Microstructure and Mechanical Behavior of Al-Mg Composites Synthesized by Reactive Sintering

Lightweight metal matrix composites are synthesized from elemental powder mixtures of aluminum and magnesium using pressure-assisted reactive sintering. The effect of the reaction between aluminum and magnesium on the microstructure and mechanical properties of the composites due to the formation of β-Al3Mg2 and γ-Al12Mg17 intermetallics is investigated. The formation of the intermetallic compounds progressively consumes aluminum and magnesium and induces strengthening of the composites: the yield and compressive strengths increase with the increase of the content of intermetallic reinforcement at the expense of the plastic deformation. The yield strength of the composites follows the iso-stress model when the data are plotted as a function of the intermetallic content.


Introduction
Developments in the production of energy efficient and environmentally friendly vehicles have shifted the interest of the transport industry from conventional structural materials to high-strength lightweight components [1][2][3][4][5]. The reduction of the weight of vehicles can enhance the fuel efficiency and limit greenhouse gas emissions, ultimately improving environmental protection [4][5][6]. Because of their remarkable properties, including a low density, high specific strength and stiffness, and improved deformability [7][8][9][10][11][12], aluminum and magnesium-based alloys and composites are evolving as potential replacements for conventional materials in the automobile, ship and aerospace industries; electronic products; and industrial equipment [13][14][15][16][17][18].
Metallic reinforcements have several advantages over conventional reinforcing ceramic particulates, including easy and economical availability [35,36], improved recyclability of MMCs [37], better compatibility with a metallic matrix in terms of coefficient of thermal expansion and wettability [38][39][40], which may create a stronger interface between matrix and reinforcement [41]. The reaction between matrix and reinforcement can generate new phases during processing or post-processing heat treatment, thereby affecting the mechanical properties of the composite depending upon the volumetric changes associated with the transformation, nature of the new phases and their bonding with the surrounding material [8,42]. Such transformations may induce strengthening in MMCs, as, for example, reported in composites consisting of an aluminum matrix reinforced with Al-Cu-Fe quasicrystals, where the reaction generates a new phase (the ω-Al 7 Cu 2 Fe) and increases the strength of the material [39,43]. In another example, the reaction between the Fe 3 Al reinforcement and aluminum matrix produces two new phases (Al 5 Fe 2 and Al 13 Fe 4 ) that significantly enhance the strength of the composites [44]. Phase transformations due to the reaction between matrix and particulate reinforcements in the case of aluminum-magnesium matrix composites lead to the in-situ formation of several intermetallics, e.g., Al 12 Mg 17 , Al 3 Mg 2 , Al 2 MgC 2 , MgAl 2 O 4 , Al 3 Ti, Ti 3 B 4 [12,41,45], which has been reported as an effective way of composite strengthening along with other strengthening contributions.
Composite manufacturing through melting routes can lead to the heterogeneous distribution of the reinforcement due to agglomeration and segregation [46,47]. Furthermore, undesired and uncontrolled interfacial reactions between matrix and reinforcement can deteriorate the properties of the composites. Therefore, solid-state powder metallurgy represents a good choice among other processing techniques to produce near-net-shape products with uniformly dispersed reinforcements [48][49][50][51]. Powder metallurgy gives excellent control over the size, shape, distribution and volume fractions of the reinforcement and matrix, which, in turn, enables the creation of composites with desired microstructures and properties [52,53]. Furthermore, the extent of the interfacial reaction between matrix and reinforcement and the type of the reaction product(s) can be governed by properly controlling the process temperature [44,54,55].
The reaction between matrix and reinforcement achieved by powder metallurgy can be considered a type of reactive sintering, where the proper mixture of starting powders is sintered with or without the assistance of pressure in order to produce a variety of materials ranging from porous intermetallics or ceramics, crystallographic phase transformations, the formation of solid solutions and chemical compounds to nickel, and titanium aluminides and composites [56][57][58][59][60][61][62][63][64]. In this work, the reactive sintering of aluminum and magnesium elemental powder mixtures has been used to synthesize Al-Mg composites reinforced with Al 3 Mg 2 and Al 12 Mg 17 intermetallics formed in-situ during processing. The aim is to analyze the effect of the initial volume percent of magnesium on the microstructural modifications, in-situ formation of the intermetallic reinforcements and the resulting mechanical properties of the Al-Mg lightweight composites.

Materials and Methods
Commercially available aluminum and magnesium powders (-100 mesh) were blended by ball milling under an argon atmosphere for 1 h at a rotational speed of 100 rpm using a Retch PM 400 planetary ball mill (Retsch GmbH, Haan, Germany) with hardened steel balls and vials. To avoid a temperature rise during the blending process, milling was carried out for 15 min followed by rest intervals of 15 min. Powder mixtures with different amounts of Mg (V = 10, 20, 40, 60, 80, and 90 vol %) were hot-pressed under the argon atmosphere at 673 K by applying a uniaxial pressure of 640 MPa for 10 min. Additionally, to investigate the phase formation induced by the reaction between Al and Mg at different temperatures, aluminum, and magnesium bilayer specimens containing 50 vol % Mg were produced by hot pressing at 573, 623, and 673 K, while keeping the other hot pressing parameters constant (640 MPa pressure and 10 min holding time). The phases formed in the blended powders and hot-pressed composites were identified by X-ray diffraction (XRD) using a D3290 PANalytical X'pert PRO X-ray diffractometer (PANalytical, Almelo, The Netherlands) with Co Kα radiation (λ = 0.179 nm). A scanning electron microscope GEMINI 1530 (Carl Zeiss NTS GmbH, Oberkochen, Germany) equipped with an energy dispersive X-ray spectrometer (EDX, Bruker AXS Microanalysis GmbH, Berlin, Germany) was used for microstructural characterization. Fifty to sixty scanning electron microscopy (SEM) micrographs were examined for every composite to estimate the volume fraction of the different phases by using the image analysis software: ImageJ (1.48v, National Institutes of Health, Bethesda, MD, USA). ImageJ counts the pixels of the highlighted phase to calculate its percentage based on a grayscale. The 'threshold' function is used to select the required phase by adjusting the upper and lower limits of the grayscale, then the 'analysis' function is used to count the number of pixels present in the selected phase. The density of the hot-pressed composites was measured using the Archimedes method, giving an average relative density of the composites of 99 ± 0.3 g/cm 3 . Cylindrical samples of 3 mm diameter and 6 mm length were prepared from the hot-pressed composites and then tested at room temperature under quasistatic compression using an Instron 5869 universal testing machine (strain rate = 8 × 10 −5 s −1 , Instron, Norwood, MA, USA). A Fiedler laser-extensometer (Fiedler Optoelektronik GmbH, Lützen, Germany) was used to measure the sample strain. To ensure the reproducibility of the results, a minimum of four samples for each composite was tested. The compression tests were deliberately stopped at 20% strain for the composite containing 10 vol % Mg because of its high deformability. was measured using the Archimedes method, giving an average relative density of the composites of 99 ± 0.3 g/cm 3 . Cylindrical samples of 3 mm diameter and 6 mm length were prepared from the hot-pressed composites and then tested at room temperature under quasistatic compression using an Instron 5869 universal testing machine (strain rate = 8 × 10 −5 s −1 , Instron, Norwood, MA, USA). A Fiedler laser-extensometer (Fiedler Optoelektronik GmbH, Lützen, Germany) was used to measure the sample strain. To ensure the reproducibility of the results, a minimum of four samples for each composite was tested. The compression tests were deliberately stopped at 20% strain for the composite containing 10 vol % Mg because of its high deformability. Figure 1a shows the XRD patterns for the Al-Mg powder mixtures after 1 h of blending by ball milling. The patterns indicate the presence of pure aluminum and magnesium only, which illustrates the absence of any reaction during the mechanically-induced mixing process carried out in the present milling conditions. Aluminum and magnesium react during the hot consolidation of the powder mixtures at 673 K to create the β-Al3Mg2 and γ-Al12Mg17 intermetallics, as shown in Figure  1b. The XRD patterns for the consolidated composites with an initial content of 10 and 20 vol % Mg (V = 10 and 20) exhibit the presence of Al3Mg2 along with residual aluminum, whereas Al12Mg17 and magnesium are not observed. The composites with V = 40 and 60 contain both Al3Mg2 and Al12Mg17 accompanied by residual magnesium and aluminum. In contrast, the composites V = 80 and 90 show only the presence of Al12Mg17 and magnesium, while the Al3Mg2 and aluminum phases are not observed. The creation of Al3Mg2 and Al12Mg17 during the high-temperature processing of aluminum and magnesium composites is in compliance with the reported results [65][66][67][68][69]. The SEM images and EDX elemental maps and profiles for the consolidated composites are displayed in Figures 2-4. The composites with V = 10 and 20 show a core-shell microstructure that consists of a Mg core enclosed by two consecutive shells of Al12Mg17 and Al3Mg2 which are finally covered by a rather continuous residual aluminum ( Figure 2). The formation of multiphase in-situ core-shell microstructures has been reported in several composite systems [41,[70][71][72][73][74][75] and can be attributed to the incomplete transformation reactions under the current processing parameters, like temperature and pressure [44]. In the composite with V = 40, the continuity of the aluminum phase has been disrupted by the substantial growth of the intermetallics, resulting from the higher initial volume percentage of magnesium. Here, the Al3Mg2 phase is continuous throughout the The SEM images and EDX elemental maps and profiles for the consolidated composites are displayed in Figures 2-4. The composites with V = 10 and 20 show a core-shell microstructure that consists of a Mg core enclosed by two consecutive shells of Al 12 Mg 17 and Al 3 Mg 2 which are finally covered by a rather continuous residual aluminum ( Figure 2). The formation of multiphase in-situ core-shell microstructures has been reported in several composite systems [41,[70][71][72][73][74][75] and can be attributed to the incomplete transformation reactions under the current processing parameters, like temperature and pressure [44]. In the composite with V = 40, the continuity of the aluminum phase has been disrupted by the substantial growth of the intermetallics, resulting from the higher initial volume percentage of magnesium. Here, the Al 3 Mg 2 phase is continuous throughout the microstructure, the magnesium cores are surrounded by the Al 12 Mg 17 phase, and only a limited amount of isolated Al is observed. In the composite V = 60 ( Figure 3), the size of the isolated aluminum regions further decreases, while the size of the magnesium cores increases and both the Al 3 Mg 2 and Al 12 Mg 17 intermetallic phases are rather continuous throughout the microstructure. Conversely, in the composite V = 80, pure aluminum is not found and the Al 3 Mg 2 phase appears as cores encapsulated by the Al 12 Mg 17 phase, which is further covered by a continuous residual magnesium phase. The microstructure of the composite V = 90 consists of islands of Al 12 Mg 17 phase surrounded by residual magnesium. Phase identification in the composites is supported by the EDX elemental concentration profiles shown in Figure 4, which are in good agreement with the XRD results in Figure 1b. microstructure, the magnesium cores are surrounded by the Al12Mg17 phase, and only a limited amount of isolated Al is observed. In the composite V = 60 ( Figure 3), the size of the isolated aluminum regions further decreases, while the size of the magnesium cores increases and both the Al3Mg2 and Al12Mg17 intermetallic phases are rather continuous throughout the microstructure. Conversely, in the composite V = 80, pure aluminum is not found and the Al3Mg2 phase appears as cores encapsulated by the Al12Mg17 phase, which is further covered by a continuous residual magnesium phase. The microstructure of the composite V = 90 consists of islands of Al12Mg17 phase surrounded by residual magnesium. Phase identification in the composites is supported by the EDX elemental concentration profiles shown in Figure 4, which are in good agreement with the XRD results in Figure  1b.   microstructure, the magnesium cores are surrounded by the Al12Mg17 phase, and only a limited amount of isolated Al is observed. In the composite V = 60 (Figure 3), the size of the isolated aluminum regions further decreases, while the size of the magnesium cores increases and both the Al3Mg2 and Al12Mg17 intermetallic phases are rather continuous throughout the microstructure. Conversely, in the composite V = 80, pure aluminum is not found and the Al3Mg2 phase appears as cores encapsulated by the Al12Mg17 phase, which is further covered by a continuous residual magnesium phase. The microstructure of the composite V = 90 consists of islands of Al12Mg17 phase surrounded by residual magnesium. Phase identification in the composites is supported by the EDX elemental concentration profiles shown in Figure 4, which are in good agreement with the XRD results in Figure  1b.     Figure 5 illustrates the volume percentages of the different phases in the consolidated composites, estimated from the SEM images, as a function of the initial volume content of magnesium (V). The quantity of Al3Mg2 progressively increased from ~15 vol % for V = 10 to ~55 vol % for V = 40. Afterward, it gradually declined to ~1.5 vol % for V = 80 and, finally, it was not present anymore in the composite with V = 90. The amount of Al12Mg17 slowly rose to ~12 vol % for V = 40, then sharply increased to 50 and 53 vol % for V = 60 and 80, respectively, and finally decreased to about 25 vol % for V = 90. The amount of aluminum decreased continuously from ~85% for V = 10 to ~3% for V = 60 and then it diminished to zero in the composites V = 80 and 90. The quantity of residual magnesium for V = 40 was ~2 vol % and then it steadily increased to ~75 vol % for V = 90. Figure 5 divides the consolidated composites into two sections separated by vertical lines on the basis of the volume profiles intersections. At about V = 55, the amount of aluminum and magnesium were equal. Additionally, the volume profiles of Al12Mg17 and Al3Mg2 also intersected at the same point. This suggests that the microstructure of the composites from V = 10 to 55 consist of aluminum as the major soft phase along with Al3Mg2 and Al12Mg17 as dominant reinforcements, whereas for V > 55, magnesium becomes the main soft phase predominantly reinforced by Al12Mg17, except for the composite with V = 60, where both Al12Mg17 and Al3Mg2 are present in high percentages.    Figure 5 illustrates the volume percentages of the different phases in the consolidated composites, estimated from the SEM images, as a function of the initial volume content of magnesium (V). The quantity of Al3Mg2 progressively increased from ~15 vol % for V = 10 to ~55 vol % for V = 40. Afterward, it gradually declined to ~1.5 vol % for V = 80 and, finally, it was not present anymore in the composite with V = 90. The amount of Al12Mg17 slowly rose to ~12 vol % for V = 40, then sharply increased to 50 and 53 vol % for V = 60 and 80, respectively, and finally decreased to about 25 vol % for V = 90. The amount of aluminum decreased continuously from ~85% for V = 10 to ~3% for V = 60 and then it diminished to zero in the composites V = 80 and 90. The quantity of residual magnesium for V = 40 was ~2 vol % and then it steadily increased to ~75 vol % for V = 90. Figure 5 divides the consolidated composites into two sections separated by vertical lines on the basis of the volume profiles intersections. At about V = 55, the amount of aluminum and magnesium were equal. Additionally, the volume profiles of Al12Mg17 and Al3Mg2 also intersected at the same point. This suggests that the microstructure of the composites from V = 10 to 55 consist of aluminum as the major soft phase along with Al3Mg2 and Al12Mg17 as dominant reinforcements, whereas for V > 55, magnesium becomes the main soft phase predominantly reinforced by Al12Mg17, except for the composite with V = 60, where both Al12Mg17 and Al3Mg2 are present in high percentages.

Phase Formation During the Reaction of Al and Mg
In order to further examine the phase formation induced by the reaction between Al and Mg, aluminum and magnesium bilayer specimens were produced by hot pressing at different temperatures (573, 623, and 673 K). The corresponding SEM images and EDX elemental concentration profiles across the interface between the Al and Mg layers are shown in Figure 6. In a homogenous composite, phase transformations occurring at an Al-Mg interface may have an influence on the diffusion and reaction fields arising from its neighboring similar interfaces. The synthesis of the bilayer composites helps to analyze the sequence of the intermetallics formation at the Al-Mg interface in the absence of any reactional caused by the neighboring interfaces. The Al-Mg bilayer composite hot-pressed at 573 K does not display any reaction product (Figure 6a-c). Because the diffusion coefficient of aluminum is higher than that of magnesium [76][77][78], it is plausible to assume that at a higher processing temperature (623 K) the aluminum atoms diffuse into the magnesium layer and form Al 12 Mg 17 as the first reaction product (Figure 6d-f). The Al 3 Mg 2 phase may then be generated by the additional diffusion of aluminum into the Al 12 Mg 17 intermetallic and by the diffusion of magnesium into the Al matrix [79,80]. At this stage, the reaction is rather discontinuous and a rough layer of Al 12 Mg 17 and Al 3 Mg 2 intermetallics is formed at the Al-Mg interface. The interfacial layer is composed of a 3 ± 1 µm thick Al 12 Mg 17 layer next to magnesium and a 4 ± 3 µm thick Al 3 Mg 2 layer next to aluminum. The existence and position of both intermetallic layers at the Al-Mg interface observed here are in accordance with previous studies [81][82][83][84][85]. A further increase of temperature to 673 K smoothens the interface and enlarges the widths of the intermetallic layers to 5 ± 1 µm thick Al 12 Mg 17 and 14 ± 1 µm thick Al 3 Mg 2 . The limited growth of the Al 12 Mg 17 layer can be attributed to its smaller value of the inter-diffusion coefficient in comparison to Al 3 Mg 2 [80,86].

Phase Formation During the Reaction of Al and Mg
In order to further examine the phase formation induced by the reaction between Al and Mg, aluminum and magnesium bilayer specimens were produced by hot pressing at different temperatures (573, 623, and 673 K). The corresponding SEM images and EDX elemental concentration profiles across the interface between the Al and Mg layers are shown in Figure 6. In a homogenous composite, phase transformations occurring at an Al-Mg interface may have an influence on the diffusion and reaction fields arising from its neighboring similar interfaces. The synthesis of the bilayer composites helps to analyze the sequence of the intermetallics formation at the Al-Mg interface in the absence of any reactional caused by the neighboring interfaces. The Al-Mg bilayer composite hot-pressed at 573 K does not display any reaction product (Figure 6a-c). Because the diffusion coefficient of aluminum is higher than that of magnesium [76][77][78], it is plausible to assume that at a higher processing temperature (623 K) the aluminum atoms diffuse into the magnesium layer and form Al12Mg17 as the first reaction product (Figure 6d-f). The Al3Mg2 phase may then be generated by the additional diffusion of aluminum into the Al12Mg17 intermetallic and by the diffusion of magnesium into the Al matrix [79,80]. At this stage, the reaction is rather discontinuous and a rough layer of Al12Mg17 and Al3Mg2 intermetallics is formed at the Al-Mg interface. The interfacial layer is composed of a 3 ± 1 μm thick Al12Mg17 layer next to magnesium and a 4 ± 3 μm thick Al3Mg2 layer next to aluminum. The existence and position of both intermetallic layers at the Al-Mg interface observed here are in accordance with previous studies [81][82][83][84][85]. A further increase of temperature to 673 K smoothens the interface and enlarges the widths of the intermetallic layers to 5 ± 1 μm thick Al12Mg17 and 14 ± 1 μm thick Al3Mg2. The limited growth of the Al12Mg17 layer can be attributed to its smaller value of the inter-diffusion coefficient in comparison to Al3Mg2 [80,86].  A closer look at the interface of the bilayer hot-pressed at 623 K (Figure 7a) reveals additional features characterizing the surface of the aluminum layer. Every area on the aluminum surface undergoing the reaction with Mg is concave, whereas the intermetallics are hump-shaped islands with a similar morphology throughout the specimen. As already mentioned, the diffusion coefficient of aluminum is one to two orders of magnitude higher than that of magnesium [77,87]. The aluminum atoms diffuse through the interface with a higher flux than that of magnesium and when the composition of the γ-Al 12 Mg 17 intermetallic phase is locally reached, it nucleates and progressively grows into the magnesium layer [77,79,83]. The preferential formation of the intermetallic phases at the concave zones of the aluminum surface indicates that the stoichiometric concentration of aluminum atoms to form the intermetallics is reached faster in the concave zone than at the convex or straight zones. This aspect is schematically depicted in Figure 7b, which shows the linear paths of the diffusion of aluminum atoms in the Al-Mg bilayer. The diffusion line density in the concave zones is larger than that in the convex zones. As a result, the stoichiometric concentration in the concave zones is reached faster than in any other zone. A closer look at the interface of the bilayer hot-pressed at 623 K (Figure 7a) reveals additional features characterizing the surface of the aluminum layer. Every area on the aluminum surface undergoing the reaction with Mg is concave, whereas the intermetallics are hump-shaped islands with a similar morphology throughout the specimen. As already mentioned, the diffusion coefficient of aluminum is one to two orders of magnitude higher than that of magnesium [77,87]. The aluminum atoms diffuse through the interface with a higher flux than that of magnesium and when the composition of the γ-Al12Mg17 intermetallic phase is locally reached, it nucleates and progressively grows into the magnesium layer [77,79,83]. The preferential formation of the intermetallic phases at the concave zones of the aluminum surface indicates that the stoichiometric concentration of aluminum atoms to form the intermetallics is reached faster in the concave zone than at the convex or straight zones. This aspect is schematically depicted in Figure 7b, which shows the linear paths of the diffusion of aluminum atoms in the Al-Mg bilayer. The diffusion line density in the concave zones is larger than that in the convex zones. As a result, the stoichiometric concentration in the concave zones is reached faster than in any other zone. The bilayer composite hot-pressed at 673 K (Figure 8a) shows thicker layers of intermetallics. The Arrhenius diffusion equation predicts a higher atomic and vacancy mobility at higher temperatures [88], which, in turns, explains the larger thickness of the intermetallic phases compared with the material hot-pressed at 623 K. An oxide layer is present along the aluminum layer facing the edge of β-Al3Mg2. This oxide layer represents the initial contact surface between the aluminum and magnesium particles [77] and its existence near aluminum indicates that the phase transformations occur within the magnesium region [77]. The presence of the boundaries of the pre-existing magnesium particle in the γ-Al12Mg17 phase supports this conclusion (compare Figure 8a,b). The interfaces between Mg and γ-Al12Mg17 and between γ-Al12Mg17 and β-Al3Mg2 do not show porosity, while the interface between β-Al3Mg2 and Al exhibits porosity that can be designated as Kirkendall's porosity [89]; this aspect can be ascribed to the higher mobility of aluminum atoms in comparison to magnesium and to the coalescence of vacancies towards aluminum [76,77,90]. The bilayer composite hot-pressed at 673 K (Figure 8a) shows thicker layers of intermetallics. The Arrhenius diffusion equation predicts a higher atomic and vacancy mobility at higher temperatures [88], which, in turns, explains the larger thickness of the intermetallic phases compared with the material hot-pressed at 623 K. An oxide layer is present along the aluminum layer facing the edge of β-Al 3 Mg 2 . This oxide layer represents the initial contact surface between the aluminum and magnesium particles [77] and its existence near aluminum indicates that the phase transformations occur within the magnesium region [77]. The presence of the boundaries of the pre-existing magnesium particle in the γ-Al 12 Mg 17 phase supports this conclusion (compare Figure 8a,b). The interfaces between Mg and γ-Al 12 Mg 17 and between γ-Al 12 Mg 17 and β-Al 3 Mg 2 do not show porosity, while the interface between β-Al 3 Mg 2 and Al exhibits porosity that can be designated as Kirkendall's porosity [89]; this aspect can be ascribed to the higher mobility of aluminum atoms in comparison to magnesium and to the coalescence of vacancies towards aluminum [76,77,90].  Figure 9a,b display the room temperature compression stress-strain curves for the composites and the summary of the corresponding mechanical data. The yield strength (0.2% offset) of the composites increases from 126 ± 7 MPa for the material with V = 10 to 562 ± 23 MPa for the composite V = 60 and then decreases to 207 ± 20 MPa for the composite V = 90. Similarly, the compressive strength (the maximum stress that a material can stand under compression loading [91]) increases from 238 ± 7 MPa for the composites with V = 10 to 562 ± 23 MPa for V = 60, and afterward it decreases to 385 ± 20 MPa for the composites with V = 90. Conversely, the plastic deformability of the composites drastically decreases from 17% ± 0.2% to 1.2% ± 0.1% for the composites with V = 20 and 60, then it increases again to 11 ± 0.3% for the composite V = 90. The composite with V = 10 does not show failure up to the 20% strain, where the test was stopped. The yield strength of the composites cannot be directly correlated to the initial volume percentage of magnesium, as shown in Figure 9c. The in-situ formed Al3Mg2 and Al12Mg17 intermetallics are the reinforcing phases in the present composites. Therefore, the yield strength data are plotted in Figure 9d as a function of the cumulative amount of both intermetallics along with the strength values predicted by the iso-stress (red dashed lines) and iso-strain (blue dot-dashed line) models of the rule of mixture (RoM) [92][93][94]. The yield strength of the composites increases with increasing the total amount of intermetallics. The values of strength are closer to the iso-stress model, which assumes that the matrix and reinforcement experience the same stress during deformation, while the iso-strain model, which assumes that all components of a composite undergo the same strain during deformation [93], overestimates the values. Under compressive loading, the reinforcement particles move in the loading direction to get closer to each other if the volume fraction of reinforcement is high and the distance between the reinforcement particles is small [95]. In such cases, the reinforcement particles may undergo deformation due to load transfer and the strength of the composites follows the iso-strain model [95]. On the other hand, a smaller volume fraction of reinforcement leads to a larger distance between reinforcement particles and, in such cases, deformation takes place mainly in the matrix. The iso-stress model describes the strength behavior of this type of composite [95][96][97]. Therefore, the deviation of the yield strength data for the composites with V = 40, 60 and 80 could be attributed to the high volume fraction of the reinforcement and complex microstructure. The iso-stress behavior for aluminum matrix composites and the deviation of the yield strength from it have been observed in previous studies [44,92,97,98].  Figure 9a,b display the room temperature compression stress-strain curves for the composites and the summary of the corresponding mechanical data. The yield strength (0.2% offset) of the composites increases from 126 ± 7 MPa for the material with V = 10 to 562 ± 23 MPa for the composite V = 60 and then decreases to 207 ± 20 MPa for the composite V = 90. Similarly, the compressive strength (the maximum stress that a material can stand under compression loading [91]) increases from 238 ± 7 MPa for the composites with V = 10 to 562 ± 23 MPa for V = 60, and afterward it decreases to 385 ± 20 MPa for the composites with V = 90. Conversely, the plastic deformability of the composites drastically decreases from 17% ± 0.2% to 1.2% ± 0.1% for the composites with V = 20 and 60, then it increases again to 11 ± 0.3% for the composite V = 90. The composite with V = 10 does not show failure up to the 20% strain, where the test was stopped. The yield strength of the composites cannot be directly correlated to the initial volume percentage of magnesium, as shown in Figure 9c. The in-situ formed Al 3 Mg 2 and Al 12 Mg 17 intermetallics are the reinforcing phases in the present composites. Therefore, the yield strength data are plotted in Figure 9d as a function of the cumulative amount of both intermetallics along with the strength values predicted by the iso-stress (red dashed lines) and iso-strain (blue dot-dashed line) models of the rule of mixture (RoM) [92][93][94]. The yield strength of the composites increases with increasing the total amount of intermetallics. The values of strength are closer to the iso-stress model, which assumes that the matrix and reinforcement experience the same stress during deformation, while the iso-strain model, which assumes that all components of a composite undergo the same strain during deformation [93], overestimates the values. Under compressive loading, the reinforcement particles move in the loading direction to get closer to each other if the volume fraction of reinforcement is high and the distance between the reinforcement particles is small [95]. In such cases, the reinforcement particles may undergo deformation due to load transfer and the strength of the composites follows the iso-strain model [95]. On the other hand, a smaller volume fraction of reinforcement leads to a larger distance between reinforcement particles and, in such cases, deformation takes place mainly in the matrix. The iso-stress model describes the strength behavior of this type of composite [95][96][97]. Therefore, the deviation of the yield strength data for the composites with V = 40, 60 and 80 could be attributed to the high volume fraction of the reinforcement and complex microstructure. The iso-stress behavior for aluminum matrix composites and the deviation of the yield strength from it have been observed in previous studies [44,92,97,98]. These results are in line with the data reported for Al-based composites reinforced with β-Al3Mg2 particles [101] and indicate that lightweight composites with a wide range of strength and plastic deformation can also be successfully synthesized by the reactive sintering of the proper elemental powder mixtures. This processing route represents an additional degree of freedom in the development of lightweight materials with designed mechanical and physical properties able to meet specific requirements.

Conclusions
The reactive hot consolidation of Al-Mg powder mixtures has been used to synthesize lightweight composites reinforced with β-Al3Mg2 and γ-Al12Mg17 intermetallics formed in-situ during consolidation. Detailed microstructural characterization and comprehensive phase analysis have been employed to examine the morphology of the phases in the composites and to achieve their identification and quantification. The formation of core-shell microstructures is observed in the Alrich composites, where the Mg cores are surrounded by a double-shell of β-Al3Mg2 and γ-Al12Mg17 intermetallic phases. In the Mg-rich composites, aluminum has been fully consumed except in the composite V = 60. Microstructural examinations of bilayer composites suggest that γ-Al12Mg17 is the  [101] and indicate that lightweight composites with a wide range of strength and plastic deformation can also be successfully synthesized by the reactive sintering of the proper elemental powder mixtures. This processing route represents an additional degree of freedom in the development of lightweight materials with designed mechanical and physical properties able to meet specific requirements.

Conclusions
The reactive hot consolidation of Al-Mg powder mixtures has been used to synthesize lightweight composites reinforced with β-Al 3 Mg 2 and γ-Al 12 Mg 17 intermetallics formed in-situ during consolidation. Detailed microstructural characterization and comprehensive phase analysis have been employed to examine the morphology of the phases in the composites and to achieve their identification and quantification. The formation of core-shell microstructures is observed in the Al-rich composites, where the Mg cores are surrounded by a double-shell of β-Al 3 Mg 2 and γ-Al 12 Mg 17 intermetallic phases. In the Mg-rich composites, aluminum has been fully consumed except in the composite V = 60. Microstructural examinations of bilayer composites suggest that γ-Al 12 Mg 17 is the first reaction product that originates at the interface between Al and Mg, followed by the formation of the β-Al 3 Mg 2 phase.
The mechanical behavior of the composites has been examined using room-temperature compression tests. The formation of the intermetallic phases increases the yield and compressive strengths of the composites, but decreases their plastic deformation. The strengthening effect depends on the amount of intermetallics formed during the consolidation step and, as a result, the yield strength of the composites follows the iso-stress model when the data are plotted as a function of the cumulative volume percentage of the intermetallic reinforcement.
Author Contributions: R.N.S. and S.S. designed the experiments, analyzed and interpreted the data, and wrote the paper; R.N.S. synthesized and characterized the specimens.