4.2. T6 Heat-Treated Condition
a,b show that Ni has a clear detrimental effect on both the YS and the UTS of T6 heat-treated alloys. The reason for this behavior can be explained by the composite material approach and the mechanism of load transfer from the primary Al matrix to the reinforcement phases, suggested by Moffat [30
] and reproposed by Casari at al. [25
]. This approach takes into account the empirical relation between hardness (H
) and ultimate tensile strength (UTS
), reported in Equation (1):
In addition, it assumes that rigid phases, i.e., eutectic silicon particles and Ni-bearing phases, have a linear elastic behavior, thus the deformation of fracture εfr
can be expressed by Equation (2):
is Young’s modulus, in GPa. Since an absolute relationship is not required, relative fracture strain f(εfr
) can be obtained by Equation (3), combining Equations (1) and (2):
Assuming homogeneity of the strain distribution for each phase, the values of H
, found in literature and reported by Moffat [30
] and Song et al. [31
], can be used in Equation (3) to determine the values of f(εfr
), reported in Table 4
. It is worth noting that Ni-rich compounds show a lower value of relative fracture strain compared to eutectic Si particles and, thus, have a higher tendency to fracture.
These findings suggest the following fracture mechanism of the material: when an external stress is applied, the load is mainly born by the reinforcing network of eutectic Si and Ni aluminides. These phases show the highest tendency to fracture [25
], leading to the formation of micro-cracks in the material. The specimen, thus, reaches YS for lower stress values. Furthermore, since micro-cracks can easily propagate across brittle phases, the fracture rapidly propagates when there is a high interconnectivity of the reinforcing 3D network. This leads to the formation of extended de-cohesion regions that reduce the dimension of the resistant section. For this reason, UTS is reached at lower stress values.
In comparison with the mechanical properties of the heat-treated A356 base alloy (Figure 1
a,b), the addition of 0.5 wt % Ni has a detrimental influence on the YS of the alloy, since it yields a decrease of about 40 MPa owing to the early fracture of brittle Ni-based compounds. Likewise, a further decrease of 100 MPa at Ni contents up to 2 wt % can be clearly related to the increased area fraction of such intermetallics (Figure 8
As far as UTS is concerned, both the base A356 alloy and the alloy with 0.5 wt % Ni exhibit a similar value. Apparently, this Ni content does not enable to maintain the 3D network after T6 heat treatment. Such loss of interconnectivity limits the ease of crack propagation by involving the more ductile Al phase in the fracture process, reasonably explaining the higher UTS average value than those of the alloys with 1 and 2 wt % Ni. In contrast, higher Ni concentrations limit the fragmentation and spheroidization of eutectic Si particles during T6 heat treatment [9
]. As a result, the 3D network maintains a higher level of interconnectivity, fostering crack propagation through a much easier path and, thus, leading to lower UTS values.
These hypotheses seem also in line with fractographic observations of alloys with Ni contents of 0.5, 1 and 2 wt % (Figure 5
b,d,f, respectively). In Figure 5
b, the heat-treated alloy with low Ni content presents not only a certain number of cleavage regions and fractured brittle phases (mainly eutectic Si particles) but also plastic deformation of the Al matrix. On the other hand, in Figure 5
d,f, a lower amount of plastic deformation can be detected and it is possible to distinguish an increased number of either cleavage propagation or debonding sites.
It is worth noting that the morphology of brittle phases influences the mechanical properties of the investigated alloys. Statistical evaluations of geometrical parameters highlighted that the presence of Ni leads to a reduction of the spheroidizing effect of heat treatment on eutectic Si particles. ED and C increase with Ni contents of 1 and 2 wt % compared to the same parameters of the alloy with 0.5 wt % of Ni, as can be observed in Figure 6
a,b. With 1 and 2 wt % of Ni, eutectic Si particles maintain an irregular morphology after heat treatment (Figure 6
a), which eases the localized concentration of stresses. This then leads to crack formation at lower load values. Additionally, a higher average dimension of eutectic Si particles (Figure 6
b) indicates that the network fragmentation during heat treatment is limited when Ni contents are 1 and 2 wt %. Eutectic Si particles of considerable dimension promote fracture propagation once it is nucleated, explaining the decrease in YS and UTS average values depicted in Figure 1
shows the detrimental effect of Ni content on Vickers microhardness of the α-Al matrix after T6 heat treatment. It is well-known that the T6 heat treatment leads to the precipitation of the M-Mg2
Si phase in nanometric particles, which act as a reinforcement and increase the microhardness of the α-Al matrix. Nevertheless, microhardness is found to decrease with further Ni additions up to 2 wt % (Figure 2
). This result, in agreement with the outcome of tensile tests, may indicate that the presence of Ni exerts a negative effect on the precipitation of reinforcing intermetallic compounds to some extent.
In summary, Ni aluminides cause a remarkable decrease in the room temperature YS, UTS, and Vickers microhardness of the A356 heat-treated alloy. Despite additions of Ni up to 2 wt % hindering spheroidization effects of eutectic Si particles during T6 heat treatment, they also promote the formation of a higher number of brittle phases that easily promote fracture propagation.