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Effect of fabrication process on microstructure of high Nb containing TiAl alloy
Journal of Alloys and Compounds 458 (2008) 313–317
Effect of fabrication process on microstructure of high Nb containing TiAl alloy Y.H. Wang a , J.P. Lin a,∗ , X.J. Xu a , Y.H. He b , Y.L. Wang a , G.L. Chen a a
State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, PR China b State Key Laboratory for Powder Metallurgy, Central South University, Changsha 410083, PR China Received 22 May 2006; received in revised form 23 October 2006; accepted 20 February 2007 Available online 6 April 2007
Abstract Ti–45Al–8.5Nb–(W,B) (at.%) alloys were fabricated by ingot metallurgy (IM) and elemental powder metallurgy (EPM) processes, respectively. The effect of fabrication process on microstructure of high Nb containing TiAl alloy was investigated. The results showed that IM alloy has near lamellar (NL) microstructure composed of coarse lamellar colonies with a lamellar spacing of about 700 nm and equiaxed ␥ phase existing at the boundaries of lamellar colonies. While EPM alloy has fully lamellar (FL) microstructure with a lamellar spacing of about 180 nm. A small amount of  phase exists either at the boundaries of the lamellar colonies or in the lamellar colonies in platelet or blocky morphology in IM alloy. However, the micro-segregation does not exist in EPM microstructure. Borides appear in both the microstructures, but their sizes are smaller in EPM microstructure. HIP treatment can significantly eliminate pores and decrease the  phase in IM alloy. It is regretful that HIP treatment can only decrease porosity to some extent, but cannot completely eliminate pores in EPM alloy. © 2007 Elsevier B.V. All rights reserved. Keywords: Intermetallics; Fabrication process; Microstructure
1. Introduction High Nb containing TiAl alloys have a great potential for high temperature applications in aerospace and automotive fields because of their high strength, low density, high modulus, good resistance to creep and high temperature oxidation [1–4]. High Nb addition can significantly improve high temperature properties of TiAl-based alloys, but it also enhances the difficulties of their fabrication [5]. These alloys are mostly fabricated by ingot metallurgy (IM) up to date [6–8]. However, IM process usually encounters with segregations. Furthermore, workability of high Nb containing TiAl alloys is poor due to its inherent brittleness. Therefore, many advanced processes are attempted to improve their microstructural homogeneity and workability, such as forging, HIP and extrusion. Unfortunately, these problems are still overcome unsuccessfully [5–8]. It is significant to pursue a novel approach for achieving a fine and homogeneous microstructure. Presently, powder metallurgy technique appears to be more attractive since high degrees of ∗
Corresponding author. Tel.: +86 10 62332192; fax: +86 10 62332508. E-mail address: [email protected] (J.P. Lin).
0925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2007.02.163
chemical homogeneities can be obtained and large-scale segregations are avoided. Among the powder metallurgy techniques, elemental powder metallurgy (EPM) has drawn intensive attention because of its low cost and convenient addition of alloy elements. Recently, numerous work has been conducted on the microstructures and properties of EPM TiAl-based alloys [9–11]. The objective of the present paper was to investigate the effect of fabrication process on microstructure of high Nb containing TiAl alloy. 2. Experimental procedures For the present study, the alloy with nominal composition Ti–45Al– 8.5Nb–(W,B) (at.%) was fabricated by two different processes, namely IM and EPM. IM alloy was fabricated in an induction skull melting furnace. In EPM process, for elemental Ti and Al powders, the mean particle sizes were 43 and 38 m, respectively, and the purity was higher than 99.5%. For Nb, W and B, the mean particle sizes were smaller than 25 m, and the purity was higher than 99.9%. The elemental powders were mixed for 6 h using a V-blender without any additive. Subsequently, powder mixture was die-pressed under 350 MPa to green compacts. The compacts were then put into a graphite die and hot-pressed for 1 h in an argon protective atmosphere under 25 MPa at 1400 ◦ C. Subsequently, both of the alloys were hot isostatic-pressed at 1200 ◦ C/4 h/200 MPa to eliminate porosity.
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Fig. 1. Micrographs of IM Ti–45Al–8.5Nb–(W,B) alloy: (a) optical microstructure, (b) BSE microstructure, (c) boride morphology in deep-etched sample and (d) TEM microstructure. Alloys fabricated by the two different processes were cut, polished and etched using Kroll’s solution. Microstructural observation was carried out by optical microscopy (OM), scanning electron microscopy using back scattering electron imaging (BSE) and energy dispersive spectroscopy (EDS). Constituent phases were characterized by X-ray diffraction (XRD). Lamellar spacing was determined by transmission electron microscopy (TEM). The lamellar spacing given in this paper was arithmetic average value measured without taking account of types of lamellar boundaries. TEM foils were prepared by twin jet electropolished in a solution of 60% (vol.%) methanol, 35% butyl alcohol and 5% perchloric acid at 15 V and −30 ◦ C. The lamellar colony size was determined by the intersection linear method.
3. Results
mean lamellar spacing of the lamellar colonies is about 700 nm (Fig. 1d). 3.2. Microstructure of EPM alloy Fig. 3 shows the microstructures of EPM Ti–45Al– 8.5Nb–(W,B) alloy before HIP treatment. A typical FL microstructure containing fine and homogeneous lamellar colonies ␣2 /␥ was successfully developed out (Fig. 3a). This FL microstructure is similar to that of IM alloy, but finer. The lamellar colony size is approximately 45 m. From XRD pattern of EPM alloy, as shown in Fig. 4, only ␥ and ␣2 phases were detected. Thus, it is deduced that Nb element diffuses into
3.1. Microstructure of IM alloy The IM Ti–45Al–8.5Nb–(W,B) alloy after HIP treatment has a homogenous NL microstructure mainly consisting of ␣2 /␥ lamellar structure (Fig. 1a). The mean lamellar colony size is about 85 m. It can be seen that there is evidence of ␥ phase with black contrast and another phase with white contrast at colony boundaries or inside colonies, but no dendritic core region enriched with W and Nb from BSE microstructure of the IM alloy (Fig. 1b). XRD pattern of the IM alloy confirms the presence of these phases, namely, ␥, ␣2 and  (Fig. 2). EDS analysis result indicates that the white phase contains higher Ti, Nb and W than the nominal composition of the alloy. Therefore, we can draw a conclusion that the white phase is  (B2) phase. Many boride particles or rods with bright contrast are also observed either at colony boundaries or in the colonies. Closer examination of borides in deeply etched sample reveals that they are needles or rods longer than 50 m with aspect ratios of 20–100 (Fig. 1c). TEM observation shows that the
Fig. 2. X-ray diffraction pattern of phases in IM Ti–45Al–8.5Nb–(W,B) alloy.
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315
Fig. 3. Micrographs of hot-pressed Ti–45Al–8.5Nb–(W,B) alloy: (a) optical microstructure, (b) BSE microstructure, (c) particulate boride morphology in deep-etched sample and (d) acicular boride morphology in deep-etched sample.
matrix by solid solution in reactive hot pressing. BSE image indicates that elemental Nb diffuses completely, and pore nests form in situ. The deformation and agglomeration phenomena of pores are observed in some pore nests (Fig. 3b). Furthermore, few particulate or acicular borides appear in matrix, but their sizes are smaller than that in the IM alloy (Fig. 3 c and d). HIP treatment can only decrease porosity to some extent, but cannot eliminate pores completely (Fig. 5a). Agglomeration phenomenon of pores is more evident. However, compression deformations of pores basically disappear. TEM observation shows that the mean lamellar spacing of the lamellar colonies is about 180 nm (Fig. 5b).
Fig. 4. X-ray diffraction pattern of phases in hot-pressed Ti–45Al–8.5Nb–(W,B) alloy.
4. Discussion 4.1. Effect of fabrication process on microstructure According to Ti–Al binary phase diagram, Ti–45Al– 8.5Nb–(W,B) IM alloy evolves the following path way during solidification: L → L +  →  →  + ␣ → ␣ → ␣ + ␥ → lamella (␣2 + ␥) + ␥ [2]. As the IM alloy solidifies through the  phase zone,  phase will be the first production. The alloy does not appear evident columnar character when it solidifies [12]. The reason is that although the 1 0 0 axis of the  phase is the preferential direction of crystal growth during solidification, there are three equivalent directions for 1 0 0, namely, [1 0 0], [0 1 0] and [0 0 1]. During cooling after solidification,  crystals are transformed to ␣, following the orientation relationship of {1 1 0} //(0 0 0 1)␣ and 1 1 1 //1 1 2 0␣ . The transformation of  to ␣, will not consume all the  phase owing to the low diffusivity of these elements and quick solidification rate. As a result, twelve variants of ␣ with different orientations are formed. During subsequent cooling, the ␥ phase precipitates from each ␣ variant in terms of a strict orientation relationship: {1 1 1}␥ //(0 0 0 1)␣ and 1 1 0␥ //[1 1 2 0]␣ . Finally, the microstructure composed of lamellar colony structure as the major constituent and ␥ and  as the minor was obtained in IM alloy. It is worth noting that the presence of  phase in IM TiAlbased alloys with alloying additions of  stabilizers, such as Nb, Cr, Mo and W has been reported previously [6,13]. It is also found that in the alloy with Nb addition above 9.5 at.%, there exists three phases (␣2 + ␥ + ) field in the high Nb containing TiAl alloy [2]. Thus, it is suspected that the  phase precipitates again from the ␣ phase in the regions rich in Nb.
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Fig. 5. Micrographs of as-HIP Ti–45Al–8.5Nb–(W,B) alloy: (a) BSE microstructure and (b) TEM microstructure.
Therefore, apart from the residual  phase presenting at colony boundaries and triple junctions,  phase also exists in the lamellar colonies. However,  phase does not form in EPM high Nb containing TiAl alloy, even if Nb and W powders, especially high Nb addition, exists in present alloy. It is related with their different fabrication processes. Reactive sintering of Ti, Al elemental powders is a diffusioncontrolled process including the formation of transient phases, such as TiAl3 and TiAl2 [9]. At the beginning of hot pressing, TiAl3 forms by the solid–solid and solid–liquid reactions between Ti and Al powders. After the reaction is completed, pure Ti and TiAl3 are present. During subsequent hot pressing, further reaction between Ti and TiAl3 leads to the formation of the intermetallic compounds Ti3 Al, TiAl and TiAl2 . In addition, Ti3 Al can react with TiAl2 to form TiAl. Finally, only Ti3 Al and TiAl are present. However, niobium aluminides does not form in the reactive hot pressing. This is due to the different diffusivity of these elements. Generally, Al seems to be the fastest specie, with Ti having a mobility close to that of Al, and Nb being the lowest mover in Ti–Al–Nb system [14,15]. It indicates that the reaction ability between Ti and Al powders is dominant in the ternary system, restraining the reaction between Nb and Al powders. It is deduced that, therefore, elemental Nb dissolves into the Ti–Al matrix by diffusion. Borides are observed in the two different microstructures, but their sizes are smaller in the EPM microstructure. The effect of borides on refining microstructure is given elsewhere [16,17]. Their formation mechanisms are different, which are similar to those of Ti–Al matrix. One forms by a solidification process, the other by diffusion reaction process. For the studied alloys, the crystal structure of the particulate or acicular borides is not known yet, but from the hexahedral prism appearance of the particles it can be concluded that the boride phase is TiB2 [18]. Therefore, it is thought that TiB2 forms via the solid–solid reaction in situ between Ti powder and B powder, namely, Ti(s) + 2B(s) → TiB2 (s). 4.2. Effect of HIP on microstructure It is well known that HIP treatment can eliminate pores significantly. For IM alloy, the densification effect is more evident because the melting process was in vacuum. In addition, HIP can significantly decrease the  phase, which partly transforms
to ␣ and ␥, but the  phase cannot be eliminated completely [8]. However, for EPM alloy, pores are not eliminated completely through HIP treatment. The main reason is that most of the pores in the alloy are closed pores, including pore nests formed in situ after Nb powders diffusion as well as Kirkendall pores generated from the transient phases during the reaction between Al and Ti powders. The closed pores are filled with the protective gas (Ar) at a certain pressure. During HIP-compaction, these pores shrink until their inner gas pressure equals the applied HIP-pressure (200 MPa in this case). Therefore, it indicates that the compression of closed pores is of practical importance to obtain dense alloys. Presently, many efficient methods have been reported in other alloys, such as using finer elemental powders, higher force in hot pressing and extrusion [9,19]. In addition, the process by replacing elemental Nb powders with Nb–Al alloy powders is favorable to improve the diffusion of Nb. We also found that agglomeration phenomenon of pores is more intensive and compression deformations of pores disappear, which are related with the applied uniform force in the HIP treatment. 5. Conclusions (1) The IM high Nb containing TiAl alloy has NL microstructure composed of coarse lamellar colonies with a lamellar spacing of about 700 nm and equiaxed ␥ phase existing at the boundaries of lamellar colonies, while EPM alloy has FL microstructure with a lamellar spacing of about 180 nm. Borides appear in both the microstructures, but their sizes are smaller in the IM alloy. (2) A small amount of  phase exists either at the boundaries of the lamellar colonies or in the lamellar colonies in platelet or blocky morphology in IM alloy. However, the microsegregation does not exist in EPM alloy. (3) For IM alloy, HIP treatment can significantly eliminate pores and decrease the  phase. For EPM alloy, HIP treatment can only decrease porosity to some extent, but cannot completely eliminate pores. Acknowledgements This research was sponsored by the Key Grant Project of Chinese Ministry of Education (No. 704008) and by the program
Y.H. Wang et al. / Journal of Alloys and Compounds 458 (2008) 313–317
from New Century Excellent Talents in University (No. NCET04-01017). References [1] G. Chen, W. Zhang, Y. Wang, G. Wang, Z. Sun, in: R. Darolia, J.J. Lewandowski, C.T. Liu, P.L. Martin, D.B. Miracle, M.V. Nathal (Eds.), Structural Intermetallics, TMS, Warrendale, PA, 1993, pp. 319– 324. [2] G. Chen, W. Zhang, Z. Liu, S. Li, Y.W. Kim, in: Y.W. Kim, D.M. Dimiduk, M.H. Loretto (Eds.), Gamma Titanium Aluminides, TMS, Warrendale, PA, 1999, pp. 371–380. [3] G. Chen, L. Zhang, Mater. Sci. Eng. 329–331A (2002) 163–170. [4] Y.Q. Yan, Z.Q. Zhang, Z.G. Zhen, K.G. Wang, L. Zhou, Mater. Sci. Eng. 280A (2000) 187–191. [5] W.J. Zhang, F. Appel, Mater. Sci. Eng. A 329–331 (2002) 649–652. [6] W.J. Zhang, G.L. Chen, F. Apple, T.G. Nieh, S.C. Deevi, Mater. Sci. Eng. A 315 (2001) 250–253.
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[7] Z.C. Liu, J.P. Lin, S.J. Li, G.L. Chen, Intermetallics 10 (2002) 653–659. [8] X.J. Xu, L.H. Xu, J.P. Lin, Y.L. Wang, Z. Lin, G.L. Chen, Intermetallics 13 (2005) 337–341. [9] G.X. Wang, M. Dahms, Powder Metall. Int. 24 (1992) 219–225. [10] Y. Wu, S.K. Hwang, Acta Mater. 50 (2002) 1479–1493. [11] S.H. Yang, W.Y. Kim, M.S. Kim, Intermetallics 11 (2003) 849–855. [12] Q. Xu, M.C. Chaturvedi, N.L. Richards, N. Goel, in: M.V. Nathal, R. Dorolia, C.T. Liu, P.L. Martin, D.B. Miracle, R. Wagner, M. Yamaguchi (Eds.), Structural Intermetallics, TMS, Warrendale, PA, 1997, pp. 323–329. [13] L.M. Hsiung, T.G. Nieh, Mater. Sci. Eng. A 364 (2004) 1–10. [14] V.A.R. Henriques, H.R.Z. Sandim, G.C. Coelho, C.R.M. Silva, Mater. Sci. Eng. 347A (2003) 315–324. [15] F.J.J. Van Loo, G.D. Rieck, Acta Metall. 21 (1973) 61–71. [16] D. Hu, Intermetallics 10 (2002) 851–858. [17] L. Huang, C.T. Liu, P.K. Liaw, B.Y. Huang, JOM 56 (2004) 361–364. [18] A.A. Abdel-Hamid, S. Hamar Thibault, R. Hamar, J. Cryst. Growth 71 (1985) 744–750. [19] Z. Witczak, P. Witczak, R. Jemielniak, Mater. Sci. Forum 475–479 (2005) 759–762.
Effect of fabrication process on microstructure of high Nb containing TiAl alloy Y.H. Wang a , J.P. Lin a,∗ , X.J. Xu a , Y.H. He b , Y.L. Wang a , G.L. Chen a a
State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, PR China b State Key Laboratory for Powder Metallurgy, Central South University, Changsha 410083, PR China Received 22 May 2006; received in revised form 23 October 2006; accepted 20 February 2007 Available online 6 April 2007
Abstract Ti–45Al–8.5Nb–(W,B) (at.%) alloys were fabricated by ingot metallurgy (IM) and elemental powder metallurgy (EPM) processes, respectively. The effect of fabrication process on microstructure of high Nb containing TiAl alloy was investigated. The results showed that IM alloy has near lamellar (NL) microstructure composed of coarse lamellar colonies with a lamellar spacing of about 700 nm and equiaxed ␥ phase existing at the boundaries of lamellar colonies. While EPM alloy has fully lamellar (FL) microstructure with a lamellar spacing of about 180 nm. A small amount of  phase exists either at the boundaries of the lamellar colonies or in the lamellar colonies in platelet or blocky morphology in IM alloy. However, the micro-segregation does not exist in EPM microstructure. Borides appear in both the microstructures, but their sizes are smaller in EPM microstructure. HIP treatment can significantly eliminate pores and decrease the  phase in IM alloy. It is regretful that HIP treatment can only decrease porosity to some extent, but cannot completely eliminate pores in EPM alloy. © 2007 Elsevier B.V. All rights reserved. Keywords: Intermetallics; Fabrication process; Microstructure
1. Introduction High Nb containing TiAl alloys have a great potential for high temperature applications in aerospace and automotive fields because of their high strength, low density, high modulus, good resistance to creep and high temperature oxidation [1–4]. High Nb addition can significantly improve high temperature properties of TiAl-based alloys, but it also enhances the difficulties of their fabrication [5]. These alloys are mostly fabricated by ingot metallurgy (IM) up to date [6–8]. However, IM process usually encounters with segregations. Furthermore, workability of high Nb containing TiAl alloys is poor due to its inherent brittleness. Therefore, many advanced processes are attempted to improve their microstructural homogeneity and workability, such as forging, HIP and extrusion. Unfortunately, these problems are still overcome unsuccessfully [5–8]. It is significant to pursue a novel approach for achieving a fine and homogeneous microstructure. Presently, powder metallurgy technique appears to be more attractive since high degrees of ∗
Corresponding author. Tel.: +86 10 62332192; fax: +86 10 62332508. E-mail address: [email protected] (J.P. Lin).
0925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2007.02.163
chemical homogeneities can be obtained and large-scale segregations are avoided. Among the powder metallurgy techniques, elemental powder metallurgy (EPM) has drawn intensive attention because of its low cost and convenient addition of alloy elements. Recently, numerous work has been conducted on the microstructures and properties of EPM TiAl-based alloys [9–11]. The objective of the present paper was to investigate the effect of fabrication process on microstructure of high Nb containing TiAl alloy. 2. Experimental procedures For the present study, the alloy with nominal composition Ti–45Al– 8.5Nb–(W,B) (at.%) was fabricated by two different processes, namely IM and EPM. IM alloy was fabricated in an induction skull melting furnace. In EPM process, for elemental Ti and Al powders, the mean particle sizes were 43 and 38 m, respectively, and the purity was higher than 99.5%. For Nb, W and B, the mean particle sizes were smaller than 25 m, and the purity was higher than 99.9%. The elemental powders were mixed for 6 h using a V-blender without any additive. Subsequently, powder mixture was die-pressed under 350 MPa to green compacts. The compacts were then put into a graphite die and hot-pressed for 1 h in an argon protective atmosphere under 25 MPa at 1400 ◦ C. Subsequently, both of the alloys were hot isostatic-pressed at 1200 ◦ C/4 h/200 MPa to eliminate porosity.
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Fig. 1. Micrographs of IM Ti–45Al–8.5Nb–(W,B) alloy: (a) optical microstructure, (b) BSE microstructure, (c) boride morphology in deep-etched sample and (d) TEM microstructure. Alloys fabricated by the two different processes were cut, polished and etched using Kroll’s solution. Microstructural observation was carried out by optical microscopy (OM), scanning electron microscopy using back scattering electron imaging (BSE) and energy dispersive spectroscopy (EDS). Constituent phases were characterized by X-ray diffraction (XRD). Lamellar spacing was determined by transmission electron microscopy (TEM). The lamellar spacing given in this paper was arithmetic average value measured without taking account of types of lamellar boundaries. TEM foils were prepared by twin jet electropolished in a solution of 60% (vol.%) methanol, 35% butyl alcohol and 5% perchloric acid at 15 V and −30 ◦ C. The lamellar colony size was determined by the intersection linear method.
3. Results
mean lamellar spacing of the lamellar colonies is about 700 nm (Fig. 1d). 3.2. Microstructure of EPM alloy Fig. 3 shows the microstructures of EPM Ti–45Al– 8.5Nb–(W,B) alloy before HIP treatment. A typical FL microstructure containing fine and homogeneous lamellar colonies ␣2 /␥ was successfully developed out (Fig. 3a). This FL microstructure is similar to that of IM alloy, but finer. The lamellar colony size is approximately 45 m. From XRD pattern of EPM alloy, as shown in Fig. 4, only ␥ and ␣2 phases were detected. Thus, it is deduced that Nb element diffuses into
3.1. Microstructure of IM alloy The IM Ti–45Al–8.5Nb–(W,B) alloy after HIP treatment has a homogenous NL microstructure mainly consisting of ␣2 /␥ lamellar structure (Fig. 1a). The mean lamellar colony size is about 85 m. It can be seen that there is evidence of ␥ phase with black contrast and another phase with white contrast at colony boundaries or inside colonies, but no dendritic core region enriched with W and Nb from BSE microstructure of the IM alloy (Fig. 1b). XRD pattern of the IM alloy confirms the presence of these phases, namely, ␥, ␣2 and  (Fig. 2). EDS analysis result indicates that the white phase contains higher Ti, Nb and W than the nominal composition of the alloy. Therefore, we can draw a conclusion that the white phase is  (B2) phase. Many boride particles or rods with bright contrast are also observed either at colony boundaries or in the colonies. Closer examination of borides in deeply etched sample reveals that they are needles or rods longer than 50 m with aspect ratios of 20–100 (Fig. 1c). TEM observation shows that the
Fig. 2. X-ray diffraction pattern of phases in IM Ti–45Al–8.5Nb–(W,B) alloy.
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315
Fig. 3. Micrographs of hot-pressed Ti–45Al–8.5Nb–(W,B) alloy: (a) optical microstructure, (b) BSE microstructure, (c) particulate boride morphology in deep-etched sample and (d) acicular boride morphology in deep-etched sample.
matrix by solid solution in reactive hot pressing. BSE image indicates that elemental Nb diffuses completely, and pore nests form in situ. The deformation and agglomeration phenomena of pores are observed in some pore nests (Fig. 3b). Furthermore, few particulate or acicular borides appear in matrix, but their sizes are smaller than that in the IM alloy (Fig. 3 c and d). HIP treatment can only decrease porosity to some extent, but cannot eliminate pores completely (Fig. 5a). Agglomeration phenomenon of pores is more evident. However, compression deformations of pores basically disappear. TEM observation shows that the mean lamellar spacing of the lamellar colonies is about 180 nm (Fig. 5b).
Fig. 4. X-ray diffraction pattern of phases in hot-pressed Ti–45Al–8.5Nb–(W,B) alloy.
4. Discussion 4.1. Effect of fabrication process on microstructure According to Ti–Al binary phase diagram, Ti–45Al– 8.5Nb–(W,B) IM alloy evolves the following path way during solidification: L → L +  →  →  + ␣ → ␣ → ␣ + ␥ → lamella (␣2 + ␥) + ␥ [2]. As the IM alloy solidifies through the  phase zone,  phase will be the first production. The alloy does not appear evident columnar character when it solidifies [12]. The reason is that although the 1 0 0 axis of the  phase is the preferential direction of crystal growth during solidification, there are three equivalent directions for 1 0 0, namely, [1 0 0], [0 1 0] and [0 0 1]. During cooling after solidification,  crystals are transformed to ␣, following the orientation relationship of {1 1 0} //(0 0 0 1)␣ and 1 1 1 //1 1 2 0␣ . The transformation of  to ␣, will not consume all the  phase owing to the low diffusivity of these elements and quick solidification rate. As a result, twelve variants of ␣ with different orientations are formed. During subsequent cooling, the ␥ phase precipitates from each ␣ variant in terms of a strict orientation relationship: {1 1 1}␥ //(0 0 0 1)␣ and 1 1 0␥ //[1 1 2 0]␣ . Finally, the microstructure composed of lamellar colony structure as the major constituent and ␥ and  as the minor was obtained in IM alloy. It is worth noting that the presence of  phase in IM TiAlbased alloys with alloying additions of  stabilizers, such as Nb, Cr, Mo and W has been reported previously [6,13]. It is also found that in the alloy with Nb addition above 9.5 at.%, there exists three phases (␣2 + ␥ + ) field in the high Nb containing TiAl alloy [2]. Thus, it is suspected that the  phase precipitates again from the ␣ phase in the regions rich in Nb.
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Fig. 5. Micrographs of as-HIP Ti–45Al–8.5Nb–(W,B) alloy: (a) BSE microstructure and (b) TEM microstructure.
Therefore, apart from the residual  phase presenting at colony boundaries and triple junctions,  phase also exists in the lamellar colonies. However,  phase does not form in EPM high Nb containing TiAl alloy, even if Nb and W powders, especially high Nb addition, exists in present alloy. It is related with their different fabrication processes. Reactive sintering of Ti, Al elemental powders is a diffusioncontrolled process including the formation of transient phases, such as TiAl3 and TiAl2 [9]. At the beginning of hot pressing, TiAl3 forms by the solid–solid and solid–liquid reactions between Ti and Al powders. After the reaction is completed, pure Ti and TiAl3 are present. During subsequent hot pressing, further reaction between Ti and TiAl3 leads to the formation of the intermetallic compounds Ti3 Al, TiAl and TiAl2 . In addition, Ti3 Al can react with TiAl2 to form TiAl. Finally, only Ti3 Al and TiAl are present. However, niobium aluminides does not form in the reactive hot pressing. This is due to the different diffusivity of these elements. Generally, Al seems to be the fastest specie, with Ti having a mobility close to that of Al, and Nb being the lowest mover in Ti–Al–Nb system [14,15]. It indicates that the reaction ability between Ti and Al powders is dominant in the ternary system, restraining the reaction between Nb and Al powders. It is deduced that, therefore, elemental Nb dissolves into the Ti–Al matrix by diffusion. Borides are observed in the two different microstructures, but their sizes are smaller in the EPM microstructure. The effect of borides on refining microstructure is given elsewhere [16,17]. Their formation mechanisms are different, which are similar to those of Ti–Al matrix. One forms by a solidification process, the other by diffusion reaction process. For the studied alloys, the crystal structure of the particulate or acicular borides is not known yet, but from the hexahedral prism appearance of the particles it can be concluded that the boride phase is TiB2 [18]. Therefore, it is thought that TiB2 forms via the solid–solid reaction in situ between Ti powder and B powder, namely, Ti(s) + 2B(s) → TiB2 (s). 4.2. Effect of HIP on microstructure It is well known that HIP treatment can eliminate pores significantly. For IM alloy, the densification effect is more evident because the melting process was in vacuum. In addition, HIP can significantly decrease the  phase, which partly transforms
to ␣ and ␥, but the  phase cannot be eliminated completely [8]. However, for EPM alloy, pores are not eliminated completely through HIP treatment. The main reason is that most of the pores in the alloy are closed pores, including pore nests formed in situ after Nb powders diffusion as well as Kirkendall pores generated from the transient phases during the reaction between Al and Ti powders. The closed pores are filled with the protective gas (Ar) at a certain pressure. During HIP-compaction, these pores shrink until their inner gas pressure equals the applied HIP-pressure (200 MPa in this case). Therefore, it indicates that the compression of closed pores is of practical importance to obtain dense alloys. Presently, many efficient methods have been reported in other alloys, such as using finer elemental powders, higher force in hot pressing and extrusion [9,19]. In addition, the process by replacing elemental Nb powders with Nb–Al alloy powders is favorable to improve the diffusion of Nb. We also found that agglomeration phenomenon of pores is more intensive and compression deformations of pores disappear, which are related with the applied uniform force in the HIP treatment. 5. Conclusions (1) The IM high Nb containing TiAl alloy has NL microstructure composed of coarse lamellar colonies with a lamellar spacing of about 700 nm and equiaxed ␥ phase existing at the boundaries of lamellar colonies, while EPM alloy has FL microstructure with a lamellar spacing of about 180 nm. Borides appear in both the microstructures, but their sizes are smaller in the IM alloy. (2) A small amount of  phase exists either at the boundaries of the lamellar colonies or in the lamellar colonies in platelet or blocky morphology in IM alloy. However, the microsegregation does not exist in EPM alloy. (3) For IM alloy, HIP treatment can significantly eliminate pores and decrease the  phase. For EPM alloy, HIP treatment can only decrease porosity to some extent, but cannot completely eliminate pores. Acknowledgements This research was sponsored by the Key Grant Project of Chinese Ministry of Education (No. 704008) and by the program
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