Effects of bonding temperature on microstructure and mechanical properties of diffusion-bonded joints of as-cast Mg–Gd alloy

Materials Science & Engineering A 767 (2019) 138408

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Effects of bonding temperature on microstructure and mechanical properties of diffusion-bonded joints of as-cast Mg–Gd alloy

T

Xin Tonga,1, Le Zaia,1, Guoqiang Youa,b,∗, Hong Wuc, Hengyu Wend, Siyuan Longa a

College of Materials Science and Engineering, Chongqing University, Chongqing, 400045, China National Engineering Research Center for Magnesium Alloys, Chongqing University, Chongqing, 400044, China c Shanghai Institute of Aerospace Precision Machinery, Shanghai, 201600, China d Guizhou Aerospace Xinli Casting & Forging Company Limited, Zunyi, 563003, China b

A R T I C LE I N FO

A B S T R A C T

Keywords: Mg-Gd alloy Diffusion bonding Microstructure Compound layer Tensile properties

In this study, an as-cast Mg–Gd binary alloy was bonded successfully via vacuum diffusion bonding at 500–560 °C in a closed graphite mould, and the microstructures and mechanical properties of the formed joints were analysed. It was found that Gd atoms diffuse to the bonding interface at temperatures lower than the eutectic temperature, with the resulting precipitated compounds improving the bonding strength. At temperatures higher than the eutectic temperature, the eutectic liquid produced at the interface gets squeezed into the residual voids, where it solidifies to form metallurgical bonds. Two types of grain boundary migration processes were observed at the interface in the joint formed at 550 °C. With an increase in the bonding temperature, the tensile strength of the joints first increased and then decreased, with the joint formed at 550 °C exhibiting a maximum efficiency of 88.3%. Casting defects in the base metals were eliminated during the bonding process with the use of a closed mould; this was also the reason that the joint produced at 550 °C showed a higher elongation than the as-cast alloy. However, the amount of the interfacial eutectic liquid increased sharply with the increase in the temperature. Thus, the joint formed at 560 °C showed a low elongation of 0.4% because a large number of microcracks formed in it owing to the solidification and shrinkage of the liquid phase.

1. Introduction Mg-rare earth (Mg-RE) alloys have attracted considerable attention in the aerospace industry, where reducing the weight of components is of great importance. This is because these alloys are lightweight structural materials with high specific strengths, good electromagnetic shielding properties, and desirable damping characteristics [1–4]. In particular, the Mg–Gd system, which includes Mg-Gd-Y-Zr [5,6], MgGd-Al-Zn [7,8], Mg-Gd-Zn-Zr [9,10], and Mg-Gd-Ag-Zr alloys [11,12], has been reported to show both high strength and reasonable elongation as well as high creep resistance at elevated temperatures [13]. Intricately structured parts cast from these alloys are used widely as aerospace components [14,15]. At the same time, some structures inevitably require the joining of as-cast Mg–Gd alloys in order to reduce the manufacturing difficulty. Thus, various technologies for welding Mg–Gd alloys have also been developed, such as arc welding [16], laser welding [17,18], and friction stir welding [19,20]. It is worth noting that the element Gd exhibits an even stronger

tendency to form oxide inclusions than Mg [13]. The high temperatures during conventional fusion welding can lead to a significant oxidationrelated burning loss of Gd as well as the development of a wide heataffected zone, resulting in the deterioration of the properties of the thus-formed joints. Similarly, a major disadvantage of friction stir welding is that it cannot be used with complicated structures [21]. On the other hand, vacuum diffusion bonding is an advanced solid-state joining technology that results in the minimum deformation of the components and involves relatively low and controllable process temperatures. One of the primary applications of this technology is in the fabrication of complex structures in a single step, such as lattice truss structures [22], honeycomb-like structures [23], heat exchange devices [24], and blanket structures [25]. Thus, Vacuum diffusion bonding shows unique prospects to promote the further application of as-cast Mg–Gd alloys in the aerospace field. However, there are no works focusing on the diffusion bonding of Mg–RE alloys – much less as-cast Mg–Gd alloys – a limited number of studies on the diffusion bonding of deformed AZ and ZK series Mg

∗

Corresponding author. College of Materials Science and Engineering, Chongqing University, Chongqing, 400045, China. E-mail address: [email protected] (G. You). 1 Xin Tong and Le Zai contributed equally to this work. https://doi.org/10.1016/j.msea.2019.138408 Received 26 August 2019; Received in revised form 8 September 2019; Accepted 9 September 2019 Available online 10 September 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.

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machining process [34,38,39]. More importantly, the effects of the extruded liquid on the joint quality are still not clear. Given these facts, in the present study, the diffusion bonding of an as-cast Mg–Gd binary alloy was performed without any interlayers at four different bonding temperatures. In order to prevent the interface liquid from being squeezed out and the material deformation, the bonding was performed in a closed graphite mould. The bonding mechanism at the different temperatures is discussed in terms of the microstructural observations and mechanical properties of the formed joints. We believe that these results will improve our understanding of the diffusion bonding of Mg-RE alloys and further its applicability.

Table 1 Actual chemical composition of the samples determined by XRF analysis. Bonding temperature (°C)

500

540

550

560

Chemical composition (wt %)

Mg-8.15Gd

Mg-8.07Gd

Mg-8.13Gd

Mg-8.06Gd

alloys can be found in the literature [21,26,27]. Sun et al. [26] reported the transient-liquid-phase (TLP) diffusion bonding of the alloy AZ31 using interlayers of pure Al. The shear strength of the thus-formed joints was as high as 92.4% of that of the parent metals. Similarly, Jin et al. [27] could form TLP bonding joints of AZ31 with a shear strength of 36 MPa using pure Ni interlayers. Based on these results, Xu et al. [21] successfully performed TLP bonding on the alloy AZ31B under ultrasonic conditions using pure Zn interlayers. The bonding time was shortened to 1 s, and the maximum shear strength of the resulting joints was 42 MPa. When the heating temperature is higher than the eutectic point, the interlayers (of Al, Ni, and Zn, among other elements) react with α-Mg to form the lowmelting-point eutectic liquid. Subsequently, the interlayer elements diffuse to the parent materials on either side, inducing the isothermal solidification of the liquid phase and resulting in joint formation [27]. However, the incorporation of the intermediate layer not only increases the complexity of the joint-forming process but also affects the composition and corrosion resistance of the formed joints. Moreover, it results in differences in colour and thus reduces the aesthetics of the final product. When the heating temperature is lower than the eutectic temperature, more time is needed to form sound joints that do not exhibit the original bonding line, owing to the grain boundary (GB) migrations caused by elemental diffusion [28–30]. Somekawa [28,29] was able to form bonding joints of AZ31 by direct diffusion in the temperature range of 250–400 °C (the eutectic temperature of Mg–Al alloys is 473 °C) without using any interlayers. The maximum ratio of the lap shear strength was 0.85 at 400 °C, and the original bond line disappeared after bonding for 3 h. This could be ascribed to the slow diffusion of the elements and the migration of the GBs. These results indicate that the temperature has a significant effect on the diffusionbonding mechanism. Thus, it is necessary to study the diffusion-based formation of joints of as-cast Mg–Gd alloys at temperatures higher and lower than their eutectic temperature. Because the process of preparing as-cast Mg–Gd alloys is primarily based on nonequilibrium solidification [31–33], the resulting nonequilibrium eutectic structure and the eutectic secondary phases can also undergo a reverse eutectic reaction during the heating procedure. Thus, in theory, the diffusion welding of as-cast Mg–Gd alloys can be performed without any interlayers. In addition, moulds are not essential for diffusion welding, as evidenced by several recent studies [34–37]. Hence, when the heating temperature or bonding pressure is too high, the eutectic liquid phase is squeezed out, resulting in the irregular deformation of the parent alloy. This would affect the subsequent

2. Experimental procedure The cylindrical samples (Ф 12 mm × H 15 mm) used for diffusion bonding were machined from Mg-8wt%Gd ingots. The alloy ingots were prepared in an electromagnetic induction furnace, which was protected by a CO2 + 1%SF6 (volume fraction) mixed atmosphere. The raw materials were commercially available pure Mg (99.99 wt%) and Mg30 wt%Gd. After melting pure Mg, the Mg-30 wt%Gd master alloy preheated to 250 °C was added to the melt at 780 °C. Finally, the molten alloy was casted into a water-cooled copper mould pre-heated to 200 °C. Detailed information regarding the preparation of the as-cast Mg–Gd alloy can be found elsewhere [31]. A fluorescence analyser (XRF-1800 CCDE) was applied to test the actual chemical compositions of the samples used for bonding at different temperatures, the results are listed in Table 1. The microstructure of the as-cast Mg–8Gd alloy before bonding is shown in Fig. 1. Prior to bonding, the mating surfaces of the samples were ground with 100#−2000# grit SiC papers and polished to a mirror-like finish. Next, the polished specimens were ultrasonically degreased in acetone for 20 min, and the specimens were stored in ethanol till further use. For the bonding process, two thus-prepared specimens were stacked in a graphite mould under a uniaxial pressure of 30 MPa. As shown in Fig. 2a and b, the diffusion bonding was performed in a vacuum hotpressing furnace at a vacuum degree of 2–5 Pa, and Fig. 2c displays the size of the tensile specimen. Considering that the eutectic temperature of the Mg–Gd binary alloy is 548 °C [40], the specimens were bonded at 500, 540, 550, and 560 °C for 90 min at an induction heating rate of 20 °C/min. To decrease the damage caused by the residual thermal stress, the uniaxial load was removed first, upon the completion of the bonding process. The samples were then slowly cooled to room temperature within the furnace. Metallographic samples for microstructural observations were cut perpendicular to the bonded surface, then mechanically ground, polished, and etched in a 4 vol% natal solution for 8 s. The characteristics of the bonding interface were analysed using an optical microscopy (OM) system (Carl Zeiss, Axivoert 40). A scanning electron microscopy (SEM) system (JEOL, JSM-7800F Prime) equipped with an energy dispersive X-ray spectroscopy (EDS) attachment was used for detailed morphological characterisation and microchemistry analysis. Electron

Fig. 1. Optical (a) and SEM (b) micrographs of as-cast Mg–8Gd alloy. (c) XRD spectra of the as-cast Mg–8Gd alloy. 2

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Fig. 2. (a) Experimental apparatus used and (b) schematic of vacuum diffusion-bonding process. (c) The size of the tensile specimen.

Fig. 3. Optical micrographs of joints formed at (a) 500, (b) 540, (c) 550, and (d) 560 °C.

hardness tester (MH-60) using a load of 50 g and dwelling time of 10 s. The tensile strengths of the diffusion-bonded joints were measured using an electronic universal testing machine (Sans, CMT 5105) at room temperature at an initial strain rate of 1 mm/min. Three specimens were tested from each temperature group in order to prevent random errors. After the tensile tests, the phases of the fracture surfaces of the

back scattering diffraction (EBSD) analysis was performed to accurately investigate the migration of the GBs along the bonding interface. The EBSD specimens were prepared by mechanical grinding, which was followed by electrolytic polishing in an AC2 solution for 100 s at a voltage of 10 V and current of 0.02 A. The microhardness profiles of the joints were measured using crosssections of the joints. The measurements were performed with a Vickers 3

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Fig. 4. SEM micrographs corresponding to different magnifications of specimens bonded at (a–b) 500, (d–e) 540, (g–i) 550, and (j–k) 560 °C. Images in (c), (f), and (l) are magnified views with their corresponding EDS maps of (b), (e), and (k).

microstructures of the joints bonded by heating in the temperature range of 500–560 °C for 90 min. As can be seen from Fig. 3a–c, the Mg–8Gd alloy could be bonded well in this temperature range, with neither obvious defects such as pores nor areas with incomplete bonding being present. However, large voids attributable to the absence of bonding were seen in the joint formed at the higher bonding temperature of 560 °C (see Fig. 3d). As the bonding temperature was increased, the bonding line of the joints became wider. Moreover, the grain size of the α-Mg matrix in the joints also increased significantly with the increasing temperature, whereas the number of secondary-phase particles dispersed within the matrix decreased gradually. It is worth noting that a few continuous secondary phases were observed along the GBs when the bonding temperature was higher than the eutectic temperature of Mg–Gd alloys

tensile specimens were determined through X-ray diffraction (XRD) analysis (Rigaku, Ultima IV), which was performed at a scan rate of 2°/ min. 3. Results 3.1. Microstructural analysis As shown in Fig. 1a, the as-cast Mg–8Gd alloy fabricated for diffusion bonding process exhibited a coarse, dendritic microstructure. The α-Mg matrix exhibited uneven grain sizes of approximately 50–250 μm. According to Fig. 1c, a large number of fine particles (sizes of approximately 1–5 μm) proved to be the secondary phase of Mg5Gd were distributed uniformly within the matrix [2]. Fig. 3 shows the 4

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Fig. 5. EBSD analysis results for specimens bonded at (a) 500, (b) 540, (c) 550, and (d) 560 °C; (e) is a magnified view of the rectangular region in (c) and (f) is a magnified view of the rectangular region in (e).

Fig. 6. SEM micrographs of parent alloy after bonding at different temperatures: (a) 500, (b) 540, (c) 550, and (e) 560 °C; (d) and (f) are magnified views with corresponding EDS maps of the rectangular region in (c) and (e), respectively.

where no compound layer was present. As shown in Fig. 4h, GB migration occurred primarily at the triple junctions, which formed wedge among three grains (Type 1). On the other hand, the localised interfacial GBs tended to migrate to form spherical caps, as shown in Fig. 4i (Type 2). As the bonding temperature was increased to 560 °C, a continuous Gd-rich compound layer with a width of 2.1 μm appeared along the bonding interface (see Fig. 4j and l). However, as shown in Fig. 4k, many microcracks were present in the compound layer, in contrast to the case for the joints bonded at the lower temperatures. EBSD analysis of the joints bonded at the different temperatures are shown in Fig. 5. As the bonding temperature was increased, the grain

(548 °C). Finally, a large number of subgrains were observed in the joint formed at 560 °C. SEM images and corresponding EDS maps of the joints are shown in Fig. 4. In the case of the joint fabricated at 500 °C, an area free of bonding (width of area ~390 nm) was observed at larger magnifications, as shown in Fig. 4b and c. In addition, the EDS surface analysis results showed that the element Gd was concentrated along this unbonded interface (see Fig. 4c). Sound joints free of incompletely bonded areas were formed at 540 and 550 °C, and the width of the Gd-rich compound layer increased gradually with the increasing temperature, as shown in Fig. 4d and g. Interestingly, in the joint formed at 550 °C, two types of GB migrations were observed at the bonding interface, 5

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Fig. 7. (a) Vickers hardness profiles and (b) indentation distributions of joints formed at different temperatures.

Fig. 8. (a) Engineering stress-strain curves and (b) tensile properties of various joints.

interface was subject to a higher amount of deformation. The OM observation of sub-grains in the joint fabricated at 560 °C (Fig. 3d) was confirmed by the differences in colour within grains in Fig. 5d, which indicate the existence of sub-structures with different orientations. Fig. 6 shows the microstructural evolution of the parent materials after bonding at the different temperatures. The sizes of the secondary phases after the bonding process were significantly smaller than those in the as-cast state; this was in keeping with the results shown in Figs. 3 and 4. Moreover, in the joints formed at 500 and 540 °C, the local matrix around the secondary-phase particles became significantly brighter (see regions within ovals in Fig. 6a and b). This phenomenon can also be observed in Fig. 4a and d; the solid solubility of Gd in local matrix was increased owing to re-dissolution of the Mg5Gd phases during the bonding process, which affected the SEM contrast. As can be seen from the EDS map in Fig. 6d, continuous Gd-rich secondary phases appeared along the GB (see Fig. 6c) when bonding was performed at a temperature higher than the eutectic temperature. Moreover, further coarsening of the secondary phases occurred (see Fig. 6e) when the bonding temperature was increased to 560 °C.

Fig. 9. Photographs showing fracture locations of tensile specimens corresponding to different joints.

size of α-Mg matrix also increased. The migration behaviour of the interfacial GB was observed in the joint formed at 550 °C (see Fig. 5e and f). The boxed and oval regions in Fig. 5e show GB migrations of Types 1 and 2, respectively. It can be seen that the distance of migration of the GBs at the triple junctions without existence of compound layer was ~850 nm after bonding for 90 min. Twinning was observed only in the joint fabricated at 560 °C (see Fig. 5d), indicating that this bonding

3.2. Microhardness and tensile properties Fig. 7a shows the microhardness profiles as determined across the bonding interface of the various specimens. The hardness of the original 6

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Fig. 10. SEM micrographs corresponding to different magnifications of tensile fracture surfaces of (a) as-cast parent materials and joints formed at (b–c) 500, (e–f) 540, (h–i) 550, and (j) 560 °C. (d) and (g) are EDS maps for samples in (c) and (e), respectively.

and then decreased to 57.2 HV when the bonding temperature used was 560 °C. As shown in Fig. 8a, in order to investigate the effects of the bonding temperature on the tensile properties of the diffusion-bonded joints of the as-cast Mg–8Gd alloy, tensile tests were performed on the joints. The 0.2% tensile yield strength (YS), ultimate tensile strength (UTS), and elongation (EL) values of these joints are shown in Fig. 8b. With the increase in the bonding temperature, the UTS first increased from 85.2 MPa at 500 °C to the maximum value of 151 MPa at 550 °C with the highest joint efficiency of 88.3%. It then decreased to 95.9 MPa at 560 °C. The YS and EL values also showed similar trends. Further, the EL value of the joints fabricated at 550 °C was 10.6%, which was unexpectedly higher than that of the original as-cast alloy by 19.1%. However, when the bonding temperature was increased to 560 °C, the

as-cast Mg–8Gd alloy, which was 67.8 HV, is shown as a horizontal line for comparison. Fig. 7b shows the distributions of the indentations made in the various joints. It can be seen that the microhardness of the joints was always lower than that of the as-cast alloy, with the trend in the microhardness curves varying with the bonding temperature. With the increasing bonding temperature, the microhardness of the interface of the joint formed at 500 °C was 60.1 HV, and gradually increased to the maximum value of 62.2 HV at 550 °C, which was 91.7% of that of the as-cast state. However, when the temperature was increased further to 560 °C, the hardness decreased sharply by 25%–46.2 HV. On the other hand, the hardness trend for the base metal away from the bonding interface was different. With the increase in the bonding temperature, the hardness of the base metal after bonding at 500 °C was 59.6 HV, reached its maximum of 64.6 HV in the joint formed at 540 °C, 7

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Fig. 11. Optical micrographs of longitudinal sections of fracture surfaces of (a) as-cast alloy and joints bonded at (b) 500, (c) 540, (d) 550, and (e) 560 °C.

river-like pattern was observed in the lower-right area (see magnified image in Fig. 10f). When viewed with the higher Gd content, it suggests that failure here is attributable to the brittle Gd-rich compound layer that formed along the bonding interface. Fig. 10h shows that the joint formed at 550 °C failed in the area with the parent materials away from the bonding interface; a magnified SEM image of the boxed region is shown in Fig. 10i. The tensile fracture surface consisted of a few cleavage planes, coarse dimples, tear ridges, and torn boundaries. In addition, a few fine cubic particles were also present at the bottom of the dimples in Fig. 10i, suggesting the main fracture characteristic of quasicleavage. Compared with the fracture morphology of as-cast Mg–8Gd alloy in Fig. 10a, it can be concluded that the bonding process at 550 °C have changed the fracture mode of the parent alloy, as it increased the material plasticity. The fracture surface of the joint formed at 560 °C showed a completely brittle cleavage feature with many rough river patterns. Based on this observation and the failure location of the joint (see Fig. 9), the fracturing of this joint can also be attributed to the brittle Gd compound layer formed along the bonding interface. Fig. 11 shows OM images of longitudinal sections of the ruptured joints. It can be seen from Fig. 11a that the cracks in the as-cast Mg–8Gd alloy, which underwent transgranular fracturing, were initiated within the grain interiors. Further, as can be seen from Fig. 11b and e, in the case of the joints formed at 500 and 560 °C, fracturing occurred along the bonding interface. Further, their failure surfaces were relatively smooth, indicating that the strength of their bonding interfaces was low. The joint formed at 540 °C also initially failed at the bonding interface. However, the initiated crack propagated toward the base metal near the interface, suggesting that the local bonding strength of the interface in this case was higher than that of the base metal. Although the joint formed at 550 °C failed in the base metal, as mentioned above, both transgranular and intergranular cracks were observed in it (see Fig. 11d), with its fracture mode being different from that of the original as-cast state (see Fig. 11a).

EL value was drastically reduced to only 0.4%.

3.3. Fracture behaviour Fig. 9 shows that the failure locations of the tensile specimens corresponding to the joints formed at the different bonding temperatures. It can be observed that, with the exception to the joint fabricated at 550 °C, all the other joints failed at the bonding interface. This indicates that the strength of the bonding interface of the joints formed at 550 °C was higher than that of the parent materials. Further, while the joint prepared at 500 °C fractured with a smooth failure surface, the fracture surfaces of the samples bonded at 540, 550 and 560 °C were relatively rougher. Fig. 10a shows an SEM image of the typical fracture features of the as-cast Mg–8Gd alloy at room temperature. It can be seen that the tensile fracture surface primarily consists of a large number of cleavage planes, indicating that the sample had undergone brittle cleavage fracture. Fig. 10b–j shows SEM fractographic images of the diffusionbonded joints formed at the four different temperatures. The images indicate that the bonding temperature had a pronounced effect on the fracture behaviour of the joints. Fig. 10b shows the fracture surface of the joint bonded at 500 °C. The surface consists of two distinct regions, which are separated by a yellow dotted line. River-like patterns can be observed on the left part of the fracture surface, suggesting that a torn compound layer was probably present in this region. The right part of the fracture surface was relatively flat, and a few fine subgrains were observed in the region at a larger magnification (see Fig. 10c), indicating that the fracture failure here was attributable to the absence of bonding, thus the surface ground and polished prior to the bonding process was exposed. Surprisingly, elemental Gd was concentrated along the subgrain boundaries of the unbonded surface, as is evident from the EDS map in Fig. 10d. Fig. 10e and f shows the fracture morphology of the joint formed at 540 °C. Again, the fracture surface can be divided into two distinct regions with different fracture characteristics. EDS analysis (Fig. 10g) revealed that the lower-right region of the fracture surface in Fig. 10e had a higher Gd content than the upper-left. The upper-left area of the fracture surface in Fig. 10e was primarily characterised by cleavage planes, indicating that this area was broken because of the transgranular fracturing of the matrix near the bonding interface. However, a

4. Discussion Fig. 12 shows a schematic of the mechanism of the diffusion bonding of the as-cast Mg–8Gd alloy. As can be seen from Fig. 12a, the base metals were locally contacted at room temperature under the bonding pressure; with a few macroscopic voids also being present. At 8

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Fig. 12. Schematics showing (a) changes in alloy microstructure after application of bonding pressure at room temperature and microstructural evolution of joints formed at (b–c) 500, (d–e) 540, (f–g) 550, and (h–i) 560 °C.

could fill into the voids to achieve the metallurgical bonding. This is the reason no unbonded areas were observed in this joint (see Fig. 4e and f). As shown in Fig. 12f, the Mg–Gd phases in the matrix underwent a reverse eutectic reaction with α-Mg to produce the eutectic liquid phase at 550 °C [31]. The liquid phase at the bonding interface was squeezed into the residual voids present nearby at the high bonding pressure and eventually solidified to form the Gd-rich compound layer during the subsequent cooling process. If this hypothesis is correct, Gd-rich compound layers should also be present at the contact interfaces between the graphite mould and the alloy ingot. As shown in Fig. 13a and c, Gdrich compound layers were indeed found at these interfaces. Moreover, the two types of GB migrations shown in Fig. 4h and i are generally attributed to the differences in the roughness of the initial interface [30,42,43]. In addition to the formation of intermetallic compounds, GB migrations across the initial interface because of atomic diffusion also result in metallurgical bonding. The amount of eutectic liquid phase at the interface was much greater when bonding was conducted at 560 °C as opposed to 550 °C. However, this liquid phase was different from the low-melting-point eutectic liquid phase that is formed by the incorporation of an

the bonding temperature of 500 °C, the thermal plasticity of the base materials was insufficient, causing some voids to remain at the bonding interface, which would hinder the atomic diffusion required for metallurgical bonding. Therefore, an area free of bonding and having a width of ~390 nm existed in the joint after the bonding process, as shown in Fig. 4c. At the same time, Gd atoms diffused spontaneously and aggregated to the unbonded surfaces during the heating process, because there are many vacancies presented in the crystal defects (e.g., surfaces, GBs, and subgrain boundaries) with the relative higher energy [41]. The Mg5Gd phase would then precipitate at the unbonded surface once the Gd content exceeds the solubility limit of the matrix. However, the amount of the produced compounds was too small to fully fill the unbonded voids. Hence, as shown in Fig. 4c, a thin Gd-rich compound layer was observed along the unbonded interface, which was also confirmed from the fracture images in Fig. 10c and d. Increasing the bonding temperature from 500 to 540 °C improved the thermal plasticity of the base metal and enhanced atom diffusivity. The number and size of the residual voids were significantly reduced at 540 °C. Thus, Gd atoms subsequently diffused to the local unbonded surface during heating process, and the resulting Gd-rich compound was formed and 9

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Fig. 13. (a, c) SEM micrographs of different contact interfaces between graphite mould and alloy ingot. (b) EDS map of (a). (d) Magnified view and corresponding EDS map of boxed region in (c).

As can be seen from Fig. 7, the hardness of the bonding interface formed at 500 °C was lower than that of the bonds formed at 540 °C and 550 °C, owing to incomplete bonding. Increasing the bonding temperature increased the volume fraction of the interfacial compound layers and hence the hardness of the bonding interface [46,47]. As a result, the interfacial hardness was the highest at 62.2 HV in the case of the joint formed at 550 °C, but then decreased drastically by 25%–46.2 HV for the joint formed at 560 °C because of the formation of a large number of shrinkage cracks in this joint. However, in the 540 °C joint, the base materials showed the highest hardness, which decreased with the increase in the bonding temperature. This trend may be attributed to the following factors: the solution-strengthening effect was the strongest at 540 °C. Further, overburning and severe grain coarsening were observed at 550 and 560 °C. Several studies [48,49] have stated that precipitation strengthening is the most significant factor in the strengthening of Mg–Gd alloys. However, the degree of precipitation strengthening was significantly reduced by the dissolution of the secondary phases during the bonding process, because of which the overall hardness of the joints was always lower than that of the original as-cast alloy. As per the SEM microstructural image shown in Fig. 4a, distinct unbonded areas existed along the bonding interface of the 500 °C joint; these greatly reduced the mechanical properties of the joint. This is the reason the elongation of the joint was only 1.8% as well as why it failed at the bonding interface during the tensile test. Similarly, in the 560 °C joint, a large number of microcracks were observed at the interfacial Gd-rich compound layer. This explains why this joint showed the lowest elongation of 0.4% and a poor tensile strength of 94 MPa. Compared to the case for the 500 °C joint, since the thermal plasticity of the base metal was higher at 540 °C, which contribute to improvement of the joint quality by eliminating the unbonded defects. Thus, the UTS and EL of the joint were increased to 132 MPa and 6.2%, respectively. Many studies have shown that brittle intermetallic compounds when present at the bonding interface can act as crack initiators during tensile tests [50,51], because they adversely affect the ductility of the alloy. This is in keeping with the tensile failure positions shown in Fig. 9. For this reason, the tensile properties of the joint formed at 540 °C were poorer than those of the original as-cast alloy, as shown in Fig. 8b.

interlayer [21,26,27]. In this experiment, the gradient in the concentration of Gd between the interfacial eutectic liquid phase and the base metal was relatively low. Moreover, the atomic radius of Gd is much larger than that of Mg. Hence, the Gd atoms in the interfacial liquid did not diffuse readily to the base metals and there was no isothermal solidification during the bonding process. Only when the temperature was increased, that is, when heating was performed could the interfacial eutectic liquid phase solidify during the subsequent furnace cooling stage, resulting in metallurgical bonding. It is worth noting that, when the temperature was increased to 560 °C, the amount of the eutectic liquid phase produced along the interface was the highest. Further, in this case, the closed mould could prevent the liquid phase from being squeezed out, as shown in Fig. 12h. During the cooling process after removing the uniaxial pressure, the eutectic liquid phase at the interface solidified and shrank, resulting in the generation of a large number of microcracks in the compound layer, as shown in Fig. 4k and l. A few subgrains were also observed in the microstructures of the joint formed at 560 °C and this was because of the softening of the base metals and the breaking of the grains at the elevated temperature. Moreover, no GB migrations were observed in the joint bonded at 560 °C, owing to the suppression effects of the large number of continuous compounds formed along the bonding interface. The heating performed during the diffusion-bonding process also had a pronounced effect on the base materials. The secondary phases dissolved into the α-Mg matrix when the bonding temperature was lower than the eutectic temperature [17,44]. In other words, the base materials underwent a solid-solution treatment, as shown in Fig. 12c and e. Thus, as the bonding temperature was increased, the size and volume fraction of the secondary phases were effectively reduced, and the saturation degree of the alloy elements in the matrix increased, as shown in Fig. 6a and b. As shown in Fig. 12f and h, when the bonding temperature was higher than the eutectic temperature, the eutectic secondary phases in the matrix, and especially those at the GBs, would react with the surrounding α-Mg to form local eutectic liquid phases, resulting in overburning. It can be seen from Fig. 6d that there were continuous Gd-rich secondary phases distributed along the triple junctions of the GBs and this is a typical characteristic of overburning defects [45]. 10

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Fig. 14. Schematics of fracture mechanisms of joints formed at different temperatures.

interface. Thus, the tensile test results were reflective of the mechanical properties of the base metals after the heating step of the diffusionbonding process but not those of the bonding interface. On the other hand, the diffusion bonding process increased the EL of the as-cast base metals but reduced their tensile strengths. As mentioned above, the Gdrich compound layer formed at the interface facilitated metallurgical bonding when the bonding temperature was higher than the eutectic temperature. However, the local eutectic liquid phase may also form within the matrix of the base metal and lead to overburning, which would significantly hinder intergranular combination. Therefore, the UTS of this joint was lower compared with that of the original as-cast alloy. In general, typical casting defects, such as pores, shrinkage pores, shrinkage cavities, and cold shuts can sharply lower the tensile properties, especially the EL of castings. Since the bonding experiments were performed in a closed graphite mould, the three-dimensional compressive stress state of the specimens during the bonding process could effectively eliminate these casting defects. This was the reason the EL of this joint was higher than that of the original as-cast alloy. Fig. 14 shows schematics of the fracture mechanisms of the joints bonded at the different temperatures. As the continuous unbonded area in the 500 °C joint significantly reduced the effective load-bearing area of the joint, the tensile crack inevitably propagated along the unbonded area and resulted in failure owing to the stress concentration. The fracture surface with a distinct unbonded area can be seen in Fig. 10b. Similarly, as shown in Fig. 14, the shrinkage crack in the compound

Fig. 15. XRD spectra of tensile fracture surfaces of various joints.

The tensile strength of the 550 °C joint was the highest, with its UTS and EL being 88.3% and 119.1%, respectively, of those of the original as-cast alloy. Based on the fracture positions shown in Fig. 9, it can be said that the joint failed at the base metal and not at the bonding 11

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phase reduced the mechanical properties of the joint. 5. Although increasing the bonding temperature aided GB migration, because it resulted in a higher atomic diffusion rate, the continuous compound layer generated at the interface at temperatures higher than the eutectic temperature inhibited GB migration, with overburning and severe grain coarsening occurring simultaneously.

layer formed at the bonding interface of the 560 °C joint was also a relatively weaker region in the joint. The tensile cracks were first initiated at the voids created by the solidification-related shrinkage and then continued to expand along the interfacial compound layer which had a continuous and coarse morphology. Eventually, the joint failed at the bonding interface; the river-like pattern, which is considered to be the typical fracture morphology of the brittle compound layer [38,46,52], can be seen in Fig. 10j. The fact that this joint exhibited the lowest EL (1.8%) proved that the hard and brittle compound layer at the interface adversely affected the mechanical properties of the joint, especially its plasticity. Fig. 15 shows the XRD results of the tensile fracture surfaces of the various joints. Compared to case for the as-cast base alloy, the XRD pattern of the joint fabricated at 560 °C contained a greater number of peaks with higher intensities related to the Mg5Gd phase, indicating that the volume fraction of the Mg5Gd phase in the fracture surface was higher. As can be seen from Fig. 11c, tensile cracks initiated at the interface and then propagated into the base metal near the interface. The morphologies of the transgranular fracture and compound layer fracture all can be seen in Fig. 10e. Since Gd atoms diffused to the bonding surface to form compounds before a complete bond had formed, the number of diffraction peaks related to Mg5Gd in the XRD pattern of the joint fabricated at 540 °C was also relatively higher. As shown in Fig. 14, GB migration was observed at the interface of the 550 °C joint, and the resulting jagged interface improved the bonding quality by hindering the compound layers from being formed continuously along the interface. Although local overburning had an adverse effect on the mechanical properties, the application of a bonding pressure at the high temperature significantly eliminated the casting defects in the base metal. The competition between these two factors eventually increased the material density and improved the plasticity of the joint. As shown in Fig. 10h, the 550 °C joint exhibited a quasi-cleavage fracture with a few coarse dimples, tear ridges, and torn boundaries while the as-cast base alloy exhibited a complete cleavage fracture (see Fig. 10a). This was the reason the EL of the joint bonded at 550 °C was higher than that of the original as-cast alloy.

Acknowledgments This research was supported by a Key Project of the National Key Research & Development Program of China (No. 2016YFB0301100), the Natural Science Foundation of Chongqing (No. cstc2018jcyjAX0430), and the Dongguan Key Project of Core Technology (No. 2019622134013). In addition, special thanks to Dr. Yichang Wang for her company all this time. References [1] L. Jiang, W. Liu, G. Wu, W. Ding, Effect of chemical composition on the microstructure, tensile properties and fatigue behavior of sand-cast Mg-Gd-Y-Zr alloy, Mater. Sci. Eng., A 612 (2014) 293–301. [2] T. Zhao, Y. Hu, B. He, C. Zhang, T. Zheng, F. Pan, Effect of manganese on microstructure and properties of Mg-2Gd magnesium alloy, Mater. Sci. Eng., A 765 (2019) 138292. [3] Y. Zhang, W. Rong, Y. Wu, L. Peng, J. Nie, N. Birbilisc, A comparative study of the role of Ag in microstructures and mechanical properties of Mg-Gd and Mg-Y alloys, Mater. Sci. Eng., A 731 (2018) 609–622. [4] X. Tong, G. You, Y. Ding, H. Xue, Y. Wang, W. Guo, Effect of grain size on lowtemperature electrical resistivity and thermal conductivity of pure magnesium, Mater. Lett. 229 (2018) 261–264. [5] Y. Wang, F. Zhang, Y. Wang, Y. Duan, K. Wang, W. Zhang, J. Hu, Effect of Zn content on the microstructure and mechanical properties of Mg-Gd-Y-Zr alloys, Mater. Sci. Eng., A 745 (2019) 149–158. [6] Y. Chi, C. Xu, X. Qiao, M. Zheng, Effect of trace zinc on the microstructure and mechanical properties of extruded Mg-Gd-Y-Zr alloy, J. Alloy. Comp. 789 (2019) 416–427. [7] B. Pourbahari, H. Mirzadeh, M. Emamy, Toward unraveling the effects of intermetallic compounds on the microstructure and mechanical properties of Mg-Gd-AlZn magnesium alloys in the as-cast, homogenized, and extruded conditions, Mater. Sci. Eng., A 680 (2017) 39–46. [8] B. Pourbahari, H. Mirzadeh, M. Emamy, R. Roumina, Enhanced ductility of a finegrained Mg-Gd-Al-Zn magnesium alloy by hot extrusion, Adv. Eng. Mater. 20 (2018) 1701171. [9] J. Zhang, W. Zhang, L. Bian, W. Cheng, X. Niu, C. Xu, S. Wu, Study of Mg-Gd-Zn-Zr alloys with long period stacking ordered structures, Mater. Sci. Eng., A 585 (2013) 268–276. [10] W. Rong, Y. Zhang, Y. Wu, Y. Chen, T. Tang, L. Peng, D. Li, Fabrication of highstrength Mg-Gd-Zn-Zr alloys via differential-thermal extrusion, Mater. Char. 131 (2017) 380–387. [11] R. Li, H. Shafqat, J. Zhang, R. Wu, G. Fu, L. Zong, Y. Su, Cold-working mediated converse age hardening responses in extruded Mg-14Gd-2Ag-0.5Zr alloy with different microstructure, Mater. Sci. Eng., A 748 (2019) 95–99. [12] Y. Zhang, Y. Wu, L. Peng, P. Fu, F. Huang, W. Ding, Microstructure evolution and mechanical properties of an ultra-high strength casting Mg-15.6Gd-1.8Ag-0.4Zr alloy, J. Alloy. Comp. 615 (2014) 703–711. [13] K. Luo, L. Zhang, G. Wu, W. Liu, W. Ding, Effect of Y and Gd content on the microstructure and mechanical properties of Mg-Y-RE alloys, J. Magn. Alloys 7 (2019) 345–354. [14] W. Liu, L. Jiang, L. Cao, J. Mei, G. Wu, S. Zhang, L. Xiao, S. Wang, W. Ding, Fatigue behavior and plane-strain fracture toughness of sand-cast Mg-10Gd-3Y-0.5Zr magnesium alloy, Mater. Des. 59 (2014) 466–474. [15] Q. Wang, L. Xiao, W. Liu, H. Zhang, W. Cui, Z. Li, G. Wu, Effect of heat treatment on tensile properties, impact toughness and plane-strain fracture toughness of sandcast Mg-6Gd-3Y-0.5Zr magnesium alloy, Mater. Sci. Eng., A 705 (2017) 402–410. [16] D. Meng, B. Zhou, D. Wu, Y. Ma, R. Chen, P. Li, Parameter optimization of gas tungsten-arc repair welding technique in Mg-6Gd-3Y-0.5Zr alloy, Int. J. Metalcast. 13 (2019) 345–353. [17] L. Wang, J. Huang, J. Dong, K. Feng, Y. Wu, P. Chu, Microstructure evolution in the fusion zone of laser-welded Mg-Gd-Y-Zr alloy during solution and aging treatment, Mater. Char. 118 (2016) 486–493. [18] L. Wang, J. Huang, Y. Peng, Y. Wu, Precipitates evolution in the heat affected zone of Mg-Gd-Y-Zr alloy in T6 condition during laser welding, Mater. Char. 154 (2019) 386–394. [19] C. Luo, X. Li, D. Song, N. Zhou, Y. Li, W. Qi, Microstructure evolution and mechanical properties of friction stir welded dissimilar joints of Mg-Zn-Gd and Mg-AlZn alloys, Mater. Sci. Eng., A 664 (2016) 103–113. [20] Q. Yang, B. Xiao, Z. Ma, Enhanced superplasticity in friction stir processed Mg-GdY-Zr alloy, J. Alloy. Comp. 551 (2013) 61–66. [21] Z. Xu, Z. Li, L. Peng, J. Yan, Ultra-rapid transient liquid phase bonding of Mg alloys within 1 s in air by ultrasonic assistance, Mater. Des. 161 (2019) 72–79.

5. Conclusions The vacuum diffusion bonding of an as-cast Mg–Gd binary alloy was performed at four different temperatures in a closed graphite mould, and the microstructures and mechanical properties of the joints were investigated. The conclusions of the study can be summarised as follows: 1. With an increase in the bonding temperature from 500 to 560 °C, the mechanical properties of the diffusion-bonded joints of the as-cast Mg–8Gd alloy first increased and then decreased, exhibiting a maximum tensile strength of 151 MPa and joint efficiency of 88.3% at 550 °C. Further, the elongation was 119.1% of that of the original as-cast alloy. 2. At bonding temperatures lower than the eutectic temperature, Gd atoms diffused to the interface, and the resulting precipitated compounds improved the metallurgical bonding process. However, unbonded areas remained in the joint formed at 500 °C owing to the low thermal plasticity of the base metals. 3. The Mg–Gd phases formed at the bonding interface undergo a reverse eutectic reaction with α-Mg to form a eutectic liquid at temperatures higher than the eutectic temperature. This interfacial liquid phase was squeezed into the residual voids because of the bonding pressure and solidified, resulting in metallurgical bonding. 4. The casting defects in the base alloy could be eliminated during the diffusion bonding process, which was performed using a closed mould. However, the eutectic liquid phase was not squeezed from the bonding interface in the 560 °C bond. Thus, the microcracks formed because of the solidification and shrinkage of the liquid 12

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