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Cast structure and property variability in gamma titanium aluminides
PII: SO966-9795(98)00042-9 ELSEVIER
Intermetallics 6 (1998) 629-636 0 1998 Published by Elsevier Science Limited Printed in Great Britain. All rights reserved 09669795/98/$-+x front matter
Cast structure and property variability in gamma titanium aluminides L. L. Rishel, N. E. Biery, R. Raban, V. Z. Gandelsman, T. M. Pollock* & A. W. Cramb Department of Materials Science and Engineering, Carnegie Mellon University, Pittsburgh, PA 15213, USA
(Received 20 April 1998; accepted 27 April 1998)
Aspects of cast structure that influence mechanical property variability, and in particular tensile ductility, have been studied in a Ti-Q8Al-2Cr-2Nb titanium aluminide alloy. Macrostructure, porosity, microstructure and tensile properties have been characterized over a range of casting processing conditions. Local solidification times and subsequent solid-state cooling rates during casting have been characterized via local thermal measurements in combination with solidification modeling. Large variations in cooling rate during casting dramatically influence the initial cast structure as well as the distribution of defects such as porosity. Variations in as-cast structure persist through subsequent thermal and hot isostatic pressing cycles and contribute to the variability of tensile ductility. The current understanding of the relationship of tensile ductility to processinginduced changes in structure will be briefly discussed. 0 1998 Published by Elsevier Science Limited. All rights reserved Key words: A. titanium aluminides, based on TiAl, B. ductility, mechanical properties at ambient temperature, C. casting, D. microstructure, E. microscopy.
1 INTRODUCTION
there is a need for a better understanding of the relationship between casting parameters, the range of cooling rates that can be achieved during solidification, and their effect on the structure and properties of the cast material. Recent experiments9yi0 have measured cooling rates in the mold during investment casting of gamma titanium aluminide plates. Here, we briefly review the results of those experiments, along with results of solidification modeling studies of investment cast plates. Variations in cast structure and tensile ductility, which are influenced by changes in cooling rate during casting, are shown. Issues which appear to be in need of further study are identified.
Gamma titanium aluminides are among the most promising intermetallics for high temperature structural applications. These materials may be produced via either cast or wrought processing techniques. However, casting processes have been of greater interest for components such as turbine blades and automobile exhaust valves, where component cost is a barrier to implementation.14 As with any casting, it is expected that the final properties of the material will be sensitive to process parameters which influence as-cast microstructure or casting defects such as porosity. This is particularly critical for gamma titanium aluminides, where some degree of tensile ductility is important for tolerance to foreign object damage and for the design of components which contain stress concentrations.S8 Control of the casting process is constrained by the need for cold crucible processes that maintain the chemistry of these highly reactive alloys during melting. Given these constraints,
2 INVESTMENT CASTING OF GAMMA TITANIUM ALUMINIDE PLATES Cooling rates during casting were measured in a series of experiments designed to cover a wide range of variations in investment casting process parameters. The experiments utilized investment molds that contained four rectangular plates measuring 10.2x 10.2cm with the following thicknesses: 12.7, 3.8, 6.3 and 3.8 mm. The composition
*To whom correspondence should be addressed. Fax: 001 412268-7596; e-mail: tplm + @andrew.cmu.edu 629
630
L. L. Rishel et al.
of the material studied in the cooling rate experiments was Ti48*OAl-2Cr-2Nb-O.O 15Fe-O. 180 (at%, avg. composition based on measurements on five plates). The casting process consisted of vacuum arc remelting (VAR) followed by pouring into investment molds. Since variations in the superheat of the molten metal cannot easily be achieved in such cold crucible processes, control of the solidification process for a given alloy must be achieved through the control of heat extraction from the investment mold. Accordingly, a range of mold preheat temperatures and mold insulation conditions were utilized. Mold preheat temperatures were 315, 815 or 1204”C, while molds were wrapped with varied amounts of Kaowool insulation. Thermocouples were attached in the mold near the surface of each of the four plates to obtain thermal data during mold preheating, pouring and solidification. The measured variations in mold temperature for each of the experiments were input as boundary conditions into the UES ProCAST simulation package. Given the arrangement of the plates in the cluster, the heat extraction at the plate centers was essentially one dimensional. Accordingly, one dimensional heat transfer analyses of the process were performed to obtain estimates of solidification times and subsequent cooling rates through the plate thickness.‘O Following casting, the structure and porosity distribution in these plates was characterized in the as-cast state and after hot isostatic pressing and heat treatment. Additional details on the experimental set-up and modeling are given elsewhere.9*t0 A second set of 12.7 mm plates from a heat of material with essentially the same composition, Ti-47.9A l-2Cr-2Nb-O.007Fe-O. 140, were also utilized along with the plates from the cooling rate study- to characterize the variability of tensile ductility.
3 COOLING
RATES, MICROSTRUCTURE DEVELOPMENT AND DUCTILITY
The results of solidification simulations for two different plate thicknesses with five different combinations of casting conditions are shown in Table 1. Local solidification times varied from the surface to the centerline of the plates; the calculeited range is given for each set of casting conditions. At the centerlines of the plates, local solidifitition times varied by a factor of more than 25 X, ‘from about 5 s for the 3.8 mm plates solidified with a mold preheat of 3 15°C to 140 s for the
12.7 mm thick plate cast with a fully insulated mold and a preheat of 1204°C. It is apparent that, in spite of superheat limitations, a fairly wide range of solidification times can be achieved, particularly with higher mold preheats. Also shown in Table 1 is the total time at temperatures above the eutectoid temperature (approximately 1125°C in the binary phase diagram) following solidification. This is an important parameter, since it influences the character of the initial microstructure of the casting, as will be discussed further. The wide range of cooling rates during solidification, Table 1, result in significant variation in ascast structure and porosity distribution. Figure 1 shows that the as-cast macrostructures at high and low cooling rates exhibit markedly different columnar grain widths. Figure 2 shows the distribution of porosity in the 3.8 mm thick as-cast plates. At higher cooling rates, the bulk of the porosity is present at the centerline of the plate. Conversely, at lower cooling rates (high degree of mold preheat and wrap), the porosity is preferentially located in near-surface regions, apparently due to reactions with the mold material. Since these materials are typically subjected to hot isostatic pressing (HIP) following casting, the distribution of porosity following a HIP cycle of 1260°C 170 MPa/4 h is also shown in Fig. 2(c) and (d) for comparison. The porosity remaining after HIP indicates the degree to which it was connected to the surface in the as-cast state. Near-surface porosity associated with very high mold preheats, Fig. 2(b), is retained through HIP’ing. Centerline porosity of the type shown in Fig. 2(a) is typically largely eliminated by HIP, Fig. 2(c). However, due to the fact that most of the porosity is interdendritic and highly interconnected, small pockets of micro-porosity often persist through the HIP cycle. The dendrite ,arm spacings, grain size and grain morphology in near-surface regions are sensitive to mold preh&ttP and this is likely to strongly influence the degree to which the porosity connects to the surface. The widely varying as-cast microstructures observed over the range of cooling rates studied responded differently to thermal cycles employed during HIP and accompanying heat treatments. The microstructures of Plates 2-l and 5-l following a HIP cycle of 126O”C/l70 MPa/4 h are shown in Fig. 3(a) and (b). These plates were cast under the same conditions as the plates shown in Figs. 1 and 2, with relatively high and low cooling rates, respectively. Additionally, Fig. 3(c) shows the microstructure of a plate which was subject to
Cast structure and property variability in gamma titanium aluminides Table 1. (hating coditkm
Plate number l-l 2-1 3-l 4-1 5-l 1-4 2-4 3-4 4-4 5-4
8ml cakdatd
raagea of local solidification times md cdbg
thnes to the eatectd
Nominal plate thickness (mm)
Casting conditions (mold preheat/wrap condition)
Range of local solidification time (s)*
12.7 12.7 12.7 12.7 12.7 3.8 3.8 3.8 3.8 3.8
3 1W/Full wrap 3 1S”C/Partial wrap 8 1S”C/No wrap 1204”C/Partial wrap 1204”C/Full wrap 3 15”C/Full wrap 31 SC/Partial wrap 815”C/No wrap 1204”C/Partial wrap 1204”C/Full wrap
1l-26 lo-25 13-30 2&44 186222 45 5-6 5.8-6.1 79-7.6 22-23
631 teqerrrtrae
Time to 1125”C(s)*~+ 54-39 5639
72-55 191-178 1124-1097 9-8 10-9 12-3-l 1.8 186-18.5 1123-1121
*Ranges indicate times at surface-centerline, respectively. +Time between the non-equilibrium solidus temperature and the binary eutectoid temperature of 1125°C.
(a> Fig. 1. As-cast macrostructure
lmm
Go
following casting with a high cooling rate, Plate 2-4 of Table 1 (a) and low cooling rate, Plate 5-4 of Table 1 (b).
cooling rates during solidification similar to the plate shown in Fig. 3(a), except it was subjected to a HIP and heat treatment cycle of 1093°C 2 h + 1205”C/ 170 MPa/4 h + 1205”C/2 h, followed by rapid cooling. Clearly there are wide variations in the final microstructure. At very slow cooling rates, large, fully lamellar y-a2 grains developed due to the long solidification times and subsequent slow cooling rate through the high temperature (II and Q + y phase fields, Fig. 3(b). From Table 1 it is apparent that some castings remained above the 1125°C eutectoid temperature for times of the order of 20min. This type of fully lamellar structure produced by slow cooling during casting is not easily changed with subsequent HIP and heat treatment cycles. lo Faster cooled plates possess an unusual fully lamellar structure in the as-cast state,‘O which is devoid of the ~2 phase and decomposes during HIP to produce equiaxed gamma grains. This results in a duplex structure, Fig. 3(a) and (c). Pre-HIP heat treatments such as the 1093”C/2 h cycle maximize the volume fraction of gamma grains prior to HIP in the a + y phase field. This results in a more equiaxed gamma grain
structure. Across the range of cooling rates studied, the equiaxed gamma grains were typically smaller in size than the original columnar grain width in the as-cast material, Fig. 3(c). The distribution of the brightly contrasting CQ phase at the grain boundaries (Fig. 3(a) and (c)), as well as the multivariant Q plates within the grains (Fig. 3(c)), is sensitive to heating rates during HIP and heat treatment. *%‘* These variations in casting conditions and HIP/heat treatment cycles have also been shown to influence the texture of the gamma phase in the cast plates. lo Note also that significant dendritic microsegregation remains after HIP and heat treatment, Figs. 3(aXc). Figure 4 shows the results of 16 room temperature tension tests on samples taken from five different plates. All plates, except 4-1, were subjected to a HIP and heat treatment cycle of 1093”C/ 2 h + 1205”C/ 170 MPa/4 h + 1205”C/2 h, resulting in a microstructure similar to that shown in Fig. 3(c). Plate 4-1 was subject to cooling rates during casting that were intermediate to those of Plates 2-1 and 5-1, Table 1. This plate was given a HIP and heat treatment cycle of 126O”C/170 MPa/
Rishel et al.
W
w
2mm
Fig. 2. Distribution of as-cast porosity for low mold preheat and partial wrap, Plate 2-4 of Table 1, (a), and high mold preheat and full wrap, Plate 5-4 of Table 1, (b). Distribution of porosity following HIP and heat treatment for plates cast with low mold preheat and partial wrap (c) and high mold preheat and full wrap, (d). Plates (a) and (c) were cast in the same investment mold, while (b) and (d) were also cast together in a single mold. White regions indicate porosity.
4 h, resulting
in a final microstructure that was predominantly lamellar, intermediate to those shown for Plate 2-1 in Fig. 3(a) and Plate 5-1 in Fig. 3(b). Samples V2-V8 were sheet samples with gage dimensions of 2.7 x 5.8 x 32 mm, machined from successive slices from a single plate. Sample V2 was located near the centerline of the plate, while V8 was from the more rapidly cooled outer edge of the plate. Samples Al-R9 were 4mm diameter cylindrical samples with a gage length of 2 1 mm removed from three different 12.7 mm thick plates at locations that were subject to similar
(cl
Fig. 3. Microstructure of Plate 2-l with local solidification time of 10s and time to 1125°C of approximately 1 min (a). Plate 5-l with local solidification time of 190s and time to 1125°C of approximately 20min (b). Plate 2-l and 5-l were both HIPed at 126O”C/170 MPa/4 h. Microstructure of 12.7mm thick plate with a local solidification time of approximately 20 s and time to 1125°C of approximately 40 s, followed by a HIP and heat treatment cycle of 1093°C 2 h+ 1205”C/170MPa/4 h+ 1205”C/2 h, (c).
cooling rates. Although these plates were not instrumented with thermocouples during casting, based on the structure of the plates and modeling results shown in Table 1, we estimate a local solidification time of the order of 25s at the plate
Cast structure and property variability in gamma
centerlines. Finally, samples 4- l-l-4 l-4 were cylindrical samples of the same geometry as AlR9, removed from successive slices from Plate 4-l. Sample 4- 1- 1 was located near the outer edge of the plate and 4-l-4 near the more slowly cooled centerline region of the plate. Calculated ranges of local solidification times are indicated. From Fig. 4 it is apparent that samples removed from faster cooled regions of the plates exhibited higher ductilities, usually in the range of 1.5-2%. Overall, the samples which came from slower cooled regions near the centers of the plates exhibited lower and more variable ductilities, ranging from 0 to 0.6%. It is interesting to note that the maximum tensile ductilities of the equiaxed near gamma structures compared with the highest ductility of the more lamellar microstructure of Plate 4-1, but only at higher cooling rates. On average the tensile ductilities of faster cooled plates (AlR9) were higher than those of more slowly cooled plates such as 4-1, for a 6xed test specimen geometry. With the Ti47*9Al-2Cr-2Nb heat of material, additional tension tests have been conducted on a range of notched and unnotched specimen configurations.6*‘3 These studies have shown that as the strained volume of material increases, tensile strains to failure decrease on average, but exhibit a lower degree of variability,6 consistent with increasing probability of sampling detrimental structural features within the casting. Similarly, in Fig. 4, the sheet samples V2-V4, with larger gage sections exhibited lower strains to failure, compared to samples Al-R9, for local solidification times of the order of 25 s.
titanium aluminides
633
Detailed studies of fracture origins of the samples shown in Fig. 4 have been conducted. The lowest ductilities were generally associated with the presence of porosity. Figure 5(a) shows the fracture surface of sample 4-l-3, which failed with only about 0.1% plastic strain. Large amounts of porosity are present, in spite of the fact that the plate was subjected to a HIP cycle. Porosity was also observed in a number of the other lower ductility specimens, but generally in lower amounts than shown in Fig. 5. Enhanced local plastic straining within gamma grains, apparent from the intersection of slip bands with the fracture surface, Fig. 5(b), was also commonly observed at the fracture origins of the equiaxed gamma microstructures. In these microstructures, a large equiaxed gamma grain size was associated with lower tensile ductility. Previous in situ straining studies on samples from the 47.9Al heat of material (microstructure shown in Fig. 3(c)) have shown that crack initiation occurs at boundaries between equiaxed gamma grains where slip bands of the most highly strained grains impinge on the boundary.13 Thus large gamma grains are apparently detrimental due to early crack initiation that arises from incompatible deformation with their neighboring grains. Enhanced local plastic straining of multiple gamma grains al& sample surfaces was also occasionally observed, suggesting residual plastic strains from machining. However, electropolishing prior to testing has shown no clear improvement in ductility.
*ate ‘ion
Sample Fig. 4. Tensile ductilities of samples of Ti-47.9Al-2Cr-2Nb
ID
samples removed from five different 12.7 mm thick cast plates.
L. L. Rishel et al.
634
6)
Fig. 5. Failure origin for sample 4-l-3 (0.1% ductility), showing local porosity, (a), and typical failure origin for higher ductility
samples showing intense local straining of gamma grains, (b).
4 DISCUSSION CONCLUDING
AND REMARKS
As mentioned previously, tensile ductility is considered to be an important material property for cast gamma titanium aluminides for several reasons. First, since ductility is related to the degree of damage suffered in a high speed impact,’ it is required for resistance to foreign object damage. Second, plastic strains to failure in the range of l2% permit significant reduction of high elastic stress concentrations at notched features in components.536 Finally, ductility is also needed for resistance to damage during manufacturing operations, such as insertion of blades into turbine disks.@
Cast Ti-48Al-2Cr-2Nb type alloys are known to exhibit variable tensile ductility over the composition range of commercial interest, 4549% Al. For example, ductility has been reported to vary by a factor of approximately 10, with minimum plastic strains to failure ranging from 0.25% for a 455% Al alloy to 2.6% for a 47.4% Al alloy.8 Over this range of composition, studies of fracture surfaces by Austin et al* have revealed some preference for fracture initiation in interdendritic regions. This seems to indicate an important role of cooling rate and/or segregation during solidification. Nevertheless, it is still not clear which aspect of the Al-induced changes in structure most strongly influences tensile ductility. This is hindered by an incomplete understanding of phase
Cast structure and property variability in gamma titanium aluminides
equilibria in these multicomponent alloys, particularly at high temperatures, as well as by a lack of information on the influence of cooling rate on microstructure evolution. Changes in Al have the potential to change the solidification path, microsegregation, proportion and distribution of the y and CQphases, as well as the intrinsic deformation properties of the constituent phases. Such features of cast structures, although undoubtedly important to properties, are difficult to isolate and quantify. It has been shown in this study that substantial variations in microstructure, macrostructure and porosity distribution result from varying cooling rates during casting. Associated with this is a substantial variability in tensile ductility,for a constant composition.
Considering that variations in processing parameters have the same degree of influence on tensile ductility as fairly substantial changes in Al content, it is apparent that a more detailed understanding of the evolution of cast microstructure and its influence on properties is needed. The results of the experiments conducted in this study are significant since they show that variations in casting cooling rates not only influence as-cast structure, but also its evolution. The features most sensitive to cooling rates include porosity distribution, microsegregation, scale of columnar structure and the character of the initial lamellar structure. At high cooling rates, the porosity is interdendritic and arises primarily due to insufficient feeding of solidification shrinkage. At low cooling rates (high mold preheats) the porosity is localized to a greater degree to the surface of the casting, apparently due to reactions with the mold. Surface connected porosity as well as other characteristics of the as-cast structure persist through subsequent HIP and heat treatment cycles typically employed for these materials. For example, as shown in Figs. 1 and 3, slow cooling rates during casting result in coarse columnar structures which remain fully lamellar after HIP and heat treatment, due to the initially slow cooling rates of the casting through the a and CY + y phase fields. Additionally, as shown in Fig. 2, slow cooling rates produce surface connected porosity which is not eliminated by HIP. Conversely, castings cooled at moderate to high rates develop an equiaxed near-gamma microstructure for the same series of HIP and heat treatment cycles. Also, with regard to thermal cycles following casting, heat treatments (or slow heating/cooling rates) at temperatures near the binary eutectoid temperature of 1125°C produce a variety of microstructures. At these temperatures
635
the volume fraction of the y phase is maximized, growth of gamma grains can occur, and changing distributions of the cr2, phase develop due to the y to cytransformation at higher temperatures in subsequent thermal cycles. 11yi2 Finally, it is worth noting that variations in cooling rate and tensile ductility within a single cast plate can also be significant. This further implies that cooling rate variations in components of more complex shape could be substantial, resulting in a wide range of microstructure, macrostructure, porosity distribution and properties that vary with section thickness. The results obtained thus far suggest that higher cooling rates during casting result in higher ductility and reduced variability in this property. Slower cooling rates are detrimental due to the development of greater amounts of surface connected porosity and larger grains in the as-cast state that persist through the standard HIP and heat treatment cycles. However, further work is clearly needed to establish quantitative relationships between cooling rates during casting, porosity distribution, microstructure, macrostructure and the resultant ductility. The results outlined here, where casting cooling rates have been quantified and related to variations in structure evolution and tensile ductility, provide a first step in this direction.
ACKNOWLEDGEMENTS
This research was funded by the AFOSR/PRET Program on Gamma Titanium Aluminides under Grant F49620-95-1-0359. The authors would like to acknowledge D. Larsen, K. Foran, S. Salter and M. Szczerbinski of the Howmet Corporation for providing access to casting facilities and property modeling data, and assistance with casting experiments. We also wish to acknowledge S. McLeod of Pratt and Whitney Aircraft for providing thermophysical property data and D. Banerjee of UES for supplying a copy of ProCAST. The authors are also grateful for discussions with C.M. Austin, T.J. Kelly, K. Muraleedharan, M. De Graef and J. Beuth. REFERENCES Austin, C. M. and Kelly, T. J., in Superalfoys 1996. TMS, Warrendale, PA, USA, 1996, p. 539. Hartfield-Wunsch, S. E., Sperling, A. A., Morrison, R. S., Dowling, W. E. and Allison, J. E., in Gamma Titanium Alurninides, TMS, Warrendale, PA, USA, 1995, p. 53. Jones, P. E. and Porter, W. J. III, Eylon, D. and Colvin, G., in Gamma Titanium Aluminides, TMS, Warrendale, PA, USA, 1995, p. 41.
L. L. Rishel et al 4. Isobe, S. and Nada, T., in Structural Intermetallics 1997. TMS, Warrendale, PA, USA, 1997, p, 427. 5. Wright, P. K., in Structural Intermetallics, TMS, Warrendale, PA, USA, 1993, p, 885. 6. Knaul, D. A., Beuth, J. L. and Milke, J. G., Metall. Mater. Trans, in press. 7. Steif, P. S., Jones, J. W., Harding, T., Rubal, M. P., Gandelsman, V. Z., Biery, N. and Pullock, ‘F. M., in Structural Intermetallics 1997. TMS, Warrendale, PA, USA, 1997, p. 435. 8. Austin, C. M., Kelly, T. J., McAllister, K. G. and Chesnutt, J. C., in Structural IntermetaIIics 1997. TMS, Warrendale, PA, USA, 1997, p. 413. 9. Rishel, L. L., Pollock, T. M., Cramb, A. W. and Larsen, D.
E., in Proceedings of the International Symposium on Liquid Metal Processing and Casting. Vacuum Metallurgy Division of AVS, 1997, p. 214. 10. Muraieedharan, K., Rishel, L. L., De Graef, M., Cramb, A. W., Pollock, T. M. and Gray, G. T. III, in Structural Intermetallics 1997, TMS, Warrendale, PA, USA, 1997, p. 215. Il. Ott, E. A. and Pollock, T. M., Metall. Mater. Trans., 1998,29A, p. 965. 12. Wang, P. and Vasudevan, V. K., in Solid-Solid Phase Transformations, TMS, Warrendale, PA, USA, 1994, p. 303. 13. Pollock, T. M., Mumm, D. R., Muraleedharan, K. and Martin, P. L., Scripta Metall. Mater. 1996, 35, p. 1311.
Intermetallics 6 (1998) 629-636 0 1998 Published by Elsevier Science Limited Printed in Great Britain. All rights reserved 09669795/98/$-+x front matter
Cast structure and property variability in gamma titanium aluminides L. L. Rishel, N. E. Biery, R. Raban, V. Z. Gandelsman, T. M. Pollock* & A. W. Cramb Department of Materials Science and Engineering, Carnegie Mellon University, Pittsburgh, PA 15213, USA
(Received 20 April 1998; accepted 27 April 1998)
Aspects of cast structure that influence mechanical property variability, and in particular tensile ductility, have been studied in a Ti-Q8Al-2Cr-2Nb titanium aluminide alloy. Macrostructure, porosity, microstructure and tensile properties have been characterized over a range of casting processing conditions. Local solidification times and subsequent solid-state cooling rates during casting have been characterized via local thermal measurements in combination with solidification modeling. Large variations in cooling rate during casting dramatically influence the initial cast structure as well as the distribution of defects such as porosity. Variations in as-cast structure persist through subsequent thermal and hot isostatic pressing cycles and contribute to the variability of tensile ductility. The current understanding of the relationship of tensile ductility to processinginduced changes in structure will be briefly discussed. 0 1998 Published by Elsevier Science Limited. All rights reserved Key words: A. titanium aluminides, based on TiAl, B. ductility, mechanical properties at ambient temperature, C. casting, D. microstructure, E. microscopy.
1 INTRODUCTION
there is a need for a better understanding of the relationship between casting parameters, the range of cooling rates that can be achieved during solidification, and their effect on the structure and properties of the cast material. Recent experiments9yi0 have measured cooling rates in the mold during investment casting of gamma titanium aluminide plates. Here, we briefly review the results of those experiments, along with results of solidification modeling studies of investment cast plates. Variations in cast structure and tensile ductility, which are influenced by changes in cooling rate during casting, are shown. Issues which appear to be in need of further study are identified.
Gamma titanium aluminides are among the most promising intermetallics for high temperature structural applications. These materials may be produced via either cast or wrought processing techniques. However, casting processes have been of greater interest for components such as turbine blades and automobile exhaust valves, where component cost is a barrier to implementation.14 As with any casting, it is expected that the final properties of the material will be sensitive to process parameters which influence as-cast microstructure or casting defects such as porosity. This is particularly critical for gamma titanium aluminides, where some degree of tensile ductility is important for tolerance to foreign object damage and for the design of components which contain stress concentrations.S8 Control of the casting process is constrained by the need for cold crucible processes that maintain the chemistry of these highly reactive alloys during melting. Given these constraints,
2 INVESTMENT CASTING OF GAMMA TITANIUM ALUMINIDE PLATES Cooling rates during casting were measured in a series of experiments designed to cover a wide range of variations in investment casting process parameters. The experiments utilized investment molds that contained four rectangular plates measuring 10.2x 10.2cm with the following thicknesses: 12.7, 3.8, 6.3 and 3.8 mm. The composition
*To whom correspondence should be addressed. Fax: 001 412268-7596; e-mail: tplm + @andrew.cmu.edu 629
630
L. L. Rishel et al.
of the material studied in the cooling rate experiments was Ti48*OAl-2Cr-2Nb-O.O 15Fe-O. 180 (at%, avg. composition based on measurements on five plates). The casting process consisted of vacuum arc remelting (VAR) followed by pouring into investment molds. Since variations in the superheat of the molten metal cannot easily be achieved in such cold crucible processes, control of the solidification process for a given alloy must be achieved through the control of heat extraction from the investment mold. Accordingly, a range of mold preheat temperatures and mold insulation conditions were utilized. Mold preheat temperatures were 315, 815 or 1204”C, while molds were wrapped with varied amounts of Kaowool insulation. Thermocouples were attached in the mold near the surface of each of the four plates to obtain thermal data during mold preheating, pouring and solidification. The measured variations in mold temperature for each of the experiments were input as boundary conditions into the UES ProCAST simulation package. Given the arrangement of the plates in the cluster, the heat extraction at the plate centers was essentially one dimensional. Accordingly, one dimensional heat transfer analyses of the process were performed to obtain estimates of solidification times and subsequent cooling rates through the plate thickness.‘O Following casting, the structure and porosity distribution in these plates was characterized in the as-cast state and after hot isostatic pressing and heat treatment. Additional details on the experimental set-up and modeling are given elsewhere.9*t0 A second set of 12.7 mm plates from a heat of material with essentially the same composition, Ti-47.9A l-2Cr-2Nb-O.007Fe-O. 140, were also utilized along with the plates from the cooling rate study- to characterize the variability of tensile ductility.
3 COOLING
RATES, MICROSTRUCTURE DEVELOPMENT AND DUCTILITY
The results of solidification simulations for two different plate thicknesses with five different combinations of casting conditions are shown in Table 1. Local solidification times varied from the surface to the centerline of the plates; the calculeited range is given for each set of casting conditions. At the centerlines of the plates, local solidifitition times varied by a factor of more than 25 X, ‘from about 5 s for the 3.8 mm plates solidified with a mold preheat of 3 15°C to 140 s for the
12.7 mm thick plate cast with a fully insulated mold and a preheat of 1204°C. It is apparent that, in spite of superheat limitations, a fairly wide range of solidification times can be achieved, particularly with higher mold preheats. Also shown in Table 1 is the total time at temperatures above the eutectoid temperature (approximately 1125°C in the binary phase diagram) following solidification. This is an important parameter, since it influences the character of the initial microstructure of the casting, as will be discussed further. The wide range of cooling rates during solidification, Table 1, result in significant variation in ascast structure and porosity distribution. Figure 1 shows that the as-cast macrostructures at high and low cooling rates exhibit markedly different columnar grain widths. Figure 2 shows the distribution of porosity in the 3.8 mm thick as-cast plates. At higher cooling rates, the bulk of the porosity is present at the centerline of the plate. Conversely, at lower cooling rates (high degree of mold preheat and wrap), the porosity is preferentially located in near-surface regions, apparently due to reactions with the mold material. Since these materials are typically subjected to hot isostatic pressing (HIP) following casting, the distribution of porosity following a HIP cycle of 1260°C 170 MPa/4 h is also shown in Fig. 2(c) and (d) for comparison. The porosity remaining after HIP indicates the degree to which it was connected to the surface in the as-cast state. Near-surface porosity associated with very high mold preheats, Fig. 2(b), is retained through HIP’ing. Centerline porosity of the type shown in Fig. 2(a) is typically largely eliminated by HIP, Fig. 2(c). However, due to the fact that most of the porosity is interdendritic and highly interconnected, small pockets of micro-porosity often persist through the HIP cycle. The dendrite ,arm spacings, grain size and grain morphology in near-surface regions are sensitive to mold preh&ttP and this is likely to strongly influence the degree to which the porosity connects to the surface. The widely varying as-cast microstructures observed over the range of cooling rates studied responded differently to thermal cycles employed during HIP and accompanying heat treatments. The microstructures of Plates 2-l and 5-l following a HIP cycle of 126O”C/l70 MPa/4 h are shown in Fig. 3(a) and (b). These plates were cast under the same conditions as the plates shown in Figs. 1 and 2, with relatively high and low cooling rates, respectively. Additionally, Fig. 3(c) shows the microstructure of a plate which was subject to
Cast structure and property variability in gamma titanium aluminides Table 1. (hating coditkm
Plate number l-l 2-1 3-l 4-1 5-l 1-4 2-4 3-4 4-4 5-4
8ml cakdatd
raagea of local solidification times md cdbg
thnes to the eatectd
Nominal plate thickness (mm)
Casting conditions (mold preheat/wrap condition)
Range of local solidification time (s)*
12.7 12.7 12.7 12.7 12.7 3.8 3.8 3.8 3.8 3.8
3 1W/Full wrap 3 1S”C/Partial wrap 8 1S”C/No wrap 1204”C/Partial wrap 1204”C/Full wrap 3 15”C/Full wrap 31 SC/Partial wrap 815”C/No wrap 1204”C/Partial wrap 1204”C/Full wrap
1l-26 lo-25 13-30 2&44 186222 45 5-6 5.8-6.1 79-7.6 22-23
631 teqerrrtrae
Time to 1125”C(s)*~+ 54-39 5639
72-55 191-178 1124-1097 9-8 10-9 12-3-l 1.8 186-18.5 1123-1121
*Ranges indicate times at surface-centerline, respectively. +Time between the non-equilibrium solidus temperature and the binary eutectoid temperature of 1125°C.
(a> Fig. 1. As-cast macrostructure
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following casting with a high cooling rate, Plate 2-4 of Table 1 (a) and low cooling rate, Plate 5-4 of Table 1 (b).
cooling rates during solidification similar to the plate shown in Fig. 3(a), except it was subjected to a HIP and heat treatment cycle of 1093°C 2 h + 1205”C/ 170 MPa/4 h + 1205”C/2 h, followed by rapid cooling. Clearly there are wide variations in the final microstructure. At very slow cooling rates, large, fully lamellar y-a2 grains developed due to the long solidification times and subsequent slow cooling rate through the high temperature (II and Q + y phase fields, Fig. 3(b). From Table 1 it is apparent that some castings remained above the 1125°C eutectoid temperature for times of the order of 20min. This type of fully lamellar structure produced by slow cooling during casting is not easily changed with subsequent HIP and heat treatment cycles. lo Faster cooled plates possess an unusual fully lamellar structure in the as-cast state,‘O which is devoid of the ~2 phase and decomposes during HIP to produce equiaxed gamma grains. This results in a duplex structure, Fig. 3(a) and (c). Pre-HIP heat treatments such as the 1093”C/2 h cycle maximize the volume fraction of gamma grains prior to HIP in the a + y phase field. This results in a more equiaxed gamma grain
structure. Across the range of cooling rates studied, the equiaxed gamma grains were typically smaller in size than the original columnar grain width in the as-cast material, Fig. 3(c). The distribution of the brightly contrasting CQ phase at the grain boundaries (Fig. 3(a) and (c)), as well as the multivariant Q plates within the grains (Fig. 3(c)), is sensitive to heating rates during HIP and heat treatment. *%‘* These variations in casting conditions and HIP/heat treatment cycles have also been shown to influence the texture of the gamma phase in the cast plates. lo Note also that significant dendritic microsegregation remains after HIP and heat treatment, Figs. 3(aXc). Figure 4 shows the results of 16 room temperature tension tests on samples taken from five different plates. All plates, except 4-1, were subjected to a HIP and heat treatment cycle of 1093”C/ 2 h + 1205”C/ 170 MPa/4 h + 1205”C/2 h, resulting in a microstructure similar to that shown in Fig. 3(c). Plate 4-1 was subject to cooling rates during casting that were intermediate to those of Plates 2-1 and 5-1, Table 1. This plate was given a HIP and heat treatment cycle of 126O”C/170 MPa/
Rishel et al.
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Fig. 2. Distribution of as-cast porosity for low mold preheat and partial wrap, Plate 2-4 of Table 1, (a), and high mold preheat and full wrap, Plate 5-4 of Table 1, (b). Distribution of porosity following HIP and heat treatment for plates cast with low mold preheat and partial wrap (c) and high mold preheat and full wrap, (d). Plates (a) and (c) were cast in the same investment mold, while (b) and (d) were also cast together in a single mold. White regions indicate porosity.
4 h, resulting
in a final microstructure that was predominantly lamellar, intermediate to those shown for Plate 2-1 in Fig. 3(a) and Plate 5-1 in Fig. 3(b). Samples V2-V8 were sheet samples with gage dimensions of 2.7 x 5.8 x 32 mm, machined from successive slices from a single plate. Sample V2 was located near the centerline of the plate, while V8 was from the more rapidly cooled outer edge of the plate. Samples Al-R9 were 4mm diameter cylindrical samples with a gage length of 2 1 mm removed from three different 12.7 mm thick plates at locations that were subject to similar
(cl
Fig. 3. Microstructure of Plate 2-l with local solidification time of 10s and time to 1125°C of approximately 1 min (a). Plate 5-l with local solidification time of 190s and time to 1125°C of approximately 20min (b). Plate 2-l and 5-l were both HIPed at 126O”C/170 MPa/4 h. Microstructure of 12.7mm thick plate with a local solidification time of approximately 20 s and time to 1125°C of approximately 40 s, followed by a HIP and heat treatment cycle of 1093°C 2 h+ 1205”C/170MPa/4 h+ 1205”C/2 h, (c).
cooling rates. Although these plates were not instrumented with thermocouples during casting, based on the structure of the plates and modeling results shown in Table 1, we estimate a local solidification time of the order of 25s at the plate
Cast structure and property variability in gamma
centerlines. Finally, samples 4- l-l-4 l-4 were cylindrical samples of the same geometry as AlR9, removed from successive slices from Plate 4-l. Sample 4- 1- 1 was located near the outer edge of the plate and 4-l-4 near the more slowly cooled centerline region of the plate. Calculated ranges of local solidification times are indicated. From Fig. 4 it is apparent that samples removed from faster cooled regions of the plates exhibited higher ductilities, usually in the range of 1.5-2%. Overall, the samples which came from slower cooled regions near the centers of the plates exhibited lower and more variable ductilities, ranging from 0 to 0.6%. It is interesting to note that the maximum tensile ductilities of the equiaxed near gamma structures compared with the highest ductility of the more lamellar microstructure of Plate 4-1, but only at higher cooling rates. On average the tensile ductilities of faster cooled plates (AlR9) were higher than those of more slowly cooled plates such as 4-1, for a 6xed test specimen geometry. With the Ti47*9Al-2Cr-2Nb heat of material, additional tension tests have been conducted on a range of notched and unnotched specimen configurations.6*‘3 These studies have shown that as the strained volume of material increases, tensile strains to failure decrease on average, but exhibit a lower degree of variability,6 consistent with increasing probability of sampling detrimental structural features within the casting. Similarly, in Fig. 4, the sheet samples V2-V4, with larger gage sections exhibited lower strains to failure, compared to samples Al-R9, for local solidification times of the order of 25 s.
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Detailed studies of fracture origins of the samples shown in Fig. 4 have been conducted. The lowest ductilities were generally associated with the presence of porosity. Figure 5(a) shows the fracture surface of sample 4-l-3, which failed with only about 0.1% plastic strain. Large amounts of porosity are present, in spite of the fact that the plate was subjected to a HIP cycle. Porosity was also observed in a number of the other lower ductility specimens, but generally in lower amounts than shown in Fig. 5. Enhanced local plastic straining within gamma grains, apparent from the intersection of slip bands with the fracture surface, Fig. 5(b), was also commonly observed at the fracture origins of the equiaxed gamma microstructures. In these microstructures, a large equiaxed gamma grain size was associated with lower tensile ductility. Previous in situ straining studies on samples from the 47.9Al heat of material (microstructure shown in Fig. 3(c)) have shown that crack initiation occurs at boundaries between equiaxed gamma grains where slip bands of the most highly strained grains impinge on the boundary.13 Thus large gamma grains are apparently detrimental due to early crack initiation that arises from incompatible deformation with their neighboring grains. Enhanced local plastic straining of multiple gamma grains al& sample surfaces was also occasionally observed, suggesting residual plastic strains from machining. However, electropolishing prior to testing has shown no clear improvement in ductility.
*ate ‘ion
Sample Fig. 4. Tensile ductilities of samples of Ti-47.9Al-2Cr-2Nb
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samples removed from five different 12.7 mm thick cast plates.
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6)
Fig. 5. Failure origin for sample 4-l-3 (0.1% ductility), showing local porosity, (a), and typical failure origin for higher ductility
samples showing intense local straining of gamma grains, (b).
4 DISCUSSION CONCLUDING
AND REMARKS
As mentioned previously, tensile ductility is considered to be an important material property for cast gamma titanium aluminides for several reasons. First, since ductility is related to the degree of damage suffered in a high speed impact,’ it is required for resistance to foreign object damage. Second, plastic strains to failure in the range of l2% permit significant reduction of high elastic stress concentrations at notched features in components.536 Finally, ductility is also needed for resistance to damage during manufacturing operations, such as insertion of blades into turbine disks.@
Cast Ti-48Al-2Cr-2Nb type alloys are known to exhibit variable tensile ductility over the composition range of commercial interest, 4549% Al. For example, ductility has been reported to vary by a factor of approximately 10, with minimum plastic strains to failure ranging from 0.25% for a 455% Al alloy to 2.6% for a 47.4% Al alloy.8 Over this range of composition, studies of fracture surfaces by Austin et al* have revealed some preference for fracture initiation in interdendritic regions. This seems to indicate an important role of cooling rate and/or segregation during solidification. Nevertheless, it is still not clear which aspect of the Al-induced changes in structure most strongly influences tensile ductility. This is hindered by an incomplete understanding of phase
Cast structure and property variability in gamma titanium aluminides
equilibria in these multicomponent alloys, particularly at high temperatures, as well as by a lack of information on the influence of cooling rate on microstructure evolution. Changes in Al have the potential to change the solidification path, microsegregation, proportion and distribution of the y and CQphases, as well as the intrinsic deformation properties of the constituent phases. Such features of cast structures, although undoubtedly important to properties, are difficult to isolate and quantify. It has been shown in this study that substantial variations in microstructure, macrostructure and porosity distribution result from varying cooling rates during casting. Associated with this is a substantial variability in tensile ductility,for a constant composition.
Considering that variations in processing parameters have the same degree of influence on tensile ductility as fairly substantial changes in Al content, it is apparent that a more detailed understanding of the evolution of cast microstructure and its influence on properties is needed. The results of the experiments conducted in this study are significant since they show that variations in casting cooling rates not only influence as-cast structure, but also its evolution. The features most sensitive to cooling rates include porosity distribution, microsegregation, scale of columnar structure and the character of the initial lamellar structure. At high cooling rates, the porosity is interdendritic and arises primarily due to insufficient feeding of solidification shrinkage. At low cooling rates (high mold preheats) the porosity is localized to a greater degree to the surface of the casting, apparently due to reactions with the mold. Surface connected porosity as well as other characteristics of the as-cast structure persist through subsequent HIP and heat treatment cycles typically employed for these materials. For example, as shown in Figs. 1 and 3, slow cooling rates during casting result in coarse columnar structures which remain fully lamellar after HIP and heat treatment, due to the initially slow cooling rates of the casting through the a and CY + y phase fields. Additionally, as shown in Fig. 2, slow cooling rates produce surface connected porosity which is not eliminated by HIP. Conversely, castings cooled at moderate to high rates develop an equiaxed near-gamma microstructure for the same series of HIP and heat treatment cycles. Also, with regard to thermal cycles following casting, heat treatments (or slow heating/cooling rates) at temperatures near the binary eutectoid temperature of 1125°C produce a variety of microstructures. At these temperatures
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the volume fraction of the y phase is maximized, growth of gamma grains can occur, and changing distributions of the cr2, phase develop due to the y to cytransformation at higher temperatures in subsequent thermal cycles. 11yi2 Finally, it is worth noting that variations in cooling rate and tensile ductility within a single cast plate can also be significant. This further implies that cooling rate variations in components of more complex shape could be substantial, resulting in a wide range of microstructure, macrostructure, porosity distribution and properties that vary with section thickness. The results obtained thus far suggest that higher cooling rates during casting result in higher ductility and reduced variability in this property. Slower cooling rates are detrimental due to the development of greater amounts of surface connected porosity and larger grains in the as-cast state that persist through the standard HIP and heat treatment cycles. However, further work is clearly needed to establish quantitative relationships between cooling rates during casting, porosity distribution, microstructure, macrostructure and the resultant ductility. The results outlined here, where casting cooling rates have been quantified and related to variations in structure evolution and tensile ductility, provide a first step in this direction.
ACKNOWLEDGEMENTS
This research was funded by the AFOSR/PRET Program on Gamma Titanium Aluminides under Grant F49620-95-1-0359. The authors would like to acknowledge D. Larsen, K. Foran, S. Salter and M. Szczerbinski of the Howmet Corporation for providing access to casting facilities and property modeling data, and assistance with casting experiments. We also wish to acknowledge S. McLeod of Pratt and Whitney Aircraft for providing thermophysical property data and D. Banerjee of UES for supplying a copy of ProCAST. The authors are also grateful for discussions with C.M. Austin, T.J. Kelly, K. Muraleedharan, M. De Graef and J. Beuth. REFERENCES Austin, C. M. and Kelly, T. J., in Superalfoys 1996. TMS, Warrendale, PA, USA, 1996, p. 539. Hartfield-Wunsch, S. E., Sperling, A. A., Morrison, R. S., Dowling, W. E. and Allison, J. E., in Gamma Titanium Alurninides, TMS, Warrendale, PA, USA, 1995, p. 53. Jones, P. E. and Porter, W. J. III, Eylon, D. and Colvin, G., in Gamma Titanium Aluminides, TMS, Warrendale, PA, USA, 1995, p. 41.
L. L. Rishel et al 4. Isobe, S. and Nada, T., in Structural Intermetallics 1997. TMS, Warrendale, PA, USA, 1997, p, 427. 5. Wright, P. K., in Structural Intermetallics, TMS, Warrendale, PA, USA, 1993, p, 885. 6. Knaul, D. A., Beuth, J. L. and Milke, J. G., Metall. Mater. Trans, in press. 7. Steif, P. S., Jones, J. W., Harding, T., Rubal, M. P., Gandelsman, V. Z., Biery, N. and Pullock, ‘F. M., in Structural Intermetallics 1997. TMS, Warrendale, PA, USA, 1997, p. 435. 8. Austin, C. M., Kelly, T. J., McAllister, K. G. and Chesnutt, J. C., in Structural IntermetaIIics 1997. TMS, Warrendale, PA, USA, 1997, p. 413. 9. Rishel, L. L., Pollock, T. M., Cramb, A. W. and Larsen, D.
E., in Proceedings of the International Symposium on Liquid Metal Processing and Casting. Vacuum Metallurgy Division of AVS, 1997, p. 214. 10. Muraieedharan, K., Rishel, L. L., De Graef, M., Cramb, A. W., Pollock, T. M. and Gray, G. T. III, in Structural Intermetallics 1997, TMS, Warrendale, PA, USA, 1997, p. 215. Il. Ott, E. A. and Pollock, T. M., Metall. Mater. Trans., 1998,29A, p. 965. 12. Wang, P. and Vasudevan, V. K., in Solid-Solid Phase Transformations, TMS, Warrendale, PA, USA, 1994, p. 303. 13. Pollock, T. M., Mumm, D. R., Muraleedharan, K. and Martin, P. L., Scripta Metall. Mater. 1996, 35, p. 1311.