## 3 Results and Discussion
### Tribological Properties
#### 3.3.3 Surface Wear Mechanism Evolution with Temperature  
Figure 8 shows SEM images of the wear track surfaces for Alloy A and Alloy B under various temperature conditions (RT, 150 \({}^{\circ}\)C, 250 \({}^{\circ}\)C), illustrating the temperature-dependent transitions in dominant wear mechanisms.  
Figure 7: Three-dimensional optical profilometry of wear track morphology of Alloy A and Alloy B at different temperatures: (**a**) Alloy A at RT, (**b**) Alloy B at RT, (**c**) Alloy A at 150 \({}^{\circ}\)C, (**d**) Alloy B at 150 \({}^{\circ}\)C, (**e**) Alloy A at 250 \({}^{\circ}\)C, and (**f**) Alloy B at 250 \({}^{\circ}\)C.  
At room temperature, Alloy A primarily experienced abrasive wear (Figure 8a), with the wear tracks displaying uniformly distributed shallow grooves and fine scratch marks. This behavior is linked to the increased shear resistance and load distribution facilitated by the lattice distortion from the uniformly dissolved Cu and the dispersion of thermally stable \(\alpha\)-Al(FeMn)Si phases. In contrast, Alloy B (Figure 8b) exhibited a combination of adhesive and delamination wear, characterized by adhered debris, plastic flow features, spalling pits, and layered delamination tracks. These patterns suggest that the stress concentration induced by agglomerated Mg\({}_{2}\)Si and \(\beta\)-AlFeSi phases accelerated the surface failure.  
Figure 8: SEM images of wear track surfaces after sliding wear tests at different temperatures: (**a**) Alloy A at RT, (**b**) Alloy B at RT, (**c**) Alloy A at 150 \({}^{\circ}\)C, (**d**) Alloy B at 150 \({}^{\circ}\)C, (**e**) Alloy A at 250 \({}^{\circ}\)C, and (**f**) Alloy B at 250 \({}^{\circ}\)C.  
At 150 \({}^{\circ}\)C, both alloys exhibited transitions in wear mechanisms due to thermal softening. Alloy A (Figure 8c) continued to display abrasive wear as the dominant mechanism, but crater-like features, delamination zones, and plastic deformation bands began to emerge, indicating localized deformation in the softened matrix. Owing to its more stable microstructure, Alloy A retained a considerable wear resistance. Conversely, Alloy B (Figure 8d) remained dominated by adhesive and delamination wear. While clustered Mg\({}_{2}\)Si particles may have temporarily acted as pseudo-lubricants, their thermal instability at elevated temperatures led to a rapid decline in the tribological performance.  
At 250 \({}^{\circ}\)C, both alloys transitioned predominantly to adhesive wear. In Alloy A (Figure 8e), softened matrix zones formed near the wear surface, and localized delamination cracks and craters became more evident. Abrasive features diminished, while the plastic flow-induced surface deformation increased. Alloy B (Figure 8f) exhibited a complex mixture of spalling, delamination, adhesion, and plastic deformation. Coarsened \(\beta\)-AlFeSi and agglomerated Mg\({}_{2}\)Si phases fragmented under cyclic loading, destabilizing the tribofilm and facilitating crack initiation. As a result, the wear track morphology in Alloy B became highly irregular, with an increased wear depth and width compared to Alloy A.  
Across all temperature conditions, the appearance of spalling pits and scuffing traces indicated that repeated frictional loading, in conjunction with the oxide film breakdown and particle detachment, contributed to the wear track instability. Wear debris accumulated within craters is indicative of tribofilm breakdown and reformation cycles, which are closely correlated with fluctuations in the friction coefficient during sliding.