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Solidification reactions in undercooled alloys
52
Materials Science and Engineering, A 179/A 180 (1994) 52-56
Solidification reactions in undercooled alloys J. H. Perepezko Department of Materials Science and Engineering, University of Wisconsin, 1509 UniversityA venue, Madison, W153706 (USA)
Abstract It is now recognized that rapid solidification conditions can be achieved with slow cooling rates provided that the liquid is undercooled substantially prior to nucleation. In fact, many of the novel metastable microstructures produced by rapid solidification require the consideration of an undercooled liquid for analysis. In general, rapid solidification techniques involve either constrained growth in which the solid phase formation is limited by the rate of heat extraction or delayed nucleation of the solid followed by unconstrained growth. With delayed nucleation methods such as the droplet emulsion technique direct measurement of undercooling is available for analysis of metastable phase formation. In fine droplet samples an effective nucleation isolation allows for undercoolings of about 0.3 Tmwith a limit that is usually set by heterogeneous nucleation. Processing variables can be used to control the undercooling and produce a transition in solidification reactions. In this case the use of metastable phase diagrams is important for the analysis of product structures and pathways during solidification and solid state treatments. A key to the understanding of structural evolution is the consideration of competitive nucleation and growth kinetics and thermal history, which can also provide a model for control of solidification reactions as demonstrated in selective alloys.
1. Introduction The most commonly used methods for generating high liquid undercoolings prior to solidification involve a rapid quenching from the melt. Indeed, the main attention in rapid solidification processing (RSP) methods has been focused upon the attainment of high cooling rates in the range 1 0 4 - 1 0 8 ° C s -1. With these methods small liquid volumes are quenched either by contact with the substrate as in melt spinning and directed energy beam processing or by an inert gas as in atomization. However, the most important impact of RSP can be related directly to the solidification structures. The wide range of reported microstructural variations encompass not only highly refined equilibrium phase mixtures but also novel microstructures. While the development of refined structures during RSP can be viewed in terms of the decrease in local solidification time, the understanding of novel microstructures representing non-equilibrium phases requires a consideration of the level undercooling at the onset of nucleation in addition to rapid growth kinetics [1-3]. In fact, the available free energy for non-equilibrium phase formation is directly related to the amount of undercooling. In most work on RSP cooling rate is used as a convenient process variable. However, in solidification studies of undercooled samples it has been demonstrated that undercooling can also be 0921-5093/94/$7.00 SSDI 0921-5093(93)05685-I
treated in a meaningful way as a process variable. In many respects undercooling is a more fundamental parameter than cooling rate. From this experience an examination of solidification reactions in undercooled liquids can provide the necessary basis for the prediction of microstructural transitions that can be applied to guide alloy design in processing to optimize the yield of the preferred microstructure. Some selected highlights can serve to illustrate the variety of pathways for microstructural development that are available in undercooled alloys.
2. Development of undercooling Crystallization of a large continuous bulk liquid metal will be catalyzed by the most potent nucleation site present, so that a means of circumventing the effect of catalytic sites such as oxides or container walls is necessary in order to observe extensive undercooling. Following from this basis, an effective approach to obtain large undercoolings is to disperse the bulk liquid into a collection of fine droplets. This effectively isolates potent nucleation sites into a small fraction of the droplet population as illustrated in Fig. I(A). As a stabilized liquid droplet emulsion is cooled, those drops containing potent nucleants will freeze at low undercoolings, but the majority will not freeze until © 1994 - Elsevier Sequoia. All rights reserved
J. H. Perepezko
/
Solidification reactions in undercooled alloys
53
(a) ATOMIZATION
o®0 0 000(90 0 o°®oo0o ® °0®0 ° oOoq ®
L
(b) TzI
(c)
,
Tz I
~
T m .T. i. I. ~. .-. . . . . . . . . . . . . .~. . ., ,
X=
×
i
\L*i
i
Z~ -
i
° , -.. ~""==.. . .L. . . . l /
WHEEL MELT SPINNING
/T E
SURFACE MELTING
Fig. l. Schematic representation of three common methods of rapid solidification processing involving (A) powder atomization, (B) melt spinning and (C) surface melting of thin layers. In B and C the position of the liquid-solid interface is shown relative to the alloy liquidus for conditions that yield liquid undercooling. Some of the vertical (Z) distances are exaggerated for clarity.
reaching the maximum nucleation undercooling level, which can range from 0.3 Tm to 0.4 Tm [4]. While nucleant isolation accounts for the development of undercooled melts during powder atomization, it is worthwhile to consider how undercooled melts can develop during other common RSP methods. For example, during the rapid freezing of melt steams, as in melt spinning shown in Fig. I(B), undercooled conditions will develop if the liquid exists for some time or distance in contact with the wheel without nucleation. It is evident that a combination of high undercooling and high heat transfer coefficient promotes the most rapid solidification conditions [3]. Alternatively, during surface melting with a rapidly moving or pulsed heat source an undercooled melt zone can develop in front of the resolidification interface if the rate of growth of the unmelted substrate is insufficient to keep up with the rate of heat extraction to the substrate. The development of metastable crystalline phases as well as amorphous products during both melt spinning and surface melting is clear evidence for the existence of an undercooled melt [5].
3. Solidification reactions
One of the first demonstrations of the microstructural benefits of RSP was the development of extensions of primary phase solid solubility [6]. This has often involved eutectic alloy systems where the limit to supersaturation can be analyzed effectively in terms of
the T0 curves to reveal the thermodynamic bounds on supersaturation, including the composition range where glass formation may become kinetically favored
[3]. The experience in solubility extension in peritectic systems is limited compared with eutectics, but recent results on Sn-Sb peritectic alloys (Fig. 2) serve to illustrate some new behavior [7]. With droplet samples and undercoolings in excess of 100°(3 below the primary phase liquidus, supersaturated fl-Sn solutions have been produced for solute levels up to 20 at.% Sb in agreement with previous observations by rapid quenching methods. It is of interest to consider the limit to supersaturation. In Fig. 2 the To curve is also indicated and reveals that the extended liquidus and solidus for the fl solution converge to a congruent maximum in the metastable domain. Following the congruent point, the metastable liquidus and solidus drop to much lower temperatures as a consequence of the melting temperature of Sn in the Sb structure. As a result, the required undercooling below the stable liquidus to reach the To temperature increases rapidly. The behavior is unique to peritectic systems and reveals a new limitation to solubility extension. In addition to extension of solid solubility, the synthesis of metastable intermediate phases has also been an important capability of RSP. Since a metastable intermediate phase will also have a melting temperature below that of the equilibrium intermediate phase, the available free energy for metastable phase formation will be less favorable than that for equilibrium
J. H, Perepezko
54
/
Solidification reactions in undercooled alloys
260 L÷ •
~
vo
Lo o
w250 ~Z
k-
240 el_
B÷zr
W
23o
P / . . . .
0
i
.
,
,
I
,
,
,
,
t
,
5 I0 15 ATOMIC PERCENT Sb
,
.
,
20
l
Fig. 2. Tin-rich portion of the Sn-Sb phase diagram showing metastable extensions of the fl liquidus and solidus as well as T0R-L (chain curve). A metastable melting point maximum is indicated in the vicinity of 15 at.% Sb [71.
phase freezing. Thus the successful development of a metastable phase requires faster nucleation and growth kinetics than for the equilibrium phase. Moreover, the maintenance of the kinetic preference favoring the metastable intermediate phase is limited by the thermal history during recalescence following nucleation at high undercooling. In the Mn-AI system for near-equiatomic alloys the equilibrium phase is an e (h.c.p.) structure, but a metastable r (L10) phase has been reported to form from e during solid state treatments [8]. By containerless processing and splat cooling, the metastable r phase has recently been reported to form directly from the melt [9]. From other thermodynamic measurements the melting temperature of the r phase is approximately 87 K below the melting temperature of e [10]. As indicated by the microstructures in Fig. 3, r phase formation initiates on the cooling substrate, but, during rapid growth, latent heat is released during recalescence. As a consequence, a decreasing interfacial undercooling and interface velocity develop with the progress of solidification to allow for the heterogeneous nucleation of the equilibrium e phase at the moving ~-L interface to yield a final freezing of the central portion of the splat to the e phase.
4. Thermal-history-modified microstructures The sample thermal history during RSP is known to play a central role in the overall solidification kinetics. It has been included in a number of models of RSP [10, 11] but there are some further issues relating to the influence of recalescence behavior on microstructure development. One of the consequences of the recal-
(b)
' l l l l l
Fig. 3. (a) Scanning electron micrograph of the cross-section of splat-quenched Mn0.ssA10.43~Cil.lll7 alloy. (b) Optical micrograph of magnetic powder patterns on the cross-section of splatquenched Mn0ssAl~.~33C0.m7 alloy. Note the confinement of the colloid magnetic particles over the ferromagnetic r phase (dark) and their absence on the non-ferromagnetic e phase (bright)19].
escence history is illustrated in Fig. 3, where a change in product structure and/or microstructure size scale develops owing to the effect of recalescence. During surface melting with directed energy beams another microstructural feature has been identified due to the interplay of solidification kinetics and latent heat release. In an equiatomic Fe-V alloy a high temperature b.c.c, a phase forms from the melt under equilibrium conditions and then transforms to a stable low temperature o phase. If a single-phase substrate of o is subjected to pulse laser annealing, rapid heating bypasses the o ~ a transition so that o melts and places the liquid at an undercooling of 54 K [ 12]. As shown in Fig. 4 when the laser pulse is turned off, the o contains a layer of liquid. During cooling, the a phase nucleates at the free surface and begins to grow, releasing latent heat which causes propagation of the liquid further into
J. H. Perepezko Schematic
of
o->a
in F e V
a~
I
Iil o
a
55
SOLIDIFICATION PATHWAYS
-~
o
a)
a
/ Solidification reactions in undercooled alloys
1
b) a
c)
7"
e
Martensite
0
__d)
e)
Fig. 4. Schematic diagram showing the process by which the latent-heat-assisted transformation of o ~ a is believed to occur
in Fe-V [13].
the o phase and allows for the development of additional surface liquid. Associated with the latent heat release is a slowing of the a phase growth until further development of an intermediate a layer occurs to yield a bimodal microstructure of a phase, i.e. equiaxed grains in contact with o and columnar grains extending outward to the free surface. The net effect is the enhancement of the transformed depth due to the latent heat released by solidification of the a phase beyond that due to the initial laser energy. This phenomenon can be understood by recognizing that the scaling of the melt depth is enhanced by a factor related to the difference in the heat of fusion between the a and the o phase [13]. Another factor of importance in defining the influence of sample thermal history on the final microstructure development during RSP is related to the influence of post-solidification solid state transformations. A clear illustration of this aspect can be developed by considering the microstructure development in undercooled Fe-Ni alloys [14]. Over the range from Fe-10wt.%Ni to Fe-30wt.%Ni large millimeter drops were determined to undercool by about 160 °C during processing in a drop tube apparatus. At this level of undercooling competitive nucleation and growth of either a stable f.c.c. 7 phase or a metastable b.c.c. 6 phase is possible as noted in Fig. 5. With the metastable path two options are available: either retention of the b.c.c, phase or transformation in the solid state to the f.c.c, phase which upon further cooling will form martensite. The stable 7 product will also develop martensite upon further cooling. Clearly the selection of the final transformation structure is a function of the solid state kinetics. With sufficiently rapid cooling in the solid state, which is promoted by a fine sample size,
Martensite
Fig. 5. Schematic diagram of two possible solidification pathways and various solid state reactions for drop-tube-processed samples [14].
retention of the b.c.c, phase is possible. However, in large samples with slow solid state cooling the b.c.c. phase can still develop but must be studied closely to determine the pathway based upon the characteristics of the transformation structure. By examining the microstructural development in samples covering seven orders of magnitude in sample volume, a clear trend has been defined as indicated in Fig. 6 to mark the operation of various solidification and solid state transformation reactions on a microstructure development map. In fact, if the boundary separating the b.c.c. phase retention and decomposition in Fig. 6 is rate limited by the solid state nucleation of the f.c.c, phase from the b.c.c, solidification product, then during cooling at a constant rate T, r,, a
J J.(T)dT = 1
(1)
T06-L
where a is the heterogeneous nucleation site area, J . ( T ) is the nucleation rate and T. is the nucleation temperature. Since J~( T ) is a steep function of T, to a good approximation J. a A T ~F- = 1
(2)
where A T= To - Tn. If it is considered also that along the retention-decomposition boundary J~ is constant, then for powders where a oc d 2 and ~oc d- 2, d 4 A T ~ constant
(3)
J. H. Perepezko
56
/
Solidification reactions in undercooled alloys
This relation is obeyed in the power law dependence illustrated in Fig. 6. The unified description of the overall solidification microstructure development in undercooled Fe-Ni alloys that has been developed indicates that in both bulk samples and powders a b.c.c. phase develops from the melt when sufficient undercooling is available to place the melt below the metastable b.c.c, phase boundary upon nucleation.
coherent guide to the control of processing that will permit an optimization of alloy design and processing in order to achieve a given structure. Recent progress has yielded an improved comprehension of the solidification reaction path, including the influence of competing metastable phase reactions and of thermal history effects on microstructural evolution during RSP, and provides new opportunities for microstructure synthesis.
5. Summary Acknowledgments Rapid solidification has been and continues to remain an important processing approach for materials. One of the main attractions is the flexibility that RSP offers for new approaches to material design and structure synthesis and the fabrication of components. These attributes allow for the production of premium quality material. As the result of numerous studies on the rapid solidification of undercooled liquids, it has become clear that processing conditions play a major role in determining the selection of specific solidification phases and microstructural morphologies. Indeed, the rich variety of solidification microstructures accessible represents an important strength of the approach. At the same time the flexibility in microstructure selection represents a key challenge to implementing a
\ \
o, C ~
o
(..)
k~Y
Imm
\
B ~ i,_
L ' ~ --'?'-o '
O~--
xc~ ~ 00= -(D ~
I00/~
t_~3
o
,6
zb
3b
wt °/oNi Fig. 6. Summary of the structural evolution in the drop tube processing of Fe-Ni alloy processed in an He-2%H 2 environment
[14].
The support of the US Army Research Office (DAAL-03-90-G-0042 and DAAH-04-93-G-0296) and NASA (NAGW-2841) are gratefully acknowledged.
References 1 M. Cohen, B. H. Kear and R. Mehrabian, in R. Mehrabian, B. H. Kear and M. Cohen (eds.), Rapid Solidification Processing: Principles and Technologies H, Claitors, Baton Rouge, LA, 1980, p. I. 2 J.H. Perepezko, Mater. Sci. Eng., 65 (1984) 125. 3 W.J. Boettinger and J. H. Perepezko, in S. K. Das, B. H. Kear and C, M. Adams (eds.), Rapidly Solidified Crystalline Alloys', TMS-AIME, Warrendale, PA, 1985, p. 21. 4 J. H. Perepezko, B. A. Mueller and K. O. Ohsaka, in E. W. Collings and C. C. Koch (eds.), Undercooled Alloy Phases, TMS-AIME, Warrendale, PA, 1987, p. 289. 5 J. H. Perepezko and W, J. Boettinger, Su@zce Alloying by Ion, Electron and Laser Beams, ASM, Metals Park, OH, 1987, p. 51. 6 G. Falkenhagen and W. Hoffman, Z. Metall., 43 (1952) 69. 7 W. E Allen and J. H. Perepezko, Metall. Trans. A, 22 ( 1991 ) 753. 8 H. Kono, J. Phys. Soc. Jpn., 13(1958)1444. 9 Y.J. Kim and J. H. Perepezko, J. Appl. Phys., 71 (1992) 676. 10 Y.J. Kim and J. H. Perepezko, Mater. Sci. Eng. A, 163 (1993) 127. 11 C. G. Levi and R. Mehrabian, Metall. Trans. A, 13 (1982) 221 and 13. 12 T.W. Clyne, Metall. Trans. B, 158 (1984) 369. 13 D. M. Follsteadt, E S. Peercy and J. H. Perepezko, MRS Symp. Proc., 100(1988) 513. 14 D.J. Thoma and J. H. Perepezko, Metall. Trans. A, 23 (1992) 1347.
Materials Science and Engineering, A 179/A 180 (1994) 52-56
Solidification reactions in undercooled alloys J. H. Perepezko Department of Materials Science and Engineering, University of Wisconsin, 1509 UniversityA venue, Madison, W153706 (USA)
Abstract It is now recognized that rapid solidification conditions can be achieved with slow cooling rates provided that the liquid is undercooled substantially prior to nucleation. In fact, many of the novel metastable microstructures produced by rapid solidification require the consideration of an undercooled liquid for analysis. In general, rapid solidification techniques involve either constrained growth in which the solid phase formation is limited by the rate of heat extraction or delayed nucleation of the solid followed by unconstrained growth. With delayed nucleation methods such as the droplet emulsion technique direct measurement of undercooling is available for analysis of metastable phase formation. In fine droplet samples an effective nucleation isolation allows for undercoolings of about 0.3 Tmwith a limit that is usually set by heterogeneous nucleation. Processing variables can be used to control the undercooling and produce a transition in solidification reactions. In this case the use of metastable phase diagrams is important for the analysis of product structures and pathways during solidification and solid state treatments. A key to the understanding of structural evolution is the consideration of competitive nucleation and growth kinetics and thermal history, which can also provide a model for control of solidification reactions as demonstrated in selective alloys.
1. Introduction The most commonly used methods for generating high liquid undercoolings prior to solidification involve a rapid quenching from the melt. Indeed, the main attention in rapid solidification processing (RSP) methods has been focused upon the attainment of high cooling rates in the range 1 0 4 - 1 0 8 ° C s -1. With these methods small liquid volumes are quenched either by contact with the substrate as in melt spinning and directed energy beam processing or by an inert gas as in atomization. However, the most important impact of RSP can be related directly to the solidification structures. The wide range of reported microstructural variations encompass not only highly refined equilibrium phase mixtures but also novel microstructures. While the development of refined structures during RSP can be viewed in terms of the decrease in local solidification time, the understanding of novel microstructures representing non-equilibrium phases requires a consideration of the level undercooling at the onset of nucleation in addition to rapid growth kinetics [1-3]. In fact, the available free energy for non-equilibrium phase formation is directly related to the amount of undercooling. In most work on RSP cooling rate is used as a convenient process variable. However, in solidification studies of undercooled samples it has been demonstrated that undercooling can also be 0921-5093/94/$7.00 SSDI 0921-5093(93)05685-I
treated in a meaningful way as a process variable. In many respects undercooling is a more fundamental parameter than cooling rate. From this experience an examination of solidification reactions in undercooled liquids can provide the necessary basis for the prediction of microstructural transitions that can be applied to guide alloy design in processing to optimize the yield of the preferred microstructure. Some selected highlights can serve to illustrate the variety of pathways for microstructural development that are available in undercooled alloys.
2. Development of undercooling Crystallization of a large continuous bulk liquid metal will be catalyzed by the most potent nucleation site present, so that a means of circumventing the effect of catalytic sites such as oxides or container walls is necessary in order to observe extensive undercooling. Following from this basis, an effective approach to obtain large undercoolings is to disperse the bulk liquid into a collection of fine droplets. This effectively isolates potent nucleation sites into a small fraction of the droplet population as illustrated in Fig. I(A). As a stabilized liquid droplet emulsion is cooled, those drops containing potent nucleants will freeze at low undercoolings, but the majority will not freeze until © 1994 - Elsevier Sequoia. All rights reserved
J. H. Perepezko
/
Solidification reactions in undercooled alloys
53
(a) ATOMIZATION
o®0 0 000(90 0 o°®oo0o ® °0®0 ° oOoq ®
L
(b) TzI
(c)
,
Tz I
~
T m .T. i. I. ~. .-. . . . . . . . . . . . . .~. . ., ,
X=
×
i
\L*i
i
Z~ -
i
° , -.. ~""==.. . .L. . . . l /
WHEEL MELT SPINNING
/T E
SURFACE MELTING
Fig. l. Schematic representation of three common methods of rapid solidification processing involving (A) powder atomization, (B) melt spinning and (C) surface melting of thin layers. In B and C the position of the liquid-solid interface is shown relative to the alloy liquidus for conditions that yield liquid undercooling. Some of the vertical (Z) distances are exaggerated for clarity.
reaching the maximum nucleation undercooling level, which can range from 0.3 Tm to 0.4 Tm [4]. While nucleant isolation accounts for the development of undercooled melts during powder atomization, it is worthwhile to consider how undercooled melts can develop during other common RSP methods. For example, during the rapid freezing of melt steams, as in melt spinning shown in Fig. I(B), undercooled conditions will develop if the liquid exists for some time or distance in contact with the wheel without nucleation. It is evident that a combination of high undercooling and high heat transfer coefficient promotes the most rapid solidification conditions [3]. Alternatively, during surface melting with a rapidly moving or pulsed heat source an undercooled melt zone can develop in front of the resolidification interface if the rate of growth of the unmelted substrate is insufficient to keep up with the rate of heat extraction to the substrate. The development of metastable crystalline phases as well as amorphous products during both melt spinning and surface melting is clear evidence for the existence of an undercooled melt [5].
3. Solidification reactions
One of the first demonstrations of the microstructural benefits of RSP was the development of extensions of primary phase solid solubility [6]. This has often involved eutectic alloy systems where the limit to supersaturation can be analyzed effectively in terms of
the T0 curves to reveal the thermodynamic bounds on supersaturation, including the composition range where glass formation may become kinetically favored
[3]. The experience in solubility extension in peritectic systems is limited compared with eutectics, but recent results on Sn-Sb peritectic alloys (Fig. 2) serve to illustrate some new behavior [7]. With droplet samples and undercoolings in excess of 100°(3 below the primary phase liquidus, supersaturated fl-Sn solutions have been produced for solute levels up to 20 at.% Sb in agreement with previous observations by rapid quenching methods. It is of interest to consider the limit to supersaturation. In Fig. 2 the To curve is also indicated and reveals that the extended liquidus and solidus for the fl solution converge to a congruent maximum in the metastable domain. Following the congruent point, the metastable liquidus and solidus drop to much lower temperatures as a consequence of the melting temperature of Sn in the Sb structure. As a result, the required undercooling below the stable liquidus to reach the To temperature increases rapidly. The behavior is unique to peritectic systems and reveals a new limitation to solubility extension. In addition to extension of solid solubility, the synthesis of metastable intermediate phases has also been an important capability of RSP. Since a metastable intermediate phase will also have a melting temperature below that of the equilibrium intermediate phase, the available free energy for metastable phase formation will be less favorable than that for equilibrium
J. H, Perepezko
54
/
Solidification reactions in undercooled alloys
260 L÷ •
~
vo
Lo o
w250 ~Z
k-
240 el_
B÷zr
W
23o
P / . . . .
0
i
.
,
,
I
,
,
,
,
t
,
5 I0 15 ATOMIC PERCENT Sb
,
.
,
20
l
Fig. 2. Tin-rich portion of the Sn-Sb phase diagram showing metastable extensions of the fl liquidus and solidus as well as T0R-L (chain curve). A metastable melting point maximum is indicated in the vicinity of 15 at.% Sb [71.
phase freezing. Thus the successful development of a metastable phase requires faster nucleation and growth kinetics than for the equilibrium phase. Moreover, the maintenance of the kinetic preference favoring the metastable intermediate phase is limited by the thermal history during recalescence following nucleation at high undercooling. In the Mn-AI system for near-equiatomic alloys the equilibrium phase is an e (h.c.p.) structure, but a metastable r (L10) phase has been reported to form from e during solid state treatments [8]. By containerless processing and splat cooling, the metastable r phase has recently been reported to form directly from the melt [9]. From other thermodynamic measurements the melting temperature of the r phase is approximately 87 K below the melting temperature of e [10]. As indicated by the microstructures in Fig. 3, r phase formation initiates on the cooling substrate, but, during rapid growth, latent heat is released during recalescence. As a consequence, a decreasing interfacial undercooling and interface velocity develop with the progress of solidification to allow for the heterogeneous nucleation of the equilibrium e phase at the moving ~-L interface to yield a final freezing of the central portion of the splat to the e phase.
4. Thermal-history-modified microstructures The sample thermal history during RSP is known to play a central role in the overall solidification kinetics. It has been included in a number of models of RSP [10, 11] but there are some further issues relating to the influence of recalescence behavior on microstructure development. One of the consequences of the recal-
(b)
' l l l l l
Fig. 3. (a) Scanning electron micrograph of the cross-section of splat-quenched Mn0.ssA10.43~Cil.lll7 alloy. (b) Optical micrograph of magnetic powder patterns on the cross-section of splatquenched Mn0ssAl~.~33C0.m7 alloy. Note the confinement of the colloid magnetic particles over the ferromagnetic r phase (dark) and their absence on the non-ferromagnetic e phase (bright)19].
escence history is illustrated in Fig. 3, where a change in product structure and/or microstructure size scale develops owing to the effect of recalescence. During surface melting with directed energy beams another microstructural feature has been identified due to the interplay of solidification kinetics and latent heat release. In an equiatomic Fe-V alloy a high temperature b.c.c, a phase forms from the melt under equilibrium conditions and then transforms to a stable low temperature o phase. If a single-phase substrate of o is subjected to pulse laser annealing, rapid heating bypasses the o ~ a transition so that o melts and places the liquid at an undercooling of 54 K [ 12]. As shown in Fig. 4 when the laser pulse is turned off, the o contains a layer of liquid. During cooling, the a phase nucleates at the free surface and begins to grow, releasing latent heat which causes propagation of the liquid further into
J. H. Perepezko Schematic
of
o->a
in F e V
a~
I
Iil o
a
55
SOLIDIFICATION PATHWAYS
-~
o
a)
a
/ Solidification reactions in undercooled alloys
1
b) a
c)
7"
e
Martensite
0
__d)
e)
Fig. 4. Schematic diagram showing the process by which the latent-heat-assisted transformation of o ~ a is believed to occur
in Fe-V [13].
the o phase and allows for the development of additional surface liquid. Associated with the latent heat release is a slowing of the a phase growth until further development of an intermediate a layer occurs to yield a bimodal microstructure of a phase, i.e. equiaxed grains in contact with o and columnar grains extending outward to the free surface. The net effect is the enhancement of the transformed depth due to the latent heat released by solidification of the a phase beyond that due to the initial laser energy. This phenomenon can be understood by recognizing that the scaling of the melt depth is enhanced by a factor related to the difference in the heat of fusion between the a and the o phase [13]. Another factor of importance in defining the influence of sample thermal history on the final microstructure development during RSP is related to the influence of post-solidification solid state transformations. A clear illustration of this aspect can be developed by considering the microstructure development in undercooled Fe-Ni alloys [14]. Over the range from Fe-10wt.%Ni to Fe-30wt.%Ni large millimeter drops were determined to undercool by about 160 °C during processing in a drop tube apparatus. At this level of undercooling competitive nucleation and growth of either a stable f.c.c. 7 phase or a metastable b.c.c. 6 phase is possible as noted in Fig. 5. With the metastable path two options are available: either retention of the b.c.c, phase or transformation in the solid state to the f.c.c, phase which upon further cooling will form martensite. The stable 7 product will also develop martensite upon further cooling. Clearly the selection of the final transformation structure is a function of the solid state kinetics. With sufficiently rapid cooling in the solid state, which is promoted by a fine sample size,
Martensite
Fig. 5. Schematic diagram of two possible solidification pathways and various solid state reactions for drop-tube-processed samples [14].
retention of the b.c.c, phase is possible. However, in large samples with slow solid state cooling the b.c.c. phase can still develop but must be studied closely to determine the pathway based upon the characteristics of the transformation structure. By examining the microstructural development in samples covering seven orders of magnitude in sample volume, a clear trend has been defined as indicated in Fig. 6 to mark the operation of various solidification and solid state transformation reactions on a microstructure development map. In fact, if the boundary separating the b.c.c. phase retention and decomposition in Fig. 6 is rate limited by the solid state nucleation of the f.c.c, phase from the b.c.c, solidification product, then during cooling at a constant rate T, r,, a
J J.(T)dT = 1
(1)
T06-L
where a is the heterogeneous nucleation site area, J . ( T ) is the nucleation rate and T. is the nucleation temperature. Since J~( T ) is a steep function of T, to a good approximation J. a A T ~F- = 1
(2)
where A T= To - Tn. If it is considered also that along the retention-decomposition boundary J~ is constant, then for powders where a oc d 2 and ~oc d- 2, d 4 A T ~ constant
(3)
J. H. Perepezko
56
/
Solidification reactions in undercooled alloys
This relation is obeyed in the power law dependence illustrated in Fig. 6. The unified description of the overall solidification microstructure development in undercooled Fe-Ni alloys that has been developed indicates that in both bulk samples and powders a b.c.c. phase develops from the melt when sufficient undercooling is available to place the melt below the metastable b.c.c, phase boundary upon nucleation.
coherent guide to the control of processing that will permit an optimization of alloy design and processing in order to achieve a given structure. Recent progress has yielded an improved comprehension of the solidification reaction path, including the influence of competing metastable phase reactions and of thermal history effects on microstructural evolution during RSP, and provides new opportunities for microstructure synthesis.
5. Summary Acknowledgments Rapid solidification has been and continues to remain an important processing approach for materials. One of the main attractions is the flexibility that RSP offers for new approaches to material design and structure synthesis and the fabrication of components. These attributes allow for the production of premium quality material. As the result of numerous studies on the rapid solidification of undercooled liquids, it has become clear that processing conditions play a major role in determining the selection of specific solidification phases and microstructural morphologies. Indeed, the rich variety of solidification microstructures accessible represents an important strength of the approach. At the same time the flexibility in microstructure selection represents a key challenge to implementing a
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[14].
The support of the US Army Research Office (DAAL-03-90-G-0042 and DAAH-04-93-G-0296) and NASA (NAGW-2841) are gratefully acknowledged.
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