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Microstructure and properties of a TiAl alloy prepared by mechanical milling and subsequent reactive sintering
Materials Science and Engineering A 403 (2005) 186–190
Microstructure and properties of a TiAl alloy prepared by mechanical milling and subsequent reactive sintering Fang Wenbin ∗ , Hu Lianxi, He Wenxiong, Wang Erde, Li Xiaoqing School of Material Science and Engineering, Harbin Institute of Technology, Harbin 150001, PR China Accepted 29 April 2005
Abstract A TiAl-based alloy with nominal composition of Ti–36 wt.% Al was prepared by a mechanical milling and reactive sintering route consisting of three basic steps: (1) mechanical milling of blended elemental Ti and Al powder mixture, (2) densification and consolidation of as-milled Ti/Al composite powder compact via warm extrusion, and (3) subsequent reactive sintering of as-extruded powder billet. The microstructure of the as-sintered TiAl alloy samples was observed by optical microscopy (OM) and transmission electron microscopy (TEM), respectively. The mechanical properties were measured by compression tests performed at room and elevated temperatures up to 800 ◦ C. The results show that mechanical milling before reactive sintering is very effective in microstructure refining of reactively synthesized TiAl alloy. The TiAl alloy sintered from 1 h milled Ti/Al composite powders presented a homogeneous ␥/(␣2 + ␥) duplex microstructure with very fine grains of 5–10 m and good mechanical properties. On the basis of microstructure observation and compressive stress–strain relationship, the deformation mechanism at different temperature was discussed. © 2005 Elsevier B.V. All rights reserved. Keywords: Mechanical milling; Reactive sintering; TiAl alloy; Microstructure and mechanical properties
1. Introduction TiAl-based alloys are strong candidate materials for hightemperature structural applications where weight reduction is of great importance due to their good thermal stability up to 600–800 ◦ C, low densities of 3.9–4.1 g/cm3 , and high strength retention at elevated temperatures [1–3]. Below 700 ◦ C, however, TiAl alloys have intrinsically very poor plasticity and workability, with their elongation in the range of 0–4%. At high temperature above 1100 ◦ C, despite plasticity improvement, they are still rather difficult to be shaped because their flow stress is usually as high as 200 MPa even with such a low strain rate as 1 × 10−3 S−1 [4–6]. Indeed, the poor workability is one of the major obstacles for their application. Therefore, it is of great significance to develop effective techniques for preparation and forming of TiAl alloys. So far, many researches have been focused on this topic in
∗
Corresponding author. Tel.: +86 451 8641 8733; fax: +86 451 5362 7384. E-mail address: [email protected] (F. Wenbin).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.04.049
an effort to accelerate the application of TiAl alloys as light structural materials for high-temperature services [7–19]. Recently, elemental powder metallurgy route has been gaining more and more attention because near-net shape TiAl alloy products can be fabricated by the consolidation and forming of blended Ti and Al elemental powders followed by a subsequent reactive sintering process [7–10]. However, due to the large difference between partial diffusion coefficients of Ti and Al and the immobility of Ti atoms for Ti/Al solid state reaction, the synthesis of TiAl alloy via reactive sintering takes the mechanism that Al atoms move into Ti lattice, thus leading to the formation of Kirkendall diffusion pores [20–22]. Although hot isostatic pressing (HIP) has been reported to be effective in eliminating porosity of reactively sintered TiAl alloys, the high cost and low production efficiency make it not suitable for commercial use. It has been reported that the microstructure of reactively sintered TiAl-based alloys can be refined and the porosity reduced by refining the reactant Ti and Al phases through cold deformation, such as extrusion and rolling, of the Ti and Al powder mixture billet prior to reactive sintering [7–9]. In
F. Wenbin et al. / Materials Science and Engineering A 403 (2005) 186–190
an effort to minimize the porosity due to Kirkendall diffusion and to improve the microstructure of the sintered TiAl-based alloys, we have recently proposed to use mechanical milling to refine the Ti and Al constituents prior to reactive sintering. In the present paper, we report the fabrication of a fully densified TiAl alloy by a cost effective elemental powder metallurgy route consisting of mechanical milling, warm extrusion, and subsequent reactive sintering.
187
cal microscopy (OM) and transmission electron microscopy (TEM), respectively. The mechanical properties of the assintered TiAl alloy samples were measured by isothermal compression tests at room and elevated temperatures. For all tests, cylindrical samples with a dimension of Ø 4 mm × 5 mm were used, and the strain rate for compression was 1.67 × 10−3 S−1 .
3. Results and discussion 2. Experimental 3.1. Microstructure Ti/Al composite powders with a nominal composition of Ti–36 wt.% Al were prepared by mechanical milling of the blended powder mixture of elemental titanium powders (99.5%, <50 m) and aluminum powders (99.9%, <80 m). The milling process was carried out under protection of pure argon using an attritor with a water cooling system. The milling vial and balls were made of stainless steel and GCr15 hardened steel, respectively. All the balls used were of the same size of 6 mm in diameter. The ball:powder weight ratio was 20:1, and the attritor shaft rotation was 400 rev/min. For each batch of milling, the weight of powders processed was 0.5 kg. The as-milled Ti/Al composite powders were cold compacted under a specific pressure of 1000 MPa, and the compacts were then consolidated into dense powder billets by warm extrusion at 250 ◦ C with an extrusion ratio of 16:1. The reactive sintering of the as-extruded Ti/Al composite powder billets was performed in a sintering furnace under vacuum better than 6.67 × 10−3 Pa. The sintering was performed following a two-step procedure, i.e., pre-sintering at 630 ◦ C for 2 h and subsequent final sintering at 1250 ◦ C for 4 h. The microstructure of various samples was observed by opti-
Fig. 1 shows the microstructure of TiAl alloy samples obtained by reactive sintering of powder billets consolidated from Ti/Al composite powders prepared by mechanical milling for various times up to 3 h. Though all as-sintered TiAl alloy samples were characterized by a ␥/(␣2 + ␥) duplex microstructure, differences in grain size and ␥:(␣2 + ␥) phase ratio were obvious. With increasing milling time, the assintered TiAl alloy presented a finer microstructure, and also a higher volume fraction of ␥ phase, suggesting a higher nucleation rate for both ␣ and ␥ phases, and accelerated diffusion kinetics to phase equilibrium, during the reactive sintering process. The high nucleation rate and accelerated diffusion kinetics can be attributed to the refining of the Ti/Al alternating lamellar structure with increasing milling. Indeed, our previous work [18,19] revealed that the lamellar spacing of the reactant Ti and Al phases in the mechanically synthesized elemental Ti/Al composite powders was reduced to less than 0.5 m after milling only for 1 h, and prolonged milling led to formation of nanocrystalline or partially amorphous Ti/Al composite powders with still smaller lamellar spacing. However, the compactibility of the Ti/Al composite
Fig. 1. Microstructure of TiAl alloy samples sintered from Ti/Al composite powders milled for various times: (a) 0 h; (b) 0.5 h; (c) 1 h; (d) 3 h.
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powders became deteriorated with prolonged milling. Therefore, it is very important to optimize the milling parameters to obtain as-milled Ti/Al composite powders with adequately refined Ti/Al lamellar spacing as well as acceptable compactibility. In our opinion, it appears that 1 h milling under conditions specified in the present study is the optimal milling process. As shown in Fig. 1(c), the microstructure of the TiAl alloy sample reactively sintered from extruded billet of 1 h milled Ti/Al composite powders is quite fine and homogenous, with grain size being 5–10 m on average. In comparison, the microstructure of the TiAl alloy sample reactively sintered from starting elemental Ti and Al powders is rather heterogeneous, with some grains being about 10 m while others as coarse as more than 50 m in size (Fig. 1(a)). This suggests that proper milling processing of Ti and Al blended powders before reactive sintering results in not only effective refining of grain size but also significant improvement of microstructure homogeneity of the as-sintered TiAl alloy. Fig. 2 shows a comparison in the lamellar structure of the (␣2 + ␥) phase between two TiAl alloy samples which were sintered from blended Ti/Al elemental powders and 1 h milled Ti/Al composite powders, respectively. It was observed that the average thickness of the ␣2 /␥ lamellae in the TiAl alloy sintered from blended elemental Ti/Al powders was about 250 nm, as shown in Fig. 2(a). In comparison, the ␣2 and ␥ lamellae formed in the TiAl alloy by reactive sintering of the 1 h milled Ti/Al composite powders were much finer, with their average thickness being about only one-third of those in the TiAl alloy sintered from the elemental Ti and Al powder mixture, as shown in Fig. 2(b). This can be explained by the fact that, when as-milled Ti/Al composite powders were used instead of blended elemental Ti/Al powders, the grains of the ␣ phase formed during reactive sintering at temperature above ␣ → ␣2 + ␥ eutectoid transition were much smaller due to a much higher nucleation rate, which in turn resulted in the refinement of the final ␣2 /␥ lamellae structure formed by the
Fig. 3. Compressive stress–strain curves of TiAl alloy sintered from 1 h milled Ti/Al composite powders.
eutectoid transformation of ␣ to ␣2 + ␥ during cooling after sintering. 3.2. Properties It is well known that one major advantage of using the elemental powder metallurgy route for synthesizing of TiAl alloy is the good formability of the precursory material prior to reactive sintering. Indeed, it was found in our present study that the Ti/Al composite powders milled up to 3 h could be readily consolidated to nearly full density by warm hydrostatic extrusion at 250 ◦ C with a moderate extrusion ratio of 16:1 and a specific extrusion pressure about 1000 MPa. By reactive pre-sintering at 630 ◦ C for 2 h, and then final sintering at 1250 ◦ C for 4 h, near fully densified TiAl alloy samples, with a density of about 3.92 g/cm3 , were obtained. Since, as discussed earlier, the TiAl alloy sintered from the 1 h milled Ti/Al composite powders presented the best microstructure, its mechanical properties were further investigated.
Fig. 2. Comparison in (␣2 + ␥) lamellar structure between two TiAl alloy samples sintered from: (a) blended Ti/Al elemental powders and (b) 1 h milled Ti/Al composite powders.
F. Wenbin et al. / Materials Science and Engineering A 403 (2005) 186–190
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Fig. 4. TEM images of the TiAl alloy deformed to a compressive strain of 0.2 at 800 ◦ C showing: (a) recrystallization and twinning in ␥ phase region and (b) recrystallization in (␣2 + ␥) phase region.
Fig. 3 shows the compressive stress–strain curves at room and elevated temperatures up to 800 ◦ C of the TiAl alloy sintered from 1 h milled Ti/Al composite powders. At room temperature, its yield strength is about 560 MPa and the compressive strain to failure is below 0.05. Unlike conventional metal materials, the yield strength of the TiAl alloy increases with increasing temperature in a certain temperature range, and achieves a maximum of about 820 MPa at 600 ◦ C. Then, with further increase of the temperature, the yield strength decreases. On the other hand, the compressive strain before failure increases steadily with increasing temperature. Indeed, the compressive strain to failure at 800 ◦ C is as high as 0.25, suggesting significant improvement of plasticity or appearance of superplastic deformation behavior at this temperature. The mechanical behavior shown in Fig. 3 can be interpreted according to the deformation mechanism of TiAl alloy. As suggested by Chen and Lin [23], for TiAl alloys with duplex microstructure, the deformation behavior below 600 ◦ C is dominated by the movement and pinning of the 1¯ 1 0 screw dislocations. At relatively higher temperature, the 1¯ 1 0 screw dislocations tend to shift on to the {1 1 1} plane by cross-slipping, and then back to the starting plane, forming non-spiral type kinks which pin dislocation movement. When the temperature is enhanced from room temperature to 600 ◦ C, the density of pinning sites steadily increases, while the distance between pinning points decreases. Besides, ¯ and 1/21 1 2] ¯ type tend to super dislocations of 1 0 1] decompose through thermal activation at relatively higher temperature to form dislocation dipoles which also present pinning effect. Therefore, a higher shear stress is needed for dislocation movement at a higher temperature. This leads to the increase of the yield strength with increasing temperature up to 600 ◦ C. On the other hand, at temperature above 600 ◦ C, the mobility of both the 1/21 1 0] dislocations and the 1/61 1 2] twinning dislocations is significantly improved due to thermal activation. Meanwhile, the pinning effect of the kinks by cross-slipping is reduced, and grain boundary
sliding or grain rotating is also possible as alternative deformation modes. Therefore, the yield strength above 600 ◦ C decreases with increasing temperature. Fig. 4 shows TEM images of the TiAl alloy deformed at 800 ◦ C to a compressive strain of 0.2. Evidence of dynamic recrystallization in both ␥ and (␣2 + ␥) phase regions during deformation was obvious, as shown in Fig. 4(a and b), respectively. Besides, evidence of deformation twinning was also observed (Fig. 4(a)). The dynamic recrystallization during compression can be attributed to two major factors: (1) the high deformation temperature, 800 ◦ C (1073 K), which is well above 0.55Tm , and (2) the low strain rate, 1.67 × 10−3 S−1 , which is favorable for recrystallization nucleation. It is believed that dynamic recrystallization is the major softening mechanism of TiAl alloy during hightemperature deformation with low strain rate, which accounts well for the steady flow stress after being compressed at 800 ◦ C to a compressive strain more than 0.05, as shown in Fig. 3. Due to the very fine microstructure of the assintered TiAl alloy, it is quite possible that grain boundary sliding and grain rotating occur during deformation at such a temperature as 800 ◦ C. Therefore, both dynamic recrystallization and alternative deformation modes including grain boundary sliding, grain rotating, and twinning at high temperature should be responsible for the significant improvement of plasticity of TiAl alloy for compression at 800 ◦ C.
4. Conclusions (1) Mechanical milling of blended elemental Ti and Al powders before reactive sintering is very effective in microstructure homogeneity improvement and grain size refining of reactively synthesized TiAl alloy prepared by powder metallurgy route using elemental powders as starting material. In particular, the TiAl alloy sintered from 1 h milled Ti/Al composite powders presented
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a homogeneous ␥/(␣2 + ␥) duplex microstructure with very fine grains of 5–10 m in size. (2) At room temperature, the compressive yield strength of the TiAl alloy sintered from 1 h milled Ti/Al composite powders is about 560 MPa. With the increase of temperature, the yield strength at first increases and then decreases, with a maximum value of about 820 MPa achieved at 600 ◦ C. At temperature above 600 ◦ C, the plasticity of the TiAl alloy is improved significantly, with a failure strain as high as 0.25 for compression at 800 ◦ C. (3) Dynamic recrystallization is the major softening mechanism of TiAl alloy during high-temperature deformation with low strain rate. Both dynamic recrystallization and possible alternative deformation modes, such as grain boundary sliding, grain rotating, and twinning, are responsible for the significant improvement of plasticity of TiAl alloy at 800 ◦ C.
[4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17]
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K. Equo, M. Yoshinao, Light Met. 11 (1991) 595–600. H. Keizo, Light Met. 11 (1994) 609–613. A. Yoshio, Light Met. 11 (1994) 604–608. G.X. Wang, M. Dahms, JOM 5 (1993) 52–56. M. Dahms, F. Schmeizer, J. Seeger, B. Wildhagen, Mater. Sci. Technol. 8 (4) (1992) 359–362. S. Kazuhisa, F. Tsutomu, M.S. Kim, K. Masaki, Powder Powder Metall. 41 (2) (1994) 146–151. J.B. Yang, K.W. Teoh, W.S. Hwang, Mater. Sci. Technol. 13 (8) (1997) 695–701. R. Bohn, T. Klassen, R. Bormann, Intermetallics 9 (2001) 559–569. G. Fanta, R. Bohn, M. Dahms, Intermetallics 9 (2001) 45–49. M.R. Shagiev, O.N. Senkov, G.A. Salishchev, J. Alloys Compd. 313 (2000) 201–208. O.N. Senkov, M. Cavusoglu, F.H. Froes, Mater. Sci. Eng. A300 (2001) 85–93. W. Yu, D.L. Lin, C.C. Law, Trans. Nonferrous Met. Soc. Chin. 9 (Suppl. 1) (1999) 290–297. H.W. Ouyang, Y. Liu, Y.H. He, B.Y. Huang, Chin. J. Nonferrous Met. 2 (2002) 334–337. X.Q. Li, L.X. Hu, E.D. Wang, Trans. Nonferrous Met. Soc. Chin. 4 (2002) 621–624. X.Q. Li, L.X. Hu, E.D. Wang, Mater. Sci. Technol. 1 (2002) 14–18. E.D. Wang, L.X. Hu, X.Q. Li, Powder Metall. Technol. 5 (2003) 259–263. G.X. Wang, M. Dahms, Powder Metall. 24 (4) (1992) 219–224. F.J.J. Van Loo, G.D. Rieck, Acta Metall. 21 (1) (1973) 61. M. Dahms, Mater. Sci. Eng. A 110 (1989) L5–L8. G.L. Chen, J.P. Lin, Ordered Intermetallics Structural Materials, Metallurgical Industry Press, Beijing, China, 1999.
Microstructure and properties of a TiAl alloy prepared by mechanical milling and subsequent reactive sintering Fang Wenbin ∗ , Hu Lianxi, He Wenxiong, Wang Erde, Li Xiaoqing School of Material Science and Engineering, Harbin Institute of Technology, Harbin 150001, PR China Accepted 29 April 2005
Abstract A TiAl-based alloy with nominal composition of Ti–36 wt.% Al was prepared by a mechanical milling and reactive sintering route consisting of three basic steps: (1) mechanical milling of blended elemental Ti and Al powder mixture, (2) densification and consolidation of as-milled Ti/Al composite powder compact via warm extrusion, and (3) subsequent reactive sintering of as-extruded powder billet. The microstructure of the as-sintered TiAl alloy samples was observed by optical microscopy (OM) and transmission electron microscopy (TEM), respectively. The mechanical properties were measured by compression tests performed at room and elevated temperatures up to 800 ◦ C. The results show that mechanical milling before reactive sintering is very effective in microstructure refining of reactively synthesized TiAl alloy. The TiAl alloy sintered from 1 h milled Ti/Al composite powders presented a homogeneous ␥/(␣2 + ␥) duplex microstructure with very fine grains of 5–10 m and good mechanical properties. On the basis of microstructure observation and compressive stress–strain relationship, the deformation mechanism at different temperature was discussed. © 2005 Elsevier B.V. All rights reserved. Keywords: Mechanical milling; Reactive sintering; TiAl alloy; Microstructure and mechanical properties
1. Introduction TiAl-based alloys are strong candidate materials for hightemperature structural applications where weight reduction is of great importance due to their good thermal stability up to 600–800 ◦ C, low densities of 3.9–4.1 g/cm3 , and high strength retention at elevated temperatures [1–3]. Below 700 ◦ C, however, TiAl alloys have intrinsically very poor plasticity and workability, with their elongation in the range of 0–4%. At high temperature above 1100 ◦ C, despite plasticity improvement, they are still rather difficult to be shaped because their flow stress is usually as high as 200 MPa even with such a low strain rate as 1 × 10−3 S−1 [4–6]. Indeed, the poor workability is one of the major obstacles for their application. Therefore, it is of great significance to develop effective techniques for preparation and forming of TiAl alloys. So far, many researches have been focused on this topic in
∗
Corresponding author. Tel.: +86 451 8641 8733; fax: +86 451 5362 7384. E-mail address: [email protected] (F. Wenbin).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.04.049
an effort to accelerate the application of TiAl alloys as light structural materials for high-temperature services [7–19]. Recently, elemental powder metallurgy route has been gaining more and more attention because near-net shape TiAl alloy products can be fabricated by the consolidation and forming of blended Ti and Al elemental powders followed by a subsequent reactive sintering process [7–10]. However, due to the large difference between partial diffusion coefficients of Ti and Al and the immobility of Ti atoms for Ti/Al solid state reaction, the synthesis of TiAl alloy via reactive sintering takes the mechanism that Al atoms move into Ti lattice, thus leading to the formation of Kirkendall diffusion pores [20–22]. Although hot isostatic pressing (HIP) has been reported to be effective in eliminating porosity of reactively sintered TiAl alloys, the high cost and low production efficiency make it not suitable for commercial use. It has been reported that the microstructure of reactively sintered TiAl-based alloys can be refined and the porosity reduced by refining the reactant Ti and Al phases through cold deformation, such as extrusion and rolling, of the Ti and Al powder mixture billet prior to reactive sintering [7–9]. In
F. Wenbin et al. / Materials Science and Engineering A 403 (2005) 186–190
an effort to minimize the porosity due to Kirkendall diffusion and to improve the microstructure of the sintered TiAl-based alloys, we have recently proposed to use mechanical milling to refine the Ti and Al constituents prior to reactive sintering. In the present paper, we report the fabrication of a fully densified TiAl alloy by a cost effective elemental powder metallurgy route consisting of mechanical milling, warm extrusion, and subsequent reactive sintering.
187
cal microscopy (OM) and transmission electron microscopy (TEM), respectively. The mechanical properties of the assintered TiAl alloy samples were measured by isothermal compression tests at room and elevated temperatures. For all tests, cylindrical samples with a dimension of Ø 4 mm × 5 mm were used, and the strain rate for compression was 1.67 × 10−3 S−1 .
3. Results and discussion 2. Experimental 3.1. Microstructure Ti/Al composite powders with a nominal composition of Ti–36 wt.% Al were prepared by mechanical milling of the blended powder mixture of elemental titanium powders (99.5%, <50 m) and aluminum powders (99.9%, <80 m). The milling process was carried out under protection of pure argon using an attritor with a water cooling system. The milling vial and balls were made of stainless steel and GCr15 hardened steel, respectively. All the balls used were of the same size of 6 mm in diameter. The ball:powder weight ratio was 20:1, and the attritor shaft rotation was 400 rev/min. For each batch of milling, the weight of powders processed was 0.5 kg. The as-milled Ti/Al composite powders were cold compacted under a specific pressure of 1000 MPa, and the compacts were then consolidated into dense powder billets by warm extrusion at 250 ◦ C with an extrusion ratio of 16:1. The reactive sintering of the as-extruded Ti/Al composite powder billets was performed in a sintering furnace under vacuum better than 6.67 × 10−3 Pa. The sintering was performed following a two-step procedure, i.e., pre-sintering at 630 ◦ C for 2 h and subsequent final sintering at 1250 ◦ C for 4 h. The microstructure of various samples was observed by opti-
Fig. 1 shows the microstructure of TiAl alloy samples obtained by reactive sintering of powder billets consolidated from Ti/Al composite powders prepared by mechanical milling for various times up to 3 h. Though all as-sintered TiAl alloy samples were characterized by a ␥/(␣2 + ␥) duplex microstructure, differences in grain size and ␥:(␣2 + ␥) phase ratio were obvious. With increasing milling time, the assintered TiAl alloy presented a finer microstructure, and also a higher volume fraction of ␥ phase, suggesting a higher nucleation rate for both ␣ and ␥ phases, and accelerated diffusion kinetics to phase equilibrium, during the reactive sintering process. The high nucleation rate and accelerated diffusion kinetics can be attributed to the refining of the Ti/Al alternating lamellar structure with increasing milling. Indeed, our previous work [18,19] revealed that the lamellar spacing of the reactant Ti and Al phases in the mechanically synthesized elemental Ti/Al composite powders was reduced to less than 0.5 m after milling only for 1 h, and prolonged milling led to formation of nanocrystalline or partially amorphous Ti/Al composite powders with still smaller lamellar spacing. However, the compactibility of the Ti/Al composite
Fig. 1. Microstructure of TiAl alloy samples sintered from Ti/Al composite powders milled for various times: (a) 0 h; (b) 0.5 h; (c) 1 h; (d) 3 h.
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powders became deteriorated with prolonged milling. Therefore, it is very important to optimize the milling parameters to obtain as-milled Ti/Al composite powders with adequately refined Ti/Al lamellar spacing as well as acceptable compactibility. In our opinion, it appears that 1 h milling under conditions specified in the present study is the optimal milling process. As shown in Fig. 1(c), the microstructure of the TiAl alloy sample reactively sintered from extruded billet of 1 h milled Ti/Al composite powders is quite fine and homogenous, with grain size being 5–10 m on average. In comparison, the microstructure of the TiAl alloy sample reactively sintered from starting elemental Ti and Al powders is rather heterogeneous, with some grains being about 10 m while others as coarse as more than 50 m in size (Fig. 1(a)). This suggests that proper milling processing of Ti and Al blended powders before reactive sintering results in not only effective refining of grain size but also significant improvement of microstructure homogeneity of the as-sintered TiAl alloy. Fig. 2 shows a comparison in the lamellar structure of the (␣2 + ␥) phase between two TiAl alloy samples which were sintered from blended Ti/Al elemental powders and 1 h milled Ti/Al composite powders, respectively. It was observed that the average thickness of the ␣2 /␥ lamellae in the TiAl alloy sintered from blended elemental Ti/Al powders was about 250 nm, as shown in Fig. 2(a). In comparison, the ␣2 and ␥ lamellae formed in the TiAl alloy by reactive sintering of the 1 h milled Ti/Al composite powders were much finer, with their average thickness being about only one-third of those in the TiAl alloy sintered from the elemental Ti and Al powder mixture, as shown in Fig. 2(b). This can be explained by the fact that, when as-milled Ti/Al composite powders were used instead of blended elemental Ti/Al powders, the grains of the ␣ phase formed during reactive sintering at temperature above ␣ → ␣2 + ␥ eutectoid transition were much smaller due to a much higher nucleation rate, which in turn resulted in the refinement of the final ␣2 /␥ lamellae structure formed by the
Fig. 3. Compressive stress–strain curves of TiAl alloy sintered from 1 h milled Ti/Al composite powders.
eutectoid transformation of ␣ to ␣2 + ␥ during cooling after sintering. 3.2. Properties It is well known that one major advantage of using the elemental powder metallurgy route for synthesizing of TiAl alloy is the good formability of the precursory material prior to reactive sintering. Indeed, it was found in our present study that the Ti/Al composite powders milled up to 3 h could be readily consolidated to nearly full density by warm hydrostatic extrusion at 250 ◦ C with a moderate extrusion ratio of 16:1 and a specific extrusion pressure about 1000 MPa. By reactive pre-sintering at 630 ◦ C for 2 h, and then final sintering at 1250 ◦ C for 4 h, near fully densified TiAl alloy samples, with a density of about 3.92 g/cm3 , were obtained. Since, as discussed earlier, the TiAl alloy sintered from the 1 h milled Ti/Al composite powders presented the best microstructure, its mechanical properties were further investigated.
Fig. 2. Comparison in (␣2 + ␥) lamellar structure between two TiAl alloy samples sintered from: (a) blended Ti/Al elemental powders and (b) 1 h milled Ti/Al composite powders.
F. Wenbin et al. / Materials Science and Engineering A 403 (2005) 186–190
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Fig. 4. TEM images of the TiAl alloy deformed to a compressive strain of 0.2 at 800 ◦ C showing: (a) recrystallization and twinning in ␥ phase region and (b) recrystallization in (␣2 + ␥) phase region.
Fig. 3 shows the compressive stress–strain curves at room and elevated temperatures up to 800 ◦ C of the TiAl alloy sintered from 1 h milled Ti/Al composite powders. At room temperature, its yield strength is about 560 MPa and the compressive strain to failure is below 0.05. Unlike conventional metal materials, the yield strength of the TiAl alloy increases with increasing temperature in a certain temperature range, and achieves a maximum of about 820 MPa at 600 ◦ C. Then, with further increase of the temperature, the yield strength decreases. On the other hand, the compressive strain before failure increases steadily with increasing temperature. Indeed, the compressive strain to failure at 800 ◦ C is as high as 0.25, suggesting significant improvement of plasticity or appearance of superplastic deformation behavior at this temperature. The mechanical behavior shown in Fig. 3 can be interpreted according to the deformation mechanism of TiAl alloy. As suggested by Chen and Lin [23], for TiAl alloys with duplex microstructure, the deformation behavior below 600 ◦ C is dominated by the movement and pinning of the 1¯ 1 0 screw dislocations. At relatively higher temperature, the 1¯ 1 0 screw dislocations tend to shift on to the {1 1 1} plane by cross-slipping, and then back to the starting plane, forming non-spiral type kinks which pin dislocation movement. When the temperature is enhanced from room temperature to 600 ◦ C, the density of pinning sites steadily increases, while the distance between pinning points decreases. Besides, ¯ and 1/21 1 2] ¯ type tend to super dislocations of 1 0 1] decompose through thermal activation at relatively higher temperature to form dislocation dipoles which also present pinning effect. Therefore, a higher shear stress is needed for dislocation movement at a higher temperature. This leads to the increase of the yield strength with increasing temperature up to 600 ◦ C. On the other hand, at temperature above 600 ◦ C, the mobility of both the 1/21 1 0] dislocations and the 1/61 1 2] twinning dislocations is significantly improved due to thermal activation. Meanwhile, the pinning effect of the kinks by cross-slipping is reduced, and grain boundary
sliding or grain rotating is also possible as alternative deformation modes. Therefore, the yield strength above 600 ◦ C decreases with increasing temperature. Fig. 4 shows TEM images of the TiAl alloy deformed at 800 ◦ C to a compressive strain of 0.2. Evidence of dynamic recrystallization in both ␥ and (␣2 + ␥) phase regions during deformation was obvious, as shown in Fig. 4(a and b), respectively. Besides, evidence of deformation twinning was also observed (Fig. 4(a)). The dynamic recrystallization during compression can be attributed to two major factors: (1) the high deformation temperature, 800 ◦ C (1073 K), which is well above 0.55Tm , and (2) the low strain rate, 1.67 × 10−3 S−1 , which is favorable for recrystallization nucleation. It is believed that dynamic recrystallization is the major softening mechanism of TiAl alloy during hightemperature deformation with low strain rate, which accounts well for the steady flow stress after being compressed at 800 ◦ C to a compressive strain more than 0.05, as shown in Fig. 3. Due to the very fine microstructure of the assintered TiAl alloy, it is quite possible that grain boundary sliding and grain rotating occur during deformation at such a temperature as 800 ◦ C. Therefore, both dynamic recrystallization and alternative deformation modes including grain boundary sliding, grain rotating, and twinning at high temperature should be responsible for the significant improvement of plasticity of TiAl alloy for compression at 800 ◦ C.
4. Conclusions (1) Mechanical milling of blended elemental Ti and Al powders before reactive sintering is very effective in microstructure homogeneity improvement and grain size refining of reactively synthesized TiAl alloy prepared by powder metallurgy route using elemental powders as starting material. In particular, the TiAl alloy sintered from 1 h milled Ti/Al composite powders presented
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a homogeneous ␥/(␣2 + ␥) duplex microstructure with very fine grains of 5–10 m in size. (2) At room temperature, the compressive yield strength of the TiAl alloy sintered from 1 h milled Ti/Al composite powders is about 560 MPa. With the increase of temperature, the yield strength at first increases and then decreases, with a maximum value of about 820 MPa achieved at 600 ◦ C. At temperature above 600 ◦ C, the plasticity of the TiAl alloy is improved significantly, with a failure strain as high as 0.25 for compression at 800 ◦ C. (3) Dynamic recrystallization is the major softening mechanism of TiAl alloy during high-temperature deformation with low strain rate. Both dynamic recrystallization and possible alternative deformation modes, such as grain boundary sliding, grain rotating, and twinning, are responsible for the significant improvement of plasticity of TiAl alloy at 800 ◦ C.
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