## 3 Results and Discussion
### Effects of Heat Treatment on the Interfacial Microstructures  
Figure 2 shows the interfacial microstructure of the Al/Fe alloys hybrid structure heat treated for 1 h at different solution temperatures, followed by furnace cooling. At a solution temperature of 500 \({}^{\circ}\)C (Figure 2a), the interface of the Al/Fe alloys hybrid structure demonstrated a similar microstructure to the as-cast condition without chemical bonding, and no intermetallic phase was formed in the Al matrix. When the solution temperature increased to 530 \({}^{\circ}\)C, the diffusion of atoms at the interface resulted in an intermetallic layer (Figure 2b). The nickel diffused into the Al matrix, and this led to formation of the intermetallic phase, due to the low solubility of nickel in aluminum [23]. There were no intermetallic phases in the Al matrix; however, as in the as-cast and 500 \({}^{\circ}\)C solution temperature conditions. Therefore, the temperature at which the inter-diffusion of atoms begins to occur at the interface is approximately 530 \({}^{\circ}\)C. When the solution temperature increased to 550 \({}^{\circ}\)C, the intermetallic phases at the interface formed a continuous layer and the thickness of the intermetallic layer increased significantly (Figure 2d). The chemical  
Figure 1: (a) SEM micrographs of the Al/Fe alloys hybrid structure’s interface with nickel coating; (b) XRD patterns of nickel coated steel before and after heat treatment at 300 \({}^{\circ}\)C; (c) chemical element quantitative maps of (a): Fe, Ni, Al, and Si. bonding that formed at the interface under this particular heat treatment condition was superior. The parabolic law correlates the solution temperature with the growth of the intermetallic layer [24; 25; 26], and is expressed as follows:  
\[X=Kt^{1/2} \tag{1}\]  
\[K=K_{0}\text{exp}(-\frac{Q}{RT}) \tag{2}\]  
where \(X\) is the intermetallic layer thickness, \(t\) is the heat treatment time, \(K\) is the coefficient of diffusion, \(K_{0}\) is the diffusion constant, \(Q\) is the activation energy for intermetallic phase growth, \(T\) is the absolute solution temperature, and \(R\) is the gas constant. According to Equations (1) and (2), the solution temperature and diffusion coefficient are proportional, implying that an increase in solution temperature promotes the diffusion of atoms at the interface. Figure 2c,e show the SEM micrographs of the intermetallic layers created at the interface, at different solution temperatures. Table 3 shows the results of the compositional analysis of the intermetallic layers formed at 530 and 550 \({}^{\circ}\)C. The compositions of the intermetallic layer after heat treatment were Al\({}_{3}\)CuNi / Al\({}_{3}\)Ni at 530 \({}^{\circ}\)C, and NiAl / Al\({}_{3}\)Ni at 550 \({}^{\circ}\)C. Thus, composition of the intermetallic layer varied with solution treatment temperature. The types of intermetallic layers formed through heat treatment were related to the ratio of diffused atoms. The higher solution temperature caused greater diffusion at the interface, affecting the types of intermetallic phase, which were determined by the Ni:Cu ratio in the Cu-containing Al matrix. As the Ni:Cu ratio increased, a nickel-bearing phase formed in the sequence of \(\gamma\)-Al\({}_{7}\)Cu\({}_{4}\)Ni, \(\delta\)-Al\({}_{3}\)CuNi, and \(\varepsilon\)-Al\({}_{3}\)Ni [27]. Therefore, depending on the solution temperature, the intermetallic layer at the interface primarily comprised \(\delta\)-Al\({}_{3}\)CuNi (at 530 \({}^{\circ}\)C), or \(\varepsilon\)-Al\({}_{3}\)Ni (at 550 \({}^{\circ}\)C). Figure 3a shows the changes in thickness of the intermetallic layer, at different solution treatment temperatures and annealing times. Heat treatment at 500 \({}^{\circ}\)C for 1 h and 3 h did not generate an intermetallic layer. The thickness of the layer formed after 8 h was approximately 10 \(\upmu\)m, which was thinner than that for other solution treatment temperatures.