- Email: [email protected]
Microstructural control of TiAl–Mo–B alloys by directional solidification
PII:
Acta mater. Vol. 46, No. 18, pp. 6529±6540, 1998 # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain S1359-6454(98)00310-3 1359-6454/98 $19.00 + 0.00
MICROSTRUCTURAL CONTROL OF TiAl±Mo±B ALLOYS BY DIRECTIONAL SOLIDIFICATION D. R. JOHNSON, K. CHIHARA, H. INUI and M. YAMAGUCHI{ Department of Materials Science and Engineering, Kyoto University, Sakyo-ku, Kyoto 606-01, Japan (Received 6 July 1998; accepted 17 August 1998) AbstractÐDirectional solidi®cation experiments were performed to see if the g/a2 lamellar boundaries could be aligned parallel to the growth direction for an ingot containing columnar lamellar grains. From microstructural analysis of arc-melted ingots, a projection of the partial liquidus surface for the Ti±Al±Mo system was constructed and this was used with an analytical description of the phase diagram to calculate the solidi®cation paths. Following this analysis, Ti±Al±Mo(+B) ingots were directionally solidi®ed at growth rates between 40 and 100 mm/h. By controlling the solidi®cation path, a possible processing window where a reasonably well aligned g/a2 microstructure can be produced was identi®ed. As-processed microstructures exhibited a compressive yield strength greater than 500 MPa at 8008C and 2% ductility at room temperature. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.
1. INTRODUCTION
Gamma titanium aluminides are candidate materials for lightweight structural applications at temperatures near 8008C. Of the numerous microstructures that can be developed in TiAlbased alloys [1, 2], the fully lamellar microstructure consisting of TiAl (g) and Ti3Al (a2) possesses the best combination of room temperature toughness and high temperature strength [1±5]. However, as detailed by studies using polysynthetically twinned (PST) crystals, where the whole ingot consists of a single lamellar grain, mechanical properties are found to be extremely anisotropic with respect to the lamellar orientation [5±7]. For example, the optimum combination of strength, toughness, and ductility can be obtained if the lamellar microstructure is aligned parallel to the test axis [6, 7]. However, if the lamellar boundaries are perpendicular to the tensile axis, the material is extremely brittle [6]. Therefore, if the lamellar boundaries can be aligned parallel to the ingot's axis by an appropriate processing technique, then the resulting material would have a good combination of strength and toughness. One way to produce such materials is by working TiAl alloys within the a phase temperature range [8, 9] to produce the desired texture. However, this paper will focus on the problem of forming a favorable texture directly upon solidi®cation. For this case, the target microstructure is one consisting of columnar grains that have transformed to the fully lamellar microstructure upon cooling with all the g/a2 boundaries parallel to the growth {To whom all correspondence should be addressed.
direction. In addition, each columnar lamellar grain should be rotated about its longitudinal axis to provide reasonable o-axis fracture toughness as shown in Fig. 1. For clarity, only a few columnar grains are shown in Fig. 1. However, for increased ductility [4] and a greater dynamic fracture toughness [10], the columnar grain size should be as small as possible. Unfortunately, simple casting operations often produce the worst case scenario with the lamellar boundaries all perpendicular to the heat ¯ow direction [3, 11]. To produce materials with a balance of properties, thermomechanical-processing routes are then required to remove the solidi®cation texture and to re®ne the lamellar colony size [1, 3, 11]. This unfavorable solidi®cation texture is a result of the preferred growth of the a phase parallel to the [0001] direction. Upon cooling, the lamellar microstructure would then form from the parent a phase following the Blackburn orientation relationship [12] of 0001a ==f111gg and h1120ia ==h110ig resulting in the observed lamellar orientation. As was recently shown, the orientation of the a phase can be controlled during solidi®cation by using a seed from the TiAl±Si system [13]. However, for this case, the resulting ingots consist of a single lamellar grain whereas the target microstructure should consist of numerous columnar grains. Another approach is to change the solidi®cation pathway such that the bulk of the ingot solidi®es as the Ti-based b.c.c. solid solution (b phase). This can be accomplished by lowering the Al content below 45 at.% [2] for binary Ti±Al alloys. Typically such alloys form an equiaxed lamellar microstructure upon casting. However, if directionally solidi-
6529
6530
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION Table 1. Compositions and processing conditions of directionally solidi®ed ingots with the results of EDS analysis enclosed in parentheses Ingot number
Fig. 1. Target microstructure for a directionally solidi®ed TiAl-based alloy.
®ed, the resulting ingot can consist of large columnar grains with the lamellar boundaries oriented between 0 and 458 to the growth direction [14]. If the b phase dendrites have a preferred growth direction parallel to h001i, then the orientation of the lamellar boundaries can be explained by the Burgers orientation relationship [15] between the a and b phases. For this case, a would nucleate from the b phase with the (0001)a plane being inclined by 458 or parallel to the growth direction [14]. The lamellar microstructure would then form upon further cooling as previously mentioned resulting in the observed lamellar orientation. Unfortunately, directionally solidi®ed ingots with a composition near 45 at.% Al tend to be brittle [2]. This can be attributed to the larger volume fraction of the a2 phase associated with the lower Al content, the very large columnar grain size (only a few columnar grains are found in a 12 mm diameter ingot [14]), and the poor alignment of the microstructure since the majority of the lamellar grains tend to be oriented 458 to the growth direction (as would be expected from Burgers orientation relationship). Therefore, to improve the mechanical properties, the lamellar grain size must be signi®cantly reduced and the lamellar orientation needs to be well aligned with the growth direction. To explore this possibility, an investigation of the solidi®cation behavior of TiAl±Mo and TiAl±Mo±B ingots was undertaken. Additions of Mo were used to promote b solidi®cation [16±18] as well as to provide some solid solution strengthening [18, 19]. In addition, boron additions were used to re®ne the solidi®cation microstructure [3, 20, 21]. With these
Fig. 2. Schematic drawing of the directional solidi®cation furnace.
1 2 3 4 5 6
Composition (at.%)
Growth rate (mm/h)
Ti±45.5Al±1.5Mo (Ti±43.4Al±3.8Mo)$ Ti±46.5Al±1.0Mo (Ti±44.8Al±2.1Mo)$ Ti±46.5Al±1.5Mo±0.5B (Ti±45.5Al±1.6Mo±B)% Ti±46.5Al±1.5Mo±0.6B (Ti±47.5Al±1.5Mo±B) Ti±46.5Al±1.5Mo±0.7B (Ti±46.9Al±1.5Mo±B) Ti±46.5Al±1.5Mo±0.7B (Ti±47.0Al±1.4Mo±B)
40 40 100 100 100 100
$Mo pickup from melting on a Mo hearth, all other ingots were directly melted in CaO crucibles. %Boron content was not measured for these ingots.
alloying additions, the purpose of this investigation was to determine the conditions that are needed to develop a well aligned solidi®cation texture in a cast or directionally solidi®ed ingot. 2. EXPERIMENTAL PROCEDURES
Master ingots for directional solidi®cation were prepared by arc-melting in puri®ed argon at 1 atm. Directional solidi®cation was performed in a laboratory scale ram-type furnace using 16 mm diameter CaO crucibles as schematically drawn in Fig. 2. For the ®rst few ingots processed, the initial charge was melted in a Mo hearth and then dropped into a CaO crucible. However, during this procedure, there was signi®cant Mo pickup from the hearth. Therefore, to maintain the target composition, all other ingots were melted directly in the crucible and then directionally solidi®ed at velocities between 40 and 100 mm/h. The compositions of all six ingots processed are listed in Table 1 and these will be referred to by number in the proceeding text. A partial liquidus surface for the Ti±Al±Mo system was determined by comparing the microstructures of arc-melted ingots. In addition, the primary dendritic phase was also identi®ed in TiAl±Mo alloys containing small additions of boron (0.6 and 0.8 at.%). Small arc-melted buttons (about 15 g each) were made from high purity elements (99.99 at.%). Each arc-melted ingot was melted and turned at least ®ve times to promote homogeneity. A scanning electron microscope (SEM) equipped with an energy-dispersive spectrometer (EDS) as well as an optical microscope (OM) were all used to characterize the microstructures of the arc-melted and directionally solidi®ed ingots. Backscattered electron (BSE) imaging was used to help identify the phases present and to locate the dendrite cores. Specimens for SEM observations were electropolished in a solution of 5% HClO4±35% N-butanol± 60% methanol by volume at ÿ658C. Specimens viewed by optical microscopy were polished and etched using a solution of 5% HF±5% HNO3±90% H2O by volume. Room temperature tensile tests were performed on selected alloys using ¯at specimens 20 mm in
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
Fig. 3. Projected view of the partial liquidus surface for the Ti±Al±Mo system.
length with a gage length of 5 mm 2 mm 1 mm. Specimens were electropolished and tested in vacuum of 5 10ÿ5 Torr at a strain rate of 2 10ÿ4/s. In addition, compression tests between 25 and 11008C were also performed on specimens with dimensions of 2.5 mm 2.5 mm 5 mm. These specimens were also electropolished and tested in vacuum of 1 10ÿ4 Torr. For all tests, the tensile or compression axis was parallel to the growth direction. 3. RESULTS
3.1. Microstructures from arc-melted ingots The solidi®cation sequence was initially determined by examining the dendrite morphologies found in arc-melted buttons containing either 2 or 5 at.% Mo. From these observations, a partial liquidus surface was constructed as shown in Fig. 3. The basic result is that Mo additions shift both the primary a and b regions towards greater Al contents. In addition, the average composition within the prior dendritic regions was found to be enriched in Mo when compared to the original alloy composition, Table 2, suggesting that Mo additions increase the melting temperature of both the a and g phases as noted by the direction of the arrows in Fig. 3. 3.1.1. Ti±Al±5Mo series. Typical microstructures for alloys containing 5 at.% Mo are shown in Fig. 4. For this series of alloys, b was found to be the primary solidi®cation phase for alloys with an aluminum content less than about 52 at.% as Table 2. Average composition of dendrites as measured by EDS analysis Alloy composition (at.%)
b dendrites Ti±48Al±2Mo Ti±48Al±5Mo a dendrites Ti±53Al±2Mo Ti±55Al±5Mo
Average composition Ti
Al
Mo
51.2 47.9
46.5 46.5
2.3 5.6
47.3 41.7
50.4 52.6
2.3 5.7
6531
shown in Fig. 4(a)±(d). The phase with bright contrast within the prior dendritic regions was found to be rich in Mo and lean in Al when compared to the ingot composition. For example for the Ti±48Al± 5Mo alloy (all compositions are given in atomic percentages), a composition of Ti±47.7Al±9.8Mo was measured for this phase, which can be inferred from the literature as being the B2 phase [16, 17, 22± 24]. The remaining matrix was found to consist of the g/a2 lamellar microstructure. The microstructure for the Ti±46Al±5Mo button, Fig. 4(a) and (b), was essentially equiaxed and an obvious aluminum-rich interdendritic region was not observed suggesting that the complete ingot solidi®ed as the b phase. The prior b grains are also found to be subdivided into a number of colonies due to the formation of WidmanstaÈtten plates of a [16]. A similar microstructure for a Ti±44Al±2Mo alloy has been characterized by Li and Loretto both in terms of microstructure and fracture behavior [23]. In contrast, remnants of the prior dendritic morphology are clearly visible in all the alloys containing 5 at.% Mo with compositions between 47 and 52 at.% Al, and typical microstructures are shown in Fig. 4(c) and (d). The primary phase is identi®ed as b by noting that the dendrite arms are orthogonal to one another. Within the interdendritic region, the microstructure is a result of a peritectic cascade consisting of the L + b 4 a and the L + a 4 g reactions. The dark regions surrounding the dendrites in Fig. 4(c) are found to consist of the lamellar microstructure, Fig. 4(d), indicating it was originally the peritectic a phase. Similarly, regions of peritectic g are found within the prior peritectic a regions [as indicated by the arrow in Fig. 4(d)], completing the cascade. Increasing the Al content simply increases the amount of peritectic g while decreasing the amount of the other phases. For higher aluminum contents in the range of 53±57 at.%, a is identi®ed as the primary phase. The hexagonal symmetry of the dendrite arms can be seen for typical microstructures in Fig. 4(e) and (f). The prior dendritic regions for the a phase are also found to be rich in Mo and lean in Al as was the case for the b phase, Table 2. In addition, a bright phase can be seen surrounding the a phase in Fig. 4(e) and (f), which may be the result of b (or B2) precipitation at intermediate temperatures. Once the aluminum content was increased above 57 at.% Al, a drastic change in the dendrite morphology was observed, Fig. 4(g) and (h), and this suggests a transition to the solidi®cation of primary g. A bright phase with a needle-like morphology, having a composition roughly of Ti±55.4Al±6.9Mo, is observable in the g phase. Since, a clear cubic dendrite morphology cannot be observed, these needles are most likely the remnants of a very small volume fraction of a phase dendrites. A possible side arm oriented 608 to the dendrite spine is indi-
6532
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
Fig. 4. SEM backscattered electron images of microstructures taken from arc-melted Ti±xAl±5Mo ingots showing (a) and (b) the Ti±46Al±5Mo alloy, (c) and (d) prior b dendrites of Ti±48Al±5Mo alloy with interdendritic g marked by an arrow, (e) and (f) prior a dendrites of Ti±56Al±5Mo alloy, and (g) and (h) Ti±58Al±5Mo alloy with an arrow marking a dendrite side arm.
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
6533
Fig. 5. SEM backscattered electron images of microstructures taken from arc-melted Ti±xAl±2Mo ingots showing prior b dendrites for (a) the Ti±46Al±2Mo alloy and (b) the Ti±50Al±2Mo alloy.
cated by the arrow in Fig. 4(h) and provides further evidence for this assumption. These prior a dendrites would then be sites for the formation of the b or B2 phase at intermediate temperatures. Also an eutectic-like microstructure is visible in Fig. 4(h) and consists of a ®brous bright phase (Ti±57.9Al± 7.7Mo) within a g matrix (Ti±58.8Al±4.4Mo). From their compositions, it would appear that the eutectic phases are b and g. From the partial isothermal sections proposed by Das et al. [22] and Singh and Banerjee [17], both the Ti±56Al±5Mo and Ti±58Al±5Mo alloys should consist of the g and b (or B2) phases once equilibrated below 014008C. Hence, the bright contrast in Fig. 4(e)±(h) should result from the formation of the B2 phase. Furthermore, the eutectic microstructure consisting of the g and b phases in Fig. 4(g) and (h), provides evidence for the L + a 4 b + g invariant reaction as suggested by Singh and Banerjee [16]. From the above arguments and by using the compositions measured from the eutectic g and b phases, the L + a 4 b + g reaction is sketched in the projected view of the liquidus surface shown in Fig. 3. 3.1.2. Ti±Al±2Mo series. The solidi®cation behavior of alloys containing 2 at.% Mo was essentially the same as those containing 5 at.% Mo. The composition range where a distinct cubic (b) dendritic morphology was visible was found to be between 46 and 50 at.% Al as shown in Fig. 5(a) and (b). The main dierence between these ®gures is the relative amounts of the phases that formed from the liquid. In Fig. 5(a), the prior b dendrites contain veins of the B2 phase. The dark regions surrounding the dendrites were found to consist of the g/a2 lamellar microstructure indicating a prior peritectic a phase. In contrast, the dark regions in Fig. 5(b) consist of the g phase. For this microstructure, the amount of B2 phase is much less due to a smaller volume fraction of primary b dendrites. Surrounding these b dendrites are lamellar regions that formed from the peritectic a phase and can be identi®ed by the lack
of the B2 phase. Hence, during solidi®cation, the Ti±50Al±2Mo alloy cascaded through two peritectic reactions while for the Ti±46Al±2Mo alloy only the L + b 4 a peritectic reaction occurred. 3.2. Solidi®cation paths from thermodynamic calculations In order to determine the solidi®cation paths and to determine the degree of microsegregation, an analytical description of the phase diagram would be useful. For example, if the tie lines of the ternary system are known then a Scheil type analysis can be used to determine the solidi®cation path. For the Ti±Al±Mo system, a simple approach was taken to estimate the Gibbs energies of the ternary phases (liquid, b, and a) from the Gibbs energies of the binary phases [25±27]. For this approach, Toop's method [28] was used in conjunction with the experimentally determined liquidus surface. The form of Toop's equation used along with the analytical description of the binary phases from the literature are listed in Table 3. Since the melting temperature of the a phase was found to increase with Mo additions, the maximum liquidus temperature should locate within the ternary system. For this case, Toop's equation alone is insucient to calculate the ternary Gibbs energies. Hence, a ternary term was added to the equation using a polynomial of the type proposed by Pelton and Bale [29]. The coecients of the polynomial were estimated by ®rst varying the interaction parameter of the a phase in the binary Mo±Al system and noting the response in the ternary system. The coecients of the ternary term were then adjusted to reproduce this response with the a interaction parameter set to zero in the Mo±Al system. This was done by an iterative approach and the coecients were not optimized. Once a satisfactory ®t was obtained, the chemical potentials were numerically calculated and from these, the equilibrium concentrations were found. A partial isothermal section at 1773 K that was calculated is shown in
6534
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION Table 3. Analytical description for the excess free energy of the liquid, b and a phases in the Ti±Al±Mo system (units are in J/mol and J/mol K). For the respective elements Ti, Al, or Mo the standard states are: a refers to f.c.c.; b refers to b.c.c.; and liquid refers to liquid Lattice stabilities of the pure elements [25±27] ÿ 8GLTi ÿ14146:0 7:2880 T 8Gb:c:c: Ti ÿ 8GLTi ÿ18318:0 10:8984 T 8Gh:c:p: Ti f:c:c: ÿ 8GLTi ÿ12316:0 10:7984 T 8GTi ÿ 8GLAl ÿ628:0 6:6598 T 8Gb:c:c: Al ÿ 8GLAl ÿ5151:4 9:6178 T 8Gh:c:p: Al f:c:c: ÿ 8GLAl ÿ10711:0 11:4728 T 8GAl L 8Gb:c:c: Mo ÿ 8GMo ÿ24267:0 8:368 T L 8Gh:c:p: Mo ÿ 8GMo ÿ15899:0 8:368 T f:c:c: ÿ 8GLMo ÿ13807:0 8:996 T 8GMo Ti±Mo binary [26] Liquid: DGxs 14943:0 XTi XMo b phase: DGxs 15657:0 5238:0 1 ÿ 2 XMo XTi XMo a phase: DGxs 36647:0 XTi XMo Mo±Al binary [27] Liquid: DGxs ÿ32024:0 XMo XAl b phase: DGxs ÿXMo XAl 30543:0 XMo 16736:0 XAl ÿ 12:552 XMo XAl XAl T Ti±Al binary [25] Liquid: DGxs ÿ112570:2 41:11378 T ÿ 7950:8 XTi ÿ XAl XTi XAl b phase: DGxs ÿ129396:7 40:06310 T XTi XAl a phase: DGxs ÿ123788:5 33:20902 T ÿ16034:9 12:18272 T XTi ÿ XAl XTi XAl Ti±Al±Mo Form of Toop's equation used with additional ternary term xs xs DGxs XMo Gxs TiMo XAl GTiAl = 1 ÿ XTi 1 ÿ XTi 1 ÿ XTi GMoAl P XTi XAl XMo Liquid: P = 0 b phase: P = 0 a phase: P ÿ76500:0 ÿ0:5 ÿ 0:9 XTi 5:9 XAl
Fig. 6. In this ®gure, the liquidus composition at either of the three phase tie triangles matches quite well with what would be expected from the experimentally determined liquidus projection, Fig. 3. While the above analysis is quite simple, it is useful for predicting the range of possible microstructures that can form. The limiting cases range from equilibrium to the case were diusion in the solid is neglected. For the latter, the solidi®cation paths were found using an approach that couples a thermodynamic model of the system with the assumptions of the Scheil model as described by Brody and Flemings [30] and by Chen et al. [31]. The basic approach is that the equations describing the mass balance during solidi®cation are solved numerically so that the partition coecients (or tie lines) can be calculated from the thermodynamic model at each iteration. From this Scheil type analysis, the amount of interdendritic g was calculated for the arc-melted buttons and compared to those measured from metallographic analysis. These results are shown in Fig. 7. The predicted values compare well with
those measured from the arc-melted ingots containing 2 and 5 at.% Mo. In addition, for alloys containing 2 at.% Mo, the calculated relative amounts of each phase that solidi®ed as a function of ingot composition are shown in Fig. 8.
Fig. 6. Estimated partial isothermal section at 1773 K for the Ti±Al±Mo system.
Fig. 7. Calculated and measured amounts of interdendritic g for Ti±Al±Mo alloys.
3.3. Directional solidi®cation of TiAl±Mo alloys For the ®rst set of directionally solidi®ed ingots, the initial charge was melted in a Mo hearth and then dropped into a CaO crucible. Unfortunately, during this process some of the Mo hearth was melted and the resulting ingots were found to have a much greater Mo content than the target composition as listed in Table 1. However, some interesting microstructures developed from these ingots. As the molten charge was dropped on a water cooled chill plate, these ingots consist of a chill region and then a directionally solidi®ed region.
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
Fig. 8. Relative amounts of each phase solidi®ed in Ti± Al±Mo alloys as calculated by the Scheil model.
Microstructures for a typical case are shown in Fig. 9 for a Ti±43.4Al±3.8Mo ingot (ingot 2). Macroscopically, columnar grains consisting of the g/a2 lamellar microstructure are found for the complete length of the ingot [Fig. 9(a)]. The microstructure for the directionally solidi®ed section consists of b grains that have a columnar shape as shown in Fig. 9(b). However, the microstructure appears equiaxed within the chill zone as shown in Fig. 9(c). The presence of the B2 phase within the cored regions suggests a prior structure of equiaxed b grains. From the previous calculations of the solidi®cation paths, the minimum amount of b that should form upon solidi®cation is nearly 80%. Thus, the columnar lamellar grains are probably the result of directional growth of a via the solid state transformation of b to a.
6535
Although this ingot solidi®ed through the b phase, the microstructure did not approach the target microstructure as speci®ed in Fig. 1 since the grain size was large and the lamellar microstructure was not aligned with the growth direction. Therefore, solidifying through the b phase alone is not sucient to produce the desired microstructure and a dierent approach must be taken. Since the lamellar microstructure is formed from the a phase, better microstructural control may result by changing the morphology of the secondary a phase that forms from the liquid. To investigate this problem, boron was added to try to change the solidi®cation behavior of the TiAl±Mo ingots. From the microstructure of arcmelt ingots, the addition of boron does not signi®cantly change the compositional region where primary b and a form as shown in Fig. 10. Thus, the volume fraction of the phases calculated earlier (Fig. 8) should roughly be the same as that for alloys containing small boron additions. Using this assumption, a baseline composition of Ti±46.5Al± 1.5Mo±0.6B was chosen for directional solidi®cation experiments since about half the ingot should solidify as b and half as a. 3.4. Directional solidi®cation of TiAl±Mo±B alloys The addition of boron greatly reduced the lamellar grain size as shown in Fig. 11(a) for ingot 3. The lamellar orientation measured along the length of this ingot was mixed and lamellar grains could be found at orientations from parallel to nearly per-
Fig. 9. Microstructures from a directionally solidi®ed Ti±43.4Al±3.8Mo ingot showing (a) the macroscopic view of the longitudinal section, and SEM backscattered electron images from (b) the directionally solidi®ed portion of the ingot, and (c) the chill zone.
6536
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
Fig. 10. Projected view of the Ti±Al±B liquidus surface as determined by Hyman et al. [20] plotted together with the results from alloys containing 1.5 at.% Mo (these compositions are denoted by the open circles and squares). The solidi®cation path for Ti±46.5Al±1.5Mo±0.6B (denoted by the small closed circle) is also outlined by the numbers ``1'' to ``3''.
pendicular to the growth direction. However, with a slight increase in boron content, portions of ingots 4±6 had a well aligned lamellar microstructure. The lamellar grain structures along the length of these ingots were similar and that for ingot 4 is shown in Fig. 11(b). The ®rst half of this ingot consists of columnar lamellar grains with mixed lamellar orientations, and this section is labeled ``M'' in Fig. 11(b). However, within the second half, a columnar to equiaxed transition occurred (labeled as ``E'') that was followed again by columnar grains but now with a much smaller grain size. These short columnar grains (the region labeled as ``A'') had a very well aligned lamellar microstructure as shown in Fig. 11(c). Note that a number of these grains are oriented such that the lamellar plane is
Fig. 11. Microstructures of directionally solidi®ed TiAl±Mo±B ingots showing (a) the longitudinal section of a Ti±46.5Al±1.5Mo±0.5B ingot, (b) the longitudinal section of a Ti±46.5Al±1.5Mo±0.6B ingot with aligned ``A'', equiaxed ``E'', and mixed ``M'' lamellar regions marked, and (c) the enlarged view of the microstructure of region ``A'' with arrows marking the lamellar colonies rotated such that (111)g is almost parallel with the plane.
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
6537
the unetched microstructure of ingot 4 is shown in Fig. 12. The bright contrast consists of the boride phase and associated pullout during polishing. The rest of the microstructure consists of b phase dendrites with the main dierence between each section being the dendrite spacing which is somewhat ®ner within the aligned section of the ingot. Since no major changes are found in Fig. 12, the change in microstructure as observed in Fig. 11(b) must be due to changes occurring in the interdendritic a phase and this will be considered later in the text. 3.5. Compression and tensile test results Tensile specimens were machined from the aligned region in ingots 4±6 and compression specimens were also taken from ingot 6 within the same region. From Fig. 13, a compressive yield strength greater than 500 MPa is maintained up to 8008C after which it decreases to 400 MPa at 10008C. Any further increase in temperature causes a very rapid decrease in yield strength. The results from the tensile tests are listed in Table 4 and these basically re¯ect the dierences in the lamellar orientation of the test specimens. Plastic strains up to 2% were measured from specimens taken from ingots 4±6 where the lamellar microstructure was aligned with the growth direction. However, specimens taken from the ®rst half of these ingots displayed little ductility when tested. A dierence in the lamellar orientation is found when observing the fracture surfaces as shown in Fig. 14. Regions of fast fracture indicated by large cleavage surfaces are found in all the tested specimens machined from the ®rst half of ingots 4 and 5 with a typical fracture surface shown in Fig. 14(a). In contrast, the fracture surfaces associated with specimens machined from the portion of the ingot with the aligned lamellar microstructure displayed a composite or splitting type fracture surface without the regions of fast fracture, Fig. 14(b).
Fig. 12. Photographs taken from the unetched ingot shown in Fig. 14 for (a) the mixed orientation, (b) the equiaxed, and (c) aligned lamellar regions.
roughly parallel to the plane of view as indicated by the coarse spacing of the lamellar trace for regions marked in Fig. 11(c). Thus, the microstructure within the aligned portion of the ingot approaches that of the target microstructure with the lamellar grains rotated with respect to each other along the ingot's axis. Somewhat surprising is that when ingots 4±6 are polished and viewed unetched by eye, little change in microstructure is observable along the entire length of the ingots. What is seen is well-de®ned columnar texture outlined by the interdendritic boride phase. For example, the macroscopic view of
Fig. 13. Compressive yield strength of a Ti±46.5Al± 1.5Mo±0.7B alloy processed at 100 mm/h as a function of temperature.
6538
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
Table 4. Tensile properties of directionally solidi®ed Ti±46.5Al± 1.5Mo±(0.6±0.7)B (at.%) alloys with a well aligned lamellar microstructure and with a mixed lamellar orientation Specimen
Yield stress (MPa)
Elongation (%)
Ingot 4: Ti±46.5Al±1.5Mo±0.6B Aligned region 440 370 430 Mixed region Ð Ð Ingot 5: Ti±46.5Al±1.5Mo±0.7B Aligned region 440 400 400 Mixed region 360 280 Ingot 6: Ti±46.5Al±1.5Mo±0.7B Aligned region 500 500
2.9 1.2 1.2 0 0 2.5 0.4 1.2 0.3 0.5 0.4 0.5
Growth rate = 100 mm/h.Tested in vacuum of 5 10ÿ5 Torr.
4. DISCUSSION
The microstructures of the directionally solidi®ed TiAl±Mo±B alloys only approached the target microstructure for a very short portion of the ingots [Fig. 11(c)]. Since the portion with the aligned lamellar microstructure was found after a columnar to equiaxed transition, the qualitative analysis described by Hunt [32] on this matter may be helpful. Hunt's analysis considers the growth of equiaxed grains ahead of a columnar front during directional solidi®cation and the ®nal results of the analysis are that fully equiaxed growth occurs when 3 3 G<0:617N 1=3 0 f1 ÿ DTN = DTc gDTc
where G is the average thermal gradient, N0 the total number of heterogeneous nucleation sites per volume, DTN the undercooling at the heterogeneous nucleation temperature, and DTc the undercooling at the solid±liquid interface, which is proportional to the square root of the growth velocity [32]. However, for the directionally solidi®ed TiAl± Mo±B alloys, little change was found in the b dendrite morphology as columnar b growth was found for the complete ingot (Fig. 12). Hence, the microstructural changes found in these alloys [Fig. 11(b)] must result from a change in the secondary a morphology. Therefore, both the b and a phases must be considered when applying the above model to the columnar to equiaxed transition and this is schematically shown in Fig. 15. The columnar to equiaxed transitions depend upon the eciency of the nucleating agent, DTN, and by the number of heterogeneous nucleation sites, N0. By increasing the nucleation eciency (decreasing DTN) the curves in Fig. 15 are shifted down towards lower velocities. Similarly, by increasing N0, the curves are shifted right towards higher values of G. Under the same processing conditions (G and V), if the a and b phases are considered separately and if the set of values for DTN and N0 are similar, then the columnar to equiaxed
Fig. 14. Fracture surfaces of tensile specimens machined from (a) the mixed orientation, and (b) the aligned lamellar regions of ingot number 4.
transition will always occur ®rst for the b phase for the same processing conditions (e.g. G and V). This is due to the smaller freezing range of a as given by the phase diagram which, when compared to b, results in smaller values of DTc [33]. Furthermore, since the growth of the peritectic a phase occurs behind the primary b dendrites, the solute undercooling for the a phase should be even less since the b dendrite spacing will adjust itself to minimize the constitutional supercooling within the interdendritic liquid. For the TiAl±Mo ingots, boron was added to change the above behavior by providing an active catalyst for a nucleation. The solidi®cation path of the Ti±46.5Al±1.5Mo±0.6B ingots can be described by using Fig. 10. For this composition (labeled as
Fig. 15. Schematic representation of the columnar/ equiaxed transition for the a and b phases in a TiAl±Mo± B alloy. For the processing conditions marked by the ®lled circle, columnar growth of b would be expected while, for the a phase, equiaxed growth would be expected.
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
``1'' in Fig. 10), growth of b would ®rst occur causing the interdendritic liquid to become enriched in boron. Boride particles would then nucleate once the L 4 TiB2 + b monovariant line is met (composition ``2'' in Fig. 10). However, the a phase would not nucleate until the liquid composition reaches the L + B 4 a + TiB2 invariant reaction at composition ``3'' as labeled in Fig. 10. The important point is that the boride particles can nucleate within the interdendritic liquid before the solidi®cation of the a phase. Thus, if the boride particles are an active catalyst for a nucleation, then the columnar to equiaxed transition curve for the a phase can be shifted towards higher G values and possibly past that for the b phase as is schematically drawn in Fig. 15. Hence, it may be possible to have equiaxed a growth even though the growth of b is columnar, and Figs 11 and 12 provide evidence that this can indeed be the case. The growth morphology for this case is sketched in Fig. 16. If a grains are continuously nucleated at the boride particles, then growth can occur only over short distances before impingement upon more recently nucleated grains occurs. During this period, there will be competitive growth between the dierently oriented a grains. Consequently, growth parallel to h1010i may be possible. Therefore, the aligned lamellar microstructure shown in Fig. 11(c) is suggested to occur from competitive growth of h1010i oriented a grains behind a columnar b dendritic front and from the nucleation of a from prior [100] oriented b dendrites following the Burgers orientation relationship. Lastly, only a small section of the directionally solidi®ed Ti±Al±Mo±B ingots had an aligned g/a2 lamellar microstructure as shown in Fig. 11(b). From Fig. 15, the columnar to equiaxed transition is very sensitive to changes in the thermal gradient at higher growth velocities, and thus the processing window where columnar growth of b and equiaxed growth of a can occur together is very small. Therefore, the variation of the microstructure along
Fig. 16. Schematic view of columnar growth of b and columnar plus equiaxed growth of a.
6539
the length of the ingot may result from the change in the thermal gradient during processing. For example, initially there is a high thermal gradient due to the water cooled chill plate used, resulting in columnar growth of both the a and b phases [region ``M'' in Fig. 11(b)]. However, away from the chill plate, the thermal gradient decreases and equiaxed growth of a occurs [regions ``A'' and ``E'' in Fig. 11(b)]. The question of whether steady state conditions can be reached and maintained within this narrow processing window must be answered. However, the present work shows that such a possibility may exist and work is now under way to produce larger size ingots. 5. CONCLUSIONS
The solidi®cation paths for Ti±Al±Mo alloys were calculated from an experimentally determined liquidus projection and from an analytical description of the phase diagram. By controlling the solidi®cation path, a possible processing window where a reasonably well aligned g/a2 microstructure can be produced was identi®ed. This is accomplished by alloying with boron such that boride particles will nucleate behind the primary b dendrites before the nucleation of the a phase, resulting in columnar growth of the b phase followed by equiaxed growth of the a phase. Such microstructures exhibited a compressive yield strength of more than 500 MPa at 8008C and 2% ductility at room temperature. AcknowledgementsÐThis work was supported by ``Research for the Future'' Program (JSPS-RFTF 96R12301), The Japan Society for the Promotion of Science, and in part by a research grant from the Research and Development Institute of Metals and Composites for Future Industries. Funding from the Ministry of Education, Science and Culture, Japan (grant No. 09750772) is also gratefully acknowledged. REFERENCES 1. Kim, Y.-W., JOM, 1994, 46, 30. 2. Huang, S. C., in Structural Intermetallics, ed. R. Darolia, J. J. Lewandowski, C. T. Liu, P. L. Martin, D. B. Miracle and M. V. Nathal. TMS, Warrendale, Pennsylvania, 1993, p. 299. 3. Wagner, R., Appel, F., Dogan, B., Ennis, P. J., Lorenz, U., Mullauer, J., Nicolai, H. P., Quadakkers, W., Singheiser, L., Smarsly, W., Vaidya, W. and Wurzwallner, K., in Gamma Titanium Aluminides, ed. Y.-W. Kim, R. Wagner and M. Yamaguchi. TMS, Warrendale, Pennsylvania, 1995, p. 387. 4. Liu, C. T., Schneibel, J. H., Maziasz, P. J., Wright, J. L. and Easton, D. S., Intermetallics, 1996, 4, 429. 5. Yamaguchi, M. and Inui, H., in Structural Intermetallics, ed. R. Darolia, J. J. Lewandowski, C. T. Liu, P. L. Martin, D. B. Miracle and M. V. Nathal. TMS, Warrendale, Pennsylvania, 1993, p. 127. 6. Inui, H., Kishida, K., Misaki, M., Kobayashi, M., Shirai, Y. and Yamaguchi, M., Phil. Mag. A, 1995, 72, 1609. 7. Yakoshima, S. and Yamaguchi, M., Acta mater., 1996, 44, 873.
6540
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
8. Kim, Y.-W. and Dimiduk, D. M., in Structural Intermetallics 1997, ed. M. V. Nathal et al. TMS, Warrendale, Pennsylvania, 1997, p. 531. 9. Jin, Z., Gray, G. T. III and Kim, Y.-W., Mater. Sci. Engng, 1997, A239-240, 729. 10. Enoki, M., Fujikawa, S. and Kishi, T., J. Japan Inst. Metals, 1994, 58, 424. 11. Mecking, H. and Hartig, Ch., in Gamma Titanium Aluminides, ed. Y.-W. Kim, R. Wagner and M. Yamaguchi. TMS, Warrendale, Pennsylvania, 1995, p. 525. 12. Blackburn, M. J., in The Science, Technology, and Application of Titanium, ed. R. I. Jaee and N. E. Promisel. Pergamon Press, Oxford, 1970, p. 663. 13. Johnson, D. R., Masuda, Y., Inui, H. and Yamaguchi, M., Acta metall. mater., 1997, 45, 2523. 14. Johnson, D. R., Masuda, Y., Shimada, Y., Inui, H. and Yamaguchi, M., in Structural Intermetallics 1997, ed. M. V. Nathal et al. TMS, Warrendale, Pennsylvania, 1997, p. 287. 15. Burgers, W. G., Physica, 1934, 1, 561. 16. Singh, A. K. and Banerjee, D., Metall. Mater. Trans. A, 1997, 28A, 1735. 17. Singh, A. K. and Banerjee, D., Metall. Mater. Trans. A, 1997, 28A, 1745. 18. Naka, S., Thomas, M., Sanchez, C. and Khan, T., in Structural Intermetallics 1997, ed. M. V. Nathal et al. TMS, Warrendale, Pennsylvania, 1997, p. 313. 19. Oh, M. H., Inui, H., Nakamura, A. and Yamaguchi, M., Acta metall. mater., 1992, 40, 167.
20. Hyman, M. E., McCullough, C., Levi, C. G. and Mehrabian, R., Metall. Trans. A, 1991, 22A, 1647. 21. Godfrey, A. B. and Loretto, M. H., Intermetallics, 1996, 4, 47. 22. Das, S., Mishurda, J. C., Allen, W. P., Perepezko, J. H. and Chumbley, L. S., Scripta metall. mater., 1993, 28, 489. 23. Li, Y. G. and Loretto, M. H., Acta metall. mater., 1994, 42, 2913. 24. Morris, M. A., Li, Y. G. and Leboeuf, M., Scripta metall. mater., 1994, 31, 449. 25. Kattner, U. R., Lin, J.-C. and Chang, Y. A., Metall. Trans. A, 1992, 23A, 2081. 26. Murray, J. L. (ed.), Phase Diagrams of Binary Titanium Alloys. ASM International, Metals Park, Ohio, 1987, pp. 169±175. 27. Kaufman, L. and Nesor, H., CALPHAD, 1978, 2, 325. 28. Toop, G. W., Trans. Am. Inst. Min. Engrs, 1965, 233, 233. 29. Pelton, A. D. and Bale, C. W., CALPHAD, 1977, 1, 253. 30. Brody, H. D. and Flemings, M. C., Trans. TMSAIME, 1966, 236, 615. 31. Chen, S.-W., Chuang, Y.-Y., Chang, Y. A. and Chu, M. G., Metall. Trans. A, 1991, 22A, 2837. 32. Hunt, J. D., Mater. Sci. Engng, 1984, 65, 75. 33. Kurz, W. and Fisher, D. J., Fundamentals of Solidi®cation. Trans Tech Publications, Rockport, Massachusetts, 1984.
Acta mater. Vol. 46, No. 18, pp. 6529±6540, 1998 # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain S1359-6454(98)00310-3 1359-6454/98 $19.00 + 0.00
MICROSTRUCTURAL CONTROL OF TiAl±Mo±B ALLOYS BY DIRECTIONAL SOLIDIFICATION D. R. JOHNSON, K. CHIHARA, H. INUI and M. YAMAGUCHI{ Department of Materials Science and Engineering, Kyoto University, Sakyo-ku, Kyoto 606-01, Japan (Received 6 July 1998; accepted 17 August 1998) AbstractÐDirectional solidi®cation experiments were performed to see if the g/a2 lamellar boundaries could be aligned parallel to the growth direction for an ingot containing columnar lamellar grains. From microstructural analysis of arc-melted ingots, a projection of the partial liquidus surface for the Ti±Al±Mo system was constructed and this was used with an analytical description of the phase diagram to calculate the solidi®cation paths. Following this analysis, Ti±Al±Mo(+B) ingots were directionally solidi®ed at growth rates between 40 and 100 mm/h. By controlling the solidi®cation path, a possible processing window where a reasonably well aligned g/a2 microstructure can be produced was identi®ed. As-processed microstructures exhibited a compressive yield strength greater than 500 MPa at 8008C and 2% ductility at room temperature. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.
1. INTRODUCTION
Gamma titanium aluminides are candidate materials for lightweight structural applications at temperatures near 8008C. Of the numerous microstructures that can be developed in TiAlbased alloys [1, 2], the fully lamellar microstructure consisting of TiAl (g) and Ti3Al (a2) possesses the best combination of room temperature toughness and high temperature strength [1±5]. However, as detailed by studies using polysynthetically twinned (PST) crystals, where the whole ingot consists of a single lamellar grain, mechanical properties are found to be extremely anisotropic with respect to the lamellar orientation [5±7]. For example, the optimum combination of strength, toughness, and ductility can be obtained if the lamellar microstructure is aligned parallel to the test axis [6, 7]. However, if the lamellar boundaries are perpendicular to the tensile axis, the material is extremely brittle [6]. Therefore, if the lamellar boundaries can be aligned parallel to the ingot's axis by an appropriate processing technique, then the resulting material would have a good combination of strength and toughness. One way to produce such materials is by working TiAl alloys within the a phase temperature range [8, 9] to produce the desired texture. However, this paper will focus on the problem of forming a favorable texture directly upon solidi®cation. For this case, the target microstructure is one consisting of columnar grains that have transformed to the fully lamellar microstructure upon cooling with all the g/a2 boundaries parallel to the growth {To whom all correspondence should be addressed.
direction. In addition, each columnar lamellar grain should be rotated about its longitudinal axis to provide reasonable o-axis fracture toughness as shown in Fig. 1. For clarity, only a few columnar grains are shown in Fig. 1. However, for increased ductility [4] and a greater dynamic fracture toughness [10], the columnar grain size should be as small as possible. Unfortunately, simple casting operations often produce the worst case scenario with the lamellar boundaries all perpendicular to the heat ¯ow direction [3, 11]. To produce materials with a balance of properties, thermomechanical-processing routes are then required to remove the solidi®cation texture and to re®ne the lamellar colony size [1, 3, 11]. This unfavorable solidi®cation texture is a result of the preferred growth of the a phase parallel to the [0001] direction. Upon cooling, the lamellar microstructure would then form from the parent a phase following the Blackburn orientation relationship [12] of 0001a ==f111gg and h1120ia ==h110ig resulting in the observed lamellar orientation. As was recently shown, the orientation of the a phase can be controlled during solidi®cation by using a seed from the TiAl±Si system [13]. However, for this case, the resulting ingots consist of a single lamellar grain whereas the target microstructure should consist of numerous columnar grains. Another approach is to change the solidi®cation pathway such that the bulk of the ingot solidi®es as the Ti-based b.c.c. solid solution (b phase). This can be accomplished by lowering the Al content below 45 at.% [2] for binary Ti±Al alloys. Typically such alloys form an equiaxed lamellar microstructure upon casting. However, if directionally solidi-
6529
6530
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION Table 1. Compositions and processing conditions of directionally solidi®ed ingots with the results of EDS analysis enclosed in parentheses Ingot number
Fig. 1. Target microstructure for a directionally solidi®ed TiAl-based alloy.
®ed, the resulting ingot can consist of large columnar grains with the lamellar boundaries oriented between 0 and 458 to the growth direction [14]. If the b phase dendrites have a preferred growth direction parallel to h001i, then the orientation of the lamellar boundaries can be explained by the Burgers orientation relationship [15] between the a and b phases. For this case, a would nucleate from the b phase with the (0001)a plane being inclined by 458 or parallel to the growth direction [14]. The lamellar microstructure would then form upon further cooling as previously mentioned resulting in the observed lamellar orientation. Unfortunately, directionally solidi®ed ingots with a composition near 45 at.% Al tend to be brittle [2]. This can be attributed to the larger volume fraction of the a2 phase associated with the lower Al content, the very large columnar grain size (only a few columnar grains are found in a 12 mm diameter ingot [14]), and the poor alignment of the microstructure since the majority of the lamellar grains tend to be oriented 458 to the growth direction (as would be expected from Burgers orientation relationship). Therefore, to improve the mechanical properties, the lamellar grain size must be signi®cantly reduced and the lamellar orientation needs to be well aligned with the growth direction. To explore this possibility, an investigation of the solidi®cation behavior of TiAl±Mo and TiAl±Mo±B ingots was undertaken. Additions of Mo were used to promote b solidi®cation [16±18] as well as to provide some solid solution strengthening [18, 19]. In addition, boron additions were used to re®ne the solidi®cation microstructure [3, 20, 21]. With these
Fig. 2. Schematic drawing of the directional solidi®cation furnace.
1 2 3 4 5 6
Composition (at.%)
Growth rate (mm/h)
Ti±45.5Al±1.5Mo (Ti±43.4Al±3.8Mo)$ Ti±46.5Al±1.0Mo (Ti±44.8Al±2.1Mo)$ Ti±46.5Al±1.5Mo±0.5B (Ti±45.5Al±1.6Mo±B)% Ti±46.5Al±1.5Mo±0.6B (Ti±47.5Al±1.5Mo±B) Ti±46.5Al±1.5Mo±0.7B (Ti±46.9Al±1.5Mo±B) Ti±46.5Al±1.5Mo±0.7B (Ti±47.0Al±1.4Mo±B)
40 40 100 100 100 100
$Mo pickup from melting on a Mo hearth, all other ingots were directly melted in CaO crucibles. %Boron content was not measured for these ingots.
alloying additions, the purpose of this investigation was to determine the conditions that are needed to develop a well aligned solidi®cation texture in a cast or directionally solidi®ed ingot. 2. EXPERIMENTAL PROCEDURES
Master ingots for directional solidi®cation were prepared by arc-melting in puri®ed argon at 1 atm. Directional solidi®cation was performed in a laboratory scale ram-type furnace using 16 mm diameter CaO crucibles as schematically drawn in Fig. 2. For the ®rst few ingots processed, the initial charge was melted in a Mo hearth and then dropped into a CaO crucible. However, during this procedure, there was signi®cant Mo pickup from the hearth. Therefore, to maintain the target composition, all other ingots were melted directly in the crucible and then directionally solidi®ed at velocities between 40 and 100 mm/h. The compositions of all six ingots processed are listed in Table 1 and these will be referred to by number in the proceeding text. A partial liquidus surface for the Ti±Al±Mo system was determined by comparing the microstructures of arc-melted ingots. In addition, the primary dendritic phase was also identi®ed in TiAl±Mo alloys containing small additions of boron (0.6 and 0.8 at.%). Small arc-melted buttons (about 15 g each) were made from high purity elements (99.99 at.%). Each arc-melted ingot was melted and turned at least ®ve times to promote homogeneity. A scanning electron microscope (SEM) equipped with an energy-dispersive spectrometer (EDS) as well as an optical microscope (OM) were all used to characterize the microstructures of the arc-melted and directionally solidi®ed ingots. Backscattered electron (BSE) imaging was used to help identify the phases present and to locate the dendrite cores. Specimens for SEM observations were electropolished in a solution of 5% HClO4±35% N-butanol± 60% methanol by volume at ÿ658C. Specimens viewed by optical microscopy were polished and etched using a solution of 5% HF±5% HNO3±90% H2O by volume. Room temperature tensile tests were performed on selected alloys using ¯at specimens 20 mm in
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
Fig. 3. Projected view of the partial liquidus surface for the Ti±Al±Mo system.
length with a gage length of 5 mm 2 mm 1 mm. Specimens were electropolished and tested in vacuum of 5 10ÿ5 Torr at a strain rate of 2 10ÿ4/s. In addition, compression tests between 25 and 11008C were also performed on specimens with dimensions of 2.5 mm 2.5 mm 5 mm. These specimens were also electropolished and tested in vacuum of 1 10ÿ4 Torr. For all tests, the tensile or compression axis was parallel to the growth direction. 3. RESULTS
3.1. Microstructures from arc-melted ingots The solidi®cation sequence was initially determined by examining the dendrite morphologies found in arc-melted buttons containing either 2 or 5 at.% Mo. From these observations, a partial liquidus surface was constructed as shown in Fig. 3. The basic result is that Mo additions shift both the primary a and b regions towards greater Al contents. In addition, the average composition within the prior dendritic regions was found to be enriched in Mo when compared to the original alloy composition, Table 2, suggesting that Mo additions increase the melting temperature of both the a and g phases as noted by the direction of the arrows in Fig. 3. 3.1.1. Ti±Al±5Mo series. Typical microstructures for alloys containing 5 at.% Mo are shown in Fig. 4. For this series of alloys, b was found to be the primary solidi®cation phase for alloys with an aluminum content less than about 52 at.% as Table 2. Average composition of dendrites as measured by EDS analysis Alloy composition (at.%)
b dendrites Ti±48Al±2Mo Ti±48Al±5Mo a dendrites Ti±53Al±2Mo Ti±55Al±5Mo
Average composition Ti
Al
Mo
51.2 47.9
46.5 46.5
2.3 5.6
47.3 41.7
50.4 52.6
2.3 5.7
6531
shown in Fig. 4(a)±(d). The phase with bright contrast within the prior dendritic regions was found to be rich in Mo and lean in Al when compared to the ingot composition. For example for the Ti±48Al± 5Mo alloy (all compositions are given in atomic percentages), a composition of Ti±47.7Al±9.8Mo was measured for this phase, which can be inferred from the literature as being the B2 phase [16, 17, 22± 24]. The remaining matrix was found to consist of the g/a2 lamellar microstructure. The microstructure for the Ti±46Al±5Mo button, Fig. 4(a) and (b), was essentially equiaxed and an obvious aluminum-rich interdendritic region was not observed suggesting that the complete ingot solidi®ed as the b phase. The prior b grains are also found to be subdivided into a number of colonies due to the formation of WidmanstaÈtten plates of a [16]. A similar microstructure for a Ti±44Al±2Mo alloy has been characterized by Li and Loretto both in terms of microstructure and fracture behavior [23]. In contrast, remnants of the prior dendritic morphology are clearly visible in all the alloys containing 5 at.% Mo with compositions between 47 and 52 at.% Al, and typical microstructures are shown in Fig. 4(c) and (d). The primary phase is identi®ed as b by noting that the dendrite arms are orthogonal to one another. Within the interdendritic region, the microstructure is a result of a peritectic cascade consisting of the L + b 4 a and the L + a 4 g reactions. The dark regions surrounding the dendrites in Fig. 4(c) are found to consist of the lamellar microstructure, Fig. 4(d), indicating it was originally the peritectic a phase. Similarly, regions of peritectic g are found within the prior peritectic a regions [as indicated by the arrow in Fig. 4(d)], completing the cascade. Increasing the Al content simply increases the amount of peritectic g while decreasing the amount of the other phases. For higher aluminum contents in the range of 53±57 at.%, a is identi®ed as the primary phase. The hexagonal symmetry of the dendrite arms can be seen for typical microstructures in Fig. 4(e) and (f). The prior dendritic regions for the a phase are also found to be rich in Mo and lean in Al as was the case for the b phase, Table 2. In addition, a bright phase can be seen surrounding the a phase in Fig. 4(e) and (f), which may be the result of b (or B2) precipitation at intermediate temperatures. Once the aluminum content was increased above 57 at.% Al, a drastic change in the dendrite morphology was observed, Fig. 4(g) and (h), and this suggests a transition to the solidi®cation of primary g. A bright phase with a needle-like morphology, having a composition roughly of Ti±55.4Al±6.9Mo, is observable in the g phase. Since, a clear cubic dendrite morphology cannot be observed, these needles are most likely the remnants of a very small volume fraction of a phase dendrites. A possible side arm oriented 608 to the dendrite spine is indi-
6532
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
Fig. 4. SEM backscattered electron images of microstructures taken from arc-melted Ti±xAl±5Mo ingots showing (a) and (b) the Ti±46Al±5Mo alloy, (c) and (d) prior b dendrites of Ti±48Al±5Mo alloy with interdendritic g marked by an arrow, (e) and (f) prior a dendrites of Ti±56Al±5Mo alloy, and (g) and (h) Ti±58Al±5Mo alloy with an arrow marking a dendrite side arm.
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
6533
Fig. 5. SEM backscattered electron images of microstructures taken from arc-melted Ti±xAl±2Mo ingots showing prior b dendrites for (a) the Ti±46Al±2Mo alloy and (b) the Ti±50Al±2Mo alloy.
cated by the arrow in Fig. 4(h) and provides further evidence for this assumption. These prior a dendrites would then be sites for the formation of the b or B2 phase at intermediate temperatures. Also an eutectic-like microstructure is visible in Fig. 4(h) and consists of a ®brous bright phase (Ti±57.9Al± 7.7Mo) within a g matrix (Ti±58.8Al±4.4Mo). From their compositions, it would appear that the eutectic phases are b and g. From the partial isothermal sections proposed by Das et al. [22] and Singh and Banerjee [17], both the Ti±56Al±5Mo and Ti±58Al±5Mo alloys should consist of the g and b (or B2) phases once equilibrated below 014008C. Hence, the bright contrast in Fig. 4(e)±(h) should result from the formation of the B2 phase. Furthermore, the eutectic microstructure consisting of the g and b phases in Fig. 4(g) and (h), provides evidence for the L + a 4 b + g invariant reaction as suggested by Singh and Banerjee [16]. From the above arguments and by using the compositions measured from the eutectic g and b phases, the L + a 4 b + g reaction is sketched in the projected view of the liquidus surface shown in Fig. 3. 3.1.2. Ti±Al±2Mo series. The solidi®cation behavior of alloys containing 2 at.% Mo was essentially the same as those containing 5 at.% Mo. The composition range where a distinct cubic (b) dendritic morphology was visible was found to be between 46 and 50 at.% Al as shown in Fig. 5(a) and (b). The main dierence between these ®gures is the relative amounts of the phases that formed from the liquid. In Fig. 5(a), the prior b dendrites contain veins of the B2 phase. The dark regions surrounding the dendrites were found to consist of the g/a2 lamellar microstructure indicating a prior peritectic a phase. In contrast, the dark regions in Fig. 5(b) consist of the g phase. For this microstructure, the amount of B2 phase is much less due to a smaller volume fraction of primary b dendrites. Surrounding these b dendrites are lamellar regions that formed from the peritectic a phase and can be identi®ed by the lack
of the B2 phase. Hence, during solidi®cation, the Ti±50Al±2Mo alloy cascaded through two peritectic reactions while for the Ti±46Al±2Mo alloy only the L + b 4 a peritectic reaction occurred. 3.2. Solidi®cation paths from thermodynamic calculations In order to determine the solidi®cation paths and to determine the degree of microsegregation, an analytical description of the phase diagram would be useful. For example, if the tie lines of the ternary system are known then a Scheil type analysis can be used to determine the solidi®cation path. For the Ti±Al±Mo system, a simple approach was taken to estimate the Gibbs energies of the ternary phases (liquid, b, and a) from the Gibbs energies of the binary phases [25±27]. For this approach, Toop's method [28] was used in conjunction with the experimentally determined liquidus surface. The form of Toop's equation used along with the analytical description of the binary phases from the literature are listed in Table 3. Since the melting temperature of the a phase was found to increase with Mo additions, the maximum liquidus temperature should locate within the ternary system. For this case, Toop's equation alone is insucient to calculate the ternary Gibbs energies. Hence, a ternary term was added to the equation using a polynomial of the type proposed by Pelton and Bale [29]. The coecients of the polynomial were estimated by ®rst varying the interaction parameter of the a phase in the binary Mo±Al system and noting the response in the ternary system. The coecients of the ternary term were then adjusted to reproduce this response with the a interaction parameter set to zero in the Mo±Al system. This was done by an iterative approach and the coecients were not optimized. Once a satisfactory ®t was obtained, the chemical potentials were numerically calculated and from these, the equilibrium concentrations were found. A partial isothermal section at 1773 K that was calculated is shown in
6534
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION Table 3. Analytical description for the excess free energy of the liquid, b and a phases in the Ti±Al±Mo system (units are in J/mol and J/mol K). For the respective elements Ti, Al, or Mo the standard states are: a refers to f.c.c.; b refers to b.c.c.; and liquid refers to liquid Lattice stabilities of the pure elements [25±27] ÿ 8GLTi ÿ14146:0 7:2880 T 8Gb:c:c: Ti ÿ 8GLTi ÿ18318:0 10:8984 T 8Gh:c:p: Ti f:c:c: ÿ 8GLTi ÿ12316:0 10:7984 T 8GTi ÿ 8GLAl ÿ628:0 6:6598 T 8Gb:c:c: Al ÿ 8GLAl ÿ5151:4 9:6178 T 8Gh:c:p: Al f:c:c: ÿ 8GLAl ÿ10711:0 11:4728 T 8GAl L 8Gb:c:c: Mo ÿ 8GMo ÿ24267:0 8:368 T L 8Gh:c:p: Mo ÿ 8GMo ÿ15899:0 8:368 T f:c:c: ÿ 8GLMo ÿ13807:0 8:996 T 8GMo Ti±Mo binary [26] Liquid: DGxs 14943:0 XTi XMo b phase: DGxs 15657:0 5238:0 1 ÿ 2 XMo XTi XMo a phase: DGxs 36647:0 XTi XMo Mo±Al binary [27] Liquid: DGxs ÿ32024:0 XMo XAl b phase: DGxs ÿXMo XAl 30543:0 XMo 16736:0 XAl ÿ 12:552 XMo XAl XAl T Ti±Al binary [25] Liquid: DGxs ÿ112570:2 41:11378 T ÿ 7950:8 XTi ÿ XAl XTi XAl b phase: DGxs ÿ129396:7 40:06310 T XTi XAl a phase: DGxs ÿ123788:5 33:20902 T ÿ16034:9 12:18272 T XTi ÿ XAl XTi XAl Ti±Al±Mo Form of Toop's equation used with additional ternary term xs xs DGxs XMo Gxs TiMo XAl GTiAl = 1 ÿ XTi 1 ÿ XTi 1 ÿ XTi GMoAl P XTi XAl XMo Liquid: P = 0 b phase: P = 0 a phase: P ÿ76500:0 ÿ0:5 ÿ 0:9 XTi 5:9 XAl
Fig. 6. In this ®gure, the liquidus composition at either of the three phase tie triangles matches quite well with what would be expected from the experimentally determined liquidus projection, Fig. 3. While the above analysis is quite simple, it is useful for predicting the range of possible microstructures that can form. The limiting cases range from equilibrium to the case were diusion in the solid is neglected. For the latter, the solidi®cation paths were found using an approach that couples a thermodynamic model of the system with the assumptions of the Scheil model as described by Brody and Flemings [30] and by Chen et al. [31]. The basic approach is that the equations describing the mass balance during solidi®cation are solved numerically so that the partition coecients (or tie lines) can be calculated from the thermodynamic model at each iteration. From this Scheil type analysis, the amount of interdendritic g was calculated for the arc-melted buttons and compared to those measured from metallographic analysis. These results are shown in Fig. 7. The predicted values compare well with
those measured from the arc-melted ingots containing 2 and 5 at.% Mo. In addition, for alloys containing 2 at.% Mo, the calculated relative amounts of each phase that solidi®ed as a function of ingot composition are shown in Fig. 8.
Fig. 6. Estimated partial isothermal section at 1773 K for the Ti±Al±Mo system.
Fig. 7. Calculated and measured amounts of interdendritic g for Ti±Al±Mo alloys.
3.3. Directional solidi®cation of TiAl±Mo alloys For the ®rst set of directionally solidi®ed ingots, the initial charge was melted in a Mo hearth and then dropped into a CaO crucible. Unfortunately, during this process some of the Mo hearth was melted and the resulting ingots were found to have a much greater Mo content than the target composition as listed in Table 1. However, some interesting microstructures developed from these ingots. As the molten charge was dropped on a water cooled chill plate, these ingots consist of a chill region and then a directionally solidi®ed region.
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
Fig. 8. Relative amounts of each phase solidi®ed in Ti± Al±Mo alloys as calculated by the Scheil model.
Microstructures for a typical case are shown in Fig. 9 for a Ti±43.4Al±3.8Mo ingot (ingot 2). Macroscopically, columnar grains consisting of the g/a2 lamellar microstructure are found for the complete length of the ingot [Fig. 9(a)]. The microstructure for the directionally solidi®ed section consists of b grains that have a columnar shape as shown in Fig. 9(b). However, the microstructure appears equiaxed within the chill zone as shown in Fig. 9(c). The presence of the B2 phase within the cored regions suggests a prior structure of equiaxed b grains. From the previous calculations of the solidi®cation paths, the minimum amount of b that should form upon solidi®cation is nearly 80%. Thus, the columnar lamellar grains are probably the result of directional growth of a via the solid state transformation of b to a.
6535
Although this ingot solidi®ed through the b phase, the microstructure did not approach the target microstructure as speci®ed in Fig. 1 since the grain size was large and the lamellar microstructure was not aligned with the growth direction. Therefore, solidifying through the b phase alone is not sucient to produce the desired microstructure and a dierent approach must be taken. Since the lamellar microstructure is formed from the a phase, better microstructural control may result by changing the morphology of the secondary a phase that forms from the liquid. To investigate this problem, boron was added to try to change the solidi®cation behavior of the TiAl±Mo ingots. From the microstructure of arcmelt ingots, the addition of boron does not signi®cantly change the compositional region where primary b and a form as shown in Fig. 10. Thus, the volume fraction of the phases calculated earlier (Fig. 8) should roughly be the same as that for alloys containing small boron additions. Using this assumption, a baseline composition of Ti±46.5Al± 1.5Mo±0.6B was chosen for directional solidi®cation experiments since about half the ingot should solidify as b and half as a. 3.4. Directional solidi®cation of TiAl±Mo±B alloys The addition of boron greatly reduced the lamellar grain size as shown in Fig. 11(a) for ingot 3. The lamellar orientation measured along the length of this ingot was mixed and lamellar grains could be found at orientations from parallel to nearly per-
Fig. 9. Microstructures from a directionally solidi®ed Ti±43.4Al±3.8Mo ingot showing (a) the macroscopic view of the longitudinal section, and SEM backscattered electron images from (b) the directionally solidi®ed portion of the ingot, and (c) the chill zone.
6536
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
Fig. 10. Projected view of the Ti±Al±B liquidus surface as determined by Hyman et al. [20] plotted together with the results from alloys containing 1.5 at.% Mo (these compositions are denoted by the open circles and squares). The solidi®cation path for Ti±46.5Al±1.5Mo±0.6B (denoted by the small closed circle) is also outlined by the numbers ``1'' to ``3''.
pendicular to the growth direction. However, with a slight increase in boron content, portions of ingots 4±6 had a well aligned lamellar microstructure. The lamellar grain structures along the length of these ingots were similar and that for ingot 4 is shown in Fig. 11(b). The ®rst half of this ingot consists of columnar lamellar grains with mixed lamellar orientations, and this section is labeled ``M'' in Fig. 11(b). However, within the second half, a columnar to equiaxed transition occurred (labeled as ``E'') that was followed again by columnar grains but now with a much smaller grain size. These short columnar grains (the region labeled as ``A'') had a very well aligned lamellar microstructure as shown in Fig. 11(c). Note that a number of these grains are oriented such that the lamellar plane is
Fig. 11. Microstructures of directionally solidi®ed TiAl±Mo±B ingots showing (a) the longitudinal section of a Ti±46.5Al±1.5Mo±0.5B ingot, (b) the longitudinal section of a Ti±46.5Al±1.5Mo±0.6B ingot with aligned ``A'', equiaxed ``E'', and mixed ``M'' lamellar regions marked, and (c) the enlarged view of the microstructure of region ``A'' with arrows marking the lamellar colonies rotated such that (111)g is almost parallel with the plane.
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
6537
the unetched microstructure of ingot 4 is shown in Fig. 12. The bright contrast consists of the boride phase and associated pullout during polishing. The rest of the microstructure consists of b phase dendrites with the main dierence between each section being the dendrite spacing which is somewhat ®ner within the aligned section of the ingot. Since no major changes are found in Fig. 12, the change in microstructure as observed in Fig. 11(b) must be due to changes occurring in the interdendritic a phase and this will be considered later in the text. 3.5. Compression and tensile test results Tensile specimens were machined from the aligned region in ingots 4±6 and compression specimens were also taken from ingot 6 within the same region. From Fig. 13, a compressive yield strength greater than 500 MPa is maintained up to 8008C after which it decreases to 400 MPa at 10008C. Any further increase in temperature causes a very rapid decrease in yield strength. The results from the tensile tests are listed in Table 4 and these basically re¯ect the dierences in the lamellar orientation of the test specimens. Plastic strains up to 2% were measured from specimens taken from ingots 4±6 where the lamellar microstructure was aligned with the growth direction. However, specimens taken from the ®rst half of these ingots displayed little ductility when tested. A dierence in the lamellar orientation is found when observing the fracture surfaces as shown in Fig. 14. Regions of fast fracture indicated by large cleavage surfaces are found in all the tested specimens machined from the ®rst half of ingots 4 and 5 with a typical fracture surface shown in Fig. 14(a). In contrast, the fracture surfaces associated with specimens machined from the portion of the ingot with the aligned lamellar microstructure displayed a composite or splitting type fracture surface without the regions of fast fracture, Fig. 14(b).
Fig. 12. Photographs taken from the unetched ingot shown in Fig. 14 for (a) the mixed orientation, (b) the equiaxed, and (c) aligned lamellar regions.
roughly parallel to the plane of view as indicated by the coarse spacing of the lamellar trace for regions marked in Fig. 11(c). Thus, the microstructure within the aligned portion of the ingot approaches that of the target microstructure with the lamellar grains rotated with respect to each other along the ingot's axis. Somewhat surprising is that when ingots 4±6 are polished and viewed unetched by eye, little change in microstructure is observable along the entire length of the ingots. What is seen is well-de®ned columnar texture outlined by the interdendritic boride phase. For example, the macroscopic view of
Fig. 13. Compressive yield strength of a Ti±46.5Al± 1.5Mo±0.7B alloy processed at 100 mm/h as a function of temperature.
6538
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
Table 4. Tensile properties of directionally solidi®ed Ti±46.5Al± 1.5Mo±(0.6±0.7)B (at.%) alloys with a well aligned lamellar microstructure and with a mixed lamellar orientation Specimen
Yield stress (MPa)
Elongation (%)
Ingot 4: Ti±46.5Al±1.5Mo±0.6B Aligned region 440 370 430 Mixed region Ð Ð Ingot 5: Ti±46.5Al±1.5Mo±0.7B Aligned region 440 400 400 Mixed region 360 280 Ingot 6: Ti±46.5Al±1.5Mo±0.7B Aligned region 500 500
2.9 1.2 1.2 0 0 2.5 0.4 1.2 0.3 0.5 0.4 0.5
Growth rate = 100 mm/h.Tested in vacuum of 5 10ÿ5 Torr.
4. DISCUSSION
The microstructures of the directionally solidi®ed TiAl±Mo±B alloys only approached the target microstructure for a very short portion of the ingots [Fig. 11(c)]. Since the portion with the aligned lamellar microstructure was found after a columnar to equiaxed transition, the qualitative analysis described by Hunt [32] on this matter may be helpful. Hunt's analysis considers the growth of equiaxed grains ahead of a columnar front during directional solidi®cation and the ®nal results of the analysis are that fully equiaxed growth occurs when 3 3 G<0:617N 1=3 0 f1 ÿ DTN = DTc gDTc
where G is the average thermal gradient, N0 the total number of heterogeneous nucleation sites per volume, DTN the undercooling at the heterogeneous nucleation temperature, and DTc the undercooling at the solid±liquid interface, which is proportional to the square root of the growth velocity [32]. However, for the directionally solidi®ed TiAl± Mo±B alloys, little change was found in the b dendrite morphology as columnar b growth was found for the complete ingot (Fig. 12). Hence, the microstructural changes found in these alloys [Fig. 11(b)] must result from a change in the secondary a morphology. Therefore, both the b and a phases must be considered when applying the above model to the columnar to equiaxed transition and this is schematically shown in Fig. 15. The columnar to equiaxed transitions depend upon the eciency of the nucleating agent, DTN, and by the number of heterogeneous nucleation sites, N0. By increasing the nucleation eciency (decreasing DTN) the curves in Fig. 15 are shifted down towards lower velocities. Similarly, by increasing N0, the curves are shifted right towards higher values of G. Under the same processing conditions (G and V), if the a and b phases are considered separately and if the set of values for DTN and N0 are similar, then the columnar to equiaxed
Fig. 14. Fracture surfaces of tensile specimens machined from (a) the mixed orientation, and (b) the aligned lamellar regions of ingot number 4.
transition will always occur ®rst for the b phase for the same processing conditions (e.g. G and V). This is due to the smaller freezing range of a as given by the phase diagram which, when compared to b, results in smaller values of DTc [33]. Furthermore, since the growth of the peritectic a phase occurs behind the primary b dendrites, the solute undercooling for the a phase should be even less since the b dendrite spacing will adjust itself to minimize the constitutional supercooling within the interdendritic liquid. For the TiAl±Mo ingots, boron was added to change the above behavior by providing an active catalyst for a nucleation. The solidi®cation path of the Ti±46.5Al±1.5Mo±0.6B ingots can be described by using Fig. 10. For this composition (labeled as
Fig. 15. Schematic representation of the columnar/ equiaxed transition for the a and b phases in a TiAl±Mo± B alloy. For the processing conditions marked by the ®lled circle, columnar growth of b would be expected while, for the a phase, equiaxed growth would be expected.
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
``1'' in Fig. 10), growth of b would ®rst occur causing the interdendritic liquid to become enriched in boron. Boride particles would then nucleate once the L 4 TiB2 + b monovariant line is met (composition ``2'' in Fig. 10). However, the a phase would not nucleate until the liquid composition reaches the L + B 4 a + TiB2 invariant reaction at composition ``3'' as labeled in Fig. 10. The important point is that the boride particles can nucleate within the interdendritic liquid before the solidi®cation of the a phase. Thus, if the boride particles are an active catalyst for a nucleation, then the columnar to equiaxed transition curve for the a phase can be shifted towards higher G values and possibly past that for the b phase as is schematically drawn in Fig. 15. Hence, it may be possible to have equiaxed a growth even though the growth of b is columnar, and Figs 11 and 12 provide evidence that this can indeed be the case. The growth morphology for this case is sketched in Fig. 16. If a grains are continuously nucleated at the boride particles, then growth can occur only over short distances before impingement upon more recently nucleated grains occurs. During this period, there will be competitive growth between the dierently oriented a grains. Consequently, growth parallel to h1010i may be possible. Therefore, the aligned lamellar microstructure shown in Fig. 11(c) is suggested to occur from competitive growth of h1010i oriented a grains behind a columnar b dendritic front and from the nucleation of a from prior [100] oriented b dendrites following the Burgers orientation relationship. Lastly, only a small section of the directionally solidi®ed Ti±Al±Mo±B ingots had an aligned g/a2 lamellar microstructure as shown in Fig. 11(b). From Fig. 15, the columnar to equiaxed transition is very sensitive to changes in the thermal gradient at higher growth velocities, and thus the processing window where columnar growth of b and equiaxed growth of a can occur together is very small. Therefore, the variation of the microstructure along
Fig. 16. Schematic view of columnar growth of b and columnar plus equiaxed growth of a.
6539
the length of the ingot may result from the change in the thermal gradient during processing. For example, initially there is a high thermal gradient due to the water cooled chill plate used, resulting in columnar growth of both the a and b phases [region ``M'' in Fig. 11(b)]. However, away from the chill plate, the thermal gradient decreases and equiaxed growth of a occurs [regions ``A'' and ``E'' in Fig. 11(b)]. The question of whether steady state conditions can be reached and maintained within this narrow processing window must be answered. However, the present work shows that such a possibility may exist and work is now under way to produce larger size ingots. 5. CONCLUSIONS
The solidi®cation paths for Ti±Al±Mo alloys were calculated from an experimentally determined liquidus projection and from an analytical description of the phase diagram. By controlling the solidi®cation path, a possible processing window where a reasonably well aligned g/a2 microstructure can be produced was identi®ed. This is accomplished by alloying with boron such that boride particles will nucleate behind the primary b dendrites before the nucleation of the a phase, resulting in columnar growth of the b phase followed by equiaxed growth of the a phase. Such microstructures exhibited a compressive yield strength of more than 500 MPa at 8008C and 2% ductility at room temperature. AcknowledgementsÐThis work was supported by ``Research for the Future'' Program (JSPS-RFTF 96R12301), The Japan Society for the Promotion of Science, and in part by a research grant from the Research and Development Institute of Metals and Composites for Future Industries. Funding from the Ministry of Education, Science and Culture, Japan (grant No. 09750772) is also gratefully acknowledged. REFERENCES 1. Kim, Y.-W., JOM, 1994, 46, 30. 2. Huang, S. C., in Structural Intermetallics, ed. R. Darolia, J. J. Lewandowski, C. T. Liu, P. L. Martin, D. B. Miracle and M. V. Nathal. TMS, Warrendale, Pennsylvania, 1993, p. 299. 3. Wagner, R., Appel, F., Dogan, B., Ennis, P. J., Lorenz, U., Mullauer, J., Nicolai, H. P., Quadakkers, W., Singheiser, L., Smarsly, W., Vaidya, W. and Wurzwallner, K., in Gamma Titanium Aluminides, ed. Y.-W. Kim, R. Wagner and M. Yamaguchi. TMS, Warrendale, Pennsylvania, 1995, p. 387. 4. Liu, C. T., Schneibel, J. H., Maziasz, P. J., Wright, J. L. and Easton, D. S., Intermetallics, 1996, 4, 429. 5. Yamaguchi, M. and Inui, H., in Structural Intermetallics, ed. R. Darolia, J. J. Lewandowski, C. T. Liu, P. L. Martin, D. B. Miracle and M. V. Nathal. TMS, Warrendale, Pennsylvania, 1993, p. 127. 6. Inui, H., Kishida, K., Misaki, M., Kobayashi, M., Shirai, Y. and Yamaguchi, M., Phil. Mag. A, 1995, 72, 1609. 7. Yakoshima, S. and Yamaguchi, M., Acta mater., 1996, 44, 873.
6540
JOHNSON et al.: DIRECTIONAL SOLIDIFICATION
8. Kim, Y.-W. and Dimiduk, D. M., in Structural Intermetallics 1997, ed. M. V. Nathal et al. TMS, Warrendale, Pennsylvania, 1997, p. 531. 9. Jin, Z., Gray, G. T. III and Kim, Y.-W., Mater. Sci. Engng, 1997, A239-240, 729. 10. Enoki, M., Fujikawa, S. and Kishi, T., J. Japan Inst. Metals, 1994, 58, 424. 11. Mecking, H. and Hartig, Ch., in Gamma Titanium Aluminides, ed. Y.-W. Kim, R. Wagner and M. Yamaguchi. TMS, Warrendale, Pennsylvania, 1995, p. 525. 12. Blackburn, M. J., in The Science, Technology, and Application of Titanium, ed. R. I. Jaee and N. E. Promisel. Pergamon Press, Oxford, 1970, p. 663. 13. Johnson, D. R., Masuda, Y., Inui, H. and Yamaguchi, M., Acta metall. mater., 1997, 45, 2523. 14. Johnson, D. R., Masuda, Y., Shimada, Y., Inui, H. and Yamaguchi, M., in Structural Intermetallics 1997, ed. M. V. Nathal et al. TMS, Warrendale, Pennsylvania, 1997, p. 287. 15. Burgers, W. G., Physica, 1934, 1, 561. 16. Singh, A. K. and Banerjee, D., Metall. Mater. Trans. A, 1997, 28A, 1735. 17. Singh, A. K. and Banerjee, D., Metall. Mater. Trans. A, 1997, 28A, 1745. 18. Naka, S., Thomas, M., Sanchez, C. and Khan, T., in Structural Intermetallics 1997, ed. M. V. Nathal et al. TMS, Warrendale, Pennsylvania, 1997, p. 313. 19. Oh, M. H., Inui, H., Nakamura, A. and Yamaguchi, M., Acta metall. mater., 1992, 40, 167.
20. Hyman, M. E., McCullough, C., Levi, C. G. and Mehrabian, R., Metall. Trans. A, 1991, 22A, 1647. 21. Godfrey, A. B. and Loretto, M. H., Intermetallics, 1996, 4, 47. 22. Das, S., Mishurda, J. C., Allen, W. P., Perepezko, J. H. and Chumbley, L. S., Scripta metall. mater., 1993, 28, 489. 23. Li, Y. G. and Loretto, M. H., Acta metall. mater., 1994, 42, 2913. 24. Morris, M. A., Li, Y. G. and Leboeuf, M., Scripta metall. mater., 1994, 31, 449. 25. Kattner, U. R., Lin, J.-C. and Chang, Y. A., Metall. Trans. A, 1992, 23A, 2081. 26. Murray, J. L. (ed.), Phase Diagrams of Binary Titanium Alloys. ASM International, Metals Park, Ohio, 1987, pp. 169±175. 27. Kaufman, L. and Nesor, H., CALPHAD, 1978, 2, 325. 28. Toop, G. W., Trans. Am. Inst. Min. Engrs, 1965, 233, 233. 29. Pelton, A. D. and Bale, C. W., CALPHAD, 1977, 1, 253. 30. Brody, H. D. and Flemings, M. C., Trans. TMSAIME, 1966, 236, 615. 31. Chen, S.-W., Chuang, Y.-Y., Chang, Y. A. and Chu, M. G., Metall. Trans. A, 1991, 22A, 2837. 32. Hunt, J. D., Mater. Sci. Engng, 1984, 65, 75. 33. Kurz, W. and Fisher, D. J., Fundamentals of Solidi®cation. Trans Tech Publications, Rockport, Massachusetts, 1984.