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Brazing TiAl intermetallics using TiNi–V eutectic brazing alloy
Materials Science and Engineering A 551 (2012) 133–139
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Brazing TiAl intermetallics using TiNi–V eutectic brazing alloy X.G. Song a,b , J. Cao a,∗ , H.Y. Chen a , Y.F. Wang a , J.C. Feng a a b
State Key Lab of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, China School of Materials Science and Engineering, Harbin Institute of Technology at WeiHai, WeiHai 264209, China
a r t i c l e
i n f o
Article history: Received 3 September 2011 Received in revised form 14 January 2012 Accepted 1 May 2012 Available online 11 May 2012 Keywords: Intermetallics Brazing TiNi–V eutectic brazing alloy Interfacial microstructure Mechanical properties
a b s t r a c t A novel TiNi–V25 (at.%) eutectic brazing alloy, which was composed of TiNi intermetallics and V, was designed and fabricated. Subsequently, the brazing alloy was used to braze TiAl intermetallics (Ti–42.5Al–9V–0.3Y (at.%)) and reliable joints were obtained. The typical interfacial microstructure of TiAl brazed joints was characterized by employing SEM, EDS and XRD, and the brazing mechanism was analyzed in detail. The brazed joints mainly consisted of B2 phase and 3 -Al3 NiTi2 intermetallics. Interfacial morphology of TiAl joints varied with the increase of brazing temperature, while the phase constitution in joints did not change sound TiAl joints with the highest average shear strength of 196 MPa were obtained when the specimens were brazed at 1220 ◦ C for 10 min. The presence of B2 phase and 3 Al3 NiTi2 intermetallics in joints deteriorated the joining properties due to their brittleness, which led to cleavage-dominated fracture after shear test. The crack propagation path and fracture location were mainly determined by the content and distribution of 3 -Al3 NiTi2 phase. © 2012 Elsevier B.V. All rights reserved.
1. Introduction TiAl intermetallics have been considered as the most promising alternative materials to replace traditional heat-resistant steels and superalloys in automobile industry for significant weight saving [1,2]. The excellent properties such as high specific strength, excellent creep strength and good oxidation resistance make them particularly appropriate for next-generation aircraft turbine engines [3–5]. During the last 30 years, an enormous research activity has been devoted to developing TiAl intermetallics all around the world [6–8]. However, extensive applications have not been achieved due to the materials’ intrinsic brittleness and the difficulties in manufacturing as well as high cost [9]. It is well known that joining techniques will play an important role in the applications of TiAl alloys. Considerable investigations on joining TiAl alloys have been carried out along with their development. Conventional fusion welding is unsuitable although Smarsly et al. [10] reported that sound TiAl joints could be produced with careful process controlling involving preheating at 700 ◦ C and pulsing slow-cooling after welding. Whereas, the solidification cracks caused by the limited ductility and fracture toughness of TiAl alloys cannot be easily avoided during the cooling process. Solid state bonding has also been used to join TiAl intermetallics and high quality joints can be obtained [11–13]. However, a long-time exposure
∗ Corresponding author. Tel.: +86 451 86418882; fax: +86 451 86418146. E-mail address: cao [email protected] (J. Cao). 0921-5093/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.05.002
of TiAl substrate at elevated temperature in bonding process can change its microstructure and thus degrade its mechanical properties. Furthermore, the use of these methods will be limited when the shape of components is complicated. Brazing, as the most feasible and economical method, has received special attentions in joining TiAl alloys to themselves or to other materials. BAlSi-4 brazing alloy was used by Wu et al. [14] to join TiAl alloy and the corresponding relationship between interfacial microstructure and joints properties was analyzed in detail. Infrared brazing of TiAl alloy using pure Ag and BAg-8 brazing alloys was investigated by Shiue et al. [15,16]. The interfacial microstructural evolution, joints properties and reaction kinetics were comprehensively evaluated. However, the joints brazed using Al, Ag or Au based brazing alloys have poor thermal resistance and are usually used below 500 ◦ C. As a result, the high-temperature properties of TiAl alloys could not be taken full advantage of. Therefore, the high temperature brazing alloys (e.g. Ti and Ni based brazing alloys) showed their advantage for high temperature applications of brazed joints. Ti–Ni–Cu brazing alloys were used by Wallis et al. [17] and Wu et al. [18] to braze TiAl alloys and high-quality brazed joints with good room-temperature and high-temperature properties were obtained. It is well known that the development of high temperature brazing alloys is very important for the practical applications of TiAl alloys. Therefore, in the present study, a novel high temperature TiNi–V eutectic brazing alloy was designed and fabricated to braze TiAl alloy. Furthermore, the typical interfacial microstructure was identified, and the effects of brazing temperature on interfacial microstructure and shear strength of TiAl brazed joints were investigated.
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Table 1 Properties of the joining materials. Materials
0.2 (MPa)
b (MPa)
ı (%)
TiAl TiNi–V
420–462 510–540
520–586 722–766
0.5–0.65 3.4–5.6
2. Experimental In this study, the TiAl alloy with nominal composition of Ti–42.5Al–9V–0.3Y (at.%) was fabricated by induction skull melting using high purity (99.99 wt.%) Ti, Al, Al–V and Al–Y interalloys. Subsequently, the ingot was treated at 900 ◦ C/48 h in air and hot isostatic-pressed at 1250 ◦ C/4 h/175 MPa in argon atmosphere to eliminate segregation and shrinkage porosity respectively. The TiNi–V brazing alloy with its nominal composition of Ti37.5Ni37.5V25 (at.%) was prepared by vacuum arc remelting using high purity (99.99 wt.%) Ti, Ni and V rods [19]. All the raw materials were cleaned by 1.5HF–2.5HNO3 –5HCl–91H2 O (ml) and saturated in NaOH solution (cNaOH = 2 mol/l) before vacuum arc remelting. Alloy ingots were turned and remelted at least six times to be homogenized; the final weight loss of the master alloy was less than 0.1 wt.%. A diffusion-annealing treatment was done at 1000 ◦ C for 4 h after vacuum arc remelting. The main properties of the TiAl alloy and TiNi–V brazing alloy are listed in Table 1. The TiNi–V alloy was cut into foils with a thickness of 400 m by spark cutting. Both surfaces of TiNi–V foil were ground on SiC grit papers until its thickness was about 200 m. The TiAl base material was also cut into blocks with two different sizes (5 mm × 5 mm × 2.5 mm and 20 mm × 10 mm × 2.5 mm). Prior to brazing, the brazing surfaces of TiAl specimens were ground on SiC grit papers and polished using diamond pastes. All samples were cleaned ultrasonically in acetone for 10 min and dried by air blowing. The TiNi–V foil was placed between two TiAl blocks. At the beginning of brazing process, the furnace was heated to 1100 ◦ C at a rate of 30 ◦ C/min, then to the brazing temperatures at a rate of 10 ◦ C/min. Subsequently, the furnace temperature were held for 10 min and then cooled down to 200 ◦ C at a rate of 5 ◦ C/min, and finally cooled down spontaneously to room temperature. During the brazing process, the vacuum was kept at (1.3–2.0) × 10−3 Pa and a pressure of 20 kPa was applied to ensure a proper contact.
The specimens for metallographic observation were crosssectioned, perpendicular to the brazed interface. The interfacial microstructure was characterized employing scanning electron microscopy (SEM). Componential analysis of various phases in the joints was carried out using an energy dispersive spectrometer (EDS) with the operation voltage of 15 kV and minimum spot size of 1 m. The shear tests were performed at a constant speed of 0.5 mm/min by a universal testing machine (Instron1186) at room temperature. A schematic of the shear test could be seen in Ref. [20]. For each set of experimental data, at least five samples were used to average the joints strength. After shear test, the fracture of brazed joints was inspected by SEM. In order to identify the interfacial phases accurately, the analysis was performed using an X-ray diffraction (XRD, JDX-3530M) spectrometer equipped with Cu K␣ radiation on the fracture surfaces. In addition, differential thermal analysis (DTA) was performed in order to determine the melting point of TiNi–V brazing alloy. 3. Results and discussion 3.1. Microstructure of TiAl alloy and TiNi–V brazing alloy Fig. 1(a) shows the microstructure of TiAl substrate in backscattered electron (BSE) mode. It is seen that the alloy was homogenous, fine-grained nearly lamellar structure, and the lamellar colony size is about 40–80 m. Fig. 1(b) shows the XRD patterns measured on the surface of the alloy. The result indicates that the alloy was primarily composed of ␥-TiAl phase, ␣2 -Ti3 Al phase and B2 phase. Each phase could be identified according to the EDS results taken from the selected spots in Fig. 1(a). It is shown that V concentration in B2 phase (marked by B) was much higher than that in other phases. Relatively high V concentration in B2 phase is consistent with the viewpoint that element V could stabilize  phase in TiAl alloy [21]. The resultant microstructure was characterized by the near lamellar colonies surrounded by B2 + ␥ structure. The white particles (marked by C) could be determined as YAl2 intermetallic compound according to literature [22]. However, diffraction peaks of YAl2 phase cannot be found in the XRD patterns, for the reason that the phase was too sparse to be detected by XRD. It is noted that these YAl2 particles were mainly distributed in ␥-TiAl phase. The solidification and transformation pathway has been reported as the follows: L → L +  →  → ␣ +  → ␣ +  + ␥ → Lamellar +B2 + ␥ [23].
Fig. 1. Backscattered electron image (BEI), XRD patterns and EDS analysis results in atomic percent of TiAl alloy. (a) BEI of TiAl alloy, (b) XRD patterns of TiAl alloy.
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Fig. 2. (a) BEI, (b) XRD patterns and (c) DTA result of TiNi–V brazing alloy.
The typical microstructure of TiNi–V brazing alloy is shown in Fig. 2(a). It is clearly seen that eutectic structure was obtained after vacuum arc remelting. In order to further investigate its phase constitution, X-ray diffraction analysis was carried out. The XRD patterns in Fig. 2(b) indicate that the alloy consisted of TiNi intermetallics and V. According to the DTA result shown in Fig. 2(c), it can be indicated that the melting point of TiNi–V brazing alloy is about 1142 ◦ C. 3.2. Interfacial microstructure of TiAl/TiNi–V/TiAl brazed joints Fig. 3 shows the typical interfacial microstructure of TiAl/TiNi–V/TiAl joint brazed at 1220 ◦ C for 10 min. It clearly demonstrates that a sound brazing seam was obtained, and its thickness is about 400 m. An increasing in thickness indicates that intensive interactions occurred between TiAl substrate and molten TiNi–V brazing alloy during the brazing process. For the sake of convenience, the brazing seam of TiAl joint is artificially divided into three zones: zone I, II and III respectively, as shown in Fig. 3. The distribution of elements was measured by EDS along the white base line. A relatively high concentration of Al in brazing seam indicates that both the amount of dissolved Al and the diffusion rate of Al in molten brazing alloy were high. The concentration profiles of Ti and V show that their distributions were heterogeneous, which is probably due to their low diffusion rates in molten brazing alloy. In order to further investigate the interfacial microstructure, more details are presented in a larger visual field in BSE mode, as shown in Fig. 4, and the EDS results taken from different phases are also included. Fig. 4(a) shows the microstructure of TiAl substrate after brazing. Compared with the original microstructure shown in Fig. 1, it is observed that there was no obvious change except that the lamellar structure leveled more smoothly. According to the phase diagram drawn by Takeyama et al. [24], during the heating process in brazing, the solid-state phase transformation
Fig. 3. BEI of TiAl joint brazed using TiNi–V brazing alloy at 1220 ◦ C for 10 min and the elemental distributions of Al, Ti, V and Ni along the white base line.
of Lamellar + B2 + ␥ → ␣ +  + ␥ occurred in TiAl substrate, as well as the B2 →  disordered transformation which occurred above 1100 ◦ C. However, in the present study, the transformations of TiAl substrate were incomplete due to a short holding time at high temperature. Therefore, during cooling process, the fresh Lamellar (␥ + ␣2 ) and B2 phase could precipitate based on the remnant Lamellar and B2 phase respectively, which resulted in a similar microstructure to that of original TiAl alloy. In addition, the lower cooling rate during brazing is beneficial to the transformation of ␣ +  + ␥ → Lamellar + B2 + ␥, and thus a more smooth lamellar structure was obtained finally. The microstructure of zone I is shown in Fig. 4(b). The gray phase (marked by D) had the average composition of 42.44Ti, 38.72Al, 16.27V, 2.13Ni and 0.44Y (at.%), which could be identified as B2 phase. It is evident that the amount of B2 phase increased in this zone. The formation process of B2 phase can be described as follows: TiNi–V brazing alloy began to melt when the brazing temperature exceeded its melting point. Subsequently, the dissolution of TiAl to molten brazing alloy and the diffusion of molten brazing alloy to TiAl substrate occurred simultaneously, which led to an increase of V content in the solid–liquid interfacial area. Thus, the amount of B2 phase in the joints increased after brazing because element V could stabilize  phase in TiAl alloy [20]. In addition, the massive ␥ phase in zone I (marked by C) was not dissolved completely, and it was alloyed with small amounts of Ni (1.04 at.%) due to the interdiffusion between the brazing alloy and TiAl substrate. Furthermore, a new phase (marked by E) came forth, which has an average composition of 31.53Ti, 47.57Al, 2.52V, 17.58Ni and 0.8Y (at.%). According to Rogl’s research [25], the phase could be identified as 3 -Al3 NiTi2 with the MgZn2 -type structure. Fig. 4(c) shows the typical microstructure in zone II. It can be clearly seen that the zone mainly consisted of B2 phase (marked by G) and 3 -Al3 NiTi2 phase (marked by F). It is important to note that the B2 phase colonies were surrounded by 3 -Al3 NiTi2 intermetallics, and the white YAl2 particles were mainly distributed in 3 -Al3 NiTi2 phase. The microstructure in zone II is very helpful to deducing the reaction mechanisms. Chen et al. [22,26] reported that element Y promoted ␥ phase nucleation, which resulted in more Y in ␥ phase. The same results were also found in this study, as above analysis of the microstructure of TiAl (shown in Fig. 1). The redistribution of YAl2 from ␥ phrase to 3 -Al3 NiTi2 phase after brazing indicates that most of the massive ␥ phase at grain boundaries dissolved into the molten brazing alloy during the brazing process. Thus, the lamellar colony was gradually surrounded by the molten brazing alloy which dissolved Al, and Ti. The diffusion of V into lamellar colony or ␥ single phase led to the transformation from lamellar colony (␥ + ␣2 ) or ␥ single phase to B2 phase [27]. In addition, B2 phase could also form by the solidification of molten brazing alloy. Due to the slow cooling rate, the solidification could be considered as an equilibrium or quasi-equilibrium process, in which impurities (e.g. YAl2 ) were driven from the crystallization surfaces, finally located in 3 -Al3 NiTi2 phase, as shown in Fig. 4(c).
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Fig. 4. High magnification BEIs of each zone marked in Fig. 3. (a) TiAl substrate, (b) zone I, (c) zone II and (d) zone III.
The microstructure in zone III is shown in Fig. 4(d). It can be seen that the microstructure of this zone was mainly composed of two phases. The location marked by H with an average composition of 32.05Ti, 33.18Al, 31.66V, 2.74Ni and 0.37Y (at.%) could be determined as B2 phase because of the high V content. In addition, the other location marked by I had a similar chemical composition with locations E and F, thus, the phase could also be identified as 3 Al3 NiTi2 intermetallics. However, compared with the 3 -Al3 NiTi2 phase in zone II (shown in Fig. 4(c)), which was located at the boundaries of massive B2 phase, the distribution of 3 -Al3 NiTi2 phase was changed greatly, as shown in Fig. 4(d). According to above analysis, it can be concluded that the brazed joints mainly consisted of B2 phase and 3 -Al3 NiTi2 intermetallics. During the brazing process, TiAl substrate was partially dissolved into the molten brazing alloy. Thus, the contents of Al and Ti in molten brazing alloy increased, which led to the formation of ternary intermetallic compound 3 -Al3 NiTi2 . The diffusion of V to TiAl substrate also played an important role, which resulted in the transformation of lamellar colony (␥ + ␣2 ) to B2 phase. During the cooling process, B2 phase directly solidified from the molten brazing alloy, while element Ni was enriched in the remnant molten brazing alloy and finally reacted with Ti and Al to form 3 -Al3 NiTi2 phase at the boundaries of massive B2 phase. Further cooling, the 3 -Al3 NiTi2 phase precipitated in the B2 phase at the center of brazing seam, probably because solubility of element Ni in B2 phase decreased as did brazing temperature.
3.3. Effects of brazing temperature on the interfacial microstructure and shear strength of the brazed joints The interfacial microstructure of TiAl joints brazed at different temperatures is shown in Fig. 5. The interfacial morphology changed obviously with the increase of brazing temperature. In other words, brazing temperature had a great influence on the interfacial microstructure. When TiAl alloy was brazed at lower temperatures (e.g. 1180 ◦ C and 1200 ◦ C), a continuous layer formed between TiAl substrate and brazing seam, as shown in Fig. 5(a) and (b). A small amount of remnant brazing alloy was found in the central region of the joint brazed at 1180 ◦ C, which was caused by poor melting at lower brazing temperatures. Another style of interfacial microstructure was obtained when the brazing temperature exceeded 1220 ◦ C (e.g. 1240 ◦ C and 1260 ◦ C), as shown in Fig. 5(c) and (d). Continuous reaction layer disappeared and the microstructure was similar to that observed in Fig. 3. Fig. 5(e)–(g) shows the high magnification BEIs of the rectangular regions in Fig. 5(a)–(c) respectively, where the EDS chemical analysis results taken at the selected locations are also included. The interfacial microstructure between TiAl and brazing seam obtained at lower temperatures can be seen in Fig. 5(e) and (f). According to the BEIs and EDS results, the continuous layer close to TiAl was determined as B2 phase, as marked by A and D in Fig. 5(e) and (f) respectively. The massive phase marked by B and F was also B2 phase due to the similar composition to that of the continuous layer. The locations marked by C and E which had relatively high
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Fig. 5. Interfacial microstructure of TiAl joints brazed at different temperatures. (a) 1180 ◦ C, (b) 1200 ◦ C, (c) 1240 ◦ C, (d) 1260 ◦ C, (e) high magnification BEIs of the rectangular regions in (a), (f) high magnification BEIs of the rectangular regions in (b), and (g) high magnification BEIs of the rectangular regions in (c).
Ni content could be inferred as 3 -Al3 NiTi2 . Based on the analysis above, it can be concluded that the interfacial morphologies of TiAl joints brazed at different temperatures were different from each other. However, the interfacial phase constitution was insensitive to brazing temperature. Fig. 5(g) shows the high magnification BEI of the rectangular region in Fig. 5(c). A colony of B2 + ␥ surrounded by 3 -Al3 NiTi2 strips was observed, and its size was about
50–60 m. It can be inferred that the colony was transformed from original lamellar colony. The typical microstructure was very similar to that shown in Fig. 4(c), and it confirmed that the transformations stated above were reasonable. Fig. 6 illustrates the average shear strength of TiAl joints brazed at different temperatures. The specimen brazed at 1220 ◦ C for 10 min demonstrated the highest average shear strength of
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Fig. 6. Effect of brazing temperature on joining shear strength.
Fig. 8. XRD patterns of the fracture surfaces after shear test for specimen brazed at 1220 ◦ C for 10 min.
196 MPa. In addition, the shear strength decreased slightly with increasing brazing temperature. In order to investigate the relationship between joints properties and interfacial microstructure, fracture analysis was performed by SEM, EDS and XRD. The typical fractographs are shown in Fig. 7(a) and (b). Cleavage facets were widely observed in the fracture surfaces. XRD result of the fracture surfaces is illustrated in Fig. 8. Combining the EDS results of selected locations and X-ray analysis, it is concluded that the 3 -Al3 NiTi2 phase and B2 phase dominated in the facture surfaces, as well as some lamellar colonies, as shown in Fig. 7(b). Actually, the presence of B2 phase and 3 Al3 NiTi2 phase deteriorated the joints properties due to their nature
of brittleness. Therefore, cleavage-dominated fracture formed after shear test. In addition, the XRD result confirms that the above analysis about interfacial microstructure and phases is accurate. Fig. 7(c)–(e) show cross-section BEIs of brazed joints after shear test. It is obvious that the joints always crack in the brazed seam region while the fracture location varied with the brazing temperature. When brazed at low temperatures (e.g. 1180 ◦ C), the crack initiated and propagated in a continuous reaction layer during shear test, as shown in Fig. 7(c). According to the analysis of interfacial microstructure shown in Fig. 5(a), it is found that the reaction layer was 3 -Al3 NiTi2 phase. When brazing temperature was increased to 1200 ◦ C, fracture location changed from
Fig. 7. Fractographs of TiAl joint brazed at 1220 ◦ C for 10 min and the EDS analysis results of selected locations. (a) Low magnification SEM image, (b) high magnification SEM image, (c) cross-section BEI of TiAl joint after shear test brazed at 1180 ◦ C, (d) cross-section BEI of TiAl joint after shear test brazed at 1200 ◦ C and (e) at 1220 ◦ C.
X.G. Song et al. / Materials Science and Engineering A 551 (2012) 133–139
3 -Al3 NiTi2 layer to the brazing seam because the continuous 3 Al3 NiTi2 layer disappeared, as shown in Fig. 5(b). Fig. 7(e) shows the cross-section BEI of TiAl joint brazed at 1220 ◦ C, which exhibited the highest shear strength. It can be seen that the crack propagated along the interface (zone II) between TiAl substrate and brazing seam. By comparing the interfacial microstructure and fracture locations in different joints, it is concluded that the crack preferred to initiate and propagate at the locations where the 3 -Al3 NiTi2 content was relatively high. The result also can be confirmed in Fig. 7(b) where the fracture surfaces contained lots of 3 -Al3 NiTi2 phases. So it can be inferred that 3 -Al3 NiTi2 phase had poor properties and its content and distribution determined the crack propagation path and fracture location, as well as joining properties, which is consistent with our previous results from the investigation of brazing TiAl alloy using TiNi–Nb eutectic brazing alloy [28]. 4. Conclusions Interfacial microstructure and joining properties of TiAl joints brazed using TiNi–V eutectic brazing alloy were investigated in this study. Primary conclusions are summarized as follows. (1) A TiNi–V eutectic brazing alloy was designed and developed by vacuum arc remelting. The alloy had typical eutectic microstructure which mainly consisted of TiNi intermetallic compound and V. The melting point of this filler metal was about 1142 ◦ C. (2) Brazing of TiAl intermetallics using TiNi–V brazing alloy was also achieved. The joint was mainly composed of B2 phase and 3 -Al3 NiTi2 intermetallics. During the brazing process, intensive reactions occurred between TiAl substrate and molten brazing alloy. The diffusion of V to TiAl led to two transformations as following: Lamellar (␥ + ␣2 ) + V → B2 and ␥ + V → B2. (3) The interfacial morphologies of TiAl joints brazed at different temperatures were different from each other. However, all the TiAl joints had the same phase constitution, i.e. B2 phase and 3 -Al3 NiTi2 phase, which was detrimental to the joining properties. The highest average shear strength of 196 MPa was obtained when the brazing condition was 1220 ◦ C/10 min. Fracture analysis showed that cleavage-dominated fracture formed during shear test and the crack always initiated and propagated
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at the locations where the 3 -Al3 NiTi2 content was relatively high. Acknowledgments The authors gratefully acknowledge the financial support from Project 2011CB605505 supported by the National Basic Research Program of China. The authors also acknowledge the financial support from the National Natural Science Foundation of China under Grant No. 50805038 and the Fundamental Research Funds for the Central Universities under Grant No. HIT.NSRIF.2011036. References [1] K. Maruyama, M. Yamaguchi, G. Suzuki, H.L. Zhu, H.Y. Kim, M.H. Yoo, Acta Mater. 52 (2004) 5158. [2] E.A. Loria, Intermetallics 8 (2000) 1339. [3] S.L. Draper, D. Krause, B. Lerch, I.E. Locci, B. Doehnert, R. Nigam, G. Das, P. Sickles, B. Tabernig, N. Reger, K. Rissbacher, Mater. Sci. Eng. A 464 (2007) 330. [4] T. Tetsui, Mater. Sci. Eng. A 329–331 (2002) 582–588. [5] S.L. Draper, D. Krause, B. Lerch, Mater. Sci. Eng. A 464 (2007) 330. [6] X. Wu, Intermetallics 14 (2006) 1114–1122. [7] G.L. Chen, X.J. Xu, Z.K. Teng, Y.L. Wang, J.P. Lin, Intermetallics 15 (2007) 625. [8] G.L. Chen, L.C. Zhang, Mater. Sci. Eng. A 329–331 (2002) 163. [9] A. Lasalmonie, Intermetallics 14 (2006) 1123. [10] W. Smarsly, H. Cramer, A.W.E. Nentwig, Presented at Werkstoffwoche 98, Munich, 12–15 October, 1998 (in German). [11] G. Cam, H. Clemens, R. Gerling, M. Kocak, Intermetallics 7 (1999) 1025. [12] A.I. Ustinov, Y.V. Falchenko, A. Ya Ishchenko, G.K. Kharchenko, T.V. Melnichenko, A.N. Muraveynik, Intermetallics 16 (2008) 1043. [13] G.X. Luo, G.Q. Wu, Z. Huang, Z.J. Ruan, Scr. Mater. 57 (2007) 521. [14] R.K. Shiue, S.K. Wu, S.Y. Chen, Intermetallics 11 (2003) 661. [15] R.K. Shiue, S.K. Wu, S.Y. Chen, Intermetallics 12 (2004) 929. [16] R.K. Shiue, S.K. Wu, S.Y. Chen, Acta Mater. 51 (2003) 1991. [17] I.C. Wallis, H.S. Ubhi, M.-P. Bacos, P. Josso, J. Lindqvist, D. Lundstrom, A. Wisbey, Intermetallics 12 (2004) 303. [18] S.J. Lee, S.K. Wu, R.Y. Lin, Acta Mater. 46 (1998) 1283. [19] Villars, Prince, Okamoto, Handbook of Ternary Alloy Phase Diagrams, ASM International, 1995. [20] X.G. Song, J. Cao, Y.F. Wang, J.C. Feng, Mater. Sci. Eng. A 528 (2011) 5135. [21] Y. Jin, J.N. Wang, J. Yang, Y. Wang, Scr. Mater. 51 (2004) 113. [22] Y.Y. Chen, B.H. Li, F.T. Kong, J Alloy Compd. 457 (2008) 265. [23] B.H. Li, Y.Y. Chen, Z.Q. Hou, F.T. Kong, J Alloy Compd. 473 (2009) 123. [24] M. Takeyama, S. Kobayashi, Intermetallics 13 (2005) 993. [25] B. Huneau, P. Rogl, K. Zeng, R. Schmid-Fetzer, M. Bohn, J. Bauer, Intermetallics 7 (1999) 1337. [26] F.T. Kong, Y.Y. Chen, B.H. Li, Mater. Sci. Eng. A 499 (2009) 53. [27] Z.W. Huang, W. Voice, P. Bowen, Scr. Mater. 48 (2003) 79. [28] X.G. Song, J. Cao, Y.Z. Liu, J.C. Feng, Intermetallics 22 (2012) 136.
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Brazing TiAl intermetallics using TiNi–V eutectic brazing alloy X.G. Song a,b , J. Cao a,∗ , H.Y. Chen a , Y.F. Wang a , J.C. Feng a a b
State Key Lab of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, China School of Materials Science and Engineering, Harbin Institute of Technology at WeiHai, WeiHai 264209, China
a r t i c l e
i n f o
Article history: Received 3 September 2011 Received in revised form 14 January 2012 Accepted 1 May 2012 Available online 11 May 2012 Keywords: Intermetallics Brazing TiNi–V eutectic brazing alloy Interfacial microstructure Mechanical properties
a b s t r a c t A novel TiNi–V25 (at.%) eutectic brazing alloy, which was composed of TiNi intermetallics and V, was designed and fabricated. Subsequently, the brazing alloy was used to braze TiAl intermetallics (Ti–42.5Al–9V–0.3Y (at.%)) and reliable joints were obtained. The typical interfacial microstructure of TiAl brazed joints was characterized by employing SEM, EDS and XRD, and the brazing mechanism was analyzed in detail. The brazed joints mainly consisted of B2 phase and 3 -Al3 NiTi2 intermetallics. Interfacial morphology of TiAl joints varied with the increase of brazing temperature, while the phase constitution in joints did not change sound TiAl joints with the highest average shear strength of 196 MPa were obtained when the specimens were brazed at 1220 ◦ C for 10 min. The presence of B2 phase and 3 Al3 NiTi2 intermetallics in joints deteriorated the joining properties due to their brittleness, which led to cleavage-dominated fracture after shear test. The crack propagation path and fracture location were mainly determined by the content and distribution of 3 -Al3 NiTi2 phase. © 2012 Elsevier B.V. All rights reserved.
1. Introduction TiAl intermetallics have been considered as the most promising alternative materials to replace traditional heat-resistant steels and superalloys in automobile industry for significant weight saving [1,2]. The excellent properties such as high specific strength, excellent creep strength and good oxidation resistance make them particularly appropriate for next-generation aircraft turbine engines [3–5]. During the last 30 years, an enormous research activity has been devoted to developing TiAl intermetallics all around the world [6–8]. However, extensive applications have not been achieved due to the materials’ intrinsic brittleness and the difficulties in manufacturing as well as high cost [9]. It is well known that joining techniques will play an important role in the applications of TiAl alloys. Considerable investigations on joining TiAl alloys have been carried out along with their development. Conventional fusion welding is unsuitable although Smarsly et al. [10] reported that sound TiAl joints could be produced with careful process controlling involving preheating at 700 ◦ C and pulsing slow-cooling after welding. Whereas, the solidification cracks caused by the limited ductility and fracture toughness of TiAl alloys cannot be easily avoided during the cooling process. Solid state bonding has also been used to join TiAl intermetallics and high quality joints can be obtained [11–13]. However, a long-time exposure
∗ Corresponding author. Tel.: +86 451 86418882; fax: +86 451 86418146. E-mail address: cao [email protected] (J. Cao). 0921-5093/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.05.002
of TiAl substrate at elevated temperature in bonding process can change its microstructure and thus degrade its mechanical properties. Furthermore, the use of these methods will be limited when the shape of components is complicated. Brazing, as the most feasible and economical method, has received special attentions in joining TiAl alloys to themselves or to other materials. BAlSi-4 brazing alloy was used by Wu et al. [14] to join TiAl alloy and the corresponding relationship between interfacial microstructure and joints properties was analyzed in detail. Infrared brazing of TiAl alloy using pure Ag and BAg-8 brazing alloys was investigated by Shiue et al. [15,16]. The interfacial microstructural evolution, joints properties and reaction kinetics were comprehensively evaluated. However, the joints brazed using Al, Ag or Au based brazing alloys have poor thermal resistance and are usually used below 500 ◦ C. As a result, the high-temperature properties of TiAl alloys could not be taken full advantage of. Therefore, the high temperature brazing alloys (e.g. Ti and Ni based brazing alloys) showed their advantage for high temperature applications of brazed joints. Ti–Ni–Cu brazing alloys were used by Wallis et al. [17] and Wu et al. [18] to braze TiAl alloys and high-quality brazed joints with good room-temperature and high-temperature properties were obtained. It is well known that the development of high temperature brazing alloys is very important for the practical applications of TiAl alloys. Therefore, in the present study, a novel high temperature TiNi–V eutectic brazing alloy was designed and fabricated to braze TiAl alloy. Furthermore, the typical interfacial microstructure was identified, and the effects of brazing temperature on interfacial microstructure and shear strength of TiAl brazed joints were investigated.
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Table 1 Properties of the joining materials. Materials
0.2 (MPa)
b (MPa)
ı (%)
TiAl TiNi–V
420–462 510–540
520–586 722–766
0.5–0.65 3.4–5.6
2. Experimental In this study, the TiAl alloy with nominal composition of Ti–42.5Al–9V–0.3Y (at.%) was fabricated by induction skull melting using high purity (99.99 wt.%) Ti, Al, Al–V and Al–Y interalloys. Subsequently, the ingot was treated at 900 ◦ C/48 h in air and hot isostatic-pressed at 1250 ◦ C/4 h/175 MPa in argon atmosphere to eliminate segregation and shrinkage porosity respectively. The TiNi–V brazing alloy with its nominal composition of Ti37.5Ni37.5V25 (at.%) was prepared by vacuum arc remelting using high purity (99.99 wt.%) Ti, Ni and V rods [19]. All the raw materials were cleaned by 1.5HF–2.5HNO3 –5HCl–91H2 O (ml) and saturated in NaOH solution (cNaOH = 2 mol/l) before vacuum arc remelting. Alloy ingots were turned and remelted at least six times to be homogenized; the final weight loss of the master alloy was less than 0.1 wt.%. A diffusion-annealing treatment was done at 1000 ◦ C for 4 h after vacuum arc remelting. The main properties of the TiAl alloy and TiNi–V brazing alloy are listed in Table 1. The TiNi–V alloy was cut into foils with a thickness of 400 m by spark cutting. Both surfaces of TiNi–V foil were ground on SiC grit papers until its thickness was about 200 m. The TiAl base material was also cut into blocks with two different sizes (5 mm × 5 mm × 2.5 mm and 20 mm × 10 mm × 2.5 mm). Prior to brazing, the brazing surfaces of TiAl specimens were ground on SiC grit papers and polished using diamond pastes. All samples were cleaned ultrasonically in acetone for 10 min and dried by air blowing. The TiNi–V foil was placed between two TiAl blocks. At the beginning of brazing process, the furnace was heated to 1100 ◦ C at a rate of 30 ◦ C/min, then to the brazing temperatures at a rate of 10 ◦ C/min. Subsequently, the furnace temperature were held for 10 min and then cooled down to 200 ◦ C at a rate of 5 ◦ C/min, and finally cooled down spontaneously to room temperature. During the brazing process, the vacuum was kept at (1.3–2.0) × 10−3 Pa and a pressure of 20 kPa was applied to ensure a proper contact.
The specimens for metallographic observation were crosssectioned, perpendicular to the brazed interface. The interfacial microstructure was characterized employing scanning electron microscopy (SEM). Componential analysis of various phases in the joints was carried out using an energy dispersive spectrometer (EDS) with the operation voltage of 15 kV and minimum spot size of 1 m. The shear tests were performed at a constant speed of 0.5 mm/min by a universal testing machine (Instron1186) at room temperature. A schematic of the shear test could be seen in Ref. [20]. For each set of experimental data, at least five samples were used to average the joints strength. After shear test, the fracture of brazed joints was inspected by SEM. In order to identify the interfacial phases accurately, the analysis was performed using an X-ray diffraction (XRD, JDX-3530M) spectrometer equipped with Cu K␣ radiation on the fracture surfaces. In addition, differential thermal analysis (DTA) was performed in order to determine the melting point of TiNi–V brazing alloy. 3. Results and discussion 3.1. Microstructure of TiAl alloy and TiNi–V brazing alloy Fig. 1(a) shows the microstructure of TiAl substrate in backscattered electron (BSE) mode. It is seen that the alloy was homogenous, fine-grained nearly lamellar structure, and the lamellar colony size is about 40–80 m. Fig. 1(b) shows the XRD patterns measured on the surface of the alloy. The result indicates that the alloy was primarily composed of ␥-TiAl phase, ␣2 -Ti3 Al phase and B2 phase. Each phase could be identified according to the EDS results taken from the selected spots in Fig. 1(a). It is shown that V concentration in B2 phase (marked by B) was much higher than that in other phases. Relatively high V concentration in B2 phase is consistent with the viewpoint that element V could stabilize  phase in TiAl alloy [21]. The resultant microstructure was characterized by the near lamellar colonies surrounded by B2 + ␥ structure. The white particles (marked by C) could be determined as YAl2 intermetallic compound according to literature [22]. However, diffraction peaks of YAl2 phase cannot be found in the XRD patterns, for the reason that the phase was too sparse to be detected by XRD. It is noted that these YAl2 particles were mainly distributed in ␥-TiAl phase. The solidification and transformation pathway has been reported as the follows: L → L +  →  → ␣ +  → ␣ +  + ␥ → Lamellar +B2 + ␥ [23].
Fig. 1. Backscattered electron image (BEI), XRD patterns and EDS analysis results in atomic percent of TiAl alloy. (a) BEI of TiAl alloy, (b) XRD patterns of TiAl alloy.
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Fig. 2. (a) BEI, (b) XRD patterns and (c) DTA result of TiNi–V brazing alloy.
The typical microstructure of TiNi–V brazing alloy is shown in Fig. 2(a). It is clearly seen that eutectic structure was obtained after vacuum arc remelting. In order to further investigate its phase constitution, X-ray diffraction analysis was carried out. The XRD patterns in Fig. 2(b) indicate that the alloy consisted of TiNi intermetallics and V. According to the DTA result shown in Fig. 2(c), it can be indicated that the melting point of TiNi–V brazing alloy is about 1142 ◦ C. 3.2. Interfacial microstructure of TiAl/TiNi–V/TiAl brazed joints Fig. 3 shows the typical interfacial microstructure of TiAl/TiNi–V/TiAl joint brazed at 1220 ◦ C for 10 min. It clearly demonstrates that a sound brazing seam was obtained, and its thickness is about 400 m. An increasing in thickness indicates that intensive interactions occurred between TiAl substrate and molten TiNi–V brazing alloy during the brazing process. For the sake of convenience, the brazing seam of TiAl joint is artificially divided into three zones: zone I, II and III respectively, as shown in Fig. 3. The distribution of elements was measured by EDS along the white base line. A relatively high concentration of Al in brazing seam indicates that both the amount of dissolved Al and the diffusion rate of Al in molten brazing alloy were high. The concentration profiles of Ti and V show that their distributions were heterogeneous, which is probably due to their low diffusion rates in molten brazing alloy. In order to further investigate the interfacial microstructure, more details are presented in a larger visual field in BSE mode, as shown in Fig. 4, and the EDS results taken from different phases are also included. Fig. 4(a) shows the microstructure of TiAl substrate after brazing. Compared with the original microstructure shown in Fig. 1, it is observed that there was no obvious change except that the lamellar structure leveled more smoothly. According to the phase diagram drawn by Takeyama et al. [24], during the heating process in brazing, the solid-state phase transformation
Fig. 3. BEI of TiAl joint brazed using TiNi–V brazing alloy at 1220 ◦ C for 10 min and the elemental distributions of Al, Ti, V and Ni along the white base line.
of Lamellar + B2 + ␥ → ␣ +  + ␥ occurred in TiAl substrate, as well as the B2 →  disordered transformation which occurred above 1100 ◦ C. However, in the present study, the transformations of TiAl substrate were incomplete due to a short holding time at high temperature. Therefore, during cooling process, the fresh Lamellar (␥ + ␣2 ) and B2 phase could precipitate based on the remnant Lamellar and B2 phase respectively, which resulted in a similar microstructure to that of original TiAl alloy. In addition, the lower cooling rate during brazing is beneficial to the transformation of ␣ +  + ␥ → Lamellar + B2 + ␥, and thus a more smooth lamellar structure was obtained finally. The microstructure of zone I is shown in Fig. 4(b). The gray phase (marked by D) had the average composition of 42.44Ti, 38.72Al, 16.27V, 2.13Ni and 0.44Y (at.%), which could be identified as B2 phase. It is evident that the amount of B2 phase increased in this zone. The formation process of B2 phase can be described as follows: TiNi–V brazing alloy began to melt when the brazing temperature exceeded its melting point. Subsequently, the dissolution of TiAl to molten brazing alloy and the diffusion of molten brazing alloy to TiAl substrate occurred simultaneously, which led to an increase of V content in the solid–liquid interfacial area. Thus, the amount of B2 phase in the joints increased after brazing because element V could stabilize  phase in TiAl alloy [20]. In addition, the massive ␥ phase in zone I (marked by C) was not dissolved completely, and it was alloyed with small amounts of Ni (1.04 at.%) due to the interdiffusion between the brazing alloy and TiAl substrate. Furthermore, a new phase (marked by E) came forth, which has an average composition of 31.53Ti, 47.57Al, 2.52V, 17.58Ni and 0.8Y (at.%). According to Rogl’s research [25], the phase could be identified as 3 -Al3 NiTi2 with the MgZn2 -type structure. Fig. 4(c) shows the typical microstructure in zone II. It can be clearly seen that the zone mainly consisted of B2 phase (marked by G) and 3 -Al3 NiTi2 phase (marked by F). It is important to note that the B2 phase colonies were surrounded by 3 -Al3 NiTi2 intermetallics, and the white YAl2 particles were mainly distributed in 3 -Al3 NiTi2 phase. The microstructure in zone II is very helpful to deducing the reaction mechanisms. Chen et al. [22,26] reported that element Y promoted ␥ phase nucleation, which resulted in more Y in ␥ phase. The same results were also found in this study, as above analysis of the microstructure of TiAl (shown in Fig. 1). The redistribution of YAl2 from ␥ phrase to 3 -Al3 NiTi2 phase after brazing indicates that most of the massive ␥ phase at grain boundaries dissolved into the molten brazing alloy during the brazing process. Thus, the lamellar colony was gradually surrounded by the molten brazing alloy which dissolved Al, and Ti. The diffusion of V into lamellar colony or ␥ single phase led to the transformation from lamellar colony (␥ + ␣2 ) or ␥ single phase to B2 phase [27]. In addition, B2 phase could also form by the solidification of molten brazing alloy. Due to the slow cooling rate, the solidification could be considered as an equilibrium or quasi-equilibrium process, in which impurities (e.g. YAl2 ) were driven from the crystallization surfaces, finally located in 3 -Al3 NiTi2 phase, as shown in Fig. 4(c).
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Fig. 4. High magnification BEIs of each zone marked in Fig. 3. (a) TiAl substrate, (b) zone I, (c) zone II and (d) zone III.
The microstructure in zone III is shown in Fig. 4(d). It can be seen that the microstructure of this zone was mainly composed of two phases. The location marked by H with an average composition of 32.05Ti, 33.18Al, 31.66V, 2.74Ni and 0.37Y (at.%) could be determined as B2 phase because of the high V content. In addition, the other location marked by I had a similar chemical composition with locations E and F, thus, the phase could also be identified as 3 Al3 NiTi2 intermetallics. However, compared with the 3 -Al3 NiTi2 phase in zone II (shown in Fig. 4(c)), which was located at the boundaries of massive B2 phase, the distribution of 3 -Al3 NiTi2 phase was changed greatly, as shown in Fig. 4(d). According to above analysis, it can be concluded that the brazed joints mainly consisted of B2 phase and 3 -Al3 NiTi2 intermetallics. During the brazing process, TiAl substrate was partially dissolved into the molten brazing alloy. Thus, the contents of Al and Ti in molten brazing alloy increased, which led to the formation of ternary intermetallic compound 3 -Al3 NiTi2 . The diffusion of V to TiAl substrate also played an important role, which resulted in the transformation of lamellar colony (␥ + ␣2 ) to B2 phase. During the cooling process, B2 phase directly solidified from the molten brazing alloy, while element Ni was enriched in the remnant molten brazing alloy and finally reacted with Ti and Al to form 3 -Al3 NiTi2 phase at the boundaries of massive B2 phase. Further cooling, the 3 -Al3 NiTi2 phase precipitated in the B2 phase at the center of brazing seam, probably because solubility of element Ni in B2 phase decreased as did brazing temperature.
3.3. Effects of brazing temperature on the interfacial microstructure and shear strength of the brazed joints The interfacial microstructure of TiAl joints brazed at different temperatures is shown in Fig. 5. The interfacial morphology changed obviously with the increase of brazing temperature. In other words, brazing temperature had a great influence on the interfacial microstructure. When TiAl alloy was brazed at lower temperatures (e.g. 1180 ◦ C and 1200 ◦ C), a continuous layer formed between TiAl substrate and brazing seam, as shown in Fig. 5(a) and (b). A small amount of remnant brazing alloy was found in the central region of the joint brazed at 1180 ◦ C, which was caused by poor melting at lower brazing temperatures. Another style of interfacial microstructure was obtained when the brazing temperature exceeded 1220 ◦ C (e.g. 1240 ◦ C and 1260 ◦ C), as shown in Fig. 5(c) and (d). Continuous reaction layer disappeared and the microstructure was similar to that observed in Fig. 3. Fig. 5(e)–(g) shows the high magnification BEIs of the rectangular regions in Fig. 5(a)–(c) respectively, where the EDS chemical analysis results taken at the selected locations are also included. The interfacial microstructure between TiAl and brazing seam obtained at lower temperatures can be seen in Fig. 5(e) and (f). According to the BEIs and EDS results, the continuous layer close to TiAl was determined as B2 phase, as marked by A and D in Fig. 5(e) and (f) respectively. The massive phase marked by B and F was also B2 phase due to the similar composition to that of the continuous layer. The locations marked by C and E which had relatively high
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Fig. 5. Interfacial microstructure of TiAl joints brazed at different temperatures. (a) 1180 ◦ C, (b) 1200 ◦ C, (c) 1240 ◦ C, (d) 1260 ◦ C, (e) high magnification BEIs of the rectangular regions in (a), (f) high magnification BEIs of the rectangular regions in (b), and (g) high magnification BEIs of the rectangular regions in (c).
Ni content could be inferred as 3 -Al3 NiTi2 . Based on the analysis above, it can be concluded that the interfacial morphologies of TiAl joints brazed at different temperatures were different from each other. However, the interfacial phase constitution was insensitive to brazing temperature. Fig. 5(g) shows the high magnification BEI of the rectangular region in Fig. 5(c). A colony of B2 + ␥ surrounded by 3 -Al3 NiTi2 strips was observed, and its size was about
50–60 m. It can be inferred that the colony was transformed from original lamellar colony. The typical microstructure was very similar to that shown in Fig. 4(c), and it confirmed that the transformations stated above were reasonable. Fig. 6 illustrates the average shear strength of TiAl joints brazed at different temperatures. The specimen brazed at 1220 ◦ C for 10 min demonstrated the highest average shear strength of
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Fig. 6. Effect of brazing temperature on joining shear strength.
Fig. 8. XRD patterns of the fracture surfaces after shear test for specimen brazed at 1220 ◦ C for 10 min.
196 MPa. In addition, the shear strength decreased slightly with increasing brazing temperature. In order to investigate the relationship between joints properties and interfacial microstructure, fracture analysis was performed by SEM, EDS and XRD. The typical fractographs are shown in Fig. 7(a) and (b). Cleavage facets were widely observed in the fracture surfaces. XRD result of the fracture surfaces is illustrated in Fig. 8. Combining the EDS results of selected locations and X-ray analysis, it is concluded that the 3 -Al3 NiTi2 phase and B2 phase dominated in the facture surfaces, as well as some lamellar colonies, as shown in Fig. 7(b). Actually, the presence of B2 phase and 3 Al3 NiTi2 phase deteriorated the joints properties due to their nature
of brittleness. Therefore, cleavage-dominated fracture formed after shear test. In addition, the XRD result confirms that the above analysis about interfacial microstructure and phases is accurate. Fig. 7(c)–(e) show cross-section BEIs of brazed joints after shear test. It is obvious that the joints always crack in the brazed seam region while the fracture location varied with the brazing temperature. When brazed at low temperatures (e.g. 1180 ◦ C), the crack initiated and propagated in a continuous reaction layer during shear test, as shown in Fig. 7(c). According to the analysis of interfacial microstructure shown in Fig. 5(a), it is found that the reaction layer was 3 -Al3 NiTi2 phase. When brazing temperature was increased to 1200 ◦ C, fracture location changed from
Fig. 7. Fractographs of TiAl joint brazed at 1220 ◦ C for 10 min and the EDS analysis results of selected locations. (a) Low magnification SEM image, (b) high magnification SEM image, (c) cross-section BEI of TiAl joint after shear test brazed at 1180 ◦ C, (d) cross-section BEI of TiAl joint after shear test brazed at 1200 ◦ C and (e) at 1220 ◦ C.
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3 -Al3 NiTi2 layer to the brazing seam because the continuous 3 Al3 NiTi2 layer disappeared, as shown in Fig. 5(b). Fig. 7(e) shows the cross-section BEI of TiAl joint brazed at 1220 ◦ C, which exhibited the highest shear strength. It can be seen that the crack propagated along the interface (zone II) between TiAl substrate and brazing seam. By comparing the interfacial microstructure and fracture locations in different joints, it is concluded that the crack preferred to initiate and propagate at the locations where the 3 -Al3 NiTi2 content was relatively high. The result also can be confirmed in Fig. 7(b) where the fracture surfaces contained lots of 3 -Al3 NiTi2 phases. So it can be inferred that 3 -Al3 NiTi2 phase had poor properties and its content and distribution determined the crack propagation path and fracture location, as well as joining properties, which is consistent with our previous results from the investigation of brazing TiAl alloy using TiNi–Nb eutectic brazing alloy [28]. 4. Conclusions Interfacial microstructure and joining properties of TiAl joints brazed using TiNi–V eutectic brazing alloy were investigated in this study. Primary conclusions are summarized as follows. (1) A TiNi–V eutectic brazing alloy was designed and developed by vacuum arc remelting. The alloy had typical eutectic microstructure which mainly consisted of TiNi intermetallic compound and V. The melting point of this filler metal was about 1142 ◦ C. (2) Brazing of TiAl intermetallics using TiNi–V brazing alloy was also achieved. The joint was mainly composed of B2 phase and 3 -Al3 NiTi2 intermetallics. During the brazing process, intensive reactions occurred between TiAl substrate and molten brazing alloy. The diffusion of V to TiAl led to two transformations as following: Lamellar (␥ + ␣2 ) + V → B2 and ␥ + V → B2. (3) The interfacial morphologies of TiAl joints brazed at different temperatures were different from each other. However, all the TiAl joints had the same phase constitution, i.e. B2 phase and 3 -Al3 NiTi2 phase, which was detrimental to the joining properties. The highest average shear strength of 196 MPa was obtained when the brazing condition was 1220 ◦ C/10 min. Fracture analysis showed that cleavage-dominated fracture formed during shear test and the crack always initiated and propagated
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at the locations where the 3 -Al3 NiTi2 content was relatively high. Acknowledgments The authors gratefully acknowledge the financial support from Project 2011CB605505 supported by the National Basic Research Program of China. The authors also acknowledge the financial support from the National Natural Science Foundation of China under Grant No. 50805038 and the Fundamental Research Funds for the Central Universities under Grant No. HIT.NSRIF.2011036. References [1] K. Maruyama, M. Yamaguchi, G. Suzuki, H.L. Zhu, H.Y. Kim, M.H. Yoo, Acta Mater. 52 (2004) 5158. [2] E.A. Loria, Intermetallics 8 (2000) 1339. [3] S.L. Draper, D. Krause, B. Lerch, I.E. Locci, B. Doehnert, R. Nigam, G. Das, P. Sickles, B. Tabernig, N. Reger, K. Rissbacher, Mater. Sci. Eng. A 464 (2007) 330. [4] T. Tetsui, Mater. Sci. Eng. A 329–331 (2002) 582–588. [5] S.L. Draper, D. Krause, B. Lerch, Mater. Sci. Eng. A 464 (2007) 330. [6] X. Wu, Intermetallics 14 (2006) 1114–1122. [7] G.L. Chen, X.J. Xu, Z.K. Teng, Y.L. Wang, J.P. Lin, Intermetallics 15 (2007) 625. [8] G.L. Chen, L.C. Zhang, Mater. Sci. Eng. A 329–331 (2002) 163. [9] A. Lasalmonie, Intermetallics 14 (2006) 1123. [10] W. Smarsly, H. Cramer, A.W.E. Nentwig, Presented at Werkstoffwoche 98, Munich, 12–15 October, 1998 (in German). [11] G. Cam, H. Clemens, R. Gerling, M. Kocak, Intermetallics 7 (1999) 1025. [12] A.I. Ustinov, Y.V. Falchenko, A. Ya Ishchenko, G.K. Kharchenko, T.V. Melnichenko, A.N. Muraveynik, Intermetallics 16 (2008) 1043. [13] G.X. Luo, G.Q. Wu, Z. Huang, Z.J. Ruan, Scr. Mater. 57 (2007) 521. [14] R.K. Shiue, S.K. Wu, S.Y. Chen, Intermetallics 11 (2003) 661. [15] R.K. Shiue, S.K. Wu, S.Y. Chen, Intermetallics 12 (2004) 929. [16] R.K. Shiue, S.K. Wu, S.Y. Chen, Acta Mater. 51 (2003) 1991. [17] I.C. Wallis, H.S. Ubhi, M.-P. Bacos, P. Josso, J. Lindqvist, D. Lundstrom, A. Wisbey, Intermetallics 12 (2004) 303. [18] S.J. Lee, S.K. Wu, R.Y. Lin, Acta Mater. 46 (1998) 1283. [19] Villars, Prince, Okamoto, Handbook of Ternary Alloy Phase Diagrams, ASM International, 1995. [20] X.G. Song, J. Cao, Y.F. Wang, J.C. Feng, Mater. Sci. Eng. A 528 (2011) 5135. [21] Y. Jin, J.N. Wang, J. Yang, Y. Wang, Scr. Mater. 51 (2004) 113. [22] Y.Y. Chen, B.H. Li, F.T. Kong, J Alloy Compd. 457 (2008) 265. [23] B.H. Li, Y.Y. Chen, Z.Q. Hou, F.T. Kong, J Alloy Compd. 473 (2009) 123. [24] M. Takeyama, S. Kobayashi, Intermetallics 13 (2005) 993. [25] B. Huneau, P. Rogl, K. Zeng, R. Schmid-Fetzer, M. Bohn, J. Bauer, Intermetallics 7 (1999) 1337. [26] F.T. Kong, Y.Y. Chen, B.H. Li, Mater. Sci. Eng. A 499 (2009) 53. [27] Z.W. Huang, W. Voice, P. Bowen, Scr. Mater. 48 (2003) 79. [28] X.G. Song, J. Cao, Y.Z. Liu, J.C. Feng, Intermetallics 22 (2012) 136.