## 4 Discussion  
The analysis of the as-cast structure of the studied alloy revealed Mg\({}_{2}\)Si and Cu-enriched Al\({}_{2}\)Cu and Q phases. The non-equilibrium Cu-rich phases dissolved during homogenization annealing whereas the Mg\({}_{2}\)Si particles were fragmented and spherodzied. The spherodzied and fragmentation of phases of eutectic origin during annealing positively influenced the structural homogeneity and particles distribution after thermomechanical treatment [54; 55]. The homogenization annealing of the alloy studied also provided for the formation of the nanoscale L1\({}_{2}\) phase dispersoids, which were homogeneously distributed in the Al-based matrix. The dispersoids provided a strong Zener pinning effect during MDF [56; 57; 58; 59] and hot deformation inhibiting the static and dynamic grain growth. It should be noted that the Si atoms can also dissolve in the Al\({}_{3}\)(Sc,Zr) phase, increasing its thermal stability due to the formation of an (Al,Si)\({}_{3}\)(Sc,Zr) phase [60; 61]. In the present study, dispersoids exhibited a size of \(\sim\)10 nm after a second high temperature homogenization step that can be a result of partial Al substitution by Si atoms, which increase the thermal stability of dispersoids. This is in agreement with our earlier studies; dispersoids with similar parameters were found after two-step homogenization in 6 xxx-type alloys [42; 62]. The increase in cumulative strain and MDF temperature resulted in an increase in particle distribution homogeneity. The homogeneously distributed microscale particles can contribute to grain refinement during thermomechanical treatment and improve deformation behavior due to the particles' stimulated nucleation effect (PSN) [46; 51; 63; 64; 65]. During MDF, the Mg\({}_{2}\)Si phase was noticeably fragmented, the result of both the temperature and strain effect. The Mg\({}_{2}\)Si particle size after MDF was in the range of 1.3-1.5 \(\upmu\)m. Despite a small volume fraction, these particles can provide for the PSN effect and additionally refine grain in the alloys studied. The positive effect of Mg\({}_{2}\)Si particles with similar parameters on grain refinement was previously discussed in refs. [7; 42]. The Sc-rich type-A and type-B phases observed in the as-forged sample exhibited a solidification origin. It was suggested that the Cu-bearing type-A phase corresponded to the W(AlCuSc) phase [66; 67]. The Si atoms may dissolve in W-phase by analogy with the (Al,Si)\({}_{3}\)(Sc,Zr) phase. The Cu-free type-B phase should correspond to the V(AlSc\({}_{2}\)Si\({}_{2}\)) phase discussed in refs. [60; 66]. The Cu-enriched microscale particles with bright contrast observed in the as-forged at 250-350 \({}^{\circ}\)C samples were formed due to the ageing that occurred during MDF. These particles corresponded to the Q-(AlMgSiCu)-phase [68; 69] or/and the Al\({}_{2}\)Cu phase. The MDF significantly refined grain structure in the alloy studied. The grain-refinement effect accompanied by a decrease in stress value with increasing cumulative strain reflects recrystallization duringac MDF. The decrease in temperature of MDF to 150 \({}^{\circ}\)C resulted in a decrease in grain/subgrain size. A high-volume fraction of LAGBs was observed after forging at 150 \({}^{\circ}\)C. This phenomenon suggests the continuous dynamic recrystallization [46] as the dominant recrystallization mechanism occurred during MDF of the alloy studied [70; 71]. The continuous dynamic recrystallization is the result of LAGBs transforming into HAGBs induced by the strain [46]. A similar process can be observed during the heating of a pre-strained alloy [71; 72; 73]. We suggest that the dynamic recrystallization is the dominant process at 150 \({}^{\circ}\)C, whereas at higher temperatures of 250-350 \({}^{\circ}\)C, post-dynamic and static recrystallization can also occur. The static recrystallization may occur when the sample is heated to a treatment temperature before starting the forging.