Joining of Zirconia and Ti-6Al-4V Using a Ti-based Amorphous Filler

J. Mater. Sci. Technol., 2011, 27(7), 653-658. Joining of Zirconia and Ti-6Al-4V Using a Ti-based Amorphous Filler Yuhua Liu 1), Jiandong Hu 1), Yaping Zhang 1), Zuoxing Guo 1) and Yue Yang 2) 1) Key Lab. of Automobile Materials, Ministry of Education, College of Materials Science and Engineering, Jilin University, Changchun 130025, China 2) Key Lab. of Advanced Structural Materials, Ministry of Education, Changchun University of Technology, Changchun 130012, China [Manuscript received February 24, 2011, in revised form April 24, 2011] Polycrystalline Zr 2 3 mol.%y 2 3 was brazed to Ti-6Al-4V by using a 7 Zr 28 Cu 14 Ni 11 (at.%) amorphous ribbon at 1123 1273 K in a high vacuum. The influences of brazing temperature on the microstructure and shear strength of the joints were investigated. The interfacial microstructures can be described as Zr 2 /Ti+Ti 2 +Cu 2 +Ni 2 /α-ti+(ti,zr) 2 (Cu,Ni) eutectic/acicular Widmanstäten structure/ti 6Al 4V alloy. With the increase in the brazing temperature, the thickness of the Ti+Ti 2 +Cu 2 +Ni 2 layer reduced, the content of the α-ti+(ti,zr) 2 (Cu,Ni) eutectic phase decreased, while that of the coarse α- Ti phase gradually increased. The shear strength of the joints did not show a close relationship with the thickness of the Ti+Ti 2 +Cu 2 +Ni 2 layer. However, when the coarse (Ti,Zr) 2 (Cu,Ni) phase was non-uniformly distributed in the α-ti phase, or when α-ti solely situated at the center of the joint, forming a coarse block or even connecting into a continuous strip, the shear strength greatly decreased. KEY WRDS: Brazing; Amorphous filler; Interface; Microstructure; Shear strength 1. Introduction Ti-6Al-4V is one of the most important titanium alloys and widely used in aerospace industries [1]. This alloy has excellent fracture toughness and corrosion resistance. Also, it can be readily welded, forged and machined. n the other hand, zirconia is an important structural ceramic because of its high strength and fracture toughness [2]. However, its brittleness and lack of flexibility make the fabrication of complexshaped and/or large-sized components difficult [3,4]. Joining of zirconia can overcome this drawback to a great extent. ver the past decades, many investigations have been carried out on the joining of zirconia to itself and to metals [1 3,5 9]. Various methods such as brazing, transient liquidphase bonding and diffusion bonding could be used for ceramic-metal or ceramic-ceramic joining [1,2]. However, the most popular method is active brazing, par- Corresponding author. Tel: +86 431 85095813; Fax: +86 431 85095876. E-mail address: guozx@jlu.edu.cn (Z.X. Guo). ticularly using the Ag-Cu-Ti alloys in eutectic or close to eutectic compositions [1 3,5, 6]. Nevertheless, consumption of a large amount (60 70 wt%) of the noble metal, Ag, in the brazing filler is the main disadvantage of these alloys. In recent decades, efforts have been taken to develop novel brazing fillers with low cost and equivalent or even superior bonding properties as compared to the Ag-Cu-Ti alloys. Recently, amorphous alloys have been used as brazing fillers. They were reported to accelerate atomic diffusion and surface reaction during brazing process [10,11]. In addition, they are expected to decrease brazing temperature so as to reduce residual stress developed in the joint and hence to increase the joint strength. They also have superior wettability than atomized powders or paste formulations that require large gaps at the joint for filling. Furthermore, the amorphous fillers [12 15] have significant advantages such as unique of foil form combined with outstanding ductility and flexibility, wide range of width and thickness available. In particular, the

654 Y.H. Liu et al.: J. Mater. Sci. Technol., 2011, 27(7), 653 658 amorphous fillers containing active Ti and Zr elements with good glass forming ability may replace the traditional Ag-Cu-Ti filler alloys in some cases. The Ti-based amorphous fillers offer a combination of low density, high specific strength and relativly low cost [16]. Zou et al. [17] joined the Si 3 N 4 ceramic to itself using a Ti-Zr-Ni-Cu amorphous foil, and obtained much larger bending strength of the joint as compared with that using a crystalline filler. However, to our knowledge, there is no report on the joining of the Zr 2 /metal assemblies using amorphous fillers. In this study, the Zr 2 ceramic was brazed to the Ti-6Al-4V alloy using a 7 Zr 28 Cu 14 Ni 11 (at.%) amorphous filler. The microstructural development in the joints was examined. The influences of brazing temperature on the interfacial microstructure and the joint shear strength were investigated. Fig. 1 Schematic of assembling parts for the brazing experiments 2. Experimental 2.1 Starting materials The base metal used in the experiments was Ti 6Al 4V plate with the dimensions of 20 mm 5 mm 5 mm, and its chemical composition in weight percent was 5.76 Al, 4.03 V, 0.28 Fe, 0.06 C and balance Ti. The sizes of the sintered Zr 2 pieces were also 20 mm 5 mm 5 mm. The 7 Zr 28 Cu 14 Ni 11 amorphous foil in a thickness of 25 µm was cut into dimensions of 15 mm 5 mm. The melting temperature of the 7 Zr 28 Cu 14 Ni 11 foil was measured to be 1116 K using a differential scanning calorimeter (DSC, NETZSCH STA 409 PC, Germany). 2.2 Joint sample preparation Both the Ti 6Al 4V metal and the Zr 2 piece were polished with different sizes of diamond pastes. The 7 Zr 28 Cu 14 Ni 11 amorphous foil was gently ground with 1000# SiC sand paper. Before assembling, the Zr 2, Ti 6Al 4V and 7 Zr 28 Cu 14 Ni 11 foil were ultrasonically cleaned in acetone and then they were assembled into a sandwich structure, as schematically shown in Fig. 1. Brazing was performed at isothermal temperatures between 1123 K and 1273 K in a high vacuum of 5 10 4 Pa. The samples were heated at a rate of 20 K min 1, holding at the brazing temperature for 30 min, then cooled at 5 K min 1 to 600 K and finally furnace-cooled. 2.3 Sample characterization Selected samples were sectioned and polished for microstructural observation using a field emission scanning electron microscope (FESEM, JSM-6700F, Japan) equipped with an energy dispersive spectrometer (EDS). The phases at the interfaces were identified by X-ray diffraction (XRD, D/2500PC Rigaku, Fig. 2 Schematic of the assemblies for the shear strength test Japan) using a Cu target with Kα radiation. XRD pattern was performed on the fractured surfaces after the shear test. The shear strength of the joint was measured by using a MTS tester (MTS-810, U.S.A.) at a constant speed of 0.1 mm/min. The samples were fixed in a special device, as illustrated in Fig. 2. The reported values were average of three tests. Finally, the fractured surfaces were examined by FESEM. 3. Results and Discussion 3.1 Microstructures Figures 3(a) and (b) show the back-scattered electron (BSE) image and elemental mapping of the Zr 2 /Ti 6Al 4V joint brazed at 1123 K for 30 min, respectively. It is obvious that four layers, with a total thickness of about 45 µm, were formed at the interface (see Fig. 3(b)). The gray layer adjacent to

Y.H. Liu et al.: J. Mater. Sci. Technol., 2011, 27(7), 653 658 655 Intensity / a.u. (c) -Ni 2 -(Ti,Zr) 2 (Cu,Ni) -Ti -Cu 2 - -Ti -Ti 2 20 30 40 50 60 70 80 2 / deg. Fig. 3 (a) Back-scattered electron image of the Zr 2/Ti 6Al 4V interface zone, (b) elemental mapping for the Zr 2/Ti 6Al 4V interface and (c) XRD pattern for the interfacial reaction product in the joint brazed at 1123 K for 30 min the Zr 2 ceramic is called layer 1, whose thickness is about 10 µm. In the intermediate part, there are layers 2 and 3. Layer 3 consists of coarse white phase while layer 2 mainly of white phase intermixed with uniformly distributed black phase. The zone adjacent to the Ti 6Al 4V alloy is layer 4, whose composition is more or less close to that in layer 2. The elemental mapping images show that a considerable amount of the titanium atoms moved to the interface. n the contrary, Cu, Ni, Zr and a low content of the Ti atoms were assembled in layer 3. Layer 1 contained a large number of Cu, Ni, Zr, and Ti atoms. The detailed elemental compositions in the interfacial reaction layer of the Zr 2 /Ti 6Al 4V joint were investigated by the quantitative EDS analysis under high magnifications and the results are given in Table 1. The compositions were normalized excluding oxygen due to its insufficient accuracy by EDS analysis. As indicated, point A in Fig. 3(a) is rich in Ti with small contents of Cu, Ni and Zr. Point B is also composed of the Ti rich phase with low contents of Cu, Ni and Zr. Point C in composition is close to the original Ti Zr Cu Ni amorphous filler with the incorporation of some amount of Al. Point D is composed of the Ti rich phase with the low contents of Cu, Ni, Zr, Al and V as well. In order to identify the reaction phases in the joint, XRD analysis was performed on the interfacial

656 Y.H. Liu et al.: J. Mater. Sci. Technol., 2011, 27(7), 653 658 Intensity / a.u. c b a (d) a 1173 K b 1223 K c 1273 K Ni 2 -(Ti,Zr) 2 (Cu,Ni) -Ti -Cu 2 - -Ti -Ti 2 20 30 40 50 60 70 80 2 / deg. Fig. 4 Back-scattered electron images of the Zr 2/Ti 6Al 4V interfaces in the joints brazed at (a) 1173 K, (b) 1223 K, (c) 1273 K and (d) XRD patterns for the interfacial reaction products produced at different brazing temperatures for 30 min Table 1 Chemical compositions of various phases observed in Figs. 3 and 4 (at.%) Location Ti Ni Cu Zr Al V A 65.6 8.9 10.7 14.8 B 80.0 3.6 4.5 11.9 C 37.1 12.0 15.4 26.8 8.7 D 73.2 2.9 5.5 6.7 6.6 5.2 E 64.9 6.9 7.4 18.3 2.5 F 87.3 5.5 7.2 H 66.0 9.3 8.5 16.2 I 84.5 2.1 1.8 8.7 2.9 J 68.2 7.3 7.9 14.6 2.0 reaction layer after brazing at 1123 K for 30 min. As shown in Fig. 3(c), the reaction products are quite complex, consisting of Ti, Ti 2, (Ti,Zr) 2 (Cu,Ni), α-ti, Cu 2 and Ni 2 phases. In combination with Table 1, we can conjecture that the reaction products in layer 1 are mainly Ti, Ti 2, Cu 2 and Ni 2. Because Ti has a strong affinity for, the Ti atoms actually aggregated at the zirconia/filler interface. Subsequently, reactions between the atoms in zirconia and the Ti atoms in the filler occurred, creating a dark, -deficient Zr 2 region [2]. Because Cu and Ni have a strong affinity for Ti, it is easy to form stable Ti 2 Ni(Cu) intermetallic compounds [18] and then react with, leading to the formation of the Cu 2 +Ni 2 phases. The composition of layer 3 is almost identical with that of the original filler metal. Through the EDS and XRD analyses, the white block phase (marked by C in layer 3) was confirmed to be (Ti,Zr) 2 (Cu,Ni) and layer 2 consisted mainly of (Ti,Zr) 2 (Cu,Ni) that was uniformly distributed in the α-ti phase. The two phases formed the eutectic-like structure. These phases resulted from the reaction of the dissolved elements from the base alloy with Cu, Ni and Zr in the liquid filler. Hong and Koo [19] also reported the similar eutectic microstructure at the joint interface after brazing a niobium alloy and Ti-6Al-4V using a Ti-20Cu-20Ni-20Zr (wt%) filler metal. In layer 4, an acicular Widmanstäten structure appeared which was adjacent to the Ti 6Al 4V alloy. Figure 4 shows the BSE images and the XRD patterns for the interfacial reaction layers of the joints after brazing at 1173 1273 K for 30 min. Different reaction layers can be clearly distinguished from the contrast. As shown in Figs. 4(a) (c), there is a distinct gray layer in all the brazed joints adjacent to the Zr 2 ceramic. In combination with Table 1 and Fig. 4(d), the reaction products in this layer marked by points H, E and J are also Ti+Ti 2 +Cu 2 +Ni 2. The thickness of this layer decreases with increasing brazing temperature, which was approximately 7 µm,

Y.H. Liu et al.: J. Mater. Sci. Technol., 2011, 27(7), 653 658 657 Shear strength / MPa 80 70 60 50 40 30 20 10 Holding time 30 min 0 1100 1150 1200 1250 1300 T / K Fig. 5 Variation in the shear strength of the joints with the brazing temperature 5 µm and 4 µm, respectively. As the brazing temperature increased, the content of the α- Ti+(Ti,Zr) 2 (Cu,Ni) eutectic phase decreased, while that of the coarse α-ti phase gradually increased and even connected into a continuous strip. It is expected that the higher brazing temperature would enhance interfacial reaction resulting from the increased dissolution and diffusion. With increasing brazing temperature, atomic diffusion rate increased as seen from the disappearance of a large amount of (Ti,Zr) 2 (Cu,Ni) phase. Cu, Ni and Zr in the liquid filler diffused to the base metal, forming white (Ti,Zr) 2 (Cu,Ni) phase adjacent to the Ti alloy (see Figs. 4(b) and (c)). The residual Ti in the liquid filler together with that diffused from the base metal aggregated at the interface. Consequently, the higher temperature gave rise to the larger thickness of the black α-ti phase. The existence of a large number of the brittle α-ti phase gave rise to cracks (see Fig. 4(c)). Based on the above analysis, it can be concluded that the interfacial microstructure of the Zr 2 /Ti 6Al 4V joint can be described as Zr 2 /Ti+Ti 2 +Cu 2 +Ni 2 /α-ti+(ti, Zr) 2 (Cu, Ni) eutectic/acicular Widmanstäten structure/ti 6Al 4V alloy. 3.2 Shear strength Figure 5 shows the effect of the brazing temperature on the shear strength of the joints. As can be seen, the shear strength of the joints decreased with increasing brazing temperature. The maximum shear strength was 63 MPa, and the average was about 60 MPa for the joints brazed at 1123 K for 30 min. Figure 6 shows the cross-sections of the fractured joints. All the fracture occurred in the joints, but not at the reaction layers. It was reported that the formation of a continuous interfacial reaction layer is the most important factor in determining the joint strength [2,20]. We presume that the formation of the Ti+Ti 2 +Cu 2 +Ni 2 reaction layer in this study is critical for the joint strength. However, the Fig. 6 SEM images showing the cross-section microstructures of the brazed joints after shear tests: (a) 1123 K, (b) 1173 K, (c) 1223 K, (d) 1273 K

658 Y.H. Liu et al.: J. Mater. Sci. Technol., 2011, 27(7), 653 658 thickness of the reaction layer does not seem to have a significant effect on the joint strength when it was in the range of 4 10 µm. As shown in Fig. 6(a), fracture occurred at the white block (Ti,Zr) 2 (Cu,Ni) phase. The (Ti,Zr) 2 (Cu,Ni) phase is an intermetallic compound, which is brittle. As the brazing temperature increased to 1173 K (Fig. 6(b)), the location of the fracture moved to the α-ti+(ti,zr) 2 (Cu,Ni) eutectic phases. The non-uniformly distributed (Ti,Zr) 2 (Cu,Ni) phase intermixed with the α-ti phase, forming the coarse structure, resulted in its low strength. As the brazing temperature further increased, as shown in Figs. 6(c) and (d), the fracture occurred at the black α-ti phase. The brittle nature of this phase is harmful to the joint strength [21]. The presence of a high residual stress at the interface may lead to the formation of cracks (Fig. 4(c)). The coefficient of thermal expansion of Ti alloy (10.8 10 6 C 1 ) is similar to that of zirconia (10.0 10 6 C 1 ), thus reducing the thermal mismatch between the ceramic and metal, ultimately minimizing the residual stresses developed upon cooling of the joined component from the relatively high brazing temperature. However, with the thickening of the brittle product, the residual stress in the joint might also accumulate. As a result, the joint strength was lowered down with increasing brazing temperature. 4. Conclusion Zirconia and Ti 6Al 4V alloy were successfully brazed using the 7 Zr 28 Cu 14 Ni 11 (at.%) amorphous foil. The interfacial microstructure can be described as Zr 2 /Ti+Ti 2 +Cu 2 +Ni 2 /α- Ti+(Ti,Zr) 2 (Cu,Ni) eutectic/ acicular Widmanstäten structure/ti 6Al 4V. The shear strength of the joints decreases with increasing brazing temperature as a result of the overgrowth of the brittle α-ti phase. And the presence of residual stress plays a role in the thickening of the brittle phase. Also, when (Ti,Zr) 2 (Cu,Ni) was non-uniformly distributed in the α-ti phase, forming the coarse or block structure, it is harmful to the joint strength. The maximum shear strength is 63 MPa for the joint brazed at 1123 K for 30 min. Acknowledgements This work was supported by 2009 pen Foundation of the Key Lab of Automobile Materials, Jilin University, from Natural Scientific Basic Research Fund for Platform and Base Construction (Grant No. 09-421060352467) and by the Department of Science & Technology of Jilin Province (Grant No. 20100545). REFERENCES [1 ] R. Roger, E.W. Collings and G. Welsch: Materials Properties Handbook: Titanium Alloys, Materials Park, H, ASM International, 1994. [2 ] W.B. Hanson, K.I. Ironside and J.A. Fernie: Acta Mater., 2000, 48, 4673. [3 ] H. Hao, Y. Wang, Z. Jin and X. Wang: J. Am. Ceram. Soc., 1995, 78, 2157. [4 ] X.H. Wang and Y.C. Zhou: J. Mater. Sci. Technol., 2010, 26(5), 385. [5 ] M. L. Muolo, E. Ferrera, L. Morbelli and A. Passerone: Scripta Mater., 2004, 50, 325. [6 ]. Smorygo, J.S. Kim, M.D. Kim and T.G. Eom: Mater. Lett., 2007, 61, 613. [7 ] A.V. Durov, B.D. Kostjuk, A.V. Shevchenko and Y.V. Naidich: Mater. Sci. Eng. A, 2000, 290, 186. [8 ] A.V. Durov, Y.V. Naidich and B.D. Kostyuk: J. Mater. Sci., 2005, 40, 2173. [9 ] M. Singh, T.P. Shpargel and R. Asthana: Int. J. Appl. Ceram. Technol., 2007, 4, 119. [10] B.A. Kalin, V.T. Fedotov and A.E. Grigoriew: Fusion Eng. Des., 1995, 28, 119. [11] D. Szewieczek and J. Tyrlik: J. Mater. Process. Technol., 1995, 53, 405. [12] A. Rabinkin: Sci. Technol. Weld. Join., 2004, 9, 181. [13] M. Naka and I. kamoto: Trans. JWRI, 1985, 14, 185. [14] M. Singh, R. Asthana and T.P. Shpargel: Mater. Sci. Eng. A, 2008, 498, 19. [15] M. Singh and R. Asthana: Mater. Sci. Eng. A, 2007, 460,153. [16] Y.C. Kim, W.T. Kim and D.H. Kim: Mater. Sci. Eng. A, 2004, 375, 127. [17] J.S. Zou, Z.G. Jiang, Q.Z. Zhao and Z. Chen: Mater. Sci. Eng. A, 2009, 507, 155. [18] Z.Y. Wu, R.K. Shiue and C.S. Chang: J. Mater. Sci. Technol., 2010, 26(4), 311. [19] I.T. Hong and C.H. Koo: Int. J. Refarct. Met. H., 2006, 24, 247. [20] D. Sciti, A. Bellosi and L. Esposito: J. Eur. Ceram. Soc., 2001, 21, 45. [21]. Botstein and A. Rabinkin: Mater. Sci. Eng. A, 1994, 188, 305.