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4/27/12 Max Planck Institute for Metals Research - Department Mittemeijer: Research Areas

1/14www.mf.mpg.de/en/abteilungen/mittemeijer/english/research/researchareas.htm

Areas of Scientific Work

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1. Phase Transformation Kinetics

1.1 Analysis and modelling of solid state phase transformation kinetics

1.2 Abnormal and normal kinetics; the austenite-ferrite phase transformation

1.3 Computer simulation of phase transformation kinetics and microstructure

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2. Phase transformations in Laves phases

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3. Reactions and phase transformations at surfaces and interfaces

3.1 Initial oxide-film growth on bare metals and alloys

3.2 Nitriding and nitrocarburizing

3.2.1 Nitride precipitation and coarsening; coherent and incoherent diffraction effects

3.2.2 Microstructure and residual stress

3.2.3 Unexpected formation of iron nitride by gas nitriding of nanocrystalline iron films

3.2.4 Unusual and extreme elastic anisotropy of γ'-Fe4N and of cementite, Fe3C

3.3 Diffusion-induced phase transformation in thin films systems

3.3.1 Metal-induced crystallization; tailoring the crystallization temperature

3.3.2 Driving force for Sn-whisker growth in the system Cu-Sn

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4. Central Scientific Facility 'X-ray diffraction' (CSFX)

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Introduction

The approach to understanding phase transformations has to be of both micro- and macro-scopical character. Apart from understanding transformation mechanisms on an atomicscale, the challenge is to use the detailed microstructural knowledge in building'mesoscopical' transformation models. Computer simulations of diffusion and phasetransformations will be indispensable.

The department is "experiment driven". Research of the type as described here requirespowerful experimental methods to characterise the (changing) microstructure of materials.The department excels in the application of X-ray diffraction methods, surface analyticaland calorimetric and dilatometric methods. Methodological research will also always be anim-portant part of our research activity.

1. Phase Transformation Kinetics

Scientists: Prof. Dr F. Sommer, Dr C. Bos, Prof. Dr F. Liu, Prof. Dr Y.C. Liu.Doctoral Students: R. Bauer, E.Jägle

1.1. Analysis and modelling of solid state phase transformation kinetics

On the basis of a general modular model for the quantitative description of solid-state phase



transformation kinetics it has been shown that, even if the kinetic parameters (e.g. growthexponent, activation energy) distinctly vary during the course of transformation, the trans-formation mechanism can be the same [Intern. Mater. Reviews 52 (2007) 573; J. Mater.Sci. 42 (2007) 573]. The applicability of the models proposed has been tested forisothermally and isoch-ronally conducted phase transformations both by numericalsimulations and by application to the crystallization kinetics of amorphous alloys and thekinetics of precipitation in supersatu-rated solid solutions.

1.2. Abnormal and normal kinetics; the austenite-ferrite phase transformation

Abnormal transformation kinetics, characterized by the presence of two maxima in thetrans-formation rate curve (Fig. 1), was recognized upon isochronal austenite(γ) -ferrite(α)trans-formation of an ultra-low-nitrogen (0.005 at.%N) Fe-N alloy. The main (later), normaltrans-formation stage of this transformation can be well described by a model based on sitesatura-tion, interface-controlled growth and an appropriate impingement correction.

Fig. 1: The ferrite-formation rate, dfα /dt, as a function of ferrite fraction, fα ,

for Fe- 0.005at.%N at various cooling rates.

The abnormal part of the transformation is due to autocatalytic nucleation of ferritic grainsin advance of the migrating γ/&alpha: interface, which is caused by the accommodation ofmisfit stain energy [Metall. Mater. Trans. A39 (2008) 2306]. The kinetics of the γĺαtransformation is greatly influenced by the development of misfit stress, in flagrant contrastwith usual concepts based on only chemical driving forces. An uncommon transition ofdiffusion-controlled to interface-controlled transformation was found and explained for theγĺα transformation in ultra-low-carbon Fe-C alloys upon isochronal cooling [Int. J. Mater.Res. 99 (2008) 925]; upon isothermal transformation subsequent interface-controlled(extremely fast) and diffusion-controlled (sluggish) growth modes could be identified andmodelled successfully adopting site-saturation for both stages [Acta Mater. 56 (2008)3833].



1.3. Computer simulation of phase transformation kinetics and microstructure

The movement of interphase boundaries during the interface-controlled phasetransformation from austenite to ferrite in iron was studied by atomistic simulations. It couldbe concluded that (i) a single atomic jump at the interface is not the rate determining step,but, instead, the transformation rate is controlled by a series of unfavourable jumps bygroups of atoms at the interface, and that (ii) the amount of excess volume associated withan interface as well as its distribution has a significant influence on interface mobility andboundary self diffusion and that the effect is stronger for the latter. [Phil. Mag. 87 (2007)2245].If a valid description of phase transformation kinetics is available, the challenge is to de-scribe the evolving microstructure upon phase transformation. The aim of ongoing work is topredict the grain-size distribution and by comparison with corresponding experimentallydetermined data, to identify the underlying mechanisms of the phase transformation. Near-future work will investigate the effect of residual stresses (e.g. caused by microstructuralchanges) on interface movement. This work is carried out as part of the Cluster ofExcellence "Simulation Technology" at the University of Stuttgart.

2. Phase transformations in Laves phases

Scientists: Dr A. LeineweberDoctoral Students: J. Aufrecht, W. Baumann

Within the Max Planck Society the inter-institutional research initiative "The nature ofLaves phases" was founded in 2005 (with groups from MPI-CPfS, Dresden, MPIE,Düsseldorf, MPI-FKF, Stuttgart and MPI-MF, Stuttgart) to create a platform for experimen-tal and theoretical investigations of intermetallic phases with strong interdisciplinary charac-ter. Within this initiative we study the mechanism and kinetics of polytypic phase transitionsin Laves phases of the binary systems Ti-Cr, Hf-Cr and Nb-Cr. It was found that Laves-phase material having previously experienced a polytypic phase transformation contain spe-cific types of stacking faults, which can be characterised using X-ray powder diffraction andhigh-resolution transmission electron microscopy (Fig. 2).



Fig. 2: Stacking fault in C36-TiCr after phase transition from C14-TiCr2.

The type of stacking faults in C36-TiCr2 and C36-NbCr2 is indicative of a partial-dislocation

pair mechanism driving the phase transition starting hfrom a parent C14 phase.

3. Reactions and phase transformations at surfaces andinterfaces

Scientists: Dr L.P.H. Jeurgens, Dr A. Leineweber, Dr R. Schacherl, Dr J.Y. Wang, Dr Z.Wang, Dr U. Welzel.Doctoral Students: : G. Bakradze, T. Gressmann, M. Nikolussi, E. Panda, F. Reichel, J.F.Sheng, M. Sobiech, M.S. Vinodh, M. Wohlschlögel, T. Wöhrle

3.1 Initial oxide-film growth on bare metals and alloys

We have shown that ultrathin (< 5 nm) amorphous oxide films (Fig. 3), as well as ultrathin(< 5 nm) crystalline oxide films with an exceptionally high lattice mismatch with respect tothat of the substrate [Acta Mater. 55 (2007) 6027], can be stabilized by relatively lowsurface and interface energies. Only above some critical oxide-film thickness, atransformation into an, according to bulk thermodynamics, more stable oxide structure isrealized [Acta Mater. 56 (2008) 659].Employing a model developed in our group [Phys. Rev. B 74 (2006) 144103], thethermodynami-cally preferred microstructures of thin oxide overgrowths on their metalsubstrates were pre-dicted, in very good agreement with experimental data, as function ofgrowth temperature and orientation of the substrate (Fig. 4) [Acta Mater. 56 (2008) 659; J.Appl. Phys. 103 (2008) 093515; Surf. Interf. Anal. 40 (2008) 259; Acta Mater. 56 (2008)4621].



Fig. 3:HR-TEM micrograph of anamorphous Al2O3 oxide film grown on

a bare Al{111} substrate by thermaloxidation for 6000 s at 373 K and pO2

= 1×10-4 Pa.

Fig. 4:Critical thickness up to whichan amorphous oxide overgrowth isthermodynamically preferred asfunction of the growth temperature.

3.2 Nitriding and nitrocarburizing

Nitriding and nitrocarburizing are widely used thermochemical surface treatments to greatlyimprove mechanical and chemical properties of iron-based materials. These treatments in-volve the uptake of the initially interstitially dissolved elements nitrogen and/or carbon andlead to the development of interstitial compounds as layers at the surface and/orprecipitates in the matrix. This project focuses on unusual, hitherto not understood,fundamental aspects of the phase transformations involved, with an emphasis on the rolesof misfit stress/distortion and the interplay of thermodynamics and kinetics.

3.2.1 Nitride precipitation and coarsening; coherent and incoherent diffraction effects

The formation of alloying element nitrides leads to a long range strain field in the ferritematrix surrounding the precipitate. Diffraction analysis of nitrided samples has revealed thatseparate reflections of inner nitrides are not observed and that the ferrite reflexions are ex-tremely broadened and distorted. HRTEM investigations (see Fig. 5) show that the VN plate-lets are continuous, but that the atomic planes are slightly bent, and defects asdislocations can be discerned in the platelets. Because of the strong tetragonallyanisotropic misfit be-tween precipitates and matrix, the coherently diffracting domain,comprising the precipitate and the surrounding matrix, can be considered as a b.c.t. "phase"(see Fig. 6). Thus the X-ray diffraction patterns obtained from such specimens originatefrom diffraction by a b.c.t. "phase" (VN platelets plus the surrounding tetragonally distortedferrite) and by a b.c.c. "phase" (cubic ferrite), see Fig 6 [Acta Mater. 56 (2008) 4137].



Fig. 5: HR-TEM for a nitrided specimen of Fe-2.23at.%V alloy. (a) VN plateletwith an indication of defects (see arrows). (b) Detail of a single platelet (seearrow), revealing the bending of the lattice planes at the location of theplatelet.

Fig. 6: (a) Schematic view of a VN platelet coherent with the surroundingferrite matrix. (b) Schematic view of the contributions of the tetragonallydistorted b.c.t. "phase" (VN platelet plus surrounding ferrite) and the cubicb.c.c. "phase" (cubic ferrite) to the total diffraction profile.

3.2.2 Microstructure and residual stress

The precipitation of alloying-element nitrides as CrN in the matrix of iron-based substrates isusually thought to lead to compressive residual stress in the surface region, which is consid-ered responsible for the pronounced increase of the fatigue resistance. Detailed XRD analy-sis of the residual stress depth profiles in nitrided Fe-Cr alloys has shown that this picturecan be much too simple: highly complicated dependencies of stress on depth can occurwith, unusually alternating, regions of tensile and compressive stress.



Fig. 7: (a) Residual stress depth-profile measured for a specimen of nitridedFe-4 wt.% Cr alloy. (b) Schematic representation of stress development uponnitriding.

The residual stress depth profile development could successfully be modelled as the conse-quence of stress build up and relaxation processes in the microstructure: see Fig. 7 [ActaMa-ter. 56 (2008) 1196].

3.2.3 Unexpected formation of ε iron nitride by gas nitriding of nanocrystalline α ironfilms

It was surprisingly found that formation of ε-Fe3N1+x occurs upon nitriding of

nanocrystalline iron at thermodynamic conditions at which pure γ'-Fe4N1-x should form

according to bulk thermodynamics: see the double tangent constructions pertaining to the(α/gas,) γ'/gas and ε/gas equilibria, on the basis of the Gibbs energies as function ofnitrogen content, in Fig. 8.



Fig. 8:Gibbs energies as functions of the atom fraction nitrogen, xN, for α-Fe,

γ' and ε at T = 550°C, p = 1 atm and r = 1 atm-1/2.

Now, due to the curvature of a (spherical) nanocrystalline particle an extra pressure Δp =2γ/r, with γ as the interfacial energy and r as the radius of the particle, acts on it givingrise to an increase in Gibbs energy (Gibbs-Thomson effect). It followed that a radius of

40nm and an interfacial energy of 1.6 J/m2 for the γ' crystallites suffice tothermodynamically stabilize the ε phase. [Appl. Phys. Lett. 91 (2007) 141901].

3.2.4 Unusual and extreme elastic anisotropy of γ'-Fe4N and of cementite, Fe3C

Many basic physical properties of the transition-metal nitrides are unknown, due to the dif-ficulty to grow single crystals or to prepare these materials in massive solid form. By use oftotal-energy calculations and by comparing the results with X-ray-diffractionmeasurements, single-crystal elastic constants could be determined for ' iron nitride and forcementite. The results show that the interstitial atoms dissolved in the metal matrix candramatically change the elastic properties of the phase. For γ'-Fe4N a reversal of the elastic

anisotropy occurs with respect to fcc iron (austenite; Fig. 9a) [Acta Mater. 55 (2007)5833] and cementite was shown to be extremely elastically anisotropic (Fig. 9b) [ScriptaMater. 59 (2008) 814].



Fig. 9:(a)Elastic anisotropy of γ'-Fe4N. Reversal of the elastic

anisotropy of γ'-Fe4N in comparison

with austenite (fcc; dotted line).Fig. 9:(b) Elastic anisotropy ofcementite.

3.3 Diffusion-induced phase transformation in thin films systems

3.3.1 Metal-induced crystallization; tailoring the crystallization temperature

The transformation of solid silicon and germanium from its amorphous to its crystalline formis one of the most important processes in semiconductor technologies. This crystallizationprocess is obstructed at low temperatures due to the strong, covalent bonding between theSi and Ge atoms in the amorphous solid: amorphous Si and Ge (a-Si, a-Ge) can only becrystal-lized at high temperatures around 700ºC. Yet, crystallization can be induced at lowtempera-ture if the a-Si or a-Ge is in contact with crystalline Al. We have demonstrated,both theo-retically and experimentally, that the thickness of a thin-enough Al film (< 20 nm)on top of a-Si provides an extremely sensitive tool to tailor the crystallization temperatureof a-Si, down to as low as 150ºC (Figs. 10 and 11).

Fig. 10:The low-temperature crystallization of a-Si in the presence of anultrathin aluminum overlayer. Initially, grain boundaries in the Al over-layer get'wetted' by a-Si. Beyond a critical thickness of the wetting a-Si film,crystallization can initiate at these wetted Al grain boundaries.



Fig. 11:(a) Dependence of the crystallization temperature of amorphous siliconon the aluminum overlayer thickness: (b) In-situ observation by HRTEM of thecrystallization and growth of a Si grain within an original Al grain boundaryduring annealing of an a-Si/Al bilayer at 150ºC.

The temperature at which the crystallization occurs is governed by a delicate balance ofthe crystallization energy and the crystallization-induced changes of the interface andsurface energies. This energy balance can be adjusted through a variation of the Aloverlayer thick-ness, leading to a striking dependence of the crystallization temperature ofa-Si on the Al overlayer thickness (Fig. 11a). This work demonstrates in general howsurface and interface energetics control diffusion, wetting and phase transformations inlow-dimensional systems. [Phys. Rev. Lett. 100 (2008) 125503; Phys. Rev. B 77 (2008)045424].

3.3.2 Driving force for Sn-whisker growth in the system Cu-Sn

Sn coatings on Cu substrates are prone to the formation of Sn whiskers on the Sn surface(Fig. 12). The Cu6Sn5 formation, occurring due to preferential Cu diffusion along the Sn

grain boundaries intersecting the Cu/Sn interface, is accompanied by a volume expansionand as a consequence, residual compressive stresses can be generated in the Sn thin films(Fig. 13). Near-surface stress-depth profiles for specimens which do exhibit and which donot exhibit Sn-whisker growth as obtained upon ageing at ambient temperature for speci-mens after one day (before whisker formation for the specimen exhibiting whisker growth)and after about 8 months (during whisker growth for the specimen exhibiting whiskergrowth), investigated by laboratory and synchrotron X-ray diffraction stress measurements,revealed, for the first time, a vertical stress gradient as driving force for spontaneous Sn-whisker formation (Fig. 14) [Appl. Phys. Lett. 93 (2008) 011906.].



Fig. 12:Scanning electronmi-croscopy images of Sn-whiskers growing on anelectrodeposited Sn coat-ing.

Fig. 13:Illustration of theprocesses taking place inSn thin films (thicknessesof several microns)deposited on Cu substratesupon ageing at ambienttemperatures.

Fig. 14:Near-surfacestress-depth profiles forwhisker-prone and whisker-resistant specimens for anageing time of 8 months.

4. Central Scientific Facility 'X-ray diffraction' (CSFX)

Head: Dr U. Welzel Doctoral Students: J. Chakraborty, A. Kumar, Y. Kuru, M. Sobiech, M. Wohlschlögel

The central scientific facility X-ray diffraction (CSFX) is a laboratory for powder diffraction,i.e. the diffraction analysis of polycrystalline materials. Apart from qualitative and quantita-tive phase analysis, a broad variety of dedicated analyses can be performed: textureanalysis, (macro)stress measurements, analysis of the crystalline imperfection (i.e. analysisof the diffraction line broadening), non-ambient X-ray diffraction measurements, structuredeter-mination from powder patterns, and - for thin film investigations - high resolutiondiffracto-metry and reflectometry.The CSFX, in an integrated way with the department Mittemeijer, is a centre of excellencein the following areas:

The CSFX is world-leading in the diffraction analysis of mechanical stresses [seereview J. Appl. Cryst. 38 (2005) 1], the investigation of elastic grain interaction inpolycrystalline materials [see review JOM 59 (2007) 66] and line broadening analysisfor the investigation of the crystalline imperfection [see review Z. Kristallogr. inpress].Diffraction instrumentation: Use of new optical components; parallel beam diffractionem-ploying laboratory diffractometers and non-ambient diffraction measurements [J.Appl. Cryst. 41 (2008) 124].Non-ambient, in-situ diffraction: In-situ investigation of stress evolution, interdiffusionand phase formation in thin film systems [J. Appl. Cryst. 41 (2008) 428; Thin Solid Film516 (2008) 7615].



Fig. 15:(a) Real-time stress evolutions in (a) the Al phase and (b) thecrystallized Ge phase as function of temperature during thermal cycling of ana-Ge/c-Al bilayer specimen (in situ XRD analysis).

An example of the impact of our methodological research is the following (cf. alsoSection 3.3.3.1): The real-time evolution of stress during metal-induced crystallization(MIC) and subsequent layer exchange in amorphous a-Ge/Al bilayers was explored byin situ X-ray diffraction (XRD) analysis in the temperature range of 25 ºC to 250 ºC[Acta Mater. 2008, doi:10.1016/j.actamat. 2008.06.026] (see Fig. 15). Lateral graingrowth of c-Ge at the location of the original Al sublayer occurs at the expense of thesmaller nanograins of c-Ge nucleated at the location of the original a-Ge sublayer.This Ostwald ripening process at anomaly low temperatures leads to a gradual



exchange of the positions of the Al and crystallized Ge sublayers in association withthe observed stress change.

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