3.1. The Impact of the Strain Rate on Diffusion Thickness
Figure 1 displays the SEM images capturing the interfaces of diverse Mg/Al composite sheets. This refers to the present hot-pressing technique with a thickness reduction of 40% at 350 °C, which can generate a good connection between 1060 and AZ31. According to reports [
28,
29], the bonding of the composite sheets relies on two crucial factors: the deformation temperature and the amount of strain. Diffusion is slower at low temperatures, thus affecting the quality of the metallurgical bond. When the strain amount falls below a critical value, the oxide layer present on the metal surface proves challenging to break, thus restricting the exposure of fresh metal. As a result, the composite plates are unable to generate strong metallurgical bonding. Liu et al. [
10] discovered that the roll bonding threshold strain for the AZ31/1060 composite sheets is 20%. Thus, a thickness reduction of 40% at 350 °C can generate an effective metallurgical bond. Additionally, it has been observed that the strain rate exerts a significant impact on the characteristics of the interface. When subjected to hot pressing with strain rates varying from 1.0 × 10
−3 s
−1 to 1.0 × 10
−2 s
−1, a flawless interface layer forms, characterized by the absence of cracks and pores. After the strain rate is further reduced to 3.3 × 10
−4 s
−1, the interface layer undergoes fragmentation. Moreover, numerous white particles can be found on the entire cross-section. EDS analysis indicates that these particles may be rich in O, C, Si, Mn, etc. It indicates that some particles may be the polishing residues from SiC sandpapers.
In order to delve deeper into the interface characteristics,
Figure 2 presents magnified SEM images along with area-scanning images. It is evident that a distinct interface layer is visibly formed between AZ31 and 1060 during the hot-pressing process. The element Mg and element Al exhibit full inter-diffusion across the entire interface. This further proves that direct hot pressing in an atmospheric environment can generate an effective bond at the interface of AZ31 and 1060. As the strain rate decreases, the thickness of the inter-diffusion layer gradually increases.
Figure 2d shows that the interface layer of the P4 sample is obviously broken. This is primarily the underlying cause for the presence of numerous cracks and pores on the interface of the P4 sample. As the strain rate decreases, a noticeable trend emerges: the diffusion layer progressively thickens. The element line-scanning spectrums of different samples are shown in
Figure 3. Due to the occurrence of inter-diffusion, the concentrations of Mg and Al show a gradual change at the interface. Within this diffusion layer, the concentration of Al decreased from the Al alloy side to the Mg alloy side, while the concentration distribution of Mg exhibited the opposite trend. SEM studies revealed no evidence of intermetallics in the bonding structure. Moreover, no significant single-phase layer is found at the interface for the P1, P2, and P3 samples. For the P4 sample, significant fluctuations in the concentrations of the elements at the interface occur, which are related to the fragmentation of the compounds. Through the analysis of these spectrums, the thickness of the diffusion layer is determined and presented in
Figure 4. Specifically, when the hot-pressing strain rate decreased from 1.0 × 10
−2 s
−1 to 3.3 × 10
−4 s
−1, there is a steady progression in the thickness of the diffusion layer from 5.7 μm to 38.7 μm.
In the previous work, a linear relationship between the thickness of the compound and the annealing time was reported [
30]. Thus, the diffusion mechanism governs the growth of the compounds. In general, the thickness of the intermetallic compound layer increases linearly with the square root of the holding time. This can be expressed by the following empirical law:
where k
p is the growth constant (m
2/s). Wang et al. [
31] reported that the growth constant of the diffusion layer at 350 °C is 1.679 × 10
−13 m
2/s for a rolled Mg/Al sheet. Based on this, the predicted thickness with holding time is also shown in
Figure 4. For the present work, the hot-pressing temperature and strain are constant. Thus, the hot-pressing rate determines the hot-pressing time, as shown in
Table 1. It is considered that the hot-pressing time will be the important factor influencing the thickness of the diffusion layer.
Figure 4 indicates that with a decreasing hot-pressing rate, the thickness of the diffusion layer gradually deviates from the predicted value. When the strain rate is 3.3 × 10
−4 s
−1 (corresponding to a hot-pressing time of 30 min), the thickness of the diffusion layer (38.7 μm) is much higher than the predicted value (17.4 μm). Unlike the annealing process, hot pressing imposes high stresses and generates a large strain (40%) during the holding process. Plastic strain might reduce the thickness of the diffusion layer [
2]. This indicates that high stress will largely promote inter-diffusion at the interface [
32]. Firstly, high bonding pressure increases the contact sites, promoting the diffusion rate and interface reaction [
33]. Secondly, crystal defects (e.g., dislocations) via plastic deformation will reduce the activation of the energy of diffusion and enhance the volume diffusion [
10,
34,
35]. Thus, both initial dislocations and plastic strain might accelerate inter-diffusion during hot pressing.
Moreover, previous studies have shown that both high temperatures and prolonged holding times can lead to the formation of Kirkendall holes on the Al side [
32]. It is reported that when the annealing temperatures exceed 250 °C, conspicuous Kirkendal voids can be seen at the interface [
10]. The emergence of these voids is expected to diminish the overall properties and bonding strength. The primary factor underlying this phenomenon is the discrepancy in diffusion rates between Mg and Al at the interface. Generally, Al exhibits a swifter inter-diffusion rate compared to Mg [
36]. Consequently, the surplus vacancies on the Al side eventually facilitate the generation of voids. In this study, no voids appeared at the interface. This may be related to the pressure applied by the hot-pressing process. Hot-pressing deformation can promote inter-diffusion and cause plastic flow of the material, which, in turn, eliminates the voids.
3.2. Phase Compositions at the Interface
Figure 5a displays the positions of EDS points for the P1 samples. The chemical compositions highlighted in
Figure 5a can be found in
Table 2. Referring to the Al-Mg binary phase diagram (see
Figure 5d) at a temperature of 350 °C, the maximum solution of Al in Mg was about ~5 at.%, and the maximum solid solution of Mg in Al was ~10 at.%. Moreover, Mg
2Al
3 (β-phase) and Mg
17Al
12 (γ-phase) can be formed at the interface. According to the stoichiometric proportion, it is inferred that position 1 and position 2 are the α-Mg and position 9 is the α-Al. Moreover, abundant LMCs are present between them. Adjacent to the Mg side (position 3–5), the microstructure might be the mixture of α-Mg and Mg
17Al
12 LMCs. However, next to the side of the Al (position 8), the mixture of α-Al and Mg
2Al
3 LMCs might be the dominant microstructure. The stoichiometric proportion at position 7 is closer to Mg
2Al
3 LMCs. It also infers that position 6 might contain a mixture of Mg
17Al
12 and Mg
2Al
3. Thus, for the P1 sample, the interface is mostly a multi-phase mixture, and there is no obvious single-phase LMC layer.
Table 3 provides the chemical compositions obtained from various positions at the interface of the P3 sample, as indicated in
Figure 5b. Next to the side of the Mg, position 4 has a close chemical composition with Mg
17Al
12 LMCs. However, the composition remains at a very limited width (<1.5 μm). This indicates that the layer of Mg
17Al
12 is very narrow. In contrast, the interface of the P3 sample contains a relatively thicker single-phase Mg
2Al
3 layer (~2.4 μm, see positions 6–8). Thus, with decreasing hot-pressing strain, the single-phase LMC layer will be generated at the interface, and the Mg
2Al
3 layer exhibits faster growth than Mg
17Al
12.
Table 4 provides the chemical compositions obtained from various positions at the interface of the P4 sample (see
Figure 5c). The Mg
17Al
12 layer is still very narrow. However, the width of the Mg
2Al
3 layer reaches 35 μm (see positions 3–5). Moreover, the EDS results indicate that the broken phase at the interface is mainly Mg
2Al
3 LMCs.
Clearly, during the hot-pressing process, extensive inter-diffusion takes place at the interface between Mg and Al. First, Mg and Al supersaturated solid solutions are formed on both sides of the interface. Then, two types of intermetallic compounds, Mg
17Al
12 and Mg
2Al
3, precipitate within the supersaturated solid solutions of Mg and Al. In the subsequent reaction process, a large number of intermetallic compounds continuously precipitate, forming a single-phase layer. Many studies have proven that the growth constant of Mg
2Al
3 is higher than Mg
17Al
12 [
38,
39,
40]. Thus, both annealing and thermal processing will generate a thick Mg
2Al
3 layer and a thin Mg
2Al
3 layer. In this work, long-time hot pressing also exhibits a similar phenomenon.
Figure 4 indicates that a thick Mg
2Al
3 layer fragmentation occurs during hot pressing. According to the ref. [
41], the intermetallic compound layer of the Mg/Al composite sheet is very brittle. The brittle nature of the Mg
2Al
3 phase results in a difficultly of accommodating deformation, and fragmentation occurs during subsequent deformation.
3.3. Shear Strength and Fracture Analysis
For the Mg/Al composite plates as a structural material, preserving structural integrity is of utmost importance. The occurrence of interfacial delamination during operation can result in complete failure. It is reported that a high bonding strength can enhance ductility and tensile strength [
10]. Furthermore, when the interface between Mg and Al becomes separated, the Mg matrix loses the protection provided by the Al alloy cladding. Hence, interfacial bond strength serves as a significant gauge for evaluating the quality of composite plates. The interfacial bond strength of composite plates is typically assessed using two methods: the shear test and roller peel test. Here, the shear test using the ASTM D4896 standard [
27] was used to evaluate the shear strength.
Table 5 displays the shear strengths of the hot-pressed plates. As the strain rate decreases, the shear strength shows an initial increase followed by a decrease. This suggests that the P3 sample demonstrates the maximum shear strength (4.39 MPa).
This indicates that the hot-pressing rate can largely influence the shear strength. Hot pressing predominantly creates two types of bonding: mechanical bonding and diffusion bonding [
10]. Diffusion bonding is a time-intensive process affected by temperature and stress, while mechanical bonding generates rapidly. Therefore, hot pressing leads to the formation of robust mechanical bonds with uniform flat interfaces and a consistent fixed plastic strain across all samples, and thus the impact of mechanical bonding is believed to be uniform. The restricted hot-pressing time limits extensive inter-diffusion, and a slow hot-pressing speed may encourage heightened interfacial inter-diffusion. As illustrated in
Figure 5, a slow hot-pressing rate results in increased compound thickness. Consequently, differences in shear strength are mainly ascribed to discrepancies in metallurgical bond quality. As shown in
Figure 2, significant elemental inter-diffusion occurred at the interface between the substrates and the diffusion layer, and no voids or cracks were generated. This demonstrates that the current hot-pressing process produces effective metallurgical bonding independent of the strain rate. The influence of the strain rate on shear strength might be mainly attributed to the changes in structure, composition, and thickness of the diffusion layer.
As the hot-pressing rate decreases, the thickness of the diffusion layer increases, and the composition of the interface gradually changes from a mixed zone of Mg
17Al
12 and Mg
2Al
3 to two single-phase zones. It is reported that a slight thickening of the diffusion layer improves the bonding strength, and then the bonding strength deteriorates when the thickness of the diffusion layer is greater than 10 μm [
10]. As shown in
Figure 4, the thickness of the diffusion layer is 5.7 μm, 6 μm, and 8.5 μm for the P1, P2, and P3 samples, respectively. This might be the reason why the P3 sample has higher shear strength. For the P4 sample, the thickness of the diffusion layer is 38.7 μm, and a thick Mg
2Al
3 single-phase layer can be found. It is reported that the Vickers hardness of the Mg
2Al
3 phase (~275 HV) is always higher than the Mg
17Al
12 phase (~250 HV) [
42]. Thus, the Mg
2Al
3 phase exhibits higher hardness and brittleness. The debonding during tension and shear usually occurs at the interface between the Mg
2Al
3 single-phase layer and the Al alloy side [
11]. The high brittleness also causes the Mg
2Al
3 single-phase layer to be more susceptible to breakage during processing.
The SEM analysis depicted in
Figure 6 provided insights into the shear fracture surfaces of different samples at the Al side. Additionally, the composition of phases on these surfaces was determined through EDS analysis, as outlined in
Table 6. The identification of possible phases on the shear fracture surfaces was obtained based on the results of the EDS analysis. The shear fracture interface features of the P1 sample are shown in
Figure 6a, and there are distinct granular protrusions at the fracture surface. This infers that the micro-cracks might extend along the grain boundaries during shearing. According to EDS results, the Al side fracture contains a mixed region of the α-Al phase and Mg
2Al
3 phase. As shown in
Figure 3 and
Table 2, no single-phase layer is formed in the P1 sample. Thus, it infers that the shear fracture took place within the mixed region of Mg
2Al
3 and Mg
17Al
12, which exists between the Mg and Al components.
For the P3 sample, the predominant cleavage river and lamellar structures were prevalent on the fracture surface. The EDS results show the phases on the shear fracture surfaces may be the Mg2Al3 phase according to the atomic ratio. This morphology may occur due to the formation of a layer-like single-phase region of intermetallic compounds during the hot-pressing process. Stress is deflected along the Al/Mg2Al3 interface. Since the matrix and the diffusion layer have a strong metallurgical bond, the interface will not fail. Therefore, the stress is further transmitted to the diffusion layer, causing cleavage fracture. Although cleavage fracture is a brittle fracture, the crack has to extend through the interior of the compound grains. Therefore, this type of fracture requires a higher stress than an intergranular fracture. This may explain the high shear strength of the P3 sample.
The shear fracture interface morphology of the P4 sample is shown in
Figure 6c, where the fracture surfaces of the Al sides exhibit a mixed feature of granular and fluvial. Moreover, some large bumps and pits can also be found on the fracture surface. EDS results indicate a large amount of Mg
2Al
3 in the phase on both the Mg and Al sides. The results indicate that the fracture occurred in the single-phase Mg
2Al
3 region. As shown in
Figure 5, during the hot-pressing process of sample P4, under high temperatures, a thick Mg
2Al
3 layer occurs a certain degree of brittle fracture, and the broken Mg
2Al
3 layer could lead to a decrease in the bonding strength.
To validate the accuracy of the previously mentioned shear fracture position prediction, XRD analysis was conducted on the shear fracture of the P4 sample, as shown in
Figure 7. The Al matrix side shear fracture surface (
Figure 7a) exhibited the presence of the Al phase and Mg
2Al
3 phase. Similarly, the Mg matrix side shear fracture surface (
Figure 7b) revealed the detection of the Mg phase, Mg
2Al
3 phase, and Mg
17Al
12 phase. These XRD results are in line with the findings from the EDS analysis, further confirming that the shear fracture position typically occurs at the Mg
2Al
3/Al interface and the Mg
2Al
3 layer. The fractured behavior at this position is greatly influenced by the interface structure and the properties of the diffusion layer formed at the Al/Mg interface during hot pressing. It is important to note that intermetallic compounds, such as Mg
2Al
3 and Mg
17Al
12, possess significantly higher Vickers hardness and lower plasticity compared to matrix materials. Additionally, studies have reported that the Vickers hardness of the Mg
2Al
3 phase is higher than Mg
17Al
12 [
43]. Consequently, a brittle fracture is more likely to occur in the Mg
2Al
3 phase layer during shear. The thickness of intermetallic compounds also plays a crucial role in the fracture behavior and shear strength [
44], where a thickness exceeding 10 μm results in reduced shear strength. The smooth cleavage plane observed on the shear fracture surface corresponds to the higher shear strength achieved, which explains the highest shear strength exhibited by the P3 sample.
It has been shown that rapid hot pressing for a short duration (1 min) and a high strain (40%) under normal conditions can effectively bond AZ31 and 1060. This is contrary to previous methods that required a longer hot-pressing time (>1 h) to achieve sufficient diffusion at the interface. Rapid hot pressing promotes faster diffusion compared to annealing and vacuum diffusion welding [
45]. However, a prolonged hot-pressing time (>30 min) may result in the thickening and fragmentation of compounds at the interface, which adversely affects bond strength. Thus, rapid hot pressing prevents compound thickening and enhances processing efficiency. The lack of thickness of the diffusion layer due to rapid hot pressing can also be compensated by subsequent heat treatment. Therefore, rapid hot pressing has the advantage of high efficiency and low cost compared to slow hot pressing.
In addition to the strain rate, the hot-pressing temperature, time, and strain amount also affect the final bond quality and the structure of the interface. Generally, both a high temperature and a long period of time can lead to the thickening of the diffusion layer [
10]. The strain amount involves the creation of new surfaces and the evolution of matrix defects, which might also influence the diffusion process [
10]. Moreover, the hot-pressing parameters also have a significant influence on the microstructure and texture of the substrates [
46]. The interfacial bonding quality and the microstructure of the matrix will jointly affect the overall mechanical properties of the composite sheets [
11]. Therefore, the next research phase will focus on developing and optimizing the rapid hot-pressing process. The combined impact of interfacial structure and matrix microstructure on the overall sheet’s mechanical properties will be investigated.