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Effects of melt treatment on the cast structure of M963 superalloy
Scripta Materialia 48 (2003) 425–429 www.actamat-journals.com
Effects of melt treatment on the cast structure of M963 superalloy F.S. Yin a
a,b,*
, X.F. Sun a, J.G. Li a, H.R. Guan a, Z.Q. Hu
a
Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, P.R. China b College of Mechanical Engineering, Shandong University of Technology, Zibo 255012, P.R. China Received 6 June 2002; received in revised form 2 September 2002; accepted 4 September 2002
Abstract The cast structure of M963 superalloy is greatly affected by the melt superheating treatment. Melt superheating treatment (below 1923 K) coarsens the grain structure and transforms the blocky MC carbide into Chinese-script one that mostly distributed in dendrite. However, when the melt superheating temperature reaches 2123 K, the grain structure becomes very fine equiaxed, and the morphology of MC carbide changes back to the blocky appearance. Ó 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Cast; Nickel; Melt treatment; Carbide
1. Introduction The grain structure and primary MC carbide morphology have a very important effect on the mechanical properties of cast nickel-base superalloys [1,2]. They can be controlled by optimizing the process variables, such as melt treatment, including melt superheating temperature and the holding time at that temperature, as well as casting parameters including mold preheating temperature and pouring temperature of the melt. The effect of casting parameters has been extensively studied for several superalloys [3–7]. Liu et al. [7] have reported that the variation of melt superheat (below * Corresponding author. Address: Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, P.R. China. Tel.: +86-24-238-43531/55432; fax: +86-24-238-91320. E-mail address: [email protected] (F.S. Yin).
1823 K), i.e. the temperature above the liquidus of the alloy, produces a significant effect on the cast microstructure of IN738LC superalloy, and low melt superheating temperature causes significant reduction in the grain size. There is no report on the effect of high melt superheating temperature up to 2123 K on the cast structure of superalloy. This paper presents the results and mechanism of high temperature treatment of melt.
2. Experimental procedures The superalloy used in this study is M963 alloy having the following composition (wt. pct): 0.15C, 8.8Cr, 9.6Co, 10.5W, 1.6Mo, 1.0Nb, 5.8Al, 2.4Ti, 0.04Zr, 0.03B and rest Ni. The liquidus temperature of the alloy is 1623 K determined by DTA. Samples were prepared in a VIM-25F vacuum induction furnace. Approximately 4.8 kg charge
1359-6462/03/$ - see front matter Ó 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 6 2 ( 0 2 ) 0 0 4 4 6 - 3
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per heat was melted and superheated to different temperatures of up to 2123 K for 5 min and poured into a preheated ceramic mold, 15 mm in diameter. The mold preheating temperature and pouring temperature were kept at 1123 and 1723 K, respectively, in all cases. The vacuum during melt treatment was 0.1 Pa. The specimens for microstructure examination cut from the cast ingots were ground, polished and chemically etched with 40 ml HCl þ 40 ml C2 H5 OH þ 20 ml H2 O þ 1:5 g CuSO4 solution. The grain structure and primary MC carbide morphology were examined with standard metallographic techniques. An ISM-6301F scanning electron microscope (SEM) was used to reveal further microstructural details. Electron-probe microanalysis (EPMA) was used to determine the phase composition.
3. Results Fig. 1 shows the grain structures of the alloy cast with different melt superheating temperatures.
In Fig. 1, one can see that melt-superheating temperature has a significant effect on the grain size. Without any melt treatment, the grain structure is almost equiaxed and fine. Below 1873 K the grain size increases with increasing the melt superheating temperature, which is in accordance with the results of Liu et al. [8]. The grain structure is refined slightly when the melt superheating temperature is over 1873 K. It is very unexpected that the superhigh temperature (2123 K) treatment of melt produces a very fine equiaxed grain structure. In order to verify this result, the same experiment was repeated and a similar result was obtained. In the as-cast condition, castings of M963 superalloy solidified under different melt superheating temperatures all consist of c solid solution matrix, c0 precipitate, (c þ c0 ) eutectic and MC carbide (see Fig. 2). The morphology and distribution of the MC carbide was varied by melt treatment. Without any melt superheating treatment, the MC carbide usually has a blocky appearance and distributes mainly in the interdendritic regions (Fig. 2a). In fact, this blocky carbide has an octahedral morphology in three-dimension
Fig. 1. Grain structures of M963 superalloy under the conditions of different melt superheating treatment. Superheating temperature of melt treatment: (a) without any melt superheating treatment, (b) 1823 K, (c) 1873 K, (d) 1923 K, (e) 2023 K, (f) 2123 K.
F.S. Yin et al. / Scripta Materialia 48 (2003) 425–429
427
Fig. 2. MC carbide morphology of M963 superalloy under the conditions of different melt superheating treatment. Superheating temperature of melt treatment: (a) without any melt superheating treatment, (b) 1823 K, (c) 1923 K, (d) 2123 K.
[8]. Under the condition of low melt superheating temperature (1823 K), the MC carbide becomes a highly developed Chinese-script type MC carbide (Fig. 2b). With further increases in the melt superheating temperature, the Chinese-script MC carbide becomes finer (Fig. 2c). When the melt superheating temperature reaches a superhigh temperature (2123 K), the blocky MC carbide is found again (Fig. 2d), but the size is smaller and more homogeneous than that without any melt superheating treatment. Observed in the SEM,
Table 1 Composition of the carbide with a core (in at pct) Element
Ti
Cr
Ni
Nb
W
Central part Periphery
58.1 46.7
2.6 2.17
5.26 7.21
16.7 23.0
17.3 20.9
many of the fine blocky carbide particles have a cored structure (Fig. 3). The composition of each part of the blocky carbide with a core was determined by EPMA. The results that are listed in Table 1 and shown in Fig. 4 indicate that the core is mainly a carbonitride. The gas analysis results show that nitrogen content in the sample cast with 2123 K melt superheat is 32 ppm, which is five times higher than that in master alloy (6 ppm).
4. Discussion
Fig. 3. A carbide particle present in the cast sample with 2123 K melt superheating treatment.
It is well known that the grain size is related to heterogeneous nucleation and the undercooling of the melt. The mechanism of grain refinement without any or low temperature (1823 K) superheating treatment of the melt can be attributed to
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F.S. Yin et al. / Scripta Materialia 48 (2003) 425–429
Fig. 4. Results of EDS of cored carbide particles indicate that the core is carbonitride: (a) central part, (b) periphery.
the presence of undissolved primary carbide particles in the melt [7]. Liu et al. has found that the MC carbide is very stable in IN738LC alloy melt above the liquidus temperature [7]. The results of a differential thermal analysis of the heating curves of IN738 alloy provided an indication of a reaction above the alloy melting point, between 1768 and 1868 K that is related to the dissolution of primary carbide [9]. By quenching the melt from 1768 K, they obtained a larger MC carbide size than for quenching from 1868 K, indicating that MC car-
bide will not completely dissolve below 1768 K. When the melt superheating temperature reaches 1873 K, the contribution of undissolved primary carbide to heterogeneous nucleation disappears and so the grain size becomes larger. With further increases in the melt superheating temperature (1873–2073 K), the distribution of alloying elements will be homogeneous due to the thermal diffusion. This will clearly affect the subsequent solidification process. It has been found that the undercooling of the melts will be increased about 40 K after the melt superheating treatment in nickel-base superalloys [10]. Therefore, the slight grain refinement after the melt superheating temperature over 1873 K can be attributed to the increment of the undercooling of the melt. Grain refinement due to the melt superheat has also been found in a Sb–4.6wt.%Bi alloy [11]. The unexpected grain refinement is very difficult to understand when the melt superheating temperature reaches 2123 K. This might be related to the increment of both the undercooling of melt and nitrogen content in the melt. The nitrogen combines with Ti to produce TiN, which can act as the nucleator of the grain nucleation. Numerous investigations [12–16] have characterized the formation sequence, composition and morphological evolution of MC carbides from slow to rapid solidification conditions. The results indicate that the processing parameters such as sample growth rate, thermal gradient, solidification interface morphology and chemical composition are important factors governing MC carbide growth morphology, size, distribution, composition and growth mechanism. It is now well understood that the near-equilibrium growth morphology of the MC carbide in superalloys is octahedral blocks, which gradually transform to a Chinese-script morphology as the carbide growth rate increases. As mentioned above, the primary carbide is very stable in superalloy melts above the liquidus temperature. Under the condition of no melt superheating treatment, there are many undissolved primary carbide particles in the melt, which are beneficial to the formation of MC carbide. In this condition, the MC carbide grows in the nearequilibrium mode into an octahedral blocky
F.S. Yin et al. / Scripta Materialia 48 (2003) 425–429
morphology (Fig. 2a). When the alloy melt is treated by high temperature (1823–2023 K) superheat, the primary carbide gradually dissolves into the melt and distribution of alloying elements becomes more homogeneous, therefore, the formation of MC carbide is postponed. In this condition, MC carbide forms by a (MC þ c) eutectic reaction and has the script-like shape (Figs. 3b and 2c). Moreover, the script-like MC carbide becomes finer due to the increment of the undercooling of melt with increasing melt superheating temperature. It is very unexpected that the script-like MC carbide transforms to the octahedral blocky one after the superhigh temperature (2123 K) treatment of melt. The results from SEM and EPMA (Fig. 3 and Table 1) indicate that the blocky MC carbide has a cored structure and the core is mainly TiN. The gas analysis results show that the nitrogen content in the alloy is increased greatly by the superhigh temperature treatment of melt. According to these results, the following assumption could be made: At superhigh temperature (2123 K), the alloy melt would absorb nitrogen from environment, because the solubility of nitrogen (N2 ) increases exponentially with increasing the temperature of melt [17]. Nitrogen combines with Ti in the alloy melt to produce TiN, which can act as the nucleator of nucleation of MC carbide. The MC carbide grows in a near-equilibrium mode again into an octahedral blocky morphology (Fig. 3d).
5. Conclusion The cast structures including grain structure and MC carbide morphology in M963 superalloy are greatly affected by the melt superheating treatment. Without any melt treatment, the grain structure is almost equiaxed and fine, and MC carbide usually has an octahedral block appearance distributed at interdendritic region. Melt superheating treatment coarsens the grain structure and transforms the blocky MC carbide into
429
Chinese-script one that mostly distributed in dendrite. When the melt superheating temperature reaches 2123 K, the grain structure becomes very fine equiaxed, and the morphology of MC carbide changes back to the blocky appearance.
Acknowledgements The authors are grateful to Prof. C.Z. Zhang, Northeast University, Prof. Y. Yu, Q. Zheng and G.C. Hou, Institute of Metal Resarch, Chinese Academy of Sciences, for their useful discussion and contributions.
References [1] Sims CT, Stoloff NS, Hagel WC, et al. In: Superalloy II [M]. New York: John Wiley & Sons, Inc.; 1987. p. 24. [2] Kotva PS, Venables JD, Calder RW. Metall Trans 1972;3:453–8. [3] Chou JC. AFS Trans 1987;87–168:734–7. [4] Chang E, Chou JC, Yin V. AFS Trans 1988;88–124:47–54. [5] Woulds M, Beson H. In: Gell M et al. (Eds.), Superalloys. The Metallurgical Society of AIME, 1984, 3. [6] Zhen BL, Liu L, Zhong ZY, Banerji A, Reif W. Metallurgy 1994;48:118–22. [7] Liu L, Zhen BL, Banerji A, Reif W, Sommer F. Scripta Metall 1994;30:593–8. [8] Liu L, Sommer F, Fu HZ. Scripta Metall 1994;30:587–91. [9] Bouse GK, Mihalisin JR. In: Superalloys, supercomposites and superceramics. Boston: Academic Press, 1989, 125. [10] Gui ZL. Aero Manuf Eng 1995;4:12–4 (in Chinese). [11] Chen G. Ph.D. Thesis. Northwestern Polytechnical University, 1995 (in Chinese). [12] Fernandez R, Lecomte JC, Kattamis TZ. Metall Trans 1978;9A:1381–6. [13] Liu L, Fu HZ, Shi ZX. Acta Metall Sinica (English edition) Series A, 1990;3:46–51. [14] Wang HM, Shang JH, Tang YJ, Hu ZQ. Mater Sci Eng A 1992;156:109–16. [15] Sun WR, Lee JH, Seo SM, Choe, Hu SJ. Mater Sci Eng A 1999;271:143–9. [16] Chen J, Lee JH, Jo CY, Choe SJ, Lee YT. Mater Sci Eng A 1998;247:113–25. [17] Winkle O, Bakisb R. In: Vacuum Metallurgy. Amsterdam, London, New York: Elsevier Publishing Company; 1971. p. 15.
Effects of melt treatment on the cast structure of M963 superalloy F.S. Yin a
a,b,*
, X.F. Sun a, J.G. Li a, H.R. Guan a, Z.Q. Hu
a
Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, P.R. China b College of Mechanical Engineering, Shandong University of Technology, Zibo 255012, P.R. China Received 6 June 2002; received in revised form 2 September 2002; accepted 4 September 2002
Abstract The cast structure of M963 superalloy is greatly affected by the melt superheating treatment. Melt superheating treatment (below 1923 K) coarsens the grain structure and transforms the blocky MC carbide into Chinese-script one that mostly distributed in dendrite. However, when the melt superheating temperature reaches 2123 K, the grain structure becomes very fine equiaxed, and the morphology of MC carbide changes back to the blocky appearance. Ó 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Cast; Nickel; Melt treatment; Carbide
1. Introduction The grain structure and primary MC carbide morphology have a very important effect on the mechanical properties of cast nickel-base superalloys [1,2]. They can be controlled by optimizing the process variables, such as melt treatment, including melt superheating temperature and the holding time at that temperature, as well as casting parameters including mold preheating temperature and pouring temperature of the melt. The effect of casting parameters has been extensively studied for several superalloys [3–7]. Liu et al. [7] have reported that the variation of melt superheat (below * Corresponding author. Address: Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, P.R. China. Tel.: +86-24-238-43531/55432; fax: +86-24-238-91320. E-mail address: [email protected] (F.S. Yin).
1823 K), i.e. the temperature above the liquidus of the alloy, produces a significant effect on the cast microstructure of IN738LC superalloy, and low melt superheating temperature causes significant reduction in the grain size. There is no report on the effect of high melt superheating temperature up to 2123 K on the cast structure of superalloy. This paper presents the results and mechanism of high temperature treatment of melt.
2. Experimental procedures The superalloy used in this study is M963 alloy having the following composition (wt. pct): 0.15C, 8.8Cr, 9.6Co, 10.5W, 1.6Mo, 1.0Nb, 5.8Al, 2.4Ti, 0.04Zr, 0.03B and rest Ni. The liquidus temperature of the alloy is 1623 K determined by DTA. Samples were prepared in a VIM-25F vacuum induction furnace. Approximately 4.8 kg charge
1359-6462/03/$ - see front matter Ó 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 6 2 ( 0 2 ) 0 0 4 4 6 - 3
426
F.S. Yin et al. / Scripta Materialia 48 (2003) 425–429
per heat was melted and superheated to different temperatures of up to 2123 K for 5 min and poured into a preheated ceramic mold, 15 mm in diameter. The mold preheating temperature and pouring temperature were kept at 1123 and 1723 K, respectively, in all cases. The vacuum during melt treatment was 0.1 Pa. The specimens for microstructure examination cut from the cast ingots were ground, polished and chemically etched with 40 ml HCl þ 40 ml C2 H5 OH þ 20 ml H2 O þ 1:5 g CuSO4 solution. The grain structure and primary MC carbide morphology were examined with standard metallographic techniques. An ISM-6301F scanning electron microscope (SEM) was used to reveal further microstructural details. Electron-probe microanalysis (EPMA) was used to determine the phase composition.
3. Results Fig. 1 shows the grain structures of the alloy cast with different melt superheating temperatures.
In Fig. 1, one can see that melt-superheating temperature has a significant effect on the grain size. Without any melt treatment, the grain structure is almost equiaxed and fine. Below 1873 K the grain size increases with increasing the melt superheating temperature, which is in accordance with the results of Liu et al. [8]. The grain structure is refined slightly when the melt superheating temperature is over 1873 K. It is very unexpected that the superhigh temperature (2123 K) treatment of melt produces a very fine equiaxed grain structure. In order to verify this result, the same experiment was repeated and a similar result was obtained. In the as-cast condition, castings of M963 superalloy solidified under different melt superheating temperatures all consist of c solid solution matrix, c0 precipitate, (c þ c0 ) eutectic and MC carbide (see Fig. 2). The morphology and distribution of the MC carbide was varied by melt treatment. Without any melt superheating treatment, the MC carbide usually has a blocky appearance and distributes mainly in the interdendritic regions (Fig. 2a). In fact, this blocky carbide has an octahedral morphology in three-dimension
Fig. 1. Grain structures of M963 superalloy under the conditions of different melt superheating treatment. Superheating temperature of melt treatment: (a) without any melt superheating treatment, (b) 1823 K, (c) 1873 K, (d) 1923 K, (e) 2023 K, (f) 2123 K.
F.S. Yin et al. / Scripta Materialia 48 (2003) 425–429
427
Fig. 2. MC carbide morphology of M963 superalloy under the conditions of different melt superheating treatment. Superheating temperature of melt treatment: (a) without any melt superheating treatment, (b) 1823 K, (c) 1923 K, (d) 2123 K.
[8]. Under the condition of low melt superheating temperature (1823 K), the MC carbide becomes a highly developed Chinese-script type MC carbide (Fig. 2b). With further increases in the melt superheating temperature, the Chinese-script MC carbide becomes finer (Fig. 2c). When the melt superheating temperature reaches a superhigh temperature (2123 K), the blocky MC carbide is found again (Fig. 2d), but the size is smaller and more homogeneous than that without any melt superheating treatment. Observed in the SEM,
Table 1 Composition of the carbide with a core (in at pct) Element
Ti
Cr
Ni
Nb
W
Central part Periphery
58.1 46.7
2.6 2.17
5.26 7.21
16.7 23.0
17.3 20.9
many of the fine blocky carbide particles have a cored structure (Fig. 3). The composition of each part of the blocky carbide with a core was determined by EPMA. The results that are listed in Table 1 and shown in Fig. 4 indicate that the core is mainly a carbonitride. The gas analysis results show that nitrogen content in the sample cast with 2123 K melt superheat is 32 ppm, which is five times higher than that in master alloy (6 ppm).
4. Discussion
Fig. 3. A carbide particle present in the cast sample with 2123 K melt superheating treatment.
It is well known that the grain size is related to heterogeneous nucleation and the undercooling of the melt. The mechanism of grain refinement without any or low temperature (1823 K) superheating treatment of the melt can be attributed to
428
F.S. Yin et al. / Scripta Materialia 48 (2003) 425–429
Fig. 4. Results of EDS of cored carbide particles indicate that the core is carbonitride: (a) central part, (b) periphery.
the presence of undissolved primary carbide particles in the melt [7]. Liu et al. has found that the MC carbide is very stable in IN738LC alloy melt above the liquidus temperature [7]. The results of a differential thermal analysis of the heating curves of IN738 alloy provided an indication of a reaction above the alloy melting point, between 1768 and 1868 K that is related to the dissolution of primary carbide [9]. By quenching the melt from 1768 K, they obtained a larger MC carbide size than for quenching from 1868 K, indicating that MC car-
bide will not completely dissolve below 1768 K. When the melt superheating temperature reaches 1873 K, the contribution of undissolved primary carbide to heterogeneous nucleation disappears and so the grain size becomes larger. With further increases in the melt superheating temperature (1873–2073 K), the distribution of alloying elements will be homogeneous due to the thermal diffusion. This will clearly affect the subsequent solidification process. It has been found that the undercooling of the melts will be increased about 40 K after the melt superheating treatment in nickel-base superalloys [10]. Therefore, the slight grain refinement after the melt superheating temperature over 1873 K can be attributed to the increment of the undercooling of the melt. Grain refinement due to the melt superheat has also been found in a Sb–4.6wt.%Bi alloy [11]. The unexpected grain refinement is very difficult to understand when the melt superheating temperature reaches 2123 K. This might be related to the increment of both the undercooling of melt and nitrogen content in the melt. The nitrogen combines with Ti to produce TiN, which can act as the nucleator of the grain nucleation. Numerous investigations [12–16] have characterized the formation sequence, composition and morphological evolution of MC carbides from slow to rapid solidification conditions. The results indicate that the processing parameters such as sample growth rate, thermal gradient, solidification interface morphology and chemical composition are important factors governing MC carbide growth morphology, size, distribution, composition and growth mechanism. It is now well understood that the near-equilibrium growth morphology of the MC carbide in superalloys is octahedral blocks, which gradually transform to a Chinese-script morphology as the carbide growth rate increases. As mentioned above, the primary carbide is very stable in superalloy melts above the liquidus temperature. Under the condition of no melt superheating treatment, there are many undissolved primary carbide particles in the melt, which are beneficial to the formation of MC carbide. In this condition, the MC carbide grows in the nearequilibrium mode into an octahedral blocky
F.S. Yin et al. / Scripta Materialia 48 (2003) 425–429
morphology (Fig. 2a). When the alloy melt is treated by high temperature (1823–2023 K) superheat, the primary carbide gradually dissolves into the melt and distribution of alloying elements becomes more homogeneous, therefore, the formation of MC carbide is postponed. In this condition, MC carbide forms by a (MC þ c) eutectic reaction and has the script-like shape (Figs. 3b and 2c). Moreover, the script-like MC carbide becomes finer due to the increment of the undercooling of melt with increasing melt superheating temperature. It is very unexpected that the script-like MC carbide transforms to the octahedral blocky one after the superhigh temperature (2123 K) treatment of melt. The results from SEM and EPMA (Fig. 3 and Table 1) indicate that the blocky MC carbide has a cored structure and the core is mainly TiN. The gas analysis results show that the nitrogen content in the alloy is increased greatly by the superhigh temperature treatment of melt. According to these results, the following assumption could be made: At superhigh temperature (2123 K), the alloy melt would absorb nitrogen from environment, because the solubility of nitrogen (N2 ) increases exponentially with increasing the temperature of melt [17]. Nitrogen combines with Ti in the alloy melt to produce TiN, which can act as the nucleator of nucleation of MC carbide. The MC carbide grows in a near-equilibrium mode again into an octahedral blocky morphology (Fig. 3d).
5. Conclusion The cast structures including grain structure and MC carbide morphology in M963 superalloy are greatly affected by the melt superheating treatment. Without any melt treatment, the grain structure is almost equiaxed and fine, and MC carbide usually has an octahedral block appearance distributed at interdendritic region. Melt superheating treatment coarsens the grain structure and transforms the blocky MC carbide into
429
Chinese-script one that mostly distributed in dendrite. When the melt superheating temperature reaches 2123 K, the grain structure becomes very fine equiaxed, and the morphology of MC carbide changes back to the blocky appearance.
Acknowledgements The authors are grateful to Prof. C.Z. Zhang, Northeast University, Prof. Y. Yu, Q. Zheng and G.C. Hou, Institute of Metal Resarch, Chinese Academy of Sciences, for their useful discussion and contributions.
References [1] Sims CT, Stoloff NS, Hagel WC, et al. In: Superalloy II [M]. New York: John Wiley & Sons, Inc.; 1987. p. 24. [2] Kotva PS, Venables JD, Calder RW. Metall Trans 1972;3:453–8. [3] Chou JC. AFS Trans 1987;87–168:734–7. [4] Chang E, Chou JC, Yin V. AFS Trans 1988;88–124:47–54. [5] Woulds M, Beson H. In: Gell M et al. (Eds.), Superalloys. The Metallurgical Society of AIME, 1984, 3. [6] Zhen BL, Liu L, Zhong ZY, Banerji A, Reif W. Metallurgy 1994;48:118–22. [7] Liu L, Zhen BL, Banerji A, Reif W, Sommer F. Scripta Metall 1994;30:593–8. [8] Liu L, Sommer F, Fu HZ. Scripta Metall 1994;30:587–91. [9] Bouse GK, Mihalisin JR. In: Superalloys, supercomposites and superceramics. Boston: Academic Press, 1989, 125. [10] Gui ZL. Aero Manuf Eng 1995;4:12–4 (in Chinese). [11] Chen G. Ph.D. Thesis. Northwestern Polytechnical University, 1995 (in Chinese). [12] Fernandez R, Lecomte JC, Kattamis TZ. Metall Trans 1978;9A:1381–6. [13] Liu L, Fu HZ, Shi ZX. Acta Metall Sinica (English edition) Series A, 1990;3:46–51. [14] Wang HM, Shang JH, Tang YJ, Hu ZQ. Mater Sci Eng A 1992;156:109–16. [15] Sun WR, Lee JH, Seo SM, Choe, Hu SJ. Mater Sci Eng A 1999;271:143–9. [16] Chen J, Lee JH, Jo CY, Choe SJ, Lee YT. Mater Sci Eng A 1998;247:113–25. [17] Winkle O, Bakisb R. In: Vacuum Metallurgy. Amsterdam, London, New York: Elsevier Publishing Company; 1971. p. 15.