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Vacuum brazing of TiAl-based intermetallics with Ti–Zr–Cu–Ni–Co amorphous alloy as filler metal
Intermetallics 57 (2015) 7e16
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Vacuum brazing of TiAl-based intermetallics with TieZreCueNieCo amorphous alloy as filler metal Xiaoqiang Li, Li Li*, Ke Hu, Shengguan Qu National Engineering Research Center of Near-net-shape Forming for Metallic Materials, South China University of Technology, Guangzhou 510640, People's Republic of China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 26 May 2014 Received in revised form 24 September 2014 Accepted 26 September 2014 Available online
An amorphous Ti41.7eZr26.7eCu14.7eNi13.8eCo3.1 (wt%) ribbon fabricated by melt spinning was used as filler to vacuum braze Tie48Ale2Nbe2Cr (at%) intermetallics. The influences of brazing temperature and time on the microstructure and strength of the joints were investigated. It is found that intermetallic phases of Ti3Al and g-Ti2Cu/Ti2Ni form in the brazed joints. The tensile strength of the joint first increases and then decreases with the increase of the brazing temperature in the range of 900e1050 C and the brazing time varying from 3 to 15 min. The maximum tensile strength at room temperature is 316 MPa when the joint is brazed at 950 C for 5 min. Cleavage facets are widely observed on all of the fracture surfaces of the brazed joints. The fracture path varies with the brazing condition and cracks prefer to initiate at locations with relatively high content of g-Ti2Cu/Ti2Ni phases and propagate through them. © 2014 Elsevier Ltd. All rights reserved.
Keywords: A. Titanium aluminides, based on TiAl A. Amorphous metals B. Bonding B. Fracture C. Brazing D. Microstructure
1. Introduction Within the family of titanium aluminides, g-TiAl alloys are considered to be an ideal material for high temperature applications, such as moving parts of advanced automobile and aero engine components [1e3]. Compared with traditional titanium alloys, g-TiAl alloys possess lower density, higher strength-to-weight ratio, better corrosion and oxidation resistance, and higher mechanical properties at elevated temperature [3e5]. However, the effective utilization of g-TiAl alloys is delayed by its high cost, complex manufacturing process and intrinsic properties such as brittleness and poor workability [5,6]. Development of reliable techniques to join g-TiAl alloys to themselves or other materials is indispensable to extend the application of these alloys in real components. Fusion welding [7], friction welding [8], diffusion bonding [9] and brazing [10] have been used to join g-TiAl alloys. Among those, brazing as the most feasible and economical bonding technique, has received special attention in joining g-TiAl alloys. Selection of an appropriate filler metal to braze g-TiAl alloys plays a crucial role in obtaining satisfactory joints. Ag-based and Albased filler metals feature low liquidus temperature and are
* Corresponding author. Tel./fax: þ86 20 87112111. E-mail address: [email protected] (L. Li). http://dx.doi.org/10.1016/j.intermet.2014.09.010 0966-9795/© 2014 Elsevier Ltd. All rights reserved.
successfully applied to braze TiAl alloys. However, Ag-based and Albased filler metals suffer from insufficient bonding strength, weak corrosion and oxidation resistance, and low creep strength. Especially, the mechanical stability of the brazed joints badly degrades at temperature higher than 400 C [3,11e13]. Due to the good compatibility between Ti-based filler and TiAl substrate, the joints brazed by Ti-based filler commonly possess high bonding strength and good corrosion resistance at ambient and elevated temperature [5,14]. Nevertheless, titanium of the Ti-based filler metal has a high affinity to many interstitial elements such as oxygen and nitrogen, and easily forms oxide and nitride [15]. Therefore, the joining of TiAl alloys is always performed under inert gas or high-vacuum atmosphere. The most widely applied Ti-based brazing filler are TieCueNi alloys, in which Cu and Ni are added as melting point depressants. Brazing temperature of TieCueNi filler can be further lowered with the addition of Zr [16,17]. Recently, some amorphous Ti-based filler metals have been developed for brazing TiAl or Ti alloys to other materials, owing to their advantage of accelerated atomic diffusion, enhanced surface reaction, low brazing temperature and superior wettability [5,18,19]. Amorphous filler metals can facilitate the capillarity and shorten the gaps [20,21]. In such a case, the formation of voids or shrinkage cavities and residual stress in the joints is reduced, and the bonding strength of the joints will improve. He et al. brazed a Ti3Al-based alloy with amorphous Ti35eZr35eCu15eNi15 filler metal and obtained the maximum
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X. Li et al. / Intermetallics 57 (2015) 7e16
Fig. 1. Microstructure of the Tie48Ale2Nbe2Cr alloy.
shear joint strength of 259.6 MPa by brazing at 1323 K for 300 s [22,23]. However, no research work has been reported on the joining of TiAl alloys to themselves by using amorphous foils of the TieZreCueNi system. The present work attempts to develop an amorphous TieZreCueNieCo filler with low melting temperature for the brazing of g-TiAl alloys, and the focus is placed on the microstructure and mechanical properties of the brazed joints.
2. Experimental procedure 2.1. Materials The joined base-metal was a g-TiAl alloy with a nominal composition of Tie48Ale2Nbe2Cr (at%) (Supplied by Institute of Metal Research, Chinese Academy of Science). The microstructure
Fig. 3. (a) XRD pattern and (b) Bright-field TEM image inset with the corresponding SAED pattern of the TieZreCueNieCo filler foil.
of the g-TiAl alloy was fully lamellar, as shown in Fig. 1. The filler metal was a Ti41.7eZr26.7eCu14.7eNi13.8eCo3.1 (wt%) ribbon. It was prepared as follows: 1) first, arc-melting was done in a watercooled copper mold under high purity argon atmosphere; 2) the ingot was then homogenized at 950 C for 6 h to reduce segregation; 3) after induction melting in a quartz tube under Ar gas protection, rapid solidification was carried out by the single-roller melt spinning technique with a rotating Cu wheel at circumferential speed of 20.5 m/s. The resulting filler metal was flexible and had ribbon shape in 4e5 mm width and 50e60 mm thickness. The amorphous filler ribbon was characterized by an x-ray diffractometer (XRD) with Cu-Ka radiation and transmission
Fig. 2. (a) Schematic of assembly for brazing, (b) timeetemperature profile of the brazing cycle, and (c) Schematic of specimen for tensile test (mm).
X. Li et al. / Intermetallics 57 (2015) 7e16
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timeetemperature profile of the brazing cycle is shown in Fig. 2(b). Brazing temperature varied from 900 to 1050 C and holding time was 3e15 min. After brazing, the joints were longitudinally cut and the crosssections of the joints were examined by scanning electron microscopy (SEM). The elemental distribution in the brazed seam was detected by energy dispersive spectroscopy (EDS). For tensile test of brazed joint, the national standard of China (GB/T 11363-2008) was used [24] and the test pieces were processed from the brazed specimens, as shown in Fig. 2(c). The tensile strength of the joints was evaluated at room temperature using a Shimadzu AG-10 universal testing machine with a constant loading rate of 0.1 mm/min. After tensile test, SEM observation was used to examine the fracture surfaces and fracture path. EDS and XRD was used to identify the interfacial compounds on the fractured surfaces. 3. Results and discussion Fig. 4. DSC curve of the filler foil.
electron microscopy (TEM). The thermal behavior was measured by differential scanning calorimetry (DSC) at a heating rate of 20 K/ min. 2.2. Brazing Prior to brazing, the brazing surfaces of the Tie48Ale2Nbe2Cr specimens and the TieZreCueNieCo filler foil were both polished, and then degreased ultrasonically for 1200 s in a bath of acetone. After polishing, the thickness of the filler foil decreased to about 20 mm. Subsequently, the filler foil was sandwiched between two Tie48Ale2Nbe2Cr alloy rods with a TC4 clamp, as shown in Fig. 2(a). The assembly was then placed into a flat-bottomed corundum crucible. Furnace brazing was performed in an HP12 12 12 furnace with a vacuum of ~5 104 Pa. The
3.1. Microstructure of filler metal The XRD pattern of the TieZreCueNieCo filler foil is shown in Fig. 3(a). The diffraction pattern only consists of broad diffuse diffraction peak at a diffraction angle 2q of about 40 . No diffraction peaks corresponding to crystallization can be seen, indicating that the foil is in amorphous state. EDS analysis of ten random points clearly shows that the foil has good composition homogeneity. Besides, the SEM micrograph of the cross section of the foil reveals a featureless contrast in an etched state using a solution of HF. The detailed microstructure of the filler foil was further investigated by TEM, as shown in Fig. 3(b). No crystalline particles can be discovered. The entirely amorphous structure is affirmed by the selectedarea electron diffraction (SAED) pattern comprised only of a series of diffuse halo rings. Fig. 4 shows the DSC curve of the TieZreCueNieCo filler foil. The glass transition temperature (Tg) of 423 C and the
Fig. 5. Backscattered electron images (BEIs) of the joint brazed at 950 C for 3 min (a) interfacial microstructure, (b) and (c) high magnification BEIs of zone I and zone II in (a), respectively.
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Table 1 EDS results of the spots in Fig. 5(b) and (c). Position Element (at%) Ti A B C D E
Al
Nb
Phase Cr
Zr
Cu
Ni
61.16 26.58 1.73 1.02 2.74 3.38 2.93 60.22 22.27 1.16 0.62 6.43 4.30 4.21 42.02 6.82 0.23 0.26 18.83 15.40 13.77 62.01 4.78 0.23 0.22 10.63 10.24 8.73 57.54 2.45 0.10 0.27 10.08 8.77 16.16
slowly preheated to 500 C with holding time of 5 min, and then rapidly heated up to the scheduled brazing temperature, as shown in Fig. 2(b).
Co 0.46 0.79 2.67 3.16 4.63
Ti3Al Ti3Al (Ti, Zr)2(Cu, Ni) a-Ti þ (Ti, Zr)2(Cu, Ni) (Ti, Zr)2(Cu, Ni) þ a-Ti
crystallization temperature Tx (onset temperature of the first exothermic peak) of 575 C are marked by arrows on the trace. The solidus temperature (Ts) and the liquidus temperature (Tl) are 851 C and 875 C, respectively. On account of the ambiguity in melting temperature determination, the Tl is assumed to represent the melting temperature. Therefore, the brazing temperature in this study was chosen in the range of 900e1050 C, which was 25e175 C higher than the melting point of the filler metal. The transition behavior from the amorphous solid to the supercooled liquid will occur at near Tg. To minimize the impact of crystallization on the melting point of the filler foil, all the specimens were
3.2. Interfacial microstructure of the brazed joints Fig. 5 shows a typical backscattered electron image (BEI) of the joints brazed at 950 C for 3 min. It is clearly seen that the g-TiAl substrates are tightly bonded by the TieZreCueNieCo filler foil. No pores and cracks exist in the brazed seam. The width of the brazed seam was about 30 mm, thicker than the amorphous TieZreCueNieCo brazing foil, demonstrating that a strong interaction (including dissolution, diffusion and reaction) between the TiAl substrate and the molten brazed alloy occurred during brazing. Based on the microstructural morphology and the chemical composition, all the brazed joints can be classified into two reaction zones (zone I and zone II), as shown in Fig. 5(a). Fig. 5(b) and (c) display the highly magnified microstructures of zone I and zone II, respectively. Based on the difference in contrast, zone I mainly consists of two phases (as shown by marks of A and B), whereas zone II is a three-phase mixed region (marked by C, D and E for the corresponding phases/spots). Both the element pairs of Ti and Zr,
Fig. 6. Interfacial microstructures of the joints brazed at (a) 900 C, (b) 950 C, (c) 1000 C and (d) 1050 C for 15 min, high magnification BEIs of joints brazed at (e) 950 C and (f) 1000 C.
X. Li et al. / Intermetallics 57 (2015) 7e16
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Table 2 EDS results of the central brazed layer zone II in Fig. 6. Brazing temperature ( C)
Element (at%) Ti
900 950 1000 1050
Al
50.60 53.11 50.55 48.76
± ± ± ±
0.22 1.99 0.31 1.22
Nb
7.05 8.67 13.37 17.46
± ± ± ±
0.18 0.49 0.90 0.98
0.32 0.53 0.82 1.24
Cr ± ± ± ±
0.09 0.11 0.34 0.17
Zr
0.36 0.45 1.08 1.06
± ± ± ±
0.09 0.05 0.19 0.19
Cu
12.68 13.10 12.49 11.38
± ± ± ±
0.39 0.24 0.36 0.16
12.59 10.36 10.16 9.51
Ni ± ± ± ±
1.04 0.15 0.43 0.16
13.83 11.13 9.74 8.90
Co ± ± ± ±
1.62 0.37 0.41 0.27
2.54 2.65 1.79 1.69
± ± ± ±
0.20 0.11 0.17 0.07
Table 3 EDS results of the reaction layer zone I in Fig. 6. Brazing temperature ( C)
Element (at%) Ti
900 950 1000 1050
58.58 59.44 56.77 58.63
Phase Al
± ± ± ±
0.48 1.41 1.03 0.84
25.15 27.70 28.87 29.67
Nb ± ± ± ±
0.47 0.62 0.34 0.46
3.18 2.91 4.07 3.49
Cr ± ± ± ±
0.08 0.11 0.12 0.12
and Cu and Ni are not only chemically compatible (having similar atomic radii and crystal structure), but also completely soluble to each other [16]. Therefore, Zr and Ni in a certain sense can be regarded as Ti and Cu, respectively. The Zr2Cu and Zr2Ni can be regarded as Ti2Cu and Ti2Ni, even the four phases can be collectively referred to as (Ti, Zr)2(Cu, Ni). Table 1 shows the EDS analysis results of all the spots in Fig. 5(b) and (c). According to the elemental contents, Ti(Zr)eAleNi(Cu) and Ti(Zr)eCueNi ternary phase diagrams [25] and previous research [16,19], the phases in zone I are lamellar phase Ti3Al (spot A) and intermittent black blocky phase Ti3Al (spot B). Zone II consists of light grey phase (Ti, Zr)2(Cu, Ni) (spot C), mixed color blocky a-Ti phase with (Ti, Zr)2(Cu, Ni) (spot D) and dark grey phase (Ti, Zr)2(Cu, Ni) with a-Ti phase (spot E).
1.03 1.06 1.77 1.74
Zr ± ± ± ±
0.09 0.24 0.04 0.09
4.07 3.10 1.80 1.28
Cu ± ± ± ±
0.22 0.75 0.49 0.13
4.13 3.12 3.47 2.61
Ni ± ± ± ±
0.44 0.24 0.35 0.17
3.37 2.19 2.82 2.13
Co ± ± ± ±
0.13 0.18 0.54 0.36
0.49 0.48 0.43 0.45
± ± ± ±
0.08 0.08 0.09 0.10
Ti3Al Ti3Al Ti3Al Ti3Al
The characteristic microstructures are formed by atomic diffusion during the brazing, including diluting effect (especially during initial melting of the brazing alloy), isothermal solidification and solid-state diffusion between the TiAl substrate and each zone of the brazing seam. At the early stage of brazing, the TiAl substrate dissolves considerably into the molten TieZreCueNieCo filler, and the diffusion transport of Al, Nb and Cr elements from the base material into the brazed joint is driven by the concentration gradient. Therefore, the liquid pool initially grows wider, and then the width of the brazed seam is larger than the thickness of the filler foil. During brazing, the successive diffusion of Ti and Al atoms from the base-metal into the molten brazing alloy dilutes the concentration of Zr, Cu and Ni in the liquid pool. In addition, the diffusion of Zr, Cu and Ni elements of the filler into the TiAl
Fig. 7. Interfacial microstructures of the joints brazed at 950 C for (a) 3 min, (b) 5 min, (c) 7 min and (d) 10 min, respectively.
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substrate also reduce the concentration of Zr, Cu and Ni in the liquid pool. According to TieCu, TieNi and TieAl binary phase diagrams [26], the reduction of Cu and Ni and the increase of Al and Ti in the liquid pool lead to an increase in melting temperature of the liquid phase. So, the remaining liquid phase with relatively high Al content and low Cu and Ni concentrations initially re-solidifies even during the isothermal step. The maximum solubilities of Cu and Ni in b-Ti (13.5 and 10 at%, respectively) are much higher than those in a-Ti [14]. The content reduction of Cu and Ni as b-Ti stabilizers is favorable to the phase transformation of b / a-Ti during the cooling step. When the temperature is lower than the aeb transformation temperature, the eutectoid decomposition of b-Ti to a-Ti and g phases (b-Ti / a-Ti þ (Ti, Zr)2(Cu, Ni)) occurs, resulting in the absence of b-Ti phase in the brazed joint. Fig. 6 shows the BEIs of the joints brazed at different temperatures for 15 min. With the increase of brazing temperature, the microstructure of the joint markedly changes, which means that temperature plays an important role in brazing. When the TiAl alloy is brazed at 900 C and 950 C, a regular brazing seam can be obtained and a continuous fine lamellar interface reaction layer forms between the TiAl substrate and the brazed seam, as shown in Fig. 6(a) and (b). With the brazing temperature up to 1000 C and 1050 C, the intermittent black Ti3Al layer (marked as B in Fig. 5) disappears and the continuous lamellar reaction layer (marked as A in Fig. 5) becomes messy. Moreover, zone II transfers to a two-phase mixture (a2-Ti3Al and high Al% g-(Ti, Zr)2(Cu, Ni) phase), as shown in Fig. 6(c) and (d). EDS analysis results of the chemical composition in the central brazed layer zone II were listed in Table 2. In zone II, the content of melting point depressing elements (Ni and Cu) decreases with the increase of brazing temperature, whereas the content of Al increases. According to the studies of Lee et al. [27], the successive diffusion of Al atoms from the TiAl substrate into the joint was found to be the main controlling factor pertaining to the microstructure of the brazing seam. Therefore, remarkable change of microstructure takes place in the brazing seams, when the brazing temperature is 1000 C and 1050 C. Fig. 6(e) and (f) show the high magnification backscattered electron images of the reaction layer. Compared with the interface microstructure at 950 C (Fig. 6(e)), the morphology of the reaction layer is similar to the colony structure in the TiAl substrate when the brazing temperature is 1000 C (Fig. 6(f)). The phases and their composition in the reaction layer were indentified by EDS (results were listed in Table 3). As shown in Table 3, the phase in the reaction layer is Ti3Al regardless of the brazing temperature, and the phase composition is almost the same. We thus deduce the dissolution of lamellar colonies rather than different phase composition in the reaction layer leads to the irregular interface morphology when the brazing temperature is 1000 C and 1050 C. Based on the previous study in brazing of Tie45Ale5Nb-(W, B, Y) alloys by using TiNieNb braze alloy [6], a flat interface formed between the TiAl based alloy and the brazed seam, due to the dissolution of the TiAl based alloy into the molten filler metal in the form of single atomic layers at low temperature. However, at high brazing temperature, massive dissolution took place not only in the form of single atomic
Fig. 8. Effect of brazing temperature on tensile strength of the joints brazed at different temperatures for 15 min.
layers but also in the form whole lamellar colonies dissolving, leading to the irregular interface morphology [6]. This is in accordance with our experimental observation in the present work. Fig. 7 shows the microstructure of the joint brazed at 950 C for different time. The width of the brazed seam increases with the brazing time. Actually, within the brazed seam, the width of the central brazed layer zone II increases with the brazing time, whereas the thickness of the interfacial Ti3Al reaction layer Zone I basically remains unchanged regardless of the brazing time. Table 4 shows the EDS analysis of the chemical composition in the central brazed layer zone II. In zone II, the content of melting point depressing elements (Ni and Cu) decreases with the increase of brazing time, whereas the content of Al increases. In such a case, more interdiffusion and interaction between the braze zone and the TiAl substrate occur with the increase of the brazing time. The width of the zone II thus increases with the brazing time. On the other hand, the content of the light grey (Ti, Zr)2(Cu, Ni) phase increases with the brazing time, but the content of the mixed color blocky a-Ti phase with (Ti, Zr)2(Cu, Ni) decreases. The TiAl substrate is readily dissolved into the molten brazing alloy and thus there is no interfacial a2-Ti3Al layer to form at this time. It is documented that the formation of interfacial Ti3Al phase primarily results from the cooling cycle of brazing due to the limited solubility of Al in the a-Ti at low temperature [4]. Accordingly, the interfacial Ti3Al layer can not be removed by increasing the brazing time. Once the continuous Ti3Al phase layer forms in the joint interface, further atomic interdiffusion between the base-metal and the filler-metal will be obstructed due to the limited solubilities of Cu and Ni in the a2-Ti3Al ordered structure [27,28]. Hence, the thickness of the continuous Ti3Al layer is insensitive to the brazing time. In addition, there is no change in the microstructure of the TiAl substrate with increase of brazing time, i.e., the parent TiAl alloy still exhibits a lamellar structure without grain growth after brazing.
Table 4 EDS results of the central brazed layer zone II in Fig. 7. Brazing time (min)
Element (at%) Ti
3 5 7 10
51.99 51.43 53.72 53.88
Al ± ± ± ±
1.33 1.02 1.58 1.37
7.21 7.56 7.68 8.12
Nb ± ± ± ±
0.11 0.09 0.18 0.13
0.37 0.47 0.48 0.47
Cr ± ± ± ±
0.19 0.11 0.11 0.17
0.44 0.43 0.44 0.49
Zr ± ± ± ±
0.04 0.03 0.02 0.05
13.43 13.86 12.68 12.95
Cu ± ± ± ±
0.40 0.24 0.28 0.37
11.97 11.64 10.47 9.80
Ni ± ± ± ±
0.25 0.30 0.32 0.28
11.93 11.86 11.65 11.60
Co ± ± ± ±
0.18 0.20 0.29 0.19
2.66 2.75 2.88 2.69
± ± ± ±
0.11 0.16 0.06 0.14
X. Li et al. / Intermetallics 57 (2015) 7e16
3.3. Mechanical strength and fracture morphology of the brazed joint Fig. 8 shows the room temperature tensile strength of the joints brazed at different temperature for 15 min. For the brazing temperature of 900 C, the tensile strength is very low (only 140 MPa). The main cause is the low amount of atomic diffusion and the insufficient dissolution reaction between the TiAl substrate and the molten TieZreCueNieCo filler metal. It is documented that
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increasing the brazing temperature can improve dissolution and diffusion, leading to an enhanced metallurgical reaction [19]. Moreover, increasing brazing temperature could decrease the amount of residual brazing filler. Therefore, the maximum tensile strength of 217 MPa is obtained when the brazing temperature rises up to 950 C. Nevertheless, further increasing the brazing temperature again decreases the bonding strength, due to the different microstructure involved in the brazing seam. At the higher brazing temperature (1000 C and 1050 C), the formation of brittle Ti3Al
Fig. 9. Tensile fracture morphology and fracture paths of the joints brazed at (a) and 900 C, (c) and (d) 950 C, (e) and (f) 1000 C, (g) and (h) 1050 C for 15 min, respectively.
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X. Li et al. / Intermetallics 57 (2015) 7e16
Table 5 EDS results of the fracture surfaces in Fig. 9. Position
A B C D
Element (at%)
Phase
Ti
Al
Nb
Cr
Zr
Cu
Ni
Co
59.50 31.55 49.53 63.55
2.46 14.78 5.95 24.78
0.38 0.49 2.78 0
0.61 0.81 0.48 0.81
8.31 21.59 12.44 5.67
8.33 16.52 12.29 1.67
16.28 11.21 13.45 2.50
4.13 3.05 3.08 1.02
and the disappearance of plastic a-Ti in zone II (as shown in Fig. 6(c) and (d)) deteriorates the joints properties. Moreover, microcracks nucleation along the brittle phase boundaries (as shown by the arrows in Fig. 6(d)) resulted from residual thermal stress accumulation is also detrimental to the bonding strength. Therefore, the tensile strength decreases to the lowest value of 120 MPa when the brazing temperature up to 1050 C. In a word, the tensile strength first increases and then decreases with the increase of the brazing temperature. The same phenomenon was also found in previous research when TiAl alloy was brazed using TiNieNb eutectic braze alloy [6]. Fractographic analysis is a useful method to study the crack nucleation and propagation during tensile test. To reveal the fracture path from the tensile test, the tested specimens were reassembled and the cross section of the joint was observed by metallographic examination, as described in Ref. [29]. Fig. 9 shows the fracture paths and fracture morphology of the joints brazed at different brazing temperature for 15 min after the tensile test. At low brazing temperature (<1000 C), cracks primarily initiate and propagate in the central brazed layer (zone II), especially along the Ti2Cu þ Ti2Ni intermetallic compound clusters, as shown in Fig. 9(a) and (c). However, at the higher brazing temperature (1000 C and 1050 C), cracks nucleate in zone II, then deflect and penetrate into the boundary between the Ti3Al reaction layer (zone I) and the central brazed layer (Zone II) (see Fig. 9(e) and (g)). Typical brittle cleavage fracture characterized by cleavage facets can be observed on the fracture surfaces of all the TiAl/TieZreCueNieCo/TiAl joints. Based on the EDS chemical analysis results shown in Table 5, the Ti2Cu þ Ti2Ni intermetallic compounds occupy most of the fracture surfaces. XRD pattern of the fracture surfaces (as shown in Fig. 10) indicates that the peak intensities of the Ti2Ni and Ti2Cu phases are much higher than the others, which further confirms the results of Table 5. Hence, cracks are expected to nucleate in zone II and then
Fig. 10. XRD patterns of the fracture surfaces after tensile test for specimen brazed at 950 C for 15 min.
(Ti, Zr)2(Cu,Ni) þ a-Ti High Al% (Ti, Zr)2(Cu,Ni) (Ti, Zr)2(Cu, Ni) Ti3Al
grow on the boundary of the intermetallic compounds Ti2Cu þ Ti2Ni. Fig. 11 shows the room temperature tensile strength of the joints brazed at 950 C. Short brazing time (3 min) leads to residual brazing filler and insufficient atomic diffusion between the melting filler metal and the TiAl substrate, and thus the tensile strength is low (only 232 MPa). The tensile strength reaches the maximum value of 316 MPa when the brazing time is 5 min. Further prolonging the brazing time, the increase in the thickness of the brittle Ti2Cu þ Ti2Ni layer in zone II (seen in Fig. 7) leads to a decrease of tensile strength. The formation of excessive Ti2Cu þ Ti2Ni intermetallic compounds embrittle the joints and deteriorate the bonding strength. Moreover, with the increase of brittle product, the residual stress in the brazed joint accumulates, inducing the formation of microcracks. As a result, the joint strength decreases with the brazing time. In the literature about vacuum brazing of Ti3Al with TieZr35eNi15eCu15 [22], the thickness of the brittle Ti2Cu þ Ti2Ni layer increased with the brazing time, and the shear strength of the brazed joint initially increased and turned to decrease with the increase of the thickness of the brittle Ti2Cu þ Ti2Ni layer. The tensile fracture paths of the TiAl joints brazed at 950 C change with brazing time, as shown in Fig. 12. When brazing time is 3 min, the fracture occurs in the central brazed layer (zone II), as shown in Fig. 12(a). With the brazing time increasing, although the crack initially nucleates in zone II, it then deflects and propagates to the continuous reaction layer (zone I) and penetrates into the TiAl substrate, as shown in Fig. 12(c), (e) and (g). The mechanical properties of TiAl joint depend noticeably on the microstructure of the brazed joint. The distribution and content of the Ti2Cu þ Ti2Ni intermetallic phases in brazing seam affect the bonding strength, a moderate amount and uniform distribution is benefit to gain high bonding strength. TiAl materials' intrinsic brittleness and low constant bonding strength between the stable brittle Ti3Al reaction
Fig. 11. Effect of brazing time on tensile strength of the joints brazed at 950 C.
X. Li et al. / Intermetallics 57 (2015) 7e16
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Fig. 12. Fracture morphology and fracture paths of the joints after tensile test brazed at 950 C for (a) and (b) 3 min, (c) and (d) 5 min, (e) and (f) 7 min, (g) and (h) 10 min, respectively.
layer (zone I) and the TiAl substrate result in the crack deflection to the TiAl substrate when the joint is robust. An increase of the content of brittle g-Ti2Cu/Ti2Ni in zone II with brazing time and embrittles zone II and then causes a drop in mechanical strength again. Further prolonging the brazing time to 15 min, the amount of hard and brittle intermetallic compounds (Ti2Cu and Ti2Ni) becomes so high that the fracture nucleates and propagates along the zone II again (Fig. 9(c)). By comparing the interfacial microstructure and fracture locations in the joints brazed under different
conditions, it is concluded that the crack prefers to initiate at the locations with relatively high g-Ti2Cu/Ti2Ni content and spread through them, owing to the weaker strength. 4. Conclusion An amorphous TieZreCueNieCo alloy was designed and fabricated by vacuum arc remelting and rapid solidification, and then used as filler metal to vacuum braze the Tie48Ale2Nbe2Cr
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alloy. The influences of the brazing temperature and time on the microstructure and bonding strength of the brazed joints were investigated. Conclusions are summarized as follows: (1) Sound brazed joints of the TiAl alloy can be obtained with the amorphous TieZreCueNieCo filler foil. The typical brazed seam consists of reaction layer zone I and central brazed layer zone II. The intermetallic phases in zone I and zone II are Ti3Al and Ti2Cu þ Ti2Ni, respectively. (2) The interfacial morphology of the joints is determined by the dissolution way of the TiAl substrate into the molten brazing alloy, which is greatly affected by the brazing temperature. Great change takes place in the interfacial morphology when the brazing temperature exceeds 1000 C. (3) The bonding strength of the brazed joints first increases and then decreases with the increase of brazing temperature in the range of 900e1050 C and brazing time varying from 3 to 15 min. The maximum tensile strength at room temperature is 316 MPa when the joint is brazed at 950 C for 5 min. (4) The fracture morphology of the joints treated under different brazing conditions is characterized as brittle cleavage. The cracks prefer to initiate at the locations with relatively high g-Ti2Cu/Ti2Ni content and spread through them. Acknowledgments This topic of research was financed by the Research Project of Special Furnishment and Part (Grant No. XZJQ-B1120680), the Research Foundation of State Key Laboratory of Advanced Welding and Joining (Grant No. AWJ-Z14-02). References [1] Lin TS, Li HX, He P, Wei HM, Li L, Feng JC. Microstructure evolution and mechanical properties of transient liquid phase (TLP) bonded joints of TiAl intermetallics. Intermetallics 2013;37:59e64. [2] Li YL, Liu W, Sekulic DP, He P. Reactive wetting of AgCuTi filler metal on the TiAl-based alloy substrate. Appl Surf Sci 2012;259:343e8. [3] Shiue RK, Wu SK, Chen SY. Infrared brazing of TiAl using Al-based braze alloys. Intermetallics 2003;11:661e71. [4] Shiue RK, Wu SK, Chen SY, Shiue CY. Infrared brazing of Ti50Al50 and Tie6Ale4V using two Ti-based filler metals. Intermetallics 2008;16:1083e9. [5] Dong HG, Yang ZL, Yang GS, Dong C. Vacuum brazing of TiAl alloy to 40Cr steel with Ti60Ni22Cu10Zr8 alloy foil as filler metal. Mater Sci Eng A 2013;561: 252e8. [6] Song XG, Cao J, Liu YZ, Feng JC. Brazing high Nb containing TiAl alloy using TiNi-Nb eutectic braze alloy. Intermetallics 2012;22:136e41.
[7] Noda T, Shimizu T, Okabe M, Iikubo T. Joining of TiAl and steels by induction brazing. Mater Sci Eng A 1997;239-240:613e8. [8] Lee WB, Kim YJ, Jung SB. Effects of copper insert layer on the properties of friction welded joints between TiAl and AISI 4140 structural steel. Intermetallics 2004;12:671e8. [9] Cam G, Clemens H, Gerling R, Kocak M. Diffusion bonding of g-TiAl sheets. Intermetallics 1999;7:1025e31. [10] Song XG, Cao J, Chen HY, Wang YF, Feng JC. Brazing TiAl intermetallics using TiNieV eutectic brazing alloy. Mater Sci Eng A 2012;551:133e9. [11] He P, Feng JC, Xu W. Mechanical property of induction brazing TiAl-based intermetallics to steel 35CrMo using AgCuTi filler metal. Mater Sci Eng A 2006;418:45e52. [12] Song XG, Cao J, Chen HY, Wang HQ, Feng JC. Contact reactive brazing of Ti53311S alloy using Cu foil as interlayer: interfacial microstructure and joining properties. Mater Des 2013;46:895e901. [13] Liaw DW, Wu ZY, Shiue RK, Chang CS. Infrared vacuum brazing of Tie6Ale4V and Nb using the Tie15Cue15Ni foil. Mater Sci Eng A 2007;454-455:104e13. [14] Ganjeh E, Sarkhosh H, Bajgholi ME, Khorsand H, Ghaffari M. Increasing Tie6Ale4V brazed joint strength equal to the base metal by Ti and Zr amorphous filler alloys. Mater Charact 2012;71:31e40. [15] Chang CT, Shiue RK. Infrared brazing Tie6Ale4V and Mo using the Tie15Cue15Ni braze alloy. Int J Refract Met Hard Mater 2005;23:161e70. [16] Lee MK, Lee JG. Mechanical and corrosion properties of Tie6Ale4V alloy joints brazed with a low-melting-point 62.7Zre11.0Tie13.2Cue9.8Nie3.3Be amorphous filler metal. Mater Charact 2013;81:19e27. [17] Chang CT, Wu ZY, Shiue RK, Chang CS. Infrared brazing Tie6Ale4V and SP-700 alloys using the Tie20Zre20Cue20Ni braze alloy. Mater Lett 2007;61:842e5. [18] Liu YH, Hu JD, Shen P, Han XH, Li JC. Microstructural and mechanical properties of jointed ZrO2/Tie6Ale4V alloy using Ti33Zr17Cu50 amorphous brazing filler. Mater Des 2013;47:281e6. [19] Hong IT, Koo CH. Microstructural evolution and shear strength of brazing C103 and Tie6Ale4V using Tie20Cue20Nie20Zr (wt%) filler metal. Int J Refract Met Hard Mater 2006;24:247e52. [20] Hong IT, Koo CH. Vacuum-furnace brazing of C103 and Tie6Ale4V with Tie15Cue15Ni filler-metal. Mater Sci Eng A 2005;398:113e27. [21] Singh M, Asthana R. Joining of zirconium diboride-based ultra hightemperature ceramic composites using metallic glass interlayers. Mater Sci Eng A 2007;460-461:153e62. [22] He P, Feng JC, Zhou H. Microstructure and strength of brazed joints of Ti3Albase alloy with TiZrNiCu filler metal. Mater Sci Eng A 2005;392:81e6. [23] He P, Feng JC, Zhou H. Microstructure and strength of brazed joints of Ti3Albase alloy with different filler metals. Mater Charact 2005;54:338e46. [24] GB/T 11363-2008. Test method of the strength of brazed and soldered joint (ISO 5187:1985, NEQ). Beijing: China Standard Press; 2008. [25] Villars P, Prince A, Okamoto H. Handbook of ternary alloy phase diagrams. 1st ed. U.S: ASM International; 1995. [26] Massalski TB. Binary alloy phase diagrams. Materials Park (OH): ASM International; 1990. [27] Lee SJ, Wu SK. Infrared joining strength and interfacial microstructures of Ti48Al-2Nb-2Cr intermetallics using Ti-15Cu-15Ni foil. Intermetallics 1999;7: 11e21. [28] Shiue RK, Wu SK, Chen YT. Strong bonding of infrared brazed a2-Ti3Al and Tie6Ale4V using TieCueNi fillers. Intermetallics 2010;18:107e14. [29] Ganjeh E, Sarkhosh H. Microstructural, mechanical and fractographical study of titanium-CP and Tie6Ale4V similar brazing with Ti-based filler. Mater Sci Eng A 2013;559:119e29.
Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Vacuum brazing of TiAl-based intermetallics with TieZreCueNieCo amorphous alloy as filler metal Xiaoqiang Li, Li Li*, Ke Hu, Shengguan Qu National Engineering Research Center of Near-net-shape Forming for Metallic Materials, South China University of Technology, Guangzhou 510640, People's Republic of China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 26 May 2014 Received in revised form 24 September 2014 Accepted 26 September 2014 Available online
An amorphous Ti41.7eZr26.7eCu14.7eNi13.8eCo3.1 (wt%) ribbon fabricated by melt spinning was used as filler to vacuum braze Tie48Ale2Nbe2Cr (at%) intermetallics. The influences of brazing temperature and time on the microstructure and strength of the joints were investigated. It is found that intermetallic phases of Ti3Al and g-Ti2Cu/Ti2Ni form in the brazed joints. The tensile strength of the joint first increases and then decreases with the increase of the brazing temperature in the range of 900e1050 C and the brazing time varying from 3 to 15 min. The maximum tensile strength at room temperature is 316 MPa when the joint is brazed at 950 C for 5 min. Cleavage facets are widely observed on all of the fracture surfaces of the brazed joints. The fracture path varies with the brazing condition and cracks prefer to initiate at locations with relatively high content of g-Ti2Cu/Ti2Ni phases and propagate through them. © 2014 Elsevier Ltd. All rights reserved.
Keywords: A. Titanium aluminides, based on TiAl A. Amorphous metals B. Bonding B. Fracture C. Brazing D. Microstructure
1. Introduction Within the family of titanium aluminides, g-TiAl alloys are considered to be an ideal material for high temperature applications, such as moving parts of advanced automobile and aero engine components [1e3]. Compared with traditional titanium alloys, g-TiAl alloys possess lower density, higher strength-to-weight ratio, better corrosion and oxidation resistance, and higher mechanical properties at elevated temperature [3e5]. However, the effective utilization of g-TiAl alloys is delayed by its high cost, complex manufacturing process and intrinsic properties such as brittleness and poor workability [5,6]. Development of reliable techniques to join g-TiAl alloys to themselves or other materials is indispensable to extend the application of these alloys in real components. Fusion welding [7], friction welding [8], diffusion bonding [9] and brazing [10] have been used to join g-TiAl alloys. Among those, brazing as the most feasible and economical bonding technique, has received special attention in joining g-TiAl alloys. Selection of an appropriate filler metal to braze g-TiAl alloys plays a crucial role in obtaining satisfactory joints. Ag-based and Albased filler metals feature low liquidus temperature and are
* Corresponding author. Tel./fax: þ86 20 87112111. E-mail address: [email protected] (L. Li). http://dx.doi.org/10.1016/j.intermet.2014.09.010 0966-9795/© 2014 Elsevier Ltd. All rights reserved.
successfully applied to braze TiAl alloys. However, Ag-based and Albased filler metals suffer from insufficient bonding strength, weak corrosion and oxidation resistance, and low creep strength. Especially, the mechanical stability of the brazed joints badly degrades at temperature higher than 400 C [3,11e13]. Due to the good compatibility between Ti-based filler and TiAl substrate, the joints brazed by Ti-based filler commonly possess high bonding strength and good corrosion resistance at ambient and elevated temperature [5,14]. Nevertheless, titanium of the Ti-based filler metal has a high affinity to many interstitial elements such as oxygen and nitrogen, and easily forms oxide and nitride [15]. Therefore, the joining of TiAl alloys is always performed under inert gas or high-vacuum atmosphere. The most widely applied Ti-based brazing filler are TieCueNi alloys, in which Cu and Ni are added as melting point depressants. Brazing temperature of TieCueNi filler can be further lowered with the addition of Zr [16,17]. Recently, some amorphous Ti-based filler metals have been developed for brazing TiAl or Ti alloys to other materials, owing to their advantage of accelerated atomic diffusion, enhanced surface reaction, low brazing temperature and superior wettability [5,18,19]. Amorphous filler metals can facilitate the capillarity and shorten the gaps [20,21]. In such a case, the formation of voids or shrinkage cavities and residual stress in the joints is reduced, and the bonding strength of the joints will improve. He et al. brazed a Ti3Al-based alloy with amorphous Ti35eZr35eCu15eNi15 filler metal and obtained the maximum
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X. Li et al. / Intermetallics 57 (2015) 7e16
Fig. 1. Microstructure of the Tie48Ale2Nbe2Cr alloy.
shear joint strength of 259.6 MPa by brazing at 1323 K for 300 s [22,23]. However, no research work has been reported on the joining of TiAl alloys to themselves by using amorphous foils of the TieZreCueNi system. The present work attempts to develop an amorphous TieZreCueNieCo filler with low melting temperature for the brazing of g-TiAl alloys, and the focus is placed on the microstructure and mechanical properties of the brazed joints.
2. Experimental procedure 2.1. Materials The joined base-metal was a g-TiAl alloy with a nominal composition of Tie48Ale2Nbe2Cr (at%) (Supplied by Institute of Metal Research, Chinese Academy of Science). The microstructure
Fig. 3. (a) XRD pattern and (b) Bright-field TEM image inset with the corresponding SAED pattern of the TieZreCueNieCo filler foil.
of the g-TiAl alloy was fully lamellar, as shown in Fig. 1. The filler metal was a Ti41.7eZr26.7eCu14.7eNi13.8eCo3.1 (wt%) ribbon. It was prepared as follows: 1) first, arc-melting was done in a watercooled copper mold under high purity argon atmosphere; 2) the ingot was then homogenized at 950 C for 6 h to reduce segregation; 3) after induction melting in a quartz tube under Ar gas protection, rapid solidification was carried out by the single-roller melt spinning technique with a rotating Cu wheel at circumferential speed of 20.5 m/s. The resulting filler metal was flexible and had ribbon shape in 4e5 mm width and 50e60 mm thickness. The amorphous filler ribbon was characterized by an x-ray diffractometer (XRD) with Cu-Ka radiation and transmission
Fig. 2. (a) Schematic of assembly for brazing, (b) timeetemperature profile of the brazing cycle, and (c) Schematic of specimen for tensile test (mm).
X. Li et al. / Intermetallics 57 (2015) 7e16
9
timeetemperature profile of the brazing cycle is shown in Fig. 2(b). Brazing temperature varied from 900 to 1050 C and holding time was 3e15 min. After brazing, the joints were longitudinally cut and the crosssections of the joints were examined by scanning electron microscopy (SEM). The elemental distribution in the brazed seam was detected by energy dispersive spectroscopy (EDS). For tensile test of brazed joint, the national standard of China (GB/T 11363-2008) was used [24] and the test pieces were processed from the brazed specimens, as shown in Fig. 2(c). The tensile strength of the joints was evaluated at room temperature using a Shimadzu AG-10 universal testing machine with a constant loading rate of 0.1 mm/min. After tensile test, SEM observation was used to examine the fracture surfaces and fracture path. EDS and XRD was used to identify the interfacial compounds on the fractured surfaces. 3. Results and discussion Fig. 4. DSC curve of the filler foil.
electron microscopy (TEM). The thermal behavior was measured by differential scanning calorimetry (DSC) at a heating rate of 20 K/ min. 2.2. Brazing Prior to brazing, the brazing surfaces of the Tie48Ale2Nbe2Cr specimens and the TieZreCueNieCo filler foil were both polished, and then degreased ultrasonically for 1200 s in a bath of acetone. After polishing, the thickness of the filler foil decreased to about 20 mm. Subsequently, the filler foil was sandwiched between two Tie48Ale2Nbe2Cr alloy rods with a TC4 clamp, as shown in Fig. 2(a). The assembly was then placed into a flat-bottomed corundum crucible. Furnace brazing was performed in an HP12 12 12 furnace with a vacuum of ~5 104 Pa. The
3.1. Microstructure of filler metal The XRD pattern of the TieZreCueNieCo filler foil is shown in Fig. 3(a). The diffraction pattern only consists of broad diffuse diffraction peak at a diffraction angle 2q of about 40 . No diffraction peaks corresponding to crystallization can be seen, indicating that the foil is in amorphous state. EDS analysis of ten random points clearly shows that the foil has good composition homogeneity. Besides, the SEM micrograph of the cross section of the foil reveals a featureless contrast in an etched state using a solution of HF. The detailed microstructure of the filler foil was further investigated by TEM, as shown in Fig. 3(b). No crystalline particles can be discovered. The entirely amorphous structure is affirmed by the selectedarea electron diffraction (SAED) pattern comprised only of a series of diffuse halo rings. Fig. 4 shows the DSC curve of the TieZreCueNieCo filler foil. The glass transition temperature (Tg) of 423 C and the
Fig. 5. Backscattered electron images (BEIs) of the joint brazed at 950 C for 3 min (a) interfacial microstructure, (b) and (c) high magnification BEIs of zone I and zone II in (a), respectively.
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X. Li et al. / Intermetallics 57 (2015) 7e16
Table 1 EDS results of the spots in Fig. 5(b) and (c). Position Element (at%) Ti A B C D E
Al
Nb
Phase Cr
Zr
Cu
Ni
61.16 26.58 1.73 1.02 2.74 3.38 2.93 60.22 22.27 1.16 0.62 6.43 4.30 4.21 42.02 6.82 0.23 0.26 18.83 15.40 13.77 62.01 4.78 0.23 0.22 10.63 10.24 8.73 57.54 2.45 0.10 0.27 10.08 8.77 16.16
slowly preheated to 500 C with holding time of 5 min, and then rapidly heated up to the scheduled brazing temperature, as shown in Fig. 2(b).
Co 0.46 0.79 2.67 3.16 4.63
Ti3Al Ti3Al (Ti, Zr)2(Cu, Ni) a-Ti þ (Ti, Zr)2(Cu, Ni) (Ti, Zr)2(Cu, Ni) þ a-Ti
crystallization temperature Tx (onset temperature of the first exothermic peak) of 575 C are marked by arrows on the trace. The solidus temperature (Ts) and the liquidus temperature (Tl) are 851 C and 875 C, respectively. On account of the ambiguity in melting temperature determination, the Tl is assumed to represent the melting temperature. Therefore, the brazing temperature in this study was chosen in the range of 900e1050 C, which was 25e175 C higher than the melting point of the filler metal. The transition behavior from the amorphous solid to the supercooled liquid will occur at near Tg. To minimize the impact of crystallization on the melting point of the filler foil, all the specimens were
3.2. Interfacial microstructure of the brazed joints Fig. 5 shows a typical backscattered electron image (BEI) of the joints brazed at 950 C for 3 min. It is clearly seen that the g-TiAl substrates are tightly bonded by the TieZreCueNieCo filler foil. No pores and cracks exist in the brazed seam. The width of the brazed seam was about 30 mm, thicker than the amorphous TieZreCueNieCo brazing foil, demonstrating that a strong interaction (including dissolution, diffusion and reaction) between the TiAl substrate and the molten brazed alloy occurred during brazing. Based on the microstructural morphology and the chemical composition, all the brazed joints can be classified into two reaction zones (zone I and zone II), as shown in Fig. 5(a). Fig. 5(b) and (c) display the highly magnified microstructures of zone I and zone II, respectively. Based on the difference in contrast, zone I mainly consists of two phases (as shown by marks of A and B), whereas zone II is a three-phase mixed region (marked by C, D and E for the corresponding phases/spots). Both the element pairs of Ti and Zr,
Fig. 6. Interfacial microstructures of the joints brazed at (a) 900 C, (b) 950 C, (c) 1000 C and (d) 1050 C for 15 min, high magnification BEIs of joints brazed at (e) 950 C and (f) 1000 C.
X. Li et al. / Intermetallics 57 (2015) 7e16
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Table 2 EDS results of the central brazed layer zone II in Fig. 6. Brazing temperature ( C)
Element (at%) Ti
900 950 1000 1050
Al
50.60 53.11 50.55 48.76
± ± ± ±
0.22 1.99 0.31 1.22
Nb
7.05 8.67 13.37 17.46
± ± ± ±
0.18 0.49 0.90 0.98
0.32 0.53 0.82 1.24
Cr ± ± ± ±
0.09 0.11 0.34 0.17
Zr
0.36 0.45 1.08 1.06
± ± ± ±
0.09 0.05 0.19 0.19
Cu
12.68 13.10 12.49 11.38
± ± ± ±
0.39 0.24 0.36 0.16
12.59 10.36 10.16 9.51
Ni ± ± ± ±
1.04 0.15 0.43 0.16
13.83 11.13 9.74 8.90
Co ± ± ± ±
1.62 0.37 0.41 0.27
2.54 2.65 1.79 1.69
± ± ± ±
0.20 0.11 0.17 0.07
Table 3 EDS results of the reaction layer zone I in Fig. 6. Brazing temperature ( C)
Element (at%) Ti
900 950 1000 1050
58.58 59.44 56.77 58.63
Phase Al
± ± ± ±
0.48 1.41 1.03 0.84
25.15 27.70 28.87 29.67
Nb ± ± ± ±
0.47 0.62 0.34 0.46
3.18 2.91 4.07 3.49
Cr ± ± ± ±
0.08 0.11 0.12 0.12
and Cu and Ni are not only chemically compatible (having similar atomic radii and crystal structure), but also completely soluble to each other [16]. Therefore, Zr and Ni in a certain sense can be regarded as Ti and Cu, respectively. The Zr2Cu and Zr2Ni can be regarded as Ti2Cu and Ti2Ni, even the four phases can be collectively referred to as (Ti, Zr)2(Cu, Ni). Table 1 shows the EDS analysis results of all the spots in Fig. 5(b) and (c). According to the elemental contents, Ti(Zr)eAleNi(Cu) and Ti(Zr)eCueNi ternary phase diagrams [25] and previous research [16,19], the phases in zone I are lamellar phase Ti3Al (spot A) and intermittent black blocky phase Ti3Al (spot B). Zone II consists of light grey phase (Ti, Zr)2(Cu, Ni) (spot C), mixed color blocky a-Ti phase with (Ti, Zr)2(Cu, Ni) (spot D) and dark grey phase (Ti, Zr)2(Cu, Ni) with a-Ti phase (spot E).
1.03 1.06 1.77 1.74
Zr ± ± ± ±
0.09 0.24 0.04 0.09
4.07 3.10 1.80 1.28
Cu ± ± ± ±
0.22 0.75 0.49 0.13
4.13 3.12 3.47 2.61
Ni ± ± ± ±
0.44 0.24 0.35 0.17
3.37 2.19 2.82 2.13
Co ± ± ± ±
0.13 0.18 0.54 0.36
0.49 0.48 0.43 0.45
± ± ± ±
0.08 0.08 0.09 0.10
Ti3Al Ti3Al Ti3Al Ti3Al
The characteristic microstructures are formed by atomic diffusion during the brazing, including diluting effect (especially during initial melting of the brazing alloy), isothermal solidification and solid-state diffusion between the TiAl substrate and each zone of the brazing seam. At the early stage of brazing, the TiAl substrate dissolves considerably into the molten TieZreCueNieCo filler, and the diffusion transport of Al, Nb and Cr elements from the base material into the brazed joint is driven by the concentration gradient. Therefore, the liquid pool initially grows wider, and then the width of the brazed seam is larger than the thickness of the filler foil. During brazing, the successive diffusion of Ti and Al atoms from the base-metal into the molten brazing alloy dilutes the concentration of Zr, Cu and Ni in the liquid pool. In addition, the diffusion of Zr, Cu and Ni elements of the filler into the TiAl
Fig. 7. Interfacial microstructures of the joints brazed at 950 C for (a) 3 min, (b) 5 min, (c) 7 min and (d) 10 min, respectively.
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X. Li et al. / Intermetallics 57 (2015) 7e16
substrate also reduce the concentration of Zr, Cu and Ni in the liquid pool. According to TieCu, TieNi and TieAl binary phase diagrams [26], the reduction of Cu and Ni and the increase of Al and Ti in the liquid pool lead to an increase in melting temperature of the liquid phase. So, the remaining liquid phase with relatively high Al content and low Cu and Ni concentrations initially re-solidifies even during the isothermal step. The maximum solubilities of Cu and Ni in b-Ti (13.5 and 10 at%, respectively) are much higher than those in a-Ti [14]. The content reduction of Cu and Ni as b-Ti stabilizers is favorable to the phase transformation of b / a-Ti during the cooling step. When the temperature is lower than the aeb transformation temperature, the eutectoid decomposition of b-Ti to a-Ti and g phases (b-Ti / a-Ti þ (Ti, Zr)2(Cu, Ni)) occurs, resulting in the absence of b-Ti phase in the brazed joint. Fig. 6 shows the BEIs of the joints brazed at different temperatures for 15 min. With the increase of brazing temperature, the microstructure of the joint markedly changes, which means that temperature plays an important role in brazing. When the TiAl alloy is brazed at 900 C and 950 C, a regular brazing seam can be obtained and a continuous fine lamellar interface reaction layer forms between the TiAl substrate and the brazed seam, as shown in Fig. 6(a) and (b). With the brazing temperature up to 1000 C and 1050 C, the intermittent black Ti3Al layer (marked as B in Fig. 5) disappears and the continuous lamellar reaction layer (marked as A in Fig. 5) becomes messy. Moreover, zone II transfers to a two-phase mixture (a2-Ti3Al and high Al% g-(Ti, Zr)2(Cu, Ni) phase), as shown in Fig. 6(c) and (d). EDS analysis results of the chemical composition in the central brazed layer zone II were listed in Table 2. In zone II, the content of melting point depressing elements (Ni and Cu) decreases with the increase of brazing temperature, whereas the content of Al increases. According to the studies of Lee et al. [27], the successive diffusion of Al atoms from the TiAl substrate into the joint was found to be the main controlling factor pertaining to the microstructure of the brazing seam. Therefore, remarkable change of microstructure takes place in the brazing seams, when the brazing temperature is 1000 C and 1050 C. Fig. 6(e) and (f) show the high magnification backscattered electron images of the reaction layer. Compared with the interface microstructure at 950 C (Fig. 6(e)), the morphology of the reaction layer is similar to the colony structure in the TiAl substrate when the brazing temperature is 1000 C (Fig. 6(f)). The phases and their composition in the reaction layer were indentified by EDS (results were listed in Table 3). As shown in Table 3, the phase in the reaction layer is Ti3Al regardless of the brazing temperature, and the phase composition is almost the same. We thus deduce the dissolution of lamellar colonies rather than different phase composition in the reaction layer leads to the irregular interface morphology when the brazing temperature is 1000 C and 1050 C. Based on the previous study in brazing of Tie45Ale5Nb-(W, B, Y) alloys by using TiNieNb braze alloy [6], a flat interface formed between the TiAl based alloy and the brazed seam, due to the dissolution of the TiAl based alloy into the molten filler metal in the form of single atomic layers at low temperature. However, at high brazing temperature, massive dissolution took place not only in the form of single atomic
Fig. 8. Effect of brazing temperature on tensile strength of the joints brazed at different temperatures for 15 min.
layers but also in the form whole lamellar colonies dissolving, leading to the irregular interface morphology [6]. This is in accordance with our experimental observation in the present work. Fig. 7 shows the microstructure of the joint brazed at 950 C for different time. The width of the brazed seam increases with the brazing time. Actually, within the brazed seam, the width of the central brazed layer zone II increases with the brazing time, whereas the thickness of the interfacial Ti3Al reaction layer Zone I basically remains unchanged regardless of the brazing time. Table 4 shows the EDS analysis of the chemical composition in the central brazed layer zone II. In zone II, the content of melting point depressing elements (Ni and Cu) decreases with the increase of brazing time, whereas the content of Al increases. In such a case, more interdiffusion and interaction between the braze zone and the TiAl substrate occur with the increase of the brazing time. The width of the zone II thus increases with the brazing time. On the other hand, the content of the light grey (Ti, Zr)2(Cu, Ni) phase increases with the brazing time, but the content of the mixed color blocky a-Ti phase with (Ti, Zr)2(Cu, Ni) decreases. The TiAl substrate is readily dissolved into the molten brazing alloy and thus there is no interfacial a2-Ti3Al layer to form at this time. It is documented that the formation of interfacial Ti3Al phase primarily results from the cooling cycle of brazing due to the limited solubility of Al in the a-Ti at low temperature [4]. Accordingly, the interfacial Ti3Al layer can not be removed by increasing the brazing time. Once the continuous Ti3Al phase layer forms in the joint interface, further atomic interdiffusion between the base-metal and the filler-metal will be obstructed due to the limited solubilities of Cu and Ni in the a2-Ti3Al ordered structure [27,28]. Hence, the thickness of the continuous Ti3Al layer is insensitive to the brazing time. In addition, there is no change in the microstructure of the TiAl substrate with increase of brazing time, i.e., the parent TiAl alloy still exhibits a lamellar structure without grain growth after brazing.
Table 4 EDS results of the central brazed layer zone II in Fig. 7. Brazing time (min)
Element (at%) Ti
3 5 7 10
51.99 51.43 53.72 53.88
Al ± ± ± ±
1.33 1.02 1.58 1.37
7.21 7.56 7.68 8.12
Nb ± ± ± ±
0.11 0.09 0.18 0.13
0.37 0.47 0.48 0.47
Cr ± ± ± ±
0.19 0.11 0.11 0.17
0.44 0.43 0.44 0.49
Zr ± ± ± ±
0.04 0.03 0.02 0.05
13.43 13.86 12.68 12.95
Cu ± ± ± ±
0.40 0.24 0.28 0.37
11.97 11.64 10.47 9.80
Ni ± ± ± ±
0.25 0.30 0.32 0.28
11.93 11.86 11.65 11.60
Co ± ± ± ±
0.18 0.20 0.29 0.19
2.66 2.75 2.88 2.69
± ± ± ±
0.11 0.16 0.06 0.14
X. Li et al. / Intermetallics 57 (2015) 7e16
3.3. Mechanical strength and fracture morphology of the brazed joint Fig. 8 shows the room temperature tensile strength of the joints brazed at different temperature for 15 min. For the brazing temperature of 900 C, the tensile strength is very low (only 140 MPa). The main cause is the low amount of atomic diffusion and the insufficient dissolution reaction between the TiAl substrate and the molten TieZreCueNieCo filler metal. It is documented that
13
increasing the brazing temperature can improve dissolution and diffusion, leading to an enhanced metallurgical reaction [19]. Moreover, increasing brazing temperature could decrease the amount of residual brazing filler. Therefore, the maximum tensile strength of 217 MPa is obtained when the brazing temperature rises up to 950 C. Nevertheless, further increasing the brazing temperature again decreases the bonding strength, due to the different microstructure involved in the brazing seam. At the higher brazing temperature (1000 C and 1050 C), the formation of brittle Ti3Al
Fig. 9. Tensile fracture morphology and fracture paths of the joints brazed at (a) and 900 C, (c) and (d) 950 C, (e) and (f) 1000 C, (g) and (h) 1050 C for 15 min, respectively.
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X. Li et al. / Intermetallics 57 (2015) 7e16
Table 5 EDS results of the fracture surfaces in Fig. 9. Position
A B C D
Element (at%)
Phase
Ti
Al
Nb
Cr
Zr
Cu
Ni
Co
59.50 31.55 49.53 63.55
2.46 14.78 5.95 24.78
0.38 0.49 2.78 0
0.61 0.81 0.48 0.81
8.31 21.59 12.44 5.67
8.33 16.52 12.29 1.67
16.28 11.21 13.45 2.50
4.13 3.05 3.08 1.02
and the disappearance of plastic a-Ti in zone II (as shown in Fig. 6(c) and (d)) deteriorates the joints properties. Moreover, microcracks nucleation along the brittle phase boundaries (as shown by the arrows in Fig. 6(d)) resulted from residual thermal stress accumulation is also detrimental to the bonding strength. Therefore, the tensile strength decreases to the lowest value of 120 MPa when the brazing temperature up to 1050 C. In a word, the tensile strength first increases and then decreases with the increase of the brazing temperature. The same phenomenon was also found in previous research when TiAl alloy was brazed using TiNieNb eutectic braze alloy [6]. Fractographic analysis is a useful method to study the crack nucleation and propagation during tensile test. To reveal the fracture path from the tensile test, the tested specimens were reassembled and the cross section of the joint was observed by metallographic examination, as described in Ref. [29]. Fig. 9 shows the fracture paths and fracture morphology of the joints brazed at different brazing temperature for 15 min after the tensile test. At low brazing temperature (<1000 C), cracks primarily initiate and propagate in the central brazed layer (zone II), especially along the Ti2Cu þ Ti2Ni intermetallic compound clusters, as shown in Fig. 9(a) and (c). However, at the higher brazing temperature (1000 C and 1050 C), cracks nucleate in zone II, then deflect and penetrate into the boundary between the Ti3Al reaction layer (zone I) and the central brazed layer (Zone II) (see Fig. 9(e) and (g)). Typical brittle cleavage fracture characterized by cleavage facets can be observed on the fracture surfaces of all the TiAl/TieZreCueNieCo/TiAl joints. Based on the EDS chemical analysis results shown in Table 5, the Ti2Cu þ Ti2Ni intermetallic compounds occupy most of the fracture surfaces. XRD pattern of the fracture surfaces (as shown in Fig. 10) indicates that the peak intensities of the Ti2Ni and Ti2Cu phases are much higher than the others, which further confirms the results of Table 5. Hence, cracks are expected to nucleate in zone II and then
Fig. 10. XRD patterns of the fracture surfaces after tensile test for specimen brazed at 950 C for 15 min.
(Ti, Zr)2(Cu,Ni) þ a-Ti High Al% (Ti, Zr)2(Cu,Ni) (Ti, Zr)2(Cu, Ni) Ti3Al
grow on the boundary of the intermetallic compounds Ti2Cu þ Ti2Ni. Fig. 11 shows the room temperature tensile strength of the joints brazed at 950 C. Short brazing time (3 min) leads to residual brazing filler and insufficient atomic diffusion between the melting filler metal and the TiAl substrate, and thus the tensile strength is low (only 232 MPa). The tensile strength reaches the maximum value of 316 MPa when the brazing time is 5 min. Further prolonging the brazing time, the increase in the thickness of the brittle Ti2Cu þ Ti2Ni layer in zone II (seen in Fig. 7) leads to a decrease of tensile strength. The formation of excessive Ti2Cu þ Ti2Ni intermetallic compounds embrittle the joints and deteriorate the bonding strength. Moreover, with the increase of brittle product, the residual stress in the brazed joint accumulates, inducing the formation of microcracks. As a result, the joint strength decreases with the brazing time. In the literature about vacuum brazing of Ti3Al with TieZr35eNi15eCu15 [22], the thickness of the brittle Ti2Cu þ Ti2Ni layer increased with the brazing time, and the shear strength of the brazed joint initially increased and turned to decrease with the increase of the thickness of the brittle Ti2Cu þ Ti2Ni layer. The tensile fracture paths of the TiAl joints brazed at 950 C change with brazing time, as shown in Fig. 12. When brazing time is 3 min, the fracture occurs in the central brazed layer (zone II), as shown in Fig. 12(a). With the brazing time increasing, although the crack initially nucleates in zone II, it then deflects and propagates to the continuous reaction layer (zone I) and penetrates into the TiAl substrate, as shown in Fig. 12(c), (e) and (g). The mechanical properties of TiAl joint depend noticeably on the microstructure of the brazed joint. The distribution and content of the Ti2Cu þ Ti2Ni intermetallic phases in brazing seam affect the bonding strength, a moderate amount and uniform distribution is benefit to gain high bonding strength. TiAl materials' intrinsic brittleness and low constant bonding strength between the stable brittle Ti3Al reaction
Fig. 11. Effect of brazing time on tensile strength of the joints brazed at 950 C.
X. Li et al. / Intermetallics 57 (2015) 7e16
15
Fig. 12. Fracture morphology and fracture paths of the joints after tensile test brazed at 950 C for (a) and (b) 3 min, (c) and (d) 5 min, (e) and (f) 7 min, (g) and (h) 10 min, respectively.
layer (zone I) and the TiAl substrate result in the crack deflection to the TiAl substrate when the joint is robust. An increase of the content of brittle g-Ti2Cu/Ti2Ni in zone II with brazing time and embrittles zone II and then causes a drop in mechanical strength again. Further prolonging the brazing time to 15 min, the amount of hard and brittle intermetallic compounds (Ti2Cu and Ti2Ni) becomes so high that the fracture nucleates and propagates along the zone II again (Fig. 9(c)). By comparing the interfacial microstructure and fracture locations in the joints brazed under different
conditions, it is concluded that the crack prefers to initiate at the locations with relatively high g-Ti2Cu/Ti2Ni content and spread through them, owing to the weaker strength. 4. Conclusion An amorphous TieZreCueNieCo alloy was designed and fabricated by vacuum arc remelting and rapid solidification, and then used as filler metal to vacuum braze the Tie48Ale2Nbe2Cr
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X. Li et al. / Intermetallics 57 (2015) 7e16
alloy. The influences of the brazing temperature and time on the microstructure and bonding strength of the brazed joints were investigated. Conclusions are summarized as follows: (1) Sound brazed joints of the TiAl alloy can be obtained with the amorphous TieZreCueNieCo filler foil. The typical brazed seam consists of reaction layer zone I and central brazed layer zone II. The intermetallic phases in zone I and zone II are Ti3Al and Ti2Cu þ Ti2Ni, respectively. (2) The interfacial morphology of the joints is determined by the dissolution way of the TiAl substrate into the molten brazing alloy, which is greatly affected by the brazing temperature. Great change takes place in the interfacial morphology when the brazing temperature exceeds 1000 C. (3) The bonding strength of the brazed joints first increases and then decreases with the increase of brazing temperature in the range of 900e1050 C and brazing time varying from 3 to 15 min. The maximum tensile strength at room temperature is 316 MPa when the joint is brazed at 950 C for 5 min. (4) The fracture morphology of the joints treated under different brazing conditions is characterized as brittle cleavage. The cracks prefer to initiate at the locations with relatively high g-Ti2Cu/Ti2Ni content and spread through them. Acknowledgments This topic of research was financed by the Research Project of Special Furnishment and Part (Grant No. XZJQ-B1120680), the Research Foundation of State Key Laboratory of Advanced Welding and Joining (Grant No. AWJ-Z14-02). References [1] Lin TS, Li HX, He P, Wei HM, Li L, Feng JC. Microstructure evolution and mechanical properties of transient liquid phase (TLP) bonded joints of TiAl intermetallics. Intermetallics 2013;37:59e64. [2] Li YL, Liu W, Sekulic DP, He P. Reactive wetting of AgCuTi filler metal on the TiAl-based alloy substrate. Appl Surf Sci 2012;259:343e8. [3] Shiue RK, Wu SK, Chen SY. Infrared brazing of TiAl using Al-based braze alloys. Intermetallics 2003;11:661e71. [4] Shiue RK, Wu SK, Chen SY, Shiue CY. Infrared brazing of Ti50Al50 and Tie6Ale4V using two Ti-based filler metals. Intermetallics 2008;16:1083e9. [5] Dong HG, Yang ZL, Yang GS, Dong C. Vacuum brazing of TiAl alloy to 40Cr steel with Ti60Ni22Cu10Zr8 alloy foil as filler metal. Mater Sci Eng A 2013;561: 252e8. [6] Song XG, Cao J, Liu YZ, Feng JC. Brazing high Nb containing TiAl alloy using TiNi-Nb eutectic braze alloy. Intermetallics 2012;22:136e41.
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