## 3 Results and Discussion
### Mechanical Properties  
The stress-strain curves of B\({}_{4}\)C-Al composites prepared by high-energy ball milling and low-energy ball milling processes are shown in Figure 7. Figures 8 and 9 illustrate the yield and tensile strength of the composites, respectively. The yield strength values of the 15 wt % and 20 wt % B\({}_{4}\)C-Al composites prepared by low-energy ball milling process are 130 MPa and 141 MPa, respectively, which are \(\sim\)20% and \(\sim\)31% higher than those of the pure Al samples prepared at the same condition. Values of the yield strength of the 15 wt % and 20 wt % B\({}_{4}\)C-Al composites prepared by high-energy ball milling process are 184 MPa and 213 MPa, respectively, which are \(\sim\)70% and \(\sim\)97% higher than those of the pure Al samples. Moreover, when increasing the ratio of the B\({}_{4}\)C particles, the yield strength of the composite prepared by the low-energy ball milling increases slightly, whereas that of the composites prepared by high-energy ball milling increases dramatically. As shown in Figure 9, the tensile strength of the composites prepared by the low-energy ball milling process (191 MPa and 194 MPa for 15 wt % and 20 wt % B\({}_{4}\)C-Al composites, respectively) was almost the same as those of the pure Al samples (191 MPa) prepared using the same process. On the contrary, after high-energy ball milling, the tensile strength of 15 wt % and 20 wt % B\({}_{4}\)C-Al composites  
Figure 6: Neutron radiography of B\({}_{4}\)C-Al composites: (**a**) 15 wt % B\({}_{4}\)C-Al prepared by low-energy ball milling; (**b**) 20 wt % B\({}_{4}\)C-Al prepared by low-energy ball milling; (**c**) 15 wt % B\({}_{4}\)C-Al prepared by high-energy ball milling; (**d**) 20 wt % B\({}_{4}\)C-Al prepared by high-energy ball milling. Figure 10 shows the SEM and the concomitant BSE images (BSTL) of the typical fracture surface of the specimens. For the sintered pure Al sample (Figure 10a), multiple shallow dimples with an average diameter of 0.8 \(\upmu\)m are observed, which are formed by the joining of micro voids along shear slip bands. On the contrary, although the stress-strain curves of the B4C-Al composites indicates apparent brittle fracture in a macroscopic scale without apparent necking (see Figure 7), the fracture surface morphology observed from SEM shows a ductile nature in microscopic scale with fine dimples of -0.5 \(\upmu\)m (Figure 10b-e). Addition of the B4C reinforced particles reduces the macroscopic ductility, which can be verified from the formation of smaller dimples compared to those without addition of B4C. The particle-matrix de-cohesion resulting dimples are observed, indicating that the tensile fracture of the composites is controlled by void nucleation and growth [1; 25]. Figure 10: _Cont._  
Figure 9: The tensile strength of pure Al plate and B4C-Al composites prepared by high-energy ball milling and low-energy ball milling processes. Figure 10: Fracture surfaces of B\({}_{4}\)C-Al composite: (**a**) Al without B\({}_{4}\)C; (**b**) 15 wt % B\({}_{4}\)C-Al prepared by low-energy ball milling; (**c**) 20 wt % B\({}_{4}\)C-Al prepared by low-energy ball-milling; (**d**) 15 wt % B\({}_{4}\)C-Al prepared by high-energy ball milling; (**e**) 20 wt % B\({}_{4}\)C-Al prepared by high-energy ball milling. For the composites reinforced with 15 and 20 wt % B\({}_{4}\)C particles prepared by the low-energy ball milling process, and BSTL images reveal an inhomogeneous, stripe-like distribution of the B\({}_{4}\)C particles (with a length of 15 \(\upmu\)m\(\sim\)30 \(\upmu\)m and a width of 4 \(\upmu\)m\(\sim\)10 \(\upmu\)m) in the Al matrix (Figure 10b,c). The agglomeration of the B\({}_{4}\)C particles inside the Al matrix failed to transfer the shear and tensile stresses during tensile tests, blocked the movement of the slip bands, and thus lead to a de-cohesion failure of particle-matrix interfaces and fast initiation and propagation of cracks [3; 25].