Transient Liquid Phase Bonding of Nickel-Base Single Crystal Alloy with a Novel Ni-Cr-Co-Mo-W-Ta-Re-B Amorphous Interlayer

Introduction

Nickel-base single crystal superalloys have been one of the candidate materials in the manufacturing of hot section components of aircraft engines and power generation turbines components due to their excellent properties at elevated temperatures [1]. However, it is difficult to weld nickel-base single crystal superalloys by conventional fusion welding methods, such as argon tungsten-arc welding, metal-inert gas welding, etc. because of their sensitivity to hot cracking during welding and post-weld heat treatment [2]. An ideal joining method would be transient liquid phase (TLP) bonding, also known as diffusion brazing. TLP bonding is a hybrid process that combines the beneficial features of brazing welding and solid-sate diffusion welding, gaining attention in the joining of nickel-base superalloys [3, 4].

The selection of interlayer alloy is crucial for different nickel-base superalloys during TLP bonding process. This interlayer alloy is inserted between the mating surfaces to wet the base metal, and subsequently isothermally solidified at the bonding temperature [5]. At the same time, the formation of eutectic structures would degrade the mechanical properties of parent materials unless sufficient time is provided for the isothermal solidification to reduce such eutectic structures. A post-weld heat treatment (PWHT) is also required to fully homogenize the joint in terms of composition and microstructure [6, 7].

A lot of work has been done on TLP bonding of the nickel-base crystal superalloys [4, 8, 9, 10, 11]. The influence of bonding time and temperature on isothermal solidification of the bond has been studied by Lin et al. [9]. Pouranvari et al. have welded GTD-111 nickel-based superalloy with Ni–Si–B amorphous interlayer. The microstructure evolution in the bonded region with different holding time and bonding temperature were studied [12]. However, only few reports have focused on the design and preparation of those interlayer alloys.

In this work, a nickel-base amorphous foil consisting of Ni-Cr-Co-W-Mo-Ta-Re-B was prepared as bonding interlayer for the TLP bonding of nickel-base single crystal superalloy (Rene’ N5). The microstructure and wetting characteristics of the interlayer alloy on the Rene’ N5 material were characterized and analyzed. TLP bonding experiments were conducted at 1,240 ℃. Furthermore, the effect of PWHT on the interfacial microstructures and mechanical properties of Rene’ N5 joints were investigated.

Experimental procedures

Nickel, chromium, cobalt, tungsten, molybdenum, tantalum, rhenium and nickel-boron alloy, with mass fractions more than 99 %, were melted and refined using a high frequency induction furnace (NEW-M04C), and casted into ingots with a weight of 30 g in each. The chemical nominal compositions of the alloy are shown in Table 1. It is noted that boron here was added to the fillers as a melting point depressant for TLP bonding. Aluminum could form very stable interface phases during TLP bonding [5]. Hence aluminum was excluded from the interlayer composition and was diffused into the joint in this work.

Ni-Cr-Co-W-Mo-Ta-Re-B ingots were fabricated into amorphous foils with a thickness of 40 μm by single-roll rapid solidification equipment (NEW-A05). During rapid solidification, the molten alloy was sprayed on the cooling copper roller rotating at high speed to obtain amorphous structure. The solidus and liquidus temperatures of the amorphous foil were determined by differential scanning calorimetry (DSC, STA499F3). The wettability of the alloy on nickel-base material was tested through sessile drop method. The spherical interlayer alloys were dropped on the Rene’ N5 substrate and heated from 1,000 ℃ to 1,140 ℃ at 10 ℃∙min⁻¹ with a vacuum of 5×10⁻⁴ Pa. A series of photographs were taken by a high speed camera system during heating and cooling. Consequently, the contact angles between the interlayer alloy and the substrate were measured from the photographs.

The chemical compositions of Rene’ N5 single crystal superalloy used in this work are listed in Table 1. The average shear strength of Rene’ N5 is about 654.5 MPa tested with three samples. Before TLP bonding, Rene’ N5 was sectioned into 5 mm×4 mm×4 mm blocks using electro-discharge machine for microstructure and bonding strength analysis. The bonding surfaces of the specimens were polished with 800 grade SiC paper and then ultrasonically cleaned in acetone for 20 min. The sandwiched structure of the interlayer alloy between Rene’ N5 base metals was prepared. Then the assembly was placed into a vacuum furnace and TLP bonded at 1,240 ℃ for 12 h. The post-weld heat treatment (PWHT) procedure was carried out at 1,305 ℃ for 4 h, followed by 1,125 ℃ for 4 h, then held at 900 ℃ for 4 h, and finally cooled to room temperature in furnace. The whole PWHT process was done in argon gas atmosphere (99.999 % Ar).

The cross-sections of samples were cut from the TLP bonding joint and etched in a solution of CuSO₄ (20 g)+HCL (100 ml)+H₂O (100 ml). The microstructures and phase constitute of the TLP joint were characterized by scanning electron microscopy (SEM, CS3400) equipped with energy dispersive spectrometer (EDS) and X-ray diffraction (XRD，D/max2500). Vickers microhardness cross the joining interface of the joint was measured by indenter (DHV-1000), in which a load of 100 g and a dwell time of 15 s were used. The shear strength of the joint was tested by a universal testing machine with a cross-head speed of 0.5 mm/min. The shear fixture for the shear strength test is shown in Figure 1.

Figure 1:

Schematic of shear test fixture.

Results and discussion

Characterization of Ni-Cr-Co-Mo-W-Ta-Re-B alloy

The chemical compositions and physical characteristics of the interlayer materials are of vital importance for TLP bonding. Reasonable chemical composition similar to the base metal is the prerequisite to form a homogeneous joint after bonding. As displayed in Figure 2(a) and (b), the Ni-Cr-Co-Mo-W-Ta-Re-B alloy consists of the M₂₃B₆ (phase A) and plate γ (phase B) based on the EDS result (Table 2) and XRD result of this alloy (Figure 3). The element boron could not be detected due to its lower concentration and the difficulty in quantifying light elements through EDS. The XRD result of the Ni-Cr-Co-Mo-W-Ta-Re-B foil is shown in Figure 4. There is no obvious diffraction peak in the curve but only a typical diffuse scattering peak around 45^o indicating an amorphous structure.

Figure 2:

SEM morphology of Ni-Cr-Co-Mo-W-Ta-Re-B ingot (a) lower magnification (b) higher magnification.

Table 2:

Chemical composition (mass %) of Ni-Cr-Co-Mo-W-Ta-Re-B ingot alloy.

Location	Cr	Co	Ni	Mo	Ta	W	Re
A(White area)	11.60	11.77	25.78	11.34	4.77	26.25	8.48
B(Gray area)	5.51	11.06	66.81	1.00	11.42	3.25	0.95

Figure 3:

XRD pattern of Ni-Cr-Co-Mo-W-Ta-Re-B ingot.

Figure 4:

XRD pattern of Ni-Cr-Co-Mo-W-Ta-Re-B ribbon alloy.

Figure 5 shows the DSC analysis curve of the Ni-Cr-Co-Mo-W-Ta-Re-B ribbon alloy. It illustrates that the alloy appeared anexothermic peak at 380 ℃ approximately indicating its crystallization reaction. The solidus and liquidus temperatures of the alloy foil were 1,065 ℃ and 1,075 ℃, respectively.

Figure 5:

DSC curve of Ni-Cr-Co-Mo-W-Ta-Re-B ribbon alloy.

The wetting of molten filler alloy on base metal is important for forming a joint, thus the wetting of molten interlayer on Rene’ N5 was quantified by contact angles at various temperatures as shown in Figure 6. A series of contact angles show almost a liner decline with the rise of temperature in the range of 1,070–1,120 ℃ for 3 min, while decreases slowly after 1,120 ℃ (Figure 6(a) and (b)). Higher temperature reduces the viscosity of molten interlayer alloy. An excellent wettability (contact angle < 20°) of the molten interlayer alloy on Rene’ N5 superalloy was achieved when the temperature is above 1,110 ℃, and the contact angle reached about 1° at 1,140 ℃ (Figure 6(a) and (b)). Figure 6(c) shows the spreading process of the interlayer alloy on Rene’ N5 at 1,100 ℃, and the contact angel vs. time curve is displayed in Figure 6(d). The spreading process could be categorized into two stages as indicated in Figure 6(d). In the initial stage, as the holding time was less than 10 s, the contact angel decreased rapidly due to the viscous flow of the melting alloy. When the holding time was 10 s, the contact angel was about 6°. Then the contact angle declined slowly and finally reached equilibrium (stage II). During this stage, the interface reaction and diffusion between the liquid alloy and solid base metal are the main influence factors for spreading process. The experimental results show that this designed interlayer alloy has a good wettability on Rene’ N5 when the temperature is higher than 1,110 ℃.

Figure 6:

Variation of contact angles of molten Ni-Cr-Co-Mo-W-Ta-Re-B alloy on Rene’ N5 substrate (a) macro morphology vs. temperature (for 3 min) (b) contact angle vs. temperature(for 3 min) (c) contact angle vs. time(at 1,100 ℃) (d) contact angle vs. heating time (at 1,100 ℃).

Microstructures of the TLP joint

The interfacial microstructures of the TLP joint bonded at 1,240 ℃ for 12 h without PWHT process are shown in Figures 7(a) and (b). No pores and cracks were observed at the interface of the TLP joint. Few white discontinues lines could be found at the centerline of the joint as shown in Figure 7(a). The magnified SEM image indicated these were the precipitates, containing Ni rich phase (Ni > 50 mass %, and Co, Cr, W elements) according to EDS results, which is similar to the reported results in the literatures [5, 8, 10, 13]. These white precipitates were mainly nickel-rich boride phases due to the low solubility of boron in the nickle-base superalloy [5, 10]. It is reported that such boride precipitates were easily formed in the grain boundary [9]. This is due to that the grain boundary could act as a preferred route with high diffusivity for B element in the interlayer. However, the diffusion is slow in base metal, since the high diffusivity paths (grain boundary) is absent in single crystal material. It is worth noting that these borides are brittle and detrimental to the mechanical properties of the joint because they could introduce the nonuniform microstructure in the joint [7, 12].

Figure 7:

Effect of PWHT on the microstructure of joint bonded at 1,240 ℃ for 12 h (a),(b)without PWHT; (c),(d)after PWHT.

These boride precipitates decreased gradually with the increase of the holding temperature and holding time. PWHT, as an effective process to uniform the chemical elements cross the joining interface, is employed to improve microstructure and mechanical properties of the joint. Solution heat treatment and precipitation heat treatment are usually recommended for nickel-base superalloy TLP processing [14, 15]. Considering the microstructure and mechanical properties of the base metal, the PWHT parameters adopted in this study was similar to that of base metal. As can be seen obviously in Figure 7(c) and (d), a homogeneous microstructure without white precipitates at the joining interface is attained after PWHT. Note that it is easy for B to diffuse into the nickel-base superalloy at higher temperature due to its small atomic size. Moreover, a significant amount of γ´ phase formed in the bonding region. This was resulted from the diffusion of Al from base metal into the bonding region and γ´ phase formed consequently during PWHT. Similar results have been reported in references [12, 16].

Hardness and shear strength of the TLP joint

Microhardness across the TLP joint is shown in Figure 8. It can be seen from the profile that the hardness increased in the middle of the TLP joint without PWHT, around 50 Hv above that of base metal. This phenomenon was caused by the higher content of B element with insufficient diffusion [3], causing some intermetallics (M_xB_y) precipitated in this region. After PWHT process, relatively uniform chemical element distribution was attained by element diffusion and no obvious variation of microhardness was observed cross the joint, as shown in Figure 7(c) and (d).

Figure 8:

Microhardness across the bonding region before and after PWHT.

The shear strengths of the TLP joints before and after PWHT were 437.2 MPa and 533.4 MPa, which were 66.8 % and 81.5 % of the base metal respectively as shown in Figure 9. The joint strength had much higher value after PWHT than that without PWHT. Similar results have also been confirmed in other literature [6]. This can be attributed to the elimination of brittle boride as well as the increase of volume fraction of γ´ phase in bonding region after PWHT, since the nickle-base superalloy is γ´ strengthened [12].

Figure 9:

The shear strength of joints.

Conclusion

The Ni-Cr-Co-W-Mo-Ta-Re-B alloy was prepared as interlayer alloy for the TLP bonding of nickle-base superalloy (Rene’ N5). The effect of PWHT on the interfacial microstructure and mechanical properties of the joints were investigated. From this investigation, following conclusions can be drawn:

1. The new Ni-Cr-Co-W-Mo-Ta-Re-B interlayer alloy prepared for bonding was amorphous, which contained plate γ and M₂₃B₆ phases. The molten Ni-Cr-Co-W-Mo-Ta-Re-B alloy exhibited an excellent wettability on the nickle-base superalloy (Rene’ N5).

2. A homogeneous TLP joint can be obtained at 1,024 ℃ for 12 h followed by PWHT. PWHT eliminated the intermetallic compounds and promoted the formation of γ´ precipitates in the bonding region.

3. The microhardness profile of TLP joint became more uniform after PWHT. The shear strength of the joint significantly increased to 533.4 MPa after PWHT (437.2 MPa without PWHT).