary system (the σ phase disappears at significantlylower temperatures for all Cr contents) and the results of [1957Bec] are not consistent with the Fe-Mo bin-ary system, as the samples probably were not in thermodynamic equilibrium. The isopleths published by[1957Bec] were redrawn by [1988Ray] to correct the obvious discrepancies with the binary diagram. Never-theless, the results of [1957Bec] do not agree with the newer isothermal sections from [1986Liu, 1988And,1988Liu, 1989Liu] at higher temperatures. No high-temperature σ phase above 1000°C was found by[1957Bec] contrary to experimental studies [1986Liu, 1988And, 1988Liu, 1989Liu] and therefore, differentphase fields were proposed in the isopleths. Also, the upper temperature limit for the low-temperature σphase close to the Cr-Fe binary system found by [1957Bec] at 70 mass% Fe does not correspond to thenewer results [1986Liu, 1988And, 1988Liu, 1989Liu]. On the other hand, they confirm the lower tempera-ture limit of the R phase to be above 950°C.Only the isopleth for 16 mass% Mo published by [1986Liu] in the range 700-1000°C is available in therecent literature. This isopleth corresponds reasonably well with the accepted isothermal sections. Theexception is the lower temperature limit of the R phase (about 930°C) compared with [1988And]. For thatreason, this isopleth indicates a different invariant reaction than that accepted in this evaluation. Therefore,no temperature-composition section is published here.The isopleth at 30 at.% Mo was calculated by [1983Sun].

Thermodynamics

The specific heat of chromium rich Cr-Ni and Cr-Fe-Mo alloys was measured by [1971Bau] in the tempera-ture range 1.3-4.2 K. Measurements were made for compositions of 20 mass% Mo and 0-20 mass% Fe.This experimental program was intended primarily for the determination of the electronic band structureof 3d transition elements. It is known that Cr–Fe alloys exhibit unusual electron specific heat coefficientsand an abnormally low Debye temperature, which is attributed to a complex magnetic structure of thesealloys. The addition of 20 at.% Mo may avoid such complications, as the alloy becomes paramagnetic atliquid He temperatures. The results of the measurements are shown in Table 4. The temperature dependenceof the specific heat is expressed as C = 234 R (T/Θ)3 + γ0T, where R is the gas constant and Θ is the appar-ent Debye temperature. The authors concluded that the values of the specific heat coefficients and the Debyetemperature are rather unusual (a low value of the apparent Debye temperature). They cannot be caused by

Cr–Fe–Mo 5



specific magnetic properties of these alloys (antiferromagnetism), but they are probably caused by the pre-cipitation of fine particles in the alloys. An X-ray diffraction study and metallographic examination revealedfine needle-like precipitates in the matrix. The nature of these precipitates was not determined.


There have been few systematic studies of materials properties of the ternary Cr-Fe-Mo alloys despite theirimportance for high temperature applications. Only a few hardness measurements have been carried out inthe past; by [1957Age, 1957Wes, 1964Alf]. They studied the influence of precipitating phases on the over-all hardness of alloys [1957Age, 1964Alf] and the hardness of the intermetallic phases themselves[1957Wes]. It was found by [1957Age, 1964Alf] that the precipitation of the σ phase generally increasedthe hardness of the material, but a decrease in the hardness was observed in some cases after a longerhigh-temperature annealing. The hardness measurements of the σ and τ1 phases [1957Wes] showed thatboth phases are significantly harder than typical high temperature alloys at room and elevated temperatures.The σ phase is generally harder and softens more slowly with increasing temperature than the ternary τ1phase does.The oxidation resistance of Cr-Fe-Mo alloys was studied by [1964Alf]. It was found that an increase in theMo content in the Cr rich solid solution does not influence it significantly, but the oxidation resistancedecreases abruptly when the σ phase starts to precipitate.The compositions of surface layers of Cr-Fe-Mo alloys were investigated under various conditions by[1979Mat, 1984Goe, 1997Ked]. An examination of the influence of Mo on the corrosion behavior of Cr-Fe-Mo alloys or (Cr,Mo) steels was the driving force for these studies. The influence of Cr and the polar-ization potential in hydrochloric acid on the surface composition was established by [1984Goe]. The resultsobtained showed that active dissolution leads to a Mo enrichment of the surface layer, which consequentlyinhibits further active dissolution.A new monolayer-sensitive technique - optical second harmonic generation - was applied by [1989Ham] tostudy the interfaces at ambient pressure and buried solid/solid and buried liquid/liquid interfaces. It providesremote sensing of surface composition at elevated temperatures. The authors proved that Cr/N segregationaffected the second harmonic generation from the 18Cr-79Fe-3Mo alloy and the intensity is related to theatomic fraction of segregating atoms on the surface.As mentioned in the Introduction, attention is currently given more to the modern Cr, Mo steels and otheradvanced materials containing Cr and Mo. For example, the spinodal decomposition of α phase into Fe andCr rich α phases in modern duplex steels is very important in relation to mechanical properties (thermal age-ing embrittlement near 480°C). It was studied by Monte Carlo modeling by [2000Hon, 2003Got, 2003Iwa,2003Kuw].There is also extensive literature on the properties of modern high alloyed (9-12Cr,1-2Mo) steels, where theknowledge of the ternary Cr-Fe-Mo phase diagram is very important. These steels are widely used at ele-vated temperatures (600-650°C) in power stations and other important industrial applications where longterm structural stability is necessary in order to guarantee safe and reliable service. The study of the mechan-ical properties of these steels confirmed the important role of Fe-Mo based intermetallic phases, namely theλ (Laves) phase, for the long term development of materials properties [1997Str, 2000Cho, 2001Par,2003Hal, 2003Nak, 2005Mal]. Many studies carried out on these steels have provided important informa-tion relating to the behavior of the λ phase, which is not otherwise available for the ternary system, becauselong annealing times were used and the structures are supposed to be close to thermodynamic equilibrium.Even though the composition of the steels is very complex and the equilibria are influenced by other ele-ments, the λ phase dissolves mainly W and Si, other elements are not significant. As many steels do notcontain W, the information about phase relations and composition of the λ phase may be important forthe ternary system.Among other interesting applications of materials based on the Cr-Fe-Mo system, the increasing use of Cr,Mo sintered steels [1999Mol1, 1999Mol2, 2006Cam] may be mentioned, as well as the development of Fe,Al/Cr,Mo steel composites [2004Mas, 2006Mas].

6 Cr–Fe–Mo


Miscellaneous

[1962Gri] reported the existence of five allotropic modifications of pure chromium at normal pressures in aCr rich region of the Cr-Fe-Mo system. According to [1962Gri], the allotropic Cr structures extend into theternary phase diagram, and phase boundaries of two-phase fields containing the two neighboring allotropicmodifications dependant on Cr or Mo were shown. The sequence of bcc, fcc, bcc, hcp and bcc structureswas found to exist with increasing temperature up to the melting point. There is no other information aboutsuch behavior of pure Cr, and therefore, this paper is not further considered here.The miscibility gap in the ternary Cr-Fe-Mo system was studied by [1971Yam], who found Mo to decreaseits critical temperature compared with that in the binary Cr-Fe system.The phase relations between the α and γ phases in the Fe rich corner were theoretically modeled by[1992Qiu] and [1988Kum]. The results of [1992Qiu] were compared with available experimental resultsof [1931Wev, 1967Bun] and good agreement was found.As mentioned above, the theoretical assessment of the Cr-Fe-Mo system using mainly high temperature datawas realized by [1988And]. This assessment forms part of a number of commercial thermodynamic data-bases and is widely used (along with other thermodynamic data) for predictions of phase stability in com-plex Cr, Mo steels. Certain disagreements between the prediction and experimental results have beensubsequently found for 9Cr-1Mo steels, where the predicted temperature stability of the λ phase was foundto be too low [1997Nat]. The experimental data gave it as approximately 600°C in these steels, but thiscould not be reproduced by the modeling. The thermodynamic description of the Laves and μ phaseswas reassessed by [2002Kro] and [2003Bal] using this information together with first principle calculations.The agreement above 900°C is still very good and the results for lower temperatures, where the λ phaseplays a significant role, are improved.

Table 1. Experimental Studies of the Phase Relations in the Cr-Fe-Mo System


[1931Wev] Thermal analysis on coolingcurves

900-1400°C, Cr/Mo ratio - 2:3, 2:1, 5:1, Fe 90-100 mass%, α+γ region

[1949And] X-ray diffraction (XRD) Cr, Mo, Ni steels

[1950Duw] XRD 650,1375°C, Cr 30-45 mass%, Fe 30-45 mass%,Mo 10-40 mass%, σ phase

[1950Put] Thermal analysis, chemicalanalysis

1100-1300°C, Cr 10-70 mass%, Fe 10-70 mass%, Mo10-60 mass%, σ phase

[1951Bae] XRD, microscopy 650°C, 180 alloys in the whole phase diagram, allphases

[1951Put] Thermal analysis, chemicalanalysis

Melting temperatures, Cr 10-70 mass%,Fe 10-70 mass%, Mo 10-60 mass%, liquidus

[1952Gol] XRD 650°C, whole diagram, α, ε, σ, N phases

[1953Koh] XRD, metallography 650-870°C, Cr 25 mass%, Fe 65 mass%,Mo 10 mass%, τ1 phase

[1954Kas] XRD, neutron diffraction 815, 900°C, Cr 16.8 mass%, Fe 56.2 mass%,Mo 26.9 mass%, τ1 phase

[1954McM] XRD, metallography 815, 900°C, Cr 5-35 mass%, Fe 45-87 mass%,Mo 1-49 mass%, all phases

(continued)

Cr–Fe–Mo 7




[1957Age] XRD 750-1050°C, Cr 60 mass%, Fe 15-25 mass%, Mo25-15 mass%, α, σ phases

[1957Bec] XRD, long term annealing,metallography

815-1205°C, Cr 0-30 mass%, Fe 70 mass%, Mo 0-30mass%, all phases

[1957Gol] XRD 900-1400°C, Cr 13.5-25 mass%, Fe 62-73.5 mass%,Mo 13-23.5 mass%, τ1 phase

[1957Tak] Thermal analysis melting temperatures, whole phase diagram, liquidus

[1962Gri] Thermal analysis, XRD 1400-1700°C, Cr 55 mass%, Fe:Mo=3:1, α phase

[1964Alf] Microstructural analysis, DTA,XRD

from 700°C to melting temperature, Cr 15-90 mass%,Fe 10-55 mass%, Mo 10-30 mass%

1966Kim] Mössbauer spectroscopy 4, 77, 300 K, Cr21Fe62Mo17, τ1 phase

[1971Bau] Specific heat, XRD 1.3-4.2 K, Cr0.75−xFexMo0.25, x = 0-20

[1971Yam] DTA, Mössbauer spectroscopy,optical microscopy

480°C, Cr 28-45 at.%, Fe 50-90 at.%,Mo 0-10 at.%, miscibility gap in α phase

[1976Kie] Microscopy, EDX, ElectronMicroprobe Analysis (EMPA)

500-1000°C, Cr 24-28 at.%, Fe 68-87 at.%,Mo 2-5 at.%, α, τ1, σ phases

[1990Ere] Lattice parameter measurement,XRD

1350°C, Cr21Fe62Mo17, τ1 phase

[1981Zha] Diffusion couple, EMPA,microhardness

1100, 1200°C, whole phase diagram, α, σ phases

[1986Liu] XRD, EMPA, microscopy 850, 1050, 1250°C, Cr 0-25 mass%,Fe 40-95 mass%, Mo 5-60 mass%, R, σ, phases

[1988And] Diffusion couple, EMPA, EDS 950-1200°C, whole phase diagram, all phases

[1988Liu] XRD, EMPA 850-1250°C, Cr 0-25 mass%, Fe 40-95 mass%,Mo 5-60 mass%, all phases

[1989Liu] XRD, EMPA 1250°C, Cr 0-25 mass%, Fe 40-95 mass%,Mo 5-60 mass%, all phases

[1990Liu] XRD, EMPA 850-1250°C, Cr 0-25 mass%, Fe 40-95 mass%,Mo 5-60 mass%, R phase

8 Cr–Fe–Mo




PearsonSymbol/Space Group/Prototype


Comments/References

α, Cr1–x–yFexMoy

(Cr)< 1863(αFe)< 912(δFe)1538 - 1394(Mo)< 2623

cI2Im�3mW

a = 288.48

a = 286.65

a = 293.15

a = 314.70

a = 312.5

0 < x <1, at y = 0 in the range 1513-1394°C and 846-830°C;0 < y < 1 at x = 0 in the range 1820-880°C;at x + y = 1, 0 < x < 0.244; 0.687 < x < 1

at x = 0, y = 0 [Mas2]

at x = 1, y = 0, (αFe) at 25°C, [Mas2](δFe) [Mas2]at x = 0, y = 1 (Mo) [Mas2]

at approximately x = 0.11, y = 0.87[1951Bae]

γ, (γFe)1394 - 846

(γFe)1394 - 912

cF4Fm�3mCu

a = 364.67

dissolves up to 1.7 at.% Mo and up to12 at.% Cr, 7 at.% Cr at 846°C [Mas2]

γFe at 915°C, [Mas2]

σ, Cr1–x–yFexMoy

σ, CrFe830 - 472σ (FeMo)1611 - 1235

tP30P42/mnmCrFe

a = 879.66 ± 0.06c = 455.82 ± 0.03a = 921.8 ± 0.2c = 481.3 ± 0.2a = 879.9c = 456.6

at y = 0, 0.44 < x < 0.505;at x+y = 1, 0.433 < x < 0.571; [Mas2]

at y = 0. x < 0.505 [V-C2]

at x + y = 1, x = 0.5 [V-C2]

at x = 0.615, y = 0.035 [1976Kie]

R, (Fe1–xCrx)3+yMo2–y~1650 - 950

R, Fe3Mo21488 - 1200

hR159R�3Co5Cr2Mo3

a = 1091.0 ± 0.3c = 1935.4 ± 0.5

0.0075 < y < 0.0305 at x = 0 [Mas2]0 < x ≲ 0.21 at y = 0.3 [1988And]down to 930°C [1986Liu, 1989Liu]at x = 0, y = 0.1 [V-C2]

λ, (Fe1–xCrx)2Mo

λ, Fe2Mo< 927

hP12P63/mmcMgZn2

a = 474.4 ± 0.2c = 772.5 ± 0.7a = 474.5c = 773.4

Stoichiometric at x = 0 [Mas2]dissolves ~20 at.% Cr [1988And]Cr content unknown [1957Bec]at x = 0 [1997Vil]

(continued)

Cr–Fe–Mo 9






Comments/References

μ, (Fe1–xCrx)7+yMo6–y

μ, Fe7Mo6< 1370

hR39R�3mFe7W6

a = 475.46 ± 0.05c = 2571.6 ± 0.3

0.28 < y <0.93 at x = 0 [Mas2]dissolves ~20 at.% Cr [1988And]Cr content unknown [1957Bec]at x = 0 [V-C2]

* τ1, Cr1–x–yFexMoy~1180 - 550

cI58I�43mαMn

a = 892.0a = 889a = 885.4a = 895.9 ± 0.5

defined as Cr6Fe18Mo5 by [1954Kas];0.49 < x <0.64 at y = 0.18, 0.35 < x < 0.55at x + y = 0.73 [1988And]at x = 0.62, y = 0.17 [1954Kas]at x = 0.606, y = 0.11 [1953Koh]at x = 0.62, y = 0.12 [1976Kie]for 21Cr-62Fe-17Mo (at.%)quenched from 1350°C [1990Ere]



Cr Fe Mo

L + σ ⇌ α + R* 1471 U1 LσαR

12.816.813.614.1

71.361.370.455.1

15.921.916.030.8

σ ⇌ α + R + τ1* 1026 E1 σαRτ1

10.09.3

10.013.4

63.278.263.261.2

26.812.526.825.4

R ⇌ α + μ + τ1* 991 E2 Rαμτ1

9.310.57.6

16.4

61.082.354.467.4

29.77.2

38.016.2

α + μ ⇌ λ + τ1* 899 U2 αμτ1λ

11.57.0

18.75.4

83.454.566.462.2

5.138.514.932.4

*Reaction prognosed by the theoretical modeling

10 Cr–Fe–Mo


Table 4. Specific Heat of Selected Cr-Fe-Mo Alloys

Alloy Temperature Range Property [mJ·[mol·K]–1] Comments

Cr0.57Fe0.18Mo0.20 1.3-4.2 K C = 234R(T/325)3 + 8.8 T [1971Bau]

Cr0.57Fe0.18Mo0.20 1.3-4.2 K C = 234R(T/144)3 + 10.8 T [1971Bau]

Fig. 1. Cr-Fe-Mo. Calculated liquidus surface projection

Cr–Fe–Mo 11



Fig. 2. Cr-Fe-Mo. Probable partial reaction scheme

12 Cr–Fe–Mo


Fig. 3. Cr-Fe-Mo. Contours for onset

Cr–Fe–Mo 13



Fig. 4. Cr-Fe-Mo. Isothermal section for 1900°C

14 Cr–Fe–Mo


Fig. 5. Cr-Fe-Mo. Isothermal section for 1600°C

Cr–Fe–Mo 15



Fig. 6. Cr-Fe-Mo. Experimental isothermal section at 1250°C

16 Cr–Fe–Mo


Fig. 7. Cr-Fe-Mo. Calculated isothermal section at 1200°C

Cr–Fe–Mo 17




18 Cr–Fe–Mo



Cr–Fe–Mo 19




20 Cr–Fe–Mo



Cr–Fe–Mo 21



Fig. 12. Cr-Fe-Mo. Experimental isothermal section at 850°C

22 Cr–Fe–Mo


Fig. 13. Cr-Fe-Mo. Experimental isothermal section at 700°C α-γ reaction

Cr–Fe–Mo 23



References[1931Wev] Wever, F., Heinzel, A., “Two Instances of Ternary Systems of Iron with a Closed γ Space”

(in German),Mitt. K.-W.-Inst. Eisenforschung, 13, 193–197 (1931) (Morphology, Phase Dia-gram, Phase Relations, Experimental, 14)

[1949And] Andrews, K.W., “A New Intermetallic Phase in Alloy Steels”, Nature, 164, 1015–1015(1949) (Crys. Structure, Experimental, 2)

[1949Jae] Jaenecke, E., “Cr-Fe-Mo” (in German) in “Kurzgefasstes Handbuch aller Legierungen”,Winter Verlag, Heidelberg, 616–617 (1949) (Phase Diagram, Review, 2)

[1950Put] Putman, J.W., Grant, N.J., Bloom, D.S., “σ Phase in Chromium-Molybdenum Alloys withIron or Nickel” in “Symposium on the Nature, Occurrence, and Effects of Sigma Phase”,American Society for Testing Materials, Philadelphia, PA, 61–68 (1950) (Morphology,Phase Diagram, Experimental, 6)

[1950Duw] Duwez, P., Baen, S.R., “X-Ray Study of the σ Phase in Various Alloy Systems” in “Sympo-sium on the Nature, Occurrence, and Effects of Sigma Phase”, American Society for TestingMaterials, Philadelphia, PA, 48–60 (1950) (Crys. Structure, Experimental, Morphology,Phase Diagram, Phase Relations, 17)

[1951And] Andrews, K.V.W, Brookes, P.E., “χ Phase in Alloy Steels”, Metal Treatment and Drop For-ging, 301–311 (1951) (Crys. Structure, Experimental) as quoted in [1953Koh]

[1951Bae] Baen, S.R., Duwez, P., “Constitution of Iron-Chromium-Molybdenum Alloys at 1200F”,Trans. Amer. Inst. Min. Met. Eng., 191, 331–335 (1951) (Crys. Structure, Morphology, PhaseDiagram, Experimental, 13)

[1951Put] Putman, J.W., Potter, R.D., Grant, N.J., “The Ternary System Chromium-Molybdenum-Iron”, Trans. Amer. Soc Met., 43, 824–845 (1951) (Crys. Structure, Morphology, Phase Dia-gram, Experimental, 14)

[1952Bue] Bueckle, H., Plansee Proc., 186 (1952) (Phase Diagram, Experimental)[1952Gol] Goldschmidt, H.J., “Phase Diagrams of the Ternary Systems Fe-Cr-W and Fe-Cr-Mo at Low

Temperature”, Iron Steel Inst. Special Rep., 43, 249–257 (1952) (Phase Diagram, PhaseRelations, Experimental, 27)

[1953Koh] Koh, P.K., “Occurrence of χ Phase in Molybdenum-Bearing Stainless Steels”, Trans. AIME,197, 339–343 (1953) (Crys. Structure, Morphology, Experimental, Mechan. Prop., 3)

[1954Kas] Kasper, J.S., “The Ordering of Atoms in the χ-Phase of the Iron-Chromium-MolybdenumSystem”, Acta Metall., 2, 456–461 (1954) (Crys. Structure, Experimental, Electronic Struc-ture, 4)

[1954McM] McMullin, J.G., Reiter, S.F., Ebeling, D.G., “Equilibrium Structures in Fe-Cr-Mo Alloys”,Trans. Amer Soc. Metals, 46, 799–811 (1954) (Crys. Structure, Morphology, Phase Diagram,Phase Relations, Experimental, 4)

[1957Age] Ageev, N.V., Guseva, L.N., Markovich, K.P., “Phase Transformations in Cr-Rich Cr-Mo-FeAlloys” (in Russian), Izvest. Akad. Nauk SSSR, (Tekhn.), (4), 23–32 (1957) (Phase Relations,Morphology, Experimental, Mechan. Prop., 6)

[1957Bec] Bechtoldt, C.J., Vacher, H.C., “Phase Diagram Study of Alloys in the Fe-Cr-Mo-NiSystem”, J. Res. Nat. Bur. Stand., 58(1), 7–19 (1957) (Phase Diagram, Experimental, Crys.Structure, 29)

[1957Gol] Goldschmidt, H.J., “Occurrence of the β-Mn structure in Transition-Metal Alloys and SomeObservation on χ Phase Equlibria”, Metallurgia, 56, 17–26 (1957) (Crys. Structure, Mor-phology, Phase Diagram, Phase Relations, Experimental, 17)

[1957Tak] Takeda, S., Yukawa, N., “Studies on Cr-Mo-Fe Super Heat-Resistant Alloys”, J. Jpn. Inst.Met., 21, 275–279 (1957) (Phase Diagram, Phase Relations, Experimental) as quoted in[1984Ray]

[1957Wes] Westbrook, J.H., “Temperature-Dependence of the Hardness of Secondary Phases Commonin Turbine Bucket Alloys”, Trans. AIME, 209, 898–904 (1957) (Morphology, Experimental,Mechan. Prop., 68)

[1961Eng] English, J.J., “Molybdenum-Chromium-Iron Sistem” in “Binary and Ternary Phase Dia-grams of Columbium, Molybdenum, Tantalum and Tungsten”, Defense Metals Information

24 Cr–Fe–Mo


Center, Batelle Memorial Institute, Columbus 1, Ohio, 152, 138–141 (1961) (Review, PhaseDiagram, 2)

[1962Gri] Grigor’ev, A.T., Sokolovskaya, E.M., Sokolova, I.G., Maksimova, M.V., “PolymorphicTransformations of Cr and the Constitution of the Cr Solid Solution in the Cr-Fe-Mo TernarySystem” (in Russian), Issled. po Zharoprochn. Splavam, Akad. Nauk SSSR, 8, 42–46 (1962)(Morphology, Phase Diagram, Phase Relations, Experimental, 14)

[1964Alf] Alfintseva, R.A., Dmitrieva, G.P., Korobeinikova, V.G., Pan, V.M., Svechnikov, V.N.,Shurin, A.K., “Study of the Cr-Fe-Mo and Cr-Fe-W Alloys” (in Russian), Sborn. Nauchn.Trudov Inst. Metallofiz., 20, 108–124 (1964) (Crys. Structure, Phase Diagram, Phase Rela-tions, Experimental, Mechan. Prop., 10)

[1966Kim] Kimball, C.W., Phillips, W.C., Nevitt, M.V., Preston, R.S., “Magnetic Hyperfine Interactionsand Electric Quadrupolar Coupling in Alloys of Iron with the α-Manganese Structure”, Phys.Rev., 146, 375–378 (1966) (Crys. Structure, Experimental, Electr. Prop., Magn. Prop., 18)

[1967Bun] Bungardt, K., Kunze, E., Horn, E., Arch. Eisenhuettenwesen, 38, 309 (1967) (Phase Rela-tions, Experimental) as quoted in [1992Qiu]

[1971Bau] Baum, N.P., Schroder, K., “Specific Heat of Chromium-Rich Chromium-Nickel andChromium-Iron-Molybdenum Alloys Between 1.3 and 4.2 K”, Phys. Rev. B: Solid State,3, 3847–3851 (1971) (Thermodyn., Experimental, 29)

[1971Wes] West, D.R.F., “The Constitution of Fe-Rich Alloys of the Fe-Cr-Co-Mo System - a Review”,Kobalt, 51, 68–80 (1971) (Crys. Structure, Phase Diagram, Phase Relations, Review, 104)

[1971Yam] Yamaguchi, M., Umakoshi, Y., Mima, G., “On the Miscibility Gap in the Fe-Cr-X (X = Cu,Mn, Mo, Ni, V, Si and Al) System”, Proc. Int. Conf. Sci. Technol. Iron Steel Tokyo, 11,1015–1019 (1971) (Electronic Structure, Experimental, Phase Relations, 35)

[1975Kau] Kaufman, L., Nesor, H., “Calculation of Superalloy Phase Diagrams: Part IV”,Metall. Trans.A, 6, 2123–2131 (1975) (Phase Diagram, Phase Relations, Calculation, Assessment, 38)

[1976Kie] Kiesheyer, H., Brandis, H., “Investigation of Phase Equilibria in the Ternary System Fe-Cr-Mo in the Solid State” (in German), Z. Metallkd., 67, 258–263 (1976) (Phase Diagram,Experimental, #, 24)

[1979Mat] Mathieu, H.J., Landolt, D., “Influence of Sputtering on the Surface Composition ofFe-Cr-Mo Alloys”, Appl. Surf. Sci., 3, 348–355 (1979) (Morphology, Experimental, Mechan.Prop., 16)

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Cr–Fe–Mo 25



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[1990Ere] Eremenko, V.N., Velikanova, T.Ya., Khar’kova, A.M., Velikanova, T.A., “Intersticial Phaseson the Base of Intermetallic Compounds in Ternary Systems of Transition Metals with Car-bon” (in Russian), Ukrain. Khim. Zhur., 56, 356–363 (1990) (Crys. Structure, Experimental,Phase Relations, Review, 11)

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26 Cr–Fe–Mo


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[2004Mas] Masahashi, N., Hanada, S., “Microstructure Evolution Mechanism on Iron Aluminides/CrMo Steel Composite Prepared by Solid State Bonding”, ISIJ Int., 44, 878–885 (2004),,(Phase Relations, Experimental, Kinetics, 19)

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Cr–Fe–Mo 27



Chromium – Iron – Nitrogen

Pierre Perrot

Introduction

Chromium is widely used in the elaboration of high nitrogen steels, because it is one of the rare elementswhich increase the nitrogen solubility in iron. [1929Bli, 1934Eri] recognized the easy formation of nitrogencompounds by reaction of pure chromium or ferrochromium with nitrogen. The first Cr-Fe-N tentative dia-gram, drawn by [1934Eri] at 400 and 700°C differs from the now accepted diagrams only by the position ofthe (Cr,Fe)-CrN-Cr2N three-phase field. Later, [1942Kra] investigated the evolution of the γ loop in the Cr-Fe system with the nitrogen content, and presented a diagram which is reproduced in the review of[1949Jae]. One of the most complete experimental work was carried out by [1967Ima1, 1967Ima2]. Anextensive review of the phase equilibria may be found in [1987Rag, 1993Rag]. More recent experimentalinvestigations on phase equilibria and thermodynamics are gathered in Table 1.

Binary Systems

The well known Cr-Fe system is accepted from [Mas2]. The Fe-N phase diagram in the solid state isaccepted from the review of [1987Wri]. The Calphad assessment carried out by [1991Fri] and justifiedby the model proposed by [1994Fer] gives an insight on the phase equilibria under high nitrogen pressures.The Cr-N phase diagram, accepted by [Mas2], has been thermodynamically assessed by [1991Fri].The nitrogen solubility in solid and liquid chromium is given by the following expressions [1976Jeh]:In solid chromium: log10(c/mass%) = 0.1 – (1410 / T) + ½ log10(p/Pa) (900–1400°C)In liquid chromium: log10(c/mass%) = –3.14 + (2670 / T) + ½ log10(p/Pa) (1600–1900°C).

Solid Phases

The solid phases are presented in Table 2. Iron and chromium nitrides Fe4N, Fe2N, Cr2N and CrN are stable.However, they are observed only under high nitrogen potentials, obtained for instance with H2-NH3 atmo-spheres at temperatures not higher than 800°C [1967Ima1]. Below 800°C, NH3 decomposes in two steps,the first step being the formation of N + 1.5 H2; N gives a nitrogen potential high enough to nitridize themetals. Above 800°C, the life time of N is too low and the metals are not nitridized.


The liquidus surface, calculated by [1990Fri] and accepted by [1993Rag] is shown in Fig. 1. The E and Uinvariant points were calculated at 1404 and 1463°C, respectively.

Isothermal Sections

The nitrogen solubility in various Cr-Fe alloys [1999Ust] under 0.1 MPa N2 is shown in Fig. 2. In the liquidstate, the nitrogen solubility increases with Cr content of the alloy up to 4.1 mass% N in pure Cr at 1900°Cunder 0.1 MPa of N2 pressure. The nitrogen solubility in a Cr-Fe alloy (18 mass% Cr) at 1600°C under0.1 MPa N2 is accepted at 0.25 to 0.30 mass% [1999Ma2]. Up to 0.1 MPa N2, the nitrogen solubility inliquid alloys obeys the Sievert’s law. At higher nitrogen pressures, a departure is observed (the solubilitygrows more slowly than the square root of the nitrogen pressure), which means that nitrogen has a negativeinfluence on its own solubility [1991Sat]. The departure from the Sievert’s law decreases as the temperaturerises [1984Pom]. A (Fe60Cr40) alloy may dissolve up to 3 mass% N at 1600°C under a nitrogen pressure of2.5 MPa [1976Ras, 1984Pom]. The nitrogen solubility in pure liquid Fe increases with the temperaturewhereas it decreases for pure Cr. The nitrogen solubility does not depend on temperature for Fe95Cr5 alloys[1968Wad, 1972Tor, 1973Lak]. For Fe34Cr67 liquid alloys, the nitrogen solubility varies from 3 mass% N at

Cr–Fe–N 1



1450°C to 0.5 mass% at 1600°C [1972Men]; for Fe70Cr30 liquid alloys, the nitrogen solubilities at the sametemperatures are 0.75 and 0.55 mass%, respectively [1973Sur]. [1974Rab] measures 2 mass% N at 1700°Cfor the Fe30Cr70 alloy under 0.1 MPa N2. The same trend has been observed by [1971Tav, 1975Gra] in γalloys. The nitrogen solubility at 1250°C under 0.1 MPa N2 was measured at 0.6 and 1.1 mass% N for20 and 35 mass% Cr in the alloy, respectively.The isothermal section at 700°C, from [1987Rag] is shown in Fig. 3. In the solid state, isothermal sectionsare characterized by a general trend of the tie lines from Fe to CrN and Cr2N, which is the result of a nitro-gen affinity greater for Cr than for Fe. Isothermal sections of the iron rich corner at 900, 1000, 1100, 1200and 1300°C taken from [1987Rag, 1990Fri, 1993Rag] are shown in Figs. 4 to 8.

Thermodynamics

The interaction coefficients of Cr and N in liquid iron at 1600°C, (< 2 mass% Cr) were first evaluated by[1969Niz] as eN

(Cr) = {∂ log10 fN / ∂ (mass% Cr)} = –0.039 at 1600°C, with fN = {(mass% N in pure Fe) /(mass% N in alloy)} and by [1977Wad] as eN

(Cr) = –0.046 at the same temperature. This value is acceptedby [1990Siw, 1999Ma1] to modelize the nitrogen solubility in liquid Cr-Fe-Mn alloys, whereas [1991Sat]accepts eN

(Cr) = –0.048 while introducing a second order interaction coefficient. The first order interactioncoefficient decreases when temperature rises as shown by [1972Tor, 1973Lak] which measures –0.042,–0.032 and –0.024 at 1600, 1800 and 2000°C, respectively and proposes expressions for the second andthird order interaction coefficients. Such a result contradicts the value eN

(Cr) = –0.045 measured at 2140-2240°C by [1968Uda]. It must be pointed to that [1972Fis] measured eN

(Cr) = –0.044 at 1600°C for aFe-17 % Cr alloy. The presence of Cr has for effect to decrease the activity coefficient of nitrogen in ironand thus, to increase its solubility [1968Nem]. As the Gibbs energy of formation of chromium nitrides isnot very negative, CrN and Cr2N do not precipitate in the alloy and high nitrogen alloys (> 3 mass% N)may easily be synthesized under pressure. The interaction coefficient eN

(N) = 0.13 at 1600°C [1991Sat,1996Hor] describes the nitrogen interaction on itself and cannot be neglected under high nitrogen pressure.In the γ(Fe,Cr) alloys, [1963Mor] measures the interaction coefficient eN

(Cr) = –0.098, –0.121 and –0.136at 1050, 1150 and 1200°C respectively.The interaction coefficients of Cr and N in (γFe) at 1200°C, (< 2 % mass Cr) were first evaluated by[1963Sch] as eN

(Cr) = –0.20 at 1200°C, confirming thus that Cr increases the solubility of nitrogen in iron.[1975Kik] proposed the following expression for the nitrogen interaction coefficient in the γ phase:εN

(Cr) = {∂ log10 γN / ∂ xCr} = 31.0 –(80 600 / T) between 900 and 1400°C, withγN = {(xN in pure Fe) / (xN in alloy)}.The nitrogen solubility in α(Fe,Cr) (18.6 mass% Cr) was measured at 0.09 mass% N at 1485°C under 1 bar(100 kPa) of nitrogen pressure [1964Flo]. This value is in good agreement with more recent measurementsof [1996Kun] which reported a nitrogen solubility at 1400°C under 91 kPa N2 of 0.072 and 0.182 mass% Nfor a chromium content in the α(Fe,Cr) alloy of 14.5 and 24 mass% Cr, respectively. The interactionparameter eN

(Cr) was measured at –0.078 and –0.070 at 1250 and 1350°C respectively [1980Tar].[1996Kun] proposed the following expression for the nitrogen interaction coefficient in δ phase:εN

(Cr) = 3.649 –(30 990 / T) between 1250 and 1500°CThe enthalpy of dissolution of nitrogen in Cr-Fe liquid alloys at 1900°C [1968Wad] is shown in Table 3.Positive for pure iron (nitrogen solubility increases with temperature), it decreases with the Cr content ofthe alloys, goes through zero for Fe95Cr5, presents a minimum at Fe50Cr50, then increases up to pure Cr.A similar behavior is observed in α(Fe,Cr) solid solutions at 1200°C [1975Sch] as shown in Table 3.The Gibbs energy of dissolution of nitrogen presents a minimum towards 71 at.% Cr.Theoretical discussion of the nitrogen solubility in alloys may be found in [1970Kun] which proposes gen-eral equations for the iron rich solutions and models of solid solutions are presented in [1975Gra, 1975Sch,1983Ko, 2001Ma]. Models of liquid solutions are presented in [1986Lin, 1991Sat, 1994Tan, 1999Ma1,1999Ma2, 2001Ma].

2 Cr–Fe–N



The main experimental investigations are reported in Table 4. Nitrogen is known to improve the mechanicalproperties of steels by increasing its hardness and tensile strength [1961Tur, 1999Nak] without any dete-rioration of the toughness [2003Gav]. Cr allows the formation of high nitrogen steels owing the low stabi-lity of chromium nitrides. For the same reason, the presence of Mn in Cr-Fe alloys increases the nitrogensolubility [1969Niz]. There is a wide spread tendency towards the full or partial replacement of C by Nin steels justified by the fact that N provides a better combination of mechanical characteristics, corrosionresistance and weldability. Replacement of C by N in high chromium steels avoids the precipitation ofM23C6 carbides. Nitrogen is also a strong γ stabilizer and may be used instead of Ni. N is considered asthe only element which stabilizes austenite in Cr-Fe alloys [1999Ust]. It is however hard to produce an aus-tenitic structure in high chromium steels by introducing nitrogen, because Cr is an element which contractsthe γ field in iron based diagrams [1994Ust]. An austenitic single phase may be obtained from a nitrided Fe-23 mass% Cr if the steel is rapidly cooled [1999Nak, 1999Tak2] to avoid the eutectoid decomposition ofaustenite into ferrite + Cr2N. On the other hand, [1999Tak1] suggests an addition of 5 mass% Mn to over-come the loss of ductility by delaying the nitride precipitation. The possibility to produce high nitrogen mar-tensitic steels by using N2 injected in the melt or by using a Cr-Fe-N master alloy has been explored by[1996Hor]. The mechanical properties and corrosion resistance of the obtained steel are comparable to thoseof a commercial AISI 420 steel.Thermal analysis [1963Oka] shows that the equilibrium between N2 gas under 1 bar and Cr-Fe alloys (12 to60 mass% Cr) is reached in one hour at 1250°C. Chromium increases the diffusion coefficient of nitrogen inliquid iron [1981Ers] from 5.4 · 10–5 cm2·s–1 for pure iron to 7.6 · 10–5 cm2·s–1 for iron with 7 mass% Cr at1600°C and the rate of nitrogen dissociation and absorption in liquid iron [1986Gla]. The Cr-N bonds arestronger than the Fe-N bonds, which leads to the formation of Cr-N clusters [1999Sum] confirmed by neu-tron diffraction. These clusters brake the movement of the dislocations, but do not change the random Cr-Fedistribution through the matrix. Nitrogen has also been proposed as an inexpensive alternative for argon inthe stainless steel refining process. As nitrogen is absorbed in the melt, it must be flushed out by argon in alater stage.Cr-Fe-N alloys undergos various transient morphologies according to the heat treatment. The precipitationof metastable α”Fe16N2 may be observed by annealing around 400°C an alloy with 2.2 mass% Cr and0.3 mass% N [1996Num]. By quenching an alloy with 18 mass% Cr and 1.35 mass% N, the followingmorphologies may be observed: austenite only, austenite + martensite, austenite + CrN or ferrite + CrN[1996Ust]. CrN precipitated from austenite or martensite has a tetragonally distordered NaCl structure.Stable (Cr,Fe)2N nitride may be precipitated via several intermediate phases such as ε nitride or cubic(Cr,Fe)N solution during tempering [1997Ber].

Miscellaneous

Kinetics measurements of nitrogen absorption in liquid alloys at 1600°C show that Cr decreases slightly theabsorption rate of nitrogen [1967Cho]. The transfer of N across the gas/metal interface is controlled by thetransport in the metal.The nitrogen diffusion was investigated in solid [1972Per1, 1972Per2] and liquid alloys [1973Kun]. Thepresence of Cr decreases the nitrogen diffusion coefficient. For instance, at 1600°C,DN = 14.9 · 10–5 cm2·s–1

for pure iron, and DN = 8.9 · 10–5 cm2·s–1 for Fe85Cr15 alloy.Nitriding Cr-Fe alloys yields the best fatigue resistance of any surface treatment. Several techniques may beused. The precipitation of iron and chromium nitrides α”Fe16N2, γ’Fe4N, CrN and Cr2N is observed[1980Mit] by nitriding Cr-Fe alloys (4 mass% Cr) at 545°C with NH3 prepared by hydrolysis of calciumcyanamide CaCN2. A discontinuous precipitation is observed by nitriding Cr-Fe alloys (< 3.6 mass% Cr)under H2-NH3 (90-10) atmospheres at 560°C [1985Hek]. High nitrogen contents may be obtained by useof a H2-NH3 atmospheres [2003Sch] or N2-NH3-N2O atmospheres [2004Sen1]. Surface segregation withthe formation of a two-dimensional CrN compound was observed [1995Ueb, 1998Mue] by nitriding Cr-Fe alloys (15 mass% Cr) at 1200°C under H2-N2 atmospheres.

Cr–Fe–N 3



Nitrogen plasma offers an attractive means to nitrogenize Cr-Fe alloys [1993Sin, 2000Alv, 2006Goe] inview of rapid absorption to high nitrogen content, up to 0.3 mass%. Plasma nitriding steels after chromiza-tion [1999Cha] increases the microhardness due to chromium nitrides formation.Mechanical alloying of mixtures γ’Fe4N + Cr at room temperature leads to the formation of metastable εphase in which Cr enters in a substitutional solid solution [1989Roc]. Annealing at 190°C gives againthe stable γ’Fe4N. Mechanical alloying of Fe+Cr mixtures under N2 atmospheres gives nanosized grains[1993Ogi, 1995Fuk, 2002Cis] which exhibits up to 12 at.% N in the case of a mixture Fe50Cr50. The nitro-gen content grows exponentially with the milling time. Grains produced under Ar have a bcc structure,whereas N2 produces an amorphization. Interstitial N occupies octahedral sites. Chromium nitride Cr2N pre-cipitate after annealing at 575°C. Phases α, γ, Cr2N and CrN may be formed, depending on the annealingtemperature [2002Cis]. Cr2N instead of N2 may be used to incorporate nitrogen in Cr-Fe alloys by mechan-ical alloying [2002Tsu, 2003Kar].The NGAS (Nitrogen Gas Absorption Sintering) process described by [1996Nak] allows to introduce asmuch as 4 mass% N in an alloy with 23 mass% Cr by sintering the alloy powder under nitrogen atmosphereat 1000°C. The amount of nitrogen absorbed decreases down to 1 mass% N when the sintering temperatureraises up to1200°C. Indeed, the time needed to reach the saturation increases (from 30 to 108 ks) when thetemperature decreases (from 1200 to 1000°C). Mechanical alloying under ammonia of an AISI 304 steel(18 mass% Cr and 10 mass% Ni) allows a rapid uptake up to 7 mass% N [1998Cal], mainly under the formof εFe3N. Annealing at 1000°C precipitates chromium nitrides in the following order: Cr2N, Cr2N+CrN andCrN.Nitrogen increases the absolute coefficient of the thermoelectromotive force [1999Kap] in the stabilitydomain of austenite. This tendency increases sharply near the austenite-martensite transition.

Table 1. Investigations of the Cr-Fe-N Phase Relations, Structures and Thermodynamics


[1958Tur] Equilibrium under H2-NH3 atmospheres, X-rayanalysis

800-1400°C, < 14 mass% Cr,< 0.1 mass% N, 1 bar N2

[1960Hum] Solubility measurements in liquid alloys,Sieverts method

1500-1800°C, Cr-Fe liquid alloys,< 1 mass% N, 1 bar N2

[1960Peh] Solubility measurements in liquid alloys,Sieverts method

1606°C, < 12 mass% Cr, < 1 mass% N,< 1 bar N2

[1961Tur] Solubility measurements in austenitic alloys,Sieverts method

1200-1370°C, < 30 mass% Cr,< 1.6 mass% N, < 1 bar N2

[1963Mor] Solubility measurements in (Fe,Cr) alloys,quenching method

1050-1250°C, < 15 mass% Cr,< 1 mass% N, < 1 bar N2

[1963Oka] Nitrogen absorption measurements in solid (Fe,Cr) alloys

1250°C, 12-60 at.% Cr, < 0.2 mass% N,1 bar N2

[1963Sch] Solubility measurements in solid (Fe,Cr) alloys 700-1200°C, < 2.4 at.% Cr,< 0.01 mass% N, 1 bar N2

[1964Flo] Solubility measurements in solid (Fe,Cr) alloys 1485°C, 18.6 mass% Cr, 0.09 mass% N,1 bar N2

[1967Ima1] Reaction of (Fe,Cr) alloys under H2-NH3

atmospheres800°C, < 40 mass% Cr, < 1 mass% N

(continued)

4 Cr–Fe–N



[1967Ima2] Isothermal sections, X-ray analysis 700-1300°C, < 40 mass% Cr,< 1 mass% N

[1968Nem] Solubility measurements 1600°C, 21-47 mass% Cr,< 0.7 mass% N, 1 bar N2

[1968Uda] Solubility measurements in arc-melted andlevitation molten alloys

2140-2240°C, < 10 mass% Cr,< 5 kPa N2

[1968Wad] Solubility measurements in molten alloys 1780-2200°C, 0-100 % Cr, < 50 kPa N2

[1969Niz] Solubility measurements in molten alloys 1600°C, 20 mass% Cr, 0 to15 mass% Mn, 1 bar N2

[1971Tav] Solubility measurements in ferritic andaustenitic alloys

1250°C, 15 to 35 mass% Cr,< 1.6 mass% N, 1 bar N2

[1972Fis] Solubility measurements in molten alloys 1550-1650°C, 17 mass% Cr,< 0.1 mass% N, < 100 torr (14 kPa) N2

[1972Men] Solubility measurements in molten alloys bythe Sievert’s method

1500-1700°C, 58 to 72 mass% Cr,< 3 mass% N, < 0.1 MPa N2

[1972Tor,1973Lak]

Solubility measurements in molten alloys 1600-2100°C, 0-100 % Cr, 0.1 MPa N2

[1973Sur] Solubility measurements in molten alloys,volumetric method

1550-1700°C, < 30 mass% Cr,< 0.1 MPa N2

[1974Rab] Solubility measurements in molten alloys 1700-1900°C, 59-71 mass% Cr,< 2 mass% N, < 0.1 MPa N2

[1975Gra] Solubility measurements in solid alloys,chemical analysis

900-1200°C, < 16 mass% Cr,< 0.6 mass% N, < 0.1MPa N2

[1975Kik] Solubility and activities measurements in γsolid alloys

900-1400°C, < 2.5 mass% Cr,< 0.1 MPa N2

[1975Sch] Solubility measurements in α solid alloys,weight gains measurements

1200°C, 21 to 99 at.% Cr,< 0.4 mass% N, < 0.1 MPa N2

[1976Rab] Solubility measurements in γ solid alloys 1300-1500°C, < 16 mass% N,0.1 MPa N2

[1976Ras] Solubility measurements in molten alloysunder pressure

1600°C, < 40 mass% Cr, < 3 mass% N,< 2.5 MPa N2

[1977Wad] Solubility measurements in molten alloys,Sievert’s method

1600°C, < 40 mass% Cr,< 0.22 mass% N, < 0.1 MPa N2

[1980Tar] Solubility measurements in α solid alloys 1200-1350°C, 16 to 32 mass% Cr,< 0.2 mass% N, < 0.1 MPa N2

[1984Pom] Solubility measurements in molten alloysunder pressure

1800-2000°C, < 40 mass% Cr,< 3 mass% N, < 1 MPa N2

(continued)

Cr–Fe–N 5




[1985Leb] XRD, chemical analysis of N by Kjeldahlmethod

805-1085°C, < 26.5 mass% Cr,< 0.1 MPa of N2

[1990Siw] Solubility measurements in molten alloysunder pressure

1600°C, 18 mass% Mn, 18 mass% Cr,0.125 to 0.9 MPa N2

[1991Sat] Solubility measurements in molten alloysunder pressure

1600°C, < 35 mass% Cr, < 4 mass% N,0.4 to 10 MPa N2

[1996Kun] Solubility measurements in δ solid alloys 1250-1500°C, 6 to 24 mass% Cr,< 0.2 mass% N, < 91 kPa N2





Comments/References

α, (δFe,αFe,Cr)(δFe)< 1538(αFe)< 912(Cr)< 1983

cI2Im�3mW

a = 293.78

a = 286.65

a = 288.48

pure Fe at 1480°C [V-C, Mas2]

pure Fe at 20°C [Mas2, V-C2](A2 structure)dissolves up to 0.4 at.% N at 590°Cpure Cr at 25°C [Mas2]

(γFe)1394-590

cF4Fm�3mCu

a = 293.16 at 915°C [Mas2, V-C2].Dissolves up to 10.3 at.% N at 650°C[1987Rag]

σCrFe830-440

tP30P42/mnmσCrFe

a = 879.4c = 455.2

44.5-50 at.% Cr[Mas2, 1987Rag]

α”Fe16N2 tI*I4/mmn

a = 572c = 629

ordered fcc structure, metastable[1987Rag]

γ’Fe4N< 680

cP5Pm�3mFe4N

a = 378.7 19.4 to 20.6 at.% N.Ordered fcc structure[1987Rag]

εFe3N< 580

hP3P63/mmlFeN3

a = 259.7b = 434.1a = 278.7b = 444.8

15.8 to 33.2 at.% N [1987Rag]εFe5.23N

εFe2N

ζFe2N< 500

oP12PbcnFe2N

a = 551.2b = 482.0c = 441.6

at 25°C [1987Rag]

(continued)

6 Cr–Fe–N





Comments/References

Cr2N< 1800

hP9P�31mCr2N

a = 478.0 ± 0.1c = 446.8 ± 1.1

30.0 to 33.3 at.% N [1987Rag]

CrN cF8Fm�3mNaCl

a = 414.3 Decomposes at 1000°C under 0.1 MPa N2

[1987Rag]



Quantity, per mole of atoms[J, mol, K]

Comments

½ N2 ⇌ {N} (in liquid Fe) 1500-18001900

ΔrG° = 3650 + 24.06 TΔrH° = + 12000

[1960Peh][1968Wad]

½ N2 ⇌{N} (in liquid Fe70Cr30) 1550-17001800-2000

ΔrG° = – 50000 + 31.34 TΔrH° = – 64000 ± 6000

[1973Sur][1984Pom]

½ N2 ⇌ {N} (in liquid Fe50Cr50) 1900 ΔrH° = – 61000 [1968Wad]

½ N2 ⇌ {N} (in liquid Cr) 1900 ΔrH° = – 50000 [1968Wad]

½ N2 ⇌ {N} (in α Fe) 1200 ΔrG° = + 59000 Extrapolated

½ N2 ⇌ {N} (in α Fe79Cr21) 1200 ΔrG° = + 24450 [1975Sch]

½ N2 ⇌ {N} (in α Fe50Cr50) 1200 ΔrG° = 0 [1975Sch]

½ N2 ⇌ {N} (in α Fe29Cr71) 1200 ΔrG° = – 6950 [1975Sch]

½ N2 ⇌ {N} (in α Fe1Cr99) 1200 ΔrG° = – 1800 [1975Sch]

Table 4. Investigations of the Cr-Fe-N Materials Properties

Reference Method / Experimental Technique Conditions

[1961Tur] Hardness, elongation, tensile strengthmeasurements

1200-1370°C, < 30 mass% Cr, < 1.6 mass% N,< 1 bar N2

[1967Cho] Kinetics of nitrogen absorption 1600°C, < 10 mass% Cr

[1972Per1,1972Per2]

Microscopy, diffusion 400-1000°C, < 13 mass% Cr, < 1.9 mass% N

[1973Kun] Diffusion measurements, steady-statemethod

1550-1700°C, < 15 mass% Cr

[1975Kan] Decarburization and denitrogenizationby plasma

1600°C, 60 mass% Cr, < 0.2 mass% N

(continued)

Cr–Fe–N 7




[1980Mit] Powder nitridizing, X-ray andmetallographical examination

545°C, Cr and Fe nitrides precipitation

[1981Ers] Nitrogen diffusion measurements inliquid Cr-Fe alloys

1600°C, < 7 mass% Cr, < 0.1 MPa N2

[1983Pul] SEM and Auger investigation ofnitrided alloys

500°C, < 3 at.% Cr, N2-NH3 atmospheres,Nitrogen segregation

[1985Hek] SEM, optical, X-ray and hardnessmeasurements

560°C, < 3.6 at.% Cr, N2-NH3 atmospheres

[1986Gla] Kinetics of absorption, isotopicexchange technique

1600°C, < 20 mass% Cr

[1989Roc] X-ray diffraction, Mössbauerspectroscopy

γ’Fe4N + Cr, mechanical alloying, annealing at190°C

[1993Ogi] X-ray diffraction, TEM, DSC,Mössbauer spectroscopy

Fe + Cr, mechanical alloying under N2

atmosphere, annealing at 575°C

[1993Sin] X-ray diffraction, X-ray fluorescenceanalysis

Nitrogenation of Cr-Fe alloys by arc currentplasma, 1700-2100°C

[1993Ueb] SEM, Auger electron spectroscopy 15 mass% Cr, < 0.03 mass% N, surfacesegregation of nitrides

[1994Ust] X-ray diffraction, SEM, hardnessmeasurement

18 mass% Cr, 0.72 mass% N, heat treatments,quenching from < 1250°C

[1995Fuk] X-ray and neutrons diffraction,chemical analysis

Mechanical alloying of Fe30Cr70 under Ar andN2 atmospheres

[1995Ueb] Auger electron spectroscopy, lowenergy electron diffraction

15 mass% Cr, < 0.03 mass% N, surfacesegregation of nitrides

[1996Hor] SEM, TEM observations, hardnessmeasurements

12 to 16 mass% Cr, 0.16 to 0.19 mass% N

[1996Nak] X-ray diffraction, optical observations 23 mass% Cr, < 4.2 mass% N. Nitrogenationby powder sintering

[1996Num] Internal friction, anelastic andmagnetic relaxation

< 2.2 mass% Cr, < 0.3 mass% N, heat treatment

[1996Ust] X-ray, electron microscopy andhardness measurements

18 mass% Cr, < 1.3 mass% N, quenched from< 1250°C

[1997Ber] TEM, Mössbauer spectroscopy,hardness measurements

15 mass% Cr, 0.6 mass% N, heat treatment

[1998Cal] X-ray diffraction, thermal analysis AISI 304 Steel (Fe18Cr10Ni), < 7 mass% N(ball milled under NH3)

[1998Mue] Auger electron spectroscopy, lowenergy electron diffraction

15 mass% Cr, < 0.019 mass% N, surfacesegregation of nitrides

(continued)

8 Cr–Fe–N



[1999Cha] SEM, EDX, X-ray diffraction,microhardness

Steels chromized (1200-1300°C) then nitrided(530°C)

[1999Kap] X-ray diffraction, thermo-electromotive force

18 mass% Cr, 0.42 to 1.48 mass% N

[1999Nak] Microscopic examination, ductilitymeasurements

12 to 23 mass% Cr, 1200°C under 0.1 MPa ofN2

[1999Sum] Neutron diffraction 15 to 24 mass% N, 1 to 3 mass% N

[1999Tak1] SEM, X-ray diffraction 23 % Cr, 5 % Mn, 1 % N, (mass%)

[1999Tak2] Optical microscopy, X-ray diffraction 1000-1300°C, 12 to 23 mass% Cr,< 1 mass% N, powder nitriding

[1999Ust] Chemical analysis, X-ray diffraction 15 to 19 mass% N, 0.9 to 1.3 mass% N

[2000Alv] X-ray diffraction, microhardness < 20 mass% Cr, plasma nitriding (H2-N2

atmospheres)

[2002Cis] SEM, X-ray diffraction 18 Cr, 11 Mn, < 2.5 N, (mass%), ball milling

[2002Kos] SEM, hardness measurements 15 to 24 mass% Cr, 0.4 to 1.3 mass% N, hotrolling

[2002Tsu] SEM, TEM, X-ray diffraction 23 mass% Cr, 1 mass% N, mechanical alloying(Fe-Cr + Cr2N)

[2003Gav] X-ray diffraction, Mössbauer 15 Cr, 1 Mo, 0.62 N (mass%)

[2003Kar] Dilatometry, cyclic polarizationcurves, optical micrography

18 mass% Cr, < 3.3mass% N, mechanicalalloying (Fe-Cr + Cr2N)

[2003Sch] SEM, optical and electron microprobeanalysis

< 20 mass% Cr, < 0.5 mass% N, (H2-NH3

atmospheres)

[2003Wie] TEM, dilatometry, hardnessmeasurements

10 mass% Cr, 0 or 1 mass% Mo,< 0.24 mass% N

[2004Sch] SEM, optical and electron microprobeanalysis

< 20 mass% Cr, < 0.5 mass% N, (H2-NH3

atmospheres)

[2004Sen1,2004Sen2]

SEM, TEM, electron backscatteringdiffraction

1 to 3 mass% Cr, CrN precipitationinvestigations

[2006Goe] Fatigue tests, microhardness profilesdetermination

Plasma nitrided chromium steels

[2006Kam] Diffusion profile measurements Nitridized ferritic steels

Cr–Fe–N 9



Fig. 1. Cr-Fe-N. Partial liquidus surface projection

10 Cr–Fe–N


Fig. 2. Cr-Fe-N. Nitrogen solubility (0.1 MPa N2) in various Cr-Fe alloys (Compositions in mass%)

Cr–Fe–N 11



Fig. 3. Cr-Fe-N. Partial isothermal section at 700°C

12 Cr–Fe–N



Cr–Fe–N 13




14 Cr–Fe–N



Cr–Fe–N 15




16 Cr–Fe–N



Cr–Fe–N 17



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18 Cr–Fe–N


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Cr–Fe–N 19



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20 Cr–Fe–N


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[1995Fuk] Fukunaga, T., Ishikawa, E., Koyano, T., Mizutani, U., “Nitrogen Induced AmorphizationObserved by Neutron Diffraction in the Cr-Fe System”, Physica B, B213-214, 526–528(1995) (Crys. Structure, Phase Relations, Experimental, 7)

[1995Ueb] Uebing, C., Viefhaus, H., Grabke, H.J., “Surface Cosegregation on Fe-3% V-C, Fe-3% V-C,Nand Fe-15%Cr-N (110) Single Crystals”, Fresenius. Z. Anal. Chem., 353(3-4), 254–258(1995) (Interface Phenomena, Experimental, 16)

[1996Hor] Horovitz, M.B., Neto, F.B., Garbogini, A., Tschiptschin, A.P., “Nitrogen Bearing MartensiticStainless Steels - Microstructure and Properties”, ISIJ Int., 36(7), 840–845 (1996) (Experi-mental, Mechan. Prop., Morphology, 14)

[1996Kun] Kunze, J., Broz, P., Stloukal, L., “Thermodynamic Analysis of the Delta Ferrite in the Sys-tems Fe-Cr-N and Fe-Cr-Mn-N”, Steel Research, 67(7), 279–284 (1996) (Phase Relations,Thermodyn., Experimental, 18)

[1996Nak] Nakamura, N., Takaki, S., “Structural Control of Stainless Steel by Nitrogen Absorption inSolid State”, ISIJ Int., 36(7), 922–926 (1996) (Experimental, Mechan. Prop., Morphology,Phase Diagram, Phase Relations, 17)

[1996Num] Numakura, H., Miura, M., Matsumoto, H., Koiwa, M., “Nitrogen Trapping to Chromium in αIron Studied by Internal Friction and Magnetic After-Effect Techniques”, ISIJ Int., 36(3),290–299 (1996) (Crys. Structure, Experimental, Magn. Prop., 40)

[1996Ust] Ustiovshikov, Y., Ruts, A., Bannykh, O., Blinov, V., “The Microstructure of Fe-18%CrAlloys With High N Contents”, Acta Mater., 44(3), 1119–1125 (1996) (Experimental,Mechan. Prop., Morphology, Phase Diagram, Phase Relations, 6)

[1997Ber] Berns, H., Duz, V.A., Ehrhardt, R., Gavriljuk, V.G., Petrov, Yu.N., Tarasenko, A.V., “Preci-pitation During Tempering of Chromium-Rich Iron-based Martensite Alloyed with Carbonand Nitrogen”, Z. Metallkd., 88(2), 109–116 (1997) (Experimental, Mechan. Prop., Morphol-ogy, 16)

[1998Cal] Calka, A., Wexler, D., Zhou, J., Dunne, D., “Nitrogenation During Ball Milling of StainlessSteel”, Mater. Sci. Forum, 269–272, 265–270 (1998) (Experimental, Phase Relations, 21)

Cr–Fe–N 21



[1998Mue] Muehller, C., Uebing, C., Kottcke, M., Rath, C., Hammer, L., Heinz, K., “The Structure ofthe Surface Compound CrN Formed by Cosegregation on a Fe 15%Cr-N(100) Single CrystalSurface”, Surf. Sci., 400(1-3), 87–94 (1998) (Experimental, Phase Relations, Interface Phe-nomena, 29)

[1999Cha] Chang, D.Y., Lee, S.Y., Kang, S.G., “Effect of Plasma Nitriding on the Surface Properties ofthe Chromium Diffusion Coating Layer in Iron-base Alloys”, Surf. Coat. Technol., 116–119,391–397 (1999) (Experimental, Mechan. Prop., Morphology, Phase Relations, 16)

[1999Kap] Kaputkina, L.M., Sumin, V.V., Bazaleeva, K.O., “Influence of Nitrogen on the Tendency toPacking Defect Formation and the Temperature Dependence of the ThermoelectromotiveForce in Cr-Fe Alloys”, Tech. Phys. Lett., 25(12), 992–993 (1999), translated from Pis’maZh. Tekh. Fiz, 25(12), 50–54 (1999) (Experimental, Crys. Structure, Electr. Prop., 7)

[1999Ma1] Ma, Z., Janke, D., “Oxygen and Nitrogen Reactions in Fe-X and Fe-Cr-Ni-X Melts”, SteelResearch, 70(10), 395–402 (1999) (Calculation, Thermodyn., Review, 69)

[1999Ma2] Ma, Z., Janke, D., “Activities of Carbon and Nitrogen in Cr-Fe and Fe-Cr-Ni Melts”, SteelRes., 70(12), 491–495 (1999) (Calculation, Phase Relations, Thermodyn., 25)

[1999Nak] Nakamura, N., Tsuchiyama, T., Takaki, S., “Effect of Structural Factors on the MechanicalProperties of High Nitrogen Austenitic Steels”, Mater. Sci. Forum, 318–320, 209–214(1999) (Experimental, Mechan. Prop., Morphology, Phase Relations, 15)

[1999Sum] Sumin, V.V., Chimid, G., Rashev, Ts., Saryivanov, L., “The Neutron-Spectroscopy Proof ofthe Strong Cr-N Interactions in Nitrogen Stainless Steels”, Mater. Sci. Forum, 318–320,31–40 (1999) (Crys. Structure, Electronic Structure, Experimental, 11)

[1999Tak1] Takaki, S., Nakamura, N., Goto, H., “Alloy Design for Suppressing Eutectoid Reaction inHigh Nitrogen Austenitic Steels”,Mater. Sci. Forum, 318–320, 249–254 (1999) (Experimen-tal, Mechan. Prop., Phase Relations, 8)

[1999Tak2] Takaki, S., Nakamura, N., “Metallurgy of Nitrogen Absorption Process in High Cr Steels”,Mater. Sci. Forum, 318–320, 723–732 (1999) (Experimental, Mechan. Prop., Morphology,Kinetics, 9)

[1999Ust] Ustinovshikov, Y., Ruts, A., Bannykh, O., Blinov, V., Kostina, M., “Microstructure and Prop-erties of the High-Nitrogen Cr-Fe Austenite”, Mater. Sci. Eng. A, 262(1-2), 82–87 (1999)(Experimental, Mechan. Prop., Morphology, Phase Relations, 16)

[2000Alv] Alves, C., Rodrigues, J.D., Martinelli, A.E., “Growth of Nitrided Layers on Cr-Fe Alloys”,Mater. Sci. Eng. A, 279(1-2), 10–15 (2000) (Experimental, Mechan. Prop., Morphology,Interface Phenomena, 11)

[2001Ma] Ma, Z., “Thermodynamic Description for Concentrated Metallic Solutions Using InteractionParameters”, Metall. Trans. B, 32B, 87–103 (2000) (Calculation, Theory, Phase Relations,Thermodyn., 70)

[2002Cis] Cisneros, M.M., Lopez, H.F., Mancha, H., Vazquez, D., Valdes, E., Mendoza, G.,Mendez, M., “Development of Austenitic Nanostructures in High-Nitrogen Steel PowdersProcessed by Mechanical Alloying”, Metall. Mater. Trans. A, 33A(7), 2139–2144 (2002)(Experimental, Phase Relations, Interface Phenomena, 24)

[2002Kos] Kostina, M.V., Dymov, A.V., Blinov, V.M., Bannykh, O.A., “Effect of Plastic Deformationon the Structure and Properties of High-Nitrogen Alloys of the Cr-Fe System”,Met. Sci. HeatTreat., 44(1-2), 9–14 (2002), translated fromMetalloved. i Therm. Obrab. Metallov (1), 8–13(2002) (Experimental, Mechanical Prop., Morphology, 9)

[2002Tsu] Tsuchiyama, T., Uchida, H., Kataoka, K., Takaki, S., “Fabrication of Fine-Grained HighNitrogen Austenitic Steels through Mechanical Alloying Treatment”, ISIJ Int., 42(12),1438–1443 (2002) (Experimental, Mechan. Prop., Morphology, Phase Relations, 17)

[2003Gav] Gavriljuk, V., Rawers, J., Shanina, B., Berns, H., “Nitrogen and Carbon in Austenitic andMartensitic Steels: Atomic Interaction and Structural Stability”, Mater. Sci. Forum,426–432, 943–950 (2003) (Calculation, Experimental, Phase Diagram, Phase Relations, 11)

22 Cr–Fe–N


[2003Kar] Karsokas, N., Rodrigues, D., Ambrosio, F., Toro, A., Tschiptschin, A.P., “Sintered HighNitrogen Stainless Steel Obtained Using Pre-mixed Powders”, Mater. Sci. Forum,416–418, 269–274 (2003) (Experimental, Morphology, Phase Relations, Phys. Prop., 8)

[2003Sch] Schacherl, R.E., Graat, P.C.J., Mittemeijer, E.J., “Modelling the Nitriding Kinetics of Iron-Chromium Alloys”, Mater. Sci. Forum, 426–432, 1047–1052 (2003) (Experimental,Kinetics, Morphology, 19)

[2003Wie] Wiedermann, J., Zalecki, W., Malec, M., “The Influence of Nitrogen on the Structure andProperties of Fe-10Cr-N and Fe-10Cr-1Mo-N Steels After Tempering in the TemperatureRange of 650-750°C”, J. Mat. Proc. Tech., 133(1-2), 225–229 (2003) (Experimental,Mechan. Prop., Morphology, Phase Diagram, 7)

[2004Sch] Schacherl, R.E., Graat, P.C.J., Mittemejer, E.J., “The Nitriding Kinetics of Iron-ChromiumAlloys; The Role of Excess Nitrogen: Experiment and Modelling”, Metall. Mater. Trans. A,35(11), 3387–3398 (2004) (Interface Phenomena, Transport Phenomena, Kinetics, Morphol-ogy, Phase Relations, 47)

[2004Sen1] Sennour, M., Jouneau, P.H., Esnouf, C., “TEM and EBSD Investigation of Continuous andDiscontinuous Precipitation of CrN in Nitrided Pure Cr-Fe Alloys”, J. Mater. Sci., 39(14),4521–4531 (2004) (Experimental, Interface Phenomena, Morphology, Phase Relations, 20)

[2004Sen2] Sennour, M., Jacq, C., Esnouf, C., “Mechanical and Microstructural Investigations ofNitrided Cr-Fe Layers”, J. Mater. Sci., 39(14), 4533–4541 (2004) (Experimental, Mechan.Prop., Morphology, 11)

[2006Goe] Goelgeli, B., Genel, K., “Fatigue Strength Improvement of a Hard Chromium Plated AISI4140 Steel Using a Plasma Nitriding Pre-treatment”, Fatigue Fract. Eng. Mater. Struct., 29(2), 105–111 (2006) (Experimental, Mechan. Prop., Morphology, 32)

[2006Kam] Kamminga, J.-D., Janssen, G.C.A.M., “Calculation of Nitrogen Depth Profiles in NitridedMulti-component Ferritic Steel”, Surf. Coat. Technol., 200(20-21), 5869–5901 (2006)(Experimental, Calculation, Transport Phenomena, Mechan. Prop., 22)



Cr–Fe–N 23



Chromium – Iron – Niobium

Kostyantyn Korniyenko

Introduction

Interest in the chromium-iron-niobium system is due mainly to its alloys being the basis of ferritic heatresisting steels. These steels are based on a tempered martensite matrix, strengthened by solution hardeningwith an element such as niobium. Knowledge of the phase relations in the Cr-Fe-Nb ternary system is ofgreat importance with regard to the optimization of alloy compositions in the preparation of these materials.The available information is presented in the literature in the form of a number of isothermal sections[1980Abr, 1981Abr, 1987Kal, 1988Kal]. Phase contents of the alloys and crystal structures of the inter-mediate phases have been studied by [1972Jaf, 1973Vow, 1980Abr, 1981Abr, 1987Kal, 1988Kal,1993Gru, 1999Zhu, 2002Yam]. The experimental techniques used and the temperature and compositionranges investigated are shown in Table 1. A review of the phase stability of the NbCr2-based Laves phasealloys was presented in [1997Zhu]. However, further studies are necessary, in particular concerning the con-stitution of the quasibinary NbCr2-NbFe2 system, the liquidus, solidus and solvus surfaces, as well as thereaction scheme for the whole range of compositions and alloy thermodynamics.

Binary Systems

The Cr-Fe system is accepted from [Mas2]. The Cr-Nb system is accepted according to [1993Oka] (Fig. 1),based on the work of [1992Tho]. Invariant reaction temperatures and solid solution ranges have beenamended with respect to earlier work [1986Ven, Mas2]. The Fe-Nb system is taken from [1993Bej] (Fig. 2).The earlier data [1986Pau, Mas2] were reinvestigated experimentally, in particular, the homogeneity rangeof the ε phase, which was found to be far narrower. In addition, the composition of the μ phase was given asNb19Fe21 rather than NbFe. The last phase does not form congruently, but peritectically at 1520°C throughthe reaction L + ε⇌ μ. On the basis of experimental studies, [1993Bej] concluded that the ω (Nb-Fe) phasewith the approximate composition of Nb3Fe2 is metastable, forming only during certain cooling conditions.The observation of thermal events around 1490°C and 1460°C in DTA studies of some samples in this com-position range was attributed to the metastable peritectoid formation and eutectoid decomposition of thisphase, respectively. The corresponding invariant horizontal lines and approximate homogeneity range ofthe ω phase are plotted in Fig. 2 by dotted lines.

Solid Phases

Crystallographic data relating to the unary, binary and ternary phases are listed in Table 2. The maximumsolubility of the third component in the binary Cr-Fe, Cr-Nb and Fe-Nb phases is observed in the ε phase.Both at 1000 and at 700°C, the solubility of chromium in the ε phase is about 60 at.%. One ternary phasehaving a crystal structure different from any of the unary and binary phases was found; τ, Nb(Cr,Fe)2[1993Gru]. It was observed in a Nb67.20Cr18.47Fe14.33 alloy, annealed at 1150°C for 96 h and water-quenched. The ternary phase appeared together with (Nb) and the ε phase. Probably, this phase exists onlyover a limited temperature range. For example, under equilibrium conditions at 1000 and at 700°C, thisphase was not observed by [1980Abr, 1981Abr, 1987Kal, 1988Kal, 1999Zhu].

Quasibinary Systems

The βNbCr2-NbFe2 section seems to be quasibinary because both limiting phases have the same crystalstructure and form congruently in the binary systems. However, the constitution of this quasibinary systemis not known.[1972Jaf] reported quasibinary morphology for the eutectic in an alloy with composition Nb13.3Cr23Fe63.7(mass%, in at.% - Nb8.29Cr25.62Fe66.09). This alloy solidified at a temperature of 1275 ± 10°C. It contained

Cr–Fe–Nb 1



approximately 22 vol% of the intermetallic phase (Nb48.11Cr10.43Fe41.46 in at.%, hexagonal close-packedcrystal structure with a = 480.0 pm, c = 763.57 pm) in a Cr-Fe solid solution (α) matrix(Nb0.27Cr26.37Fe73.36 in at.%). The lattice parameters of the latter phase are listed in Table 2. It was supposedthat the intermetallic phase is not a simple modification of the ε phase in spite of their similar structures.However, the tie line joining these two phases cross the NbCr2-NbFe2 section, which is supposed to be qua-sibinary.

Isothermal Sections

The isothermal section for 1000°C for the whole range of compositions is presented in Fig. 3, based on theexperimental data of [1980Abr, 1981Abr, 1987Kal]. Slight amendments have been made to maintain con-sistency with the accepted binary phase diagrams. The section shows a significant homogeneity region forthe solid solution based on the hexagonal Laves phase ε, NbFe2. A considerable part of the isotherm is occu-pied by two and three-phase equilibria. The reported results agree well with [1999Zhu] relating to theNbCr2-NbFe2 section and the e/a correlation with phase stability.Figure 4 shows the isothermal section for 700°C for the whole range of compositions, based on the experi-mental data of [1988Kal]. The extent of the α solid solution region in the Cr-Fe binary system given in[1988Kal] is somewhat different to that in the accepted binary diagram, and hence, the tie lines of theε + α + σ from the α corner are denoted by dashed lines. The character of the phase equilibria over the rangeof compositions 30 to 100 at.% Nb at a temperature of 700°C is similar to that at 1000°C. The positions ofthe three-phase regions (Nb) + μ + ε and ε + (Nb) + αNbCr2 in Figs. 3 and 4 are plotted on the basis ofmicroprobe results listed in [1987Kal] and [1988Kal], respectively. These data were preferred because theirpositions presented by authors in the isothermal sections contradict the tabulated data.

Thermodynamics

Experimental investigations of the thermodynamic properties Cr-Fe-Nb alloys have not been reported. Ther-modynamic data for liquid, ferrite and austenite of some iron-based systems were optimized by [1998Mie]in order to give a more accurate representation of phase equilibria between these phases in multicomponentsteels. But for the Cr-Fe-Nb system, the author has noted that a careful examination of the system is neces-sary in order to get a reliable prediction of the partitioning tendency of niobium.


Alloys of the Cr-Fe-Nb system are the basis of ferritic heat resisting steels that are widely used in powergenerating plants for high creep resistance and superior oxidation resistance [2002Yam]. Body-centeredcubic solid solutions based on niobium (in particular, Cr-Fe-Nb) were proposed by [2001Esa] as potentialmaterials for hydrogen storage. Moreover, the investigation of alloys of this system is expedient in connec-tion with the need to develop technological processes for obtaining new magnetically hard materials[1988Kal]. Literature data concerning investigation of the materials properties of chromium-iron-niobiumalloys are listed in Table 3.

Table 1. Investigations of the Cr-Fe-Nb Phase Relations, Structures and Thermodynamics

Reference Method / Experimental Technique Temperature / Composition / Phase RangeStudied

[1972Jaf] Metallography, SEM, X-ray diffraction ≤ 1200°C, Nb8.29Cr25.69Fe66.02

[1973Vow] Chemical analysis, TEM, X-raydiffraction(Debye-Scherrer technique)

≤ 1380°C, Nb1.08Cr16.05Fe82.87,Nb0.48Cr9.63Fe89.89

(continued)

2 Cr–Fe–Nb



[1980Abr] X-ray diffraction, optical microscopy 1000°C, whole range of compositions

[1981Abr] X-ray diffraction, optical microscopy 1000°C, whole range of compositions

[1987Kal] Neophot-2 optical microscopy, powderX-ray diffraction, Cameca MS-46microprobe analysis

1000°C, whole range of compositions

[1988Kal] Neophot-2 optical microscopy, powderX-ray diffraction, Cameca MS-46microprobe analysis

700°C, whole range of compositions

[1993Gru] X-ray diffraction (room andhigh-temperature)

≤ 1150°C; 78 mass% Nb, 0 to 22 at.% Cror Fe

[1999Zhu] X-ray diffraction (Scintag XD200diffractometer)

≤ 1000°C; 33.3 at.% Nb, 0 to 66.7 at.% Fe

[2002Yam] Optical microscopy, SEM, TEM ≤ 1100°C; 10 at.% Cr, 0.5 to 3.5 at.% Nb

[2005Yam] TEM 600°C, 700°C; Nb1Cr10Fe89





Comments/References

α, (Cr1–xFex)

(Cr)< 1863

(δFe) (h2)1538 - 1394(αFe) (r)< 912

cI2Im�3mW

a = 288.4a = 288.15a = 288.51a = 293.15

a = 286.64a = 287.29

a = 287.93

a = 287.52

a = 287.57

a = 287.32

0 ≤ x ≤ 1. Dissolves 5.6 at.% Nb at x = 0, 1668°C[1992Tho, 1993Oka]at x = 1 dissolves 3.27 at.% Nb at 1370°C [1993Bej],0.73 at.% Nb at 960°C [1993Bej]pure Cr, T = 27°C [V-C2][1958Ere][1992Tho]pure Fe, T = 1394°C [V-C2]

pure Fe, T = 20°C [V-C2]in the Nb8.29Cr25.62Fe66.09 as-cast alloy, together with asecond phase [1972Jaf]in the Nb1.08Cr16.05Fe82.87 alloy solution treated at1380°C [1973Vow]in the Nb1.08Cr16.05Fe82.87 alloy solution treated at1380°C and aged at 630°C [1973Vow]in the Nb0.48Cr9.63Fe89.89 alloy solution treated at1380°C [1973Vow]in the Nb0.48Cr9.63Fe89.89 alloy solution treated at1380°C and aged at 630°C [1973Vow]dissolves 3 at.% Nb at 1000°C

(continued)

Cr–Fe–Nb 3






Comments/References

(γFe) (h1)1394 - 912

cF4Fm�3mCu

a = 364.68 pure Fe, T = 912°C [V-C2]dissolves 11.9 at.% Cr at ~1100°C[Mas2]dissolves 1 at.% Nb at 1190°C [1993Bej]

(εFe) hP2P63/mmcMg

a = 246.8c = 396

T = 25°C [Mas2]high pressure phase > 1.3·105 bar

(Nb)< 2469

cI2Im�3mW

a = 330.04a = 330.62

pure Nb, T = 25°C [Mas2][1992Tho]dissolves 24.4 at.% Cr at 1703°C [1992Tho, 1993Oka]

σ, CrFe830 - 440

tP30P42/mnmCrFe

a = 879.95c = 454.42

50.0 to 55.5 at.% Fe [Mas2][V-C2]dissolves 10 at.% Nb at 700°C [1988Kal]

βNbCr21730 - ~ 1585

hP12P63/mmcMgZn2 a = 493.01

c = 811.67

~30 to ~39 at.% Nb[1992Tho, 1993Oka][1961Pan]

αNbCr2≲ 1625

cF24Fd�3mMgCu2 a = 698.54

a = 699.71

a = 699.13

a = 699

~30 to ~39 at.% Nb[1992Tho, 1993Oka][1958Ell]at 34.29 at.% Nb, specimen annealed at 1400°C during100 h [1992Tho]at 32.88 at.% Nb, specimen annealed at 1400°C during100 h [1992Tho][1993Gru]dissolves 4 at.% Fe at 1000°C [1987Kal]

ε, NbFe2< 1630

hP12P63/mmcMgZn2

a = 483.39c = 788.02a = 482.5c = 787.9a = 483.8c = 788.9a = 483.3c = 784.2

a = 480 to 495c = 787 to 810

32 to 37 at.% Nb [1993Bej][1958Ell]

[1973Vow]

[1992Rag]

[1993Gru]

dissolves 60 at.% Cr at 1000°C [1987Kal]at 0 to 60 at.% Cr,T = 1000°C [1987Kal]

(continued)

4 Cr–Fe–Nb





Comments/References

μ, Nb19Fe21< 1520

hR39R�3mW6Fe7 a = 492.8

c = 2683

48 to 54 at.% Nb [1993Bej]

[1992Rag]

dissolves 7 at.% Cr at 1000°C [1987Kal] and 700°C[1988Kal]

ω (Nb-Fe) - - metastable [1993Bej]

*τ, Nb(Cr,Fe)2 hP24P63/mmcMgNi2

a = 486c = 1531

at 1150°C [1993Gru]

Table 3. Investigations of the Cr-Fe-Nb Materials Properties


[1972Jaf] TF-KM model Instron tensile testingmachine tests

Microhardness, ultimate tensile strength,ductility

[1973Vow] Aging of the specimens, hardnessmeasurements

Hardness, aging-time dependence

[1981Abr] Mechanical properties tests Hardness, microhardness

[1987Kal] Vickers hardness (PT type tester) andmicrohardness (PMT 3 type tester)measurements; magnetic measurements

Hardness, microhardness, magneticsusceptibility

[1988Kal] Vickers hardness (PT type tester) andmicrohardness (PMT 3 type tester)measurements

Hardness; microhardness

[1992Wol] Tensile deformation tests Tensile ductility; toughness

[1999Zhu] Mechanical properties tests Vickers hardness; fracture toughness

[2002Yam] Tensile tests Elongation to fracture; tensile 0.2 % flowstress

Cr–Fe–Nb 5



Fig. 1. Cr-Fe-Nb. The Cr-Nb phase diagram

6 Cr–Fe–Nb


Fig. 2. Cr-Fe-Nb. The Fe-Nb phase diagram

Cr–Fe–Nb 7



Fig. 3. Cr-Fe-Nb. Isothermal section at 1000°C

8 Cr–Fe–Nb


Fig. 4. Cr-Fe-Nb. Isothermal section at 700°C

Cr–Fe–Nb 9



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10 Cr–Fe–Nb


[1997Zhu] Zhu, J.H., Liaw, P.K., Liu, C.T., “Effect of Electron Concentration on the Phase Stability ofNbCr2-based Laves Phase Alloys”, Mater. Sci. Eng. A, 239–240, 260–264 (1997) (PhaseRelations, Review, Electronic Structure, 30)

[1998Mie] Miettinen, J., “Approximate Thermodynamic Solution Phase Data for Steels”, Calphad, 22(2), 275–300 (1998) (Phase Diagram, Phase Relations, Thermodyn., Assessment, Calcula-tion, 83)

[1999Zhu] Zhu, J.H., Liu, C.T., Liaw, P.K., “Phase Stability and Mechanical Behavior of NbCr2-basedLaves Phases”, Intermetallics, 7, 1011–1016 (1999) (Crys. Structure, Phase Relations,Experimental, Electronic Structure, Mechan. Prop., *, 25)

[2001Esa] Esayed, A.Y., “Metal Hydrides {Hydrogen Storage}”, Energy Sources, 23(3), 257–265(2001) (Crys. Structure, Phase Relations, Review, Phys. Prop., 13)

[2002Yam] Yamamoto, K., Kimura, Y., Wei, F-G., Mishima, Y., “Design of Laves Phase StrengthenedFerritic Heat Resisting Steels in the Fe-Cr-Nb (-Ni) System”, Mater. Sci. Eng. A, 329–331,249–254 (2002) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan.Prop., 8)

[2005Yam] Yamamoto, K., Kimura, Y., Mishima, Y., “Precipitation Behavior and Phase Stability of Inter-metallic Phases in Fe-Cr-W-Co Ferritic Alloys”, Mat. Sci. Forum, 475–479, 845–848 (2005)(Morphology, Phase Relations, Experimental, Electronic Structure, 14)



Cr–Fe–Nb 11



Chromium – Iron – Nickel

Hans Leo Lukas, Pavel Agraval

Introduction

The Cr-Fe-Ni system is of primary interest for ferrous and nonferrous metallurgy. In ferrous metallurgy itsimportance is in the production and application of stainless steels. The basic types are austenitic (~18 mass% Cr, 8 mass% Ni and ~0.1 mass% C), austenitic-ferritic (13-25 mass% Cr, 9-37 mass% Ni and ~0.1 mass%C), ferritic-martensitic (13-30 mass% Cr and up to 1.5 mass% C) and martensitic (13-18 mass% Cr and 1-4mass% Ni) stainless steels with excellent corrosion resistance and high temperature strength. Cr is one of themain alloying elements in Ni based alloys with improved heat resistance, as Incolloy- and Hastelloy typealloys. The ferromagnetism near the composition FeNi3 leads to alloys with very small thermal expansion(Invar-effect). Cr based alloys are not widely used because generally they are very brittle, but brittleness canbe overcome by purification [2002His]. Alloys of the Cr-Fe-Ni system also are used to produce very softmagnetic and electrotechnical materials. Purely austenitic steels have special interest as nonmagnetic steels.The first experimental investigations of ternary phase diagram data are from [1927Bai], [1927Che] and[1930Jen]. [1931Wev] investigated the system above 1100°C and constructed the liquidus surface and sev-eral vertical sections showing the phases liquid, ferrite (α) and austenite (γ). The existence of the σ phase,already found by [1927Bai], was confirmed. These authors classified their work as preliminary due to prob-ably insufficient annealing to get true equilibrium. [1937Jen] studied the area with less than 50 mass% Crand got similar results, but assumed a minimum of the (liquid + α + γ) three-phase equilibrium near the Cr-Ni binary system. [1939Sch] studied the stability range of the σ phase and constructed isothermal sectionsat 650 and 800°C. Homogeneity range of the σ phase was investigated by many authors [1941Bra,1950Pug, 1952Coo, 1994Sop, 1996Sop]. Several papers report measurements of the stability range of theσ phase in commercial steel or multicomponent alloys containing additionally C, Si or Mn [1940Mon,1952Lis, 1952Nic, 1953Tal, 1956Pry]. Extrapolations of these values support the results measured at pureternary Cr-Fe-Ni alloys.The (α + γ) field was the topic of investigations at 1300°C by [1959Pri], from 816 to 1260°C, below 40mass% Fe by [1972Sch] and at 1200 to 1350° in 50°C steps by [1983Mun1]. The tie lines of solid (α orγ)-liquid equilibria in the Fe corner (more than 70 at.% Fe) were measured by annealing in the two-phasefields and analyzing the two phases after quenching [1977Sch1, 1977Sch2, 1988Kun, 1994Sch, 1994Sop].To get a planar interface between solid and liquid [1988Kun] used directional solidification. Besides theseequilibrium determinations the kinetics of solidification was extensively studied by several researchers[1976Her, 1982Fre, 1995Ohs, 1996Bob, 1996Ume, 1997Moi, 1997Vol1, 1997Vol2, 1998Kim]. Thekinetics of the α to γ transformation after solidification is the topic of several papers [1983Mun2,1993Kaj, 1995Kaj, 1995Vit, 1998Var, 1999Fuk, 2000Fuk]. It is reported that from undercooled melts itis possible to get metastable primary α, where γ is the stable equilibrium phase and vice versa [1994Loe,1996Kos, 1996Loe, 1996Vol, 2001Vit].The γ phase has a tricritical point in the Fe-Ni binary system, occurring very close to where the Curie linemeets the (α+γ)/γ boundary. The corresponding miscibility gap was studied by [1989Cha] using meteoritespecimen and samples heavily irradiated by electrons. It is reported that the demixed samples do not showthe Invar-effect associated with the second-order change of magnetization near the Curie temperature.Thermodynamic experimental studies were mainly concerning activities of components in molten alloys,mostly using Knudsen-cells connected with mass spectrometry [1958Lyu, 1969Gil, 1976Vre, 1979Pro,1980Skr, 1995Bro, 1998Tom, 1998Vre]. Enthalpies of mixing of solid phases (α and γ) were determinedby [1967Kub] using an adiabatic direct reaction calorimeter. Enthalpy of mixing of liquid was determinedby [1998Thi] using levitation calorimetry. [1984Der] measured low temperature heat capacities of certainalloys.All papers containing original experimental values of phase diagram or thermodynamic data are summar-ized in Table 1.

Cr–Fe–Ni 1



There are many reports on thermodynamic modelling and calculation of the Cr-Fe-Ni phase diagram[1972Cou, 1974Kau, 1977Kau, 1980Cha, 1985Her, 1987Chu, 1988Yam, 1989Cha, 1990Hil, 1990Yam,1993Lee1, 1999Mie, 2004Tom].A critical review of the ternary Cr-Fe-Ni system was published by Raynor and Rivlin [1981Ray1,1988Ray]. Two updates of this review are due to Raghavan [1994Rag, 2003Rag].

Binary Systems

The binary systems are accepted from [Mas2]. The temperature of decomposition of the σ phase wasaccepted by [Mas2] for the binary Cr-Fe system to be about 440°C. It is most probably somewhat higher[1981Ray2]. In thermodynamic calculations with recently assessed binary Cr-Fe datasets [1985Her,1987And, 2004Tom] this temperature appears between 509 and 512°C.Phase diagrams presented in this report are generated using thermodynamic calculations using Gibbs energyfunctions of various phases. The required functions are taken from [1993Lee1], that uses binary datasets of[1993Lee2] (Cr-Fe), [1992Lee] (Cr-Ni) and [1993Lee2] (Fe-Ni). The calculated binary phase diagramsusing these datasets agree well with the binary systems given in [Mas2]. The rationale for selecting thisdataset from many reported in the literature ([1972Cou, 1974Kau, 1977Kau, 1980Cha, 1985Her,1987Chu, 1988Yam, 1989Cha, 1990Hil, 1990Yam, 1993Lee1, 1999Mie, 2004Tom]) is given below.As ferromagnetism is important in the Fe rich α phase and in the Ni rich γ phase, for temperatures below1000°C, a thermodynamic modeling that takes into account Gibbs energy contribution from ferromagnetismis preferred. This contribution is described by the present state of the art of the Calphad method that useseither the formalism of [1976Ind], modified by [1978Hil], or a similar formalism established by[1986Chu]. Such an approach is used in most of the datasets published beyond 1985, except in[1988Kun] and [2004Tom]. For the liquidus-solidus equilibria all these datasets give reasonably goodagreement with experimental results. The Ni rich (liquid+γ) field of the binary Fe-Ni system calculatedusing the dataset of [2004Tom] is claimed to give better agreement with experimental data than those cal-culated using the datasets by [1990Hil, 1993Lee2, 1999Mie]. However, this is difficult to decide, as the dif-ference is within the scatter of the experimental data. In this context the dataset from [1993Lee1] seems tobe the most reasonable to be accepted for phase diagram calculations of the ternary system.

Solid Phases

The crystallographic data of all solid phases are summarized in Table 2. There are no ternary phasesreported. The γ phase forms two different superstructures, a cubic Cu3Au type at FeNi3 and an orthorhom-bic MoPt2 type at CrNi2. Both superstructures extend only very little into the ternary composition range[1981Gom]. For the CrNi2 superstructure no lattice parameters are reported. Those given in Table 2 are cal-culated from a = 356.14 pm, given by [1949Ree] for the γ-substructure cell of Cr32Ni68, neglecting a pos-sible deformation of the cubic substructure cell.The α phase at lower temperatures shows a large miscibility gap between Cr rich and Fe rich compositions.The critical temperature of this miscibility gap is masked by equilibria of the σ phase.


A four-phase equilibrium σ ⇌ α1 + α2 + γ, Table 3, follows from the eutectoid decomposition of the σphase in the Cr-Fe binary system. Its temperature is not experimentally determined. Calculations with thedatasets of [1985Her, 1990Hil, 1993Lee1, 1999Mie, 2004Tom] give values between 460 and 490°C.Another four-phase equilibrium γ1 ⇌ γ2 + α1 + α2 follows from the tricritical miscibility gap of the γphase. Calculations with the above datasets except [2004Tom] find this equilibrium between 340 and380°C. The dataset of [2004Tom] does not find this four-phase equilibrium, which is a result of ferromag-netism of the system. The reaction scheme in Fig. 1 is drawn only for temperatures above 400°C, as at lowertemperatures equilibrium cannot be reached in reasonable time and thus calculated equilibria cannot beproved experimentally. Two of the datasets [1990Hil, 1999Mie] give a minimum of the three-phase equili-bria (L + α + γ) near the Cr-Ni binary eutectic (1345°C), at 1333 or 1311°C respectively, whereas the otherthree datasets show no such minimum. The value 1333°C is comparable with the experimental data of

2 Cr–Fe–Ni


[1937Jen], whereas 1311°C may be considered too low. A three-phase field without minimum agrees withthe data reported by [1931Wev].

Liquidus and Solidus Surfaces

The liquidus surface projection, calculated using the dataset of [1993Lee1] is shown in Fig. 2. In the Fe richcorner the dataset of [1999Mie] gives best fit to the measurements of Schürmann et al. [1977Sch1,1994Sch], Kundrat and Elliott [1988Kun], Yamada and Umeda [1986Yam, 1988Yam], but elsewhere theagreement is poor. The liquidus isotherms in the Ni rich part are not well established, they differ slightlyin the reported experimental results [1931Wev, 1937Jen]. The solidus surface projection, calculated withthe same dataset as Fig. 2, is shown in Fig. 3. It contains boundaries of the single-phase fields α and γas well as the two-phase field (α + γ).

Isothermal Sections

Several isothermal sections are shown in Figs. 4 to 9, calculated with the dataset of [1993Lee1]. Betweenabout 1345 and about 940°C the only stable phases are α and γ. Below 940°C the extension of equilibriacontaining the σ phase is of interest, as already small amounts of this phase considerably may embrittle analloy. At 450°C and below the calculated isothermal sections show the miscibility gap of γ due to ferromag-netism, for the existence of which [1989Cha] gave experimental evidence. Near the tricritical point theGibbs energy vs composition is so flat, that it is impossible to calculate the miscibility gap until it closes.The formation of the ordered phases FeNi3 and CrNi2 of the γ phase at lower temperatures is not taken intoaccount in Figs. 4 to 11.


Figures 12 and 13 show two calculated vertical sections, both passing through the nominal composition ofthe most commonly used stainless steel (18 mass% Cr, 8 mass% Ni).

Thermodynamics

Enthalpies of formation of solid phases at 1292°C (α and γ phase) were determined by direct reactioncalorimetry [1967Kub]. Enthalpies of mixing of the liquid phase were measured calorimetrically[1998Thi] by dropping pieces of solid Cr into an electromagnetically levitated Fe-Ni melt near 1625°C from0 to 40 at.% Cr. Chemical potentials of all three elements in the liquid and solid state were determined byKnudsen cell effusion. The evaporation rates were determined either by analysis of the condensate on a tar-get [1958Lyu, 1969Gil, 1976Vre, 1980Skr] or by mass spectrometry [1969Gil, 1979Pro, 1983Vre,1995Bro, 1996Vre, 1998Tom, 1998Vre]. Differences between the measurements are somewhat larger thanthe expected experimental scatter.Several assessed thermodynamic datasets of Cr-Fe-Ni covering the whole system are reported in the litera-ture [1972Cou, 1974Kau, 1985Her, 1990Hil, 1993Lee1, 1999Mie, 2004Tom]. Some assessed datasetsare intended to describe only selected parts of the system: the liquid-solid equilibria in the Fe rich corner[1988Kun, 1988Yam, 1990Yam, 1999Mie], equilibria containing the σ phase [1977Kau, 1980Cha], orequilibria influenced by the contribution of ferromagnetism to the Gibbs energy [1987Chu, 1989Cha].The Cr-Fe-Ni system contains large phase fields of ferromagnetic phases. The temperature dependence ofthe contribution of ferromagnetic ordering to the total Gibbs energy cannot be assessed by using only poly-nomial terms to describe the excess Gibbs energy. The integration of the shaped Cpvs T dependence isnecessary as proposed by [1976Ind] and modified by [1978Hil] or similarly by [1987Chu]. This type oftreatment is considered only in the datasets of [1985Her, 1987Chu, 1989Cha, 1990Hil, 1993Lee1,1999Mie]. The other ones may be used for calculations far above any Curie temperature, like all equilibriacontaining the liquid phase, but in the vicinity of a Curie temperature they should be avoided.Enthalpies and Gibbs energies calculated employing different datasets disagree slightly more than the esti-mated accuracies of the experimental data on which they are based. The enthalpies of mixing of liquid from[1998Thi] are best represented by the datasets of [1990Hil, 1993Lee1], whereas the excess Gibbs energies

Cr–Fe–Ni 3



of liquid derived from the measurements of [1996Vre, 1998Vre] are better represented by [1999Mie,2004Tom]. For the liquidus and solidus surfaces the datasets of [1993Lee1, 2004Tom] give best fit.In the Fe rich corner [1999Mie] may be somewhat better than [1993Lee1], but at the expense of accuracyoutside this composition area. The statement of [2004Tom], that the dataset of [1993Lee1] is worse thanhis own one is not correct, it may be caused by a conversion error of Lee's data from the Redlich-Kister-Muggianu description into the TAP series used by [2004Tom], as indicated by the original figure 3 of[2004Tom] and original figure 7a of [1998Tom].


One of the technically important problems of austenitic stainless steels is the problem of cracking. Usuallycracking can be largely avoided when certain amount of δ ferrite forms. Consequently, in the modern litera-ture great attention is given to the problem of α-γ [1998Var, 2005Yan] or δ-γ [1995Ino, 1996Bob,1996Kos, 1997Moi, 1997Vol1, 1997Vol2] transformations. Mechanical properties of austenitic steels andsuperalloys are subjects of investigation in [1987Kli, 1994Mar, 1999Har, 2003Ale, 2003Cai, 2005Ye].Metastable formation of various phases was investigated by [1987Mal, 1996Kos, 1996Kuw, 1997Moi,1997Vol1, 1997Vol2, 1999Fuk, 2000Fuk]. Physical and magnetic properties of stainless steels and alloyswere reported by [1986Ric, 1987Bol, 1991Nat, 1994Mar, 1997Sin, 1998Tak, 1999Kak, 2000Ban].The experimental works devoted to study of materials properties in the Cr-Fe-Ni system are listed in Table 4.[2000Mie] obtained fitted equations for liquidus temperatures of steels from calculated results based on amodel of interdendritic solidification. [2002His] summarized literature data on mechanical properties ofhigh-chromium iron-base alloys and chromium-base alloys. This review covers data obtained up to 2002.Literature data on elastic constants of seven alloy compositions are given in [2004Li].Austenitic and austenitic-ferritic (duplex) stainless steels. Linear expansion of Cr-Fe-Ni alloys [1987Bol]has no anomalies at the Curie point, indicating a large temperature interval of the transition from paramag-netic to ferromagnetic state. The temperature dependences of elastic moduli of Cr-Fe-Ni alloys duringvarious magnetic transformations were investigated by [1987Kli]. [1994Mar] investigated the variationsin lattice parameter, electrical resistivity and microhardness in Cr-Fe-Ni alloys. Solidification and ferrite-austenite transformation in austenitic stainless steels were studied by [1995Ino, 1996Bob]. Kinetics of α-γ phase transformation was investigated by [1998Var]. The thermal residual elastic strains in α and γ phaseshave been measured by [1999Har]. The influence of hydrogen charging time on the microstructure of α andγ phases was investigated by [2003Glo]. [2005Ram] calculated multicomponent phase diagrams that weresuccessfully used to explain the behavior of grain growth during welding of stainless steels. Microstructuralevolutions of austenitic stainless steel and the flow stress in semi-solid state are investigated by [2005Yan].Low-cycle fatigue tests have been conducted in air at room temperature on polycrystalline austenitic stain-less steel samples [2005Ye].Iron-based superalloys. The influence of solution treatment temperature on the microstructure, fracture sur-face, composition and amount of precipitates, tensile properties of hot-rolled specimens of Cr-Fe-Ni alloyswere studied by [2003Cai]. It is reported that optimum mechanical properties can be obtained by solution-treatment at 1000°C followed by conventional aging.Ni-based superalloys. Temperature dependence of specific heat capacity of Inconel 600 alloy was investi-gated by [1986Ric]. Caustic stress corrosion cracking behavior of thermally treated and sensitized samplesof Ni-base superalloys have been studied by [2000Kim]. It was found that resistance to stress corrosioncracking of these alloys can be improved by increasing of the grain boundary chromium concentrationand of the grain size, and by formation of intergranular chromium carbide.The temperature dependence of magnetoresistance has been measured by [1991Nat, 1997Sin]. A study ofthe electrical resistance and thermopower of fcc alloys is reported in [1998Tak] and a H-T magnetic phasediagram of the Cr20Fe64Ni16 alloy is presented. Effects of a magnetic field on the athermal and isothermalmartensitic transformations have been examined in [1999Kak]. It is reported that the austenitic state in theCr0.5Fe68.1Ni31.4 alloy is ferromagnetic, whereas that in the Cr4Fe71Ni25 alloy is spin glass. Magnetic phasetransitions have been studied by examining the temperature dependence of average hyperfine field[2000Ban]. This study confirmed the existence of a nonequilibrium ferromagnetic state below the Curietemperature.

4 Cr–Fe–Ni


A general thermodynamic model for calculating the energy of stacking faults is presented and applied toCr-Fe-Ni alloys in [1998Fer]. This model is useful for understanding of the deformation characteristicsand the mechanical behavior of materials. The martensitic transformations in Cr-Fe-Ni alloys from the iso-thermal mode to the burst mode were studied by [2000Rag]. Addition of Cr increases appreciably the driv-ing force atMb. In [2001Jha] the thermal stability of chromised coatings on iron alloys, containing Ni or Cr,has been studied. It was found that the coatings on Fe-5% Ni steel are more stable than such layers on Fe-1% Ni. The grain boundary character distribution of 2 alloy samples was investigated as a function ofannealing time at 360°C by [2003Ale]. Hardness measurements were performed at coincident site latticeand high angle boundaries. Results show that the hardness is increased with the fraction of contiguous coin-cident site lattice boundaries. Phase transformations between phases during gaseous nitriding in cold rolledfoils were studied by [2004Che]. Grain refinement and reorientation take place at 400°C. The reduction oftexture is caused by presence of CrN precipitates.Compositions of typical commercial alloys are indicated in Fig. 14.

Miscellaneous

[1987Mal] observed the formation of α phase in dc sputter deposits of C-Cr-Fe-Ni alloys with carbon con-tent less then 3 mass%. Deposits with larger carbon contents become wholly amorphous. In all crystallinedeposits metastable α phase coexists with the stable equilibrium γ phase. Heating the deposits at 500°Cleads to the formation of M23C6 carbides (in crystalline deposits) or M7C3 carbides (in amorphous deposits).The solidification mode in five undercooled alloys was investigated by [1996Kos]. The primary phase atequilibrium solidification is fcc in three hypoeutectic alloys with lower Cr content and bcc in two hypereu-tectic alloys with higher Cr content. The primary phase to solidify in the undercooled hypoeutectic alloys isbcc. At lower Cr contents the additional fcc phase forms by a solid-state transformation of the metastablebcc phase. CrFeNi alloys were investigated in [1997Moi] by electromagnetic levitation technique. The equi-librium structure of the studied alloys is fcc, but, when undercooled sufficiently, a bcc phase nucleates.At lower undercooling the fcc phase nucleates and at higher undercooling the bcc phase. The critical under-cooling increases with increasing nickel content. Measured growth velocities of fcc and bcc phases arecompared with calculated ones by nucleation theories. Details on calculations are given in [1997Vol1].In [1996Kuw] spinodal decomposition of Cr-Fe-Ni alloys is investigated by detailed analysis of themagnetic-hyperfine-field distribution.The diffusional mobilities of C and N in the α phase are assessed by [1994Joe]. Interdiffusion in differentmultiphase diffusion couples was studied experimentally in [1995Eng]. Literature data on the diffusionalmobilities of Cr, Fe and Ni in bcc CrFeNi alloys are assessed by [1995Joe]. The dissolution kinetics at1100°C of the α phase in the γ phase of the CrFeNi system was studied using a diffusion couple techniqueby [1995Kaj]. [1999Fuk] reviewed works devoted to the problem of metastable formation of the γ phasefrom the δ phase. The crystallographic orientation relationships between δ and γ during unidirectional soli-dification of Cr-Fe-Ni alloys were evaluated in [1999Fuk, 2000Fuk]. The transition from primary γ to δ indissimilar alloy welds by laser experiments was thought to be controlled by nucleation of ferrite [2000Fuk].[2003Nam] determined lattice parameters by X-ray diffraction for CrxFe75Ni25–x with x = 0, 0.05, 0.10 and0.25 and for CrxFe65Ni35–x with x = 0, 0.05, 0.10, 0.15 and 0.20. It was found, that the alloys CrxFe75Ni25–xhave a dominant α phase and a minor γ phase, whereas the CrxFe65Ni35–x alloy system have γ phase. Thesefindings obviously are valid for samples not annealed after preparation by arc-melting. Mössbauer spectrashow that all CrxFe75Ni25–x alloys are magnetically ordered. The magnetic behavior is attributed to the pre-sence of α phase.

Cr–Fe–Ni 5



Table 1. Investigations of the Cr-Fe-Ni Phase Equilibria, Crystal Structures and Thermodynamics

Reference Method / Experimental Technique Temperature / Composition / PhaseRanges Studied

[1927Che] Dilatometry, magnetometry Curie temperatures, Fe-Ni, 0-15 mass%Cr

[1930Jen] “Freezing temperatures”, method not specified Less than 50 mass% Cr, 0-100 mass%Fe, Ni

[1931Wev] Thermal analysis, optical microscopy, X-raydiffraction, dilatometry

Fe-Ni-Cr2Ni-CrFe-Fe

[1937Jen] Thermal analysis, optical microscopy ofannealed alloys

Less than 50 mass% Cr, 0-100% Fe, Ni;800°C to melting temperature

[1939Sch] Micrography of long-time annealed samples 650°C, 800°C, 200 h and 1000 hannealed

[1941Bra] X-ray diffraction Annealed 2d at 1100-1300°C, slowlycooled, filed and powder annealed at800-900°C

[1949Ree] Metallography, X-ray diffraction 650-800°C, 0-50 mass% Cr, 40-80mass% Fe, 0-60 mass% Ni

[1950Pug] Dilatometry, hardness, tensile strength, opticalmicroscopy α + γ field

1100, 800, 400°C, whole compositiontriangle

[1952Coo] Metallography, X-ray diffraction 550, 650, 800°C; 0-35 mass% Cr,0-40 mass% Ni

[1953Tal] Metallography 650-900°C, 15-30 at.% Cr, 20-35 at.%Ni

[1958Lyu] Composition of vapor above liquid phase 1360, 1409, 1464°C; xCr = 0.0495,0.259, 0.372

[1959Pri] Metallography 1300°C, whole α + γ field

[1967Kub] Adiabatic direct reaction calorimetry Integral enthalpy at 1292°C, wholecomposition range, α, γ, α + γ

[1968Art] X-ray diffraction, electric resistance,Youngs modulus

(Fe,Cr)Ni3, 20-800°C

[1969Gil] Knudsen cell effusion, mass spectrometry orchemical analysis of target

1600°C, 5, 10, 20, 30 at.% Cr, from Cr-Fe to Cr-Ni

[1971Yam] DTA, Mössbauer spectroscopy Miscibility gap of α, around 600°C

[1972Sch] Metallography, microprobe 816-1260°C, α + γ field, 0-40 mass%Fe

[1974Ul'y] Magnetometry Martensite formation at –196°C

[1976Vre] Knudsen cell effusion, analysis of target bymicroprobe

1227°C, 11.6 and 28 at.% Fe

(continued)

6 Cr–Fe–Ni



[1977Sch1] Isothermal annealing of liquid + solid two- phasesamples and quenching, optical microscopy,microprobe

Around 1500°C, > 65 at.% Fe

[1977Sch2] Same as [1977Sch1], some Mo or V added Around 1500°C, > 65 at.% Fe

[1979Pro] Knudsen cell mass spectrometry Two alloys: Cr13Fe16Ni71 at 1220 and1620°C; Cr14Fe26Ni60 at 1235 and1635°C

[1980Skr] Knudsen cell effusion, analysis of target bymicroprobe

1227°C, 21.5 at.% Cr, 0-77.3 at.% Ni

[1981Gom] Superstructure formation by neutron diffraction Along section Ni3Fe-Ni2Cr

[1983Mun1] Isothermal annealing, microprobe for phasecompositions, quantitative optical microscopyfor phase amounts

1200 to 1350°C, α + γ two-phase field

[1983Vre] Knudsen cell effusion, analysis of target

[1984Der] Cp by adiabatic calorimetry Fe65Ni35–xCrx (x = 6, 10, 20), 5-90 K

[1986Kun][1988Kun]

Liquid-solid tie lines: directional solidificationinterrupted by quenching, analyzing of bothsides of interface; DTA for temperaturebelonging to liquidus

> 75 mass% Fe, 1477-1538°C

[1987Gom] Ordering parameter by neutron diffraction Along section Ni3Fe-Ni2Cr

[1988Yam] Annealing of liquid-solid two-phase samples andquenching; DTA

> 70 at.% Fe, 1460-1530°C

[1989Cha] Lattice parameter vs temperature Tricritical miscibility gap near FeNi

[1994Sch] Isothermal saturation of melt with Fe andanalyzing

0-22 mass% Cr, 0-22 mass% Ni, 1450-1530°C

[1994Sop] Isothermal long time annealing and analyzing bymicroprobe

ca. 30 mass% Cr + 7.7, 17.6,31.2 mass% Ni; 1000, 900, 795°C

[1995Bro] Knudsen cell mass spectrometry Fe1–x(Ni0.86Cr0.14)x, 0 < x < 1; ca.1400°C (always solid state γ phase)

[1995Mar] Long-time annealing (< 32000 h), electricresistance, SEM, TEM

450-600°C; from Ni2Cr to Ni3Cr,up to 50 mass% Fe

[1996Kuw] Mössbauer spectroscopy Cr30Fe65Ni5 and Cr16Fe81Ni5; 450°C

[1996Sop] Scanning electron microscopy, TEM 795, 880, 900, 1000°C,Fe61.9Cr30.9Ni7.2,Fe51.6Cr31.9Ni16.5,Fe38.6Cr31.9Ni29.5

[1996Vre] Knudsen cell mass spectrometry Fe1–x(Cr1/6Ni5/6)x, 1577°C

(continued)

Cr–Fe–Ni 7




[1998Roe] Levitation drop calorimetry 70 at.% Fe; 14, 15, 18.5 at.% Ni; 1455-1463°C

[1998Var] Mössbauer spectroscopy Cr12Fe84Ni4, 500-650°C

[1998Vre] Knudsen cell mass spectrometry xCr/xNi = 0.46 or 0.19; 0 < xFe < 1;1677°C



Pearson Symbol/ SpaceGroup/Prototype


Comments/References

γ, (Ni1–x–yFexCry)

(Ni)<1453(γFe)1394 - 912

cF4Fm�3mCu

a = 354.7

a = 352.40

a = 364.67

0 < x < 1, 0 < y < 0.5alloy Cr32Ni69, 20°C[1949Ree]at 25°C [Mas2]

at 915°C [Mas2]

α, (Cr1–x–yFexNiy)

(αCr)< 1863(δFe)1538 - 1394(αFe)< 912

cI2Im�3mW

a = 288.48

a = 293.15

a = 286.65

0 < x < 1, 0 < y < 0.32α1 - Cr base solid solutionα2 - Fe base solid solutionat 25°C [Mas2]

[Mas2]

at 25°C [Mas2]

σ, (Fe,Ni)10Cr4(Cr,Fe,Ni)16

tP30P4nmU

a = 879.66c = 455.82a = 879.61c = 456.05

Fe50.5Cr49.5 [V-C2]

Fe52.2Cr47.8 [V-C2]

FeNi3 cP4Pm�3mCu3Au

a = 354.7

a = 356.0

[1968Art] room temperature

at 300°C [1968Art]

CrNi2 oI6ImmmMoPt2

a = (252)a

b = (755)a

c = (356)a

-

acalculated from parameter of disordered phase (γ) neglecting a possible deformation

8 Cr–Fe–Ni




Cr Fe Ni

α + γ ⇌ σ 940 ± 40 p2max σαγ

515526

393750

108

24

σ ⇌ α1 + α2 + γ 475 ± 15 E1 σα1

α2

γ

5294.4158

435.5

7662

5<0.19

30

Table 4. Investigations of the Cr-Fe-Ni Materials Properties


[1986Ric] Adiabatic high temperature calorimetry Specific heat capacity of Cr15.2Fe8.3Ni75.2alloy between 50 and 1000°C

[1987Bol] Dilatometry Thermal expansion of Cr18Fe82–xNix alloyswith x = 45, 50, 60 and 80 mass% Nibetween 140 and 300 K (–133°C-27°C)

[1987Kli] Resonance method Elastic moduli of CrFeNi alloys with 17.5mass% Cr, 50-75 mass% Fe and 6-31 mass%Ni between 4.2 and 300 K

[1991Nat] Magnetoresistance and magnetizationmeasurements

Magnetoresistance of Cr20Fe80–xNix alloyswithx = 21, 23 and 30 at.% Ni between 11 and300 K

[1994Mar] X-ray diffraction, TEM, resistance andmicrohardness measurements

Lattice parameters, electrical resistivity andmicrohardness of CrFeNi alloys with 17-32at.% Cr, 1-51 at.% Fe and 30-67 at.% Nibetween 450 and 600°C

[1995Ino] Gas tungsten arc welding, SEM, EPMA Phase transformations on weld alloyCr9Fe21Ni60

[1996Bob] Directional solidification, opticalmicroscopy, DTA

Solidification path of CrxFe100–x–yNiy mass%Cr 17.6 < x < 21.4, mass% Ni 5.8 < y < 14.4

[1997Sin] X-ray and neutron diffraction Low field magnetoresistance, Cr20Fe80–xNix,at.% 19 < x < 30; from 4 to 90 K

[1998Tak] X-ray, electrical resistivity and thermopowermeasurements, dilatometry

Electrical resistivity, thermopower,Cr20FexNi80–x, at.% 44 < x < 70; from 5 to60 K

[1998Var] Mössbauer spectroscopy, X-ray diffraction Kinetics of α-γ phase transition,Cr12Fe84Ni14, 500, 550, 600, 650°C

(continued)

Cr–Fe–Ni 9




[1999Har] Dilatometry, X-ray and neutron diffraction,optical microscopy

Thermal residual elastic strains,CrxFe100–x–yNiy, mass% 22 < x < 36; 4 < y< 18, 1000°C

[1999Kak] Optical microscopy, X-ray, electricalresistivity measurements, magnetic fieldsusceptibility measurements

Effect of magnetic field on martensitictransformation, 0.5 and 4,0 mass% Cr, 31.4and 25 mass% Ni, from 4.2 to 293 K

[2000Ban] Mössbauer spectroscopy Magnetic properties (hyperfine fielddistribution), Cr20Fe54Ni26, from 13 to 295 K

[2000Kim] Modified Huey and potentiodynamic tests,optical microscopy, SEM

Stress corrosion cracking of alloy 600,alloy 690, Cr10Fe10Ni80

[2001Jha] Optical microscopy, EPMA, SEM, X-ray,microhardness measurements

Cr diffusion-coatings on low-alloy steels(0.014 mass% C, 1 and 5 mass% Ni, 0.6 and4 mass% Si), 830, 1000 and 1120°C

[2003Ale] Nano-hardness measurements, TEM High temperature deformation at 16 mass%Cr, 7.9 and 9.1 mass% Fe, 360°C

[2003Cai] TEM, SEM, EDS analyze, X-ray Mechanical properties of 11.6 mass% Cr,23.4 mass% Ni, 1.1 mass% Mo, 3.1 mass%Ti, 1000, 1050, 1100 and 1150°C

[2003Glo] TEM The influence of hydrogen charging on themicrostructure of 6 mass% Ni, 26.5 mass%Cr, 25°C

[2004Che] X-ray analysis, transmission Mössbauerspectroscopy, TEM

Microstructure of 3 at.% Cr, 4 at.% Ni, from300 to 500°C

[2005Yan] Optical microscopy, EPMA Microstructure evolution, mechanicalproperties (strain-stress) at 18 mass% Cr,9 mass% Ni, 1412 and 1400°C

[2005Ye] Optical microscopy, TEM, atomic forcemicroscopy

Hardness measurements, mechanicalproperties of a 18 mass% Cr, 9 mass% Nialloy at 1040°C

10 Cr–Fe–Ni


Fig.1.

Cr-Fe-Ni.

Reactionscheme

Cr–Fe–Ni 11



Fig. 2. Cr-Fe-Ni. Liquidus surface projection

12 Cr–Fe–Ni


Fig. 3. Cr-Fe-Ni. Solidus surface projection

Cr–Fe–Ni 13



Fig. 4. Cr-Fe-Ni. Isothermal section at 1400°C

14 Cr–Fe–Ni



Cr–Fe–Ni 15




16 Cr–Fe–Ni



Cr–Fe–Ni 17




18 Cr–Fe–Ni



Cr–Fe–Ni 19




20 Cr–Fe–Ni



Cr–Fe–Ni 21



Fig. 12. Cr-Fe-Ni. Vertical section at xCr:xNi = 2.54

22 Cr–Fe–Ni


Fig. 13. Cr-Fe-Ni. Vertical section at 74 at.% Fe

Cr–Fe–Ni 23



Fig. 14. Cr-Fe-Ni. Compositions of typical commercial alloys

24 Cr–Fe–Ni


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Cr–Fe–Ni 25



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[1976Ind] Inden, G., “The Role of Magnetism in the Calculation of Phase Diagrams”, presented at the6th Calphad-meeting, Düsseldorf, June 1976, published in Physica, 103B+C, 82–100(1981) (Calculation, Theory, 31)

[1976Vre] Vrestal, J., Pokorna, A., Rek, A., “Determination of Thermodynamic Activities of Compo-nents in the Nickel Rich Corner of the γ-Solid Solution Fe-Ni-Cr at 1500°C”, KovoveMat., 14, 481–489 (1976) (Experimental, Thermodyn., 6)

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[1980Cha] Chart, T., Putland, F., Dinsdale, A., “Calculated Phase Equilibria for the Cr-Fe-Ni-SiSystem - I Ternary Equilibria”, Calphad, 4(1), 27–46 (1980) (Assessment, Calculation,Phase Diagram, Phase Relations, 75)

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26 Cr–Fe–Ni


[1981Ray1] Raynor, G.V., Rivlin, V.G., “The Cr-Fe-Ni (Chromium-Iron-Nickel) System”, Bull. AlloyPhase Diagrams, 2(1), 89–99 (1981) (Phase Diagram, Phase Relations, Review, 46)

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[1983Mun2] Mundt, R., Hoffmeister, H., “The Isothermal δ-γ Transformation of Ferritic-Austenitic Iron-Chromium-Nickel Alloys”, Arch. Eisenhuettenwes., 54(7), 291–294 (1983) (Experimental,Kinetics, 19)

[1983Vre] Vrestal, J., Rek, A., “Determination of Thermodynamic Activities of Components in theSections xCr=0.11 and xCr=0.33 of Solid Solution Cr-Fe-Ni at 1500 K” (in Czech),Kov. Mater.-Met. Mater., 21(2), 97–104 (1983) (Experimental, Thermodyn., 13)

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[1985Her] Hertzman, S., Sundman, B., “AThermodynamic Analysis of the Fe-Cr-Ni System”, Scand.J. Metall., 14, 94–102 (1985) (Assessment, Calculation, Phase Diagram, Phase Relations,Thermodyn., 47)

[1986Chu] Chuang, Y.-Y., Chang, Y.A., “A Thermodynamic Model for Ternary σ Phase”, Scr. Metal.,20(8), 1115–1118 (1986) (Assessment, Calculation, Phase Diagram, Phase Relations, 7)

[1986Kun] Kundrat, D.M., Elliott, J.F., “Phase Relationships in the Iron Rich Fe-Cr-Ni-C System atSolidification Temperatures”, Metall. Trans. A, 17A, 1461–1469 (1986) (Experimental,Phase Relations, Calculation, 16)

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[1987Bol] Bolshutkin, D.N., Derevyanchenko, M.V., Pecherskaya, V.I., Sirenko, V.A., “ThermalExpansion of Concentrated Fe-Cr-Ni Alloys near the Curie Point”, Phys. Met. Metall., 64(5), 190–192 (1987), translated from Fiz. Met. Metalloved., 64(5), 1029-1031, 1987 (Experi-mental, Mechan. Prop., Phys. Prop., 6)

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[1987Mal] Malavasi, S., Oueldennaoua, A., Foos, M., Frantz, C., “Metastable Amorphous and Crystal-line (α,σ) Phase in Physical Vapor Deposited Fe-(Cr)-Ni-(C) Deposites”, J. Vac. Sci. Tech-nol. A, A5(4), 1888–1891 (1987) (Experimental, Morphology, 13)

[1988Ray] Raynor, G.V., Rivlin, V.G., “Cr-Fe-Ni” in “Phase Equilibria in Iron Ternary Alloys”, Inst.Metals, London, 406, 316–332 (1988) (Review, Assessment, Phase Diagram, 41)

[1988Kun] Kundrat, D.M., Elliott, J.F., “Phase Relationships in the Iron Rich Fe-Cr-Ni System at Soli-dification Temperatures”, Metall. Trans. A, 19, 899–908 (1988) (Experimental, Phase Rela-tions, Calculation, 35)

[1988Yam] Yamada, A., Umeda, T., “Determination of Liquidus and Solidus Surfaces for the Iron richCorner of the Fe-Cr-Ni System and its Application to Calculating Phase Diagram” in “Soli-dif. Proc. 87”, Proc. 3rd Int. Conf., Sheffield, Sept. 21-24 1987, 485–487 (1988) (Experi-mental, Phase Relations, Calculation, 6)

[1989Cha] Chang, Y.A., “Magnetic-Induced Tricritical Point in Alloys and the Low-Temperature Fe-Niand Fe-Ni-Cr Phase Diagrams”, Bull. Alloy Phase Diagrams, 10(5), 513–521 (1989)(Experimental, Magn. Prop., Phase Diagram, Phase Relations, 24)

[1990Hil] Hillert, M., Qiu, C., “A Reassessment of the Cr-Fe-Ni System”, Met. Trans. A, 21(6),1673–1680 (1990) (Assessment, Calculation, Phase Diagram, Phase Relations, 40)

[1990Yam] Yamada, A., Umeda, T., Kimura, Y., “Calculation of Liquidus and Solidus Surfaces of theIron Rich Corner of the Fe-Cr-Ni System”, J. Iron Steel Inst. Jpn., 76(12), 2137–2143(1990) (Calculation, Phase Diagram, Phase Relations, 25)

[1991Nat] Nath, T.K., Majumdar, A.K., “Magnetoresistance in Fe80–xNixCr20 (21≤x≤30) Alloys”,J. Appl. Phys., 70(10), 5828–5830 (1991) (Experimental, Magn. Prop., 6)

[1992Lee] Lee, B.-J., “On the Stability of Cr Carbides”, Calphad 16(2), 121–149 (1992) (Assessment,Thermodyn., Phase Diagram, Phase Relations, 83)

[1993Kaj] Kajhara, M., Kikuchi, M., “Numerical Analysis of Dissolution of α Phase in γ/α/γ DiffusionCouples of the Fe-Cr-Ni System”, Acta Metall. Mat., 41(7), 2045–2059 (1993) (Experimen-tal, Theory, 27)

[1993Lee1] Lee, B.-J., “AThermodynamic Evaluation of the Fe-Cr-Ni System” (in Korean), J. KoreanInst. Met., 31(4), 480–489 (1993) (Assessment, Calculation, Phase Diagram, Phase Rela-tions, 68)

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[1994Joe] Joensson, B., “Ferromagnetic Ordering and Diffusion of Carbon and Nitrogen in bcc Cr-Fe-Ni Alloys.”, Z. Metallkd., 75(7), 498–501 (1994) (Assessment, Kinetics, 34)

[1994Loe] Löser, W., Volkmann, T., Herlach, D.M., “Nucleation and Metastable Phase Formation inUndercooled Fe-Cr-Ni Melts”,Mater. Sci. Eng. A, 178(1-2), 163–166 (1994) (Experimental,Morphology, Kinetics, 13)

[1994Mar] Marucco, A., “Atomic Ordering in the Ni-Cr-Fe System”, Mater. Sci. Eng. A, 189(1-2),267–276 (1994) (Experimental, Crys. Structure, 32)

[1994Rag] Raghavan, V., “Cr-Fe-Ni (Chromium-Iron-Nickel)”, J. Phase Equilib., 15(5), 534–538(1994) (Phase Diagram, Phase Relations, Review, 32)

[1994Sch] Schuermann, E., Durdevic, M., Degen, T., “Melting Equilibria in the Iron rich Corner of theFe-Ni-Cr System”, Steel Research, 65(12), 517–522 (1994) (Experimental, Phase Dia-gram, Phase Relations, 12)

[1994Sop] Sopousek, J., Vrestal, J., “Phase Equilibria in the Fe-Cr-Ni and Fe-Cr-C Systems”,Z. Metallkd., 85(2), 111–115 (1994) (Experimental, Phase Diagram, Phase Relations, 20)

[1995Bro] Broz, P., Vrestal, J., Tomiska, J., “Computer-Aided Thermodyn. of Solid TernaryFe1–x(Ni0.86Cr0.14)x Alloys by Knudsen Cell Mass Spectrometry”, Ber. Bunsen-Ges. Phys.Chem., 99(6), 802–806 (1995) (Experimental, Thermodyn., 11)

28 Cr–Fe–Ni


[1995Eng] Engstroem, A., “Interdiffusion in Multiphase Fe-Cr-Ni Diffusion Couples”, Scand. J. Metall.,24, 12–20 (1995) (Experimental, Interface Phenomena, Morphology, Phase Diagram, PhaseRelations, 15)

[1995Ino] Inoue, H., Koseki, T., Ohkita, S., Tanaka, T., “Effect of Solidification on Subsequent Ferrite-to-Austenite Massive Transformation in an Austenitic Stainless Steel Weld Metal”, ISIJ Int., 35(10), 1248–1257 (1995) (Calculation, Experimental, Morphology, Phase Diagram, Phase Rela-tions, 23)

[1995Joe] Joensson, B., “Assessment of the Mobilities of Cr, Fe and Ni in fcc Cr-Fe-Ni Alloys”,Z. Metallkd., 86, 686–692 (1995) (Assessment, Calculation, Phys. Prop., Review, 46)

[1995Kaj] Kajihara, M., Kikuchi, M., “Analysis of Dissolution of α Phase in γ/α/γ Diffusion Couplesof the Fe-Cr-Ni System using Analytical Solutions for Semi-Infinite Diffusion Couples”,Acta Metall. Mat., 43(2), 807–820 (1995) (Calculation, Experimental, Kinetics, 17)

[1995Kaj] Kajihara, M., Kikuchi, M., “Analysis of Dissolution of α Phase in γ/α/γ Diffusion Couplesof the Fe-Cr-Ni System Using Analytical Solutions for Semi-Infinite Diffusion Couples”,Acta Metall. Mat., 43(2), 807–820 (1995) (Calculation, Experimental, Kinetics, 17)

[1995Mar] Marucco, A., “Phase Transformations During Long-term Aging of Ni-Fe-Cr Alloys in theTemperature Range 450-600°C”, Mat. Sci. Eng. A, A194, 225–233 (1995) (Experimental,Crys. Structure, Kinetics, 26)

[1995Ohs] Ohsasa, K., Nakaue, S., Kudoh, M., Narita, T., “Analysis of Solidification Path of Fe-Cr-NiTernary Alloy”, ISIJ Int., 35(6), 629–636 (1995) (Calculation, Kinetics, 15)

[1995Vit] Vitek, J.M., Vitek, S.A., David, S.A., “Numerical Modeling of Diffusion-controlled PhaseTransformations in the Ternary Systems and Application to the Ferrite/Austenite Transfor-mation in the Fe-Cr-Ni System”, Metall. Mater. Trans. A, 26(8), 2007–2025 (1995) (Calcu-lation, Kinetics, 28)

[1996Bob] Bobadilla, M., Lacaze, J., Lesoult, G., “Effect of Cooling Rate on the Solidification Beha-viour of Austenitic Steels”, Scand. J. Metall., 25, 2–10 (1996) (Calculation, Experimental,Morphology, Kinetics, 30)

[1996Kos] Koseki, T., Flemings, M.C., “Solidification of Undercooled Fe-Cr-Ni Alloys: Part II. Micro-structural Evolution”, Met. Mater. Trans. A., 27(10), 3226–3240 (1996) (Experimental,Kinetics, Morphology, 19)

[1996Kuw] Kuwano, H., Nakamura, Y., Ito, K., Yamada, T., “Spinodal Decomposition of Fe-Cr-NiAlloys Studied by Mössbauer Spectroscopy”, Nuovo Cimento, 18D, 259–262 (1996) (Cal-culation, Experimental, Phase Relations, 5)

[1996Loe] Loeser, W., Volkmann, T., “Metastable Phase Formation in Undercooled Stainless SteelAlloys”, Mater. Sci. Forum, 225–227, 27–32 (1996) (Crys. Structure, Experimental, PhaseRelations, 18)

[1996Sop] Sopousek, J., Kruml, T., “σ Phase Equilibria and Nucleation in Fe-Cr-Ni Alloys at HighTemperature”, Scr. Mater., 35(6), 689–693 (1996) (Experimental, Phase Relations, 18)

[1996Ume] Umeda, T., Okane, T., Kurz, W., “Phase Selection During Solidification of PeritecticAlloys”, Acta Mater., 44(10), 4209–4216 (1996) (Calculation, Morphology, Kinetics, 39)

[1996Vol] Volkmann, T., Löser, W., Herlach, D.M., “Verification of Metastable Phase Diagrams by insitu Measurements on Undercooled Melts”, Int. J. Thermophys., 17(5), 1217–1226 (1996)(Experimental, Kinetics, 11)

[1996Vre] Vrestal, J., Broz, P., Tomiska, J., “Thermodynamic Parameters of Liquid Ternary Fe1–x(Ni5/6Cr1/6)x Alloys”, Monatsh. Chem., 127, 135–142 (1996) (Experimental, Thermodyn., 8)

[1997Moi] Moir, S.A., Herlach, D.M., “Observation of Phase Selection from Dendrite Growth inUndercooled Fe-Ni-Cr Melts”, Acta Mater., 45(7), 2827–2837 (1997) (Experimental,Kinetics, Morphology, 34)

[1997Sin] Sinha, G., Barnard, R.D., Majumdar, A.K., “Low-field Magnetoresistance in Diverse Phasesof γ-Fe80–xNixCr20 Alloys”, Phys. Rev. B, 55(14), 8982–8989 (1997) (Crys. Structure,Experimental, 17)

Cr–Fe–Ni 29



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[1998Fer] Ferreira, P.J., Mullner, P., “A Thermodynamic Model for the Stacking-Fault Energy”, ActaMater., 46(13), 4479–4484 (1998) (Calculation, Theory, Thermodyn., 23)

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[1998Roe] Roesner-Kuhn, M., Matson, D.M., Drewes, K., Thiedemann, U., Kuppermann, G., Flem-ings, M.C., Frohberg, M.G., “Enthalpies and Heat Capacities of Liquid Fe-Cr-Ni Alloyswith the Focus on Pure Liquid Chromium”, Thermochim. Acta, 314, 123–129 (1998)(Experimental, Thermodyn., 43)

[1998Tak] Takzei, G.A., Sych, I.I., Cherepov, S.V., “Transport Properties of Antiferromagnetic fcc-Iron Alloys”, Phys. Solid State, 40(1), 89–93 (1998), translated from Fiz. Tverd. Tela(St. Petersburg), 40, 101-105, (1998) (Experimental, Transport Phenomena, 20)

[1998Thi] Thiedemann, U., Roesner-Kuhn, M., Matson, D.M., Kuppermann, D., Drewes, K.,Flemings, M.C., Frohberg, M.G., “Mixing Enthalpy Measurements in the Liquid TernarySystem Iron-Nickel-Chromium and its Binaries”, Steel Res., 69, 3–7 (1998) (Experimental,Thermodyn., 36)

[1998Tom] Tomiska, J., Vrestal, J., “Computation of Phase Equilibria in the Fe-Ni-Cr System Basedupon Mass Spectrometric Investigations”, Thermochim. Acta, 314, 155–157 (1998) (Calcu-lation, Experimental, Phase Relations, 40)

[1998Var] Varga, I., Kuzmann, E., Vertes, A., “Kinetics of α-γ Phase Transition of Fe-12Cr-4Ni AlloyAged between 500-650°C”, Hyperfine Interact., 112, 169–173 (1998) (Experimental, PhaseRelations, 8)

[1998Vre] Vrestal, J., Theiner, J., Broz, P., Tomiska, J., “Mass-Spectrometric Determination of theThermodynamic Mixing Behavior of Liquid Ternary Fe-Ni-Cr Alloys”, Thermochim. Acta,319(1-2), 193–200 (1998) (Experimental, Thermodyn., 11)

[1999Fuk] Fukumoto, S., Kurz, W., “Solidification Phase and Microstructure Selection Maps for Fe-Cr-Ni Alloys”, ISIJ Int., 39(12), 1270–1279 (1999) (Experimental, Morphology, Kinetics, 52)

[1999Fuk] Fukumoto, S., Kurz, W., “Solidification Phase and Microstructure Selection Maps for Fe-Cr-Ni Alloys”, ISIJ Int., 39(12), 1270–1279 (1999) (Calculation, Experimental, Morphology,Phase Diagram, Phase Relations, 52)

[1999Har] Harjo, S., Tomota, Y., Ono, M., “Measurements of Thermal Residual Elastic Strains in Fer-rite-austenite Fe-Cr-Ni Alloys by Neutron and X-ray Diffractions”, Acta Mater., 47(1),353–362 (1999) (Experimental, Phase Relations, 28)

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[1999Mie] Miettinen, J., “Thermodynamic Reassessment of Fe-Cr-Ni System with Emphasis on theIron Rich Corner”, Calphad, 23(2), 231–248 (1999) (Assessment, Phase Diagram, PhaseRelations, Thermodyn., 39)

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30 Cr–Fe–Ni


[2000Fuk] Fukumoto, S., Okane, T., Umeda, T., Kurz, W., “Crystallographic Relationships Betweenδ-Ferrite and γ-Austenite During Unidirectional Solidification of Fe-Cr-Ni Alloys”, ISIJ Int.,40, 677–684 (2000) (Experimental, Phase Relations, 40)

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[2003Rag] Raghavan, V., “Cr-Fe-Ni (Chromium-Iron-Nickel)”, J. Phase Equilib., 24(3), 261–264(2003) (Review, Crys. Structure, Phase Diagram, Phase Relations, 20)

[2004Che] Chezan, A.R., Craus, C.B., Chechenin, N.G., Vystavel, T., Niesen, L., De Hosson, J.T.M.,Boerma, D.O., “On the Formation of Ultra-Fine Grained Fe-Base Alloys via Phase Transfor-mations”, Mater. Sci. Eng. A, 367(1-2), 176–184 (2004) (Crys. Structure, Experimental,Kinetics, Morphology, Phase Relations, 13)

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Cr–Fe–Ni 31



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32 Cr–Fe–Ni


Chromium – Iron – Oxygen

Pierre Perrot

Introduction

The Cr-Fe-O system is of a great technological interest. It has been intensively investigated with respect tomelting, refining and corrosion of stainless steels and to the refractory properties of oxides. Oxygen, as anon metallic element, exert a strong influence on the production process and service properties. The firstspinel solid solution was prepared by [1912Gro], even if it was not recognized. The main investigationson the ternary system concern the solubility measurements of oxygen in liquid (Fe,Cr) alloys and the oxy-gen pressure measurements at equilibrium between alloy, spinel and corundum solid solutions. The mainexperimental investigations on phase equilibria and thermodynamics are gathered in Table 1. A Calphadassessment which agrees well with the known experimental data below 1900°C was carried out by[1993Tay]. A tentative diagram is drawn at 2300 K (2127°C).

Binary Systems

The well known Cr-Fe system is accepted from [Mas2]. The Fe-O system accepted by [Mas2], mainly fromthe fundamental work of [1945Dar, 1946Dar] has been carefully assessed by [1991Sun, 1995Kow]. The par-tial Cr-O diagram given in [Mas2] does not present Cr3O4 as a stable phase, which contradicts the formationof Cr3O4 observed by [1986Gho] at equilibrium with liquid Cr-Fe alloys containing more than 8 mass% Cr.A Calphad assessment carried out by [1995Kow] showed that Cr3O4 is stable between 1550 and 1707°C, astability domain in agreement with the evaluation of [1993Tay, 1996Deg] which proposed 1650-1705°C.The diagram proposed by [1995Kow] seems more probable, because Cr3O4 was shown to be stable at1600°C [1960Koc, 1986Gho].

Solid Phases

The solid phases are presented in Table 2. Iron chromite FeCr2O4 is a normal spinel which presents twothermal anomalies at 75 and 135 K [1944Sho]. The second anomaly is λ shaped and corresponds to the2nd order transition between a tetrahedrically distorted spinel (below 135 K) and a normal spinel[1964Shi] whereas magnetite Fe3O4 is an inverse spinel. Cr

III may replace FeIII in the spinel structure. How-ever, CrII can not replace FeII in the composition range Fe3O4-FeCr2O4. The spinel solid solution Fe(Cr,Fe)2O4 may be obtained by mixing Fe2O3 and Cr2O3 at 950°C under a (CO2 + 1 % CO) atmospherewhereas the mixing under pure CO2 leads to the corundum (Cr,Fe)2O3 solid solution. The lattice parameterfor the Fe3O4-FeCr2O4 solid solution measured by [1995All] is shown in Fig. 1. It presents two extrema,because the magnetite rich part of the solution is inverse whereas the chromite rich part of the solution isdirect. Other properties of the solid solution follow the same scheme. Above 1600°C, the Fe3O4-FeCr2O4

solid solution seems extend beyond the chromite composition FeCr2O4, up to Fe0.4Cr2.6O4 at 1700°C,which supposes the presence of CrII in the spinel solid solution FeCr2O4-Cr3O4. The spinel solid solution(Fe,Cr)3O4 cannot be stable up to Cr3O4 because the spinel structure of Cr3O4 is unstable. However, a meta-stable solid solution Fe0.6Cr2.4O4 with a distorted spinel structure was obtained [2000Lei, 2001Sch] byplasma spraying of a Cr-Fe alloy. A miscibility gap due to cation ordering in tetrahedral and octahedral sitesis forecast [2003Zie] below 600°C in the Fe(Fe1–xCrx)2O4 solid solution. At 500°C, two phases coexist for0.05 < x < 0.70.The stable α(Fe,Cr)2O3 solid solution exists in the whole composition range. It may be prepared at roomtemperature by high energy milling [1998Mic] or at 1000°C by conventional solid state reaction underair or CO2 atmosphere. A miscibility gap has been forecast by [1978Kau] below 859°C. However, thisresult is quite improbable because it implies a positive excess Gibbs energy of mixing, which contradictsthe calculations of [1979Pel, 1993Tay] and the strong negative Gibbs energy of mixing which is observed

Cr–Fe–O 1



in the spinel solid solution. It is generally accepted that FeIII and CrIII ions in an oxygen lattice presentattractive interactions, which precludes the existence of a miscibility gap. A good agreement is observerbetween the various determination of the crystal parameters of the solid solution [1959Dic, 1982Sch,1995All]. Optical, Mössbauer [1996Mus] and spectrometric measurements [2003Mur] show the existenceof a metastable miscibility gap below 940°C. At 600°C, both phases exist between 35 and 80 mol% Fe2O3.Metastable γ(Fe,Cr)2O3 solid solution, with a defect spinel structure may be prepared by oxidation at 400-500°C of the spinel solid solution FeII(FeIII2–3xCr

III3x)O4 (0.5 < x < 0.67) [1984Gil, 1985Cha]. Their stabi-

lity increases with the Cr content. γ(Fe,Cr)2O3 may also be obtained by oxidizing (Fe,Cr) alloys (< 12 mass% Cr) below 250°C [1986Tha].Wuestite Fe1–xO may dissolve up to 1 mass% Cr at 1000°C [1966Lev, 1968Fuj] and ~4 mass% at 1300-1400°C [1964Kat, 1964Rib]. The maximum Cr content decreases from the iron rich side (Fe0.95O) to theoxygen rich side (Fe0.88O) of the wuestite.

Isothermal Sections

In the solid state above 800°C, the Cr-Fe-O system presents three solid solution: the sesquioxide(Fe1–xCrx)2O3 (0 < x <1), the spinel (Fe1–xCrx)3O4 stable for x < 0.67 and the fcc alloy (Fe1–xCrx)(0 < x <1). The wuestite phase Fe1–xO (0.87 < x < 0.95 dissolves only a limited amount of Cr, up to 2.5mass% at 1100°C and 5 mass% at 1300°C [1972Vik]. Each tie-line in the sesquioxide-spinel two-phasedomain is characterized by an oxygen pressure at equilibrium depending on the temperature. The isothermalsection proposed at 1000°C by [1966Ada] was not taken into account because it contradicts every other dia-grams. At fixed oxygen pressure, the temperature increases with the chromium content. Under pO2 =0.21 bar, Fe2O3 loose oxygen to give Fe3O4 at 1380°C, (Fe0.9Cr0.1)2O3, (Fe0.8Cr0.2)2O3 and (Fe0.7Cr0.3)2O3

loose oxygen to give a spinel solid solution at 1400, 1450 and 1500°C, respectively. At higher chromiumcontent, the oxygen loss, more progressive, is observed between 1500 and 1600°C for (Fe0.6Cr0.4)2O3

and between 1600 and 1700°C for (Fe0.5Cr0.5)2O3 [1954Ric]. The equilibrium between alloy and spinelhas been investigated in [1965Dah] between 1000 and 1600°C for the whole composition range. The Crcontent of the metal phase in equilibrium with FeCr2O4 + Cr2O3 is 1 mass%, 2 mass%, 4 mass% and 6mass% at 1000, 1300, 1500 and 1600°C respectively. At higher Cr content, Cr2O3 is the only oxide in equi-librium with the alloy [1965Dah, 1970Iwa]. The isothermal section at 900°C, investigated by [1969Sch] byequilibrating solutions (Fe1–xCrx)2O3 under various gaseous atmospheres CO-CO2 is shown in Fig. 2. Thevalues (79 in the triangle Fe-wuestite-spinel and 90 in the triangle Fe-Cr2O3-FeCr2O4) represent the COcontent at equilibrium in the corresponding triangles. At 900°C, wuestite Fe0.95O is reduced in Fe undergaseous mixtures containing more than 79 mol% CO; A CO-CO2 atmosphere is never reducing enoughto reduce Cr2O3 in metallic Cr. Below 1000°C, Cr2O3 is the only oxide formed in the oxidation scales ofCr-Fe alloys [1999Bry]. The isothermal section at 1200°C is given in Fig. 3 with log10(pO2). The diagramproposed by [1975Kat, 1975Sne] has been slightly modified to take into account the presence of Cr inthe alloy in equilibrium with Cr2O3 and FeCr2O4. This diagram is in good agreement with the iso-thermal section drawn at 1200°C [1980Lah] from diffusion couple measurements. Figure 4 presents theFeO-Fe2O3-Cr2O3 isothermal section at 1300°C, from [1964Kat], which is in good qualitative agreementwith the isothermal section at 1300°C drawn previously by [1960Sey] and at 1370°C by [1965Dah]. Eachtie line and each three phased field is characterized by an oxygen pressure at equilibrium. Values oflog10(pO2/bar) are given in the figure. The Figs. 5 and 6 present the phase equilibria calculated by[1993Tay] at 1700 and 2027°C respectively.In the liquid state, the oxygen solubility decreases with the addition of Cr up to a minimum which isobserved around 2 mass% Cr in the alloy at 1600°C. The existence of such a minimum is a common obser-vation since [1953Lin, 1959Ave, 1960Koc] and has been explained theoretically by [1968StP, 1974Fel,1998Lee]. Isothermal sections investigated by [1955Hil] at three temperatures near the Cr-Fe side are givenin Fig. 7. At low Cr content, liquid alloy saturated in oxygen is in equilibrium with a liquid oxide FeO hav-ing dissolved a small amount of Cr. In the middle Cr range (between ~2 and 10 mass% Cr), the liquid alloyis in equilibrium with a spinel solid solution. At highest Cr content, liquid alloy is in equilibrium with Cr2O3

having dissolved a small amount of Fe2O3 [1954Tur, 1968StP]. The nature of the crucible in equilibriumwith the liquid Cr-Fe alloys is important, which may explain the controversies about the nature of the oxide

2 Cr–Fe–O


in equilibrium with the liquid alloy. At higher chromium content (> 8 mass%), the liquid alloy saturatedwith O is in equilibrium with Cr3O4 according to [1960Koc, 1986Gho] and with Cr2O3 according to[1991Gel]. However, it is not excluded that Cr2O3 introduced in the cell by [1991Gel] transforms intoCr3O4 when the Cr content of the alloy exceeds 8 mass%. [1973Nak] measured a solubility which is low-ered by the use of a chromite crucible (MgCr2O4 or CaCr2O4), but the existence of a minimum in the oxy-gen solubility is always observed [1986Gho]. A careful investigation [1991Tok] shows that, at 1600°C thespinel solid solution in equilibrium with alloy have a composition from approximately Fe1.2Cr1.8O4 toFe0.7Cr2.3O4 and, at 1700°C, a composition from approximately Fe1.2Cr1.8O4 to Fe0.4Cr2.6O4. In both cases,the interval includes the “chromite” FeCr2O4 composition.


The vertical section Fe3O4-Cr2O3, calculated by [1978Kau] is shown in Fig. 8 whereas the same verticalsection calculated by [1993Tay] is shown in Fig. 9. It has been slightly modified to take into account theaccepted melting point of Cr2O3. The difference between both diagrams lies in the fact that [1978Kau] doesnot take into account the oxygen pressure. The oxygen potential at equilibrium in a two-phase field dependson the temperature. On the other hand, the vertical section calculated by [1993Tay] supposes a potentialoxygen fixed at 21 kPa of oxygen pressure (air atmosphere). This diagram agrees well with the experimentalone drawn by [1960Mua]. The vertical section FeO-Cr2O3, investigated by [1964Rib] under a CO2-H2 (50-50) atmosphere is shown in Fig. 10. The diagram has been modified to take into account the accepted melt-ing points of Cr2O3 and FeCr2O4. The liquidus and solidus temperatures of pure wuestite are raisedfrom ~1380 to 1420°C at the peritectic point Fe1–xO - FeCr2O4 - Liquid.

Thermodynamics

The Gibbs energy of formation of FeCr2O4 from its elements and from its oxides are given in Table 3. Sev-eral expressions are found into the literature [1975Jac, 1978Kau, 1979Pel, 1981Lev]. The proposed expres-sions represent a compromise which match the best with the experimental equilibrium measurements of[1969Sch, 1975Jac, 1987Mar]. The heat content, enthalpies and entropies of CrFe2O4 have been carefullymeasured below room temperature by [1944Nay] and above room temperature up to 1500°C by [1944Sho].Their data, summarized in Table 4 are still accepted in the most recent databases and Calphad assessmentsof the Cr-Fe-O system.The activities of Fe3O4 and FeCr2O4 in the spinel solid solution have been evaluated at 900°C from equili-brium under various CO-CO2 gaseous atmosphere [1969Sch], at 1227°C from equilibrium measurements insolid phase [1975Kat] and at 1400°C [1982Pet] from equilibrium under CO-CO2 atmosphere. The spinelsolid solution presents a strong negative departure towards ideality and may be described with a regularmodel characterized by an excess Gibbs energy:ΔmixG

xs/ J = – 33500 x (1 – x) [1969Sch][1993Tyu] calculates theoretically a semi regular model which gives an excess Gibbs energy of mixing lessnegative than observed one. At 1227°C, [1975Kat] proposes a slight positive departure towards ideality, butthis result seems unlikely. The reason is the use of solid electrolyte cell which covers a wide interval of oxy-gen pressure, but with a lower precision than the use of CO-CO2 atmospheres.The excess Gibbs energy of mixture of the solid solution Fe2O3-Cr2O3 has been evaluated from equilibriumbetween the spinel and corundum solid solution [1979Pel]. The corundum solid solution presents negativedeparture towards ideality and may be described with a regular model:ΔmixG

xs/ J = – 7500 x (1 – x)Pressure-composition diagrams have been drawn at 1200°C by [1996Ina], 1300°C by [1960Sey, 1979Pel,2002Bal], 1600°C by [1991Tok, 2002Dav], 1700°C by [1991Tok]. Figure 11 shows the potential drawnat 1600°C by [2002Dav]. Two-dimensional potential diagrams may be constructed, such as the diagram cal-culated at 1200°C by [1996Ina] and shown in Fig. 12. This kind of diagram is useful for visualizing theoverall features of interfacial reactions.Cr is not a strong deoxidizer of steels and the interaction coefficients of Cr and O in liquid iron at 1600°C,(< 2 mass% Cr) are evaluated by [1967Buz] as εCr

(Cr) = –1.00 (eCr(Cr) = –0.005) and eO

(Cr) = –0.06, by

Cr–Fe–O 3



[1968StP] at εO(Cr) = – 8.8. At higher Cr content (5 mass% Cr) where the liquid alloy is in equilibrium with

a spinel phase, eO(Cr) = – 0.03 [1973Buz]. A more precise expression is proposed by [1973Nak]:

eO(Cr) = 0.151 – 380 / T (eO

(Cr) = – 0.05 at 1600°C, which is coherent with commonly accepted values andwith the value of –0.043 measured later by [1976Jan]). [1982Ska] measures eO

(Cr) = –0.044, –0.039,–0.037 and –0.029 at 1550, 1600, 1650 and 1700°C respectively. More recent determinations [1995Dim]by electromotive force measurements lead to slightly lower values: eO

(Cr) = –0.0546, –0.0486, and–0.0428 at 1550, 1600 and 1650°C respectively. Interaction coefficients in multicomponent systems basedon liquid Fe-Cr-O alloys may be found in [1999Ma]. It must be pointed out that the JSPS (Japan Society forthe Promotion of Science) recommends eO

(Cr) = – 0.055 at 1600°C [2000Ito].The interaction coefficient of Cr and O in (δFe) at 1500°C, is evaluated by [1970Nis] from the distributioncoefficient of oxygen between L and δ phases, as eO

(Cr) = – 0.23 (< 0.5 mass% Cr in the alloy).Cr-Cr2O3 oxygen buffer has been used in solid electrolyte cells to investigate the oxygen solubility in mol-ten steel [1978Apt]. FeCr2O4 rather than Cr2O3 was shown to be the main desoxidation product of Cr-Fealloys.[1993Tay] carried out a realistic thermodynamic assessment using an ionic liquid model, a sublattice modelfor the solid phases, with 4, 3 and 2 sublattices for the spinel, corundum and wuestite phases, respectively.Thermodynamic models were also developed for the oxidation of alloys in air [2004Gon], and for the des-oxidation equilibria in molten steels [2004Jun].


The main experimental works regarding properties are summarized in Table 5. FeCr2O4 and Fe3O4 are anti-ferromagnetic with a Neel temperature of 90 and 119 K respectively [1985Tay]. The magnetic moment ofFeCr2O4 is 0.65 μB measured at 4 K [1964Shi]. The spinel solid solutions Fe3–xCrxO4 (x < 1) undergoes aferrimagnetic-paramagnetic transition measured at 379, 290 and 158°C for x = 0.6, 0.8 and 1.0 respectively[1982Ina]. The ferrimagnetic-paramagnetic transition of magnetite (x = 0) is at 593°C. Thermal expansioncoefficients of FeCr2O4 are reviewed in [1985Tay].Cr-Cr2O3 oxygen buffer has been used in solid electrolyte cells to investigate the oxygen solubility in mol-ten steel [1978Apt].The deviation from stoichiometry δ in (Cr,Fe)3–δO4 was measured at 1200°C and was found lower than0.01. Two defects may be present: cation vacancies which governs iron diffusion at low oxygen activitiesand cations intersticials which governs iron diffusion at high oxygen activities.

Miscellaneous

Monocrystals of (FexCr1–x)2O4 (0.1 < x < 0.4) which are antiferromagnetic with interesting spiral magneticstructures have been prepared by chemical vapor transport [1980Hay] using FeCl3 as transporting agent. Forthese compounds, the transition to the helical magnetic structure lies between 150 and 300 K.Superficial oxidation of Fe0.84Cr0.16 alloy under low oxygen pressure at 600-800°C [1992Lin] leads to anoxide film which is richer in Cr at low oxygen exposure and became increasingly richer in Fe with increas-ing exposure. Cr in the oxidation layer is always in the CrIII state.

Table 1. Investigations of the Cr-Fe-O Phase Relations, Structures and Thermodynamics


[1944Nay] Heat capacity measurement 100-1500°C, FeCr2O4

[1944Sho] Heat capacity measurement 52-298 K, FeCr2O4

[1953Lin] Oxygen solubility in liquid Cr-Fe alloys 1625-1710°C, < 22 mass% Cr, H2-H2Oatmospheres

(continued)

4 Cr–Fe–O



[1954Ric] Equilibrium sesquioxide-spinel measurements 1300-1650°C, (Fe1–xCrx)2O3 (x < 0.5), airoxygen pressure

[1954Tur] Equilibrium alloy-oxides under H2-H2Oatmospheres

1550-1700°C, Cr-Fe liquid alloy + Cr2O3

[1955Hil] XRD, metallography, chemical analysis,potentiometry

1550-1650°C, < 50 at.% Cr, oxygensolubility measurements

[1955Woo] Equilibrium sesquioxide-spinel measurements 500-1600°C, Fe2O3-Cr2O3-FeO, airoxygen pressure

[1960Koc] Oxygen solubility in liquid Cr-Fe alloys 1550-1650°C, < 20 mass% Cr, < 0.2mass% O

[1960Mua] Thermal analysis, X-ray and microscopicexamination

1300-2100°C, Fe2O3-Cr2O3, airatmosphere

[1960Sey] X-ray and electron diffraction, metallographicobservations

1000-1300°C, solid phase equilibria

[1964Kat] Equilibria measurements under CO-CO2 andH2-CO2 atmospheres

1300°C, FeO-Fe2O3-Cr2O3, 0.21 to1.5E–11 bar of oxygen pressure

[1964Rib] X-Ray analysis and microscopic examinations 1350-1500, FeO-Cr2O3 system under H2-CO2 atmospheres

[1965Dah] Equilibrium spinel-alloy measurements 1000-1600°C, Cr-Fe-Fe2O3-Cr2O3

[1966Ada] X-Ray analysis and solid state reactions 1000°C, FeCr2O4-Fe2O3-Cr2O3

[1968Fuj] Solubility or Cr oxides in wuestite measuredunder CO-CO2.

1000°C, Fe0.88O-Fe0.95O, up to 2 mass%Cr

[1968StP] Oxygen solubility in liquid Cr-Fe alloys 1550-1650°C, < 50 mass% Cr, < 0.2mass% O

[1969Fru] Oxygen solubility in liquid (Fe,Cr) alloys 1600°C, < 50 mass% Cr

[1969Sch] Equilibria measurements under CO-CO2

atmospheres900°C, Cr-Fe-Fe2O3-Cr2O3

[1970Iwa] Equilibria measurements under H2-CO2

atmospheres1600°C, Fe-Fe2O3-Cr2O3

[1970Nis] Oxygen solubility in δ(Fe,Cr) alloys 1520-1534°C, < 10 mass% Cr

[1972Vik] X-Ray diffraction, equilibria measurements 790-1100°C, Wuestite-Alloy-Spinelequilibria

[1973Buz] Oxygen solubility in liquid (Fe,Cr) alloys 1600°C, < 10 mass% Cr

[1973Nak] Oxygen solubility in liquid (Fe,Cr) alloys 1606-1823°C, < 30 mass% Cr,equilibrium with (Mg,Ca)Cr2O4

[1975Jac] Electromotive force measurements with ZrO2-CaO solid electrolyte

750-1600°C, Fe-FeCr2O4-Cr2O3

(continued)

Cr–Fe–O 5




[1975Kat] Activities by solid electrolyte cellmeasurements in solid phase

1227°C, Cr-Fe-Fe2O3-Cr2O3

[1975Oud] Electrochemistry 1200°C, Cr-Fe alloys (< 23 mass% Cr)under H2-H2O mixtures

[1975Sne] Isothermal reduction under controlled oxygenfugacities

1000-1200°C, sesquioxide-spinelleequilibria

[1976Jan] Activities by solid electrolyte cellmeasurements in liquid phase

1600°C, < 45 mass% Cr

[1978Apt] Oxygen solubility in Cr-Fe alloys by solidelectrolyte cell

1600°C, < 6 mass% Cr

[1980Lah] Diffusion couple, microscopy, electron probemicroanalysis

1200°C, (Fe,Cr)-Fe2O3-Cr2O3 isothermalsection

[1982Ina] Heat capacity and magnetic transitionmeasurements

200-850 K, Fe3–xCrxO4

(x = 0.6, 0.8, 1)

[1982Pet] Magnetite activity measurement byequilibrium with CO-CO2 mixtures

1200-1400°C, Fe3O4-FeCr2O4

[1982Ska] Oxygen activity measurements in liquid Cr-Fealloys

1550-1700°C, Cr-Fe (< 7 mass% Cr)

[1991Gel] Chromium activity measured by solidelectrolyte cell

1600°C, Cr-Fe alloys in equilibrium withCr2O3

[1991Tok] Spinel-alloy equilibrium measurements H2-CO2 mixtures

1600-1825°C, Cr rich part of the systemCr-Fe-Fe3O4-Cr3O4

[1995Dim] Oxygen activity measurements in liquid Cr-Fealloys

1550-1650°C, < 50 at.% Cr

[2000Ito] Oxygen activity measurements in liquid Cr-Fealloys

1550-1650°C, < 50 at.% Cr

[2002Hin,2002Kim]

Oxygen activity measurements in liquid Cr-Fealloys

1550-1650°C, < 8 at.% Cr

6 Cr–Fe–O






Comments/References

α, (Fe1–xCrx)(δFe)1538 - 1394(αFe)< 912(Cr)< 1863

cI2Im�3mW

a = 293.15

a = 286.65

a = 288.48

0 < x < 1pure Fe at 1360°C [V-C2]

pure Fe at 20°C [Mas2, V-C2](A2 structure)pure Cr at 25°C [Mas2]

(γFe)1394 - 912

cF4Fm�3mCu

a = 293.16 at 915°C [Mas2, V-C2]

σCrFe830 - 440

tP30P42/mnmσCrFe

a = 879.4c = 455.2

44.5-50 at.% Cr[Mas2, 1989Rag]

Fe1–xO (Wuestite)1422 - 569

cF8Fm�3mNaCl

a = 431.0a = 429.3

0.05 < x < 0.12 [1991Sun]x = 0.05x = 0.12

Fe3O4 I< 580

oP56PbcmFe3O4 I

a = 1186.8b = 1185.1c = 1675.2

[V-C2]

SpinelFe3O4 (h)(Magnetite)1597 - 580FeCr2O4

(“Chromite”)< 1757°C

cF56Fd�3mMgAl2O4

(Spinel)

a = 839.6a = 854.5

a = 837.8

at 25°Cat 1000°C [V-C2]inverse spinelat 25°C [1995All]direct spinel

CorundumαFe2O3 (Hematite)< 1451

αCr2O3

< 2330

hR30R�3cαAl2O3 (Corundum)

a = 503.42c = 1374.83a = 499.4 ± 0.2c = 1361.4 ± 0.2a = 500.1c = 1361.7a = 496.07c = 1359.9

melts at 1892°C under O2 pressure

α(Fe0.835Cr0.165)2O3

[1995All]

α(Fe0.4Cr0.6)2O3 [1982Sch]Cr2O3 [1989Rag]

βFe2O3 cI80Ia�3Mn2O3

a = 939.3 metastable phase[V-C2]

γFe2O3 (maghemite) cF56Fd�3mMgAl2O4

a = 834

a = 826.0

metastable phase[1989Rag]γ(Fe0.5Cr1.5)O3 [1985Cha]

(continued)

Cr–Fe–O 7






Comments/References

Cr3O4

1707 - 1550tI28I41/amdMn3O4

a = 620.26c = 753.86

[1989Rag, 1995Kow]distorted spinel

CrO2 tP6P42/mnmTiO2 (rutile)

a = 441.90c = 291.54

[1989Rag]

CrO3

< 470oC16Ama2

a = 574.3b = 855.7c = 478.9

[1989Rag]decomposes under air at 470°C



Quantity, per mole of atoms[J, mol, K]

Comments

2 Fe + 2 Cr2O3 + O2 ⇌2 FeCr2O4

750-15361536-170013001400

ΔrG° = – 633400 + 145.5 TΔrG° = – 661000 + 160.5 TΔrG° = – 405080ΔrG° = – 390560

[1987Mar][1975Jac][1982Pet][1982Pet]

Fe0.95O + 0.05 Fe + Cr2O3 ⇌ FeCr2O4 750-130025-1500

ΔrG° = – 52100 + 7.24 TΔrG° = – 51630 + 8.013 T

[1975Jac]accepted

Cr2O3 + Fe3O4 ⇌ FeCr2O4 + Fe2O3 1300 ΔrG° = – 5210 [1982Pet]

Cr2O3 ⇌ 2 {Cr}Fe +3 {O}Fe

1500-1700 ΔrG° = 1022700 – 438.80 T(Standard state: 1 mass%)ΔrG° = 843150 – 371.79 T(Standard state: 1 mass%)

< 5 mass% Cr[2000Ito]> 5 mass% Cr[2000Ito]

Table 4. Thermodynamic Properties of Single Phases

Phase Temperature Range[°C]

Property, per mole of atoms[J, mol, K]

Comments

(1/7)(FeCr2O4) 25-1500 Cp = 23.293 + 3.193E–3 T– 4.56E5 T – 2

[1944Nay], accepted value

(1/7)(FeCr2O4) 25 S° = 20.86S° = 20.846

[1944Sho, 1988Bag]accepted value [1998Cha]

(1/7)(FeCr2O4) 25 ΔfH° = – 232000ΔfH° = – 207250ΔfH° = – 206000

[1984Kaz], calculatedaccepted value[1998Cha]

(1/7)(FeCr2O4) 251700

ΔfG° = – 192800ΔfG° = – 113700

accepted value[1991Tok]

8 Cr–Fe–O


Table 5. Investigations of the Cr-Fe-O Materials Properties

Reference Method / Experimental Technique Type of Property/Conditions

[1956Lot] Magnetization, magnetic susceptibilitymeasurements

FeCr2O4, 0-1000 K

[1959Dic] X-ray diffraction Crystal parameter of the solid solution(FexCr1–x)2O3 (0 < x < 1)

[1964Shi] Magnetic moment measurements FeCr2O4, 0-100 K.

[1966Lev] X-ray diffraction 1000°C, crystal parameter of Fe1–xO doped withCr (< 1 mass% Cr)

[1968Mor] Gravimetry, metallography 650-1000°C, Cr and Fe50Cr50

[1978Aki] Oxidation kinetics Fe-Cr alloys, 20°C

[1980Hay] Chemical vapor transport andmagnetic measurements

(FexCr1–x)2O3, 0.1 < x < 0.4, 770-1040°C

[1982Sch] X-Ray diffraction Crystal parameter of the solid solution(FexCr1–x)2O3 (0 < x < 1)

[1984Gil] Infra-red spectrometry Vacancy ordering in spinel and γ(Fe,Cr)2O3 solidsolutions

[1985Cha] X-ray diffraction, magneticmeasurements

Preparation of γ(Fe,Cr)2O3 solid solutions withspinel structure

[1986Tha] Raman spectroscopy, ESCA,chronopotentiometry

Cr-Fe alloys (< 12 mass% Cr) oxidized at 250-260°C

[1992Lin] XPS photoelectron spectroscopy 600-800°C, surface oxidation of Fe0.84Cr0.16 alloy

[1995All] X-ray diffraction Crystal parameters of the spinel and corundumsolid solutions

[1995Top] Thermogravimetry, diffusionmeasurements

1200°C, δ, deviation from stoichiometry on(Fe,Cr)3–δO4

[1996Mus] Fourier transform IR, Mössbauer,optical spectroscopy

< 1200°C, (Fe,Cr)2O3 solid solution

[1998Mic] X-ray diffraction, EDX, EDS, TEMobservations

Preparation of α(Fe,Cr)2O3 solid solutions bymechanical alloying

[1999Bry] Thermogravimetry under H2-H2Oatmospheres

750-950°C, oxide scale development on Cr-Fealloys

[1999Mor] X-ray diffraction, coercivity,permeability, Magnetization

Cr-Fe-O thin films cosputtered using Fe and Cr2O3

dual targets

[2000Lei] X-ray diffraction, X-ray fluorescence,Mössbauer

Cr-Fe-O particles obtained by plasma spraying of aFe-13 mass% Cr alloy

[2001Liu] Magnetoresistance Cr-Fe-O thin films cosputtered using Fe and Cr2O3

dual targets

(continued)

Cr–Fe–O 9



Reference Method / Experimental Technique Type of Property/Conditions

[2001Sch] X-ray diffraction, Mössbauer Cr-Fe-O particles obtained by plasma spraying of aCr-Fe alloy

[2002Ast] Auger electron spectroscopy Analysis of a (Fe,Cr)2O3 oxide scale on anoxidized stainless steel

[2002Muk] Sessile drop technique, surfacetension, wettability

1550°C, low oxygen activity, 16 mass% Cr

[2003Mik] Microstructure, oxidation kinetics 900°C, Ar-H2-H2O atmospheres

[2003Mur] Fourier transform IR, X-ray analysis 600-940°C, metastable miscibility gap in thecorundum solid solution

[2003Zie] Cation disordering calculations < 600°C, metastable miscibility gap in the spinelsolid solution

[2005Lee] Surface tension measurements 1550°C, Cr-Fe-O, < 30 % Cr

Fig. 1. Cr-Fe-O. Lattice parameter a of the spinel solid solution

10 Cr–Fe–O


Fig. 2. Cr-Fe-O. Isothermal section at 900°C

Cr–Fe–O 11



Fig. 3. Cr-Fe-O. Isothermal section at 1200°C (values give log10(pO2/bar) at equilibrium)

12 Cr–Fe–O


Fig. 4. Cr-Fe-O. Isothermal section of the FeO - Fe2O3 - Cr2O3 subsystem at 1300°C (values give log10 (pO2/bar))

Cr–Fe–O 13



Fig. 5. Cr-Fe-O. Calculated isothermal section at 1700°C

14 Cr–Fe–O


Fig. 6. Cr-Fe-O. Calculated isothermal section at 2027°C

Cr–Fe–O 15



Fig. 7. Cr-Fe-O. Solubility of oxygen in liquid (Fe,Cr) alloys at 1550, 1600 and 1650°C

16 Cr–Fe–O


Fig. 8. Cr-Fe-O. The Fe3O4-Cr2O3 vertical section under a variable oxygen pressure

Cr–Fe–O 17



Fig. 9. Cr-Fe-O. Phase relations in the system iron oxides - Cr2O3 under air atmosphere

18 Cr–Fe–O


Fig. 10. Cr-Fe-O. The FeO-Cr2O3 vertical section under CO-H2 (50-50) atmosphere

Cr–Fe–O 19



Fig. 11. Cr-Fe-O. Pressure-composition diagram at 1600°C

20 Cr–Fe–O


Fig. 12. Cr-Fe-O. Two-dimensional potential diagram at 1200°C

Cr–Fe–O 21



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22 Cr–Fe–O


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Cr–Fe–O 23



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24 Cr–Fe–O


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Cr–Fe–O 25



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26 Cr–Fe–O


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Cr–Fe–O 27



Chromium – Iron – Phosphorus

Kostyantyn Korniyenko, Hans Leo Lukas

Introduction

Phase relations in the Cr-Fe-P system are of great interest because of the importance of chromium and phos-phorus as alloying elements in corrosion-resistant amorphous alloys. They are also important for enhancingthe sintering behavior of 316L stainless steel and improving the magnetic iron-phosphorus alloys. But up tonow, information regarding phase equilibria is quite sparse and that available was obtained long ago.[1939Vog] reported liquidus and solidus projections for the composition range 0-30 mass% P, a series oftemperature-composition sections as well as a partial isothermal section at 1140°C in the region adjacentto the Fe rich Cr-Fe binary edge of the system. A partial isothermal section at 800°C was later presentedby [1965Kan1]. Experimentally determined crystal structure data were published by [1948Now, 1969Fru,1969Rog, 1973Mae, 1977Got, 1978Sel, 1981Maa, 1991Myu, 1999Kum, 1999Lit, 2004Kum]. Thermody-namic properties of the Cr-Fe-P system have been reported by [1970Had, 1979Yam, 1983Yam, 1984Ban,1993Din, 1997Zai, 1998Mie, 1998Zai1, 1998Zai2, 1999Mie, 2005Sho]. The experimental techniques usedalong with the temperature and composition ranges studied are listed in Table 1. Reviews of the literaturedata include information concerning phase equilibria and crystal structures [1949Jae, 1988Rag], classifica-tion of phase correlations in related systems [1965Kan1, 1965Kan2], thermodynamics [1974Sig, 1979Yam,1984Ban] and crystal structure types [1948Now].Further study of the phase equilibria is desirable, in particular the constitution of temperature-compositionsections for the phosphorus contents of 50 at.% and higher, as well phase relations for different temperaturesin the homogenized state.

Binary Systems

The Cr-Fe and Cr-P systems are accepted from [Mas2].The Fe-P system is accepted from [2002Per].

Solid Phases

Crystallographic data for the known unary and binary phases as well as for the single ternary phase (τ) arecompiled in Table 2. Three phases are reported to form continuous series of solid solutions between iso-structural binary phases existing over the whole composition range and wide temperature ranges, namelythe M3P phase [1939Vog] (Ni3P type structure), the M2P phase [1939Vog] (Fe2P type) and the MP phase[1978Sel] (MnP type). In pure FeP [1972Sel] determined that the z-value of the position of the P atomsis 0.296 and thus considerably shifted from the value 0.25 demanded by the mirror plane in space groupPnma, Thus the mirror plane disappears and the space group transforms to the subgroup Pn21a with thesame symmetry elements except for this mirror plane. The remaining coordinates of the atomic positionsare not changed and thus the finding of [1972Sel] does not contradict the presence of a continuous solidsolution between CrP and FeP. The ternary τ phase was found in alloys along the Cr2P-Fe2P section con-taining between 10 and 36.67 at.% Cr [1969Fru, 1969Rog]. The crystal structure type of the τ phase differsfrom both the βCr2P and the Fe2P (I) phases, but it is identical to that of the high-pressure binary λ, Fe2P(II) phase.

Quasibinary Systems

The phase diagram of the Cr2P-Fe2P quasibinary system is presented in Fig. 1, based mainly on the data of[1939Vog]. The temperature of the transformation between the high- and low-temperature modifications ofthe Cr2P compound is not yet well known. [1948Now, 1969Fru, 1969Rog] prepared alloy specimens of this

Cr–Fe–P 1



quasibinary system and the structure investigations clearly verify the continuous series of solid solutions ofthe M2P phase.


The invariant reactions of the binary Cr-P system between the liquid and the phases (Cr), Cr3P, Cr2P, CrPform monovariant three-phase equilibria in the ternary decreasing smoothly to the corresponding invariantequilibria of the binary Fe-P system [1939Vog]. No invariant equilibrium was found in the ternary system inthe region Cr-CrP-FeP-Fe.


Figure 2 shows the partial liquidus and solidus surface projections in the composition range up to about40 at.% P. It is derived from the data of [1939Vog], the compositions are transformed from mass% to at.%. Some small corrections were necessary to ensure agreement with the accepted binary systems.

Isothermal Sections

A partial isothermal section for 1140°C in the Fe rich corner is presented in Fig. 3, showing the α + γ field.A correction was made to the original diagram of [1939Vog] to make it consistent with the accepted Cr-Fephase diagram: the α + γ two-phase field does not close when approaching the binary Cr-Fe edge. An iso-thermal section for 800°C in the range Cr-Cr2P-Fe2P-Fe was plotted by [1965Kan1]. They assumed that atthis temperature, a miscibility gap appears in the M3P phase. [1965Kan2] measured the P solubility in the αphase from 0 to 18 mass% (19.08 at.%) Cr at 700, 800, 900 and 1000°C. The solubility decreases sharplywith increasing Cr content.


[1939Vog] constructed eight temperature-composition sections at constant Cr/Fe ratios and one from Fe2Pto Cr and one from Cr3P to Fe3P. Three of these sections are drawn in Figs. 4 to 6, calculated from the data-set of [1999Mie]. They deviate to some extent from those drawn by [1939Vog], but this deviation seems tobe in the same magnitude as the uncertainty of the system. Calculated by the dataset of [1999Mie] the qua-sibinary two-phase field L+(Cr,Fe)2P shows a minimum near the Fe2P side, contrary to the diagram of[1939Vog], shown in Fig. 1. The three-phase field L+(Cr,Fe)2P+(Cr,Fe)3P shows a minimum in both dia-grams, of [1939Vog] and [1999Mie], only the curvatures of some lines in Vogel’s diagram are very unlikely.The experimental temperatures of the three-phase fields given by [1939Vog] may suffer from segregationduring primary crystallization and be too low.

Thermodynamics

[1970Had] related interaction coefficients applicable to conditions of constant chemical potential of one of thesolute elements to the usual interaction coefficients at infinite dilution. A statistical technique was applied tothe data of the Cr-Fe-P system. The analysis indicates that a formal relationship can be used where the chemi-cal potential of one component is fixed. [1974Sig] collected and assessed literature data on the thermodynamicbehavior of dilute solute elements (including chromium and phosphorus) in liquid iron as the solvent. Theactivity interaction parameters between alloying elements (in particular, chromium) and phosphorus in liquidiron at 1873 K (1600°C) were determined by [1979Yam, 1983Yam] using a mass spectrometry technique, andlater - by [1993Din] using thermodynamic calculation. The calculated values agree satisfactorily with theexperimental data. The vapor pressure of phosphorus above liquid Cr-Fe-P alloys has been measured by[1984Ban] at 1673 K (1400°C). The activities of the components of the Cr-Fe-P melt were determined by[1997Zai, 1998Zai1, 1998Zai2] over wide temperature and concentration ranges. The associated-solutionmodel was used to represent the thermodynamic properties of the melt, assuming that several types of ternarycomplexes can exist in the liquid solution, along with all binary associates of the subsystems Fe-P and Cr-P.The types of ternary associative complexes and their thermodynamic functions were determined by treatment

2 Cr–Fe–P


of the complete file of experimental data by an optimizing procedure. It was shown that only one ternary com-plex (CrFeP), should be taken into account. The thermodynamic functions of the binary complexes werefound to coincide with those established from the data for the Cr-P and Fe-P systems. Extrapolation of theexperimental data to the binary Cr-Fe side resulted in good agreement with the thermodynamic informationavailable and with the phase diagram. Based on the data of [1997Zai, 1998Zai1, 1998Zai2] and their ownexperimental results, [2005Sho] obtained an empirical expression for the solubility of phosphorus in the Cr-Fe melt given as a function of temperature and phosphorus activity.Thermodynamic optimization of the Cr-Fe-P system has been carried out by [1998Mie, 1999Mie] usingclassical thermodynamic models to describe the properties of the individual phases of the system. The ther-modynamic parameters of the Cr-Fe and Fe-P binaries were taken from earlier assessments and the para-meters of the Cr-P and Cr-Fe-P systems were optimized by using experimental data on activity andGibbs energy [1997Zai, 1998Zai1, 1998Zai2] as well as experimental phase equilibrium data [1939Vog,1965Kan1, 1965Kan2]. The description is valid for P contents of up to 33 at.%. The optimization resultedin good agreement between calculations and experimental data.


A series of corrosion-resistant iron-based amorphous alloys contain both chromium as a film former and phos-phorus. The FeP-, Fe2P- and Fe3P-based alloys with chromium additions possess specific magnetic proper-ties. Literature data concerning investigations of the material properties of chromium-iron-phosphorusalloys are listed in Table 3. Among them, information about magnetic characteristics is presented from[1969Fru, 1969Rog, 1977Got, 1978Sel, 1999Kum, 1999Lit, 2004Kum]. In particular, [1978Sel] found nosigns of ferri- or ferromagnetic impurities in the alloys with 50 at.% P. However, the inverse magnetic suscept-ibility χ–1(T ) curves for these alloys are nonlinear and convex towards the temperature axis. [1999Kum]reported that the (Cr0.03Fe0.97)2P alloy is paramagnetic at temperatures above ~ 180 K with persistent shortrange ferromagnetic order. At lower temperatures three different regions of magnetic behavior have been iden-tified. [1999Lit] determined that magnetico-rheological suspensions (MRS) based on Cr6Fe82P12 amorphousribbons possess high magneto-rheological behavior. According to [2004Kum], the nature of the magnetiza-tion-temperature curve for this CrFeP alloy is suggestive of antiferromagnetic nature with a Neel temperature(TN) of 280 ± 10 K. The passivation behavior of phosphorus-implanted alloys in a Cl-containing electrolyte asa function of Cr concentration and P implantation has been investigated by [1986Sor]. Amorphous and nano-crystalline Cr-Fe-P electrodeposits were studied by [1998Kun, 1999Kun] and [2000Kun, 2004Kun], respec-tively. Their electrochemical measurements revealed an increased corrosion resistance in substrates coatedwith these deposits. [1999Pre] found that additions of Fe3P led to improved sintering densities in 316L stain-less steel and to a reduction in the total and interconnected porosity. Hardness increases with increasing con-tent of Fe3P. A tensile strength of around 300 MPa was achieved in 316L steel sintered with 8, 10 and12 mass% Fe3P, which is comparable with traditionally sintered 316L, but much lower than expected for afully dense 316L stainless steel. The cause was assumed to be the presence of brittle grain boundary eutecticmaterial which severely reduces the ductility of the alloys.

Miscellaneous

The grain-boundary segregation of phosphorus and chromium has been studied by [1981Erh] for aCr2.2Fe97.752P0.048 (mass%, in at.% - Cr2.358Fe97.555P0.087) alloy annealed at 500°C. The authors establishedthat both P and Cr are enriched at the grain boundaries of this alloy. The grain-boundary concentration ofchromium is only slightly higher than the bulk concentration. The segregation of P and Cr to grain bound-aries in Cr-Fe-P ferritic alloys was studied by [1981Mat] in relation to the quenching temperature and agingconditions using Auger electron spectroscopy (AES). The heat of phosphorus segregation increases from14.5 to 71 kJ·mol–1 with an increase in the bulk chromium concentration from 0 to 5 at.%. It was estab-lished that the equilibrium grain boundary segregation of Cr is slightly increased by the presence of P.The heat of chromium segregation is small. The grain boundary concentration of Cr decreases rapidly

Cr–Fe–P 3



(negative segregation) by short time aging, but it increases by prolonged aging (normal segregation). Morerecently, [1986Yuq] studied surface segregation in a ferritic alloy containing 1.04 mass% Cr and 0.034 mass% P as well as small amounts (0.005 mass%) of carbon, sulfur and nitrogen. The AES investigations did notshow evidence of Guttmann-type cosegregation of Cr + P, however, evidence was found for Cr + N andCr + S cosegregation. These observations support the previous contention that the Cr + P interaction isunimportant for the temper embrittlement of low-alloy steels. The role of chromium in the intergranularfracture of high purity Fe-0.2 mass% P-(1 or 2) mass% Cr alloys containing also small amounts of carbonwas studied by [1994Liu] using small-scale Charpy impact tests and SEM on fracture surfaces. ScanningAES, tensile tests at low temperatures and optical microscopy were performed to clarify the mechanismof the effect of chromium on intergranular fracture. The experimental results show that the addition of chro-mium to Fe-P and C-Fe-P alloys within its solubility limit has no effect on the segregation of phosphorous,but reduces the susceptibility of these alloys to intergranular fracture. The effect of chromium is attributed toan increased grain boundary cohesion caused by chromium segregated to the grain boundaries and/or toreduced deleterious effect of phosphorus, resulting from the attractive interaction between segregated chro-mium and phosphorus.

Table 1. Investigations of the Cr-Fe-P Phase Relations, Structures and Thermodynamics


[1939Vog] X-ray diffraction, thermal analysis,optical microscopy

< 1600°C, 0-30 mass% P

[1948Now] X-ray diffraction The Cr3P-Fe3P section

[1965Kan1] X-ray diffraction, chemical analysis 800°C; ≤ 33.3 at.% P, ≤ 39.1 at.% Cr

[1965Kan2] X-ray diffraction, chemical analysis ≤ 1000°C

[1969Fru] X-ray diffraction (Seeman-Bohlincamera), Mössbauer spectroscopy

The Cr2P-Fe2P section

[1969Rog] X-ray diffraction (Seeman-Bohlincamera), Mössbauer spectroscopy

The Cr2P-Fe2P section

[1973Mae] X-ray powder diffraction, Mössbauerspectroscopy

900°C; CrFeP

[1977Got] X-ray diffraction The Cr3P-Fe3P section

[1978Sel] X-ray diffraction The CrP-FeP section

[1979Yam] Knudsen-cell mass-spectrometry 1600°C

[1981Maa] X-ray powder diffraction, Mössbauerspectroscopy

Cr0.8Fe1.2P

[1983Yam] Knudsen cell-mass spectrometry 1600°C

[1984Ban] Transportation method 1400°C

[1991Myu] X-ray diffraction, differential scanningcalorimetry (DSC)

17 at.% P, 7 to 10 at.% Cr

[1997Zai] Knudsen effusion technique 1403-1826 K (1130-1553°C); ≤ 31.8 at.% P,≤ 79.8 at.% Cr

(continued)

4 Cr–Fe–P



[1998Zai1][1998Zai2]

Knudsen effusion technique 1403-1821 K (1130-1548°C); ≤ 31.8 at.% P,≤ 79.8 at.% Cr

[1999Kum] X-ray diffraction (Cr0.03Fe0.97)2P

[1999Lit] X-ray diffraction, DTA Cr1Fe79P20, Cr3Fe77P20, Cr7Fe73P20, Cr6Fe82P12

[2004Kum] X-ray diffraction and neutrondiffractiontechniques

CrFeP

[2005Sho] Knudsen effusion technique 1400-1800 K (1127-1527°C); ≤ 50 at.% P





Comments/References

α, (Cr1–xFex)

(Cr)< 1863(δFe) (h2)1538 - 1394(αFe) (r)< 912

cI2Im�3mW

a = 288.4

a = 293.15

a = 286.64

0 ≤ x ≤ 1, dissolves 5 at.% P at x = 1,1048°C [1990Oka, Mas2]

pure Cr, at 27°C[V-C2]pure Fe, at 1360°C[V-C2]pure Fe, at 20°C [V-C2]

γ, (γFe) (h1)1394 - 912

cF4Fm�3mCu

a = 364.68 pure Fe, at 912°C[V-C2]dissolves 11.9 at.% Cr at ~1100°C[Mas2]dissolves 0.56 at.% P at ~1150°C[1962Lor, 1990Oka, Mas2]

(εFe)> 1.3·105 bar

hP2P63/mmcMg

a = 246.8c = 396

at 25°C [Mas2]high pressure phase

(P) (red)< 417

c*66 a = 1131 sublimation at 1 bar. Stable form of P.Triple point at 576°C, > 36.3 bar; triplepoint at 589.6°C at 1 atm [Mas2, V-C2]

(P) (white)< 44.14

c**-P (white)

a = 718 common form of P [Mas2, V-C2]

(P) (black) oC8CmcaP (black)

a = 331.36b = 1047.8c = 437.63


(continued)

Cr–Fe–P 5






Comments/References

σ, CrFe830 - 440

tP30P42/mnmCrFe

a = 879.95c = 454.42

50.0 to 55.5 at.% Fe[Mas2][V-C2]

M3P,(CrxFe1–x)3P

Cr3P< 1510Fe3P< 1166

tI32I�4Ni3P

a = 919c = 456a = 911c = 445.5a = 910c = 451.5

x = 0 to 1 [1977Got]

at x = 1, T = 800°C [1977Got]

at x = 0, T = 800°C [1977Got]

at x = 0.5,T = 800°C [1977Got]

M2P,(CrxFe1–x)2P

βCr2P (h)≲ 1640

Fe2P (I)< 1370

hP9P�62mFe2P

a = 598c = 344a = 584.0c = 348.5a = 583.6c = 345.3a = 593.0c = 345.3a = 586.7c = 345.6a = 586.8c = 346.0

x = 0 to 1 [1969Rog]

~ 33 at.% P [Mas2]at x = 1 [1948Now]at x = 0.1 [1969Rog]

at x = 0.03, T = 27°C[1999Kum]33.3 to 34 at.% P [1990Oka, Mas2][1948Now]at 33.3 at.% P [1969Fru]

[1969Rog]

αCr2P (r)< 1640 (?)

oP18PmmmCr2P (r)

a = 633.2b = 1033.9c = 329.9

~ 33 at.% P [Mas2][1988Rag]

γ, Cr12P7 (orCr1.7P)

hP20P63/mTh7S12

a = 896.6c = 331.1

~ 37 at.% P [Mas2]single crystals [1981Maa]

(continued)

6 Cr–Fe–P





Comments/References

MP, (Cr1–xFex)P

CrP≲ 1360

FeP≲ 1300

oP8PnmaMnP

a = 536.2b = 311.3c = 601.8a = 535 to 519b = 310 to 309c = 600 to 580a = 519.3b = 309.9c = 579.2

0≤ x≤ 1, 50 at.% P [1939Vog, 1978Sel]

decomposes at ~1360°C, according tocrude estimation [1962Rip][1988Rag]0 to 50 at.% Fe, room temperature[1978Sel]

[1972Sel], symmetry diminished tospace group Pn21a by slight shift of theP atoms

ε, Cr2P3 (?) - - 60 at.% P (?) [Mas2]

η, CrP2 mC12C2/mOsGe2

a = 821.3b = 303.4c = 709.8β = 119.47°

~ 67 at.% P [Mas2][1988Rag]

κ, CrP4 mC20C2/cMoP4

a = 519.14b = 1076.00c = 577.12β = 110.648°

~ 80 at.% P [Mas2]prepared by reaction of Cr powder andred P in boron nitride crucibles between900 and 1200°C, at pressures of 15 to65 kbar [1972Jei]

λ, Fe2P (II) oP12PnmaCo2Si a = 519.3 to 524

b = 310 to 320.6c = 579 to 584a = 520 to 529b = 310 to 309c = 580 to 592

33.3 at.% P, high-pressure phase[1990Oka, Mas2]at T = 25°C to 1000°C [1978Sel]

0 to 25 at.% Cr, room temperature[1978Sel]

ν, FeP2 oP6PnnmFeS2 (marcasite)

a = 497.29b = 565.68c = 272.30

66 to 67 at.% P [1990Oka, Mas2]at 66.7 at.% P [1969Dah, 1990Oka]

θ, FeP4 (I) mP30P21/c a = 461.9

b = 1367.0c = 700.2β = 101.48°

80 at.% P [1990Oka, Mas2][1978Jei]

θ, FeP4 (II) oC20C2221

a = 500.5b = 1021.2c = 553.0

80 at.% P, high-pressure phase,synthesized at 60 kbar in a cubic anvildevice [1978Sug]

(continued)

Cr–Fe–P 7






Comments/References

χ (Fe-P) o**

a = 359b = 401c = 432

< 20 at.% P, metastable, labelled as “Fe4+P” [1990Oka]after 25 h aging at T = 500°C [1961Hor]

* τ,(CrxFe1–x)2P

oP12PnmaCo2Si or ZrFeP a = 580 to 583

b = 352 to 351.7c = 663.5 to 668.5a = 582.6b = 351.5c = 665a = 583.3b = 356.9c = 665.8

x = 0.15 to 0.55 [1969Fru, 1969Rog]

x = 0.2 to 0.55 [1969Rog]

x = 0.5 [1969Rog]

x = 0.5, T = 27°C [2004Kum]

Table 3. Investigations of the Cr-Fe-P Materials Properties


[1969Fru] Thermomagnetic balance technique Magnetic moments; Curie temperatures;magnetic-transition temperatures

[1969Rog] Thermomagnetic balance technique Magnetic moments; Curie temperatures;magnetic-transition temperatures

[1977Got] Thermomagnetic balance, samplevibrating magnetometer techniques

Magnetic transition point; saturation magneticmoment

[1978Sel] Magnetic susceptibility measurements Magnetic susceptibility

[1986Sor] Electrochemical implantation ofphosphorus

Potentiodynamic polarization curves;potentiostatic current density dependences

[1991Myu] Differential scanning calorimetry (DSC);thermomechanical analysis (TMA)

Glass transition temperatures; viscosity

[1998Kun] Conventional electrodeposition;electrochemical measurements

Corrosion resistance

[1999Kum] SQUID magnetometry Magnetization; magnetic moment

[1999Kun] Conventional electrodeposition;electrochemical measurements


[1999Lit] VSM 7.300 magnetometer techniqueapplied for the magneto-rheologicalsuspensions (MRS)

Saturation magnetization; ratio betweenremanent magnetization and saturationmagnetization; coercivity; magneticsusceptibility

(continued)

8 Cr–Fe–P



[1999Pre] Specimens dimensions and volumesdetermination; Vickers microhardnessand hardness tests; tensile tests

Sintered density; interconnected porosity;volume shrinkage; microhardness andhardness; tensile strength

[2000Kun] Nanocrystalline electrodeposition;electrochemical measurements


[2004Kum] Vibrating sample magnetometer,SQUID magnetometer techniques

Magnetization

[2004Kun] Nanocrystalline electrodeposition;electrochemical measurements


Fig. 1. Cr-Fe-P. The quasibinary Cr2P-Fe2P system

Cr–Fe–P 9



Fig. 2. Cr-Fe-P. Partial projection of the liquidus and solidus surfaces

10 Cr–Fe–P


Fig. 3. Cr-Fe-P. Partial isothermal section at 1140°C

Cr–Fe–P 11



Fig. 4. Cr-Fe-P. Calculated temperature-composition section Cr-Fe2P

12 Cr–Fe–P


Fig. 5. Cr-Fe-P. Calculated temperature-composition section Cr/Fe = 1/1

Cr–Fe–P 13



Fig. 6. Cr-Fe-P. Calculated temperature-composition section Cr3P-Fe3P

14 Cr–Fe–P


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[1969Dah] Dahl, E., “Refined Crystal Structures of PtP2 and FeP2”, Acta Chem. Scand., 23(8),2677–2684 (1969) (Crys. Structure, Experimental) as quoted by [1990Oka]

[1969Fru] Fruchart, R., Roger, A., Senateur, J.P., “Crystallographic and Magnetic Properties of SolidSolutions of the Phosphides M2P, M = Cr, Mn, Fe, Co, and Ni”, J. Appl. Phys., 40(3),1250–1257 (1969) (Crys. Structure, Experimental, Magn. Prop., 45)

[1969Rog] Roger, A., Senateur, J.-P., Fruchart, R., “Crystallographic and Magnetic Properties of SolidSolutions Among the Phosphides Ni2P - Co2P - Fe2P - Mn2P and Cr2P” (in French),Ann. Chim. (Paris), 4(2), 79–91 (1969) (Crys. Structure, Experimental, Magn. Prop., 44)

[1970Had] Hadrys, H.G., Frohberg, M.G., Elliott, J.F., Lupis, C.H.P., “Activities in the Liquid Fe-Cr-C(sat), C-Fe-P (sat) and Fe-Cr-P Systems at 1600°C”, Metall. Trans., 1, 1867–1874 (1970)(Thermodyn., Calculation, 23)

[1972Jei] Jeitschko, W., Donohue, P.C., “The High Pressure Synthesis, Crystal Structure, and Proper-ties of CrP4 and MoP4”, Acta Cryst., Sect. B., 28(6), 1893–1898 (1972) (Crys. Structure,Experimental, Phys. Prop., 30)

[1972Sel] Selte, K., Kjekshus, A., “Structural and Magnetic Properties of FeP”, Acta Chem. Scand., 26(3), 1276–1277 (1972) (Crys. Structure, Experimental, Magn. Prop.) as quoted by [1990Oka]

[1973Mae] Maeda, Y., Takashima, Y., “Mössbauer Studies of FeNiP and Related Compounds”, J. Inorg.Nucl. Chem., 35(6), 1963–1969 (1973) (Crys. Structure, Experimental, Electronic Structure,12)


[1977Got] Goto, M., Tange, H., Tokunaga, T., Fujii, H., Okamoto, T., “Magnetic Properties of the (Fe1–xMx)3P Compounds”, Jpn. J. Appl. Phys., 16(12), 2175–2179 (1977) (Crys. Structure,Experimental, Magn. Prop., 16)

[1978Jei] Jeitschko, W., Braun, D.J., “Synthesis and Crystal Structure of the Iron Polyphosphide FeP4”,Acta Cryst. B, 34, 3196–3201 (1978) (Crys. Structure, Experimental, 30)

[1978Sel] Selte, K., Birkeland, L., Kjekshus, A., “On the Structural and Magnetic Properties of Cr1–-tFetP, Mn1–tCotP and Fe1–tCotP”, Acta Chem. Scand., Ser. A, 32(8), 731–735 (1978) (Crys.Structure, Experimental, Magn. Prop., 21)

Cr–Fe–P 15



[1978Sug] Sugitani, M., Kinomura, N., Koizumi, M., “Preparation and Properties of a New Iron Phos-pide FeP4”, J. Solid State Chem., 26(2), 195–201 (1978) (Crys. Structure, Experimental,Electr. Prop., Magn. Prop., 14)

[1979Yam] Yamada, K., Kato, E., “Mass Spectrometric Determination of Activities of Phosphorus inLiquid Fe-P-Si, Al, Ti, V, Cr, Co, Ni, Nb and Mo Alloys” (in Japanese), Tetsu to Hagane(J. Iron Steel Inst. Jap.), 65(2), 273–280 (1979) (Thermodyn., Calculation, Experimental,Review, 40)

[1981Erh] Erhart, H., Grabke, H.J., “Equilibrium Segregation of Phosphorus at Grain Boundaries ofFe-P, Fe-C-P, Fe-Cr-P, and Fe-Cr-C-P Alloys”, Met. Sci., 15(9), 401–408 (1981) (Morphol-ogy, Experimental, Electronic Structure, 17)

[1981Maa] Maaref, S., Madar, R., “Crystal Chemistry of M12P7 Phases in Relation with the M2P Phos-phides”, J. Solid State Chem., 40, 131–135 (1981) (Crys. Structure, Experimental, 11)

[1981Mat] Matsuyama, T., Suto, H., “Grain Boundary Segregation in Fe-P-Cr Ferritic Alloys”, J. Jpn.Inst. Met., 45(3), 233–241 (1981) (Morphology, Experimental, Interface Phenomena, 32)

[1983Yam] Yamada, K., Kato, E., “Effect of Dilute Concentrations of Si, Al, Ti, V, Cr, Co, Ni, Nb andMo on the Activity Coefficient of P in Liquid Iron”, Trans. Iron Steel Inst. Jpn., 23(1),51–55 (1983) (Thermodyn., Calculation, Experimental, 16)

[1984Ban] Ban-Ya, S., Maruyama, N., Kawase, Y., “Effects of Ti, V, Cr, Mn, Co, Ni, Cu, Nb, Mo and Won the Activity of Phosphorus in Liquid Iron” (in Japanese), Tetsu to Hagane, 70(1), 65–72(1984) (Thermodyn., Calculation, Experimental, Review, 21)

[1986Sor] Sorensen, N.R., Diegle, R.B., Picraux S.T., “The Effect of P Implantations on Passivity ofFe-Cr Alloys in Acidic Electrolytes”, J. Mater. Res., 1(6), 752–757 (1986) (Morphology,Experimental, Electrochemistry, 21)

[1986Yuq] Yu-Qing, W., McMahon, C.J., “Surface Segregation in an FeCrP Alloy”, Scr. Metall., 20(1),19–23 (1986) (Morphology, Experimental, Interface Phenomena, 12)

[1988Rag] Raghavan, V., “The Cr-Fe-P System” in “Phase Diagrams of Ternary Iron Alloys”, IndianInst. Metals, Calcutta, 3, 60–67 (1988) (Crys. Structure, Phase Diagram, Review, #, 12)

[1990Oka] Okamoto, H., “The Fe-P (Iron-Phosphorus) System”, Bull. Alloy Phase Diagrams, 11(4),404–412 (1990) (Crys. Structure, Phase Diagram, Thermodyn., Assessment, Review, Magn.Prop., #, 88)

[1991Myu] Myung, W.-N., Yang, S.-J., Kim, H.-G., “Glass Transition and Viscous Flow Behavior ofAmorphous Fe-M-P (M = Cr, V or Mo) Alloys”, Mater. Sci. Eng., A133, 418–422 (1991)(Crys. Structure, Morphology, Experimental, Kinetics, Phys. Prop., 9)

[1993Din] Ding, X., Wang, W., Han, Q., “Thermodynamic Calculation of Fe-P-j System Melt”,Acta Metall. Sin. (China), 29(12), B527-B532 (1993) (Thermodyn., Calculation, Theory, 7)

[1994Liu] Liu, C., Abiko, K., Tanino, M., “Role of Chromium in the Intergranular Fracture of High Pur-ity Fe-P-Cr Alloys with Small Amounts of Carbon”,Mater. Sci. Eng.,A176, 363–369 (1994)(Morphology, Experimental, Kinetics, Interface Phenomena, 24)

[1997Zai] Zaitsev, A.I., Shelkova, N.E.,Mogutnov, B.M., “Thermodynamic Properties of Iron-Chromium-Phosphorus Melts”, Russ. J. Inorg. Chem., 42(10), 1567–1573 (1997) (Thermodyn., Calcula-tion, Experimental, 40)

[1998Kun] Kunioshi, C.T., Correa, O.V., de Lima, N.B., Ramanathan, L.C., “The Development ofAmorphous Fe-Cr-P Electrodeposits” in “Proceedings and Fabrication of Advanced Materi-als VII”, Proceedings of TMS Symposium, TMS - Miner. Metals & Mater. Soc., Warrendale,PA, USA, 475–486 (1998) (Morphology, Experimental, Electrochemistry, Interface Phenom-ena, 24)

[1998Mie] Miettinen, J., “Approximate Thermodynamic Solution Phase Data for Steels”, Calphad, 22(2), 275–300 (1998) (Phase Diagram, Phase Relations, Thermodyn., Assessment, Calcula-tion, 98)

[1998Zai1] Zaitsev, A.I., Litvina, A.D., Shelkova, N.E., Mogutnov, B.M., “Association in TernaryMetallic Melts Fe-Mn-Si, Fe-Cr-P and Fe-Mn-P”, Thermochim. Acta, 314, 307–315 (1998)(Phase Relations, Thermodyn., Calculation, Experimental, 31)

16 Cr–Fe–P


[1998Zai2] Zaitsev, A.I., Shelkova, N.E., Mogutnov, B.M., “Thermodynamic Properties of the Fe-Cr-PLiquid Solution”, Metall. Mater. Trans. B, 29B, 155–161 (1998) (Thermodyn., Calculation,Experimental, 39)

[1999Kum] Kumar, S., Paranjpe, S.K., Srivastava, B.P., Krishnamurthy, A., Sahni, V.C., “MagneticStructure of (Fe0.97Cr0.03)2P”, Pramana - J. Phys., 52(1), 111–120 (1999) (Crys. Structure,Experimental, Magn. Prop., 18)

[1999Kun] Kunioshi, C.T., Correa, O.V., de Lima, N.B., Ramanathan, L.C., “Development of Electrode-posited Fe-Cr-P Amorphous Metallic Alloys”, Surf. Eng., 15(5), 395–400 (1999) (Morphol-ogy, Experimental, Electrochemistry, 27)

[1999Lit] Lita, M., Nicoara, M., “The Use of Amorphous and Quasi-Amorphous Fe-Cr-P Powders forFabrication of Magneto-Rheological Suspensions”, J. Magn. Magn. Mater., 201, 49–52(1999) (Crys. Structure, Experimental, Magn. Prop., 3)

[1999Mie] Miettinen, J., “Thermodynamic Description of Cr-P and Fe-Cr-P Systems at Low PhosphorusContents”, Calphad, 23(1), 141–154 (1999) (Phase Relations, Thermodyn., Calculation, 28)

[1999Pre] Preusse, H., Bolton, J.D., “Use of Phosphide Phase Additions to Promote Liquid Phase Sin-tering in 316L Stainless Steels”, Powder Metall., 42(1), 51–62 (1999) (Morphology, Experi-mental, Mechan. Prop., Phys. Prop., 53)

[2000Kun] Kunioshi, C.T., Correa, O.V., de Lima, N.B., Ramanathan, L.C., “Effect of Processing Para-meters on Structure and Properties of Nanocrystalline Fe-Cr-P Electrodeposits” in “UltrafineGrained Materials”, Proceedings of a Symposium, TMS - Miner. Metals Mater. Soc.,99–110 (2000) (Morphology, Experimental, Electrochemistry, 24)

[2002Per] Perrot, P., Batista, S., Xing, X., “Fe-P (Iron-Phosphorus)”, Diagrams as Published in MSITWorkplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stutt-gart; Document ID: 20.16107.1.20, (2002) (Phase Diagram, Phase Relations, Crys. Structure,23)

[2004Kum] Kumar, S., Krishnamurthy, A., Srivastaya, B.K., Das, A., Paranjpe, S.K., “Magnetization andNeutron Diffraction Studies on FeCrP”, Pramana - J. Phys., 63(2), 199–205 (2004) (Crys.Structure, Experimental, Magn. Prop., 7)

[2004Kun] Kunioshi, C.T., de Lima, N.B., Correa, O.V., Castro, N.A., Ramanathan, L.C., “Effect of Pro-cessing Parameters on Properties of Nanocrystalline FeCrP Electrodeposits” in “MPMD FifthGlobal Innovations Proceedings. Surfaces and Interfaces in Nanostructured Materials &Trends in LIGA, Miniaturization, and Nanoscale Materials”, TMS, Warrendale, PA, USA,173–181 (2004) (Morphology, Experimental, Electrochemistry, Interface Phenomena, 20)

[2005Sho] Shohoji, N., Monteiro Dias, M.C., “Empirical Expression of Phosphorus Solubility in MoltenFe1–yCry Given as Functions of Temperature and Phosphorus Activity”, ISIJ Int., 45(9),1226–1231 (2005) (Phase Relations, Thermodyn., Experimental, 12)



Cr–Fe–P 17



Chromium – Iron – Sulfur

Vasyl Tomashik

Introduction

The Cr-Fe-S ternary system is one of the basic constituent systems of relevance to stainless steels. The firstattempt to determine the phase diagram was carried out in 1938 by [1938Vog]. Later, it was assessed by[1949Jae, 1988Rag] with subsequent updates [1998Rag, 2004Rag], the latter review taking into considera-tion the literature up to the year 2000. [2000Oik] carried out a thermodynamic evaluation of this system.Some literature data concerning the investigation of the Cr-Fe-S ternary system were missing in the reviewsof [1988Rag, 1998Rag, 2000Oik, 2004Rag], but all available data are taken into account in the presentassessment.At temperatures above 700°C, Cr1–xS and Fe1–xS form a continuous series of solid solutions [1938Vog,1963Kan, 1978Moh, 1978Rao, 1996Nog1, 1997Nog, 2000Oik]. This solid solution splits into an Fe richand a Cr rich sulfide solution at lower temperatures [1951Vog, 1969Gor2, 1973Nis, 1978Moh,1978Rao]. The mutual solubility of the two monosulfides decreases on cooling below 700°C [1969Gor2].[1982Ind2] determined the phase diagram of the FeS-Cr2S3 system. A peritectic reaction between liquid andCr2S3 at 1440°C yields Cr2FeS4, and a eutectic reaction yielding Cr2FeS4 and Fe1–xS occurs at 1090°C. Apolymorphic transition of the ternary compound Cr2FeS4 was observed at 1060°C. But these results areinconsistent with the constructed liquidus surface and isothermal sections [1938Vog, 1969Gor2,1969Gor3, 1978Moh, 1979Cha, 1988Rag, 2000Oik].A ternary compound (mineral name daubreelite) Cr2FeS4 was confirmed to form in this system [1944Lun,1956Lot, 1965Alb, 1965Bou, 1969Gib1, 1969Gib2, 1969Gor1, 1969Gor2, 1969Shi, 1972Wat, 1981Rie].The liquidus surface of the Fe-FeS-CrS-Cr region was constructed by [1938Vog]. Isothermal sections of theCr-Fe-S ternary system at 1300, 1090, 950, 800, 700 and 600°C were constructed by [1965Dah, 1969Gor2,1969Gor3, 1978Moh, 1979Cha, 1988Fuj]. [2000Oik] calculated isothermal sections of the Fe-FeS-CrS-Crpart of this ternary system at 1600, 1500, 1370, 1300, 1090 and 950°C. Some polythermal sections weredetermined experimentally and calculated by [1938Vog, 2000Oik, 2002Mit].Thermodynamic properties of solid and liquid alloys in the Cr-Fe-S ternary system were investigated bothexperimentally and theoretically by [1960Gri, 1969Ban, 1970Don, 1971Don, 1972Bur, 1973Buz, 1974Sig,1974Tre, 1976Kes, 1976Tre, 1977Tre1, 1977Tre2, 1979Shv, 1981Rug, 1982Zhu, 1987Nar, 1988Fuj,1998Fed, 1999Kau, 2000Oik]. At temperatures between 662 and 949°C, the sulfur is present in the vaporphase above Cr2FeS4 [1982Zhu].All of the details of the experimental studies are given in Table 1.

Binary Systems

The Cr-Fe, Cr-S and Fe-S binary systems are accepted from [Mas2].

Solid Phases

Crystallographic data of all unary, binary and ternary phases are listed in Table 2.Only one ternary compound, τ (Cr2FeS4), is formed in the Cr-Fe-S ternary system [1944Lun, 1956Lot,1965Alb, 1965Bou, 1969Gib1, 1969Gib2, 1969Gor1, 1969Gor2, 1969Shi, 1972Wat, 1981Rie]. It is stableat ambient pressure and temperatures up to 1350°C [1969Gor2]. Meanwhile, [1969Shi] indicated that Cr2FeS4decomposes on heating above 750°C, and according to the data of [1976Tre], it decomposes at 1020°C. Thecrystal structure of this compound is cubic at temperatures up to 4.2 K [1964Shi, 1969Gib2], and at 1040°C(at 1060°C [1982Ind2]) it transforms into the NiAs structure type [1985Sok]. An excess of Fe moves the tem-perature of the polymorphic transformation to 670°C. At high pressures, it transforms to a NiAs type hexago-nal [1965Alb, 1968Tre1, 1969Gor2, 1998Fed] or the ordered monoclinic structure [1967Bou, 1968Tre1,

Cr–Fe–S 1



1968Tre2, 1969Gor2, 1978Moh, 1998Fed]. These high pressure modifications can be considered as ternarysolid solutions based on the corresponding binary chromium sulfide [1988Rag]. According to the data of[1967Bou], the Cr2FeS4 compound retains its monoclinic structure when heated to 450°C at ambient pressure.At 600°C, however, it reverts to the spinel starting material. The difference in the density of the low- and high-pressure Cr2FeS4 polymorphs is equal to 6.6 % [1968Tre1]. A reversible distortion of the spinel cubic cell wasdetermined within the temperature interval from 370 to 770°C [1979Los]. At 1.0 GPa and 1000°C, decompo-sition of the spinel was observed by [1998Fed].A complete series of solid solutions exists in the Cr1–xS-Fe1–xS system at high temperatures [1938Vog,1963Kan, 1978Moh, 1978Rao, 1996Nog1, 1997Nog, 2000Oik]. A nonlinear variation of the lattice para-meters was reported by [1973Nis]. Nonstoichiometry of these solid solutions increases with increasing Crcontent and sulfur pressure, while it decreases with increasing temperature at a constant composition andS pressure [1996Nog1, 1997Nog]. Cation vacancies are the predominant lattice defects in the Cr1–xS-Fe1–xSsolid solutions.The solubility of sulfur in solid Cr-Fe alloys has been determined by [1968Jos, 1976Bar, 1976Gro, 1977Pet,1981Oud, 1989Bar]. [1968Jos] indicates that the solubility of S in the γ phase decreases by an order ofmagnitude from 0.0132 mass% for pure Fe to 0.0013 mass% for an Fe + 10 mass% Cr alloy at 1000°C.The results of [1976Bar, 1976Gro, 1981Oud, 1989Bar] were determined by equilibration of such alloysin an H2/H2S atmosphere. In some cases, the limiting solubility corresponding to the precipitation of chro-mium sulfide was also determined [1976Gro]. For a constant H2/H2S ratio, the solubility in the α phaseincreases with increasing Cr content, but no systematic variation could be seen in the results of [1976Bar,1981Oud, 1989Bar] (the sulfur solubility was expressed as a function of H2S/H2 ratio and temperature).In the H2S partial pressure range investigated, the sulfur content was always fixed, and the temperaturewas strictly proportional to the pressure following Sieverts’ law [1981Oud].

Quasibinary Systems

The CrS-FeS section was recognized as quasibinary by [1951Vog, 2000Oik], showing complete miscibilityin the liquid and solid state. The CrS-FeS phase diagram was calculated by [2000Oik] for temperaturesabove 700°C, and is shown in Fig. 1. Considering the fact that both compounds CrS and FeS have homo-geneity ranges in the binary systems, and that below 700°C the (Cr1–xFex)S solid solution decomposes intotwo phases, there are doubts whether the CrS-FeS is a true quasibinary system across the whole compositionand temperature range.The CrS-Fe vertical section was called “pseudobinary” by [1999Kau, 2000Oik], although it shows severalthree-phase regions, which should not appear in a quasibinary (pseudobinary) system. This section is anordinary vertical section and is reported below under “Temperature-Composition Sections”.It is possible that the Cr2S3 - FeS section is also quasibinary, but the constructed phase diagram for this sys-tem [1982Ind2] is inconsistent with the liquidus surface and isothermal sections given by [1938Vog,1969Gor2, 1969Gor3, 1978Moh, 1979Cha, 1988Rag, 2000Oik].


A tentative reaction scheme is given in Fig. 2 [1988Rag]. The transition reaction U4 at 1050°C has beendetermined by [1938Vog]. According to the data of [2004Rag], the temperature of this reaction is estimatedfrom the computed vertical section of [2000Oik] to be 1325°C (at 19.5 mass% S), which is 275°C higher.The transition reactions U6 and U7 are those given by [1969Gor3]. The other postulated reactions yield thecorrect three-phase equilibria in the isothermal sections at 700 and 600° [1969Gor2, 1969Gor3]. The S richliquid (L3) is assumed to solidify through the ternary eutectic reaction E just below 115°C [1988Rag]. Thisinvariant equilibrium is also included in Table 3.


The liquidus surface of the Fe-FeS-CrS-Cr region is shown in Fig. 3. The liquidus line enclosing the mis-cibility gap L1+L2 has a lower critical point c at 1390°C and ~5 at.% Cr, ~25 at.% S. Isotherms at 100°C

2 Cr–Fe–S


intervals between 1800 and 1100°C have been drawn by [1988Rag] using the data of [1938Vog] and[1965Dah]. Two pairs of coexisting compositions of the metallic and sulfide liquids are shown bydotted lines.No experimental data are available for the CrS-S-FeS region. Based on the binary data and Cr2FeS4 beingstable up to at least 1350°C, [1988Rag] presented a hypothetical liquidus projection, as shown in Fig. 3.The dependence of S concentration on Cr content in liquid Fe can be expressed as [1972Bur]:1.0 – log {mass% Cr} + 0.020 {mass% Cr} = log {mass% S} – 0.06 {mass% S}.

Isothermal Sections

The isothermal sections for 1300, 950, 700 and 600°C are shown in Figs. 4 to 7, as derived from [1969Gor2,1969Gor3, 1978Moh, 1988Fuj, 1995Smi]. The bottom part of the section at 950°C (Fig. 5) is based on theresults of [1963Kan]. The upper part has been drawn schematically by [1988Rag] using their reactionscheme (Fig. 2). The isothermal section at 600°C, determined by [1995Smi], is in broad agreement withthat assessed by [1988Rag]. The main finding is that the homogeneity range of Cr2FeS4 is larger than pre-viously believed, and is 16.2 to 23.7 mass% Fe, 33.8 to 39.7 mass% Cr and 42.4 to 44.9 mass% S.The isothermal sections at 1600 and 1550°C were calculated by [2000Oik]. Both isothermal sections are notcompatible with the assumed liquidus surface and, therefore, are not accepted in the present evaluation. Theisothermal sections at 1600 and 1550°C were calculated in [2000Oik] where the monotectic temperature inthe Cr-CrS part of the Cr-S system was estimated to be above 1800°C while in the liquidus surface this tem-perature is assumed to be 1550°C in agreement with [Mas2].


Five vertical sections of the Cr-Fe-S system at 19.5 mass% S and at Fe/Cr weight ratios of 1 : 1, 4 : 1, 9 : 1and 39 : 1 were constructed by [1938Vog]. [2000Oik] compared the isothermal sections, constructed by[1938Vog], with vertical sections calculated at 19.5 mass% S and at Cr/Fe = 4 : 1. It is interesting to notethat the phase boundaries calculated by [2000Oik] agree with the thermal analysis data points given in[1938Vog], although the configuration of the phase fields is quite different. Figure 8 summarizes the partialvertical sections for 0.001 to 0.1 mass% S in the vicinity of the γ loop of the Cr-Fe system, andshows the change in the solubility of sulfur in the bcc(α) and the fcc(γ) phases in equilibrium with eitherthe (Cr1–xFex)1–yS or the L2 phase [2000Oik]. It is shown that the sulfur solubility limit increases withincreasing temperature. For the same S content, the solubility limit is at a higher temperature in the fcc(γ) phase than in the bcc(α) phase. Calculated phase boundaries for the γ+α+(Cr1–xFex)1–yS(λ),γ+α+(Cr1–xFex)1–yS(λ)+L, γ + α + L equilibria for various S contents are given in Fig. 9 [2000Oik].The CrS-Fe vertical section presented by [2000Oik] is shown in Fig. 10.The vertical sections at 0.05, 1 and 13 mass% Cr were also calculated by [2002Mit].

Thermodynamics

The heat of S dissolution in Cr-Fe alloys as a function of alloy composition is characterized by an unex-pected variation [1981Oud, 1989Bar]: it passes through a minimum between 23 and 30 mass% Cr[1976Gro]. The free energy change for the formation of the Cr2FeS4 spinel from a monoclinic Cr3S4, hex-agonal (Cr1–xFex)1–yS and sulfur vapor is given by the relationship ΔG = –1523 + 1.09 T kJ·mol–1

[1987Nar].Cr has a pronounced negative effect on the activity coefficient of sulfur in liquid Fe up to about 20 mass%Cr (eCrS = –0.019) [1960Gri] (eCrS = –0.0107 and εCrS = –2.23 at 1550°C [1969Ban], eCrS = –0.012 andεCrS = –2.64 at 1600°C and 2.4-21 mass% S [1973Buz], at 1300°C eCrS = –0.031 and –0.0085 for alloyscontaining less than 3.3 and 9.6-48 mass% Cr [1988Fuj]). According to the data of [1970Don, 1971Don,1974Sig], the first order interaction coefficients are eCrS = –94.2/T + 0.040 and εCrS = –20.200/T + 8.64for temperatures between 1525 and 1755°C (εSCr = –153/T + 0.062 [1974Sig]). The first order enthalpyand entropy interaction coefficients are found to be hCrS = –430 ± 70 and sCrS = –0.183 ± 0.007(ηCrS = –40000 ± 6500 and σCr

S = –17.11 ± 0.65). The second order interaction coefficient rCrS is equalto zero [1974Sig].

Cr–Fe–S 3



The activity of S in a melt containing 16 mass% Cr is one-half of that to be expected in a Cr-free alloy[1960Gri]. The entropy change owing to the high-pressure transformation of Cr2FeS4 is equal to23.0 ± 8.4 J·(mol·K)–1 [1965Alb]. According to thermodynamic calculations, an excess sulfide phase inthe Cr-Fe-S alloys must be CrS [1979Shv]. The activation energy of sulfur volume diffusion in Fe-18mass% Cr alloys at temperatures between 850 and 1200°C has been determined to be 184.2 kJ·mol–1

[1981Rug].Some thermodynamic data concerning the Cr-Fe-S system are given in the Tables 4 and 5.

Notes on Materials Properties and Application

CrS is commonly present in the form of non-metallic inclusions in stainless steels [2000Oik]. This sulfideinclusion modifies the corrosion resistance of free-cutting stainless steels.Hydrogen sulfide corrosion of Cr-Fe alloys is an important problem in the petrochemical, fossil fuel com-bustion, coal conversion and other energy related industries [1985Nar, 1995Smi]. The corrosion resistanceof the Cr-Fe alloys decreases significantly in the presence of S or H2S occurring in the atmosphere of hightemperature industrial processes [1993Cie, 1994Cie]. The results of [1995Smi] suggest that there is a criti-cal Cr composition (≈ 51 mass%) above which outward Cr diffusion promotes interfacial layer formation.Cr-Fe alloys below this critical composition would be more resistant to sulfidation corrosion. According tothe data of [1996Smi], the optimum Cr composition range for sulfidation resistance at 600°C is between 20and 40 mass% Cr.The addition of sulfur to Fe-1 mass% Cr alloy drastically increases the hardness, and the hardness of Fe-8mass% Cr and that of Fe-8 mass% Cr-0.05 mass% S alloys show the same level [2003Mit].The presence of a monolayer of adsorbed S on the surface of an Fe-18 mass% Cr alloy prevents the high-temperature depletion of Cr [1980Oud]. But it causes an enrichment of Cr if the surface has been previouslydepleted. It is suggested that the presence of the adsorbed sulfur layer hinders the vaporization of Cr at hightemperatures.Cr2FeS4 is ferrimagnetic and has semiconducting properties (ΔEg = 0.038 eV at room temperature)[1965Bou, 1966Rac, 1970Rob, 1975Goe, 1997Ram, 1998Bou, 1999Che] (ΔEg = 0.09 eV [1998Fed]).According to the data of [1968Tre2, 1999Che], this compound is paramagnetic at 27°C. The high pressurevariant of Cr2FeS4 is antiferromagnetically ordered below the Neel temperature [1968Tre2]. According tothe preliminary results of [1993Sok], quenched Cr0.5Fe0.5S samples are antiferromagnets. The Cr2FeS4compound is characterized by the existence of colossal magnetoresistance [1997Ram, 1999Che, 2002Nat,2005Mer] and does not exhibit double-exchange electron transfer nor Jahn-Teller distortion [2002Nat].There are no Fe3+ and Cr2+ ions in Cr2FeS4 [1999Che, 2000Tsu, 2001Kim2].The maximum resistance ratio for Cr2Fe0.92S4 was observed at 183 K about 12 % under 1.6 T [2001Kim1].Cr1–xFexS4 solid solutions possess a record-high Curie temperature (TC = 940 K) [1980Pet].

Miscellaneous

In steels, containing 2.45 mass% Cr and 0.059 or 0.91 mass% S, a solid solution of (Cr,Fe)1–xS is formed inthe inclusions [1968Ska]. The sulfidation kinetics of Cr-Fe alloys shows that the formation of a sulfide scaleproceeds according to a parabolic law, irrespective of temperature and Cr content in the alloy [1968Mro1,1968Mro3, 1996Nog2, 1998Cie1, 2006Sro] (at less than 60 at.% Cr [1982Nar]), but at 1000°C and a sulfurpressure of 10–5 Pa, it follows a linear rate law [2006Sro]. The parabolic rate constant increased withincreasing Cr content and sulfur partial pressure, while they showed an abnormal temperature dependenceat p(S2) = 10–3 Pa [1996Nog2]. This indicates that the reaction rate is determined by the diffusion ofreagents in the reaction product. Three ranges of concentration of alloying elements with different sulfida-tion mechanisms may be distinguished. For low-Cr alloys, the scale formed is single phase Fe1–xS[1968Mro2, 1998Cie1, 1998Cie2]. At chromium contents between 4 and 40 at.%, heterophase scales com-posed of two layers of distinctly different morphological structures are formed. The compact outer layercontains the Fe1–xS phase, whereas the inner layer is a heterophase mixture of Fe1–xS and Cr2FeS4. Accord-ing to the data of [1998Cie2], CrxFe1–xS solid solutions are formed on sulfidation of Cr-Fe alloys containingmore than 10 at.% Cr. Sulfidation of alloys containing more than 40 at.% Cr yielded monophase scales built

4 Cr–Fe–S


up from Cr2S3 [1968Mro2, 1968Mro3, 1982Nar]. The scales on quenched specimens were composed ofeither (Cr,Fe)5S6 or (Cr,Fe)1–xS plus small amounts of (Cr,Fe)3S4, at p(S2) = 10–2 Pa and of (Cr,Fe)5S6 atp(S2) = 10–5 Pa [1982Nar]. Internal sulfidation is most pronounced in high Cr content alloys and at low sul-fur pressures. The parabolic sulfidation rate at p(S2) = 10–2 Pa is large and of a practically constant valueirrespective of alloy composition. The kinetics at p(S2) = 10–5 Pa, however, increase from a very low valueby three orders of magnitude as the Cr content of the alloy is increased from 5 to 60 at.% [1982Nar]. A tri-plex scale (Cr,Fe)1–xS / (Cr,Fe)3S4 / (Fe,Cr)1–xS is formed at high S pressures, a single-phase (Fe,Cr)1–xS orduplex (Cr,Fe)1–xS / (Fe,Cr)1–xS scale at an intermediate S pressure range and a single-phase (Fe,Cr)1–xSscale at low S pressure [1984Nar1, 1984Nar2]. The critical sulfur pressures are 10–2 and 10–5 Pa at 700°C, 1 and 10–4 Pa at 800°C and 102 and 10–2 Pa at 900°C. Temperature has a much stronger effect than expo-sure time on the growth of the sulfide layer [1985Nar]. It was shown that Cr in the scale enriches near analloy/scale interface and then decreases gradually toward a scale/gas interface, and these cation profilesbecome steeper with increasing sulfur pressure [1996Nog2]. According to the data of [2006Sro], a mono-layer scale with a composition corresponding to a mixture of the chromium sulfides Cr7S8 and Cr5S6is formed on the surface of a Fe-46 at.% Cr alloy. The addition of Sn increases the sulfidation rate ofsuch alloy.Solid-state diffusion couples of low Cr-Fe alloys and Fe0.95S produce internal sulfide precipitates within thelow Cr-Fe alloy microstructure [1996Smi]. The first sulfide to precipitate is Cr1–xS. Subsequent Cr and Fediffusion produces a phase change to Cr2FeS4 and then to Fe1–xS. The internal sulfide precipitate morphol-ogy changes from small, spherical particles near the diffusion front to large faceted precipitates near the ori-ginal alloy surface. The thinnest internal sulfidation zone occurred in alloys in the composition range 20.5 to51.1 mass% Cr.For CrxFe100–x (x ≤ 45.5) samples treated at 800°C in an H2/H2S atmosphere, the partial pressure of sulfurbeing 10–3 Pa, the sulfidation was selective [1993Cie, 1994Cie, 1998Cie3]. At x ≤ 3-5, the metallic phasebecome enriched in Cr, while for x ≥ 3-5 it become enriched in Fe.Figures 11 and 12 show phase relations in the CrS-FeS section at 800 and 900°C in the S-pressure ranges of10–2-10–6 and 100-10–5 Pa respectively [1987Nar]. With decreasing sulfur pressure, the Cr2FeS4 spineltransforms into the monoclinic and hexagonal sulfides. The sulfur pressure gap in the two-phase field ofthe (Cr,Fe)3S4 and (Fe,Cr)1–xS decreases with decreasing temperature. The temperature at which the two-phase field disappears is estimated to be 763°C.Chromium reduces the sulfur diffusion mobility in the Cr-Fe-S system [1974Mok].The diffusion coefficients of S in Cr-Fe alloys containing 18, 26 and 34 mass% Cr could be expressed by:D = 0.166 exp(–184/RT), D = 3.06·10–2 exp(–171/RT) and D = 6.18·10–2 exp(–177/RT) respectively[1982Fil]. The sulfur diffusion in these alloys seems to proceed via vacant sites rather than through inter-stices.A structural anomaly below 60 K, which could be interpreted in terms of a triclinic distortion within crystal-lographic domains, was found for the Cr2FeS4 compound by [2005Mer].Between 9.7 and 10 K, the Mössbauer spectrum of the Cr2FeS4 shows an abrupt change in the electric fieldgradient, which indicates a low-temperature transition [1972Spe]. It is possible that a-sites, which are occu-pied by Fe2+ ions, distort tetragonally with c/a < 1 below 10 K.The spinel-monoclinic transition of Cr2FeS4 was observed at sulfur pressures of 10–3-10–3.5 Pa at 800°Cand at 100-100.5 Pa at 900°C [1987Nar]. The results of high-temperature XRD measurements and the phasecomposition data indicate the occurrence of a reversible solid-phase reaction Cr2FeS4 + Fe ⇌ 4Cr0.5Fe0.5S,coexisting with the first-order structural transition in Cr2FeS4 [1993Sok].Cr2FeS4 is characterized by significant nonstoichiometry, which is the result of metallic point defects[1986Bal, 1988Rag, 1995Smi, 1999Bou]. The nonstoichiometric compound undergoes a phase transition“metal-dielectric” at ~670°C, which is accompanied by the lattice distortion and magnetic transition ofthe “order-disorder” type [1979Los].Materials in the system Fe2+Fe3+xCr2–xS4 have been prepared between x = 0 and x = 0.5 [1970Rob,1981Rie]. All of the compositions form the spinel structure. Magnetic data suggest an ionic distributionof Fe2+1–xFe

3+x{Fe

2+xCr2–x}S4.

Cr–Fe–S 5



Neutron diffraction studies of the Cr2FeS4 spinel show the degree of inversion to be zero and the sulfurparameter to be u = 0.2608 [1964Col] (u = 0.384 ± 0.002 [1964Shi]). Low temperature results suggest asimple ferrimagnetic model for the compound.Daubreelite (Cr2FeS4) is a common constituent of iron meteorites [1944Lun, 1965Alb, 1969Gor1,1969Gor2, 1978Moh, 1986Bal].The Curie temperature of Cr2FeS4 is within the temperature interval of 167 to 195 K [1969Gib2, 1969Shi,1975Goe, 1979Los, 1997Ram, 1998Fed, 1999Che, 2001Kim2] (170, 169 and 200 K for Cr2Fe0.96S4,Cr2Fe0.92S4, Cr1.9Fe1.1S4 respectively [2000Yan2, 2001Kim2]), the Neel temperature is 180 K (near 200K [1968Tre2], 172 K [2002Kim2]) with σ0 = 1.59 μB [1964Shi] and Θp = –290 K [1969Gib1]. The abso-lute magnetization at 0 K was found to be 1.79 μB [1969Gib2] (1.6 μB at 4.2 K [1969Shi]). For the materialsin the system Fe2+Fe3+xCr2–xS4 the Curie temperature increases with increasing Fe3+ content, from 180 Kfor Cr2FeS4 to 302 K for Fe2+Fe3+0.5Cr1.5S4 [1970Rob] (it is 220 K for Fe2+Fe3+0.2Cr1.8S4 [1975Bab]).The magnetic moment of Cr2FeS4 is 1.52 μB (1.8 μB [2000Tsu]) and 1.79 μB for Fe2+Fe3+0.3Cr1.7S4. Inthe Fe2+Fe3+xCr2–xS4 ions, Fe

2+ lie in the octahedral positions with the Cr3+ ions [1975Bab]. The excessFe ions in Cr1.9Fe1.1S4 enter the lattice [2000Yan2].Neutron diffraction of Cr2FeS4 above 10 K shows that there is no crystallographic distortion and revealsantiferromagnetic ordering with the magnetic moment of Fe2+ (–3.52 μB) aligned antiparallel to Cr3+

(2.72 μB) [2002Kim1, 2002Kim2]. With increasing Fe deficiency, the maximum magnetoresistance tem-perature increases steadily [2002Kim1]. Electron-spin-resonance measurements (100-290 K) reveal thatthe paramagnetic-ferrimagnetic transition is incomplete and that a paramagnetic phase coexists with a fer-rimagnetic phase over a certain temperature range below TC [2000Yan1].The enthalpy and entropy of the magnetic transition in the Cr1.8Fe1.2S4 alloy are 2.1 ± 0.8 kJ·mol–1 and10.3 ± 0.4 J·(mol·K)–1 respectively [1977Shc, 1977Tre2]. The Debye temperature for this alloy is395 ± 10 K [1977Tre1].For Cr2Fe0.97S4, a cooperative antiferro-distortive Jahn-Teller transition exists at 9.25 K [1975Lot].Cr2FeS4 begins to lose weight only at 800°C and begins to oxidize in an open system at 550°C formingFe2O3 and Cr2O3 (Cr1.8Fe1.2S4 begins to oxidize in an open system at 400°C forming also Fe2O3 andCr2O3) [1977Tre2].Single crystals of the Cr2FeS4 compound can be grown by a chemical transport reaction [1969Gib1,1969Shi, 1972Wat, 1998Fed, 2000Tsu] using CrCl3 as the transport agent [1972Wat]. They were alsogrown by spontaneous crystallization [1998Fed]. Films of this compound can be obtained by “explosive”vacuum condensation of the Cr2FeS4 meal powder [1982Ind1]. Polycrystalline samples of Cr2FeS4 wereobtained from FeSO4·7H2O and Cr2(SO4)3·8H2O [1998Bou].The morphology of the sulfide in Cr-Fe-S ternary alloys was found to change from a cell wall type to a glob-ular type with increasing Cr content [2002Mit, 2003Mit]. Accompanying the phase transformation ofthe matrix from the δ phase to the γ phase, two types of transgranular fine particle sulfide were formed.One is a fine spherical sulfide formed from the FeS rich liquid phase through the remelting reaction,δa→ .γ + L in less than 5 mass% Cr alloys, and the other is a fine rod-like sulfide formed through the eutec-toid reaction of δ → γ + sulfide in 5 to 13 mass% Cr alloys.

Table 1. Investigations of the Cr-Fe-S Phase Relations, Structures and Thermodynamics

Reference Method / Experimental Technique Temperature / Composition / PhaseRange Studied

[1938Vog] DTA, XRD, metallography Up to 1600°C / Fe-FeS-CrS-Cr

[1944Lun] XRD Room temperature / Cr2FeS4

[1956Lot] XRD, magnetic measurements 190 K - room temperature / Cr2FeS4

[1960Gri] Pressure measurements 1600-1760°C / (Cr-Fe-S) + H2/H2S

(continued)

6 Cr–Fe–S



[1963Kan] XRD, chemical analysis 950°C / CrS-FeS

[1964Col] Neutron diffraction study 4.2 and 300 K / Cr2FeS4

[1964Shi] XRD, neutron diffraction study 4.2 and 300 K / Cr2FeS4

[1965Alb] XRD 27-1030°C / Cr2FeS4

[1965Bou] XRD, electrical measurements 90-420°C / Cr2FeS4

[1966Rac] XRD Room temperature / Cr2FeS4

[1967Bou] XRD, metallography Up to 1000°C / Cr2FeS4

[1968Jos] Metallography, chemical analysis 1000°C / (Fe + 10 mass% Cr) + S

[1968Mro1,1968Mro3]

Gravimetry 700-1000°C / Cr-Fe alloys up to 74at.% Cr

[1968Mro2] XRD, EPMA, metallography 700-1000°C / Cr-Fe alloys up to 74at.% Cr

[1968Ska] EPMA Fe + 2.45 mass% Cr + 0.059 or 0.91mass% S

[1968Tre1] XRD 100-900°C / Cr2FeS4

[1968Tre2] XRD, magnetic measurements 1.5-300 K / Cr2FeS4

[1969Ban] Chemical analysis 1550°C / Cr-Fe-S

[1969Gib1] XRD Room temperature / Cr2FeS4

[1969Gib2] Magnetic measurements 200-1100 K / Cr2FeS4

[1969Gor1] XRD, high-pressure DTA, EPMA Up to 820°C / Cr2FeS4

[1969Gor2] XRD 600 and 700°C / Cr2FeS4

[1969Shi] XRD, magnetic measurements Up to 750°C / Cr2FeS4

[1970Rob] Magnetic and electrical propertiesmeasurements

4.2-300 K / Fe2+Fe3+xCr2–xS4

[1970Don,1971Don]

Levitation melting technique, chemicalanalysis

1525-1755°C / Cr-Fe-S up to 40mass% Cr

[1972Spe] Mössbauer spectroscopy 2-19 K / Cr2FeS4

[1972Wat] XRD Room temperature / Cr2FeS4

[1973Buz] XRD, EPMA 1600°C / Cr-Fe-S

[1973Nis] XRD Up to 1100°C / Cr-Fe-S

[1974Mok] Radiochemical method 950-1150°C / Fe + 7.5 mass% Cr +0.008 mass% S

[1974Tre] Calvet type calorimetry Room temperature / Cr2FeS4

(continued)

Cr–Fe–S 7




[1975Bab] Magnetic measurements, neutrondiffraction study

4.2-300 K / Cr1.8Fe1.2S4

[1975Goe] XRD, magnetic measurements 4.2-300 K / Cr2FeS4

[1975Lot] Mössbauer spectroscopy 42-300 K / Cr2FexS4 (x = 0.87, 0.92,0.97 and 1.00)

[1976Bar] Radiochemical method 985-1200°C / (Cr-Fe) + H2/H2S

[1976Gro] Radiochemical method 985-1200°C / (Cr-Fe) + H2/H2S

[1976Kes] XRD, Calvet type calorimetry, chemicalanalysis

Up to 700°C / Cr2FeS4

[1976Tre] XRD, DTA, Calvet type calorimetry Up to 1200°C / Cr2FeS4

[1977Los] XRD, DTA, electrical resistivitymeasurement

77-1000 K / Cr1–xFexS (0 < x < 0.5)

[1977Pet] Radiochemical method, chemical analysis 950, 100 and 1050°C / Cr-Fe alloyscontaining 0.53, 1.00, 1.93, 7.92 and18.7 mass% Cr

[1977Shc] Adiabatic calorimetry 173-673 K / Cr1.8Fe1.2S4

[1977Tre1] Calvet type calorimeter 25°C / Cr2FeS4 and Cr1.8Fe1.2S4

[1977Tre2] DTA, DTG, XRD, chemical analysis Up to 800°C / Cr2FeS4 andCr1.8Fe1.2S4

[1978Moh] XRD, metallography 700°C / Cr-Fe-S

[1978Rao] Gravimetry, metallography, SEM, EPMA,XRD, energy dispersive X-ray analysis

640-1100°C / CrS-FeS

[1979Cha] XRD, EPMA, metallography 800°C / Cr-Fe-S

[1979Los] XRD Up to 1060°C / Cr2FeS4

[1980Oud] Radioisotope method 850-1200°C / (Fe-18 mass% Cr) +H2/H2S

[1980Pet] XRD, DTA, magnetic electric resistivitymeasurements, Mössbauer spectroscopy

Up to 1100°C / Cr1–xFexS

[1981Rie] XRD, Mössbauer spectroscopy 77-300 K / Cr2–xFe1+xS4(0 ≤ x ≤ 0.5)

[1981Rug] Radioisotope method 850-1200°C/(Fe+18, 26 and 34mass% Cr)+ H2/H2S

[1982Fil] Radioisotope method 850-1200°C / (Fe-18 mass% Cr) +H2/H2S

[1982Ind1] XRD Room temperature / Cr2FeS4

[1982Ind2] XRD, DTA Up to 1600°C / Cr2S3-FeS

(continued)

8 Cr–Fe–S



[1982Zhu] Vapor pressure measurements 662-949°C / Cr2FeS4

[1982Nar] XRD, SEM, EPMA, metallography 800°C / (Cr-Fe) + H2/H2S

[1984Nar1,1984Nar2]

XRD, SEM, metallography 700, 800 and 900°C / (Fe-26.6 at.%Cr) + H2/H2S

[1985Nar] Auger spectrometry 200-300°C / (Fe + 25 mass% Cr) +H2S/Ar

[1985Sok] DTA, high-temperature XRD 30-1100°C / Cr2FeS4-Fe

[1986Bal] XRD, metallography, X-ray microanalysis 550, 660, 745 and 840°C / Cr2FeS4

[1987Nar] XRD, EPMA, metallography 800 and 900°C / Cr-Fe-S

[1988Fuj] XRD, EPMA, metallography, chemicalanalysis

1300°C / (Cr-Fe) + H2/H2S

[1993Cie, 1994Cie,1998Cie1, 1998Cie2,1998Cie3]

Mössbauer spectroscopy, XRD, EPMA,metallography

800°C / (CrxFe100–x at x ≤ 45.5) +H2/H2S

[1993Sok] XRD, Mössbauer spectroscopy Up to 1000°C / Cr2FeS4

[1995Smi] SEM, EPMA, diffusion couple technique 600°C / Cr-Fe-S

[1996Nog1,1997Nog]

Thermogravimetry 700-900°C / (Cr,Fe)1–xS

[1996Nog2] EPMA, thermogravimetry 750, 800 and 850°C / (Fe+25 at.%Cr) + S

[1996Smi] EPMA, metallography 600°C / Cr-Fe alloys containing11.3-51.1 mass% Cr + Fe0.95S

[1998Fed] XRD, vapor pressure measurements 900°C / Cr2FeS4

[1999Bou] Electrical resistivity measurements 600-1000°C / Cr2FeS4

[1999Che] XRD, TEM. XRPS, electron-spinresonance, Mössbauer spectroscopy

77-300 K / Cr2FeS4

[2000Kim,2005Nam]

XRD, Mössbauer spectroscopy 77-600 K / Cr0.025Fe0.975S

[2000Tsu] XRD, XRPS Room temperature / Cr2FeS4

[2000Yan1] XRD, magnetic and transport propertiesmeasurements

4.2-400 K / Cr2FeS4

[2000Yan2] XRD, electron-spin resonance, magneticmeasurements

77-300 K / Cr2–xFe1+xS4 (x = 0 and0.1)

[2001Kim1] Mössbauer spectroscopy, XRPS, magneticmeasurements

18-300 K / Cr2Fe0.92S4

(continued)

Cr–Fe–S 9




[2001Kim2,2002Kim1,2002Kim2]

Mössbauer spectroscopy, XRPS, neutronpowder diffraction, magneticmeasurements

10-300 K / Cr2Fe1–xS4 (x = 0.0,0.04, 0.08)

[2002Mit, 2003Mit] XRD, SEM, metallography Up to 1400°C / Fe + (0.3-18 mass%Cr) + (0.05-0.3) mass% S

[2002Nat] Mössbauer spectroscopy 90-200 K / Cr2FeS4

[2005Mer] XRD, TEM, EPMA, selected area electrondiffraction

27-70 K / Cr2FeS4

[2006Sro] Thermogravimetry, SEM, EDS, XRD,EPMA, metallography, chemical analysis,Mössbauer spectroscopy

800, 900 and 1000°C / (Fe-46 at.%Cr) + H2/H2S





Comments/References

α, (αCr,αFe,δFe)< 1863

(αCr)< 1863(δFe)1538 - 1394(αFe)< 912

cI2Im�3mW

a = 288.48

a = 293.15

a = 286.65

continuous solid solution [Mas2]

pure Cr at 25°C [Mas2]

pure Fe [Mas2]


γ, (γFe)1394 - 845

cF4Fm�3mCu

a = 364.67dissolves up to 11.9 at.% Cr [Mas2]


(βS)115.22 - 95.5

mP64P21/cβS

a = 1102b = 1096c = 1090β = 96.7°

[Mas2]

(αS)< 95.5

oF128FdddαS

a = 1046.4b = 1286.60c = 2448.60

pure S at 25°C [Mas2]

CrFe (σ sphase)830-440

tP30P42/mnmCrFe

a = 879.66 ± 0.06c = 455.82 ± 0.03

[Mas2, V-C2]

Cr1.03S< 597

mC8C2/c

a = 382.6b = 591.3

[Mas2, V-C2]

(continued)

10 Cr–Fe–S





Comments/References

Cr1.03S c = 608.9β = 101.6°

βCr2S3< 1250

hR30R�3βCr2S3

a = 593.84 ± 0.04c = 1666.83 ± 0.10a = 592.44 ± 0.04c = 1663.89 ± 0.11a = 592.24 ± 0.04c = 1664.08 ± 0.12

at 300 K [Mas2, V-C2]

at 80 K [Mas2, V-C2]

at 4.2 K [Mas2, V-C2]

αCr2S3< 1200

hP20P�31cαCr2S3

a = 593.9 ± 0.2c = 1119.2 ± 0.3

[Mas2, V-C2]

Cr3S4< 1152

mC14C2/mCr2S3

a = 1257.8b = 342.24c = 595.54β = 116.74°

[Mas2, V-C2]

Cr5S6< 327

hP22P�31cCr5S6

a = 596.2 ± 0.1c = 1150.6 ± 0.5

[Mas2, V-C2]

Cr5S8(I) mC52C2/mFeTi4S8

a = 1178.3 ± 1.0b = 678.6 ± 0.6c = 1106.3 ± 0.8β = 90.82 ± 0.02°

high-pressure phase [Mas2, V-C2]

Cr7S8< 317

hP4P�3m1LiTiTe2

a = 346.4 ± 0.1c = 576.3 ± 0.2

[Mas2, V-C2]

γFeS1188 - 315

hP4P63/mmcNiAs

a = 344.36 ± 0.05c = 587.59 ± 0.05

[V-C2, Mas2]

βFeS315 - ~60

Cr0.025Fe0.975S

hP24P�62cβFeS

a = 596.3 ± 0.1c = 1175.4 ± 0.1a = 586.1c = 1157.7 ± 0.1a = 599.8 ± 1.1c = 1171 ± 1a = 597.9 ± 0.2c = 1173 ± 5

at 21°C [V-C2, Mas2]

at 21°C and 3.33 GPa [V-C2, Mas2]

at 120°C [V-C2, Mas2]

[2000Kim, 2005Nam]

αFeS< 138

hP6P63/mmcαFeS

a = 345.59 ± 0.05c = 577.89 ± 0.05

[V-C2, Mas2]

FeS (I) oP8Pnma

a = 582.5 ± 0.2b = 346.8 ± 0.1

at 190°C [V-C2]

(continued)

Cr–Fe–S 11






Comments/References

MnP c = 693.5 ± 0.6a = 571.6 ± 0.9b = 334.7 ± 0.3c = 669.4 ± 0.9a = 565 ± 1b = 331.6 ± 0.3c = 663.1 ± 0.8

at 21°C and 4.15 GPa [V-C2]

at 21°C and 6.35 GPa [V-C2]

FeS (mackinawite) tP4P4/nmmPbO

a = 376.8c = 503.9

mineral mackinawite [V-C2]

βFeS2(h)743 - 444.5

cP12Pa�3FeS2(pyrite)

a = 541.79 ± 0.11a = 534.8 ± 0.2a = 529.3 ± 0.2a = 525.5 ± 0.2

mineral pyrite [V-C2, Mas2]at 1.57 GPa [V-C2]at 2.87 GPa [V-C2]at 3,85 GPa [V-C2]

αFeS2(r)< 444.5

oP6PnnmFeS2(marcasite)

a = 444.1b = 542.5c = 338.7a = 446.4b = 544c = 339

mineral marcasite [V-C2, Mas2]

at 327°C [V-C2]

Fe2S3 tP80P43212

a = 1053c = 1001

[V-C2]

Fe3S4(smythite)

hR21R�3mFe3S4

a = 347 ± 2c = 3450 ± 20

mineral smythite [V-C2]

Fe3S4 cF56Fd�3mAl2MgO4

a = 987.6 ± 0.2 [V-C2]

Fe7S8 hP45P3121Fe7S8

a = 686.52 ± 0.06c = 1704.6 ± 0.2

metastable phase, mineralpyrrhotite-3C [V-C2]

(continued)

12 Cr–Fe–S





Comments/References

λ, (Cr1–xFex)1–yS< 1595

λ1, γFeS(h)≲ 1188

λ2, CrS≲ 1565

hP4P63/mmcNiAs

a = 344.36 ± 0.05c = 587.59 ± 0.05

a = 341.9c = 555.0

0 ≤ x ≤ 10 ≤ y ≤ 0.18

dissolves up to ~5 at.% S and~22 at.% Cr at 700°C [1988Rag]at x = 1, y = 0 [V-C2, Mas2]

dissolves up to ~9 at.% S and ~ 5 at.% Feat 700°C [Mas2, 1988Rag]at x = 0, y = 0 [Mas2, V-C2, 2000Oik]

* τ, Cr2FeS4< 1350

cF56Fd�3mMgAl2O4

a = 997.5a = 999.8a = 999.5 ± 0.2

a = 998.3 ± 0.1a = 999 ± 1a = 999.6 ± 0.2a = 998.93 ± 0.08a = 998.9 ± 0.1a = 996.9 ± 0.4a = 999.3 ± 0.4a = 999.41 ± 0.07a = 999.25 ± 0.02a = 999.8 ± 0.1a = 997.56 ± 0.03a = 998.13 ± 0.03

[1944Lun][1956Lot][1964Shi, 1965Bou, 1966Rac, 1972Wat,1975Goe, 1976Tre, 1980Pet,1982Ind1, 1982Ind2][1968Tre2][1969Gib1][1969Gor1][1969Shi][1975Bab][1977Los][1979Los, 2000Tsu][1999Che, 2000Yan1, 2000Yan2]for Fe0.92Cr2S4 [2001Kim1][2001Kim2]at 10 K [2002Kim1, 2002Kim2]at 175 K [2002Kim1, 2002Kim2]

* τ´, Cr2FeS4 (I) h** a = 340c = 570a = 344.4c = 574.9

high pressure phase [1965Alb][1993Sok]

* τ´´, Cr2FeS4 (II) mC14C2/mCr3S4

a = 1127.1 ± 0.2b = 343.0 ± 0.1c = 594.8 ± 0.2β = 91.21 ± 0.02°a = 1147 ± 5b = 344 ± 1c = 594 ± 1β = 90.85 ± 0.05°

at 6.5 GPa [1967Bou]

at 5.5 GPa [1968Tre1, 1968Tre2]

Cr–Fe–S 13





Cr Fe S

L2 + λ ⇌ Cr2FeS4 + L3 ~ 1235 U1 Cr2FeS4 28.57 14.29 57.14

L2 + Cr2FeS4 ⇌ λ + L3 ~ 1225 U2 Cr2FeS4 28.57 14.29 57.14

λ + L3 ⇌ Cr2FeS4 + Cr2S3 ~ 1070 U3 Cr2FeS4Cr2S3

28.5740.00

14.290

57.1460.00

α + L2 ⇌ λ + γ ? U4 - - - -

λ + L3 ⇌ Cr2FeS4 + FeS2 ~ 740 U5 Cr2FeS4FeS2

28.570

14.2933.33

57.1466.67

λ1 + λ2 ⇌ Cr2FeS4 + α ~ 650 U6 Cr2FeS4 28.57 14.29 57.14

L3 + Cr2FeS4 ⇌ Cr2S3 + FeS2 ~ 500 U7 Cr2FeS4Cr2S3FeS2

28.5740.000

14.290

33.33

57.1460.0066.67

L3 ⇌ (S) + Cr2S3 + FeS2 ~ 114 E (S)Cr2S3FeS2

040.000

00

33.33

~ 10060.0066.67



Quantity, per mole of atoms[kJ, mol, K]

Comments

2Cr + Fe + 4S ⇌ Cr2FeS42Cr + Fe + 4S ⇌ Cr2FeS4

2525

ΔH°= (4070 ± 100)·ΔH°FeS–(360830 ± 8370)ΔH° = –457300 ± 7900

[1974Tre, 1976Tre][1976Kes, 1977Tre1,1977Tre2]

1.8Cr + 1.2Fe + 4S ⇌Cr1.8Fe1.2S4

25 ΔH° = –498550 ± 8790 [1977Tre1, 1977Tre2]

Table 5. Vapor Pressure Measurements

Phase(s) Temperature [°C] Pressure [Pa] Comments

Cr2FeS4 662 - 949 (–6210 ± 315)/T + (8.28 ± 0.30) [1982Zhu, 1998Fed]

14 Cr–Fe–S


Fig. 1. Cr-Fe-S. The CrS-FeS quasibinary section

Cr–Fe–S 15



Fig.2.

Cr-Fe-S.

Reactionscheme.λ=(Fe,Cr)1-xS

16 Cr–Fe–S


Fig. 3a. Cr-Fe-S. Liquidus surface projection

Cr–Fe–S 17



Fig. 3b. Cr-Fe-S. Probable liquidus surface projection in the S corner (schematic)

18 Cr–Fe–S


Fig. 4. Cr-Fe-S. Isothermal section at 1300°C

Cr–Fe–S 19




20 Cr–Fe–S



Cr–Fe–S 21




22 Cr–Fe–S


Fig. 8. Cr-Fe-S. Calculated vertical sections at 0.001-0.1 mass% S

Cr–Fe–S 23



Fig. 9. Cr-Fe-S. Calculated vertical sections at 0.5-20 mass% S

24 Cr–Fe–S


Fig. 10. Cr-Fe-S. The CrS-Fe vertical section

Cr–Fe–S 25



Fig. 11. Cr-Fe-S. Phase relations and stability fields at 800°C as a function of sulfur pressure

26 Cr–Fe–S


Fig. 12. Cr-Fe-S. Phase relations and stability fields at 900°C as a function of sulfur pressure

Cr–Fe–S 27



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28 Cr–Fe–S


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[1975Bab] Babaev, G.Yu., Kocharov, A.G., Ptasevich, Kh., Yamzin, I.I., Vinnik, M.A., Saksonov, Yu.G., Alferov, V.A., Gordeev, I.V., Tretyakov, Yu.D., “Magnetic and Neutron-Diffraction Stu-dies of the Sulphospinels Cu0.2Fe0.8Cr2S4 and Fe1.2Cr1.8S4”, Sov. Phys.-Crystallogr. (Engl.

Cr–Fe–S 29



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30 Cr–Fe–S


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[1984Nar1] Narita, T., Smeltzer, W.W., “Sulfidation Properties of an Fe-26.6 at.% Cr Alloy at Tempera-tures of 973-1173 K in H2S-H2 Atmospheres at Sulfur Pressures 104-10–6 Pa”, Oxid. Met., 21(1-2), 39–55 (1984) (Experimental, Kinetics, Morphology, Phase Relations, 8)

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Cr–Fe–S 31



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[1998Cie1] Cieslak, J., Dubiel, S.M., Zurek, Z., “Investigation of Scales Resulted from a High-TemperatureSulphidation of Fe-Cr Alloys”, J. Alloys Compd., 265(1-2), 297–304 (1998) (Experimental,Kinetics, Interface Phenomena, 12)

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32 Cr–Fe–S


[2000Oik] Oikawa, K., Mitsui, H., Ohtani, H., Ishida, K., “Thermodynamic Calculations of Phase Equi-libria in the Fe-Cr-S System”, ISIJ Int., 40(2), 182–190 (2000) (Calculation, Phase Relations,Thermodyn., #, *, 37)

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[2003Mit] Mitsui, H., Oikawa, K., Ohnuma, I., Ishida, K., “Microstructural Evolution of Sulfide inFe-Cr-S Alloys”, Mater. Sci. Forum, 426–432, 993–998 (2003) (Experimental, Mechan.Prop., Morphology, Phase Relations, 20)

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[2005Mer] Mertinat, M., Tsurkan, V., Samusi, D., Tidecks, R., Haider, F., “Low-Temperature StructuralTransition in FeCr2S4”, Phys. Rev. B: Condens. Matter, 71(10), 100408-1-100408–4 (2005)(Experimental, Crys. Structure, Morphology, Phase Relations, 33)

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[2006Sro] Sroda, S., Zurek, Z., Dubiel, S.M., “Effect of Tin on High-Temperature Sulphidation of Fe-Cr Alloys in H2/H2S Atmospheres”, Corros. Sci., 48(3), 727–740 (2006) (Experimental,Kinetics, Morphology, Phase Relations, 14)



Cr–Fe–S 33



Chromium – Iron – Silicon

Andy Watson, Marina Bulanova, Jean-Claude Tedenac

Introduction

This system has an industrial and technological importance for applications in commercial alloyed steels(iron based) and special applications in high temperature thermoelectric materials. This system has beenreviewed on five occasions, [1981Nis, 1987Rag, 1993Rag, 2003Rag, 2004Rag]. In these reviews, theauthors give a reaction scheme in the Cr-Fe rich part of the system and some isothermal sections for the tem-perature range 427-1127°C where the solid state transformations take place. Since the first experimentalwork of [1926Den], many publications have presented results of phase equilibria determined by thermalanalysis, X-Ray diffraction and micrograph studies [1936And, 1936Jet, 1937Sve, 1938And, 1942Kur,1949Kur, 1949Jae, 1956Gla1, 1956Gla2, 1957Aro].The phase equilibria have been calculated by [1980Cha] based on earlier experimental phase diagram stu-dies. This work was extended by [1985Ans]. Finally, according to these results [1993Rag] revised the reac-tion scheme previously reported in [1987Rag]. [1988Har] computed the (αFe)(bcc)/(γFe)(fcc) equilibriumat 6 temperatures in the iron rich region and these results are consistent with [1985Ans]. [1997Lin] has reas-sessed this ternary using unpublished new data. [1998Mie, 1999Mie] remodeled the solution phases usingnew thermodynamic data.Details of experimental studies of phase equilibria and thermodynamics are given in Table 1.

Binary Systems

A Calphad assessment of the Fe-Si system was performed by [1991Lac] incorporating B2 ordering for thebcc phase. The results were very good and agreed with the phase diagram given by [Mas2]. However, D03 isalso a feature of this phase but was not included in the modeling. [1998Mie] made some adjustments to themodeling of the liquid, fcc and bcc phases in order to improve the fit of experimental data for the C-Fe-Sisystem, and these improved data were used in a modeling of the Cr-Fe-Si ternary system a little later. As theversion of the Fe-Si is most complete in [Mas2], this is accepted for the purposes of this assessment.The Cr-Fe system has been assessed by [1987And]. However, there are three major differences between theassessed phase diagram and that given in [Mas2]. The Cr rich liquidus and solidus are at much higher tem-peratures in the diagram of [1987And]. This is explained by the fact that the experimental work on whichthis part of the diagram is based suffered from problems of oxidation and Cr vaporization. This is underlinedby the fact that the melting point for pure Cr given in [Mas2] (1863°C) is over 40°C lower than that in[1987And] (1907°C). This part of the phase diagram was improved in the assessment performed by[1993Lee], which also brought the melting minimum closer to the experimentally determined value. Theremaining discrepancies are associated with the location of the σ phase field. The lower composition limitof the σ phase is different in the assessments, 49 at.% Cr in [Mas2] and 51.4 at.% Cr in [1987And]. Also,the limiting temperature differs, 440°C in [Mas2], 512°C in [1987And]. There is a difference in the upperlimit of the σ phase, the composition of the congruent transformation to bcc. In [Mas2], this is given as47 at.% Cr, whereas in [1987And] this point is given as ~45.5 at.% Cr. The reason for the shift in the com-position range of the σ phase in the Calphad assessment was due to the requirement to describe the experi-mental heat of transformation from the σ to the bcc phase. As the liquidus and solidus of the system arebetter described in [1993Lee], the preferred phase diagram is a composite of this and the assessment by[1987And]. This version of the phase diagram is given in Fig. 1. The Cr-Si system is taken from the assess-ment provided by the MSIT binary assessment program [2007Leb].

Solid Phases

The crystal structures of the relevant phases in the Cr-Fe-Si ternary system are presented in Table 2.

Cr–Fe–Si 1



[1936Jet, 1936And, 1938And] studied the phase repartition of the α/σ phases on samples after quenchingand annealing. They concluded that the brittleness of the samples with Si contents of over 3 mass% was dueto σ phase precipitation.The σ phase in the Cr-Fe-Si was investigated by [1957Aro] who studied ten alloy compositions afterannealing at different temperatures followed by rapid quenching. The lattice parameters were reported asa function of Si content. This phase is tetragonal with 30 atoms per unit cell, lattice constants are summar-ized in Table 2.The kinetic of precipitation of the σ phase induced by deformation was investigated by [1980Sau], a modelis deduced from previous published experimental results. Later [1981Sau1, 1981Sau2] studied experimen-tally the kinetics of precipitation using metallography for samples treated at 650-800°C. In this study thevolume fraction of precipitates is given as a function of aging time. It is shown that after 100 h the σ contentremains constant.The composition (Cr1–xFex)3Si with x = ±0.2 falls in the ordered D03 region of the α phase field. [1975Pic]studied these two compositions by NMR and concluded that Cr substitutes for iron in with 8 Fe nearestneighbors in the D03 lattice of Fe3Si. A new ternary compound, (Cr0.26Fe0.74)3Si2 has been prepared by[1994Mot] using a vapor deposition process, the crystallographic characteristics have been determinedand presented Table 2. Electronic structures of chromium doped ξL, FeSi2 phase were investigated by[2002Tan] using first principles pseudopotential calculations (GGA).The CrSi and FeSi phases are isostructural having the cubic B20 structure. They have been found to form acomplete series of solid solutions [1936And, 1942Kur, 1949Kur, 1956Gla1, 1956Gla2, 1962Bur, 1966Gla,1968Sid, 1980Kav]. Figure 2 shows the lattice parameter variation of this phase as a function of Cr content.

Quasibinary Systems

The section FeSi-CrSi has been shown to be quasibinary, through lattice parameter measurements along thissection [1936And, 1942Kur, 1949Kur, 1956Gla1, 1956Gla2, 1962Bur, 1966Gla, 1968Sid, 1980Kav].However, this section can only be partially quasibinary as there is no maximum on the monovariant liquidusline (L + Cr5Si3 ⇌ ε) that crosses this section. Alternatively, [1961Dub, 1962Dub] found the section FeSi2-CrSi2 to be quasibinary with a eutectic reaction at 18 mol% CrSi2 and 1150°C. But again, this can only bepartially quasibinary, as the high temperature variant of the FeSi2 phase dissociates into the lower tempera-ture variant and silicon below 937°C. Also, according to [Mas2], the composition of the congruent meltingof FeSi2 is approximately 1 at.% richer in Si than the composition of the eutectoid decomposition of thephase. The section presented by [1962Dub] is reproduced in Fig. 3.


Only limited experimental studies of the invariant equilibria and liquidus surface have been made[1926Den, 1973Wet]. Based on these data and invariant reaction temperatures in the binary systems, a ten-tative reaction scheme was constructed by [1987Rag], shown in Figs. 4a and 4b, although this was laterupdated in [1993Rag] based on the calculations by [1985Ans]. The binary Cr-Si phase diagram used by[1987Rag] gave quite a different temperature for the peritectic reaction L + αCr5Si3 ⇌ CrSi from the valuein the currently accepted diagram (1475°C in [1987Rag], 1424°C in [2007Leb]). This leads to a change ofreaction type in the ternary system for the reaction U2, from transition to ternary peritectic (P1), the acceptedbinary invariant reaction temperature being lower than that for the ternary invariant reaction. In any case, itshould be stressed that this section is only tentative. The invariant reactions are given in Table 3. Approx-imate liquid compositions are taken from the liquidus surface, the compositions of the solid phases areunknown and hence are omitted from the Table 3.


There have only been two experimental studies of the liquidus surface. [1926Den] studied 18 compositionsin the Fe-30 at.% Cr-30 at.% Si region of the system by thermal analysis. Electrolytic Fe, Cr (98.23% Cr,(Fe+Al) 1.12%) and Si (97.63% Si, 1.82% (Fe+Al)) were used. Later, [1973Wet] studied the liquidus tem-peratures of 21 ternary alloys in the region 20-50 mass% Cr, 10-30 mass% Fe, 25-55 mass% Si by DTA.

2 Cr–Fe–Si


A classical quenching technique was also employed, which involved quenching samples from above andbelow the liquidus and examining the samples microscopically. The experimental results of both studieswere combined in a tentative liquidus surface by [1987Rag] (Fig. 5a), modifications were made to ensureagreement with the binary systems, particularly the Fe-Si system where it was necessary to add a tentativeliquidus line corresponding to the peritectic formation of αCr5Si3 at 1519°C. Also, as discussed above, thebinary Cr-Si phase diagram used by [1987Rag] has quite a different temperature for the peritectic reactionL + αCr5Si3 ⇌ CrSi from the value in the currently accepted diagram leading to a change of reaction type inthe ternary system, from transition U2 to ternary peritectic P1. Therefore, the monovariant line drawn fromthe Cr-Si binary edge rises with respect to temperature as it extends into the ternary system, rather than fall-ing as given in [1987Rag].More recently, [1997Lin] calculated a liquidus surface of part of their Calphad assessment of the system(Fig. 5b). The main feature of this work was to model the σ phase and the extension of the Fe5Si3 phasein the ternary system following experimental work conducted by [1995Hay]. As a consequence, the liquidussurface featured a primary solidification surface for the σ phase, and also the Fe5Si3 phase. There is noexperimental justification for these features. The later work of [1999Mie], which involved a more thoroughmodeling of the liquid, bcc and fcc phases gave similar results for the liquidus surface as [1997Lin]. But thisis not so surprising, as the modeling of the other phases was taken from the earlier work.

Isothermal Sections

Figures 6 to 10 show isothermal sections for 427, 527, 900, 1047 and 1150°C. The earliest comprehensivestudy of the phase equilibria of the system was undertaken by [1936And], who studied the phase composi-tion of 70 alloys in the region up to 50 at.% Cr and 35 at.% Si and at temperatures of 600, 800 and 1000°C.Samples were melted under vacuum in a high frequency furnace in alundum crucibles. The samples werehomogenized in evacuated silica tubes before grinding to powder for heat treatment. This took place, again,in evacuated silica tubes in a vertically mounted furnace that allowed the samples to be quenched after heattreatment. The duration of the annealing was between 16-48 h. Following chemical analysis and XRD stu-dies, phase boundaries were determined for the three temperatures in the Fe corner of the phase diagram.However, on their composite isothermal section for the temperatures, there is a phase field denotedα3+(Fe,Cr)3Si2, suggesting the possibility of complete miscibility between Fe3Si2 and Cr3Si2. However,these phases do not exist in the binary systems. The phases closest to these compositions are Fe5Si3 andCr5Si3, but these have different crystal structures precluding complete mutual solid solubility. Other thanthis feature, the α+σ phase boundaries are in good agreement with later work.[1966Gla] prepared 120 alloys of different compositions and determined an isothermal section for 900°C.Carbonyl Fe, electrolytic Cr and Si (all 99.9%) were melted under helium before annealing for 400 h at900°C. XRD revealed complete solubility between FeSi and CrSi, but only partially between Fe5Si3 andCr5Si3, and between CrSi2 and FeSi2. The section was produced in [1987Rag]. Later, these results wereused in a ternary assessment of the system by [1980Cha] (for Si contents less than ~43 at.%), their sectionfor 900°C being in very good agreement with the earlier work. The major difference between the calculatedand experimental section being the size of the α+Cr3Si phase field. Figure 8 presents the isothermal sectionfor 900°C. It comprises the calculated diagram of [1980Cha] along with equilibria for Si contents greaterthan 43 at.% taken from [1987Rag], based on the work of [1966Gla]. [1980Cha] used their modeling to cal-culate further partial isothermal sections. These are shown in Figs.6 (527°C) and 7 (427°C).Unpublished experimental data for the stability of the σ phase [1995Hay] were used by [1997Lin] to assessthe system. The assessment used all available phase equilibrium data together with experimental thermody-namic data from the literature. Isothermal sections for 1047 and 947°C were calculated and agree well withthe experimental data from [1995Hay], although the fit of the calculated enthalpies of mixing do not fit theexperimental data of [1975Igu] particularly well.The section for 927°C agrees well with Fig. 8. The section for 1047°C from [1997Lin] is shown in Fig. 9.In each of the presented isothermal sections, amendments were made to ensure agreement with the acceptedbinary phase diagrams. There was a significant shift in the phase boundaries associated with the Cr-Fe sys-tem, particularly with respect to Figs. 6 to 8. Also, the B2 and D03 ordering were ignored in all of the sec-tions. In Fig 9, it was necessary to add tentative equilibria involving the Fe2Si phase. At 1047°C, this phase

Cr–Fe–Si 3



is just stable in the accepted Fe-Si phase diagram given by [Mas2], whereas in the calculated phase diagramused by [1997Lin] to produce this section, it is just unstable. Also, the FeSi phase is treated here as stoichio-metric, whereas in reality it has a small homogeneity range of approximately 1 at.%. It was also necessary toadd phase boundaries for the ‘γ-loop’, which had been omitted in the original work. Figure 10 shows a cal-culated section for 1150°C, also from [1997Lin]. There are no experimental data to support this section, andit is interesting to note that not only the liquid phase becomes stable at this temperature, but the Fe5Si3 phaseis predicted to be quite stable well into the ternary system, even though it is not stable in the binary Fe-Sisystem at this temperature. This is a feature that requires further investigation as it would seem to be a resultof the modeling and may not reflect reality.It was necessary to add a line to this section to create a two-phase region, Cr3Si+η to maintain consistencywith the vertical section taken from the same publication that is reproduced here in Fig. 17.


[1937Sve] studied phase equilibria in the region of the γ loop using a combination of thermal analysis andmetallography. Partial isopleths for Cr contents of 1, 4, 7.3 and 10.2 mass%, and for Si contents of 0.5 and 1mass% Si were produced (Figs. 11 to 16). [1942Kur, 1949Kur] studied sections at 15 and 25 mass% Fe.Alloys were melted from electrolytic Fe and Cr along with relatively impure Si (97%) in a high frequencyfurnace and thermal arrests detected on cooling. The liquidus points are in reasonable agreement with theaccepted liquidus surface, but the sections themselves are far from complete, and are therefore not repro-duced here. Figure 17 shows an isopleth for xFe:xCr = 2:3 as calculated by [1997Lin] (this section is mis-takenly labelled as xFe:xCr = 3:2 in the original article). It should be noted that in this section, there is nodistinction between the α and β forms of Cr5Si3 as this was not modeled in the Cr-Si binary description usedin the calculation.

Thermodynamics

Experimental thermodynamic data for ternary Cr-Fe-Si alloys are limited. [1968Che] report the activity ofsilicon in Cr-Fe-Si melts along sections with 12, 18 and 25 % Cr at 1600°C. It was shown that the activityexhibits a negative deviation from ideal behavior. The dependence of log10γSi on chromium concentration islinear, and extrapolation of the relationship to zero chromium content yields a value of about 0.0027.[1975Igu] studied the partial molar enthalpies of the components and the integral molar enthalpies of mixingin ternary Cr-Fe-Si alloys at 1600°C by calorimetry. Isoenthalpy of mixing lines were constructed. In[1978Pet], these were calculated for this temperature based on thermodynamic data for the binary systems.The results of the two studies are in good agreement. According to [1975Igu], a decrease in the xFe/xCr ratioresults in a decrease in the energy of interaction between the Si atoms and the matrix. Partial enthalpies ofmixing of Si in liquid Fe-Cr alloys were determined by [1977Ost] by high temperature calorimetry. Theenthalpy of mixing of a single composition of Fe and Cr in liquid Si was measured by [1960Gel] using hightemperature calorimetry.The activity of Si in liquid alloys was modeled by [1990Dre] and the whole system by [1997Lin] and laterby [1999Mie]. The latter work incorporated improved modeling for the liquid, bcc and fcc phases andresulted in a better agreement with the experimental data given by [1975Igu] than found with the work of[1997Lin]. Figures 18 to 21 show the experimental integral enthalpies of mixing of liquid alloys along withthe calculated curves for xFe/xCr = 4.55, 1.86, 0.93 and 0.464, respectively, at 1600°C taken from[1999Mie]. The modeling by [1998Mie] was used in the prediction of liquidus curves in steels by[2000Mie].[1977Si] studied the specific heat of Cr-Fe-Si alloys with 3 at.% Si at temperatures between 1.4 and 4.2 K.The electronic specific heat coefficient γ was shown to depend on electronic concentration, suggesting theelectrons transfer from Si to the d-metals.[1969Bau] has shown that the wetting angle and the work of adhesion at 1700°C with respect to Si concen-tration have a maximum and a break, respectively, at 45-50 at.% Si.

4 Cr–Fe–Si



The FeSi2 intermetallic compound has been studied for thermoelectric applications at high temperatures.Owing to its peritectoid transformation at 982°C, applications are limited to that temperature with a maxi-mum of efficiency at ~550°C. The introduction of chromium as a substitute for iron in the structure leadsto a p-type semiconductor. [2003Kim] studied the compositions FexCr1–xSi2 (with x = 0.01-0.1).[1999Tan] measured the Hall effect and electrical conductivity for compositions of Fe1-xCrxSi2 ranging fromx = 0.01 to 0.05 at low temperatures and obtained a temperature dependence for the carrier mobility, thebehavior of which was explained by using the small polaron model. The understanding of such physicalproperties has been enhanced by first principles calculations made by [2002Tan] using the DFT methodin a GGA approximation, the DOS curves as a function of energy are given along with the position ofthe Fermi level.[2004Aru] reported on the magnetic properties on Cr doped (~1 at.%) FeSi2 concluding spin glass behavior.Nanocrystallization has been reported in this system by [2005Tur]. Samples of the compositionFe78.5Si20Cr1.5 were studied by NMR and X-Ray diffraction in order to determine the grain size of the nano-crystals.The low temperature specific heat of ten ternary alloys with 3 at.% of Si have been investigated by [1977Si],the chromium content ranging from 0.049 to 0.97 at.%. The Debye temperature was obtained by interpola-tion between the two limits of the solid solution.[1995Zho1, 1995Zho2] prepared Fe-Si/Cr multilayers by RF sputtering and studied the magneto-optic prop-erties of these multilayers.The magnetic susceptibility of alloys having compositions Cr2FeSi2 and Cr2FeSi3 have been measured by[1978Hed]. The susceptibility was found to follow a Curie law.The mechanical properties of sintered Cr-Fe-Si alloys were studied by [1974Sie]. The materials were pre-pared by in-situ decomposition of Si3N4, Cr2N and CrSi2 in an iron matrix, all taking place in either aNH3 or H2 atmosphere, for temperatures between 1000 and 1300°C for times between 1 and 8 h. Tensilestrength, elongation and hardness were measured as a function of tempering time, temperature and theresulting sintered density. The tensile strength was found to be in the region of 525 MPa with an elongationto fracture of 3 % at a sintered density of 6700 kg/m3.Owing to the brittle nature of the σ phase, its formation in steels is of significant importance. [1965Sch] stu-died the stability of the σ phase in ferritic chromium steels containing 13, 18 and 24 mass% Cr with siliconand aluminium at temperatures up to 900°C. The deformation behavior was studied with respect to theamount and distribution of the σ phase in a Fe-30 mass% Cr-3.5 mass% Si alloy, [1982Sau]. The volumefraction of σ phase in such an alloy was in the region of 47%, and the presence of the σ phase resultedin hard material with a small impact energy. This was found to decrease, with a corresponding increasein strength, as the σ content in the material increases on heating. It was also found that the strength increasewas less for finer distributions of the σ phase.Details of studies of materials properties are given in Table 4.

Miscellaneous

Oxidization and surface morphologies are reported in [1981Min]. Using air, H2/H2O, and H2/H2O/H2Senvironments at a fixed temperature (1200°C, 24 h) and partial pressure (p(O2) = 1.0-1.8 MPa, p(S2) =1.0-6.0 MPa) they were able to explain the role of silicon in the formation of the oxide layer.Measuring carbon activity by means of diffusion couples, [1991Nis] calculated the interaction parameters ofC in the quaternary system C-Cr-Fe-Si. Samples were prepared at 900°C and held for 7 h. Equilibrium wasmaintained over three different heat treatment cycles (from 4 days at 1100°C to 2 weeks at 900°C). Theresulting activities were: aC = 0.14-0.24, (900°C), 0.01 (1000°C), 0.043(1100°C).Corrosion studies of the CrFeSi compound in an industrial galvanizing bath (containing aluminium) wereconducted by [2005Liu]. During these experiments it was shown that a layer of Fe-Cr-Al alloys was formedat the surface conferring good resistance against corrosion.The interaction of the alloys with different gases has been investigated. The solubility of hydrogen in liquidalloys has been determined over the temperature range 1500-1700°C. It follows Sievert’s law, [1971Blo].

Cr–Fe–Si 5



The “Siliconization” of the surface of stainless steel at 1100°C entails the growth of tube-shaped columnarCr-Fe-Si compounds of thicknesses in the region of 50-100 μm, at a growth temperature of 1000°C for30 min. Analysis suggests a composition of (Cr1–xFex)5Si3, [1989Mot].An experimental study on the oxidation of alloys containing 0-9 mass% Cr and 0-1 mass% Si in carbondioxide at 500°C has been reported by [1982Mos]. Several experimental techniques were used for thisinvestigation (X-Ray, TEM, SEM, Photoelectron microscopy), they studied the silicon distribution in thealloys and compared the results with theoretical predictions by an oxidation model. An attempt to growmaterials with a eutectic composition in this ternary system by directional solidification was made by[1978Hao]. However, the alloys were found to have no eutectic.High Cr cast irons are reviewed briefly in [1996Rey].

Table 1. Investigations of the Cr-Fe-Si Phase Relations, Structures and Thermodynamics


[1926Den] Thermal analysis all temperature range/ subternary Fe-30 at.% Cr-30at.% Si

[1936And] X-Ray diffraction/powder method Normal temperature, 17.44 - 23.36 mass% Cr,68.04 - 61.52 mass% Fe, 14.52 - 15.12 mass% Si

[1936Jet] X-Ray diffraction/powder method Room temperature, 0.006 - 0.086 mass% Si

[1937Sve] Metallography, thermal analysis. Isopleths in the γ loop region at 1, 4, 7.3, 10.2mass% Cr, and 0.5, 1 mass% Si

[1938And] Optical micrography, Brinell hardness Room temperature after quenching at 100°C adageing at 610-820°C, 0.05 - 1.51 mass% Si,2.28 - 57.28 mass% Cr

[1942Kur] Thermal analysis, metallography as-cast, 10 - 25 at.% Fe, 25 - 70 at.% Si

[1949Kur] Thermal analysis, metallography 1000°C, 1215°C, 10 - 25 at.% Fe, 25 - 70 at.% Si

[1956Gla1,1956Gla2]

Microscopy, X-ray diffraction FeSi2-CrSi2 Cr5Si3-Fe5Si3 sections

[1957Aro] Crystal structure/powder diffraction After quenching from 950°C/(Cr,Fe)Si0.68/ two-phase sample: D88+T1

[1958Aro] Crystal structure/powder diffraction Room temperature/FeCr and FeCrSi

[1961Dub] Microscopy/X-ray diffraction Mutual solubility in the FeSi2-CrSi2 section

[1962Gla,1966Gla]

X-ray diffraction/ powder method/lattice constant measurements

Along the section CrSi-FeSi

[1968Sid] X-ray diffraction/ powder method Along the section CrSi-FeSi

[1971Yam] DTA experiment, Mössbauerabsorption

Arc melted/annealed 80 days at 480°C

[1973Wet] DTA, quenching of samples fromtemperatures above and belowliquidus

Liquidus temperatures, 20 - 50 mass% Cr,10 - 30 mass% Fe, 25 - 55 mass% Si

[1975Pic] Neutron diffraction Fe2.88Cr0.12Si

[1977Ost] High temperature calorimetry Partial enthalpies of mixing of Si in Fe-Cr alloys

(continued)

6 Cr–Fe–Si



[1978Hed] Magnetic susceptibility 0 - 400 K/ Cr2FeSi2 and Cr2FeSi3

[1980Sau] Metallography/optical microscopy Precipitation of σ phase

[1981Min] Oxidation/ air furnace 1200°C, 24 h, 18 mass% Cr, 0.5 - 2 mass% Si

[1981Sau1,1981Sau2]

Metallography/optical microscopy Precipitation of σ phase between 650-800°C/3samples

[1986Vos] Dilatometric analysis, XRD Lattice parameter and linear expansion coefficientsof alloys with up to 9 at.% Cr and 12 at.% Si

[1989Mot]. Vapor deposition/Si2Cl6 /CVD (Cr1−xFex)5Si3, (CrFe)3Si2

[1991Nis] Diffusion couples/carbon analysis Treated at 900°C during 7h, 3 mass% Si, 0.37,0.69, 1.04 mass% Cr

[1994Mot] Vapor deposition/Si2Cl6 /CVD (CrFe)3Si2

[1995Zho1,1995Zho2]

Multilayers deposition-magneto-optic/RF sputtering-Kerr effect

Normal temperature / Fe-Si/Cr layers

[2004Wan] Microscopy TEM, microprobeanalysis, EDS profiles

Grain boundary precipitation in a 12.64 mass%and 2.8 mass% Si

[2004Yam] Microscopy (SEM, TEM),microprobe analysis

1173K, α/D03 partition in the iron rich part





Comments/References

α, (δFeαFe,Cr)

(Cr)< 1907

(δFe)1538 - 1394(αFe)< 912

cI2Im�3mW a = 288.48

a = 288.47 ± 0.05a = 288.09 ± 0.02a = 288.0a = 293.15

a = 286.65a = 284

Pure Cr at 25°C [Mas2]4.5 ± 0.2 at.% Si at 1200°C [2007Leb]6.7 ± 0.2 at.% Si at 1400°C [2007Leb]10 at.% Si [V-C2]Pure Fe at 1390°C [Mas2]

Pure Fe at 25°C [Mas2]78.5Fe-1.5Cr-20Si at 576°C [2005Tur]Dissolves up to 19.5 at.% Si.

(γFe)1394 - 912

cF4Fm�3mCu

a = 364.67 at 915°C [Mas2, V-C2]Dissolves up to 11.9 at.% Cr and 3.8 at.%Si [Mas2].

(continued)

Cr–Fe–Si 7






Comments/References

εFe hP2P63/mmcMg

a = 246.8c = 396.0


(Si)< 1414

cF8Fd�3mC (diamond)

a = 543.06 at 25°C [Mas2]

σCrFe830 - 512

(CrFe)1–xSix

tP30P42/mnmCrFe

a = 879.66 ± 0.06c = 455.82 ± 0.03a = 879.68 ± 0.05c = 455.85 ± 0.03a = 879.6 ± 0.04c = 456.05 ± 0.03a ≈ 876c ≈ 458

[Mas2]

[1987And]49.5 at.% Cr [V-C2]

48.2 at.% Cr [V-C2]

47.8 at.% Cr [V-C2]

x = 0.15 at 690°C [1957Aro].Values are taken from the figure.Stable to at least 1047°Cin the ternary system [1997Lin].

Cr3Si< 1780

cP8Pm�3nCr3Si

a = 456.5 to 455.2

a = 456.27 ± 0.04a = 456.46 ± 0.02a = 456.67 ± 0.02a = 456.65 ± 0.03

[V-C2]20.8 - 25.3 at.% Si22.5 ± 0.4 at.% Si at 1200°C21.5 ± 0.4 at.% Si at 1400°C20.8 ± 0.4 at.% Si at 1600°Cas solidified [2007Leb]

βCr5Si31666 - 1488

- - 37.5 - 37.7 at.% Si [2007Leb]

αCr5Si3< 1488

tI32I4/mcmW5Si3

a = 917.0 ± 0.6c = 463.6 ± 0.3

[V-C2]

ε, Cr1–x,FexSi

CrSi< 1424FeSi< 1410

cP8P213FeSi a = 462.2 ± 0.1

a = 450.0a = 451.7 ± 0.5

0 ≤ x ≤ 1

[V-C2]

at 25°C [V-C2]at 300°C [V-C2]

CrSi2< 1438 ± 2

hP9P6222CrSi2

a = 442.83 ± 0.01c = 636.80 ± 0.09

66.3 – 68 at.% Si [2007Leb]

(continued)

8 Cr–Fe–Si





Comments/References

α2

< 1280cP2Pm�3mCsCl

Ordered B2 structure. ~10-22 at.% Si[Mas2]

α1, Fe3Si< 1150

cF16Fm�3mBiF3

a = 565.0a = 564.4

[V-C2]Ordered D03 structure. ~10-30 at.% Si[Mas2]

β, Fe2Si (h)1212 - 1040

hP6P�3m1Fe2Si

a = 405.2 ± 0.2c = 508.55 ± 0.03

[V-C2]

η, Fe5Si3 (h)1060 - 825

hP16P63/mcmMn5Si3

a = 674.16 ± 0.06c = 470.79 ± 0.06a = 675.9 ± 0.5c = 472.0 ± 0.5

[V-C2]

ξH, Fe0.92Si2 (h)1220 - 937

tP3P4/mmmFe0.92Si2

a = 269.01 ± 0.01c = 513.4 ± 0.2

[V-C2]

ξL, FeSi2 (l)< 982

(Fe1–xCrx)Si2

oC48CmcaFeSi2

a = 986.3 ± 0.7b = 779.1 ± 0.6c = 783.3 ± 0.6a = 988.2b = 780.5c = 783.2a = 988.6b = 780.4c = 783.5a = 988.4b = 780.6c = 783.2 ± 0.6a = 984.3b = 781.8c = 787.6

[V-C2]

[1999Tan] x = 0.01Cr-doped FeSi2



[2002Tan] x = 0.125Ab-initio calculations

*τ, (Cr,Fe)3Si2

metastable

P63/mcmD88Mn5Si3

a = 680c = 471

[1994Mot] for (Cr0.26Fe0.74)3Si2composition. Obtained innon-equilibrium conditions.

Cr–Fe–Si 9





Cr Fe Si

L + Cr3Si ⇌ α + βCr5Si3 ~1625 U1 L 57 7 36

L +βCr5Si3 + α ⇌ ε ~1430 P1 L 31.7 20.7 47.6

L ⇌ α + ε + Fe2Si ~1170 E1 L 4.8 60 35.2

L ⇌ ξH + CrSi2 ~1150 e9(max) L 6 23.5 70.5

L ⇌ ξH + CrSi2 + ε ~1125 E2 L 7.1 27.7 65.2

L ⇌ ξH + CrSi2 + (Si) ~1120 E3 L 6.6 20.8 72.6

ε + ξH ⇌ α + Fe5Si3 ~1050 U2 - - - -

ε + ξH ⇌ ξL + CrSi2 ~960 U3 - - - -

ξH + CrSi2 ⇌ ξL + (Si) ~950 U4 - - - -

ε + α ⇌ Fe5Si3 + Cr5Si3 ? U5 - - - -

α + Cr5Si3 ⇌ Fe5Si3 + Cr3Si >900 U6 - - - -

Fe5Si3 + Cr5Si3 ⇌ ε + Cr3Si <900 U7 - - - -

Fe5Si3 + Cr3Si ⇌ ε + α ? U8 - - - -

Table 4. Investigations of the Cr-Fe-Si Materials Properties

Reference Method / ExperimentalTechnique

Type of Property

[1974Sie] Rockwell hardness, Stahl-Eisen-Prufblatt tensile testing

Mechanical properties of sintered compacts

[1977Si] Adiabatic calorimetry wit liquidhelium bath

Low-T specific heat. Alloys with 3 at.% Si.

[1978Hed] Faraday method Magnetic susceptibility of Cr2FeSi2 and Cr2FeSi3 at 4.3< T < 340 K

[1982Sau] Vickers hardness measurement,compression testing

Mechanical properties of Fe-30%Cr-3.5%Si (by mass)alloy with respect to σ phase content and distribution

[1995Zho1,1995Zho2]

Magnetooptical spectrometryand ellipsometry

Magnetic and magneto-optical properties of sputteredFe-Si/Cr multilayers

[1999Tan] Van der Pauw technique Hall effect and electrical resistivity of Cr doped βFeSi2

[2002Tan] Ab initio calculation – DFTmethod, GGA approximation

Electronic structure and density of states of Cr dopedβFeSi2

(continued)

10 Cr–Fe–Si


Reference Method / ExperimentalTechnique

Type of Property

[2003Kim] DC four probe and laser flashtechniques

Thermoelectric power, electrical and thermalconductivity

[2004Aru] SQUID magnetometry at 5 < T< 300 K and H up to 8 kOe.

Magnetization, magnetic susceptibility and resistivity ofCr doped βFeSi2

[2005Tur] NMR and XRD Nanostructure and magnetic properties

Fig. 1. Cr-Fe-Si. Cr-Fe binary system

Cr–Fe–Si 11



Fig. 2. Cr-Fe-Si. The lattice parameter variation of the ε phase as a function of Cr content

12 Cr–Fe–Si


Fig. 3. Cr-Fe-Si. FeSi2-CrSi2 quasibinary section

Cr–Fe–Si 13



Fig.4a

.Cr-Fe-Si.Tentativereactio

nscheme.Part1

14 Cr–Fe–Si


Fig.4b

.Cr-Fe-Si.Tentativereactio

nscheme.Part2

Cr–Fe–Si 15



Fig. 5a. Cr-Fe-Si. Tentative liquidus surface projection

16 Cr–Fe–Si


Fig. 5b. Cr-Fe-Si. Calculated liquidus surface projection

Cr–Fe–Si 17



Fig. 6. Cr-Fe-Si. Isothermal section at 427°C

18 Cr–Fe–Si



Cr–Fe–Si 19




20 Cr–Fe–Si


Fig. 9. Cr-Fe-Si. Calculated isothermal section at 1047°C

Cr–Fe–Si 21



Fig. 10. Cr-Fe-Si. Calculated isothermal section at 1150°C

22 Cr–Fe–Si


Fig. 11. Cr-Fe-Si. Partial isopleth for 1 mass% Cr, plotted in at.%

Cr–Fe–Si 23



Fig. 12. Cr-Fe-Si. Partial isopleth for 4 mass% Cr, plotted in at.%

24 Cr–Fe–Si


Fig. 13. Cr-Fe-Si. Partial isopleth for 7.3 mass% Cr, plotted in at.%

Cr–Fe–Si 25



Fig. 14. Cr-Fe-Si. Partial isopleth for 10.2 mass% Cr, plotted in at.%

26 Cr–Fe–Si


Fig. 15. Cr-Fe-Si. Partial isopleth for 0.5 mass% Si, plotted in at.%

Cr–Fe–Si 27



Fig. 16. Cr-Fe-Si. Partial isopleth for 1 mass% Si, plotted in at.%

28 Cr–Fe–Si


Fig. 17. Cr-Fe-Si. Calculated isopleth for xFe:xCr = 2:3

Cr–Fe–Si 29



Fig. 18. Cr-Fe-Si. Experimental integral enthalpies of mixing of liquid alloys along with the calculated curves for xFe:xCr = 4.55 at 1600°C

30 Cr–Fe–Si



Cr–Fe–Si 31




32 Cr–Fe–Si



Cr–Fe–Si 33



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154, 178–185 (1926) (Phase Diagram, Phase Relations, Experimental, 13)[1936And] Andersen, A.G.H., Jette, E.R., “X-Ray Investigation of the Iron-Chromium-Silicon Phase

Diagram”, Trans. Amer. Soc. Metals, 24, 375–419 (1936) (Crys. Structure, Phase Diagram,Phase Relations, Experimental, 37)

[1936Jet] Jette, E.R., Foote, F., “The Fe-Cr Alloy System”, Metals and Alloys, 7, 207–210 (1936)(Crys. Structure, Phase Relations, Experimental, 16)

[1937Sve] Svechnikov, V.N., Alferova, N.S., “Investigation of the Alloys of the Fe-Cr-Si System”(in Russian), Teor. Prakt. Metall., (3), 60–69 (1937) (Experimental, Morphology, Phase Dia-gram, Phase Relations, 27)

[1938And] Andersen, A G.H., Jette, E.R., “Notes on Microstructure and Hardness of Alloys ConsistingEssentially of Iron, Chromium and Silicon”, Trans. Amer. Inst. Min. Met. Eng., 131, 318–326(1938) (Phase Relations, Experimental, Mechan. Prop., 9)

[1942Kur] Kurnakov, N.N., “Investigation of the System Chromium-Silicon-Iron in the Region of Sili-cochrome”, Compt. Rend. (Doklady) Acad. Sci. URSS., 34(6), 158–159 (1942) (Phase Dia-gram, Phase Relations, Experimental, 8)

[1949Jae] Jaenecke, E., “Cr-Fe-Si” (in German) in “Kurzgefasstes Handbuch aller Legierungen”,Winter Verlag, Heidelberg, 614–615 (1949) (Phase Relations, Experimental, 1)

[1949Kur] Kurnakov, N.N., “Investigation of the Silicochromuim” (in Russian), Izv. Sekt. Fiziko-Khimich. Analiza, 17, 209–219 (1949) (Experimental, Morphology, Phase Diagram, PhaseRelations, 29)

[1956Gla1] Gladyshevskii, E.I., Cherkahsin, E.E., “Solid Solutions on the Basis of Metallic Com-pounds”, Russ. J. Inorg. Chem., 1(6), 288–295 (1956) (Crys. Structure, Phase Diagram,Phase Relations, Review, 4)

[1956Gla2] Galdyshevsky, E.I., Cherkashin, E.E, “Solid Solutions Based on Metallic Compounds”(in Russian), Zh. Neorg. Khim., 1(6), 1394–1401 (1956) (Crys. Structure, Experimental, 4)

[1957Aro] Aronsson, B., Lundstrom, T., “Investigations on σ-FeCrSi.”, Acta Chem. Scand., 11,365–372 (1957) (Crys. Structure, Phase Diagram, Experimental, 31)

[1958Aro] Aronsson, B., “An Investigation of the Me3Si3-MeSi Region of the Mn-Fe-Si and someRelated Systems”, Acta Chem. Scand., 12, 308–313 (1958) (Crys. Structure, Phase Diagram,Experimental, 22)

[1960Gel] Geld, P.V., Gertman, Y.M., “About the First Heat of Solubility of Liquid Transition Metals inLiquid Silicon” (in Russian), Fiz. Met. Metalloved., 10, 299–300 (1960) (Thermodyn.,Experimental, 8)

[1961Dub] Dubrovskaya, L.B., Geld, P.V., “The Quasibinary System α-Leboite-CrSi2” (in Russian),Trudy Uralsk. Politekhn. Inst., 114, 151–153 (1961) (Experimental, Mechan. Prop., PhaseRelations, 2)

[1962Bur] Burger, K.O., Wittmann, A., Nowotny, H., “The Crystal Structure of Co3Al3Si4 and Co2AlSi2and the Formation of some Monosilicidsystems of Transition Metals”, Monatsh. Chem.,93(1), 914 (1962) (Crys. Structre, Experimental, 8)

[1962Dub] Dubrovskaya, L.B., Geld, P.V., “Structure of the α-Leboite Chromium Disilicide Pseudobin-ary System”, Russ. J. Inorg. Chem. (Engl. Transl.), 7(1), 73–76 (1962), translated from Zhur.Neorg. Khim., 7, 145–148, (1962) (Crys. Structure, Phase Relations, Experimental, 16)

[1962Gla] Gladyshevskii, E.I., “Crystal Structure of Compounds and Phase Equilibria in Ternary Sys-tems of Two Transition Metals and Silicon”, Sov. Powder Metall. Met. Ceram. (Engl. Transl.),(4), 262–265 (1962) (Crys. Structure, Experimental, Phase Diagram, 17)

[1965Sch] Schueller, H.J., “Position of the Temperature Range of σ-Phase in Ferritic Cr Steels”(in German), Arch. Eisenhuettenwes., 36, 513–516 (1965) (Crys. Structure, Phase Relations,Experimental, 25)

[1966Gla] Gladyshevsky, E.I., Borusevich, L.K., “The Cr-Fe-Si Ternary System” (in Russian), Izv. Akad.Nauk SSSR, Met., 1, 159–164 (1966) (Crys. Structure, Phase Diagram, Experimental, #, 21)

34 Cr–Fe–Si


[1968Che] Cherkasov, P.A., Averin, V.V., Samarin, A.M., “Activities of Silicon and Titanium in MoltenIron, Cobalt, and Nickel Containing Chromium”, Russ. J. Phys. Chem. (Engl. Transl.), 42(3),401–404 (1968), translated from Zh. Fiz. Khim, 42(3), 767 (1968) (Experimental, Phase Dia-gram, Phase Relations, Thermodyn., 19)

[1968Sid] Sidorenko, A.F., Dmitriev, E.A., Apasova, E.A., “Crystal Lattice Constants in Some Quasi-binary Alloys Based on FeSi” (in Russian), Trudy Uralsk. Politekhn. Inst., (167), 124–127(1968) (Crys. Structure, Experimental, 3)

[1969Bau] Baum, B.A., Pavars, I.A., Geld, P.V., “Density and Surface Energy of Liquid Chromium-Iron-Silicon Alloys”, Russ. J. Phys. Chem. (Engl. Transl.), 43(10), 1379–1382 (1969)(Experimental, 9)

[1971Blo] Blossey, R.G., Pehlke, R.D., “Solubility of Hydrogen in Liquid Fe-Co-Ni Alloys”, Metall.Trans., 2, 3157–3161 (1971) (Phase Relations, Experimental, 26)

[1971Yam] Yamaguchi, M., Umakoshi, Y., Mima, G., “On the Miscibility Gap in the Fe-Cr-X (X = Cu,Mn, Mo, Ni, V, Si and Al) System”, Proc. Int. Conf. Sci. Technol. Iron Steel, Tokyo, 11,1015–1019 (1971) (Electronic Structure, Experimental, Phase Relations, 35)

[1973Wet] Wethmar, J.C.M., Howat, D.D., Jochens, P.R., Strydom, O.A.W., “Liquidus Temperatures inthe Cr-Fe-Si System in the Composition Range Representative of Ferrochromium-SilicideProduced in South Africa”, J. S. African Inst. Min. Met., 73, 181–183 (1973) (Experimental,Phase Diagram, Phase Relations, 11)

[1974Sie] Sienel, D., Thummler, F., Zapf, G., “Sintered Iron-Base Alloys with High Strength UsingSelected Nitrides and Silicides”, Powder Met., 17(33), 206–226 (1974) (Experimental,Kinetics, Mechan. Prop., 5)

[1975Igu] Igushev, V.F., Tolstoguzov, N.V., Rudenko, V.A., “Investigation of the Enthalpies of LiquidFe-Cr-Si Alloys” (in Russian), Izv. Vuz. Chern. Metallurg., 6, 46–50 (1975) (Experimental,Thermodyn., 9).

[1975Pic] Pickart, S., Litrenta, T., Burch, T., Budnick, J.I., “Site Preference of Dilute Transition MetalSolutes in Fe3Si”, Phys. Letters A, 53A(4), 321–323 (1975) (Experimental, Crys. Structure, 4)

[1977Ost] Ostrovskii, O.I., Stomakhin, A.Ya., Dietrich, E., Grigoryan, V.A., “Heats of Solution of Alu-minium, Silicon and Titanium in Iron-Chromium and Nickel-Chromium Melts” (in Russian),Vses. Konf. Kalorim., (Rasshir. Tezisy Dokl.), 59–63 (1977) (Experimental, Thermodyn., 11)

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Cr–Fe–Si 35



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[1995Hay] Hayes, F.H., Bin Awais, H., Unpublished work, as quoted by [1997Lin]

36 Cr–Fe–Si


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Cr–Fe–Si 37



Mater., 294(2), e151-e154 (2005) (Crys. Structure, Experimental, Magn. Prop., Phase Rela-tions, 6)

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38 Cr–Fe–Si


Chromium – Iron – Titanium

Volodymyr Ivanchenko, Tetyana Pryadko

Introduction

The earliest report on the liquidus surface is that of [1940Vog]. These results were reproduced in the reviewby [1949Jae]. Because [1940Vog] used Ti of only 95% purity for making their alloys and also assumed acompound Ti3Cr2 (which was not confirmed by the later studies), their results are not considered in the pre-sent evaluation. The phase diagram of the Ti rich portion of the Cr-Fe-Ti system was studied by [1953Thy]and isothermal sections at 900, 800, 750, 700, 650, and 550°C, as well as isopleths at 2 at.% Cr, 6 at.% Cr,10 at.% Cr, and 96 at.% Ti, 94 at.% Ti, 90 at.% Ti were constructed by studying arc-cast alloys. However,the main results from the study of phase equilibria in Cr-Fe-Ti alloys have been obtained by [1959Bor,1961Bor, 1964Bor, 1964Kor], who presented isothermal sections at 550, 800, and 1000°C, a projectionof the liquidus surface and the TiFe2-TiCr2 quasibinary section. Dilatometry studies were carried out inorder to determine the Ac3 transformation temperature. These results, both [1952Roe, 1953Thy] and[1959Bor, 1961Bor, 1964Bor, 1964Kor] were the basis for the review presented by [1987Rag].[1975Kau] calculated isothermal sections at 1427, 1227, and 1000°C. A comparison of the calculatedand experimental [1964Bor] sections at 1000°C showed good general agreement except for two features.Firstly, the observed section does not show the fcc field arising from the stable structure of iron at 1000°C. Secondly, the calculations did not take into account the ternary phase τ. [1978Hao] studied the Cr-Fe-Tisystem in order to identify a new Fe based eutectic suitable for the development of materials for the fabrica-tion of directionally solidified turbine blades with operating temperatures up to 1150°C. The thermody-namic properties of Ti in Cr-Fe melts at 1600°C were studied by [1968Che, 1977Ost, 2005Cho] usingdifferent experimental techniques. Reoptimized parameters for liquid, bcc and fcc phases of the Cr-Fe-Tisystem were presented by [1998Mie].There is only one ternary phase in the Cr-Fe-Ti system. It has an αMn cubic structure [1959Bor, 1960Bor,1964Bor, 1965Hug] and was originally named in the literature as the χ phase. It is designated in Table 2as τ. Some peculiarities of its crystal structure have been studied by [1966Kim] and [1971Sny, 1972Sny].A summary of experimental studies on phase relations, structure and thermodynamics is given in Table 1.

Binary Systems

There is a number of Calphad assessments of the Fe-Ti system in the literature, but none of them shows ahomogeneity range for the TiFe2 Laves phase. Therefore the phase diagram for this system is taken from[Mas2]. The Cr-Ti system is taken from the assessment of [2000Zhu]. The Cr-Fe phase diagram is takenfrom [1993Itk].

Solid Phases

The details of the crystal structures and lattice parameters of the solid phases are listed in Table 2.The σ phase in the Cr-Fe system is tetragonal with 30 atoms per unit cell. Although there is no experimentalevidence for the extension of the σ phase into the ternary system, any extension would be severely limitedby the equilibria involving (γFe), αβδ, λ1 and τ. The extension of the αβδ phase into the ternary system islimited at ~ 1360°C by the precipitation of the intermetallic phase λ1.TiCr2 exists in all the three Laves phase modifications. The γTiCr2 phase with the MgZn2 type crystal struc-ture forms from the (βTi,αCr) solid solution at 1359°C. A polymorphic transformation of the MgZn2 typestructure to the MgNi2 type structure leads to a peritectoid reaction at 1271°C on the Ti rich side of theβTiCr2 homogeneity region, and a eutectoid reaction at 1269°C on the Cr rich side. This phase decomposeseutectoidally at 804°C to give αTiCr2 with the MgCu2 type structure and the (βTi,αCr) solid solution.αTiCr2 decomposes peritectoidally at 1223°C into βTiCr2 and the (βTi,αCr) solid solution. According to

Cr–Fe–Ti 1



[1953Thy] and [1964Kor], there should be a complete miscibility between γTiCr2 and the isostructuralTiFe2 phase depending on the temperature, owing to the similar atomic sizes of Fe and Cr. The γTiCr2 phaseis stable over a limited temperature range in the Cr-Ti binary system, 1359-1269°C, whereas the TiFe2 phaseis stable up to its melting point of 1424°C. [1953Thy] determined lattice parameters across the phase dia-gram from the TiFe2 composition towards FeCr2 for a number of temperatures and found that completesolubility occurred at 1300°C. The lattice parameters increase continuously with increasing substitutionof Cr for Fe. The existence of the λ1, Ti(Fe1–xCrx)2 solid solution indicates that Fe most likely substitutesfor Cr. The variation is not quite linear, but data from both studies show good agreement. The solubilityof Fe in all modifications of the Laves phase is limited.There is only one ternary phase formed in this system. The ideal formula is Ti5Cr7Fe17. It corresponds totwo formula units per unit cell. In the cubic crystal structure of αMn, 58 atoms occupy four crystallographi-cally independent sets of positions. There are two equivalent Mn I sites (coordination number (CN) 16), 8Mn II sites (CN 16), 24 Mn III sites (CN 13) and 24 Mn IV sites (CN 12). This complex structure is similarto that of electronic intermetallic compounds in which the occupation of the sites is defined by the space-filling of different sized atoms. Larger atoms tend to occupy the CN 16 sites with which relatively largevolumes are associated, while smaller atoms favor the sites of lesser volume, CN 13 and CN 12.[1966Kim] used the Mössbauer technique to study hyperfine interactions in an alloy of the Ti17Cr24Fe59composition and obtained a set of data consistent with the assumption that of the four available sets of sites,Fe occupies only sites III and IV. Domain structures, related to the deviation from stoichiometry have beenobserved in the τ phase by [1972Sny]. Domain boundaries, named as “dissociated anti-phase boundaries”[1971Sny], are non conservative phase boundaries dissociated into several components with decreasingenergy.The metastable ω phase is formed by quenching from the αβδ region.The effect of mechanical grinding on the phase transformation in Cr-Ti alloys with nominal compositions ofTiCr2–x (x = 0, 0.2, 0.5) was studied by [2003Tak]. The intermetallic TiFe2 phase is unstable upon mechan-ical milling producing a mixture of super-saturated α-Fe(Ti) solid solution and FeTi intermetallic phase.This solid solution decomposes for prolonged milling ejecting a large fraction of the Ti atoms that were pre-viously dissolved in the αFe structure. The instability of compound could be related to the accumulation ofa many stacking faults on prolonged milling which favors a hcp - bcc structure transition [2004Gue].


Table 3 lists the invariant reactions occurring in the system. This should be considered in conjunction withFigs. 2, 3, 4 and 5. The reaction scheme is presented in Fig. 1 according to [1987Rag]. However thetwo maxima presented by [1987Rag] as peritectic reactions have been corrected here for eutectic reactionse2 and e3.


The liquidus surface was constructed by [1987Rag] based on the results of [1964Bor]. Using the binary dataand the liquidus temperatures from vertical sections of [1940Vog] as a guide, approximate isotherms at 100°C intervals ranging from 1100 to 1800°C were presented. The peritectic formation of the ternary τ phasegiven by [1987Rag] taken from [1964Bor] is in contradiction with the nature of peritectic reactions. Forexample, the Ti content in e2 is lower then in the τ phase whereas it should be higher. And similarly, theTi content in p3 is higher than in the τ phase, whereas it should be lower. For this reason, it was assumedthat the τ phase melts incongruently and, as a consequence, forms eutectics with αβδ and λ1 phases.The modified liquidus surface is presented in Fig. 2.

Isothermal Sections

The isothermal sections at 1000, 800 and 550°C are redrawn from [1987Rag] based on the work of[1964Bor]. They are presented in Figs. 3 to 5.A tentative phase boundary relating to the α/γ phase transformation for 1070°C was presented by[1952Roe] following dilatometric studies of an alloy containing 7.01 mass% Cr and 1.29 mass% Ti and

2 Cr–Fe–Ti


the corresponding phase boundaries in the Cr-Fe and Fe-Ti binary systems. However, as only a single tern-ary experimental measurement was used, this result has not been reproduced here.


The TiCr2-TiFe2 section was presented by [1964Kor] but it is omitted here owing to the low accuracy of thesolidus temperature determination. For example, the measured melting temperature of TiFe2 was given asapproximately 1500°C, whereas the assessed value [1981Mur, 1998Dum] is 1427°C. The temperature com-position sections of the Ti rich region of the phase diagram at 2, 6 and 10 mass% Cr were derived by[1953Thy] based on their isothermal section determinations. However, in their sections, there was no dis-tinction between the three types of the Laves phase present in the Cr-Ti system. Therefore they are in dis-agreement with the accepted binary systems and the ternary isothermal sections from [1987Rag]. Hence,they are not reproduced here.

Thermodynamics

Experimental data of Cr-Fe-Ti ternary alloys are limited to the results presented in [1968Che, 1977Ost,2005Cho]. In [1968Che] and [1977Ost], the dependence of the activity coefficient log10�0Ti on the chro-mium concentration at 1600°C is linear, and extrapolation of the relationship to zero chromium concentra-tion yields for the activity coefficient of titanium a value of 0.008. The interaction parameter betweenchromium and titanium in Fe-based melts was calculated to be eCrTi= 0.024. The calculated value of the activ-ity coefficient of Ti in a Fe-30%Cr melt is 0.039. [1977Ost] obtained for a Fe-Ti melt �0Ti= 0.007 and 0.032for Fe-30%Cr. The measured values of �0Ti and eCrTi [2005Cho] are slightly higher (0.018 and 0.032) thanthose presented earlier. Thermodynamic data based on the few experiments were optimized for liquid, fer-rite and austenite for some iron-based systems that had not been assessed earlier [1998Mie]. The data wereshown to improve the correlation between calculated and measured liquidus temperature, solute partitionand primary phase data in typical stainless steels.


In an attempt to reduce the use of expensive alloying elements such as nickel, steel makers have developedless expensive ferrite grades of stainless steels. It is shown in [2004Elk] that ultra-high purity 18% Cr fer-ritic stainless steel exhibits superior corrosion resistance, such as resistance to chloride stress corrosioncracking, corrosion in oxidizing aqueous media and pitting and crevice corrosion in chloride media.[2000Ike1, 2000Ike2] showed that Cr-Fe-Ti alloys containing 1.6 to 7.1 mass% Fe and 3.3 to 14.1 mass% Cr, having the same number of electrons per atom (~ 4.26), with higher Cr content had a smaller volumefraction of the athermal ω phase then alloys with higher Fe content. The tensile strength decreased withincreasing Cr content, while elongation and reduction of cross section are significantly increased. The bal-ance between tensile strength and ductility (elongation and reduction in area) improved in the alloys withsubstitution of Cr for Fe. Therefore, no negative influences of ferro-chromium alloys on mechanical proper-ties was observed in these studies. [2002Ike] studied the influence of cooling rate on phase constitution andmechanical properties for two alloys (Ti-4.9Cr-2.9Fe and Ti-6.9Cr-4.2Fe (mass%)). The hardness for aTi-4.9Cr-2.9Fe (mass%) alloy cooled at 25 K·s–1, was about HV=400, and for Ti-6.9Cr-4.2Fe (mass%)HV was around 320. It was shown that the last alloy has a lower sensitivity to quenching conditions. Thetensile strength and reduction of area obtained with the Ti-4.9 mass% Cr-2.9 mass% Fe alloy were about1200 MPa and 45%, respectively. A tensile strength of 1050 MPa and a reduction in cross sectional areaof 60% were obtained with a Ti-6.9 mass% Cr-4.2 mass% Fe alloy. [1998Che] showed that addition ofFe to TiCr2, which partitions to both the A and B sublattices, can improve the toughness. Stabilization ofthe cubic C15 crystal structure also resulted in higher toughness values. [1992Wol] studied the influenceof microalloying additions (0.05 mass% Ti) on the plastic deformation of a Fe-40Cr base alloy. Titaniumappears to be active in changing the local strain parameter which governs the stress required to nucleatea void, apparently by modifying the second phase particles. The C14 Laves phase TiCr2 is an antiferromag-net with a Neel temperature of about 12°C. The magnetic properties and the magnetic phase diagram of Crwith substitutions for Fe have been investigated by [2003Yam]. In the Cr substituted Ti(CrxFe1–x)2 alloy,

Cr–Fe–Ti 3



antiferromagnetism is maintained up to around x = 0.3, while spontaneous magnetization appears at0.075 < x < 0.2 at low temperatures under antiferromagnetic order. A summary of experimental investiga-tions of properties is given in Table 4.

Miscellaneous

The influence of Ti on the kinetics of σ phase formation promoted by an isothermal annealing was studiedby [2000Bla1, 2000Bla2, 2001Bla] using 57Fe Mössbauer spectroscopy. It was found that the kinetics couldbe well described in terms of the Johnson-Avrami-Mehl equation. The addition of Ti (up to 3 at.%) wasrevealed to affect the transformation kinetics in the following way: for xTi ≤ 1.5 at.%, its presence acceler-ates the process with the highest transformation rate. For xTi = 0.3 at.%, and xTi ≥ 1.5 at.%, titanium retardsthe formation of the σ phase. Quantitatively, the effect of Ti on the kinetics was described in terms of achange in the effective activation energy [2000Bla1, 2000Bla2]. The study of σ phase formation inFe53.8Cr46.2 and Fe53.8Cr46.2-0.3 at.%Ti by isothermal annealing in the temperature interval of 650-730°Cshowed that average value of the activation energy was found to be equal to 196±2 kJ·mol–1 for theFe53.8Cr46.2 sample and to 153±2 kJ·mol–1 for a sample doped with Ti [2001Bla]. Phase transformationsof the TiCr1.8 and TiCr1.7Fe0.1 intermetallic compounds during their interaction with hydrogen at pressuresfrom 1 to 200 MPa and temperatures of 195-295 K were studied by [1999Kly]. For TiCr1.8 formation, threetypes of hydride phases were described: (I) hexagonal Laves phase preserving the MgZn2 structure of theinitial intermetallic, (II) fcc phase of the CaF2 structure type which can be considered as a solid solutionin the quasibinary TiH2-CrH2 system and the (III) bcc phase with the W structure type - a high temperatureβ-solution in the Cr-Ti system formed on decomposition of the (II) phase during hydrogen desorption overthe temperature interval from 100 to 150°C and a hydrogen pressure of 0.01 MPa. The substitution of 5 %of the Cr for Fe completely suppressed the formation of the fluoride type phase.[2002Boz] used the Bozzolo-Ferrante-Smith (BFS) method, one of the family of quantum approximatemethods, to analyze site substitution behavior of Cr additions to the binary TiFe compound. The decompo-sition of βTi in step-quenched Ti base alloys was studied by [1952Phi1] and discussed in [1952Phi2].A nondiffusive transition into the intermediate state was observed by [1990Wir] at a temperature of 450±5°C on heating the nonequilibrium bcc Ti60Cr32Fe8 solid solution obtained by quenching bulk material at amoderate cooling rate. This transformation is known as “inverse melting”, i.e. the thermodynamic transitionfrom the quenched β phase into the frozen liquid; even though the resulting phase exhibits configurationallong-range order.Ultra-high purity 18% Cr stainless steel exhibits superior corrosion resistance. But, grain enlargement islikely to occur during their production and heat treatment, and this causes practical drawbacks in termsof surface qualities and mechanical properties. It was found that simultaneous additions of B and Ti leadto a coarser grain structure than with single additions of Ti [2004Elk]. When B is added with Ti, coarseM23(C,B)6 was found to precipitate primarily on TiN and hence the total density of the precipitates in aTi+B steel becomes lower than that in a steel with only Ti additions containing both TiN and TiC precipi-tates producing the pinning effect.The sulphidation properties of Fe-25Cr alloyed with Ti were studied by [1993Qi].[1992Wol] studied the influence of microalloying with Ti on the plastic deformation of Fe-40Cr alloys. Astudy of surface mechanical properties of Ti+ implanted Cr-Fe alloys has been performed by [1988Sas].

4 Cr–Fe–Ti


Table 1. Investigations of the Cr-Fe-Ti Phase Relations, Structures and Thermodynamics


[1940Vog] Light microscopy, X-ray diffraction, TA Up to 1700°C; Fe- Ti33.3Fe66.7-Ti60Cr40-Cr;phase diagram

[1952Phi1] Light microscopy, X-ray diffraction,hardness

Up to 650°C, α, βTi,martensite under decomposition of (βTi)

[1952Roe] Dilatometry 850-1088°C, 2.46-8.2 mass% Cr and 0.06-2.71mass% Ti, α→γ transformation

[1953Thy] Light microscopy, X-ray diffraction,incipient melting studies, Vickershardness

Up to 1300°C; Ti-Ti60Cr40-Ti60Fe40,TiCr2-TiFe2

[1959Bor] Light microscopy, X-ray diffraction,density

Quenched from 1000°C, TiFe2-CrFe section

[1960Bor] Light microscopy, X-ray diffraction,Vickers hardness

Homogenization at 1100°C for 100 h,following quenching from 1000°C (100 hsoaking), annealing at 550°C for 500 h;Cr-TiCr2-TiFe2-Fe

[1961Bor] Light microscopy, X-ray diffraction,Vickers hardness, microhardness of theindividual phases

Homogenization at 1100°C, followingquenching from 1000°C (100 h soaking) andquenching after annealing at 550°C for 500 h;Cr-TiCr2-TiFe2-Fe

[1964Bor] Light microscopy, X-ray diffraction,thermal analysis, microhardness of theindividual phases

Homogenization at 1100-1200°C (100 h),annealing at 1000°C for 100 h, 800°C for 300 hand quenching, annealing at 600°C for 400 hand 550°C for 1000 h; Cr-Fe-Ti

[1965Hug] X-ray diffractometry Annealing at 850°C for 120 h,Ti10Cr12Fe36, τ phase

[1966Kim] Mössbauer spectroscopy 4, 77 and 300 K,Ti17Cr24Fe59, τ phase

[1968Che] Equilibration of melt with steam/hydrogenmixture

Activities of Ti in Cr-Fe-Ti melts at 1600°Cwith 5, 18, and 25 mass% Cr and up to 2 mass%Ti

[1972Sny] Bright field and dark field TEM, electronmicrodiffraction

Fe36Cr22–xTix, 8 < x < 10

[1975Kau] Calculations 1427, 1227, and 1000°C; Cr-Fe-Ti

[1977Ost] Isoperibolic calorimetry 1600°C, Fe-30 at.% Cr, liquid

[1978Hao] Light metallography, SEM, EDAX (60Fe-5Cr-35Ti)-(60Fe-15Cr-25Ti)-(80Fe-5Cr-15Ti) (mass%) composition region

(continued)

Cr–Fe–Ti 5




[1990Wir] DSC, dilatometry, X-ray diffraction Up to 500°C, Ti60Cr32Fe8 cooled from theliquid state with the rate 103 K·s–1

[2005Cho] Galvanic cell technique Activity of Ti in Cr-Fe-Ti melts at 1600°C,4.09 - 30.98 mass% Cr,0.018-0.19 mass% Ti





Comments/References

αβδ, (Ti1–xCrx)Fey

(αCr)< 1863(δFe)1538 - 1394(αFe)< 912(βTi)1670 - 882

cI2Im�3mW

a = 288.28a = 288.14a = 287.97a = 287.79a = 287.67a = 287.53a = 287.41a = 287.23a = 287.16a = 286.55a = 286.33a = 325.95a = 325.34a = 324.45a = 323.55a = 322.94a = 321.40a = 288.48

a = 293.22

a = 286.65

a = 330.65

x = 1, y = 0.043x = 1, y = 0.11x = 1, y = 0.208x = 1, y = 0.306x = 1, y = 0.408x = 1, y = 0.470x = 1, y = 0.580x = 1, y = 0.698x = 1, y = 0.833x = 1, y = 0.935x = 1, y = 0.979 [V-C2]x = 0.06, y = 0x = 0.08, y = 0x = 0.10, y = 0x = 0.12, y = 0x = 0.14, y = 0x = 0.16, y = 0 [V-C2]pure Cr at 25°C [Mas2]

pure Fe at 1394°C [V-C2]


pure Ti at 25°C [Mas2]

(γFe)1394 - 912

cF4Fm�3mCu

a = 364.97 at 915°C [V-C2, Mas2]

(αTi)< 882

hP2P63/mmcMg

a = 295.06c = 468.35

pure Ti at 25°C [Mas2]dissolves 0.6 at.% Cr at 667°C[1987Mur]

(continued)

6 Cr–Fe–Ti





Comments/References

(εFe) hP2P63/mmcMg

a = 246.8c = 396.0


(ωTi)HP hP3P6/mmmωTi

a = 462.5c = 281.3

at 25°C [Mas2]

σ, CrFe tP30P42/mnmCrFe

a = 879.66 ± 0.06c = 455.82 ± 0.03

a = 879.68 ± 0.05c = 455.85 ± 0.03a = 879.61 ± 0.04c = 456.05 ± 0.03

49.5 at.% Cr, [V-C2]

48.2 at.% Cr, [V-C2]

47.8 at.% Cr, [V-C2]

λ1, Ti(Fe1–xCrx)2

γTiCr2 (h2)1359 - 1269TiFe2< 1424

hP12P63/mmcMgZn2

a = 493.2c = 800.5a = 480.4 to 477.7c = 784.9 to 780.7a = 486.7c = 795.3

x = 0.85 at 1000°C.C14 structure.63.57-66.31 at.% Cr at 1300°C [1998Che]at 25°C [V-C2]

65.2-72.8 at.% Fe, annealing at 1000°C[1981Mur]Ti2Cr3.1Fe0.78 [V-C2]. Composition isoff-stoichiometry owing to the homogeneity range

λ2, αTiCr2< 1223

cF24Fd3mMgCu2 a = 693.2

C15 structure. 63.5-65.2 at.% Cr at 1000°C[1998Che]at 25°C TiCr1.9 [V-C2].Dissolves ~3 at.% Cr at 1000°C

λ3, βTiCr21271 - 804

hP24P63/mmcMgNi2 a = 493.2

c = 1601.0

C36 structure. 63.57-66.31 at.% Cr at 1200°C[1998Che]at 25°C Ti1.12Cr2 [V-C2].Dissolves ~3 at.% Cr at 1000°C

TiFe< 1317

cP2Pm�3mCsCl

a = 298.8 to 297.7 48-50.2 at.% Fe [1981Mur]

ω(Ti,Cr)< 450

hP3P3m1ω (Ti,Cr)

a = 461.6c = 282.7

3-9 at.% Cr [1987Mur]at 4.6 at.% Crmetastable

(continued)

Cr–Fe–Ti 7






Comments/References

ω(Ti,Fe) h**P6/mmm-

- The composition approaches ~ 4 at.% Fe, formsduring quenching from αβδ-region (10-16 at.%Fe) or during aging between 300 and 500°C[1981Mur]

* τ, Ti5Fe17 Cr7 cI58P63/mmmαMn

a = 894.7

a = 892.2

at Ti5Fe18Cr6annealed at 900°C [V-C2]~17 at.% Ti, 19-34 at.% Cr,49-69 at.% Feat Ti5Cr7Fe17 [1961Bor]



Cr Fe Ti

l ⇌ λ1 + αβδ 1550 e1 l 65 8 27

L ⇌ τ > 1290 congruent L, τ ~ 24.1 ~ 58.7 ~ 17.2

l ⇌ αβδ + τ ~ 1290 e2 lτ

~ 24.1~ 24.1

~ 62.9~ 58.7

~ 13~ 17.2

l ⇌ τ + λ1 ~ 1290 e3 Lτ

~ 24.1~ 24.1

~ 58.5~ 58.7

~ 17.4~ 17.2

L ⇌ αβδ + τ + λ1 ~ 1270 E1 L ~ 37.8 ~ 45 ~ 17.2

L ⇌ αβδ + τ + λ1 ~ 1270 E2 L ~ 9.3 ~ 76.7 ~ 14

L + λ1 ⇌ αβδ + FeTi ~ 1200 U1 L 10.5 28.5 61

αβδ + λ3 ⇌ λ1 + λ2 > 1000 U2 - - - -

λ1 + λ3 ⇌ λ2 + αβδ ~ 804 U3 - - - -

αβδ + λ2 ⇌ (αTi) + λ1 ~ 650 U4 - - - -

αβδ ⇌ (αTi) + λ1 + TiFe 540 E3 αβδ ~ 8? ~ 13 ~ 79

8 Cr–Fe–Ti


Table 4. Investigations of the Cr-Fe-Ti Materials Properties


[1953Thy] Vickers hardness Hardness

[1988Sas] Surface hardness, friction coefficient Surface mechanical properties ofTi+-implanted Cr-Fe alloys

[1992Wol] Specimens were tensile tested to failure at an initialstrain rate of 10–4·s–1 (cross head speed = 0.15mm·min–1) at RT

Mechanical properties

[1993Qi] Quartz spring thermal balance, SEM, EDAX, EPMA,XRD

Sulphidation of Fe-25Cr-4Ti inH2-H2S mixtures at 800°C

[1998Che] Vickers hardness, EPMA, XRD Resistant to fracture of Ti(Fe,Cr)2ternary alloys

[2000Ike1][2000Ike2][2002Ike]

Vickers hardness, tensile test with a cross head speed =3 mm·min–1, light microscopy, XRD, SEM, electricalresistivity at room temperature and liquid nitrogentemperature

Mechanical properties

[2003Yam] Magnetization measurements were carried out from 4.2to 300 K and up to 10.3 kOe by a torsion magneticbalance

Magnetic phase diagram of the Ti(CrxFe1–x)2 system with x < 0.5

Cr–Fe–Ti 9



Fig.1.

Cr-Fe-Ti.

Reactionscheme

10 Cr–Fe–Ti


Fig. 2. Cr-Fe-Ti. Liquidus surface projection

Cr–Fe–Ti 11



Fig. 3. Cr-Fe-Ti. Isothermal section at 1000°C

12 Cr–Fe–Ti



Cr–Fe–Ti 13




14 Cr–Fe–Ti


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[1965Hug] Hughes, H., “Precipitation in Alloy Steels Containing Cr, Ni, Al and Ti”, J. Iron Steel Inst.,203, 1019–1023 (1965) (Crys. Structure, Phase Relations, Experimental, 29)

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[1968Che] Cherkasov, P.A., Averin, V.V., Samarin, A.M., “Activities of Silicon and Titanium in MoltenIron, Cobalt, and Nickel Containing Chromium”, Russ. J. Phys. Chem. (Engl. Transl.), 42(3),401–404 (1968), translated from Zh. Fiz. Khim, 42(3), 767 (1968) (Experimental, Phase Dia-gram, Thermodyn., 19)

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[1975Kau] Kaufman, L., Nesor, H., “Calculation of Superalloy Phase Diagrams. Part III”, Metall. Trans.A, 6(11), 2115–2122 (1975) (Phase Diagram, Phase Relations, Thermodyn., Calculation, 35)

[1977Ost] Ostrovskii, O.I., Stomakhin, A. YA., Dietrich, E., Grigoryan, V.A., “Heats of Solution of Alu-minium, Silicon and Titanium in Iron-Chromium and Nickel-Chromium Melts” (in Russian),

Cr–Fe–Ti 15



in “Vses. Konf. Kalorim., (Rasshir. Tezisy Dokl.)”, 59–63 (1977) (Experimental, Thermo-dyn., 11)

[1978Hao] Haour, G., Mollard, F., Lux, B., Wright, G., “New Eutectics Based on Fe, Co or Ni”,Z. Metallkd., 69(1), 26–32 (1978) (Phase Relations, Phase Diagram, Morphology, Experi-mental, 24)

[1981Mur] Murray, J.L., “The Fe-Ti (Iron-Titanium) System”, Bull. Alloy Phase Diagrams, 2(3)320–334 (1981) (Crys. Structure, Phase Diagram, Thermodyn., Assessment, *, 124)

[1987Mur] Murray, J.L., “The Cr-Ti (Chromium-Titanium) System” in “Phase Diagrams of Binary Tita-nium Systems”, ASM International, Metals Park, Ohio, 68–78 (1978) (Crys. Structure, PhaseDiagram, Thermodyn., Assessment, *, 60)

[1987Rag] Raghavan, V., “Section I. The Cr-Fe-Ti (Chromium-Iron-Titanium) System” in “Phase Dia-grams of Ternary Iron Alloys. Part 1”, Ind. Inst. Techn., Delhi, 1, 43–54 (1987) (Crys. Struc-ture, Phase Diagram, Phase Relations, Review, #, 15)

[1988Sas] Sasaki, J., Iwaki, M., “Surface Mechanical Properties and Micro-Characteristics of Ti+-Implanted Fe-Cr Alloys Depending on Cr Concentration” in “Fundamentals of Beam-SolidInteractions and Transient Thermal Processing”, Symposium. Mater. Res. Soc., 185-190,(1988) (Phase Relations, Mechan. Prop., Experimental, 8)

[1990Wir] Wirz, Ch., Blatter, A., Baltzer, N., von Allmen, M., “Transformations Preceding Amorphiza-tion in Cr-Ti and Cr-Ti-Fe β Phases”, Phys. Rev. B, 42(11), 6993–6999 (1990) (Crys. Struc-ture, Phase Relations, Thermodyn., Experimental, Kinetics, 37)

[1992Wol] Wolff, I.M., Ball, A., “Substitutional Alloying and Deformation Modes in High ChromiumFerritic Alloys”, Metall. Trans. A, 23A(2), 627–638 (1992) (Crys. Structure, Experimental,Interface Phenomena, Mechan. Prop., Morphology, 37)

[1993Itk] Itkin, V.P., “Cr-Fe (Chromium-Iron)” in “Phase Diagram of Binary Iron Alloys”, Okamoto,H. (Ed.), ASM International, Materials Park, OH (1993) (Crys. Structure, Phase Diagram,Assessment, #, 33)

[1993Qi] Qi, H.B., Zhu, R.Z., He, Y.D., “The Sulphidation Properties of Titanium-, Manganese- andNiobium-Bearing Fe-25Cr Alloys in H2-H2S Mixtures at 800°C”, Corros. Sci., 35(5-8),1099–1103 (1993) (Crys. Structure, Experimental, Kinetics, Morphology, 10)

[1998Che] Chen, K.C., Allen, S.M., Livingston, J.D., “Factors Affecting the Room-TemperatureMechanical Properties of TiCr2-Base Phase Alloys”, Mater. Sci. Eng. A, A242, 162–173(1998) (Crys. Structure, Experimental, Mechan. Prop., 51)

[1998Dum] Dumitrescu, L.F.S., Hillert, M., Saunders, N., “Comparison of Fe-Ti Assessments”, J. PhaseEquilib., 19(5), 441–448 (1998) (Phase Diagram., Thermodyn., Assessment, #, 48)

[1998Mie] Miettinen, J., “Approximate Thermodynamic Solution Phase Data for Steels”, Calphad, 22(2), 275–300 (1998) (Assessment, Calculation, Phase Diagram, Phase Relations, Thermo-dyn., 98)

[1999Kly] Klyamkin, S.N., Kovriga, A.Yu., Verbetsky, V.N., “Effect of Substitution on F.C.C. and B.C.C. Hydride Phase Formation in the Ti-Cr2-H2 System”, Int. J. Hydrogen Energy, 24, 149–152(1999) (Crys. Structure, Experimental, Phase Relations, 4)

[2000Bla1] Blachowski, A., Cieslak, J., Dubiel, S.M., Sepiol, B., “Effect of Titanium on the Kinetics ofthe σ phase Formation in a Coarse-Grained Fe-Cr Alloy”, Intermetallics, 8, 963–966 (2000)(Experimental, Kinetics, 10)

[2000Bla2] Blachowski, A., Cieslak, J., Dubiel, S.M., Zukrowski, J., “Effect of Titanium on the Kineticsof the σ Phase Formation Ion a Small Grain Fe-Cr Alloy”, J. Alloys Compd., 308, 189–192(2000) (Experimental, Kinetics, 9)

[2000Ike1] Ikeda, M., Komatsu, S., Inoue, K., Shiota, H., Imose, T., “Effect of Chromium Content onElectrical Resistivity and Tensile Properties of Ti-Fe-Cr Alloys”, Mater. Sci. Technol., 16(6), 605–608 (2000) (Experimental, Mechan. Prop., Morphology, 20)

[2000Ike2] Ikeda, M., Komatsu, S., Inoue, K., Hiroyuki, S., Imose, T., “Effect of Cr Addition on TensileProperties of Ti-Fe-Cr β Alloys Quenched from β Single Phase Region at 1173K”

16 Cr–Fe–Ti


(in Japanese), J. Jpn. Inst. Met., 64(5), 279–282 (2000) (Experimental, Mechan. Prop., Mor-phology, 12)

[2000Zhu] Zhuang, W., Shen, J., Liu, Yu., Ling, L., Shang, S., Du, Y., Schuster, J.C., “ThermodynamicOptimization of the Cr-Ti System”, Z. Metallkd., 91 (2), 121-127 (Crys. Structure, Phase Dia-gram, Thermodyn., Assessment, #, 53)

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[2002Ike] Ikeda, M., Komatsu, S.-Y., Imose, T., Inoue, K., “The Effect of Average Cooling Rate from aTemperature within β Single Phase Region on Phase Constituiton and Tensile Properties ofTi-Fe-Cr Alloys” (in Japanese), J. Jpn. Inst. Met., 66(3), 131–134 (2002) (Experimental,Mechan. Prop., Morphology, 8)

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Cr–Fe–Ti 17



Chromium – Iron – Vanadium

Volodymyr Ivanchenko, Tetyana Pryadko

Introduction

The constitution of many ferritic steels includes chromium and vanadium as basic alloying elements.Knowledge of the phase diagram and thermodynamic properties of the Cr-Fe-V system is essential to under-stand the behavior of such iron-based alloys.The first experimental work on the phase equilibria of the Cr-Fe-V system was undertaken by [1952Mar]who investigated the phase boundaries between the bcc and σ phases at 700°C, and found that the σ phaseextends across the ternary diagram from the Fe-V to the Cr-Fe binary side. These results were confirmedexperimentally by [1954Kor, 1957Kor1, 1957Kor2] as well as the calculations of [1977Bra1, 1977Bra2],who used a cluster model. The isothermal section presented by [1957Kor1] was almost the same as thatof [1952Mar]. The effect of vanadium on the miscibility gap was investigated by [1967Mim, 1968Mim,1970Mim, 1971Yam]. They reported that the miscibility gap in the Cr-Fe system extends into the Cr-Fe-V ternary system and takes the form of a mountain with a peak.The thermodynamic properties of some solid ternary alloys were measured by [1973Mal] and [1987Che].The enthalpies of formation of the αδ solid solution as well as enthalpies of transformation from the σ phaseto the bcc structure were determined. The activities of Fe and V in liquid Cr-Fe-V alloys were reported by[1975Fur].The composition dependence of the lattice parameters of the αδ and σ phases were presented by [1988Vas1,2002Ham] and [1952Mar].Thermodynamic computations of the phase equilibria were carried out [1973Spe, 1988Kum, 1992Lee].[1973Spe] calculated isothermal sections illustrating the αδ/σ phase relationships at 427, 700 and 900°C,and compared them with the experimental results of [1952Mar]. The results of [1973Spe] had certain dis-crepancies with respect to the results of [1952Mar], in particular with regard to the relative positions ofthe tie lines. [1973Spe] pointed out that errors could arise from the methods used to fit the experimental dataand in the summation of the three binary expressions for calculation of the ternary Gibbs energy values,mostly for the σ phase region. Later, [2001Spe] showed that calculations of the Cr-Fe-V system using dif-ferent interaction models gave different results in the σ phase domain. [1988Kum] computed the (γ/γ+αδ)and (γ+αδ/γ) phase boundaries near the Fe corner at 1350, 1250, 1150, 950, and 900°C. A single-latticemagnetic solution model was used. Isothermal sections of the Cr-Fe-V system at 1727, 1552, 1527,1027, 700, and 337°C were calculated by [1989Kau]. [1992Lee] made a detailed thermodynamic analysisof this system. The binary description for the Cr-Fe was taken from [1987And] and that for Fe-V from[1991Hua]. The Cr-V system was computed by [1992Lee]. In the last case, the depression in the Cr-Vfusion diagram was ignored. The ternary interaction parameters were obtained by optimization using experi-mental data. A three-sublattice model was used to express the Gibbs energy of the σ phase. The computedresults were presented as isothermal sections at 700 and 480°C, as a metastable miscibility gap at 480°C, avertical section at 50 at.% Fe, and solidus and liquidus isotherms at 1900, 1800, 1700, 1600 and 1500°C.Metastable vertical sections at 76 and 60 mass% Fe were also presented by suspending the σ phase inthe calculation. They show that the addition of V to Cr-Fe binary alloys increases the critical temperatureof the miscibility gap. These results are in good qualitative agreement with experimental data [1967Mim,1968Mim, 1970Mim, 1971Yam]. The results of the calculations of [1992Lee] are in better agreement withthe experimental phase equilibrium data than the work of [1973Spe] and [1989Kau], and for this reasonthey are preferred. Reviews of the Cr-Fe-V system have been presented by [1986Ban], [1988Ray] and[1994Rag]. Investigations of phase relations, structures and thermodynamics are listed in Table 1. Physicaland mechanical properties of some Cr-Fe-V alloys were studied theoretically and experimentally by[1974Rus, 1979Kat, 1981Kon, 1986Gal, 1988Vas1, 1988Vas2, 1995Nic, 1996Cer, 1997Gal, 1998Gal1,1998Gal2, 1998Gal3, 2000Som]. The experimental techniques as well as measured properties are listedin Table 4.

Cr–Fe–V 1



Binary Systems

The Fe-V system is taken from the thermodynamic assessment of [1991Hua]. At temperatures above 600°Cit is close to the assessed phase diagram of [1984Smi]. The Cr-Fe phase diagram is taken from the assess-ment of [1987And], which has no essential differences from those of [Mas2] and [1993Itk]. The Cr-V sys-tem is accepted from [1982Smi].

Solid Phases

A continuous series of bcc solid solutions αδ are formed at high temperatures, and a continuous solid solu-tion exists between the σ phases in the Fe-V and Cr-Fe systems. The variation in the lattice constants of theσ phase and the αδ solid solution in Cr-Fe-V alloys with a constant Fe content 50 at.% was presented by[1952Mar]. The lattice parameters of the σ phase rise smoothly from CrFe to VFe with a slow negativedeviation from Vegard’s law (Fig. 1a). [1952Mar] presented an isothermal section at 700°C showing linesof constant lattice parameter value for the bcc phase. The lattice parameter of the αδ phase in alloys ofcomposition VxCr1–xFe2 as a function of V concentration is shown in Fig. 1b. [2002Ham] studied theVxCr1–xFe2 alloys using XRD and Mössbauer spectroscopy. X-ray diffraction patterns showed a singlebcc phase, indicating that the alloys are atomically disordered. The lattice parameter increases with increas-ing V content, changing from about 288 pm for CrFe2 to 291 pm for VFe2, demonstrating a 1% latticeexpansion. This expansion was attributed to the replacement of Cr atoms by V atoms, which have a largeratomic radius. [1975Hag] reported a CsCl type order-disorder transformation in the Cr-Fe-V ternary systembased on DTA studies, but these results were not confirmed and are thus omitted here. The crystal structuresof phases existing in the Cr-Fe-V system are presented in Table 2.


The first determination of the liquidus surface of the Cr-Fe-V system was by [1949Jae], based on phaseequilibria in the bordering binary systems. However, as the binary fusion diagrams, as well as the meltingpoints of pure elements were known with only limited accuracy, this work is of only historical interest. Inspite of the industrial importance of the Cr-Fe-V system, no systematic study of the liquidus surface hasbeen undertaken. Liquidus and solidus isotherms have been calculated by [1992Lee]. These results are pre-sented in Fig. 2a. However, more close to reality may be the liquidus projection constructed by [1988Ray],who took into account the existence of the temperature depression in the fusion diagram of the Cr-V systemin accordance with the assessment by [1982Smi]. It is presented in Fig. 2b. But, the contours suggested areonly speculative and have not been established experimentally. The (γ/γ+αδ) and (γ+αδ/αδ) solvus sur-faces can be reconstructed from the results of thermodynamic calculations produced by [1988Kum]. Theyare presented in Fig. 3.

Isothermal Sections

As indicated above, parts of isothermal sections in the Fe corner of phase diagram are presented in Fig. 3.The isothermal section of the Cr-Fe-V system at 700°C calculated by [1992Lee] is given in Fig. 4. This iso-thermal section is in good agreement with experimental data [1952Mar]. The stable isothermal section at480°C, presented by [1992Lee], is shown in Fig. 5. The metastable miscibility gap at 480°C calculatedby [1992Lee] is shown in Fig. 6. This was calculated by suspending the σ phase during the calculation.


The temperature-composition section of the Cr-Fe-V system at 50 at.% Fe calculated by [1992Lee] is pre-sented in Fig. 7. The results of the calculations are in excellent agreement with the results of the experimen-tal study by [1957Kor1].

2 Cr–Fe–V


Thermodynamics

Some thermodynamic properties of the αδ and σ phase were measured by [1973Mal] and [1987Che]. Theenthalpy of formation of V0.3Cr0.2Fe0.5 at 1127°C (1400 K) is –1250±500 J·mol–1. The enthalpy of trans-formation from σ to αδ at a composition of V0.3Cr0.2Fe0.5 is 2300±80 J·mol–1 at 1087±10°C (1360±10 K)[1973Mal]. The enthalpy of formation of bcc solid solutions with constant Fe content (53 at.%), as well asthe heat of transformation from the σ phase to the αδ bcc solid solutions were measured by [1987Che]. Theresults are presented in Table 3. The activities of Fe and Vat 1600°C were measured by [1975Fur]. They areas follows: log γV = –0.42 + 0.70 xCr, log γFe = –0.04 – 0.18 xCr (xV = 0.2, xCr ≤ 0.2) and log γV = –0.28 +0.48 xCr, log γFe = –0.08 – 0.20 xCr (xV = 0.3, xCr ≤ 0.2). These results were used by [1992Lee] to evaluatethe interaction parameters in the system; however they could not be included in the optimization proceduredirectly.


The magnetic susceptibility and thermal expansion of Cr based alloys containing 0.77, 0.78, 1.08, and1.63 at.% V were studied by [1981Kon] over the temperature range from –196 to 127°C (77 to 400 K).Analysis of the results shows that at low concentrations, atoms of V have a local magnetic momentof ~ 1 Bohr m. in the Cr matrix. The concentration relationships of the volume modulus of elasticity forFe1–x(Cr0.5V0.5)x were derived from the alloys. It was shown that particularities in the composition depen-dences of the modulus and the probability of a Mössbauer effect are linked with the effects of a strong d-d-correlation but not with the formation of the Fe3Cr and Fe3V compounds [1988Vas1]. [1995Nic] shows thatowing to high magneto-mechanical hysteresis, some Cr-Fe-V alloys can provide a higher level of dampingthan pure iron and C-Fe alloys. The pressure dependence of the resistivity ρ(T ) in spin-density-wave(SDW) Cr based alloys of the Cr-Fe-V system were studied by [1979Kat, 1997Gal]. The behavior of theparamagnetic phase was believed to be due to the Kondo effect. The Kondo effect in Cr-Fe-V alloys wasstudied by [1974Rus]. The temperature dependence of the magnetic susceptibility of (SDW) Cr-2.7 at.%Fe alloys with different V contents (from 0.07 to 20 at.%) were measured over the range 5-400 K in a fieldof 100 Oe by [1998Gal1, 1998Gal2, 1998Gal3]. It conforms to the Curie-Weiss law, with different para-meters above and below the Neel temperature in the SDW alloys. The formation of the σ phase-based con-tinuous solid solutions, as well as the existence of a miscibility gap owing to its eutectoidal decomposition,lead to embrittlement in Cr-Fe-V alloys.

Miscellaneous

The position of the resonance level was calculated and the magnetic phase diagram for (Cr-2.7Fe)1–yVy

(0.0 ≤⊊y ≤⊊0.0501) was presented by [1986Gal]. Alloying with V destroys the antiferromagnetism ofCr in Cr-Fe alloys. TN rapidly decreases with increasing V concentration. The magnetic phase diagramsfor Cr1–xFex, V0.01Cr0.99–xFex, and VyCr1–x–yFex for y = 0.01, 0.04, 0⊊≤ x⊊≤⊊0.24 as well as the compo-sition dependence of the magnetoresistance in a field of 4T were presented by [2000Som]. The diffusioncoefficient of Cr in a sample of composition Fe-15 mass% Cr and Fe-1.7 mass% V is the same and equalto D = 2.0 exp {–13623/RT}, [1968Pav]. The time of transformation from the bcc solid solution to CrFeand VFe σ phases at 700°C is equal 350 h and 11 h, respectively [1957Kor2].

Cr–Fe–V 3



Table 1. Investigations of the Cr-Fe-V Phase Relations, Structures and Thermodynamics


[1952Mar] X-ray diffraction 700°C, Cr-Fe-V, αδ, αδ+σ, σ

[1954Kor] DTA, light microscopy, magneticproperties, Vickers hardness

Up to 1500°C, annealing at 1350°C followed byquenching in ice water, annealing at 1200°Cfollowed by tempering at 700°C,FeCr-FeV, αδ, αδ+σ, σ

[1957Kor1] DTA, X-ray diffraction, light microscopy,specific electrical resistance, Vickershardness

Up to 1500°C, annealing at 1350°C followed byquenching in iced water, annealing at 1200°Cfollowed by tempering at 700°C,along the composition lines with Cr/V=1/3, 1/1,and 1/3, αδ, αδ+σ, σ

[1970Mim] DTA, Mössbauer spectroscopy, TEM,simple bend testing and tensile test atroom temperature

Annealing at 1150°C followed by quenching iniced water, and aged isothermally at 480 or510°C, (Fe-25Cr)-(Fe-50Cr)-(Fe-35V),αδ, αδ+σ, σ, αδ(1)+αδ(2)-miscibility gap

[1971Yam] DTA, Mössbauer spectroscopy Annealing at 1150°C followed by quenching iniced water, and aged isothermally at 480 or510°C, (Fe-25Cr)-(Fe-50Cr)-(Fe-35V),αδ, αδ+σ, σ, αδ(1)+αδ(2) –miscibility gap

[1973Mal] High-temperature adiabatic calorimetry 727-1377°C (1000-1650 K)Cr0.2V0.3Fe0.5, αδ⇌σ transformation

[1973Spe] Thermodynamic calculation 427, 700, and 900°CCr-Fe-V, αδ, αδ+σ, σ,

[1975Fur] Knudsen cell effusion method with amass spectrometer

1600°CxV ≤ 0.3, xCr ≤ 0.2, liquid phase

[1975Hag] DTA, thermodynamic calculation Up to 800°C, (Fe-40V-20Cr)-(Fe-50V-22Cr)-(Fe-60V-10Cr)-(Fe-60V-5V)-(Fe-55V-5Cr)-(Fe-50V-8Cr), at.%CsCl type order-disorder transformation

[1987Che] Calorimetric investigation, based on DTAtechnique

954-1176°C (1227-1449 K),xFe = 0.53, xCr = 0, 0.047, 0.094, 0.157, 0.235,0.313, xv=1–xCr–xFe

[1988Kum] Thermodynamic calculation 900, 950, 1050, 1150, 1250, and 1350°C,Up to 2 at.% V and up to 15 at.% Crαδ+γ

[1988Vas1] X-ray diffraction, Mössbauerspectroscopy

Annealing at 1250°C followed by quenching inwater, Fe1–x(Cr0.5V0.5)x, 0.01 ≤ x ≤ 0.45,αδ bcc solid solution

[1988Vas2] Calculations, Mössbauer spectroscopy Stability of σ phase

(continued)

4 Cr–Fe–V



[1992Lee] Thermodynamic evaluation The stable isothermal sections at 480°C(753 K), and 700°C (973 K), the metastablemiscibility gap at 480°C (753 K), vertical sec-tion at 50 at.% Cr, liquidus and solidus surfaces

[2002Ham] X-ray diffraction, Mössbauerspectroscopy

As cast, Fe2Cr1–xVx, x = 0.0, 0.1, 0.2, 0.4, 0.6,0.8, and 1.0, αδ bcc solid solution





Comments/References

(α’Cr)HP tI2I4/mmmα’Cr

a = 288.2c = 288.7

at 25°C, [Mas2]

αδ, (V,αCr,αδFe)αδ,(V0.5Cr0.5)xFe1–x

αδ, CrxFe1–xαδ, VxCr1–xαδ, VxFe1–x1538 - ~1270

α1δ, VxFe1-x< 650α2δ, VxFe1–x< 650(αCr)< 1863(δFe)1538 - 1394(αFe)< 912(V)< 1910

cI2Im�3mW

a= 289.6a= 287.2a= 288.0a= 288.6a= 289.0a= 288.6a= 290.0a= 291.1a= 291.3a= 294.0a = 289.8 to 290.7a= 288.4 to 302.6a = 288.8a = 290.8a = 292.0a = 293.3a = 296.8

a = 288.48

a = 293.15

a = 286.65

a = 302.4

[1988Vas1]x = 0.02 annealed at 1400°Cx = 0.1x = 0.2x = 0.3x = 0.4x = 0.5x = 0.6x = 0.7x = 0.8x = 0.90 ≤ x ≤ 1 annealed at 800°C [V-C2]0 ≤ x ≤ 1 [1982Smi]x = 0.2 at 25°C, [1984Smi]x = 0.4x = 0.5x = 0.6x = 0.80 < x ≲ 0.18

0.75 ≲ x ≤ 1

pure Cr at 25°C [Mas2]

[Mas2]


pure V at 25°C [Mas2]

(continued)

Cr–Fe–V 5






Comments/References

(εFe) hP2P63/mmcMg

a = 246.8c = 396.0


γ, (γFe)1394 - 912

cF4Fm�3mCu

a = 364.67 at 915°C [V-C2, Mas2]

σ, (V,Cr)Feσ, CrFe< 830

σ, VFe~ 1200 - ~ 650

tP30P42/mnmσCrFe

a = 879.66 ± 0.06c = 455.82 ± 0.03a = 879.68 ± 0.05c = 455.85 ± 0.03a = 879.61 ± 0.04c = 456.05 ± 0.03a = 886.5 to 901.5c= 460.5 to 464.2

49.5 at.% Cr, [V-C2]

48.2 at.% Cr, [V-C2]

47.8 at.% Cr, [V-C2]at 25°C29.6-60.1 at.% V, [1984Smi][2005Ust]



Quantity, per mole ofatoms [kJ, mol, K]

Comments

0.47V+0.53Fe ⇌ V0.47Fe0.53 1170 –5.261 bcc solid solutionformation [1987Che]

0.423V+0.047Cr+0.53Fe⇌V0.423Cr0.047Fe0.53

1176 –2.695 bcc solid solutionformation [1987Che]

0.376V+0.0.094Cr+0.53Fe⇌V0.376Cr0.094Fe0.53


0.313V+0.157Cr+0.53Fe⇌V0.313Cr0.157Fe0.53


0.235V+0.235Cr+0.53Fe⇌V0.235Cr0.235Fe0.53

1129 1.819 bcc solid solutionformation [1987Che]

0.157V+0.313Cr+0.53Fe⇌V0.157Cr0.313Fe0.53


0.094V+0.376Cr+0.53Fe⇌V0.094Cr0.376Fe0.53


0.047V+0.423Cr+0.53Fe⇌V0.047Cr0.423Fe0.53


0.47Cr+0.53Fe⇌ Cr0.47Fe0.53 954 5.582 bcc solid solutionformation [1987Che]

σV0.47Fe0.53⇌ αδ V0.47Fe0.53 1145 3.183 [1987Che]

(continued)

6 Cr–Fe–V



Quantity, per mole ofatoms [kJ, mol, K]

Comments

σV0.423Cr0.047Fe0.53⇌αδ V0.423Cr0.047Fe0.53

1131 2.830 [1987Che]

σV0.376Cr0.094Fe0.53⇌αδ V0.376Cr0.094Fe0.53

1111 2.676 [1987Che]

σV0.313Cr0.157Fe0.53⇌αδ V0.313Cr0.157Fe0.53

1075 2.314 [1987Che]

σV0.235Cr0.235Fe0.53⇌αδ V0.235Cr0.235Fe0.53

1023 1.858 [1987Che]

σV0.157Cr0.313Fe0.53⇌αδV0.157Cr0.313Fe0.53

944 1.216 [1987Che]

Table 4. Investigations of the Cr-Fe-V Materials Properties

References Method / Experimental Technique Type of Property

[1957Kor2] Magnetic saturation measurement Kinetics of phase transformation from αδbcc solid solution to σ (V,Cr)Fe phase

[1968Pav] Radioisotopic tracer technique Diffusivity of Cr

[1974Rus] Electrical resistivity from 0.3 to 40 K Kondo effect

[1979Kat] Electrical resistivity from 4.2 K to roomtemperature at high pressures

The dependence of resistivity minimum onpressure

[1981Kon] Dilatometry, magnetic measurements Magnetic susceptibility, thermal expansion

[1986Gal] Electroresistivity at temperatures from4.2 to 373 K

Magnetic phase diagram

[1988Vas1] Measuring of the resonance ultrasonicfrequency excited by piezoelectric vibrator

Volume elasticity

[1989Vas] Measuring of electromotive force Electronic structure, d-state density functionnear the Fermi level

[1995Nic] Internal friction, coercitivity,magnetostriction

Magnetoelastic properties

[1997Gal] The pressure dependence of ρ(T ) Kondo effect in paramagnetic phase

[1998Gal1] Magnetic susceptibility χ(T ) in thetemperature range 5-400 K

Magnetic behavior in Cr-base spin-density-waves alloys


Magnetic behavior in Cr-base spin-density-waves and paramagnetic alloys


Magnetic behavior in Cr-base spin-density-waves alloys

[2000Som] Squid magnetometry and Faraday balance Magnetization and magnetoresistance

Cr–Fe–V 7



Fig. 1a. Cr-Fe-V. Lattice parameters of the σ phase solid solution vs composition

8 Cr–Fe–V


Fig. 1b. Cr-Fe-V. The composition dependencies of lattice parameters of the αδ VxCr1–xFe2 solid solution vsComposition

Cr–Fe–V 9



Fig. 2a. Cr-Fe-V. Liquidus/solidus surface projections calculated by [1992Lee]

10 Cr–Fe–V


Fig. 2b. Cr-Fe-V. Liquidus surface

Cr–Fe–V 11



Fig. 3. Cr-Fe-V. Isothermal sections in the Fe corner

12 Cr–Fe–V


Fig. 4. Cr-Fe-V. Calculated isothermal section at 700°C projection

Cr–Fe–V 13



Fig. 5. Cr-Fe-V. Calculated isothermal section at 480°C

14 Cr–Fe–V


Fig. 6. Cr-Fe-V. Calculated metastable isothermal section at 480°C

Cr–Fe–V 15



Fig. 7. Cr-Fe-V. Calculated temperature - composition section at 50 at.% Fe

16 Cr–Fe–V


References[1949Jae] Jaenecke, E., “Fe-Cr-V” (in German) in “Kurzgefasstes Handbuch aller Legierungen”, Win-

ter Verlag, Heidelberg, 612–613 (1949) (Phase Diagram, Phase Relations, 1)[1952Mar] Martens, H., Duwez, P., “Phase Relationships in the Iron-Chromium-Vanadium System”,

Trans. Amer. Soc. Metals, 44, 484–494 (1952) (Crys. Structure, Phase Diagram, Phase Rela-tions, Experimental, #, 10)

[1954Kor] Kornilov, I.I., Matveeva, N.M., “The Continues Solid Solutions of the FeCr and FeV MetallicCompounds” (in Russian), Dokl. Akad. Nauk SSSR, 98, 787–790 (1954) (Morphology, PhaseRelations, Phase Diagram, Mechan. Prop., Experimental, #, 9)

[1957Kor1] Kornilov, I.I., Matveeva, N.M., “Phase Transformations in the Iron-Chromium-VanadiumSystem”, Russ. J. Inorg. Chem., 2(2), 196–216 (1957), translated from Zh. Neorg. Chem. 2(2), 355–366 (1957) (Crys. Structure, Morphology, Phase Relations, Experimental, Electr.Prop., Mechan. Prop., #, 9)

[1957Kor2] Kornilov, I.I., Matveeva, N.M., “The Rate of Transformation of the α Solid Solution into theσ - Phase in the System Fe-Cr-V”, Russ. J. Inorg. Chem., 2, 275–287 (1957),, translated fromZh. Neorg. Khim., 2(6), 1383–1391 (1975) (Phase Relations, Experimental, Magn. Prop.,Kinetics, 11)

[1967Mim] Mima, G., Yamaguchi, M., Takahashi, J., “Ageing Behaviour of Ternary Iron-Chromium-Vanadium Alloys Around 500°C” (in Japanese), J. Jpn. Inst. Met., 31, 470–475 (1967) (Crys.Structure, Morphology, Phase Diagram, Phase Relations, Experimental, 12)

[1968Mim] Mima, G., Yamaguchi, M., Takahash, J., “Aging of Iron-Chromium and Iron-Chromium-Vanadium Alloys at about 500°C”, Trans. Jpn. Inst. Met., 9, Suppl. S, 407–413 (1968) (Crys.Structure, Morphology, Phase Diagram, Phase Relations, Experimental, 9)

[1968Pav] Pavlinov, L.V., Isadzhanov, E.A., Smirnov, V.P., “Diffusion of Chromium in Alloys of Ironwith Vanadium and Chromium”, Phys. Met. Metallogr. USSR, 25(5), 206–207 (1968), trans-lated from Phyz. Met. Metalloved., 25(5), 959–960 (1967) (Experimental, Transport Phenom-ena, 4)

[1970Mim] Mima, G., Yamaguchi, M., “Constitutional Investigations on the Miscibility Gap in Iron-Chromium and Iron-Chromium-Vanadium Systems”, Trans. Jpn. Inst. Met., 11, 239–244(1970) translated from J. Jpn. Inst. Met., 33, 1308–1313 (1969) (Crys. Structure, Morphol-ogy, Phase Diagram, Phase Relations, Experimental, Thermodyn., 21)

[1971Yam] Yamaguchi, M., Umakoshi, Y., Mima, G., “On the Miscibility Gap in the Fe-Cr-X (X = Cu,Mn, Mo, Ni, V, Si and Al) System” in “Proc. Int. Conf. Sci. Technol. Iron Steel Tokyo”, 11,1015–1019 (1971) (Crys. Structure, Morphology, Phase Relations, Experimental, 35)

[1973Mal] Malinsky, I., Claisse, F., “Thermodynamic Properties of α- and σ-Phases in V + Fe andV + Cr + Fe”, J. Chem. Thermodyn., 5, 911–916 (1973) (Thermodyn., Experimental, 7)

[1973Spe] Spencer, P.J., Counsell, J.F., “A Thermodynamic Calculation of the Iron-Chromium-Vanadium Equilibrium Diagram”, Z. Metallkd., 64, 662–665 (1973) (Calculation, Phase Dia-gram, Phase Relations, Thermodyn., 15)

[1974Rus] Rusby, R.L., Coles, B.R., “Resistance Minima in (V-Cr)Fe Alloys”, J. Phys. F - Metal Phy-sics, 4(6), L161-L166 (1974) (Experimental, Electr. Prop., 22)

[1975Fur] Furukawa, T., Kato, E., “Thermodynamic Properties of the Fe-V, Fe-V-Cr Alloys at 1600°Cby Mass Spectrometry” (in Japanese), Tetsu-To-Hagane, 61, 3050–3059 (1975) (Thermo-dyn., Experimental, 9)

[1975Hag] Hagiwara, M., Seki. J.-I., Suzuki, T., “CsCl Type Superlattice Formation in V-Mn-Fe, Fe-V-Cr, and V -Mn-Cr Ternary Alloys” (in Japanese), J. Jpn. Inst. Met., 39, 402–408 (1975)(Phase Diagram, Phase Relations, Thermodyn., Calculations, Experimental, 27)

[1977Bra1] Brauwers, M., “Occurrence of the σ Phase Computed from a Cluster Model”, J. Phys. F: Met.Phys., 7(6), 921–927 (1977) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn.,Calculation, 17)

[1977Bra2] Brauwers, M., Brouers, F., “On the Occurrence of the σ Phase in Transition-Metal Alloys”,Philos. Mag., 35(4), 1105–1109 (1977) (Crys. Structure, Phase Diagram, Phase Relations,Thermodyn., Calculation, 7)

Cr–Fe–V 17



[1979Kat] Katano, S., Mori, N., “Pressure Effect on the Resistivity Minimum in (Cr92V8)-1.0 at.% FeAlloy”, J. Phys. Soc. Jpn., 46(2), 691–692 (1979) (Electr. Prop., Experimental, 7)

[1981Kon] Kondorskii, E.I., Kostina, T.I., Medvedchikov, V.P., “Magnetic Susceptibility of Cr-Co-VandCr-Fe-V Alloys of Chromium”, Moscow University Physics Bulletin, USA, 36(5), 25–30(1981), translated from Vestn. Mosk. Univ., Ser. 3 (Fizika, Astronomiya), 36(5), 22–26(1981) (Thermodyn., Experimental, Magn. Prop., 8)

[1982Smi] Smith, J.F., Bailey, D.M., Carlson, O.N., “The Cr-V (Chromium-Vanadium) System”, Bull.Alloy Phase Diagrams, 2(4), 469–473 (1982) (Crys. Structure, Phase Diagram, Phase Rela-tions, Review, Thermodyn., #,17)

[1984Smi] Smith, J.F., “The Fe-V (Iron-Vanadium) System”, Bull. Alloy Phase Diagr., 5(2), 184–193(1984) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Assessment, #, 99)

[1986Ban] Bannykh, O.A., Drits, M.E., “Iron-Vanadium-Chromium” (in Russian) in “Phase Diagramsof Binary and Multicomponent Systems Based on Iron, Reference Book”, Metallurgia, Mos-cow, 189–191 (1986) (Phase Diagram, Phase Relations, Review, 6)

[1986Gal] Galkin, V.Yu., Tugushev, V.V., Tugusheva T.E., “Resonance Impurity Scattering in Dilute Cr-Fe-M (M = Mn, V) Alloys”, Sov. Phys. - Solid State, USA, 28(8), 1282–1287 (1986), trans-lated from Fizika Tverdogo Tela, USSR, 28(8), 2290–2298 (1986) (Phase Relations, Calcu-lation, Experimental, Electronic Structure, Magn. Prop.,16)

[1987And] Andersson, J.O., Sundman, B., “Thermodynamic Properties of the Cr-Fe System”, Calphad,11(1), 83–92 (1987) (Crys. Structure, Phase Diagram, Phase Relations, Review, 51)

[1987Che] Cheng, S., Fan, M., “A Calorimetric Study of Solid Cr-Fe-V Alloys”, Z. Metallkd., 78(11),815–817 (1987) (Experimental, Thermodyn., 7)

[1988Kum] Kumar, K.C.H., Raghavan, V., “BCC-FCC Equilibrium in Ternary Iron Alloys”, J. AlloyPhase Diagrams, 4(1), 53–71 (1988) (Experimental, Phase Relations, Phase Diagram, Ther-modyn., 27)

[1988Ray] Raynor, G.V., Rivlin, V.G., “Cr-Fe-V” in “Phase Equilibria in Iron Ternary Alloys”, Inst.Metals, London, 406, 332–341 (1988) (Phase Diagram, Phase Relations, Thermodyn.,Review, 10)

[1988Vas1] Vasman, G.I., Demidenko, V.S., Muslov, S.A., “Volume Modulus of Elasticity and the Prob-ability of the Mössbauer Effect in Cr-Fe-VAlloys” (in Russian), Metallofizika, 10(6), 68–73(1988) (Crys. Structure, Phase Relations, Electronic Structure, Phys. Prop., Experimental, 15).

[1988Vas2] Vasman, G.I., Kalyanov, A.P., Demidenko, V.S., “Role of Electronic Structure in the Forma-tion of the Properties and Characteristics of Interatomic Interaction in Alloys of Iron withChromium and Vanadium”, Sov. Phys. J., 31(12), 1021–1026 (1988), translated from Izv.Vyss. Uchebn. Zaved., Fiz., 31(12), 77–83 (1988) (Phase Relations, Calculation, Experimen-tal, Electronic Structure, 17).

[1989Kau] Kaufman, L., “Computer Bases Thermochemical Modeling of Multicomponent Phase Dia-grams” in “Alloys Phase Stability”, Stocks, G.M., Gonis, A. (Eds.), Kluwer Acad. Publ.,145–175 (1989) (Phase Diagram, Phase Relations, Experimental, 43)

[1989Vas] Vasman, G.I., Demidenko, V.S., “The Connection of the Thermopower with Pecularities ofElectronic-Structure of Concentrated Alloys Fe-Cr-V” (in Russian), Izv. Vyss. Uchebn.Zaved., Fiz., 32(5), 115–117 (1989) (Experimental, Phys. Prop., Electronic Structure, 7)

[1991Hua] Huang, W., “A Thermodynamic Evaluation of the Fe-V-C System”, Z. Metallkd., 82(5),391–401 (1991) (Calculation, Phase Diagram, Phase Relations, Thermodyn., Experimen-tal, #, 54)

[1992Lee] Lee, B.J., “A Thermodynamic Evaluation of the Cr-Fe-V System”, Z. Metallkd., 83(5),292–299 (1992) (Calculation, Phase Diagram, Phase Relations, Thermodyn., #, 34)

[1993Itk] Itkin, V.P., “Cr-Fe (Chromium-Iron)” in “Phase Diagrams of Binary Iron Alloys”, Okamoto,H. (Ed.), ASM International, Materials Park, OH (1993) (Review, Phase Diagram, PhaseRelations, 17)

[1994Rag] Raghavan, V., “Cr-Fe-V (Chromium-Iron-Vanadium)”, J. Phase Equilib., 15(5), 538–539(1994) (Phase Diagram, Phase Relations, Review, 8)

18 Cr–Fe–V


[1995Nic] Nichipuruk, A.P., Potekhin, B.A., Gerasimov, Y.G., Potekhin, A.B., Mansurov, S.V., “Damp-ing Decrement of Mechanical Oscillations and Magnetic Properties of Iron and Fe-Cr,Cr-Fe-VAlloys” (Russian), Fiz. Met. Metalloved., 80(1), 49–53 (1995) (Experimental, Magn.Prop., Mechan. Prop., 5)

[1996Cer] Cermak, J., Ruizickova, J., Pokorna, A., “Tracer Diffusion of V-48 Along High-DiffusivityPaths in Cr-Fe Ferritic Alloys”, Scr. Mater., 34(3), 429–433 (1996) (Phase Relations, Experi-mental, Transport Phenomena, 9)

[1997Gal] Galkin, V.Y., de Camargo, P.C., Budko, S.L., Saitovitch, E.B., Fawcett, E., “Kondo Effect inthe Paramagnetic Phase of CrFe Alloy with the Spin-Density-Wave Suppressed by Pressureor V Doping”, J. Appl. Phys., 81(8), Part 2A, 4176–4178 (1997) (Phase Relations, Experi-mental, Electr. Prop., Magn. Prop.,14)

[1998Gal1] Galkin, V.Y., Ortiz, W.A., Fawcett, E., “Frustration in the Paramagnetic Phase of Spin-Den-sity-Wave CrFeVAlloys”, J. Appl. Phys, 83(11), Part 2, 7384–7386 (1998) (Phase Relations,Experimental, Magn. Prop.,9)

[1998Gal2] Galkin, V.Y., Ali, N., Fawcett, E., de Camargo, P.C., “The Moment of Fe in Cr1–xVx Host: I.The Paramagnetic Phase”, J. Phys.-Condes. Matter., 10(22), 4901–4909 (1998) (Phase Rela-tions, Experimental, Magn. Prop.,2)

[1998Gal3] Galkin, Y., Ortiz, W.A., Fawcett, E., Ali, N., de Camargo, P.C., “The Moment of Fe in aCr1–xVx Host: II. Effect of Magnetic Field in the Spin-Density-Wave Phase”, J. Phys.-Con-des. Matter., 10(22), 4911–4917 (1998) (Phase Relations, Experimental, Magn. Prop., 14)

[2000Som] Somsen, Ch., Acet, M., Nepecks, G., Wassermann, E.F., “The Effect of Magnetic Orderingon the Giant Magnetoresistance of Cr-Fe-V and Cr-Fe-Mn”, J. Magn. Magn. Mater., 208,191–206 (2000) (Experimental, Magn. Prop., Phase Relations, 26)

[2001Spe] Spenser, P.J., “Computational Thermochemistry: from its Early Calphad Days to a Cost-Effective Role in Materials Development and Processing”, Calphad, 25(2), 163–174(2001) (Calculation, Phase Relations, Thermodyn., 31)

[2002Ham] Hamasha, K.M., Al-Omari, I.A., Mahmood, S.H., “Mössbauer and Structural Studies of theFe2Cr(1–x)Vx Alloy System”, Physica B, 321(1-4), 154–158 (2002) (Crys. Structure, Experi-mental, Magn. Prop., 8)

[2005Ust] Ustinovshikov, Y., Pushkarev, B., Sapegina, I., “Phase Transformations in Alloys of the Fe-VSystem”, J. Alloys Compd., 398, 133–138 (2005) (Crys. Structure, Morphology, Phase Dia-gram, Phase Relations, Experimental, #, 9)



Cr–Fe–V 19



Chromium – Iron – Zirconium

Gabriele Cacciamani

Introduction

Several experimental investigations of the Cr-Fe-Zr phase equilibria are available in literature: they are sum-marized in Table 1. However, no isothermal section or liquidus surface was investigated in the completecomposition range. Monovariant liquidus lines, isothermal section at 700°C and a reaction scheme extendedto the whole composition range were extrapolated in the assessment compiled by [1992Rag] on the basis ofthe partial experimental information available in 9 papers appeared between 1959 and 1986.Typical purities of the elements employed for sample preparation were: iodide zirconium (99.96 mass%),electrolytic iron and chromium (99.96 mass%). In a few cases metal purities were not indicated, while ina few others N, O, Si and selected metal impurities were analyzed and reported. Alloys were generally pre-pared by arc melting.

Binary Systems

The Cr-Zr system is accepted according to the assessment by [2003Per]. Fe-Zr phase diagram is taken fromthe recent re-investigation by [2002Ste]: according to this paper, in particular, Zr6Fe23 (or ZrFe3 accordingto some authors) and the cubic form of Zr2Fe are metastable phases stabilized by oxygen contamination. Cr-Fe is accepted from [Mas2].

Solid Phases

Structure data are summarized in Table 2. No ternary phase has been reported in the Cr-Fe-Zr system. How-ever binary Laves phases show extended solubility ranges along the ZrFe2-ZrCr2 quasibinary section. Thevariation of the lattice parameters with composition along the section are reported in Fig. 1 according to[1970Kan].

Quasibinary Systems

The ZrFe2-ZrCr2 is an approximately (due to the solubility range of the end member phases) quasibinarysection first investigated by [1959Mos] and [1962Sve] and subsequently revised by [1972Sve] and reportedalso in the review paper by [1973Sve]. It is shown in Fig. 2.


Polythermal equilibria involving α- β- and γ-Fe have been investigated by [1970Sve]. In particular, invar-iant equilibria between these phases and liquid have been identified. As in other papers by the same authors,however, an η phase (presumably Zr6Fe23) is reported as a stable phase, in contradiction with the acceptedFe-Zr phase diagram. For this reason it is not possible to confirm and report here the proposed invariantequilibria.


The liquidus surface in the Fe-Cr-ZrCr2-ZrFe2 region has been reported by [1963Sve]. It looks incompletebecause only one monovariant line (from the Fe rich to the Cr rich binary eutectics) is reported. It has beencompleted here and adapted to the accepted binary and quasibinary systems in order to be shown in Fig. 3.Notice however that it has to be considered poorly reliable. [1978Hao] confirmed that no ternary eutectic ispresent in the Fe rich corner of the phase diagram.

Cr–Fe–Zr 1



Isothermal Sections

The Fe rich corner of the 1000°C isothermal section has been investigated by [1963Sve]: it is reported inFig. 4 with corrections in order to agree with the accepted Fe-Zr binary system. Three partial isothermal sec-tions at 875, 800 and 700°C have been investigated by [1979Mal] in the Zr-ZrCr2-ZrFe2 composition range.The section at 700°C is shown in Fig. 5 slightly modified in order to be consistent with the quasibinary sec-tion ZrFe2-ZrCr2.


A temperature-composition partial section at 2 at.% Cr has been determined by [1963Sve] in two versions:at equilibrium and in metastable conditions respectively. In the equilibrium section however an η phase(presumably Zr6Fe23) and related equilibria are present, in contradiction with the accepted Fe-Zr phase dia-gram. Then it has been considered unreliable and it is not reported here.


Cr-Fe-Zr is one of the ternary systems relevant to Zircaloys, which are widely used as fuel cladding materialin nuclear industry due to their low neutron-capture cross-section, high mechanical strength, high thermalconductivity and good corrosion resistance. In these alloys Fe and Cr, which are essentially insoluble inZr at temperatures lower than about 600°C, are present in the form of Zr(Fe,Cr)2 precipitates. This is oneof the reasons why several studies have been performed on properties and characteristics of the Zr(Fe,Cr)2 phase.ZrCr2 and Zr(Fe,Cr)2, moreover, exhibit high hydrogen absorption capacity in both hexagonal and cubicstructures. Hydrogen concentrations so far reported are about 3 H atoms per formula unit while the theore-tical hydrogen capacity reaches 6.3 H atoms per formula unit.

Miscellaneous

Pure Zr and Cr-Fe-Zr alloys were used as test samples for the accurate determination of phase boundarytemperatures from DTA measurements elaborated by means of an extrapolation equation proposed by[1991Zhu].Crystal structure, twinning, dislocation structures and orientation relationships of MgCu2 type precipitatesin Zr rich alloys have been studied by [1986Men1]. Polytypic structures assumed by Zr(Fe,Cr)2 Lavesphases have been studied by [1986Men2, 1991Bur] both in bulk materials and in precipitates included inZn rich alloys. Samples were analyzed by [1991Bur] in as cast conditions as well as after heat treatmentat 850 or 900°C. The variety of polytypic structures observed in precipitates was attributed to the latticemisfit and associated stresses between precipitates and the αZr matrix.Several experimental investigations have been carried out on the Zr-~1.15mass% Cr-~0.1mass% Fe alloyused as a fuel sheathing for water cooled nuclear reactors: mechanical properties in cold-worked and irra-diated alloys were studied by [1974Hol1, 1974Hol2]; microstructure and precipitation kinetics of Lavesphase precipitates have been investigated by [1983Dos] by electron microscopy; thermoelectric power inrelation to dislocation density for a cold rolled alloy was studied by [1993Sun, 1994Sun1, 1994Sun2] onthe basis of X-ray line broadening analysis; abnormal grain growth during ageing of quenched sampleswas investigated by [1997Nor].Thermal expansion curves and magnetic properties of amorphous CrxFe90–xZr10 alloys have been investi-gated by [1994Lu] providing experimental evidence that the invar effect of these alloys is also related tothe instability of the ferromagnetism.Polarization curves for ZrCr2, ZrFe2 and Zr(Fe0.67Cr0.33)2 have been measured at room temperature[1995Mur] and 250°C [1996Mur]. Polarization curves of the ternary alloy at high temperature resulted morestable and reproducible than those at room temperature. In general Cr-containing alloys had lower anodiccurrent than the Cr-free ones.

2 Cr–Fe–Zr


X-ray diffraction and Mössbauer spectroscopic studies have been performed by [1999Coa] on Cr-Fe-ZrLaves phases and their hydrides while magnetic and spin-glass properties of the MgZn2 type Laves phasehave been investigated by [2001Coa1, 2001Coa2].Phase stabilities of metastable solid solutions (amorphous, bcc, hcp) and stable Zr(Fe,Cr)2 phase have beencalculated by [2002Rod] using Miedema's model in order to explain alloy amorphization and Fe depletionof the Zr(Fe,Cr)2 precipitates in Zircaloy irradiated at intermediate temperatures.Volume changes upon hydrogenation in ZrCr2, Zr(Fe,Cr)2 and other ZrMe2 based alloys has been studied by[2003Dor]. They found that the volume change per hydrogen atom tends to decrease with hydrogen concen-tration (in the hexagonal structure faster than in the cubic one) and suggested that different hydrogen sitesand their occupation are responsible for such volume change. Hydrogenation effects on Cr-Fe-Zr alloyshave also been studied by [2003Sor] by Mössbauer spectroscopy.High temperature oxidation process of the hexagonal Laves phase oxidized in an open furnace has been stu-died by [2004Boz]. Oxidation was found to proceeds through the increasing presence of α-Fe2O3 and thestructural evolution of Zr oxides.The grain refinement process during equal channel angular pressing in commercial purity Zr7O2 alloy(including Hf, Fe, Cr, H, N, C and O) was investigated by [2005Cao] in view of an improvement of themechanical properties of this material for nuclear applications.

Table 1. Investigations of the Cr-Fe-Zr Phase Relations, Structures and Thermodynamics


[1959Mos] DTA, XRD, LOM ZrCr2-ZrFe2 temperature/composition section

[1962Sve] DTA, XRD, chemical analysis ZrCr2-ZrFe2 temperature/composition section

[1963Sve] DTA, XRD, microstructureanalysis, dilatometry,magnetometry

Liquidus surface, 1000°C isothermal section and oneT/C vertical section in the Cr-ZrCr2-ZrFe2-Fecomposition range

[1970Kan] XRD, magnetometry Crystal structure and magnetic moment in the ZrCr2-ZrFe2 section

[1970Sve] DTA, XRD, microstructureanalysis, dilatometry,magnetometry

Polythermal equilibria in the Fe rich corner

[1972Sve] DTA, XRD, chemical analysis ZrCr2-ZrFe2 temperature/composition section

[1979Mal] XRD, microstructure analysis Isothermal sections at 875, 800 and 700°C in the Zr-ZrCr2-ZrFe2 composition range

[1985Yu] XRD Zr(Cr1–xFex)2 at x = 0.5

[1986Uch] XRD Laves phase at ZrCrFe1.6

[1995Sou] neutron diffraction Zr(Cr1–xFex)2 Laves phase

Cr–Fe–Zr 3



Table 2. Crystallographic Data of Solid Phases Cr-Fe-Zr




Comments/References

αδ, (αCr,δFe,αFe)(αCr)< 1863(δFe)1538 - 1394(αFe)< 912

cI2Im�3mW a = 288.48

a = 293.15

a = 286.65

< 0.6 at.% Zr at 1592°C [Mas2, V-C2]at 25°C [Mas2]pure Fe at 1360°C [Mas2, V-C2]


(α’Cr) tI2I4/mmmα’Cr

a = 288.2c = 288.7

at 25°C, HP [Mas2]

(γFe)1394 - 912

cF4Fm�3mCu

a = 364.67a = 357.3

[Mas2]at 25°C [V-C2]

(βZr)1855 - 831

cI2Im�3mW

a = 360.90< 8 at.% Cr at 1332°C[Mas2]

(αZr)< 863

hP2P63/mmcMg

a = 323.16c = 514.75

< 0.5 at.% Cr at 836°Cat 25°C [V-C2]

(ωZr) hP3P6/mmmωTi

a = 503.6c = 310.9

< 4 at.% Crat 25°C, HP, at 1 atm [Mas2]

σ, CrFe830 - 440

tP30P42/mnmCrFe

a = 879.66c = 455.82

43 to 48 at.% Cr [Mas2, V-C2]

γZr(Cr1–xFex)21677 - 1625

hP12P63/mmcMgZn2

a = 510.2c = 828.9a = 511.1c = 834.1a = 511.3c = 830.9a = 508.7c = 831.2a = 505.7c = 829.7a = 502.9c = 824.1

C14 structure. 0 < x ≲ 0.8at 20°C [1997Kur], x = 0

at 300°C [1997Kur], x = 0

neutron diffraction, refinement [1995Sou], x = 0

at x = 0.1 (neutron diffraction, refinement)[1995Sou]at x = 0.2 (neutron diffraction, refinement)[1995Sou]at x = 0.5 [1985Yu]

(continued)

4 Cr–Fe–Zr





Comments/References

a = 500.1c = 818.0a = 495.8c = 813.3

at ZrCrFe1.6 [1984Pou] quoted in [1986Uch]

at ZrCrFe1.6 (sample contaminated by 0.3 at.% O)[1986Uch]

βZr(Cr1–xFex)21625 - 1546

hP24P63/mmcMgNi2

a = 510.2c = 1662

C36 structure. Dissolves ~1 at.% Feat x = 0.025 and 1500°C [1972Sve]

αZr(Cr1–xFex)2< 1560

cF24Fd�3mMgCu2 a = 720.4

C15 structure. Dissolves ~7 at.% Fe at 800°C[1972Sve]at x = 0 [1995Sou]

ZrFe2< 1673

cF24Fd�3mMgCu2 a = 702 to 709

C15 structure. Dissolves ~9 at.% Cr[1972Sve]at 0.275 ≤ x(Zr) ≤ 0.344 [2002Ste]

ZrFe2(h)1345 - 1240

hP24P63/mmcMgZn2

a = 495c = 1614

0.265 ≤ x(Zr) ≤ 0.27 [2002Ste]C14 structure.

Zr2Fe951 - 780

tI12I4/mcmAl2Cu

a = 638c = 560

0.667 ≤ x(Zr) ≤ 0.672 [2002Ste]

Zr3Fe< 851

oC16CmcmBRe3

a = 332b = 1100c = 882

0.748 ≤ x(Zr) ≤ 0.754 [2002Ste]

Zr6Fe23(metastable)

cF116Fm�3mTh6Mn23

a = 1172 metastable, oxygen stabilized0.206 ≤ x(Zr) ≤ 0.216 [2002Ste]

Zr2Fe (metastable)951 - 780

cF96Fd�3mTi2Ni

a = 1221 metastable, oxygen stabilized [2002Ste]

Cr–Fe–Zr 5



Fig. 1. Cr-Fe-Zr. Laves phase lattice parameters versus composition across the ZrFe2-ZrCr2 quasibinary section

6 Cr–Fe–Zr


Fig. 2. Cr-Fe-Zr. ZrFe2-ZrCr2 quasibinary section

Cr–Fe–Zr 7



Fig. 3. Cr-Fe-Zr. Partial liquidus surface projection

8 Cr–Fe–Zr


Fig. 4. Cr-Fe-Zr. Partial isothermal section at 1000°C

Cr–Fe–Zr 9



Fig. 5. Cr-Fe-Zr. Partial isothermal section at 700°C

10 Cr–Fe–Zr


References[1959Mos] Moss, S.C., “Structural Relationships in the Pseudobinary System ZrFe2-ZrCr2”, Thesis,

Mass. Inst. Technology (1959) (Experimental, Phase Diagram)[1962Sve] Svechnikov, V.N., Spektor, A.Ts., “Phase Diagram of the ZrCr2-ZrFe2 System” (in Russian),

Voprosy Fiz. Met. Metalloved., Akad. Nauk Ukr. SSR, Inst. Metallofiziki, 16, 145–152 (1962)(Experimental, Phase Diagram, 6)

[1963Sve] Svechnikov, V.N., Spektor, A.Ts., “Phase Diagram of the Ternary System Cr-Fe-Zr” (in Rus-sian), Voprosy Fiz. Met. Metalloved., Akad. Nauk Ukr. SSR, Inst. Metallofiziki, 17, 181–186(1963) (Experimental, Phase Diagram, 11)

[1970Kan] Kanematsu, K., Fujita, Y., “Magnetic Moment in Laves Phase Compound. I. Zr(Fe1–xVx)2and Zr(Fe1–xCrx)2”, J. Phys. Soc. Jpn., 29(4), 864–868 (1970) (Crys. Structure, Experimen-tal, Magn. Prop., 9)

[1970Sve] Svechnikov, V.N., Spektor, A.Ts., “Effect of Zr on α, γ, δ Range of Polymorphism in theTernary System” (in Russian), Akad. Nauk Ukr. SSR, Metallofizika, 32, 33–38 (1970)(Experimental, Phase Diagram, 11)

[1972Sve] Svechnikov, V.N., Markiv, V.Ya., Pet’kov, V.V., “Interaction of the Intermetallide ZrCr2with Certain Compounds of Zirconium with Iron, Cobalt and Nickel” (in Russian), Akad.Nauk Ukr. SSR, Metallofizika, 42, 112–117 (1972) (Crys. Structure, Experimental, PhaseDiagram, 10)

[1973Sve] Svechnikov, V.N., Kocherzhinsky, Yu.A., Markiv, V.Ya., Pet’kov, V.V., “Laves Phases inTransition Metal Systems of the IV-VII Groups of Periodic Systems” (in Russian), Akad.Nauk Ukr. SSR, Metallofizika, 46, 35–45 (1973) (Experimental, Phase Diagram, Review, 74)

[1974Hol1] Holt, R.A., “The Effect of Metallurgical Condition and Irradiation on Strength and Ductilityof Zr-1.14% Cr-0.1% Fe”, J. Nucl. Mater., 50, 207–214 (1974) (Experimental, Mechan.Prop., 17)

[1974Hol2] Holt, R.A., “Work Hardening and Ductility of Zr-1.14% Cr-0.1% Fe”, J. Nucl. Mater., 51,309–320 (1974) (Experimental, Mechan. Prop., 23)

[1978Hao] Haour, G., Mollard, F., Lux, B., Wright, G., “New Eutectics Based on Fe, Co or Ni”,Z. Metallkd., 69(1), 26–32 (1978) (Experimental, Interface Phenomena, Mechan. Prop.,Morphology, Phase Diagram, Phase Relations, Phys. Prop., 24)

[1979Mal] Malakhova, T.O., “Study of Phase Diagrams of the Zirconium Part of the Zr-Fe, Zr-Cr-Feand Zr-Cr-Cu Systems” (in Russian), Splavy At. Energ., 123–130 (1979) (Experimental,Phase Diagram, 19)

[1983Dos] Dosen, K., Robinson, J.W., Northwood, D.O., “Electron Metallographic Examination ofZrCr2 Precipitation in Zr-1.15 mass% Cr-0.1 mass% Fe” in “Microstructural Science” Pro-ceedings of the Fifteenth Annual Technical Meeting of the International MetallographicSociety, Elsevier, New York, USA, 11, 235–242 (1983) (Experimental, Morphology, 15)

[1984Pou] Pourarian, F., Wallace, W.E., “Hydrogenation Characteristics of the HyperstoichiometricZrCrFeTx System (T = Mn, Fe, Co, Ni or Cu)”, J. Solid State Chem., 55, 181 (1984) asquoted by [1986Uch]

[1985Yu] Yu, G.Y., Pourarian, F., Wallace, W.E., “The Crystallographic, Thermodynamic and KineticProperties of the Zr1–xTixCrFe-H2 System”, J. Less-Common Met., 106, 79–87 (1985) (Crys.Structure, Experimental, Kinetics, Thermodyn., 16)

[1986Men1] Meng, X.Y., Northwood, D.O., “A Study of the Structure of Zr(CrFe)2 Laves Phase Precipi-tates in a Zr-Cr-Fe Alloy”, J. Nucl. Mater., 137(3), 217–226 (1986) (Crys. Structure, Experi-mental, 12)

[1986Men2] Meng, X.Y., Northwood, D.O., “Polytype Structures in Zr-Cr-Fe Laves Phase”,J. Less-Common Met., 125, 33–44 (1986) (Crys. Structure, Experimental, 25)

[1986Uch] Uchida, M., Bjurstroem, H., Suda, S., Matsubara, Y., “On the Equilibrium Properties ofSome ZrMn2-Related Hydride-Forming Alloys”, J. Less-Common Met., 119, 63–74 (1986)(Crys. Structure, Experimental, Thermodyn., 16)

[1991Bur] Burany, X.M., Northwood, D.O., “Polytypic Structures in Close-Packed Zr(FeCr)2 LavesPhases”, J. Less-Common Met., 170, 27–35 (1991) (Crys. Structure, Experimental, 23)

Cr–Fe–Zr 11



[1991Zhu] Zhu, Y.T., Devletian, J.H., “Precise Determination of Isomorphous and Eutectoid Transfor-mation Temperatures in Binary and Ternary Zr Alloys”, J. Mater. Sci., 26(22), 6128–6122(1991) (Experimental, Phase Relations, Thermodyn., 13)

[1992Rag] Raghavan, V., “The Cr-Fe-Zr (Chromium-Iron-Zirconium) System” in “Phase Diagrams ofTernary Iron Alloys”, Indian Inst. Metals, Calcutta, 6B, 711–721 (1992) (Crys. Structure,Phase Diagram, Phase Relations, Review, 9)

[1993Sun] Sun, X.-C., Do, N., “Effect of Substructure on the Thermoelectric Power of Deformed Zr-Cr-Fe Sheet”, Mater. Charact., 30(2), 97–105 (1993) (Crys. Structure, Experimental, 25)

[1994Lu] Lu, Z.C., Xianyu, Z., Shen, B.G., Liu, J., “Magnetovolume Effect of Amorphous Fe-(Co,Cr)-Zr Alloys”, Mat. Sci. Eng. A, 182, 1001–1003 (1994) (Experimental, Magn. Prop., 5)

[1994Sun1] Sun, X.-C., “The Thermoelectric Power (TEP) Characteristics of a Zr-1.14mass%Cr-0.10mass% Fe Alloy”, Diss. Abstr. Int. B, 54(9), 4873 (1994) (Abstract)

[1994Sun2] Sun, X-C., Northwood, D.O., “Effect of the Lattice Strain on the Thermoelectric Power in aβ-Quenched Zr-1.14 mass%Cr-0.1 mass%Fe Alloy”, Mater. Charact., 33(1), 3–9 (1994)(Experimental, Phase Relations, Thermodyn., 12)

[1995Mur] Murai, T., Isobe, T., Mae, Y., “Polarization Curves of Precipitates in Zirconium Alloys”,J. Nucl. Mater., 226, 327–329 (1995) (Experimental, 6)

[1995Sou] Soubeyroux, J.L., Bououdina, M., Fruchart, D., Pontonnier, L., “Phase Stability and NeutronDiffraction Studies of Laves Phases Zr(Cr1–xMx)2 with M = Mn, Fe, Co, Ni, Cu and 0 < x <0.2 and their Hydrides”, J. Alloys Compd., 219, 48–54 (1995) (Crystal Structure, Experi-mental, 20)

[1996Mur] Murai, T., Isobe, T., Mae, Y., “Polarization Curves of Precipitates in Zirconium Alloys 2”, J.Nucl. Mater., 230, 178–180 (1996) (Experimental, 2)

[1997Kur] Kuranaka, S., Gamo, T., Morita, Y., “Powder X-ray Diffraction under a High PressureHydrogen Atmosphere for Zr-Cr Based Laves Phase Alloys”, J. Alloys Compd., 253–254,268–271, (1997) (Crys. Structure, Experimental, 13)

[1997Nor] Northwood, D.O., Dosen, K., “A Metallographic Study of Abnormal Grain GrowthDuring Aging of a β-Quenched Zr-1.14 mass% Cr-0.08 Mass% Fe Alloy”, Mater. Charact.,39(2–5), 381–398 (1997) (Crys. Structure, Experimental, Morphology, Phase Relations, 15)

[1999Coa] Coaquira, J.A.H., Rechenberg, H.R., Filho, J.M., “Structural and Mössbauer SpectroscopicStudy of Hexagonal Laves-phase Zr(FexCr1–x)2 Alloys and their Hydrides”, J. AlloysCompd., 288, 42–49 (1999) (Crys. Structure, Experimental, 19)

[2001Coa1] Coaquira, J.A.H., Rechenberg, H.R., “Spin-Glass Behavior of Zr(FexCr1–x)2 Compounds”,J. Magn. Magn. Mater., 226–230, 1306–1308 (2001) (Crys. Structure, Experimental, Magn.Prop., 7)

[2001Coa2] Coaquira, J.A.H., Rechenberg, H.R., “Magnetic Properties of Hexagonal Laves-Phase Zr(Cr1–xFex)2 Compounds”, J. Phys.: Condens. Matter, 13(36), 8415–8434 (2001) (Crys.Structure, Experimental, Magn. Prop., 37)

[2002Rod] Rodriguez, C., Barbiric, D.A., Pepe, M.E., Kovacs, J.A., Alonso, J.A., Hojvat de Tendler, R.,“Metastable Phase Stability in the Ternary Zr-Fe-Cr System”, Intermetallics, 10, 205–216(2002) (Calculation, Phase Relations, Thermodyn., 51)

[2002Ste] Stein, F., Sauthoff, G., Palm, M., “Experimental Determination of Intermetallic Phases,Phase Equilibria, and Invariant Reaction Temperatures in the Fe-Zr System”, J. Phase Equi-lib., 23(6), 480–494 (2002) (Crys. Structure, Experimental, Phase Relations, 88)

[2003Dor] Dorogova, M., Hirata, T., Filipek, S.M., “Hydrogen-Induced Volume Changes in ZrCr2 andPseudo-Binary Compounds of ZrCr2, ZrMn2 and ZrV2”, Phys. Status Solidi A, 198(1),38–42 (2003) (Crystal Structure, Experimental, 13)

[2003Per] Perrot, P., “Cr - Zr (Chromium - Zirconium)”, MSIT Binary Evaluation Program, in MSITWorkplace, Effenberg, G. (Ed.), Materials Science International Services GmbH, Stuttgart;Document ID: 20.15393.1.20 (2003) (Phase Diagram, Assessment, Crys. Structure, 14)

12 Cr–Fe–Zr


[2003Sor] Sorescu, M., Pourarian, F., Brand, “Mössbauer Study of Hydrohenation Effect in Iron RichIntermetallics”, J. Mater. Sci. Lett., 22(22), 1569–1572 (2003) (Crys. Structure, Experimen-tal, 8)

[2004Boz] Bozzano, P.B., Ramos, C., Saporiti, F., Vazquez, P.A., Versaci, R.A., Saragovi, C., “Oxida-tion of the Hexagonal Zr(Cr0.4Fe0.6)2 Laves Phase”, J. Nucl. Mater., 328(2–3), 225–231(2004) (Crys. Structure, Experimental, Interface Phenomena, Kinetics, Phase Relations, 27)

[2005Cao] Cao, W.Q., Yu, S.H., Chun, Y.B., Yoo, Y.C., Lee, C.M., Shin, D.H., Hwang, S.K.,“Strain Path Effects on the Microstructure Evolution and Mechanical Properties of Zr7O2”,Mater. Sci. Eng. A, 395(1–2), 77–86 (2005) (Experimental, Mechan. Prop., Morphology, 32)



Cr–Fe–Zr 13



Cesium – Iron – Oxygen

Pierre Perrot

Introduction

Basic oxides such as Cs2O are known to stabilize the highest oxidation states (IV and VI) of iron. There is agrowing interest to the ferrate (VI) compounds due to their potential as powerful oxidizing agents[2001Ded]. It has also been suggested than dry extraterrestrial environments promote the formation ofhigher oxidation states of iron. Investigations on ferrates (IV and VI) of Cs has been carried out by Möss-bauer spectroscopy [1994Kop, 1995Ran, 1999Kul, 2001Ded], which is one of the most appropriateapproach to characterize iron containing materials. The Cs-Fe-O system could be important to oxide fueland cladding interactions in the LMFBR (Liquid Metal Fast Breeding Reactor) systems [1981Lin]. Thelatest experimental works carried out on the Cs-Fe-O system are summarized in Table 1.

Binary Systems

Cs and Fe present no mutual solubility in the solid and in the liquid state [Mas2]. The Fe-O systemis accepted from thermodynamic assessment of [1991Sun]. This diagram is in a very good agreement withevaluation of [Mas2] mainly based on the fundamental work of [1945Dar, 1946Dar]. The Cs-O system hasnever been thermodynamically assessed. There data on phase equilibria in the Cs-O system are scarce andcontradictory. The latest phase diagram is from [1979Kni], where it was claimed that phase diagram databetween Cs7O and Cs3O were not correct. According to [1979Kni] Cs4O melts incongruently at 53°C thatis higher than 10.5°C given by [Mas2]. According to [Mas2] Cs3O has homogeneity range, while [1979Kni]show this phase as stoichiometric. The Cs2O3 phase [V-C2] was not found in [1979Kni]. Probably thisphase is metastable. The accepted Cs-O phase diagram from [1979Kni] is presented in the evaluation ofthe Cs-Mo-O system in the present volume. Crystallographic data are from [V-C2].

Solid Phases

Crystallographic data of all unary phases and binary and ternary oxides are listed in Table 2. The mixed oxi-des of Cs and Fe lie mainly on the joint Cs2O-Fe2O3. All these phases are easily synthesized by solid-statereaction between Fe2O3 and a cesium salt such as Cs2CO3 or CsNO3. The compound CsFe11O17 whosestructure is the “magnetoplumbite” type is decomposed above 800°C by heating under N2, air or O2

[1987Ito]. It is also easily reduced and gives Fe3O4 + CsFeO2.The mixed oxide Cs2FeO4 in which iron is in the oxidation state VI can be synthesized via the oxidation ofFe(OH)3 in concentrated alkaline media by Cl2 or using the interaction between solid Fe2O3 and CsO2 indry oxygen flow at high temperatures [1995Kul]. It may also be obtained by reacting a mixture (Fe +CsO2) with the ratio Cs/Fe = 4, at 200°C during 10 hours under an oxygen atmosphere [1999Kul]. Cs2FeO4

decomposes above 500 - 600°C with the formation of CsFeO2.5 in which iron is in the state of oxidation IV[1992Kop]. Mössbauer studies [1995Kul] detected iron IV in Cs2FeO4, which is explained by the result ofelectron capture decay of 57Co in the solid matrix of Cs2FeO4. A slight non stoichiometry has been pro-posed for CsFeO2.5 which is sometimes labelled CsxFeO2+0.5x (with x ~1). However, Mössbauer investiga-tion and magnetic measurements carried out by [1994Kop] failed to detect iron with an oxidation statedifferent from IV. Further lattice of CsFeO2.5 is indexed in a P cubic structure perowskite like. A superstruc-ture may be observed, which is due to the ordering of the oxygen vacancies. By taking into account thesuperstructure, CsFeO2.5 may be indexed in a F cubic structure with a parameter twice. The ferrate CsFeO2

in which iron is in the oxidation state III may be obtained by decomposition of the mixed oxalate Cs3Fe(C2O4)3·2H2O in a shorter time and a lower temperature (700°C) than by the conventional ceramic method[1995Ran].

Cs–Fe–O 1



The compound Cs8Fe2O7 in which Fe has a mean oxidation state of 2.33 is obtained by reacting Fe2O3 withmetallic Cs under Ar at 500°C [2004Fri]. It is probable that the phases Cs8Fe2O7 and αCs5FeO4 with thesame structure P21/c and whose compositions are very close belong to the same solid solution P21/c.

Isothermal Section

The Cs-Fe-O diagram in the solid state given in Fig. 1 is taken from the informations given in [1981Lin]related to ferrates in which iron is in a state of oxidation lower than 3. First, they pointed out the fact thatno reference to the compounds Cs4FeO3 and Cs2FeO2 could be found. Nevertheless, they assumed the exis-tence of these compounds and evaluated their properties. In the Fig. 1, the compound Cs4FeO3 has beenidentified with Cs4FeO3.5 (or Cs8Fe2O7) and the compound Cs2FeO2 (or Cs6Fe3O6) have been identifiedwith Cs5Fe3O6, both compounds being described by [2004Fri]. The main characteristics of the diagram isthe existence of the triangle Cs-Fe-Cs5Fe3O6, which means that the mixture Cs+Fe oxidize more easily thanpure Cs, that is under oxygen pressures lower than the oxygen pressure at Cs-Cs2O equilibrium. This dia-gram may be used in the temperature range between 600 and 1000°C; below 570°C, FeO is never stable andthe tie lines between FeO and CsFeO2 have to be deleted. The ferrate (VI) Cs2FeO4 stable in highly basicmedium [2004Lic] loses its oxygen by heating and thus, is not shown in Fig. 1.

Thermodynamics

The thermodynamic properties of some cesium ferrates given in Table 3 are mainly taken from the evalua-tion of [1981Lin]. However, as seen above, the compounds Cs4FeO3 and Cs2FeO2 whose existence wasassumed by these authors were identified respectively with the compounds Cs8Fe2O7 and Cs5Fe3O6

described by [2004Fri]. The entropies were estimated by the same authors with the hypothesis that theentropy of formation of ferrates from pure oxides equals zero, which is fairly unrealistic. The entropiesgiven in Table 2 are thus corrected by using a more probable value of 3 J per mole of metal observed withanalogous chromium compounds. [1981Lin] propose a tentative Ellingham (pO2-T) diagram comparing thestability domains of iron oxides, Cs5Fe3O6 and CsFeO2. However, this diagram is very unlikely, because theauthors assumed that FeO does not exists below 200°C, while this temperature is commonly accepted to be570°C.


The ferrite CsFe11O17 is an excellent ionic conductor. However its conductivity, electronic and ionic, mea-sured between 300 and 600°C by [1987Ito] is lower than that of analogous compounds RbFe11O17 andKFe11O17 because of the higher ionic radius of the Cs+ ion. The ferrate VI Cs2FeO4 present a high solidstate stability, is hardly soluble in concentrated KOH and may be used as an alkali cathode able to sustaincurrent densities similar to that of conventional MnO2 cathodes [2004Lic].

Miscellaneous

Mössbauer spectra are given for Cs2FeO4 [1999Kul], CsFeO2 [1995Ran] and CsFeO2.5 [1994Kop].CsFeO2.5 presents a simple symmetrical line with an isomer shift δ = 0.15 mm·s–1 corresponding to FeIV

in an octahedra oxygen environment. The ferrate (VI) presents also a simple line with an isomer shiftδ = –0.59 mm·s–1 corresponding to FeVI in an octahedral environment and an antiferromagnetic transitionat 2.8 K.

2 Cs–Fe–O


Table 1. Recent Investigations of the Cs-Fe-O System


[1986Mor] X-Ray diffraction analysis 20°C, CsFe11O17

[1987Ito] Ionic conductibility measurements 300 - 600°C, CsFe11O17

[1992Kop] Thermal analysis, X-Ray determination 20 - 600°C, Cs2FeO4-CsFeO2.5

[1994Kop] Thermal analysis, Mössbauer 20 - 600°C, Cs2FeO4-CsFeO2.5

[1995Kul] Synthesis, Mössbauer –196°C, +20°C, Cs2FeO4

[1995Ran] Synthesis, Mössbauer 20 - 700°C, CsFeO2

[1999Kul] Synthesis, Mössbauer 200°C, Cs2FeO4

[2001Ded] Mössbauer, Magnetic measurements 2 - 6 K, 20°C, Cs2FeO4

[2004Fri] Synthesis, X-Ray diffraction analysis 500°C, Cs8Fe2O7, Cs5Fe3O6, CsFeO2

[2004Lic] Electrochemistry, IR measurement 20°C, Cs2FeO4

[2005Fri] X-Ray diffraction analysis 20°C, Cs5FeO4, Cs7Fe2O8





Comments/References

(Cs)< 28.39

cI2Im�3mW

a = 614.1 at 25°C [Mas2]

(αFe)1538 - 1394, 912

cI2Im�3mW

a = 286.65 at 25°C [Mas2]

(γFe)1394 - 912

cF4Fm�3mCu

a = 364.67 at 25°C [Mas2]

Cs7O< 3

hP24P�6m2Cs7O

a = 1639.3c = 919.3


Cs4O< 53

oP*Pna21

a = 1682.3b = 2052.5c = 1237.2

[1979Kni, V-C2]

(continued)

Cs–Fe–O 3






Comments/References

Cs7O2

< –10.5mP56P21/cCs11O3

a = 1761.0b = 921.8c = 2404.7β = 100.14°

[1979Kni, V-C2] labelled Cs11O3 in [Mas2]

Cs3O< 164

- - 23 to 25 at.% O [Mas2]

Cs2O< 495

hR9R�3mWN2

a = 425.6c = 1899.2

[1979Kni, 2005Gem]

CsO< 590

oI8ImmmCsO

a = 432.2b = 751.7c = 643.0

[1979Kni, V-C2]

Cs2O3

< 502cI28I�43dTh3P4

a = 988 [Mas2, V-C2]. Not given in [1979Kni]. Probablymetastable

αCsO2 (r)< 200

tI6I4/mmmCaC2

a = 446.2c = 732.6

[Mas2, V-C2]

βCsO2 (h)450 - 200

cF8Fm�3mNaCl

a = 662 [Mas2, V-C2]

CsO3

< 70mP16P21/cRbO3

a = 670.9b = 624.4c = 899.7

[1979Kni, V-C2]

Fe1–xO (wüstite)1422 - 569

cF8Fm�3mNaCl

a = 431.0a = 429.3

0.05 < x < 0.12 [1991Sun]x = 0.05x = 0.12

Fe3O4 (r)< 580

oP56PbcmFe3O4 (r)

a = 1186.8b = 1185.1c = 1675.2

[V-C2]

Fe3O4 (h)(mgnetite)1597 - 580

cF56Fd�3mMgAl2O4

a = 839.6a = 854.5

at 25°Cat 1000°C [V-C2]

αFe2O3 (hematite)< 1451

hR30R�3cAl2O3

a = 503.42c = 1374.83

at 600°C [Mas2, V-C2]


a = 939.3 metastable phase [V-C2]

(continued)

4 Cs–Fe–O





Comments/References

γFe2O3

(maghemite)tP60P41212Mn5Si2 (?)

a = 833.96c = 832.21

metastable phase [V-C2]

* CsFe11O17 hP64P63/mmcNaAl11O17

(β alumina)

a = 592.3c = 2415.8

[1986Mor]

* CsFeO2.5 cP36Pm�3mCaTiO3

(perowskite)

a = 419.9 [1992Kop, 1994Kop]

* CsFeO2 cF32Fd�3mβ cristobalite

a = 839.2 [2004Fri]

* Cs5Fe3O6 oP56P212121Cs5Fe3O6

a = 867.1b = 872.9c = 1670.1

at – 22°C [2004Fri]

* Cs2FeO4 oP28PnmaβK2SO4

a = 843.43b = 629.23c = 1112.74

[1992Kop]Complete conversion of K2FeO4 by CsOH[2004Lic]

* Cs7Fe2O8 mP68P21/cRb7Fe2O8

a = 666.0b = 1097.4c = 2156.6β = 92.83°

at 22°C [2005Fri]

* Cs8Fe2O7 mP68P21/cCs8Fe2O7

a = 722.32b = 1789.0c = 733.9β = 118.98°

at –28°C [2004Fri]

* αCs5FeO4 (r) mP40P21/cαCs5FeO4

a = 1133.9b = 1269.5c = 725.05β = 99.07°

at 22°C [2005Fri]

* βCs5FeO4 (l) mP40P21/cβCs5FeO4

a = 880.78b = 1067.4c = 1115.7β = 97.35°

at – 25°C [2005Fri]

Cs–Fe–O 5




Phase Temperature Range [°C] Property, per mole of atoms[J, mol, K]

Comments

1/17 (Cs8Fe2O7) 25 S° = 47.2 ± 1ΔfH° = – 138 500 ± 500

see text[1981Lin]

1/14 (Cs5Fe3O6) 25 S° = 44.4 ± 1ΔfH° = – 148 500 ± 500

see text[1981Lin]

1/4 (CsFeO2) 25 S° = 32.3 ± 1ΔfH° = – 175 000 ± 500

see text[1981Lin]

Fig. 1. Cs-Fe-O. The equilibria in the solid state

6 Cs–Fe–O


References[1945Dar] Darken, L.S., Gurry, G.W., “The System Iron-Oxygen - I - The Wuestite Field and Related

Equilibria”, J. Am. Chem. Soc., 67, 1398–1412 (1945) (Experimental, Phase Diagram, Ther-modyn., *, #, 26)

[1946Dar] Darken, L.S., Gurry, G.W., “The System Iron-Oxygen - II - Equilibrium and Thermody-namics of Liquid Oxides and other Phases”, J. Am. Chem. Soc., 68, 798–816 (1946) (Experi-mental, Phase Diagram, Phase Relations, Thermodyn., *, #, 24)

[1979Kni] Knight, C.F., Phillips, B.A., “The Cs-O system; Phase Diagram and Oxygen Potentials”,J. Nucl. Mater., 84, 196–206 (1979) (Phase Diagram, Phase Relations, Review, 25)

[1981Lin] Lindemer, T.B., Besman, T.M., Johnson, C.E., “Thermodynamic Review and Calculations -Alkali Metal Oxide Systems with Nuclear Fuels, Fission Products, and Structural Materials”,J. Nucl. Mater., 100, 178–226 (1981) (Phase Diagram, Thermodyn., Review, 280)

[1986Mor] Morgan, P.E.D., Miles, J.A., “Magnetoplumbite-Type Compounds: Further Discussion”,J. Am. Ceram. Soc., 69(7), 157–159 (1986) (Crys. Structure, Experimental, Review, 27)

[1987Ito] Ito, S., Kubo, N., Nariki, S., Yoneda, N., “Ion Exchange in Alkali Layers of Potassium β-Ferrite((1+x)K2O·11Fe2O3) Single Crystals”, J. Am. Ceram. Soc., 70(2), 874–879 (1987) (Crys. Struc-ture, Electr. Prop., Experimental, 29)

[1991Sun] Sundman, B., “An Assessment of the Fe-O System”, J. Phase Equilib., 12(1), 127–140(1991) (Phase Diagram, Thermodyn., Assessment, 53)

[1992Kop] Kopelev, N.S., Val’kovskii, M.D., Popov, A.I., “The Thermal Decomposition of Caesium Fer-rate(VI)”, Russ. J. Inorg. Chem., 37(3), 267–268 (1992), translated from Zh. Neorg. Khim.,37(3), 540–543, (1992) (Crys. Structure, Phase Relation, Experimental, 6)

[1994Kop] Kopelev,N.S., Popov,A.I.,Val’kovskii,M.D., “Properties of theProducts ofCsxFeO2+0.5xThermalDecomposition”, J. Radioanal. Nucl. Chem., 188(2), 99–108 (1994) (Crys. Structure, Magn.Prop., Phase Relations, 24)

[1995Kul] Kulikov, L.A., Perfil’ev, Y.D., Kopelev, N.S., “The Iron Charge State in Solid Cesium Ferrate(VI) Deduced from Mössbauer Absorption and Emission Spectroscopy”, J. Phys. Chem.Solids, 56(8), 1089–1094 (1995) (Experimental, Crys. Structure, Electronic Structure, 44)

[1995Ran] Randhawa, B.S., “Mössbauer Study on Thermal Decomposition of Cesium Tris (Oxalato)Ferrate (III) Dihydrate”, J. Radioanal. Nucl. Chem., 201(1), 57–63 (1995) (Experimental,Electronic Structure, Phase Relations, 20)

[1999Kul] Kulikov, L.A., Yurchenko, A.Y., Perfil’ev, Y.D., “Preparation of Cesium Ferrate (VI) fromMetallic Iron” (in Russian), Vestn. Mosk. Univ., Ser. 2: Khim., 40(2), 137–138 (1999) (Elec-tronic Structure, Experimental, Phase Relations, 12)

[2001Ded] Dedushenko, S.K., Perfiliev, Yu.D., Goldfeld, M.G., Tsapin, A.I., “Mössbauer Study of Hex-avalent Iron Compounds”, Hyperfine Interact., 136/137, 373–377 (2001) (Magn. Prop., Elec-tronic Structure, Experimental, 19)

[2004Fri] Frisch, G., Roehr, C., “A5{Fe3O6} (A=Rb,Cs), Cs{FeO2} and Cs8{Fe2O7}: New Oxoferratesof the Heavy Alkaline Metals” (in German), Z. Naturforsch. B, 59, 771–781 (2004) (Crys.Structure, Experimental, 40) (Summary with 5 ref. published in Z. Kristallogr., 21, 156 (2004)

[2004Lic] Licht, S., Naschitz, V., Rozen, D., Halperin, N., “Cathodic Charge Transfer and Analysis ofCs2FeO4, K2FeO4 and Mixed Alkali Fe (VI) Ferrate Super-Irons”, J. Electrochem. Soc.,151(8), A1147-A1151 (2004) (Crys. Structure, Electr. Prop., Experimental, 15)

[2005Fri] Frisch, G., Roehr, C., “New Orthoferrates of Rubidium and Cesium: α-, βCs5[FeIIIO4] and

AI7[(Fe

IVO4)(FeVO4)] (AI =Rb, Cs)” (in German), Z. Anorg. Allg. Chem., 631(2-3),

507–517 (2005) (Crys. Structure, Experimental, 30)[2005Gem] Gemmings, S., Seifert, G., Muehle, C., Jansen M., Abu-Yaron, A., Arad, T., Tenne, R., “Elec-

tron Microscopy, Spectroscopy and First Principles Calculations of Cs2O”, J. Solid-StateChem., 178(4), 1190–1196 (2005) (Crys. Structure, Experimental, Calculation, 21)



Cs–Fe–O 7



Copper – Iron – Hydrogen

Suray Bhan, Nataliya Bochvar, Boris Kasper, Ortrud Kubaschewski, Pierre Perrot, Peter Rogl

Introduction

Experimental information on the Cu-Fe-H system, gathered in Table 1 and reviewed in [1974Sig] is limitedto the solubility of hydrogen in Cu-Fe alloys under 0.1MPa H2. The majority of experiments were per-formed using the conventional Sievert's technique in the temperature range from 1083 to 1672°C. However[1974Boo] used the constant volume method. Considering the experimental difficulties, i.e. reaction withcrucible and contamination of the starting materials, there is acceptable agreement among the authors of[1963Wei] (1592°C for 0-12 mass% Cu), [1965Bur] (1600°C for 0-100 mass% Cu), [1967Ban] (1548,1610 and 1672°C for 0-15 mass% Cu), [1969Kat, 1970Kat] (1150°C for 0-2.5 mass% Fe). Heat and entropyof solution data were presented by [1972Deg]. Solubility of hydrogen was also reported for a selection oftechnically used multicomponent copper base alloys [1976Ger].

Binary Systems

The Cu-Fe system, accepted from [1994Swa] has been critically evaluated by [1995Che]. It presents a totalsolubility in the liquid state and a limited solubility in the solid state. No stable intermetallic compoundis known.The Fe-H system, given by [Mas2] at 0.1 MPa has been assessed by [1990San] which presents also a dia-gram at 40 MPa. The solubility of H in pure liquid iron under 0.1 MPa is fitted by the function validbetween 1530 and 1800°C:log10 (H/Fe) = ½ log10 (pH2/bar) –1.893 – 1825 / T, where H/Fe represents an atomic ratio.At 1600°C, (H/Fe) = 1.357·10–3, which has to be compared with direct measurement of [1963Wei,1974Boo]: 27.7±1.3 cm3 H2 measured at 0°C and 101325 Pa, dissolved in 100 g Fe, leading to H/Fe = (1.38±0.06)·10–3.The Cu-H system is given by [Mas2] at 50 MPa. The solubility of H in liquid copper, given by [1970Kat]obeys the Sievert's law up to 1 bar:log10 (H/Cu) = ½ log10 (pH2/bar) – 1.809 – 2370/TThe relation is valid between 1100°C and 1300°C. However, an extrapolation at 1500°C gives H/Cu = 7.71·10–4 under 1 bar, whereas the direct measurement of [1972Deg] gives H/Cu = 7.15·10–4 inthe same conditions.CuH may be synthesized by reduction of aqueous CuSO4 with hypophosphorous acid H3PO2 [2002Now]. Itis metastable, but decomposes slowly at room temperature so that its thermodynamic properties have beeninvestigated [2003Bur]. The hydrogen pressure at equilibrium Cu/CuH has been evaluated at 106 TPa atroom temperature [2000Bur].

Solid Phases

The solid phases are presented in Table 2. No ternary hydrides are known.

Isothermal Sections

A number of investigations on the solubility of H in liquid Fe [1963Wei, 1965Bur] report first order inter-action coefficient to be positive, indicating that Cu increases the activity coefficient of H in the melt and thusdecreases its solubility at a given H2 pressure. On the contrary, the solubility of hydrogen in liquid Cu isincreased by the presence of iron [1967Ban, 1969Kat, 1970Kat, 1976Ger, 1983Str], a result which disagreewith the first measurements of [1963Wei, 1965Bur].Whereas the afore-mentioned authors arrived at essentially linear dependencies of the dissolved volume ofH versus composition, the results of [1970Deg, 1972Deg] revealed striking deviations from linearity such as

Cu–Fe-H 1



a minimum at 2.6 mass% Fe in the temperature range from 1083 to 1325°C, which, in the late paper, wasattributed to oxygen contamination, and a pronounced maximum at 60-70 mass% Fe at 1500°C. Employingthe constant volume technique, [1974Boo] found practically no concentration dependence for alloys up to13 mass% Cu in iron at 1600°C.Figure 1 displays the isothermal sections at 1150°C [1969Kat], 1420°C [1983Str] and 1600°C [1965Bur,1983Str]. According to [1974Boo], copper, which does not form compounds with iron in the solid stateand shows lower solubility of hydrogen than iron, has very little effect on the solubility of hydrogen in iron.

Thermodynamics

The Wagner first order interaction parameter eH(Cu) has first been experimentally determined at 1592°C by

[1963Wei] and at 1600°C by [1965Bur]. It is defined by eH(Cu) = {∂ log10 f H / ∂ (% Cu)} where (% Cu)

represents the copper content of the alloy expressed in mass% and fH the activity coefficient of H in the alloydefined by fH = (% H in pure iron) / (% H in the alloy). It is calculated from the solubility measurements atconstant temperature and hydrogen pressure and its most probable value is eH

(Cu) = –0.0004 at 1600°C[1974Boo] for less than 12 mass% Cu in the alloy.In the Cu rich alloys, eH

(Fe) = {∂ log10 f H / ∂ (% Fe)} where (% Fe) represents the iron content of the alloyexpressed in mass% and fH the activity coefficient of H in the alloy defined by fH = (% H in pure copper) /(% H in the alloy). It was measured by [1972Deg] and its most probable value is eH

(Fe) = –0.0184 at 1250°Cand –0.0131 at 1500°C, indicating that Fe increases the H solubility in liquid Cu. This result agrees with thevalue eH

(Fe) = –0.015 proposed by [1969Kat, 1970Kat] at 1300°C. [1970Deg] is the only author to proposea positive value, so that the minimum of H solubility proposed at ~3 mass% Fe in liquid Cu is doubtful.


The microstructure of fine grains of Fe-Cu alloys prepared by mechanical alloying change critically[2001Rad] after 2 min. of annealing at 1130°C under H2 atmosphere. A very fast coarsening of the micro-structure is observed. High hardness and good conductivity is observed after hardening.

Table 1. Investigations of the Cu-Fe-H Phase Relations, Structures and Thermodynamics


[1963Wei] Hydrogen solubility by the Sievert’smethod, Fe rich alloys

1592°C, < 12 mass% Cu, < 0.1 MPa H2

[1965Bur] Hydrogen solubility by the Sievert’smethod, Cu-Fe alloys

1600°C, 0-100 mass% Cu, 0.1 MPa H2

[1967Ban] Hydrogen solubility by the Sievert’smethod, Fe rich alloys

1548-1672°C, < 10 mass% Cu, 0.1 MPa H2

[1969Kat,1970Kat]

Hydrogen solubility, Sievert’s andsampling methods

1100-1300°C, < 2.5 mass% Fe, 0.1 MPa H2

[1970Deg] Hydrogen solubility by the Sievert’smethod, Cu rich alloys

1180-1325°C, < 5 mass% Fe, 0.1 MPa H2

[1972Deg] Hydrogen solubility by the Sievert’smethod, Cu-Fe alloys

1250°C, < 12 mass% Fe, 0.1 MPa H2 1500°C,< 80 mass% Fe, 0.1 MPa H2

[1974Boo] Hydrogen solubility, constant volumemethod, Fe rich alloys

1600°C, < 13 mass% Cu, 0.1 MPa H2

(continued)

2 Cu–Fe-H



[1974Lin] Hydrogen solubility by the Sievert’smethod, Cu rich alloys

1150-1500°C, < 80 at.% Fe, 0.1 MPa H2

[1976Ger] Hydrogen solubility by the Sievert’smethod, Cu rich alloys

1150-1300°C, < 10 at.% Fe, 0.1 MPa H2

[1983Str] Hydrogen solubility by the Sievert’smethod, Cu rich alloys

1150-1300°C, < 6 mass% Fe, 0.1 MPa H2

[2001Rad] Micrography, density, electron probemicroanalysis

< 1130°C, Fe-30 mass% Cu composites sinteredunder H2





Comments/References

(δFe)1538 - 1394

(αFe)< 912

cI2Im�3mW

a = 293.15

a = 286.65

dissolves 15 at.% Cu at 1440°C [2007Tur] and0.1 mass% H at 1538°C [2002Ant]

pure Fe at 20°C [Mas2, V-C2] (A2 structure). Dissolves1.6 at.% Cu at 847°C

(γFe)1394 - 912

cF4Fm�3mCu

a = 293.16 at 915°C [Mas2, V-C2].Dissolves 15 at.% Cu at 1440°C

(Cu)< 1084.62

cF4Fm�3mCu

a = 361.46 at 25°C [Mas2, V-C2].Dissolves 5 at.% Fe at 1095°C

CuH hP4P63mcZnS(Würtzite)

a = 289c = 462a = 293c = 468

CuH [2000Bur]actually CuH0.8

CuD [2005Tka]

Cu–Fe-H 3



Fig. 1. Cu-Fe-H. Isothermal sections at 1150, 1420 and 1600°C under 0.1 MPa of hydrogen pressure

4 Cu–Fe-H


References[1963Wei] Weinstein, M., Elliott, J.F., “Solubility of Hydrogen in Liquid Iron Alloys”, Trans. Met. Soc.

AIME, 227, 382–393 (1963) (Calculation, Experimental, Thermodyn., 27)[1965Bur] Burylev, B.P., “Solubility of H in Liquid Fe Alloys” (in Russian), Izvest. Vyssh. Ucheb.

Zaved., Chern. Met., (2), 17–22 (1965) (Experimental, 13)[1967Ban] Ban-ya, S., Fuwa, T., Ono, K., “Solubility of Hydrogen in Liquid Iron Alloys” (in Japanese),

Tetsu to Hagane, 53(2), 101–116 (1967) (Experimental, Phase Relations, Thermodyn., 18)[1969Kat] Kato, E., Orimo, T., “Solubility of Hydrogen in Liquid Copper and Certain Copper Alloys

Measured by Sieverts Method” (in Japanese), J. Jpn. Inst. Met., 33, 1165–1170 (1969)(Experimental, Phase Relations, Thermodyn., 20)

[1970Deg] Degtyarev, Yu.V., Linchevsky, B.V., Chursin, V.M., “Solubility and Activity of Hydrogen inLiquid Copper and its Alloys with Manganese, Iron and Nickel”, Russ. Metall., (4), 30–33(1970), translated from Izv. Akad. Nauk SSSR, Met., (4), 42–46 (1970) (Experimental, Inter-face Phenomena, Thermodyn., 6)

[1970Kat] Kato, E., Ueno, H., Orimo, T., “Solubility of Hydrogen in Liquid Copper Alloys”, Trans. Jpn.Inst. Met., 11, 351–358 (1970) (Experimental, Phase Relations, 19)

[1972Deg] Degtyarev, Yu.V., Linchevsky, B.V., Chursin, V.M., Yudin, A.F., “The Solubility and Activityof Hydrogen in Copper-Base Alloys”, Russ. Metall., (5), 31–35 (1972), translated from Izv.Akad. Nauk SSSR, Met., (5), 38-42, (1972) (Experimental, Phase Relations, 6)

[1974Boo] Boorstein, W.M., Pehlke, R.D., “Measurement of Hydrogen Solubility in Liquid Iron AlloysEmploying a Constant Volume Technique”, Metall. Trans., 5, 399–405 (1974) (Calculation,Experimental, Phase Relations, 29)

[1974Lin] Linchevsky, B.V., Degtyarev, Yu.V., Chursin, V.M., “Solubility of H in Binary Cu-BaseAlloys” (in Russian) in “Kinet. Termodin. Vzaimodeistviya Gazov Zhidk. Met. Mater. Simp.1972”, Nauka, Moscow, 15–20 (1974) (Experimental, Phase Relations, Thermodyn., 5)

[1974Sig] Sigworth, G.K., Elliott, J.F., “The Thermodynamics of Dilute Liquid Copper Alloys”, Can.Metall. Quart., 13(3), 455–461 (1974) (Review, Thermodyn., 122)

[1976Ger] Gershkovich, V.K., Strel’tsov, F.N., Kunin, L.L., “Solubility of Hydrogen in Copper BasedAlloys”, Russ. Metall., (2), 55–57 (1976), translated from Izv. Akad. Nauk SSSR, Met., (2),72–74 (1976) (Experimental, Phase Relations, 7)

[1983Str] Strel’tsov, F.N., “Solubility of Hydrogen in Copper Melts”, Russ. Metall., (6), 46–49 (1983),translated from Izv. Akad. Nauk SSSR, Met., (6), 55–58 (1983) (Experimental, Phase Rela-tions, 10)

[1990San] San-Martin, A., Manchester, F.D., “The Fe-H (Iron-Hydrogen) System”, Bull. Alloys PhaseDiagrams, 11(2), 173–184 (1990) (Phase Diagram, Review, 86)

[1994Swa] Swartzendruber, L.J., “Cu-Fe (Copper-Iron)” in “Phase Diagrams of Binary Copper Alloys”,Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM International, MaterialsPark, OH, 167–172 (1994) (Phase Diagram, Review, 102)

[1995Che] Chen, Q., Jin, Z., “The Fe-Cu System: A Thermodynamic Evaluation”, Metall. Trans. A, 26,417–426 (1995) (Thermodyn., 55)

[2000Bur] Burtovyy, R., Utzig, E., Tkacz, M., “Study of the Thermal Decomposition of the CopperHydride”, Thermochim. Acta, 363, 157–163 (2000) (Crys. Structure, Thermodyn., Experi-mental, 16)

[2001Rad] Radchenko, O.G., Getman, O.I., “Microstructure Evolution of Fine-Grain Fe-Cu CompositesDuring Heat Treatment in Hydrogen”, Int. J. Hydrogen Energy, 26(5), 489–491 (2001)(Experimental, Morphology, 3)

[2002Ant] Antonov, V.E., Baier, M., Dorner, B., Fedotov, V.K., Grosse, G., Kolesnikov, A.I., Ponya-tovsky, E.G., Schneider, G., Wagner, F.E., “High-Pressure Hydrides of Iron and its Alloys”,J. Phys.: Condens. Matter, 14, 6427–6445 (2002) (Crys. Structure, Experimental, PhaseRelations, Review, 46)

[2002Now] Nowak, B., Burtovyy, R., Tkacz, M., “2D and 63Cu NMR Study in Copper Hydride”, J. AlloysCompd., 384, 71–75 (2002) (Crys. Structure, Phys. Prop., Experimental, 16)

Cu–Fe-H 5



[2003Bur] Burtovyy, R., Wlosewicz, D., Czopnik, A., Tkacz, M., “Heat Capacity of Copper Hydride”,Thermochim. Acta, 400, 121–129 (2003) (Thermodyn., Experimental, 21)

[2005Tka] Tkacz, M., Burtovyy, R., “Isotope Effect in a Cu-H(D) System with Hexagonal HydridePhase”, J. Alloys Compd., 404/406, 368–371 (2005) (Crys. Structure, Thermodyn., Experi-mental, 14)

[2007Tur] Turchanin, M., Agraval, P., “Cu-Fe (Copper-Iron)”, MSIT Binary Evaluation Program, inMSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH,Stuttgart; Document ID: 20.11107.1 (2007) (Phase Diagram, Crys. Structure, Thermodyn.,Assessment, 31)



6 Cu–Fe-H


Copper – Iron – Manganese

Nathalie Lebrun, Pierre Perrot

Introduction

The Cu-Fe-Mn has raised much attention because of its technological importance due to the fact that Cu-Febased alloys are of interest in practical development of high strength, high electrical conductive Cu richalloys and age-hardened Fe rich alloys. The general shape of the liquidus lines was known since[1909Jae, 1913Par, 1925Ost] and the existence of a solid solution with a miscibility gap for alloys with lessthan 50 mass% Mn was recognized by [1913Par, 1958Top]. The first schematic diagram showing few liqui-dus lines was proposed in [1949Jae] and reproduced in [1969Gue, 1979Dri, 1983Riv]. Experimental worksdone on the ternary system since the first important, although early work of [1913Par] completed by that of[1977Has] are gathered in Table 1 and a comprehensive review may be found in [1988Ray] updated by[1994Rag, 2002Rag]. Calphad assessments were carried out by [1981Nis, 1989Kum, 2003Mie]. A calcula-tion of the miscibility gaps, stable or metastable in the α, γ and liquid solutions is also presented in[2004Wan2].

Binary Systems

Cu-Fe and Cu-Mn are accepted from [2007Tur] and [2006Tur], respectively. In the Cu-Fe system, a meta-stable miscibility gap in the liquid phase exists. Using new experimental thermodynamic data, the Fe-Mnsystem has been recently reassessed and the Calphad description was updated by [2004Wit]. This new ree-valuation of the phase equilibria leads to consistently better fits to the available experimental data. There ishowever, a typographical error in [2004Wit] in that the Mn rich invariant reaction involving the liquid phaseis given as a peritectic type reaction in the table of invariants. This reaction should be denoted as eutectic, asconfirmed by [2007Wit], and have been written as L ↾ (δMn) + (γMn,γFe). Consequently, the Fe-Mn sys-tem is accepted from [2004Wit].

Solid Phases

Solid phases are presented in Table 2. At high temperature, γ(Mn,Fe) alloys present a continuous solid solu-tion which is stabilized by the presence of Cu. At low temperature, the γ alloy (23 mass% Mn, 4 mass% Cu)undergoes a tetrahedral distortion (around 150°C) then a rhombic distortion (around 100°C) [1981Dem1],which may be considered as second order transitions. A γ alloy (10 mass% Cu, 1 mass% Fe) undergoes amartensitic transformation whose Ms lies around 160°C [1981Dem2]. Cooling below Ms brings an increasein tetragonality (1 – c/a) from 0.02 at 160°C to 0.04 at 120°C. A tetragonality of 0.03 at room temperaturehas also been measured by [2005Zha] on a γ alloy (5 at.% Cu, 9.5 at.% Fe).


The liquidus proposed by [1913Par] is well defined and fits well with the accepted binaries. More recently,[2003Mie] computed the liquidus surface over the entire ternary system. A general good agreement isobserved between its calculation and the experimental data of [1913Par] assessed by [1988Ray]. Neverthe-less, calculation deviates from experimental data in the Fe-Mn region up to 30 mass% Cu, especially in theMn rich region. Consequently, the experimental liquidus curves of [1913Par] reproduced by [1988Ray]have been preferred and are shown in Fig. 1. Good agreement has been also observed between the experi-mental data of [1977Has] and the calculated isotherm at 1050°C done by [2003Mie]. This isothermal line,accepted from [2003Mie] has been added to the drawing in Fig. 1.The saturation lines in the Fe and Mn rich corners have been deduced from the binary reactions and areschematic lines (indicated as dashed lines in the drawing). In their review, [1969Gue] indicated the exis-tence of a critical point at 10Fe-55-Cu-35Mn (mass%) and 880°C. It was later shown in the review of

Cu–Fe–Mn 1



[1988Ray]. Its composition has been moved to lower Fe content in [1969Gue, 1979Cha, 1979Dri]. The highFe composition of the critical point seems to be doubtful since at around 900°C the boundary of the liquidphase is very close to the Cu-Mn binary edge. Moreover, there is a big disagreement with the experimentalisotherm of the liquidus surface. The composition calculated by [2003Mie] seems to be more realistic and isin agreement with the reviews of [1969Gue, 1979Cha, 1979Dri]. Consequently the final composition of thecritical point is estimated to be 0.8 mass% Fe, 35.3 mass% Mn [2003Mie] at 880°C [1969Gue], see Fig. 1.Further investigations are needed to determine experimentally the exact composition of the critical point.A well defined experimental solidus map has been proposed by [1913Par] and is reproduced in Fig. 2 withsmall modification along the binary edge in agreement with the accepted binary systems. Moreover, the sol-vus lines in the Fe rich corner have been modified in order to be in agreement with the miscibility gap.[1925Ost] stipulated that this miscibility could be favoured by the presence of carbon. The single phasefields have been indicated as dashed lines in Fig. 2.A metastable miscibility gap is observed in the liquid phase along the Cu-Fe binary edge. The addition ofMn decreases the critical temperature as shown in Fig. 3 reproduced from the Calphad assessment of[2004Wan2]. The miscibility gap at 927°C was not included in the drawing since the accepted Cu-Fe binarysystem presents only a metastable miscibility gap above 1107°C. A metastable solid miscibility gap islandof the γ phase was predicted by Calphad method [2004Wan2] as depicted in Fig. 4.

Isothermal Sections

Several isothermal sections are available in the literature including calculations and experimental investiga-tions from 1300 to 550°C. The agreement between experimental data and calculations is quite reasonablethough at high Mn contents at 1300 and 1100°C, the calculated solubility of the iron rich (γFe) phase inthe Cu rich liquid is higher than indicated by the experimental data. These discrepancies are mainly dueto some disagreements between the experimental data of [1977Has, 1997Oht]. [1966Sal] investigated a par-tial isothermal section at 1250°C and concluded that addition of about 10 mass% Mn to pure Fe leads tosmall effect on the solubility of copper.Only schematic isothermal sections at 1150°C have been reported in the literature: these are the calculationsdone by [1981Nis] and reproduced by [1994Rag] with a slight modification near the Mn corner to accountfor the presence of (δMn). According to the accepted binary system Fe-Mn, the (βMn) should also be pre-sent and has been added in the isothermal section at 1150°C shown in Fig. 7.The isothermal section at 900°C is based on [1988Ray] with modification of the liquid boundary and thethree-phase field (Cu)+L+(γFe) in agreement with the accepted liquidus surface.A solid miscibility gap is present between the (γFe) and (Cu) phases at 850°C. Calculations [1977Has,1981Nis, 2003Mie] reproduced well the experimental data except along the Fe-Mn binary edge wherethe calculated data present higher Fe composition compared with the experimental data of [1977Has].The calculations have been accepted here since the experimental data did not fit well with the acceptedCu-Fe binary system in the Fe rich corner. Moreover, the three-phase field (Cu)+(αFe)+(γFe) has not beenretained since no experimental data were reported on it. The calculation done by [1989Kum] has been takeninto account to describe the boundaries between (γFe) and (αFe) in the Fe rich corner of the ternary system.The accepted isothermal sections are reported in Fig. 5 (1300°C, accepted from [1994Rag]), Fig. 6 (1200°C,accepted from [2004Wan1], Fig. 7 (1150°C) accepted from [1988Ray] Fig. 8 (1050°C accepted from[1977Has, 2003Mie]), Fig. 9 (900°C) and Fig. 10 (850°C) both accepted from [1988Ray].[1989Kum] computed the phase boundaries in the Fe rich corner and found the existence of the three-phasefield (Cu)+(αFe)+(γFe) which is shifted towards the higer Fe composition with decreasing temperaturefrom 750 to 550°C, as shown in Fig. 11 (750°C) and Fig. 12 (550°C). This three-phase field was not takeninto account in the calculated isothermal section at 600°C determined by [1977Has, 1981Nis]. Only thesolid miscibility gap between (Cu) and (γFe) has been reported with good agreement between both calcula-tions. With the addition of Mn, the gap increases inside the ternary. The phase separation extends from theCu-Fe to the Cu-Mn and Fe-Mn with decreasing temperature. Minor modifications have been done to agreewith the accepted binary data and the isothermal lines L / L+(γFe) determined experimentally by [1913Par,1977Has].

2 Cu–Fe–Mn


Thermodynamics

The excess Gibbs energy in a phase φ, in which φ represents one of the α, γ and liquid solutions was eval-uated by [1997Oht, 2003Mie] using the following model:Gxs(φ) = Gxs(Cu,Fe,φ) + Gxs(Cu,Mn,φ) + Gxs(Fe,Mn,φ) + Gxs(Cu,Fe,Mn,φ)The binary terms for a φ phase are given by a Redlich-Kister expansion:Gxs(i, j, φ) = xixj ΣkL

(k,φ)(xi – xj)k

The ternary term is given by Gxs(Cu,Fe,Mn,φ) = xCu xjFe xMn A(φ). The ternary interaction parameters A(φ)proposed by [2003Mie] are:A(α) = 30000A(γ) = –68000 + 50 TA(L) = (115000 – 60 T) xCu + 13000 xFe + 10000 xMn.


Experimental works on materials properties are gathered in Table 3. γ(Mn,Cu) solid solutions are distin-guished by a very high electrical resistivity but a low thermal stability. Fe additions cause a segregationof the (αMn) phase and a strong embrittlement of the alloy and a sharp decrease of resistivity [1958Top].During the rolling process of mild steels, surface cracks or fissures may appear, which can lead to the scrap-ping of considerable quantity of steel. Cu is the main detrimental element associated with this phenomenon[1966Sal] known as ‘hot shortness’. Non magnetic Cu-Fe-Mn alloys (20-55 mass% Cu, < 37.5 mass% Fe)exhibit the Elinvar properties with high corrosion resistivities, which makes these alloys suitable for preci-sion instruments [1973Mas]. The studies of evaporation using the levitation melting technique [1974Fis]offer useful informations to optimize experimental parameters for the removal of undesired Cu in liquidsteels. The interest of copper precipitation in steels is receiving much attention with respect to precipitationhardening [2003Des, 2006Zha] which provides a good combination of strength, roughness and weldability.This result is attributed to the precipitation of coherent Cu particles before peak hardening is achieved.

Miscellaneous

Cu has been used to stabilize γ(FexMn1–x) alloys [1981Dem1]. Discontinuities have been observed in physicalproperties such as Neel temperature, electrical resistivities between x = 0.4 and x = 0.5 [1974Ish]. The Neeltemperature presents a maximum around x = 0.4 and 200°C for the composition (Fe0.45Mn0.55)0.95Cu0.05. Sucha maximum exists probably for the alloy Fe0.45Mn0.55 around 220°C. Cu and Mn diffusion in γFe has beeninvestigated by [1975Lar] at 1100°C. Cu enters in austenite grains by penetration of a liquid rather than bysolid state diffusion; Mn is governed by a substantially slower grain boundary diffusion mechanism.

Table 1. Investigations of the Cu-Fe-Mn Structures, Phase Equilibria and Thermodynamics


[1913Par] Micrography, thermal analysis 880-1450°C, diagrams established in the wholecomposition domain

[1977Has] Electron Microprobe Analysis (EMPA),diffusion couple

600-1350°C, diagrams established in the wholecomposition domain

[1997Oht] Diffusion couples, scanning electronmicroscopy

1100-1300°C, solid-liquid equilibria

[2004Wan1] Micrography, metastable miscibility gapin the liquid

< 1800°C, 48 masss% Fe, 48 mass% Cu,4 mass% Mn

Cu–Fe–Mn 3







Comments/References

α,(δFe,δMn,αFe)(δFe)1538 - 1394(δMn)1246 - 1138

(αFe) (Ferrite)< 912

cI2Im�3mW

a = 293.15

a = 308.0

a = 286.65

at 1394°C. Dissolves 10 at.% Mn at1474.5°C [2004Wit]. Dissolves 5.6at.% Cu at 1487°C [2007Tur][Mas2]. Dissolves 12.2 at.% Fe at1236.6°C [2004Wit]. Dissolves12.9 at.% Cu at 1097°C [2006Tur]pure Fe at 20°C [Mas2, V-C2].Dissolves 5 at.% Mn at 527°C[2004Wit].Dissolves 1.4 at.% Cu at 847°C[2007Tur]

(εFe) hP2P63/mmcMg

a = 246.8c = 396.0

at 25°C, > 13 GPa [Mas2]

γ, (γMn,γFe,Cu)(γFe) (Austenite)1394 - 912

(γMn)1138 – 1087

(Cu)< 1084.62

cF4Fm�3mCu

a = 364.67

a = 386.0

a = 361.46

at 915°C [V-C2, Mas2]. Dissolves14.2 at.% Cu at 1425°C [2007Tur];complete solubility with (γMn)[Mas2]. Complete solubility with(γFe) and (Cu)pure Cu at 25°C [Mas2]. Dissolves5 at.% Fe at 1095°C; complete solu-bility with (γFe) and (γMn)

(βMn)1087 - 707

cP20P4132βMn

a = 631.52 [Mas2]

(αMn)< 707

cI58I�43mαMn

a = 891.26 pure Mn at 25°C [Mas2]

γ1, MnCu5≤ 410

cF* ordered γ phase[2006Tur]

γ2, MnCu3≤ 450


γ3≤ 700


(Cu,γMn)’ t** metastable [2006Tur]

4 Cu–Fe–Mn


Table 3. Investigations of the Cu-Fe-Mn Materials Properties

Reference Method / Experimental Technique Conditions / Type of Property

[1966Sal] Micrography, Electron MicroprobeAnalysis (EMPA)

800-1250°C, < 10 at.% Cu, < 12 at.% Mn

[1969End] Magnetic and electric properties,Mössbauer

4.2-600 K, Mn0.95Cu0.05 and(Mn0.95Cu0.05)0.99Fe0.01

[1970Win,1971Win]

Magnetic properties, Mössbauer 4.2-300 K, (MnxCu1–x)0.99Fe0.01

[1973Mas] Hardness, Young’s modulus, rigiditymodulus

25°C, 20-55 at.% Cu, 2-30 at.% Fe

[1974Fis] Vaporization measurements by levitationmelting

1760-1930°C, < 0.1 MPa of Ar

[1974Ish] Neel temperature, electrical conductivitiesmeasurements

< 220°C, (FexMn1–x)0.95Cu0.05(0 < x < 1)

[1975Lar] EMPA, diffusivity measurements 1100°C, Cu and Mn in γFe

[1978Nik] XRD, electrostriction 77-500 K, 17-20 mass% Mn, 2-5 mass%Cu

[1980Dem] XRD, diffuse scattering, twinning < 200°C, 10 mass% Fe, 3 mass% Cu

[1980Tka] Dilatometry, strain-stress curves < 900°C, 7 mass% Mn, < 2 mass% Cu

[1981Dem1] XRD, cubic -tetragonal- orthorhombictransitions

< 250°C, 4 mass% Cu, 10-26 mass% Fe

[1981Dem2] Mössbauer, cubic-tetragonal transition <250°C, 10 mass% Cu, 1 mass% Fe

[2003Des] Transition Electron Microscopy (TEM),small angles diffraction

500°C, 0.78 mass% Cu, < 0.35 mass% Mn,precipitation strengthening

[2005Res] XRD, Mössbauer, grain size measurements Cu70Fe15Mn15, ball milled

[2005Zha] XRD, TEM, DSC < 350°C, Cu5.0Fe9.5Mn85.5, magnetic shapememory effect

Cu–Fe–Mn 5



Fig. 1. Cu-Fe-Mn. Liquidus surface projection

6 Cu–Fe–Mn


Fig. 2. Cu-Fe-Mn. Solidus surface projection

Cu–Fe–Mn 7



Fig. 3. Cu-Fe-Mn. Calculated metastable miscibility gap in the liquid phase

8 Cu–Fe–Mn


Fig. 4. Cu-Fe-Mn. Calculated metastable miscibility gap in the γ phase

Cu–Fe–Mn 9



Fig. 5. Cu-Fe-Mn. Isothermal section at 1300°C

10 Cu–Fe–Mn



Cu–Fe–Mn 11




12 Cu–Fe–Mn



Cu–Fe–Mn 13




14 Cu–Fe–Mn



Cu–Fe–Mn 15



Fig. 11. Cu-Fe-Mn. Partial isothermal section at 750°C

16 Cu–Fe–Mn


Fig. 12. Cu-Fe-Mn. Partial isothermal section at 550°C

Cu–Fe–Mn 17



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[1975Lar] Larsson, L.-E., “Note on the Grain Boundary Penetration of Cu and Mn in γ-Fe at 1100°CDuring Liquid-Solid Powder Alloying”, Mater. Sci. Eng., 19, 241–244 (1975) (Experimen-tal, Kinetics, Transport Phenomena, 10)

[1977Has] Hasebe, M., Nishizawa, T., “Analysis and Synthesis of Phase Diagrams of the Fe-Cr-Ni, Fe-Cu-Mn, and Fe-Cu-Ni”, NBS Spec. Publication, 496(I), 911–954 (1977) (Experimental,Phase Diagram, Phase Relations, Thermodyn., #, 109)

[1978Nik] Nikolin, B.I., Makogon,, Yu.N., “ε’-Martensite in Carbon-Free Iron Manganese Alloys withCopper” (in Russian), Akad. Nauk Ukr. SSR, Metallofizika, 74, 103–105 (1978) (Electr.Prop., Crys. Structure, Experimental, 8)

[1979Cha] Chang, Y.A., Neumann, J.P., Mikula, A., Goldberg, D., “Cu-Fe-Mn” in “INCRA MonographSeries. 6. Phase Diagrams and Thermodynamic Properties of Ternary Copper-Metal Sys-tems”, Uni. Wisconsin-Milwaukee, USA, 461–463 (1979) (Crys. Structure, Phase Relations,Review, 4)

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18 Cu–Fe–Mn


[1980Dem] Demin, S.A., Ustinov, A.I., Chuistov, K.V., “Formation of a Twin Structure in Alloys Baseon γ -Manganese During the Faced-Centred Cubic to Face-Centred Tetragonal Transforma-tion” Phys. Met. Metallogr., 50(3), 86–91 (1980), translated from Fiz. Met. Metalloved., 50(3), 553–559 (1980) (Crys. Structure, Experimental, 11)

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[1983Riv] Rivlin, V.G., Raynor, G.V., “Phase Equilibriums in Iron Ternary Alloys 9: Critical Review ofConstitution of Ternary Systems Fe-Mn-X (X = Ti, V, Cr, Co, Ni, Cu)”, Int. Met. Rev., 28(1),23–64 (1983) (Phase Diagram, Phase Relations, Review, 68)

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Cu–Fe–Mn 19



[2005Zha] Zhang, J.H., Peng, W.Y., Chen, S., Hsu, T.Y., “Magnetic Shape Memory Effect in an Anti-ferromagnetic γ-Mn-Fe(Cu) Alloy”, Appl. Phys. Lett., 86(2), 022506-1–3 (2005) (Crys.Structure, Experimental, Magn. Prop., 20)

[2006Zha] Zhang, C., Eanomoto, M., “Study of the Influence of Alloying Elements on Cu Precipitationin Steel by Non-Classical Nucleation Theory”, Acta Mater., 54(16), 4183–4191 (2006) (Cal-culation, Morphology, Phase Relations, 32)

[2006Tur] Turchanin, M., Agraval, P., Gröbner, J., Matusch, D., Turkevich, J., “Cu-Mn (Copper-Manganese)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.),MSI, Materials Science International Services GmbH, Stuttgart; Document ID:20.14136.1.20 (2006) (Phase Diagram, Phase Relations, Crys. Structure, #, 25)


[2007Wit] Witusiewicz, V.T., private communication to MSI, (2007)[Mas2] Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International,

Metals Park, Ohio (1990)[V-C2] Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic


20 Cu–Fe–Mn


Copper – Iron – Molybdenum

Nataliya Bochvar, Lazar Rokhlin

Introduction

Cu rich alloys having high strength and high electrical conductivity, and age-hardenable Fe rich alloyscan be developed through consideration of the Cu-Fe-Mo phase diagram. Molybdenum is added to steelstraditionally in order to improve their mechanical properties, in particular ductility, since molybdenumpromotes grain refinement. The additions of copper to steels increase their corrosion resistance and improveformability.First investigation of the Cu-Fe-Mo phase diagram was carried out by [1938Dan], who studied Fe richalloys (up to 25 mass% Cu and Mo) by thermal and microscopic analyses. Eight vertical sections, at 5,10, 15, 20 and 25 mass% Mo, and at 5, 10 and 15 mass% Cu, were constructed. The schematic liquidus sur-face and the reaction scheme were proposed. According to [1938Dan], the Cu-Fe-Mo phase diagram ischaracterized by a wide liquid miscibility gap which spreads from the binary Cu-Mo system into the ternarysystem. [1938Dan] presented eight four-phase equilibria and determined experimentally temperatures forthree of them. However, [1938Dan] assumed only two binary compounds in the Fe-Mo system, μ and σ,because two other compounds, R and λ were still not discovered. A new version of the Cu-Fe-Mo phasediagram taking into account the existence of R and λ was suggested by [1992Rag], who corrected the sche-matic liquidus projection in the Cu corner and the reaction scheme.Phase equilibria in the Cu-Fe-Mo system in the temperature range of 800 to 1300°C were studied by[2000Wan] using optical metallography and SEM/energy-dispersive X-ray analysis. Using the limitedexperimental data and the binary interaction parameters, [2000Wan] performed a thermodynamic calcula-tion of four isothermal sections at 900, 1100, 1300 and 1500°C and of two vertical sections at 5 and10 mass% Mo. However, the calculated vertical sections do not agree with the experimental vertical sectionsconstructed by [1938Dan]. [2000Wan] presented also calculated miscibility gap of the liquid phase between2727 and 1227°C considering it as metastable.[2004Wan1] calculated the limits of the miscibility gap of the liquid phase along the (Fe-4 mass% Mo) -(Cu-4 mass% Mo) section.There are several reviews on the Cu-Fe-Mo phase diagram. [1949Jae] discussed the data of [1938Dan] andpresented the liquidus surfaces with the miscibility gap in the whole concentration range, the liquidus sur-face in Fe rich alloys being after [1938Dan]. [1979Cha] presented the liquidus surface of the phase diagramand the four-phase equilibria after [1938Dan], however, noting that additional four-phase equilibria wouldoccur in the ternary system because of the presence of two new binary phases discovered later. [2002Rag]discussed the results of [2000Wan] and constructed three isothermal sections of the phase diagram at 1300,1100 and 900°C, based on [2000Wan], but slightly modified them to agree with the accepted binary data. Asit was noted by [2002Rag], the isothermal section at 1500°C thermodynamically calculated by [2000Wan]needs experimental confirmation.Literature on the experimental and thermodynamic data on phase relations in the Cu-Fe-Mo system is sum-marized in Table 1.

Binary Systems

The binary Cu-Fe and Cu-Mo phase diagrams are accepted from [2007Tur] and [2007Lys], respectively. Thebinary Fe-Mo phase diagram is accepted from [1982Gui, Mas2].

Solid Phases

There are no ternary compounds in the Cu-Fe-Mo system. The unary and binary phases discussed in thisassessment are listed in Table 2.

Cu–Fe–Mo 1




The reaction scheme is shown in Fig. 1 after [1938Dan] and [1992Rag] in the Fe rich and in the Cu richcorners, respectively. Eleven four-phase and two three-phase invariant equilibria were established in theCu-Fe-Mo system. The types of some reactions in the Cu corner were corrected here for compatibilityof the reaction scheme with the liquidus surface. The temperatures of the invariant equilibria in the pointse2(max), E1, E2, E3 and U5 were experimentally determined by [1938Dan]. The temperature in the pointU3 is estimated to be above 1300°C taking into account the isothermal section at 1300°C [2000Wan]. Simi-larly the temperature in the point U6 is estimated to be above 900°C taking into account the isothermal sec-tion at 900°C [2000Wan]. The temperatures of all other invariant reactions were assumed speculativelytaking into account data of [1992Rag, 2000Wan]. The compositions of the phases participating in the invar-iant reactions e2(max), E1 and E2 were estimated from the figures in [1938Dan]. They are presented inTable 3. The compositions of other phases were not established.

Liquidus Surface

In the Cu-Fe-Mo system the miscibility gap in the liquid state is well established. There are two liquidphases, Mo rich (L′) and Cu rich (L″), in a wide composition range. A critical point c was identified onthe monovariant line bordering the miscibility gap on the Cu-Fe side. Its position is at about 1450°C, about4 mass% Cu and 35 to 40 mass% Mo. These data were compiled from [1979Cha] and [1992Rag]. The liqui-dus surface of the Cu-Fe-Mo system is shown in Fig. 2 according to [1938Dan] without details of the Cucorner. The liquidus surface in the Cu corner is shown schematically in Fig. 3 according to [1992Rag].The calculated metastable miscibility gap in the liquid-phase region of the Cu-Fe-Mo system [2000Wan] isshown in Fig. 4. The tie lines in the miscibility gap are located along the radial lines from the Cu corner tothe Fe-Mo side.

Isothermal Sections

The calculated isothermal sections at 1300, 1100 and 900°C are shown in Figs. 5, 6 and 7, respectively.They are given according to [2002Rag], based, in turn, on data of [2000Wan], with minor correctionsaccording to the accepted binary systems.


In Fig. 8 a partial vertical section at 5 mass% Mo is presented. The section is constructed using the respec-tive vertical section of [1938Dan] with major corrections taking into account the isothermal sections at 900,1100 and 1300°C [2002Rag]. In Fig. 9 the calculated vertical section of the miscibility gap along the sectionFe-4 mass% Mo - Cu-4 mass% Mo is shown after [2004Wan1].


In [1991Par] the cast iron containing small contents of Cu and Mo was studied aiming to improve mechan-ical properties. Phase transformations during heat treatment and their influence on strength, ductility andfracture toughness were established.[2004Rad] studied the effect of added molybdenum powder on properties of sintered fine-grained Cu-Fealloys. The mechanical alloying was used for the preparation of the powder mixtures. The specimens weresintered in the range 600 to 1130°C. Additions of Mo powder accelerated the sintering process.

Miscellaneous

[1971Laz] examined the volume and boundary diffusion of Cu in the Fe-Mo alloy with 0.7 mass% Mo inthe (γFe) region. It was established that Mo did not effect the mobility of Cu atoms in the (γFe) phasevolume, but increased diffusion mobility of Cu atoms along the grain boundaries.

2 Cu–Fe–Mo


[1994Bah] constructed the isothermal transformation diagrams for a ductile cast iron alloyed with Mn, Moand Cu. The isothermal transformations proceeded over the temperature range of 275 to 650°C for timesbetween 0.5 and 120 min after austenitization at 870 and 920°C.[2004Wan1] compiled a thermodynamic database of the Cu-Fe-X (X: Al, Co, Cr, Mn, Mo, Ni, V) systemsby the CALPHAD method. According to [2004Wan1], on the basis of this database much information con-cerning stable and metastable phase equilibria, isothermal and vertical sections, molar fractions of constitu-ent phases, liquidus surface projection, etc., could be predicted.[2004Wan2] experimentally studied the effect of alloying elements (C, Cr, Mo, Nb, Si, V) on the macro-scopic morphologies of the Cu-Fe base alloys. It was shown that the core-type macroscopic morphologywas formed at addition of Mo to the Cu-Fe base alloy. Formation of such morphology was assumed tobe due to the stable miscibility gap in the liquid phase in the Cu-Fe-Mo system. Interfacial energy betweentwo liquid phases was said to play a key role in the formation of the core-type macroscopic morphology.

Table 1. Investigations of the Cu-Fe-Mo Phase Relations, Structures and Thermodynamics


[1938Dan] Thermal and microscopic analysis.Electrolytic Cu, Swedish charcoal Fe, wireof Mo (purities not stated). Melting in ahigh frequency induction furnace.Theoretical generalization based on thebinary systems.

Fe corner, alloys with Cu and Mo up to 25mass% each. Fe rich part of the liquidussurface; vertical section at 5, 10, 15, 20, 25mass% Mo; 5, 10, 15 mass% Cu.Schematic liquidus surface and reactionscheme.

[1992Rag] Theoretical modification of [1938Dan] toaccount for the existence of two newcompounds in the Fe-Mo system

Schematic reaction scheme and liquidusprojection in the Cu corner

[2000Wan] Electrolytic Co (99.99%), Fe (99.99%), Mo(99.5%). Melting in alumina crucibles in ahigh induction furnace under an Agatmosphere. Hot-rolling at 800°C. Opticalmicroscopy, SEM/energy dispersive X-rayanalysis.Thermodynamic calculations with theRedlich-Kister model and thermodynamicparameters evaluated with the PARROTsoftware.

The alloys of the compositions near Cu/Fe:50/50 with Mo from 0 to 6 mass%.Annealing at 800 to 1300°C for 3 to 1680 h.Experimentally determined compositions ofthe phases in equilibrium at 1300, 1200,1100, 1000, 900, 800°C.Calculated isothermal sections at 1500,1300, 1100, 900°C; vertical sections at 5and 10 mass% Mo and up to 30 mass%Cu; metastable miscibility gap of the liquidphase.

[2004Wan1] Thermodynamic calculation usingCALPHAD method.

Vertical section at 10 mass% Mo and up to30 mass% Cu; miscibility gap along the Fe-4 mass% Mo - Cu-4 mass% Mo section

Cu–Fe–Mo 3







Comments/References

(Cu)< 1084.62

cP4Fm�3mCu

a = 361.46

dissolves up to 5 at.% Fe at 1095°C[2007Tur] and up to 0.061 at.% Mo at1083.4°C [2007Lys]at 25°C [Mas2]

(δFe)1538 - 1394

cI2Im�3mW a = 293.15

dissolves up to 5.8 at.% Cu at 1487°C[2007Tur][Mas2]

(γFe)< 1394 - 912

cP4Fm�3mCu

a = 364.67

dissolves ~13 at.% Cu at 1450°C[2007Tur] and ~4 at.% Mo at 1100°C[Mas2]at 915°C [Mas2, V-C2]

(αFe)< 912

cI2Im�3mW

a = 286.65

dissolves up to 1.6 at.% Cu at 847°C[2007Tur] and ~24 at.% Mo at 1449°C[Mas2]at 25°C [Mas2]

(Mo)< 2623

cI2Im�3mW

a = 314.7

dissolves 1.6 at.% Cu at 2515°C[2007Lys] and 32 at.% Fe at 1611°C[Mas2]at 25°C [Mas2]

λ, MoFe2< 927

hP12P63/mmcMgZn2

a = 474.4c = 772.5

33.3 at.% Mo [Mas2][1982Gui]

R1488 - 1200

hR159R�3Co5Cr2Mo3

a = 1095.6c = 1935.3

33.9 to 38.5 at.% Mo [Mas2][V-C2]

μ, Mo6Fe7< 1370

hR39R�3mFe7W6

a = 475.1c = 256.8

39.0 to 44.0 at.% Mo [1982Gui, Mas2][1982Gui]

σ, MoFe1611 - 1235

tP30P42 /mnmσCrFe

a = 921.8c = 481.3

42.9 to 56.7 at.% Mo[1982Gui, Mas2][1982Gui]

4 Cu–Fe–Mo




Cu Fe Mo

l′ ⇌ l″ + (αFe) 1450 e2(max) l′ 14.9 77.6 7.5

L′ ⇌ L″ + (αFe) + (γFe) 1420 E1 L′ 4.6 70.1 25.3

L′ ⇌ L″ + (αFe) + R 1420 E2 L′ 8.8 76.2 15.0

Cu–Fe–Mo 5



Fig.1.

Cu-Fe-Mo.

Reactionscheme

6 Cu–Fe–Mo


Fig. 2. Cu-Fe-Mo. Liquidus surface projection

Cu–Fe–Mo 7



Fig. 3. Cu-Fe-Mo. Enlarged scheme of the liquidus surface in the Cu corner

8 Cu–Fe–Mo


Fig. 4. Cu-Fe-Mo. Calculated metastable miscibility gap. Solid thick lines: isotherms of the miscibility gap in theliquid state. Dashed lines: the tie lines between liquid Cu and the Fe-Mo alloys at different Fe/Mo ratios. Solid thin lines:the tie lines between L′ and L″ at 2727°C

Cu–Fe–Mo 9



Fig. 5. Cu-Fe-Mo. Calculated isothermal section at 1300°C

10 Cu–Fe–Mo



Cu–Fe–Mo 11




12 Cu–Fe–Mo


Fig. 8. Cu-Fe-Mo. Calculated vertical section at 5 mass% Mo, plotted in at.%

Cu–Fe–Mo 13



Fig. 9. Cu-Fe-Mo. Calculated stable miscibility gap of the liquid phase along the Fe-4 mass% Mo - Cu-4 mass% Mosection, plotted in at.%

14 Cu–Fe–Mo


References[1938Dan] Dannoehl, W., “The Alloys of Iron, Copper and Molybdenum” (in German), Wiss. Veroeff.

Siemens-Werk, 17, 113–125 (1938) (Phase Diagram, Phase Relations, Morphology, Experi-mental, 5)

[1949Jae] Jaenecke, E., “Cu-Fe-Mo” (in German) in “Kurzgefasstes Handbuch aller Legierungen”,Winter Verlag, Heidelberg, 661–664 (1949) (Phase Diagram, Phase Relations, Review, 1)

[1971Laz] Lazarev, V.A., Golikov, V.M., “Diffusion of Copper in Iron and Iron-Boron and Iron-Molybdrnum Alloys”, Phys. Met. Metallogr., 31(4), 213–214 (1971), translated from Fiz.Met. Metalloved., 31(4), 885–886 (1971) (Experimental, 8)

[1979Cha] Chang, Y.A., Neumann, J.P., Goldberg, D., “Cu-Fe-Mo” in “INCRA Monograph Series. 6.Phase Diagrams and Thermodynamic Properties of Ternary Copper-Metal Systems”, Uni.Wisconsin-Milwaukee, USA, 464–467 (1979) (Phase Diagram, Phase Relations, Review, 2)

[1982Gui] Guillermet, F., “The Fe-Mo (Iron-Molybdenum) System”, Bull. Alloy Phase Diagrams, 3(3),359–366 (1982) (Phase Diagram, Phase Relations, Crys. Structure, Thermodyn., Review, 40)

[1991Par] Park, Hyun-Ku,, Park, Seung-Ho,, MLee, Won-Sik, Kang, In-Chan, “Fracture Properties onthe Austempering Control Times of Mo-Cu Ductile Cast Iron” (in Korean), J. Korean Inst.Met., 29(7), 665–670 (1991) (Experimental, Mechan. Prop., Morphology, 11)

[1992Rag] Raghavan, V., “Cu-Fe-Mo (Copper-Iron-Molybdenum) System” in “Phase Diagram of Tern-ary Iron Alloys”, Ind. Inst. Metal, Calcutta, 722–725 (1992) (Phase Diagram, Phase Rela-tions, Review, 3)

[1994Bah] Bahmani, N., Elliott, R., “Isothermal Transformation Diagrams for Alloyed Ductile Cast-Iron”, Mater. Sci. Technol., 10(12), 1050–1056 (1994) (Experimental, Morphology, 9)

[2000Wan] Wang, C.P., Liu, X.J., Ohnuma, I., Kainuma, P., Hao, S.M., Ishida, K., “Phase Equlibria inthe Cu-Fe-M- and Cu-Fe-Nb Systems”, J. Phase Equilib., 21(1), 54–62 (2000) (Phase Dia-gram, Phase Relations, Thermodyn., Calculation, Experimental, 22)

[2002Rag] Raghavan, V., “Cu-Fe-Mo (Copper-Iron-Molybdenum)”, J. Phase Equilib., 23(3), 260–262(2002) (Phase Diagram, Phase Relations, Review, 6)

[2004Rad] Radchenko, P.Ya, Radchenko, O.G., Panichkina, V.V., Skorokhod, V.V., “Theory, Technol-ogy of Sintering, Heat and Chemical Heat-Treatment Processes. Effect of MolybdenumAdditions on Compaction During Sintering of Fine-Grained Iron-Copper Pseudoalloys”,Powder Metall. Met. Ceram., 43(3-4), 132–137 (2004), translated from Poroshk. Metall.,43(3-4), 26–32 (2004) (Experimental, Electr. Prop., Mechan. Prop., Kinetics, 7)

[2004Wan1] Wang, C.P., Liu, X.J., Ohnuma, I., Kainuma, P., Ishida, K., “Thermodynamic Database of thePhase Diagram in Cu-Fe Base Ternary Systems”, J. Phase Equilib., Diffus., 25(4), 320–328(2004) (Phase Diagram, Phase Relations, Thermodyn., Calculation, 40)

[2004Wan2] Wang, C.P., Liu, X.J., Takaku, Y., Ohnuma, I., Kainuma, P., Ishida, K., “Formation of Core-Type Macroscopic Morphologies in Cu-Fe Base Alloys with Liquid Miscibility Gap”,Metall. Mater. Trans. A., 35A(4), 1243–1253, (2004) (Experimental, Morphology, PhaseRelations, Phase Diagram, Thermodyn., 31)

[2007Lys] Lysova, E., Rokhlin, L., “Cu-Mo (Copper-Molybdenum)”, MSIT Binary Evaluation Pro-gram, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Ser-vices GmbH, Stuttgart; to be published (2007) (Phase Diagram, Crys. Structure,Thermodyn., Assessment, 31)




Cu–Fe–Mo 15



Copper – Iron – Niobium

Nataliya Bochvar, Lazar Rokhlin

Introduction

Cu rich alloys having high strength and high electrical conductivity, and age-hardenable Fe rich alloys canbe developed through consideration of the Cu-Fe-Nb phase diagram. However, very little attention has beenpaid to phase equilibria in this system. There is only one experimental work [2000Wan], where the coordi-nates of the three-phase field L’+ε+(γFe) at 1300, 1200 and 1100°C were determined experimentally. Usingthese data, [2000Wan] thermodynamically calculated the sections at 1500, 1300 and 1100°C, and[2004Wan1] calculated the isothermal section at 1200°C in the whole concentration interval.[2000Wan] established the existence of the metas and stable miscibility gaps in the liquid-phase region ofthe Cu-Fe-Nb system and presented calculated isothermal section of the phase diagram at 1500°C withthe stable miscibility gap and isotherms of the metastable miscibility gap of the liquid phase between2727 and 1227°C.[2004Wan1] calculated also the temperature-concentration limits of the miscibility gap of the liquid phasealong the (Fe-4 mass% Nb) - (Cu-4 mass% Nb) section.[2002Rag] reviewed the results of [2000Wan] and slightly modified the isothermal sections of the phase dia-gram to meet the accepted binary data. [2002Rag] postulated four-phase invariant reactions in the Cu-Fe-Nbsystem based on the calculated isothermal sections of [2000Wan].

Binary Systems

The binary Cu-Fe and Cu-Nb phase diagrams are accepted after [2007Tur] and [2002Rom],respectively. The binary Fe-Nb phase diagram is accepted after [2000Tof]. The Fe-Nb phase diagramassessed thermodynamically by [2000Tof] is shown in Fig. 1. [2002Oka] remarks that this diagram appearsto be a better presentation of the experimental phase boundary data, including the ε and μ phases.

Solid Phases

There are no ternary compounds in the Cu-Fe-Nb system. The unary and binary phases discussed in thisassessment are listed in Table 1.


Seven four-phase invariant equilibria were postulated by [2002Rag] in the Cu-Fe-Nb system. Six of themoccur with the participation of the liquid phase (points U1 to U4, E1, E2) and one equilibrium (point E3)occurs in the solid state. The reaction scheme is shown in Fig. 2 according to [2002Rag] with some correc-tions. The corrections concernthe point c3 indicated in the reaction scheme by [2002Rag]. Two monovariant lines L’’+γ+ε and L’+L’’+εconverge at this point with falling temperature and one monovariant line L’+γ+ε descends from it, suggest-ing the existence of a four-phase transition reaction with L’, L’’, γ and ε. However, no invariant reactionwas suggested by [2002Rag].The invariant temperatures were accepted speculatively after [2002Rag], taking into account, however,the phase equilibria established by [2000Wan] for 1500, 1300 and 1100°C. The reaction scheme suggestedby [2002Rag] is speculative at the temperatures below 1100°C because of the absence of the experimentaldata at these temperatures. However, taking into account the existence of the certain three-phase regions:(Nb)+(Cu)+μ, μ+(Cu)+ε, ε+(Cu)+(αFe) at room temperature, the part of the reaction scheme below1100°C is quite probable.

Cu–Fe–Nb 1



Liquidus Surface

The liquidus surface of the Cu-Fe-Nb system was not studied experimentally. Thermodynamic calculation[2000Wan] results in a miscibility gap in the liquid state, the thermodynamic parameters being evaluated onthe basis of own experimental results.The schematic liquidus surface shown in Fig. 3 in the present assessment is constructed speculatively basingon the invariant equilibria in the Cu-Fe-Nb system (Fig. 2). There are two liquid phases in the system. Theliquid phase bordering with (Nb) is denoted as L’’ and the liquid phase near the Cu corner is denoted as L’.Two critical points, c1 and c2, exist on the monovariant line outlining the miscibility gap where the L’ + L’’field contacts with L’’ + (Nb) and L’ + ε fields, respectively. As indicated by [2002Rag], temperatures of c1and c2 are about 1700°C.The calculated metastable miscibility gap (see above) in the liquid-phase region of the Cu-Fe-Nb system isshown in Fig. 4 [2000Wan]. A miscibility gap island exists inside the ternary system. The tie-lines in theimmiscible region lie along the radial lines from the Cu corner to the Fe-Nb side. In Fig. 4 they are shownby dashed lines for different temperatures at different Fe/Nb ratios. Solid lines in Fig. 4 show the tie-lines at1827°C.

Isothermal Sections

Only compositions of the coexisting phases in the three-phase region L+(γFe)+ε at 1100, 1200 and 1300°Cwere studied experimentally by [2000Wan] be means of optical microscopy and SEM/energy-dispersive X-ray analysis. The alloys of the compositions near Cu/Fe = 50/50 with 1, 2, 4 and 6 mass% Nb were preparedby melting in alumina crucibles in a high-frequency induction furnace under Ar atmosphere. Hot rolling of theingots at 800°C followed by a solution treatment at 900°C for 24 h was used to obtain strips for preparation ofthe specimens. The specimens were sealed in quartz capsules and equilibrated at various temperatures in therange from 1100 to 1300°C for up to 1660 hours followed by quenching into iced water. On the basis of theseexperimental data [2000Wan] evaluated the thermodynamic parameters for the Cu-Fe-Nb system and calcu-lated the isothermal section at 1500, 1300 and 1100°C in all concentration range. The calculated isothermalsections at 1300 and 1100°C are shown in Figs. 5 and 6, respectively. They are reproduced from[2002Rag] who modified the versions of [2000Wan] to agree with the modern accepted binary data. Some cor-rections were made according to the binary systems accepted in the present assessment. The section at 1500°C[2000Wan] was not confirmed by the experimental data and, therefore, is not reproduced in the presentevaluation.


Figure 7 presents the limits of the miscibility gap along the section Fe-4Nb-Cu-4Nb (mass%) of the Cu-Fe-Nb system calculated by [2004Wan1].

Miscellaneous

[1990Toy] measured the Meissner effect in Cu-clad Nb wires doped with Fe down to 40 mK. They foundthe thickness of the Meissner region in Cu to increase in proportion to T–1/2, as in case of the normal impur-ity presence. The thickness was determined mainly by the electron mean free path in Cu and the interfacebetween Cu and Nb. The T –1/2 dependence of thickness suggests the pair breaking effect weakens linearlywith decreasing temperature below 1 K due to magnetic impurities.In [1996Yav] a model based on the kinetic effect of diffusion double-layers of larger atoms of Nb, and smallerand faster diffusing atoms of Cu was presented. These atoms were rejected into the surrounding amorphousmatrix at the αFe nanocrystallization front. It was shown that the effect of Cu addition on grain size of nano-crystallized FeB based amorphous precursors was explainable through the formation of diffuse interfacesbetween the Fe rich amorphous phase and amorphous Cu rich clusters on which αFe nanocrystals could nucle-ate with reduced interfacial energy. The stronger nucleating effect of Cu addition in Nb-containing FeB basedalloys was explained in terms of CuFe metastable phase diagram that indicated the Cu-enriched clusters inthese alloys should crystallize into bcc structure. Addition of larger atoms of Nb that were rejected together

2 Cu–Fe–Nb


with B atoms at the αFe nanocrystallization interface was found to generate diffusion double-layers with sharpconcentration gradient. The concentration gradient had an additional thermodynamic stabilization effect on theamorphous interlayers by reducing the driving force for the intermetallic formation.[2004Wan1] compiled a thermodynamic database of the Cu-Fe-X (X: Al, Co, Cr, Mn, Mo, Ni, Nb, V) sys-tems by the CALPHAD method. According to [2004Wan1], on the basis of this database much informationconcerning stable and metastable phase equilibria, isothermal and vertical sections, molar fractions of con-stituent phases, liquidus surface projection, etc., could be predicted.[2004Wan2] experimentally studied the effect of alloying elements (C, Cr, Mo, Nb, Si, V) on the macro-scopic morphologies of the Cu-Fe base alloys. It was shown that core-type macroscopic morphology wasformed by the addition of Nb in the Cu-Fe base alloy. Formation of such morphology was assumed to bedue to the stable miscibility gap in the liquid phase in the Cu-Fe-Nb system. Interfacial energy betweentwo liquid phases was said to play a key role in the formation of the core-type macroscopic morphology.





Comments/References

(Cu)< 1084.62

cP4Fm�3mCu

a = 361.46

dissolves up to 5 at.% Fe at 1095°C [2007Tur] and up to 0.2at.% Nb at 1091°C [2002Rom]

at 25°C [Mas2]

(δFe)1538 - 1394

cI2Im�3mW

a = 293.15

dissolves up to 5.8 at.% Cu at 1487°C [2007Tur]and 3.1 at.% Nb at 1394°C [2000Tof]

[Mas2]

(γFe)1394 - 912

cP4Fm�3mCu

a = 364.67

dissolves ~13 at.% Cu at 1450°C [2007Tur] and 1.15 at.%Nb at 1199°C [2000Tof]

at 915°C [Mas2, V-C2]

(αFe)< 912

cI2Im�3mW

a = 286.65

dissolves up to 1.6 at.% Cu at 847°C [2007Tur] and 0.026at.% Nb at 942°C [2000Tof]

at 25°C [Mas2]

(Nb)< 2477

cI2Im�3mW

a = 314.7

dissolves 1.2 at.% Cu at 109°C [2002Rom] and 7.7 at.% Feat 1484°C [2000Tof]

at 25°C [Mas2]

(continued)

Cu–Fe–Nb 3






Comments/References

ε, NbFe2< 1629

hP12P63 /mmcMgZn2

a = 483.63c = 789.19a = 484.06c = 789.66a = 485.04c = 790.93

30.5 to 35.0 at.% Nb [2000Tof]at 32.3 at.% Nb, quenched from 1000°C [1999Zhu]

at 33.3 at.% Nb, quenched from 1000°C [1999Zhu]

at 34.3 at.% Nbquenched from 1000°C [1999Zhu]

μ, Nb6Fe7< 1520

hR39R�3mFe7W6

a = 493.2c = 268.1

45.5 to 49.4 at.% Nb [2000Tof]at 48 at.% Nb [1986Pau]

Fig. 1. Cu-Fe-Nb. Fe-Nb phase diagram

4 Cu–Fe–Nb


Fig.2.

Cu-Fe-Nb.

Assum

edreactio

nscheme

Cu–Fe–Nb 5



Fig. 3. Cu-Fe-Nb. Speculative scheme of the liquidus surface projection

6 Cu–Fe–Nb


Fig. 4. Cu-Fe-Nb. Calculated metastable miscibility gap of the liquid phase

Cu–Fe–Nb 7



Fig. 5. Cu-Fe-Nb. Calculated isothermal section at 1300°C

8 Cu–Fe–Nb


Fig. 6. Cu-Fe-Nb. Calculated isothermal section at 1100°C

Cu–Fe–Nb 9



Fig. 7. Cu-Fe-Nb. Calculated stable miscibility gap of the liquid phase along the Fe-4 mass% Nb - Cu-4 mass% Nbsection, plotted in at.%

10 Cu–Fe–Nb


References[1986Pau] Paul, E., Swartzendruber, L.J., “The Fe-Nb (Iron-Niobium) System”, Bull. Alloy Phase Dia-

grams, 7(3), 248–254 (1986) (Phase Diagram, Phase Relations, Crys. Structure, Thermo-dyn., Review, 54)

[1990Toy] Toyoda, H., Sumiyama, A., Oda, Y., Asayama, K., “Superconducting Proximity Effect inCu-Clad Nb Wires Doped with Fe, Co and Ni”, J. Phys. Soc. Jpn, 59(12), 4215–4218(1990) (Experimental, Phys. Prop., 10)

[1996Yav] Yavari, A.R., Drbohlav, O., “Mechanisms of Nanocrystallization of Fe- and Al-Based Amor-phous Precursors”, Mater. Sci. Forum, 225–227, 295–304 (1996) (Experimental, Phys.Prop., 40)

[1999Zhu] Zhu, J.H., Pike, L.M., Liu, C.T., Liaw, P.K., “Point Defects in Binary Laves Phase Alloys”,Acta Mater., 47(7), 2003–2018 (1999) (Experimental, Crys. Structure, 38)

[2000Tof] Toffolon, C., Servant, C., “Thermodynamic Assessment of the Fe-Nb System”, Calphad, 24(2), 97–112 (2000) (Phase Diagram, Phase Relations, Thermodyn., Calculation, Assessment,40)

[2000Wan] Wang, C.P., Liu, X.J., Ohnuma, I., Kainuma, P., Hao, S.M., Ishida, K., “Phase Equlibria inthe Cu-Fe-M- and Cu-Fe-Nb Systems”, J. Phase Equilib., 21(1), 54–62 (2000) (Phase Dia-gram, Phase Relations, Thermodyn., Calculation, Experimental, 22)

[2002Oka] Okamoto, H., “Fe-Nb (Iron-Niobium)”, J. Phase Equilib., 23(1), 112 (2002) (Phase Dia-gram, Review, 3)

[2002Rag] Raghavan, V., “Cu-Fe-Nb (Copper-Iron-Niobium)”, J. Phase Equilib., 23(3), 263–266(2002) (Phase Relations, Phase Diagram, Review, 5)

[2002Rom] van Rompaey, T., “Cu-Nb (Copper-Niobium)”, MSIT Binary Evaluation Program, in MSITWorkplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stutt-gart; Document ID: 20.12479.1.20, (2002) (Phase Diagram, Crys. Structure, Thermodyn.,Assessment, 16)

[2004Wan1] Wang, C.P., Liu, X.J., Ohnuma, I., Kainuma, P., Ishida, K., “Thermodynamic Database of thePhase Diagram in Cu-Fe Base Ternary Systems”, J. Phase Equilib., Diffus., 25(4), 320–328(2004) (Phase Diagram, Phase Relations, Thermodyn., Calculation, 40)

[2004Wan2] Wang, C.P., Liu, X.J., Takaku, Y., Ohnuma, I., Kainuma, P., Ishida, K., “Formation of Core-Type Macroscopic Morphologies in Cu-Fe Base Alloys with Liquid Miscibility Gap”,Metall. Mater. Trans. A., 35A(4), 1243-1253, (2004) (Experimental, Morphology, PhaseRelations, Phase Diagram, Thermodyn., 31)




Cu–Fe–Nb 11



Copper – Iron – Nickel

Nathalie Lebrun, Pierre Perrot, Mikhail Turchanin, An Serbruyns

Introduction

Cu-Fe based alloys are of interest in practical development of high strength, high electrical conductive Curich alloys and age-hardened Fe rich alloys. The Cu-Fe-Ni is an important system for producing compositematerials with high electrical and thermal conductivity and low thermal expansion. The combination ofthese properties makes the Cu-Fe-Ni an interesting candidate for the development of new data storage facil-ities [2006Dan]. Moreover Cu-Fe-Ni alloys present attractive hard magnetic properties. Their high conduc-tivity provides the necessary flexibility to fabricate soft magnets [2007Mon].The Cu-Fe-Ni ternary system has been widely studied using various experimental methods and thermody-namic optimization. Although the general shape of the liquidus was forecast by [1909Jae, 1910Vog], andthe miscibility gap in the γ phase was well known by [1935Dah], the first important experimental workon the phase equilibria was carried out by [1935Koe]. The first theoretical investigation on the miscibilitygap was carried out by [1942Hir]. Details on methods employed are indicated in Table 1.The thermodynamic properties of ternary Cu-Fe-Ni alloys were investigated by [1971Cou, 1977Vel,1978Kon, 1986Ric] for the γ phase and [1985Tse, 1997Fuj, 2002Mor] for liquid phase. In [1989Kum,1991Jia, 1996Ron, 2004Mor, 2006Mor] the thermodynamic properties of phases were calculated. Thesystem was thermodynamically assessed by [1985Chu, 1985Mos, 1985Spe, 1987Jan, 1996Ron,2001Ser1, 2004Wan1]. Numerous reviews are available in the literature [1949Jae, 1969Gue, 1975Cha,1979Cha, 1979Dri, 1990Gup, 2004Rag].

Binary Systems

The Cu-Fe, Cu-Ni and Fe-Ni phase diagrams are accepted from [2007Tur], [2002Leb] and [2007Kuz]respectively. In the Cu-Fe and Cu-Ni systems, metastable miscibility gaps in the liquid phase exist. How-ever no liquid separation occurs in the Cu-Ni system.

Solid Phases

Solid phases are presented in Table 2. [2001Ser2] has shown an evidence of a new phase inside the ternarysystem using electron micrographs on thin film. This phase would be of cubic symmetry (a = 1080 pm with108 atoms in the cell). The composition of this phase, around Cu13Fe20Ni67 is not clear since large discre-pancies are found in the composition between measurements and calculations. This phase seems to derivefrom the binary ordered FeNi3 phase and it forms during aging at 500°C. Thermodynamic calculationshowed that this phase dissolves about 5 at.% Cu whereas the experimental composition was found twicelarger [1970Gom2, 2001Ser2]. The thermal treatment used is doubtful since the sample was aged onlyfor 2 h at 500°C after one week at 1150°C. Further experimental results are needed to determine exactlythe composition range of this phase.The order-disorder temperature of γ’, FeNi3 is lowered by the addition of Cu [1951Che]. The Ni3Fe long-range order is destroyed at 9-10 at.% Cu [1970Gom1, 1970Gom2].The lattice parameter for the γ phase in alloys quenched from above 1050°C decreases as Ni contentincreases [1941Bra].


No invariant reaction exists inside the ternary system. A critical tie-line has been observed at 1220°C[1935Koe]. At this temperature the L+γ1+γ2 tie-line triangle degenerates to a straight line. The end pointsof this line are the critical point c and the upper point pc of the monovariant peritectic line pc-p3. The

Cu–Fe–Ni 1



composition of the γ phase at the point c is 37.5Cu-35.1Fe-27.4Ni (mass%) [1935Koe]. The composition ofthe liquid and the γ phase at the point pc is 67Cu-8.8Fe-24.2Ni (mass%) [1935Koe].


The binary peritectic located in the Cu rich region of the Cu-Fe binary system extends into the ternary sys-tem, terminating in a critical tie-line at 1220°C [1935Koe]. The experimental data of [1977Has] are fairly inagreement with the monovariant line estimated by [1935Koe]. The liquidus projection shown in Fig. 1 isreproduced from the review done by [1987Gup] and based on the experimental data of [1935Koe]. Goodagreement is observed with the accepted binary systems except the isothermal line at 1450°C in the Ni richcorner. To be in accordance with the Fe-Ni binary the end line at the Fe-Ni edge has been moved from 10 to19 at.% Fe.

Isothermal Sections

The basic features of the Cu-Fe-Ni ternary diagram were established by [1910Vog, 1929Che1, 1929Che2,1933Rol] where considerable disagreement was noticed regarding the extent of the solid miscibility gap inthe ternary system. [1935Koe] first developed a complete diagram of the ternary system. With improvementof purity of metals and experimental and calculation methods, the isothermal sections were later determinedin numerous experimental and thermodynamic works (see Table 1).There is no molten copper rich phase formed at 1000°C [1989Zou]. A wide miscibility gap region of the γphase was first established by [1910Vog]. Indeed, the high temperature iron (γ1 phase), which seems toundergo a spinodal unmixing during aging, gives rise to two coherent fcc phases: one enriched in copper(called γ3) and the other enriched in nickel (called γ2) with similar crystal structures. With the additionof Ni, the gap increases inside the ternary system over a wide temperature range. The γ phase gap extendsfrom the Cu-Fe into the ternary system with decreasing temperature.Phase equilibria proposed by [1929Che1, 1929Che2] agree broadly with the diagram proposed by[1910Vog] but disagree regarding the extent of the miscibility gap. [1990Gup] reviewed the shape of thisstable miscibility gap between 1050 and 600°C omitting, the Fe rich region, where the bcc (αFe) phasebecomes stable below 912°C. From the thermodynamic properties of the phases γ1 and γ3 as well the liquidphase, curves of the miscibility gap were calculated between 777 and 1477°C [1975Cha, 1985Chu,1985Mos, 1986Chu, 1987Hoc, 1987Jan, 1989Kum, 1996Ron, 2004Wan1]. These computed curves are ingood agreement with the experimental data of [1935Koe, 1977Has, 1990Gan, 1992Gan]. Nevertheless,the gaps calculated by [1985Chu, 1986Chu, 1987Hoc, 1987Jan] are systematically wider than the experi-mental data and tend to be larger at higher Ni content in both γ1 and γ3 equilibrium phases. Using diffusionmeasurements, [1996Ron, 2000Qin] obtained a better agreement between their experimental data and thecalculated curve of [1977Has] at 1050, 1000 and 800°C. The compositions of the Cu and Ni rich phasesobtained in the experimental work of [1993Lop] are in agreement with the boundary of the miscibilitygap calculated by [1977Has]. The miscibility gaps measured by [1941Bra, 1971Cou] tend to be smallercompared to the other experimental works cited above. Moreover the miscibility gap measured by[1941Bra] shows a peculiar bulge in the phase boundary near the Cu rich corner which was not observedin later experimental works, e.g. [1952Pal]. Due to more reliable experimental data, the miscibility gap mea-sured by [1977Has, 1990Gan, 1992Gan, 1996Ron, 2000Qin] is accepted in this assessment and is in agree-ment with the Cu-Fe binary edge. Results are reported in Fig. 2. The Darken stability function for ternaryalloys is discussed in [1985Sch1, 1986Chu] and calculated for the Cu-Fe-Ni alloys at 950°C. Gibbs energysurfaces were calculated and drawn by [1985Sch2] at 877°C for the γ phase and at 1577°C for the liquidphase. No stable miscibility gap has been reported in the liquid phase. The Cu-Fe binary system presentsa metastable liquid miscibility gap whose critical point is lowered by 80°C with the addition of 4 mass%Ni [2004Wan2].Large discrepancies have been observed between the calculations available in the literature regarding theextension of the L+γ region [1977Has, 1985Spe, 1986Chu] at 1300 and 1400°C. [1977Has] systematicallyfound a larger liquid region than [1985Spe]. Since no experimental data are available at that temperatures,no conclusion can be made.

2 Cu–Fe–Ni


Using their own experimental data and the available literature data on thermodynamic optimization,[1997Oht] computed an isothermal section at 1200°C. Results are of the same nature as the sections pre-viously computed at 1150, 1200 and 1250°C by [1985Chu, 1985Spe, 1986Chu, 1987Jan]. The L + γ regiondetermined by [1977Has] tends to be smaller with a lower maximum Ni composition. The calculatedboundary γ / γ+L by [1985Chu, 1985Spe, 1986Chu] differs from the experimental data of [1977Has]. Theircalculated curves are systematically at lower Ni content than the experimental data. On the contrary, calcu-lations done by [1987Jan] are in agreement with [1977Has]. Recently, [2004Wan1] also undertook a precisethermodynamic assessment which is in good agreement with the ones computed at 1150°C by [1985Chu,1986Chu, 1987Jan] and at 1200°C by [1997Oht]. The more recent computed data at 1200°C[2004Wan1] have been shown in Fig. 3. Slight modification (less than 2 at.%) has been done to match withthe composition of the boundaries along the Cu-Ni and Cu-Fe binaries.The isotherms at 850, 750, 650 and 550°C [1989Kum], at 600°C [1987Jan], at 450, 500 and 550°C[2001Ser2], at 500°C [1990Tay] are nearly of the same nature showing a miscibility gap of the γ phase.General agreement has been observed with the experimental data of [2000Qin]. Nevertheless, the three-phase field region (αFe)+γ1+γ2 is systematically larger in the calculated isothermal sections than in theexperimental ones. Moreover, the calculated miscibility gap is larger at Ni rich compositions. The experi-mental results of [2000Qin] at 800 and 600°C based on diffusion measurements have been preferred andare shown in Figs. 4 and 5. Slight modifications have been done regarding the solid miscibility curve tobe in agreement with the binary edge Cu-Fe. No experimental data are available at temperatures lower than600°C. Consequently the most recent reliable thermodynamic assessment [2001Ser1] has been chosen tak-ing into account the order-disorder transformation in the γ-(Cu,Ni)3Fe and the corresponding isothermalsection at 500°C is reported in Fig. 6.The phase boundaries involving (γFe)-(αFe) transformations are uncertain at very low temperaturesbecause of the strong hysteresis effects. A comparison with the accepted binary system Fe-Ni indicates thatthe (αFe) region shown in the isothermal section at 20°C available in the literature [1935Koe, 1941Bra]should extend to only about 1 at.% Ni. Due to the lack of data at this temperature regarding all the acceptedbinaries, no definitive isotherm can be drawn at 20°C. A metastable solid miscibility gap exists in the αphase [2004Wan1] and the computed curves are shown in Fig. 7.


Thirteen isopleths were constructed by [1935Koe]. The phase boundaries involving (γFe)-(αFe) transfor-mations are uncertain at lower temperatures because of strong hysteresis effects. Consequently results upto 600°C have not been retained in this assessment. [1952Pal] determined the solvus curves in isoplethsat 5 and 10 mass% Ni from 0 up to 5 mass% Fe showing a solubility limit of the γ phase which increaseswith temperature. [2000Qin] measured larger solubility limits and these more recent results have been pre-ferred. Moreover, Ni increases Cu solubility in molten iron [1952Pal]. Two vertical sections have beendepicted in Figs. 8 and 9. The boundaries involving (αFe) and the γ phases have been modified in the ver-tical sections to be in agreement with the more recent isothermal sections at 600 and 800°C as well as themeasured solid miscibility gap at 900, 1000 and 1100°C.

Thermodynamics

Thermodynamic activities of iron in γ solid Cu-Fe-Ni alloys were determined at 800-1100°C by [1978Kon]and those of copper and nickel in liquid Cu-Fe-Ni alloys were investigated in [1997Fuj] at 1350°C andxCu > 0.3. The activities demonstrate a positive deviation from the ideal behavior in the studied compositionrange and the deviation for copper increases with increasing iron content as shown in Fig. 10.Partial pressures of the components above liquid alloys were investigated in [1985Tse] by gas transportmethod for 5 ternary compositions at < 6.3 mass% Cu and 1520°C. The result are given in Table 3 as activ-ity coefficients of the components fMe. In [2004Mor] thermodynamic activities in liquid alloys were calcu-lated from the Gibbs energies of the mixture modelled in [2002Mor]. In [2006Mor] the results of thecalculations are compared with the experimental data of [1997Fuj] and a reasonable agreement wasobserved.

Cu–Fe–Ni 3



Mixing enthalpy of components in the γ phase was measured in [1971Cou] at 1050°C for nine ternaryalloys, most of which were nickel rich. These values were fitted with binary mixing enthalpies and integralmixing enthalpy ΔHM was described by an equation. Excess entropy values ΔSE for the three binary sys-tems were calculated from the mixing enthalpies and the observed phase equilibria at 850 and 1050°C.The final equations used to represent the binary and ternary values are given in Table 3. These data weretaken into account by [1985Mos, 1985Spe, 1987Jan, 2001Ser1] for the thermodynamic assessment of thesystem.The thermodynamic activities of the components in the Ni/Fe = 3 section were determined at 1247°C andxCu = 0.185-0.241 for the alloys consisting of the γ phase [1977Vel]. It has been shown that copper behavesalmost ideally with respect to the remaining components in the section; nickel and iron show negative andpositive deviations from the ideal solution behaviour.The thermodynamic activities of Fe in the γ phase have been determined at 800 to 1000°C, xFe < 0.5 andxCu < 0.5 [1978Kon]. The additions of Cu raise the aFe. The negative deviations from the ideal behaviorobserved in the Ni-Fe system change to the positive ones in the ternary system. A temperature increaseshifts the system closer to ideality. So there is a contradiction between the results of [1977Vel] and[1978Kon] regarding the definition of the copper influence on the activities of the components in the γphase. The advantage can by given to [1978Kon] because these data were obtained as a result of a moredetailed study. Moreover, positive deviations observed in the ternary γ phase are consistent with the phaseseparation that occurs with increasing Cu content in the phase. Isothermal thermodynamic activities of thecomponents in γ phase at 1000°°C are shown in Figs. 11a, 11b and 11c according to the calculations of[1996Ron]The results of [1978Kon] were taken into account in the thermodynamic assessments of the system[1985Spe, 1987Jan, 2001Ser1]. The assessment of [2001Ser1] is based on [1987Jan], with a particularattention given to the ordering in the phase (Cu,Ni)3Fe. It leads to a good agreement between the calculatedand experimental boundaries of the miscibility gap of the γ phase.The specific heat capacity of Cu-1.6Fe-9.9Ni (mass%) and Cu-0.8Fe-30.9Ni (mass%) alloys was measuredby [1986Ric] in the temperature interval from 59 to 946°C. Both alloys show nearly the same specific heatcapacity as copper. In Cu-1.6Fe-9.9Ni (mass%) alloy an additional contribution, due to the precipitation ofnickel-iron particles is found. It should be noted that investigated alloys contained up to 0.03 mass% C and0.81 mass% Mn.


The main experimental works on material properties are gathered in Table 4. The Cu-Fe-Ni alloys aremainly known for their magnetic properties [1938Bum1]. Permanent magnets are encountered in the com-position range 18-35 mass% Ni and 5-25 mass% Fe. Materials with high magnetic permeability and lowcoercitive force, are encountered around the composition 20 mass% Fe, 20 mass% Ni; more generally,materials with high permeability and low hysteresis losses lie in the domain 10-20 mass% Fe, 0-40 mass%Cu. These materials, known as Permalloys®, widely used in telecommunications are amongst the softestmagnetic materials known [2005Ost]. Their properties, dependent on the heat treatment [1950Jos,2005Ost], are attributed to the order-disorder transition observed around the composition Ni3Fe (at 517°Cfor pure Ni3Fe). A conductive wire such as Cu (20 µm of diameter) electroplated with Permalloy® presentsa giant magnetoimpedance [2003Li, 2003Nag, 2007Yi] which attracts much interest because of its potentialapplications in magnetic sensors.Addition of 5 mass% Cu to Ni3Fe alloys was shown to decrease by about 20°C the ordering ability of Ni3Fe[1951Che, 1952Jos]. The long range order parameter decreases from unity for Ni3Fe to 0.4 for(Ni0.75Fe0.25)0.95Cu0.05 [1970Gom1, 1970Gom2]. The alloy Cu52Ni38Fe10 annealed at 460°C during850 h gives superlattice reflections attributed to the formation of a Ni3Fe type ordered structure[1984Aal]. The Cu-Fe-Ni system presents a line of zero linear magnetostriction [1969Ash] going fromthe Cu-Ni side (55 mass% Cu) to the Fe-Ni side (18 mass% Fe) and a line of zero extraordinary Hall coeffi-cient going from the Cu-Ni side (50 mass% Cu) to the Fe-Ni side (12 mass% Fe).The addition of 1 to 2 mass% Fe was shown to improve the corrosion resistance: impingement attack,deposit attack, pitting, of copper alloys containing 10 mass% Ni [1951Bai]. Copper precipitation has been

4 Cu–Fe–Ni


widely investigated in alloys containing 10 or 30 mass% Ni added with less than 2 mass% Fe, [1983Ric,1984Bri, 1988Lug, 1989Wac] as well as in alloys in which Cu/Ni ~ 1 added with less than 8 at.% Fe, alloyscharacterized by a high corrosion resistance [1980Kat, 1993Lop, 1997Gud, 1997Kas]. There is a strongexperimental evidence [1987Wag, 1991Wol] that precipitates dissolve under heavy ions or neutrons irradia-tion. Different models of Cu precipitation during thermal ageing were presented in [2004Chr, 2004Sek,2006Zha] and compared with experimental observations.Copper is considered as a “tramp” element in steels and cannot be deleted easily, because it is less oxidablethan iron. It is the main cause of the “Hot Shortness Effect” [1966Sal], that is the appearance of surfacecracks or fissures which can reach tenth of millimeters and lead to the scrapping of a considerable quantityof steel. Nickel was shown to restrict the hot shortness effect. Copper may also be used instead of Ni toobtain Maraging steels after austenitizing at 1050°C, quenching at room temperature, then annealing at485°C [1968Wil]. Ni was shown to improve hot workability and hardenability of Cu maraging steels.It is possible to prepare, in the two-phase domain, composite alloys consisting of a high conductivity copperrich phase and a low expansivity Invar rich phase [1998Cot] provided that the ratio Fe/Ni be maintainednear by 64/36 (in mass%) because the Invar effect occurs only in a narrow composition range.

Miscellaneous

Diffusion in Cu-Fe-Ni alloys has been widely investigated [2006Dan] and fundamental atomistic informa-tion was extracted from measured interdiffusion coefficients by [2005Bel]. One of the first theoretical andexperimental studies was made by [1948Dan] who investigated the “uphill” (against the concentration gra-dient) diffusion presented by the Cu4FeNi3 alloy which is single-phase above 800°C and two-phase owingto the miscibility gap below that temperature. A similar investigation was carried out on the Cu10Fe3Ni7alloys by [1949Har] whose observations may be clearly interpreted as a proof of a spinodal decomposition.A Mössbauer investigation on the same alloy [1967Nag] shows a probably more complex decompositionmechanism, even in the spinodal region, but this result may be explained by uncertainties on the positionof the binodal and spinodal curves. Mössbauer seems to be an efficient tool [1970Ben] to follow thekinetics of the decomposition γ ↾ γ1 + γ2. Later investigation by [1979Lib] shows that spinodal decomposi-tion occurs during annealing so that two coherently linked tetragonal phases which differ in tetragonality(c/a < 1 and c/a > 1) are formed from the original γ phase. Elastic stress have a large effect on spinodaldecomposition. Small Angle Neutron Scattering has been successfully applied [1984Aal, 1986Poe1,1986Poe2, 1988Lyo] in order to put into evidence the spinodal decomposition and to have informationon the size and orientation of precipitates. Large Angle X-Ray Diffraction [2001Lyo, 2005Lyo] gives, onthe other hand, information on the distortion of the lattice created by the precipitates. The coherent spinodaltemperature was evaluated by [1993Lop] at 525±25°C for the Cu50Ni46Fe4 alloy, and at 625±25°C for theCu44Ni48Fe8 alloy. Similar results were also obtained from electrical resistivity measurements carried out by[1996Lop]. The coherent spinodal temperature was measured at 517±20°C (instead of 525±25°C) for theCu50Ni46Fe4 alloy and at 317±20°C for the Cu30Ni70 alloy. The coherent phase seems constituted of a meta-stable bcc (αCu) phase which precipitates into a stable fcc (Cu) phase upon reaching a critical size[1994Osa]. An alloy (20 mass% Cu, 20 mass% Ni) spinodally decomposed shows a dramatic improvementof the magnetoresistance [1994Che] when uniaxially deformed if the size scale of each line is small, of theorder of 1.5 nm.The density and surface tension of liquid Cu-Fe-Ni alloys have been measured over temperature interval1127-1577°C including the undercooled regime [2005Egr, 2006Bri1, 2006Bri2]. The investigation wasundertaken for two sections - CuxFe0.6(1–x)Ni0.4(1–x) and 20 at.% Cu. The results of investigation of densitypointed out on the positive deviations from ideality. The equation, describing excess molar volume ΔVE ofliquid Cu-Fe-Ni alloys is given in Table 3. The established values of surface tension are shown in Figs. 12and Fig. 13 along two sections. The surface tension was correctly predicted by the Butler equation from thethermodynamic potentials of the binary systems.Viscosity of ternary Cu-Fe-Ni metallic melts was estimated by [2002Wan] on the base of a new model,based on Seetherman’s viscosity model and Chou’s geometrical thermodynamic model. Predicted viscosityof liquid alloys along the section Fe/Ni = 1 is shown in Fig. 14.

Cu–Fe–Ni 5



Table 1. Investigations of the Cu-Fe-Ni Phase Relations, Structures and Thermodynamics


[1929Che1] Dilatometry, resistivity, microscopicexamination

< 1000°C, 34-35 mass% Ni, < 20 mass%Cu

[1929Che2] Dilatometry, microscopic examination < 900°C, < 25 mass% Cu, determination oftwo-phase equilibria

[1929Kus] Magnetic coercitivity and hardnessmeasurements

< 800°C, < 50 mass% Ni, determination ofphase relations

[1933Rol] Optical microscopy, copper solubility in Fe-Ni alloys

Cooled from the liquid state, < 35 mass%Ni, < 25 mass% Cu

[1935Dah] Electric conductivity, hardnessmeasurements

450-1050°C, < 65 mass% Ni, miscibilitygap in the γ phase

[1935Koe] Electric conductivity, magnetometricmeasurements

20, 600 and 800°C, the whole diagram,miscibility gap in the γ phase

[1937Leg] Electric conductivity, magnetometricmeasurements

< 1250°C, < 25 mass% Ni, < 25 mass% Cu

[1941Bra] XRD, crystal parameters measurements 750-1300°C, the whole diagram, tie linesinside the miscibility gap

[1951Che] Dilatometry, Curie temperaturemeasurement, electrical resistivity

< 500°C, FeNi3 + 5 mass% Cu, Cuinfluence on the ordering of Ni3Fe

[1952Pal] Micrographic observation, hardnessmeasurements

400-1000°C, < 10 % Ni, < 2 %Fe (mass%), Fe solubility in Cu-Ni alloys

[1971Cou] High temperature calorimetry, integralGibbs energy of a mixture

1050°C, γ phase < 33 at.% Cu, < 33 at.%Fe

[1977Has] Electron Microprobe Analysis (EMPA),diffusion couples

400-1300°C, phase equilibria in the wholecomposition range

[1977Vel] Knudsen effusion method 1247°C, γ phase, Ni/Fe = 3 section at 18.5to 24.1 at.% Cu

[1978Kon] EMF measurements, stabilized thoria assolid electrolyte

800-1000°C, < 50 at.% Cu, < 50 at.% Fe,iron activities

[1985Mos] EMF measurements, stabilized zirconia assolid electrolyte

900-1100°C, Cu-Ni and Fe-Ni binaryalloys.

[1985Tse] Vapor pressure measurements by a gastransport method

1520°C, liquid phase at < 6.3 mass% Cu,activities calculation.

[1986Ric] Heat capacity measurements by calorimetry < 1000°C, 1.6 % Fe, 9.9 % Ni and 0.8 %Fe, 30.9 % Ni (in mass%)

[1990Gan,1992Gan]

EMPA, hardness measurements, diffusioncouples

900°C, < 50 at.% Ni, miscibility gapdetermination

(continued)

6 Cu–Fe–Ni



[1996Ron] Optical microscopy, SEM, EMPA,diffusion coefficients

1000°C, < 50 at.% Ni, diffusion coupleswith tracers

[1997Fuj] Mass spectrometry, double Knudsen cell 1350°C, < 50 at.% Ni, activitiesmeasurements in liquid phase

[1997Oht] Energy Dispersive X-ray analysis (EDX) 1100-1300°C, < 20 mass% Ni, solid-liquiddiffusion couples

[2000Qin] EPMA, ternary diffusion couples 600-1050°C, < 60 at.% Ni

[2000Uga] Optical microscopy, SEM, EMPA,Kirkendall effect

1100°C, < 70 at.% Ni, diffusion couples

[2001Ser2] TEM, high resolution TEM Cu80Fe3Ni17 annealed at 500°C

[2004Wan2] Optical microscopy CuFe + 4 mass% Ni, 1800°C then cooled





Comments/References

(αFe) (Ferrite)< 912

(δFe)1538 - 1394

cI2Im�3mW

a = 286.65

a = 293.15

pure Fe at 20°C [Mas2, V-C2]. Dissolves 4.6 at.% Ni at 495°C[2007Kuz]. Dissolves 1.4 at.% Cu at 847°C [2007Tur]

at 1394°C. Dissolves 3.8 at.% Ni at 1517°C [2007Kuz].Dissolves 5.6 at.% Cu at 1487°C [2007Tur]

(εFe) hP2P63/mmcMg

a = 246.8c = 396.0

at 25°C, > 13 GPa [Mas2]

γ, (γFe,Ni,Cu)(Austenite)

γ1, (γFe)a)

< 1394 – 912γ3, (Ni)

a)

< 1455γ2, (Cu)< 1084.62

cF4Fm�3mCu

a = 364.67

a = 352.40

a = 361.46

complete solubility range

at 915°C [V-C2, Mas2]; dissolves 14.2at.% Cu at 1425°C [2007Tur]pure Ni at 25°C [Mas2]

pure Cu at 25°C [Mas2]melting point [2002Leb]

(continued)

Cu–Fe–Ni 7






Comments/References

γ’, FeNi3< 517

cP4Pm�3mAuCu3-I

a = 355.23 63 to 85 at.% Ni [2007Kuz]

γ”, FeNimetastable

tP4P4/mmmAuCu

a = 357.9 [V-C2]metastable ordering temperature320°C at 51.2 at.% Ni [2007Kuz]

a)Below critical temperature: γ1 - paramagnetic, Fe enriched, γ2 - ferromagnetic, Ni enriched


Phase TemperatureRange[°C]


Comments

Liquid 1520 fCu = 6.3, fFe = 0.66, fNi = 0.35;fCu = 3.2, fFe = 0.54, fNi = 0.45;fCu = 2.42, fFe = 0.99, fNi = 0.84;fCu = 2.09, fFe = 1.82, fNi = 1.67;fCu = 2.86, fFe = 1.05, fNi = 0.97;

Cu 0.91, Ni 24.50, Fe 74.59 mass%;Cu 0.92, Ni 74.48, Fe 24.60 mass%;Cu 4.82, Ni 23.93, Fe 71.25 mass%;Cu 6.29, Ni 63.91, Fe 29.80 mass%;Cu 5.17, Ni 46.88, Fe 47.95 mass%;[1985Tse]

γ phase 1050 ΔHM = xCuxNi (3.586xCu+12.937xNi)++ xNixFe (–10.826 xNi–2.021 xFe) ++ xFexCu (54.124xFe+42.288xCu) ++ xCuxNi xFe(–8.644) kJ·mol–1;ΔSE = xCuxNi (–4.89xCu–2.04xNi) ++ xNixFe (–13.10 xNi+1.93 xFe) ++ 12.47xFexCu–24.94 xCuxNi xFeJ· (mol·K)–1

[1971Cou]

Liquid 1500 ΔVE = 0.6xFexCu – 0.85xCuxNi ++ 11.5 xCuxNi xFe cm

3·mol–1[2006Bri1, 2006Bri2]

Table 4. Investigations of the Cu-Fe-Ni Materials Properties


[1934Dah] Magnetic hysteresis, electrical conductivitymeasurements

Laminated alloys, precipitationhardening,< 20 % Cu, < 80 % Fe (mass%)

[1937Neu] Magnetization, remanence, coercitivitymeasurements

900-1050°C, < 50 mass% Ni plasticallydeformed

(continued)

8 Cu–Fe–Ni



[1938Bum1] Permeability, coercitive force, conductivitymeasurements

600-800°C, the whole diagram

[1938Bum2] XRD, permeability, coercitive force,conductivity measurements

600-800°C, 20-30 mass% Ni,12 mass% Al

[1938Bum3] XRD, permeability, conductivitymeasurements

< 1000°C, < 13 mass% Cu

[1938Dan] Coercitive force, remanence, hardnessdepending on heat treatment

< 1200°C, < 25 mass% Cu,< 25 mass% Ni

[1941Lif] Permeability, coercitive force, conductivitymeasurements

500-1100°C, 20 mass% Fe, 20 mass%Ni, various heat treatment

[1948Dan] XRD, diffusion measurements bymicrophotometry

Cu4FeNi3, 500-800°C, Uphill diffusioninvestigation

[1949Har] XRD 550-800°C, Cu10Fe3Ni7,decomposition mechanisminvestigation

[1950Jos] Dilatometry, thermomagnetic analysis < 1000°C, 74.5 mass% Ni, 5.6 mass%Cu, heat treated samples

[1951Bai] Corrosion resistance, hardness, magneticpermeability

< 1000°C, < 30 % Ni, < 2 %Fe (mass%), various heat treatments

[1952Smi] Emissivity of liquid alloys 1535°C, the whole composition range

[1957Rav] Thermomagnetic, electrical conductivitymeasurements

< 1000°C, Ni3Fe + 5 mass% Cu, order-disorder transition in Ni3Fe.

[1966Sal] Micrographic analysis, Electron MicroprobeAnalysis (EMPA)

900-1250°C, < 15 % Cu, < 15 %Ni (mass%), hot shortness effect

[1967Nag] Mössbauer 500-800°C, Cu10Fe3Ni7 decompositionmechanism investigation

[1968Wil] Optical microscopy, hardness measurements < 1250°C, < 8% Cu, < 3% Ni (mass%),hot rolled, austenitized

[1969Ash] Magnetostriction, Hall coefficient, resistivitymeasurements

20 K, > 60 mass% Ni

[1970Ben] Mössbauer, EMPA 25°C, < 53 at.% Ni, < 8 at.% Fe

[1970Gom1][1970Gom2]

Neutron diffraction Ni3Fe + Cu (< 6 at.% Cu), long rangeorder parameter

[1971Wat] Density measurement by the maximumbubble pressure method

1650°C, Cu-Fe and Cu-Ni binary liquidmixtures

[1971Win] Mössbauer 27°C, < 10 at.% Cu, < 1 at.% Fe

[1975Dob] XRD, diffuse diffraction rings analysis 20 mass% Cu, 20 mass% Fe, 1050°C,quenched, then annealed at 650°C

(continued)

Cu–Fe–Ni 9




[1977Dah] Yield strength measurements 550-800°C, < 15 at.% Fe, 35 to 52 at.%Ni, age hardened alloys

[1977Kok] Superparamagnetism 77-573 K, 6 mass% Fe, 14 mass% Ni,early stages decomposition

[1979Lib] Saturation magnetization in a field ofcompression stress

20 % Fe, 20 % Ni and 30 % Fe, 30 %Ni (mass%), annealed 500-700°C

[1979Ric,1980Ric]

Dilatometry, electron microscopy < 900°C, 2 mass% Cu, < 7 mass% Ni,Austenite ⇌ Ferrite transformation

[1980Buh] Optical micrography, EPMA < 400°C, 1.5 Cu, 10 % Ni (mass%),creep-rupture embrittlement

[1980Kat] XRD, XPS, SEM, ESCA 9.4% Ni, 1.7% Fe (mass%), corrosionin aqueous NaCl

[1981Ric] High speed dilatometry, optical microscopy,TEM

< 900°C, 2 mass% Cu, 2 mass% Ni,Austenite ⇌ Ferrite transformation

[1982Kok1,1982Kok2]

XRD, magnetic susceptibility measurements,superparamagnetism

9 to 22 at.% Fe, 17 to 40 at.% Ni,annealed at 610°C

[1983Mec] Magnetic susceptibility by the Faraday’smethod

< 1000°C, pure Cu with 0.1 mass% Niand < 0.3 mass% Fe as impurities

[1983Ric] Mössbauer, superparamagnetism, SEM < 850°C, (Cu0.7Ni0.3)0.98Fe0.02 and(Cu0.9Ni0.1)0.98Fe0.02

[1984Aal] Neutrons scattering 400-800°C, 37 to 40 at.% Ni, 5 to 13 at.% Fe, superlattice observation

[1984Bri] TEM, TTT (Temperature-Time-Transformation) curves

< 800°C, 1.3 % Fe, 9.5 % Ni and 0.5 %Fe, 30 % Ni (in mass%)

[1984Szy] Coercive force, remanent induction 750°C, 20 at.% Fe, 20 at.% Ni, uniaxialcompression stress applied

[1986Cha1] Electrical resistivity < 600°C, 11-16 at.% Cu, < 64 at.% Fe,samples annealed at 500°C

[1986Cha2] Magnetic moment, spontaneousmagnetization, Curie point

< 600°C, 11-16 at.% Cu, < 64 at.% Fe,samples annealed at 500°C

[1986Kok] Magnetoresistance 600-1050°C, 10-25 at.% Fe, 7-15 at.%Ni

[1986Poe1,1986Poe2]

Small Angle Neutrons Scattering (SANS) < 750°C, 47 at.% Ni, < 8 at.% Fe,kinetics of short range ordering

[1986Sch,1987Sch]

SANS, TEM. Decomposition underirradiation (electrons, heavy ions)

500°C, 46 at.% Ni, 4.0 at.% Fe and47.8 at.% Ni, 8.0 at.% Fe

[1987Lo] SEM, hardness measurements 10 mass% Ni, 8 mass% Fe rapidlysolidified, then annealed 550-800°C

(continued)

10 Cu–Fe–Ni



[1987Lyo,1988Lyo]

Anomalous small-angle X-ray scattering Cu0.43Ni0.42Fe0.15 aged at 500°C.Partial factor structure determination

[1987Wag] SANS, decomposition under irradiation(electrons)

500°C, 46 to 48 at.% Ni, < 8 at.% Fe

[1988Lug] TEM, creep rupture tests 400°C, 10 mass% Ni, 1.5 mass% Fe

[1989Wac] SEM, coercitive force, saturation polarization 400-600°C, 10 mass% Ni, 0.8 to 1.3mass% Fe, precipitation kinetics

[1989Zou] SEM, diffusion couple experiment 930-1000°C, Cu transport in austenite

[1991Oth] TEM, High Resolution Electron Microscopy(HTEM)

1.28% Cu, 1.43% Ni (mass%) aged at550°C. Cu precipitates

[1991Sim] Anomalous small-angle X-ray scattering < 25 at.% Fe, 2-41 at.% Ni, annealed at500°C. Structure factors determination

[1991Wol] SEM Cu45Fe8Ni47. Structure underirradiation

[1992Lyo] SEM, anomalous small-angle X-rayscattering

Cu44Fe14Ni42 aged at 550°C. Partialfactor structure determination

[1993Lop] Atom Probe Field Ion Microscopy (AP-FIM) Cu50Ni46Fe4 and Cu44Fe8Ni48,annealed 400-650°C. Coherentspinodal

[1994Che] Magnetoresistance 4.2 and 298 K, 20 mass% Ni, 20 mass%Fe, spinodally decomposed

[1994Osa] TEM, SANS, AP-FIM, precipitationmechanism

500°C, 1.4 at.% Cu, 0.27 at.% Ni,0.25 at.% Mn

[1995Nav] Mössbauer, saturation magnetization Cu70Fe30 and Cu70Fe25Ni5, isochronalthermal treatment 400-900°C

[1996Lop] Electrical resistivity, spinodal decomposition Cu30Ni70 and Cu50Fe4Ni46, annealed150-550°C

[1997Gud] TEM, hardness and yield strengthmeasurements

Cu44Fe8Ni46, annealed 450-650°C

[1997Kas] Tomographic atom probe Cu45Fe8Ni47, annealed 500°C

[1998Cot] Expansivity, electrical conductivity Cu-Invar composites (Fe/Ni=64/36 inmass%)

[2001Li] XRD, optical microscopy, SEM, TEM,hardness measurements

Cu45Fe25Ni30, aged at 600-900°C andcold-rolled

[2001Lyo,2005Lyo]

SEM, small angles and large angles X-rayscattering

Cu42Fe1-6Ni42, aged at 500-550°C,surface energies calculations

[2002Kin] Atomic volume, magnetization, Mössbauer < 1000°C, 50 and 60 at.% Fe preparedby mechanical alloying

(continued)

Cu–Fe–Ni 11




[2003Li] SEM, magnetoresistance Cu wires (20µm diameter) electro-plated with Permalloy (Ni80Fe20)

[2003Nag] High angle XRD, X-ray reflectivity,magnetoresistance

Cu-NiFe composites. Paramagneticinterfacial layer investigation

[2005Egr,2006Bri1,2006Bri2]

Image analyses of magnetically levitatedsamples. Surface tension and density of theliquid

1065-1577°C, Cu0.2(Fe1–xNix)0.8 andCux(Fe0.4Ni0.6)1–x (0 < x < 1)

[2007Mon] XRD, DSC, magnetization Nanocrystalline alloys prepared bymechanical alloying

[2007Yi] Atomic force microscopy, magnetization Cu sputtered on Fe21Ni79 substrate.Giant magnetoimpedance effect

Fig. 1. Cu-Fe-Ni. Liquidus surface projection

12 Cu–Fe–Ni


Fig. 2. Cu-Fe-Ni. Solid miscibility gap of the γ phase

Cu–Fe–Ni 13



Fig. 3. Cu-Fe-Ni. Isothermal section at 1200°C

14 Cu–Fe–Ni



Cu–Fe–Ni 15




16 Cu–Fe–Ni



Cu–Fe–Ni 17



Fig. 7. Cu-Fe-Ni. Metastable solid miscibility gap of the α phase

18 Cu–Fe–Ni


Fig. 8. Cu-Fe-Ni. Vertical section Ni-Fe10Cu90 (in mass%), plotted in at.%

Cu–Fe–Ni 19



Fig. 9. Cu-Fe-Ni. Vertical section Ni-Fe50Cu50 (in mass%), plotted in at.%

20 Cu–Fe–Ni


Fig. 10. Cu-Fe-Ni. Isoactivity lines of copper in liquid alloys at 1350°C

Cu–Fe–Ni 21



Fig. 11a. Cu-Fe-Ni. Isoactivity lines of copper in the γ phase at 1000°C

22 Cu–Fe–Ni


Fig. 11b. Cu-Fe-Ni. Isoactivity lines of nickel in the γ phase at 1000°C

Cu–Fe–Ni 23



Fig. 11c. Cu-Fe-Ni. Isoactivity lines of iron in the γ phase at 1000°C

24 Cu–Fe–Ni


Fig. 12. Cu-Fe-Ni. Surface tension of liquid alloys along the CuxFe0.6(1–x)Ni0.4(1–x) section at 1527°C

Cu–Fe–Ni 25



Fig. 13. Cu-Fe-Ni. Surface tension of liquid alloys along the section at 20 at.% Cu and 1527°C

26 Cu–Fe–Ni


Fig. 14. Cu-Fe-Ni. Predicted viscosity of liquid alloys along the section Fe/Ni = 1 at 1550 and 1650°C

Cu–Fe–Ni 27



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30 Cu–Fe–Ni


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[1983Ric] Richter, F., Pepperhoff, W., “Mössbauer Investigations on Magnetically Ordered Precipi-tates in Cu-Ni-Fe” (in German), Z. Metallkd., 74(8), 500–503 (1983) (Experimental, Magn.Prop., Electronic Structure, 18)

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[1985Chu] Chuang, Y.Y., Schmid, R., Chang, Y.A., “Calculation of the Equilibrium Phase Diagramsand the Spinodally Decomposed Structures of the Cu-Fe-Ni System”, Acta Metall., 33(8),1369–1380 (1985) (Calculation, Phase Diagram, Phase Relations, Thermodyn., Assess-ment, 24)

[1985Mos] Moser, Z., Zakulski, W., Spencer, P., Hack, K., “Thermodynamic Investigations of SolidCopper-Nickel and Iron-Nickel Alloys and Calculation of the Solid State Miscibility Gapin the Copper-Iron-Nickel System”, Calphad, 9(3), 257–269 (1985) (Phase Diagram, PhaseRelations, Experimental, Thermodyn., Calculation, 44)

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32 Cu–Fe–Ni


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Cu–Fe–Ni 33



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34 Cu–Fe–Ni


[2004Chr] Christien, F., Barbu, A., “Modelling of Copper Precipitation in Iron during Thermal Agingand Irradiation”, J. Nucl. Mater., 324(2-3), 90–96 (2004) (Calculation, Electronic Structure,Morphology, Phase Relations, 15)

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36 Cu–Fe–Ni


Copper – Iron – Oxygen

Pierre Perrot, Sander Arnout, Jan Vrestal

Introduction

The Cu-Fe-O system is of a great technological interest because of its importance in the metal-slag equili-bria encountered in the copper metallurgy. The easy formation of CuFeO2 in the slags causes the loss ofcopper [1981Zin1], so, this system is of prime importance in understanding the roasting, smelting, refiningand leaching process in the extraction of copper.The first credible assemblage of phases in the ternary system was proposed by [1964Yun] and comprehensivereview may be found in [1989Rag]. The main experimental works on the phase equilibria in the Cu-Fe-Osystem are gathered in Table 1.

Binary Systems

The Cu-Fe system is accepted from [2007Tur]. The Fe-O and Cu-O systems are accepted from [Mas2]. TheFe-O diagram comes from the fundamental work of [1945Dar, 1946Dar] and has been carefully assessed by[1991Sun, 1995Kow]. The Cu-O diagram has been recently assessed by [2003Hal, 2004Cla].

Solid Phases

The solid phases are presented in Table 2. τ,CuFeO2 is paramagnetic at room temperature with a Neel tem-perature of 14 K [1979Dim, 2004ElA]. Copper ferrite CuFe2O4 which is ferrimagnetic at room temperaturepresents a Curie temperature given at 507°C [1988Nai], 490°C [1988Upa] or 441°C [1988San]. It under-goes a cubic-tetragonal transformation which has been reported as low as 360°C [1959Ohn, 1969Mex]and as high as 960°C. The most probable transition temperature is 760°C [1961Yun]. CuFe2O4 developsa defect structure above 960°C, thoroughly investigated by [1972Tre], which may explain in part thereported variation in its transformation temperature. As pointed out by [1989Tan], pure stoichiometricCuFe2O4 does not exist and the mixture CuO+Fe2O3 gives Cu1–ηFe2+ηO4–δ with CuO in excess. The hightemperature modification, βCuFe2O4 give a spinel solid solution which is in the whole domain range above975°C. The reduction of CuFe2O4 by hydrogen [1984Qia, 1992Xu] begins at 215°C to form Cu and Fe withthe formation of a spinel solid solution as an intermediate phase.The lattice parameters of the Fe3O4-CuFe2O4 solid solutions, shown in Fig. 1 presents a maximum for thecomposition CuFe5O8 [1981Sap, 1987Lis], which is explained by the presence of monovalent copper in thesolid solution. The maximum content of Cu+ occurs for the composition CuFe5O8 [1969Haa, 1969Len,1981Dhu, 1981Sap]. However, a large uncertainty remains [1981Zin3, 1985Han] about the value of theratio Cu+/Cu2+ in the solid solution. The CuFe5O8 spinel ferrite has the structure of LiFe5O8. It is a disor-dered spinel with an inversion degree of 0.56 [1970Cer], which means that 56% of Cu+ ions are in octahe-dral sites. It is ferromagnetic with a Curie temperature reported at 344°C [1962The]. A more probable Curietemperature was measured at 470°C after slow cooling and at 445°C after quenching [1968Len]. The spinelsolid solution has been shown as nonstoichiometric [1968Ber, 1981Zin2] in air, which is, according to[1983Zin], is a characteristic of spinels with multivalent cations. CuFe5O8 around 1300°C can absorb oxy-gen up to CuFe5O8.08. Above 977°C, CuFe5O8 belongs to the continuous solid solution between Fe3O4 andCuFe2O4. Below that temperature, CuFe5O8 undergoes a eutectoidal transformation and decomposes intoCuFeO2 + 2Fe2O3 [1969Yam, 1971Sim].The low temperature modification αCuFe2O4 presents tetragonal structure with low c/a which decreaseswhen Fe3O4 enters in solid solution, as shown in Fig. 2. When αCuFe2O4 has dissolved about 7 mol%Fe3O4, c/a = 1, and the αCuFe2O4 phase takes the spinel structure of the β phase [1970Sch].The τ,CuFeO2 rhombohedral ferrite, known as delafossite, presents an incongruent melting at 1180°C[1996Zha2, 1997Zha2]. However, it must be pointed out that its stability domain is strongly dependent

Cu–Fe–O 1



on the surrounding atmosphere. A first evaluation from [1962The] shows its decomposition intoCuFe5O8 + Cu2O at low temperatures. Actually, more precise investigations of [1966Bui, 1966Gad,1981Zin1] show that CuFeO2 is stable between 1015°C and 1090°C under air atmosphere, a result in agree-ment with the observation of [2003Yan] which does not detect the formation of CuFeO2 by reactingCuO + Fe2O3 at 960°C under air. At 1090°C it presents incongruent melting into spinel + liquid. At1100°C, the spinel has the composition Cu0.63Fe2.37O4 and, at 1240°C, the composition Cu0.5Fe2.5O4

[1981Zin1]. Below 1015°C it decomposes into CuO + spinel. Under high oxygen pressure CuFeO2 oxidizesinto CuO + CuFe2O4; under low oxygen pressure (for instance pure Ar or CO2), CuFeO2 reduces intoCu + Fe3O4 [1996Zha2, 1997Zha2]. At 1000°C, CuFeO2 reduces into Cu + (Fe304)0.7(Cu0.5Fe2.5O4)0.3when pO2 < 0.06 Pa [1969Zal2]. A thorough investigation [1997Zha3] shows that the τ phase is non stoi-chiometric and may be represented by CuFeO2+δ with –0.11 < δ < +0.08 depending on the oxygen pressure.In stoichiometric CuFeO2, Cu and Fe are in the oxidation state +1 and +3, respectively [2000Suk].Wuestite Fe1–xO in equilibrium with metallic Fe may dissolve a small amount of Cu (0.35 at.% at 1000°C[1999Kat]. The copper solubility in wuestite increases slightly with the oxygen pressure. Fe1–xO in equili-brium with Fe3O4 may dissolve 0.34 at.% Cu at 800°C and 0.77 at.% Cu at 1000°C. Cu2O may dissolve upto 3 mol% CuFeO2 at 1000°C. The copper oxide in equilibrium with CuFeO2 has the compositionCu63Fe2O35 [1964Sch]. Ball milling was shown to enhance considerably the solubility of Fe2O3 in CuO[1998Ste, 2001Ste] and the solubility of CuO in Fe2O3 [2005Jia]. Small particles of CuFe2O4 form afterannealing at 575°C the metastable solid solution obtained by mechanical alloying.

Quasibinary Systems

Quasibinary system CuO-Fe2O3 has been investigated by [1966Gad, 1966Yam, 1988Zin]. These diagramscontradict the accepted oxygen pressure at equilibria Cu-Cu2O and Cu2O-CuO. Indeed CuO decomposesinto Cu2O at 1031°C in air. [1966Gad, 1966Yam, 1988Zin] give for this process 1050, 1080, and 1020°C, respectively. The difference seems not very big, but the main discrepancies lie in the nature of the phasesin equilibrium with τ,CuFeO2. The Fig. 3 presents quasibinary section CuO-FeO1.5 which agree with allexperimental observations. Under air, CuO is reduced into Cu2O at 1031°C, Fe2O3 into Fe3O4 at 1380°Cand the stability domain of CuFeO2 lies between 1015 and 1085°C. Delafossite CuFeO2 may be in equili-brium with CuO, Cu2O, a liquid phase and a spinel solid solution.


In the solid state, the 4 phases Cu, Fe2O3, Fe3O4 and CuFe2O4 present an invariant equilibrium. Cu + Fe2O3

are stable together below 675 ± 25°C whereas Fe3O4 + CuFeO2 are stable together above that temperature.A second invariant equilibrium in the solid state is observed at 977°C between the 4 phases CuFeO2, Fe2O3,Sp1,Fe3O4 and Sp2,CuFe2O4. Below 977°C, CuFeO2 and Fe2O3 are stable together and the two spinelphases present only a partial miscibility; above that temperature, the two spinel phases present a completemiscibility and the two phases CuFeO2 and Fe2O3 are no more stable together. 977°C is the temperatureunder which the spinel CuFe5O8 undergoes a eutectoid decomposition according to:CuFe5O8 ⇌ CuFeO2 + Fe2O3.


A liquidus surface was proposed by [1989Rag] on the basis of the works of [1966Gad, 1976Lur]. This dia-gram does not take into account the existence of an equilibrium between CuFeO2 and the liquid phase andthe tie lines at the liquid metal - slag equilibria are not realistic [1986Acu]. On the other hand, this figuresrests on very few experimental results so that it has not been reproduced in this report. The liquidus surfaceat 1200 and 1500°C in the copper rich corner is reproduced in Fig. 4. At 1200°C, the oxygen solubility inCu-Fe liquid alloys presents a minimum at ~1 at.% Fe in Cu [1973Kul]. At 1500°C, it presents a minimumof 0.03 mass% O around a composition of 2 mol% Fe [1989Ois].

2 Cu–Fe–O


Isothermal Sections

The phase equilibria in the solid state at 650°C, investigated by [1964Yun], presented in Fig. 5, are charac-terized by an equilibrium between metallic copper and the oxides FeO, Fe3O4, Fe2O3 and τ,CuFeO2. No tie-line appears between the two spinel phases (Sp1, Fe3O4) and (Sp2, βCuFe2O4), neither between Fe3O4 andτ,CuFeO2. Above 675 ± 25°C, Cu and Fe2O3 are no more stable together and the following transformationis observed: Cu + 2 Fe2O3 ⇌ CuFeO2 + Fe3O4. The phase equilibria in the solid state at 800°C, investigatedby [1961Yun, 1964Kul], presented in Fig. 6, show the presence of a stable equilibrium between delafossiteand magnetite, without any tie line between the spinels (Sp1,Fe3O4) and (Sp2, βCuFe2O4). The oxidation ofa two-phase alloy (Fe50Cu50) forms delafossite [1979Nan]. Although [1966Bui, 1972Ono, 1979Dri] see nomiscibility between Fe3O4 and CuFe2O4 below 1000°C, actually, [1966Yam] shows notable solubilitywhich is more coherent with the complete solubility observed above 977°C, and confirmed by[1967Yam] at 1080°C. The Fig. 7 shows the phase relations at 927°C, mainly from [1976Fre]. Fe3O4

may dissolve up to 30 mol% CuFe2O4 to give the spinel Sp1 solid solution CuxFe3–xO4 (x < 0.3) whichis in equilibrium, in the oxidized side, with Fe2O3. In the reduced side, Sp1, CuxFe3–xO4 is in equilibriumwith the metallic Cu for x < 0.13 and with τ,CuFeO2 for 0.13 < x < 0.30. At 927°C, the oxygen pressureat equilibrium in the triangle Cu-CuFeO2-Cu0.13Fe2.87O4 is pO2 = 10–2.49 Pa; the oxygen pressure at equili-brium in the triangle CuFeO2-Fe2O3-Cu0.3Fe2.7O4 is pO2 = 10–0.47 Pa; at the same temperature, Fe2O3

is in equilibrium with the spinel solid solution Cu0.13Fe2.87O4 under an oxygen pressure of 10–1.42 Pa.At 1000°C, the miscibility gap between magnetite and copper ferrite has disappeared as observed by[1964Sch] and definitively confirmed by [1970Sch, 1977Jac]. Figure 8 shows the phase equilibria in thesolid state at 1000°C. The τ,CuFeO2-Fe2O3 tie line is no more present. In the oxidized side, the spinel solidsolution CuxFe3–xO4 (0 < x < 1) is in equilibrium with Fe2O3; in the reduced side, it is in equilibrium withCu (x < 0.09) or with CuFeO2 (0.09 < x < 1) [1964Sch]. Above 1085°C, the assemblage (Cu2O + CuFeO2)presents a eutectic and a liquid domain spreads across the phase diagram as shown in Fig. 9 which presentisothermal sections at 1200°C [1986Acu]. The iron distribution between metal and slag, investigated by[1976Bur] at 1300°C shows a copper enrichment of the metallic phase, in fair agreement with the higheraffinity of oxygen for iron [1976Lur]. The slag-metal equilibria at 1500°C is shown in Fig. 10. The oxygenpressures for each metal-oxide tie line is also given [1989Ois]. As shown by [2000Ris], the ferrite ceramicsCuxFe3–xO4 may be prepared by solid reactions between CuO and Fe2O3 or Fe3O4 at high temperatures,which is in agreement with the phase equilibria observed above 1000°C.


The solubility of oxygen in liquid Cu which is of ~14 mass% under 0.1 MPa of oxygen pressure at 1200°Cdecreases quickly when Fe content increases up to a minimum of 22 ppm (0.0020 mass%) for 1.1 at.% Fe at1200°C in the alloy [1973Kul]. The vertical section from [2003Kat] is shown in Fig. 11. Under 977°C, thespinel solid solution CuFe5O8 decomposes into CuFeO2 + Fe2O3. The three-phase domains are character-ized by an oxygen pressure depending on the temperature:The oxygen pressure in the triangle Sp1 (Fe3O4 rich)-Fe2O3-CuFeO2 is given by:log10 (pO2 / Pa) = 23.46 – (29068 / T)The oxygen pressure in the triangle Sp2 (CuFe2O4 rich)-Fe2O3-CuFeO2 is given by:log10 (pO2 / Pa) = – 14.22 + (18060 / T)Both curves meet at 1250 K (977°C). Above that temperature, the spinel solid solution Fe3O4-CuFe2O4 isstable in the whole concentration range.

Thermodynamics

The interaction coefficients of oxygen in liquid alloys, defined, in the iron rich alloys, as ε(Cu)O = ∂ lnγO /∂ xCu where γO = (xO in pure Fe)/ (xO in alloy), has been evaluated as ε(Cu)O = –2.63 at 1600°C by[1966Sch] for xCu < 0.13 (15 mass% Cu). The interaction coefficients in the same conditions definedas e(Cu)O = ∂ log10 fO /∂ (mass% Cu), where fO = (mass% O in pure Fe) / (mass% O in alloy), has been eval-uated as e(Cu)O = –0.0095 at 1600°C. This value may be considered to be in fair agreement with the valuee(Cu)O = –0.016 from [1970Tan2] at 1550°C and –0.013 accepted by [1970Fis, 1971Fis, 1974Sig2], at

Cu–Fe–O 3



1600°C. In the copper rich alloys, the interaction coefficient e(Fe)O was determined [1970Fis, 1971Fis] ase(Fe)O = –0.27 (0 to 3 mass% Fe) and by [1970Tan1] as e(Fe)O = – 0.226 at 1550°C. Other interaction para-meters are available in the literature. Values of ε(Fe)O between –300 and –500 around 1350°C, are given by[1973Bis, 1976Ois]. Although confirmed by [1973Kul], and accepted by [1974Sig1, 1983Ani], they are fartoo negative to be seriously taken into consideration. It is probable that the huge oxygen solubility in liquidcopper (nearly 40 at.% at 1200-1400°C) and its quick decreases with iron content explains the discrepanciesbetween the various oxygen activities determination. A more credible expression is proposed by [1980Chu]:ε(Fe)O= 4.4 – (8900/T)The integral Gibbs energy of mixing for reaction O2 (gas) ⇌ 2 [O] (in liquid alloys) at 1600°C isΔG{O} = –248 kJ·mol–1, for pure iron. It decreases with the copper content down to a minimum of–260 kJ·mol–1 for 40 mass% Cu, then increases up to – 108 kJ·mol–1 for pure Cu [1970Fis, 1971Fis].The minimum value ΔGo = – 260 kJ·mol–1 is to be compared to the value of –268.8 kJ·mol–1, determinedin [1971Tan] and to the values shown in Table 3.Activities of Fe3O4 in the solid solution Fe3O4-CuFe2O4 system show large negative deviation fromRaoult`s law at 1000°C below 0.8 mole fraction Fe3O4 [1980Kat]. The thermodynamic values for reactionsin Cu-Fe-O are given in Table 3.The standard Gibbs energy of formation of CuFeO2 from its elements has been measured by [1975Pau] withpotentiometric method. Between 727 and 1027°C the following expression may be used:ΔfG

o (CuFeO2) = – 389215 + 86.89 T lnT – 541.8 TOther thermodynamic data for pure phases are given in Table 4. The Gibbs energy of the reaction4CuFe2O4+O2 ⇌ 2Cu+2CuO+2CuFe2O4 has been chosen from [1964Sch] because of the more credibleslope of the curve which is closer to the entropy of O2. CuFe2O4 is easily reduced [1972Shc] (the oxygenpressure at equilibrium is of the order of 1 kPa between 900 and 1200°C) with the formation of CuFeO2 + aspinel solid solution. The Gibbs energy of formation of CuFe2O4 determined by emf using stabilized zirco-nia [1976Eri] may be expressed by the following expression:ΔfG

o (CuFe2O4) = – 749230 + 156.0 T lnT – 945.2 T[1977Gro] investigates the Gibbs energy of the reaction Cu2O+3FeO → Fe3O4+2Cu, but proposes valueswhich are 30 kJ lower than the values calculated from more recent databases, such as Thermo-Calc. For thesame reason, values measured by [1977Jac] reported in Table 3 are preferred to those given by [1975Pau].Activities of Fe3O4 in the solid solutions Fe3O4-CuFe2O4 have been measured from reduction isotherms[1977Zal] and from emf measurements [1977Jac, 1980Kat]. They present a marked negative departure fromideal behavior. Potential diagrams of the Cu-Fe-O system are given in [1979Sek, 1980Ros, 1996Ina]. Thepotential diagram shown in Fig. 12 (potential oxygen-Cu and Fe activities) is from [1996Ina].


Main experimental works carried out on the properties are reported in Table 5. The spinel solid solutions areused in the synthesis of materials for absorber of electromagnetic energy and for storage elements in com-puter technology [1981Zin2]. The magnetoresistance of the spinel CuFe2O4 was shown to be negative[1968Rez]: the electrical resistance decreases under the influence of a magnetic field. The electrical resistiv-ity of CuFe2O4 which evolve with time and temperature according to its conditions of preparation[1972Rez, 1988San] was measured as 8.3 · 10 –3 Ω –1.cm –1, at 27°C on a sample sintered at 1000°C.CuFeO2 has been used as a cathode material [2000Suk] in Li batteries for its ability to intercalate Li+ ions.CuFe2O4 used as sensor was shown to have a good response to alcohol and reducing gases [2000Tao], whichis attributed to reactions between these gases and the absorbed oxygen. CuFe2O4 is also strongly recom-mended as a non consumable and green electrode [2003Sel] for substituting the C anode in Hall-Heroult cellsused in Al production. Indeed, CuFe2O4 electrodes present a good electrical conductivity and release O2

instead of CO2.Single crystals of CuFeO2 may be prepared by a flux method [1988Dor]. Cu crucibles with LiBO2 as a fluxproduce n type whereas Pt crucibles with Cu2O as a flux produce p type crystals. n type crystals exhibit aweak anisotropic conductivity with large activation energy and small mobilities; p type crystals, less aniso-tropic, present low activation energy with higher mobilities. CuFeO2 presents two magnetic transitions at 10

4 Cu–Fe–O


and 16 K [1991Mit]. A more recent investigation made on a single crystal [1995Zha] gives these transitionsat 9.5 and 13.5 K, respectively. In both magnetic structures, spins are collinear and parallel to the c axis.

Miscellaneous

Cu-Fe-O alloys (1.5 to 3.5 mass% Fe, 0.05 to 0.5 mass% O) have been proposed [1976Wue] for solderingalloys. Internal oxidation of Cu-Fe alloys, theoretically investigated by [1994Fra] forms simultaneously twotypes of iron oxides, both with a spinel structure, and gives plates on different habit planes of the matrix.The first type, Fe3O4 forms plates on the {111} plane of the copper lattice; the second type, close toγ-Fe2O3 forms plates on the {100} planes [1994Par].The magnetic moment at saturation of the spinel CuFe2O4 was calculated at 1.2 μB at 0 K [1975Oud], valueto be compared to the measured ones (1.2 and 1.66 μB). The magnetic moment measured by [1989Bre1] is1.37 μB at 300°C and 2.36 μB at 900°C. The saturation magnetization of ball milled CuFe2O4 dependsstrongly of the milling time, because of the decomposition of CuFe2O4 with the initial formation ofFe2O3, then of Fe3O4 [1998Goy1]. Fe2O3 decreases, whereas Fe3O4 increases the saturation magnetizationof ball milled samples. The decomposition reaction under ball milling was shown to be reversible[1998Goy2]. Equal phase composition was observed after 420 h of mechanical alloying CuFe2O4 and(CuO+Fe2O3).Copper, which is known as a tramp element in recycled steel produced from steel scraps, was found to besegregated by heating at 800°C [1997Suz]. The surface segregation seems to increase the thickness of theoxide layer formed at room temperature.

Table 1. Investigations on the Cu-Fe-O Phase Relations, Structures and Thermodynamics


[1959Ohn] Phase transition determinationX-ray diffraction (XRD),

320-790°C, CuFe2O4

[1961Yun,1964Yun]

Microscopic examination, XRD 200-800°C, Cu-Fe-O equilibria

[1964Sch] pO2 measurement, XRD 1000-1200°C, Cu2O-Fe3O4-Fe2O3-CuO region,pO2 < 0.1 MPa

[1966Bui,1966Gad]

Thermogravimetry, XRD 800-1600°C, Cu2O-Fe3O4-Fe2O3-CuO region, pO2< 0.1 MPa

[1966Yam,1967Yam]

Microscopy, XRD 1000-1600°C, Cu2O-Fe3O4-Fe2O3-CuO region,pO2 = 21 kPa (air atmosphere)

[1967Flo] pO2 measurement, XRD 700-1200°C, Cu2O-Fe3O4-Fe2O3-CuOregion, pO2 < 0.1 MPa,

[1968Ber] Thermogravimetry, XRD, dilatometry 1100 -1350°C, Cu2Fe5O8+δ, δa� < 0.08

[1968Nav] Solution calorimetry in aqueous HF 700°C, CuFe2O4

[1969Yam] Micrography, eutectoiddecomposition

450-1000°C, CuFe5O8

[1969Zal1] Equilibrium in H2/H2O atmosphere 700-1000, reduction of CuFeO2

[1969Zal2] XRD, equilibrium measurements 1000°C, Cu-Fe-O, < 0.1 MPa O2

[1970Fis] EMF, stabilized zirconia 1600°C, 0-100% Fe

(continued)

Cu–Fe–O 5




[1970Sch] XRD, chemical analysis, equilibriumunder N2/O2

900-975°C, 10 Pa to 50 kPa of oxygenpressure

[1970Tan1] Equilibrium in H2/H2O atmosphere 1100-1600°C, < 11 at.% Fe

[1970Tan2] Equilibrium in H2/H2O atmosphere 1500-1700°C, < 20 at.% Cu

[1971Fis] EMF, stabilized zirconia 1600°C, 0-100% Fe

[1971Sim] XRD, DTA 500-1200°C, CuFe5O8

[1971Tan] Equilibrium in H2/H2O atmosphere 1450-1700°C, 0-100% Fe

[1972Ono] Equilibrium in CO/CO2 atmosphere,thermogravimetry

500-1200°C, Cu2O-Fe3O4-Fe2O3-CuOregion

[1972Shc] Oxygen pressure at equilibrium 900-1200°C, reduction of CuFe2O4

[1972Tre] EMF, stabilized zirconia, oxygenpressure measurements

900-1000°C, Cu1–xFe2+xO4+γ

(γ < 0.04, x < 0.45)

[1973Bis] Equilibrium in H2/H2O atmosphere,thermogravimetry

1250-1350°C, 0-10.2 mass% Fe

[1973Kul] EMF, oxygen solubility measurement 1100-1300°C, < 10 at.% Fe

[1975Pau] EMF, stabilized zirconia 700 - 1025°C, CuFe2O4

[1976Bur] Isotopic analysis with radioactiveCo 60

1300°C, copper distribution betweenmetal liquid and slag

[1976Eri] EMF, stabilized zirconia 700 - 1050°C, CuFe2O4

[1976Fre] EMF, stabilized zirconia, XRD 825-1025°C, Cu-CuFeO2-Fe3O4-Fe2O3

domain

[1976Ois] EMF, stabilized zirconia, oxygenactivity measurements

1100-1200°C, < 0.5 at.% Fe

[1977Gro] Isobaric-Isothermal equilibrium 1100-1400°C, Cu2O+3FeO reaction

[1977Jac] EMF, stabilized zirconia, activitiesand oxygen potential measurements

1000°C, Cu-Fe-O phase relations, Fe3O4-CuFe2O4

measurements

[1977Zal] Activity measurements, equilibriummethod

1000°C, Cu0.5Fe2.5O4-Fe3O4 solid solution

[1978Rez] Calorimetry 25-527°C, CuFe5O8

[1979Nan] XRD 700-1000°C, oxidation of Fe50Cu50

[1980Chu] Interaction parameters measurements 1200-1500°C, < 10 mass% Fe

[1980Kat] EMF, stabilized zirconia 1000-1100°C, CuFe2O4-Fe3O4

[1981Zin1] XRD, thermogravimetry 1000-1300°C, CuFeO2, p(O2) = 21 kPa

[1981Zin2] XRD, thermogravimetry 800-1330°C, CuFe2O4-Fe3O4,p(O2) = 21 kPa

(continued)

6 Cu–Fe–O



[1986Acu] XRD, chemical analysis, slag-metalequilibrium

1100-1300°C, < 50 mass% Fe,< 20 mass% O

[1987Lis] XRD Fe3O4-CuFe2O4 solid solution

[1988Zin] XRD, thermogravimetry 800-1400°C, Fe2O3-Cu2O-CuO-Fe2O3 under21 kPa and 0.1 MPa of pO2

[1989Ois] EMF, stabilized zirconia 1500-1600°C, O in (Fe-Cu) alloys

[1992Xu] XRD, DTA, thermogravimetry 100-800°C, CuO-Fe2O3 reactions, reduction by H2

[1999Kat] EMF, stabilized zirconia, potentialoxygen measurements

800-1000°C, Fe1–xCuyO (0.05 < x < 0.12,y < 0.008)

[2000Ris] XRD, Fourier transformation infra-red (FTIR), Mössbauer

< 1350°C, CuO+FeOx mixtures (x = 1.33 or 1.5),Cu/Fe < 0.5

[2003Kat] EMF, stabilized zirconia, potentialoxygen measurements

800-1000°C, Fe3O4-CuFe2O4

[2003Yan] XRD 960°C, CuO-Fe2O3 in air

[2005Zin] XRD 800-1100°C, CuFe2O4-Cu0.5Fe2.5O4

under H2





Comments/References

(δFe)1538 - 1394

(αFe)< 912

cI2Im�3mW

a = 293.15

a = 286.65

pure Fe at 1360°C [Mas2, V-C2]dissolves 15 at.% Cu at 1440°C[2007Tur]pure Fe at 20°C [Mas2, V-C2](A2 structure). Dissolves 1.6 at.% Cuat 847°C

(γFe)1394 - 912

cF4Fm�3mCu

a = 293.16 at 915°C [Mas2, V-C2]. Dissolves15 at.% Cu at 1440°C

(Cu)< 1084,62

cF4Fm�3mCu

a = 361.46 at 25°C [Mas2, V-C2] dissolves 5 at.%Fe at 1095°C

Fe1–xO (Wuestite)1422 - 569

cF8Fm�3mNaCl

a = 431.0a = 429.3

0.05 < x < 0.12 [1991Sun]x = 0.05x = 0.12

(continued)

Cu–Fe–O 7






Comments/References

Fe3O4 I< 580

oP56PbcmFe3O4 I

a = 1186.8b = 1185.1c = 1675.2

[V-C2]

spinelSp1, Fe3O4 (h)(Magnetite)1596 - 580CuFe5O8

Sp2, βCuFe2O4 (h)1085 - 360

cF56Fd�3mMgAl2O4

a = 839.6a = 854.5

a = 841.6a = 838.2

at 25°Cat 1000°C [V-C2]inverse spinel[1969Haa][1969Mex, 1982Sch] disordered spinel

αFe2O3 (Hematite)< 1451

hR30R�3cαAl2O3

(Corundum)

a = 503.42c = 1374.83

melts at 1892°C under O2 pressure


a = 939.3 metastable phase[V-C2]

γFe2O3 (Maghemite) cF56Fd�3mMgAl2O4

a = 834 metastable phase[1989Rag]

Cu2O (Cuprite)< 1229

cP6Pn�3mAg2O

a = 421.7 [1989Rag]

CuO (Tenorite)<1230

mC8C2/cCrS

a = 466.2b = 341.6c = 511.8β = 99.49°

[1989Rag], decomposes at 1125°Cunder 0.1 MPa O2

αCuFe2O4 (r)< 360

tI56I41/amd

a = 821.6c = 870.9

[1968Eva, 1989Rag]dissolves up to 4 mol% Fe3O4

[1970Sch]

*τ, CuFeO2

(Delafossite)< 1180

hR*R�3mCuFeO2

a = 303.6 ± 0.1c = 1716.9 ± 0.3

a = 295.1c = 1703.0

[1995Zha, 1997Zha2] incongruentmelting under 21 kPa O2.

under 10 GPa of hydrostatic pressure[1997Zha1]

8 Cu–Fe–O




Quantity,per mole of atoms[J, mol, K]

Comments

4 CuFeO2 + O2 ⇌2 CuFe2O4 + 2 CuO 850 - 1000 ΔrG° = 221150 – 160.2 T [1964Sch,1970Sch]

CuO + Fe2O3 ⇌ CuFe2O4 700 - 1000 ΔrG° = 17150 – 17.99 T [1977Jac]

Cu2O + Fe2O3 ⇌2 CuFeO2

700 - 1000 ΔrG° = – 20160 + 8.76 T [1977Jac]

3 Cu + Fe3O4 + O2 ⇌3 CuFeO2

700 - 1025 ΔrG° = – 408030 + 194.5 T [1977Jac]

O2 (gas) ⇌ 2 {O} (in liquid Cu) 1100 - 1600 ΔrG° = – 141300 – 3.67 T [1970Tan1]{O} in at.%

O2 (gas) ⇌ 2 {O} (in liquidCu0.89Fe0.11)

1100 - 1600 ΔrG° = – 204400 – 29.1 T [1970Tan1]{O} in at.%


1500 - 1700 ΔrG° = – 238200 – 16.4 T [1970Tan2]{O} in mass%


1500 - 1700 ΔrG° = – 237700 – 17.8 T [1971Tan]{O} in mass%

O2 (gas) ⇌2 {O} (in liquidCu0.10Fe0.90)

1500 - 1700 ΔrG° = – 238700 – 11.9 T [1970Tan2]{O} in mass%

O2 (gas) ⇌ 2 {O} (in liquid Fe) 1500 - 1700 ΔrG° = – 240200 – 4.60 T [1971Tan]{O} in mass%


Phase Temperature Range[°C]


Comments

(1/14) CuFe5O8 25527

Cp = 19.6Cp = 25.4

[1978Rez]

(1/4) τ,CuFeO2 25700 - 1000

Cp = 20.0S° = 22.2ΔfG° = – 129700 + 45.7 T

[1988Bag]

[1969Zal1]

(1/7) CuFe2O4 25700 - 1000

Cp = 21.2S° = 20.14ΔfG° = – 137100 + 46.5 T

[1988Bag]

[1976Eri]

Cu–Fe–O 9



Table 5. Investigations of the Cu-Fe-O Materials Properties


[1968Rez] Magnetoresistance at 25°C CuFe2O4, sintered at 1000°C

[1968Eva] XRD, Mössbauer 4.2 - 1000 K, CuFe2O4

[1968Len] XRD, Curie point measurement CuFe2O4, below 800°C

[1969Haa] XRD CuFe5O8 crystal parameters

[1969Len] XRD, Curie point measurement CuFe5O8

[1969Mex] XRD, thermogravimetry CuFe2O4-CuFe5O8 solid solution

[1970Cer] XRD, spinel inversion CuFe5O8

[1970Kam] Voltamperometry 77 - 295 K, CuFe2O4

[1972Rez] Electrical conductivity 150 - 315°C, CuFe2O4

[1978Sid] XRD, BET, metallography Fe3O4 doped with Cu precipitation

[1979Dim] Mössbauer CuO-Fe3O4 and Cu2O-Fe3O4

reactions

[1979Has] EXAFS Cu+75 ppm Fe, internal oxidation

[1980Nan] XRD, EDXA, optical microscopy 700 - 1000°C, Fe + 4.5 mass% Cu,oxidation kinetics

[1980Wil] Mössbauer 76-340 K, Cu + 0.2 at.% Fe oxidizedat 850°C. p(O2) < 0.02 Pa

[1981Dhu] ESCA, Cu oxidation state Fe3O4-CuFe2O4

[1981Sap] Voltamperometry, Cu oxidation state Fe3O4-CuFe2O4

[1984Qia] XRD, thermogravimetry < 800°C, Cu-Fe-CuFe2O4, reductionisotherms, p(H2) = 15 kPa

[1985Han] High temperature XRD, Mössbauer 500 - 1200°C, CuFe2O4

[1986Pru,1987Pru]

AES (Auger Electron Spectroscopy), LEED (LowEnergy Electron Diffraction)

< 327°C, Fe on Cu surface,p(O2) < 0.1 MPa, oxygen uptake

[1988Dor] Electrical conductivity 820 - 1180°C, CuFeO2, n- and p-type

[1988Nai] XRD, magnetic susceptibility < 1000°C, αCuFe2O4

[1988San] XRD, resistivity 25 - 650°C, CuFe2O4

[1988Upa] XRD, Mössbauer, magnetic susceptibility 80 - 300 K, CuFe2O4

[1989Bre1] XRD, magnetic moment 300 - 900°C, CuFe2O4

[1989Bre2] XRD, magnetic moment CuFeO2 prepared by aqueousprecipitation, then annealed

[1989Tan] XRD, thermogravimetry 550 - 920°C, CuFe2O4

[1991Mit] Neutron diffraction < 50 K, CuFeO2

(continued)

10 Cu–Fe–O



[1994Fra,1994Par]

XRD, small-angle X-ray scattering Cu-Fe, internal oxidation

[1995Zha,1996Zha1]

XRD, Mössbauer, magnetic susceptibility CuFeO2 single crystal grown by afloating zone technique

[1996Zha2] Metallography, SEM, electron microprobeanalysis

CuFeO2 single crystal grown underAr, CO2 and Ar+0.5 % O2

atmospheres

[1997Suz] X-ray photoelectron, Auger spectroscopy, surfaceanalysis

< 800°C, Fe + 0.82 at.% Cu, surfacesegregation of Cu under oxidation

[1997Zha1] XRD, high pressure (diamond anvil) 25°C, < 10 GPa, CuFeO2

[1997Zha2] SEM, electron microprobe analysis CuFeO2 under Ar, CO2 and Ar+O2

atmospheres (< 2 % O2)

[1997Zha3] SEM, electron microprobe analysis CuFeO2+δ under CO2 and O2 (< 2 kPaO2), –0.11 < δ < +0.08

[1998Goy1] XRD, Mössbauer, magnetization 4.2 and 300 K, ball milled CuFe2O4

[1998Goy2] XRD, Mössbauer, magnetization 300 K, ball milled CuFe2O4 and(CuO+ Fe2O3) mixtures

[1998Ste] XRD, Mössbauer CuO + 0.25 mol% Fe2O3 ball milled,then annealed (< 727°C)

[1999Cro] XRD, SEM, EDXA, Fourier TransformationInfra-Red (FTIR)

CuFe2O4, self propagating hightemperature synthesis

[2000Suk] XRD, Mössbauer, EXAFS CuFeO2, 800 - 1000°C

[2000Tao] DTA, thermogravimetry, conductivity CuFe2O4 prepared by the sol-gelprocess, < 800°C

[2001Ste] XRD, TEM, Mössbauer, magnetization CuO + Fe2O3 ball milled

[2003Sel] XRD, FTIR, SEM, electrical conductivity CuFe2O4 prepared by combustionsynthesis, < 1000°C

[2004ElA] XRD, Mössbauer, magnetic susceptibility CuFeO2, 4.2 and 293 K

[2005Jia] XRD, Mössbauer CuO + Fe2O3 ball milled

[2006Gin] XRD, IR and UV spectroscopy, magneticsusceptibility

CuFe2O4, nanoparticle preparation bysoft chemistry methods

Cu–Fe–O 11



Fig. 1. Cu-Fe-O. Lattice parameter of the Fe3O4-CuFe2O4 spinel solid solution

12 Cu–Fe–O


Fig. 2. Cu-Fe-O. Lattice parameters of the Fe3O4-αCuFe2O4 and Fe3O4-βCuFe2O4 solutions on quenched samples

Cu–Fe–O 13



Fig. 3. Cu-Fe-O. Fe2O3-CuO quasibinary system under air

14 Cu–Fe–O


Fig. 4. Cu-Fe-O. Liquidus lines at 1200°C and 1500°C near the copper rich corner

Cu–Fe–O 15



Fig. 5. Cu-Fe-O. Isothermal section at 650°C

16 Cu–Fe–O



Cu–Fe–O 17




18 Cu–Fe–O



Cu–Fe–O 19




20 Cu–Fe–O


Fig. 10. Cu-Fe-O. Isothermal section at 1500°C showing L1 + L2 tie lines and log10(pO2 / bar)

Cu–Fe–O 21



Fig. 11. Cu-Fe-O. CuFe2O4-Fe3O4 vertical section

22 Cu–Fe–O


Fig. 12. Cu-Fe-O. Potential diagram at 1000°C

Cu–Fe–O 23



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26 Cu–Fe–O


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[1988Bag] Bagdavadze, D.I., Tsagareishvili, D.Sh., Tskhadaya, R.A., Gvelesiani, G.G., “Method ofComputation of Enthalpy Increment of Crystalline Inorganic Compounds at 0-298.15 KTemperature Range” (in Russian), Izv. Akad. Nauk Gruz. SSR, Ser. Khim, 14(3), 199–206(1988) (Calculation, Review, Thermodyn., 8)

[1988Dor] Dordor, P., Chaminade, J.P., Wichainchai, A., Marquestaut, E., Doumerc, J.P., Pouchard, M.,Hagenmuller, P., Ammar, A., “Crystal Growth and Electrical Properties of CuFeO2 SingleCrystals”, J. Solid State Chem., 75, 105–112 (1988) (Crys. Structure, Electr. Prop., 10)

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[1988San] Sankpal, A.M., Sawant, S.R., Vaingankar, A.S., “Electrical Resistivity and Lattice Para-meters Variation in Copper-Nickel Ferrite”, Indian J. Pure Appl. Phys., 26, 459–461(1988) (Experimental, Crys. Structure, Electr. Prop., 21)

[1988Upa] Upadhyay, R.V., Kulkarni, R.G., “Magnetic Properties of the Cu1+xTixFe2–2xO4 System”,J. Magn. Magn. Mater., 74(3), 327–329 (1988) (Crys. Structure, Experimental, Magn.Prop., 9)

[1988Zin] Zinovik, M.A., “The Phase Equilibrium Diagram of the CuO-Cu2O-Fe2O3-Fe3O4 System”,Russ. J. Inorg. Chem., 33(10), 1543–1545 (1988), translated from Zh. Neorg. Khim., 33,2689–2691 (1988) (Experimental, Phase Diagram, Phase Relations, 5)

[1989Bre1] Brezeanu, M., Patron, L., Contescu, A., “Structure and Properties of Copper Ferrite. I. Gen-eral Considerations” (in Romanian), Rev. Chim. (Bucharest), 40(5), 404–407 (1989) (Crys.Structure, Experimental, Magn. Prop., 36)

[1989Bre2] Brezeanu, M., Patron, L., Contescu, A., Mincu, V., Segal, E., “Structure and Properties ofCopper Ferrite. II. New Possibilities of Synthesis of CuFe2O4 with a Cubic Structure”(in Romanian), Rev. Chim. (Bucharest), 40(6), 510–513 (1989) (Crys. Structure, Experimen-tal, Phys. Prop., 18)

[1989Ois] Oishi, T., Nakagawa, K., Ono, K., “Phase Diagram and Thermodynamics of the MoltenCu-Fe-O System at 1773 K” (in Japanese), J. Jpn. Inst. Met., 53(7), 692–697 (1989) (Inter-face Phenomena, Phase Diagram, Thermodyn., 15)

28 Cu–Fe–O


[1989Rag] Raghavan, V., “Cu-Fe-O System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Inst.Metals, Calcutta, 5, 105–117 (1989) (Crys. Structure, Phase Diagram, Phase Relations,Review, 25)

[1989Tan] Tang, X., Xia, M.A., Goodenough, J.B., “Copper Ferrite Revisited”, J. Solid State Chem., 79(2), 250–262 (1989) (Crys. Structure, Experimental, Theory, 26)

[1991Mit] Mitsuda, S., Yoshizawa, H., Yaguchi, N., Mekata, M., “Neutron Diffraction Study ofCuFeO2”, J. Phys. Soc. Jpn., 60(6), 1885–1890 (1991) (Crys. Structure, Magn. Prop., 8)

[1991Sun] Sundman, B., “An Assessment of the Fe-O System”, J. Phase Equilib., 12(1), 127–140(1991) (Phase Diagram, Phase Relations, Thermodyn., Assessment, 53)

[1992Xu] Xu, J., Wang, W.-X., “Interaction in Cupric Oxide-Ferric Oxide System” (in Chinese), Cui-hua Xuebao, 13(6), 420–424 (1992) (Experimental, Thermodyn., 6)

[1994Fra] Fratzl, P., Paris, O., “Internal Oxidation of Cu-Fe-II. The Morphology of Oxide Inclusionsfrom the Minimization of Elastic Misfit Energy”, Acta Metall. Mat., 42(6), 2027–2033(1994) (Theory, 15)

[1994Par] Paris, O., Fratzl, P., Langmayr, F., Vogl, G., Haubold, H.G., “Internal Oxidation of Cu-Fe-I.Small Angle X-ray Scattering Study of Oxide Precipitation”, Acta Metall. Mat., 42(6),2019–2026 (1994) (Experimental, Crys. Structure, Physical Prop., 36)

[1995Kow] Kowalski, M., Spencer, P.J., “Thermodynamic Revaluation of the Cr-O, Fe-O and Ni-O Sys-tems: Remodelling the Liquid, BCC and FCC Phases”, Calphad, 19(3), 229–243 (1995)(Assessment, Phase Diagram, Phase Relations, Thermodyn., Review, 47)

[1995Zha] Zhao, T.-R., Hasegawa, M., Takei, H., “Growth and Characterization of CuFeO2 SingleCrystals”, J. Cryst. Growth, 154, 322–328 (1995) (Crys. Structure, Electronic Structure,Experimental, 11)

[1996Ina] Inaba, H., Yokokawa, H., “Analysis of Interfacial Reactions by the Use of Chemical Poten-tial Diagrams”, J. Phase Equilib., 17(4), 278–289 (1996) (Thermodyn., Calculations, 42)

[1996Zha1] Zhao, T.-R., Hasegawa, M., Takei, H., “Crystal Growth and Characterization of Cuprous Fer-rite (CuFeO2)”, J. Cryst. Growth, 166, 408–413 (1996) (Crys. Structure, Electronic Struc-ture, Experimental, 15)

[1996Zha2] Zhao, T.-R., Hasegawa, M., Takei, H., “Phase Equilibrium of the Cu-Fe-O System Under Ar,CO2 and Ar+0.5% O2 Atmospheres during CuFeO2 Single-Crystal Growth”, J. Mater. Sci.,31(21), 5657–5663 (1996) (Experimental, Morphology, Phase Relations, 16)

[1997Suz] Suzuki, S., Waseda, Y., “Copper Segregated Layer and Oxide Layer Formed on the Surfaceof an Fe-0.8 at.% Cu Alloy by Angle Resolved XPS”, Scr. Mater., 36(8), 915–920 (1997)(Interface Phenomena, Experimental, 15)

[1997Zha1] Zhao, T.R., Hasegawa, M., Kondo, T., Yagi, T., Takei, H., “X-ray Diffraction Study of Cop-per Iron Oxide (CuFeO2) under Pressures up to 10 GPa”, Mater. Res. Bull., 32(2), 151–157(1997) (Crys. Structure, Experimental, 33)

[1997Zha2] Zhao, T.R., Takei, H., “Study of the Oxidation and Reduction Kinetics of Copper Iron Oxide(CuFeO2) in the Cu-Fe-O System”, Mater. Res. Bull., 32(10), 1377–1393 (1997) (PhaseRelations, Experimental, 20)

[1997Zha3] Zhao, T.R., Hasegawa, M., Takei, H., “Oxygen Nonstoichiometry in Copper Iron Oxide(CuFeO2+δ) Single Crystals”, J. Cryst. Growth, 181, 55–60 (1997) (Crys. Structure, PhaseRelations, Experimental, 16)

[1998Goy1] Goya, G.F., Rechenberg, H.R., Jiang, J.Z., “Structural and Magnetic Properties of BallMilled Copper Ferrite”, J. Appl. Phys., 84(2), 1101–1108 (1998) (Crys. Structure, Experi-mental, Magn. Prop., Electronic Structure, 33)

[1998Goy2] Goya, G.F., Rechenberg, H.R., “Reversibility of the Synthesis-Decomposition Reaction inthe Ball-Milled Cu-Fe-O System”, J. Phys.-Condens. Matter., 10(50), 11829–11840 (1998)(Experimental, Magn. Prop., Phase Relations, 35)

[1998Ste] Stewart, S.J., Borzi, R.A., Punte, G., Mereader, R.C., “Phase Stability and Magnetic Beha-vior of Fe-doped CuO Powders”, Phys. Rev. B, 57(9), 4983–4986 (1998) (Experimental,Magn. Prop., 24)

Cu–Fe–O 29



[1999Cro] Cross, W.B., Affleck, L., Kuznetsov, M.V., Parkin, I.P., Pankhurst, Q.A., “Self-PropagatingHigh-Temperature Synthesis of Ferrites MFe2O4 (M = Mg, Ba, Co, Ni, Cu, Zn); Reactionsin an External Magnetic Field”, J. Mater. Chem., 9, 2545–2552 (1999) (Crys. Structure,Experimental, Magn. Prop., 26)

[1999Kat] Katkov, A.E., Lykasov, A.A., “Wuestite Solid Solutions in the Fe-Cu-O System”, Inorg.Mater., 35(7), 706–708 (1999), translated from Neorg. Mater., 35(7), 836–839 (1999)(Experimental, Phase Relations, Thermodyn., 6)

[2000Ris] Ristic, M., Hannoyer, B., Popovic, S., Music, S., Bajraktaraj, N., “Ferritization of CopperIons in the Cu-Fe-O System”, Mater. Sci. Eng. B, B77, 73–82 (2000) (Crys. Structure,Experimental, Electronic Structure, Phase Relations, 31)

[2000Suk] Sukeshini, A.M., Kobayashi, H., Tabuchi, M., Kageyama, H., “Physicochemical Character-ization of CuFeO2 and Lithium Intercalation”, Solid State Ionics, 128, 33–41 (2000) (Crys.Structure, Experimental, Optical Prop., 18)

[2000Tao] Tao, S., Gao, F., Liu, X., Sorensen, O.T., “Preparation and Gas-sensing Propertiesof CuFe2O4 at Reduced Temperature”, Mater. Sci. Eng. B, B77, 172–176 (2000) (Crys.Structure, Experimental, Electr. Prop., 17)

[2001Ste] Stewart, S.J., Borzi, R.A., Punte, G., Mercader, R.C., Garcia, F., “Microstructural and Mag-netic Characterization of Nanostructured α-Fe2O3 and CuO Mixtures Obtained by BallMilling”, J. Phys.-Condens. Matter., 13(8), 1743–1757 (2001) (Experimental, Magn.Prop., 33)

[2003Hal] Hallstedt, B., Gauckler, L.J., “Revision of the Thermodynamic Description of the Cu-O,Ag-O, Ag-Cu-O, Bi-Sr-O, Bi-Ca-O, Bi-Cu-O, Sr-Cu-O, Ca-Cu-O and Sr-Ca-Cu-O Sys-tems”, Calphad, 27, 177–191 (2003) (Phase Diagram, Phase Relations, Thermodyn., Assess-ment, 57)

[2003Kat] Katkov, A.E., Lykasov, A.A., “Spinel Phase Relations in the Fe3O4-CuFe2O4 System”,Inorg. Mater., 39(2), 171–174 (2003), translated from Neorg. Mater., 39(2), 223–226(2003) (Experimental, Thermodyn., Phase Relations, 10)

[2003Sel] Selvan, R.K., Augustin, C.O., Berchmans, L.J., Saraswathi, R., “Combustion Synthesis ofCuFe2O4”, Mater. Res. Bull., 38, 41–54 (2003) (Electr. Prop., Morphology, Transport Phe-nomena, 35)

[2003Yan] Yang, L.T., Liang, J.K., Song, G.B., Chang, H., Rao, G.H., “Compounds and Phase Rela-tions in the SrO-Fe2O3 -CuO, SrO-Fe2O3-Gd2O3 and Gd2O3-Fe2O3-CuO Ternary Systems”,J. Alloys Compd., 353, 301–306 (2003) (Crys. Structure, Experimental, Phase Relations, 48)

[2004Cla] Clavaguera-Mora, M.T., Touron, L., Rodriguez-Viejo, J., Clavaguera, N., “ThermodynamicDescription of the Cu-O System”, J. Alloys Compd., 377, 8–16 (2004) (Phase Diagram,Assessment, Thermodyn., #, 28)

[2004ElA] El Ataoui, Kh., Doumerc, J.-P., Ammar, A., Fournes, L., Wattiaux, A., Grenier, J.-C., Pou-chard, M., “Mössbauer Study and Magnetic Properties of CuFe1–xGaxO2”, J. Alloys Compd.,368, 79–83 (2004) (Crys. Structure, Experimental, Magn. Prop., Electronic Structure, 8)

[2005Jia] Jiang, J.S., Yang, X.L., Gao, L., Guo, J.K., “Nanostructured CuO- αFe2O3 Solid SolutionObtained by High-Energy Ball Milling”, Mater. Sci. Eng. A, 392, 179–183 (2005) (Crys.Structure, Electronic Structure, Experimental, Phase Relations, 15)

[2005Zin] Zinovik, E.V., Zinovik, M.A., “Hydrogen Reduction of Cu-Mn-Fe-O Spinel SolidSolutions”, Inorg. Mater., 41(3), 272–278 (2005), translated from Neorg. Mater., 41(3),332–338, (2005) (Crys. Structure, Experimental, Phase Diagram, Thermodyn., 14)

[2006Gin] Gingasu, D., Mindru, I., Patron, L., Carp, O., Matei, D., Neagoe, C., Balint, I., “Copper Fer-rite Obtained by Two “Soft Chemistry” Routes”, J. Alloys Compd., 425, 357–361 (2006)(Crys. Structure, Experimental, Magn. Prop., Optical Prop., 34)


30 Cu–Fe–O




Cu–Fe–O 31



Copper – Iron – Phosphorus

Kostyantyn Korniyenko

Introduction

The constitution of the ternary Cu-Fe-P system is of great interest in the first instance because its alloys arecharacterized by dispersional hardening and a good combination of heat and electrical conductivity withmechanical properties. At the same time, they are sufficiently technologically convenient in metallurgicalindustry. In particular, they may be melted practically in any furnace unit, which favorably distinguishes themfrom such well-known mass-produced low-alloy copper bronzes as chrome bronze and chromo-zirconiumbronze which contain readily oxidizing alloying elements [1987Kum]. Table 1 lists the publications dealingwith experimental studies of phase relations, crystal structures and thermodynamics of the Cu-Fe-P systemas well as the applied techniques.Crystal structures of the phases present in phosphorzus-containing ternary alloys, in particular, along theFe2P-Cu2P and Fe3P-Cu3P sections, were reported by [1948Now]. Slightly later phase relations in the par-tial Fe-Fe2P-Cu3P-Cu system were investigated in detail by [1950Vog]. As result, a liquidus surface projec-tion, a reaction scheme as well as a series of temperature-composition sections were proposed. Studies ofsimultaneous solubility of iron and phosphorus in copper at various temperatures are reflected in[1967Gla, 1968Gla, 1987Kum]. [1987Sau] used vapor co-deposited thin films possessing concentrationgradients for the evaluation of phase equilibria in the Fe-Fe2P-Cu3P-Cu partial system at 600°C. Phasesformed during sintering in high speed steels (HSS) at different temperatures were identified by[1990Bol]. Phase relations of precipitates formed during heat treatment of the commercial EFTEC5 alloywere reported by [1995Fuj1].Data on thermodynamic properties, mainly on the activity of phosphorus in liquid copper-iron alloys wereexperimentally obtained by [1959Sch, 1983Yam, 1984Ban, 2004Kai]. Reviews of literature data deal withCu-Fe-P phase equilibria [1972Gla, 1979Cha, 1979Dri, 1988Rag, 1998Rag], crystal structures [1979Cha,1988Rag] and thermodynamics [1974Sig]. Thermodynamic calculations of isothermal sections and a liqui-dus surface projection were carried out by [1990Gus]. Activity interaction parameters between copper andphosphorus in liquid iron were assessed by [1974Sig, 1993Din].As a whole, phase equilibria in the ternary Cu-Fe-P system are studied insufficiently. In particular, the char-acter of phase relations in the partial Fe-Fe2P-Cu3P-Cu system should be verified by modern techniques ofphysico-chemical investigations. The gap in data on phase equilibria in the partial Fe2P-Cu3P-P systemshould be filled up. Acquiring new information would extend practical applications of Cu-Fe-P alloys.

Binary Systems

The Cu-Fe, Cu-P and Fe-P systems are accepted from [2007Tur], [2002Per1] and [2002Per2], respectively.

Solid Phases

Crystallographic data on the known unary and binary Cu-Fe-P phases and their concentration and tempera-ture ranges of stability are listed in Table 2. Among the pure components, iron forms an extended fieldof solid solutions, in particular, solubility of copper in (γFe) reaches 13.5 at.% at 1417°C in the binaryCu-Fe boundary system. It should be noted that this solubility of Cu in Fe is of retrograde character. Withtemperature decreasing from the solidus solubility of copper first increases. Visible solubilities of the thirdcomponent in the binary phases were not found. The temperature of mutual transformation between the(βCu3P) and (αCu3P) modifications is not yet established clearly, therefore this compound in the text andfigures is marked as (Cu3P). No ternary phases were found in the system.

Cu–Fe–P 1



Quasibinary Systems

A phase diagram of the Fe2P-Cu3P quasibinary system was constructed by [1950Vog] based on thermal ana-lysis, chemical analysis and optical microscopy. It is shown in Fig. 1 with certain corrections according to theconstitution of the accepted Fe-P and Cu-P binary systems (temperatures of congruent melting of the Fe2P andCu3P phases are shifted from 1350 and 1023°C to 1370 and 1022°C, respectively). The temperature of theeutectic was reported by [1950Vog] as 1013°C, measured by thermal analysis, but it seems not to be precise,because the experimental points of the eutectic temperature drawn by [1950Vog] scatter between slightlyhigher than 1000°C at the Cu3P-side and about 960°C at the Fe2P-side (at 15 at.% Cu). Thermal curves arenot shown in the article, therefore the temperature of the invariant reaction Le ⇌ δ + (Cu3P) needs further ver-ification. This temperature is assumed in our assessment to be ~ 1013°C as reported in the review of [1979Dri](misprinted as 1073°C). [1950Vog] did not find visible solubilities of the third component, neither Cu in theFe2P, nor Fe in the Cu3P phase. [1990Gus], however, in his thermodynamic calculation assessed a solubility ofabout 3.3 at.% Fe in Cu3P without referencing a corresponding experimental determination. The eutectic pointwas calculated by [1990Gus] to be at about 4.7 at.% Fe (or about 69.8 at.% Cu).As a whole, the quasibinary character of the Fe2P-Cu3P section allows to triangulate the ternary Cu-Fe-Psystem along this section and to consider the phase equilibria of the Fe-Fe2P-Cu3P-Cu partial system inde-pendently of the more P-rich part of the system.[1967Gla, 1968Gla], based on electrical conductivity and microhardness measurements, expressed anopinion, that also the Cu-Fe2P and Cu-Fe3P sections would be quasibinary, but this assumption contradictsthe constitution of the Fe-Fe2P-Cu3P-Cu partial system determined by [1950Vog].


Temperatures, types of reactions and available compositions of the phases taking place in the invariant equi-libria in the partial Fe-Fe2P-Cu3P-Cu system are listed in Table 3. Compositions were converted from massto atomic per cents in some cases. These data and the partial reaction scheme in Figs. 2a and 2b are basedmainly on [1950Vog] concerning invariant reactions and the liquidus surface. The reaction scheme was ear-lier presented in the review of [1988Rag], the invariant reactions were also listed in the reviews of[1979Cha, 1979Dri]. Some corrections have been carried out in Fig. 2 according to the constitution ofthe binary boundary systems. Seven four-phase and three three-phase invariant reactions with participationof liquid, listed in Fig. 2 and in Table 3, demonstrate the character of stable ternary liquid-solid equilibria inthe Fe-Fe2P-Cu3P-Cu partial system. This aspect, however, is reported mainly in only one publication,[1950Vog], which was written more than fifty years ago, it is desirable to verify it by modern techniquesof physico-chemical analysis. Also in the phosphorus rich region, there is a gap of knowledge on phase rela-tions, which is desirable to be filled up.An attempt of thermodynamic calculation of temperatures and co-ordinates of the four-phase equilibria inthe partial Fe-Fe2P-Cu3P-Cu system was carried out by [1990Gus] (namely, U2, U3 and E1), but the resultsquite essentially differ from the experimental data reported by [1950Vog]. In particular, the temperature ofthe invariant reaction L’ + (γFe) ⇌ (αFe) + L’’ was calculated to be 1210°C, in contrast to the value of1094°C, experimentally determined by [1950Vog]. After the calculation also Cu3P and Fe3P are assumedto be in mutual equilibrium below the solidus temperatures, whereas [1950Vog] and [1987Kum] experimen-tally found Fe2P and (Cu) to be in mutual equilibrium. As the calculated Gibbs energies of these four phasesdo not significantly depend on ternary terms, that means already the binary Cu-P and Fe-P datasets needrevised assessments.[1988Rag] postulated a quasibinary monotectic at the place, where [1950Vog] assumed the critical point c2between L’ and L” of the three-phase equilibrium L’ + L” + Fe3P. The calculations, however, show at least,that there is no reason to assume such a quasibinary monotectic.


Figure 3 shows the liquidus surface projection of the Fe-Fe2P-Cu3P-Cu partial system. It is based mainly onthe data of [1950Vog] and reproduced similarly in the reviews of [1979Cha, 1988Rag]. Occurrence of a

2 Cu–Fe–P


large miscibility gap closed in itself is the characteristic feature of this liquidus surface. Locations of tie linesin the miscibility gap and compositions of the c1 (1405°C, Fe57.5Cu41.4P1.1 in at.%) and c2 (1210°C,Fe50Cu37P13 in at.%) points were established by [1950Vog]. Three four-phase invariant equilibria contain-ing two different liquids take place, and one of them is the ternary monotectic equilibrium E1: L’ ⇌ L’’ +(αFe) + Fe3P at 1030°C. The liquidus surface projection of the copper rich corner is shown schematically.[1950Vog] gave compositions of the equilibrium phases in the Cu rich corner, however, at least the reportedcompositions touching the surface of primary crystallization of (Cu) are thermodynamically very unlikely.As in this range liquid and (Cu) can be treated as dilute solutions, for which Henry’s law is approximatelyvalid, the slopes of the surface of primary crystallization of (Cu) in the directions Cu-Fe and Cu-P cannotdeviate much from those in the binary Cu-Fe and Cu-P systems, respectively. [1950Vog] has drawn thepoints of the equilibria called U3, E2, U4 in Fig. 2 in this sequence on a line within less than 1% distance.Thus the temperatures 1090, 1028 and 1070°C should follow in very short distance demanding an extre-mely wavy surface of primary crystallization of (Cu), which is thermodynamically very unlikely. Thereforefurther investigations in this range of compositions are necessary.A calculated liquidus surface projection of the partial Fe-Fe2P-Cu3P-Cu system was reported by [1990Gus].The presented arrangement of the monovariant curves on the Fe rich side is not very different from thosereported by [1950Vog], but at the Cu rich side the saturated liquid shows much larger Fe- and P-solubilitiesthan that of [1950Vog] avoiding the unlikely steep slopes. But the dataset of [1990Gus] contradicts thestable equilibrium between (Cu) and Fe2P, experimentally well established by [1950Vog] and [1987Kum].Some isotherms of the solidus surface projection in the Cu rich corner are shown in Fig. 4 as reported by[1987Kum]. Like the solidus line in the binary Cu-P system, the solidus temperature steeply falls withincreasing phosphorus content, whereas the slope in direction of the Fe content is slightly positive likethe solidus line in the Cu-Fe binary system.In general, information on the liquid-solid equilibria is incomplete, especially on the solidus and solvus sur-faces and further experimental studies of these aspects are necessary.

Isothermal Sections

Simultaneous solubilities of iron and phosphorus in copper at the temperatures 950, 700, 600 and 200°Cwere studied by [1967Gla]. Those at 950 and 700°C are also reported by [1968Gla]. The maximal solubi-lities at each temperature are listed in Table 2.Isothermal sections of the Cu-Fe-P system in the Cu rich corner at 1000, 900 and 700°C are shown in Figs. 5,6 and 7, respectively, as reported by [1987Kum]. These diagrams contradict those of Glazov et al. [1967Gla,1968Gla], but they seem to be the more reliable ones. They show decreasing Fe-solubility and increasing P-solubility in the copper-based solid solution with decreasing temperature like in the binary boundary sys-tems. Both additions, phosphorus to the Cu-Fe alloys and iron to the Cu-P alloys intensively reduce thesolubility of the other alloying element in the copper solid solution. Solubilities of iron and phosphorusin copper along the Cu-Fe2P section at different temperatures are listed in Table 2.Vapor co-deposited Cu-Fe-P alloy thin films possessing concentration gradients were used by [1987Sau] forthe rapid assessment of phase equilibria. According to the data obtained from diffraction patterns taken fromthe film deposited at about 600°C, three three-phase fields, namely (Cu) + (αFe) + γ, (Cu) + δ + γ and (Cu)+ (Cu3P) + δ were clearly defined. This agree well with the data about phase relations obtained by[1950Vog] and [1987Kum]. The same three-phase fields appear in a schematic isothermal section of theFe-Fe2P-Cu3P-Cu partial system at room temperature reported by [1950Vog] based on experimental data.The room temperature solubilities of the other components in the (Cu) and (αFe) solid solutions as wellas of the third component in the binary compounds are shown to be negligible.[1990Gus] thermodynamically calculated isothermal sections of the Cu-Fe-P system from the iron corner upto 8 mass% (13.55 at.%) P and up to 50 mass% (46.78 at.%) Cu at 1250 and 1120°C as well as up to10 mass% (16.7 at.%) P at 1100°C. The obtained diagrams agree well with the accepted Cu-Fe and Fe-Pbinary systems. In the presented ranges of compositions and temperatures equilibria of liquid, iron and cop-per occur with solid (αFe) and (γFe). Using the Cu-P, Fe-P and Cu-Fe binary phase diagrams and the data of[1950Vog] on phase relations in the Cu-Fe-P system, [2004Kai] deduced participation of the same phases inequilibria near the Cu-Fe edge at 1200°C.

Cu–Fe–P 3




[1950Vog] constructed from his experimental results several of temperature-composition sections, namelyfrom Cu towards Fe2P up to 3 mass% Fe2P (or Cu-Fe2.66Cu96.01P1.33 in at.%), the whole range fromFe3P towards Cu3P, as well as the sections with Fe:Cu ratios of 9:1 and 1:1 up to 21 and 18 mass% P, respec-tively. The last two sections after conversion to at.% correspond to Fe91.1Cu8.9-Fe61.3Cu6.0P32.7 andFe53.2Cu46.8-Fe37.4Cu32.9P29.7, respectively. As a whole, displacement of the primary crystallization regionsin all these sections agree satisfactorily with the determined constitution of the liquidus surface but for ver-ification of the character of the phase equilibria at lower temperatures it is desirable to determine more pre-cisely the solidus and solvus compositions.

Thermodynamics

The effect of copper on the activity coefficient of phosphorus in liquid iron alloys was studied by[1983Yam, 1984Ban, 2004Kai]. The Knudsen cell – mass spectrometer combination measurements triedat 1600°C by [1983Yam] show that the vapor pressure of copper in liquid iron alloy is high, and the insidewall of the ion source was severally contaminated by this vapor and this contamination made accurate deter-minations impossible. Later the vapor pressure of phosphorus in liquid Cu-Fe-P alloys containing up to10 at.% Cu was measured by [1984Ban] using the transportation method at 1400°C. The results were trea-ted by the model of interstitial solution proposed by J. Chipman, and the effect of copper on the activitycoefficient of P in liquid Fe was determined by assuming the copper to dissolve substitutionally. A valueof εP

Cu = – 9.42 ± 2.86 at 1400°C was obtained for the interaction parameter. In order to determine theactivities of phosphorus and iron in liquid Cu-Fe-P alloys, the subject specimens were brought by[2004Kai] into equilibrium with Cu-Fe-P solid solutions and a mixture of (Al2O3) + (AlPO4) + (FeAl2O4)at temperatures of 1253 and 1143°C. The equilibrium oxygen partial pressures were determined with the aidof an electrochemical technique involving MgO-stabilized zirconia as electrolyte and a Mo + MoO2 refer-ence electrode. The activity coefficient of phosphorus referred to 1 mass% solution in pure liquid copperwas determined as log fP

o = (4.46 ± 0.40) – (8.710 ± 770)/T (temperature in K). The iron activities referredto pure solid Fe were obtained to be log aFe = (0.37 ± 0.12) + (500 ± 200)/T (temperature in K).Thermodynamic properties of copper and phosphorus in liquid iron have been evaluated by [1974Sig]and [1993Din]. The following expressions were obtained by [1974Sig]: lnγoCu(Fe) = 4026/T andlnγoP(Fe) = –7645/T – 4.39 (temperature in Kelvin). The interaction parameter was obtained by[1993Din] at 1400°C as εP

Cu = 0.36. A thermodynamic database of the Cu-Fe-P system was assessedby [1990Gus], Using this dataset in the computer program Thermo-Calc a partial liquidus surface projec-tion, the Fe2P-Cu3P quasibinary section as well as partial isothermal sections at 1250, 1120 and 1100°Cwere calculated. To improve this dataset first the binary datasets Cu-P and Fe-P need re-assessment, asthe chemical potential of P in the two-phase field Fe3P + Fe2P must be smaller than in the two-phase fieldCu3P + (Cu) in order to reproduce equilibrium between the two phases (Cu) and Fe2P in the ternary system.Precise measurements of P in these binary two-phase fields would be of very high value.


In general, copper-based alloys with excellent strength and electrical conductivity are widely used for manyapplications, such as lead frames, connectors and conducting springs, sliding contacts, crystal tube for con-tinuous casting and electrodes for resistance welding. Precipitation hardening is the main method of devel-oping copper-based high-strength and high-conductivity alloys because it can effectively improve thestrength of the alloys with little harm on the electrical conductivity. Many precipitation hardening copperalloys have been developed and commercially used, such as Cu-Fe. Addition of phosphorus to them is alsoperspective because it allows to form Fe2P precipitates during the aging process. However, additions of ironand phosphorus to copper must be optimal because these components decrease its conductivity [2006Lu].Among the important practically useful alloys are the CuFe2P alloy named as “Wieland-K65” for connectorapplications in the automotive industry [1986Zei, 1987Puc, 1988Puc, 1993Iwe, 1994Boe] and the EFTEC5Cu-based alloy containing 1 mass% Fe and a trace amount of phosphorus used as a semiconducting material[1991Fuj, 1995Fuj1].

4 Cu–Fe–P


Other important aspects are phosphide phase (Cu3P and Fe3P) additions, applied to promote liquid phasesintering in 316L stainless steel because they reduce porosity and increase hardness [1999Pre].The applied experimental techniques and investigated types of properties of the Cu-Fe-P alloys are general-ized in Table 4. Reviews on various properties of these materials are presented in [1972Gla, 1982Sak,1983Tak, 1986Zei, 1987Puc, 1990Seg, 1993Iwe, 1999Mor].[1990Wat, 1996Tyl] (reviewed in [1999Mor]) studied conditions of gas filled porosity, present afterhigh spray deposition of Cu-Fe-P alloys. Such porosity deteriorates mechanical and electrical properties.No spray deposition condition could be found to diminish this porosity. [1996Tyl] observed that such porescan be reduced in size by thermo-mechanical working but they did not reopen during long high-temperaturetreatments. Investigations of [1990Wat] demonstrated that small additions of sufficiently strong nitrideforming elements (for example, zirconium) may lead to avoid the gas porosity problem.

Miscellaneous

[1979Mij] investigated physical properties of some Cu-Fe-P alloys and developed a new mold material“KT-1” for continuous casting; they also presented a model of heat transfer from molten steel to coolingwater. Oxidability of the Fe1.2Cu98.4P0.4 (in mass%, Fe1.4Cu97.8P0.8 in at.%) alloy at 600°C during 2, 4, 6and 10 hours was determined by [1979Ste]. They showed that additions of cerium to this alloy reducethe oxidability. [1987Sau] studied Cu-Fe-P vapor co-deposited thin films sputtered onto heated substrates.Because all the relevant equilibrium three-phase fields were observed in a single film, the authors consid-ered about the potential usefulness of this technique for rapid assessment of phase equilibria in ternary sys-tems. Mechanisms of sintering in HSS, alloyed with phosphorus additions, were studied by [1990Bol].They showed that densification occurs in distinct stages, due to successive formation of a series of liquidphases at various sintering temperatures. An activated sintering model for densification by the formationof a phosphide rich grain boundary liquid was developed. [1999Pre] established that phosphide powdersof Cu3P and Fe3P enhance the sintering behavior of 316L stainless steel by the formation of a phosphideliquid at temperatures around 1050°C. The optimum amount of Cu3P as well as Fe3P in terms of sinteringtemperature, dimensional stability, interconnected porosity and density was assessed to be approximately 10mass%. Improvements in sintering behavior are caused by an activated liquid phase sintering reaction due toa complex eutectic reaction between austenite and iron-chromium phosphide phases, which may consist ofboth M3P and M2P type phosphides. In [1972Gla] aspects of interaction between alloying components inalloys are presented, in particular for Cu-Fe-P alloys. It was shown that singular maxima of microhardnessand electric conductivity on the compositional dependences of properties as well as sharp curvatures ofsolubility isotherms correspond to certain ratios between alloying components.

Table 1. Investigations of the Cu-Fe-P Phase Relations, Structures and Thermodynamics

Reference Method / Experimental Technique Temperature/ Composition/ Phase RangesStudied

[1948Now] X-ray diffraction The Fe3P-Cu3P and Fe2P-Cu2P sections

[1950Vog] Melting in Pythagor’s crucibles, thermalanalysis, optical microscopy, chemicalanalysis

≥ 1520°C, the Fe-Fe2P-Cu3P-Cu partialsystem

[1967Gla] Melting in Tamman furnace (graphitecrucibles), chemical analysis, opticalmicroscopy, annealing followed by waterquenching, electrical conductivity andmicrohardness measurements

950, 800, 700, 600, 300, 200°C,≥ 7 at.% Fe,≥ 7 at.% P

(continued)

Cu–Fe–P 5



Reference Method / Experimental Technique Temperature/ Composition/ Phase RangesStudied

[1968Gla] Melting in Tamman furnace (graphitecrucibles), chemical analysis, annealingfollowed by water quenching, electricalconductivity and microhardnessmeasurements

950, 700°C, ≥ 4 at.% Fe, ≥ 4 at.% P

[1983Yam] Melting, Knudsen cell-mass spectrometry εPCu in liquid phase, 1600°C

[1984Ban] Inert gas transportation method εPCu in liquid phase, 1400°C,

< 10 mass% Cu

[1987Kum] X-ray diffraction, thermomagneticanalysis(TMA-2 apparatus technique),metallography

1000-450°C, ≥ 3 mass% Fe,≥ 3 mass% P

[1987Sau] Simultaneous secondary iron beamsputterco-deposition, X-ray diffraction, EMPA

600, 500, 400°C, the Fe-Fe2P-Cu3P-Cu partialsystem

[1990Bol] Sintering, optical microscopy, SEM,wave-length dispersive analysis, EDX, X-raydiffraction, TEM, DTA

≥ 1180°C, high-speed steels (HSS) with Cu-Por Fe-P alloys powders addition

[1995Fuj1] Heat treatment with precipitation,chemicalanalysis, SEM, EMPA

≥ 980°C, the EFTEC5 alloys

[2004Kai] Emf, X-ray diffraction 1253, 1143°C, Cu-Fe-P liquid alloyssaturated with Cu-Fe-P solid solutions





Comments/References

(P) (red)< 417

c*66 a = 1131 sublimation at 1 bar, triple point at576°C, > 36.3 bar, triple point at589.6°C, 1.013 bar [Mas2, V-C2]

(P) (white)< 44.14

c**-P (white)

a = 718 at 25°C [Mas2]

(P) (black) oC8CmcaP (black)

a = 331.36b = 1047.8c = 437.63


(continued)

6 Cu–Fe–P





Comments/References

(αδFe)

(αFe) (r) (ferrite)< 912(δFe) (h2)1538 - 1394

cI2Im�3mW a = 286.65

a = 293.15

at 25°C [Mas2]

at > 1394°C [Mas2]

dissolves 5.8 at.% Cu at 1487°C[2007Tur]dissolves 1.6 at.% Cu at 847°C[2007Tur]dissolves 4.55 at.% P at 1048°C[2002Per2]

(γFe) (h1) (austenite)1394 - 912

cF4Fm�3mCu

a = 364.67 at 915°C [V-C2, Mas2]

dissolves ~ 14 at.% Cu at 1425°C[2007Tur]dissolves 0.56 at.% P at ~1150°C[Mas2]

(Cu)< 1084.62

FexCu1–x–yPy

cF4Fm�3mCu

a = 361.46 at 25°C [Mas2]dissolves 5 at.% Fe at 1095°C, 1 at.%Fe at 847°C [2007Tur]dissolves 3.5 at.% P at 714°C[2002Per1]0 < x < 0.037,0 < y < 0.018, 1000°C [1987Kum]0 < x < 0.035,0 < y < 0.017, 950°C [1987Kum]0 < x < 0.024,0 < y < 0.012, 950°C [1967Gla,1968Gla]0 < x < 0.026,0 < y < 0.013, 900°C [1987Kum]0 < x < 0.017,0 < y < 0.009, 800°C [1987Kum]0 < x < 0.001,0 < y < 0.0005, 700°C [1987Kum]0 < x < 0.017,0 < y < 0.008, 700°C [1967Gla,1968Gla]0 < x < 0.014,0 < y < 0.007, 600°C [1967Gla]0 < x < 0.005,0 < y < 0.01, 200°C [1967Gla]

(continued)

Cu–Fe–P 7






Comments/References

βCu3P< 1022

hP8P3m1βCu3P a = 409.2

c = 718.6

25 to 31 at.% P [Mas2, V-C2]

congruent melting at 1.013 barat 560°C [V-C2]

αCu3P hP24P63cmαCu3P a = 695.93

c = 714.3a = 699.2c = 717.0

25 to 31 at.% P (?) [Mas2, V-C2]low-temperature phase[V-C2]

at 25 at.% P [1965Man, 1988Rag,1994Sub]

α, CuP2< 891

mP12P21/cCuP2 a = 580.04

b = 480.63c = 752.63β = 112.70°

~ 66.7 at.% P [Mas2, 2002Per1]

congruent melting at 15.2 bar[1982Mol, Mas2,V-C2]

β, Cu2P7< 850

mC72C2/mCu2P7

a = 1265.8b = 725.6c = 1463.0β = 107.46°

77.78 at.% P [1994Sub, 2002Per1][1982Mol, 1994Sub]

γ, Fe3P< 1166

tI32I�4Ni3P

a = 910.8c = 445.5

25 at.% P [Mas2, V-C2]

δ, Fe2P (I)< 1370

hP9P�62mFe2P

a = 586.4c = 346.0

33.3 to 34 at.% P [Mas2]at 33.3 at.% P [Mas2, V-C2]

δ’, Fe2P (II) oP12PnmaFe2P

a = 577.5b = 357.1c = 664.1

33.3 at.% P, high-pressure phase [Mas2]at 800°C and 80 kbar [Mas2, V-C2]

ε, FeP≲ 1370

oP8Pna21MnP

a = 519.3b = 579.3c = 309.9

50 at.% P [Mas2][1972Sel, 1990Oka]

η, FeP2 oP6PnnmFeS2 (marcasite)

a = 497.29b = 565.68c = 272.30

66 to 67 at.% P [Mas2]

at 66.7 at.% P [1969Dah, 1990Oka]

(continued)

8 Cu–Fe–P





Comments/References

κ, FeP4 (I) mP30P21/cFeP4

a = 461.9b = 1367.0c = 700.2β = 101.48°

80 at.% P [Mas2][1978Jei]

κ', FeP4 (II) oC20C2221FeP4

a = 500.5b = 1021.2c = 553.0

80 at.% P, high-pressure phase,synthesized at 60 kbar in a cubic anvildevice [1978Sug]

χ (Fe-P) o**

a = 359b = 401c = 432

< 20 at.% P, metastable, labelled as“Fe4+P” [1990Oka]after 25 h aging at 500°C [1961Hor]



Cu Fe P

L’ + δ ⇌ L’’ + γ 1103 U1 L ~ 13 ~ 66.5 ~ 20.5

L’ + (γFe) ⇌ (αδFe) + L’’ 1094 U2 L’ ~ 12.0 ~ 75.5 ~ 12.5

L’’ + (γFe) ⇌ (Cu) + (αδFe) 1090 U3 - - - -

L’’ ⇌ (Cu) + δ 1071 e2 - - - -

L’’ + δ ⇌ γ + (Cu) 1070 U4 - - - -

L’ ⇌ L’’ + (αδFe) + γ 1030 E1 L’ ~ 11 ~ 74 ~ 15

L’’ ⇌ (αδFe) + (Cu) + γ 1028 E2 L’’ - - -

L ⇌ δ + Cu3P ~ 1013 e4 L ~ 72 ~ 3 ~ 25

L’’ ⇌ δ + Cu3P + (Cu) ~ 713 E3 L’’ ~ 0.5 ~ 84.5 ~ 15

Table 4. Investigations of Cu-Fe-P Materials Properties


[1967Gla] Microhardness measurements (PMT-3type tester)

Microhardness

[1968Gla] Microhardness (PMT-3 type tester) andelectrical conductivity measurements

Microhardness, electrical conductivity

(continued)

Cu–Fe–P 9




[1976Wue] Soldering tests Soldering properties

[1979Mij] Thermal behavior tests, mechanicalproperties tests (incl. Vickers and Brinellhardness techniques), electricalconductivity measurements

Thermal conductivity,thermal expansion, tensile strength, hardness,electrical conductivity, Young’s modulus,fatigue behavior

[1979Ste] Electrical conductivity measurements,mechanical (Vickers hardness andplasticity) tests

Electrical conductivity, plasticity, hardness

[1979Sve] Partial prealloying of powder, SEM,laboratory-scale tensile tests

Dimensional change, tensile strength

[1980Tak] TEM, hardness tests Hardness

[1982Ste] Mechanical and electrical properties tests Hardness, electrical conductivity

[1987Kum] Mechanical and electrical properties tests,thermomagnetic analysis (TMA-2apparatus)

Strength, yield point, electrical resistivity

[1987Puc] Mechanical and electrical properties tests Hardness, electrical conductivity

[1988Puc] Physical properties tests Solderability, surface roughness, brightness,bondability

[1989Miy] Mechanical and electrical properties tests,TEM

Tensile strength, yield strength, elongation,conductivity, resistance to heat-softening

[1990Pie] Mechanical tests Plasticity

[1991Fuj] SEM, optical microscopy, TEM,electrical properties tests, Vickershardness measurements

Electrical resistivity and conductivity, hardness

[1991Oha] Mechanical and electrical properties tests Tensile strength, elongation, cyclic bendingstrength, electrical conductivity

[1992Boe] Vickers hardness tests Hardness, breaking strength, electricalconductivity, residual voltage

[1994Boe] Mechanical tests with Larson-Millerparameters applying, conductivitymeasurements

Stress relaxation behavior, hardness,electrical conductivity

[1994Kam] SEM, magnetic measurements Coercive force

[1995Fuj1] Vickers hardness tests, electricalresistivitymeasurements

Hardness, electrical resistivity

[1995Fuj2] Mechanical tests Hardness

[1999Pre] Specimens dimensions and volumesdetermination, Vickers microhardnessand hardness tests, tensile tests

Sintered density, interconnected porosity,volume shrinkage, microhardness andhardness, tensile strength

(continued)

10 Cu–Fe–P



[2000Kit] Thermal and mechanical tests Heat resistance, strength

[2002Don] Mechanical tests, conductivitymeasurements

Microhardness, electrical conductivity

[2006Lim] Mechanical testing Tensile strength, tensile elongation

[2006Lu] Electropolishing method, TEM, tensiletesting (WES-100 testing machine),conductivity testing (ZY9987 digitmicrometer)

Tensile and yield strength, elongation,conductivity

Fig. 1. Cu-Fe-P. The Fe2P-Cu3P phase diagram

Cu–Fe–P 11



Fig.2a

.Cu-Fe-P.

Reactionscheme,part1

12 Cu–Fe–P


Fig.2b

.Cu-Fe-P.

Reactionscheme,part2

Cu–Fe–P 13



Fig. 3. Cu-Fe-P. Partial liquidus surface projection

14 Cu–Fe–P


Fig. 4. Cu-Fe-P. Partial solidus surface projection, Cu rich corner

Cu–Fe–P 15



Fig. 5. Cu-Fe-P. Partial isothermal section at 1000°C

16 Cu–Fe–P



Cu–Fe–P 17




18 Cu–Fe–P


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Cu–Fe–P 19



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20 Cu–Fe–P


[1990Oka] Okamoto, H., “The Fe-P (Iron-Phosphorus) System”, Bull. Alloy Phase Diagrams, 11(4),404–412 (1990) (Crys. Structure, Phase Diagram, Thermodyn., Assessment, Review, Magn.Prop., *, 88)

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[1993Din] Ding, X., Wang, W., Han, Q., “Thermodynamic Calculation of Fe-P-j System Melt”, ActaMetall. Sin. (China), 29(12), B527-B532 (1993) (Thermodyn., Calculation, Theory, 7)

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Cu–Fe–P 21



[2000Kit] Kita, K., Kobayashi, K., Monzen, R., “Dispersions of Particles and Heat Resistance in a Cu-Fe-P Alloy” (in Japanese), J. Soc. Mater. Sci., Jpn., 49(5), 482–487 (2000) (Morphology,Experimental, Mechan. Prop., Phys. Prop., 6) cited from abstract

[2002Don] Dong-Feng, W., Bu-Xi, K., Bao-Hong, T., Ping, L., Jin-Lang, H., Dong-Mei, Zh., “Effect ofAging Treatment on Microhardness and Electrical Conductivity of Cu-Fe-P Alloy” (in Chi-nese), J. Luoyang Inst. Technol., China, 23(3), 10–12 (2002) (Morphology, Experimental,Electr. Prop., Mechan. Prop., 6) cited from abstract

[2002Per1] Perrot, P., Batista, S., Xing, X., “Cu-P (Copper-Phosphorus)”, MSIT Binary Evaluation Pro-gram, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International ServicesGmbH, Stuttgart, Document ID: 20.16300.1.20, (2002) (Crys. Structure, Phase Diagram,Phase Relations, Thermodyn., Assessment, Phys. Prop., #, 7)

[2002Per2] Perrot, P., Batista, S., Xing, X., “Fe-P (Iron-Phosphorus)”, MSIT Binary Evaluation Program,in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International ServicesGmbH, Stuttgart, Document ID: 20.16107.1.20, (2002) (Crys. Structure, Phase Diagram,Phase Relations, Thermodyn., Assessment, Phys. Prop., #, 23)

[2004Kai] Kaida, Y., Hasegawa, M., Kikuchi, Y., Wakimoto, K., Iwase, M., “Activities of Phosphorus inCopper-Iron Liquid Alloys Saturated with Copper-Iron Solid Solutions”, Steel Res., 75 (6),393–398 (2004) (Phase Diagram, Thermodyn., Experimental, *, 12)

[2006Lim] Lim, C.Y., Han, S.Z., Lee, S.H., “Formation of Nano-Sized Grains in Cu and Cu-Fe-PAlloysby Accumulative Roll Bonding Process”, Met. Mater.-Int., 12(3), 225–230 (2006) (Morphol-ogy, Experimental, Mechan. Prop., 10) cited from abstract

[2006Lu] Lu, D.-P., Wang, J., Zeng, W.-J., Liu, Y., Lu, L., Sun, B.-D., “Study on High-Strength andHigh-Conductivity Cu-Fe-PAlloys”,Mater. Science Eng. A., 421, 254–259 (2006) (Morphol-ogy, Experimental, Electr. Prop., Mechan. Prop., 16)

[2007Tur] Turchanin, M., Agraval, P., “Cu-Fe (Copper-Iron)”, MSIT Binary Evaluation Program, inMSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH,Stuttgart, Document ID: 20.11107.1 (2007) (Crys. Structure, Phase Diagram, Phase Rela-tions, Thermodyn., Assessment, Phys. Prop., #, 36)



22 Cu–Fe–P


Copper – Iron – Platinum

Vasyl Tomashik, Andy Watson

Introduction

A diagram of primary crystallization in the ternary system Cu-Fe-Pt was proposed for the first time by[1909Jae], but the earliest experimental studies were by [1941Nem] and [1943Nem2]. They used thermalanalysis, optical microscopy hardness measurement, electrical resistivity and temperature coefficient deter-mination. The thermoelectric power with respect to Pt, the extension and the tensile strength of the alloyswere also determined.Four isopleths were produced which suggested that all Cu-Fe-Pt alloys form solid solutions at high tempera-tures except for Pt poor alloys with less than 20 at.% Pt. The compound FePt was found to separate out ofthe solid solution and penetrates deep into the ternary system with a suggestion of the formation of a ternarycompound FePt2Cu at about 1200°C. This ternary compound was seen on investigation of the section Pt-(50at.% Cu + 50 at.% Fe), forming an unbroken series of solid solutions with FePt and limited solutions withPtCu [1941Nem, 1943Nem2]. These data were included in the reviews [1951Sol], [1969Gue] and[1979Cha], but as stated in [1979Cha], by that time, the existence of the ternary compound hadn’t been con-firmed.Interest in this system and the proposed compound continued [1973Cab]. Tulameenite is a naturally occur-ring deposit of the composition FePt2Cu adding weight to the hypothesis that this ternary compound exists[1973Cab, 1991Cab, 2002Tol, 2002Coo, 2002Gar, 2004Dis]. [1973Cab] conducted an extensive XRD,EPMA and optical microscopical study of tulameenite, but it was not until the work of [1986Sha] thatthe true nature of the material came to light. Through XRD, EPMA, TEM and DTA studies, they determinedthe relationship between FePt2Cu and the rest of the ternary system. Arc-melted samples were homogenizedat 1200°C for 3 months before quenching into water, followed by a further heat treatment at 1200°C for1 week, followed by slow cooling (10°C·d–1) to 1000°C where they were held for a further 3 months. Somesamples were cooled to 600°C and held for 6 months before quenching. They postulated that the singlephase region of FePt2Cu extends all the way to FePt at 1000°C and the FePt2Cu composition lies in atwo phase field at 600°C. This would suggest that the ternary compound is in fact an extension of the binarycompound. Later work [1986Sha] confirmed that this is the case with the binary and “ternary” compoundsbeing isomorphous.Little thermodynamic study has taken place on this system. [1989Par] measured the activity of Cu alongcomposition lines of constant Fe/Pt ratios by Knudsen cell effusion. [2004Wil] used electronic structure cal-culations to calculate the Gibbs energy of mixing along the Fe1–xCuxPt section, suggesting phase separationwith a critical temperature of 450°C.Experimental investigations of phase equilibria and thermodynamics are listed in Table 1.

Binary Systems

Binary systems Cu-Fe, Cu-Pt and Fe-Pt are accepted from [2006Tur], [2006Kuz] and [Mas2], respectively.

Solid Phases

The L10-FePt ordered phase extends into the ternary system by just over 30 at.% Cu at 1000°C and includesthe composition FePt2Cu. Thus, there is no ternary compound in this system. The ordering temperature ofthe L10 phase falls with increasing Cu content [2002Mae]. At the FePt2Cu composition, the ordering tem-perature falls to 1150°C [1986Sha]. It proved to be impossible to retain the high-temperature disordered fccstructure by quenching [1986Sha].In rapidly cooled polycrystalline material the grains are subdivided into twin-related domains with {101}habit to compensate for the stress set up by the tetragonal distortion. Colonies of small twins are found

Cu–Fe–Pt 1



within larger twins on three levels of scale, and it is suggested that these smaller twins are formed by suc-cessive re-ordering of a previously ordered matrix during cooling [1986Sha]. Cu substitutes for Fe in theFePt lattice [1986Sha, 2004Wil]. Early crystallographic work indicates 4 atoms per unit cell in the structureof the L10 compound, but later work by [1990Bay] did not reveal any additional reflections to indicate Feand Cu ordering. No information is available on the solubility of Cu in either Fe3Pt or FePt3.PtCu dissolves about 4.3 mass% Fe [1979Cha], PtCu3 also dissolves a small quantity of iron [1943Nem1].Crystallographic data for solid phases in the system are given in Table 2.

Isothermal Sections

Isothermal sections of the Cu-Fe-Pt ternary system at 1000 and 600°C, based on the work of [1985Sha] areshown in Figs. 1 and 2. At 1000°C there is a single phase region based on FePt extending far into the ternarysystem. Modification to the section presented in [1986Sha] is necessary, however, in order to take intoaccount equilibria involving FePt3. At 600°C, equilibria involving the PtCu and PtCu3 binary phases arealso included, although dissolution of the third element in these compounds is fairly small [1979Cha,1941Nem].


Two vertical sections CuPt - FePt and Pt - 50Cu50Fe (at.%) were constructed by [1941Nem, 1943Nem2]and included in the review [1969Gue]. However, there are inconsistencies between these sections and thebinary systems, particularly Cu-Pt. In fact, their measurements of the CuPt-FePt section disagree with theirown determination of the liquidus and solidus of the Cu-Pt binary. For these reasons, the sections have notbeen included here. Electronic structure calculations were used by [2004Wil] to calculate the Gibbs freeenergy of mixing of disordered Fe1–xPtCux alloys. This enabled the calculation of a temperature-composi-tion section revealing phase separation with a critical temperature of approximately 177°C. The section isgiven in Fig. 3. It should be noted that this behavior has yet to be seen experimentally.

Thermodynamics

The partial pressures of Cu in the system Cu-Fe-Pt in the temperature range 1240 to 1360°C have been mea-sured by the Knudsen effusion technique and the thermodynamic properties of this system at 1300°C havebeen derived [1989Par]. The activities of Fe in solid solutions at 1300°C were calculated by Gibbs-Duhemintegration of the Cu activities. The experimental alloys were prepared from Cu (99.999 mass%), Fe(99.999 mass%) and Pt (99.99 mass%) by induction melting in an alumina crucible under an Ar atmosphere.The alloy buttons were then homogenized in a H2 atmosphere for 5 to 30 days at 900 to 1300°C.The iso-aCu lines and the iso-aFe lines in the Cu-Fe-Pt system are shown in Figs. 4 and 5, respectively (thedashed lines in these figures are the estimated solidus and liquidus lines at 1300°C).


The alloys of the Cu-Fe-Pt enriched by Pt have practical interest as materials for electrical contacts[1943Nem1]. Ordered alloys of this system are good candidate materials for ultrahigh recording densityin terms of both high magnetic anisotropy energy and preparation without high-temperature treatment[2004Kai]. According to the data of [2004Wil], the magnetocrystalline anisotropy falls smoothly withincreasing Cu content, both for small and large additions of Cu. The calculations point to Cu as a promisingmaterial with which it is possible to reduce the prohibitively high ordering temperature of L10 FePt whilenot adversely affecting the magnetic properties. Furthermore, experimental data on the lower bound forthe ordering temperature of the L10 phase in FePt alloyed with Cu may aid in the future commercial viabilityof this material.The catalytic activity of Cu-Fe-Pt ternary alloy was investigated using an electrochemical method in a poly-mer electrolyte fuel cell by [2000Shi]. It was established that the electrode prepared using a Cu-Fe-Pt alloycatalyst showed higher cell performance than unalloyed Pt.Brief detals of some studies of materials properties in Cu-Fe-Pt alloys are given in Table 3.

2 Cu–Fe–Pt


Miscellaneous

The ‘compound’ FePt2Cu with minor substitution of Pd, Rh, Ir and Ni is a naturally occurring minor con-stituent of platinum ores and placers (mineral tulameenite) [1973Cab, 1976Ura, 1985Sha, 1990Bay,1991Cab, 2000Joh, 2002Coo, 2002Tol, 2004Dis]. An alloy with the composition FePt2Cu commonlyoccurs as prismatic and needle-shaped crystals at the outer margin of altaite (PbTe) [2002Coo]. This mineralis ferromagnetic [1973Cab]. According to [1973Cab], tulameenite has a range of composition with minoramounts of Ir replacing Pt and with Ni replacing Cu and Fe. Minor amounts of Sb was also detected in somegrains. Studies along the FePt2Cu-FePt join by [1985Sha] indicate that the maximum substitution of Fe forCu in tulameenite is up to between Fe1.68 Pt2Cu0.32 and Fe1.28 Pt2Cu0.72. Comparison of the experimentalresults with data for the tulameenite mineral suggests that some observed compositions may be metastablypreserved [1985Sha].Tulameenite is white in reflected light in oil and air [1973Cab]. No bireflectance could be observed and themineral is very weakly anisotropic.The Cu-Fe-Pt films show a high coercivity (HC) of around 5 kOe at 300°C, whereas the HC of Fe-Pt films isstill several hundred oersted at this temperature and starts increase to around 4.5 kOe at 400°C [2002Mae,2005Che]. Cu-Fe-Pt films show a large HC value at around 15 at.% Cu.The Mössbauer spectra of alloys of 0.1 and 1 at.% 57Fe in Cu with 5 at.% Pt have been obtained by[1971Win]. Analysis shows electric field gradients and changes in the isomer shift with Pt neighbors thatcan be explained by the increased density of states due to the virtual bound state on the Pt impurity.

Table 1. Investigations of the Cu-Fe-Pt Phase Relations, Structures and Thermodynamics


[1941Nem,1943Nem2]

Thermal analysis, hardness measurements,electrical resistance

Isopleths Fe0.5Pt0.5-Pt0.5Cu0.5,Pt-Fe0.5Cu0.5

[1973Cab] XRD, EPMA Pt2FeCu composition and alloys of thePtFe- Pt2FeCu section

[1985Sha] XRD, TEM, EPMA, DTA Equilibria involving Pt2FeCu

[1986Sha] XRD, TEM, EPMA, DTA Crystal structure and morphology ofPt2FeCu

[1989Par] Weight loss Knudsen cell Activity of Cu in ternary alloys from 1239- 1360°C

[2002Mae] XRD Structure of PtFe with added Cu

[2004Wil] Total energy calculations - Korringa-Kohn-Rostoker method

Phase separation in the Fe1–xCuxPt section





Comments/References

(εFe) hP2P63/mmcMg

a = 246.8c = 396.0


(continued)

Cu–Fe–Pt 3






Comments/References

(δFe)1538 - 1394

cI2Im�3mW

a = 293.15 [Mas2]

(γFexPt(1–x–y)Cuy)

(γFe)1394 - 912(Pt)< 1769.0(Cu)< 1084.62

cF4Fm�3mCu

a = 361.46 to 392.36a = 364.67 to 392.36

a = 364.67

a = 392.36

a = 361.46

x = 0, [2006Kuz]y = 0,

at 915°C[V-C2, Mas2]at 25°C [Mas2]

at 25°C [Mas2];melting point[1994Sub]

(αFe)< 912

cI2Im�3mW

a = 286.65 at 25°C [Mas2]

PtCu< 816

hR32R�3mCuPt

a = 1071.3c = 1319.2

[Mas2, V-C2]

PtCu3≲ 735

cP4Pm�3mAuCu3

a = 368.8 at 25 at.% Pt, 250°C [2006Kuz]

1-D LPS≲ 650

tP28P4mmCu3Pd

? [2006Kuz]; most probably - notunique phase but a set of closelyrelated phases

Fe3Pt≲ 820

cP4Fm�3mAuCu3

a = 373.0 [Mas2], [V-C2]

FePt≲ 1300

FePt2Cu< 1200

tP2P4/mmmAuCutP4P4/mmmAuCu

a = 384.1

a = 389.1 ± 0.2c = 357.7 ± 0.2a = 388.5 ± 0.1c = 358.8 ± 0.1a = 389.5c = 359.5a = 380.5 ± 0.1c = 359.5 ± 0.1a = 390.18 ± 0.07c = 358.45 ± 0.13

[Mas2][V-C2]

[1973Cab, V-C2]; mineral tulameenite[1973Cab]; synthetic compound[1986Sha, V-C2]; annealed at 1200°Cfor 3 months[1985Sha]

[1991Cab]

[1990Bay]

(continued)

4 Cu–Fe–Pt





Comments/References

tP2P4/mmm?c**??

a = 274.77 ± 0.04c = 358.70 ± 0.03

a = 379.2a = 385.0

[1985Sha]

[2000Shi]

FePt3≲ 1350

cP4Pm�3mAuCu3

a = 387.2 [V-C2]

Table 3. Investigations of the Cu-Fe-Pt Materials Properties


[1941Nem,1943Nem2]

Hardness measurements, tensile and ductilitymeasurement, electrical resistance,thermopower.

Mechanical and electrical properties withrespect to composition and temperature.

[1971Win] Mössbauer studies Spectra and hyperfine field distributionsof 1.4, 1 and 0.2 at.% Fe in Cu-5Pt

[1973Cab] Ore microscopy, hardness measurement. Reflectance, microhardness

[2002Mae] Vibrating sample magnetometry Magnetic properties of PtFe with respectto Cu addition

[2004Wil] Total energy calculations Magneto-crystalline anisotropy as afunction of Cu in Fe1–xCuxPt

[2005Che] Vibrating sample magnetometer Coercivity change diffusion of Cu intoPtFe thin film.

Cu–Fe–Pt 5



Fig. 1. Cu-Fe-Pt. Isothermal section at 1000°C

6 Cu–Fe–Pt


Fig. 2. Cu-Fe-Pt. Isothermal section at 600°C

Cu–Fe–Pt 7



Fig. 3. Cu-Fe-Pt. Calculated vertical section for FePt – PtCu

8 Cu–Fe–Pt


Fig. 4. Cu-Fe-Pt. Isoactivity lines of Cu at 1300°C

Cu–Fe–Pt 9



Fig. 5. Cu-Fe-Pt. Isoactivity lines of Fe at 1300°C

10 Cu–Fe–Pt


References[1909Jae] Jänecke, E., “Ternary Alloys of Cu, Ag, Au; Cr, Mn; Fe, Co, Ni; Pd, Pt Metals” (in German),

Z. Phys. Chem., 67, 668–688 (1909) (Experimental, Phase Diagram, 29)[1941Nem] Nemilov, V.A., Rudnickij, A.A., “Investigation of the Platinum - Iron - Copper Ternary Sys-

tem” (in Russian), Izv. Sekt. Fiz.-Khim. Anal., 14, 263–281 (1941) (Phase Diagram, Experi-mental, Mechan. Prop., Electr. Prop., 15)

[1943Nem1] Nemilov, V.A., “Alloys of Pt and Pd and their Using” (in Russian), Izv. Sekt. Platiny, 19,21–44 (1943) (Phase Diagram, Review, 43)

[1943Nem2] Nemilov, V.A., “Investigation of the Ternary Solid Solutions of Alloys of Pt and Pd” (in Rus-sian), Izv. Sekt. Fiz.-Khim. Anal., 16(1), 167–183 (1943) (Phase Diagram, Experimental, 24)

[1951Sol] Solodennikov, A.I., “The Dependences Composition - Property for Ternary Metal Systems”(in Russian), Uchen. Zap., Kirov. Gos. Pedagog. Inst., (6), 3–14 (1951) (Review, Phase Dia-gram, 18)

[1969Gue] Guertler, W., Guertler, M., Anastasiadias, E., “Copper - Iron - Platinum”, A Comp. of Const.Ternary Diagr. Met. Systems, Isr. Pro. Sci. Tr., Jerusalem, 579–583 (1969) (Review, PhaseDiagram, 1)

[1971Win] Window, B., “Mössbauer Studies of Iron in Copper Alloys”, J. Phys. F: Metal Phys., 1,533–538 (1971) (Experimental, Crys. Structure, 20)

[1973Cab] Cabri, L.G., Owens, D.R., LaFlamme, J.H.G., “Tulameenite, a New Platinum-Iron-CopperMineral form Placers in the Tulameen River Area, British Columbia”, Can. Mineral., 12,21–25 (1973) (Crys. Structure, Experimental, Optical Prop., 9)

[1976Ura] Urashima, Y., Wakabayashi, T., Masaki, T., “Osmian Ruthenium and Platinum Alloys fromHorokanai, Hokkaido, Japan” (in Japanese), Kagoschima Dai. Rika Hokoku, 25(9),165–171 (1976) (Experimental, Phase Diagram, 11)

[1979Cha] Chang, Y.A., Neumann, J.P., Mikula, A., Goldberg, D., “Cu-Fe-Pt”, in “INCRA MonographSeries 6 Phase Diagrams and Thermodynamic Properties of Ternary Copper-Metal Sys-tems”, NSRD, Washington, 490–491 (1979) (Review, Phase Diagram, 5)

[1985Sha] Shahmiri, M., Murphy, S., Vaughan, D.J., “Structural and Phase Equilibria Studies in theSystem Pt-Fe-Cu and the Occurrence of Tulameenite (Pt2FeCu)”, Mineral. Mag., 49(353),547–554 (1985) (Crys. Structure, Phase Diagram, Experimental, #, 14)

[1986Sha] Shahmiri, M., Vaughan, D.J., Murphy, S., “Observation of Twin-Related Order Domains inCuFePt2”, Phys. Status Solidi A, 95(1), 63–72 (1986) (Crys. Structure, Experimental, Mor-phology, 16)

[1989Par] Park, Y.-G., Gaskell, D.R., “The Thermodynamic Activities of Copper and Iron in the Sys-tem Copper-Iron-Platinum at 1300°C”, Metall. Trans. B, 20B(4), 127–135 (1989) (PhaseDiagram, Experimental, Thermodyn., 30)

[1990Bay] Bayliss, P., “Revised Unit-Cell Dimensions, Space Group, and Chemical Formula of SomeMetallic Minerals”, Can. Mineral., 28(4), 751–755 (1990) (Crys. Structure, Review, 32)

[1991Cab] Cabri, L.J., Genkin, A.D., “Re-Examination of Pt Alloys From Lode and Placer Deposits,Urals”, Can. Mineral., 29, 419–425 (1991) (Experimental, Crys. Structure, 16)

[1994Sub] Subramanian, P.R., “Cu (Copper)”, in “Phase Diagrams of Binary Copper Alloys”, P.R. Sub-ramanian, D.J. Chakrabarti and D.E. Laughlin (Eds.), ASM International, Materials Park,OH, 1–3 (1994) (Crys. Structure, Thermodyn., Review, 16)

[2000Joh] Johan, Z., Slansky, E., Kelly, D.A., “Platinum Nuggets from the Kompiam Area, Enga Pro-vince, Papua New Guinea: Evidence for an Alaskan-Type Complex”, Mineral. Petrol., 68,159-176 (Phase Diagram, Review, 42)

[2000Shi] Shim, J., Yoo, D.-Y., Lee, J.-S., “Characteristics for Electrocatalytic Properties and Hydrogen-Oxygen Adsorption of Platinum Ternary Alloy Catalysts in Polymer Electrolyte Fuel Cell”,Electrochim. Acta, 45, 1943–1951 (2000) (Experimental, Catalysis, Crys. Structure, 42)

[2002Coo] Cook, N.J., Ciobanu, C.L., Merkle, R.K.W., Bernhardt, H.-J., “Sobolevskite, Taimyrite, andPt2CuFe (Tulameenite?) in Complex Massive Talnakhite Ore, Noril’sk Orefield, Russia”,Can. Mineral., 40, 329–340 (2002) (Electronic Structure, Experimental, 61)

Cu–Fe–Pt 11



[2002Gar] Garuti, G., Pushkarev, E.V., Zaccarini, F., “Composition and Paragenesis of Pt Alloys fromChromitites of the Uralian-Alaskan-Type Kytlym and Uktus Complexes, Northern and Cen-tral Urals, Russia”, Can. Mineral., 40, 1127–1146 (2002) (Experimental, Morphology, 20)

[2002Mae] Maeda, T., Kikitsu, A., Kai, T., Nagase, T., Aikawa, H., Akiyama, J., “Effect of Added Cuon Disorder-Order Transformation of L1(0)-FePt”, IEEE Trans. Magn., 38(5), 2796–2798(2002) (Phase Diagram, Crys. Structure, Experimental, Magn. Prop., 12)

[2002Tol] Tolstykh, N.D., Foley, J.Y., Sidorov, E.G., Laajoki, K.V.O., “Composition of the Platinum-Group Minerals in the Salmon River Placer Deposit, Goodnews Bay, Alaska”, Can.Mineral., 40, 463–471 (2002) (Experimental, Phase Diagram, 35)

[2004Dis] Distler, V.V., Yudovskaya, M.A., Mitrofanov, G.L., Prokof´ev, V.Yu., Lishnevskii, E.N.,“Geology, Composition, and Genesis of the Sukhoi Log Noble Metals Deposit, Russia”,Ore Geol. Rev., 24, 7–44 (2004) (Review, Phase Diagram, 51)

[2004Kai] Kai, T., Maeda, T., Kikitsu, A., Akiyama, J., Nagase, T., Kishi, T., “Magnetic and ElectronicStructures of FePtCu Ternary Ordered Alloy”, J. Appl. Phys., 95(2), 609–612 (2004) (Cal-culation, Crys. Structure, Electronic Structure, Magn. Prop., 16)

[2004Wil] Willoughby, S.D., “Electronic and Magnetic Properties of Fe1–xCuxPt”, J. Appl. Phys.,95(11), 6586–6588 (2004) (Calculation, Crys. Structure, Electronic Structure, Magn.Prop., 14)

[2005Che] Chen, S.K.; Yuan, F.T.; Shiao, S.N., “Magnetic Property Modification of L10 FePt ThinFilms by Interfacial Diffusion of Cu and Au Overlayers”, IEEE Trans. Magn., 41(2),921–923 (2005) (Crys. Structure, Experimental, Magn. Prop., Transport Phenomena, 13)

[2006Kuz] Kuznetsov, V., “Cu-Pt (Copper - Platinum)”, MSIT Binary Evaluation Program, in MSITWorkplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH,Stuttgart; to be published, (2006) (Crys. Structure, Phase Diagram, Assessment, 8)

[2006Tur] Turchanin, M., Agraval, P., “Cu-Fe (Copper - Iron)”, MSIT Binary Evaluation Program, inMSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services,GmbH, Stuttgart; to be published, (2006) (Crys. Structure, Phase Diagram, Assessment, 36)



12 Cu–Fe–Pt


Selected Systems from Co-Fe-Si to Cu-Fe-Pt Landolt-Bornstein

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Transcript of Selected Systems from Co-Fe-Si to Cu-Fe-Pt Landolt-Bornstein