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					                    AO-1999-114 METCOMP
       METASTABLE SOLIDIFICATION OF COMPOSITES: NOVEL
        PERITECTIC STRUCTURES AND IN-SITU COMPOSITES

                           M. Kolbe, G. Eggeler, D. Herlach (D),
                                      T. Pusztai (H)
                                     L. Granasy (UK),
                                      A. Ludwig (A),
                                     M. Rappaz (CH),

                        Corus Technology BV (NL),
   Schwermetall GmbH&Co KG, Thyssen Krupp Stahl AG, Wieland-Werke AG (D)
                           Swissmetal SA (CH),

                                         ABSTRACT
     The aim of this project is to improve the processing of commercial peritectic alloys through
  microstructure control, e.g. through the development of a phase selection model, and of a model
  predicting the pushing/engulfment of particles by a growing dendritic front. The output of these
microstructure models will be compared to experimental data obtained in 1- and in 0-g environments.
                                                                                                                                   METCOMP MAP 114




List of Contents

I                Participants ........................................................................................................................3
      I.1     Academic Partners ............................................................................................................... 3
      I.2     Industrial Partners ................................................................................................................ 4
      I.3     Distribution List................................................................................................................... 5
II                Introduction .......................................................................................................................6
III               Reports...............................................................................................................................8
      III.1   Directional solidification experiments on metallic peritectic alloys (EPFL/Rappaz, Kohler)
                     ...................................................................................................................................... 8
      III.2   Investigations on transparent model systems of organic peritectics (MUL/Ludwig, Eck,
                    Mogeritsch) ................................................................................................................ 16
      III.3   Phase field modelling of solidification of peritectics with particles (RISSPO/Pusztai) ... 28
      III.4   Peritectic Cu-alloys and solid inclusions (DLR/Herlach, Kolbe) ..................................... 37
      III.5   Peritectic Ni- and Fe-alloys and solid inclusions (RUB/Eggeler, Lierfeld, Wu) .............. 41
      III.6   Publication list ................................................................................................................... 51
IV                  Inventory of Intellectual Property Rights ........................................................................53
V                   Evaluation of Technology ...............................................................................................53
VI                  Initiatives towards non-space industries..........................................................................54




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                                                                 METCOMP MAP 114




I     Participants

I.1     Academic Partners
Dr. Matthias Kolbe (Team co-ordinator)
German Aerospace Center (DLR)
Institut für Materialphysik im Weltraum
D - 51170 Köln, Linder Höhe, Germany
Tel: +49-2203-601-3019, Fax: +49-2203-61768
e-mail: matthias.kolbe@dlr.de
http://www.dlr.de/mp/

Prof. Dr. G. Eggeler
Ruhr-University Bochum
Institute of Materials IA 1
D - 44780 Bochum, Universitätsstr. 150, Germany
Tel: +49-234-32-23022, Fax: +49-234-32-14235
e-mail: gunther.eggeler@ruhr-uni-bochum.de
http://www.ruhr-uni-bochum.de/ww/

Prof. Dr. Andreas Ludwig
University of Leoben (MUL)
Department of Metallurgy
A - 8700 Leoben, Franz-Josef-Strasse 18, Austria
Tel.: +43-3842-402-2221, Fax: +43-3842-402-2202
e-mail: ludwig@unileoben.ac.at
http://www.smmp.at.hm

Dr. Tamás Pusztai
Research Institute for Solid State Physics and Optics (RISSPO)
Department of Experimental Solid State Physics
Konkoly Thege M. út 29-33
H - 1121 Budapest, Hungary
Tel.: +36-1-392-2222 ext 3371, Fax: +36-1-392-2219
e-mail: pusztai@szfki.hu
http://www.szfki.hu

Prof. Dr. Michel Rappaz
École Polytechnique Fédérale de Lausanne (EPFL)
School of Engineering, Institute of Materials
Computational Materials Laboratory
Station 12
CH - 1015 Lausanne, Switzerland
Tel.: +41-21-693 2844, Fax: +41-21-693 5890
e-mail: michel.rappaz@epfl.ch
http://lsmx.epfl.ch/




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                                                      METCOMP MAP 114


I.2   Industrial Partners

Dr. Huib Wouters
Corus Technology BV
Molten Aluminium Processing Group
Ijmuiden Technology Centre
PO Box 10000, The Netherlands
Tel.: +31-251-49 7553, Fax +31-251-47 0265
e-mail: huib.wouters@corusgroup.com

Dr.-Ing. Jürgen Jestrabek
Schwermetall Halbzeug-Werk GmbH&Co.KG
Breinigerberg 165
D - 52223 Stolberg, Germany
Tel.: +49-(0)2402-761213, Fax: +49-(0)2402-761210
e-mail: jestrabek@schwermetall.de

Dr. Claudio Penna
Swissmetal SA
CH - 4143 Dornach, Switzerland
Tel.: +41-61-7053220, Fax: +41-61-7053347
e-mail: claudio.penna@swissmetal.com

Dr. Jürgen Stahl
Thyssen Krupp Steel AG
Kaiser-Wilhelm-Straße 100
D - 47166 Duisburg, Germany
Tel.: +49-(0)203-52-44210, Fax: +49-(0)203-52-25721
e-mail: juergen.stahl@thyssenkrupp.com

Dr.-Ing. Wolfram Schillinger
Wieland-Werke AG
Zentrallabor und Entwicklung
Graf-Arco-Straße 36
D - 89079 Ulm, Germany
Tel.: +49-(0)731-944-6302, Fax: +49-(0)731-944-4461
wolfram.schillinger@wieland.de
http://www.wieland.de/




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                                                                                    METCOMP MAP 114




I.3    Distribution List

Dr. Olivier Minster, ESTEC, ESA
Dr-Ing. Daniela Voss, ESTEC, ESA

all participants (each 1 copy)
Prof. Dr. Gunther Eggeler, Institute of Materials, RUB
Dr. Matthias Kolbe, German Aerospace Center, DLR
Prof. Dr. Andreas Ludwig, University of Leoben, MUL
Dr. Tamás Pusztai, Research Institute for Solid State Physics and Optics (RISSPO)
Prof. Dr. Michel Rappaz, École Polytechnique Fédérale de Lausanne, EPFL


all industrial partners (each 1 copy)
Dr.-Ing. Jürgen Jestrabek, Schwermetall Halbzeug-Werk GmbH&Co.KG
Dr. Claudio Penna, Swissmetal SA
Dr.-Ing. Wolfram Schillinger, Wieland-Werke AG
Dr. Jürgen Stahl, Thyssen Krupp Stahl AG
Dr. Huib Wouters, Corus Technology BV




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                                                                                      METCOMP MAP 114




II     Introduction

Many commercially used alloys as steels and many copper-based alloys are peritectic alloys. They are
widely used and have high economic importance. The peritectic reaction follows the scheme
Lwhere a solid phase  in contact to a liquid phase L transforms into a solid phase  The
transformation is slow, since diffusion in solid is required. As a direct consequence, peritectic
reactions are in most cases far out of thermodynamic equilibrium and are very much influenced by
convection phenomena in liquid, which change heat and mass transport. More recently, new
metastable microstructures were detected in solidification of peritectics, which make them suitable for
the preparation of in situ composite materials. Nevertheless, the solidification behaviour is not well
understood in particular, because the formation of in situ composite microstructure from liquid
sensitively depends on convection that is always present under 1g-conditions.
The strength and the wear resistance of metallic materials are often improved by the introduction of
particles of a second phase. If these particles are grown by annealing from a supersaturated solution,
than the particles are thermodynamically not stable and the material has limited high temperature
stability. Ceramic particles would be desirable because of their high temperature stability, but it is
difficult to grow them by annealing. If the material is produced by casting, pushing of particles by an
advancing solid-liquid interface and sedimentation of particles due to buoyancy effects in liquid lead to
inhomogeneous distribution of the particles in the matrix resulting in low fracture toughness. The
interaction between a solid-liquid interface and ceramic particles has been investigated during
directional solidification of a planar interface both under terrestrial and reduced gravity conditions. In
general, embedding of particles is favoured for large particles and rapid growth, while pushing is
expected at small particle size and small growth velocity of the crystal. Dendritic solidification is of
much higher practical importance in casting processes, but a description of the conditions of
homogeneous distribution of particles in a metallic matrix is still lacking. In contrast to a planar
interface, particle embedding takes place in the interdendritic liquid regime and by kinetic trapping
into the stem of a rapidly propagating dendrite. This process is very complex: it is controlled by the
growth dynamics of the interface, transport phenomena in liquid due to convection, sedimentation and
mainly by the particle size. On the other hand industrial production processes require particle pushing
during casting processes in order to remove foreign phases and purify the cast material.
In 2000 the project “Metastable Solidification of Composites: Novel Peritectic Structures and In-Situ
Composites” (METCOMP) was established by ESA. The project was continued with an enlarged
consortium with starting date 1 May 2003. The scientific members are: Prof. Dr. G. Eggeler (Ruhr-
University Bochum, Germany), Dr. M. Kolbe (DLR), Prof. Dr. A. Ludwig (University of Leoben,
Austria), Dr. T. Pusztai (RISSPO, Hungary) and Prof. Dr. M. Rappaz (EPFL, Switzerland); the
industrial members are: Dr. H. Wouters (Corus Technology BV, The Netherlands), Dr.-Ing. J.
Jestrabek (Schwermetall Halbzeug-Werk GmbH & Co. KG, Germany), Dr. C. Penna (Swissmetal SA,
Switzerland), Dr. J. Stahl (Thyssen Krupp Stahl AG, Germany) and Dr.-Ing. W. Schillinger (Wieland-
Werke AG, Germany).
The partner at Ecole Polytechnique Federale de Lausanne (CH) studied peritectic solidification in Cu-
Sn and showed that bands and lamella can form in a large solidification interval alloy. It has been
demonstrated that bands can occur via an overlay, rather than nucleation-growth, mechanism. During
the lateral growth of one phase, destabilisation can lead to lamellae. Multi-phase field simulations
indicate that cooperative growth occurs with the peritectic phase slightly ahead of the primary phase,
and thus the front is not strictly isothermal. Experiments with thin capillaries for reducing convection
have been conducted. Thermal tests were performed in collaboration with ESA, EADS and


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                                                                                    METCOMP MAP 114


SOTEREM with Cu-10wt% Sn in the METCOMP cartridge for flight experiments in the Science
Reference Model of SQF. The results have been modeled with good agreement.
The partner at the University of Leoben studies transparent peritectic model systems. Oscillatory and
coupled growth of  and  phases have been observed in-situ during solidification experiments at very
low pulling speed. Investigations have been carried out to define the boundary conditions and
experimental parameters for the intended µg-experiments and support has been given to the design of
the Directional Solidification facility, DIRSOL, which is intended to be used on the ISS. Physical
properties of the two organic compounds of interest, TRIS and NPG, have been measured for NASA
safety check.
The partners at DLR and at RUB (Bochum) cooperate closely to maximise synergy. Parabolic flights
have been used frequently for experiments in the TEMPUS facility. The results show clearly the need
for micro-gravity solidification experiments, as it is possible to reduce convective effects on
microstructure considerably. Particle pushing and engulfment has been observed in Ni-Ta and Cu-Ni
alloys with Ta2O5 particles. The dependence of the engulfment fraction on alloying element Ta and on
the change in wetting conditions in Cu-Ni alloys has been measured. Experiments in the TEMPUS
facility are precursor experiments to the use of MSL-EML on ISS.
The partner at RISSPO (Hungary) is applying phase field modeling to peritectic transformations with
and without foreign particles. Depending on the nature of the foreign particles the growing dendrite is
influenced: It may be deflected by a particle or even tip splitting might occur. Modeling results
concerning the interaction of dendrite growth in the Cu-Ni system with Ta2O5 particles correlate with
findings from solidification experiments in TEMPUS during parabolic flights (DLR and RUB)
Removed




                                                  7
                                                                                                    METCOMP MAP 114




III       Reports


III.1 Directional solidification experiments on metallic peritectic alloys (EPFL/Rappaz,
Kohler)

This part of the project focuses on the competition that can occur in peritectic alloys solidified at very
low velocity, between the primary and peritectic phases, which will be called  and  in this report,
regardless of the system being investigated. For Fe-Ni and Ni-Al, which are two alloys with a very
small solidification interval, several microstructures were observed at low speed for hypoperitectic
compositions, in particular: (1) alternation of  and  bands growing parallel to the isotherms, (2)
islands of one phase embedded in the other, (3) lamellae of  and  growing as for eutectics in a
coupled (or cooperative) way with a nearly-isothermal front. Microstructures of types (1) and (2) were
interpreted as a succession of nucleation and growth events. Since neither the  nor the -planar
fronts are stable under steady-state conditions, one phase nucleates at the solidification front of the
other, then grows laterally to cover totally (1) or partially (2) the other phase. In case (1), as the new
phase cannot reach steady-state, this process starts again with an inversion of the roles. In case (2), the
new phase does not cover entirely the planar front of the other and becomes like islands embedded in a
matrix. In the coupled or cooperative growth mode (3), solute elements are rejected ahead of a planar
front made of - and -lamellae, with an average diffusion layer equal to Dℓ/V, where Dℓ is the
diffusion coefficient in the liquid and V the growth rate. That is a modulation of this average diffusion
layer by lateral exchanges of solute between the lamellae, but this is greatly reduced as compared to
eutectics. Furthermore, there is a dispute whether the temperature of the front is above or below the
peritectic temperature. For more details, see the previous METCOMP Final Report (AO 98/99 – 114,
ESTEC Contract No. 14243/00/NL/SH 2006) and references [1,2,3,4,5,6,7,8,9,10].

In order to see if similar phenomena can be observed in large solidification interval alloys, a research
project was launched to study the Cu-Sn system (bronzes), which is important from an industrial point
of view. This system exhibits two interesting properties:
         The equilibrium solidification interval of the -phase decreases with the tin concentration in
          the hypoperitectic composition range (see Figure 1).
         The equilibrium solidification interval of the primary phase in the Cu-Sn system is about 25
          times larger than in the Fe-Ni alloy.
However, several difficulties are also inherent to this system. First, tin is lighter than copper and
solutal macrosegregation is therefore expected. Second, bronzes exhibit several solid state
transformations, making the interpretation of microstructures more difficult. Third, liquid bronze does


[1] S. Dobler, T.S. Lo, M. Plapp, A. Karma, and W. Kurz, Acta Mater. 52 (2004) 2795
[2] O. Hunziker, M. Vandyoussefi and W. Kurz, Acta Mater. 46 (1998) 6325
[3] A. Karma, W.J. Rappel, B.C. Fuh, and R. Trivedi, Metall. Mater. Trans. 29A (1998) 1457
[4] R. Trivedi, Metall. and Mater. Trans. 26A (1995) 1583
[5] R. Trivedi, A. Karma, T.S. Lo, J.S. Park, and M. Plapp, Dynamic pattern formation in the two phase region of peritectic
systems. (1998) Zermatt, Switzerland (CD ROM, EPFL, Lausanne, Switzerland)
[6] R. Trivedi and J.S. Park, J. Crystal Growth 235 (2002) 572
[7] T.S. Lo, S. Dobler, M. Plapp, A. Karma, and W. Kurz, Acta Mater. 51 (2003) 599
[8] R. Trivedi, Scripta Mater. 53 (2005) 47
[9] T.S.Lo S.Dobler, M.Plapp, A.Karma, W.Kurz, Acta Mater. 52 (2004) 2795
[10] H.W. Kerr and W. Kurz, International Mater. Rev. 41 (1996) 129

                                                             8
                                                                                                METCOMP MAP 114


not wet well most crucibles. Therefore, during the previous round of the METCOMP project, the
contributions of EPFL were essentially to:
       Design the whole experimental procedure to produce the alloys, solidify specimens of reduced
        dimensions in a high thermal gradient Bridgman furnace with a liquid metal cooling bath and
        develop observation and analytical techniques (metallography, SEM, EDX, WDS, EBSD, etc).
        The geometry of the Cu-Sn specimen that was finally selected for ground experiments consists
        in an outer tube of 4/6 mm inner/outer diameter and an inner cylinder of 3 mm diameter. Such
        small dimensions reduce the solutal Grasshof number and give the opportunity to test two
        (reduced) natural convection configurations.
       Develop a Single Pan Thermal Analysis (SPTA) experiment for the characterization of the
        solidification of Cu-Sn alloys [11]. These measurements confirmed that the most reliable phase
        diagram is that recently assessed by Liu et al. [12]
       Analyze and optimize by numerical modelling the thermal conditions for the cartridge and
        furnace that should allow to solidify Cu-Sn specimens under microgravity conditions in the
        International Space Station.




                                  Figure 1:      Phase diagram of the Cu-Sn system.

During the present project, we could fully exploit the tools developed during the previous one, except
the microgravity experiments due to the problems of the space shuttle, and observe very interesting
phenomena in the Cu-Sn system under earth-gravity conditions. The results have been or will be
published in international journals (see publication list) and constitute the body of F. Kohler’s thesis
(http://library.epfl.ch/theses/?nr=4037, 2008). They are summarized below.


Bridgman solidification of Cu-Sn alloys

We report here only the main results obtained at low speed for alloys near the peritectic composition
(Cp = 21.5wt%). More details can be found in [13,14]. Figure 2(left) shows a longitudinal section of a
cylindrical Cu-20wt% specimen solidified at 0.5 m/s in a thermal gradient of 20 K/mm. The primary


[11] F. Kohler, T. Campanella, S. Nakanishi and M. Rappaz, Acta Mater. 56 (2008) 1519–1528. Corrigendum Acta Mater.
56 (2008) 3708 – 3709
[12] X.J. Liu, C.P. Wang, I. Ohnuma, R. Kainuma, K. Ishida, Metall. Mater. Trans. 35A (2004) 1641
[13] F. Kohler , L. Germond , J. Wagnière and M. Rappaz, Acta Mater. 57 (2009) 56-68
[14] F. Kohler, Peritectic solidification of Cu-Sn alloys : microstructure competition at low speed, PhD thesis, EPFL,
Lausanne, Switzerland, 4037 (2008) http://library.epfl.ch/theses/?nr=4037

                                                          9
                                                                                              METCOMP MAP 114


phase  appears in light brown, whereas the phase appears in darker bluish-brown. While the scale
is given by the 3 mm-diameter of the cylinder, the vertical axis has been transformed into a
temperature scale thanks to the temperature measurements performed during Bridgman solidification
(thermal gradient is vertical upward). As can be seen, the structure is very complex: It evolves from a
succession of - and -bands near the bottom of the specimen to lamellae, back to bands, etc. As
solidification progresses, the amount of -phase increases near the top of the specimen due to solutal
macrosegregation. In the band region, a microprobe measurement of the composition was made along
the profile labelled P1. It is shown on the right of the figure for the -phase only, since the
decomposition of  into the eutectoid (+) makes the interpretation more difficult.





Figure 2:    Solidification sequence in the 3 mm diameter cylinder of a Cu-20wt%Sn alloy directionally solidified at
             V = 0.5 μm/s (Left, Gℓ = 20 K/mm, -phase in light brown, -phase in dark brown). P1 represents the
             position of the measured composition profile shown on the right.

In each of the -bands, the composition exhibits a U-shape profile. Close to the – interfaces, the
                 
composition C*; in the various -bands increases from 16.3 wt.% Sn to 16.7 wt.% Sn, then decreases
to about 16 wt.% Sn. Since these compositions were measured in the microstructure for positions
corresponding to 500ºC at the time of the quench, they correspond fairly well to the solvus of the -
phase in equilibrium with the tin-rich phase (or with the /-phase above 520ºC, see Figure 1).
Therefore, the bands which formed during solidification have evolved in order to maintain near-
equilibrium conditions during the subsequent peritectic transformation. During band formation, the
composition at the growing interface is expected to increase as the -phase grows whereas it must
decrease once the -phase has nucleated. This explains the right portion of the U-shape of the
composition measured in the -phase. The left portion is explained by the peritectic transformation.
Due to the slope of the – solvus (Figure 1), the -phase partially shrinks in order to provide the
solute increase in the -phase, i.e. the -bands were thicker just after solidification. The composition
of -bands is not uniform because the Fourier number associated with their thickness and the time
spent between solidification and 500ºC is still fairly small (about 0.4, i.e., incomplete solute diffusion).

A similar structure is shown in Figure 3A for a Cu-21wt%Sn specimen solidified at 0.58 m/s in the
same thermal gradient, but now for the 4/6 mm tubular specimen cut along a longitudinal section (the
two sides of the tube are shown). Bands and lamellae can be clearly distinguished, especially in the top
longitudinal section. The lamellae structure can sometimes extend over several mm. Again, the
fraction of -phase increases as solidification proceeds, and at the time of the quench, only a planar
                                                       10
                                                                                               METCOMP MAP 114


front f  was present (vertical line on the extreme right, just after the position z1). The composition
profiles were measured across the lamellae along the profiles labelled P1, P2 and P3. The profile (P1)
closest to the quenched solidification front is shown in B. At this temperature, the -phase did not
transform in the solid state, except for the small peaks which can be observed near the - interface
due to the presence of a very thin layer. Otherwise, the composition is fairly uniform in both the -
and -phases. If one plots the - and -compositions measured at positions P1, P2 and P3 at the
corresponding temperature, the points shown in Figure 4 are obtained (filled squares). These measured
points are superimposed with the phase diagram. As for the bands, they follow essentially the solvus of
 and , due to the peritectic transformation that occurred after solidification. They are slightly larger,
but considering the uncertainty of the phase diagram data, this is not surprising. The extrapolation of
these points to the peritectic temperature gives the open circles, which are fairly consistent with a
cooperative growth mode. Such lamellae microstructures provide an interestsing way of probing the
phase diagram solvus in a single experiment, if one assumes that the system has enough time (low
velocity) to be in local thermodynamic equilibrium.




Figure 3:    A) Lamellar structure observed during a DS experiment in a Cu-21wt%Sn sample (4/6 mm tube, V = 0.58
             μm/s, Gℓ = 20 K/mm). The positions of the composition profiles P1, P2 and P3 are also indicated ( -phase
             in light brown, -phase in dark brown). B) Measured composition profile P1. C) EBSD analyses in the
             inverse pole figure representation ([001] direction parallel to the z-direction).




                                                        11
                                                                                                   METCOMP MAP 114




Figure 4:    The compositions and temperatures measured over the lamellae structures of Fig. 3 for positions distant
             from the quenched interface (Fo >> 1) are drawn on the Cu-Sn phase diagram (dark squares). Thick black
             lines correspond to solvus lines defined by the ASM phase diagram (Fig. 1), while the filled circles indicate
             the equilibrium concentrations. Dashed lines correspond to approximated solvus lines, whereas the
             corresponding extrapolated equilibrium concentrations C and Cp are represented with open circles.

Unfortunately, the various specimens produced during this project could never exhibit the quenched
lamellae structure directly in contact with the liquid. Therefore, it is impossible to assess whether
cooperative growth occurs below or above the peritectic temperature. Interestingly, the lamellae
structure always ended with a “hammer-like” morphology in the tubular specimens, as shown in
Figure 5. Although not fully understood, this structure is most likely due to a phenomenon similar to
that giving rise to islands. The almost facetted appearance of the - interface is again due to a partial
peritectic transformation.




Figure 5:    Hammer-like morphology of -lamellae revealed at the end of a lamellar structure observed in the 4/6 mm
             tube of a Cu-21wt%Sn sample directionally solidified at V = 0.58 μm/s (Gℓ = 20 K/mm).

Returning to Figure 3, EBSD measurements shown in (C) clearly reveal that the -bands and lamellae
belong to two grains for the upper part of the specimen (yellow and blue). This shows that: i) It is not
necessary to a have re-nucleation of one phase over the other to explain the band structures; ii)
Lamellae can directly originate from bands and conversely lamellae can be re-stabilized into bands.
This mechanism, which was not considered explicitly in previous studies on peritectics, can be
explained by the same mechanism that prevails for the onset of eutectic growth. It is shown
schematically in Figure 6. Consider the growth of one phase, say . As it grows with a planar front
morphology, the interfacial composition in the liquid has to increase. At some point, the liquid is
undercooled with respect to the -liquidus (i.e., below Tp) and the -phase must either nucleate, if it
does not exist in the whole specimen, or grow laterally from another part of the specimen where it is
present. This last mechanism is like an overlay of one phase over the other. It the lateral growth of  is
stable, this will lead to bands in cross-sections if the overlay is complete or to islands if it incomplete.
As the lateral growth of  can be at a much higher speed than the pulling speed, it can also be unstable.
This leads to cells of  which leave some room for the -phase to grow in between, thus leading to the
beginning of the lamellae cooperative growth mode. X-ray tomography experiments, which are in
progress, should allow a reconstruction of the whole microstructure with a micron-size resolution and
a confirmation of this growth mechanism.

                                                          12
                                                                                                     METCOMP MAP 114




Figure 6:      Illustration of the -phase nucleating and growing at some point on an advancing -liquid interface, The
               lateral growth of the -phase can be unstable, thus leading to the formation of cells. The -phase continues
               to grow in the thermal gradient direction in between these -cells, and an alternate sequence of - and -
               phases is therefore initiated. Reprinted from [15].


Finally, the cross section shown in Figure 7 illustrates very well the complexity of the microstructure
in the cooperative mode of peritectic growth. The - and -phases can appear under various
morphologies: Lamellae of the two phases, isolated -fibers in an -matrix or -fibers in a  matrix.
These morphologies are influenced by the local composition, itself influenced by macrosegregation,
and by the convection current itself. It was shown by numerical simulation in [14], that the tubular
specimens exhibit complex helicoidal currents with upward/downward and tangential components of
the velocity field. In this respect, microgravity experiments will really be key to observe this
cooperative growth mode in the absence of (or very limited) convection. In Figure 7, it can also be
seen that a Kurdjumov-Sachs relation exists between the two phases, i.e., {111} // {110} and <110>
// <111>.




Figure 7:      Orientation relationship between the - and -phases in the lamellar structure (same relationship observed
               in the straight lamellae and in the labyrinth-like microstructure). A) Illustration of the position of the
               diffraction analyses over the cross section. B) Pole figures: for the fcc-phase, the pole figure of <110>
               directions with a relative {111} plane are represented. On the other hand, for the bcc-phase, the pole figure
               of <111> directions with a {110} plane are shown. As can be seen, the dashed squares emphasize the exact
               orientation relationship between the primary and peritectic phases in the lamellar structures. Indeed, this
               coherency relationship corresponds to the Kurdjumov-Sachs relationship.


Phase Field Modelling

In order to better understand the cooperative growth mode of Cu-Sn lamellae, multi-phase field
simulations were performed using the model and the software of Folch and Plapp (Ecole
Polytechnique, Paris [16]). The adaptation of this code to the Cu-Sn conditions as well as the details of


[15] J. Dantzig and M. Rappaz, Solidification (EPFL-Press, Lausanne, Switzerland, 2009, to appear)
[16] R. Folch and M. Plapp, Phys. Rev. E 2005;72: 1-27

                                                            13
                                                                                                     METCOMP MAP 114


the simulations can be found in [14,17]. The simulations were restricted to two dimensions and the
thermal field was given by Bridgman conditions: T(z,t) = Tp + Gℓ (z – Vt). Among the major
assumptions of the model, we can mention: i) Equal thermal conductivities of the solid and liquid
phases; ii) No diffusion in the solid (i.e., one-sided model) and constant diffusion coefficient in the
liquid; iii) Equal interfacial energies of the -liquid, -liquid and - interfaces (i.e., angles at the
triple point are equal to 120 deg.); iv) cigar-shape phase diagram (i.e,, the solidus of both phases  and
 is parallel to the corresponding liquidus, which is a fairly good approximation for the Cu-Sn system).
The simulations were made with a mesh that is refined near the interface and that follows its
movement. Two configurations were considered: 1) two half-lamellae of  and ; 2) 4-spacings, in
both cases with a no-flux condition at the lateral boundaries.

As the number of mesh points in the diffuse interface needs to be sufficiently fine to give accurate
results, the actual spacing of the - lamellae observed in the experiments (about 80 m) had to be
reduced. Decreasing this spacing by a factor 10, it is necessary to increase the velocity by a factor 100
according to Jackson-Hunt’s law (2V = cst). In order to be below the stability limit of the planar
fronts, the gradient was also increased by a factor of at least 100 or more since Gℓ/V must remain larger
than T0/Dℓ, where T0 is the solidification interval of the -liquid interface.

The result shown in Figure 8 corresponds to four lamellae growing in a temperature gradient
Gℓ,sim = 400 Gℓ (i.e., Gℓ,sim = 8106 K/m) at a velocity Vsim = 100 V (i.e., Vsim = 58 m/s). While the
horizontal axis corresponds to the dimension, the vertical one has been converted into a temperature
scale, with indication of the peritectic temperature Tp. The average spacing set in the simulation was
 m, but as can be seen, it fluctuates slightly between the four - lamellae. More important, one can
notice two important features:

        The -lamellae are globally convex and extends slightly above Tp in the centre. Furthermore,
         the extent of the -phase above Tp is definitely correlated with the spacing: The narrower the
         -lamellar, the lower its temperature T* at the centre.
        The -phase is convex near the triple junction (indicated with a dot) in order to satisfy the
         Young-Laplace condition with the prescribed angles of 120 deg., but it becomes concave near
         the center. It is located below Tp.

In order to analyze in more details this result, several simulations were performed on two half-lamellae
with various compositions and lamellar spacings. The results are shown in Figure 9 for two alloys (20
and 21 wt%Sn) and two spacings (5.33 and 16 m), with a different representation as compared to
Figure 8. In this case, the interfacial compositions in the liquid Cℓ* and local temperatures T* are
directly superimposed with the phase diagram near the peritectic liquid composition CL. This
representation has the great advantage to outline the curvature undercooling, i.e., the difference
between the corresponding liquidus lines and the actual T*-Cℓ* points. It corresponds to the light/dark
grey shaded areas for the / phases, respectively.




[17] F. Kohler, T. Jauzein, M. Plapp and M. Rappaz, Multi-phase field simulation of directional solidification at low speed
of hypoperitectic Cu-Sn alloys, Acta Mater, (2009, to be submitted)

                                                            14
                                                                                               METCOMP MAP 114




Figure 8:    Operating temperature T*(x) of the lamellar growth front, when 4 lamellae of each solid phase are
             modelled (Grey triangles: -liquid interface; dark grey squares: -liquid interface). The position of the
             triple junctions is also represented by black dots. C0 = 19.87wt%Sn, sim = 6 μm, Gℓ,sim/Vsim = 4Gℓ/V and
             Tp = 795.7°C.

These diagrams confirm that the front is not strictly isothermal, even though the temperature
differences are small, and that the -phase is slightly ahead of the primary -phase. The centre of 
extends above Tp, while that of  is always below Tp. This gives a “compromise” solution to the
dispute regarding the operating temperature of cooperative growth in peritectics, i.e., whether T* is
above or below Tp: We find that it is “across” the peritectic temperature. Finally, the spacing does not
appear to be uniquely defined and can adapt to the conditions. 




Figure 9:    Operating temperatures T*(x) of the modelled lamellar structures drawn on the cigar-shaped Cu-Sn phase
             diagram (Gℓ,sim/Vp,sim = 4Gℓ/V): A) Cu-20wt%Sn, sim = 16 μm; B) Cu-21wt%Sn, sim = 16 μm;
             C) Cu-20wt%Sn, sim = 5.33 μm; D) Cu-21wt%Sn, sim = 5.33 μm. The shaded areas represent the
             curvature contribution Tr to the total undercooling T.




                                                        15
                                                                                        METCOMP MAP 114


Conclusion

This project has shown that bands and lamellae can exist in large solidification interval alloys, such as
Cu-Sn. It has also demonstrated that bands can occur via an overlay, rather than nucleation-growth,
mechanism. During the lateral growth of one phase, destabilisation can lead to lamellae.
Unfortunately, lamellae in contact with the liquid could not be quenched. Lamellae of  terminated
with a "hammer-like" morphology. It is important to consider the peritectic transformation when
looking at the shape and amount of phases or when measuring the compositions. Finally, multi-phase
field simulations indicate that cooperative growth occurs with the peritectic phase slightly ahead of the
primary phase, and thus the front is not strictly isothermal. Its temperature profile crosses the peritectic
temperature, and thus is neither above or below Tp.

Of course, all the experimental observations are certainly very much influenced by solutal convection.
In this respect, it is essential to carry out experiments in microgravity and/or in smaller capillaries.
Furthermore, X-ray tomography experiments should allow to see the interconnections between the two
phases and to confirm the conclusions regarding the overlay growth mechanism conjectured from
EBSD measurements in the present investigation.



III.2 Investigations on transparent model systems of organic peritectics (MUL/Ludwig, Eck,
Mogeritsch)

The team at MUL is responsible for the work package WP 2 “organic peritectics”. This work package
consists of three main tasks:

      WP 2.1      1g- research on organic peritectics
      WP 2.2      Preparation of µg-experiments on organic peritectics
      WP 2.3      Accompanying of µg- experiments on organic peritectics

In order to fully complete the work of WP2, the ESA Map Project METCOMP has been co-financed
by the ASA (Austrian Space Agency) since 2006. So far investigations have been carried out to define
the boundary conditions and experimental parameters for the intended µg-experiments. Simultaneous
to the experimental work on organic peritectics, investigations has been started to support the design of
the Directional Solidification facility, DIRSOL, which is intended to be used on the ISS. In addition,
physical properties of the two organic compounds of interest, TRIS and NPG, had to be measured for
the NASA safety check. At presence the DIRSOL facility is in its design phase (Phase B) and first
bread board tests have been completed. The DIRSOL facility is expected to be launched end of 2009.
Since WP 2.3 is directly coupled with the DIRSOL equipment, its execution is shifted to the beginning
of the corresponding experiments on ISS. In the following, the work done for WP 2.1 and WP 2.2 is
briefly described.
Material description, handling, filling and thermal stability
The transparent model system for in-situ peritectic solidification observations consists of the organic
substances TRIS(hydroxymethyl)aminoethan (TRIS) and Neopentylglycol (NPG) which reveals a
phase diagram with a non-faceted/non-faceted peritectic reaction at high temperature (Figure 10).




                                                    16
                                                                                                                                 METCOMP MAP 114




                                                              hypo-peritectic                           hyper-peritectic
                                                                                            [L]
                                                440
                                                                        [L+Cl]

                                                                                                           [L+CF]
                                                                 [Cl]



                              Temperature [K]
                                                400                         [CF+Cl]
                                                       [Cl+O]
                                                                                                                  [CF]
                                                360
                                                                            [CF+O]

                                                320                                                            [M+CF]

                                                        [O]                           [M+O]                              [M]
                                                      0.0         0.2            0.4              0.6       0.8            1.0
                                                                                            x


  Figure 10:      Phase diagram of TRIS-NPG. Both low temperature phases [O] and [M] are facetted and not miscible
                  below 310 K. Both high temperature phases [CL] and [CF] grow with non-facetted solid/liquid
                  interface and form a peritectic plateau around x = 0.5. The light shadowed area marks a region where
                  the facetted TRIS [O] phase coexists with the NPG-rich [CF] phase.

Detailed information about the phase diagram and the physical and chemical properties of the two
components are given in the Final Report of METCOMP I, 2002, and the Mid Term Report of
METCOMP II, 2004. They were gathered from literature ([18],[19],[20],[21],[22],[23],[24],[25],[26],
[27],[2],[28], [29],[30],[31],[32]). Own investigations show decompositions of TRIS close to the
boiling point. This is described in details in the Final Report of METCOMP II, 2006. Therefore, the
hot zone temperature of the used Bridgman furnace has been reduced to a minimum in order to prevent
the decomposition of TRIS.

NPG as delivered has a purity of 99%. It is very hydroscopic and thus it may contain residual water.
This fact requires in any case a pretreatment for further use. Thus, additional purification of NPG is
done by long term drying. TRIS as delivered has a purity of 99.9+% and further purification is known
to be rather complicated. In order to avoid possible decomposition of TRIS during additional
purification treatment, TRIS is used as delivered.

Alloy preparation, sample filling and purification routines were carried out in dry argon atmosphere in
a glove box. NPG has a 10 times higher vapor pressure than TRIS and thus has a tendency to


[18] M. Barrio, D.O.Lopez, J.Ll.Tamarit, P.Negrier, Y .Haget, J. Solid State Chem. 124 (1996) pp. 29-38
[19] M. Barrio, J. Font, D.O. Lopez, J. Muntasell, J.Ll. Tamarit, P.Negrier, Y. Haget, J. Phys. Chem. Solids 55 (1994) pp.
1295-1302
[20] http://www.merck.de
[21] A. Zhang, H. Zou, M. Yang, Gaodeng Xuexiao Huaxuue Suebao 9 (1988)
[22]G. Barone, G. Della, D. Ferro, V. Piacente, J. Chem. Soc. Faraday Trans. 86 (1990) pp. 75-79
[23] I. Nitta, S. Seki, M. Momotani, Proc. Japan Acad., No. 9, 26 (1950) pp. 25-29
[24] E. Murrill, L. Brees, Thermochim. Acta 1 (1970) pp. 239-246
[25] http://webbook.nist.gov
[26] DECLIC, Technical Notes on the System Definition, CNES (1999)
[27] S. Dobler, Ph.D. Thesis, EPFL Lausanne (2001)
[28] W.J. Boettinger, U.R. Kattner, Metal. Mater. Trans., No. 6, 33A (2002), pp. 1779-1794
[29] http://www.ralicks.de/deutsch/dichtungstechnik_silikon_scrintec_901.htm
[30] http://www.weidner-laboreinrichtungen.de/Programm/programm.html
[31] Prof. Paul O´Leary, Institute for Automation, University of Leoben
[32] D.Eilerman, R.Rudman, J. Chem. Ph. 72 (1980) pp. 5656

                                                                                       17
                                                                                                                 METCOMP MAP 114


sublimate. This property constrains an alloy preparation in sealed containers (firstly ceramic pans, now
sealed glass container) which is necessary in order to avoid a change in concentration during melting
(Mid Term Report of METCOMP II, 2004, and Ref. [29],[30],[31]). The accuracy of weighed
concentrations had been checked by DSC measurements. Our DSC equipment reveals an error which
is in the same order of magnitude as the published phase diagram TRIS–NPG [33]. Finally, long term
DSC measurements of different alloys show a thermal instability of TRIS after a few hours (Figure
11). The stability increases for higher NPG content or decreasing temperatures. This investigation is
published in [34] and allows the definition of suitable conditions for stable solidification experiment at
the peritectic region.

                                                                 50 mol% TRIS
                                                                 NPG
                            heat flow [mW] arbitrary units




                                                                 TRIS

                                                                                            Heat flow at 473 K



                                                                                            Heat flow at 463 K




                                                                                            Heat flow at 453 K
                                                             0    60      120       180   240     300
                                                                            time [min]
  Figure 11:      DSC traces of a peritectic alloy (50 mol% NPG), pure NPG and pure TRIS at different temperatures.
                  The graph shows heat flow versus time after reaching the selected annealing temperature. In the upper
                  line the peak of TRIS and the alloy with 50 mol% NPG shows the beginning of thermal instability.

Originally, samples of glued parallel-epiped glass plates with an inner dimension of 40x20x1 mm were
used but the high vapor pressure of the molten alloy and the weakness of the adhesive at high
temperatures makes samples unsealed. Therefore, long thin square glass capillaries with 400x400 µm
solved the problem of unsealed samples but the observation of solidification morphology was
insufficient. Now, long thin rectangle glass tubes with 2000x100 µm were used for further in-situ
observations. A detailed description of the filling process is available in the Mid Term Report 2004 of
METCOMP II.

Bridgman furnace, solid/liquid interface stability and homogenization of the samples
For the optical investigations, a Zeiss microscope is used in combination with a horizontal Bridgman
furnace (4 PT-100 temperature resistances, a ½” digital camera connected to a controlling PC). The
temperature gradient in the Bridgman furnace was measured with micro-thermocouples for the glued
parallel-epiped glass plates. This was not possible for the thin rectangle glass tubes because the micro-
thermocouples are too thick. Instead the phase transition temperature of the pure organic compounds,
NPG and TRIS, were used to determine the temperature gradient and temperatures at distinct positions.
A high temperature gradient, G, is necessary to suppress the growth of dendrites and to reach the G/V-
region where coupled and oscillatory growth is expected to occur in accessible experimental time. Due
to the aforementioned thermal instability of TRIS, the possible temperature of the hot zone is limited.
In order to still maximize the temperature gradient, the distance between the hot and cold zone was
reduced from 10 mm to 4 mm.


[33] M. Barrio, D. O. Lopez, J. Ll. Tamarit, P. Negrier, Y. Haget, J. Mater. Chem. 5 (1995) 431-439
[34] J. Mogerisch, A. Ludwig, S. Eck, M Grasser, B. McKey, Scripta Met., submitted

                                                                                 18
                                                                                                   METCOMP MAP 114




As described in the Final Report 2006 of METCOMP II (Ref. [35],[36],[37],[6],[4]), a long time
solid/liquid interface stability investigation was performed. Rectangle capillary tube samples have
been kept immobile in an established temperature gradient. In the first investigations with the square
glass tube samples, dissolving of the plastic phases [CF] and [CI] was observed and for alloys with a
concentration less then 60 mol% NPG the liquid was in direct contact with the facetted phase [O] after
approximately 8 hours. With a new furnace configuration (smaller baffle, two temperature controller)
and a reduced temperature of the hot zone (TH = 453 K), the plastic phases of these alloys contacted
the facetted phase only after 18 hours. Apart from that, alloys with peritectic concentrations showed a
stable solid/liquid interface for more than 5 hours. Furthermore, samples have been used twice. The
results of these investigations showed that the plastic phases dissolved faster than before. As a
consequence, samples can not be used more than once.
Wetting angle
The precise interpretation of the experimental studies requires knowledge of the boundary conditions
for nucleation. This includes the estimation of the contact angles between the wall and the phases.
Both plastic phases, [Cl] and [CF], show a large wetting angle with the glass wall (Figure 12), namely
approximately 145° ±10° for [Cl] and 135° ±10° for [CF]. To estimate the wetting angle between both
plastic phases, the nucleation of phase [CF] in front of the solid/liquid interface [Cl] has been
investigated. The results did not give sufficient information therefore further investigations are
required. Preliminary observations imply a wetting angle below 90°. In general, the knowledge of the
wetting angels allows the prediction of the location of nucleation sites which might influence the
solidification structure.




                             [L]                                         [L]
                                100 µm




                            [Cl]                                         [CF]
                                      a)                                       b)
Figure 12:     Wetting angel between the glass wall and the plastic phases indicated by the red arrow. a) Wetting angle
               between [Cl] in contact with [L] and the glass wall. (b) Wetting angle between [C F] in contact with [L]
               and the glass wall.

Figure 13 shows the nucleation map as described by Trivedi [8]. For TRIS-NPG alloys, the nucleation
map allows nucleation in front of the primary phase or/and at the wall. Considering the fact that in the
experiments no nucleation is observed on the wall, nucleation of the secondary or peritectic phase [C l]
might most probably occur in front of the solid/liquid interface.




[35] DIRSOL meeting at GPS in Paris, 22.05.2005
[36] http://www.waleapparatus.com/details/
[37] W. Kurz. D.J. Fisher, Fundamentals of Solidification, Trans. Tech. Publ., Aedermannsdorf, Switzerland (1989)

                                                           19
                                                                                                                                                                               METCOMP MAP 114


                                      
                                            solid/liquid
                                            nucleation
    contact angle with the wall 
                                                                                         unstable
                                    

                                                                                                                       Figure 13: Possible nucleation for the secondary
                                                  wall -solid - liquid nucleation                                   phase depending on the contact angle with the wall and
                                                                                                                       the primary phase. The red square considers the
                                                                                                                       possible state of nucleation for the alloys of the organic
                                                                                                                       compounds TRIS and NPG. The red line indicates
                                                                                                                    other possible forms of nucleation since the wetting
                                                                                                                       angle between the plastic phases is still under
                                               unstable                                 wall /liquid                   investigation.
                                                                                        nucleation
                                                                                            
                                                                contact angle 



Diffusion coefficient, interface recoil, temperature gradient, zone melting, microstructure map and
convection
To estimate the diffusion coefficient DL over the entire phase diagram, different methods have been
used. At first, the relationship between the molecular weight and the diffusion constant, published by
Polson [38], was used which gives a diffusion coefficient of approximately DL = 710-10 m2/s. Besides,
the diffusion coefficient DL in the peritectic region was measured based on the interface recoil which
occurs after starting planar growth with low withdrawal speeds. This method is fast enough not to be
influenced by any alloy decomposition far ahead of the planar front. The concentration dependence of
DL can be approximated by the regression line shown in Figure 14a. In the region, where coupled or
oscillatory growth is expected (Figure 14b), the solid/liquid interface reveals a cellular and/or dendritic
morphology and indeed shows oscillatory behavior. Details on that will be given in the next section.
                                              -10                                                                                        10
                                     6.0x10                                                                                     2.0x10
                                                                                                                                                                               possible layered struktures
                                                                                                                                              effect of undercooling
                                                                                                                                         10
                                                                                                                                1.5x10
                                              -10
                                     4.0x10
                                                                                                                  G/V [Ks/m ]




                                                                                                                                              [Cl] planar front                [CF] planar front
                                                                                                                  2
                    DL [m /s]




                                                                                                                                         10
                                                                                                                                1.0x10
               2




                                              -10
                                     2.0x10
                                                                                          DL                                              9
                                                                                                                                5.0x10

                                                                                                                                                 [Cl] cells/dendrites               [CF] cells/dendrites
                                           0.0                                                                                      0.0
                                              0.0         0.2       0.4           0.6          0.8     1.0                            0.40              0.45            0.50         0.55          0.60
                                                                              x                                                                                           x


                                                                  (a)                                                                                             (b)
 Figure 14:                                         a) Derived diffusion coefficients DL plotted versus alloy concentration c. b) Expected microstructures
                                                    depending on the applied G/V and alloy concentration c. The region where coupled and oscillatory
                                                    growth is predicted is highlighted in grey. Red lines show the effect of undercooling on that region; red
                                                    dots indicate samples where an oscillating solid/liquid interface was found.

For an unmoved sample in a given temperature gradient, the interface between a plastic phase, either
[CI] or [CF], and the liquid is usually flat and perpendicular to the sample axis. However, from time to
time the planarity disappears and the moved solid/liquid interface gets slightly curved. The cause for


[38] A. Polson, J. Phy. Chemistry (1950) pp. 649-652

                                                                                                             20
                                                                                          METCOMP MAP 114


this observation is the presence of convection. Especially in the long thin rectangle glass tubes this
happens more frequently. Figure 15 shows a bended solid/liquid interface followed by a flat interface
between the optically inactive and active phase. Although, the convection in the samples might be
relatively weak, due to the long observation time its presents can be seen as one of several reasons for
none constant solidification conditions.


                                        [L]


       450 µm


                                      [CF]             Figure 15: Curved solid/liquid interface of a hypoperi-
                                                       tectic alloy. The implemented temperature gradient is
                                                       homogeneous shown by a straight growth of the facetted
                                                       phase [O]. For comparison, both interfaces are marked
                                                       with dashed lines.

                                      [O]




In-situ observation of oscillatory peritectic growth
For an alloy with near peritectic composition, namely co = 50 mol%, solidified with a pulling velocity
V = 0.9 µm/s and a temperature gradient G = 14.2 K/mm, an oscillatory growth was observed.
Figure 16 shows the oscillation of the interface/tip position and the corresponding interface/tip
temperature as function of time. The gradual decay of the interface/tip temperature might be caused by
(i) the occurrence of macrosegregation amplified by convection and/or (ii) the decomposition of the
organic alloys within the hot zone. A typical oscillation cycle reveals a period of T  2200 s which
corresponds to a cycle length of   2 mm. Please note the similarity to Figure 2 of Section III.1.
Obviously, oscillatory growth for near peritectic compositions is found for metal as well as for organic
peritectics. In quenched metallic samples the local concentrations can be measured, which can not be
done in organics. On the other hand, in organics the dynamic of morphology formation can directly be
observed, which is quite difficult to do for metallic samples. Therefore, a comparison of the
oscillations found in both classes of materials is of utmost scientific importance.

In the following, an explanation of the oscillations found for the organic near peritectic alloy is given
based on a series of similar observations and further obvious arguments. However, it has to be
mentioned that the two plastic phases which participate at a peritectic growth, namely [Cl] and [CF],
are optically undistinguishable. On the other hand, an interpretation of the oscillatory growth can be
given if it is assumed that both phases grow simultaneously (but not necessarily at the same
temperatures).




                                                   21
                                                                                                                        METCOMP MAP 114




                                       30.412                                                   414


                                       30.271                                                   412




                                                                                                      temperature [K]
                       position [mm]
                                        30.13                                                   410


                                        29.99                                                   408


                                       29.849                                                   406


                                       29.708                                                   404
                                                0   3600   7200   10800      14400   18000   21600
                                                                  time [t]

Figure 16:   Oscillation of the interface/tip position and the corresponding interface/tip temperature as function of time
             for an alloy with near peritectic composition.

In Figure 17-20, the evolution of growth morphology for the cycle from t = 7.050 s after start pulling
to t = 9.360 s is shown. The cycle starts shortly before the minimum at t = 7.350 s occurs. The
morphology evolution at the minimum is shown in Figure 17 in more details. In Figure 17a two
cellular arrays can be seen, one slightly ahead of the other. In corresponding magnified pictures (not
shown) it became clear that the two arrays grow at different depth in the sample. In order to understand
the subsequent morphological transitions, let’s assume that the leading array consists of - and the
secondary of -cells. Figure 17a-d show that the advance of the -cells gradually disappears until both
cellular arrays grow with the same interface temperature. However, this “side-by side” growth is
obviously inherently unstable and leads to a sudden advance of smaller cells of both phases. The
reason for this instability is not yet understood.




                                                                  22
                                                                                                         METCOMP MAP 114




               300µm

                (a)           (b)                (c)                (d)              (e)                  (f)
Figure 17:   Growth of two cellular arrays close to the minimum interface/tip temperature. The corresponding tip
             temperatures were estimated to be T* = 408.6 K, 408.0 K, 407.3 K, 406.9 K, 406.9 K, and 407.3 K. Thus,
             the front instability goes along with a minimum in interface/tip temperature. The pictures are taken t =
             7.050 s after starting pulling with a time increment of Δt = 90 s.
                                                                                             Picture Nr.: 236, 239, 242, 245, 248, 251


Careful observation of morphological details makes it obvious that now the small cells of both phases
grow in close cooperation with each other until the -cells once more advances against the -cells (see
Figure 18).




                   (a)              (b)            (c)               (d)              (e)                   (f)

             300µm

                   (a)              (b)            (c)               (d)              (e)                   (f)
Figure 18:   Two examples which show that after new finer cells of both phases occur, the fine -cells are leading
             against the fine -cells. In pictures (f) and (f’)  is again the leading phase. The pictures are taken at t =
             7.380 s after starting pulling with a time increment of Δt = 60 s.
                                                                                           Picture Nr.: 247, 249, 251, 253, 255, 257


The corresponding -cells further develop into dendrites, as can be seen in Figure 19a. These isolated
dendrites grow clearly ahead of the cellular front which might now consist mostly of -cells. One
might expect that the -dendrites now gradually try to reach steady-state.




                                                           23
                                                                                                       METCOMP MAP 114




               300µm

                (a)            (b)              (c)               (d)              (e)                 (f)
Figure 19:   The leading -dendrites as well as the following -cells advance gradually until they reach the end of the
             solute boundary layer which has been formed previously. Feeling the initial melt concentration further
             acceleration of the tips occur until the tip temperatures are close to the related steady-state tip
             temperature for the fast growth. Then the tip velocity adapts to the pulling speed and tip thickening occurs.
             The pictures are taken t = 7.890 s after starting pulling with a time increment of Δt = 30 s.
                                                                                         Picture Nr.: 264, 265, 266, 267, 268, 269


However, what happens is something else. Suddenly all growing objects, -dendrites and -cells, shot
forwards, as can be seen in Figure 19b-d. This can be understood by considering the fact that a flat
growth front which consists of deep cells (as Figure 17c) does have a large solute layer similar to a
planar front. With V = 0.9 µm/s and D = 310-10 m2/s the boundary layer can be calculated to be
 =330 µm in thickness. After the sudden occurrence of the fine cellular growth, these cells have to
grow right into the solute boundary layer. Due to the fact that the tip radius of the fine cells is much
smaller than that of the former larger cells, the corresponding growth diffusion field is now more or
less localized around the tips. Therefore, the growth velocity can increase for both, the leading -
cells/dendrites as well as the following -cells. This leads to a gradually forward motion of the front
until the front reaches the end of the boundary layer. Here, the corresponding tips feel the initial melt
composition and can therefore further accelerate until they reach a temperature close to the related
steady-state tip temperature for the fast growth. This can obviously been seen in Figure 19d. However,
now the tips have to adapt to the pulling speed and therefore the tip radii gets larger and larger (Figure
19d-f). Note that the - and -tips do thicken at slightly different temperatures. The reason for that is
that due to the higher liquidus temperature of , compared to , and due to an approximately similar
tip undercooling, the tip temperature of  is higher than that of . It is important to notice that as soon
as the tips adapt to the pulling velocity, the maximum of the interface/tip temperature cycle is reached.
With the following tip thickening (see Figure 19 e-f) a change from isolated-tip-growth to a diffusion-
field-overlapping-growth occurs. In consequence the tip temperatures decrease again.




                                                          24
                                                                                                     METCOMP MAP 114




                       300µm

                        (a)             (b)             (c)                (d)              (e)
Figure 20:   After reducing the growth velocity in order to adapt the pulling speed, tip thickening results in a transition
             from dendrite-like into cell-like growth. With that the interface/tip temperature decreases. Note that the -
             cells/dendrites keep being ahead of the subsequently -cells. The Pictures are taken t = 8.040 s after
             starting pulling with a time increment of Δt = 330 s.
                                                                                            Picture Nr.: 269, 280, 291, 302, 313


The difference in tip temperatures between the -cells and -cells remains, while the interface/tip
temperature further drops. Figure 20 shows the evolution of the interface morphology from the
maximum of the interface/tip temperature to its minimum. Note that Figure 20e, taken at the end of the
corresponding cycle, is similar to Figure 17c, which represent the beginning of the cycle.

At presence, further studies with the near peritectic alloy c0 = 50 mol% are ongoing. For a pulling
velocity of V = 0.064 µm/s and a temperature gradient G = 12 K/mm evidence for a coupled peritectic
growth has been found. Figure 21 shows the corresponding solidification front, where an alternative
growth of - and -lamellar occur.




                                                              Figure 21: First evidence of a coupled peritectic growth in
                                                              the transparent peritectic system NPG-TRIS. An alternate
                                                              appearance of - and - lamella can be seen.


             100µm

                     Picture Nr.: 433



Note that detailed studies on oscillatory and coupled peritectic growth in near peritectic organic alloys
have been the major objective of WP 2. Now the experimental conditions for these studies are set.
However, two major difficulties are still unsolved:

                                                           25
                                                                                                    METCOMP MAP 114


            Possible decomposition in the TRIS-NPG system due to too high temperatures in the hot
             zone;
            Buoyancy-driven convection in the liquid part of the sample, which may lead to macroscopic
             concentration inhomogeneities.

Investigations for DIRSOL
For Phase A of the DIRSOL development the physical and chemical properties of TRIS, NPG and
their alloys have been defined. The toxic behavior has been checked by NASA and NASA attested a
low toxic level for the two organic substances and their alloys. Since the designer company requires
significant compatibility data for the construction of the DIRSOL facility, a first material compatibility
test was performed in December 2007 and a second material compatibility test in March 2008. The
results of these tests did not give any serious conflict between the substances and the material planed
for the apparatus. Since the filling apparatus of the flight sample is not directly comparable to the
experiments performed at the MUL further investigations were necessary. The sample differs in
geometry and in addition a vertical filling instead of a horizontal one is applied at DIRSOL. Therefore,
possible decomposition as a consequence of sample preparation was investigated. At high temperature
(above 400 K) an alloy of TRIS and NPG appears as plastic phase. At room temperature the alloy
decomposes to the pure substances TRIS and NPG which appear now facetted. At 310 K and under
polarized light it is possible to distinguish between NPG and TRIS because NPG is in the plastic phase
whereas TRIS is still in the facetted one. Figure 22 shows homogenously spread fine radial laminar
growth of TRIS and NPG. Based on this observation, it could be shown that significant decomposition
does not take place as an affect of the filling procedure. Another safety aspect requires the estimation
of the vapor pressures of the organic substances and there alloys. This has been investigated in summer
2008. The investigations showed that the vapor pressure of interest is below the security value of the
preferred glass cartridge.


              500 µm




                                   a)                                              b)
Figure 22:      Micrographs of an alloy with a hypoperitectic concentration (c 0 = 61 mol%) at 310 K. The facetted phase
                [M] of NPG changes to the plastic phase [CF] while TRIS is still in the facetted phase [O]. a) NPG and
                TRIS without polarization filter. Cracks occur as black lines while NPG and TRIS occur colored. b) With
                polarization filter, the plastic phase [CF] of NPG is black while the facetted phase [O] of TRIS is still
                colored.

As a conclusion the used organic material fulfills the safety requirements. So far, it seems that the
design of the DIRSOL apparatus offers the necessary system requirements for the µg-experiments of
this WP.

MUL participated on following DIRSOL meetings held during the duration of this report:

          09.05.06                 Progress meeting                   at ESTEC, NL

                                                            26
                                                                                      METCOMP MAP 114


       14.06.06              Progress meeting               web-conference
       27.06.06              Progress meeting               web-conference
       18.12.06              Progress meeting               at Campus Boucicaut, Paris, F
       07.12.07              Progress meeting               at Verheart Space, Kruibeke, NL
       18.03.08              Midterm meeting                at Verheart Space, Kruibeke, NL
       18.07.08              Progress meeting               telephone conference
       16.10.08              Progress meeting               at Verheart Space, Kruibeke, NL

Conclusion
The present report summarizes the experimental work on the peritectic TRIS-NPG alloy system. The
in-situ observations have been performed with a Bridgman furnace at a predefined G/V-ratio.
Furthermore, for the development of the DIRSOL apparatus, material properties had to be measured
and checked. Based on the results the following statements can be given:

      Under specific conditions, the plastic phase  gradually disappears until the facetted phase [O]
       gets into contact with the liquid phase. This indicates that the plastic phase of TRIS is thermal
       instable. However, the necessary observation time for a couple and/or oscillatory growth lies
       within the stable time range. After full dissolution, the plastic phase could be re-established by
       cooling the whole sample and re-starting the experiment. The results of these investigations
       have shown that the plastic phase now dissolved faster than before. As a consequence, samples
       can not be used more than once.
      The effective diffusion coefficient DL was estimated for the entire phase diagram by various
       experimental methods. The measured diffusion coefficient is in the order of DL = 310-10 m²/s.
      The contact angle between the plastic phases and the glass wall could be determined. The
       wetting angle between the phases could not be clear defined. In combination with the results of
       the solidification experiments it seems that nucleation of the peritectic phase happens only at
       the solid/liquid interface.
      Alloys with near peritectic concentration reveal oscillatory growth behavior. The mechanism
       which leads to these oscillations is a side-by-side growth of arrays of - and -cells. Both
       arrays of cells try to form a more or less planar growth front, until a sudden instability leads to
       the formation of fine cells of both types, which are now able to grow faster. The fine cells
       propagate until they reach the end of the solute boundary layer which was piled-up during
       “planar” cellular growth. After having grown beyond the boundary layer, the tips can even
       accelerate until they reach a temperature close to the related steady-state tip temperature for the
       fast growth. Afterwards, the tip thickens again and thus the interface/tip temperature starts to
       drop. Close to a “planar” growth of the two cellular arrays the oscillation starts again. The
       described mechanism for an oscillatory growth had never been observed before.
      For even lower growth velocities, a coupled growth of - and -lamellar is found. At presence,
       the dynamic of this peritectic growth mode is studied in details.
      The DIRSOL apparatus for the investigation under µg is now at Phase B. The organic
       compounds NPG and TRIS passed the necessary security checks up to now. The design of the
       DIRSOL apparatus offers the necessary system requirements for the µg-experiments of this
       WP. A first possible launch to the ISS is predicted by the end of 2009.




                                                   27
                                                                                                       METCOMP MAP 114


III.3    Phase field modelling of solidification of peritectics with particles (RISSPO/Pusztai)

(1) We have developed a phase field model of nucleation for competing crystalline phases of fcc
structure relying on a Ginzburg-Landau expanded free energy [39,40]. We have demonstrated its
applicability for the Ag-Cu eutectic system, including crystal nucleation in the metastable liquid
miscibility gap [39,40]. The respective Euler-Lagrange equations have been solved by a shooting
method outside the liquid miscibility gap, and by a relaxation method inside. Outside of the liquid
immiscibility region either an Ag rich or a Cu rich nucleus forms, depending on the composition and
temperature (Figure 23) [39]. A fairly complex behaviour has been found in the liquid immiscibility
region: At a fixed temperature below the critical point, six different types of nuclei may form
(Figure 24) [40]: two liquid-liquid nuclei; two solid-liquid nuclei; and two types of composite nuclei,
in which the crystalline core has a liquid “skirt”, whose composition falls in between the compositions
of the solid and the initial liquid phases. The phase field technique developed here is expected to be
applicable to other binary alloys, including peritectic and monotectic systems.




Figure 23: Radial phase field (light lines) and composition (heavy lines) profiles for the three types of solutions existing on
the left of the critical composition at (a) T=650 K, (b) T=750 K, and (c) T=800 K.




Figure 24: Phase selection in the Ag–Cu system, according to the minimum of the nucleation barrier at T=850 K (upmost
curve), 750 K (central curve), and 650 K (bottom curve). Note the complex behaviour below the critical point. (N Ag and NCu
stand for normal solutions that are rich in the component denoted by the subscript, while NeCu denotes the normal solution
forming on the Cu rich branch of the coexistence line.)

(2) We have developed phase field methods to incorporate walls characterised by arbitrary contact
angles into single component and binary systems. This work has been done in cooperation with J. A.
Warren, NIST, USA [41,42,43]. We have proposed three different approaches. If the normal


[39] G. I. Tóth, L. Gránásy: Phase field theory of interfaces and crystal nucleation in a eutectic system of fcc structure: I.
       Transitions in the one-phase liquid region. J. Chem. Phys., 127, 074709 (2007)
[40] G. I. Tóth, L. Gránásy: Phase field theory of interfaces and crystal nucleation in a eutectic system of fcc structure: II.
       Nucleation in the metastable liquid immiscibility region. J. Chem. Phys., 127, 074710 (2007)
[41] L. Gránásy, T. Pusztai, J.A. Warren, and D. Saylor: Phase field theory of heterogeneous nucleation. Phys. Rev. Lett.
98, 035703 (2007)
[42] T. Pusztai, G. Tegze, G.I. Tóth, L. Környei, G. Bansel, Z. Fan, and L. Gránásy: Phase-field approach to
polycrystalline solidification including heterogeneous and homogeneous nucleation. J Phys.: Condens. Matter 20, 404205
(2008)

                                                              28
                                                                                                      METCOMP MAP 114


component of the gradient of the structural order parameter is prescribed at the surface as a special
function of the order parameter (model A by RISSPO), an analogue of the classical spherical cap
model is realized. In contrast, if the value of the structural order parameter is fixed at the surface
(model B by NIST), a nonclassical behaviour can be observed, where the contact angle decreases from
its value taken at the melting point towards complete wetting at a critical undercooling, an analogue of
the surface spinodal of liquid-wall interfaces. A third possibility is to use a constant value for the
normal component of the gradient of the structural order parameter (model C by NIST & RISSPO).
To illustrate, how these three approaches work, we show the structure of heterogeneous nuclei forming
on surfaces with boundary conditions corresponding to model A, B and C under the same physical
conditions (Figure 25).




Figure 25: Radial cross sections of 3D heterogeneous nuclei forming on a foreign surface (the bottom axis). The same
contact angle of  = 60 has been realized by the respective parameters of the three different models. The rightmost panel
shows the ratio of the phase field prediction for the nucleation barrier height in model A (circles) normalized by the barrier
height for the homogeneous nucleus in the droplet model of the classical nucleation theory. For comparison, the 3D form
of catalytic potency factor f(ψ)=[23cos(ψ)+cos(ψ)3] from the spherical cap model is also shown (solid line). The
excellent agreement indicates that the classical contact angle concept is well represented by model A.

(3) Using the models described under (2), we have performed systematic studies aimed at exploring
the interaction between a growing dendrite and fixed particulate impurities in 3D. In our study, we
have varied so far the position, the size, the crystallographic orientation, and the surface properties
(contact angle, model A) of the spherical impurity particles. We have identified the micro-mechanism
of tip deflection.

(a) First, we investigated how the lateral displacement of foreign particles from the spine of the
dendrite tip influences the tip-particle interaction. In agreement with our earlier findings, the effect of
the collision with the foreign particle is maximal if the tip of the dendrite directly hits the particle; even
a small lateral displacement can hinder any effect other than simple engulfment (Figure 26).




Figure 26: Cross-sectional 2D images of the structural order parameter field (blue: solid, red: liquid, green: foreign
particle) of larger 3D simulations showing the interaction between a growing dendrite and foreign particles of the same
orientation as the dendrite. The displacement of the particles is 0, 6, 12, 18 and 24 pixels, respectively.




[43] J.A. Warren, T. Pusztai, L. Környei, L. Gránásy: Phase Field Approach to Heterogeneous Crystal Nucleation in
Alloys. Submitted to Phys. Rev. B

                                                             29
                                                                                                   METCOMP MAP 114


(b) Secondly, we have studied the effect of the size and surface properties of the foreign particles of
fixed orientation identical to that of the dendrite on the interaction between particles and the growing
dendrite.

We observed two types of behaviour: In the case of small particles and contact angles, the particle is
engulfed by the dendrite, while at large sizes and contact angles tip splitting occurs (Figure 27). Tip
splitting happens only above a critical size of the particle (comparable to the radius of the dendrite tip),
and the bigger the particle, the lower contact angle is required for tip splitting. Increasing the contact
angle beyond 150 degrees does not significantly influence the results. It can also be seen, that wetting
particles accelerate the temporary growth velocity of the dendrite, while non-wetting particles reduce
it.




Figure 27: 2D images of sections of the cross-sectional structural order parameter field (blue: solid, red: liquid, green:
foreign particle) of 3D simulations showing the interaction between the growing dendrites and foreign particles (of fixed
orientation). The columns correspond to particles of radii 2, 4, 6, 9, 12 and 15 pixels, respectively, while the rows
correspond to contact angles of 30, 60, 90, 120, 150 and 180 degrees. Snapshots displayed in the same row have been
taken at the same dimensionless time.
(c) Thirdly, we have studied the effect of the size and surface properties of the foreign particle on the
growing dendrite, when the particle is rotated by 30 degrees with respect to the dendrite (Figure 28).

                                                           30
                                                                                                      METCOMP MAP 114


The characteristic effect of the interaction in this case is tip deflection. At low contact angles and very
small particle sizes, tip deflection is only partial, i.e., the new growth direction of the tip does not fully
coincide with the orientation of the foreign particle; but except for these smallest sizes, the tip
deflection is full and pronounced. However, at larger contact angles corresponding to non-wetting
surfaces, the dendrite ignores the orientation of the impurity particle, and it continues to grow into its
original direction. The resulting growth forms are, therefore, similar to those for which the dendrite
and the foreign particles had the same crystallographic orientation (see Figure 27).




Figure 28: 2D images of the cross-sectional structural order parameter field of 3D simulations showing the interaction
between growing dendrites and foreign particles with 30 degrees misorientation. Colouring is according to the
misorientation compared with the original dendrite, the colour bar is the same as in Fig. 29. The columns correspond to
particles of radii 2, 4, 6, 9, 12 and 15 pixels, respectively, while the rows correspond to contact angles of 30, 60, 90, 120
and 150 degrees. Snapshots in the same row have been taken at the same time.

(d) Finally, we have identified the micro-mechanism of tip deflection in our simulations, a process that
is dominant in the case of wetting particles and is effectively suppressed in the case of non-wetting
particles:

When the growing dendrite tip just touches the impurity particle on its way, due to the favourable
wetting properties of the surface, the solidification front accelerates suddenly, leading to a thin solid
layer around the impurity particle (see Figure 29). Initially, this layer is only loosely connected to the
dendrite through a thin neck, but has a large common surface with the particle. Therefore, the
orientational ordering inside this layer follows the orientation of the foreign particle, and not the

                                                             31
                                                                                                      METCOMP MAP 114


orientation of the dendrite. Thus, the impurity particle ahead the dendrite tip surrounded by a thin
orientationally ordered solid layer can be considered as the growth centre of a new particle. The final
growth form looks very much as if formed by the impingement of separate particles (Figure 28).




Figure 29: Snapshots of a simulation showing the interaction between a growing dendrite (dark blue) and six impurity
particles of different sizes (green spheres semi-embedded by the dendrite tips in the snapshot on the left) and of 60 degrees
contact angle. Colouring shows the local crystallographic orientation.

(4) To investigate the apparent speedup of dendrite growth due to the interaction with a wetting
particle, we carried out systematic studies on how the foreign walls affect the dendrite growth speed
(Figure 30). We placed a solid seed at the corner formed by two planes with wetting properties set by
model A, and followed the time evolution of the dendrite tip position during growth. We found, in
accordance with Semoroz et al. [44] that the growth velocity of the dendrite is independent of the
contact angle in case of non-wetting (ψ>90°) particles, but increases with decreasing contact angle in
case of wetting particles (ψ<90°). We have also investigated the effect of local geometry on the growth
velocity. The change of growth velocity due to increased wetting is more pronounced if the dendrite
grows along a convex corner rather then on a flat surface.




Figure 30: Dendrite growth velocity predicted by phase field simulations in different geometries and at different contact
angles at the foreign surfaces. The left panel shows a snapshot of a simulation where the (invisible) foreign surfaces are
located at the back and bottom sides of the simulation box. The two types of tips labelled A and B correspond to growth
along a convex corner of 90 degrees and to growth on top of a flat surface, respectively. The right panel show the dendrite
tip speed as function of the contact angle, normalised to the value corresponding to 90°.

(5) We have, so far, treated the foreign surfaces as boundary conditions, when solving our phase field
equations (partial differential equations, PDEs). Physically, this corresponds to sharp interfaces
between the simulation volume and the foreign particle (wall) region. However, to model particle

[44] A. Semoroz, S. Henry, M. Rappaz, Metall. Mater. Trans. 31A (2000) 487

                                                             32
                                                                                       METCOMP MAP 114


pushing, this approach is generally inappropriate, since it is incompatible with a sub-pixel translation
of the foreign surfaces. To overcome this limitation, we have generalized our model A and model C
approaches to allow for diffuse walls, i.e., where the interface between the impurity particle and
simulation volume is not mathematically sharp. This has been achieved by reformulating our phase
field model so that the boundary conditions defining the foreign particles are converted to a volumetric
contribution via introducing a new phase field w(r) for the particles [43]

                                                        2              
                           F   dV   ) W  f (,c)  ( ) 2 1 W 
                                     Z(                               
                                                        2               

as opposed with the original form [3,4,5]

                                                             2      
                               F    dA Z()     
                                                   dV f (,c)  ( ) 2
                                                                 2
                                                                          ,
                                                                         

where  is the solid-liquid order parameter, Z() is the surface function of model A or model C that
determines the wetting properties of the surfaces, w(r) is the continuous “wall field” that is 1 inside
                     
the volume occupied by the impurity particles and 0 outside with a monotonic transition in a narrow
region of width  between, f (,c) and  2 2 ( ) 2 are the usual free energy density and square
gradient terms of phase field theory. It is easy to see, that in the   0 limit the generalized functional
falls back to the original functional, therefore the results obtained with the diffuse wall approach
expected to be in good agreement with the results of the original sharp wall approach.
                                
                                                          
(6) Next we have studied particle pushing by a planar solidification front in a single component
system. First, we determined the steady state configurations when the particle speed is the same as the
v advancing speed of the free solid-liquid interface. This configuration can be found as the time
independent solution of the phase field PDE expressed in a local coordinate system moving with the
same velocity v in the direction of solidification. We utilized the cylindrical symmetry of the problem,
reducing the dimensionality of the solution from 3 to 2.

Since in the time independent solutions no displacement of the impurity particle is required (the
particle moves together with the frame), both the sharp and the diffuse wall approaches can be used to
address this problem, making this scenario an ideal candidate for comparing the results obtained by the
two methods. As can be seen in Figure 31, the two solutions are almost identical outside the impurity
particle, i.e., in the physically relevant region.




                                                    33
                                                                                                      METCOMP MAP 114




Figure 31: Stationary configurations of particles pushed by a planar solidification front in Ni at an undercooling of 10K.
The panels show the half cross sections of the cylindrically symmetric 3D steady state solutions travelling with the velocity
corresponding to the velocity of the free solid-liquid interface. The transition from black to white corresponds to the
transition from the solid to liquid phase, while the contour lines corresponding to 0 = 0.1, 0.2, … , 0.9 are also plotted.
The left and right sides of the double panels show the solutions based on the sharp and diffuse wall approaches,
respectively. In the case of diffuse walls, the solution assigns  values to the interior of the particle, which, for better
visibility, is shown only by its contour lines and not by the respective colours. The diameter of the impurity particle is
40nm, while the size of the simulation box is 40  80 nm in all panels.

Next, we address the full dynamics of particle pushing. For this, we have to derive an equation of
motion for the particle. Since the spherical particle is assumed a fixed shape and size, it can be
described mathematically by a simple radial function as

                                                                                     
                                 w (x, y,z)  w  x  x0   y  y0   z  z0  .
                                                               2            2            2

                                                                                           

(In practice, a tanh profile of width  in the order of 1nm is used.) This way, all the dependence of the
free energy of the system on w can be expressed with the x0, y0 and z0 central coordinates, and the
                  
equations of motion for this coordinates is assumed to follow relaxational dynamics in the form of

                            F                
                                              
                                                         2                               x  x 
                                                                                                       
               v x  M w         M w  dV  f (,c)  ( ) 2   Z( )sign()
                                                                  w                            0
                                                                                                       ,
                            x 0              
                                              
                                                         2       
                                                                                       w    w
                                                                                               r   

where vx is the velocity of the impurity particle and Mw is the mobility of the particle that might be
related to viscosity, particle size, etc. Similar equations hold for the y and z components of the particle
  
velocity.

On the basis of the above equation of motion, we have carried out phase field simulations of particle
pushing using the same physical conditions as in the steady state case above. We have fixed Mw so that
velocity of the particle coincides with the velocity of the steady state solution for a particle with 90
degree contact angle at the solid-liquid interface. Then we varied the contact angle between 0 and 180
degrees (Figure 32). We found, that below a critical contact angle that falls between 30 and 60
degrees, the particle is engulfed, while above this critical contact angle the particle is pushed by the
advancing solidification front.


                                                             34
                                                                                                      METCOMP MAP 114




Figure 32: Particle pushing by a planar solidification front as predicted by the phase field theory. The left three panels
correspond to 30, 90 and 150 degrees contact angle and show the snapshots of the radial cross sections of the cylindrically
symmetric 3D solutions approximately 2108 s after the solidification front reached the particle. (The central snapshot
corresponds to the right side of the 90 degree panel in Figure 31.) The rightmost panel shows the velocity profiles
evaluated from the simulations. After a transient period, the particle is either engulfed by the moving solid-liquid interface
(at low contact angles) or pushed (at high contact angles). At the velocity minimum (t  1.3  108 s), the curves from
bottom to top correspond to contact angles  = 0, 30, 60, 90, 120, 150, and 180 degree, respectively.

The free energy functional specified in section (5) and the equation of motion proposed in section (6)
enable us to study not only single component systems, but binary alloys too. However, two regions
require special attention during the modelling of particle translation in a binary alloy: (a) the volume
that is freshly occupied and (b) the volume that is just left by the translating particle. We have worked
out a very simple approximate procedure to handle these regions. First, we calculate the error in the
integral of the composition field that we would make by simply shifting the particle with the required
amount. Then, we redistribute the excess solute/solvent just around the surface of the particle [in
practice, proportionally to the derivative of w(r)], which might be considered as a primitive model of
the local mass flow that restores local equilibrium in the neighbourhood of the particle.

We have performed illustrative simulations of particle pushing in a binary Cu-Ni alloy (Figure 33).
Here a solute-enriched liquid layer builds up between interface and particle. As a result, the
solid/liquid interface does not always touch the foreign particle (see e.g., the rightmost panel of
Figure 33), i.e., the wetting properties of the foreign particle are not expected to influence particle
pushing. In some specific cases, such as dendrite tip-particle interaction, however, the solid-liquid
interface may get close enough to the foreign particle so that, similarly to the single component case
(Figure 31), the wetting properties indeed matter. Indeed, under such circumstances, we have observed
both particle engulfment (at low contact angles) and particle pushing (at high contact angles).




                                                             35
                                                                                                         METCOMP MAP 114




Figure 33: Particle pushing in a binary alloy in 2D. The left two panels show the effect of the contact angle (45 and 135
degrees, respectively), while the right three panels illustrate the effect of increasing particle mobility (M w, 3Mw, 10Mw,
respectively). The foreign particle is denoted by a white circle, the trajectory of the particle by a black solid line, while the
rest of the simulation window is coloured according to the local chemical composition.


Invited talks

1.  L. Gránásy, T. Pusztai, G. Tegze, G. Bortel, J. A. Warren, J. F. Douglas:
    From needle crystals to spherulites: A phase field study. MCWASP XI., 28 May - 2 June, 2006, Opio, France
2. L. Gránásy, T. Pusztai, G. Tegze, G. Bortel, J. A. Warren, J. F. Douglas:
    Phase field modeling of polycrystalline patterns in two and three dimensions.ESF Research Conference on Solid/Fluid
    Interfaces, Complex Fluid Interfaces and Nanofluidics, 9-14 Sept., 2006, Obergurgl, Austria
3. L. Gránásy, T. Pusztai, G. Bortel, J. A. Warren, J. F. Douglas:
    Nucleation and polycrystalline freezing in two and three dimensions: A phase field study. 8th International Symposium
    on Crystallization in Glasses and Liquids, September 24 – 28, 2006, Jackson Hole, Wyoming, USA.
4. L. Gránásy, T. Pusztai, G. Tegze, G. Tóth, J. A. Warren, J. F. Douglas:
    Predicting polycrystalline patterns in 2D and 3D: A phase field approach. International Workshop on Polymorphism
    in Condensed Matter, Nov. 13-17, 2006, Dresden, Germany
5. T. Pusztai, G. Tegze, G. Tóth, L. Gránásy, J. A. Warren, J. F. Douglas:
    Phase field modeling of polycrystalline freezing in 2D and 3D: New developments. The Minerals, Metals & Materials
    Society Annual Meeting, February 25 – March 1, 2007, Orlando, Florida, USA
6. L. Gránásy:
    Phase field theory of homogeneous and heterogeneous nucleation. Workshop on Phase Field Models for the
    Evolution of Complex Structures, Institut Henri Poincaré, June 4-6, 2007, Paris, France
7. T. Pusztai:
    The orientation field in 3D: Quaternion magic. Workshop on Phase Field Models for the Evolution of Complex
    Structures, Institut Henri Poincaré, June 4-6, 2007, Paris, France
8. L. Gránásy, T. Pusztai, G. I. Tóth, G. Tegze: Phase field approach to polycrystalline solidification including
    heterogeneous and homogeneous nucleation. CODEF-II workshop, 30 March – 2 April, 2008, Bonn-Bad Godesberg,
    Germany
9. L. Gránásy, T. Pusztai, G. I. Tóth, G. Tegze, L. Környei: Phase field approach to homogeneous and heterogeneous
    crystal nucleation in alloys. SIAM Conf. on Mathematical Aspects of Materials Science, 11-14 May, 2008,
    Philadelphia, Pennsylvania, USA
10. L. Gránásy, G. Tegze, T. Pusztai, G. I. Tóth, L. Környei: Phase field modeling of self-organized polycrystalline
    structures: dendrites, spherulites, eutectics. IUCr2008, XXI Congress of the International Union of Crystallography,
    23-31 August, 2008, Osaka, Japan
11. L. Gránásy, G. Tegze, T. Pusztai, L. Környei, G. I. Tóth: Phase field modeling of complex polycrystalline
    solidification morphologies. Materials Science and Technology Conference and Exhibition, MS&T08, 5-9 October,
    2008, Pittsburgh, Pennsylvania, USA
12. L. Gránásy, G, Tegze, L. Környei, T. Pusztai: Phase field modeling of complex polycrystalline morphologies in three
    dimensions. Workshop on Phase-field simulations: Materials Science Meets Biology and Medicine. 12-14 November,
    2008, MPIPKS, Dresden, Germany



                                                              36
                                                                                                   METCOMP MAP 114




III.4   Peritectic Cu-alloys and solid inclusions (DLR/Herlach, Kolbe)
During the continuation of the ESA MAP METCOMP the work concerning the interaction of solid
particles with a dendritic solidification front, has been divided into two work packages. "Peritectic Cu
alloys and solid inclusions" - work package WP 4 - is performed by DLR and "Peritectic Ni and Fe
alloys and solid inclusions" - work package WP 5 by Ruhr-University Bochum (RUB). The
collaboration is very close in order to maximize synergy effects. This is especially true for preparation,
conduction and evaluation of parabolic flight experiments. In this period, parabolic flights have been
used for experiments in May 2006 (A), November 2006 (B), September 2007 and April 2008. The next
campaign is due in September 2009. Nevertheless separated reports are given, the line of separation
mainly defined by the different materials.

Theoretical and technological background
New developments in processing have shown the importance of the complex interaction of a
solid/liquid interface with foreign particles. Recently, semi-solid casting with ceramic particles has
been applied to produce material with high wear resistance for brake drums [45]. This means metallic
dendrites and ceramic particles moving in a mixture of metallic melt. A bad result would be if the
solidified material had all the particles located at the grain boundaries leading to low fracture
toughness of the material. The distribution of the inclusions in the as-solidified alloy is determined by
the interaction of the inclusions with the advancing solidification front:
(i) The particles are pushed forward by the solidification front and segregate with the last solidifying
liquid (particle pushing).
(ii) In the case of a dendritic solidification front, the particles may be pushed and trapped in the
interdendritic regions (particle trapping).
(iii) The solidification front engulfs the particles and they are embedded randomly into the metal
(particle engulfing).
The modes (i) and (iii) are derived from the study of the planar solidification front with particles. An
overview of experimental and theoretical work is given in [46, 47]. It has been shown that inclusions
of a given kind are engulfed, if the velocity of the moving solid/liquid interface exceeds a critical
velocity vC. Otherwise the inclusions are pushed ahead by the front [48]. Also it was shown that the
critical velocity decreases with increasing particle size. The critical velocity vC for particle engulfment
depends upon the physical properties of the melt-inclusion system, the morphology of the
solidification front and upon the geometrical properties of the inclusion (size, shape, surface
roughness).
Materials' processing relevant to application-oriented operations generally involves alloys rather than
pure metals. Solidification of alloys is accompanied with a concentration gradient in front of the
solid/liquid interface in the liquid, which can destabilise a planar interface leading to cellular or
dendritic solidification. In alloys the particles act as barriers for the solutal diffusion field ahead of the
solidification front, which can induce morphological instabilities of the interface [49]. Therefore,
destabilisation of a planar interface is inherent to the solidification of a melt containing a dispersed
phase. Besides the influence of melt convection on the incorporation of particles mentioned above



[45] G. Withers, Advanced Materials & Processes 163 (2005) 45
[46] R. Asthana and S.N. Tewari, J. Mat. Sci. 28 (1993) 5414
[47] R. Asthana and S.N. Tewari, Proc. of Adv. Mater. 3 (1993) 163
[48] D.M. Stefanescu, D. Shangguan and P. von den Brinken, Mat. Sci. Forum 77 (1991) 25
[49] J.A. Sekhar, R. Trivedi, S.A. Han: in Solidification of Metal Matrix Composites, P.K. Rohatgi (ed.) TMS (1990) 21,
      J.A. Sekhar, R. Trivedi, Mater. Sci. Eng. A147 (1990) 9

                                                           37
                                                                                                         METCOMP MAP 114


pushing of particles into the interdendritic space is favoured, if a component of the relative velocity
between particle and dendrite is tangential to the interface.
The main parameters for the pushing/embedding process are the interfacial energies between inclusion,
liquid matrix and solidification front, and the ratio of thermal conductivities of liquid (kl) and inclusion
(kp) at the given temperature [46]. Kinetic models consider the effect of the velocity of solidification
by describing the interaction of front and particle with a balance of forces acting on the particle in the
vicinity of the front. The outcome is an expression for the critical velocity of the solidification front
and a given particle radius R from which up to larger velocities the particles are engulfed. It was
shown that the critical velocity for particle engulfment is increased by higher convection levels [50].
Near the solid/liquid interface forces arise from the fluid velocity gradient because of different flow
rates on the two sides of the particle [51]. A logic extension is to consider curvature effects of the
interface. Pötschke and Rogge [52] derived an expression for the critical velocity vC, which takes into
account alloying elements and the (attractive) van der Waals force:

                                                                         1 / 2
                1.3   R  k p                                    
                               2
                                                        kp                                 C0 ml 
           vC        16    
                     a0  k l
                                                  15 
                                                              2
                                                                    
                                                                                  and   
                                                                                             k 0 GlDl
                                                                                                              *
                                                            
                                                         kl
                                                                     

The influence of alloying is incorporated in χ, pure metals are described by χ=0. The parameter
σ=σps-σpl-σsl consists of the interface tensions between particle/solid, particle/liquid and solid/liquid.
C0, ml and k0 (composition, liquidus slope and partition coefficient) are found in the phase diagram, Dl
is the diffusion coefficient, η the viscosity. Gl is the temperature gradient in the liquid ahead of the
solidification front.
Modified expressions for the critical velocity were proposed by Stefanescu, Dhindaw, Kacar and
Moitra (SDKM) [53]

                    a0      k 
          vC               2  p 
                            
                 6(n  1)R    kl 
                                                                and n = 4                                    **
                                   



and by Shangguan, Ahuja and Stefanescu (SAS) [54]:


                    a 0  k l  n  1 
                                             n
         vC                     
                3(n  1)R  k p  n 
                                                             and n = 2                                     ***

The three models are applied to the results obtained with the Ni-Ta + Ta2O5 systems and discussed in
greater detail in section III.5.

Solidification experiments with Cu-based systems in TEMPUS
In a study of the interaction of an advancing dendritic solid/liquid interface with ceramic particles one
has to overcome the experimental problem of identifying a suitable metal-matrix + ceramic particles
system. A suitable system exhibits undercoolability and chemical and physical stability: (i) the metal


[50] S. Sen, B.K. Dhindaw, D.M. Stefanescu, A. Catalina and P.A. Curreri, J. Crystal Growth 173 (1997) 574-584
[51] Q.Han, J.D. Hunt, ISIJ International 35 (1995) 693
[52] Pötschke, J., Rogge, V., J. Cryst. Growth 94 (1989) 726
[53] D.M. Stefanescu, B.K. Dhindaw, A.S. Kacar and A. Moitra, Metall. Trans. 19A (1988) 2847
[54] D. Shangguan, S. Ahuja and D.M. Stefanescu, Metall. Trans. 23A (1992) 669

                                                                     38
                                                                                                            METCOMP MAP 114


melt together with the particles can be undercooled, (ii) the particles do not react with the liquid melt
and (iii) the particles stay in the melt during processing. These criteria are fulfilled, e.g., in the Ni98Ta2
+ Ta2O5 system and, in addition, the microstructure can be observed easily by SEM due to segregation
of elemental Tantalum around the Nickel-rich dendrites.
           Cu + metallic particles                            Cu + ceramic particles


                                                   Wetting/alloying             "pure" systems

                   Cu + Nb                      Cu-(0.9-10)Ti + Al2O3            Cu + Al2O3
                                         diss
                                           .
                   Cu + W                       Cu-(0.24-10)Ti + Ta2O5           Cu + Ta2O5
                                         diss                                                    Figure 34:
                                          .
                                                                                  Cu + TiO       Overview on the investigated Cu
             Cu-(2at%)Ge + Nb                    Cu-(5at%)Ti + HfO2
                                                                                                 + particle systems. Cu-Nb,
parabolic flight 2005                            Cu-(1.1-2)Ti + Ti2O3            Cu + Ti2O3      Cu98Ge2 + Al2O3 and the Cu-Ni
                              diss
                                                                                                 alloys were processed in
parabolic flight 2006-A                         Cu-(2at%)Ge + Al2O3               Cu + TiC       parabolic flight. All "pure"
                                                                                                 systems showed bad wetting -
parabolic flight 2007                              Cu50Ni50 + Ta2O5              Cu + Nb2O5
                                                                                                 the particles left the melt. The
parabolic flight 2008                              Cu75Ni25 + Ta2O5                              indication "diss" means that the
                                                                                                 particles were dissolved during
                                                   Cu25Ni75 + Ta2O5                              processing.

                                     stable


An overview on the investigations with Cu-based systems is given in Figure 34. After finding stable
systems, the strategy was to transfer the successful approach in Ni-based alloys to Cu-based alloys in
four steps:
(i) The pure system Cu + Al2O3 particles. The basic problem is that the ceramic particles (e. g. Al2O3)
leave the bulk of the sample directly after melting and stay on the surface. This "squeezing-out" has
been analyzed and described in detail in a previous report [55]. It has been attributed to the bad
wettability of ceramics by Cu-melts.
(ii) Exchange of the ceramic particles by metallic Nb particles, which are wetted by Cu and stay inside
of the sample during parabolic flight [56]. Nb metal has a similar thermal conductivity λ (53.7 W/mK
at RT) and density ρ (8.55 g/cm3 at RT) as liquid Copper (λ=100 W/mK, ρ=8.09 g/cm3). Cu + Nb
particles is a model system for low interaction of the solidification front with the metallic particles.
(iii) Experiments with the system Cu98Ge2 + Al2O3 were conducted in order to test the influence of Ge
as a wetting agent. Ge is suitable for marking the interdendritic region in the solidified material. A
solidification experiment in TEMPUS at condition of low convection showed that the system is not
suitable for investigation: The particles were squeezed-out from the melt immediately after melting
[57].
(iv) Stable systems - under terrestrial processing conditions - are candidates for parabolic flights
(Figure 34): (a) Cu-Ni + Ta2O5, (b) Cu95Ti5 + HfO2 and (c) Cu with W particles. We investigated the
system Cu100-xNix + Ta2O5 particles (x = 25, 50, 75) in TEMPUS during parabolic flight campaigns
September 2007 and April 2008).
The idea behind this system is to combine the good wetting conditions of the Ni-based systems with
Copper. The phase diagram of Cu-Ni (Figure 35) is simple from the view point of thermodynamics, it
shows miscibility in the full range of concentration without any intermetallic phases. The solidus-


[55] Final Report ESA MAP METCOMP AO 98/99, ESA Contract No. 14243/00/NL/SH -Continuation- (2006)
[56] Mid Term Report ESA MAP METCOMP AO 98/99, ESA Contract No. 14243/00/NL/SH -Continuation- (2004)
[57] Mid Term Report ESA MAP METCOMP AO 98/99, ESA Contract No. 14243/00/NL/SH - 2nd Continuation- (2007)

                                                                         39
                                                                                                                 METCOMP MAP 114


liquidus interval of Cu-Ni is quite narrow leading to a low difference in concentration between
dendrite core and interdendritic region in the solidified material. Furthermore, Cu and Ni are neighbors
in the periodic system of elements with a similar density. This means that the contrast in the
backscattered electron images is low. Thus, microstructure analysis of the solidified sample with the
aim to determine the locations of the particles - engulfed in the dendrite or entrapped in the
interdendritic region - is more difficult than the respective analysis in the Ni-Ta + Ta2O5 systems and
requires an improved surface preparation. Remark: The hypothetic system with x = 100 is different to
the Ni-Ta + Ta2O5 systems, as it contains no elemental Tantalum.


      1600
                                                                              Figure 35: Phase diagram of Cu-Ni (after Massalski).
   T(°C)                                                                      The arrows show the compositions of the investigated
      1400                                                                    alloys. The equilibrium concentrations of dendrite core
                   liquid                                                     and interdendritic region are indicated for the alloy
      1200                                                                    Cu25Ni75.
                                      solid
                                                    Cu-Ni
      1000


       800
             0              25                50     75     at% Ni 100



All three Cu-Ni systems were suitable for processing in TEMPUS. All samples have been melted and
solidified during the 20 s of low gravity in parabolic flight. After melting clusters of the particles
appeared on the surface. The clusters stabilized the surface and suppressed oscillations of the sample
leading to a further reduction of convection in the melt. The achieved undercooling was always low
(< 5K).
Table 1: Results of solidification experiments in low gravity with Cu 100-xNix + Ta2O5 systems

           Alloy                 Parabolic flight         No. of samples           Total no. of particles      Engulfed fraction

       Cu75Ni25                       2008                      1                           50                        0

       Cu50Ni50                       2007                      1                           168                       0.10

       Cu25Ni75                       2008                      1                           575                       0.15



Post mortem microstructure analysis showed clusters of particles on the surface and in the bulk of the
material. In the cluster-free regions isolated submicron particles were found, with the tendency of a
larger density of particles in the Ni-rich systems. This fact reflects certainly the influence of wetting on
the stability of the systems: Ni-rich systems are more stable during processing. If we consider only the
isolated particles in the respective samples, we can distinguish between engulfed and entrapped
particles. The tendency in the system Cu100-xNix + Ta2O5 particles is that the fraction of engulfed
particles is larger with larger x, reflecting the improved wetting in the Ni-rich melts. The size
distributions of engulfed or entrapped, entrapped and engulfed particles are shown in Figure 36, the
results are summarized in Table 1.




                                                                         40
                                                                                                                                           METCOMP MAP 114




                          40                                                                                 120
                                                     engulfed or entrapped                                                              engulfed or entrapped
                          35
                                                                                                             100




                                                                                       Number of particles
    Number of particles




                          30                         entrapped                                                                          entrapped
                                                                                                             80
                          25                         engulfed                                                                           engulfed

                          20                                                                                 60

                          15
                                                                                                             40
                          10                               Cu50Ni50 +                                                                         Cu25Ni75 +
                                                           Ta2O5                                             20                               Ta2O5
                          5

                          0                                                                                   0
                               0,1              1                            10                                    0,1             1                            10

                                      Particle diameter (µm)                                                             Particle diameter (µm)



Figure 36: Distributions of entrapped or engulfed, entrapped and engulfed particles in Cu-Ni + Ta2O5 solidified in low
gravity. Undercooling temperature < 5K. The fraction of engulfed particles is 0.10 in Cu 50Ni50 and 0.15 in Cu25Ni75.
Scattering of the data is due to weak contrast in the backscattered electron images and the limited data base.


In section III.3 of this report Pusztai investigates particle pushing and engulfment in the system Cu-Ni
with particles, especially the influence of wetting on the interaction of solidification front and particle.
According to his results this influence may be neglected when a thick liquid layer of enriched solute is
established between particle and advancing dendrite. Under certain conditions of solidification this
layer is thin. Then he finds a transition from pushing to engulfment, which is due to the wetting
conditions (Figure 33). We have changed the wetting conditions in the system Cu-Ni + Ta2O5 particles
by replacing Ni with Cu and find a decrease of the engulfed particle fraction. Probably there exists a
correlation between these results.
We plan to repeat experiments with the Cu-Ni + Ta2O5 particles during the parabolic flight campaign
September 2009 in order to support these results and/or achieve undercooling of the melts. A variation
of melt undercooling could lead to a different liquid layer between advancing dendrite and the particles
and thus, change the influence of wetting.
In addition, it would be interesting to know the influence of elemental Tantalum on the engulfment in
the Cu-Ni systems. A link to the Ni-Ta systems would be established. Therefore we plan the
investigation of (Cu0.25Ni0.75)98Ta2 + Ta2O5 as a representative for a ternary system.


III.5                      Peritectic Ni- and Fe-alloys and solid inclusions (RUB/Eggeler, Lierfeld, Wu)
In this part we begin with a description of the state of the art of producing clean samples for the
solidification experiment with metallic matrix and ceramic particles. The procedure and the basic ideas
are given as example for Ni-Ta + Ta2O5. For all other systems the routine is similar, differences arise
mainly from different melting points of the components. Then the results from the solidification
experiments during parabolic flights are given, which were obtained by in situ high speed camera
observation and post mortem investigations.

Preparation of metal matrix + ceramic particle systems
The samples are prepared from pure Ni and Ta powders for alloying Ni98Ta2 samples with particles of
Ta2O5 in size ranging from 1 µm to 45 µm. Pure Ni is a suitable material for undercooling
experiments. Solidification of undercooled melts of pure Ni is by the growth of thermal dendrites,
which are hard to make visible in microstructure analysis of as-solidified samples. Therefore, Ta is
added to pure Ni for preparation of dilute Ni-Ta alloys. Solidification of alloys is by the growth of
                                                                                  41
                                                                                                 METCOMP MAP 114


chemical dendrites with segregation phenomena occurring which make the analysis of microstructure
by scanning electron microscopy much more easy. In such a way the position of the ceramic particles
with respect to the dendritic morphology can be determined. Ta2O5 is chosen as the addition of ceramic
particles to the metallic matrix material since it has a mass density of 8.20 g/cm3 being similar to the
mass density of pure liquid Ni of 7.93 g/cm3. Thus, sedimentation and buoyancy effects of Ta2O5
during processing of liquid samples under 1g conditions can be neglected. Ta2O5 shows a melting
temperature of 2163 K and is therefore in a stable solid state at the liquidus temperature 1717 K of
Ni98Ta2.
For sample preparation Ni and Ta powder in mass corresponding to the concentration of Ni 98Ta2 alloy
together with Ta2O5 powder is mixed under Argon atmosphere of purity of 99.998%. The powders are
milled in a ball-milling device in the enclosed high-purity atmosphere for a time of about 30 minutes.
After the milling process a homogeneous mixture of all components of the metal-ceramic composite
material is reached. Subsequently, the mixed powder is filled in a compression molding die and after
evacuation of the enclosed air the powder mix is compacted by a punch under high pressure of 500 to
600 MPa to a green body. The cylindrical or spherical material prepared in such a way is placed in a
high vacuum furnace and is sintered at temperatures between eutectic and solidus temperature of the
metallic alloy (around 1673 K) under high vacuum conditions for 4 – 8 hours. Optical and SEM
microscopy investigations show that the particles are homogeneously distributed in the as prepared
material. Subsequently spherical samples with a diameter of 6 – 8 mm are prepared by mechanical
working and surface cleaning and are finally placed in the Ultra-High-Vacuum recipient for the
levitation experiment. Figure 37 shows the procedure of sample preparation in a schematic way.




Figure 37: Procedure of sample preparation of metal-ceramic composite materials for levitation experiments. Powders of
Ni, Ta and Ta2O5 are milled, compacted at high pressure and sintered at temperature close to the solidus temperature of
Ni98Ta2 alloy. Spherical samples were produced from the as prepared material.

For undercooling experiments containerless processing by electromagnetic levitation is applied. This
technique is an effective tool for undercooling since heterogeneous nucleation on container walls is
completely avoided which otherwise act as highly efficient catalyst for crystallization limiting the
undercoolability of the melt [58]. During electromagnetic levitation of a metallic sample inside the
levitation coil eddy currents lead to heat up and eventually melt the sample. The liquid sample is
cooled and undercooled by gas cooling using He gas of high thermal conductivity streaming around
the sample. The temperature of the sample is measured contactless by a pyrometer. To determine the
growth velocity during crystallization a high-speed camera is used which enables time resolved
measurements of the rapid propagation of the solidification front through the undercooled melt [59].

The application of electromagnetic levitation on Earth is limited by the strong electromagnetic
levitation forces needed to compensate the gravitational force. This implies high heating power and
high temperatures and strong stirring of the melt by the eddy currents induced in the sample. This led


[58] D.M. Herlach, Annual Rev. Mat. Sci. 21 (1991) 23
[59] O. Funke, G. Phanikumar, P.K. Galenko, L. Chernova, S. Reutzel, M. Kolbe, and D.M. Herlach, J. Cryst. Growth 297
(2006) 211

                                                          42
                                                                                                 METCOMP MAP 114


to the complete loss of particles in our experiments. These limitations have been circumvented by the
use of TEMPUS in a reduced gravity environment. TEMPUS is an electromagnetic positioning device
developed by the German Space Agency. Different to electromagnetic levitation on Earth it is
operating with two independent coil systems one for positioning creating a quadrupole field of strong
field gradients but small magnetic field (high positioning forces, small heating effect) and another one
for heating generating a dipole field of high magnetic field and small field gradients (high power
absorption, small repulsing forces). This allows for a low convection level in the melt. Especially in
the system Ni-Ta + Ta2O5 a considerable amount of submicron particles is engulfed by the dendritic
solidification front. The microstructure of the solidified samples was analyzed by standard
metallographic techniques in a scanning electron microscope. The experiments in TEMPUS on
parabolic flights were a breakthrough for the study of pushing and engulfment behaviour and allowed
the controlled variation of experimental parameters for testing of theories.

Solidification experiments with Ni-based systems in TEMPUS

In the beginning of the project, the work had been concentrated on the selection of suitable metal
matrix / particle systems, preparation of material, processing on ground and evaluation of the solidified
microstructures by scanning electron microscopy (SEM). Solidification experiments in TEMPUS
during parabolic flight had been conducted and showed remarkable differences in microstructure
compared to terrestrial solidification. Engulfment of particles has been observed in the system Ni-Ta +
Ta2O5 particles (Figure 38). The fraction of engulfed particles of the total number of particles has been
determined and is given versus the particle diameter.




Figure 38: Scanning electron micrographs of Ni98Ta2 + Ta2O5 solidified in reduced gravity (left) and under 1g conditions
(right) at low undercooling ∆T ≈ 5 K. The dendrites (dark) are marked by interdendritic Ta enrichment (bright). Engulfed
particles are located in this area.


The results are summarized as follows:

(i) Undercooling level: No significant dependence of the fraction of engulfed particles on undercooling
level has been found in the undercooling range from 5 K to 105 K.




                                                          43
                                                                                                               METCOMP MAP 114



                                                                             Figure 39: Distribution of diameters d of engulfed
     1,0                                     1,0
                                                                             Ta2O5 particles in Ni98Ta2 (full red squares) and Ni96Ta4
                                                                             (open blue squares). The total fractions are 0.25 and




                                                    engulfed fraction
                                                    engulfed fraction
     0,8                                     0,8
                                                                             0.12-0.14 respectively. The respective undercoolings are
                                                                             84 K and 55-57 K (2 samples). The black curve
     0,6                                     0,6
                                                                             describes the engulfed or entrapped particles.
     0,4                                     0,4

     0,2                                     0,2

     0,0                                      0,0
        0,1               1                 10
                                 d / µm


(ii) Particle size: The range of particle sizes is in general from 0.25 µm to 3 µm. For a few samples it
has been extended to smaller particles down to 100 nm in diameter (nano-Ta2O5). The fraction of
engulfed particles is increasing with smaller particle size [57].
(iii) Ta content: Two different compositions of the metallic melt have been investigated, Ni98Ta2 and
Ni96Ta4. The fraction of engulfed particles decreases with increasing Ta content (Figure 39).
According to the description of particle engulfment by Pötschke and Rogge [52], one would expect the
following dependence of the critical velocity for engulfment vc on the concentration of the alloying
element C0 (here: Ta):

                                     vc ~ (1/C0)1/2

This means a lower critical velocity is expected for the system with 4 at% Ta and from this respect, a
tendency to a higher fraction of engulfed particles - in contradiction to our results.

(iv) Particle content: The volume fraction of the Ta2O5 particles has been varied from 1 vol% to
10 vol% in Ni98Ta2. At high volume fraction the particles have the tendency for clustering. Thus, a low
fraction of particles is desired for the study of particle engulfment.

A second question is whether the particles modify the process of solidification itself. For this
investigation a sample with high volume fraction is ideal. The respective results are given in an
Appendix to this section.

(v) Exchange of Ta by Hf: Ta has been replaced by Hf in the melt and the Ta2O5 particles by those of
HfO2. The system under investigation was then Ni99.5Hf0.5 + HfO2 particles. This replacement changes
the conditions for pushing and engulfment of particles and for the solidification especially by the low
partitioning coefficient (k0(NiHf) = 0.0548, k0(NiTa) = 0.8). In the model by Pötschke and Rogge the
substitution of Ta by Hf increases the expression C0ml/k0 by a factor of 8 [52]. The experiments
showed that the advantage of processing in low gravity was obvious, submicron particles of HfO2 were
engulfed. A quantitative analysis has not been carried out. The Ni/Ni5Hf eutectic appeared in the
samples and caused problems with the image analysis software [57]. This problem will be addressed in
the future.

(vi) Fe99Hf1 + HfO2: One Iron-based system has been investigated in parabolic flight. The system
contained a certain amount of Oxygen leading to Hf-Fe-O and Hf-Fe-O-C phases [57]. Future
experiments with this system in TEMPUS will be done after repeated ground experiments with
improved sample material.



                                                                        44
                                                                                                       METCOMP MAP 114


Discussion of the results

The experimental results are discussed with the models by Pötschke and Rogge (PR) [52], Stefanescu,
Dhindaw, Kacar and Moitra (SDKM) [53], and Shangguan, Ahuja and Stefanescu (SAS) [54]. All
these models are based upon the assumptions of a planar interface and steady state conditions. They
are developed by estimating various forces acting on a particle in liquid environment in front of the
solid-liquid interface. Commonly viscous drag forces, hydrodynamic buoyancy and forces due to van
der Waals interaction of particle and interface are taken into account. The three models derive the
critical velocity (formulas ,       * ** ***
                                      ,        in section III.4) from (i) a balance of forces acting on the
particle in front of the solid/liquid interface and (ii) the modification of the homogeneous temperature
field in the melt in the vicinity of particle and s/l interface. While the models are very similar in the
treatment of (i), they have considerable differences in (ii). SDKM give a qualitative description of the
heat transfer in the liquid between particle and s/l interface, which arises from the different heat
conductivities of particle and liquid. The result is a qualitative change in the shape of the s/l interface
determining the critical velocity. SAS extend this treatment by solving the heat conduction equation
explicitly and calculate the shape of the s/l interface. PR consider in addition the influence of an
alloying element on the temperature field and the shape of the s/l interface. It turns out that the
alloying element has a strong influence on the critical velocity. In the case of no alloying element, SAS
and PR yield the same formula for the critical velocity.

The critical velocities were calculated according to the models mentioned above using the materials
parameter of Ni98Ta2 + Ta2O5 as given in Table 2. The results of the values of the critical velocities as
calculated within the different models are collected in Table 3.

Table 2: Physical parameters used for calculations of critical velocities for engulfment of Ta 2O5 particles during
solidification of Ni98Ta2 alloys

       Parameter              Unit                Value            Parameter              Unit                Value
           co                 at%                   2                  Dl                 m2/s               2.5·10-9
           ml                K/at%                  5                  ∆                 N/m                  0.2
           k0                  ---                 0.8                  kp               W/mK                   4
           Gl                 K/m                 50 000                kl               W/mK                   40
                                                                        
                                  -10                                                      -3
           ao                10         m          2.32                                 10 Pa s                 5
            R                  µm                  0.1



The experimental results from samples processed in reduced gravity do not show any significant
dependence of the fraction f of particles engulfed by growing dendrites on the undercooling
temperature and therefore on the dendrite growth velocity. The growth velocity changes over three
orders of magnitude in the range of undercooling accessible in levitation experiments. We estimate an
upper limit for the critical growth velocity from the lowest achieved value of solidification velocity
(Figure 40 left) for particle engulfment as

         vc(R) < 10 cm/s       for          0.1 µm < R < 10 µm

This result is in agreement with estimations within the different models as listed in Table 3. The SAS
and SDKM models were derived for pure metals only. The PR model also regards constitutional
effects in alloys. As can be seen from the results of vc by PR and SAS as given in Table 3 the critical
velocity vc is very much reduced in Ni98Ta2 alloy compared to pure Ni. This observation could be
understood by the fact that in alloys a concentration gradient is established in addition to the negative
                                                             45
                                                                                                       METCOMP MAP 114


temperature gradient ahead the solid/liquid interface. Since the concentration gradient is much deeper
than the temperature gradient the morphology is controlled by solute redistribution rather than heat
redistribution. It leads to the formation of much more dendrites being smaller than in case of pure
thermally controlled dendrites. This may explain the difference of the estimated critical velocity vc <
0.1 m/s as found in the present work for the engulfment of Ta2O5 particles in the alloy Ni98Ta2 with
many small dendrites compared to the value of vc ≈ 8 m/s as reported for the engulfment of Ta2O5
particles in pure Ni undercooled by melt fluxing in DTA experiments with fewer but larger dendrites
[60].

Table 3: Critical velocities vc for engulfment of Ta2O5 particles during solidification of undercooled Ni98Ta2 alloy
calculated within three different models PR, SAS (n=2) and SDKM (n=4).

        Critical velocity           Unit                     PR                      SAS                   SDKM
               vc                   cm/s                   0.00125                   7.73                    0.98


Conclusions and outlook for WP 4 and WP 5

During the course of this project processing in TEMPUS during parabolic flights became a well
established experimental method and essential for the workpackages WP 4 and WP 5. In total 19
samples of various metal + particle systems have been processed in TEMPUS (5 in the second period
of the project). In nearly all of them particles were retained in the melt. Due to the reduced convection
under low gravity conditions the rate of success is much higher than in EML processing under
terrestrial conditions. In addition the solidification front has been observed with the recently integrated
high speed camera. The solidified microstructures and the observed solidification front have been
correlated. The influence of low gravity or the presence of particles on the solidification itself is
negligible in the region of the observations.

Microstructure analysis of material solidified in low gravity yields differences to results from
terrestrial experiments. The main results of low gravity experiments are:

         submicron particles are found in cluster-free regions - one fraction is engulfed, another fraction
          is entrapped

         the fractions of engulfed and entrapped particles are independent from undercooling in the
          range from <5 K to 104 K

         the fraction of engulfed particles increases with decreasing particle size

         the fraction of engulfed particles decreases with increasing amount of alloying element Ta in
          the Ni-Ta + Ta2O5 particles system

         in the Cu-Ni + Ta2O5 particles system the fraction of engulfed particles increases with
          increasing amount of Ni in the melt, which improves wetting of the particles.

In a new period of this project, we will use all available parabolic flight opportunities, because they are
the only way to get reliable data concerning pushing and engulfment of particles by a dendritic
solidification front. We plan to extend the results for the Ni-Ta + particles systems by varying the Ta
content. The investigations on solidification in the pure systems will be completed by experiments in
the low undercooling range using the high speed camera and EBSD. The results on the Cu-Ni systems


[60] G. Wilde and J.H. Perepezko, Mat. Sci. Eng. A 283 (2000) 25

                                                             46
                                                                                                                  METCOMP MAP 114


will be extended in order to establish a larger database. Fe-based systems will be investigated using a
low oxygen containing Fe powder.
It is still an open question to establish a model, which describes pushing and engulfment under non-
equilibrium conditions in undercooled metallic melts. Several of our results can be described by
existing models, for example the low critical velocities for engulfment. Other results as the dependence
on the alloying element are in contradiction. The reason is certainly that the specific non-equilibrium
conditions in the undercooled melt are not in models for planar front solidification. The importance of
especially the mass transport on the wetting layer and, thus, on pushing and engulfment behaviour has
been shown by Pusztai (section III.3). Theories for dendrite growth in undercooled melts analyze the
heat and mass transport in front of the dendrite tip and allow to calculate dendrite growth velocity vs.
undercooling.

Appendix

The velocity of the solidification front is an important parameter for the description of the interaction
of inert particles with a dendritic solidification front. In the past, we derived the velocity of the front in
the low gravity experiments by use of terrestrial measurements at the same undercooling value. This
procedure neglects low gravity effects on the solidification velocity as well as an effect of the particles
on the velocity. Recently, a high speed camera has been integrated in TEMPUS and the heat emission
of the propagating front can be observed directly in low gravity experiments. The observation in
TEMPUS has three advantages: (i) the velocity measurements are conducted with a real sample with
particles in the melt, (ii) the observations are under low gravity conditions, and (iii) the geometry of
the front itself is observed and allows the link to further microstructural analyses. Especially at low
solidification velocities, the fluid flow may have an influence on heat and mass transport at the
solidification front and, thus, on the solidification velocity. Measurements of the solidification velocity
versus undercooling temperature are given in Figure 40 for various experimental conditions.

           40                                                                     40

           35                                                                     35

           30                                                                     30

           25                                                                     25
                                                                        v (m/s)
 v (m/s)




           20                                                                     20

           15                                                                     15

           10                                                                     10                          terr. Ni96Ta4
                                       terr. Ni98Ta2
                                       PF Ni98Ta2 + Ta2O5                                                     PF Ni96Ta4 + Ta2O5
            5                                                                      5
                                       PF Ni98Ta2                                                             PF Ni96Ta4
            0                                                                      0
                0   50   100   150     200    250      300   350                       0   50   100   150    200     250      300   350
                                 TUC (K)                                                                TUC (K)


Figure 40: Velocity of solidification as function of undercooling temperature of Ni98Ta2 (left) and Ni96Ta4 (right), with and
without Ta2O5 particles solidified in reduced gravity and under 1g conditions respectively. Terrestrial measurements of
samples with particles are not shown, as a large fraction of particles leaves the melt before solidification. All values are
derived from high speed camera observations.

The velocity of solidification shows the same behaviour under terrestrial and low gravity conditions in
the pure alloy at high velocities > 5 m/s. This means that we can derive the solidification velocity from
measurement of the undercooling temperature under low gravity conditions together with terrestrial
undercooling experiments using the v vs. TUC relation. Samples with particles in TEMPUS yielded up

                                                                   47
                                                                                                   METCOMP MAP 114


to now only undercooling values < 105 K. In this region the solidification velocity is in the order of
1 m/s, which is similar to the velocity of melt convection in terrestrial levitation. This means, if we
expect an effect of melt convection on the solidification velocity, then it should appear in this low
velocity region [61]. Up to now such effect cannot be excluded, because terrestrial measurements in
this region are lacking.

The high speed camera shows the propagating solidification front due to the release of latent heat
during recalescence. As metals are not optically transparent, the pictures show the intersection of the
solidification front with the surface of the sample. The front shows a characteristic shape, which is
certainly determined by the geometrical features of solidification, e. g. an array of dendrites, which
belong to a single crystal, a few crystals or many crystals. It has been found that the solidification front
has a typical shape depending on the undercooling temperature.




               0.9 ms                1.5 ms                   2.1 ms                   2.7 ms                    3.3 ms

                                  TUC = 45 K      Figure 41: Propagation of the (bright) solidification front in
                                                  Ni98Ta2 + 10 vol% Ta2O5 particles in TEMPUS (type I). The view
                                                  on the elliptical sample is by the dark edges of the sample cage.
                                                  The full length of the sequence is 4.7 ms.
                                                  The undercooling was 45 K, the average velocity calculated by
                                                  the diameter of the sample divided by the length of the sequence
                                                  is 1 m/s. A precise determination of the velocity of the solidi-
                                                  fication front is not straightforward unless the geometrical details
               3.9 ms                  4.3 ms     of the solidified crystal are known.


Figure 41 shows a sequence of the propagating solidification front of Ni98Ta2 +10 vol% Ta2O5
particles. Remarkable is that the front shows a clear geometrical shape, which consists of a few straight
lines. This behaviour can be found for undercooling temperatures around 50 K (type I).

                                                                                      Figure 42: Two sequences of
                                                                                      solidification    (Ni96Ta4)   in
                                                                                      TEMPUS. The undercooling
                                                                                      temperature was 164 K (upper,
                                                                                      type IV) and 155 K (lower,
                                                               TUC = 164 K            type III). The time difference
                                                                                      between neighboring images is
                                                                                      0.033 ms. The spherical front of
                                                                                      the upper sequence is typical
                                                                                      for       high      undercooling
                                                                                      (>160 K).      In    a    region
                                                                                      140 K<TUC<160 K the front
                                                                                      shows features of 500 µm in
                                                                                      size.
                                                               TUC = 155 K



[61] H. Hartmann, P.K. Galenko, D. Holland-Moritz, M. Kolbe, D.M. Herlach, O. Shuleshova, J. Appl. Phys. 103 (2008)
073509

                                                         48
                                                                                                  METCOMP MAP 114


In a region 80 K<TUC<140 K the front has a pronounced zig-zag shape of millimeter size (type II). For
undercooling temperature 140 K< TUC<160 K there is a transition regime (Figure 42 lower), the front
is nearly straight, but shows small features 500 µm in size (type III). In the high undercooling range
>160 K the front is spherical (Figure 42 upper). No features are resolved (type IV). Concerning the
samples processed so far, there are no differences in the morphology of the solidification front with
respect to processing in low gravity or under terrestrial conditions.

The grain orientation of solidified samples has been analyzed by use of EBSD (electron back scatter
diffraction) in the scanning electron microscope. A cross section is given in Figure 43. It shows large
yellow and blue areas, each of them with dark and colored dots. Identical colors represent identical
crystallographic directions according to the color code. The sample consists of three large grains (blue
and yellow) inside of which we observe small grains (colored dots). The dark dots represent the
particles, which are not analyzed with respect to their crystallography. As the image is a cross section,
the two yellow grains might be connected in the bulk. A detailed analysis of the two main grain
orientations (blue and yellow) shows that these crystals are twins: They have a {111}-plane in
common and this {111}-plane is a mirror plane. This has consequences for the determination of the
solidification velocity. Usually the solidification velocity is determined for the propagating <100>
dendrite. In a {111}-plane is never a <100> direction, this means that the [100]-dendrites of the two
grains of the twins move definitely into different directions. A precise determination of the
solidification velocity should take into account the geometry of each grain. This will be done in the
future.


                                                       Figure 43: Overview cross section of the sample
                                                       described in Figure 41 (type I) in the light of an EBSD
                                                       measurement. Large yellow and blue areas are in twin
                                                       orientation. In the center is a large pore. The IPF color
                                                       code is given for Y direction.




The grain structure of a sample, which solidified as type IV at an undercooling of TUC = 201 K is given
in Figure 44. No preferential crystallographic direction is found in this grain refined structure; it is
typical for materials solidified from a highly undercooled melt. Schwarz et al. [62] explained grain
refinement by the partly remelting of dendrites during the period of recalescence when the temperature
rises to the melting point.

In the intermediate region of undercooling temperatures 140 K<TUC<160 K the solidified
microstructure is complex (Figure 45). One has the impression that the solidification started from the
left side of the sample. The size of the grains ranges from 2 mm down to 50 µm. A detailed analysis of
the grain orientations shows that several of the large grains are pairwise in twin orientation. Probably
the growth of several large grains has been observed directly with the high speed camera (Figure 42
lower).



[62] M. Schwarz, A. Karma, K.E. Eckler, D. M. Herlach, Phys. Rev. Lett. 73 (1994) 1380

                                                         49
                                                                                                 METCOMP MAP 114


The high speed camera observations and the microstructure analyses both lead to the same
interpretation. With increasing undercooling temperature the number of growing crystals is increasing.
Around TUC=50 K a single crystal or a twin is growing, leading to a solidification front with few
straight segments. In a region 80 K<TUC<140 K the front has a large zig-zag shape. The microstructure
of such solidified samples has not yet been investigated. From 140 K<TUC<160 K the front is more or
less straight, but in detail shows features of 500 µm size. These might correspond to the medium size
grains, which are observed in the solidified sample (Figure 45). The picture becomes simpler at high

                                                                 Figure 44: Grain refined structure in a sample
                                                                 solidified in TEMPUS from a highly undercooled
                                                                 melt (Ni98Ta2, TUC=201 K). The morphology of the
                                                                 solidification front was classified as type IV. The
                                                                 average grain diameter is 40 µm.




              =200 µm; GB+IPF_Y1; Step=4 µm; Grid200x150




undercooling TUC>160 K. A great number of dendrites grows in all directions from a nucleation point
leading to a featureless spherical solidification front. The corresponding solidified material is grain
refined due to partly remelting of the dendrites.

                                                            Figure 45: Overview cross section of the sample
                                                            described in Figure 42 lower (type III, Ni96Ta4,
                                                            TUC=155 K) in the light of an EBSD measurement.
                                                            Several of the large grains are pairwise in twin
                                                            orientation.




The investigation of the solidification of undercooled Ni98Ta2 and Ni96Ta4 melts with and without
particles by the use of a high speed camera and the EBSD technique gives an impression of the
complexity of non-equilibrium solidification. So far no influence of convection on the solidification
has been found in these alloys. This means that all effects, which have been found concerning particle
engulfment are effects of the reduced convection in low gravity processing on particle dendrite
interaction. Future investigations will be focused on the low undercooling region where convection
may have an influence on dendrite growth.



                                                           50
                                                                                      METCOMP MAP 114




III.6   Publication list
E. Boehm-Courjault, F. Gonzales, A. Jacot, F. Kohler, A. Mariaux, C. Niederberger, M. Salgado and
M. Rappaz, “EBSD: A Powerful Microstructure Analysis Technique in the Field of Solidification”, J.
Microscopy (2008, to appear)
S. Dobler, T. S. Lo, M. Plapp, A. Karma and W. Kurz, "Peritectic coupled growth", Acta Materialia 52
(2004) 2795-2808
S. Dobler and W. Kurz, "Phase and microstructure selection in peritectic alloys under high G-V
ratio", Z. Metallkd. 95 (2004) 592-595
S. Eck, J. Mogeritsch, A. Ludwig, "Experimental Observation of Convection during Equiaxed
Solidification of Transparent Alloys", Mater. Sci. Forum 508 (2005) 157-162
L. Gránásy, T. Pusztai, J. A. Warren, J. Phys.: Condens. Matter. 16, R1205 (2004).
L. Gránásy, T. Pusztai, T. Börzsönyi, in: Handbook of Theoretical and Computational Nanoscience,
eds. M. Rieth and W. Schrommers, American Sci. Publ., in print
L. Gránásy, T. Pusztai, J. A. Warren, J. F. Douglas, T. Börzsönyi, V. Ferreiro, "Growth of 'dizzy
dendrites' in a random field of foreign particles", Nature Materials 2 (2003) 92
L. Gránásy, T. Pusztai, T. Börzsönyi, J. A. Warren, J. F. Douglas, "A general mechanism of
polycrystalline growth", Nature Materials 3 (2004) 645-650
L. Gránásy, T. Pusztai, J.A. Warren, and D. Saylor, "Phase field theory of heterogeneous nucleation",
Phys. Rev. Lett. 98 (2007) 035703
F. Kohler, T. Campanella, S. Nakanishi and M. Rappaz, “Application of Single Pan Thermal Analysis
to Cu - Sn peritectic alloys”, Acta Mater. 56 (2008) 1519–1528. Corrigendum Acta Mater. 56 (2008)
3708 – 3709
F. Kohler, "Peritectic solidification of Cu-Sn alloys : microstructure competition at low speed", PhD
thesis, EPFL, Lausanne, Switzerland, 4037 (2008) http://library.epfl.ch/theses/?nr=4037
F . Kohler , L . Germond , J . Wagnière and M . Rappaz, “Peritectic solidification of Cu – Sn alloys:
Microstructural competition at low speed”, Acta Mater. 57 (2009) 56-68
F. Kohler, T. Jauzein, M. Plapp and M. Rappaz, "Multi-phase field simulation of directional
solidification at low speed of hypoperitectic Cu-Sn alloys", Acta Mater. (2009, to be submitted)
M. Kolbe, G. Eggeler, L. Gránásy, D. M. Herlach, A. Ludwig, M. Rappaz, (Project Partners), Corus
Technology BV (NL), Schwermetall Halbzeugwerke GmbH & Co.KG (D), Swissmetal SA (CH),
Thyssen Krupp Stahl AG (D), Wieland-Werke AG (D) (Industrial Partners), "Metastable Solidification
of Composites: Novel Peritectic Structures and In-Situ Composites", in: Microgravity Applications
Programme, A. Wilson (ed.), ESA Publications Division, Vol. 1290, ESTEC (NL), 2005, pp.50-61
M. Kolbe, X. R. Liu, T. Volkmann, R. Röstel, P. K. Galenko, G. Eggeler, B. Wei, D.M. Herlach,
"Interaction of solid ceramic particles with a dendritic solidification front", Mater. Sci. Eng. A 375-377
(2004) 524-527
M. Kolbe, T. Lierfeld, G. Eggeler, D. M. Herlach, "Experimental investigation of the interaction of a
dendritic solidification front with foreign particles", Microgravity sci. technol. XVIII-3/4 (2006) 170-
173
M. Kolbe, J. Gao T. Lierfeld, S. Schneider, G. Eggeler, D. Herlach, "The influence of electromagnetic
processing under terrestrial and parabolic flight conditions on solidifying metallic melts", in: EPM

                                                   51
                                                                                       METCOMP MAP 114


2006, eds.: S. Taniguchi, Proceedings of the 5th International Symposium on Electromagnetic Processing
of Materials, Sendai, Japan, 23-27 Oct. 2006, ISIJ, pp. 357- 362
M. Kolbe, T. Lierfeld, S. Schneider, G. F. Eggeler, D. M. Herlach, "Interaction of a dendritic
solidification front with ceramic particles", JASMA 25 (2008), Proceedings of 3rd International
Symposium on Physical Sciences in Space, 22-26 Oct 2007, Nara (Japan), pp. 463-466
T. Lierfeld, M. Kolbe, J. Gegner, G. Eggeler, D. M. Herlach, "Einbau keramischer Dispersoide in
Metalle mittels unterschiedlicher Techniken der Unterkühlung metallischer Schmelzen", Zeitschrift für
Werkstoffe, Wärmebehandlung, Fertigung 60 (2005) 64-70
T. Lierfeld, M. Kolbe, D. Herlach, G. Eggeler, "Embedding of ceramic particles in metals using
different techniques for undercooling of metallic melts", Mater. Sci. Forum 508-509 (2006) 307-312
T. Lierfeld, M. Kolbe, H. Nagai, T. Okutani, G. Eggeler, D.M. Herlach, "Experimental investigation of
the interaction of an advancing dendritic solid/liquid interface with ceramic particles", Proceedings of
the 5th Decennial International Conference on Solidification Processing, Sheffield, July 2007, pp. 57-
60
T. Lierfeld, P. Gandham, M. Kolbe, T. Schenk, H. Singer, G. Eggeler, D. M. Herlach, "Particle
incorporation in metallic melts during dendritic solidification - in-situ visualization and undercooling
experiments under reduced gravity", Mater. Sci. Eng. A 449-451 (2007) 689-692
T. Lierfeld, M. Kolbe, G. Eggeler, D. M. Herlach, "Interaction of small ceramic particles with a
dendritic solidification front", Adv. Eng. Mater. 10 (2008) 547-533
D. Lewis, T. Pusztai, L. Gránásy, J. Warren, W. Boettinger, JOM - J. Min. Met. Mat. 56 (2004) 34
T. S. Lo, S. Dobler, M. Plapp, A. Karma and W. Kurz, "Two-phase microstructure selection in
peritectic solidification: from island banding to coupled growth", Acta Materialia 51 (2003) 599-611
C. Pfeiler, M. Wu and A. Ludwig, "Influence of argon gas bubbles and non-metallic inclusions on the
flow behavior in steel continuous casting", Mater. Sci. Eng. A 413-414 (2006) 115-120
C. Pfeiler, A. Ludwig, M. Wu, "Simulation of particle and bubble motion in a steel continuous caster",
11. Int. Conf. on Modelling of Casting, Welding and Advanced Solidification Processing, Opio (F)
(2006), to be published
T. Pusztai, G. Tegze, G.I. Tóth, L. Környei, G. Bansel, Z. Fan, and L. Gránásy, "Phase-field approach
to polycrystalline solidification including heterogeneous and homogeneous nucleation", J Phys.:
Condens. Matter 20 (2008) 404205
L. Ratke, Ch.-A. Gandin, R.H. Mathiesen, M. Rappaz and S. Rex, “Materials solidification physics in
space”, Europhysics News 39 (2008) 22-24
G. I. Tóth, L. Gránásy, "Phase field theory of interfaces and crystal nucleation in a eutectic system of
fcc structure: I. Transitions in the one-phase liquid region", J. Chem. Phys. 127 (2007) 074709
G. I. Tóth, L. Gránásy, "Phase field theory of interfaces and crystal nucleation in a eutectic system of
fcc structure: II. Nucleation in the metastable liquid immiscibility region", J. Chem. Phys. 127 (2007)
074710
J. A. Warren, L. Gránásy, T. Pusztai, T. Börzsönyi, G. Tegze, J. Douglas, "The influence of foreign
particlesin the formation of polycrystalline solidification patterns", in: Solidification processes and
microstructures: A symposium in honor of Wilfried Kurz, M. Rappaz et al. (eds.), TMS Annual
Meeting 2004, Charlotte (USA), pp. 379-384
J.A. Warren, T. Pusztai, L. Környei, L. Gránásy, "Phase Field Approach to Heterogeneous Crystal
Nucleation in Alloys", Submitted to Phys. Rev. B.


                                                    52
                                                                                  METCOMP MAP 114


M. Wu, A. Ludwig, "Influence of Phase Transport Phenomena on Macrosegregation and Structure
Formation During Solidification", Adv. Eng. Mater. 5 (2003) 62-66



IV     Inventory of Intellectual Property Rights

In the original contract the claims of the partners with regard to Intellectual Property Rights made
available to the Project were described in Annex 6, Appendix 4. In the course of the project, no
inventions were made. The actual claims of the partners after the reconstruction of the consortium in
the second period are as follows:


1. Background Information Statement DLR:

DLR makes available to the project its technical and scientific knowledge on:

       (i) Electromagnetic Levitation

       (ii) Electromagnetic Positioning in Space (Patent Nr. 36 39 973, Deutsches Patentamt)


2. Background Information Statement Laboratory of physical metallurgy (LMPH-EPFL):

LMPH-EPFL makes available to the project its technical and scientific knowledge on

       (i)     Liquid Metal Cooling (LMC) Bridgman apparatus
       (ii)    Microstructure analysis (metallography, SEM, EBSD, EDX, EWX)
       (iii)   Access to SLS for X-ray tomography experiments
       (ii)    Computer modelling of directional solidification
       (iii)   Phase and microstructure selection modelling


3. Background Information Statement ThyssenKrupp Stahl AG (TKS):

TKS makes available to the project its technical and scientific knowledge on:

       (i)     Steel materials and the production of steel
       (ii)    Application of steel in various sectors
       (iii)   Tailoring materials and effects of production routes on material properties
       (iv)    Metallographic investigations as well as physical and chemical testing of materials
       (v)     Experimental techniques for the investigation of solid-solid phase transformation
               kinetics including numerical methods of data exploitation


V      Evaluation of Technology

Many technical alloys, steels and copper alloys to mention just the economically most important ones,
are peritectic alloys. Recently, new metastable microstructures were detected in solidification of
peritectics, which make them suitable for the preparation of in situ composite materials. The

                                                  53
                                                                                     METCOMP MAP 114


solidification behaviour is not well understood in particular, because the formation of in situ composite
microstructure from liquid sensitively depends on convection that is always present under 1g-
conditions. For the alloys under study in this project phase selection maps were established. They may
allow technological exploitation.
Composite materials consisting of various phases and substances are of high interest since they
combine advantageous properties of different phases within one material. Ceramic particles are often
added to metallic alloys to reinforce metallic alloys and make them applicable at high temperatures.
But for dendritic solidification that is of high importance in casting processes a description of the
conditions of homogeneous distribution of particles in a metallic matrix is still lacking. On the other
hand industrial production processes require particle pushing during casting processes in order to
remove foreign phases and purify the cast material. The conditions of pushing or embedding of
particles depend on the interaction with the solidification front. The studies undertaken in this project
may in the future be extrapolated to the conditions of industrial production processes.


VI     Initiatives towards non-space industries

The industrial consortium of this project consists of five large and medium companies (Corus
Technology BV, Schwermetall Halbzeugwerk GmbH & Co. KG, Swissmetal SA, ThyssenKrupp Steel
AG and Wieland-Werke AG), which are involved in the production of the main metallic alloys as
steel, copper-based and aluminum. No further initiatives towards non-space industry have been
undertaken.




                                                   54

				
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