## 3 Results and Discussion  
### Microstructural Characterization  
#### 3.1.1 Phase Identification via X-Ray Diffraction (XRD) Analysis  
Figure 1 presents the X-ray diffraction (XRD) patterns of Al-Si-(Cu)-Mg alloys (Alloy A and Alloy B) before and after the T6 heat treatment. All diffraction peaks were indexed with reference to ICDD PDF-4 + cards (Al [FCC] 04-0787; Si 03-016-1163; \(\alpha\)-Al(FeMn)Si 35-102; \(\beta\)-AlFeSi 65-879; Mg\({}_{2}\)Si 01-077-2101; Al\({}_{2}\)Cu 03-065-6648; Q-phase Al\({}_{5}\)Cu\({}_{2}\)Mg\({}_{5}\)Si\({}_{6}\) 35-0784). In both alloys, characteristic diffraction peaks corresponding to Al, Si, \(\alpha\)-Al(FeMn)Si, \(\beta\)-AlFeSi, and Mg\({}_{2}\)Si phases were commonly observed, indicating the dominant phases formed under the given casting and heat treatment conditions. In the as-cast condition, Alloy A exhibited distinct diffraction peaks for Cu-rich intermetallics such as Al\({}_{2}\)Cu. However, these peaks either disappeared or were significantly \begin{table}
\begin{tabular}{c c c c c c c} \hline \hline
**Material** & **Temperature** & **Counter Body** & \
\begin{tabular}{c} **Counter Body** \\ **Diameter** \\ \end{tabular} & **Load** & **Sliding Radius** &
\begin{tabular}{c} **Rotational** \\ **Speed** \\ \end{tabular} \\ \hline Alloy A, B & RT & Al\({}_{2}\)O\({}_{3}\) & 12.7 mm & 30 N & 11.5 mm & 100 rpm \\ \hline \hline \end{tabular}
\end{table} 
Table 2: Test conditions for high-temperature ball-on-disk wear experiments conducted on Alloy A and Alloy B.  
Figure 1: (**a**) X-ray diffraction (XRD) patterns of Alloy A and Alloy B in the as-cast and T6-treated conditions and (**b**) Magnified view of the region highlighted by the purple box in (**a**). reduced in intensity after the T6 heat treatment. Simultaneously, the Al peak was slightly shifted to a higher angle, as shown in the magnified region (highlighted by the purple box) in Figure 1b. According to [PERSON]'s law (n\(\lambda\) = 2d sin\(\theta\)), this peak shift suggests a decrease in the interplanar spacing (d-spacing), implying lattice contraction [19]. Upon the T6 treatment, most Al\({}_{2}\)Cu and Q-phase peaks in Alloy A collapse or diminish below the XRD detection (<0.5 wt.%), whereas Mg\({}_{2}\)Si reflections emerge strongly in Alloy B (Figure 1a). The principal Al (FCC) 111 reflection at \(\approx\) 38.5\({}^{\circ}\) 20 also changes in shape and position. A detailed view (Figure 1b) reveals a composite reflection formed by overlapping \(\alpha\)-Al (111) and Al\({}_{2}\)Cu (211). After T6, the centroid of the \(\alpha\)-Al (111) component moves from 38.41\({}^{\circ}\)\(\pm\) 0.01\({}^{\circ}\) (AC_A) to 38.47\({}^{\circ}\)\(\pm\) 0.01\({}^{\circ}\) (HT_A), corresponding to a \(-\)0.16% decrease in d\({}_{111}\) (Bragg's law), which is indicative of Cu atoms entering the Al lattice and contracting its unit cell. The Williamson-Hall analysis (Table S1) shows the FWHM of the \(\alpha\)-Al (111) peak narrows from 0.31\({}^{\circ}\) to 0.23\({}^{\circ}\) upon T6 in Alloy A (versus only 0.33\({}^{\circ}\)\(\rightarrow\)0.26\({}^{\circ}\) in Alloy B), demonstrating a more complete dislocation annihilation and strain relief in the Cu-rich matrix during solutionizing (530 \({}^{\circ}\)C, 8 h). The pseudo-Voigt fitting of the 38.0-39.0\({}^{\circ}\) window (Figure S2, Table S1) separates the two contributions: the integrated area of Al\({}_{2}\)Cu (211) falls by \(\approx\) 75% (0.18\(\rightarrow\)0.04 a.u. ), while that of \(\alpha\)-Al (111) increases slightly (1.00\(\rightarrow\)1.08 a.u.). This confirms that the post-T6 increase in the composite peak intensity arises from both the removal of the Al\({}_{2}\)Cu shoulder and the sharpening of the \(\alpha\)-Al reflection. The measured lattice constant (a) of Alloy A decreased after the heat treatment, providing quantitative evidence for the incorporation of Cu into the Al lattice and the resultant solid solution strengthening.