Evolution of Voids in Mg/Al Diffusion Bonding Process

Material and experimental procedure

The materials used in this work include pure magnesium (Mg1) and industrial aluminum (Al1060). Cylindrical-shaped specimens (Φ40 mm×10 mm) were used to prepare Mg/Al joints. The oxide film at the surface was removed by 2000 grit sandpaper, and then the materials were polished with Al₂O₃ polishing powder before diffusion bonding.

The vacuum hot pressing furnace with FVPHP-R-10FRET-40 type was used for Mg/Al diffusion bonding. The processing parameters are: the bonding temperature T=475–500 °C with a heating rate of 10 °C/min, the pressure P=1 MPa, the holding time t=80–100 min under Ar gas atmosphere. The cooling process was conducted in the vacuum chamber from bonding temperature to 100 °C.

Specimens with dimension of 10×10×10 mm, which are parallel to the load applied direction, were fabricated from the Mg/Al diffusion bonded joint region by wire cut machine. The specimens were grinded with 2000 grit SiC paper and polished by Al₂O₃ powder, and then ultrasonically cleaned in acetone.

The structure and morphology of the bonded joint were observed by the Nova Nano 230 scanning electron microscope (SEM). Surface properties were determined by the Bruker Veeco NanoManTM VS+MultiMode atomic force microscope (AFM). Distribution of Mg and Al at the bonded joint was determined by the JXA-8230 electron probe microanalysis (EPMA). The phase constitutions were identified using a RIGAKU RAPIDⅡR type selected area X-ray diffractometer (XRD). The thickness of each intermetallic compounds (IMCs) was measured by EPMA, the average thickness was calculated by averaging ten values. A VNHT type nanoindentation testing (NIT) with a load of 2 mN was used to determine the elastic modulus of IMCs. The high temperature mechanical properties of the matrixes were tested by an Instron3369 type mechanical testing machine at 300 and 475 °C, with loading rate of 1.0 mm/min and the heating rate of 2 °C/min.

Model setup

In order to truly reflect the distribution of residual stresses at joints, the present study focused on both the temperature-dependent material properties and the effect of IMCs. Simulations in this paper were performed in the ANSYS software. Al1060 and Mg1 with cylindrical shaped (Φ40 mm×10 mm) were considered for simulations and the temperature load was applied to all units with uniform radiant heating. The bonding temperature of diffusion bonding was about 0.7 times of the matrix with lower melting point (475 °C), so material properties varied with temperature should be considered in the model. Parts of the mechanical properties are given in Table 1.

To improve both the computational efficiency and accuracy, 2-D axisymmetric mode and Plane183 were applied for the residual stresses analysis. At the Mg-Al reaction layer, the regions near the matrixes and the free boundary, the mesh was refined. The geometry model of Mg/Al diffusion bonded joint and mesh finite element model are shown in Figure 1.

Figure 1:

The scheme of Mg/Al joints: (a) geometrical model; (b) meshing scheme.

Voids behaviors under plastic deformation during early stage of bonding

The surface microstructures of Al1060 and Mg1 before bonding examined by AFM are presented in Figure 2. In the initial stage of Mg/Al diffusion bonding, localized contact occurs at the connecting surface, the initial voids are generated. It can be deduced from Figure 2 that bigger voids would be obtained in the Al1060. Therefore, properties of voids are characterized by surface properties of Al1060, the maximum height (R_max=0.35 μm) and average radius (R=0.71 μm) of surface microcosmic profiles corresponding to the initial height and width of the voids respectively were determined by AFM.

Figure 2:

Surface topography before bonding obtained by AFM: (a) Al; (b) Mg.

Length of the connecting interface is small during the early stage of bonding, suggesting the pressure presented on the connecting interface is larger than the yield strength of matrix. In terms of welding pressure, if interfacial stress is greater than the yield limit of the material, microscopic plastic deformation occurs at the surface, so the initial voids become small, and the localized connecting is formed. During this course, yield strength of matrix is one of the key factors, so mechanical properties of Al1060 and Mg1 at various temperatures are tested, and the results are shown in Figure 3. The results indicate that the strength of Al1060 and Mg1 decreased from 92.3 MPa to 5.6 MPa and 69.9 MPa to 2.2 MPa respectively as the temperature is increased from 20 °C to 450 °C. So the start-up of plastic deformation mechanism becomes possible under the combined effect of temperature and pressure. It can also be observed that the yield strength of Al1060 is higher than that of Mg1 at measuring temperature of 20 °C, 300 °C and 450 °C, the softening deformation of Mg led to almost no strength at higher temperatures, and the results are in accordance with published works [8]. Since microscopic plastic deformation of Al1060 is more difficult to initiate, the yield strength of Al1060 (1.3 MPa at 475 °C) was used.

Figure 3:

Tensile properties of matrixes at various temperatures.

More accurate plastic deformation mechanism can be obtained by elliptical geometry model [9], so elliptical voids from Hill and so forth are selected [10, 11]. The roles of plastic deformation mechanism of void closure are shown in Figure 4.

Figure 4:

Hill’s geometric model: (a) the first stage; (b) the second stage.

The pressure presented on the connecting interface decreases with the increase of the connecting areas, the plastic deformation ceases when the pressure equals to the average stress. Using the theoretical analysis from Johnson to study the degree of connection by plastic deformation, the height (h′) and the length (e′) of the interface achieved can be deduced as [12]:

where a is the initial half width of void (μm), b is the initial half height of the void (μm), L is the half distance of adjacent voids (μm), P is welding pressure (MPa), g is the surface energy (J m⁻²), s_y is the yield strength (MPa).

Bonding pressure must satisfy P > g/L before the plastic deformation mechanism start [13], hence bonding pressure must be larger than 0.689 MPa in Mg1/Al1060 diffusion bonding. It can be found in our experiment that the strength of Mg1/Al1060 joint which obtained by lower bonding pressure is not enough to prepare the samples for mechanical property test.

From eqs (1) and (2), the joint-rate defined as e′/L is calculated as 24 % under the pressure of 1 MPa. In general, interface contacted by plastic deformation for traditional materials is limited (typical joint-rate is less than 10 %) [14], while for Mg1 and Al1060, the degree of connection achieved on the conditions of plastic deformation mechanism in this work is about 24 %, suggesting the plastic deformation plays an important role in the early stage of the Mg/Al bonding process.

Voids behaviors under diffusion in Mg/Al bonding

Due to the different microstructures, physical and chemical properties of Mg and Al, except the voids formed by the initial surface, some secondary voids (voids formed by acquired factors) still exist in Mg/Al diffusion bonding process.

Residual stresses are present in the joints during bonding and subsequent cooling process. Meanwhile, the formation of layered Al₃Mg₂ and Al₁₂Mg₁₇ will change the conditions of stress at the interface. Wu et al. [15, 16] showed that the nucleation sites and growth rate of voids were determined by stress gradient; the diffusion flux of voids was proportional to the stress gradient. Therefore, accurate prediction and efficient estimation of the residual stresses distribution at Mg/Al diffusion bonding joint are essential for revealing the formation of secondary voids. Recently, finite element simulations (FEM) were widely employed to predict welding temperature field, welding residual stress field and welding deformation [17]. In this paper, 2-D ANSYS finite element welding simulation procedure was used.

The formation process of IMCs at the interface during Mg/Al diffusion bonding process is very complex. In order to simplify the calculation, the interface reaction is assumed to be zero or very weak during heating and cooling process, the changes of IMCs can be ignored at this stage. So it can be concluded that the growth of IMCs occurred only at the holding period. Once the interface reaction layer has been completely formed (assuming the thickness values of Al₃Mg₂ and Al₁₂Mg₁₇ are 40 μm and 10 μm, respectively), the matrixes cannot be freely extended, due to the difference in thermal expansion coefficient. Non-uniform contraction velocity occurs across the joint when it is cooled from the bonding temperature to room temperature.

The numerical results based on the model stated above are shown in Figure 5. The maximum compressive stress of σ_x is located in the Al matrix nearby the interface; the maximum tensile stress of σ_x is presented at the Mg matrix in the vicinity of the interface. The maximum stress gradient of σ_x exists at the joint center, the partial amplification image of this region, as shown in Figure 5(b), suggesting that the interface of Al/Al₃Mg₂ has the maximum stress gradient, the stress transformed from compressive to tensile. The stress smoothly transited at the interfaces of Al₃Mg₂/Al₁₂Mg₁₇ and Al₁₂Mg₁₇/Mg. The maximum stress gradient of σ_y exists at the free-boundary of joint, but it is not obvious at the interface. The stress gradient of σ_y in IMCs is mainly in radial distribution, Al matrix and Al₃Mg₂ layers have large tensile stress and radial stress gradients of σ_y.

Figure 5:

Distribution of stress calculated by ANSYS: (a) and (b) σ_x; (c) and (d) σ_y.

Since the diffusion flux of voids is proportional to the stress gradient, the interface of Al/Al₃Mg₂ is mostly beneficial to nucleation and growth of voids. Generally, the elements with higher melting point and smaller atomic radius are easier to diffuse. The melting point of Mg and Al is almost the same, while Mg atoms have smaller atomic radius and thus easier to diffuse [18]. The diffusion of Mg atoms at the interface is accelerated by axial and radial tensile stress existed in Al₃Mg₂ layer, resulting in the aggravation of unbalance diffusion at the Al/Al₃Mg₂ interface. So secondary voids are formed easily in the Al₃Mg₂ layer nearby the Al/Al₃Mg₂ interface, with parallel arrangement to the interface, as shown in Figure 6.

Figure 6:

Voids at the Mg/Al joint interface obtained at 445 °C for 90 min.

Although residual stress at Mg/Al diffusion bonding joint is lower than the yield strength of matrix, the nucleation and growth of the secondary voids can weaken the mechanical properties of joints, joint failure would occur if the voids were connected integrally. To avoid the generation of secondary voids, the control of the formation of IMCs is essential.

Diffusion is the main behavior of the joining of Mg and Al. When the distance of Mg/Al interface lies in the range of interatomic forces, grain boundary migration and voids disappearance can be found under the action of atomic interdiffusion, the way of interface connection is transformed from metal bonding by plastic deformation to metallurgical bonding. Therefore, the diffusion is indispensable for voids closing in the Mg/Al diffusion bonding process, especially for secondary voids.

Morphology and the qualitative distribution of elements at Mg/Al bonded joint obtained at 475 °C for 90 min are shown in Figures 7 and 8, respectively. It can be claimed that a good metallurgical bonding without obvious flaws is formed. It is also clear that the interdiffusion of Mg and Al obviously exists at the interface, and a multi-layer transition zone can be seen at the joint interface. There are two regions where the quantity of elements is kept constant and thus, two different IMCs are formed here.

Figure 7:

SEM micrographs of the Mg/Al joint.

Figure 8:

The elemental distributions at the bonded joint: (a) Mg; (b) Al.

In order to identify the phases of this region, the selected area XRD was applied, as shown in Figure 9. It shows that IMCs are Al₃Mg₂ and Al₁₂Mg₁₇ respectively. At the interface of Al/Al₃Mg₂ and Mg/Al₁₂Mg₁₇, the quantity of elements transforms continuously from one side to the other side, suggesting the appearance of solid solutions.

Figure 9:

The selected area XRD results from the joint interface.

The thicknesses of IMCs from the joint interface bonded at different predominant process parameters are presented in Figure 10. Thickness of IMCs is increasing with welding temperature and holding time, the growth of IMCs layer can be attributed to the interdiffusion of Mg and Al atoms.

Figure 10:

Thickness of IMCs for different time at elevated temperatures.

Different Mg/Al welding process parameters show different diffusion bonding mechanism, assuming the diffusion is not affected by the voids at interface, the growth of the IMCs at the interface can be expressed by [19]

where Δx is the thickness of IMCs layer (μm), x0 is the initial thickness of IMCs layer (μm)，k is the growth constant (μm² s⁻¹), t is the holding time(s)，td is the latent time(s), n is the constant that is indicative of the growth mechanism. For diffusion-controlled growth of a phase with a semi-infinite boundary condition, if bulk diffusion is the controlling mechanism of growth, n ≈ 0.5, whereas if grain boundary diffusion is the controlling mechanism, n is about to 0.33 [20]. As determined by eq. (1), n can be obtained, as reported in Table 2.

The data indicate that diffusion has a significant effect on the bonding. Controlling mechanism of bulk diffusion is observed at lower bonding temperatures, while at higher bonding temperature, the controlling mechanism is grain-boundary diffusion. Bulk diffusion is the dominant diffusion mechanism of all IMCs at both 475 °C and 490 °C. The increase in the growth of IMCs at 500 °C can be attributed to the increasing interdiffusion of Mg and Al. Since diffusion coefficients of grain-boundary diffusion are larger than those of bulk diffusion, the corresponding IMCs with greater thickness can be found. Figure 11 displays the voids models by diffusion.

As shown in Figure 11(a), with the growth of IMCs, voids around them become smaller rapidly, thin voids channel is closed and grain boundaries occur. The driving force for diffusion is the difference in the excess chemical potentials formed by concentration gradient of vacancy. The bonding pressure determines the initial vacancy concentration, the excess chemical potential is originated from plastic deformation, and lower vacancy concentration can be found at the interface. It is known that the vacancy concentration at concave surface is higher than the equilibrium concentration. Due to the variations in curvature of the internal voids, atomic chemistry potentials at the voids surface are lower than those at the grain-boundary interface, vacancy diffusion or reverse atom diffusion occur driven by the different chemical potentials. Therefore, atoms diffuse to the voids along the lattice and grain boundary, accompanied by void shrinkage and the increase of connecting areas, as shown in Figure 11(b). The isolated voids at the joint interface gradually disappear as atoms diffused, and finally a high level of mechanical bonding is formed.

Figure 11:

Diffusion mechanism for void closure: (a) the initial state of voids; (b) shrinkage of voids by diffusion.