- Email: [email protected]
Effects of heat treatment on microstructures and mechanical properties of a directionally solidified cobalt-base superalloy
Materials Science and Engineering A271 (1999) 101 – 108 www.elsevier.com/locate/msea
Effects of heat treatment on microstructures and mechanical properties of a directionally solidified cobalt-base superalloy W.H. Jiang a,b,*, H.R. Guan a, Z.Q. Hu a a
Department of Superalloys, Institute of Metal Research, Academia Sinica, No. 72 Wenhua Road, Shenyang 110015, P.R. China b Department of Metal Materials Engineering, Shenyang Polytechnic Uni6ersity, Shenyang 110023, P.R. China Received 26 January 1999; received in revised form 6 April 1999
Abstract DZ40M alloy is a newly developed directionally solidified cobalt-base superalloy. The present work investigated microstructures, room-temperature tensile and stress–rupture properties at 980°C/83 MPa of the alloy in as-cast, solutionized as well as aged states. The microstructure of the DZ40M alloy can be modified by heat treatment. Solution treatment at 1280°C for 4 h dissolved the primary carbides essentially and the alloy became a single-phase supersaturated solid solution. Incorporation of aging treatment at 950°C for 12 h produced a profusion of secondary M23C6 precipitation throughout the matrix. The room-temperature mechanical properties of the alloy are mainly dependent on the microstructures of the matrix. During the stress – rupture tests, the microstructural evolution occurred in the alloy, the primary carbides dissolved sluggishly and the secondary M23C6 precipitated heavily. The precipitation hardening is the most important strengthening mechanism at high temperature for the alloy in all three states. The stress–rupture properties were dominated by both the matrix and the boundaries of grains and interdendrites. The aged alloy has a superior stress– rupture property, which is attributed to its good microstructural combination of the matrix and the boundaries of grains and interdendrites. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Cobalt-base superalloy; Directional solidification; Heat treatment; Mechanical property; Microstructure
1. Introduction Cobalt-base superalloys are widely used in industrial and aircraft turbines for vanes and combuster sections because of their intrinsic properties such as good stress–rupture parameters and excellent hot corrosion and oxidation resistance. The demand for longer lives in turbines is the driving force to improve their mechanical properties. Some efforts have been made and substantial progress has been achieved. The work has focused on composition modification, alloy development, and solidification mode [1 – 4]. However, despite the fact that precipitation hardening is their important strengthening mechanism, the potential for strengthening alloys by heat treatment has not been well developed. In industrial practice, cobalt-base superalloys are usually used in as-cast state [5]. * Corresponding author. E-mail address: [email protected] (W.H. Jiang)
As-cast cobalt-base superalloys consist of a continuous fcc matrix and a variety of carbides, mainly coarse primary M23C6, M7C3 and MC. The carbides contribute significantly to strengthening. The carbides form as alloys cooling down in shell mold. They precipitate at grain boundaries and in interdendritic regions. During service at high temperature, secondary carbides, usually M23C6, precipitate. The fine secondary carbide pins up dislocations that hardens the matrix. Obviously, the morphology, distribution and size of the secondary carbides affect precipitation hardening effect. However, little attention has been paid to control their precipitation, for heat treatment deteriorates their room-temperature ductility severely [5,6]. Lately, a modified directionally solidified X-40 (DZ40M) alloy was developed in the Institute of Metal Research, Academia Sinica [7]. The directional solidification not only eliminates transverse grain boundaries, but also produces a structure of columnar grain matrix
0921-5093/99/$ - see front matter © 1999 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 9 ) 0 0 1 8 3 - 5
W.H. Jiang et al. / Materials Science and Engineering A271 (1999) 101–108
102
Table 1 Tensile properties of DZ40M alloy at room temperature State
Ultimate tensile strength, MPa
0.2% yield strength, MPa
Elongation, %
As-cast Solutionized Aged
672 729 1039
422 455 649
19.0 36.0 9.0
with a 001 direction parallel to the specimen axis, i.e. 001 preferential orientation. The alloy has an excellent ductility. For example, its room-temperature tensile elongation reaches up to about 20%, while that of conventional equiaxed as-cast X-40 is only 8% [4]. Its excellent ductility provides a prerequisite to exploit heat treatment to increase alloy strength further. For development of heat treatment, it is essential to understand a relationship between microstructure and mechanical property. The present work investigated microstructures and mechanical properties at room and high temperatures of DZ40M alloy in as-cast, solutionized as well as aged states.
2. Experimental procedure The alloy was prepared in a conventional vacuum furnace with a mold withdrawal device. The nominal chemical composition of the alloy (in wt.%) is 25Cr, 11Ni, 7.5W, 0.45C, 0.05B, 0.8Al, 0.2Mo, 0.25Ta, 0.15Ti, 0.15Zr and the balance Co. Cylindrical rods of the alloy, 16 mm in diameter and 140 mm long, were produced at a withdrawal speed of 7 mm/min and with a thermal gradient of about 50 K/cm at the solid/liquid interface. Three typical alloy states were selected for this investigation: as-cast, solutionized as well as aged. The solution treatment was performed at 1280°C for 4 h, and followed by air cooling. The aging treatment was carried out at 950°C for 12 h, which was applied after solution treatment. The alloy microstructures were examined with scanning electron microscope (SEM) with energy dispersive X-ray analysis (EDAX). The samples for SEM observation were etched electrolytically in a solution consisting of 42 pct H3PO4, 34 pct H2SO4 and 24 pct H2O. High magnification observation was carried out using transmission electron microscope (TEM). Electropolishing of the thin foils was done with an electrolyte consisting of 5 pct perchloric acid in methanol at about − 20°C. Identification of phases was carried out by X-ray diffraction using filtered chromium radiation. An investigation was made on carbides extracted from the anodically dissolved matrix of the as-cast alloy. Room-temperature tensile and high temperature stress–rupture properties were measured along the
specimen axis. The stress–rupture tests were conducted at 980°C in air at a stress level of 83 MPa. After the stress–rupture tests, fracture surfaces and microstructures were observed with SEM.
3. Results
3.1. Mechanical properties The tensile test data obtained for the alloy in the three selected states at room temperature are given in Table 1. It can be seen that the tensile properties are closely related to the alloy states. The as-cast alloy has the lowest room-temperature strength, but a high ductility, which should be attributed to its columnar grain structure with 001 preferential orientation. The 1280°C solution treatment caused a little increase in strength, and almost doubled tensile elongation, compared to the as-cast alloy. Incorporation of 950°C/12 h aging treatment reduced tensile ductility markedly, but increased strength by 50%. It is noted that the aged DZ40M alloy has the highest room-temperature strength, and an adequate ductility which is comparative to that of the equiaxed as-cast X-40 [4]. The room-temperature mechanical property of the aged alloy is superior to that of as-cast MAR-M509, a high strength cobalt-base superalloy, whose yield strength and elongation are 519 MPa and 4.5%, respectively [6]. The stress–rupture tests at 980°C and 83 MPa were conducted for the alloy in the three states. The results are summarized in Table 2. It demonstrates that the alloy states have a more pronounced effect on high temperature stress–rupture property. The as-cast alloy possesses the shortest stress–rupture life and the highest stress–rupture ductility. Relative to the as-cast alloy, stress–rupture life was prolonged three and five times by the solution and aging treatments, respectively. However, at the same time, a severe loss in rupture Table 2 Stress–rupture properties of DZ40M alloy at 980°C/83 MPa State
Stress–rupture life, h
Elongation, %
As-cast Solutionized Aged
59.7 188.8 278.1
56.0 4.4 28.0
W.H. Jiang et al. / Materials Science and Engineering A271 (1999) 101–108
103
Fig. 1. Optical micrograph of the as-cast DZ40M alloy showing a columnar grain structure.
Fig. 3. Back-scattered electron micrograph of primary carbides in the as-cast DZ40M alloy.
ductility accompanied. The elongation of the solutionized alloy is only 4.4%, while that of the as-cast alloy is as high as 56%. The aging treatment increased the rupture ductility to 28%. Evidently, the aged alloy possesses a good combination of the stress – rupture life and ductility.
of MC is C45.10, Zr24.34, Ta11.00, Ti10.77, Co3.06, W2.84, Cr2.06, and Ni0.82. The scanning electron image (Fig. 3) shows the morphology and distribution of the carbide phases. In this backscattered electron image, the chromium-rich M7C3 carbide appears dark and the MC carbide is rather light. It is observed that M7C3 carbide is in the form of rods or irregular aggregates, while MC is present as a discrete, blocky dispersion with a well-distributed Chinese script morphology. Such a morphology of MC is also observed in MARM509 alloy [1,5]. Evidently, the formation of MC is due to the addition of reactive elements, Ta, Ti and Zr. Both the M7C3 and MC carbides are located at grain boundaries or in interdendritic regions, forming a continuous network around the columnar grained matrix. Fig. 4 is the microstructure of the solutionized alloy. It can be seen that the solution treatment at 1280°C for 4 h dissolved the primary carbides essentially and the alloy became a single-phase supersaturated solid solution. Fig. 5 reveals the microstructure of the aged alloy. The aging treatment at 950°C for 12 h gave rise to a profuse secondary precipitation throughout the alloy
3.2. Microstructural characteristic Fig. 1 shows the microstructure of the as-cast DZ40M alloy. Its matrix is a well-developed columnar grain austenite with 001 direction parallel to the growth axis. X-ray diffraction analysis (Fig. 2) and energy dispersive X-ray analysis (EDAX) indicate that the alloy contains two types of coarse primary carbides, i.e. chromium-rich M7C3 and MC which contains tantalum, titanium, zirconium, and tungsten. A typical composition of M7C3 is (in at pct) C29.94, Cr56.14, Co8.35, W3.88, Ni0.77, Mo0.67, Ti0.22 and Zr0.04, while that
Fig. 2. X-ray diffraction pattern of carbides extracted from the as-cast DZ40M alloy.
Fig. 4. Back-scattered electron micrograph showing the microstructure of the solutionized alloy.
104
W.H. Jiang et al. / Materials Science and Engineering A271 (1999) 101–108
{100}precipitate // {100}matrix 001precipitate // 001matrix
Fig. 5. SEM micrograph showing the microstructure of the aged alloy.
matrix. It can be seen that the secondary precipitates is unevenly distributed. They appeared preferentially on grain boundaries and subgrain boundaries which were decorated by them. It is noted that there are a few of the large rod-like precipitates. They seem to be sigma phase. EDAX was used to characterize their composition. From the obtained profile shown in Fig. 6, it is learnt the rod-like precipitates are chromium-rich carbide rather than sigma phase. In Fig. 6, relatively high cobalt peak resulted from the alloy matrix due to narrowness of the rod-like precipitates and a limited resolution of EDAX. TEM observation (Fig. 7) shows the blocky and rod-like precipitates. Electron diffraction analysis demonstrates that these precipitates possess the same crystal structure and their characteristic reflections are present at every one-third position of the fcc matrix reflection. This indicates that the precipitates have a fcc structure with a lattice constant that is nearly three times that of the matrix, and they also have a cube–cube orientation relationship with the matrix:
These are characteristic features of the chromium-rich M23C6 carbide [8]. Fig. 8 depicts the microstructures of the alloy in three states after the stress–rupture tests. Compared with the original ones, substantially microstructural changes have occurred in the alloy in all three states. In the as-cast specimen ruptured, a profusion of fine secondary precipitates were produced in the matrix and the primary carbides thinned (Fig. 8a). However, the coarse primary carbides were still retained, indicating that their dissolution was sluggish. It is worthy of note that the secondary precipitates are unevenly distributed. Close to the primary carbides, there is dense distribution. After the stress–rupture test, the solutionized alloy also contained a great number of fine secondary precipitates, as shown in Fig. 8b. They are evenly distributed and aligned in rows to form a grid. Evidently, during the stress–rupture test, the solutionized alloy subjected to a stress treatment. The regular arrangement of the secondary precipitates resulted from a stress effect. Furthermore, it is observed that a copious intergranular precipitation occurred, forming a continuous chain. But, the intergranular precipitates are very fine. Fig. 8c shows the microstructure of the aged alloy after the stress–rupture test. It has the most dispersed secondary precipitates. Obviously, the microstructural homogeneity was improved relative to the aged one. Coarse and fine precipitates coexist in the matrix, indicating that their coarsening and reprecipitation happened at the same time during the test. Particularly, it is noted that the rod-like M23C6 disappeared. Furthermore, it can be seen that intergranular precipitates are coarse and comparative to those in the as-cast alloy ruptured. Usually, in cast cobalt-base superalloys, secondary M23C6 carbide is more thermodynamically stable, and during service or exposure at high temperature, its precipitation is very common. In fact, in the DZ40M alloy, a heavy precipitation of secondary M23C6 carbide in the matrix occurred during high temperature low cycle fatigue, creep and thermal aging [9–11]. Therefore, it is believed that the secondary precipitates must be M23C6 in the present alloy in all three states after the stress–rupture tests.
3.3. Fracture surfaces after the stress–rupture tests
Fig. 6. EDAX spectrum of the rod-like precipitates in the DZ40M alloy aged at 950°C/12 h.
After the stress–rupture tests, the fracture surfaces of the alloy in all three states were observed with the SEM. Fig. 9a shows the fracture appearance of the as-cast alloy, where there are a great number of fine voids. This fracture is a ductile type. Microvoid coalescence is its fracture mechanism. Fig. 9b shows the fracture appearance of the solutionized alloy. Cracking
W.H. Jiang et al. / Materials Science and Engineering A271 (1999) 101–108
105
Fig. 7. (a) TEM bright field micrograph of the fine precipitates formed in the aged alloy, (b) dark field micrograph of the same region, using {111} precipitate reflection and (c) corresponding diffraction pattern showing a cube – cube matrix: carbide orientation relationship.
of longitudinal grain boundaries and interdendritic regions was observed. The fracture is transgranular and brittle. The matrix reveals very limited amount of plastic deformation. Fig. 9c reveals the fracture appearance of the aged alloy, where there are uniformly distributed voids. The fracture is a ductile type. Although its fracture mechanism also is coalescence of microvoids, the number of microvoids is less than that of the as-cast alloy.
4. Discussion Like other cast cobalt-base superalloys, the as-cast DZ40M alloy consists of a cobalt-base austenitic matrix and coarse primary carbides located at grain boundaries and interdendritic regions. However, the directional solidification developed a structure of columnar grain matrix with 001 preferential orientation. From Table 1, it is learnt that the as-cast alloy has a yield strength not higher but a little lower than that of the solutionized alloy. This demonstrates that the room-temperature strength of the as-cast alloy resulted mainly from the matrix and little from the primary carbides. At room temperature, its strengthening mech-
anism is solution hardening. At the same time, the alloy matrix with 001 preferential orientation shows an excellent ductility, for the slip system 110{111} of the fcc matrix is located at the easiest slip direction when load is applied along the specimen axis, i.e. 001 direction. At high temperature, the as-cast alloy was thermodynamically unstable. The primary carbides, both the M7C3 and MC, dissolved sluggishly and the secondary M23C6 carbide precipitated profusely. The M23C6 precipitates are unevenly distributed. There are a dense distribution around the primary carbides and a sparse distribution in the center of grains, as shown in Fig. 8a. The secondary M23C6 precipitation is closely related to the decomposition of the primary carbides. A previous work by the present authors [10] suggested a direct reaction mechanism for the M23C6 precipitation, i.e. 23M+ 6C M23C6
(1)
As a carbon reservoir, the primary carbides provided the reaction Eq. (1) with carbon atoms. Since during solidification carbon atoms were exhausted to form the primary carbides and there was little of carbon in the matrix, carbon content became a predominant factor of the precipitation reaction Eq. (1). It was the uneven
106
W.H. Jiang et al. / Materials Science and Engineering A271 (1999) 101–108
distribution of carbon that caused non-homogeneous secondary carbide precipitation. During the stress – rupture test, the microstructural evolution in the as-cast alloy brought about a precipitation hardening, which was absent at room temperature. The retained coarse primary carbides strengthened grain and interdendritic boundaries, retarding their migration. The fine secondary M23C6 carbide pinned up dislocations that hardened the matrix. Due to their relatively sparse distribution, however, the precipitation strengthening was less effective in the center of grains. The fracture surface observa-
Fig. 9. SEM micrographs showing the fracture surfaces of the alloy in all three states after the stress – rupture tests: (a) as-cast, (b) solutionized, and (c) aged.
Fig. 8. SEM micrographs showing the microstructures of the alloy in three states after the stress–rupture tests: (a) as-cast, (b) solutionized, and (c) aged.
tion shows that a substantial plastic deformation took place, which is probably related to low strength in the center of grains. The non-homogeneous microstructure and consequently uneven strength in the matrix resulted in a low stress–rupture life. At high temperature, an extraordinarily high ductility was developed by the matrix, like the situation at room temperature.
W.H. Jiang et al. / Materials Science and Engineering A271 (1999) 101–108
The solution treatment at 1280°C caused a complete dissolution of the primary carbides, both the M7C3 and MC, and no secondary M23C6 precipitation (Fig. 4). This indicates that at the solution temperature both the primary and secondary carbides are thermodynamically unstable. So, it is feasible to develop a single-phase supersaturated solid solution for the DZ40M alloy by solution treatment. This is especially significant for development of heat treatment for the alloy. Obviously, at room temperature, solution hardening is exclusively strengthening mechanism for the solutionized alloy. The solutionized alloy has a little higher strength than the as-cast alloy. This indicates that the solution treatment dissolved the primary carbides and consequently, increased alloying element content in the matrix that hardened the matrix more effectively. At the same time, this also demonstrates that the roomtemperature mechanical property of the alloy is principally dependent on the matrix, rather than the grain and interdendritic boundaries. The solution matrix with the preferential orientation exhibits an excellent ductility, which is even much higher than that of the as-cast alloy. This is attributed to absence of the primary carbides at grain boundaries and interdendritic regions, which would restrain slip. During the stress – rupture test, a great number of the secondary M23C6 carbide precipitated in the solutionized alloy. The secondary carbide distribution is in the grid form. The regular arrangement of the precipitates demonstrates the effect of stress on the precipitation. Secondary M23C6 carbide precipitated preferentially in slip bands where there was a high dislocation density [12]. Evidently, the M23C6 precipitation in the solutionized alloy also followed the direct reaction Eq. (1). The well-distributed secondary precipitation indicates that the solution treatment homogenized the matrix, and there is an even distribution of alloying elements, particularly carbon, which were released by the dissolution of the primary carbides. It is noted that the secondary precipitation also happened at grain boundaries and interdendritic regions, but the precipitates are relatively small. Differently from the solution strengthening at room temperature, the solutionized alloy was hardened effectively by M23C6 during the stress – rupture test. The effective precipitation hardening in the matrix gives the alloy a longer rupture life, relative to the as-cast alloy. Conversely to the room-temperature high ductility, however, the solutionized alloy exhibits a very low stress–rupture ductility. The fracture surface observation (Fig. 9b) indicates that fracture is a brittle and transgranular type. Unlike at room temperature, at high temperature, the boundaries of grains and interdendrites in superalloys usually are very critical to deformation and fracture. Directional solidification eliminated the transverse grain boundaries in the
107
DZ40M alloy, reducing the effect of grain boundaries. But, the migration of longitudinal boundaries of grains and interdendrites happened in the absence of coarse carbides, which would lock them. The little deformation took place in the matrix, which is strengthened effectively by M23C6 precipitates. The matrix made little contribution to the alloy ductility. It was the low strength of boundaries of grain and interdendrite that caused a premature stress rupture. The solutionized alloy does not possess a good microstructural combination of the matrix with the grain boundaries and interdendritic regions. The solution plus aging treatments produced a great number of the secondary M23C6 precipitates throughout the alloy matrix. This demonstrates that at the aging temperature, M23C6 is thermodynamically stable. The secondary precipitation is unevenly distributed. They occurred preferentially on imperfections, such as grain and subgrain boundaries. TEM observation showed that dislocations and stacking faults were preferential sites of M23C6 precipitation [12]. In other austenitic alloys, the non-homogeneous secondary M23C6 precipitation was also found [13,14]. Undoubtedly, the mechanism of M23C6 precipitation in the aged alloy can be described by the direct reaction Eq. (1). The non-homogeneous precipitation is related to an uneven distribution of alloying elements, mainly carbon. As discussed previously, carbon content is a predominant factor of the M23C6 precipitation in the matrix. It is well known that dislocations and stacking faults have a tendency to attract carbon atoms, forming carbon segregation zone, that is, the famous ‘Cottrell atmosphere’ and ‘Suzuki atmosphere’, respectively. Also, grain boundaries accommodated more carbon atoms rejected by the decomposition of the primary carbides. Subgrain boundaries are composed of dislocations [12] and consequently, tend to attract carbon atoms, forming a carbon segregation. These imperfections made a preparation in chemistry for the M23C6 formation and became preferential sites of M23C6 precipitation. The secondary M23C6 precipitation made a significant contribution to mechanical property of the aged alloy. Compared with the as-cast and solutionized alloy, the aged alloy has a higher room-temperature strength. The precipitation hardening is its most important strengthening mechanism. Furthermore, this indicates that the precipitation hardening is more effective than solution one in strengthening the DZ40M alloy at room temperature. On the other hand, the precipitation hardening restrains deformation of the matrix, reducing the alloy ductility severely. During the stress–rupture test, a microstructural evolution also occurred and the M23C6 carbide coarsened. Because of their relatively high surface energy, the large rod-like M23C6 precipitates were thermodynamically
108
W.H. Jiang et al. / Materials Science and Engineering A271 (1999) 101–108
unstable and disappeared. Their absence was beneficial for stress–rupture properties. At the same time, the intergranular precipitates became coarser. The microstructure of the alloy approached an equilibrium state. The aged alloy exhibits an excellent stress–rupture property. This should be attributed to a good microstructural combination of the matrix with the grain boundaries and interdendritic regions. Fracture surface (Fig. 9c) shows that there is an even distribution of ductile microvoids. This demonstrates that both the matrix and the boundaries of grains and interdendrites have a comparative strength, resulting in attainment of a higher strength for the alloy. The alloy matrix was strengthened by the fine precipitates, while the boundaries of grains and interdendrites were locked by the coarse precipitates. The present work shows that a propriate heat treatment could improve the microstructure of the alloy and increase its strength at both room and high temperatures substantially.
The precipitation hardening is the most important strengthening mechanism at high temperature for the alloy in as-cast, solutionized and aged states. The stress–rupture properties were dominated by the matrix and the boundaries of grains and interdendrites. The low rupture life of the as-cast alloy and low rupture ductility of the solutionized alloy are due to the unevenly distributed secondary M23C6 precipitation and absence of coarse precipitates at grain boundaries and interdendritic regions, respectively. The superior stress– rupture property of the aged alloy is attributed to its good microstructural combination of the matrix with the boundaries of grains and interdendrites. Heat treatment is an effective way to improve mechanical property of the DZ40M alloy with an excellent ductility.
References 5. Conclusions The microstructure of the DZ40M alloy can be modified by heat treatment. Solution treatment at 1280°C for 4 h dissolved the primary carbides essentially and the alloy became a single-phase supersaturated solid solution. Incorporation of aging treatment at 950°C for 12 h produced a profusion of secondary M23C6 precipitation throughout the matrix. The room-temperature mechanical properties of the alloy are mainly dependent on the microstructures of the matrix. For the alloy in the as-cast and solutionized states, solution hardening is their important strengthening mechanism, while for the aged one, the secondary M23C6 precipitation makes a significant contribution. During the stress – rupture tests, the microstructural evolution occurred in the alloy, the primary carbides dissolved sluggishly and the stable secondary M23C6 precipitated heavily.
.
[1] C.T. Sims, J. Metal. 21 (1969) 27. [2] Y. Fukui, R. Sasaki, F. Hataya, T. Kashimura, Metall. Trans. A 12A (1981) 1033. [3] A. Davin, J.M. Drapier, D. Coutsouradis, Cobalt (2) (1974) 32. [4] R.F. Vandermousen, P. Viatour, J.M. Drapier, D. Coutsouradis, Cobalt (1) (1974) 6. [5] C.T. Sims, W.C. Hagel, The Superalloy, John Wiley & Sons, New York, 1972, pp. 145 – 174. [6] J.M. Drapier, V. Leroy, C. Dupont, D. Coutsouradis, L. Habraken, Cobalt (12) (1968) 199. [7] X.D. Yao, J.H. Zhang, Z.Y. Zhang, Y.A. Li, N.R. Zhao, H.R. Guan, Z.Q. Hu, Acta Metall. Sinica 31A (1995) 320 (in Chinese). [8] H.M. Tawancy, J. Mater. Sci. 18 (1983) 2976. [9] W.H. Jiang, X.D. Yao, H.R. Guan, Z.Q. Hu, J. Mater. Sci. 34 (1999) 2859. [10] W.H. Jiang, X.D. Yao, H.R. Guan, Z.Q. Hu, Metall. Trans. A 30A (1999) 513. [11] W.H. Jiang, X.D. Yao, H.R. Guan, Z.Q. Hu, Acta Metall. Sinica, 35A (1999) 253 (in Chinese) [12] W.H. Jiang, H.R. Guan, Z.Q. Hu, Scr. Metall., (1999) submitted for publication. [13] F.R. Beckitt, B.R. Clark, Acta Metall. 15 (1969) 113. [14] L.K. Singhal, Metall. Trans. 2 (1971) 1267.
Effects of heat treatment on microstructures and mechanical properties of a directionally solidified cobalt-base superalloy W.H. Jiang a,b,*, H.R. Guan a, Z.Q. Hu a a
Department of Superalloys, Institute of Metal Research, Academia Sinica, No. 72 Wenhua Road, Shenyang 110015, P.R. China b Department of Metal Materials Engineering, Shenyang Polytechnic Uni6ersity, Shenyang 110023, P.R. China Received 26 January 1999; received in revised form 6 April 1999
Abstract DZ40M alloy is a newly developed directionally solidified cobalt-base superalloy. The present work investigated microstructures, room-temperature tensile and stress–rupture properties at 980°C/83 MPa of the alloy in as-cast, solutionized as well as aged states. The microstructure of the DZ40M alloy can be modified by heat treatment. Solution treatment at 1280°C for 4 h dissolved the primary carbides essentially and the alloy became a single-phase supersaturated solid solution. Incorporation of aging treatment at 950°C for 12 h produced a profusion of secondary M23C6 precipitation throughout the matrix. The room-temperature mechanical properties of the alloy are mainly dependent on the microstructures of the matrix. During the stress – rupture tests, the microstructural evolution occurred in the alloy, the primary carbides dissolved sluggishly and the secondary M23C6 precipitated heavily. The precipitation hardening is the most important strengthening mechanism at high temperature for the alloy in all three states. The stress–rupture properties were dominated by both the matrix and the boundaries of grains and interdendrites. The aged alloy has a superior stress– rupture property, which is attributed to its good microstructural combination of the matrix and the boundaries of grains and interdendrites. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Cobalt-base superalloy; Directional solidification; Heat treatment; Mechanical property; Microstructure
1. Introduction Cobalt-base superalloys are widely used in industrial and aircraft turbines for vanes and combuster sections because of their intrinsic properties such as good stress–rupture parameters and excellent hot corrosion and oxidation resistance. The demand for longer lives in turbines is the driving force to improve their mechanical properties. Some efforts have been made and substantial progress has been achieved. The work has focused on composition modification, alloy development, and solidification mode [1 – 4]. However, despite the fact that precipitation hardening is their important strengthening mechanism, the potential for strengthening alloys by heat treatment has not been well developed. In industrial practice, cobalt-base superalloys are usually used in as-cast state [5]. * Corresponding author. E-mail address: [email protected] (W.H. Jiang)
As-cast cobalt-base superalloys consist of a continuous fcc matrix and a variety of carbides, mainly coarse primary M23C6, M7C3 and MC. The carbides contribute significantly to strengthening. The carbides form as alloys cooling down in shell mold. They precipitate at grain boundaries and in interdendritic regions. During service at high temperature, secondary carbides, usually M23C6, precipitate. The fine secondary carbide pins up dislocations that hardens the matrix. Obviously, the morphology, distribution and size of the secondary carbides affect precipitation hardening effect. However, little attention has been paid to control their precipitation, for heat treatment deteriorates their room-temperature ductility severely [5,6]. Lately, a modified directionally solidified X-40 (DZ40M) alloy was developed in the Institute of Metal Research, Academia Sinica [7]. The directional solidification not only eliminates transverse grain boundaries, but also produces a structure of columnar grain matrix
0921-5093/99/$ - see front matter © 1999 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 9 ) 0 0 1 8 3 - 5
W.H. Jiang et al. / Materials Science and Engineering A271 (1999) 101–108
102
Table 1 Tensile properties of DZ40M alloy at room temperature State
Ultimate tensile strength, MPa
0.2% yield strength, MPa
Elongation, %
As-cast Solutionized Aged
672 729 1039
422 455 649
19.0 36.0 9.0
with a 001 direction parallel to the specimen axis, i.e. 001 preferential orientation. The alloy has an excellent ductility. For example, its room-temperature tensile elongation reaches up to about 20%, while that of conventional equiaxed as-cast X-40 is only 8% [4]. Its excellent ductility provides a prerequisite to exploit heat treatment to increase alloy strength further. For development of heat treatment, it is essential to understand a relationship between microstructure and mechanical property. The present work investigated microstructures and mechanical properties at room and high temperatures of DZ40M alloy in as-cast, solutionized as well as aged states.
2. Experimental procedure The alloy was prepared in a conventional vacuum furnace with a mold withdrawal device. The nominal chemical composition of the alloy (in wt.%) is 25Cr, 11Ni, 7.5W, 0.45C, 0.05B, 0.8Al, 0.2Mo, 0.25Ta, 0.15Ti, 0.15Zr and the balance Co. Cylindrical rods of the alloy, 16 mm in diameter and 140 mm long, were produced at a withdrawal speed of 7 mm/min and with a thermal gradient of about 50 K/cm at the solid/liquid interface. Three typical alloy states were selected for this investigation: as-cast, solutionized as well as aged. The solution treatment was performed at 1280°C for 4 h, and followed by air cooling. The aging treatment was carried out at 950°C for 12 h, which was applied after solution treatment. The alloy microstructures were examined with scanning electron microscope (SEM) with energy dispersive X-ray analysis (EDAX). The samples for SEM observation were etched electrolytically in a solution consisting of 42 pct H3PO4, 34 pct H2SO4 and 24 pct H2O. High magnification observation was carried out using transmission electron microscope (TEM). Electropolishing of the thin foils was done with an electrolyte consisting of 5 pct perchloric acid in methanol at about − 20°C. Identification of phases was carried out by X-ray diffraction using filtered chromium radiation. An investigation was made on carbides extracted from the anodically dissolved matrix of the as-cast alloy. Room-temperature tensile and high temperature stress–rupture properties were measured along the
specimen axis. The stress–rupture tests were conducted at 980°C in air at a stress level of 83 MPa. After the stress–rupture tests, fracture surfaces and microstructures were observed with SEM.
3. Results
3.1. Mechanical properties The tensile test data obtained for the alloy in the three selected states at room temperature are given in Table 1. It can be seen that the tensile properties are closely related to the alloy states. The as-cast alloy has the lowest room-temperature strength, but a high ductility, which should be attributed to its columnar grain structure with 001 preferential orientation. The 1280°C solution treatment caused a little increase in strength, and almost doubled tensile elongation, compared to the as-cast alloy. Incorporation of 950°C/12 h aging treatment reduced tensile ductility markedly, but increased strength by 50%. It is noted that the aged DZ40M alloy has the highest room-temperature strength, and an adequate ductility which is comparative to that of the equiaxed as-cast X-40 [4]. The room-temperature mechanical property of the aged alloy is superior to that of as-cast MAR-M509, a high strength cobalt-base superalloy, whose yield strength and elongation are 519 MPa and 4.5%, respectively [6]. The stress–rupture tests at 980°C and 83 MPa were conducted for the alloy in the three states. The results are summarized in Table 2. It demonstrates that the alloy states have a more pronounced effect on high temperature stress–rupture property. The as-cast alloy possesses the shortest stress–rupture life and the highest stress–rupture ductility. Relative to the as-cast alloy, stress–rupture life was prolonged three and five times by the solution and aging treatments, respectively. However, at the same time, a severe loss in rupture Table 2 Stress–rupture properties of DZ40M alloy at 980°C/83 MPa State
Stress–rupture life, h
Elongation, %
As-cast Solutionized Aged
59.7 188.8 278.1
56.0 4.4 28.0
W.H. Jiang et al. / Materials Science and Engineering A271 (1999) 101–108
103
Fig. 1. Optical micrograph of the as-cast DZ40M alloy showing a columnar grain structure.
Fig. 3. Back-scattered electron micrograph of primary carbides in the as-cast DZ40M alloy.
ductility accompanied. The elongation of the solutionized alloy is only 4.4%, while that of the as-cast alloy is as high as 56%. The aging treatment increased the rupture ductility to 28%. Evidently, the aged alloy possesses a good combination of the stress – rupture life and ductility.
of MC is C45.10, Zr24.34, Ta11.00, Ti10.77, Co3.06, W2.84, Cr2.06, and Ni0.82. The scanning electron image (Fig. 3) shows the morphology and distribution of the carbide phases. In this backscattered electron image, the chromium-rich M7C3 carbide appears dark and the MC carbide is rather light. It is observed that M7C3 carbide is in the form of rods or irregular aggregates, while MC is present as a discrete, blocky dispersion with a well-distributed Chinese script morphology. Such a morphology of MC is also observed in MARM509 alloy [1,5]. Evidently, the formation of MC is due to the addition of reactive elements, Ta, Ti and Zr. Both the M7C3 and MC carbides are located at grain boundaries or in interdendritic regions, forming a continuous network around the columnar grained matrix. Fig. 4 is the microstructure of the solutionized alloy. It can be seen that the solution treatment at 1280°C for 4 h dissolved the primary carbides essentially and the alloy became a single-phase supersaturated solid solution. Fig. 5 reveals the microstructure of the aged alloy. The aging treatment at 950°C for 12 h gave rise to a profuse secondary precipitation throughout the alloy
3.2. Microstructural characteristic Fig. 1 shows the microstructure of the as-cast DZ40M alloy. Its matrix is a well-developed columnar grain austenite with 001 direction parallel to the growth axis. X-ray diffraction analysis (Fig. 2) and energy dispersive X-ray analysis (EDAX) indicate that the alloy contains two types of coarse primary carbides, i.e. chromium-rich M7C3 and MC which contains tantalum, titanium, zirconium, and tungsten. A typical composition of M7C3 is (in at pct) C29.94, Cr56.14, Co8.35, W3.88, Ni0.77, Mo0.67, Ti0.22 and Zr0.04, while that
Fig. 2. X-ray diffraction pattern of carbides extracted from the as-cast DZ40M alloy.
Fig. 4. Back-scattered electron micrograph showing the microstructure of the solutionized alloy.
104
W.H. Jiang et al. / Materials Science and Engineering A271 (1999) 101–108
{100}precipitate // {100}matrix 001precipitate // 001matrix
Fig. 5. SEM micrograph showing the microstructure of the aged alloy.
matrix. It can be seen that the secondary precipitates is unevenly distributed. They appeared preferentially on grain boundaries and subgrain boundaries which were decorated by them. It is noted that there are a few of the large rod-like precipitates. They seem to be sigma phase. EDAX was used to characterize their composition. From the obtained profile shown in Fig. 6, it is learnt the rod-like precipitates are chromium-rich carbide rather than sigma phase. In Fig. 6, relatively high cobalt peak resulted from the alloy matrix due to narrowness of the rod-like precipitates and a limited resolution of EDAX. TEM observation (Fig. 7) shows the blocky and rod-like precipitates. Electron diffraction analysis demonstrates that these precipitates possess the same crystal structure and their characteristic reflections are present at every one-third position of the fcc matrix reflection. This indicates that the precipitates have a fcc structure with a lattice constant that is nearly three times that of the matrix, and they also have a cube–cube orientation relationship with the matrix:
These are characteristic features of the chromium-rich M23C6 carbide [8]. Fig. 8 depicts the microstructures of the alloy in three states after the stress–rupture tests. Compared with the original ones, substantially microstructural changes have occurred in the alloy in all three states. In the as-cast specimen ruptured, a profusion of fine secondary precipitates were produced in the matrix and the primary carbides thinned (Fig. 8a). However, the coarse primary carbides were still retained, indicating that their dissolution was sluggish. It is worthy of note that the secondary precipitates are unevenly distributed. Close to the primary carbides, there is dense distribution. After the stress–rupture test, the solutionized alloy also contained a great number of fine secondary precipitates, as shown in Fig. 8b. They are evenly distributed and aligned in rows to form a grid. Evidently, during the stress–rupture test, the solutionized alloy subjected to a stress treatment. The regular arrangement of the secondary precipitates resulted from a stress effect. Furthermore, it is observed that a copious intergranular precipitation occurred, forming a continuous chain. But, the intergranular precipitates are very fine. Fig. 8c shows the microstructure of the aged alloy after the stress–rupture test. It has the most dispersed secondary precipitates. Obviously, the microstructural homogeneity was improved relative to the aged one. Coarse and fine precipitates coexist in the matrix, indicating that their coarsening and reprecipitation happened at the same time during the test. Particularly, it is noted that the rod-like M23C6 disappeared. Furthermore, it can be seen that intergranular precipitates are coarse and comparative to those in the as-cast alloy ruptured. Usually, in cast cobalt-base superalloys, secondary M23C6 carbide is more thermodynamically stable, and during service or exposure at high temperature, its precipitation is very common. In fact, in the DZ40M alloy, a heavy precipitation of secondary M23C6 carbide in the matrix occurred during high temperature low cycle fatigue, creep and thermal aging [9–11]. Therefore, it is believed that the secondary precipitates must be M23C6 in the present alloy in all three states after the stress–rupture tests.
3.3. Fracture surfaces after the stress–rupture tests
Fig. 6. EDAX spectrum of the rod-like precipitates in the DZ40M alloy aged at 950°C/12 h.
After the stress–rupture tests, the fracture surfaces of the alloy in all three states were observed with the SEM. Fig. 9a shows the fracture appearance of the as-cast alloy, where there are a great number of fine voids. This fracture is a ductile type. Microvoid coalescence is its fracture mechanism. Fig. 9b shows the fracture appearance of the solutionized alloy. Cracking
W.H. Jiang et al. / Materials Science and Engineering A271 (1999) 101–108
105
Fig. 7. (a) TEM bright field micrograph of the fine precipitates formed in the aged alloy, (b) dark field micrograph of the same region, using {111} precipitate reflection and (c) corresponding diffraction pattern showing a cube – cube matrix: carbide orientation relationship.
of longitudinal grain boundaries and interdendritic regions was observed. The fracture is transgranular and brittle. The matrix reveals very limited amount of plastic deformation. Fig. 9c reveals the fracture appearance of the aged alloy, where there are uniformly distributed voids. The fracture is a ductile type. Although its fracture mechanism also is coalescence of microvoids, the number of microvoids is less than that of the as-cast alloy.
4. Discussion Like other cast cobalt-base superalloys, the as-cast DZ40M alloy consists of a cobalt-base austenitic matrix and coarse primary carbides located at grain boundaries and interdendritic regions. However, the directional solidification developed a structure of columnar grain matrix with 001 preferential orientation. From Table 1, it is learnt that the as-cast alloy has a yield strength not higher but a little lower than that of the solutionized alloy. This demonstrates that the room-temperature strength of the as-cast alloy resulted mainly from the matrix and little from the primary carbides. At room temperature, its strengthening mech-
anism is solution hardening. At the same time, the alloy matrix with 001 preferential orientation shows an excellent ductility, for the slip system 110{111} of the fcc matrix is located at the easiest slip direction when load is applied along the specimen axis, i.e. 001 direction. At high temperature, the as-cast alloy was thermodynamically unstable. The primary carbides, both the M7C3 and MC, dissolved sluggishly and the secondary M23C6 carbide precipitated profusely. The M23C6 precipitates are unevenly distributed. There are a dense distribution around the primary carbides and a sparse distribution in the center of grains, as shown in Fig. 8a. The secondary M23C6 precipitation is closely related to the decomposition of the primary carbides. A previous work by the present authors [10] suggested a direct reaction mechanism for the M23C6 precipitation, i.e. 23M+ 6C M23C6
(1)
As a carbon reservoir, the primary carbides provided the reaction Eq. (1) with carbon atoms. Since during solidification carbon atoms were exhausted to form the primary carbides and there was little of carbon in the matrix, carbon content became a predominant factor of the precipitation reaction Eq. (1). It was the uneven
106
W.H. Jiang et al. / Materials Science and Engineering A271 (1999) 101–108
distribution of carbon that caused non-homogeneous secondary carbide precipitation. During the stress – rupture test, the microstructural evolution in the as-cast alloy brought about a precipitation hardening, which was absent at room temperature. The retained coarse primary carbides strengthened grain and interdendritic boundaries, retarding their migration. The fine secondary M23C6 carbide pinned up dislocations that hardened the matrix. Due to their relatively sparse distribution, however, the precipitation strengthening was less effective in the center of grains. The fracture surface observa-
Fig. 9. SEM micrographs showing the fracture surfaces of the alloy in all three states after the stress – rupture tests: (a) as-cast, (b) solutionized, and (c) aged.
Fig. 8. SEM micrographs showing the microstructures of the alloy in three states after the stress–rupture tests: (a) as-cast, (b) solutionized, and (c) aged.
tion shows that a substantial plastic deformation took place, which is probably related to low strength in the center of grains. The non-homogeneous microstructure and consequently uneven strength in the matrix resulted in a low stress–rupture life. At high temperature, an extraordinarily high ductility was developed by the matrix, like the situation at room temperature.
W.H. Jiang et al. / Materials Science and Engineering A271 (1999) 101–108
The solution treatment at 1280°C caused a complete dissolution of the primary carbides, both the M7C3 and MC, and no secondary M23C6 precipitation (Fig. 4). This indicates that at the solution temperature both the primary and secondary carbides are thermodynamically unstable. So, it is feasible to develop a single-phase supersaturated solid solution for the DZ40M alloy by solution treatment. This is especially significant for development of heat treatment for the alloy. Obviously, at room temperature, solution hardening is exclusively strengthening mechanism for the solutionized alloy. The solutionized alloy has a little higher strength than the as-cast alloy. This indicates that the solution treatment dissolved the primary carbides and consequently, increased alloying element content in the matrix that hardened the matrix more effectively. At the same time, this also demonstrates that the roomtemperature mechanical property of the alloy is principally dependent on the matrix, rather than the grain and interdendritic boundaries. The solution matrix with the preferential orientation exhibits an excellent ductility, which is even much higher than that of the as-cast alloy. This is attributed to absence of the primary carbides at grain boundaries and interdendritic regions, which would restrain slip. During the stress – rupture test, a great number of the secondary M23C6 carbide precipitated in the solutionized alloy. The secondary carbide distribution is in the grid form. The regular arrangement of the precipitates demonstrates the effect of stress on the precipitation. Secondary M23C6 carbide precipitated preferentially in slip bands where there was a high dislocation density [12]. Evidently, the M23C6 precipitation in the solutionized alloy also followed the direct reaction Eq. (1). The well-distributed secondary precipitation indicates that the solution treatment homogenized the matrix, and there is an even distribution of alloying elements, particularly carbon, which were released by the dissolution of the primary carbides. It is noted that the secondary precipitation also happened at grain boundaries and interdendritic regions, but the precipitates are relatively small. Differently from the solution strengthening at room temperature, the solutionized alloy was hardened effectively by M23C6 during the stress – rupture test. The effective precipitation hardening in the matrix gives the alloy a longer rupture life, relative to the as-cast alloy. Conversely to the room-temperature high ductility, however, the solutionized alloy exhibits a very low stress–rupture ductility. The fracture surface observation (Fig. 9b) indicates that fracture is a brittle and transgranular type. Unlike at room temperature, at high temperature, the boundaries of grains and interdendrites in superalloys usually are very critical to deformation and fracture. Directional solidification eliminated the transverse grain boundaries in the
107
DZ40M alloy, reducing the effect of grain boundaries. But, the migration of longitudinal boundaries of grains and interdendrites happened in the absence of coarse carbides, which would lock them. The little deformation took place in the matrix, which is strengthened effectively by M23C6 precipitates. The matrix made little contribution to the alloy ductility. It was the low strength of boundaries of grain and interdendrite that caused a premature stress rupture. The solutionized alloy does not possess a good microstructural combination of the matrix with the grain boundaries and interdendritic regions. The solution plus aging treatments produced a great number of the secondary M23C6 precipitates throughout the alloy matrix. This demonstrates that at the aging temperature, M23C6 is thermodynamically stable. The secondary precipitation is unevenly distributed. They occurred preferentially on imperfections, such as grain and subgrain boundaries. TEM observation showed that dislocations and stacking faults were preferential sites of M23C6 precipitation [12]. In other austenitic alloys, the non-homogeneous secondary M23C6 precipitation was also found [13,14]. Undoubtedly, the mechanism of M23C6 precipitation in the aged alloy can be described by the direct reaction Eq. (1). The non-homogeneous precipitation is related to an uneven distribution of alloying elements, mainly carbon. As discussed previously, carbon content is a predominant factor of the M23C6 precipitation in the matrix. It is well known that dislocations and stacking faults have a tendency to attract carbon atoms, forming carbon segregation zone, that is, the famous ‘Cottrell atmosphere’ and ‘Suzuki atmosphere’, respectively. Also, grain boundaries accommodated more carbon atoms rejected by the decomposition of the primary carbides. Subgrain boundaries are composed of dislocations [12] and consequently, tend to attract carbon atoms, forming a carbon segregation. These imperfections made a preparation in chemistry for the M23C6 formation and became preferential sites of M23C6 precipitation. The secondary M23C6 precipitation made a significant contribution to mechanical property of the aged alloy. Compared with the as-cast and solutionized alloy, the aged alloy has a higher room-temperature strength. The precipitation hardening is its most important strengthening mechanism. Furthermore, this indicates that the precipitation hardening is more effective than solution one in strengthening the DZ40M alloy at room temperature. On the other hand, the precipitation hardening restrains deformation of the matrix, reducing the alloy ductility severely. During the stress–rupture test, a microstructural evolution also occurred and the M23C6 carbide coarsened. Because of their relatively high surface energy, the large rod-like M23C6 precipitates were thermodynamically
108
W.H. Jiang et al. / Materials Science and Engineering A271 (1999) 101–108
unstable and disappeared. Their absence was beneficial for stress–rupture properties. At the same time, the intergranular precipitates became coarser. The microstructure of the alloy approached an equilibrium state. The aged alloy exhibits an excellent stress–rupture property. This should be attributed to a good microstructural combination of the matrix with the grain boundaries and interdendritic regions. Fracture surface (Fig. 9c) shows that there is an even distribution of ductile microvoids. This demonstrates that both the matrix and the boundaries of grains and interdendrites have a comparative strength, resulting in attainment of a higher strength for the alloy. The alloy matrix was strengthened by the fine precipitates, while the boundaries of grains and interdendrites were locked by the coarse precipitates. The present work shows that a propriate heat treatment could improve the microstructure of the alloy and increase its strength at both room and high temperatures substantially.
The precipitation hardening is the most important strengthening mechanism at high temperature for the alloy in as-cast, solutionized and aged states. The stress–rupture properties were dominated by the matrix and the boundaries of grains and interdendrites. The low rupture life of the as-cast alloy and low rupture ductility of the solutionized alloy are due to the unevenly distributed secondary M23C6 precipitation and absence of coarse precipitates at grain boundaries and interdendritic regions, respectively. The superior stress– rupture property of the aged alloy is attributed to its good microstructural combination of the matrix with the boundaries of grains and interdendrites. Heat treatment is an effective way to improve mechanical property of the DZ40M alloy with an excellent ductility.
References 5. Conclusions The microstructure of the DZ40M alloy can be modified by heat treatment. Solution treatment at 1280°C for 4 h dissolved the primary carbides essentially and the alloy became a single-phase supersaturated solid solution. Incorporation of aging treatment at 950°C for 12 h produced a profusion of secondary M23C6 precipitation throughout the matrix. The room-temperature mechanical properties of the alloy are mainly dependent on the microstructures of the matrix. For the alloy in the as-cast and solutionized states, solution hardening is their important strengthening mechanism, while for the aged one, the secondary M23C6 precipitation makes a significant contribution. During the stress – rupture tests, the microstructural evolution occurred in the alloy, the primary carbides dissolved sluggishly and the stable secondary M23C6 precipitated heavily.
.
[1] C.T. Sims, J. Metal. 21 (1969) 27. [2] Y. Fukui, R. Sasaki, F. Hataya, T. Kashimura, Metall. Trans. A 12A (1981) 1033. [3] A. Davin, J.M. Drapier, D. Coutsouradis, Cobalt (2) (1974) 32. [4] R.F. Vandermousen, P. Viatour, J.M. Drapier, D. Coutsouradis, Cobalt (1) (1974) 6. [5] C.T. Sims, W.C. Hagel, The Superalloy, John Wiley & Sons, New York, 1972, pp. 145 – 174. [6] J.M. Drapier, V. Leroy, C. Dupont, D. Coutsouradis, L. Habraken, Cobalt (12) (1968) 199. [7] X.D. Yao, J.H. Zhang, Z.Y. Zhang, Y.A. Li, N.R. Zhao, H.R. Guan, Z.Q. Hu, Acta Metall. Sinica 31A (1995) 320 (in Chinese). [8] H.M. Tawancy, J. Mater. Sci. 18 (1983) 2976. [9] W.H. Jiang, X.D. Yao, H.R. Guan, Z.Q. Hu, J. Mater. Sci. 34 (1999) 2859. [10] W.H. Jiang, X.D. Yao, H.R. Guan, Z.Q. Hu, Metall. Trans. A 30A (1999) 513. [11] W.H. Jiang, X.D. Yao, H.R. Guan, Z.Q. Hu, Acta Metall. Sinica, 35A (1999) 253 (in Chinese) [12] W.H. Jiang, H.R. Guan, Z.Q. Hu, Scr. Metall., (1999) submitted for publication. [13] F.R. Beckitt, B.R. Clark, Acta Metall. 15 (1969) 113. [14] L.K. Singhal, Metall. Trans. 2 (1971) 1267.