Al13Fe4-Al Composites with Nanocrystalline Matrix Manufactured by Hot-Pressing of Milled Powders

The paper describes composites with the matrix containing a nanocrystalline intermetallic Al13Fe4 phase and microcrystalline aluminium. Mechanically alloyed Al80Fe20 powder, containing a metastable nanocrystalline Al5Fe2 phase, was mixed with 20, 30, and 40 vol.% of Al powder and consolidated at 750 °C under the pressure of 7.7 GPa. During the consolidation, the metastable Al5Fe2 phase transformed into a nanocrystalline Al13Fe4 phase. In the bulk samples, Al13Fe4 areas were wrapped around by networking Al regions. The hardness of the Al13Fe4-Al composites was in the range of 4.52–5.50 GPa. The compressive strength of the Al13Fe4-30%Al and Al13Fe4-40%Al composites was 805 and 812 MPa, respectively, and it was considerably higher than that of the Al13Fe4-20%Al composite (538 MPa), which failed in the elastic region. The Al13Fe4-30%Al and Al13Fe4-40%Al composites, in contrast, showed some plasticity: namely, 1.5% and 9.1%, respectively. The density of the produced composites is in the range of 3.27–3.48 g/cm3 and decreases with the increase in the Al content.


Introduction
Aluminium-based alloys are characterised by low density and good corrosion resistance. Among aluminium alloys, attention is attracted by, i.a., intermetallic phases of the Fe-Al system [1]. They have advantageous properties: a high specific strength and stiffness, good strength at intermediate temperatures, and excellent corrosion resistance at elevated temperatures [1]. Properties such as hardness and strength can be improved when iron aluminides have a nanocrystalline structure [2,3]. A well-known method of producing nanocrystalline materials is mechanical alloying (MA), followed by the consolidation of milled powders [4]. However, most consolidation techniques do not allow for obtaining of truly nanocrystalline samples that are large in size and free from artefactual defects formed during consolidation [5]. It was observed that such defects are a frequent cause of a lack of plasticity in consolidated samples [6]. These defects are often micropores between weakly bound powder particles, and they cause bulk nanocrystalline materials to crack at stresses lower than the yield strength. Therefore, the technological processes of bulk nanocrystalline materials can be treated as critical factors influencing their low ductility, and the key to its improvement is the consolidation of the material free of artefact-like defects [6].
High-temperature consolidation of mechanically alloyed powders, with a nanocrystalline structure preserved, is a task that is difficult to perform. The application of high temperatures, which is required to achieve good quality consolidation, can lead to excessive Materials 2022, 15, 4241 2 of 13 grain growth. As we previously demonstrated, application of high pressure during consolidation inhibits grain growth and maintains the nanocrystalline structure of the pressed powder [7,8].
It was reported that the presence of larger grains in a nanocrystalline matrix (bimodal grain distribution) can improve a material's ductility in comparison with samples containing only nanometric grains [9]. A grain size distribution containing a fraction of grains large enough to maintain dislocation activity enables strain hardening, and thus favours an increase in ductility [6]. This prompted us to modify the microstructure of the consolidated powders to improve the ductility of the bulk materials [10,11]. We, therefore, produced the composite materials comprising a nanocrystalline intermetallic as a matrix and aluminium as a binder [10,11]. Ductile aluminium was to be distributed in the form of a "network" between hard nanocrystalline particles. The expected effect of the presence of a soft aluminium "network" was to improve the quality of bonding nanocrystalline particles into bulk material and, being a ductile component of the composite, to improve the plasticity of material. Micrometric aluminium grains contribute to the plastic deformation, and they may release stress concentrations and, hence, retard crack initiation and propagation [12]. The introduction of aluminium lowered the final density of the composites but also resulted in a decrease in the material hardness. Powders with nanocrystalline structure intended to form the matrix were produced by mechanical alloying [10,11]. They were then mixed with Al powder and subjected to short-term low-energy milling, which ensured even distribution of Al in the powder material. Mixtures prepared in such a way were consolidated into bulk form [10,11].
There are very few publications describing bulk materials containing ultrafine and larger grains, obtained by mixing nanocrystalline and microcrystalline powders, as well as the subsequent consolidation [9,12]. Apart from those reported by our group [10,11,13], no works related to nanocrystalline powders mixed with microcrystalline aluminium were found. It is, therefore, worth noting that the nanocrystalline phase-Al composites, as well as their technology, are innovative.
Recently, we prepared Al 3 Ni 2 -Al [10] and Al 5 Fe 2 -Al [11] composites with a nanocrystalline intermetallic matrix and various content of microcrystalline Al. The increase in composites' plasticity with the increase in Al content was observed in both cases. In the present work, composites with the matrix containing a nanocrystalline intermetallic Al 13 Fe 4 phase and microcrystalline Al are being reported. The Al 13 Fe 4 intermetallic phase has the density lower than the Al 3 Ni 2 or Al 5 Fe 2 intermetallic phases, and hence, the composites produced in this work will have a lower density than those reported previously. Based on the results of the previous studies [14], the alloy with Al 80 Fe 20 stoichiometry was selected for the production of Al 13 Fe 4 phase.
The produced Al 80 Fe 20 alloy powder was blended with 20, 30, and 40 vol.% of elemental Al powder and milled for 20 min with low intensity (ball-to-powder weight ratio 3:1) to ensure uniform mixture of the two components. Low-intensity milling was carried out under argon atmosphere in a SPEX 8000 D mill.
The Al 80 Fe 20 -Al powder mixtures were consolidated using a press (ASEA, Västerås, Sweden) equipped with a toroid-type high-pressure cell (Institute for High Pressure Physics of the Russian Academy of Sciences. Troitsk, Moscow Region, Russia). The sample was placed inside a graphite tube with 1 mm thick wall serving as a resistance heater. The heater was enclosed in a toroidal container made of lithographic limestone with a 25 mm outer diameter. The assembly was compressed uniaxially between two tungsten carbide anvils using a press. The anvils' shape and the gasket's material ensure that the pressure conditions are close to isostatic. The compaction processes were performed under a pressure of 7.7 GPa for 3 min. Pressure of 7.7 GPa was achieved with a compression force of 400 Tonnes. The pressing temperature was in the range of 600-800 • C. Pellets were loaded at the rate of 0.5 GPa/min prior to heating. The heating and cooling rate was 1000 • C/min. The consolidated samples were cylindrical in shape with a diameter of 5 mm and a height of about 4 mm.
The X-ray diffraction (XRD) investigations of the powders and consolidated samples were carried out in a Rigaku MiniFlex II X-ray diffractometer (Rigaku Corporation, Tokyo, Japan) operating with CuKα radiation.
A Hitachi S-3500N scanning electron microscope (SEM) (Hitachi high-tech corporation, Tokyo, Japan), equipped with an energy dispersive spectroscopy system (EDS), was used for examinations of the low-intensity milled powder mixtures and for the determination of chemical composition. Cross-section samples for SEM investigations were prepared by embedding the powder in a conducting resin and subsequently grinding and polishing it.
Particle size distribution measurements for the low-intensity milled powder mixtures were performed using a KµK mini3D (Kamika Instruments, Warszawa, Poland) analyser.
A Zeiss AXIOVERT 40 MAT light microscope (Carl Zeiss MicroImaging GmbH, Gottingen, Germany) and a Hitachi S-3500N SEM equipped with an EDS system were used for structural investigations and chemical analysis of the consolidated material. Bulk samples for these microscopy investigations were prepared by standard polishing techniques.
EBSD examination of a consolidated sample was performed using a Quanta 3D FEG high-resolution scanning electron microscope (FEI Company, Hillsboro, OR, USA). The sample for the EBSD investigation was prepared by grinding up to 4000 SiC paper, diamond suspension polishing and polishing with 0.1 µm SiO 2 suspension.
A Zwick Roell hardness tester (ZwickRoell GmbH & Co. KG, Ulm, Germany) was employed for the measurements of Vickers hardness (HV1) of the compacted samples (average value of at least 25 indentations).
A Zwick Roell Z250 testing machine (ZwickRoell GmbH & Co. KG, Ulm, Germany) was used for uniaxial compression tests of bulk materials. Experiments were performed under displacement control at a strain rate of 10 −4 s −1 .
A Gibertini E154 (Gibertini Elettronica, Milano, Italy) balance equipped with a device for measuring the density of solids was employed for the determination of the bulk samples density (Archimedes method). The mass measurements performed allowed to calculate the open porosity of the consolidated samples.

Results and Discussion
Phase development and structural evolution in the Al 80 Fe 20 powder mixture during mechanical alloying was studied by recoding of XRD patterns of the powders after various milling times (Figure 1). In the early stage of mechanical alloying (up to 15 h), a Fe(Al) solid solution was formed. This is evidenced by the gradual decrease in the intensity of the Al diffraction peaks compared to that of the Fe peaks and by the appearance of an asymmetry of the Fe diffraction profiles, which become wider on the low-angle side. Such features of the diffraction patterns accompanying the formation of a Fe(Al) solid solution were analysed in our earlier work [15]. In the diffraction pattern taken after 20 h of milling, a few broad peaks of low intensity appear. These peaks are assigned to Al 5 Fe 2 intermetallic phase. In the pattern for 25 h-milled powder, more well-defined peaks of an Al 5 Fe 2 phase are visible. In the XRD pattern after 30 h of milling, the peaks of Al and Fe(Al) are not present, which indicates that all the Al reacted with Fe, creating Al 5 Fe 2 intermetallic, are at least partially ordered. Additionally, in the pattern of the 35 h-milled powders, only the peaks of Al 5 Fe 2 phase are present. The observed phase evolution during mechanical alloying of the Al 80 Fe 20 powder mixture is similar to that described earlier [14]. powders, only the peaks of Al5Fe2 phase are present. The observed phase evolution during mechanical alloying of the Al80Fe20 powder mixture is similar to that described earlier [14]. In the Fe-Al phase equilibrium diagram, the Al5Fe2 phase exists for the Al concentration of 70-72 at.% [16]. Thus, the produced Al80Fe20 alloy powder has a metastable phase composition. Mechanical alloying is a nonequilibrium process and the phase composition of its products most often differs from that expected from the phase equilibrium diagrams.
It was found that, during the heating of the mechanically alloyed Al80Fe20 powder in a calorimeter up to 720 °C, the Al5Fe2 phase transformed into the Al13Fe4 [14]. For the concentration of 80% Al, there is a two-phase Al13Fe4+Al area in the Fe-Al phase equilibrium diagram [16]. Hence, the heating in a calorimeter moved the phase composition of the powder towards equilibrium.
There are also reports on synthesis of an unknown Al-rich phase [17] or an amorphous phase [18,19] by mechanical alloying of an Al80Fe20 powder mixture. Even for the same composition of the initial powder mixture, products of differing phase composition and structure can be synthesised by mechanical alloying, depending on applied parameters, such as type of mill, milling intensity, etc. [4,14]. In the Fe-Al phase equilibrium diagram, the Al 5 Fe 2 phase exists for the Al concentration of 70-72 at.% [16]. Thus, the produced Al 80 Fe 20 alloy powder has a metastable phase composition. Mechanical alloying is a nonequilibrium process and the phase composition of its products most often differs from that expected from the phase equilibrium diagrams.
It was found that, during the heating of the mechanically alloyed Al 80 Fe 20 powder in a calorimeter up to 720 • C, the Al 5 Fe 2 phase transformed into the Al 13 Fe 4 [14]. For the concentration of 80% Al, there is a two-phase Al 13 Fe 4 +Al area in the Fe-Al phase equilibrium diagram [16]. Hence, the heating in a calorimeter moved the phase composition of the powder towards equilibrium.
There are also reports on synthesis of an unknown Al-rich phase [17] or an amorphous phase [18,19] by mechanical alloying of an Al 80 Fe 20 powder mixture. Even for the same composition of the initial powder mixture, products of differing phase composition and structure can be synthesised by mechanical alloying, depending on applied parameters, such as type of mill, milling intensity, etc. [4,14].
In the XRD patterns of the powders after milling time longer than 6 h, all the diffraction peaks are broad, which indicates nanometric crystallite size in the phases existing in the powders.
The produced Al 80 Fe 20 alloy powder was blended with 20, 30, or 40 vol.% of Al and subjected to short-term low-energy milling. In the XRD patterns of the processed mixtures, only the diffraction peaks of Al 5 Fe 2 and Al are visible (see Figures 2a and 3a,c). Thus, the low-energy milling processes did not cause any reactions between the phases present in the powder blends.
In the XRD patterns of the powders after milling time longer than 6 h, all the diffraction peaks are broad, which indicates nanometric crystallite size in the phases existing in the powders.
The produced Al80Fe20 alloy powder was blended with 20, 30, or 40 vol.% of Al and subjected to short-term low-energy milling. In the XRD patterns of the processed mixtures, only the diffraction peaks of Al5Fe2 and Al are visible (see Figures 2a and 3a,c). Thus, the low-energy milling processes did not cause any reactions between the phases present in the powder blends. Figure 2. XRD patterns of the Al80Fe20-30%Al powder after: (a) short-term low-energy milling, (b) short-term low-energy milling and consolidation at 600 °C, (c) short-term low-energy milling and consolidation at 700 °C, (d) short-term low-energy milling and consolidation at 750 °C, and (e) short-term low-energy milling and consolidation at 800 °C. Figure 2. XRD patterns of the Al 80 Fe 20 -30%Al powder after: (a) short-term low-energy milling, (b) short-term low-energy milling and consolidation at 600 • C, (c) short-term low-energy milling and consolidation at 700 • C, (d) short-term low-energy milling and consolidation at 750 • C, and (e) short-term low-energy milling and consolidation at 800 • C.
To analyse particle size variation and particle shape, the powders after short-term low-energy milling were subjected to particle size distribution measurements and SEM investigations. Figure 4 illustrates particle size distribution for the short-term low-energy milled Al 80 Fe 20 -40%Al powder mixture as an example. The mean equivalent diameter of powder particles for this sample is 24.8 µm. For the other short-term low-energy milled powder mixtures, similar particle size distribution was found, and the mean equivalent diameter was 23.5 µm and 22.0 µm for the Al 80 Fe 20 -30%Al and Al 80 Fe 20 -20%Al powder mixtures, respectively. Figure 5 shows SEM images of the loose powder of the short-term low-energy milled Al 80 Fe 20 -40%Al mixture as an example. In the powder mixture, two kinds of particles are visible: one, of almost spherical shape, visible as bright in BSE mode, and the other, with lamellar shape (Figure 5c), visible as grey in BSE mode (Figure 5b). Materials 2022, 15, x FOR PEER REVIEW 6 of 14 Figure 3. XRD patterns of the Al80Fe20-20%Al powder and Al80Fe20-40%Al powder after: (a,c) short-term low-energy milling, (b,d) short-term low-energy milling, and consolidation at 750 °C.
To analyse particle size variation and particle shape, the powders after short-term low-energy milling were subjected to particle size distribution measurements and SEM investigations. Figure 4 illustrates particle size distribution for the short-term low-energy milled Al80Fe20-40%Al powder mixture as an example. The mean equivalent diameter of powder particles for this sample is 24.8 μm. For the other short-term low-energy milled powder mixtures, similar particle size distribution was found, and the mean equivalent diameter was 23.5 μm and 22.0 μm for the Al80Fe20-30%Al and Al80Fe20-20%Al powder mixtures, respectively.   Figure 5 shows SEM images of the loose powder of the short-term low-energy milled Al80Fe20-40%Al mixture as an example. In the powder mixture, two kinds of particles are visible: one, of almost spherical shape, visible as bright in BSE mode, and the other, with lamellar shape (Figure 5c), visible as grey in BSE mode (Figure 5b).  . Particle size distribution for the short-term low-energy milled Al80Fe20-40%Al powder mixture. Figure 5 shows SEM images of the loose powder of the short-term low-energy milled Al80Fe20-40%Al mixture as an example. In the powder mixture, two kinds of particles are visible: one, of almost spherical shape, visible as bright in BSE mode, and the other, with lamellar shape (Figure 5c), visible as grey in BSE mode (Figure 5b). EDS investigations revealed that the spherical particles are the Al80Fe20 alloy powder, while the lamellar particles are pure Al ( Figure 6). EDS investigations revealed that the spherical particles are the Al 80 Fe 20 alloy powder, while the lamellar particles are pure Al ( Figure 6).
The alloy powder particles are mostly agglomerates, but there also exist some fine individual particles ( Figure 5). Figure 7 depicts cross-section of an alloy powder particle. The EDS analysis of the alloy powder particles cross-sections demonstrated that the particles are chemically homogeneous, with an average composition of 78.5% Al and 21.5% Fe, as well as a composition fluctuation of 1%. For this composition, there is also a two-phase Al 13 Fe 4 +Al area in the Fe-Al phase equilibrium diagram [16]. Departure from the initial composition often occurs during mechanical alloying when one of the ground metals is much more plastic than the other and may stick to the grinding media (balls or container) a bit.
The prepared Al 80 Fe 20 -Al powders were consolidated by hot-pressing. The purpose was to produce a composite built of a nanocrystalline matrix, which is an Al 13 Fe 4 phase (closer to the equilibrium state than an Al 5 Fe 2 one for the Al 80 Fe 20 composition) and an Al network that wraps around the matrix areas. The alloy powder particles are mostly agglomerates, but there also exist some fine individual particles ( Figure 5). Figure 7 depicts cross-section of an alloy powder particle. The EDS analysis of the alloy powder particles cross-sections demonstrated that the particles are chemically homogeneous, with an average composition of 78.5% Al and 21.5% Fe, as well as a composition fluctuation of 1%. For this composition, there is also a two-phase Al13Fe4+Al area in the Fe-Al phase equilibrium diagram [16]. Departure from the initial composition often occurs during mechanical alloying when one of the ground metals is much more plastic than the other and may stick to the grinding media (balls or container) a bit. The use of high temperature during the consolidation of powders facilitates obtaining high quality compaction. However, in the case of multiphase powders, the application of high temperature can result in a reaction between phases. We reported this in the case of two-phase powders containing Al [11,20,21]. Therefore, the appropriate temperature must be used: high but not high enough to cause uncontrolled reactions. To select the pressing temperature, the Al 80 Fe 20 -30%Al powder was consolidated at 600, 700, 750, and 800 • C. The melting point of Al under atmospheric pressure is 660.3 • C, and increases with increasing pressure up to about 1180 • C at 7.7 GPa [22]. Thus, under the pressure applied in our work, we did not expect the melting of Al during consolidation at 700, 750, and 800 • C.
The XRD patterns of the Al 80 Fe 20 -30%Al powder before and after consolidation are shown in Figure 2. XRD investigations revealed that: (i) after consolidation at 600 and 700 • C, besides Al, the Al 5 Fe 2 phase is still present in the material; (ii) during consolidation at 750 • C the Al 5 Fe 2 phase transformed into the Al 13 Fe 4 phase; (iii) during consolidation at 800 • C Al 13 Fe 4 phase and unidentified phases formed. Based on these results, the temperature of 750 • C was chosen for the hot-pressing of the two other powders. In the XRD patterns of both consolidated powders, the diffraction peaks of the Al 13 Fe 4 phase and Al are present (Figure 3b,d). Overlapping of the peaks of the Al 13 Fe 4 phase made it impossible to reliably estimate the crystallite size of this phase, but the width of the peaks indicates the presence of a nanocrystalline structure in this phase. The prepared Al80Fe20-Al powders were consolidated by hot-pressing. The purpose was to produce a composite built of a nanocrystalline matrix, which is an Al13Fe4 phase (closer to the equilibrium state than an Al5Fe2 one for the Al80Fe20 composition) and an Al network that wraps around the matrix areas.
The use of high temperature during the consolidation of powders facilitates obtaining high quality compaction. However, in the case of multiphase powders, the application of high temperature can result in a reaction between phases. We reported this in the case of two-phase powders containing Al [11,20,21]. Therefore, the appropriate temperature must be used: high but not high enough to cause uncontrolled reactions. To select the pressing temperature, the Al80Fe20-30%Al powder was consolidated at 600, 700, 750, and 800 °C. The melting point of Al under atmospheric pressure is 660.3 °C, and increases with increasing pressure up to about 1180 °C at 7.7 GPa [22]. Thus, under the pressure applied in our work, we did not expect the melting of Al during consolidation at 700, 750, and 800 °C.
The XRD patterns of the Al80Fe20-30%Al powder before and after consolidation are shown in Figure 2. XRD investigations revealed that: (i) after consolidation at 600 and 700 °C, besides Al, the Al5Fe2 phase is still present in the material; (ii) during consolidation at 750 °C the Al5Fe2 phase transformed into the Al13Fe4 phase; (iii) during consolidation at 800 °C Al13Fe4 phase and unidentified phases formed. Based on these results, the temperature of 750 °C was chosen for the hot-pressing of the two other powders. In the XRD patterns of both consolidated powders, the diffraction peaks of the Al13Fe4 phase and Al are present (Figure 3b,d). Overlapping of the peaks of the Al13Fe4 phase made it impossible to reliably estimate the crystallite size of this phase, but the width of the peaks indicates the presence of a nanocrystalline structure in this phase.
Light microscopy (LM) and SEM methods were used to observe the microstructure of the bulk samples and check the quality of consolidation. Examples of microstructure  [11], and TiC-Al [13] composites produced in a similar manner as the ones in this work.
Microscopy examination of the bulk samples showed that their surface is smooth, without pores, cracks, or voids.
The results of the XRD examinations of the consolidated samples, presented above, indicated the presence of a nanocrystalline structure in the Al 13 Fe 4 phase being the composites matrix. To confirm a nanocrystallinity of the Al 13 Fe 4 phase, the Al 80 Fe 20 -40%Al bulk sample was investigated in SEM by the EBSD method. Figure 10 presents the EBSD map of an area of the matrix in the bulk sample. As it can be seen, the crystallite size of the Al 13 Fe 4 phase is below 100 nm. Thus, the produced bulk composites have nanocrystalline matrix. The preservation of the nanocrystalline structure of the intermetallic phase, during consolidation at high temperatures, is due to the use of high pressing pressure, as we have shown previously [7,8]. Since diffusion is involved in grain growth, the influence of pressure apparently comes from the fact that the diffusion coefficient decreases with the increase in pressure, usually by a factor ranging from 2 to 10 for a pressure of 1 GPa [23]. Hence, the use of high pressure during consolidation can reduce the mobility of grain boundaries at high temperature.
Examples of EDS maps of the Al13Fe4-20%Al sample (Figure 9) show that the matrix areas contain Al and Fe, while in the wrapping regions, only Al is present. Microscopy observation, coupled with the XRD and EDS results, indicates that the networking regions are Al and the other areas make up the Al13Fe4 matrix. Thus, the bulk samples have Al13Fe4-Al composite microstructure. Similar microstructure, i.e., Al wrapping networking regions between regions of the nanocrystalline matrix, was obtained in the case of Al3Ni2-Al [10], Al5Fe2-Al [11], and TiC-Al [13] composites produced in a similar manner as the ones in this work.
Microscopy examination of the bulk samples showed that their surface is smooth, without pores, cracks, or voids.   Examples of EDS maps of the Al13Fe4-20%Al sample (Figure 9) show that the matrix areas contain Al and Fe, while in the wrapping regions, only Al is present. Microscopy observation, coupled with the XRD and EDS results, indicates that the networking regions are Al and the other areas make up the Al13Fe4 matrix. Thus, the bulk samples have Al13Fe4-Al composite microstructure. Similar microstructure, i.e., Al wrapping networking regions between regions of the nanocrystalline matrix, was obtained in the case of Al3Ni2-Al [10], Al5Fe2-Al [11], and TiC-Al [13] composites produced in a similar manner as the ones in this work.
Microscopy examination of the bulk samples showed that their surface is smooth, without pores, cracks, or voids.   The produced composites were also characterised by compression tests, as well as hardness, density, and open porosity measurements. The hardness, compressive strength, plastic strain, and density values are given in Table 1. talline matrix. The preservation of the nanocrystalline structure of the intermetallic phase, during consolidation at high temperatures, is due to the use of high pressing pressure, as we have shown previously [7,8]. Since diffusion is involved in grain growth, the influence of pressure apparently comes from the fact that the diffusion coefficient decreases with the increase in pressure, usually by a factor ranging from 2 to 10 for a pressure of 1 GPa [23]. Hence, the use of high pressure during consolidation can reduce the mobility of grain boundaries at high temperature. The produced composites were also characterised by compression tests, as well as hardness, density, and open porosity measurements. The hardness, compressive strength, plastic strain, and density values are given in Table 1. The hardness of the composites decreases with the increase in the networking Al regions content. However, the compressive strength of the Al13Fe4-30%Al and Al13Fe4-40%Al composites is considerably higher than that of the Al13Fe4-20%Al one. This is because, during the compression test, the Al13Fe4-20%Al sample fractured catastrophically in the elastic region, crumbling into fine pieces. This is probably due to the insufficient bonding of the powder particles during the consolidation process. The Al13Fe4-30%Al and Al13Fe4-40%Al composites, in contrast, showed some plasticity, and, after failure, remained whole. However, the composite with 30% of Al, unlike that with 40% of Al, fractured shortly after yield. Thus, the presence of more than 20% of the networking Al in the Al13Fe4-Al composites improved the interparticle bonding and the  The hardness of the composites decreases with the increase in the networking Al regions content. However, the compressive strength of the Al 13 Fe 4 -30%Al and Al 13 Fe 4 -40%Al composites is considerably higher than that of the Al 13 Fe 4 -20%Al one. This is because, during the compression test, the Al 13 Fe 4 -20%Al sample fractured catastrophically in the elastic region, crumbling into fine pieces. This is probably due to the insufficient bonding of the powder particles during the consolidation process. The Al 13 Fe 4 -30%Al and Al 13 Fe 4 -40%Al composites, in contrast, showed some plasticity, and, after failure, remained whole. However, the composite with 30% of Al, unlike that with 40% of Al, fractured shortly after yield. Thus, the presence of more than 20% of the networking Al in the Al 13 Fe 4 -Al composites improved the interparticle bonding and the ductility of the material. In addition, micrometric Al grains participate in the plastic deformation, and they can also help to release stress concentrations and, thus, retard crack initiation and propagation [12]. The compressive stress-strain curves are shown in Figure 11.
For the in-situ Al-Al 13 Fe 4 composite, prepared by mechanical alloying of Al 90 Fe 10 powder mixture followed by spark plasma sintering, the microhardness HV0.3 of 1.97 GPa and the compressive strength of 1130.1 MPa (but without any plasticity of material) were reported [24]. No other information concerning Al 13 Fe 4 -Al composites was found in the literature.
The density of the produced composites decreases with the increase in the Al content. The open porosity in all the produced composites does not exceed 0.15%. The specific compressive strength of the composites with 40% and 30% of Al is 248 kNm/kg and 235 kNm/kg, respectively, and is higher than that of Al 5 Fe 2 -Al (185-190 kNm/kg [11]) and Al3Ni2-Al (146-207 kNm/kg [10]) composites. ductility of the material. In addition, micrometric Al grains participate in the plastic deformation, and they can also help to release stress concentrations and, thus, retard crack initiation and propagation [12]. The compressive stress-strain curves are shown in Figure  11. For the in-situ Al-Al13Fe4 composite, prepared by mechanical alloying of Al90Fe10 powder mixture followed by spark plasma sintering, the microhardness HV0.3 of 1.97 GPa and the compressive strength of 1130.1 MPa (but without any plasticity of material) were reported [24]. No other information concerning Al13Fe4-Al composites was found in the literature.
The density of the produced composites decreases with the increase in the Al content. The open porosity in all the produced composites does not exceed 0.15%. The specific compressive strength of the composites with 40% and 30% of Al is 248 kNm/kg and 235 kNm/kg, respectively, and is higher than that of Al5Fe2-Al (185-190 kNm/kg [11]) and Al3Ni2-Al (146-207 kNm/kg [10]) composites.

Conclusions
Mechanically alloyed Al80Fe20 powder containing a metastable nanocrystalline Al5Fe2 phase was used as a matrix for the composites. Al80Fe20-X vol.% Al (X = 20, 30, and 40) powder mixtures were consolidated by hot-pressing. Hot-pressing tests have shown that the temperature of 750 °C is suitable for the consolidation, during which only the transformation of the Al5Fe2 phase into the nanocrystalline Al13Fe4 phase occurs, without other reactions. The microstructure of the bulk samples consisted of Al13Fe4 intermetallic matrix areas wrapped around by networking Al regions. With the increase in Al content, the hardness of the composites decreased, ranging from 4.52 to 5.50 GPa. The Al13Fe4-20%Al composite failed in the elastic region, while the composites with 30% and 40% of Al showed some plasticity: namely, 1.5% and 9.1%, respectively. The presence of more than 20% of Al in the Al13Fe4-Al composites improved their ductility. The compressive strength of the composites containing 30 and 40% of Al was 805 and 812 MPa, respectively, and it was considerably higher than that of the Al13Fe4-20%Al composite (538 MPa). The low-density materials produced have a relatively high hardness.

Conclusions
Mechanically alloyed Al 80 Fe 20 powder containing a metastable nanocrystalline Al 5 Fe 2 phase was used as a matrix for the composites. Al 80 Fe 20 -X vol.% Al (X = 20, 30, and 40) powder mixtures were consolidated by hot-pressing. Hot-pressing tests have shown that the temperature of 750 • C is suitable for the consolidation, during which only the transformation of the Al 5 Fe 2 phase into the nanocrystalline Al 13 Fe 4 phase occurs, without other reactions. The microstructure of the bulk samples consisted of Al 13 Fe 4 intermetallic matrix areas wrapped around by networking Al regions. With the increase in Al content, the hardness of the composites decreased, ranging from 4.52 to 5.50 GPa. The Al 13 Fe 4 -20%Al composite failed in the elastic region, while the composites with 30% and 40% of Al showed some plasticity: namely, 1.5% and 9.1%, respectively. The presence of more than 20% of Al in the Al 13 Fe 4 -Al composites improved their ductility. The compressive strength of the composites containing 30 and 40% of Al was 805 and 812 MPa, respectively, and it was considerably higher than that of the Al 13 Fe 4 -20%Al composite (538 MPa). The low-density materials produced have a relatively high hardness.

Conflicts of Interest:
The authors declare no conflict of interest.