## 3 Results and Discussion
### Mechanical Behavior
#### 3.2.2 Elevated Temperature Tensile Behavior  
Figure 4, which presents two representative stress-strain curves for each alloy at 25 \({}^{\circ}\)C, 150 \({}^{\circ}\)C, and 250 \({}^{\circ}\)C, and Table 5 summarize the tensile response of the hypoeutectic Al-Si-(Cu)-Mg alloys. Both materials show the expected thermal softening: the ultimate tensile strength (UTS), yield strength (YS), and elongation to fracture (EL) decrease monotonically as the temperature rises, owing to the increased matrix ductility and the coarsening of intermetallic phases that collectively lower dislocation resistance. As shown in Table 5, at room temperature, Alloy A attains 347.01 \(\pm\) 7.56 MPa (UTS), 227.74 \(\pm\) 3.12 MPa (YS), and 5.03 \(\pm\) 0.16% (EL), markedly exceeding Alloy B (277.85 \(\pm\) 4.26 MPa, 123.98 \(\pm\) 2.81 MPa, 2.12 \(\pm\) 0.11%). This difference is primarily attributed to the uniform solid solution of Cu in Alloy A, which induces lattice distortion and enhances strength via solid solution strengthening. Furthermore, the stable distribution of \(\alpha\)-Al(FeMn)Si phases in Alloy A contributes to load transfer and crack resistance. In contrast, the presence of brittle \(\beta\)-AlFeSi and agglomerated Mg\({}_{2}\)Si in Alloy B promotes stress concentration, thereby reducing both strength and ductility.  
The disparity persists at elevated temperatures. At 150 \({}^{\circ}\)C and 250 \({}^{\circ}\)C, Alloy A maintains relatively high tensile properties--343.27 \(\pm\) 5.27 MPa (UTS), 227.62 \(\pm\) 2.89 MPa (YS), and 4.26 \(\pm\) 0.21% (EL) at 150 \({}^{\circ}\)C and 306.68 \(\pm\) 4.45 MPa, 176.87 \(\pm\) 3.41 MPa, and 6.55 \(\pm\) 0.20% at 250 \({}^{\circ}\)C, demonstrating the robust retention of the mechanical integrity at elevated temperatures. This resilience is ascribed to the retention of Cu in the solid solution together with the thermal stability of the \(\alpha\)-Al(FeMn)Si skeleton, which collectively suppress strain localization. Conversely, Alloy B shows a pronounced decline  
Figure 4: Stress–strain curves of Alloy A and Alloy B obtained at RT, 150 \({}^{\circ}\)C, and 250 \({}^{\circ}\)C; two representative curves are shown for each condition.  
(150 \({}^{\circ}\)C: 249.45 \(\pm\) 7.81 MPa, 112.27 \(\pm\) 3.21 MPa, 3.12 \(\pm\) 0.12%; 250 \({}^{\circ}\)C: 229.92 \(\pm\) 4.98 MPa, 106.30 \(\pm\) 2.71 MPa, 3.43 \(\pm\) 0.18%), a consequence of the Mg\({}_{2}\)Si coarsening and the persistence of brittle \(\beta\)-AlFeSi plates, both of which increase the structural heterogeneity and local brittleness.  
The analysis of the engineering stress-strain curves (Figure 4) revealed that Alloy A maintained a noticeable strain-hardening region after yielding, indicative of uniform plastic deformation. In contrast, Alloy B showed an immediate stress saturation or sharp post-yield softening at elevated temperatures, suggesting the early onset of necking due to localized softening and the presence of brittle intermetallics.  
Both alloys exhibited an extended toe region in the initial stage of the stress-strain curve, with a more pronounced effect observed in Alloy B. This phenomenon implies a delayed load transfer and plastic deformation initiation, influenced by the alloy's initial ductility and phase distribution. While a higher ductility generally correlates with a longer toe region [15; 22], the presence of brittle phases, such as Mg\({}_{2}\)Si and \(\beta\)-AlFeSi, may induce heterogeneous deformation and localized stress shielding, thereby contributing to delayed yielding [15; 22].