**1. Introduction**

Efforts to increase aluminum alloys' strength without significant sacrifice in electrical conductivity is a current trend in industry. At the same time, the material used as an electrical conductor should demonstrate good thermal stability, which can be challenging in case of aluminum alloys [1]. With this in mind, finding ways to simultaneously improve mechanical and electrical properties is important. This combination is essential for creating new lightweight conductive materials for the electrical industry, and it can comprise ultrafine-grained Al-Fe alloys.

Aluminum alloys, particularly Al-Fe alloys, have several advantages as conductive materials. First, aluminum and iron are very common and cheap metals, which make them economically preferable. Second, the solubility of iron in aluminum for conventionally processed alloys, at temperatures ranging from room to near-melting, is close to zero [2]. This eliminates the major contribution to electrical resistance (i.e., solid solute atoms). The other contributions are grain boundaries, particles, and dislocation density [2].

Previous studies of Al-Fe alloys have found applications in conductive wires in automobiles [3–6] and as household cables. However, conventional approaches have reached the limit of increasing the mechanical strength and electrical conductivity of these materials. This is due to the absence of iron solubility in aluminum limiting the variations of precipitation morphology.

In conventionally produced aluminum alloys, there are usually two major intermetallic phases: Al13Fe4 and Al6Fe [7–9]. These phases precipitate as needle- and plate-like particles of relatively large size (microns to hundreds of microns). Conventional approaches, such as

**Citation:** Medvedev, A.; Murashkin, M.; Enikeev, N.; Medvedev, E.; Sauvage, X. Influence of Morphology of Intermetallic Particles on the Microstructure and Properties Evolution in Severely Deformed Al-Fe Alloys. *Metals* **2021**, *11*, 815. https://doi.org/10.3390/met11050815

Academic Editor: Babak Shalchi Amirkhiz

Received: 27 April 2021 Accepted: 14 May 2021 Published: 17 May 2021

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drawing/rolling etc., are not able to reduce these particles to the nanoscale, which would provide increased mechanical strength. Nanoscale particles could easily be formed from solid solution by simple heat treatment, but, as was mentioned earlier, formation of solid solution of Fe in Al is difficult to achieve via conventional methods.

Another possibility is to obtain a supersaturated solid solution thanks to the introduction of severe plastic deformation (SPD) [10,11]. As well as leading to grain size refinement, particle shredding, and an increase of defect density, SPD is also known for the formation of supersaturated solid solution in systems, which is considered very hard or nearly impossible [5,10].

A number of studies were conducted involving SPD of Al-Fe alloys. The concentration of Fe in a solid solution was increased up to 2 wt.% or more [10–25]. Such solid solutions may give an opportunity to homogeneously precipitate Al-Fe intermetallic particles in a way that was previously impossible.

The formation of strain-induced supersaturated solid solutions depends on a number of parameters. It was previously demonstrated that it could be influenced not only by the second-phase volume fraction [12], but also by the particle morphology [13]. Larger interphase areas naturally give more opportunities for solute atoms to migrate into the matrix. In addition, the second-phase morphology directly influences the fragmentation process during deformation. This effect was demonstrated in [13], where differences in the particle morphology before deformation affected microstructures and the mechanical and physical properties of the alloy after SPD. However, the underlying mechanisms were not deeply investigated, and to optimize the second-phase fragmentation process and the solid solution formation, one needs a systematic study of the influence of the primary as-cast microstructure (volume fraction, intermetallic phase, particle size, morphology, and distribution). To achieve this goal, the present study focuses on two alloys (Al-2Fe and Al-4Fe) with very different initial microstructures.

#### **2. Materials and Methods**

The samples of Al-2 wt.% Fe and Al-4 wt.% Fe alloys were prepared on the basis of A99 grade primary aluminum (GOST 11069-2001). Melting was carried out in a GRAFICARBO GF 1100 electric resistance furnace (Graficarbo S.R.L., Zorlesco, Italy) in a graphite crucible at 830 ◦C at Russia's National University of Science and Technology "MISIS". Ingots with a diameter of 20 mm and a length of 200 mm were obtained by casting into a graphite mold (at the 20 K/s cooling rate). The total amount of trace materials, including Si, did not exceed 0.1 wt.% for either of the alloys.

The cylinders were then sliced into discs (1.5 mm thick) using a wire-cutting APTA-120 machine (Delta-Test, Fryazino, Russia). HPT was carried out by placing the disc-shaped sample between two rotating anvils and applying high hydrostatic pressure to the sample, inducing a simple shear type deformation. In this study, discs of as-cast alloy were subjected to 20 revolutions of HPT performed at room temperature (RT) under a pressure of 6 GPa and a speed of 1 rpm in closed anvils. Such a processing route was chosen based on primary experiments showing that these parameters provide uniform microstructures [26,27] and, possibly, supersaturated solid solutions.

Microstructural, mechanical, and electrical data were collected from a specific location in the middle of the sample's radius 5 mm, thus rendering it comparable with other research.

Transmission electron microscopes (TEMs, JEOL, Tokyo, Japan) using different analysis modes (i.e., BF bright-field analysis, SAED-selected area electron diffraction) and STEM (scanning transmission electron microscopy-JEOL ARM 200F) were employed for the microstructure analysis. Objects for TEM were prepared by twin-jet electro-polishing on a Struers Tenupol-5 (Copenhagen, Denmark) with 20% nitric acid in methanol below −20 ◦C at an operating voltage of 20 V. To ensure statistically reliable results, at least three foils were used for each state.

Scanning electron microscopy (SEM) was performed on a ThermoFisher Helios G4 PFIB DualBeam (Waltham, MA, USA).

X-ray diffraction (XRD) analysis was conducted with a Rigaku Ultima IV (Tokyo, Japan) diffractometer using CuK α radiation (30 kV and 20 mA). Values of lattice parameter *a*, coherent scattering domain (CSD), size D, and elastic microdistortion level <sup>&</sup>lt;ε2>1/2 were calculated via the Rietveld refinement method using MAUD software (v.2.97, University of Trento, Trento, Italy) [28]. To calculate dislocation density (ρ), Equation (2) was used:

$$\rho = \frac{2\sqrt{3}\langle \epsilon^2 \rangle^{1/2}}{D \times b} \tag{1}$$

where *b* = *a* √2/2 is the Burgers vector for FCC metals, and *D* is the coherent scattering domain [29].

Tensile tests were carried out in triplicate using an Instron 5982 (Instron, Norwood, MA, USA) machine at RT and a strain rate of 10−<sup>3</sup> s<sup>−</sup>1. Yield stress ( *<sup>σ</sup>0.2*), ultimate tensile stress ( *<sup>σ</sup>UTS*), and ductility (*δ*), measured as elongation-to-failure, were obtained from the small samples with gauge dimensions of 1.0 mm × 1.0 mm × 4.0 mm, prepared by wire-cutting. The maximum error for sample dimensions was 0.02 mm.

Electrical conductivity ( ω) was determined with a ±2% error, using the eddy current method [30]. Electrical conductivity relative to annealed copper (International Annealed Copper Standard, %IACS) was calculated according to Equation (3):

$$IACS = \frac{\omega\_{Al}}{\omega\_{Cu}} \times 100\% \tag{2}$$

where *ωAl* is the measured electrical conductivity of aluminum alloy, and *ωCu* is the conductivity of annealed chemically pure copper (58 MS/m).
