Joining of C_f/SiC composite to GH783 superalloy with NiPdPtAu-Cr filler alloy and a Mo interlayer

1. Introduction

Continuous carbon fiber reinforced SiC ceramic matrix composite (C_f/SiC) has a series of good performances, such as low density, low coefficient of thermal expansion, good high-temperature strength and excellent oxidation resistance [[1], [2], [3]]. Therefore, the C_f/SiC composites are promising materials for high-temperature structural components.

Ni-based superalloys have the advantages of high strength, good corrosion resistance and good oxidation resistance, and thus they are widely used in the high-temperature environments in aerospace field [4,5]. In some cases it is highly needed for the joining of Ni-or-Co-based superalloy to C_f/SiC composite. However, there is a great difference in the coefficient of thermal expansion (CTE) for the two materials. The larger the difference of CTE is, the greater the residual thermal stresses within the dissimilar joint are. Therefore, joining of C_f/SiC composite with the Ni-or-Co-based superalloy is extremely difficult.

One of the major obstacles to the brazing of ceramic and metal is to decrease the residual stresses within the dissimilar joints. As an example, when brazed C_f/SiC composite to Ni-based superalloy, multiple interlayers (Cu/W/Cu/W/Cu) was used in order to reduce the thermal stress and improve the joint strength. The maximum joint bend strength of 102.1 MPa was obtained under the joining condition of 970 °C/10 min/34.3 MPa [6]. Similarly, for the joining of SiC ceramic to Ni-based superalloy (GH3044), triple Kovar/W/Ni interlayer was placed between the SiC ceramic and GH3044 superalloy, even so the maximum joint strength was only 62.6 MPa [7]. Another measure to relieve the thermal stress was reported when brazing the C/C composite to a Ni-based superalloy in the reference [8]. Besides an alumina ceramic interlayer was used as a buff layer, zig-zag interfacial structure between the C/C composite and the filler metal was fabricated. However, the maximum bend strength was only 73.0 MPa. Thus an appropriate technology which can achieve sound joint between C_f/SiC composite and Ni-or-Co-based superalloy with good performance is extremely lacking.

In this work, C_f/SiC composite was attempted to braze with GH783 superalloy. Between the two materials, there existed large mismatch of thermal expansion coefficient (α_Cf/SiC = 3.2 × 10^-6 K^-1, α_GH783 = 12.7 × 10^-6 K^-1, 20-100 °C). A newly-developed high-temperature NiPdPtAu-Cr filler alloy together with a Mo intermediate layer was used for decreasing the residual thermal stresses within the dissimilar joint. The considerations for the design of this new filler alloy can be concluded as the following: At first, Cr as the active element for ceramic joining was added into the filler alloy. Furthermore, AuNi-Cr system brazing filler was proposed for joining C_f/SiC composite and the corresponding joints exhibited bend strength of 154.5 MPa at room temperature [9]. Obviously, the C_f/SiC joints used AuNi-Cr filler alloy have good room-temperature strength, but half of the room-temperature strength was remained tested at 600 °C. In addition, it was reported that Pd-based and Pt-based alloys showed good thermal and oxidation resistance as well as excellent ductility [[10], [11], [12], [13]]. The addition of precious metals Pd and Pt can not only enhance the high temperature resistance but also improve the ductility of the filler alloy [14].

In order to release the large stresses within the C_f/SiC-GH783 dissimilar joint, Mo was utilized as a hard buffer due to its low CTE (4.8 × 10^-6 K^-1) [[15], [16], [17]]. Additionally, according to the previous study [14], when joining C_f/SiC composite to itself by using NiPdPtAu-Cr filler alloy, Mo participated in the interfacial reactions and improved the metallurgical behavior at the joining interface. The brittle Ni-Si compounds can transform to the refractory Mo-Ni-Si ternary compound, and thus the interfacial bonding strength increased remarkably [14].

In this study, the effects of Mo interlayer thickness on the joint microstructure and joint strength were studied. At the same time, the joint strengths at high temperatures were also tested. Furthermore, the joining mechanism for the dissimilar joint was also analyzed in detail.

2. Experimental procedure

3-D C_f/SiC composite was used as the base material to be joined, which was fabricated by Polymer Infiltration Pyrolysis (PIP) process [18]. The volume fraction of carbon fiber is about 50% within the C_f/SiC composites, while that of the porosity is about 11.4 vol.% measured by Archimedes method. The GH783 superalloy to be joined has a nominal composition of Co-28.2Ni-25.6Fe-5.65Al (wt%). It has good performances with the combination of high temperature strength and good oxidation resistance, so that it can offer long-term service at the temperature of 700-800 °C.

Two kinds of raw materials were cut into 10 mm × 10 mm × 3 mm slices for the metallographic observation and 3 mm × 4 mm × 20 mm bars for three-point bend test, respectively. And then they were ultrasonically cleaned in acetone for 10 min and dried in air. Pure Mo was chosen as the buffer layer, and four thicknesses were used including 0.2 mm, 0.5 mm, 1.0 mm and 2.0 mm. All joined surfaces were polished by SiC papers up to grit 1000. The newly-developed NiPdPtAu-Cr filler alloy was used with the nominal composition of Ni-18.0Pd-12.0Pt-10.0Au-(25.0-28.0) Cr (in at %). The fabrication method of the filler foils and the wettability result on the C_f/SiC composite were described in the previous work [14]. The butt joints of C_f/SiC composite and GH783 superalloy were prepared by inserting one Mo interlayer and two filler foils, as shown in Fig. 1.

The samples to be joined were vertically fixed in a specially designed graphite jig, and then put into a vacuum furnace. Then, they were heated to 1200 °C from ambient temperature at a heating rate of 10 °C/min, and were held at this temperature for 10 min. During the brazing process the vacuum was kept between 3.0 × 10^-3 Pa and 7.0 × 10^-3 Pa. After the high-temperature brazing the specimens were cooled down to 500 °C with a cooling rate of 5 °C/min, followed by furnace cooling.

To characterize the microstructure of the joints, the polished cross-sections of the brazed joints were examined with an optical microscope (OM) and an electron probe microanalyser (EPMA, JXA-8100) equipped with an X-ray energy-dispersive spectrometer (EDS, INCA-E305).

The joint strength was determined by three-point bend test in air at room temperature and high temperatures. At least three specimens were tested for each test temperature, and the average value was presented as the joint strength. To determine the fracture mechanism, the fracture surface after the bend test was analyzed by using a scanning electron microscope (SEM, FEI nano 450) and an X-ray diffractometer (XRD, D8 ADVANCE).

3. Results and discussion

Fig. 2 presented the microstructures of the C_f/SiC-Mo interface with different thickness of Mo interlayer. When 0.2 mm thickness of Mo interlayer was used, evident crack occurred at the interface between the C_f/SiC composite and the filler alloy, as shown in Fig. 2(a). It was reported that residual stress concentration within the dissimilar joints occurred at the reaction layer close to the ceramic surface [19]. Generally these interfacial reaction products were brittle compounds, and could not tolerate the tensile stresses. As a result, the joint cracked close to the C_f/SiC composite during the cooling process.

A sound joint was obtained with the Mo interlayer thickness of 0.5 mm. The corresponding interfacial microstructure was presented in Fig. 2(b). The interface between C_f/SiC composite and Mo interlayer showed good appearance without obvious defects. However, when the Mo thickness was increased to 1.0 mm, some transverse cracks appeared (shown by black arrows in Fig. 2(c)). Further increasing the Mo thickness to 2.0 mm, the cracks location changed into the reaction layer near the surface of the C_f/SiC composite (Fig. 2(d)).

In general, the strength of C_f/SiC-Mo-GH783 joint was mainly determined by the bonding quality of the C_f/SiC-Mo interface. Obviously Mo interlayer with a thickness of 0.5 mm offered the best bonding appearance. The whole joint microstructure (Fig. 3(a)) showed that, after brazing the Mo thickness was decreased to 375 μm from the original 500 μm. Obviously the Mo interlayer participated in the interfacial reactions. The detailed microstructure between the C_f/SiC composite and the Mo was presented in Fig. 3(b), and the corresponding compositions for some typical microzones were listed in Table 1.

At the surface of the C_f/SiC composite intensive reactions occurred and two reaction layers “1” and “2” formed. According to the area distribution map of the element Cr (shown in Fig. 4(e)), the reaction layer “1” was enriched with Cr. The XRD results in the previous work [14] have already proved that the reaction product at the interface between the NiPdPtAu-Cr filler alloy and C_f/SiC composite was Cr₃C₂ phase. During the brazing Cr element played an active role in the interfacial reaction. At the brazing temperature of 1200 °C, Cr could react with C_f/SiC composite by the following two chemical formulas [14,20]:

Cr + C→Cr₃C₂, ΔG_{f, 1200°C} =-102 kJ/mol

Cr + SiC→Cr₃C₂+Si, ΔG_{f, 1200°C} =-54 kJ/mol

From the Mo distribution map shown in Fig. 4(f), Mo also participated in the interfacial reaction at the surface of C_f/SiC composite and formed the reaction layer “2”. In this layer 26.37 at.% Mo, 23.02 at.% Ni, 21.97 at.% C, 14.73 at.% Cr and 11.38 at.% Si were detected. It can be deduced as a mixture of Mo₂C and Cr₃C₂ as well as some Ni-Si compound. As we know, Mo, a strong carbide former, has a large affinity with carbon. In the study of diffusion bonding of SiC to molybdenum, Mo₂C was an usual reaction product at the interface for joining temperatures ranging from 1200 °C to 1300 °C [21].

In the central part of the joint, according to the EDS analysis results the grey phase “3” should be Pd₂Si compound (Table 1). It has been demonstrated that Pd has a large affinity with Si [22,23], and the formation of Pd₂Si compound is commonly found when using Pd-containing filler alloy for the joining of SiC or Si₃N₄ ceramics or C_f/SiC composite [[24], [25], [26]]. The white phase “4” can be deduced as Au(Cr) solid solution. After the high-temperature brazing this ductile phase could release the residual stresses to some extent through its plastic deformation. Concerning the black phase “5” enriched with elements Mo, Ni, Cr and Si, it should be Mo-Ni(Cr)-Si complex silicides.

Adjacent to the Mo surface, a reaction layer labeled “6” formed with a thickness of 15-35 μm, containing 36.85 at.% Mo, 31.96 at.% Ni and 22.37 at.% Si. Noticeably, some scattered light grey phases “7” existed in layer “6”, composed of 50.53 at.% Mo, 29.25 at.% Ni and 15.52 at.% Si. It can be deduced as another kind of Mo-Ni-Si ternary compound. Gupta [27] has studied the phases in the Mo-Ni-Si system and found that two intermediate phases Γ₁’ and Γ₁’’ around the alloy composition Mo₃₃Ni_50.5Si_17.5 (Γ₁’) and Mo₃₂Ni₃₈Si₃₀ (Γ₁’’) extending from ∼16 to ∼21 at.% Si and ∼23 to ∼35 at.% Si, respectively. Therefore, for the typical microzones “6” and “7”, they can be confirmed as Γ₁’’(Mo-Ni-Si) and Γ₁’(Mo-Ni-Si), respectively. However, in the XRD pattern shown in Fig. 5 only one type of Mo-Ni-Si ternary compound with chemical formula of MoNiSi was detected. Thus it can be deduced the consumed Mo not only participated in the interfacial reactions with the ceramic composite, but also reacted with the filler alloy. Compared with binary molybdenum silicides, the ternary metal silicide Mo-Ni-Si should possess better toughness while still keep the inherent high-temperature stability. The formation of Mo-Ni-Si ternary compound should be favorable to the joint high-temperature strength. Finally, close to the Mo interlayer, there existed a diffusion zone “8”, which could be denoted as Mo(Ni,Si,Cr) solid solution.

The microstructure and the corresoponding elemental distribution maps for the interface between Mo interlayer and GH783 superalloy were presented in Fig. 6. The main mechanism for joining of metal to metal is the elements inter-diffusion. The microzone “1” in Fig. 6 close to the Mo interlayer was enriched with 31.83 at.% Mo (Table 2), and it can be deduced as a Mo(Ni,Co) solid solution diffusion layer.

Some precipitations labeled “2” were also observed, which was enriched with elements Pd and Al (shown in Fig. 6(d) and (i)). From the Pd to Al atomic ratio listed in Table 2 it can be seen that Pd₂Al phase formed here. There is no doubt that element Al come from the GH783 superalloy. When the filler alloy melted, element Al from GH783 superalloy would diffuse into the filler metal and react with Pd due to the large affinity of element Pd and Al (The enthalpy of Al in liquid Pd is -187 kJ mol^-1) [28]. The formation of Pd₂Al compound was also confirmed when brazed AlN ceramic with Au-Pd-Co-Ni-V filler alloy [29]. Meanwhile, white phase “3” occasionally co-existed with Pd₂Al phase, in which 30.78 at.% Au, 14.31 at.% Pd and 10.11 at.% Ni were contained. It should be the Au-based solid solution dissolved with Pd and Ni. The matrix of the central part consisted of elements Ni, Cr, Co and Fe (shown as microzone “4”), it should be the residual filler alloy dissolved with Co and Fe.

Fig. 7(a) presented the joint strength for the different thickness of Mo interlayer. When using 0.2 mm Mo interlayer, the joint strength of 5.3 MPa was given, with nearly no actual bonding. However, the joint with 0.5 mm Mo interlayer showed the maximum strength of 98.5 MPa. Further increasing the thickness of the Mo interlayer to 1.0 mm and 2.0 mm resulted in a decrease of joint strength (58.0 MPa and 68.1 MPa, respectively).

The joint appearance after the bend test was shown in Fig. 7(b). All the joints fractured at the interface between the C_f/SiC composite and the Mo interlayer, even though the incorporation of Mo interlayer could effectively improve the distribution of the thermal stresses within the joint and reduce the concentration of residual thermal stresses near the ceramic composite surface. Evidently, the residual thermal stresses in this area were not released thoroughly [30], so that this interface was still the weakest link over the whole C_f/SiC-Mo-GH783 joints.

The previous work demonstrated that the thickness of the interlayer played an important role in the joint strength [31], but limited data were valuable to the determination of the interlayer thickness for the dissimilar joint. It was assumed that the residual stresses within the joints could be decreased with increasing interlayer thickness because of the larger volume of interlay material available for plastic deformation [32,33]. While, according to another reference [34], the shear strength of the Al₂O₃/steel joints with 0.5 mm Ti interlayer was higher than that of the joints with 1.0 mm interlayer. In this study, the joint strength was increased at first and then decreased with the increase of the Mo interlayer. This was because the interlayer with 0.2 mm was too thin to effectively release the residual stresses. But when the interlayer was too thick (1.0-2.0 mm), the distribution of the residual stresses within the joint was changed and transverse microcracks were induced, and thus the joint strength was conversely decreased. This phenomenon was consistent with the results when brazing porous Si₃N₄ to Invar by using Ag-Cu-Ti/Cu/Ag-Cu multi-layer filler [35]. In this work, the joint strength first increased and then decreased as the thickness of Cu interlayer increased from 50 μm to 150 μm. And the maximum shear strength (73.0 MPa) was obtained with a 100 μm Cu interlayer.

In order to clarify the fracture mechanism, the microstructure analysis for the fracture surfaces was carried out, as shown in Fig. 8. Compositions for some typical microzoneswere presented in Table 3.

From the fractured surface morphologies at GH783 superalloy side (Fig. 8(a) and (b)) two kinds of features were included. One of them was the flat area (microzone “1”), it was mainly the Mo base material according to the EDS result. Furthermore, some particles “2” with small size were composed of 33.74 at.% Mo, 35.51 at.% Ni and 20.37 at.% Si, and they can be deduced as the Mo-Ni-Si ternary compounds. In fact MoNiSi phase was directly detectable in the XRD pattern for the fracture surface (Fig. 5). From Fig. 8(b) lots of voids were visible caused by pulling out of carbon fibers. In fact the pulling out of carbon fibers was favorable to improve the joint strength because of the good mechanical properties of carbon fiber [36,37]. In the meantime, microzone “3” beside the void-shape zone was mainly the MoNiSi compound and some Cr-riched phase. Combined with the XRD result for the fractured surface of the C_f/SiC-Mo-GH783 joint after the bend test (Fig. 5), the Cr-riched phase should be Cr₃Ni₂ compound.

Fig. 8(c) presented the fracture morphology at the C_f/SiC composite side. The micorzone “4” had a similar composition to the microzone “3”. Concerning the microzone “5”, it can be confirmed as Pd₂Si compound by the EDS analysis result together with XRD pattern.This area was in accordance with the grey phase “3” in the central part of the joint (Fig. 3(b)). Even though all joints fractured at the C_f/SiC-Mo interface (Fig. 7(b)), the microcracks propagated throughout the whole reaction zone between the C_f/SiC composite and the Mo interlayer. This fracture characteristic was illustrated in Fig. 9. It is generally believed that the C_f/SiC composite near the brazing seam will be the mostly possible fracture location owing to the thermal stresses concentration when subjected to the bending load. Indeed, further numerical simulation work is needed to reveal the distribution of residual stress within the ceramic/metal dissimilar joint. It has been proved that the residual stress in the joint was not uniform and the highest value occurs in the ceramic material near the seam [38,39]. As a result, during the bending test, cracks should originate near the interface between C_f/SiC composite and filler metal zone, where the residual stress concentration was the largest. With the bending load exerted on the brazed joint, main stress direction changed and then the cracks propagated into the filler metal zone. To sum up, for the C_f/SiC-Mo-GH783 joint, the C_f/SiC-Mo bonding interface was still the weak link over the whole joint.

For C_f/SiC-GH783 joints with 0.5 mm Mo interlayer, three-point bend tests were carried out at high temperatures, as shown in Fig. 10. The newly-developed NiPdPtAu-Cr filler alloy can offer the room-temperature bend strength of 98.5 MPa. When the test temperature was increased to 800 °C and 900 °C, the average joint strength was even elevated to 123.8 MPa and 133.0 MPa, respectively. The phenomenon that joint strengths at elevated temperatures were even higher than room-temperature strength was generally attributed to the release of the inherent stresses within the C_f/SiC-GH783 joint [40,41]. However when further increasing the test temperature to 950 °C, the joint strength was sharply decreased to only 14.4 MPa.

As previously reported, when brazed C/SiC composite to Nb with TiNiNb filler alloy, the shear strength at 800 °C was only 73.0 MPa [42]. Obviously, in this study, NiPdPtAu-Cr filler alloy used in the joining of C_f/SiC and GH783 superalloy could offer the higher joint strength at high temperature. Furthermore, the joining of C_f/SiC to Ni-or-Co-based superalloy is more difficult due to the larger mismatch of the thermal expansion coefficient. As shown in Fig. 10, the newly-developed filler alloy has a potential advantage of satisfying the requirement at service temperatures of 800-900 °C, about 300-400 °C higher than the traditional AgCu-Ti system filler alloys [[43], [44], [45], [46]].

In the previous work, when using this newly-developed filler metal, the strengths of the C_f/SiC-C_f/SiC joints with Mo interlayer were significantly improved compared with the direct brazing (without Mo interlayer). The main reason is the formation of refractory Mo-Ni-Si ternary compound rather than Ni-Si binary compound. Mo-Ni-Si ternary compound exhibits good high-temperature phase stability [47,48]. According to the partial isothermal section of Mo-Ni-Si system diagram, the refractory Mo-Ni-Si compound has a melting point at least no lower than 1280 °C [27,49], and thus the presence of Mo-Ni-Si refractory compound should contribute to the high-temperature joint strength. Furthermore, in the present work, the residual Mo interlayer can release the residual thermal stresses within the dissimilar joint as a hard buffer, and so the C_f/SiC-GH783 dissimilar joint exhibits high bonding strength.

4. Conclusions

Using a Mo interlayer together with the newly-developed NiPdPtAu-Cr filler alloy, joining of C_f/SiC composite to GH783 superalloy was studied. The following conclusions can be drawn from this study:

(1)The dissimilar joint with 0.5 mm Mo interlayer exhibited the maximum strength of 98.5 MPa. When tested at 800 °C and 900 °C the corresponding joint strength was elevated to 123.8 MPa and 133.0 MPa, respectively.

(2)During the brazing process Mo interlayer participated in the interfacial reactions. A reaction band composed of Mo₂C, Cr₃C₂ and Ni-Si compound was formed near the surface of C_f/SiC composite. Furthermore, refractory Mo-Ni-Si ternary compounds distributed in the central reaction zone between C_f/SiC composite and Mo interlayer. Whereas at the bonding interface between Mo interlayer and GH783 superalloy, good metallurgical bonding had been achieved because of the element inter-diffusion.

(3)For the C_f/SiC-Mo(0.5 mm)-GH783 joint, good high-temperature strength was attributed to the formation of refractory MoNiSi. In addition, the residual Mo interlayer, as a hard buffer, can release the residual thermal stresses within the dissimilar joint.

(4)The C_f/SiC-Mo bonding interface was still the weak link over the whole C_f/SiC-Mo-GH783 joint. After the bend test, the micocracks propagated throughout the whole reaction zone between the C_f/SiC composite and the Mo interlayer.

Acknowledgements

This work was sponsored by the National Natural Science Foundation of China (Grant Nos. 59905022, 50475160 and 51275497). We also thank Aeronautical Science Foundation of China (Grant No. 2008 ZE21005).

The authors have declared that no competing interests exist.

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