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Effect of melt treatment on carbides formation in a cast nickel-base superalloy M963
Journal of Materials Processing Technology 183 (2007) 440–444
Effect of melt treatment on carbides formation in a cast nickel-base superalloy M963 F.S. Yin a,∗ , Q. Zheng b , X.F. Sun b , H.R. Guan b , Z.Q. Hu b a
School of Mechanical Engineering, Shandong University of Technology, 12 Zhangzhou Road, Zibo 255049, PR China b Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, PR China Received 18 December 2003; received in revised form 29 May 2006; accepted 27 October 2006
Abstract The microstructures of a cast nickel-base superalloy M963 with both no melt treatment and melt treatment at 1923 K for 5 min were investigated in detail. The results show that the melt treatment has a significant role on the carbides formation in the alloy. The MC carbide tends to have blocky morphology in the alloy with no melt treatment, while it has a Chinese script-like morphology in the alloy with melt treatment at 1923 K for 5 min. During the solution heat treatment at 1248 K for 4 h followed by air-cooling, much more acicular M6 C carbide precipitates in the alloy with no melt treatment. © 2006 Elsevier B.V. All rights reserved. Keywords: Superalloy; Carbides; Melt treatment
1. Introduction The carbide phases have an important role on the mechanical properties of superalloys [1–3]. They contribute to strengthening of grain boundaries (GBs) at elevated temperature. The strengthening role of carbides depends significantly on their geometry and distribution. Three types of carbides, MC, M6 C and M23 C6 , are found in cast nickel-base superalloys. The MCtype carbide, which forms during solidification, will decompose to M6 C or/and M23 C6 during heat treatment process and service at high temperature [1]. Many investigations have characterized the composition, morphology and growth mechanism of these carbides existing in nickel-base superalloys [4–8]. However, no report was made on the effect of melt treatment on the formation of these carbides. M963 is a cast equiaxed Ni-base superalloy, which has excellent high temperature strength and temperature capability mainly due to the addition of a large amount of the refractory elements W, Mo and Nb [9]. However, this superalloy usually has low ductility, especially under the service condition of high temperature. In previous investigation [10], we have found that the melt superheating temperature has a great effect on the microstructure and high temperature stress rupture property of the M963 alloy. It ∗
Corresponding author. Tel.: +86 533 2786910. E-mail address: [email protected] (F.S. Yin).
0924-0136/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2006.11.002
was found that the melt superheating treatment at 1873–1923 K for 5 min resulted in a great improvement in stress rupture property at 1248 K under 225 MPa. In order to well understand the effect of the melt superheating treatment on the mechanical properties of this alloy, the microstructure of the alloy with two different melting schedules was examined in detail in both as-cast and heat-treated conditions. 2. Experimental The M963 master alloy used in this study was prepared by vacuum induction melting and cast into round ingots, then cut into pieces. Each piece weighs 5 kg. The master alloy has the following nominal composition (wt.%): 0.15C, 8.8Cr, 9.6Co, 10.5W, 1.6Mo, 1.0Nb, 5.8Al, 2.4Ti and rest Ni. The liquidus temperature of the alloy is about 1623 K. The pieces were remelted in a vacuum furnace and poured in a preheated ceramic mold, 15 mm in diameter. During the melting process, one heat was superheated to 1923 K and holding 5 min at that temperature, which is marked as melt treatment, while the another heat was only heated to pouring temperature and then poured without any holding at that temperature, marked as no melt treatment. Both the pouring temperature and mold preheating temperature were kept the same in the two cases. The casting bars were solution heat-treated by standard process of 1483 K for 4 h followed by air-cooling. The microstructure examination was made on JSM-6301F scanning electron microscopy (SEM) equipped with energy dispersive spectrometer (EDS) for microanalysis and Philips EM420 transmission electron microscope (TEM). The TEM foil was prepared by twin-jet thinning electrolytically in a solution of 7% perchloric acid and 93% ethanol at 253 K. The elements distribution in as-cast state was determined by electron probe microanalysis (EPMA). Three points were analyzed to get the average of compositions at interdendritic region or dendrite arm.
F.S. Yin et al. / Journal of Materials Processing Technology 183 (2007) 440–444
3. Results 3.1. As-cast microstructure The microstructure of the cast M963 superalloy consists of ␥ solid solution matrix, ␥ precipitate (␥ + ␥ ) eutectic and MC carbide [11]. The effect of melt superheating temperature up to 2123 K on the cast structure of this alloy has been reported elsewhere [11]. One of the striking differences between cast microstructures of the alloy with both no melt treatment and melt treatment is primary MC carbide morphology as shown in Fig. 1. The MC carbide tends to have blocky morphology in the alloy without melt treatment (Fig. 1a), while the MC carbide in the alloy with melt treatment at 1923 K for 5 min all has a Chinese script-like morphology no matter what the pouring and mold preheating temperature are (Fig. 1b).
441
The elements distribution determined by EPMA (Table 1) reveals that the melt treatment at 1923 K for 5 min also has effect on the segregation of the alloy. The segregation of elements Cr and Mo was reduced by the melt treatment, while the segregation of elements Ti and Nb was enhanced, i.e., the Ti and Nb have a more tendency to distribute in the interdendritic regions in the melt-treated alloy, which is unexpected. The segregation of element W was further increased by the melt treatment. The melt treatment has little effect on the segregation of the rest elements Al, Co and Ni. 3.2. Heat-treated microstructure The SEM backscatter electron micrographs of the alloys solution heat-treated at 1483 K for 4 h, followed by air-cooling are shown in Fig. 2. During the solution treatment, besides the
Fig. 1. MC carbide morphology in as-cast M963 superalloys (a) no melt treatment and (b) melt treatment at 1923 K for 5 min. Table 1 Elements distribution in as-cast M963 superalloy (wt.%) Elements Al
Ti
Nb
Cr
Mo
W
Co
Ni
No melt treatment Interdendritic region (Ci ) Dendrite arm (Cd ) Ci /Cd
4.18 4.22 0.99
2.08 1.66 1.25
0.56 0.59 0.94
12.52 8.10 1.54
2.16 1.33 1.63
8.29 12.90 0.64
10.73 10.86 0.99
59.48 60.33 0.99
Melt treatment at 1923 K for 5 min Interdendritic region (Ci ) Dendrite arm (Cd ) Ci /Cd
5.17 4.71 1.10
2.56 1.73 1.48
0.85 0.46 1.85
10.41 9.26 1.12
1.93 1.47 1.31
7.42 12.76 0.58
10.08 10.52 0.96
61.59 59.09 1.04
Fig. 2. The microstructure of the alloys heat treated at 1483 K for 4 h followed by AC; (a) no melt treatment and (b) melt treatment at 1923 K for 5 min.
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F.S. Yin et al. / Journal of Materials Processing Technology 183 (2007) 440–444
Table 2 EDS results of the carbides in M963 superalloy (wt.%) Elements Al
Ti
Nb
Cr
Mo
W
Co
Ni
Phases
No melt treatment
0.00 0.52 0.63 0.62
24.06 1.23 3.11 1.50
19.72 1.83 3.05 2.45
2.00 11.72 8.98 9.33
0.00 6.65 6.49 5.95
50.49 54.71 54.93 57.00
0.00 7.07 6.39 7.55
3.73 16.27 16.42 15.59
Blocky MC Acicular M6 C M6 C around MC M6 C in GBs
Melt treatment
0.00 0.92
17.78 1.07
18.23 0.60
3.87 6.69
5.06 2.22
42.09 57.26
0.00 5.90
3.87 25.33
Script-like MC M6 C around MC
Fig. 3. MC and M6 C carbides in heat-treated alloy with melt treatment at 1923 K for 5 min (a) dark field, (b) diffraction pattern and (c) interpretation of the pattern.
Fig. 4. Acicular M6 C in the heat treated alloy with no melt treatment; (a) bright field, (b) diffraction pattern and (c) interpretation of the pattern.
F.S. Yin et al. / Journal of Materials Processing Technology 183 (2007) 440–444
dissolution of ␥ phase and re-precipitation in finer form, the decomposition of MC carbide and secondary carbide precipitation in grain boundaries and interdendritic regions also occur [12]. Careful observation made on the heat-treated specimens reveals an interesting result that there is much more acicular phase distributed in interdendritic regions in the alloy without melt treatment (Fig. 2a), while much less acicular phase was found in the specimen with the melt treatment (Fig. 2b). The compositions of the carbides determined by EDS are summarized in Table 2. The acicular phase has the similar composition to the particular carbide precipitated both in GBs and around MC carbide particles. The selected area electron diffraction identifies that both the acicular and particular phases are all M6 C-type carbide as shown in Figs. 3 and 4. 4. Discussion The melt state of metal plays an important role in the production of castings. It is well known that liquid metals are characterized by heterogeneities that are greatly affected by melt treatment. The treatment may cause dissolution of these heterogeneities or their growth or the formation of new ones. This can be used for the control of the microstructure and the properties in the solid state. Numerous investigations have characterized the formation sequence, composition and morphological evolution of MC carbide from slow to rapid solidification conditions [4–8]. It is now well understood that the near-equilibrium growth morphology of the MC carbide in superalloy is octahedral blocks, which gradually transform to a Chinese-script morphology as carbide growth rate increases. In previous investigation on the liquid structure of the M963 superalloy by high temperature X-ray diffraction, it was found that the liquid structure of the alloy has a close relationship to the thermal history of the alloy melts [13]. There exist Ni3 (Al,Ti,Nb)-like cluster and residual MC carbide or (Ti,Nb)C cluster in the alloy with no melt treatment. When the superheating temperature reaches 1923 K, the Ni3 (Al,Ti,Nb)-like cluster and residual MC carbide are all disappeared. Under the condition of no melt treatment, the residual carbide particles in the melt might be beneficial to the formation of MC carbide. In this condition, the MC carbide grows in the near-equilibrium mode into an octahedral blocky morphology (Fig. 1a). When the alloy melt is treated by 1923 K superheat for 5 min, the primary carbide gradually
443
dissolves into the melt and distribution of alloying elements becomes more homogeneous, therefore, the formation of MC carbide is postponed and the undercooling of the solidification increases. In this condition, MC carbide is formed by a (MC + ␥) eutectic reaction and so has the script-like shape (Fig. 1b). The investigations on segregation in cast Ni-base superalloys revealed that Cr and Mo are a little rich, Ti and Nb heavily rich in interdendritic regions, while dendrite core is rich in W. In present work, it was found that the melt treatment at 1923 K for 5 min has different effects on the segregation of various elements, which can be explained by the change of liquid structure owing to the melt treatment. Under the condition of no melt treatment, Ti and Nb, existing in the form of both Ni3 (Al,Ti,Nb)-like cluster and (Ti,Nb)C cluster, have low diffusibility and are easy to be trapped during the dendrite growth as shown in Fig. 5a. When the alloy melt was treated by 1923 K superheat for 5 min, the Ni3 (Al,Ti,Nb)-like cluster and (Ti,Nb)C cluster dissolve into the melt, which increases the diffusibility of Ti and Nb, and become more difficult to be trapped during the dendrite growth as shown in Fig. 5b. Therefore, the segregation of Ti and Nb in the melt-treated alloy is higher than that in the alloy with no melt treatment. Because of the enrichment of Ti and Nb in interdendritic region, the enrichment of Cr, Mo and W in interdendritic region is inhibited. Thus, the segregation of Cr and Mo is reduced by the melt treatment. Because W is an element rich in dendrite core, its segregation is further increased by the melt treatment. It is hard to understand that why the acicular M6 C is easier to precipitate in the alloy with no melt treatment during the solution heat treatment. It is well known that the selection of phase geometry depends on both the kinetic and thermodynamic factors. Mo and W are the elements of M6 C-type carbide formation. In the specimen of alloy with no melt treatment, the concentration of Mo and W in interdendritic regions is relatively higher as shown in Table 1, which makes the M6 C carbide formation favorable on thermodynamic. However, because Mo and W all belong to heavy elements and have low diffusibility, it is difficult for M6 C carbide to precipitate in a particulate form in inderdendritic regions. The melt treatment at 1923 K for 5 min decreases the segregation tendency of Mo and W to interdendritic regions, therefore, the acicular M6 C precipitation is inhibited obviously.
Fig. 5. Sketch of dendrite growth process in alloy with (a) no melt treatment and (b) melt treatment at 1873 K for 5 min.
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5. Conclusion The microstructures of the M963 alloy with both no melt treatment and melt treatment at 1923 K for 5min were investigated in detail. The following conclusion can be drawn: (1) The MC carbide tends to have blocky morphology in the alloy with no melt treatment, while it has a Chinese scriptlike morphology in the alloy with melt treatment at 1923 K for 5 min. (2) The segregation of elements Cr and Mo was reduced by the melt treatment, while the segregation of elements Ti and Nb was enhanced in the melt-treated alloy. The segregation of element W was further increased by the melt treatment. The melt treatment has little effect on the segregation of the rest elements Al, Co and Ni. (3) During the solution heat treatment at 1248 K for 4 h followed by air-cooling, much more acicular M6 C carbide precipitates in the alloy with no melt treatment. References [1] C.T. Sims, N.S. Stoloff, W.C. Hagel, Superalloys II, John Wiley & Sons Inc., New York, 1987, p. 73.
[2] L.R. Liu, T. Jin, N.R. Zhao, X.F. Sun, H.R. Guan, Z.Q. Hu, Mater. Sci. Eng. A 361 (2003) 191–197. [3] P.S. Kotva, J.D. Venables, R.W. Calder, Metall. Trans. 3 (1972) 453– 458. [4] R. Fernandez, J.C. Lecomte, T.Z. Kattamis, Metall. Trans. 9A (1978) 1381–1386. [5] L. Liu, H.Z. Fu, Z.X. Shi. Acta Metall. Sinica (English edition) Series A, 3(1990)46–51. [6] H.M. Wang, J.H. Shang, Y.J. Tang, Z.Q. Hu, Mater. Sci. Eng. A 156 (1992) 109–116. [7] W.R. Sun, J.H. Lee, S.M. Seo, S.J. Choe, Z.Q. Hu, Mater. Sci. Eng. A 271 (1999) 143–149. [8] J. Chen, J.H. Lee, C.Y. Jo, S.J. Choe, Y.T. Lee, Mater. Sci. Eng. A 247 (1998) 113–125. [9] C. Yuan, X.F. Sun, F.S. Yin, H.R. Guan, Z.Q. Hu, Q. Zheng, Y. Yu, J. Mater. Sci. Technol. 17 (2001) 425. [10] F.S. Yin, X.F. Sun, Y.B. Li, Y. Yu, Q. Zheng, H.R. Guan, Z.Q. Hu, Acta Metall. Sin. 39 (2003) 75–78. [11] F.S. Yin, X.F. Sun, J.G. Li, H.R. Guan, Z.Q. Hu, Scripta Metall. 48 (2003) 425–429. [12] F.S. Yin, X.F. Sun, C. Yuan, H.R. Guan, Z.Q. Hu, Trans. Nonferr. Metal. Soc. China 12 (2002) 83–87. [13] F.S. Yin, X.F. Sun, H.R. Guan, Z.Q. Hu, J Alloys Compd. 364 (2004) 225–228.
Effect of melt treatment on carbides formation in a cast nickel-base superalloy M963 F.S. Yin a,∗ , Q. Zheng b , X.F. Sun b , H.R. Guan b , Z.Q. Hu b a
School of Mechanical Engineering, Shandong University of Technology, 12 Zhangzhou Road, Zibo 255049, PR China b Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, PR China Received 18 December 2003; received in revised form 29 May 2006; accepted 27 October 2006
Abstract The microstructures of a cast nickel-base superalloy M963 with both no melt treatment and melt treatment at 1923 K for 5 min were investigated in detail. The results show that the melt treatment has a significant role on the carbides formation in the alloy. The MC carbide tends to have blocky morphology in the alloy with no melt treatment, while it has a Chinese script-like morphology in the alloy with melt treatment at 1923 K for 5 min. During the solution heat treatment at 1248 K for 4 h followed by air-cooling, much more acicular M6 C carbide precipitates in the alloy with no melt treatment. © 2006 Elsevier B.V. All rights reserved. Keywords: Superalloy; Carbides; Melt treatment
1. Introduction The carbide phases have an important role on the mechanical properties of superalloys [1–3]. They contribute to strengthening of grain boundaries (GBs) at elevated temperature. The strengthening role of carbides depends significantly on their geometry and distribution. Three types of carbides, MC, M6 C and M23 C6 , are found in cast nickel-base superalloys. The MCtype carbide, which forms during solidification, will decompose to M6 C or/and M23 C6 during heat treatment process and service at high temperature [1]. Many investigations have characterized the composition, morphology and growth mechanism of these carbides existing in nickel-base superalloys [4–8]. However, no report was made on the effect of melt treatment on the formation of these carbides. M963 is a cast equiaxed Ni-base superalloy, which has excellent high temperature strength and temperature capability mainly due to the addition of a large amount of the refractory elements W, Mo and Nb [9]. However, this superalloy usually has low ductility, especially under the service condition of high temperature. In previous investigation [10], we have found that the melt superheating temperature has a great effect on the microstructure and high temperature stress rupture property of the M963 alloy. It ∗
Corresponding author. Tel.: +86 533 2786910. E-mail address: [email protected] (F.S. Yin).
0924-0136/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2006.11.002
was found that the melt superheating treatment at 1873–1923 K for 5 min resulted in a great improvement in stress rupture property at 1248 K under 225 MPa. In order to well understand the effect of the melt superheating treatment on the mechanical properties of this alloy, the microstructure of the alloy with two different melting schedules was examined in detail in both as-cast and heat-treated conditions. 2. Experimental The M963 master alloy used in this study was prepared by vacuum induction melting and cast into round ingots, then cut into pieces. Each piece weighs 5 kg. The master alloy has the following nominal composition (wt.%): 0.15C, 8.8Cr, 9.6Co, 10.5W, 1.6Mo, 1.0Nb, 5.8Al, 2.4Ti and rest Ni. The liquidus temperature of the alloy is about 1623 K. The pieces were remelted in a vacuum furnace and poured in a preheated ceramic mold, 15 mm in diameter. During the melting process, one heat was superheated to 1923 K and holding 5 min at that temperature, which is marked as melt treatment, while the another heat was only heated to pouring temperature and then poured without any holding at that temperature, marked as no melt treatment. Both the pouring temperature and mold preheating temperature were kept the same in the two cases. The casting bars were solution heat-treated by standard process of 1483 K for 4 h followed by air-cooling. The microstructure examination was made on JSM-6301F scanning electron microscopy (SEM) equipped with energy dispersive spectrometer (EDS) for microanalysis and Philips EM420 transmission electron microscope (TEM). The TEM foil was prepared by twin-jet thinning electrolytically in a solution of 7% perchloric acid and 93% ethanol at 253 K. The elements distribution in as-cast state was determined by electron probe microanalysis (EPMA). Three points were analyzed to get the average of compositions at interdendritic region or dendrite arm.
F.S. Yin et al. / Journal of Materials Processing Technology 183 (2007) 440–444
3. Results 3.1. As-cast microstructure The microstructure of the cast M963 superalloy consists of ␥ solid solution matrix, ␥ precipitate (␥ + ␥ ) eutectic and MC carbide [11]. The effect of melt superheating temperature up to 2123 K on the cast structure of this alloy has been reported elsewhere [11]. One of the striking differences between cast microstructures of the alloy with both no melt treatment and melt treatment is primary MC carbide morphology as shown in Fig. 1. The MC carbide tends to have blocky morphology in the alloy without melt treatment (Fig. 1a), while the MC carbide in the alloy with melt treatment at 1923 K for 5 min all has a Chinese script-like morphology no matter what the pouring and mold preheating temperature are (Fig. 1b).
441
The elements distribution determined by EPMA (Table 1) reveals that the melt treatment at 1923 K for 5 min also has effect on the segregation of the alloy. The segregation of elements Cr and Mo was reduced by the melt treatment, while the segregation of elements Ti and Nb was enhanced, i.e., the Ti and Nb have a more tendency to distribute in the interdendritic regions in the melt-treated alloy, which is unexpected. The segregation of element W was further increased by the melt treatment. The melt treatment has little effect on the segregation of the rest elements Al, Co and Ni. 3.2. Heat-treated microstructure The SEM backscatter electron micrographs of the alloys solution heat-treated at 1483 K for 4 h, followed by air-cooling are shown in Fig. 2. During the solution treatment, besides the
Fig. 1. MC carbide morphology in as-cast M963 superalloys (a) no melt treatment and (b) melt treatment at 1923 K for 5 min. Table 1 Elements distribution in as-cast M963 superalloy (wt.%) Elements Al
Ti
Nb
Cr
Mo
W
Co
Ni
No melt treatment Interdendritic region (Ci ) Dendrite arm (Cd ) Ci /Cd
4.18 4.22 0.99
2.08 1.66 1.25
0.56 0.59 0.94
12.52 8.10 1.54
2.16 1.33 1.63
8.29 12.90 0.64
10.73 10.86 0.99
59.48 60.33 0.99
Melt treatment at 1923 K for 5 min Interdendritic region (Ci ) Dendrite arm (Cd ) Ci /Cd
5.17 4.71 1.10
2.56 1.73 1.48
0.85 0.46 1.85
10.41 9.26 1.12
1.93 1.47 1.31
7.42 12.76 0.58
10.08 10.52 0.96
61.59 59.09 1.04
Fig. 2. The microstructure of the alloys heat treated at 1483 K for 4 h followed by AC; (a) no melt treatment and (b) melt treatment at 1923 K for 5 min.
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F.S. Yin et al. / Journal of Materials Processing Technology 183 (2007) 440–444
Table 2 EDS results of the carbides in M963 superalloy (wt.%) Elements Al
Ti
Nb
Cr
Mo
W
Co
Ni
Phases
No melt treatment
0.00 0.52 0.63 0.62
24.06 1.23 3.11 1.50
19.72 1.83 3.05 2.45
2.00 11.72 8.98 9.33
0.00 6.65 6.49 5.95
50.49 54.71 54.93 57.00
0.00 7.07 6.39 7.55
3.73 16.27 16.42 15.59
Blocky MC Acicular M6 C M6 C around MC M6 C in GBs
Melt treatment
0.00 0.92
17.78 1.07
18.23 0.60
3.87 6.69
5.06 2.22
42.09 57.26
0.00 5.90
3.87 25.33
Script-like MC M6 C around MC
Fig. 3. MC and M6 C carbides in heat-treated alloy with melt treatment at 1923 K for 5 min (a) dark field, (b) diffraction pattern and (c) interpretation of the pattern.
Fig. 4. Acicular M6 C in the heat treated alloy with no melt treatment; (a) bright field, (b) diffraction pattern and (c) interpretation of the pattern.
F.S. Yin et al. / Journal of Materials Processing Technology 183 (2007) 440–444
dissolution of ␥ phase and re-precipitation in finer form, the decomposition of MC carbide and secondary carbide precipitation in grain boundaries and interdendritic regions also occur [12]. Careful observation made on the heat-treated specimens reveals an interesting result that there is much more acicular phase distributed in interdendritic regions in the alloy without melt treatment (Fig. 2a), while much less acicular phase was found in the specimen with the melt treatment (Fig. 2b). The compositions of the carbides determined by EDS are summarized in Table 2. The acicular phase has the similar composition to the particular carbide precipitated both in GBs and around MC carbide particles. The selected area electron diffraction identifies that both the acicular and particular phases are all M6 C-type carbide as shown in Figs. 3 and 4. 4. Discussion The melt state of metal plays an important role in the production of castings. It is well known that liquid metals are characterized by heterogeneities that are greatly affected by melt treatment. The treatment may cause dissolution of these heterogeneities or their growth or the formation of new ones. This can be used for the control of the microstructure and the properties in the solid state. Numerous investigations have characterized the formation sequence, composition and morphological evolution of MC carbide from slow to rapid solidification conditions [4–8]. It is now well understood that the near-equilibrium growth morphology of the MC carbide in superalloy is octahedral blocks, which gradually transform to a Chinese-script morphology as carbide growth rate increases. In previous investigation on the liquid structure of the M963 superalloy by high temperature X-ray diffraction, it was found that the liquid structure of the alloy has a close relationship to the thermal history of the alloy melts [13]. There exist Ni3 (Al,Ti,Nb)-like cluster and residual MC carbide or (Ti,Nb)C cluster in the alloy with no melt treatment. When the superheating temperature reaches 1923 K, the Ni3 (Al,Ti,Nb)-like cluster and residual MC carbide are all disappeared. Under the condition of no melt treatment, the residual carbide particles in the melt might be beneficial to the formation of MC carbide. In this condition, the MC carbide grows in the near-equilibrium mode into an octahedral blocky morphology (Fig. 1a). When the alloy melt is treated by 1923 K superheat for 5 min, the primary carbide gradually
443
dissolves into the melt and distribution of alloying elements becomes more homogeneous, therefore, the formation of MC carbide is postponed and the undercooling of the solidification increases. In this condition, MC carbide is formed by a (MC + ␥) eutectic reaction and so has the script-like shape (Fig. 1b). The investigations on segregation in cast Ni-base superalloys revealed that Cr and Mo are a little rich, Ti and Nb heavily rich in interdendritic regions, while dendrite core is rich in W. In present work, it was found that the melt treatment at 1923 K for 5 min has different effects on the segregation of various elements, which can be explained by the change of liquid structure owing to the melt treatment. Under the condition of no melt treatment, Ti and Nb, existing in the form of both Ni3 (Al,Ti,Nb)-like cluster and (Ti,Nb)C cluster, have low diffusibility and are easy to be trapped during the dendrite growth as shown in Fig. 5a. When the alloy melt was treated by 1923 K superheat for 5 min, the Ni3 (Al,Ti,Nb)-like cluster and (Ti,Nb)C cluster dissolve into the melt, which increases the diffusibility of Ti and Nb, and become more difficult to be trapped during the dendrite growth as shown in Fig. 5b. Therefore, the segregation of Ti and Nb in the melt-treated alloy is higher than that in the alloy with no melt treatment. Because of the enrichment of Ti and Nb in interdendritic region, the enrichment of Cr, Mo and W in interdendritic region is inhibited. Thus, the segregation of Cr and Mo is reduced by the melt treatment. Because W is an element rich in dendrite core, its segregation is further increased by the melt treatment. It is hard to understand that why the acicular M6 C is easier to precipitate in the alloy with no melt treatment during the solution heat treatment. It is well known that the selection of phase geometry depends on both the kinetic and thermodynamic factors. Mo and W are the elements of M6 C-type carbide formation. In the specimen of alloy with no melt treatment, the concentration of Mo and W in interdendritic regions is relatively higher as shown in Table 1, which makes the M6 C carbide formation favorable on thermodynamic. However, because Mo and W all belong to heavy elements and have low diffusibility, it is difficult for M6 C carbide to precipitate in a particulate form in inderdendritic regions. The melt treatment at 1923 K for 5 min decreases the segregation tendency of Mo and W to interdendritic regions, therefore, the acicular M6 C precipitation is inhibited obviously.
Fig. 5. Sketch of dendrite growth process in alloy with (a) no melt treatment and (b) melt treatment at 1873 K for 5 min.
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5. Conclusion The microstructures of the M963 alloy with both no melt treatment and melt treatment at 1923 K for 5min were investigated in detail. The following conclusion can be drawn: (1) The MC carbide tends to have blocky morphology in the alloy with no melt treatment, while it has a Chinese scriptlike morphology in the alloy with melt treatment at 1923 K for 5 min. (2) The segregation of elements Cr and Mo was reduced by the melt treatment, while the segregation of elements Ti and Nb was enhanced in the melt-treated alloy. The segregation of element W was further increased by the melt treatment. The melt treatment has little effect on the segregation of the rest elements Al, Co and Ni. (3) During the solution heat treatment at 1248 K for 4 h followed by air-cooling, much more acicular M6 C carbide precipitates in the alloy with no melt treatment. References [1] C.T. Sims, N.S. Stoloff, W.C. Hagel, Superalloys II, John Wiley & Sons Inc., New York, 1987, p. 73.
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