CHARACTERIZATION OF 10MOL% Sc2O3-1MOL% CeO2-ZrO2 CERAMICS
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Transcript of CHARACTERIZATION OF 10MOL% Sc2O3-1MOL% CeO2-ZrO2 CERAMICS
CHARACTERIZATION OF 10MOL% Sc2O3-1MOL% CeO2-ZrO2 CERAMICS AS ELECTROLYTE MATERIAL FOR LOWER TEMPERATURE
SOLID OXIDE FUEL CELLS
by
Devendra Ray
A thesis submitted to the graduate faculty in partial fulfillment of the requirements for the degree of
MASTER OF SCIENCE
Department: Mechanical & Chemical Engineering Major: Mechanical Engineering Major Professor: Dr. Jag Sankar
North Carolina A&T State University Greensboro, North Carolina
2007
iii
DEDICATION
First and foremost, I would like to thank my Lord Shiva for giving me the
strength and passion to arrive at this point and for putting the right people in my life at
the right time to help guide me. I dedicate this thesis to my parents, Ramkailash Ray,
father, and late mother, Ramshakhi Devi Ray, for their love, encouragement and the
many sacrifices they made that have enabled me to arrive at this point of my life. I would
also like to dedicate this thesis to my wife, Sima Ray. Finally, I would like to thank all
my family and friends who have encouraged and supported me to reach this milestone in
my life.
iv
BIOGRAPHICAL SKETCH
Devendra Ray was born in Nepal and immigrated to the United States of America
(USA) on September 7, 2004. In 2002, he graduated from Kathmandu University, Nepal
with an undergraduate degree in Mechanical Engineering. He is a candidate for the
Master of Science degree in Mechanical Engineering at North Carolina Agricultural and
Technical State University. He has performed research on solid oxide fuel cells (SOFC)
electrolyte materials.
v
ACKNOWLEDGEMENTS
I would like to begin by thanking my major advisor, Dr. Jag Sankar, and senior
research scientist/co-advisor, Dr. Sergey Yarmolenko, for their consistent guidance and
support over the past two years. I significantly benefited from their knowledge and
valuable experience in the field of material science and its characterization. I would also
like to thank all my friends and the staff of the Center for Advanced Materials and Smart
Structures (CAMSS). I would also like to express my thanks to Dr. Z. Xu and Dr. M. T.
Saad for their suggestions. I am very grateful to my family for their support, sacrifice,
and encouragement in assisting me to achieve this goal.
vi
TABLE OF CONTENTS
LIST OF FIGURES ........................................................................................................... ix
LIST OF TABLES............................................................................................................ xii
ABSTRACT..................................................................................................................... xiii
CHAPTER 1. INTRODUCTION ........................................................................................1
1.1 Research Objective ..................................................................................................2
CHAPTER 2. LITERATURE REVIEW .............................................................................3
2.1 Background.............................................................................................................3
2.2 Function of Fuel Cells.............................................................................................3
2.3 History of Fuel Cells...............................................................................................4
2.4 Types of Fuel Cells .................................................................................................5
2.5 Basic Components of SOFC ...................................................................................7
2.5.1 Electrolyte ......................................................................................................8
2.5.2 Electrolyte for Solid Oxide Fuel Cells...........................................................8
2.6 Introduction of Zirconia........................................................................................10
2.6.1 Crystallographic Structure ...........................................................................10
2.6.2 Application of Scandia Stabilized Zirconia (ScSZ).....................................12
2.6.3 Different Electrolyte Materials ....................................................................14
CHAPTER 3. EXPERIMENTAL METHODS AND CHARACTERIZATION ..............18
3.1 Materials ...............................................................................................................18
3.2 Fabrication of Ceramic .........................................................................................18
vii
3.2.1 Powder Compaction.....................................................................................18
3.2.2 Sintering.......................................................................................................19
3.3 Sample Preparation ...............................................................................................20
3.4 Study of Mechanical Properties ............................................................................22
3.4.1 Density and Porosity Analysis .....................................................................22
3.4.2 Indentation Techniques for Hardness and Fracture Toughness Analysis....24
3.4.2.1 Nanoindentation..................................................................................24
3.4.2.2 Microindentation.................................................................................31
3.5 Microstructure Characterization ...........................................................................36
3.6 XRD Measurements..............................................................................................43
3.6.1 Coefficient of Thermal Expansion (CTE)....................................................44
3.6.1.1 Temperature Calibration .....................................................................48
CHAPTER 4. RESULTS AND DISCUSSION.................................................................50
4.1 Charaterization of Praxair and DKKK Powders...................................................50
4.1.1 Theoretical Densities and Actual Composition ...........................................50
4.1.2 XRD of Powders ..........................................................................................54
4.2 Sintering Behavior of Powders .............................................................................57
4.2.1 Density Measurements of ScCeSZ and its Porosity ....................................58
4.2.2 Microstructure and Grain Size .....................................................................63
4.3 Phases and Stability of DKKK and Praxair Ceramics ..........................................67
4.3.1 Phase Analysis in DKKK-based Samples ...................................................68
4.3.2 Phase Transition in J1600 Sample ...............................................................69
viii
4.4 Coefficients of Thermal Expansion (CTE) ...........................................................71
4.5 Hardness and Fracture Toughness ........................................................................78
4.5.1 Microindenatation of Etched and Non-Etched Samples..............................78
4.5.2 Nanoindentation...........................................................................................87
CHAPTER 5. CONCLUSION AND FUTURE WORK ...................................................89
5.1 Conclusion .............................................................................................................89
5.2 Future Work ...........................................................................................................90
REFERENCES ..................................................................................................................91
ix
LIST OF FIGURES
FIGURE PAGE
2.1 Schematic representation of fuel cell ...........................................................................4
2.2 Phase diagram of Sc2O3-ZrO2 system .......................................................................11
3.1 Illustration of changes in the surface morphology after polishing on diamond paper with grits 3µm, 1 µm, and 0.5µm (from left to right) ......................................21 3.2 A schematic representation of the indentation process showing various parameters involved in the analysis ...........................................................................26
3.3 Nanoindentation load-displacement relationship.......................................................27
3.4 Typical images collected at 500x on optical microscope for hardness study ............33
3.5 (a) Schematic figure having crack and indent length; (b) Image collected at 1000x on optical microscope (Indentation load is 200g)..................................................35
3.6 Typical SEM images for grain size study ..................................................................38
3.7 Typical AFM working principle ................................................................................39
3.8 (a) Typical AFM images for grain size study; (b) Grain size and cavitations measurements............................................................................................................42 3.9 Temperature dependent shifting of XRD peaks positions ........................................45
3.10 Constant, linear and quadratic fitting with temperature vs. log (LT) (P1600- powder) .....................................................................................................................47 3.11 Instrumentally set temperature vs. residuals (P1600-powder)..................................48
3.12 Polynomial fitting of calibrated temperature vs. instrumentally set temperature ...............................................................................................................49
4.1 SEM micrographs of powders (a-Praxair, USA) and (b-DKKK, Japan)...................53
x
4.2 XRD pattern at different temperatures.......................................................................54
4.3 XRD pattern of as-received Praxair powder..............................................................55
4.4 XRD pattern of as-received DKKK powder..............................................................56
4.5 Fracture surfaces of samples P1600-(a) and J1600-(b)..............................................59 4.6 Grain structure of samples produced from Praxair powder sintered at temperatures 1500ºC (left) and 1600ºC (right) ..........................................................60 4.7 Grain structure of samples produced from DKKK powder sintered at temperatures 1400ºC (left), 1500ºC (middle) and 1600ºC (right) ...................................................60 4.8 SEM images of polished and low angle ion-milled surface of sample P1600 (a) and thermally etched surfaces of samples P1600 (b) and J1600(c).................................61 4.9 The dependence of density on sintering temperatures...............................................62
4.10 Porosity level vs. sintering temperature....................................................................62
4.11 SEM images of DKKK samples from left to right sinter correspondingly 1200-1600°C.............................................................................................................64 4.12 SEM images of Praxair samples from left to right sinter correspondingly 1500-1600°C.............................................................................................................64 4.13 Grain size vs. sintering temperatures ........................................................................65 4.14 SEM images collected after 1hr ion milling at 30 deg (left - J1600 right - P1600,
right - P1600) ............................................................................................................66 4.15 SEM images collected after ion milling (left-J1600 milled at 9-deg for 18hrs, right - P1600 milled at 5 deg for 20hrs)....................................................................67 4.16 SEM images of P1600 collected after ion milled at 5 deg for 20hrs (right - high porous area, left - less porous) ..........................................................................67 4.17 Phase content in the sample DKKK-1200 at different temperatures from XRD (Retvield multiphase analysis) ........................................................................69 4.18 Changes in XRD pattern of powdered sample J1600 heated from 300 to 500°C with step 10°C...........................................................................................................70
xi
4.19 Phase transition in DKKK ceramic sintered at 1600°C............................................70
4.20 Instrumentally set temperature vs. residuals (Praxair powder).................................74
4.21 Constant, linear and quadratic fitting with temperature vs. log (LT) (Praxair powder) .....................................................................................................................74
4.22 Instrumentally set temperature vs. residual(DKKK powder) ...................................75
4.23 Constant, linear and quadratic fitting with temperature vs. log (LT) (DKKK powder) .......................................................................................................75 4.24 Instrumentally set temperature vs. residuals (J1500 Powder) ..................................76
4.25 Constant, linear and quadratic fitting with temperature vs. log (LT) (J1500 Powder) .....................................................................................................................76 4.26 Instrumentally set temperature vs. residuals (J1500 ceramic) ..................................77
4.27 Constant, linear and quadratic fitting with temperature vs. log (LT) (J1500 ceramic).....................................................................................................................77 4.28 Hardness vs. sintering temperature plot for etched and non-etched sample.............78
4.29 Variation in fracture toughness value with temperature ...........................................79 4.30 (a) Effect of sintering temperature on microhardness and fracture toughness at the load 1000g; (b) Micrograph of typical Vickers indents and cracks at different loads (thermally etched J1600) ..................................................................81 4.31 ISE analysis using power law (a) and Nix-Gao (b) models......................................86
4.32 Nanoindentation Young’s modulus results...............................................................88
4.33 Nanoindentation hardness results .............................................................................88
xii
LIST OF TABLES
TABLE PAGE
2.1 Description of major fuel cell types.............................................................................6
2.2 Commercially available types of electrolyte materials..............................................14
3.1 Name of sample based on powder and sintering temperature ...................................19
3.2 Calculation of actual temperature from set temperature ...........................................49
4.1 Theoretical density from the actual composition of DKKK powder.........................51
4.2 Theoretical density from the actual composition of Praxair powder.........................52
4.3 Powder characterization.............................................................................................53
4.4 Density and porosity of sintered samples ..................................................................58
4.5 Grain size and in-grain cavities at different sintering temperatures ..........................65
4.6 Coefficients of thermal expansion for Praxair and DKKK samples in ceramic and powder forms ......................................................................................................73 4.7 Hardness (H) and fracture toughness (K1C) values at different sintering temperatures and loads for non-etched samples .......................................................82 4.8 Hardness (H) and fracture toughness (K1C) values at different sintering temperatures and loads for thermally etched samples ...............................................83 4.9 Comparison of hardness (H) and fracture toughness (K1C) of non-etched and etched samples. .................................................................................................84 4.10 ISE parameters for Praxair and DKKK-based samples ............................................86
4.11 Nanoindentation values……….................................................................................87
xiii
ABSTRACT
Ray, Devendra. CHARACTERIZATION OF 10MOL% Sc2O3-1MOL% CeO2-ZrO2 CERAMICS AS ELECTROLYTE MATERIAL FOR LOWER TEMPERATURE SOLID OXIDE FUEL CELLS (Major Professor: Dr. Jag Sankar), North Carolina Agricultural and Technical State University.
Material properties such as hardness, fracture toughness, densities and
coefficients of thermal expansion are very important parameters for stability and
reliability of solid oxide fuel cell (SOFC) operation. Ceria doped Sc2O3-CeO2- ZrO2
ceramics are also very promising materials as electrolyte in solid oxide fuel cells (SOFC)
due to their high oxygen conductivity in the 800-1000°C temperature range.
In this experimental work, sintering behavior of two commercial powders with the
nominal composition 10mol% Sc2O3-1mol% CeO2-ZrO2 which were produced by Praxair
surface technologies, USA and DKKK, Japan was studied. Ceramics were made by
sintering uniaxially pressed pallets at temperature range 1100-1600°C in air. Hardness
and fracture toughness of ceramics sintered at different temperatures were studied by
microindentation method. The density measurement was done using Archimedes’
principle. Porosity level in ceramics was estimated using actual and theoretical densities
calculated using lattice parameter a=5.09 Ǻ of FCC unit cell obtained from XRD data and
actual composition of powder. Hardness of most dense ceramics reaches 15GPa and
fracture toughness is in the range of 1.8-2.5MPa×m 0.5. These results are in good
agreement. Hardness of ceramics depends on indentation depth due to indentation size
effect (ISE). Indentation size effect parameters were analyzed using Power law model.
xiv
CTEs were studied. These results can be further used for the optimal design of SOFC
layered structures as well as for determination of their reliability and durability under
operational conditions.
1
CHAPTER 1
INTRODUCTION
It is well known that production and distribution of energy affects all sectors of
the global economy. The increasing industrialization of the world requires sustainable,
highly efficient energy production. Without major technological advancement, energy
production will impact the quality of human life on Earth. Dependency on fossil fuel will
not help to maintain the global economy in the long term due to the growing consumption
and limitation of fossil fuel. For this reason, the application of fuel cell technologies may
be one of the most important technological advancements of the future.
A fuel cell operating as a sort of continuously replenished battery provides an
alternative whereby electrical energy can be made available with smaller losses.
Electrical energy is derived from the reduction/oxidation reaction of a fuel such as
hydrogen and an oxidant such as oxygen. If the fuel is clean, the effluents are in principle
only water, heat and CO2. Fuel cell plants can be modular in design, and the energy
production can be adjusted to meet the actual demand, which is a convenient feature for a
power source in a technological society [1].
The efficiency of a fuel cell is dictated by factors regarding the kinetics and
driving forces of electrochemical reactions. An electrolyte has a very important role
which is to achieve optimum efficiency. Only a few materials have survived as practical
solid oxide fuel cell (SOFC) electrolytes because many electrical, chemical, and
mechanical properties are required of a successful electrolyte. These factors can be
2
improved through the introduction of superior electrolyte materials and enhanced
processing techniques [2].
1.1 Research Objective
The objective of this research was to further the development of better electrolyte
material. This includes the study of mechanical and microstructure properties of two
powders received from Praxair surface technology, USA, and DKKK, Japan. Knowledge
and understanding of mechanical properties such as hardness (H), fracture toughness
(K1c), thermal properties and densities are very important for the design of electrolyte in
SOFC. It is equally important to know the microstructure, such as the grain size of two
powders at different sintering temperatures, which has an important role in ionic
conductivity. Samples were produced at the temperature range 1100-1600°C in air.
Specifically, the objectives can be outlined as follows:
• To study mechanical properties such as hardness, fracture toughness, and
densities of two powders having nominal composition 10mol%Sc2O3-
1mol%CeO2- ZrO2 at all sintering temperatures.
• To study the grain size measurements at different sintering temperatures of
the material.
• To study the coefficient of thermal expansion.
• To study phase stability.
• To study modulus of elasticity.
3
CHAPTER 2
LITERATURE REVIEW
2.1 Background
This chapter provides a brief review of solid oxide fuel cells (SOFC), what they
are and how they work. It also details their history, the types of fuel cells, working
principles and materials employed for SOFC production. This research work emphases
the scandia stabilized zirconia (ScSZ) electrolyte material which was selected for the
study, in particular, its various properties, research and developmental aspects.
2.2 Function of Fuel Cells
A fuel cell is a �factory� that takes fuel as input and produces electricity as output.
Furthermore, in terms of chemistry, it is an electrochemical device that directly converts
chemical energy from a reaction between a fuel and oxidant into electrical energy as long
as fuel is supplied. This is the key difference between a fuel cell and a battery. While
both rely on electrochemistry, a fuel cell is not consumed when it produces electricity. It
is really a factory, a shell, which transforms the chemical energy stored in a fuel into
electrical energy. The basic elements of a typical fuel cell, as depicted in Figure 2.1,
consists of electrolyte in intimate contact with a porous anode (negative electrode) and
porous cathode (positive electrode). The fuel and oxidant gases flow along the surface of
the anode and cathode, respectively, and then react electrochemically in the three-phase-
boundary region established at the gas-electrolyte electrode interface. A fuel cell can
4
theoretically produce electrical energy for as long as fuel and oxidant are fed to the
porous electrodes, but the degradation or malfunction of some of its components limits
the practical life span of all fuel cells.
Figure 2.1: Schematic representation of fuel cell
2.3 History of Fuel Cells
The history of the fuel cell dates back to 1839 when Sir William Grove [3] first
described its principle and demonstrated a fuel cell at room temperature using a liquid
electrolyte. In 1899, Nernst discovered the solid oxide electrolyte when using stabilized
zirconia in making filaments for electric glowers. In the middle of the 20th century the
development accelerated. Several types of fuel cells were developed in the race for
conquering space. In the eighties, focus on pollution and the demand for higher efficiency
in the exploitation of fossil resources initiated a new wave of fuel cell development. At
present, several types of fuel cells are approaching the consumer market within a limited
Anode
Electrolyte
Cathode
Fuel
Oxidant
External loadDirect current Exhaust gases and heat
e-
e-
5
number of years. The primary challenges are cost and durability, to be solved by
materials selection and design engineering [2].
2.4 Types of Fuel Cells
Currently, there are five major types of fuel cells, differentiated from one another
by their electrolyte material [1]:
1. Phosphoric acid fuel cell (PAFC)
2. Polymer electrolyte member fuel cell (PEMFC)
3. Alkaline fuel cell (AFC)
4. Molten carbonate fuel cell (MCFC)
5. Solid oxide fuel cell (SOFC).
While all five fuel types are based upon the same underlying electrochemical
principles, they operate at different temperature regimens, incorporate different materials,
and often differ in their fuel tolerance and performance characteristics, some of which are
listed in Table 2.1[1, 2, 3, and 52]. Among these, operation temperature, electrical
efficiency and demands for fuel composition are the defining prospects of each type.
Polymer electrolyte member fuel cells are well suited to transportation
applications because they provide a continuous electrical energy supply from fuel at high
levels of efficiency and power density. They also offer the the advantage of minimal
maintenance because there are no moving parts in the power generating stacks of the fuel
cell system.
6
T
able
2.1
. The
des
crip
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of m
ajor
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M
olte
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mob
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Cat
alys
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rovs
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Plat
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atin
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inum
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l com
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nts1
Cer
amic
bas
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Stai
nles
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arbo
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sed
Car
bon
base
d Fu
el
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patib
ility
1 H
2, C
H4
, CO
H
2, C
H4
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H2
H2,
Met
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App
licat
ions
2 >E
lect
ric u
tility
>T
rans
porta
tion
>Mili
tary
>Ele
ctric
util
ity
>Mili
tary
>S
pace
>E
lect
ric u
tility
>T
rans
porta
tion
>Ele
ctric
util
ity
>Por
tabl
e po
wer
>T
rans
porta
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Adv
anta
ges2
>Hig
h ef
ficie
ncy
>Fue
l fle
xibi
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>Can
use
a v
arie
ty
of c
atal
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>R
educ
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alk
alin
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ectro
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that
le
ads t
o hi
gh
perfo
rman
ce
>Can
use
impu
re
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fuel
>S
olid
ele
ctro
lyte
re
duce
s cor
rosio
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anag
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oble
ms
>Low
tem
pera
ture
Disa
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s2 >H
igh
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rate
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own
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ell
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>Hig
h te
mpe
ratu
re
enha
nce
corr
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d br
eakd
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of
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com
pone
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>Req
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s rem
oval
of
CO
2 fro
m fu
el
and
air s
tream
s, w
hich
is e
xpen
sive
>Req
uire
s pl
atin
um c
atal
yst
>Low
cur
rent
an
d po
wer
>Low
ope
ratin
g te
mpe
ratu
re
requ
ires e
xpen
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ca
taly
sts
Elec
trica
l Ef
ficie
ncy3
60%
60
%
40%
40
%
40%
7
The low temperature fuel cells (AFC, PAFC) also have a potential for propulsion
of cars, where a short heating time is needed and the efficiency has to be compared with
that of a combustion engine (~20%) given by the limitations of the Carnot cycle. In
miniature applications such as cellular phones, laptops and digital cameras, methanol is
reformed to hydrogen in the fuel cell itself (reformation reaction). For applications that
require higher output, such as cars, methanol can be reformed in a device that is located
outside the fuel cell. This reformation can be accomplished through steam reforming or
partial oxidation [2]. Recently invented micro fuel cells are smaller, lighter, cleaner,
simpler and less expensive than lithium ion batteries [2].
2.5 Basic Components of SOFC
There are different types of electrolytes that are commercially available; however,
more research and development is needed to characterize and explore new and promising
electrolytes for an intermediate temperature (IT) SOFC. The basic components of SOFC
are:
1. Cathode
2. Electrolyte
3. Anode
4. Interconnect
This thesis focuses on electrolyte material.
8
2.5.1 Electrolyte
The most effective electrolytes employed in fuel cells have ionic conductivities of
around 0.10 Ώ-1cm-1. The search for better electrolytes has led to the development of
three major candidate material classes for fuel cells: aqueous, polymer, and ceramic
electrolytes. Regardless of the class, however, any fuel cell electrolyte must meet the
following typical properties to be a better electrolyte:
• High ionic conductivity
• Low electronic conductivity
• High stability (in both oxidizing and reducing environments)
• Low fuel crossover
• Reasonable mechanical strength (if solid)
• Ease of manufacturability.
Except for the high conductivity requirement, the electrolyte stability requirement
is often the hardest to fulfill. It is difficult to find an electrolyte that is stable in both the
highly reducing environment of the anode and the highly oxidizing environment of the
cathode [1, 52].
2.5.2 Electrolyte for Solid Oxide Fuel Cells
The electrolyte is the central part of an SOFC. Within the electrolyte the oxygen
ions (O2-), which are reduced on the air electrode side (cathode), are transported and
react with, for example, hydrogen to form water on the fuel electrode side (anode).
Conversely, electrons (e-) are formed and while moving in the opposite direction, are
available for an outside current use. Nowadays, the most frequently used electrolyte
9
material is zirconia (ZrO2). Zirconia is at ambient conditions a poor ionic conductor. If
ZrO2 is heated up to temperatures above 2000°C, it becomes ionic conducting due to a
phase transformation from tetragonal to cubic structure. By adding stabilizing scandia
(ScSZ) the cubic structure is stabilized even at ambient conditions. Because of this
stabilization zirconia becomes a reasonable ionic conductor at SOFC operating
temperatures (750-1000°C) and can be used as electrolyte material. Ionic conduction
proceeds through oxygen vacancies due to insertion of a di- or trivalent element (Ca2+,
Y3+, Sc3+) instead of the tetravalent Zr4+. The lack of positive charge is balanced by free
oxygen lattice sites. Through these free sites oxygen can move through the cubic
structure.
Besides good ionic conductivity, gas tightness of the electrolyte is the most
important characteristic the material should have. If the gas tightness of the electrolyte is
insufficient, a reaction between oxygen (cathode side) and hydrogen (anode side) may
occur. Gas tightness is ensured by sintering the electrolyte at temperatures of
approximately 1400°C. Lower temperatures lead to inadequate gas tightness, but pure
yittria-stablized zirconia (YSZ) cannot be sintered to high densities below 1400°C. In
contrast to the high sintering temperature, which is necessary for the electrolyte, the
cathode tolerates only maximum temperatures of approximately 1200°C. If sintered at
higher values, the amount of triple phase boundaries reduces drastically due to enhanced
sintering. These two facts are the reason for one area of major research and development
on electrolyte materials, namely, is the reduction of the sintering temperature of the
electrolyte (goal ≤ 1300°C). Decreasing sintering temperatures can be reached by a)
10
using a nanosized starting material or b) the use of sintering additives. Additional
research and development is focused on the coating technologies for SOFC electrolytes
[4].
2.6 Introduction of Zirconia
Zirconia has become one of the most industrially important ceramic materials of
the present time. The traditional applications of ZrO2 and ZrO2-containing materials are
foundry sands and flours, refractory ceramic, and abrasives. Because of its high oxygen
ion conduction and high refractive index, it is also used in a wide range of newer
applications which include fuel cells, catalysts, oxygen sensors, and jewelry. Zirconias
have also been utilized in many mechanical applications. Along with high strength and
toughness, zirconia also possesses good hardness, wear resistance, and thermal shock
resistance. These properties have led to the use of zirconia-based components in a
number of engineering applications such as automobile engine parts, wire drawing dies
and cutting tools. The low thermal conductivity together with relatively high coefficiency
of thermal expansion makes zirconia a suitable material for thermal barrier coating on
metal components [6].
2.6.1 Crystallographic Structure
At atmospheric pressure, pure zirconia (ZrO2) has three crystalline polymorphs
with monoclinic, tetragonal and cubic structures. The monoclinic from is stable up to
1170°C when it transforms to the tetragonal modification which remains stable up to
2370°C; from 237°C to the melting point (2680°C), the cubic form is stable. The reverse
11
transformations take place on cooling. The phase transformations are martensitic in
nature. The monoclinic and tetragonal conversion is accompanied by a contraction in
volume of approximately 5% which can cause mechanical failure evident by cracking in
the ceramic. The phase transformation can be avoided and the high temperature cubic
phase can be stabilized at low temperature by substituting low-valency cations for the
zirconium (see Figure 2.2).
Figure 2.2: Phase diagram of Sc2O3-ZrO2 system
Cubic ZrO2 has the fluorite structure with the O2- ions arranged in simple cubic
packing and half the interstices in this lattice occupied by Zr4+ ions. The substitution of
12
lower-valency cations leads to O2- ion vacancies. The vacancies that stabilized the
structure also lead to high mobility in the oxygen sublattice and to behavior as a high-ion
conductor.
The elements that stabilized the cubic fluorite structure in zirconia include
scandium, lanthanides, yttrium, magnesium, calcium, manganese and indium. Scandium
has been found to give a material with a higher conductivity, which is particularly
valuable at a lower temperature. Scandia stabilized zirconia (ScSZ) is one of the solutions
to the problem of stabilizing the cubic form of zirconia [6].
2.6.2 Application of Scandia Stabilized Zirconia (ScSZ)
Fully stabilized zirconia based ceramic materials have several areas of application
such as an electrolyte in solid oxide fuel cells (SOFC), oxygen sensors and structural
applications due to their high oxygen conductivity in the 800-1000oC temperature range,
electrical and mechanical properties. Scandia stabilized zirconia (ScSZ) ceramics have
the highest oxygen conductivity among zirconia based materials [7] and, with the trend in
reducing the operation temperature of SOFCs to an intermediate temperature (IT) range
of 600-800ºC, ScSZ ceramics are very attractive materials for IT SOFCs and catalytic
membrane reactors. Important requirements for these applications are reproducibility of
the electrical and mechanical properties, high density and a low level of internal flaws.
However, most commercial ZrO2-based powders typically need high sintering
temperatures above 1400ºC in order to form dense ceramics. High sintering temperatures
create a variety of problems such as temperature induced degradation and shape change
of materials, interface reactions between zirconia-based ceramics and other components,
13
excessive grain growth as well as energy production costs. Lower processing
temperatures can limit grain growth to nanocrystalline size which has beneficial effect on
the improvement of electrical properties of solid electrolytes [8] due to specific grain
boundary conductivity. Recent studies of nanocrystalline scandia stabilized zirconia thin
films prepared by sol-gel using a polymer precursor solution showed that the electrical
transport of this material can also be enhanced [9]. Therefore the reduction of the
sintering temperature towards an SOFC�s operating temperature can improve mechanical
and electrical properties and make possible co-firing zirconia-based electrolytes with
other SOFC components and significantly lower production costs. Properties such as
hardness, fracture toughness and coefficients of thermal expansion (CTEs) are very
important parameters for stability and reliability of SOFC operation.
Extensive research has shown that a phase transition from cubic to rhombohedral
structure of 11mol%Sc2O3-89%ZrO2 occurs when the temperature decreases below
600°C [10, 11]. This phase transformation prevents the direct use of this ceramic in
SOFC applications. The aging process of scandia doped zirconia is also accompanied by
the formation of a less conductive rhombohedral phase resulting in degradation of ionic
conductivity of Sc2O3-ZrO2 electrolytes [12, 13]. Usually the cubic phase can be
stabilized over a wide temperature range by substitution of 1 mol% of scandia by another
stabilizer such as Yb2O3 [13], Y2O3 [10], Bi2O3 [14], Al2O3 [15], and so on. The addition
of a low amount of other oxides usually does not greatly deteriorate the sintering activity
and other properties of the resulting powder. It was reported recently that replacement of
1mol% Sc2O3 by CeO2 stabilized cubic phase led to only 50ºC increase in sintering
14
temperature [14]. Also it was found recently that scandia-ceria doped zirconia ceramics
have a very high ionic conductivity (0.151 S/cm at 800oC) [16].
In this experimental work the study of sintering behavior of two commercially
available powders with a nominal composition of 10mol%Sc2O3-1mol%CeO2-ZrO2
(ScCeSZ) that can be used to fabricate dense and stable ceramic electrolytes was carried
out. The crystal structure, the morphology and phase transitions of powders and ceramics
were also investigated. Mechanical properties, such as hardness and fracture toughness,
of the Sc2O3-CeO2-ZrO2 ceramics were reported. These results can be further used for the
optimal design of SOFC layered structures as well as for determination of their reliability
and durability under operational conditions.
2.6.3 Different Electrolyte Materials
Table 2.2 briefly describes all types of commercially available SOFC electrolyte
materials currently available in different forms for developing solid oxide fuel (SOFC)
[18].
Table 2.2. Commercially available types of electrolyte materials
No Name Form Description 1 Scandia
Stabilized Zirconia (10 mole %) Nanopowder
Nano Powder Scandia Stabilized Zirconia (Sc2O3)0.10(ZrO2)0.90 Nanoscale powder for use as sintering aid, catalyst support, and as a component for mixed conducting anodes and cathodes to enhance catalytic activity Surface area: >100 m2/g Particle size: softly agglomerated
15
Table 2.2. (Continued) No Name Form Description 2 Gadolinium
Doped Ceria (10% Gd) Nanopowder
Nanopowder Gadolinium Doped Ceria Gd0.10 Ce0.90 O2- Nanoscale powder for use as sintering aid, catalyst support, and as a component for mixed conducting anodes and cathodes to enhance catalytic activity. Surface area: >100 m2/g Primary crystallite size: 5-10 nanometers Secondary particle size: softly aggolomerated
3 Gadolinium Doped Ceria (10% Gd) Aqueous Suspension
Suspension Gadolinium Doped Ceria (10% Gd) aqueous suspension specially formulated suspension for colloidal deposition and spray coating methods. Suspension is stable for an extended period of time. Surface area: >20-40 m2/g (estimated) Particle size: 50-80 nm Solids loading: 10-15 volume %
4 Gadolinium Doped Ceria (10% Gd) Ceramic Powder
Ceramic Grade Powder
Gadolinium Doped Ceria Gd0.10 Ce0.90 O2- Powder suitable for tape casting, ink manufacture, pellet pressing and other non-aqueous manufacturing processes. Surface area: >5-8 m2/g Particle Size: d50-0.5 microns
5 Gadolinium Doped Ceria (20% Gd) Nanopowder
Nanopowder Gadolinium Doped Ceria Gd0.20 Ce0.80 O2- Nanoscale powder for use as sintering aid, catalyst support, and as a component for mixed conducting anodes and cathodes to enhance catalytic activity. Surface area:>100m2/g Primary crystallite size: 5-10 nanometers Secondary particle size: softly agglomerated
6 Gadolinium Doped Ceria (20% Gd) Ceramic Powder
Ceramic Grade Powder
Gadolinium Doped Ceria Gd0.20 Ce0.80 O2- Powder suitable for tape casting, ink manufacture, pellet pressing and other non-aqueous manufacturing processes. Surface area: >5-8 m2/g Particle Size: d50-0.5 microns
16
Table 2.2. (Continued) No Name Form Description 7 Samarium
Doped Ceria (15% Sm) Nanopowder
Nanopowder Samarium Doped Ceria Sm0.15 Ce0.85 O2- Nanoscale powder for use as sintering aid, catalyst support, and as a component for mixed conducting anodes and cathodes to enhance catalytic activity. Surface area: >100 m2/g Particle size: softly agglomerated
8 Samarium Doped Ceria (15% Sm) Ceramic Powder
Ceramic Grade Powder
Samarium Doped Ceria Sm0.15 Ce0.85 O2- Powder suitable for tape casting, ink manufacture, pellet pressing and other non- aqueous manufacturing processes. Surface area: >5-8 m2/g Particle Size: d50-0.5 microns
9 Yttria-Stabilized Zirconia 3 mole% Ceramic Powder
Ceramic Grade Powder
Ytrria Stabilized Zirconia (Y2O3)0.03 (ZrO2)0.97 Partially stabilized zirconia powder suitable for pellet pressing, injection molding and aqueous manufacturing processes. Can be calcined for use in non-aqueous applications. Surface area: >13-19 m2/g Particle Size: d50-0.5 microns
10 YSZ 8 mole% Engineered Coating
Suspension Yttria Stabilized Zirconia (8% Yttria) Aqueous suspension for colloidal deposition, dip coating and spray coating methods. Has engineered particle sizes to match shrinkage to typical anode materials. Can be engineered to match the shrinkage of cathode supports. Suspension is stable for an extended period of time. Surface area: Engineered Particle Size: Engineered Solid loading: 10-15 volume %
11 Yttria-Stabilized Zirconia 8 mole% Nanopowder
Nanopowder Ytrria Stabilized Zirconia (Y2O3)0.08 (ZrO2)0.92 Nanoscale powder for use as sintering aid, catalyst support, and as a component for mixed conducting anodes and cathodes to enhance catalytic activity. Surface area: >100 m2/g Particle size: softly agglomerated
17
Table 2.2. (Continued) No Name Form Description 12 Yttria-
Stabilized Zirconia 8 mole% Aqueous Suspens
Suspension
Yttria Stabilized Zirconia (8% Yttria) Aqueous Suspension specially formulated suspension for colloidal deposition and spray coating methods. Suspension is stable for an extended period of time. Surface area: 20-40 m2/g (estimated) Particle Size: 50-80 nm Solid loading: 10-15 volume %
13 Yttria-Stabilized Zirconia 8 mole% Ceramic Powder
Ceramic Grade Powder
Yttria Stabilized Zirconia (Y2O3)0.08 (ZrO2)0.92 Full stabilized powder suitable for tape casting, ink manufacture, pellet pressing and other non-aqueous manufacturing processes. Surface area: 4-8 m2/g Particle Size: d50-0.5 microns
18
CHAPTER 3
EXPERIMENTAL METHODS AND CHARACTERIZATION
3.1 Materials
Two different powders with nominal composition 10mol%Sc2O3-1mol%CeO2-
ZrO2 were studied in this research work:
• Praxair Surface Technology, Specialty Ceramics, USA
• Daiichi Kigenso Kagaku Kogyo Co., Ltd , Japan (DKKK, Japan)
3.2. Fabrication of Ceramic
3.2.1 Powder Compaction
Powder compaction was performed using the uniaxial pressing method. Uniaxial
pressing involves the compaction of powder into a rigid die by applying pressure in a
single axial direction through a rigid punch or piston. The presses were hydraulic and
needed manual operation. A 31mm steel die was used for the preparation of all samples.
The setup consisted of a steel die, a steel punch and two plugs. The powder was poured in
between the two plugs. The steel punch was placed on top of the second plug, and the
pressure was applied. The applied force was 15000N. The calculation is as follows:
)(
)(,PrAArea
FForceAppliedPessure =
4
,2dAArea ∏= = 1.1697 in2
19
Applies Force, F = 15000N = 3372.134 lb
1697.1
134.3372,Pr =Pessure
= 2882.92 Psi
3.2.2 Sintering
Sintering converts a compacted powder into a denser structure of crystallites
joined to one another by grain boundaries. Grain boundaries vary in thickness from about
100µm to over 1µm. Samples were numbered based on the company name for all
different temperatures as listed in Table 3.1. They may consist of crystalline or vitreous
second phases, or may simply be a disordered form of the major phase because of the
differing lattice orientations in neighboring grains. They impart strength to the body,
Table 3.1. Name of sample based on powder and sintering temperature
No of Samples Sintering Temperature (°C)
Name (J-Japan, P-Praxair)
1 1100 J1100 2 1200 J1200 3 1300 J1300 4 1400 J1400 5 1500 J1500 6 1600 J1600 7 1100 P1100 8 1200 P1200 9 1300 P1300
10 1400 P1400 11 1500 P1500 12 1600 P1600
20
typically in the range 70-350 MPa (cross-breaking). Grain boundaries are generally not as
dense as the crystals [6].
The sintering process of two uniaxially pressed powders was performed in a high
quality, high temperature Micropyretics Heaters International (MHI) furnace. The
sintering temperature was from 1100-1600°C in air. The sintering time was 2 hours and
heating and cooling rates were 10°/minutes.
3.3 Sample Preparation
Sample preparation is one of the most important steps during the various
experimental testing of ceramic materials. This process involves several steps to produce
a smooth surface for analysis. The main objective is to obtain an absolutely flat, highly
polished, and clean sample. This process was completed basically in two different steps
as follows:
• Polishing
• Etching.
Polishing
This is an integral part of most types of testing. The best and the fastest polishing
to produce flat and smooth surfaces was performed with the lapping tool typically used
for SEM, AFM and TEM sample preparation and polishing by diamond grinding paper
with a grain size from 30 µm to 0.5 µm. The following steps were taken during the
polishing of the samples:
• The ceramic was cut to size using a diamond saw.
21
• The sample was mounted on a lapping stage using wax and heat.
• Diamond paper was placed on a clean wet glass plate and air bubbles were
removed from beneath the paper.
• The lapping tool with the sample was slid under the self-weight of the lapping
tool on 30 micron diamond paper using water for lubrication to obtain a fully flat
surface.
• The sample was flipped over to begin final polishing on the other side
Figure 3.1 represents the polished surface obtained by different grit sized diamond
paper.
Figure 3.1: Illustration of changes in the surface morphology after polishing on diamond paper with grits 3µm, 1 µm, and .5µm (from left to right)
Etching
Etching is performed to make visible the grain boundaries for study of grain size
in material and granular microstructure. All samples were polished up to 1 micron on
diamond paper and were etched thermally in air for 30 minutes at 1300°C regardless of
22
the sintering temperature. A Thermcraft Incorporated Furnace model#TSL-2-0-6-1S-
J7409/1A was used for thermal annealing.
3.4 Study of Mechanical Properties
Mechanical properties like densities, fracture toughness and hardness were
experimentally studied for samples made at different sintering temperatures.
3.4.1 Density and Porosity Analysis
Density is the quotient of the mass and volume. The unit is g/cm3 which is
basically adopted in lab work. Density determination is frequently performed by
Archimedes� principle which states that every solid body immersed in a fluid apparently
loses weight by an amount equal to that of the fluid it displaces.
The densities of sintered specimens were determined here by the standard
Archimedes� method in acetone. For each sample 12 readings were taken. The apparatus
used for this purpose was Mettler Toledo of Model: AX105 Delta Range(R). The
procedures adopted to calculate density were as follows:
• A balance with all necessary setup (draft shield elements, bracket, equalizing
washer and platform, thermometer) for density measurements was prepared.
• The balance was initialized to zero.
• The dry weight of a sample was taken in air.
• The sample was immersed in acetone and the reading was taken after stabilization
of number displayed on a screen of the balance.
• For porous samples wet weight was taken.
23
The density ρ was calculated from the two weighings as follows
LLBAADensity ρρρρ +−−
= )(: 0
Lo
BAVVolumeρρ
α−−=:
where,
ρ = Density of sample
A = Weight of sample
B = Weight of sample in the auxiliary liquid
Lρ = Air density (0.0012 g/cm3)
0ρ = Density of the auxiliary liquid
α = Balance correction factor (0.99985), takes air buoyancy of the adjustment
weight into account [19].
Theoretical Density
Theoretical density is based on crystal structure and molar mass of material, the
formula used for theoretical density calculation is:
γβαρ
sin*sin*sin*)*(:
cbaNNMWeightAtomicDensitylTheoretica
A
=
where, M = Molar mass, N = Atomic Number, NA = Avogadro�s number, a, b, c are
lattice parameters, obtained from XRD experiments, Angle α, β, and γ = 90°, The
procedures and calculation are detailed in Chapter 4 [20].
24
3.4.2 Indentation Techniques for Hardness and Fracture Toughness Analysis
3.4.2.1 Nanonindentation
Indentation techniques are used as effective means of measuring the hardness of
bulk materials. Nanoindentation is a simple method consisting of pressing a material of
known mechanical properties such as hardness and elastic modulus into a material whose
properties are to be estimated. The material is subjected to load by pressing an indenter
through the material and the images of the indent impressions are captured. Hardness of
the material is defined as the mean pressure the material can support under load and is
calculated as H = P/A, where P is the applied load and A is the projected contact area of
the impression. Elastic modulus is a measure of the recovery of the material during
unloading and is an indicator of its elastic properties. Conventional indentation
techniques involve optical imaging of the indentations for calculation of the contact area.
Microindentation testing carried out using Rockwell and Vickers� hardness tests involve
high loads and are applicable for bulk materials. Nanoindentation, a recently developed
technique, involves indentation depths in the order of nanometers and loads of few
millinewtons.
Nanoindentation, also known as depth-sensing indentation, is a relatively new
testing technique that involves a significant expansion on the capabilities of conventional
hardness testing. Nanoindenters employ high-resolution instrumentation to continuously
control and monitor the loads and displacements of an indenter that is driven into and
withdrawn from the material being tested [22-28, 29]. Depending on the specifications of
the equipment, loads as small as 1 nN and displacements as small as 0.1nm (1 Å) can be
25
measured [30]. The first major contribution in indentation experiments was made by I.N.
Sneddon (Sneddon, 1965). Sneddon studied the relationships between the load,
displacement and contact area for different indenter geometries and derived with an
equation, P = αhm, where P is the indenter load, h is the elastic displacement of the
indenter and α and m are constants [31]. Values of the exponent m for some common
punch geometries are m=1 for flat cylinders, m=2 for paraboloids of revolution [22].
Nanoindentation technique has various advantages compared to traditional micro
indentation techniques. The load-displacement data from nanoindentation contains an
enormous amount of information and techniques are developed for characterizing a
variety of mechanical properties such as hardness, stiffness and elastic modulus. Methods
have also been developed for evaluating the yield stress and strain-hardening
characteristics of materials [32-34]. Damping and internal friction characteristics such as
storage and loss modulus, activation energy and stress exponent for creep are also being
evaluated from the nanoindentation data [35-37]. Fracture toughness is also being
estimated from the optical measurements of the cracks that have formed at the corners of
indent impressions made with sharp indenters [38-39].
Tabor used hardened spherical indenters to perform indentation tests on metals
[21]. Stillwell and Tabor noticed that at least in metals, the impression formed by the
spherical indenter is still spherical with a slightly larger radius than the indenter, and the
impression formed by a conical indenter is still conical with a large included tip angle.
Tabor showed that the shape of the entire unloading curve and the total amount of
recovered displacement could be accurately related to the elastic modulus and the size of
26
the impression for both spherical and conical indenters [40]. It is also notable that the
diameter of the contact impression on the surface formed by the conical indenter does not
recover during unloading and only the depth recovers. The indenter is loaded and
unloaded a few times before the load displacement behavior becomes perfectly
reversible. The plasticity can be dealt with by taking into account the shape of the
perturbed surface in the analysis of the elastic unloading curve. Therefore, the shape of
the unloading curve and the recovered displacement characterizes the elastic modulus of
the material.
Hardness and elastic modulus are the most frequently measured mechanical
properties using nanoindentation. A schematic of the indentation process for an
axisymmetrical indenter of arbitrary profile is shown in Figure 3.2. As the indenter is
pressed into the material, both elastic and plastic deformation processes occur, creating
an impression that conforms to the shape of the indenter of some contact depth, hc. The
Figure 3.2: A schematic representation of the indentation process showing various parameters involved in the analysis.
27
radius of the contact circle is a. As the indenter is withdrawn, only the elastic portion of
the displacement is recovered and it distinguishes the elastic properties of the material
from its plastic properties [30]. Here hs is the displacement of the surface at the perimeter
of the contact; h is the total indentation depth; hf is the final depth of the residual
hardness impression. Indentation load versus displacement relationship is shown
schematically in Figure 3.3.
Figure 3.3: Nanoindentation load-displacement relationship
The critical parameters used in the calculation of hardness and elastic modulus are
the peak load (Pmax), the maximum indentation depth (hmax), the residual or final
indentation depth after unloading (hf) and the initial unloading stiffness (S = dP/dh). S is
also known as the elastic contact stiffness, or simply contact stiffness. Hardness, the
mean pressure a material can withstand, is calculated using Equation 1,
APH = (1)
28
where P is the load and A is the projected contact area at that load. The hardness
determined by dividing the applied load by the projected contact area under the load
should not be confused with the traditional definition of hardness, the load divided by the
projected area of contact of the residual hardness impression. These two definitions yield
similar values for hardness when the deformation process is mainly dominated by a
plastic region and a fully plastic permanent hardness impression is formed. When the
contact is predominantly elastic, they exhibit different values for hardness since the
residual contact area is very small for a purely elastic contact, which results in infinite
hardness based on the traditional hardness definition [30]. This phenomenon plays a
crucial role while indenting materials with very sharp indenters at small indentation
depths. In these conditions, traditional definition of hardness yields a higher value than
that calculated using Equation 2.
Elastic modulus of the test material, E, is calculated using the equation,
i
i
r EEE
22 111 νν −+−= (2)
where Er is the reduced elastic modulus, E and ν are the elastic modulus and Poisson�s
ratio of the test material, and Ei and νi are the elastic modulus and Poisson�s ratio of the
indenter, respectively, and for a diamond indenter these values are 1141 GPa and 0.07
respectively. Reduced modulus is used to account for the elastic displacements that occur
in both the indenter and the sample. Er is related to the contact stiffness by the equation,
A
SEr βπ
2= (3)
29
where S is the contact stiffness, A is the projected contact area and β is a constant that
depends on the geometry of the indenter. This equation is derived from elastic contact
theory [41-42] and holds good for any indenter that can be described as a body of
revolution of smooth function [43]. Though the equation formally applies only to circular
contacts, it has been shown that it works well for different geometries provided a
different value of β is used [44-45]. For indenters with square cross sections β = 1.012;
for triangular cross sections like the Berkovich and cube-corner indenters, β = 1.034. It
has been shown recently that yet another correction factor has to be added to the equation
[42, 46-49].
The principal difference between conventional hardness testing techniques and
the nanoindentation technique is the manner in which the contact area is determined.
Instead of imaging of the indent impressions, the contact area is estimated from the
indentation load-displacement data. Oliver and Pharr [22] developed a method by fitting
the unloading portion of the load-displacement data to the power-law relation
mfhhAP )( −= (4)
where A and m are constants determined by a least squares fitting procedure; hf is the
final displacement after complete unloading determined from curve fit. The initial
unloading slope is found by differentiating the equation and evaluating the derivative at
the peak load and displacement. Contact stiffness is obtained by differentiating Equation
(4) and evaluating the value at the maximum indentation depth.
The contact depth, hc, is estimated using, the expression,
30
SPhhc ε−= (5)
where ε is a constant that depends on the indenter geometry. For spherical indenters ε =
0.75, for conical indenters ε = 0.72 and for Berkovich indenters ε = 0.75. Equation (5)
does not account for pile-up formed during indentation because it is assumed that the
contact is perfectly elastic and only sink-in occurs continuously [30]. The contact area is
calculated as a function of the contact depth, hc, and is written as:
A (hc) = 24.5 hc2 (6)
for a perfect Berkovich indenter. Taking into consideration the tip blunting effects, the
modified area function is written as,
128/18
64/17
32/16
16/15
8/14
4/13
2/12
11
25.24)( cccccccccc hChChChChChChChChhA ++++++++=
(7)
where C1 through C8 are constants. The area function is determined by making a series of
indentations at various depths in a calibration sample of well-known elastic properties.
The basic assumption in this process is that the elastic modulus is independent of the
indentation depth. It is also imperative that there is no pile-up. Fused silica is the widely
accepted calibration sample. The machine compliance is also calculated during the
calibration procedure [22]. Correcting for machine compliance, the load-displacement
data are reduced and used to obtain the contact stiffness (S) and the contact depths (hc).
For fused silica material, E = 72 GPa; ν = 0.17 and for diamond indenter, E = 1141 GPa;
ν = 0.07. The contact area, determined using Equation (6), is plotted against contact
depth, hc, and is fitted according to a polynomial order. A weighted fitting procedure
31
ensures that data from all depths are given equal importance. The area function
coefficients determined from the fit are used in further calculations.
3.4.2.2 Microindentation
The resistance to indentation or deformation of material is known as its hardness.
It is one of the most frequently measured properties of a ceramic. There are many
hardness scales and methods of measurements to characterize resistance to deformation,
densification, and fracture. In fact, many ceramic specifications list minimum hardness
requirements. For example, a new ASTM zirconia�s specification for surgical implants, F
1873-98, stipulates that Vickers hardness (HV) shall be no less than 11.8 GPa (1200 kg/
mm2) at a load of 9.8 N (1 kg). Although measuring and interpreting ceramic hardness
should be routine, pitfalls, controversies, and surprises abound.
The Vickers hardness test was used for hardness and fracture toughness
measurements in this research work. In most engineering and characterization
applications, approximately 60% of worldwide published ceramic hardness values are
Vickers, with loads typically in the range of a few newtons to 9.8 N (1 kg) and occasional
data for soft or high-toughness ceramics as high as 98 N (10 kg). Knoop hardness is
more frequently used for glass, glass ceramics, and ceramic white wares. About 35% are
Knoop with loads from as low as 0.98 N (100 g) to 19.6 N (2 kg). Rockwell hardness are
about 5% [21]. For research purposes, Vickers, Knoop, and Berkovich (triangular
pyramid) indenters are customary; Rockwell and Brinell indenters are rarely suitable for
ceramics research. Measurements were carried out on conventional microhardness
machines with Vickers diamond indenters.
32
A LECO Microhardness testing Machine Model: M-400-H1/H2/H3 was used for
indentations. The indentions loads were 100, 200, 300, 500, 1000g. At each load 20
indents were created. The following procedures were undertaken for indentations:
• A highly polished flat sample was placed on stage at a right location of
microhardness testing machine.
• The sample was focused and adjusted to create 20 indents at each load from 100,
200, 300, 500, 1000g.
• A loaded indentor was lowered against a polished flat surface of the sample for a
specific load time of 30 seconds and retracted automatically afterwards.
• During loading, the indentor penetrated into the ceramic and, on retracting, left a
permanent pyramidal indentation.
• A high quality impression image was collected on an Optical Microscope
adjusting the indent on the center of the lens to avoid possible errors of focusing
on Zeiss Axiovert 10 at magnification 500x (1000, 500g) and 1000x (100, 200,
300g).
• Diagonal lengths of the impression were measured by using Image Pro-Plus
software for Vickers hardness.
• Crack lengths were measured for fracture toughness analysis.
A typical impression is shown in Figure 3.4.
33
Figure 3.4: Typical images collected at 500x on optical microscope for hardness study
The Vickers hardness is calculated by using the following formula
= 28544.1
DFVHN GPa
where, F = Load in Kg, D = average diagonal length in µm, VH = kg/mm2 (1 kgf/mm2 =
9.81 MPa.)
Fracture toughness, K1C, is a measure of a ceramic part�s resistance to fracture
starting from a pre-existing crack. It is one of the most important properties of material
for virtually all design applications. If the material has a large value of fracture
toughness, it will undergo ductile fracture. Brittle fracture is a characteristic of materials
with a low fracture toughness value. There are four different types of fracture toughness,
34
KC, K1C, K11C, and K111C. KC is the fracture toughness of the sample, which is in a state
plane stress. K1C, K11C, and K111C are the fracture toughness of material under three
different modes of fracture, mode I, mode II, mode III, respectively.
Ceramic materials have a much lower K1C value than the metals. The low K1C
value shows that ceramic materials are very susceptible to cracks and undergo brittle
fracture, whereas the metals undergo ductile fracture. Thus for this study, the fracture
toughness, K1C is considered.
For fracture toughness calculation Young�s modulus was obtained from
nanoindentation data 230GPa and 210GPa for DKKK, Japan and Praxair samples
respectively.
The equations used for fracture toughness calculation is:
2/3
2/1
1 016.0C
PHEK C
=
where, E = Young�s modulus (230/210GPa), H = Vickers hardness (GPa), F =
Indentation load (mN), C = Crack length (µm). Figure 3.5 shows diagonals and cracks
taken for measurements. Young�s modulus was obtained by nanoindentation technique.
The Young�s modulus value for the DKKK based ceramic sample sintered at 1600°C was
230GPa. The same value for the Praxair based ceramic sintered at 1600°C was 210GPa.
Both values are in good agreement with the literature review of this experimental work.
35
(a)
.
(b)
Figure 3.5: (a) Schematic figure having crack and indent length; (b) Image collected at 1000x on optical microscope (Indentation load is 200g)
D
C
36
3.5 Microstructure Characterization
Microstructure observations of the Scandia Stabilized Zirconia (ScSZ) were
performed on the thermally etched samples of DKKK, Japan and Praxair materials using
a scanning electron microscope (SEM) Model: Hitachi S3000N and atomic force
microscope (AFM) respectively. A brief description of both microscopes and the
technique adopted for the collection of images follows.
Scanning Electron Microscope
The SEM uses electrons rather than light to form an image. There are many
advantages to using the SEM instead of a light microscope. The SEM has a large depth of
field, which allows a large amount of the sample to be in focus at one time. The SEM
also produces images of high resolution, which means that closely spaced features can be
examined at a high magnification. Preparation of the samples is relatively easy, since
most SEMs only require the sample to be conductive. The combination of higher
magnification, larger depth of focus, greater resolution, and ease of sample observation
makes the SEM one of the most heavily used instruments in materials research today.
The scanning electron microscope (SEM) has a very fine �probe� of electrons
with energies of up to 40 keV focused at the surface of the specimen in the microscope
and scanned across it in a �raster� or pattern of parallel lines. A number of phenomena
occur at the surface under electron impact: most important for scanning microscopy is the
emission of secondary electrons with energies of a few tens of eV and re-emission or
reflection of the high-energy backscattered electrons from the primary beam.
37
Backscattered Electron Imaging
In this particular research backscattered electron imaging was used to collect
images for microstructure characterization. The intensity of emission of both secondary
and backscattered electrons is very sensitive to the angle at which the electron beam
strikes the surface, that is, to topographical features on the specimen. When the electron
beam strikes the sample some of the electrons will interact with the nucleus of the atom.
The negatively-charged electron will be attracted to the positive nucleus but if the angle
is just right instead of being captured by the "gravitational pull" of the nucleus, it will
circle the nucleus and come back out of the sample without slowing down. These
electrons are called backscattered electrons because they come back out of the sample.
Because they are moving so fast, they travel in straight lines. In order to form an image
with BSE (backscattered electrons), a detector is placed in their path. When they hit the
detector a signal is produced which is used to form the image.
Different elements have different size of nuclei. As the size of the atom nucleus
increases, the number of BSE increases. Thus, BSE can be used to derive an image that
shows the different elements present in a sample. A typical image collected at SEM is
shown in Figure 3.6.
The magnification of this microscope is the ratio between the dimensions of the final
image display and the field scanned on the specimen. Usually, a magnification range of
SEM is between 10 to 200000X and the resolution (resolving power) is between 4 to 10
nm (40 - 100 Angstroms).
38
Figure 3.6: Typical SEM images for grain size study
There are several types of SEMs designed for specific purposes ranging from
routine morphological and microstructural studies of materials, to high-speed
compositional analyses.
Atomic Force Microscope
The AFM is also called the scanning force microscope (SFM) and was invented in
1986 by Binnig, Quate and Gerber. The AFM utilizes a sharp probe moving over the
surface of a sample in a raster scan. The probe is a tip on the end of a cantilever which
bends in response to the force between the tip and the sample. Figure 3.7 illustrates the
working of the AFM. As the cantilever flexes, the light from the laser is reflected onto the
39
split photo-diode. By measuring the difference signal (A-B), changes in the bending of
the cantilever can be measured.
Figure 3.7: AFM working principle
The interaction force between the tip and the sample can be found as the
cantilever obeys Hooke's Law for small displacements. The movement of the tip or
sample is performed by an extremely precise positioning device made from piezo-electric
ceramics, most often in the form of a tube scanner. The scanner is capable of sub-
angstrom resolution in x-, y- and z-directions. The z-axis is conventionally perpendicular
to the sample.
40
The AFM operates by measuring attractive or repulsive forces between a tip and
the sample. In its repulsive contact mode, the instrument lightly touches a tip at the end of
a cantilever to the sample. As a raster-scan drags the tip over the sample, some sort of
detection apparatus measures the vertical deflection of the cantilever, which indicates the
local sample height. Thus, in contact mode, the AFM measures hard-sphere repulsion
forces between the tip and the sample.
In non-contact mode, the tip does not touch the sample and the AFM derives
topographic images from measurements of attractive forces. AFMs can achieve a
resolution of 10 pm, and unlike electron microscopes, can image samples in air and under
liquids.
Feedback operation
The AFM can be operated in two principal modes:
• with feedback control
• without feedback control.
Once the electronic feedback is switched on, then the positioning piezo which is
moving the sample or tip up and down can respond to any changes in force which are
detected, and alter the tip-sample separation to restore the force to a predetermined value.
This mode of operation is known as constant force mode, and usually enables a fairly
faithful topographical image to be obtained.
When the feedback electronics are switched off, then the microscope is said to be
operating in constant height or deflection mode. This is particularly useful for imaging
41
very flat samples at high resolution. The following procedures were considered to obtain
high quality images for grain size study of all samples:
• The controller and the main computer were turned on and the SPM cockpit
software was launched.
• The close-contact EZ-Mode was selected and the corresponding configuration file
was loaded.
• The cantilever close-contact mode operation was installed in the AFM scanner
head.
• The cantilever tip was aligned so that the laser light was projected from the back
of the cantilever onto the photodetector.
• The AFM scanner was moved down and the surface was focused to obtain the
high quality image.
• The tip approach was carried out and a specific area on the sample was selected to
obtain better quality images with distinct grain boundaries.
• The probe was retracted from the sample.
• The above procedure was repeated with remaining samples.
• Grain size measurements were performed by the intersect method of Image Pro-
Plus software using Scanning Electron Microscope (SEM) and Atomic Force
Microscope (AFM) images (see Figure 3.8).
42
(a)
(b)
Figure 3.8: (a) Typical AFM images for grain size study; (b) Grain size and cavitations measurements
43
The measurements were performed, collecting data by moving the red lines from
top to bottom at 10 locations in the same images. Almost 20 grain intersects per line were
collected (see Figure 3.8 (b)). For each sample, the data were collected from 4/5 images
which gave 800-1000 measurements. Cavities density was collected using the threshold
method as shown in Figure 8.3 (b).
3.6 XRD Measurements
The X-ray diffraction experiment is mostly performed to characterize
crystallographic structure and chemical composition of the materials. In the present work,
the XRD experiment was used for phase identification of ceramic sintered at different
temperature, and calculation of coefficient of thermal expansion by obtaining lattice
parameters.
The experiments were performed using the AXS-Bruker D8 Discover
diffractometer with Cu Kα X-ray radiation source and Eurlean cradle. Parallel beam
optics was used to minimize errors associated with sample displacement and surface
roughness. The Anton Paar H-900 high temperature stage mounted on the XYZ-
positioning stage was used for measurements at temperatures up to 900ºC. The zero
point of detector was <0.004º, and sample height was controlled with accuracy of 0.005
mm. Ceramic samples were polished down to 0.5 mm thick tablet form with final grit
size 0.05 µm to avoid stress induced effects on XRD data.
Powder samples for high temperature measurements were deposited on Si(100)
substrate in the slurry form using colloidal silver paste. Silver matrix has been used for
44
temperature calibration. Diffraction patterns were collected at 25ºC and then between 100
and 800ºC in 100ºC intervals for CTE measurements and 10°C for phase transition
analysis. The heating rate was 10 ºC/min and the waiting time was 30 minutes before
XRD scans at every temperature for temperature stabilization. High temperature XRD
scans were performed in the interval 2θ of 28-66º with integration time 5s/point and step
0.01º. The waiting time for temperature stabilization before the XRD scan at each
temperature was 30 minutes.
3.6.1 Coefficient of Thermal Expansion (CTE)
When heat is given to a body, it normally expands. The coefficient of thermal
expansion is the relative change in a given dimension when the body is heated. Thermal
expansion of the crystalline phase means a change in the cell lattice parameters with the
temperature variation. Increase of the lattice parameter in the cubic phase with
temperature increase leads to shifting of diffraction peaks towards lower angles. An
example of this behavior is shown in Figure 3.9 for the P1600 powdered sample. To
determine coefficient of thermal expansion (CTE), cell parameters at every temperature
were obtained by Retvield refinement of at least 5 peaks in XRD data using Full-Proof
Software. The CTE is defined by the following formula:
dTdL
LCTE 1=
where, L is original length, dL is change in length and dT is temperature difference. This
demonstrates the linear expansion of solid material. The cubical or volume expansivities
45
of solids are three times the linear expansivity. The unit of linear CTE is /°C or
mm/mm/°C.
Figure 3.9: Temperature dependent shifting of XRD peaks positions
The expansivities of the majority of solid materials increase with increasing
temperature and can be fitted in three different ways:
1) Model-I:CTE- constant
obCTE =
By substitution and integration,
46
.*ln constTbL oT +=
2) Model-II: CTE-varying linearly
TbbCTE o 1+=
By substitution and integration,
.2
*ln 21 constTbTbL oT ++=
3) Model-III: CTE - varying qudratically
221 TbTbbCTE o ++=
By substitution and integration,
.32
*ln 3221 constTbTbTbL oT +++=
For all three models, coefficients bo, b1, b2 were obtained through curve fitting log
(LT) by linear, quadratic, and cubic functions respectively. An example of this fitting is
presented in Figure 3.10. Analysis of fitting residuals can help to select the appropriate
model for CTE (see Figure 3.11). By substituting value of coefficients and given
temperature, the CTE was calculated. In the case of the P1600 powdered sample
dependence Ln (LT) is not linear and the best fit is achieved by linear and quadratic
models. In cases like this, selection between linear and quadratic models is difficult
because residuals are similar. Despite the constant and the linear models, the quadratic
model is the better fitting (less residuals) model. The accuracy of coefficients of fitting
was used for selection of the appropriate model. In the case of the P1600 powdered
47
sample (see Figures 3.9, 3.10, 3.11) for the linear model b0 = (7.14±0.39)×10-6 and
b1=(4.61±0.53)×10-9, while for the quadratic model b0= (5.83±2.33)×10-6,
b1=(8.47±6.75)×10-9 and b2=-(2.68±4.67)×10-12. It is evident that statistical accuracy of
coefficients for the linear model is significantly better than for the quadratic model.
Therefore, the best model was decided based on analysis of residuals, overall standard
deviations and standard errors of parameters. To compare our data with the known value
of CTE of cubic zirconia (10.5×10-6 deg-1), we have used the results of a constant model.
300 400 500 600 700 800 900 1000 1100
1.628
1.629
1.630
1.631
1.632
1.633
1.634
1.635
1.636 Constant Linear Qudratic
Ln(L
T)
Temperature (K)
Figure 3.10: Constant, linear and quadratic fitting with temperature vs. log (LT) (P1600-powder)
48
300 400 500 600 700 800 900 1000 1100
-0.00015
-0.00010
-0.00005
0.00000
0.00005
0.00010
0.00015
Del
ta
Temperature, K
Constant Linear Quadratic
Residuals
Figure 3.11: Instrumentally set temperature vs. residuals (P1600-powder)
3.6.1.1 Temperature Calibration
It is known that the coefficient of thermal expansion is temperature dependent
parameters. The actual value of temperature is essential for the accurate calculation of
coefficient of thermal expansion. Based on literature reviews, the CTE at known
temperature of silver was taken as reference value for temperature calibration. By back
substitution and polynomial fitting using Microsoft excel and the Origin program, the
coefficient of polynomial equation was obtained for known temperature, and for each set
temperature the actual temperature was calculated as shown in Figure 3.12 and Table 3.3.
49
Figure 3.12: Polynomial fitting of calibrated temperature vs. instrumentally set temperature Table 3.2. Calculation of actual temperature from set temperature
Set Temperature(°C)
Set Temperature(K)
Cell Parameters(Å)
Calibrated Temperature(K)
100 373 4.09107 356.8 150 423 4.09618 420.2 200 473 4.10045 471.9 250 523 4.10481 523.4 300 573 4.10911 573.0 350 623 4.11372 624.9 400 673 4.11819 674.1 450 723 4.12312 727.0 500 773 4.12779 776.0 550 823 4.13246 823.8 600 873 4.13728 872.0 650 923 4.14271 925.1 700 973 4.14754 971.1 800 1073 4.15823 1069.6
200 400 600 800 1000 1200200
300
400
500
600
700
800
900
1000
1100
Actu
al T
empe
ratu
re (K
)
Set Temperature (K)
50
CHAPTER 4
RESULTS AND DISCUSSION
4.1 Characterization of Praxair and DKKK powders
Two different powders with nominal composition 10mol%Sc2O3-1mol%CeO2-
ZrO2 were studied in this research work:
• Praxair Surface Technology, Specialty Ceramics, USA
• Daiichi Kigenso Kagaku Kogyo Co., Ltd , Japan (DKKK, Japan)
4.1.1 Theoretical Dnsities and Actual Composition
The actual composition of powders (obtained by Inductively Coupled Plasma
Mass Spectroscopy at the University of St. Andrews, UK by J. Irvine and C. Savanie) and
theoretical densities were calculated using the lattice parameter obtained by XRD
experiments for both powders (see Tables 4.1, 4.2, 4.3):
DKKK Powder
The actual composition of DKKK powder is 10.07% Sc2O3 + 0.92% HfO2 +
1.03% CeO2 + 1.4% TiO2 + 86.58% ZrO2.
Praxair Powder
The actual composition of Praxair Powder is 15% Sc2O3 +0.05% HfO2 + 0.4%
CeO2 + 2% TiO2 + 82.55%ZrO2
51
Tab
le 4
.1. T
heor
etic
al d
ensi
ty fr
om th
e ac
tual
com
posi
tion
of D
KK
K p
owde
r
Com
pd.
Wei
ght
Mas
s #
of
Sc,H
f, C
e,Ti
, Z
r
Ato
mic
M
ass o
f O
xyge
n
# of
O
xyge
n A
tom
Mas
s pe
r M
etal
A
tom
Con
tent
s of
Eac
h C
ompd
.
# of
M
oles
Nor
mal
izat
ion
O
xyge
n
Ato
mic
W
eigh
t
Sc2O
3 44
.96
44.9
6 2
16
3 15
0.
1460
3 0.
1665
9 0.
2498
9 29
.960
HfO
2 17
8.49
17
8.49
1
16
2 0.
05
0.00
437
0.00
499
0.00
997
3.56
0
CeO
2 14
0.12
14
0.12
1
16
2 0.
4 0.
0059
8 0.
0068
3 0.
0136
5 3.
826
TiO
2 47
.9
47.9
1
16
2 2
0.01
752
0.01
999
0.03
998
3.83
0
ZrO
2 91
.22
91.2
2 1
16
2 82
.55
0.70
265
0.80
160
1.60
321
292.
48
Tota
l Mol
es
0.8
7655
1
Oxy
gen
1.91
670
122.
67
Sum
456.
34
Avo
gadr
o�s N
umbe
r 6.
032*
1023
Latti
ce P
aram
eter
s 5.
09*1
0-8cm
Theo
retic
al D
ensit
y 5.
74 g
m/c
m2
52
Tab
le 4
.2. T
heor
etic
al d
ensi
ty fr
om th
e ac
tual
com
posi
tion
of P
raxa
ir p
owde
r
Com
pd.
Wei
ght
Mas
s #
of S
c,
Hf,
Ce,
T
i, Zr
Ato
mic
M
ass o
f O
xyge
n
# of
O
xyge
n A
tom
Mas
s pe
r M
etal
A
tom
Con
tent
s of
Eac
h C
ompd
.
# of
M
oles
Nor
mal
izat
ion
O
xyge
n
Ato
mic
W
eigh
t
Sc2O
3 44
.96
44.9
6 2
16
3 15
0.
2175
0.
237
0.35
6 42
.749
HfO
2 17
8.49
17
8.49
1
16
2 0.
05
0.00
02
0.00
025
0.00
052
0.18
5
CeO
2 14
0.12
14
0.12
1
16
2 0.
4 0.
0023
0.
0025
3 0.
005
1.42
3
TiO
2 47
.9
47.9
1
16
2 2
0.02
50
0.02
73
0.05
4 5.
241
ZrO
2 91
.22
91.2
2 1
16
2 82
.55
0.66
99
0.73
2 1.
464
267.
14
Tota
l Mol
es
0.91
50
1
O
xyge
n
1.
881
120.
39
Sum
437.
13
Avo
gadr
o�s N
umbe
r 6.
032*
1023
Latti
ce P
aram
eter
s 5.
09*1
0-8cm
Theo
retic
al D
ensit
y 5.
49gm
/cm
2
53
Table 4.3. Powder characterization
It was found that the DKKK powder has a composition close to the nominal
composition while the Praxair powder has an excessive amount of Sc2O3 and a lower
amount of CeO2. The HfO2 is considerably less in the Praxair than in the DKKK and
there is no big difference in composition of TiO2 in both powders. The c-phases at room
temperature has almost identical lattice parameters as shown in Table 4.3. The result
shows that the DKKK powder has a composition close to the nominal while the Praxair
powder has an excessive amount of Sc2O3 and a lower amount of CeO2. From Figure 4.1,
Figure 4.1: SEM micrographs of powders (a-Praxair, USA) and (b-DKKK, Japan)
Powder Name Actual Compositions Lattice Parameter(Å)
Theorectical Density(g/cm3)
Praxair powder 15% Sc2O3 +0.05% HfO2 + 0.4% CeO2 + 2% TiO2 + 82.55%ZrO2
5.09203
5.49
DKKK powder 10.07% Sc2O3 + 0.92% HfO2 + 1.03% CeO2 + 1.4% TiO2 + 86.58% ZrO2
5.09119
5.74
54
it is very clear that Praxair powder consists of relatively large 1-2 µm pre-sintered
agglomerates of smaller ~300 nm particles and DKKK powder has significantly higher
surface area due to relatively uniform nanograins with size 60-100 nm.
4.1.2 XRD of Powders
The XRD study of as-received powders was performed at different temperature
ranges RT-600°C as shown in Figure 4.2. It was found that both powders had a mixture
of rhombohedra phase (β-phase) and cubic phases (c-phase) at room temperature.
Figure 4.2: XRD pattern at different temperatures
30 35 40 45 50 55 60
O
O
O
O
OO OO O
OO
O
OOO
X
X
X
X
X
II
II
II
I
I
I
I
I
2Theta (degree)
RT
500°C
600°C
I
O
X
I Cubic Phase (space group Fm-3m)O Rhombohedral Phase (space group R-3)X Silver (space group Fm-3m) as a reference for HT measurements
XRD of DKKK as-received powder at different temperatures
55
The result showed that with the increase of temperature the rhombohedra phase
disappeared in the XRD pattern and only remained in a cubic phase at temperatures
higher than 500°C. All samples sintered at temperatures higher than 1200°C existed in
cubic phase at room temperature.
XRD of Praxair
The following XRD patterns were obtained for the as-received Praxair powder
(see Figure 4.3).
Figure 4.3: XRD pattern of as-received Praxair powder
The XRD pattern showed that at room temperature, the powder has both
rhombohedra and cubic phases. The cubic to rhombohedral phase ratio for Praxair
powder is larger than for DKKK powder. The whole pattern was successfully fitted by
three known phases: cubic ScSZ, rhombohedral phase (b-phase) of ScSZ and FCC silver
3 0 4 0 5 0 6 0 7 0 8 0 9 0 1 0 0 1 1 0 1 2 0
0
1 0 0 0
2 0 0 0
3 0 0 0
4 0 0 0
5 0 0 0
4 9 .5 5 0 .0 5 0 .5 5 1 .0 5 1.5 5 2 .0
C /R h = 8.0 8
M u lt i-p h as e an a ly s is re su lts :A g 8 .5 7 (0 .1 4 )%C ub ic 8 1 .37 (0 .6 2 )%R ho m bo h ed ra l 1 0 .07 (0 .2 5 )%
Inte
nsity
(cou
nts)
2 T h e ta (d e g re e )
C ub ic
S ilv e r
R ho m b o he d ra l
R e tvie ld re fin em en t o f as -re ce ive d P rax a ir p o w d e r
2 T h e ta (d e g re e )
56
(used as internal standard). This powder has small amounts (~12%) of secondary b-
phase. The insert above shows details of multi-phase fitting used in the Retvield
refinement procedure: small amounts of secondary phase can be traced but quantitative
analysis cannot be done correctly. The cubic phase had lattice parameter
a=b=c=5.09203(7) Å and the corresponding theoretical density is 5.49 g/cm3 based on
the exact content of powder. The apparent crystallite size of c-phase is 68.1 nm, which
supports SEM observation: big powder particles consist of agglomerates of smaller
grains.
XRD of DKKK powder
The following XRD patterns were obtained for the as-received DKKK powder
(see Figure 4.4).
Figure 4.4: XRD pattern of as-received DKKK powder
3 0 4 0 5 0 6 0
0
2 0 0 0
4 0 0 0
6 0 0 0
8 0 0 0
1 0 0 0 0
29 .0 29 .5 30 .0 3 0 .5 3 1 .0 31 .5
R etvie ld re fin e m en t o f a s -re c e ive d D K K K p o w d e r
Inte
nsity
(cou
nts)
2 T h e ta (d e g re e )
C u bic
R ho m b o he d ra l
S ilv e r
M u lt i-p h ase an a lys is resu lts:A g 4 .2 4 (0 .0 2 )%C u bic 7 5 .6 2 (0 .4 3 )%R h o m bo h e d ra l 2 0 .1 4 (0 .3 4 )%
C /R h = 3 .75
2The ta (d egree )
57
DKKK has also both rhombohedral and cubic phases at room temperature. A
phase transition was found at a temperature higher than 500°C and cubic phases recorded
at 600°C. Whole patterns can be successfully fitted by three known phases: cubic ScSZ,
rhombohedral phase (b-phase) of ScSZ and FCC silver. This powder has also significant
amounts (~21%) of secondary β-phase. The insert above shows details of multi-phase
fitting used in the Retvield refinement procedure: the amount of secondary phase can be
quantitatively analysed. The cubic phase has a lattice parameter a=b=c=5.09119(7) Å and
its corresponding theoretical density is 5.74 g/cm3, based on the exact content of powder.
The apparent crystallite size of the c-phase is 36.8 nm, which is in good agreement with
SEM data. The grain size of DKKK powder is two times smaller than the Praxair powder
which can lead to lower sintering temperature.
4.2 Sintering Behavior of Powders
The sintering behaviors of powders were studied by density measurement and
microstructure of powder. It was found that the DKKK powder had significantly higher
sintering activity and required ~200ºC less sintering temperature than the Praxair powder.
The difference in the sintering activity of powders was a result of differences in the
powder morphology, chemical composition and levels of impurities present in the DKKK
and Praxair powder. It was also observed that the agglomeration had a significant
influence in sinteribility of behavior of as-received DKKK and Praxair powder.
58
4.2.1 Density Measurements of ScCeSZ and its Porosity
The density of sintered specimens was determined by the Archimedes� method.
The sintering temperature and density data of ScCeSZ are listed in Table 4.4.
Table 4.4. Density and porosity of sintered samples
Sample Sintering Temp(ºC)
Density (g/cm3)
Standard Dev
Porosity (%)
Standard Dev
J1100 1100 3.31 .056 42.3% 1.0% J1200 1200 4.96 .036 13.5% 0.6% J1300 1300 5.38 .008 J1400 1400 5.52 .011 J1500 1500 5.55 .026 J1600 1600 5.57 .007 P1100 1100 2.72 .044 50.2% 0.8% P1200 1200 3.11 .095 43.2% 1.7% P1300 1300 3.70 .086 32.4% 1.6% P1400 1400 4.57 .070 16.4% 1.3% P1500 1500 5.11 .041 6.6% 0.7% P1600 1600 5.27 .007 3.7% 0.1%
The results show that the porosity level in ceramics sintered from Praxair and
DKKK powders has significant sinteribility difference. It was found that the DKKK
powder was significantly more active in the sintering process and required ~200ºC lower
sintering temperature to produce ceramics of the same level of porosity as in samples
sintered from Praxair powder (Figure 4.1). Amounts of porosity had been calculated
using theoretical densities 5.49 and 5.74 g/cm3 of ScCeZrO2 ceramics Praxair and DKKK
powder respectively which were estimated using XRD data and the actual chemical
composition of the powders. Density data showed that for ceramics produced from
59
DKKK powder, density increases rapidly from 1100 to 1400ºC and at higher
temperatures become constant (Table 4.4). Similarly, in the same temperature range,
grain size in DKKK ceramics increases significantly from 200 nm to 3 µm and stabilizes
at higher temperatures. Even at 1200°C samples produced from DKKK powder are dense
enough for polishing, grain size measurements and microhardness testing while samples
sintered from Praxair powder at 1400°C are still porous for microindentation. Analysis of
grain size with increase of sintering temperature allowed for the conclusion that DKKK
powder produces dense ceramics at temperatures higher than 1400ºC while Praxair
powder is still not fully sintered at 1600ºC. This conclusion was supported by analysis of
fracture surfaces. Figure 4.5 shows typical fracture surfaces of the samples sintered at
1600ºC from Praxair powders and DKKK powder. It is evident that DKKK powder
created almost
Figure 4.5: Fracture surfaces of samples P1600-(a) and J1600-(b)
60
fully dense ceramic while the sample produced from the Praxair powder had noticeable
amounts of porosity. It should be noticed that thermal etching effects on polished
surfaces of samples produced from Praxair powders and DKKK are different (see Figures
4.6-4.8).
Figure 4.6: Grain structure of samples produced from Praxair powder sintered at temperatures 1500ºC (left) and 1600ºC (right)
Figure 4.7: Grain structure of samples produced from DKKK powder sintered at temperatures 1400ºC (left), 1500ºC (middle) and 1600ºC (right)
61
Figure 4.8: SEM images of polished and low angle ion-milled surface of sample P1600 (a) and thermally etched surfaces of samples P1600 (b) and J1600 (c)
Praxair-based samples unusually exhibit significant increase of in-grain cavities
due to thermal etching with sintering temperature while the in-grain cavitations level in
DKKK-based samples remain constant. SEM analysis reveals that all Praxair-based
samples have significant non-uniformity. Even the densest sample, P1600, has highly
dense and very porous areas with well distinguished borders between them (see Figure
4.8(a)). The difference in overall density between samples P1400, P1500 and P1600
(Table 4.4) results predominantly from different ratios between dense and porous areas in
the samples. Grain structures in Figure 4.8(b) demonstrate a higher density area of
sample P1600 while average porosity of the sample is 4.4%. Contrary to samples
(Praxair), DKKK-based samples have high uniformity, and the density of in-grain
cavities remains constant for samples J1400, J1500 and J1600 (Table 4.4). One can
conclude that DKKK powder produces dense ceramics at temperatures higher than
1400ºC while ceramics produced from Praxair powder are still not fully sintered at
1600ºC (see Figures 4.9-4.10).
62
1100 1200 1300 1400 1500 1600
2.0
2.5
3.0
3.5
4.0
4.5
5.0
5.5
6.0
Praxair Japan
Den
sity
(g/c
cm)
Temperature (C)
Figure 4.9: The dependence of density on sintering temperatures
1100 1200 1300 1400 1500 16000
5
10
15
20
25
30
35
40
45
Porous
Poro
sity
Lev
el (%
)
S intering Temperature (0C)
Praxair (I) DKKK (II)
Figure 4.10: Porosity level vs. sintering temperature
63
The difference in sintering activity of Praxair and DKKK powders is a result of
differences in powder morphology, chemical composition and impurities. Praxair powder
consists of relatively large 1-2 µm pre-sintered agglomerates of grains while DKKK
powder has a significantly higher surface area due to relatively uniform nanograins (see
Figure 4.1) which favors lower sintering temperature for DKKK powder. Also Praxair
and DKKK powders have a slightly different chemical composition that could have an
additional effect on the sintering process.
4.2.2. Microstructure and Grain Size
The microstructure properties were characterized by collecting images at SEM
and AFM of sintered ceramics. The high quality images were obtained with distinct grain
boundary and grain size measurement was performed as described in Chapter 4 using
Image Pro-Plus. It is observed from the experimental data that the sintering temperature
has a significant effect on the grain size of the ceramic and its microstructure is shown in
Figures 4.11, 4.12, and 4.13. Results of grain size and cavities are listed in Table 4.5.
J1200°C J1300°C J1400°C Grain size: 0.34µ Gain size: 0.80 Grain Size: 2.71
64
(a)
J1500°C J1600°C Grain size: 3.61µ Grain Size: 3.81µ
(b) Figure 4.11: SEM images of DKKK samples from left to right sinter correspondingly 1200-1600°C
P1500°C P1600°C Grain size: 2.48µm Grain Size: 4.42µm
Figure 4.12: SEM images of Praxair samples from left to right sinter
correspondingly 1500-1600°C
65
1100 1200 1300 1400 1500 1600-0.5
0.0
0.5
1.0
1.5
2.0
2.5
3.0
3.5
4.0
4.5
5.0
5.5
6.0
0 2 4 6 8 10
0
2
4
6
8
10
0 2 4 6 8 100
2
4
6
8
10
DKKK, Japan Praxair, USA
Gra
in s
ize
(µ)
Sintering Temperature (ºC)
Figure 4.13: Grain size vs. sintering temperatures
Table 4.5: Grain size and in-grain cavities at different sintering temperatures
Sample Sintering Temp (ºC)
Grain size (Stdev) µm
Cavities (Stdev) µm
J1100 1100 0.23 (0.10) - J1200 1200 0.34 (0.15) - J1300 1300 0.80 (0.41) - J1400 1400 2.71 (1.36) 1.0 (0.4) J1500 1500 3.61 (1.89) 1.15 (0.3) J1600 1600 3.81(1.95) 0.94(0.08) P1100 1100 - - P1200 1200 - - P1300 1300 - - P1400 1400 - - P1500 1500 2.48 (1.27) 0.68 (0.18) P1600 1600 4.82 (2.36) 1.14 (0.23)
66
The ceramic sintered at higher temperature has a large grain size while ceramic
sintered at low temperature has a smaller size. The comparative study of grain size of
ceramics J1600 and P1600 showed that the P1600 has a larger grain size. While Praxair
sintered at 1500ºC has significant difference in grain size (Table 4.5) in comparison to the
Japan powder (DKKK). It was found that below 1500ºC, ceramics from Praxair powders
is not fully sintered; on the other hand, pallets of Japan (DKKK) powders are fully
sintered even at 1400ºC. The grain size versus the sintering temperature plot reveals that
the Japan powder can produce the same dense ceramic even at 200 lower temperatures
than the Praxair powder (see Figure 4.11). The Praxair powder cannot produce fully
dense material even at 1600°C (see Table 4.4). This result was also supported by the
images collected after ion milling (see Figures 4.14, 4.15, 4.16).
Figure 4.14: SEM images collected after 1hr ion milling at 30 deg (left-J1600, right- P1600)
67
Figure 4.15: SEM images collected after ion milling (left-J1600 milled at 9-deg for
18hrs, right-P1600 milled at 5 deg for 20hrs)
Figure 4.16: SEM images of P1600 collected after ion milled at 5 deg for 20hrs (right-high porous area, left-less porous)
4.3. Phases and Stability of DKKK and Praxair Ceramics
Phase identification is very important in characterization of materials properties.
Transition of phase has a greater impact on the different mechanical, thermal, and
microstructure properties of materials. The stability of phase is crucial in the case of
68
SOFC electrolyte design. This is because of the significantly higher ionic conductivity of
the cubic phase as compared to the rhombohedral phase. To analyze the DKKK and
Praxair materials, several XRD experiments were performed. The cubic and
rhombohedral phases were identified along with their phase stability temperature and
phase transition temperature.
4.3.1 Phase Analysis in DKKK-based Samples
Ceramics sintered from DKKK powder at temperatures below 1300°C consist of
two phases at room temperature. Ceramics sintered at temperatures 1300ºC and higher
exist in cubic phase at room temperature. Phase transition from rhombohedral to cubic
phase occurs at 400-500ºC (see Figure 4.17).
(a)
100 200 300 400 500 600 7000.0
0.2
0.4
0.6
0.8
1.0
1.0
0.8
0.6
0.4
0.2
0.0
Frac
tion
of C
ubic
Pha
se
Temperature (°C)
Fra
ctio
n of
Rho
mbo
hedr
al P
hase
69
(b)
Figure 4.17: Phase content in the sample DKKK-1200 at different temperatures from XRD (Retvield multiphase analysis)
4.3.2 Phase Transition in J1600 Sample
High temperature XRD experiments were performed to study the phase transition
behavior of J1600 and P1600 samples (see Figures 4.18-4.19). Silver peaks were used for
internal temperature calibration. It was found that the DKKK based samples underwent
the slow cubic to rhombohedra phase transformation at temperatures above 300°C.
Rhombohedron to cubic phase transformation occurs at a significantly higher rate at
420°C. Results show that all Praxair-based samples exist in c-phase in the temperature
range100-850°C.
30 31 50 51
400°C 440°C 460°C 480°C 500°C 540°C
2 Theta (°)
70
Figure 4.18: Changes in XRD pattern of powdered sample J1600 heated from 300 to 500°C with step 10°C.
Figure 4.19: Phase transition in DKKK ceramic sintered at 1600°C
300 350 400 450 500-0.030
-0.025
-0.020
-0.015
-0.010
-0.005
0.0
0.2
0.4
0.6
0.8
1.0
Hea
t Flo
w (W
/g)
Temperature (°C)
1 hr
449.4 °C3.462 J/g
Fra
ctio
n of
Cub
ic P
hase
from
XR
D
71
The differential scanning calorimetry (DSC) results show that when this material
was heated at temperatures higher than 300°C DKKK samples underwent the slow cubic
to rhombohedral phase transformation. This process was not registered as a peak in slow
DSC scan. As the temperatures increased above 420°C the reverse rhombohedral to cubic
phase transformation occurred at significantly faster rate and this endothermic process
was registered by the DSC (see Figure 4.19). There was no such phase transition in the
Praxair samples and only c-phase was found in the range of RT-850°C. It is my endeavor
to study this behavior by running different samples at different heating and cooling rates
for further investigation of material properties. Further experiments are under way at the
Center for Advanced Materials and Smart Structure (CAMSS).
4.4 Coefficients of Thermal Expansion (CTE)
The CTE calculation was based on the procedure described in Chapter 4. The
results obtained for coefficient of thermal expansion of as-received powders Praxair and
DKKK ceramic samples P1600 and J1400 and powders produced from them were
obtained by high temperature XRD method in the temperature range 100-850ºC. Powders
were used in the form of slurry coating on Si (100) substrate with small amounts of silver
paint used as a binder for powders and as internal reference material for temperature
calibration. Ceramic samples with thickness equal to Si substrate were polished down to
0.5µm grit size and annealed at 850ºC for 20 minutes to release stresses on sample
surface.
72
DKKK-based powders have an almost constant CTE in the temperature range
100-850ºC while CTE of all Praxair-based samples and ceramic sample J1400 almost
linearly increase with temperature. CTE of ceramics have the highest increase in this
range (almost 2 times for the sample P1600 and 1.7 times for J1400). This difference
supports assumption of significant internal compressive stresses in ceramics at room
temperature. This also can explain formation of crater-like cavities on fracture surfaces at
room temperature due to explosive chipping of material in order to release compressive
stresses near to the fracture surface.
The three models of fittings are listed in Table 4.6, and are recognized by bold,
italic and normal fonts.
The best fitting for as-received Praxair powder, as-received DKKK powder and
J1500-powder and J1500-ceramic is shown in Figures 4.20-4.25. The best fitting model
was chosen based on an analysis of residuals, overall standard deviations and standard
errors of parameters.
73
Tab
le 4
.6. C
oeffi
cien
ts o
f the
rmal
exp
ansi
on fo
r Pr
axai
r an
d D
KK
K sa
mpl
es in
cer
amic
and
pow
der
form
s
Sam
ple
Form
C
TE
×106
(deg
-1)
SD×1
06 #
of
Tem
p.
Prax
air
as-r
ecei
ved
Pow
der
10.0
2±0.
21
(6.5
1±0.
24) +
(5.1
6±0.
35)×
10-3
T
(7.4
3±1.
19) +
(2.2
5±3.
69)×
10-3
T +
(2.1
2±2.
68)×
10-6
T2
155.
1 27
.1
28.0
9
P160
0 Po
wde
r 10
.45±
0.19
(7
.14±
0.39
) + (4
.61±
0.53
)×10
-3T
(5
.83±
2.33
) + (8
.47±
6.75
)×10
-3T
- (2.
68±4
.67)
×10-6
T2
118.
4 32
.5
35.0
8
P160
0 C
eram
ic
10.5
1±0.
23
(4.8
5±0.
43) +
(8.2
7±0.
62)×
10-3
T
(5.2
4±2.
05) +
(7.0
5±6.
37)×
10-3
T +
(0.8
9±4.
62)×
10-6
T2
244.
3 69
.9
72.2
18
DK
KK
as
-rec
eive
d Po
wde
r 9.
74±0
.16
(7.6
1±0.
91) +
(3.1
7±1.
35)×
10-3
T (4
.32±
6.06
) + (1
3.4±
18.7
)×10
-3T
- (7.
6±13
.8)×
10-6
T2
89.2
60
.0
66.0
7
J140
0 Po
wde
r 10
.35±
0.20
(7
.67±
1.13
) + (3
.80±
1.60
)×10
-3T
(11.
8±6.
8) -
(8.7
±20.
3)×1
0-3T
+ (8
.9±1
4.4)
×10-6
T2
124.
6 92
.6
99.0
8
J140
0 C
eram
ic
10.6
0±0.
29
(5.6
8±0.
36) +
(7.3
6±0.
53)×
10-3
T
(7.3
7±1.
75) +
(1.9
2±5.
52)×
10-3
T +
(4.0
4±4.
09)×
10-6
T2
212.
8 39
.9
40.0
9
J150
0 Po
wde
r 10
.56±
0.23
(4
.89±
1.51
) + (7
.27±
1.92
)×10
-3T
(22.
8±10
.2) -
(39.
8±26
.8)×
10-3
T +
(30.
2±17
.1)×
10-6
T2
115.
1 64
.2
53.9
8
J160
0 Po
wde
r 10
.45±
0.20
(8
.49±
2.29
) + (2
.47±
2.87
)×10
-3T
(24.
0±19
.0) -
(37.
6±48
.8)×
10-3
T +
(25.
2±30
.6)×
10-6
T2
121.
6 12
3.0
125.
0 13
N
ote:
Bol
d �
best
fit,
Italic
� a
ccep
tabl
e fit
, Nor
mal
� p
oor f
it.
74
Praxair as received powder The following figures represent the best fitting model for the Praxair as received
powder (see Figures 4.20-4.21).
Figure 4.20: Instrumentally set temperature vs. residuals (Praxair powder)
Figure 4.21: Constant, linear and quadratic fitting with temperature vs. log (LT) (Praxair powder)
200 300 400 500 600 700 800 900 1000 1100
-0.00015
-0.00010
-0.00005
0.00000
0.00005
0.00010
0.00015
0.00020
0.00025
0 2 4 6 8 100
2
4
6
8
10
Del
t
Temperature, K
Constant Linear Qudratic
200 400 600 800 1000
1.628
1.630
1.632
1.634
1.636
Constant Linear Quadratic
Ln(a
)
Temperature, K
75
DKKK as received powder
The following figures represent the best fitting model for the DKKK as received
powder (see Figures 4.22-4.23).
Figure 4.22: Instrumentally set temperature vs. residuals (DKKK powder)
Figure 4.23: Constant, linear and quadratic fitting with temperature vs. log (LT) (DKKK powder)
400 600 800 1000
1.628
1.630
1.632
1.634 Constant Linear Quadratic
Ln(a
)
Temperature (K)
100 200 300 400 500 600 700-0.00010
-0.00005
0.00000
0.00005
0.00010
Del
ta
Temperature,K
Constant Linear Qudratic
76
J1500 Powder
The following figures represent the best fitting model for the J1500 powder (see
Figures 4.24-4.25).
Figure 4.24: Instrumentally set temperature vs. residuals (J1500 Powder)
Figure 4.25: Constant, linear and quadratic fitting with temperature vs. log (LT) (J1500 Powder)
300 400 500 600 700 800 900 1000 1100-0.00020
-0.00015
-0.00010
-0.00005
0.00000
0.00005
0.00010
0.00015
0.00020
Res
idua
ls
Temperature (K)
Constant Linear Quadratic
200 400 600 800 1000
1.628
1.630
1.632
1.634
1.636
Constant Linear Quadratic
Ln(a
)
Temperature, K
77
J1500 Ceramic
The following figures represent the best fitting model for the J1500 ceramic (see
Figures 4.26-4.27).
Figure 4.26: Instrumentally set temperature vs. residuals (J1500 ceramic)
Figure 4.27: Constant, linear and quadratic fitting with temperature vs. log (LT) (J1500 ceramic)
200 300 400 500 600 700 800 900 1000 1100
-0.0002
-0.0001
0.0000
0.0001
0.0002
0.0003
0.0004
Ln(a
) Cal
c - Ln
(a)
Temperature (K)
Constant Linear Quadratic
200 400 600 800 1000
1.628
1.630
1.632
1.634
1.636
Constant Linear Quadratic
Ln(a
)
Temperature, K
78
4.5 Hardness and Fracture Toughness
The hardness and fracture toughness properties were characterized by collecting
high quality indent images on Optical Microscope. The measurements were performed as
described in Chapter 4 using Image Pro-Plus.
4.5.1 Microindentation of Etched and Non-Etched Samples
Vickers microindentation of mechanically polished surfaces show that hardness
increases with increase of density and consequently with increase of sintering
temperature. Fracture toughness of ceramics decreases with increase of hardness and/or
sintering temperature as shown in Figures 4.28-4.29.
1200 1300 1400 1500 16008
9
10
11
12
13
14
Har
dnes
s(G
Pa)
Sintering temperature(°C)
Non-etched(DKKK) Non-etched(Praxair) Etched(DKKK) Etched(Praxair)
Figure 4.28: Hardness vs. sintering temperature plot for etched and non-etched sample
79
1200 1300 1400 1500 1600
1.8
2.0
2.2
2.4
2.6
2.8
3.0
3.2
3.4
3.6
3.8Fr
actu
re to
ughn
ess(
MPa
·m0.
5)
Sintering temperaure(°C)
Non-etched(DKKK) Non-etched(Praxair) Etched(DKKK)Etched(Praxair)
Figure 4.29: Variation in fracture toughness value with temperature
Hardness of the samples decreases at higher loads. This widely discussed
phenomenon, namely, indentation size effect (ISE), is typical for single crystals and
metals and is almost absent in fine-grained polycrystals. Most of the mechanisms
developed are based on the dislocation nature of the ISE. ISE also can be considered as
an artifact of mechanical polishing and measurement errors. We estimated that our
mechanical polishing procedure can create damage up to 0.2-0.3 µm in depth (surface
nanoroughness range did not exceed 50 nm). To avoid the effect of mechanical polishing
we selected loads which create indents at least 1 µm or 3 times deeper than the estimated
80
threshold of damaged surface layer due to polishing. Simple evaluation of ISE can be
done by linear log H � log h relationship based on empirical power-law relation H = HI*
h-m [50], where h is the depth of indentation, m is power-law exponent of the ISE index,
and H1 is constant corresponding to hardness at the depth of 1 µm. Some typical
impressions at different loads are shown in Figure 4.30.
Load = 1000g Load = 500g Load = 300g
Magnification = 500x Magnification = 500x Magnification = 500x H = 12.69GPa H =12.92GPa H = 13.21GPa K1C = 2.12MPa·m0.5
K1C = 2.12 MPa·m0.5 K1C = 1.79 MPa·m0.5
Load = 200g Load = 100g Magnification = 1000x Magnification = 1000x H = 13.63GPa H = 13.63GPa K1C =1.65 MPa·m0.5
K1C =1.92 MPa·m0.5
(a)
81
1200 1300 1400 1500 16006
7
8
9
10
11
12
13
0
1
2
3
4
5 Praxair Hardness DKKK Hardness
Har
dnes
s (G
Pa)
Sintering Temperature (°C)
Praxair K1C
DKKK K1C Fra
ctur
e To
ughn
ess
(MP
a·m
-0.5)
(b)
Figure 4.30: (a) Effect of sintering temperature on microhardness and fracture toughness at the load 1000g and; (b) Micrograph of typical Vickers indents and cracks at different loads (thermally etched J1600)
Tables 4.7-4.8 contain the hardness and fracture toughness values for each load from
1000g-100g. The weighted value for hardness and fracture toughness was also calculated.
The results obtained for DKKK based and Praxair-based ceramics are listed in Table 4.7
(for non-etched samples), and Table 4.8 (for thermally etched samples). Table 4.9
contains hardness and fracture toughness values for both etched and non-etched samples.
82
Tab
le 4
.7. H
ardn
ess (
H) a
nd fr
actu
re (K
1C) t
ough
ness
val
ues a
t diff
eren
t sin
teri
ng te
mpe
ratu
res
and
load
s for
non
-
etch
ed sa
mpl
es
Sa
mpl
e Pr
oper
ties
1000
g 50
0g
300g
20
0g
100g
A
vg.(H
) G
Pa
Wei
ghte
d (H
) GPa
H
(StD
ev)
8.89
(0.8
7)
8.59
(1.0
3)
8.59
(1.1
7)
8.98
(1.2
8)
8.77
(1.0
7)
8.76
(0.1
7)
8.76
(0.1
7)
J120
0
K1C
(StD
ev)
3.54
(0.4
0)
3.08
(0.4
6)
3.25
-
- 3.
29(0
.23)
3.
34(0
.3)
H(S
tDev
) 10
.64(
0.82
) 10
.45(
1)
10.9
1(0.
7)
11.0
9(1)
11
.77(
1.0)
11
.97(
0.9)
10
.92(
0.4)
J1
300
K1C
(StD
ev)
2.94
(0.4
3)
3.00
(0.3
8)
2.73
(0.2
4)
2.45
2.78
(0.2
4)
2.83
(.18)
H(S
tDev
) 11
.69(
0.72
) 12
.2(0
.58)
12
.66(
0.9)
12
.78(
1.02
) 13
.4(1
.11)
12
.54(
0.6)
12
.35(
0.35
) J1
400
K1C
(StD
ev)
2.72
(0.2
6)
2.87
(0.2
4)
2.78
(0.2
6)
2.56
(0.5
1)
2.
73(0
.13)
2.
77(0
.14)
H(S
tDev
) 12
.70(
0.57
) 13
.23(
0.5)
13
.81(
0.5)
14
.63(
0.51
) 15
.28(
1.25
) 13
.93(
1.03
) 13
.70(
0.25
) J1
500
K1C
(StD
ev)
2.36
(0.1
9)
2.35
(0.3
2)
2.68
(0.2
8)
2.6(
0.32
)
2.49
(0.1
6)
2.46
(0.1
3)
H(S
tDev
) 12
.78(
0.41
) 12
.98(
0.5)
13
.07(
0.4)
13
.20(
0.53
) 13
.32(
0.61
) 13
.07(
0.48
) 13
.03(
0.2)
J1
600
K1C
(StD
ev)
2.06
(0.1
7)
1.92
(0.1
4)
2.08
(0.2
5)
1.93
(0.2
2)
1.91
(0.2
3)
1.98
(0.0
8)
1.97
(0.0
8)
H(S
tDev
) 7.
88(1
.5)
8.69
(1.8
5)
8.81
(1.6
9)
8.84
(1.4
9)
8.93
(2.3
1)
8.63
(0.4
3)
8.57
(0.7
8)
P150
0
K1C
(StD
ev)
3.53
(0.4
8)
3.23
(0.4
) 2.
52(0
.24)
2.
28(0
.18)
2.89
(0.6
) 2.
55(0
.13)
H(S
tDev
) 10
.20(
1.21
) 1.
03(0
.96)
10
.6(0
.74)
10
.86(
0.81
) 11
.70(
1.06
) 10
.74(
0.69
) 10
.74(
0.41
) P1
600
K1C
(StD
ev)
2.01
(0.3
1)
2.33
(0.3
3)
2.31
(0.3
5)
2.37
(0.2
7)
2.07
(0.3
3)
2.24
(0.1
7)
2.24
(0.1
4)
83
Tab
le 4
.8. H
ardn
ess (
H) a
nd fr
actu
re to
ughn
ess
(K1C
) val
ues a
t diff
eren
t sin
teri
ng te
mpe
ratu
res a
nd lo
ads f
or
ther
mal
ly e
tche
d sa
mpl
es
Sam
ple
Prop
erty
10
00g
500g
30
0g
200g
10
0g
Avg
. (G
Pa)
Wei
ghte
d A
vg.(G
Pa)
H(S
tDev
) 9.
05(1
.10)
8.
55(0
.87)
9.
74(0
.88)
9.
94(1
.02)
10
.56(
0.97
) 9.
57(0
.78)
9.
55(0
.43)
J1
200
K1C
(StD
ev)
3.09
(0.2
9)
2.70
(1.0
7)
1.
86(0
.79)
2.55
(0.6
3)
2.93
(0.2
6)
H(S
tDev
) 10
.7(0
.68)
11
.07(
0.54
) 11
.0(0
.64)
11
.91(
0.5)
11
.31(
0.87
) 11
.21(
0.45
) 11
.01(
0.20
) J1
300
K1C
(StD
ev)
2.75
(0.2
7)
3.08
(0.0
8)
2.88
(0.2
1)
2.50
(0.3
0)
2.
80(0
.24)
3.
00(0
.07)
H
(StD
ev)
11.9
(0.5
4)
11.9
4(0.
26)
12.1
5(0.
8)
12.2
(0.7
6)
12.1
5(0.
82)
12.0
(0.1
2)
11.9
9(0.
21)
J140
0 K
1C(S
tDev
) 2.
05(0
.14)
2.
02(0
.29)
2.
29(0
.37)
2.
16(0
.30)
1.
76(0
.42)
2.
06(0
.18)
2.
06(0
.11)
H
(StD
ev)
13.3
(0.2
8)
13.0
9(0.
54)
13.1
(0.6
4)
13.6
(0.5
5)
13.2
5(0.
85)
13.3
(0.2
1)
13.2
8(0.
28)
J150
0 K
1C(S
tDev
) 2.
00(0
.12)
2.
02(0
.16)
2.
23(0
.32)
2.
02(.1
9)
1.74
(0.0
2)
2.00
(018
) 1.
75(0
.02)
H
(StD
ev)
12.7
0.41
) 12
.92(
0.55
) 13
.2(0
.39)
13
.3(0
.45)
13
.63(
0.23
) 13
.15(
0.3)
13
.32(
0.16
) J1
600
K1C
(StD
ev)
1.91
(0.2
1)
2.12
(0.1
5)
2.11
(0.2
8)
1.79
(015
) 1.
65(0
.40)
1.
92(0
.20)
1.
95(0
.09)
H
(StD
ev)
8.81
(1.5
8)
8.44
(1.1
9)
8.88
(1.2
9)
9.49
(1.5
6)
8.85
(0.9
7)
8.89
(0.3
8)
8.84
(0.5
6)
P150
0 K
1C(S
tDev
) 2.
68(0
.53)
2.
81(0
.16)
2.
55(0
.32)
2.
29(0
.53)
1.
69(0
.5)
2.41
(0.4
4)
2.66
(0.1
3)
H(S
tDev
) 10
.471
.0)
9.93
(0.9
9)
10.2
4(0.
3)
10.9
9(1.
0)
10.2
6(0.
54)
10.3
8(0.
4)
10.2
8(0.
27)
P160
0 K
1C(S
tDev
) 2.
11(0
.21)
2.
18(0
.24)
2.
30(0
.34)
1.
70(0
.16)
1.
05(0
.52)
1.
87(0
.51)
1.
93(0
.11)
84
Table 4.9. Comparison of hardness and fracture toughness of non-etched and etched samples
Samples Sintering temperature
Non-etched sample Etched Samples
H(StDev) GPa
K1C (StDev) (MPa·m0.5)
H(StDev) GPa
K1C (StDev) (MPa·m0.5)
P1500 1500 8.57(0.78) 2.55(0.13) 8.84(0.56) 2.66(0.13) P1600 1600 10.74(0.41) 2.24(0.14) 10.28(0.27) 1.93(0.11) J1200 1200 8.76(0.17) 3.34(0.30) 9.55(0.43) 2.93(0.46) J1300 1300 10.92(0.39) 2.83(0.18) 11.01(0.20) 3.00(0.07) J1400 1400 12.35(0.35) 2.77(0.14) 11.99(0.21) 2.06(0.11) J1500 1500 13.70(0.25) 2.46(0.13) 13.28(0.28) 1.75(0.02) J1600 1600 13.03(0.20) 1.97(0.08) 13.22(0.16) 1.95(0.09)
Indentation size effect in hardness value of materials at different indentation
depths was predicted using the empirical power-law as described in Page 80. The results
of this analysis are presented in Table 4.5 and in Figure 4.10(a). Most of the samples
exhibited relatively weak ISE compared to single crystals of metals and semiconductors
(typically m from 0.12 to 0.32). ISE for polycrystals, was found to depend on the grain
size. Generally, no ISE is observed if impression covers several grains, and for large-
grained polycrystals ISE is significant if impression is localized inside one grain.
Nanocrystalline sample J1200 does not exhibit ISE. In the case of other samples,
indentation size for most loads are comparable or exceed grain size. We have noticed that
at higher loads (500 or 1000 g) hardness values are scattered significantly for samples
with large grains. These samples have less porosity. They are dense and are expected to
create more regular plastic deformation zones for every indent. This unexpected behavior
85
of large-grained samples can be explained by significant variation of ISE effect from
indent to indent due to significant scattering of the volume of stored dislocations under
indentation.
The Nix and Gao [51] model relates the ISE to geometrically necessary
dislocations (GNDs). According to this model, density of GNDs is proportional to the
inverse indentation depth. The additional hardening of GNDs is mainly based on total
line length of dislocations, necessary to form the permanent geometry of the impression,
and on the volume, in which the dislocations are stored. Using strain-gradient plasticity
theory, Nix and Gao created a model describing the contribution of GNDs to the hardness
(strain gradient effect) properties of material. Based on this model the strain gradient law
implies the following relation for the hardness:
*
0
1H hH h
= + 4.1
where H is apparent hardness, H0 is characteristic hardness arisen from the statistically
stored dislocations in the absence of geometrically generated dislocations, h indentation
depth, and h* is characteristic length (depth) at which the effect of strain gradient
becomes comparable to the effect of strain.
From the linear dependences of H2 from the reversed indentation depth (see
Figure 4.31), we estimated parameters H0 and h* as listed in Table 4.10. Characteristic
hardness of samples strongly depends on sintering temperature and material density.
Values of h* indicate that geometrically necessary dislocations are limited by grain size.
86
1 2 3 4 5 6
8
9
10
11
12
13
14H
ardn
ess
(GP
a)
Indentation Depth (µm)0.2 0.4 0.6 0.8
60
80
100
120
140
160
180
I-1500 I-1600 II-1200 II-1300 II-1400 II-1500 II-1600
H2 (G
Pa2 )
1/h (µm-1)
(a) (b)
Figure 4.31: ISE analysis using power law (a) and Nix-Gao (b) models
Table 4.10: ISE parameters for Praxair and DKKK-based samples
Power law model Nix-Gao model Sample Sintering Temp.
(ºC) m H1 (GPa) H0 (GPa) h* (µm)
P1500 1500 0.109 9.64 7.61 0.73 P1600 1600 0.116 11.86 9.35 0.71 J1200 1200 0.0 8.76 8.76 0.00 J1300 1300 0.083 11.81 9.92 0.50 J1400 1400 0.109 13.71 11.05 0.61 J1500 1500 0.072 13.97 12.25 0.32 J1600 1600 0.037 13.44 12.58 0.17
87
4.5.2 Nanoindentation
Nanoindentation tests were performed on the ceramic samples J1600 and P1600
using NanoindenterR XP in a continuous stiffness measurement (CMS) option. A three-
sided Berkovich indenter was used with a 60nm tip radius and 65.3° included half-angle.
The indentation was carried out at a depth of 1000nm for ceramic. Sixteen indents for
P1600 and sixty four indents for J1600 samples were created. The data was averaged for
each sample and is listed in Table 4.11.
Table 4.11: Nanoindentation values
Samples Young modulus(GPa) Hardness(GPa) P1600 230 16.5 J1600 210 17.99
It was observed that the hardness and modulus of DKKK ceramic is slightly
higher than the Praxair ceramic sintered at 1600°C shown in Figures 4.32-4.33. It was
also found that the density of J1600 has a slightly larger density than the P1600. The
DKKK samples were fully sintered and had a fewer number of pores and cavities in
comparison to the Praxair ceramic (see Figures 4.8-4.11).There was not much difference
in hardness and modulus data depending on the number of grains covered by the indentor
for both samples.
88
0 100 200 300 400 500
160
180
200
220
240
260
280
300 Praxair, P1600 DKKK , J1600
You
ng's
Mod
ulus
(GP
a)
Indentation Depth (nm)
Figure 4.32: Nanoindentation Young�s modulus results
0 100 200 300 400 50010
12
14
16
18
20
22
24 Praxair P1600 DKKK J1600
Har
dnes
s (G
Pa)
Indentation Depth (nm)
Figure 4.33: Nanoindentation hardness results
89
CHAPTER 5
CONCLUSION AND FUTURE WORK
5.1 Conclusion
The following conclusions can be made based on the results of the research work
outlined:
• Phase and chemical composition, morphology, and grain size of the 10 mol%
Sc2O3 � 1 mol% CeO2 - ZrO2 ceramic powders produced by Praxair Specialty
Ceramics, USA and by DKKK, Japan has been characterized.
• The sinterability of the powders has been studied by sintering at 6 different
temperatures (1100ºC -1600ºC).
• The sintering rate of the ceramic powder manufactured by DKKK is much higher
than the powder manufactured by Praxair. The DKKK powder was able to reach
its lowest porosity level at lower temperatures and shorter sintering times than the
Praxair powder.
• DKKK-based ceramics have phase instability in the region 300-500ºC.
Transformation from c- to b-phase at these temperatures is slow process and it is
required to use high heating and cooling rates to avoid formation of significant
amounts of b-phase.
• The ionic conductivity of selected samples is reported to be ~0.088 S/cm for
Praxair Ceramics and ~0.156 S/cm for DKKK ceramics at 800ºC.
90
• Coefficients of thermal expansion in the range RT-800ºC of both types of
ceramics are in the range 10.5 � 10.6×10-6 grad-1 and very close to CTE of cubic
zirconia. CTE of DKKK samples have less linear dependency on temperature than
Praxair samples.
• Young�s modulus and hardness of DKKK samples are higher than Praxair
samples.
• High temperature fracture toughness measurements confirm that these ceramics
are very brittle, especially at elevated temperatures. The highest K1C was 2.5
MPa×m.5 at 300ºC for DKKK ceramics.
5.2 Future Work
The following are recommendations for future work:
• Study of mechanical and microstructure properties of J1300 samples by changing
the dwell time and heating and cooling rate.
• Study of mechanical and microstructure properties of rhombohedral phases of
samples J1300-J1600
91
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