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Microstructure evolution of a dissimilar junction interface between an Al sheet and a Ni-coated Cu sheet joined by magnetic pulse welding
Materials Characterization 118 (2016) 142–148
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Microstructure evolution of a dissimilar junction interface between an Al sheet and a Ni-coated Cu sheet joined by magnetic pulse welding Takaomi Itoi a,⁎, Azizan Bin Mohamad a, Ryo Suzuki a, Keigo Okagawa b a b
Department of Mechanical Engineering, Chiba University, 1-33, Yayoi-cho, Inage-ku, Chiba 263-8522, Japan Department of Electrical and Electronics Engineering, Tokyo Metropolitan College of Industrial Technology, 1-10-40 Higashi ohi, Shinagawa-ku, Tokyo 140-0011, Japan
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Article history: Received 22 March 2016 Received in revised form 25 May 2016 Accepted 25 May 2016 Available online 26 May 2016 Keywords: Microstructure Amorphous alloy TEM Magnetic pulse welding Joining
a b s t r a c t An Al sheet and a Ni-coated Cu sheet were lap joined by using magnetic pulse welding (MPW). Tensile tests were performed on the joined sheets, and a good lap joint was achieved at a discharge energy of N 0.9 kJ. The weld interface exhibited a wavy morphology and an intermediate layer along the weld interface. Microstructure observations of the intermediate layer revealed that the Ni coating region consisted of a Ni–Al binary amorphous alloy and that the Al sheet region contained very fine Al nanograins. Ni fragments indicative of unmelted residual Ni from the coating were also observed in parts of the intermediate layer. Formation of these features can be attributed to localize melting and a subsequent high rate cooling of molten Al and Ni confined to the interface during the MPW process. In the absence of an oxide film, atomic-scale bonding was also achieved between the intermediate layer and the sheet surfaces after the collision. MPW utilises impact energy, which affects the sheet surfaces. From the obtained results, good lap joint is attributed to an increased contact area, the anchor effect, work hardening, the absence of an oxide film, and suppressed formation of intermetallic compounds at the interface. © 2016 Elsevier Inc. All rights reserved.
1. Introduction The development of welding methods for combinations of dissimilar materials for industrial applications has gained importance because of the economic and technical advantages of such methods. Cu, which exhibits good electrical and thermal conductivity, and Al, which exhibits low density, are two common engineering materials widely used for reducing the cost and weight of various industrial products, especially in automotive applications [1]. The joining of dissimilar metals through the conventional fusion welding method results in the formation of a brittle intermetallic compound layer, which leads to failure at the joint interface under applied external stress. Furthermore, Cu and Al cannot be joined by conventional fusion welding because of large differences in their physical and mechanical properties. Therefore, solid-state welding methods such as cold-roll welding, diffusion welding, explosive welding and friction-stir welding (FSW) have been proposed as alternative methods for joining dissimilar metals with different properties [2– 6]. Magnetic pulse welding (MPW), a solid-state welding method similar to impact welding processes such as explosive welding, is applicable to metals with high electrical conductivity. Impact energy is induced by the electromagnetic force generated by the interaction between the discharge energy (impulse current), high-density magnetic flux and eddy currents produced in the sheet. This electromagnetic force drives the ⁎ Corresponding author. E-mail address: [email protected] (T. Itoi).
http://dx.doi.org/10.1016/j.matchar.2016.05.021 1044-5803/© 2016 Elsevier Inc. All rights reserved.
flyer sheet to the parent sheet, resulting in a high-velocity collision and solid-state welding within several microseconds [7–9]. Recently, Al/Cu and Fe/Al lap joints have been successfully fabricated by using MPW [9–11]. In practical applications, metal sheets are surface-treated by coating with a layer of several metals such as nickel to increase resistance against corrosion. Welding for these sheets with metal coatings have been fabricated by solid-state welding methods such as FSW [12–14]. However, the fabrication of such welding for sheets with metal coatings by MPW has not yet been reported. During the MPW process, localized melting occurs at the interface when the discharge energy is high and intermetallic compounds partially form at the weld interface, as in the case of the interfaces in Al/Cu and Fe/Al lap joints [9–11]. Localized melting occurs because of the high pressure and temperature after the high-speed collision of both sheets. This particular phenomenon, which encompasses the formation of amorphous alloys and nanograins of intermetallic compounds at the interface created by joining dissimilar sheets, also occurs in other impact welding mechanisms such as that for explosive welding. However, the detailed microstructure and mechanism of formation of amorphous alloys at the interface during MPW is still unclear. Amorphous alloys exhibit high strength and toughness, as well as advantageous mechanical properties. The atoms in amorphous alloys are ordered arrays over short ranges, but such arrangement lacks long-range order. Thus, atomic bonding between crystalline and amorphous alloys is easily achievable. When Liu et al. joined Zr bulk metallic glass (BMG) with crystalline Al by explosive welding, they reported achieving atomic-scale bonding between these two materials [15]. To
T. Itoi et al. / Materials Characterization 118 (2016) 142–148
limit crystallisation and to suppress the formation of intermetallic compounds during joining of BMG with crystalline metallic materials, researchers have applied different welding methods, including electronbeam welding ultrasonic welding and FSW [16–18]. Results of these studies suggest that production of an amorphous alloy at the weld interface effectively suppresses surface embrittlement. However, advanced joining technology is required to control the heat input and to suppress the formation of intermetallic compounds in the formation of this amorphous alloy. The present study examined the feasibility of lap joining a Ni-coated Cu sheet and an Al sheet by MPW. The quality of the lap joints was evaluated by tensile tests at room temperature. The microstructure, especially that of an intermediate layer of the amorphous alloy formed at the Ni coating/Al bonding interface, was investigated by electron microscopy. The mechanism of formation of the amorphous alloy at the interface and the cause of the high interfacial strength of the joint are also discussed. 2. Experimental procedure Fig. 1(a) shows a schematic of the MPW circuit. A pure Cu sheet (100 × 80 × 0.6 mm3) coated with Ni (hereafter designated as Cu(Ni)) and an Al sheet (100 × 80 × 1.0 mm3) were used in the present study. Fig. 1(b) shows a scanning electron microscopy (SEM) image of the Ni coating on the Cu sheet. The thickness of the Ni coating was ~ 3 μm, and the Ni coating had a columnar structure of Ni of b1 μm in size. Both sheets were positioned over a coil, and a gap of 1.0 mm was maintained between them. The sheet closest to the coil is referred to as the ‘flyer sheet’, and the other sheet, which is firmly fixed in place, is referred to as the ‘parent sheet’. When an impulse current from the charged capacitor bank is passed through the coil, a high-density magnetic flux, which penetrates the flyer sheet, is generated around the coil. Eddy currents, which flow in the direction opposite to the impulse current in the coil, are induced in the surface layer of the flyer sheet. The generated eddy currents and high-density magnetic flux induce an upward electromagnetic force, which then drives the flyer sheet into the parent sheet, resulting in a high-velocity collision. In the present study, the Al sheet was used as the flyer sheet and Cu(Ni) was used as the parent sheet. The capacitor bank was maintained at 100 μF. The electrical energy stored in the capacitor is referred to as the ‘discharge energy’ (W), which is expressed as W = ½ CV2, where C is the capacitance of the capacitor and V is the applied voltage. The welding conditions for this experiment were controlled mainly by the discharge energy in the capacitor, which ranged from 0.7 to 2.0 kJ. Fig. 2(a) shows the macroscopic appearance of an Al/Cu(Ni) lap joint fabricated by MPW. The deformed region is indicated by the dotted area. Fig. 2(b) shows a cross-sectional view of the welded sample, where the upper part is the Cu(Ni) sheet and the lower part is the Al sheet. The width of the deformed region of the Al sheet corresponds to the width of the coil, which is approximately 5 mm. Bonding was achieved at
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two areas, as indicated by white circles in Fig. 2(b). No bonding occurred at the midsection of the deformed region. Samples for tensile tests were prepared by slicing the welded sheets using wire electrical discharge machining according to the dimensions indicated in Fig. 2(c). Tensile tests on the welded sheets were conducted at room temperature. The microstructure of the lap joint sheets was examined by SEM (SU6600, Hitachi), transmission electron microscopy (TEM), bright-field scanning transmission electron microscopy (BF-STEM) and high-angle annular dark field scanning transmission electron microscopy (HAADF-STEM) using a JEOL JEM-2100F. Indentation hardness tests of the weld interface were performed with a load of 1.96 × 10−3 N on a micro-Vickers hardness testing machine (HM-221; Mitsutoyo). 3. Results and discussions Fig. 3(a) displays a back-scattering electron image (BEI) of the lap joint sheet at a discharge energy of 2.0 kJ. The upper and lower sides correspond to the Cu(Ni) and Al sheets, respectively. As indicated by the white rectangles, two seam welded areas parallel to the side edges of the coil are formed at both sides with no bonded zone in between. Table 1 shows the weld length and weld strength of lap joint sheets welded at various discharge energies. The weld length refers to the total weld length of both sides of the bonded zones as determined by SEM. Joining was not possible at a discharge energy of 0.7 kJ or less. Therefore, the tensile test was performed only on lap joint sheets created with a discharge energy of more than 0.8 kJ. Tested specimens showing base-metal fracture and joint peeling at the joint interface are shown in Fig. 3(b) and (c), respectively. Tested specimens showing base-metal fracture or joint peeling at the joint interface are considered to indicate good or poor welding, respectively. Results of the tensile tests show that lap joint sheets prepared with a discharge energy of more than 0.9 kJ fractured at the base-metal Al sheet rather than at the bonded zone. Average failure loads of lap joints formed at a discharge energy above 0.9 kJ are N740 N, which is almost the same as the failure load of base-metal Al sheet. Thus, good welding was achieved at a discharge energy of more than 0.9 kJ. As shown in the Table 1, the weld length of the lap joint sheet tends to increase with the increase in discharge energy from 0.8 to 2.0 kJ. The weld strength of the lap joint welded at a discharge energy of 0.9 kJ, which experienced basemetal fracture, was estimated to be N120 MPa, which is the same as the tensile strength of the base-metal Al sheet. These results therefore suggest that increasing the weld length by increasing the discharge energy results in high joint strength. Thus, good welding is achieved at a discharge energy of N0.9 kJ. Fig. 4(a), (b) and (c) shows magnified BEIs of the jointed area (indicated by the white rectangle in Fig. 3(a)) formed at discharge energies of 0.8, 1.5 and 2.0 kJ, respectively. Joining with no wavy morphology at the weld interface was achieved at a discharge energy of 0.8 kJ. However, joint peeling occurred in the welded sample upon tensile testing, indicating poor welding. In contrast, samples welded at a discharge
Fig. 1. (a) Schematic of the MPW process. (b) Cross-sectional SEM image of the Ni-coated Cu sheet.
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Fig. 2. (a) Macroscopic structure of an Al/Cu(Ni) lap joint fabricated by MPW. (b) Cross-sectional view of the lap joint sheet. (c) A specimen for tensile testing.
energy of 1.5 and 2.0 kJ exhibited a wavy morphology similar to that observed in those subjected to explosive welding [5]. This wavy morphology is due to the deformation of the sheet surfaces, which is affected by the collision energy, the impact angle of the sheets and the geometry of the joint. This leads to an increased contact surface area of the sheet after MPW, resulting in strong bonding. The affected area resulting from the wavy shape is estimated to be about 10 μm from the interface in both lap joint sheets. This indicates that the bonding occurs at the outermost surface layer of these sheets. Collision speed for the Al sheet fabricated by the MPW was reported by Watanabe et al. which were used same apparatus and experimental conditions with those of the present study [19]. Their results indicated that the collision speed tend to increase with increasing discharge energy, and when the discharge energies are 1.0, 1.5, and 2.0 kJ, each collision speed indicated 98.8, 111.1 and 125.0 m/s. Furthermore, it was reported that the collision pressure is increased with increasing of collision speed, and collision pressure is estimated to be approximately 0.9 GPa when the discharge energy is 1.5 kJ. The wave amplitude and wave length of the
interfacial wave tend to increase with increasing of collision speed, because collision pressure increased. The increasing of wave amplitude brings both the increasing of contact area and the anchor effect for strong bonding between the interfaces. These results supported our results that the strong interfacial bonding is achieved by the wavy pattern formation by increasing of the discharge energy. Furthermore, intermediate layers with gray contrast were observed at the welding interface, as indicated by white arrows in Fig. 4(b) and (c). This result clearly shows that the volume fraction of the intermediate layer increases with increasing discharge energy. Formation of this intermediate layer is attributed to severe mechanical impact at the interface between the Ni coating and Al. The welding interface in impact welding methods usually exhibits two morphologies: a wavy interface and a melted layer. These morphologies have received much attention and have been thoroughly discussed [19–22]. Fig. 5(a) shows a BEI and energy-dispersive X-ray spectroscopy (EDS) mapping images of the intermediate layer of the welding interface of the lap joint sheet formed at a discharge energy
Fig. 3. (a) BEI of the cross-section of the lap joint sheet formed at a discharge energy of 2.0 kJ. Specimens after the tensile test showing (b) base-metal fracture (indicating good welding) and (c) joint peeling (indicating poor welding).
T. Itoi et al. / Materials Characterization 118 (2016) 142–148 Table 1 Weld length and weld strength of lap joint sheets welded at various discharge energies. After tensile testing of the lap joint sheets, the resulting base-metal fracture (Fig. 3(b)) or joint peeling (Fig. 3(c)) at the joint interface are considered to indicate good (denoted by ○) or poor weld strength (denoted by×), respectively. Discharge energy W (kJ) Weld length (mm) weld strength (O or×)
0.7 ̶ ̶
0.8 0.98 ×
0.9 1.31 O
1.0 1.57 O
1.5 1.75 O
2.0 2.23 O
of 2.0 kJ. As evident in the figure, the Ni coating of the Cu sheet was hammered and anchored into the Al matrix. The EDS mapping results also show that the intermediate layer consisted of a Ni/Al mixture, resembling Ni–Al binary intermetallic compounds. No clear oxide films formed on the surface of the weld interface. This finding indicates that the metal jet generated at the point of collision during the welding process removed any dust and oxide films on the sheet surfaces. This metal jet is normally achieved within a very short time, on the order of microseconds. The generated metal jet removes oxide films on both metal sheets, resulting in clean surfaces for welding [23]. Fig. 5(b) shows an SEM image of the weld interface region, a magnified image of the intermediate layer, and the micro-Vickers hardness at each position. The hardness values are indicated next to the indentations. As evident in the figure, the hardness of the Al matrix increased from 46 to 82 towards the intermediate layer. We obtained similar results for the Cu matrix, where the hardness increased from 68 to 86 towards the intermediate layer. These results indicate that the area within 10 μm from the interface underwent work hardening during the formation of the wave morphology in the welding process. This result is also supported
Fig. 4. BEIs of the joint interface of the lap joint sheet formed at discharge energies of (a) 0.8 kJ, (b) 1.5 kJ and (c) 2.0 kJ.
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by the SEM image in Fig. 4(c). The hardness value of the intermediate layer is remarkably high (348) and is greater than the hardness of both Cu and Al. Thus, the high hardness might be due to the formation of intermetallic compounds in the intermediate layer. To investigate the characteristics of the intermediate layer, a thin-film sample cut along the red line in Fig. 4(c) was machined by using a focused ion beam, and its microstructure was observed by TEM, BF-STEM and HAADF-STEM. Fig. 6(a) shows a BF-STEM image of the area around the intermediate layer. The columnar structure of the Ni coating was observed along the interface. Areas of dark contrast corresponding to the dislocation cell structure due to deformation caused by collision of the sheets were observed in both the Ni and Al grains. These results are also supported by the work hardening of the Ni coating and Al sheet around the interface (Fig. 5(b)). Fig. 6(b) shows a magnified BF-STEM image of the white square in Fig. 6(a), as well as points on the Al sheet (point 1), matrix of the intermediate layer (points 2–4), fragment with gray contrast embedded in the matrix (point 5), and Ni coating (point 6), on which EDS was conducted. Chemical compositions determined from the EDS spectra for these points are also shown in Fig. 6(b). Points 1 and 6 show that the Al sheet and Ni coating consisted almost entirely of Al and Ni, respectively. The compositions of the intermediate layer at points 2–4 are almost similar, including a Ni/Al mixture approximating Ni-35 at.% Al. On the other hand, the fragment embedded in the intermediate layer at point 5 consisted almost entirely of Ni. These fragments thus consisted of the Ni coating which has been partially dispersed in the intermediate layer. Fig. 7(a) shows an HAADF-STEM image of the intermediate layer captured from the same area of the BF-STEM image in Fig. 6(a). Metals with larger Z numbers such as Ni result in bright contrast, whereas metals with smaller Z numbers such as Al produce dark contrast. Therefore, the upper side with bright contrast is the Ni coating, while the lower side with dark contrast is the Al matrix. An intermediate layer with gray contrast was observed between the Ni coating and the Al matrix, as shown in Fig. 7(a). In addition, dispersed Ni fragments with white contrast are evident in the intermediate layer in the HAADFSTEM image. The large difference in the Ni content of the intermediate layer is due to the suppression of Ni diffusion to the matrix caused by rapid solidification; this result thus indicates rapid quenching after the melting of the Ni coating at the interface. That is, the Ni fragments observed in the intermediate layer are the melted remains of the Ni coating. Fig. 7(b) shows a magnified HAADF-STEM image of the area in the white square in Fig. 7(a). The dotted white line in Fig. 7(b) represents the contrast boundary between Al and the intermediate layer. The dark-contrast region consists of Al, and the region with gray and featureless contrast corresponds to the intermediate layer (Fig. 7(b)). Fig. 7(c) shows a BF-STEM image of the same area shown in Fig. 7(b). Fine grains are observed at the boundary between the intermediate layer and the Al region. Fig. 7(d) is a magnified BF-STEM image of the area in circle A in Fig. 7(c). The diffraction pattern taken from the finegrained area in Fig. 7(d) reveals that the fine grains with a diameter of ~10 nm consisted of Al. In contrast, the diffraction pattern in Fig. 7(e) reveals a halo ring. A high-resolution TEM (HR-TEM) image of this area shows a maze pattern near the Ni coating region, indicating the presence of an amorphous alloy. Ni3Al grains several nanometres in diameter were occasionally observed in the amorphous alloy matrix. We therefore conclude from these results that an intermediate layer consisting of Ni–Al binary amorphous alloy formed at the interface between the Ni coating and Al sheet during the MPW process. On the basis of the relationship between the mixing enthalpy (ΔHmix) and atomic radii of the components, Giessen et al. concluded that Zr- and Ni-based binary amorphous alloys prepared by a rapid quenching method exhibit GFA (glass forming ability) [24]. From their results, a contour separating GFA from non-GFA systems can be drawn. The Ni–Al binary amorphous alloy is located very close to the boundary between the GFA and non-GFA regions [23]. This proximity
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Fig. 5. (a) BEI and EDS mapping images of the intermediate layer of the welding interface. (b) Micro-Vickers hardness distribution around the weld interface region.
indicates that the Ni–Al alloy can become an amorphous alloy in the liquid state when rapidly cooled. Noya et al. investigated the GFA of Ni3Al, NiAl and NiAl3 alloys using molecular dynamics (MD) simulations [25]. Their results show that at cooling rates on the order of 1013 K/s, all three alloys can become amorphous alloys from the liquid state. Liu et al. welded Zr-based BMG to pure Al using explosive welding [15]. Their results show that atomic-scale bonding between BMG and crystalline Al can be achieved, with no indications of crystallization in the welding process. They conducted MD simulations based on a hydro-elastic–plastic model and found that in the explosive welding process, the rate of temperature increase is approximately 109 K/s and the rate of cooling is approximately 108 K/s [15]. The impact force in the MPW process is lower than that in explosive welding. However, in our experiments, melting and rapid quenching must have occurred at the weld interface between the Ni coating and Al. Conventional rapid-quenching cooling rates are approximately 105 to 106 K/s [24]. Therefore, we expect that
the cooling rate at the weld interface is greater than that for the conventional rapid quenching, consistent with the simulation results of Noya et al. [25] and Liu et al. [15] Fig. 8(a) and (b) respectively show HR-TEM images of the interfaces between the Ni coating and the amorphous alloy and between the amorphous alloy and the Al matrix. Observations of the interface indicate that atomic-scale bonding was achieved between the amorphous alloy and the Ni crystal lattice and between the amorphous alloy and the Al crystal lattice at the weld interface. The aforementioned results indicate that a strong lap joint can be achieved by MPW. The weld interface exhibited a wavy morphology, an anchor effect, a thin work-hardened layer around the intermediate layer, the absence of an oxide film, and suppressed intermetallic-compound formation by the formation of an amorphous alloy at the interface, characteristics that apparently contributed to the strong lap joint. Utilisation of the interfacial phenomena of MPW allows the formation of amorphous alloys due to localized melting and rapid cooling, which
Fig. 6. (a) BF-STEM image of the area around the weld interface. (b) Magnified BF-STEM image of the area indicated by the white square in Fig. 6(a) and EDS analysis points of the (1) Al sheet, (2–4) matrix of the intermediate layer, (5) fragment embedded in the matrix showing gray contrast, and (6) Ni coating. The beam diameter of these points was ~20 nm.
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Fig. 7. (a) HAADF-STEM image taken from the same area as in Fig. 6(a). (b) Magnified image of the area indicated by the white square in Fig. 7(a). (c) BF-STEM image of the same area shown in Fig. 7(b). (d) BF-STEM images of the area in circle A and (e) the area in circle B in Fig. 7(c).
suppresses the formation of brittle intermetallic compounds that usually form when dissimilar metals are joined. In addition, the impact force generated by MPW affects the surface structure within several micrometres of the interface. This characteristic of MPW can thus potentially be utilised for the surface treatment of dissimilar materials such as Al and Cu, through deliberate lap joining of surface-treated materials such as Ni-coated Cu to form an amorphous alloy at the intermediate layer. This layer can suppress the formation of brittle intermetallic compounds at the weld interface. According to our experimental results, MPW, which involves a high cooling rate, can produce simple and classic binary amorphous alloys without BMG components [26], which can also form at low cooling rates, at the interface of the weld junction. This approach represents a unique preparation method that may be useful in engineering applications for the welding of dissimilar metals.
4. Conclusions A good lap joint between Al and Ni-coated Cu sheets was achieved by MPW at a discharge energy of N0.9 kJ. A wavy morphology and an intermediate layer along the weld interface were observed in the lapjointed sheets. Micro-Vickers hardness results indicate that the area within 10 μm of the Al and Cu interface underwent work hardening during the formation of the wavy morphology in the welding process. SEM revealed no oxide film between the surfaces of the Al and Cu sheets and the intermediate layer. TEM, BF-STEM and HAADF-STEM observations revealed a Ni–Al binary amorphous alloy and Ni fragments embedded in the intermediate layer near the Ni-coated region. The Al sheet close to the weld interface showed Al nanograins. This result suggests that formation of an intermediate layer consisting of Ni–Al binary amorphous alloy and Al nanograins formed at the Al sheet surface are due
Fig. 8. HR-TEM images of the interfaces (a) between the Al matrix and the amorphous alloy and (b) between the amorphous alloy and the Ni coating.
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to the rapid solidification of Al and melting of the Ni coating during the MPW process. Therefore, the good joining of Al and Ni-coated Cu sheets is due to the increased contact area, anchor effect, work hardening, absence of an oxide film, and suppressed formation of intermetallic compounds at the interface. Acknowledgments This work was partly financed by Grant-aided Project (2015-2016) from the Light Metal Education Foundation of Japan. References [1] P. Bergmann, F. Petzoldt, R. Schuerer, S. Shneider, Solid-state welding to aluminium to copper, Weld. World 57 (2013) 541–550. [2] M. Abbasi, A.K. Taheri, M.T. Salehi, Growth rate of intermetallic compounds in Al/Cu bimetal produced by cold roll welding process, J. Alloys Compd. 319 (2001) 233–241. [3] C. Xia, Y. Li, U.A. Puchkov, S.A. Gerasimov, Microstructure and phase constitution near the interface of Al/Cu vacuum brazing using Al-Si filler metal, Vacuum 82 (2008) 799–804. [4] M. Asemabadi, M. Sedighi, M. Honarpisheh, Investigation of cold rolling influence on the mechanical properties of explosive weld Al/Cu bimetal, Mater. Sci. Eng. A 558 (2012) 144–149. [5] F. Findik, Recent developments in explosive welding, Mater. Des. 32 (2011) 1081–1093. [6] W.B. Lee, K.S. Bang, S.B. Jung, Effects of intermetallic compound on the electrical and mechanical properties of friction welded Cu/Al bimetallic joints during annealing, J. Alloys Compd. 390 (2005) 212–219. [7] T. Aizawa, K. Okagawa, M. Yoshizawa, N. Henmi, Impulse magnetic pressure seam welding of aluminium sheets, Proc. 4th Int. Symp. Inpact. Eng. 2001, pp. 827–831. [8] T. Aizawa, M. Kashani, K. Okagawa, Application of magnetic pulse welding for aluminum alloys and SPCC steel sheet joints, Weld. J. 86 (2007) 119–124. [9] T. Aizawa, K. Okagawa, N. Henmi, Impulse magnetic pressure seam welding of aluminum, copper and steel sheets, Adv. Technol. Plast. 2 (2002) 1687–1692.
[10] A. Stern, A. Aizenshtein, Bonding zone formation in magnetic pulse welding, Sci. Technol. Weld. Join. 7 (2002) 339–342. [11] K.J. Lee, S. Kumai, T. Arai, T. Aizawa, Interfacial microstructure and strength of steel/ aluminum alloy lap joint fabricated by magnetic pressure seam welding, Mater. Sci. Eng. A 471 (2007) 95–101. [12] Y.C. Chen, T. Komazaki, T. Tsumura, K. Nakata, Role of zinc coated in friction stir welding Al and Zn coated steel, Mater. Sci. Technol. 24 (2008) 33–39. [13] Y.C. Chen, K. Nakata, Effect of the surface state of steel on the microstructure and mechanical properties of dissimilar metal lap joints of aluminum and steel by friction stir welding, Metall. Mater. Trans. A 39 (2008) 1985–1992. [14] R.T. Lee, C.R. Liu, Y.C. Chiou, H.L. Chen, Effect of nickel coating on the shear strength of FSW lap joint between Ni-Cu alloy and steel, J. Mater. Process. Technol. 213 (2013) 69–74. [15] K.X. Liu, W.D. Liu, J.T. Wang, H.H. Yan, J. Li, Y.J. Huang, X.S. Wei, J. Shen, Atomic-scale bonding of bulk metallic glass to crystalline aluminum, Appl. Phys. Lett. 93 (2008) (081918–3). [16] Y. Kawamura, Y. Ohno, Successful electron-beam welding of bulk metallic glass, Mater. Trans. JIM 42 (2001) 2676–2678. [17] D. Wang, B.L. Xiao, Z.Y. Ma, H.F. Zhang, Friction stir welding of Zr55Cu30Al10Ni5 bulk metallic glass to Al-Zn-Mg-Cu alloy, Scr. Mater. 60 (2009) 112–115. [18] M. Maeda, Y. Takahashi, M. Fukuhara, X. Whang, A. Inoue, Ultrasonic bonding of Zr55Cu30Al10Ni5 metallic glass, Mater. Sci. Eng. B 148 (2008) 141–144. [19] M. Watanabe, S. Kumai, Interfacial morphology of magnetic pulse welded Aluminum/Aluminum and Copper/Copper lap joints, Mater. Trans. 50 (2009) 286–2292. [20] M. Acarer, B. Gulenc, F. Findik, Investigation of explosive welding parameters and their effects on microhardness and shear strength, Mater. Des. 24 (2003) 659–664. [21] N. Kahraman, B. Gulenc, F. Findik, Joining of titanium/stainless steel by explosive welding and effect on interface, J. Mater. Process. Technol. 169 (2005) 127–133. [22] S.Y. Chen, Z.W. Wu, K.X. Liu, X.J. Li, N. Luo, G.X. Lu, Atomic diffusion behavior in CuAl explosive welding process, J. Appl. Phys. 113 (2013) 044901–044906. [23] M. Watanabe, S. Kumai, High-speed deformation and collision behavior of pure aluminium plates in magnetic pulse welding, Mater. Trans. 50 (2009) 2035–2042. [24] B.C. Giessen, A two-parameter representation of rapidly glass forming binary alloy systems, Proc. 4th Int. Conf. Rapidly Quenched Metals. 1981, pp. 213–216. [25] E.G. Noya, C. Rey, L.L. Gallego, Amorphization of Ni-Al alloys by fast quenching from the liquid state: a molecular dynamics, J. Non-Cryst. Solids 298 (2002) 60–66. [26] A. Peker, W.L. Johnson, A highly processable metallic glass: Zr41·2Ti13.8Cu12·5Ni10.0Be22.5, Appl. Phys. Lett. 63 (1993) 2342–2345.
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Materials Characterization journal homepage: www.elsevier.com/locate/matchar
Microstructure evolution of a dissimilar junction interface between an Al sheet and a Ni-coated Cu sheet joined by magnetic pulse welding Takaomi Itoi a,⁎, Azizan Bin Mohamad a, Ryo Suzuki a, Keigo Okagawa b a b
Department of Mechanical Engineering, Chiba University, 1-33, Yayoi-cho, Inage-ku, Chiba 263-8522, Japan Department of Electrical and Electronics Engineering, Tokyo Metropolitan College of Industrial Technology, 1-10-40 Higashi ohi, Shinagawa-ku, Tokyo 140-0011, Japan
a r t i c l e
i n f o
Article history: Received 22 March 2016 Received in revised form 25 May 2016 Accepted 25 May 2016 Available online 26 May 2016 Keywords: Microstructure Amorphous alloy TEM Magnetic pulse welding Joining
a b s t r a c t An Al sheet and a Ni-coated Cu sheet were lap joined by using magnetic pulse welding (MPW). Tensile tests were performed on the joined sheets, and a good lap joint was achieved at a discharge energy of N 0.9 kJ. The weld interface exhibited a wavy morphology and an intermediate layer along the weld interface. Microstructure observations of the intermediate layer revealed that the Ni coating region consisted of a Ni–Al binary amorphous alloy and that the Al sheet region contained very fine Al nanograins. Ni fragments indicative of unmelted residual Ni from the coating were also observed in parts of the intermediate layer. Formation of these features can be attributed to localize melting and a subsequent high rate cooling of molten Al and Ni confined to the interface during the MPW process. In the absence of an oxide film, atomic-scale bonding was also achieved between the intermediate layer and the sheet surfaces after the collision. MPW utilises impact energy, which affects the sheet surfaces. From the obtained results, good lap joint is attributed to an increased contact area, the anchor effect, work hardening, the absence of an oxide film, and suppressed formation of intermetallic compounds at the interface. © 2016 Elsevier Inc. All rights reserved.
1. Introduction The development of welding methods for combinations of dissimilar materials for industrial applications has gained importance because of the economic and technical advantages of such methods. Cu, which exhibits good electrical and thermal conductivity, and Al, which exhibits low density, are two common engineering materials widely used for reducing the cost and weight of various industrial products, especially in automotive applications [1]. The joining of dissimilar metals through the conventional fusion welding method results in the formation of a brittle intermetallic compound layer, which leads to failure at the joint interface under applied external stress. Furthermore, Cu and Al cannot be joined by conventional fusion welding because of large differences in their physical and mechanical properties. Therefore, solid-state welding methods such as cold-roll welding, diffusion welding, explosive welding and friction-stir welding (FSW) have been proposed as alternative methods for joining dissimilar metals with different properties [2– 6]. Magnetic pulse welding (MPW), a solid-state welding method similar to impact welding processes such as explosive welding, is applicable to metals with high electrical conductivity. Impact energy is induced by the electromagnetic force generated by the interaction between the discharge energy (impulse current), high-density magnetic flux and eddy currents produced in the sheet. This electromagnetic force drives the ⁎ Corresponding author. E-mail address: [email protected] (T. Itoi).
http://dx.doi.org/10.1016/j.matchar.2016.05.021 1044-5803/© 2016 Elsevier Inc. All rights reserved.
flyer sheet to the parent sheet, resulting in a high-velocity collision and solid-state welding within several microseconds [7–9]. Recently, Al/Cu and Fe/Al lap joints have been successfully fabricated by using MPW [9–11]. In practical applications, metal sheets are surface-treated by coating with a layer of several metals such as nickel to increase resistance against corrosion. Welding for these sheets with metal coatings have been fabricated by solid-state welding methods such as FSW [12–14]. However, the fabrication of such welding for sheets with metal coatings by MPW has not yet been reported. During the MPW process, localized melting occurs at the interface when the discharge energy is high and intermetallic compounds partially form at the weld interface, as in the case of the interfaces in Al/Cu and Fe/Al lap joints [9–11]. Localized melting occurs because of the high pressure and temperature after the high-speed collision of both sheets. This particular phenomenon, which encompasses the formation of amorphous alloys and nanograins of intermetallic compounds at the interface created by joining dissimilar sheets, also occurs in other impact welding mechanisms such as that for explosive welding. However, the detailed microstructure and mechanism of formation of amorphous alloys at the interface during MPW is still unclear. Amorphous alloys exhibit high strength and toughness, as well as advantageous mechanical properties. The atoms in amorphous alloys are ordered arrays over short ranges, but such arrangement lacks long-range order. Thus, atomic bonding between crystalline and amorphous alloys is easily achievable. When Liu et al. joined Zr bulk metallic glass (BMG) with crystalline Al by explosive welding, they reported achieving atomic-scale bonding between these two materials [15]. To
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limit crystallisation and to suppress the formation of intermetallic compounds during joining of BMG with crystalline metallic materials, researchers have applied different welding methods, including electronbeam welding ultrasonic welding and FSW [16–18]. Results of these studies suggest that production of an amorphous alloy at the weld interface effectively suppresses surface embrittlement. However, advanced joining technology is required to control the heat input and to suppress the formation of intermetallic compounds in the formation of this amorphous alloy. The present study examined the feasibility of lap joining a Ni-coated Cu sheet and an Al sheet by MPW. The quality of the lap joints was evaluated by tensile tests at room temperature. The microstructure, especially that of an intermediate layer of the amorphous alloy formed at the Ni coating/Al bonding interface, was investigated by electron microscopy. The mechanism of formation of the amorphous alloy at the interface and the cause of the high interfacial strength of the joint are also discussed. 2. Experimental procedure Fig. 1(a) shows a schematic of the MPW circuit. A pure Cu sheet (100 × 80 × 0.6 mm3) coated with Ni (hereafter designated as Cu(Ni)) and an Al sheet (100 × 80 × 1.0 mm3) were used in the present study. Fig. 1(b) shows a scanning electron microscopy (SEM) image of the Ni coating on the Cu sheet. The thickness of the Ni coating was ~ 3 μm, and the Ni coating had a columnar structure of Ni of b1 μm in size. Both sheets were positioned over a coil, and a gap of 1.0 mm was maintained between them. The sheet closest to the coil is referred to as the ‘flyer sheet’, and the other sheet, which is firmly fixed in place, is referred to as the ‘parent sheet’. When an impulse current from the charged capacitor bank is passed through the coil, a high-density magnetic flux, which penetrates the flyer sheet, is generated around the coil. Eddy currents, which flow in the direction opposite to the impulse current in the coil, are induced in the surface layer of the flyer sheet. The generated eddy currents and high-density magnetic flux induce an upward electromagnetic force, which then drives the flyer sheet into the parent sheet, resulting in a high-velocity collision. In the present study, the Al sheet was used as the flyer sheet and Cu(Ni) was used as the parent sheet. The capacitor bank was maintained at 100 μF. The electrical energy stored in the capacitor is referred to as the ‘discharge energy’ (W), which is expressed as W = ½ CV2, where C is the capacitance of the capacitor and V is the applied voltage. The welding conditions for this experiment were controlled mainly by the discharge energy in the capacitor, which ranged from 0.7 to 2.0 kJ. Fig. 2(a) shows the macroscopic appearance of an Al/Cu(Ni) lap joint fabricated by MPW. The deformed region is indicated by the dotted area. Fig. 2(b) shows a cross-sectional view of the welded sample, where the upper part is the Cu(Ni) sheet and the lower part is the Al sheet. The width of the deformed region of the Al sheet corresponds to the width of the coil, which is approximately 5 mm. Bonding was achieved at
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two areas, as indicated by white circles in Fig. 2(b). No bonding occurred at the midsection of the deformed region. Samples for tensile tests were prepared by slicing the welded sheets using wire electrical discharge machining according to the dimensions indicated in Fig. 2(c). Tensile tests on the welded sheets were conducted at room temperature. The microstructure of the lap joint sheets was examined by SEM (SU6600, Hitachi), transmission electron microscopy (TEM), bright-field scanning transmission electron microscopy (BF-STEM) and high-angle annular dark field scanning transmission electron microscopy (HAADF-STEM) using a JEOL JEM-2100F. Indentation hardness tests of the weld interface were performed with a load of 1.96 × 10−3 N on a micro-Vickers hardness testing machine (HM-221; Mitsutoyo). 3. Results and discussions Fig. 3(a) displays a back-scattering electron image (BEI) of the lap joint sheet at a discharge energy of 2.0 kJ. The upper and lower sides correspond to the Cu(Ni) and Al sheets, respectively. As indicated by the white rectangles, two seam welded areas parallel to the side edges of the coil are formed at both sides with no bonded zone in between. Table 1 shows the weld length and weld strength of lap joint sheets welded at various discharge energies. The weld length refers to the total weld length of both sides of the bonded zones as determined by SEM. Joining was not possible at a discharge energy of 0.7 kJ or less. Therefore, the tensile test was performed only on lap joint sheets created with a discharge energy of more than 0.8 kJ. Tested specimens showing base-metal fracture and joint peeling at the joint interface are shown in Fig. 3(b) and (c), respectively. Tested specimens showing base-metal fracture or joint peeling at the joint interface are considered to indicate good or poor welding, respectively. Results of the tensile tests show that lap joint sheets prepared with a discharge energy of more than 0.9 kJ fractured at the base-metal Al sheet rather than at the bonded zone. Average failure loads of lap joints formed at a discharge energy above 0.9 kJ are N740 N, which is almost the same as the failure load of base-metal Al sheet. Thus, good welding was achieved at a discharge energy of more than 0.9 kJ. As shown in the Table 1, the weld length of the lap joint sheet tends to increase with the increase in discharge energy from 0.8 to 2.0 kJ. The weld strength of the lap joint welded at a discharge energy of 0.9 kJ, which experienced basemetal fracture, was estimated to be N120 MPa, which is the same as the tensile strength of the base-metal Al sheet. These results therefore suggest that increasing the weld length by increasing the discharge energy results in high joint strength. Thus, good welding is achieved at a discharge energy of N0.9 kJ. Fig. 4(a), (b) and (c) shows magnified BEIs of the jointed area (indicated by the white rectangle in Fig. 3(a)) formed at discharge energies of 0.8, 1.5 and 2.0 kJ, respectively. Joining with no wavy morphology at the weld interface was achieved at a discharge energy of 0.8 kJ. However, joint peeling occurred in the welded sample upon tensile testing, indicating poor welding. In contrast, samples welded at a discharge
Fig. 1. (a) Schematic of the MPW process. (b) Cross-sectional SEM image of the Ni-coated Cu sheet.
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Fig. 2. (a) Macroscopic structure of an Al/Cu(Ni) lap joint fabricated by MPW. (b) Cross-sectional view of the lap joint sheet. (c) A specimen for tensile testing.
energy of 1.5 and 2.0 kJ exhibited a wavy morphology similar to that observed in those subjected to explosive welding [5]. This wavy morphology is due to the deformation of the sheet surfaces, which is affected by the collision energy, the impact angle of the sheets and the geometry of the joint. This leads to an increased contact surface area of the sheet after MPW, resulting in strong bonding. The affected area resulting from the wavy shape is estimated to be about 10 μm from the interface in both lap joint sheets. This indicates that the bonding occurs at the outermost surface layer of these sheets. Collision speed for the Al sheet fabricated by the MPW was reported by Watanabe et al. which were used same apparatus and experimental conditions with those of the present study [19]. Their results indicated that the collision speed tend to increase with increasing discharge energy, and when the discharge energies are 1.0, 1.5, and 2.0 kJ, each collision speed indicated 98.8, 111.1 and 125.0 m/s. Furthermore, it was reported that the collision pressure is increased with increasing of collision speed, and collision pressure is estimated to be approximately 0.9 GPa when the discharge energy is 1.5 kJ. The wave amplitude and wave length of the
interfacial wave tend to increase with increasing of collision speed, because collision pressure increased. The increasing of wave amplitude brings both the increasing of contact area and the anchor effect for strong bonding between the interfaces. These results supported our results that the strong interfacial bonding is achieved by the wavy pattern formation by increasing of the discharge energy. Furthermore, intermediate layers with gray contrast were observed at the welding interface, as indicated by white arrows in Fig. 4(b) and (c). This result clearly shows that the volume fraction of the intermediate layer increases with increasing discharge energy. Formation of this intermediate layer is attributed to severe mechanical impact at the interface between the Ni coating and Al. The welding interface in impact welding methods usually exhibits two morphologies: a wavy interface and a melted layer. These morphologies have received much attention and have been thoroughly discussed [19–22]. Fig. 5(a) shows a BEI and energy-dispersive X-ray spectroscopy (EDS) mapping images of the intermediate layer of the welding interface of the lap joint sheet formed at a discharge energy
Fig. 3. (a) BEI of the cross-section of the lap joint sheet formed at a discharge energy of 2.0 kJ. Specimens after the tensile test showing (b) base-metal fracture (indicating good welding) and (c) joint peeling (indicating poor welding).
T. Itoi et al. / Materials Characterization 118 (2016) 142–148 Table 1 Weld length and weld strength of lap joint sheets welded at various discharge energies. After tensile testing of the lap joint sheets, the resulting base-metal fracture (Fig. 3(b)) or joint peeling (Fig. 3(c)) at the joint interface are considered to indicate good (denoted by ○) or poor weld strength (denoted by×), respectively. Discharge energy W (kJ) Weld length (mm) weld strength (O or×)
0.7 ̶ ̶
0.8 0.98 ×
0.9 1.31 O
1.0 1.57 O
1.5 1.75 O
2.0 2.23 O
of 2.0 kJ. As evident in the figure, the Ni coating of the Cu sheet was hammered and anchored into the Al matrix. The EDS mapping results also show that the intermediate layer consisted of a Ni/Al mixture, resembling Ni–Al binary intermetallic compounds. No clear oxide films formed on the surface of the weld interface. This finding indicates that the metal jet generated at the point of collision during the welding process removed any dust and oxide films on the sheet surfaces. This metal jet is normally achieved within a very short time, on the order of microseconds. The generated metal jet removes oxide films on both metal sheets, resulting in clean surfaces for welding [23]. Fig. 5(b) shows an SEM image of the weld interface region, a magnified image of the intermediate layer, and the micro-Vickers hardness at each position. The hardness values are indicated next to the indentations. As evident in the figure, the hardness of the Al matrix increased from 46 to 82 towards the intermediate layer. We obtained similar results for the Cu matrix, where the hardness increased from 68 to 86 towards the intermediate layer. These results indicate that the area within 10 μm from the interface underwent work hardening during the formation of the wave morphology in the welding process. This result is also supported
Fig. 4. BEIs of the joint interface of the lap joint sheet formed at discharge energies of (a) 0.8 kJ, (b) 1.5 kJ and (c) 2.0 kJ.
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by the SEM image in Fig. 4(c). The hardness value of the intermediate layer is remarkably high (348) and is greater than the hardness of both Cu and Al. Thus, the high hardness might be due to the formation of intermetallic compounds in the intermediate layer. To investigate the characteristics of the intermediate layer, a thin-film sample cut along the red line in Fig. 4(c) was machined by using a focused ion beam, and its microstructure was observed by TEM, BF-STEM and HAADF-STEM. Fig. 6(a) shows a BF-STEM image of the area around the intermediate layer. The columnar structure of the Ni coating was observed along the interface. Areas of dark contrast corresponding to the dislocation cell structure due to deformation caused by collision of the sheets were observed in both the Ni and Al grains. These results are also supported by the work hardening of the Ni coating and Al sheet around the interface (Fig. 5(b)). Fig. 6(b) shows a magnified BF-STEM image of the white square in Fig. 6(a), as well as points on the Al sheet (point 1), matrix of the intermediate layer (points 2–4), fragment with gray contrast embedded in the matrix (point 5), and Ni coating (point 6), on which EDS was conducted. Chemical compositions determined from the EDS spectra for these points are also shown in Fig. 6(b). Points 1 and 6 show that the Al sheet and Ni coating consisted almost entirely of Al and Ni, respectively. The compositions of the intermediate layer at points 2–4 are almost similar, including a Ni/Al mixture approximating Ni-35 at.% Al. On the other hand, the fragment embedded in the intermediate layer at point 5 consisted almost entirely of Ni. These fragments thus consisted of the Ni coating which has been partially dispersed in the intermediate layer. Fig. 7(a) shows an HAADF-STEM image of the intermediate layer captured from the same area of the BF-STEM image in Fig. 6(a). Metals with larger Z numbers such as Ni result in bright contrast, whereas metals with smaller Z numbers such as Al produce dark contrast. Therefore, the upper side with bright contrast is the Ni coating, while the lower side with dark contrast is the Al matrix. An intermediate layer with gray contrast was observed between the Ni coating and the Al matrix, as shown in Fig. 7(a). In addition, dispersed Ni fragments with white contrast are evident in the intermediate layer in the HAADFSTEM image. The large difference in the Ni content of the intermediate layer is due to the suppression of Ni diffusion to the matrix caused by rapid solidification; this result thus indicates rapid quenching after the melting of the Ni coating at the interface. That is, the Ni fragments observed in the intermediate layer are the melted remains of the Ni coating. Fig. 7(b) shows a magnified HAADF-STEM image of the area in the white square in Fig. 7(a). The dotted white line in Fig. 7(b) represents the contrast boundary between Al and the intermediate layer. The dark-contrast region consists of Al, and the region with gray and featureless contrast corresponds to the intermediate layer (Fig. 7(b)). Fig. 7(c) shows a BF-STEM image of the same area shown in Fig. 7(b). Fine grains are observed at the boundary between the intermediate layer and the Al region. Fig. 7(d) is a magnified BF-STEM image of the area in circle A in Fig. 7(c). The diffraction pattern taken from the finegrained area in Fig. 7(d) reveals that the fine grains with a diameter of ~10 nm consisted of Al. In contrast, the diffraction pattern in Fig. 7(e) reveals a halo ring. A high-resolution TEM (HR-TEM) image of this area shows a maze pattern near the Ni coating region, indicating the presence of an amorphous alloy. Ni3Al grains several nanometres in diameter were occasionally observed in the amorphous alloy matrix. We therefore conclude from these results that an intermediate layer consisting of Ni–Al binary amorphous alloy formed at the interface between the Ni coating and Al sheet during the MPW process. On the basis of the relationship between the mixing enthalpy (ΔHmix) and atomic radii of the components, Giessen et al. concluded that Zr- and Ni-based binary amorphous alloys prepared by a rapid quenching method exhibit GFA (glass forming ability) [24]. From their results, a contour separating GFA from non-GFA systems can be drawn. The Ni–Al binary amorphous alloy is located very close to the boundary between the GFA and non-GFA regions [23]. This proximity
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Fig. 5. (a) BEI and EDS mapping images of the intermediate layer of the welding interface. (b) Micro-Vickers hardness distribution around the weld interface region.
indicates that the Ni–Al alloy can become an amorphous alloy in the liquid state when rapidly cooled. Noya et al. investigated the GFA of Ni3Al, NiAl and NiAl3 alloys using molecular dynamics (MD) simulations [25]. Their results show that at cooling rates on the order of 1013 K/s, all three alloys can become amorphous alloys from the liquid state. Liu et al. welded Zr-based BMG to pure Al using explosive welding [15]. Their results show that atomic-scale bonding between BMG and crystalline Al can be achieved, with no indications of crystallization in the welding process. They conducted MD simulations based on a hydro-elastic–plastic model and found that in the explosive welding process, the rate of temperature increase is approximately 109 K/s and the rate of cooling is approximately 108 K/s [15]. The impact force in the MPW process is lower than that in explosive welding. However, in our experiments, melting and rapid quenching must have occurred at the weld interface between the Ni coating and Al. Conventional rapid-quenching cooling rates are approximately 105 to 106 K/s [24]. Therefore, we expect that
the cooling rate at the weld interface is greater than that for the conventional rapid quenching, consistent with the simulation results of Noya et al. [25] and Liu et al. [15] Fig. 8(a) and (b) respectively show HR-TEM images of the interfaces between the Ni coating and the amorphous alloy and between the amorphous alloy and the Al matrix. Observations of the interface indicate that atomic-scale bonding was achieved between the amorphous alloy and the Ni crystal lattice and between the amorphous alloy and the Al crystal lattice at the weld interface. The aforementioned results indicate that a strong lap joint can be achieved by MPW. The weld interface exhibited a wavy morphology, an anchor effect, a thin work-hardened layer around the intermediate layer, the absence of an oxide film, and suppressed intermetallic-compound formation by the formation of an amorphous alloy at the interface, characteristics that apparently contributed to the strong lap joint. Utilisation of the interfacial phenomena of MPW allows the formation of amorphous alloys due to localized melting and rapid cooling, which
Fig. 6. (a) BF-STEM image of the area around the weld interface. (b) Magnified BF-STEM image of the area indicated by the white square in Fig. 6(a) and EDS analysis points of the (1) Al sheet, (2–4) matrix of the intermediate layer, (5) fragment embedded in the matrix showing gray contrast, and (6) Ni coating. The beam diameter of these points was ~20 nm.
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Fig. 7. (a) HAADF-STEM image taken from the same area as in Fig. 6(a). (b) Magnified image of the area indicated by the white square in Fig. 7(a). (c) BF-STEM image of the same area shown in Fig. 7(b). (d) BF-STEM images of the area in circle A and (e) the area in circle B in Fig. 7(c).
suppresses the formation of brittle intermetallic compounds that usually form when dissimilar metals are joined. In addition, the impact force generated by MPW affects the surface structure within several micrometres of the interface. This characteristic of MPW can thus potentially be utilised for the surface treatment of dissimilar materials such as Al and Cu, through deliberate lap joining of surface-treated materials such as Ni-coated Cu to form an amorphous alloy at the intermediate layer. This layer can suppress the formation of brittle intermetallic compounds at the weld interface. According to our experimental results, MPW, which involves a high cooling rate, can produce simple and classic binary amorphous alloys without BMG components [26], which can also form at low cooling rates, at the interface of the weld junction. This approach represents a unique preparation method that may be useful in engineering applications for the welding of dissimilar metals.
4. Conclusions A good lap joint between Al and Ni-coated Cu sheets was achieved by MPW at a discharge energy of N0.9 kJ. A wavy morphology and an intermediate layer along the weld interface were observed in the lapjointed sheets. Micro-Vickers hardness results indicate that the area within 10 μm of the Al and Cu interface underwent work hardening during the formation of the wavy morphology in the welding process. SEM revealed no oxide film between the surfaces of the Al and Cu sheets and the intermediate layer. TEM, BF-STEM and HAADF-STEM observations revealed a Ni–Al binary amorphous alloy and Ni fragments embedded in the intermediate layer near the Ni-coated region. The Al sheet close to the weld interface showed Al nanograins. This result suggests that formation of an intermediate layer consisting of Ni–Al binary amorphous alloy and Al nanograins formed at the Al sheet surface are due
Fig. 8. HR-TEM images of the interfaces (a) between the Al matrix and the amorphous alloy and (b) between the amorphous alloy and the Ni coating.
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to the rapid solidification of Al and melting of the Ni coating during the MPW process. Therefore, the good joining of Al and Ni-coated Cu sheets is due to the increased contact area, anchor effect, work hardening, absence of an oxide film, and suppressed formation of intermetallic compounds at the interface. Acknowledgments This work was partly financed by Grant-aided Project (2015-2016) from the Light Metal Education Foundation of Japan. References [1] P. Bergmann, F. Petzoldt, R. Schuerer, S. Shneider, Solid-state welding to aluminium to copper, Weld. World 57 (2013) 541–550. [2] M. Abbasi, A.K. Taheri, M.T. Salehi, Growth rate of intermetallic compounds in Al/Cu bimetal produced by cold roll welding process, J. Alloys Compd. 319 (2001) 233–241. [3] C. Xia, Y. Li, U.A. Puchkov, S.A. Gerasimov, Microstructure and phase constitution near the interface of Al/Cu vacuum brazing using Al-Si filler metal, Vacuum 82 (2008) 799–804. [4] M. Asemabadi, M. Sedighi, M. Honarpisheh, Investigation of cold rolling influence on the mechanical properties of explosive weld Al/Cu bimetal, Mater. Sci. Eng. A 558 (2012) 144–149. [5] F. Findik, Recent developments in explosive welding, Mater. Des. 32 (2011) 1081–1093. [6] W.B. Lee, K.S. Bang, S.B. Jung, Effects of intermetallic compound on the electrical and mechanical properties of friction welded Cu/Al bimetallic joints during annealing, J. Alloys Compd. 390 (2005) 212–219. [7] T. Aizawa, K. Okagawa, M. Yoshizawa, N. Henmi, Impulse magnetic pressure seam welding of aluminium sheets, Proc. 4th Int. Symp. Inpact. Eng. 2001, pp. 827–831. [8] T. Aizawa, M. Kashani, K. Okagawa, Application of magnetic pulse welding for aluminum alloys and SPCC steel sheet joints, Weld. J. 86 (2007) 119–124. [9] T. Aizawa, K. Okagawa, N. Henmi, Impulse magnetic pressure seam welding of aluminum, copper and steel sheets, Adv. Technol. Plast. 2 (2002) 1687–1692.
[10] A. Stern, A. Aizenshtein, Bonding zone formation in magnetic pulse welding, Sci. Technol. Weld. Join. 7 (2002) 339–342. [11] K.J. Lee, S. Kumai, T. Arai, T. Aizawa, Interfacial microstructure and strength of steel/ aluminum alloy lap joint fabricated by magnetic pressure seam welding, Mater. Sci. Eng. A 471 (2007) 95–101. [12] Y.C. Chen, T. Komazaki, T. Tsumura, K. Nakata, Role of zinc coated in friction stir welding Al and Zn coated steel, Mater. Sci. Technol. 24 (2008) 33–39. [13] Y.C. Chen, K. Nakata, Effect of the surface state of steel on the microstructure and mechanical properties of dissimilar metal lap joints of aluminum and steel by friction stir welding, Metall. Mater. Trans. A 39 (2008) 1985–1992. [14] R.T. Lee, C.R. Liu, Y.C. Chiou, H.L. Chen, Effect of nickel coating on the shear strength of FSW lap joint between Ni-Cu alloy and steel, J. Mater. Process. Technol. 213 (2013) 69–74. [15] K.X. Liu, W.D. Liu, J.T. Wang, H.H. Yan, J. Li, Y.J. Huang, X.S. Wei, J. Shen, Atomic-scale bonding of bulk metallic glass to crystalline aluminum, Appl. Phys. Lett. 93 (2008) (081918–3). [16] Y. Kawamura, Y. Ohno, Successful electron-beam welding of bulk metallic glass, Mater. Trans. JIM 42 (2001) 2676–2678. [17] D. Wang, B.L. Xiao, Z.Y. Ma, H.F. Zhang, Friction stir welding of Zr55Cu30Al10Ni5 bulk metallic glass to Al-Zn-Mg-Cu alloy, Scr. Mater. 60 (2009) 112–115. [18] M. Maeda, Y. Takahashi, M. Fukuhara, X. Whang, A. Inoue, Ultrasonic bonding of Zr55Cu30Al10Ni5 metallic glass, Mater. Sci. Eng. B 148 (2008) 141–144. [19] M. Watanabe, S. Kumai, Interfacial morphology of magnetic pulse welded Aluminum/Aluminum and Copper/Copper lap joints, Mater. Trans. 50 (2009) 286–2292. [20] M. Acarer, B. Gulenc, F. Findik, Investigation of explosive welding parameters and their effects on microhardness and shear strength, Mater. Des. 24 (2003) 659–664. [21] N. Kahraman, B. Gulenc, F. Findik, Joining of titanium/stainless steel by explosive welding and effect on interface, J. Mater. Process. Technol. 169 (2005) 127–133. [22] S.Y. Chen, Z.W. Wu, K.X. Liu, X.J. Li, N. Luo, G.X. Lu, Atomic diffusion behavior in CuAl explosive welding process, J. Appl. Phys. 113 (2013) 044901–044906. [23] M. Watanabe, S. Kumai, High-speed deformation and collision behavior of pure aluminium plates in magnetic pulse welding, Mater. Trans. 50 (2009) 2035–2042. [24] B.C. Giessen, A two-parameter representation of rapidly glass forming binary alloy systems, Proc. 4th Int. Conf. Rapidly Quenched Metals. 1981, pp. 213–216. [25] E.G. Noya, C. Rey, L.L. Gallego, Amorphization of Ni-Al alloys by fast quenching from the liquid state: a molecular dynamics, J. Non-Cryst. Solids 298 (2002) 60–66. [26] A. Peker, W.L. Johnson, A highly processable metallic glass: Zr41·2Ti13.8Cu12·5Ni10.0Be22.5, Appl. Phys. Lett. 63 (1993) 2342–2345.