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Microstructural Evolution of a Ni-base Alloy DZ468 Joint Bonded with a New Co-base Filler
Available online at ScienceDirect
ScienceDirect J. Mater. Sci. Technol., 2014, 30(5), 480e486
Microstructural Evolution of a Ni-base Alloy DZ468 Joint Bonded with a New Co-base Filler Yanhong Jing, Zhi Zheng, Enze Liu*, Yi Guo Division of Superalloy, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China [Manuscript received November 10, 2012, in revised form November 20, 2012, Available online 8 January 2014]
Ni-base alloy DZ468 has been joined by transient liquid phase bonding technique with a newly developed Co-based filler. The microstructures of the Co-base filler and the joint, the effects of heat treatment on microstructure and hardness of the joint have been investigated by various experimental methods. Results show that the Co-base filler consists of g, M2B, M5B3 and M23B6 phases. Because of the interdiffusion between the base metal and the filler, g, MC, M5B3 and M23B6 phases are formed in the bonding zone. And localized liquidation of substrate occurs in the diffusion affected zone, with MC and M3B2 precipitating in this area. During heat treatment, the volume of the intermetallic phases in the bonding zone resulting from incomplete isothermal solidification decreases obviously. On the contrary, the width of the diffusion affected zone increases at the solution stage and subsequently decreases at the aging stages. KEY WORDS: Transient liquid phase (TLP) bonding; Microstructure evolution; Ni-base alloy; Co-base filler
1. Introduction DZ468 alloy is a directionally solidified superalloy that was developed by the low segregation technique and phase calculation[1e4]. It has not only superior mechanical property, but also good hot corrosion resistance[5,6]. Those excellent characteristics make DZ468 alloy a promising candidate in turbine engine components serving in marine atmosphere. However, it is very difficult to produce the complicated intrastructures just by casting. And welding is essential to produce component with complex structure. Unfortunately, the high alloying of DZ468 alloy, especially considerable contents of g0 promoting elements, reduces the weldability of this superalloy substantially[7,8]. In 1970s, a new welding method, transient liquid phase (TLP) bonding was developed by Duvall et al.[9], which is a hybrid process that combines the beneficial features of liquid phase joining and diffusion bonding techniques. This process can be successfully used to join the heat resistant alloys which are inherently susceptible to hot cracking or post-weld heat treatment cracking free from substantial pressure. Therefore, it is so suitable for the TLP bonding technology to bond DZ468 alloy. Furthermore, it is reasonable to assume that if both the bonding Corresponding author. Assoc. Prof., Ph.D.; Tel.: þ86 24 23971143; Fax: þ86 24 23971927; E-mail address: [email protected] (E. Liu). 1005-0302/$ e see front matter Copyright Ó 2014, The editorial office of Journal of Materials Science & Technology. Published by Elsevier Limited. All rights reserved. http://dx.doi.org/10.1016/j.jmst.2013.12.010 *
process and the heat treatment method of the base metal are welldesigned to accord with each other, the cost of production can be reduced greatly. However, existing fillers cannot bond DZ468 alloy well. Considering that Co-base alloy has superior hot corrosion resistance to Ni-base alloy, a Co-base filler is developed to solve the problem. Due to the great effect of microstructure on properties, a deep understanding on the microstructure evolution of the joint is very necessary. Although lots of similar researches were conducted in this area, almost no research work has been conducted for the specific alloy and the relevant filler. So in the present work, the microstructures of both the newly developed Co-base filler and its joint were characterized. In addition, the microstructure evolution and microhardness of the joint during heat treat were investigated as well.
2. Experimental DZ468 alloy used as base alloy in this work was prepared by high rate directional solidification in ZGD2 vacuum furnace. The temperature gradient was 333e353 K/cm and the withdrawal rate was 7 mm/min. The Co-base filler powder with a size less than 45 mm was produced by ultrasonic gas atomization. A flexible insert alloy cloth consisting of the alloy powder and some organism which functions to glue the loose powders was produced by thermal rolling. Their nominal compositions (weight percent) and melting ranges are shown in Table 1. DZ468 alloy was machined into disk of 415 mm 3 mm.
Y. Jing et al.: J. Mater. Sci. Technol., 2014, 30(5), 480e486
Table 1 Nominal compositions and melting ranges of DZ468 alloy and the Co-base filler (wt%) Al Mo W Co Cr Ti Ta Re
C
B
Ni
Melting range (K)
DZ468 5.2 1 5 8.5 12 0.5 5 2 0.06 0.01 Bal. 1533e1633 Filler e e 10 Bal. 22 e e e e 2 20 1366e1517
Then they are manually grinded to 600-grit, degreased with acetone in an ultrasonic bath for 15 min and cleaned with alcohol prior to joining. The flexible cloth was inserted into two pieces of DZ468 disks. And the assembly was put in a vacuum furnace with an atmosphere of 1 103e5 103 Pa at 1553 K for 5 min, and a pressure of 1 104 Pa was applied to the top disk at this time. After bonding, the specimens were furnace cooled in the vacuum chamber to room temperature. The bonded samples were heat treated sequentially to different stages according to the optimized heat treatment schedule of the DZ468 alloy. The heat treatment schedule was: 1513 K/0.5 h þ 1533 K/0.5 h þ 1553 K/2 h þ AC (solution stage), 1393 K/4 h/FC þ 1353 K/4 h/AC (1st aging stage), 1173 K/4 h þ AC (2nd aging stage) (FC: furnace cooling, AC: air cooling)[10]. The joints and the as-cast Co-base filler were sectioned transversely, mechanically polished and then etched in a 20% phosphoric acid (H3PO4) solution at 5 V for 15 s for observation by metallographic optical microscopy (OM) and field emission scanning electron microscopy (SEM). The average width and area fraction of the intermetallics in the bonding zone, and the width of the diffusion affected zone were measured at a magnification of 500. Averages of 20 measurements across the joint sections were taken from each specimen. To identify the chemical composition of each phase, the energy dispersive spectroscopy (EDS) and electron probe microanalysis (EPMA) were employed, and at least three points were analyzed from each phase. The section of the multilayer joint and the Co-base filler were polished for the X-ray diffraction (XRD) phase identification. Thin foils for transmission electron microscopy (TEM) were sectioned from the multilayer joint and the Co-base filler. The TEM foils were prepared by ion thinning method. Microhardness across the brazed region was measured using a LM247AT hardness tester with 50 g load for 15 s.
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3. Results 3.1. Microstructure of the as-cast Co-base filler Fig. 1 shows the SEM image and the XRD pattern of the ascast Co-base filler. It can be seen that there are three phases in the matrix. The constituent of the matrix g is 50.24Coe23.77Nie 23.28Cre2.71W (at.%) determined by EDS. Table 2 shows the chemical compositions of the rest of the phases, which exhibit that the dark net-like phase is (Co, Cr)-rich boride, the bright white net-like phase is (Co, Cr, W)-rich boride and the white skeleton-like phase is Cr-rich boride, respectively. XRD pattern indicates the presence of g, M5B3, M23B6 and M2B phases. Furthermore, the TEM analysis results, as shown in Fig. 2, demonstrate the crystal structures of the phases. The (Co, Cr)rich boride has the tetragonal crystal structure and can be determined to be M2B phase with the lattice parameters of a ¼ 0.51 nm and c ¼ 0.425 nm, the Cr-rich boride has the tetragonal structure and can be determined to be M5B3 with lattice parameters of a ¼ 0.544 nm and c ¼ 1.007 nm, and lattice parameter of the g phase is 0.352 nm. Combined with XRD result, it can be determined that the (Co, Cr, W)-rich boride should be M23B6. 3.2. Microstructure of as-bond joint Fig. 3 indicates the OM micrograph and the XRD pattern of the joint bonded at 1553 K for 5 min. As can be seen from Fig. 3(a), three layers are formed, which are the bonding zone, the diffusion zone and the substrate, respectively. The bonding zone is composed of isothermally solidified zone and eutectic zone. Diffusion affected zone is the region that exists at the interface between the substrate and bonding zone. XRD pattern indicates the presence of g, g0 , M3B2, M23B6, M5B3 and MC phases. Further observation on the as-bond joint is shown in Fig. 4. Seen from Fig. 4(a), four different phases are formed from the remnant liquid layer in the bonding zone during the cooling of the joint to ambient temperature. The white skeleton-like phase and the bright white net-like phase, which are always associated with the g phase, appear in the bonding zone; the small white blocky phase exists at the edge of the eutectic zone.
Fig. 1 SEM micrograph (a) and XRD pattern (b) from the as-cast Co-base filler.
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Table 2 EPMA results of the intermetallic phases in the matrix of the as-cast Co-base filler (at.%) Co
Ni
Cr
W
B
Dark net-like phase 33.009 5.651 31.353 0.389 29.598 Bright white net-like phase 30.538 6.281 21.992 11.670 29.518 White skeleton-like phase 10.792 1.859 43.884 7.408 36.058
B is detected in the bright white net-like phase and the white skeleton-like phase, and C is detected in the small white blocky phase, although their concentrations could not be determined accurately. EDS compositional analyses of the rest of the elements suggest that the bright white net-like phase should be (Co, Cr, W)-rich boride, the white skeleton-like phase should be Crrich boride, and the small white blocky phase should be (Ta, Ti)-rich carbide (Table 3). Compared with the EDS analyses of the Co-base filler, it can be deduced that the (Co, Cr, W)-rich boride is M23B6 and Cr-rich boride is M5B3. Combined with XRD analysis, (Ta, Ti)-rich carbide can be deduced to be MC carbide. Furthermore, it can be found from Fig. 3(a) that, in the diffusion affected zone, eutectic mixtures in the liquid pool are formed along substrate grain boundaries. Two kinds of eutectic mixtures can be seen from Fig. 4(b). EDS analyses indicate that the two kinds of eutectic mixtures are (Ta, Ti)-rich carbide/g eutectic and Cr-rich boride/g eutectic. Combined with XRD
result, Cr-rich boride is identified to be M3B2 and (Ta, Ti)-rich carbide is MC. EDS was employed to further confirm the Crrich phase to be M3B2. TEM analysis result suggests that M3B2 has a tetragonal structure with lattice parameters of a ¼ 0.57 nm and c ¼ 0.3135 nm (Fig. 4(c)). 3.3. Microstructural character of the joint during heat treatment process Fig. 5 shows the microstructure of the joints obtained at different heat treatment stages. Obviously, during heat treatment, the morphology of the eutectic in the bonding zone and the diffusion affected zone changes greatly, but the eutectic mixtures still exist finally. The intermetallic keeps continuous before aging treatment, but becomes discontinuous after aging stages. In the as-bond specimen, MC particles are distributed at edge of the eutectic zone (Fig. 4). After solution treatment, the morphology of MC carbide changes greatly. It becomes the script shaped morphology from blocky, and exists in the eutectic mixtures rather than at the edge of the eutectic zone. Furthermore, it is worth noting that some needle/particle phases are visible in the matrix g adjacent to the centerline eutectic at aging stages (Fig. 5(g)), which are not observed before aging treatment. EDS analysis reveals that they are Cr-rich boride. Fig. 6 shows the change of the intermetallic area fraction in the bonding zone (AFBZ I ), the intermetallic width (WI) and the width of the diffusion affected zone (WDAZ) of the joint during
Fig. 2 Morphologies of the Cr-rich phase M5B3 (a) and the (Co, Cr)-rich phase M2B (b) from the as-cast filler. The SAED patterns of M5B3 showing [110] zone axis (c), the matrix g showing [001] zone axis (d) and the M2B showing [001] zone axis (e).
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Fig. 3 OM micrograph (a) and XRD pattern (b) of the as-bond joint. DAZ: diffusion affected zone, ISZ: isothermally solidified zone, EZ: eutectic zone.
heat treatment. It can be seen that AFBZ I reduces gradually during heat treatment. In the as-bonded specimen, AFBZ is 17.76%, in I the solution treated specimen, AFBZ is 13.53%, after subsequent I 1st aging treatment, AFBZ changes to be 10.12%, and finally I decreases to 10% after the whole heat treatment. The change of WI is well accordant with AFBZ I during heat treatment. WI at each sequential heat treatment stages is 117, 71, 43 and 41 mm, respectively. Obviously, it can be observed that the volume of the intermetallics changes greatly at solution and 1st aging stage and then almost keeps constant at 2nd aging stage. Furthermore, it can be observed that WDAZ increases at solution treatment stage,
then decreases at subsequent aging stages. Solution treatment results in the increase in the width of diffusion affected zone from 329 mm in as-bond specimen to 655 mm, and then the decrease to 598 mm at 1st aging stage, and finally to 538 mm after the whole heat treatment. 3.4. Microhardness profiles in the bonded joint Microhardness profiles as a function of distance at the different stages of heat treatment are shown in Fig. 7. It can be observed that, in all conditions, the microhardness peak is
Fig. 4 Images from the joint bonded at 1553 K for 5 min: (a) SEM micrograph of the bonding zone, (b) SEM micrograph of the diffusion affected zone, (c) morphology of Cr-rich boride M3B2 in the liquid pool, the inset is the SAED pattern of the phase showing [001] zone axis, (d) micrograph of the liquid pool in the bonding zone.
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Table 3 EDS results of the phases in the as-bond joint and as-cast filler (at.%) White skeletonlike phase (joint)
M5B3 (filler)
Bright white netlike phase (joint)
M23B6 (filler)
Matrix (joint)
g (filler)
White blocky phase (joint)
e e 63.76 13.76 8.73 e 0.68 13.08
e e 70.63 13.81 2.05 e e 13.51
e 1.22 21.27 32.98 14.33 8.24 3.61 18.35
e e 29.67 41.11 8.71 e e 20.67
1.62 1.58 15.95 38.95 38.28 2.64 e 0.98
e e 23.28 50.24 23.77 e e 2.71
e 25.77 5.07 e 4.03 59.32 2.24 1.75
Al Ti Cr Co Ni Ta Mo W
present in the eutectic zone but different at each stage. In as-bond specimen, the hardness peak is 564 HV, and then increases to 737 HV after solution treatment. During subsequent 1st aging process, the hardness peak decreases to 515 HV and finally to 495 HV after the whole heat treatment. It is noted that, after aging heat treatment, a small peak appears in the matrix g adjacent to centerline eutectic. Furthermore, it can be observed that, at each heat treatment stage, the hardness in the diffusion affected zone is similar to that in the substrate, but decreases during heat treatment. The microhardness of substrate in each sequential heat treatment stage is 376, 360, 354 and 275 HV, respectively. 4. Discussion 4.1. Mechanism for the formation of new phases According to the traditional TLP model, when the assembly is held at bonding temperature, the insert filler melts and rapidly attains equilibrium with the solid base metal through the process of melt-back dissolution of the substrate. Subsequently, interdiffusion of alloying elements between the base metal and the liquid commences and causes the melting point of the filler at the solid/liquid interface to increase resulting in isothermal solidification of the liquid. Since the melting point depressing solute B diffuses continuously from the liquid into the base metal, the volume of the liquid that can be maintained at equilibrium decreases, causing solidification to proceed towards the center of the joint from the mating solid surface[11,12]. In the present experiment, due to the lack of bonding time for completing the isothermal solidification of the liquid filler, eutectic microconstituents are formed along the joint centerline. During the bonding process, the dissolution of base metal makes the alloying elements of base metal like Ta, Ti, and C enter into the liquid filler. In addition to dissolution, alloy elements from the base alloy can diffuse into the liquid, which is driven by concentration gradient. Although C is not added in the filler, the decomposition of primary carbide and C atoms in base metal act as the source to supply the free C atoms during the exposure at high temperature. Therefore, in the following solidification process, stable MC-type carbides are formed in the region. However, it is more realistic to assume that both solid/liquid equilibration and solid state diffusion processes operate simultaneously rather than sequentially. The formation of liquid pools and the needle/particle phases in the isothermal solidification matrix g adjacent to the centerline eutectic supports this
assumption. Similar phenomenon was observed in joining nickel using ternary NieSieB insert metals by Gale et al.[13]. When the bonding temperature is below the binary NieB eutectic temperature, the precipitation of Ni3B borides occurs in the substrate at the edge of the insert. In contrast, when the bonding temperature is above the binary NieB eutectic temperature, localized liquation of the substrate material takes place[13]. The formation of liquid pools was also observed during bonding IN738LC alloy using BNi-3 filler at 1473 K[14]. Guo et al.[15] investigated the effect of B on microstructure, and found that M3B2/g eutectic was inclined to form in grain boundaries once the concentration of B exceeded 0.004 at.%. It is because B atom is prone to segregate into grain boundaries, afterwards, eutectic reaction usually occurs at high temperature and produces M3B2/g eutectic. The crown morphology of the liquid pools, as shown in Fig. 4(d), along the solidification direction implies that they are formed from resolidification. It is known that B has a high diffusion rate at high temperature[16,17]. And the maximum solid solubility of B in solid Ni according NieB binary phase diagram is taken to be 0.3 at.% over the range of 1338e1373 K. Therefore, when the assemble coupon is held at 1553 K, the B element diffuses from the gap to the substrate driven by concentration gradient (the initial concentration of B in filler is about 2 wt% (w10 at.%) and DZ468 only just contains 0.01 wt% (w0.06 at.%)). The diffused B prefers to segregate along the intergranular and interphase boundaries, which decreases the melting point of these zones. Once their melting points become lower than the bonding temperature, these zones would melt. And then during furnace cooling from the bonding temperature, these liquid zones are resolidified resulting in the formation of liquid pools. It is reasonable to assume that each heat treatment stage acts as a bonding process with new filler. During aging process, the “new filler” melts or partially melts (the initial filler melting rage: 1366e1517 K) while B diffusion occurs. It is necessary to compare the relative rates of the dissolution and diffusion processes during bonding. Before the solid and liquid phases reach equilibrium, B concentration in the adjacent solid solution deriving from the centerline boride outward exceeds its solid solubility limit. That leads to the precipitation of boride phases in the substrate. 4.2. Microstructural evolution During the cooling from the bonding temperature, solidification of the filler in the gap would take place along the direction from the base metal towards the centerline, and then B is rejected into the remaining melt in the process. As a result, a concentration gradient of B is built up across the central region of the
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Fig. 5 SEM micrographs of the joints obtained from different heat treatment stages: (a) in the bonding zone and (b) the diffusion affected zone from the solution treated specimens, (c) in the bonding zone and (d) the diffusion affected zone from the solution þ 1st aging treated specimens, (e) in the bonding zone and (f) the diffusion affected zone from the completely heat treated specimens, (g) enlarged image of the box in the image (e).
brazed gap to the base metal, with the highest B concentration at the centerline. During the post-bond heat treatment, B diffuses from the central region of the gap outwards into the base metal driven by the concentration gradient of B.
At solution stage, as B diffuses from the gap into the base metal, the concentration of boride in the gap matrix decreases accordingly. On the other hand, because of more B diffusion into the adjacent base metal, more liquid pools are formed, and
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5. Conclusions
Fig. 6 Change of AFBZ I , WI and WDAZ of the joint along heat treatment progressing.
then WDAZ increases. Although B is expected to diffuse from the liquid pool outwards into the surrounding matrix, the continued supply of B from the gap, driven by the high chemical potential gradient generated by the high concentration gradient, causes more liquid pools and greater width of the diffusion affected zone. At aging stages, the decomposition of the intermetallics in the liquid pools leads to the reduction in the width of the diffusion affected zone, which can be seen from Fig. 5. 4.3. Reason for the hardness development Hardness profile is a good indicator of joint microstructure and can be used to assess the effect of the intermetallics on mechanical properties[18,19]. The hard brittle intermetallics provides a preferential low resistance path for crack initiation and/or propagation, which results in significant reduction in the strength of the joints[20e24]. The post-bond heat treatment is assumed to remove the detrimental intermetallics in the joint. During aging treatment, the decomposition of the centerline intermetallics leads to the decrease in the value of the hardness peak. And the precipitation of the needle/particle phases in the isothermal solidification matrix g adjacent to the centerline eutectic at aging stages is responsible for the presence of the small peak. Furthermore, because the precipitation of alloy elements weakens the effect of the solution strengthening, the hardness of the base metal decreases during heat treatment process (from 376 to 273 HV). Therefore, it is necessary to optimize bonding process of DZ468 alloy using the Co-base filler in order to suppress the formation of brittle phases.
Fig. 7 Microhardness profiles of the bonded joints in different heat treatment conditions.
(1) The microstructure of the as-cast Co-base filler consists of the matrix g, M5B3, M2B and M23B6 phases. (2) Due to the incomplete isothermal solidification, M5B3, M23B6 and MC are formed in the centerline of the joint. In addition, localized liquation of the substrate material occurs, with M3B2, MC formed resulting from resolidification in the liquid pools and grain boundaries. (3) In post-bond heat treatment process, more alloying elements of substrate enter into the gap. The amount of intermetallics in the bonding zone decreases steadily. Contrary to the bonding zone, the width of the diffusion affected zone increases at solution stage, and then, decreases at following aging stage. (4) The hardness peak is present in the eutectic zone and decreases gradually in the course of heat treatment, which is caused by the decomposition of the intermetallics. Acknowledgments This work was finally supported by the National Base Research Program of China. REFERENCES [1] Z.F. Yu, Z. Zheng, E.Z. Liu, Y.S. Yu, Y.X. Zhu, Acta Metall. Sin. 43 (2007) 653e658 (in Chinese). [2] Y. Zhang, N. Wanderka, G. Schumacher, R. Schneider, W. Neumann, Acta Mater. 48 (2000) 2787e2793. [3] C.M. Stander, Mater. Des. 17 (1996) 23e26. [4] U. Glatzel, M. Feller-Kniepmeier, Scripta Metall. 23 (1989) 1839e1844. [5] Z. Zheng, E.Z. Liu, Y.S. Yu, Y.X. Zhu, Mater. Sci. Forum 654e656 (2010) 554e557. [6] E.Z. Liu, Z. Zheng, J. Tong, L.K. Ning, X.R. Guan, Mater. Sci. Technol. 19 (2011) 110e115. [7] O.A. Ojo, N.L. Richards, M.C. Chaturvedi, Metall. Mater. Trans. 50 (2004) 641e646. [8] M.H. Haafkens, G.H. Matthey, Weld. J. 61 (1982) 25e30. [9] S.D. Duvall, W.A. Owczarski, D.F. Paulonis, Weld. J. 9 (1974) 203e214. [10] E.Z. Liu, S.C. Sun, G.F. Tu, Z. Zheng, J. Tong, X.R. Guan, Heat Treat. Met. 34 (2009) 84e86 (in Chinese). [11] G.O. Cook, C.D. Sorensen, J. Mater. Sci. 46 (2011) 5305e5323. [12] C.E. Campbell, W.J. Boettinger, Metall. Mater. Trans. A 31 (2000) 2835e2847. [13] W.F. Gale, E.R. Wallach, Metall. Trans. A 22 (1991) 2451e2457. [14] M. Mosallaee, A. Ekrami, K. Ohsasa, K. Matsuura, Metall. Mater. Trans. A 39 (2008) 2389e2402. [15] J.T. Guo, Acta Metall. Sin. 16 (1980) 30e37 (in Chinese). [16] K. Tokoro, N.P. Wikstrom, O.A. Ojo, M.C. Chaturvedi, Mater. Sci. Eng. A 477 (2008) 311e318. [17] L.Y. Sheng, W. Zhang, J.T. Guo, Z.S. Wang, Intermetallics 17 (2009) 572e577. [18] M. Pouranvari, A. Ekrami, A.H. Kokabi, Mater. Sci. Eng. A 490 (2008) 229e234. [19] Z. Mazur, A. Hernandez-Rossette, R. Garcia-Illescas, A. Luna-Ramirez, Eng. Fail. Anal. 15 (2008) 913e921. [20] L.Y. Sheng, L.J. Wang, T.F. Xi, Y.F. Zheng, H.Q. Ye, Mater. Des. 32 (2011) 4810e4817. [21] L.Y. Sheng, W. Zhang, J.T. Guo, H.Q. Ye, Mater. Charact. 60 (2009) 1311e1316. [22] C.Y. Su, C.P. Chou, W.J. Chang, M.H. Liu, J. Mater. Sci. Perform. 9 (2000) 663e668. [23] L.Y. Sheng, Y. Xie, T.F. Xi, J.T. Guo, Y.F. Zheng, H.Q. Ye, Mater. Sci. Eng. A 528 (2011) 8324e8331. [24] L.Y. Sheng, F. Yang, J.T. Guo, T.F. Xi, H.Q. Ye, Compos. Part B Eng. 45 (2013) 785e791.
ScienceDirect J. Mater. Sci. Technol., 2014, 30(5), 480e486
Microstructural Evolution of a Ni-base Alloy DZ468 Joint Bonded with a New Co-base Filler Yanhong Jing, Zhi Zheng, Enze Liu*, Yi Guo Division of Superalloy, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China [Manuscript received November 10, 2012, in revised form November 20, 2012, Available online 8 January 2014]
Ni-base alloy DZ468 has been joined by transient liquid phase bonding technique with a newly developed Co-based filler. The microstructures of the Co-base filler and the joint, the effects of heat treatment on microstructure and hardness of the joint have been investigated by various experimental methods. Results show that the Co-base filler consists of g, M2B, M5B3 and M23B6 phases. Because of the interdiffusion between the base metal and the filler, g, MC, M5B3 and M23B6 phases are formed in the bonding zone. And localized liquidation of substrate occurs in the diffusion affected zone, with MC and M3B2 precipitating in this area. During heat treatment, the volume of the intermetallic phases in the bonding zone resulting from incomplete isothermal solidification decreases obviously. On the contrary, the width of the diffusion affected zone increases at the solution stage and subsequently decreases at the aging stages. KEY WORDS: Transient liquid phase (TLP) bonding; Microstructure evolution; Ni-base alloy; Co-base filler
1. Introduction DZ468 alloy is a directionally solidified superalloy that was developed by the low segregation technique and phase calculation[1e4]. It has not only superior mechanical property, but also good hot corrosion resistance[5,6]. Those excellent characteristics make DZ468 alloy a promising candidate in turbine engine components serving in marine atmosphere. However, it is very difficult to produce the complicated intrastructures just by casting. And welding is essential to produce component with complex structure. Unfortunately, the high alloying of DZ468 alloy, especially considerable contents of g0 promoting elements, reduces the weldability of this superalloy substantially[7,8]. In 1970s, a new welding method, transient liquid phase (TLP) bonding was developed by Duvall et al.[9], which is a hybrid process that combines the beneficial features of liquid phase joining and diffusion bonding techniques. This process can be successfully used to join the heat resistant alloys which are inherently susceptible to hot cracking or post-weld heat treatment cracking free from substantial pressure. Therefore, it is so suitable for the TLP bonding technology to bond DZ468 alloy. Furthermore, it is reasonable to assume that if both the bonding Corresponding author. Assoc. Prof., Ph.D.; Tel.: þ86 24 23971143; Fax: þ86 24 23971927; E-mail address: [email protected] (E. Liu). 1005-0302/$ e see front matter Copyright Ó 2014, The editorial office of Journal of Materials Science & Technology. Published by Elsevier Limited. All rights reserved. http://dx.doi.org/10.1016/j.jmst.2013.12.010 *
process and the heat treatment method of the base metal are welldesigned to accord with each other, the cost of production can be reduced greatly. However, existing fillers cannot bond DZ468 alloy well. Considering that Co-base alloy has superior hot corrosion resistance to Ni-base alloy, a Co-base filler is developed to solve the problem. Due to the great effect of microstructure on properties, a deep understanding on the microstructure evolution of the joint is very necessary. Although lots of similar researches were conducted in this area, almost no research work has been conducted for the specific alloy and the relevant filler. So in the present work, the microstructures of both the newly developed Co-base filler and its joint were characterized. In addition, the microstructure evolution and microhardness of the joint during heat treat were investigated as well.
2. Experimental DZ468 alloy used as base alloy in this work was prepared by high rate directional solidification in ZGD2 vacuum furnace. The temperature gradient was 333e353 K/cm and the withdrawal rate was 7 mm/min. The Co-base filler powder with a size less than 45 mm was produced by ultrasonic gas atomization. A flexible insert alloy cloth consisting of the alloy powder and some organism which functions to glue the loose powders was produced by thermal rolling. Their nominal compositions (weight percent) and melting ranges are shown in Table 1. DZ468 alloy was machined into disk of 415 mm 3 mm.
Y. Jing et al.: J. Mater. Sci. Technol., 2014, 30(5), 480e486
Table 1 Nominal compositions and melting ranges of DZ468 alloy and the Co-base filler (wt%) Al Mo W Co Cr Ti Ta Re
C
B
Ni
Melting range (K)
DZ468 5.2 1 5 8.5 12 0.5 5 2 0.06 0.01 Bal. 1533e1633 Filler e e 10 Bal. 22 e e e e 2 20 1366e1517
Then they are manually grinded to 600-grit, degreased with acetone in an ultrasonic bath for 15 min and cleaned with alcohol prior to joining. The flexible cloth was inserted into two pieces of DZ468 disks. And the assembly was put in a vacuum furnace with an atmosphere of 1 103e5 103 Pa at 1553 K for 5 min, and a pressure of 1 104 Pa was applied to the top disk at this time. After bonding, the specimens were furnace cooled in the vacuum chamber to room temperature. The bonded samples were heat treated sequentially to different stages according to the optimized heat treatment schedule of the DZ468 alloy. The heat treatment schedule was: 1513 K/0.5 h þ 1533 K/0.5 h þ 1553 K/2 h þ AC (solution stage), 1393 K/4 h/FC þ 1353 K/4 h/AC (1st aging stage), 1173 K/4 h þ AC (2nd aging stage) (FC: furnace cooling, AC: air cooling)[10]. The joints and the as-cast Co-base filler were sectioned transversely, mechanically polished and then etched in a 20% phosphoric acid (H3PO4) solution at 5 V for 15 s for observation by metallographic optical microscopy (OM) and field emission scanning electron microscopy (SEM). The average width and area fraction of the intermetallics in the bonding zone, and the width of the diffusion affected zone were measured at a magnification of 500. Averages of 20 measurements across the joint sections were taken from each specimen. To identify the chemical composition of each phase, the energy dispersive spectroscopy (EDS) and electron probe microanalysis (EPMA) were employed, and at least three points were analyzed from each phase. The section of the multilayer joint and the Co-base filler were polished for the X-ray diffraction (XRD) phase identification. Thin foils for transmission electron microscopy (TEM) were sectioned from the multilayer joint and the Co-base filler. The TEM foils were prepared by ion thinning method. Microhardness across the brazed region was measured using a LM247AT hardness tester with 50 g load for 15 s.
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3. Results 3.1. Microstructure of the as-cast Co-base filler Fig. 1 shows the SEM image and the XRD pattern of the ascast Co-base filler. It can be seen that there are three phases in the matrix. The constituent of the matrix g is 50.24Coe23.77Nie 23.28Cre2.71W (at.%) determined by EDS. Table 2 shows the chemical compositions of the rest of the phases, which exhibit that the dark net-like phase is (Co, Cr)-rich boride, the bright white net-like phase is (Co, Cr, W)-rich boride and the white skeleton-like phase is Cr-rich boride, respectively. XRD pattern indicates the presence of g, M5B3, M23B6 and M2B phases. Furthermore, the TEM analysis results, as shown in Fig. 2, demonstrate the crystal structures of the phases. The (Co, Cr)rich boride has the tetragonal crystal structure and can be determined to be M2B phase with the lattice parameters of a ¼ 0.51 nm and c ¼ 0.425 nm, the Cr-rich boride has the tetragonal structure and can be determined to be M5B3 with lattice parameters of a ¼ 0.544 nm and c ¼ 1.007 nm, and lattice parameter of the g phase is 0.352 nm. Combined with XRD result, it can be determined that the (Co, Cr, W)-rich boride should be M23B6. 3.2. Microstructure of as-bond joint Fig. 3 indicates the OM micrograph and the XRD pattern of the joint bonded at 1553 K for 5 min. As can be seen from Fig. 3(a), three layers are formed, which are the bonding zone, the diffusion zone and the substrate, respectively. The bonding zone is composed of isothermally solidified zone and eutectic zone. Diffusion affected zone is the region that exists at the interface between the substrate and bonding zone. XRD pattern indicates the presence of g, g0 , M3B2, M23B6, M5B3 and MC phases. Further observation on the as-bond joint is shown in Fig. 4. Seen from Fig. 4(a), four different phases are formed from the remnant liquid layer in the bonding zone during the cooling of the joint to ambient temperature. The white skeleton-like phase and the bright white net-like phase, which are always associated with the g phase, appear in the bonding zone; the small white blocky phase exists at the edge of the eutectic zone.
Fig. 1 SEM micrograph (a) and XRD pattern (b) from the as-cast Co-base filler.
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Table 2 EPMA results of the intermetallic phases in the matrix of the as-cast Co-base filler (at.%) Co
Ni
Cr
W
B
Dark net-like phase 33.009 5.651 31.353 0.389 29.598 Bright white net-like phase 30.538 6.281 21.992 11.670 29.518 White skeleton-like phase 10.792 1.859 43.884 7.408 36.058
B is detected in the bright white net-like phase and the white skeleton-like phase, and C is detected in the small white blocky phase, although their concentrations could not be determined accurately. EDS compositional analyses of the rest of the elements suggest that the bright white net-like phase should be (Co, Cr, W)-rich boride, the white skeleton-like phase should be Crrich boride, and the small white blocky phase should be (Ta, Ti)-rich carbide (Table 3). Compared with the EDS analyses of the Co-base filler, it can be deduced that the (Co, Cr, W)-rich boride is M23B6 and Cr-rich boride is M5B3. Combined with XRD analysis, (Ta, Ti)-rich carbide can be deduced to be MC carbide. Furthermore, it can be found from Fig. 3(a) that, in the diffusion affected zone, eutectic mixtures in the liquid pool are formed along substrate grain boundaries. Two kinds of eutectic mixtures can be seen from Fig. 4(b). EDS analyses indicate that the two kinds of eutectic mixtures are (Ta, Ti)-rich carbide/g eutectic and Cr-rich boride/g eutectic. Combined with XRD
result, Cr-rich boride is identified to be M3B2 and (Ta, Ti)-rich carbide is MC. EDS was employed to further confirm the Crrich phase to be M3B2. TEM analysis result suggests that M3B2 has a tetragonal structure with lattice parameters of a ¼ 0.57 nm and c ¼ 0.3135 nm (Fig. 4(c)). 3.3. Microstructural character of the joint during heat treatment process Fig. 5 shows the microstructure of the joints obtained at different heat treatment stages. Obviously, during heat treatment, the morphology of the eutectic in the bonding zone and the diffusion affected zone changes greatly, but the eutectic mixtures still exist finally. The intermetallic keeps continuous before aging treatment, but becomes discontinuous after aging stages. In the as-bond specimen, MC particles are distributed at edge of the eutectic zone (Fig. 4). After solution treatment, the morphology of MC carbide changes greatly. It becomes the script shaped morphology from blocky, and exists in the eutectic mixtures rather than at the edge of the eutectic zone. Furthermore, it is worth noting that some needle/particle phases are visible in the matrix g adjacent to the centerline eutectic at aging stages (Fig. 5(g)), which are not observed before aging treatment. EDS analysis reveals that they are Cr-rich boride. Fig. 6 shows the change of the intermetallic area fraction in the bonding zone (AFBZ I ), the intermetallic width (WI) and the width of the diffusion affected zone (WDAZ) of the joint during
Fig. 2 Morphologies of the Cr-rich phase M5B3 (a) and the (Co, Cr)-rich phase M2B (b) from the as-cast filler. The SAED patterns of M5B3 showing [110] zone axis (c), the matrix g showing [001] zone axis (d) and the M2B showing [001] zone axis (e).
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Fig. 3 OM micrograph (a) and XRD pattern (b) of the as-bond joint. DAZ: diffusion affected zone, ISZ: isothermally solidified zone, EZ: eutectic zone.
heat treatment. It can be seen that AFBZ I reduces gradually during heat treatment. In the as-bonded specimen, AFBZ is 17.76%, in I the solution treated specimen, AFBZ is 13.53%, after subsequent I 1st aging treatment, AFBZ changes to be 10.12%, and finally I decreases to 10% after the whole heat treatment. The change of WI is well accordant with AFBZ I during heat treatment. WI at each sequential heat treatment stages is 117, 71, 43 and 41 mm, respectively. Obviously, it can be observed that the volume of the intermetallics changes greatly at solution and 1st aging stage and then almost keeps constant at 2nd aging stage. Furthermore, it can be observed that WDAZ increases at solution treatment stage,
then decreases at subsequent aging stages. Solution treatment results in the increase in the width of diffusion affected zone from 329 mm in as-bond specimen to 655 mm, and then the decrease to 598 mm at 1st aging stage, and finally to 538 mm after the whole heat treatment. 3.4. Microhardness profiles in the bonded joint Microhardness profiles as a function of distance at the different stages of heat treatment are shown in Fig. 7. It can be observed that, in all conditions, the microhardness peak is
Fig. 4 Images from the joint bonded at 1553 K for 5 min: (a) SEM micrograph of the bonding zone, (b) SEM micrograph of the diffusion affected zone, (c) morphology of Cr-rich boride M3B2 in the liquid pool, the inset is the SAED pattern of the phase showing [001] zone axis, (d) micrograph of the liquid pool in the bonding zone.
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Table 3 EDS results of the phases in the as-bond joint and as-cast filler (at.%) White skeletonlike phase (joint)
M5B3 (filler)
Bright white netlike phase (joint)
M23B6 (filler)
Matrix (joint)
g (filler)
White blocky phase (joint)
e e 63.76 13.76 8.73 e 0.68 13.08
e e 70.63 13.81 2.05 e e 13.51
e 1.22 21.27 32.98 14.33 8.24 3.61 18.35
e e 29.67 41.11 8.71 e e 20.67
1.62 1.58 15.95 38.95 38.28 2.64 e 0.98
e e 23.28 50.24 23.77 e e 2.71
e 25.77 5.07 e 4.03 59.32 2.24 1.75
Al Ti Cr Co Ni Ta Mo W
present in the eutectic zone but different at each stage. In as-bond specimen, the hardness peak is 564 HV, and then increases to 737 HV after solution treatment. During subsequent 1st aging process, the hardness peak decreases to 515 HV and finally to 495 HV after the whole heat treatment. It is noted that, after aging heat treatment, a small peak appears in the matrix g adjacent to centerline eutectic. Furthermore, it can be observed that, at each heat treatment stage, the hardness in the diffusion affected zone is similar to that in the substrate, but decreases during heat treatment. The microhardness of substrate in each sequential heat treatment stage is 376, 360, 354 and 275 HV, respectively. 4. Discussion 4.1. Mechanism for the formation of new phases According to the traditional TLP model, when the assembly is held at bonding temperature, the insert filler melts and rapidly attains equilibrium with the solid base metal through the process of melt-back dissolution of the substrate. Subsequently, interdiffusion of alloying elements between the base metal and the liquid commences and causes the melting point of the filler at the solid/liquid interface to increase resulting in isothermal solidification of the liquid. Since the melting point depressing solute B diffuses continuously from the liquid into the base metal, the volume of the liquid that can be maintained at equilibrium decreases, causing solidification to proceed towards the center of the joint from the mating solid surface[11,12]. In the present experiment, due to the lack of bonding time for completing the isothermal solidification of the liquid filler, eutectic microconstituents are formed along the joint centerline. During the bonding process, the dissolution of base metal makes the alloying elements of base metal like Ta, Ti, and C enter into the liquid filler. In addition to dissolution, alloy elements from the base alloy can diffuse into the liquid, which is driven by concentration gradient. Although C is not added in the filler, the decomposition of primary carbide and C atoms in base metal act as the source to supply the free C atoms during the exposure at high temperature. Therefore, in the following solidification process, stable MC-type carbides are formed in the region. However, it is more realistic to assume that both solid/liquid equilibration and solid state diffusion processes operate simultaneously rather than sequentially. The formation of liquid pools and the needle/particle phases in the isothermal solidification matrix g adjacent to the centerline eutectic supports this
assumption. Similar phenomenon was observed in joining nickel using ternary NieSieB insert metals by Gale et al.[13]. When the bonding temperature is below the binary NieB eutectic temperature, the precipitation of Ni3B borides occurs in the substrate at the edge of the insert. In contrast, when the bonding temperature is above the binary NieB eutectic temperature, localized liquation of the substrate material takes place[13]. The formation of liquid pools was also observed during bonding IN738LC alloy using BNi-3 filler at 1473 K[14]. Guo et al.[15] investigated the effect of B on microstructure, and found that M3B2/g eutectic was inclined to form in grain boundaries once the concentration of B exceeded 0.004 at.%. It is because B atom is prone to segregate into grain boundaries, afterwards, eutectic reaction usually occurs at high temperature and produces M3B2/g eutectic. The crown morphology of the liquid pools, as shown in Fig. 4(d), along the solidification direction implies that they are formed from resolidification. It is known that B has a high diffusion rate at high temperature[16,17]. And the maximum solid solubility of B in solid Ni according NieB binary phase diagram is taken to be 0.3 at.% over the range of 1338e1373 K. Therefore, when the assemble coupon is held at 1553 K, the B element diffuses from the gap to the substrate driven by concentration gradient (the initial concentration of B in filler is about 2 wt% (w10 at.%) and DZ468 only just contains 0.01 wt% (w0.06 at.%)). The diffused B prefers to segregate along the intergranular and interphase boundaries, which decreases the melting point of these zones. Once their melting points become lower than the bonding temperature, these zones would melt. And then during furnace cooling from the bonding temperature, these liquid zones are resolidified resulting in the formation of liquid pools. It is reasonable to assume that each heat treatment stage acts as a bonding process with new filler. During aging process, the “new filler” melts or partially melts (the initial filler melting rage: 1366e1517 K) while B diffusion occurs. It is necessary to compare the relative rates of the dissolution and diffusion processes during bonding. Before the solid and liquid phases reach equilibrium, B concentration in the adjacent solid solution deriving from the centerline boride outward exceeds its solid solubility limit. That leads to the precipitation of boride phases in the substrate. 4.2. Microstructural evolution During the cooling from the bonding temperature, solidification of the filler in the gap would take place along the direction from the base metal towards the centerline, and then B is rejected into the remaining melt in the process. As a result, a concentration gradient of B is built up across the central region of the
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Fig. 5 SEM micrographs of the joints obtained from different heat treatment stages: (a) in the bonding zone and (b) the diffusion affected zone from the solution treated specimens, (c) in the bonding zone and (d) the diffusion affected zone from the solution þ 1st aging treated specimens, (e) in the bonding zone and (f) the diffusion affected zone from the completely heat treated specimens, (g) enlarged image of the box in the image (e).
brazed gap to the base metal, with the highest B concentration at the centerline. During the post-bond heat treatment, B diffuses from the central region of the gap outwards into the base metal driven by the concentration gradient of B.
At solution stage, as B diffuses from the gap into the base metal, the concentration of boride in the gap matrix decreases accordingly. On the other hand, because of more B diffusion into the adjacent base metal, more liquid pools are formed, and
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5. Conclusions
Fig. 6 Change of AFBZ I , WI and WDAZ of the joint along heat treatment progressing.
then WDAZ increases. Although B is expected to diffuse from the liquid pool outwards into the surrounding matrix, the continued supply of B from the gap, driven by the high chemical potential gradient generated by the high concentration gradient, causes more liquid pools and greater width of the diffusion affected zone. At aging stages, the decomposition of the intermetallics in the liquid pools leads to the reduction in the width of the diffusion affected zone, which can be seen from Fig. 5. 4.3. Reason for the hardness development Hardness profile is a good indicator of joint microstructure and can be used to assess the effect of the intermetallics on mechanical properties[18,19]. The hard brittle intermetallics provides a preferential low resistance path for crack initiation and/or propagation, which results in significant reduction in the strength of the joints[20e24]. The post-bond heat treatment is assumed to remove the detrimental intermetallics in the joint. During aging treatment, the decomposition of the centerline intermetallics leads to the decrease in the value of the hardness peak. And the precipitation of the needle/particle phases in the isothermal solidification matrix g adjacent to the centerline eutectic at aging stages is responsible for the presence of the small peak. Furthermore, because the precipitation of alloy elements weakens the effect of the solution strengthening, the hardness of the base metal decreases during heat treatment process (from 376 to 273 HV). Therefore, it is necessary to optimize bonding process of DZ468 alloy using the Co-base filler in order to suppress the formation of brittle phases.
Fig. 7 Microhardness profiles of the bonded joints in different heat treatment conditions.
(1) The microstructure of the as-cast Co-base filler consists of the matrix g, M5B3, M2B and M23B6 phases. (2) Due to the incomplete isothermal solidification, M5B3, M23B6 and MC are formed in the centerline of the joint. In addition, localized liquation of the substrate material occurs, with M3B2, MC formed resulting from resolidification in the liquid pools and grain boundaries. (3) In post-bond heat treatment process, more alloying elements of substrate enter into the gap. The amount of intermetallics in the bonding zone decreases steadily. Contrary to the bonding zone, the width of the diffusion affected zone increases at solution stage, and then, decreases at following aging stage. (4) The hardness peak is present in the eutectic zone and decreases gradually in the course of heat treatment, which is caused by the decomposition of the intermetallics. Acknowledgments This work was finally supported by the National Base Research Program of China. REFERENCES [1] Z.F. Yu, Z. Zheng, E.Z. Liu, Y.S. Yu, Y.X. Zhu, Acta Metall. Sin. 43 (2007) 653e658 (in Chinese). [2] Y. Zhang, N. Wanderka, G. Schumacher, R. Schneider, W. Neumann, Acta Mater. 48 (2000) 2787e2793. [3] C.M. Stander, Mater. Des. 17 (1996) 23e26. [4] U. Glatzel, M. Feller-Kniepmeier, Scripta Metall. 23 (1989) 1839e1844. [5] Z. Zheng, E.Z. Liu, Y.S. Yu, Y.X. Zhu, Mater. Sci. Forum 654e656 (2010) 554e557. [6] E.Z. Liu, Z. Zheng, J. Tong, L.K. Ning, X.R. Guan, Mater. Sci. Technol. 19 (2011) 110e115. [7] O.A. Ojo, N.L. Richards, M.C. Chaturvedi, Metall. Mater. Trans. 50 (2004) 641e646. [8] M.H. Haafkens, G.H. Matthey, Weld. J. 61 (1982) 25e30. [9] S.D. Duvall, W.A. Owczarski, D.F. Paulonis, Weld. J. 9 (1974) 203e214. [10] E.Z. Liu, S.C. Sun, G.F. Tu, Z. Zheng, J. Tong, X.R. Guan, Heat Treat. Met. 34 (2009) 84e86 (in Chinese). [11] G.O. Cook, C.D. Sorensen, J. Mater. Sci. 46 (2011) 5305e5323. [12] C.E. Campbell, W.J. Boettinger, Metall. Mater. Trans. A 31 (2000) 2835e2847. [13] W.F. Gale, E.R. Wallach, Metall. Trans. A 22 (1991) 2451e2457. [14] M. Mosallaee, A. Ekrami, K. Ohsasa, K. Matsuura, Metall. Mater. Trans. A 39 (2008) 2389e2402. [15] J.T. Guo, Acta Metall. Sin. 16 (1980) 30e37 (in Chinese). [16] K. Tokoro, N.P. Wikstrom, O.A. Ojo, M.C. Chaturvedi, Mater. Sci. Eng. A 477 (2008) 311e318. [17] L.Y. Sheng, W. Zhang, J.T. Guo, Z.S. Wang, Intermetallics 17 (2009) 572e577. [18] M. Pouranvari, A. Ekrami, A.H. Kokabi, Mater. Sci. Eng. A 490 (2008) 229e234. [19] Z. Mazur, A. Hernandez-Rossette, R. Garcia-Illescas, A. Luna-Ramirez, Eng. Fail. Anal. 15 (2008) 913e921. [20] L.Y. Sheng, L.J. Wang, T.F. Xi, Y.F. Zheng, H.Q. Ye, Mater. Des. 32 (2011) 4810e4817. [21] L.Y. Sheng, W. Zhang, J.T. Guo, H.Q. Ye, Mater. Charact. 60 (2009) 1311e1316. [22] C.Y. Su, C.P. Chou, W.J. Chang, M.H. Liu, J. Mater. Sci. Perform. 9 (2000) 663e668. [23] L.Y. Sheng, Y. Xie, T.F. Xi, J.T. Guo, Y.F. Zheng, H.Q. Ye, Mater. Sci. Eng. A 528 (2011) 8324e8331. [24] L.Y. Sheng, F. Yang, J.T. Guo, T.F. Xi, H.Q. Ye, Compos. Part B Eng. 45 (2013) 785e791.