**4. Discussion**

This study shows the evolution of structural and phase transformations in the Al-Zn-Mg-Fe-Ni alloy with different histories during severe plastic deformation. An ingot after homogenization annealing (nickalin 1) and a rod after RSR (nickalin 2) served as initial specimens for HPT.

The microstructural studies, including SEM and high resolution TEM have shown that the structure formation mechanisms in common during HPT are as follows.

HPT results in an intensive refinement of the Al matrix structure to the nanocrystalline state. The formation of a dispersed structure follows two mechanisms: fragmentation at a true strain *ε* = 6.0 to 6.7 (*n* = 5, 10) and in-situ low-temperature dynamic recrystallization at *ε* = 7.1 (*n* = 15). The fragmented structure is characterized by the diffusion contrast of HABs, high dislocation density and the presence of LABs inside the grains/subgrains. The mixed structure consisting of deformation-induced fragments with LABs and HABs and recrystallized grains separated by HABs is bimodal, and this provides the best combination of strength and plastic properties. The average grain/subgrain size in HPT nickalin 1 is 100 nm, and it is 60 nm in HPT nickalin 2. Such nanostructures exhibit high hardness HV = 2000 to 2600 MPa. The strengthening mechanisms for various materials after HPT were repeatedly proved in experiments reported in [19–21,35–37]; namely, according to the Hall–Petch equation, grain-boundary strengthening increases intensely, and the abundance of structural defects is responsible for the high value of the dislocation strengthening component.

The quantitative data obtained by TEM sugges<sup>t</sup> that, during HPT, Al9FeNi eutectic aluminides crash into fragments whose size becomes commensurate with the secondary phases; the fragments are uniformly distributed in the matrix volume. If the Al9FeNi particles were micron-sized (2 to 5 μm) in the macrocrystalline analogs (the ingot and the rod), their size does not exceed 1 μm after HPT, and there appear nanoparticles when *n* = 15.

According to the TEM data, the phase composition of the alloy changes during HPT, the sequence and kinetics of the phase transformations being dependent on the history of the macrocrystalline analogs. As was noted in Section 3.2, during HPT of nickalin 1, at *ε =* 7.1, a supersaturated Al solid solution is formed due to the strain-induced dissolution of the secondary T-phase. Additional solid-solution hardening combined with grain-boundary and dislocation components provides high hardness to the material, HV = 2500 MPa.

Additionally, the formation of a supersaturated Al solid solution during HPT can be proved by the results obtained during subsequent post-deformation heat treatment [31,32]. According to this data, low-temperature annealing (T = 120–160 ◦C) is accompanied by the decomposition of the supersaturated Al solid solution, which occurs simultaneously with the thermal softening of the alloy. Dispersed particles of the T and η strengthening phases precipitate as a result of artificial aging. Due to aging, the hardening achieved under HPT remains high. The presence of the nanostructure significantly reduces the duration of the maximum effect of Al solid solution decomposition from 3–6 h to 0.5–1.0 h. Consequently, the thermal deformation processing of nickalin 1 forms a two-phase material consisting of a nanostructured supersaturated Al solid solution and micron-sized Al9FeNi aluminides of eutectic origin.

The phase transformations in nickalin 2 under HPT are more diverse. Prior to RSR deformation, the rod had a multiphase composition comprising, besides the Al matrix and Al9FeNi aluminides, the Al3Zr and MgZn2 secondary phases. The conditions of their formation are described in Section 3.1.

Under HPT these phases alter their morphologies and sizes owing to two opposite processes induced by SPD, namely the strain-induced dissolution of intermetallics and the kinetic strain aging of the Al matrix. As a result of HPT (*n* = 5, *ε* = 6.0), the MgZn2 phase dissolves with the accumulation of mobile dislocations, which, while moving, cut the particles and transfer magnesium and zinc atoms into the Al solid solution. An increase in true strain to *ε* = 6.7–7.1 (*n* = 10–15) intensifies the subsequent decomposition of the supersaturated solid solution under kinetic strain aging. This inversion of the matrix composition results in that the stick-shaped particles of the η phase disappear and that dispersed (40 nm) globular precipitates of the T-phase appear along the grain/subgrain boundaries of the Al matrix.

One of the possible reasons for the heterogeneous nucleation of secondary phases at the grain/subgrain boundaries is the appearance of abnormal grain-boundary segregations of the atoms of alloying elements under HPT, whose formation was proved theoretically [38] and experimentally [37,39]. In this connection, the precipitation of the secondary phases along the grain boundaries is easier, and they become dominant. Besides, the size of the particles and their discrete arrangemen<sup>t</sup> along the boundaries of the nanocrystalline structure does not decrease the alloy plasticity. Note that a distinctive feature of the morphology of the zinc-magnesium phases precipitating at low temperature under SPD is their globular shape.

As deformation with *n* = 10 increases (*ε* = 6.7), strain aging further develops, inside the grains/subgrains of the Al matrix there appear nanosized precipitates of the η' phase, the quantity of which increases with the amount of strain, *ε* = 7.1 (*n* = 15). The reasons for the alternation and competition of the T and η' phase precipitations under dynamic aging can be explained by referring to the results reported in [40], where it was found by atom probe tomography that, at an early stage of aging, clusters form in the Al-Zn-Mg-Cu alloy, with different Zn/Mg atomic ratios depending on Zn content in the supersaturated solid solution. With the atomic ratio Zn/Mg < 1.3, the clusters transform into the T-phase, and when Zn/Mg > 1.3 they turn into the η'-phase. As a result of the nonuniform deformation of the specimen volumes during HPT, the local regions of the strain-induced supersaturated Al solid solution have different combinations of alloying element concentrations. Accordingly, clusters appear with different Zn/Mg atomic ratios, which serve as centers for heterogeneous nucleation of strengthening phase precipitates.

The possibility of the occurrence of the above phase transformations during lowtemperature SPD stems from a several orders of magnitude increase in the atomic diffusion coefficients of the alloying elements in the Al matrix due to excess vacancy concentration [41,42]. Besides, the formation of a strain-induced supersaturated solid solution may be due to the pinning and transfer of alloying element atoms on dislocations during their migration.

The XRD analysis confirms the TEM results and based on the change in the matrix lattice parameter before and after HPT from *a* = 0.4064 ± 0.0001 nm (*n* = 0) to *a* = 0.4067 ± 0.0002 nm (*n* = 15).

The analysis of the structural and strength characteristics of the Al-Zn-Mg-Fe-Ni alloy after HPT shows their interrelation. The difference between the phase compositions of nickalin 1 and nickalin 2 causes the different mechanical characteristics of these materials due to the action of different structural strengthening mechanisms. The total strength of nickalin 1 (YS = 450 MPa, UTS = 470 MPa) is the sum of grain-boundary strengthening by the Hall–Petch mechanism, dislocation strengthening resulting from high dislocation density, and solid-solution strengthening. The higher values of the strength properties of HPT nickalin 2 (YS = 628 MPa, UTS = 640 MPa), which is a multiphase nanocomposite, are caused by the additional contribution of dispersion strengthening from the dispersed

particles of the secondary phase through the Orowan mechanism, as well as by the greatest contribution of the grain-boundary mechanism to the total strength due to the formation of smaller (half-size) grains. The comparison of the obtained strength characteristics with the corresponding properties of the Al-Zn-Mg-Fe-Ni alloy after four RSR cycles (YS = 410 MPa, UTS = 430 MPa) [33] shows the advantage of using combined deformation processing. Besides, the mechanical properties are comparable with those of other 7xxx series alloys after SPD and aging. For example, YS = 525 MPa, UTS = 547 MPa in the 7050 alloy after six ECAP passes (route Bc) [27] and YS = 520 MPa in the Al-Zn-Mg alloy after ECAP and aging [28].
