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Effect of thermal exposure on the stability of carbides in Ni–Cr–W based superalloy
Materials Science and Engineering A 528 (2011) 2339–2344
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of thermal exposure on the stability of carbides in Ni–Cr–W based superalloy Guanghai Bai, Jinshan Li, Rui Hu ∗ , Tiebang Zhang, Hongchao Kou, Hengzhi Fu State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an 710072, China
a r t i c l e
i n f o
Article history: Received 19 August 2010 Received in revised form 5 November 2010 Accepted 29 November 2010 Available online 5 December 2010 Keywords: Ni–Cr–W superalloy M6 C M23 C6 Thermal exposure Decomposition Precipitation
a b s t r a c t The microstructure evolutions of Ni–Cr–W based superalloy during thermal exposure have been investigated systematically. M6 C carbides in the alloy decompose into M23 C6 carbides at temperatures from 650 to 1000 ◦ C due to its high content of Cr. The M6 C carbides decompose dramatically from 800 to 900 ◦ C. At temperatures up to 1000 ◦ C, a few M23 C6 carbides form on the surface of M6 C carbides. The decomposition behavior of primary M6 C can be explained by the following reaction: M6 C → M23 C6 + Me (W, Ni, Cr, Mo). At temperatures below 900 ◦ C, coarse lamellar M23 C6 carbides precipitate at the grain boundaries. The carbide lamellae line almost perpendicular to the grain boundaries. While the temperature is above 1000 ◦ C, discrete M23 C6 carbides precipitate at the grain boundaries. Moreover, there are lots of small M23 C6 particles precipitated around M6 C carbides from 650 to 1000 ◦ C. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Ni–Cr–W based superalloy with excellent mechanical strength and long-term creep rupture strength has been developed recently [1,2]. It can be employed at temperatures as high as 1100 ◦ C and could be used widely in aerospace and chemical process industries with particular importance in heat exchanger tubes relating to the high temperature gas cooled reactor (HTGR). The alloy derives its strength from solid solution and primary M6 C carbides, which can prevent excessive grain coarsening and strengthen the grain boundary. Moreover, M23 C6 carbide can precipitate at the grain boundaries during the heat treatment process [3]. Because the alloy is subjected to extremely harsh environments, the effects of thermal exposure on microstructural stability need to be investigated. M6 C carbide is one of the most common carbides observed in superalloys, which can block the movement of dislocations effectively and result in the enhancement of the strength. It can form during the solidification process [4] or precipitate during thermal exposure [5–8]. Previous studies have shown that M6 C is unstable, and it can decompose into M23 C6 in certain alloys [7,9–11]. In Hastelloy X [10], M6 C carbide can decompose into M23 C6 carbide during the low cycle fatigue tests, and such phenomenon also exists in Inconel 617 [11] and Co–Ni–Fe based hard facing alloy [9]. How-
∗ Corresponding author. Tel.: +86 29 88491764; fax: +86 29 88460294. E-mail address: [email protected] (R. Hu). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.11.088
ever, systemic research about the effect of exposure temperature and time on M6 C decomposition behavior is limited. There are two different ways of M23 C6 carbide formation in superalloys. The first one is the decomposition of the carbides such as M6 C [9–11] and MC [12]. The other is that M23 C6 carbide precipitates directly from the matrix [13,14]. Different morphologies of M23 C6 carbide have been observed in superalloys, such as continuous M23 C6 carbide films [15], discrete M23 C6 particles [16,17] and nano-scaled M23 C6 particles [18]. Research shows that the morphologies of M23 C6 carbides are not only determined by the chemical compositions, but also by the heat treatment conditions [13]. In K465 alloy [13], the flower-like dendrite M23 C6 can precipitate during the heat treatment and it changes to regular polyhedron during the thermal exposure. Thus, it is necessary to understand the precipitation behavior of M23 C6 in the given superalloys. In an attempt to get a better understanding of carbides in Ni–Cr–W based superalloy, thermal exposure experiments were performed in the typical application temperatures which were from 650 to 1000 ◦ C. Microstructural evolutions, including M6 C decomposition and M23 C6 precipitation behaviors, were systematically investigated. 2. Experimental method The chemical composition (wt.%) of the material used in the present investigation was Cr, 19.82; W, 18.48; Mo, 1.24; Al, 0.46; C, 0.11; B, 0.0028; La, 0.026; P, S <0.004, Bal. Ni. The process for fabrication of the wrought alloy was as follows. Firstly, the cast material
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was prepared by vacuum induction melting (VIM) and vacuum arc re-melting (VAR). And then the ingot was homogenized at 1200 ◦ C for 24 h and furnace cooled. Finally, it was hot forged and rolled at about 1050 ◦ C to 1150 ◦ C into a 9 mm thick sheet. The specimens were cut from the sheet and solution-heated at 1260 ◦ C for 0.5 h, water quenched (WQ). Subsequently, the specimens were exposed at 650, 750, 800, 900 and 1000 ◦ C for 10–100 h, respectively. The specimens for the analysis of microstructure were prepared by standard metallographic techniques and the polished specimens were etched with aqua regia (HCl:HNO3 = 3:1) for 2 min to reveal the microstructure. The microstructures were examined by optical microscope (OM) and scanning electron microscope (SEM, JSM-6700) with energy dispersive spectrometer (EDS) analysis. Furthermore, transmission electron microscopy (TEM, Tecnai G2 F30) was used for phase identification. TEM samples were prepared by electrochemical polishing at −30 ◦ C in 8% perchloric acid carbinol solution. 3. Results and discussion 3.1. Microstructure of solution-treated Ni–Cr–W based superalloy The microstructures of Ni–Cr–W based superalloy solutiontreated at 1260 ◦ C for 0.5 h are shown in Fig. 1. It can be seen from Fig. 1a that the microstructure consists of equiaxed grains with an average grain size of 70 m. Lots of secondary phase particles (black
grains) randomly disperse inside the grains or at the grain boundaries. TEM analysis (Fig. 1b) shows that these particles are M6 C carbide with the size of 5 m in length and its calculated lattice parameter ˛M6 C is 1.118 nm. 3.2. Temperature dependence of M6 C decomposition The microstructure observation and EDS analysis of samples heat-treated at various conditions are shown in Fig. 2. Plenty of fine particles form densely on the surface of M6 C carbide at 650 ◦ C and 750 ◦ C for 10 h (Fig. 2a and b). The number of the particles increases obviously (Fig. 2c) at temperatures up to 800 ◦ C. These particles have a regular shape with the size of 0.2–0.5 m in length. The selected area electron diffraction pattern indicates that they are M23 C6 carbide with the lattice constant of ˛M23 C6 (1) = 1.064 nm (Fig. 4). The M23 C6 carbides grow larger with increasing the holding time and part of them merge together into larger ones which the grain size can reach 1 m (Fig. 2d). With increasing temperature to 900 ◦ C, most of M23 C6 particles on the surface of M6 C grow larger and abnormal M23 C6 particle coarsening happens (Fig. 2e). The Cr content of the degeneration phase (Fig. 2h) is relatively high compared with M6 C carbides (Fig. 2j). The EDS analysis further confirms that the particles are M23 C6 carbide, which is rich in Cr element. Back-scattered electron (BSE) images and EDS analyses (Fig. 3) show that most of the M6 C carbides decompose into M23 C6 . Since there is no microstructural change on the M6 C/␥ interface, it can be deduced that most of Cr atoms are from the decomposition of M6 C carbides. A few Cr atoms may probably diffuse from the matrix to the M6 C/␥ interface, and these atoms also contribute to the formation of M23 C6 carbide. At temperatures up to 1000 ◦ C, a few M23 C6 particles form at the surface of M6 C carbides (Fig. 2f and g). BSE analysis (Fig. 2g) shows that M23 C6 is encapsulated in M6 C particle indicating its transformation from M6 C carbide. Therefore, it can be concluded that the M6 C carbide decomposes dramatically from 800 to 900 ◦ C. Moreover, there are a lot of small particles which precipitate around M6 C carbides during the thermal exposure (Fig. 2a–g). These particles are M23 C6 carbides which are identified by the EDS analysis (Fig. 2i). The size of M23 C6 particles also increases with the increase of temperature until 900 ◦ C. For most nickel-based superalloy, M6 C carbide is derived from degradation of the MC carbide: MC + ␥ → M6 C + ␥ [19]. The chemical composition of M6 C is usually rich in W, Mo and poor in Cr [20]. However, Cr-rich M6 C carbide can form in the Cr-rich superalloys during the solidification process [2,9,21]. Table 1 shows the relationship between the stability of M6 C and its chemical composition. It can be deduced that when the concentration of Cr in M6 C is high enough, the degradation of M6 C would happen at certain temperatures [9–11]. In the present alloy, the concentration of Cr in M6 C carbide is as high as 14.1%, and this induces the instability of M6 C. As a result, the decomposition of M6 C happens during the thermal exposure process. Based on the previous analysis, a reaction is proposed for the decomposition of the primary M6 C carbide. During the thermal exposure ranging from 650 to 1000 ◦ C, the M6 C carbide is metastable and partly degenerates: M6 C → M23 C6 + Me(W, Ni, Cr, Mo)
Fig. 1. Microstructures of Ni–Cr–W based superalloy solution-treated at 1260 ◦ C for 0.5 h: (a) optical microstructure; (b) TEM morphology of M6 C carbide and its corresponding diffraction pattern.
During the decomposition process, the atoms such as W, Ni, Cr, Mo and C are released from M6 C carbide. These atoms contribute to the nucleation, growth and coarsening of M23 C6 carbides. Research indicates that carbon atoms segregate to dislocations, stacking faults and grain boundaries [22]. Therefore, the nucleation and subsequent growth of M23 C6 carbides are promoted by carbon segregation on the M6 C/␥ interface.
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Fig. 2. SEM micrographs showing the microstructure of M6 C decomposition behavior: (a) 650 ◦ C × 10 h; (b) 750 ◦ C × 10 h; (c) 800 ◦ C × 10 h; (d) 800 ◦ C × 100 h; (e) 900 ◦ C × 100 h; (f) 1000 ◦ C × 100 h; (g) BSE images of (f); (h) and (i) EDS analyses of M23 C6 carbides; (j) EDS analysis of M6 C carbide.
3.3. Grain-boundaries M23 C6 precipitation behavior The SEM analyses on the evolution of the grain boundary morphologies during the thermal exposure are shown in Fig. 5. It can be found that two different morphologies of the secondary phase
precipitate at the grain boundaries. TEM analysis shows that the precipitates are M23 C6 carbides and the phase connected with M23 C6 lamellae is ␥ phase as indicated in Fig. 6. At temperatures from 650 to 900 ◦ C (Fig. 5a–d), a significant amount of coarse lamellar M23 C6 carbides precipitate at the grain boundaries and the space
Table 1 The composition of M6 C carbide in superalloys (wt.%) (P – precipitation; D – MC decomposition; S – solidification). Alloy
Source
Stability
K465 [20] M963 [23] Ni–W–Co [4] Co–Ni–Fe [9]
D P S S
Yes Yes Yes No
M6 C composition Cr
Co
Ni
W
7.1 9.3 1 19.5
7.3 7.6 4.1 13.1
/ 15.6 14.3 10
52.8 57 71.4 43.3
Mo
Nb
Ti
Other elements
15.9 6 6.5 /
0.1 2.5 1.4 /
0.3 1.5 0.6 /
6.9 0.6 1.7 14.1
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Fig. 3. (a) BSE analysis of sample treated at 900 ◦ C × 100 h and (b) its EDS analysis.
between the lamellae is about 0.2 m in average. It is noteworthy that the M23 C6 carbide lamellae are nearly perpendicular to the grain boundaries. EDS analyses (Fig. 5f) show that there is only ␥ phase existed between the M23 C6 lamellae. It can be seen that the density of M23 C6 carbide lamellae decreases with increasing temperature. The grain boundaries are full of these carbides and its width is about 3–5 m as shown in Fig. 5a–d. With increasing temperature to 1000 ◦ C, the lamellar M23 C6 carbides disappear entirely. A large number of discrete M23 C6 carbides with the size of 0.2–0.5 m precipitate at the grain boundaries (Fig. 5e). This kind morphology of M23 C6 was found in many other superalloys such as Ni-based superalloy [20,24] and Fe-based alloys [25,26]. The morphological transition of M23 C6 could be due to the temperature dependence of Cr concentration in the matrix. Sahlaoui et al. studied Cr depleted-zone evolution when M23 C6 precipitated at the grain boundaries in Ni–Cr–Fe alloys, which pointed
Fig. 4. TEM analysis of M6 C decomposition processed at 800 ◦ C for 100 h: (a) microstructure observation and (b) corresponding diffraction pattern.
out that Cr concentration in the matrix increased with increasing temperature [27]. It can be deduced that for the present alloy, Cr concentration in the matrix at 1000 ◦ C is higher than that at lower temperatures, and fewer Cr atoms are supplied for the M23 C6 growth. As a result, M23 C6 growth is suppressed and discrete M23 C6 carbides are apt to form at the grain boundaries. Cr-rich M23 C6 has a face-centered cubic (fcc) structure, the same as the matrix. From the diffraction pattern analysis, the calculated lattice constant of grain boundary M23 C6 ˛M23 C6 (2) in the superalloy is 1.079 nm, which is different from ˛M23 C6 (1) . Most of the elements in M23 C6 which distribute on the surface of M6 C come from the decomposition of M6 C. Therefore, the M23 C6 carbide is high in Cr and W (Fig. 2 h). However, the grain boundary M23 C6 is high in Cr and Ni but low in W (Fig. 5f). As a result, the lattice constant change of two types of M23 C6 carbides is caused by the different chemical
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Fig. 5. SEM examinations of grain boundary precipitation in Ni–Cr–W based superalloy: (a) 650 ◦ C × 100 h; (b) 700 ◦ C × 100 h (the sample was etched for 20 s); (c) 800 ◦ C × 100 h; (d) 900 ◦ C × 100 h; (e) 1000 ◦ C × 100 h; (f) EDS of (b).
composition of the carbides. The M23 C6 carbide has a cube-cube orientation relationship with the matrix: (1 0 0)M23 C6 //(1 0 0)matrix and [1 0 0]M23 C6 //[1 0 0]matrix , as shown in Fig. 6. During the thermal exposure, M23 C6 carbides precipitate directly from the matrix. The reaction of M23 C6 carbide precipitation can be described as: 23M + 6C → M23 C6 . M is rich in Cr, also M may include some other elements such as W, Mo and so on [18,28]. All the elements in M23 C6 carbide come from the matrix. It is reported that the diffusion rate of carbon atoms is much faster than that of Cr atoms in superalloys [27]. Therefore, the growth rate of M23 C6 is controlled by diffusion rate of Cr atoms. 4. Conclusions
Fig. 6. TEM analysis of grain boundary M23 C6 processed at 800 ◦ C for 100 h.
During the thermal exposure from 650 to 1000 ◦ C, M6 C carbide in Ni–Cr–W superalloy is instable and partly decomposes into M23 C6 carbides. Cr concentration plays an important role in the stability of M6 C carbide. The M6 C carbides decompose dramatically from 800 to 900 ◦ C. And during this temperature range, the growth rate of M23 C6 on the surface of M6 C carbide is dependent on both exposure temperature and time. A few M23 C6 carbides form on the surface of M6 C at 1000 ◦ C. The primary M6 C carbide decomposes in the form of M6 C → M23 C6 + Me (W, Ni, Cr, Mo). There are two type morphologies of M23 C6 carbides which precipitate at grain boundaries. The first one is coarse lamellar M23 C6 which the lamellae line
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almost perpendicular to the grain boundaries (650–900 ◦ C), and the other is discrete M23 C6 carbides (1000 ◦ C). Moreover, there are lots of small M23 C6 particles which precipitate around M6 C carbides during the exposure process. Acknowledgements The research was supported by China New Century Excellent Talents in University (NCET-07-0690) and the 111 Project (B08040). References [1] Y. Liu, R. Hu, J.S. Li, H.C. Kou, H.W. Li, H. Chang, H.Z. Fu, Mater. Sci. Eng. A 508 (2009) 141–147. [2] G.H. Bai, J.S. Li, R. Hu, X.Y. Xue, J. Ma, S.T. Hu, H.Z. Fu, Rare Met. Mater. Eng., accepted (in Chinese). [3] Z.W. T, J.S. Li, R. Hu, Y. Liu, G.H. Bai, Rare Met. Mater. Eng. 39 (2010) 1157–1161. [4] L. Zheng, C.Q. Gu, Y.R. Zheng, Scripta Mater. 50 (2004) 435–439. [5] J.X. Yang, Q. Zheng, X.F. Sun, H.R. Guan, Z.Q. Hu, Mater. Sci. Eng. A 465 (2007) 100–108. [6] 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. [7] S. Hamar-Thibault, M. Durand-Charre, B. Andries, Metall. Mater. Trans. A 13 (1982) 545–550.
[8] F.M. Yang, X.F. Sun, W. Zhang, Y.P. Kang, H.R. Guan, Z.Q. Hu, Mater. Lett. 49 (2001) 160–164. [9] B. Djerdjare, S. Lebaili, S. Lay, Mater. Sci. Eng. A 475 (2008) 336–342. [10] G.V.P. Reddy, P. Harini, R. Sandhya, K.B.S. Rao, R.K. Paretkar, Mater. Sci. Eng. A 527 (2010) 3848–3851. [11] S. Chomette, J.M. Gentzbittel, B. Viguier, J. Nucl. Mater. 399 (2010) 266–274. [12] B.G. Choi, I.S. Kim, D.H. Kim, C.Y. Jo, Mater. Sci. Eng. A 478 (2008) 329–335. [13] X.Z. Qin, J.T. Guo, C. Yuan, J.S. Hou, H.Q. Ye, Mater. Lett. 62 (2008) 258–261. [14] L.Z. He, Q. Zheng, X.F. Sun, H.R. Guan, Z.Q. Hu, A.K. Tieu, C. Lu, H.T. Zhu, Mater. Sci. Eng. A 397 (2005) 297–304. [15] A. Koul, R. Castillo, Metall. Mater. Trans. A 19 (1988) 2049–2066. [16] T. Garosshen, G. McCarthy, Metall. Mater. Trans. A 16 (1985) 1213–1223. [17] L. He, Q. Zheng, X. Sun, G. Hou, H. Guan, Z. Hu, J. Mater. Sci. 40 (2005) 2959– 2964. [18] D. Wang, J. Zhang, L.H. Lou, Mater. Charact. 60 (2009) 1517–1521. [19] X.Z. Qin, J.T. Guo, C. Yuan, C.L. Chen, J.S. Hou, H.Q. Ye, Mater. Sci. Eng. A 485 (2008) 74–79. [20] J. Yang, Q. Zheng, X. Sun, H. Guan, Z. Hu, Mater. Sci. Eng. A 429 (2006) 341–347. [21] S. Zangeneh, H. Farhangi, Mater. Des. 31 (2010) 3504–3511. [22] W. Jiang, X. Yao, H. Guan, Z. Hu, Metall. Mater. Trans. A 30 (1999) 513–520. [23] F.S. Yin, Q. Zheng, X.F. Sun, H.R. Guan, Z.Q. Hu, J. Mater. Process. Technol. 183 (2007) 440–444. [24] K.L. Zeik, J.D. Rossi, P.R. Howell, J. Mater. Sci. Lett. 8 (1989) 732–734. [25] Y. Kadoya, B. Dyson, M. McLean, Metall. Mater. Trans. A 33 (2002) 2549–2557. [26] T.K. Kim, S.H. Kim, C.B. Lee, Ann. Nucl. Energy 36 (2009) 1103–1107. [27] H. Sahlaoui, H. Sidhom, J. Philibert, Acta Mater. 50 (2002) 1383–1392. [28] W.H. Jiang, H.R. Guan, Z.Q. Hu, Mater. Sci. Eng. A 271 (1999) 101–108.
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of thermal exposure on the stability of carbides in Ni–Cr–W based superalloy Guanghai Bai, Jinshan Li, Rui Hu ∗ , Tiebang Zhang, Hongchao Kou, Hengzhi Fu State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an 710072, China
a r t i c l e
i n f o
Article history: Received 19 August 2010 Received in revised form 5 November 2010 Accepted 29 November 2010 Available online 5 December 2010 Keywords: Ni–Cr–W superalloy M6 C M23 C6 Thermal exposure Decomposition Precipitation
a b s t r a c t The microstructure evolutions of Ni–Cr–W based superalloy during thermal exposure have been investigated systematically. M6 C carbides in the alloy decompose into M23 C6 carbides at temperatures from 650 to 1000 ◦ C due to its high content of Cr. The M6 C carbides decompose dramatically from 800 to 900 ◦ C. At temperatures up to 1000 ◦ C, a few M23 C6 carbides form on the surface of M6 C carbides. The decomposition behavior of primary M6 C can be explained by the following reaction: M6 C → M23 C6 + Me (W, Ni, Cr, Mo). At temperatures below 900 ◦ C, coarse lamellar M23 C6 carbides precipitate at the grain boundaries. The carbide lamellae line almost perpendicular to the grain boundaries. While the temperature is above 1000 ◦ C, discrete M23 C6 carbides precipitate at the grain boundaries. Moreover, there are lots of small M23 C6 particles precipitated around M6 C carbides from 650 to 1000 ◦ C. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Ni–Cr–W based superalloy with excellent mechanical strength and long-term creep rupture strength has been developed recently [1,2]. It can be employed at temperatures as high as 1100 ◦ C and could be used widely in aerospace and chemical process industries with particular importance in heat exchanger tubes relating to the high temperature gas cooled reactor (HTGR). The alloy derives its strength from solid solution and primary M6 C carbides, which can prevent excessive grain coarsening and strengthen the grain boundary. Moreover, M23 C6 carbide can precipitate at the grain boundaries during the heat treatment process [3]. Because the alloy is subjected to extremely harsh environments, the effects of thermal exposure on microstructural stability need to be investigated. M6 C carbide is one of the most common carbides observed in superalloys, which can block the movement of dislocations effectively and result in the enhancement of the strength. It can form during the solidification process [4] or precipitate during thermal exposure [5–8]. Previous studies have shown that M6 C is unstable, and it can decompose into M23 C6 in certain alloys [7,9–11]. In Hastelloy X [10], M6 C carbide can decompose into M23 C6 carbide during the low cycle fatigue tests, and such phenomenon also exists in Inconel 617 [11] and Co–Ni–Fe based hard facing alloy [9]. How-
∗ Corresponding author. Tel.: +86 29 88491764; fax: +86 29 88460294. E-mail address: [email protected] (R. Hu). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.11.088
ever, systemic research about the effect of exposure temperature and time on M6 C decomposition behavior is limited. There are two different ways of M23 C6 carbide formation in superalloys. The first one is the decomposition of the carbides such as M6 C [9–11] and MC [12]. The other is that M23 C6 carbide precipitates directly from the matrix [13,14]. Different morphologies of M23 C6 carbide have been observed in superalloys, such as continuous M23 C6 carbide films [15], discrete M23 C6 particles [16,17] and nano-scaled M23 C6 particles [18]. Research shows that the morphologies of M23 C6 carbides are not only determined by the chemical compositions, but also by the heat treatment conditions [13]. In K465 alloy [13], the flower-like dendrite M23 C6 can precipitate during the heat treatment and it changes to regular polyhedron during the thermal exposure. Thus, it is necessary to understand the precipitation behavior of M23 C6 in the given superalloys. In an attempt to get a better understanding of carbides in Ni–Cr–W based superalloy, thermal exposure experiments were performed in the typical application temperatures which were from 650 to 1000 ◦ C. Microstructural evolutions, including M6 C decomposition and M23 C6 precipitation behaviors, were systematically investigated. 2. Experimental method The chemical composition (wt.%) of the material used in the present investigation was Cr, 19.82; W, 18.48; Mo, 1.24; Al, 0.46; C, 0.11; B, 0.0028; La, 0.026; P, S <0.004, Bal. Ni. The process for fabrication of the wrought alloy was as follows. Firstly, the cast material
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was prepared by vacuum induction melting (VIM) and vacuum arc re-melting (VAR). And then the ingot was homogenized at 1200 ◦ C for 24 h and furnace cooled. Finally, it was hot forged and rolled at about 1050 ◦ C to 1150 ◦ C into a 9 mm thick sheet. The specimens were cut from the sheet and solution-heated at 1260 ◦ C for 0.5 h, water quenched (WQ). Subsequently, the specimens were exposed at 650, 750, 800, 900 and 1000 ◦ C for 10–100 h, respectively. The specimens for the analysis of microstructure were prepared by standard metallographic techniques and the polished specimens were etched with aqua regia (HCl:HNO3 = 3:1) for 2 min to reveal the microstructure. The microstructures were examined by optical microscope (OM) and scanning electron microscope (SEM, JSM-6700) with energy dispersive spectrometer (EDS) analysis. Furthermore, transmission electron microscopy (TEM, Tecnai G2 F30) was used for phase identification. TEM samples were prepared by electrochemical polishing at −30 ◦ C in 8% perchloric acid carbinol solution. 3. Results and discussion 3.1. Microstructure of solution-treated Ni–Cr–W based superalloy The microstructures of Ni–Cr–W based superalloy solutiontreated at 1260 ◦ C for 0.5 h are shown in Fig. 1. It can be seen from Fig. 1a that the microstructure consists of equiaxed grains with an average grain size of 70 m. Lots of secondary phase particles (black
grains) randomly disperse inside the grains or at the grain boundaries. TEM analysis (Fig. 1b) shows that these particles are M6 C carbide with the size of 5 m in length and its calculated lattice parameter ˛M6 C is 1.118 nm. 3.2. Temperature dependence of M6 C decomposition The microstructure observation and EDS analysis of samples heat-treated at various conditions are shown in Fig. 2. Plenty of fine particles form densely on the surface of M6 C carbide at 650 ◦ C and 750 ◦ C for 10 h (Fig. 2a and b). The number of the particles increases obviously (Fig. 2c) at temperatures up to 800 ◦ C. These particles have a regular shape with the size of 0.2–0.5 m in length. The selected area electron diffraction pattern indicates that they are M23 C6 carbide with the lattice constant of ˛M23 C6 (1) = 1.064 nm (Fig. 4). The M23 C6 carbides grow larger with increasing the holding time and part of them merge together into larger ones which the grain size can reach 1 m (Fig. 2d). With increasing temperature to 900 ◦ C, most of M23 C6 particles on the surface of M6 C grow larger and abnormal M23 C6 particle coarsening happens (Fig. 2e). The Cr content of the degeneration phase (Fig. 2h) is relatively high compared with M6 C carbides (Fig. 2j). The EDS analysis further confirms that the particles are M23 C6 carbide, which is rich in Cr element. Back-scattered electron (BSE) images and EDS analyses (Fig. 3) show that most of the M6 C carbides decompose into M23 C6 . Since there is no microstructural change on the M6 C/␥ interface, it can be deduced that most of Cr atoms are from the decomposition of M6 C carbides. A few Cr atoms may probably diffuse from the matrix to the M6 C/␥ interface, and these atoms also contribute to the formation of M23 C6 carbide. At temperatures up to 1000 ◦ C, a few M23 C6 particles form at the surface of M6 C carbides (Fig. 2f and g). BSE analysis (Fig. 2g) shows that M23 C6 is encapsulated in M6 C particle indicating its transformation from M6 C carbide. Therefore, it can be concluded that the M6 C carbide decomposes dramatically from 800 to 900 ◦ C. Moreover, there are a lot of small particles which precipitate around M6 C carbides during the thermal exposure (Fig. 2a–g). These particles are M23 C6 carbides which are identified by the EDS analysis (Fig. 2i). The size of M23 C6 particles also increases with the increase of temperature until 900 ◦ C. For most nickel-based superalloy, M6 C carbide is derived from degradation of the MC carbide: MC + ␥ → M6 C + ␥ [19]. The chemical composition of M6 C is usually rich in W, Mo and poor in Cr [20]. However, Cr-rich M6 C carbide can form in the Cr-rich superalloys during the solidification process [2,9,21]. Table 1 shows the relationship between the stability of M6 C and its chemical composition. It can be deduced that when the concentration of Cr in M6 C is high enough, the degradation of M6 C would happen at certain temperatures [9–11]. In the present alloy, the concentration of Cr in M6 C carbide is as high as 14.1%, and this induces the instability of M6 C. As a result, the decomposition of M6 C happens during the thermal exposure process. Based on the previous analysis, a reaction is proposed for the decomposition of the primary M6 C carbide. During the thermal exposure ranging from 650 to 1000 ◦ C, the M6 C carbide is metastable and partly degenerates: M6 C → M23 C6 + Me(W, Ni, Cr, Mo)
Fig. 1. Microstructures of Ni–Cr–W based superalloy solution-treated at 1260 ◦ C for 0.5 h: (a) optical microstructure; (b) TEM morphology of M6 C carbide and its corresponding diffraction pattern.
During the decomposition process, the atoms such as W, Ni, Cr, Mo and C are released from M6 C carbide. These atoms contribute to the nucleation, growth and coarsening of M23 C6 carbides. Research indicates that carbon atoms segregate to dislocations, stacking faults and grain boundaries [22]. Therefore, the nucleation and subsequent growth of M23 C6 carbides are promoted by carbon segregation on the M6 C/␥ interface.
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Fig. 2. SEM micrographs showing the microstructure of M6 C decomposition behavior: (a) 650 ◦ C × 10 h; (b) 750 ◦ C × 10 h; (c) 800 ◦ C × 10 h; (d) 800 ◦ C × 100 h; (e) 900 ◦ C × 100 h; (f) 1000 ◦ C × 100 h; (g) BSE images of (f); (h) and (i) EDS analyses of M23 C6 carbides; (j) EDS analysis of M6 C carbide.
3.3. Grain-boundaries M23 C6 precipitation behavior The SEM analyses on the evolution of the grain boundary morphologies during the thermal exposure are shown in Fig. 5. It can be found that two different morphologies of the secondary phase
precipitate at the grain boundaries. TEM analysis shows that the precipitates are M23 C6 carbides and the phase connected with M23 C6 lamellae is ␥ phase as indicated in Fig. 6. At temperatures from 650 to 900 ◦ C (Fig. 5a–d), a significant amount of coarse lamellar M23 C6 carbides precipitate at the grain boundaries and the space
Table 1 The composition of M6 C carbide in superalloys (wt.%) (P – precipitation; D – MC decomposition; S – solidification). Alloy
Source
Stability
K465 [20] M963 [23] Ni–W–Co [4] Co–Ni–Fe [9]
D P S S
Yes Yes Yes No
M6 C composition Cr
Co
Ni
W
7.1 9.3 1 19.5
7.3 7.6 4.1 13.1
/ 15.6 14.3 10
52.8 57 71.4 43.3
Mo
Nb
Ti
Other elements
15.9 6 6.5 /
0.1 2.5 1.4 /
0.3 1.5 0.6 /
6.9 0.6 1.7 14.1
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Fig. 3. (a) BSE analysis of sample treated at 900 ◦ C × 100 h and (b) its EDS analysis.
between the lamellae is about 0.2 m in average. It is noteworthy that the M23 C6 carbide lamellae are nearly perpendicular to the grain boundaries. EDS analyses (Fig. 5f) show that there is only ␥ phase existed between the M23 C6 lamellae. It can be seen that the density of M23 C6 carbide lamellae decreases with increasing temperature. The grain boundaries are full of these carbides and its width is about 3–5 m as shown in Fig. 5a–d. With increasing temperature to 1000 ◦ C, the lamellar M23 C6 carbides disappear entirely. A large number of discrete M23 C6 carbides with the size of 0.2–0.5 m precipitate at the grain boundaries (Fig. 5e). This kind morphology of M23 C6 was found in many other superalloys such as Ni-based superalloy [20,24] and Fe-based alloys [25,26]. The morphological transition of M23 C6 could be due to the temperature dependence of Cr concentration in the matrix. Sahlaoui et al. studied Cr depleted-zone evolution when M23 C6 precipitated at the grain boundaries in Ni–Cr–Fe alloys, which pointed
Fig. 4. TEM analysis of M6 C decomposition processed at 800 ◦ C for 100 h: (a) microstructure observation and (b) corresponding diffraction pattern.
out that Cr concentration in the matrix increased with increasing temperature [27]. It can be deduced that for the present alloy, Cr concentration in the matrix at 1000 ◦ C is higher than that at lower temperatures, and fewer Cr atoms are supplied for the M23 C6 growth. As a result, M23 C6 growth is suppressed and discrete M23 C6 carbides are apt to form at the grain boundaries. Cr-rich M23 C6 has a face-centered cubic (fcc) structure, the same as the matrix. From the diffraction pattern analysis, the calculated lattice constant of grain boundary M23 C6 ˛M23 C6 (2) in the superalloy is 1.079 nm, which is different from ˛M23 C6 (1) . Most of the elements in M23 C6 which distribute on the surface of M6 C come from the decomposition of M6 C. Therefore, the M23 C6 carbide is high in Cr and W (Fig. 2 h). However, the grain boundary M23 C6 is high in Cr and Ni but low in W (Fig. 5f). As a result, the lattice constant change of two types of M23 C6 carbides is caused by the different chemical
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Fig. 5. SEM examinations of grain boundary precipitation in Ni–Cr–W based superalloy: (a) 650 ◦ C × 100 h; (b) 700 ◦ C × 100 h (the sample was etched for 20 s); (c) 800 ◦ C × 100 h; (d) 900 ◦ C × 100 h; (e) 1000 ◦ C × 100 h; (f) EDS of (b).
composition of the carbides. The M23 C6 carbide has a cube-cube orientation relationship with the matrix: (1 0 0)M23 C6 //(1 0 0)matrix and [1 0 0]M23 C6 //[1 0 0]matrix , as shown in Fig. 6. During the thermal exposure, M23 C6 carbides precipitate directly from the matrix. The reaction of M23 C6 carbide precipitation can be described as: 23M + 6C → M23 C6 . M is rich in Cr, also M may include some other elements such as W, Mo and so on [18,28]. All the elements in M23 C6 carbide come from the matrix. It is reported that the diffusion rate of carbon atoms is much faster than that of Cr atoms in superalloys [27]. Therefore, the growth rate of M23 C6 is controlled by diffusion rate of Cr atoms. 4. Conclusions
Fig. 6. TEM analysis of grain boundary M23 C6 processed at 800 ◦ C for 100 h.
During the thermal exposure from 650 to 1000 ◦ C, M6 C carbide in Ni–Cr–W superalloy is instable and partly decomposes into M23 C6 carbides. Cr concentration plays an important role in the stability of M6 C carbide. The M6 C carbides decompose dramatically from 800 to 900 ◦ C. And during this temperature range, the growth rate of M23 C6 on the surface of M6 C carbide is dependent on both exposure temperature and time. A few M23 C6 carbides form on the surface of M6 C at 1000 ◦ C. The primary M6 C carbide decomposes in the form of M6 C → M23 C6 + Me (W, Ni, Cr, Mo). There are two type morphologies of M23 C6 carbides which precipitate at grain boundaries. The first one is coarse lamellar M23 C6 which the lamellae line
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almost perpendicular to the grain boundaries (650–900 ◦ C), and the other is discrete M23 C6 carbides (1000 ◦ C). Moreover, there are lots of small M23 C6 particles which precipitate around M6 C carbides during the exposure process. Acknowledgements The research was supported by China New Century Excellent Talents in University (NCET-07-0690) and the 111 Project (B08040). References [1] Y. Liu, R. Hu, J.S. Li, H.C. Kou, H.W. Li, H. Chang, H.Z. Fu, Mater. Sci. Eng. A 508 (2009) 141–147. [2] G.H. Bai, J.S. Li, R. Hu, X.Y. Xue, J. Ma, S.T. Hu, H.Z. Fu, Rare Met. Mater. Eng., accepted (in Chinese). [3] Z.W. T, J.S. Li, R. Hu, Y. Liu, G.H. Bai, Rare Met. Mater. Eng. 39 (2010) 1157–1161. [4] L. Zheng, C.Q. Gu, Y.R. Zheng, Scripta Mater. 50 (2004) 435–439. [5] J.X. Yang, Q. Zheng, X.F. Sun, H.R. Guan, Z.Q. Hu, Mater. Sci. Eng. A 465 (2007) 100–108. [6] 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. [7] S. Hamar-Thibault, M. Durand-Charre, B. Andries, Metall. Mater. Trans. A 13 (1982) 545–550.
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