Microstructure and Phase Formation of Novel Al80Mg5Sn5Zn5X5 Light-Weight Complex Concentrated Aluminum Alloys

In this work, three novel complex concentrated aluminum alloys were developed. To investigate the unexplored region of the multicomponent phase diagrams, thermo-physical parameters and the CALPHAD method were used to understand the phase formation of the Al80Mg5Sn5Zn5Ni5, Al80Mg5Sn5Zn5Mn5, and Al80Mg5Sn5Zn5Ti5 alloys. The ingots of the alloys were manufactured by a gravity permanent mold casting process, avoiding the use of expensive, dangerous, or scarce alloying elements. The microstructural evolution as a function of the variable element (Ni, Mn, or Ti) was studied by means of different microstructural characterization techniques. The hardness and compressive strength of the as-cast alloys at room temperature were studied and correlated with the previously characterized microstructures. All the alloys showed multiphase microstructures with major α-Al dendritic matrix reinforced with secondary phases. In terms of mechanical properties, the developed alloys exhibited a high compression yield strength up to 420 MPa, high compression fracture strength up to 563 MPa, and elongation greater than 12%.


Introduction
A great deal of effort has been devoted to the development of new lightweight materials during the last decades. Al-based alloys have received considerable attention, especially in the automotive and aerospace industries, due to the possibility of their light weight and consequent reduction of greenhouse gas emissions. Commercial Al alloys consist of Al as base element and small quantities of solute elements added to improve the mechanical and physical properties. In some cases, these restricted compositions lead to some limitations, such as their performance at high temperatures or low wear resistance.
One research focus with relevance to light weight in the transport industry includes the so-called high entropy alloys (HEAs) [1,2]. HEAs were defined as alloys of five or more metallic elements in equimolar or near-equimolar concentrations between 5 and 35 at.%. The main goal of this approach was to stabilize a single-phase solid solution (SS) by reducing Gibbs free energy to obtain simple crystal structures instead of brittle intermetallic compounds (IC). HEAs have received significant attention since initial reports in 2004 [3,4]. These alloys exhibit excellent mechanical properties, corrosion-oxidation resistance, and wear behavior [5][6][7][8][9], having numerous engineering applications [10].
Experimental alloys were prepared in an induction furnace VIP-I (Inductotherm Corp., Rancocas, NJ, USA) in an alumina crucible using 99.95% pure Al, Mg, Sn and Zn. Tablets of Al-Ni, Al-Mn and Al-Ti containing 75 wt.% of Ni, 80 wt.% of Mn, and 75 wt.% of Ti respectively were used. Firstly, Al was melted to guarantee a bath base where Mn, Ni or Ti was subsequently dissolved. Finally, Sn, Mg, and Zn were added to the molten alloy and held between 776 and 795 • C for 10 min to reach complete dissolution. Finally, the melt was cooled down to 700 • C and poured manually into a preheated steel mold at 200 • C.
Ingots of 200 mm × 80 mm × 40 mm were obtained for each alloy. The microstructural samples were cut from the ingots and prepared according to standard metallographic procedures, by hot mounting in conductive resin, grinding, and polishing. The X-ray diffraction (XRD) equipment used to characterize the crystal structures of the alloys was a D8 ADVANCE (BRUKER, Karlsruhe, Germany) diffractometer, with Cu Kα radiation, operated at 40 kV and 30 mA. The diffraction diagrams were measured at the diffraction angle 2θ, range from 10 • to 90 • with a step size of 0.01 • , and 1.8 s/step. The powder diffraction file (PDF) database 2008 was applied for phase identification. The microstructure, the different regions, and the averaged overall chemical composition of each sample were investigated by an optic microscope model DMI5000 M (LEICA Microsystems, Wetzlar, Germany) and a scanning electron microscope (SEM), equipped with an energy dispersive X-ray spectrometry (EDS) model JSM-5910LV (JEOL, Croissy-sur-Seine, France).
A Vickers hardness FV-700 model (FUTURE-TECH, Kawasaki, Japan) measurer was employed on the polished sample surface using a 1 kg load, applied for 10 s. At least 10 random individual measurements were made for each test. Cylindrical specimens for compression testing were machined from the ingots, with a diameter of 13 mm and a height of 26 mm, giving an aspect ratio of 1:2. Compression testing was performed using a MTS Insight 100 kN Extended Length (MTS Systems Corporation, Eden Prairie, MN, USA) at RT with a strain rate of 0.001 s −1 . For the accurate measurement of Young's modulus, a clip-on extensometer mounted on the specimens were used. At least five specimens were recorded to ensure repeatability. Finally, density measurement was conducted using the Archimedes method.

Thermo-Physical Parameters for Phase Formation in HEAs
In this work, the unexplored Al-Mg-Sn-Zn-X system (X = Ni, Mn and Ti) was selected to study the thermo-physical parameters proposed for phase formation in HEAs. The mixing enthalpy (∆H mix ) is expressed as: where n is the number of the elements in the alloy, C i is the atomic ratio of the ith element; C j is the atomic ratio of the jth element, and ∆H mix i−j is the enthalpy of mixing between ith and jth elements calculate by Miedema's model [30]. The atomic size difference (δ) is obtained by: where r i is the atomic radius of ith element and r is the composition weighted average atomic radius. The Pauling electronegativity difference is calculated as: where χ i is the Pauling electronegativity for the ith component and χ is the composition weighted average electronegativity of all the components. The Ω parameter, which can be regarded as the ratio between the enthalpy and the entropy in the liquid state, is given by: where ∆S mix = ∆S conf = −R ∑ n i C i lnC J , T m is the composition-weighted average melting point of the alloy and R is the gas constant (8.314 J·mol −1 ·K −1 ). Subsequently, the theoretical density (ρ t ) of the targeted alloys was calculated by using the assumptions of a rule of mixture: where A i is the atomic weight of the ith element; ρ i is the density of the ith element. In the preliminary study of the thermo-physical parameters for phase formation in HEAs, three different alloy systems were considered. The first system is based on the nominally AlMgSnZnX equimolar composition. The second Al 35 Mg 35 Sn 5 Zn 20 X 5 system is based on the compositional definition of HEAs given by Yeh et al. [4], and optimized to obtain the lowest possible density. Finally, the Al 80 Mg 5 Sn 5 Zn 5 X 5 system was studied, with Al as the main element, and the quantity of the other elements was kept over 5 at.% to be considered as principal element [31]. Hence, a multi-principal element alloys composed of five elements were studied with a composition of the main element of 80 at.%, which is the upper limit for an alloy consisting of five multi-principal elements. In Table 1, the calculated parameters of the defined systems are summarized. As can be observed, the very restrictive values of ∆H mix between −1 and 5 kJ/mole to avoid the formation of ICs in LW-HEAs are very difficult to satisfy, especially in equiatomic alloys. In all the alloying systems, the highest negative value of ∆H mix belongs to equiatomic compositions, followed by the Al 35 Mg 35 Sn 5 Zn 20 X 5 system, and finally, by the Al 80 Mg 5 Sn 5 Zn 5 X 5 system. Therefore, the values of the ∆H mix of the Al 80 Mg 5 Sn 5 Zn 5 X 5 system are closer to satisfying it than their equiatomic counterparts. Following the proposed criteria of δ r below 4.5% and ∆χ equal or below 0.175 to form SS, only Al 80 Mg 5 Sn 5 Zn 5 X 5 system satisfy these criteria. On the other hand, all the calculated values of Ω are very far away of those proposed for the formation of SS in LW-HEAs. The ∆H mix must be near zero and ∆S mix as high as possible to satisfy this requirement, and there are few combinations of low-density elements that satisfy it. Finally, the calculated ρ t of the Al 80 Mg 5 Sn 5 Zn 5 X 5 system also shows the benefit of this system over the other systems.
From the preliminary study of the thermo-physical parameters for phase formation in HEAs, the Al 80 Mg 5 Sn 5 Zn 5 X 5 system was selected. In this system, a microstructure predominantly formed by α-Al phase with other/s ICs is expected to be stabilized. Considering that it is very difficult to keep the ∆H mix values within the limits of SS formation, a system as close as possible to the critical values has been selected, as it is impossible to avoid the formation of ICs attempting to control their volume.

Equilibrium Phase Diagram of the Alloys
Equilibrium phase diagrams of the developed alloys are shown in Figure 1. All the alloys are expected to be composed of a mixture of five different phases at equilibrium. Hence, the alloys are expected to be composed of at least of FCC (A1), Mg 2 Si (C1) and Laves (C14) phases, also known as α-Al, Mg 2 Sn, and η-MgZn 2 phases, respectively. The Mg 2 Sn is a desirable phase to improve the mechanical properties by grain refinement and precipitation strengthening in Mg-based alloys [32,33]. The η-MgZn 2 is typical phase in Al-Zn-Mg-(Cu) alloys that plays an important role in the excellent mechanical properties of the precipitation hardenable 7xxx series.
Al80Mg5Sn5Zn5Ni5 alloy (Figure 1a) was expected to be mainly composed of FCC phase (α-Al), 77 mol % at RT. The major reinforce phase is Al3Ni (D011), which precipitates at 672 °C, and its amount of phase (10 mol %) remains stable below 557 °C. At 521 °C the Mg2Si (Mg2Sn) compound precipitates from liquidus, the amount of this phase is 4 mol % at RT. The T-phase defined by Thermo-Calc, compositionally (Al,Zn)49Mg32, precipitates from FCC matrix at 413 °C. The amount of this phase is 8 mol % at RT. Recently, the benefits of this phase in Al-Zn-Mg-Cu alloy strengthened by T-phase were reported [34]. Finally, the η-MgZn2 phase precipitates from T-(Al,Zn)49Mg32 at 112 °C, and the amount of this phase is below 1 mol %.
The Al80Mg5Sn5Zn5Mn5 alloy equilibrium phase diagram (Figure 1b), was distinctively different from a previously studied alloy containing Ni. Both alloys were expected to be composed of major α-Al matrix reinforced by T-(Al,Zn)49Mg32; Mg2Sn and MgZn2 minor phases, but there is a quantitative and qualitative transformation below 238 °C. At this temperature, the amount of phase of α-Al decreased drastically from 83 mol % to 52 mol % due to the precipitation of Al12Mn from the supersaturated α-Al phase. The amount of Al12Mn phase is 34 mol % at RT. The minor phases show a similar precipitation pattern than alloy containing Ni. The phase diagram of Al80Mg5Sn5Zn5Ti5 alloy (Figure 1c), showed that the liquidus temperature is noticeably higher than previously studied alloys due to the formation of Al3Ti (D022) intermetallic phase at 1108 °C. This phase changes to Al3Ti (LT) at 760 °C and represent the 7 mol % at equilibrium. The remaining phases at RT are the ternary Al 80 Mg 5 Sn 5 Zn 5 Ni 5 alloy (Figure 1a) was expected to be mainly composed of FCC phase (α-Al), 77 mol % at RT. The major reinforce phase is Al 3 Ni (D0 11 ), which precipitates at 672 • C, and its amount of phase (10 mol %) remains stable below 557 • C. At 521 • C the Mg 2 Si (Mg 2 Sn) compound precipitates from liquidus, the amount of this phase is 4 mol % at RT. The T-phase defined by Thermo-Calc, compositionally (Al,Zn) 49 Mg 32 , precipitates from FCC matrix at 413 • C. The amount of this phase is 8 mol % at RT. Recently, the benefits of this phase in Al-Zn-Mg-Cu alloy strengthened by T-phase were reported [34]. Finally, the η-MgZn 2 phase precipitates from T-(Al,Zn) 49 Mg 32 at 112 • C, and the amount of this phase is below 1 mol %.
The Al 80 Mg 5 Sn 5 Zn 5 Mn 5 alloy equilibrium phase diagram (Figure 1b), was distinctively different from a previously studied alloy containing Ni. Both alloys were expected to be composed of major α-Al matrix reinforced by T-(Al,Zn) 49 Mg 32 ; Mg 2 Sn and MgZn 2 minor phases, but there is a quantitative and qualitative transformation below 238 • C. At this temperature, the amount of phase of α-Al decreased drastically from 83 mol % to 52 mol % due to the precipitation of Al 12 Mn from the supersaturated α-Al phase. The amount of Al 12 Mn phase is 34 mol % at RT. The minor phases show a similar precipitation pattern than alloy containing Ni.
The phase diagram of Al 80 Mg 5 Sn 5 Zn 5 Ti 5 alloy (Figure 1c), showed that the liquidus temperature is noticeably higher than previously studied alloys due to the formation of Al 3 Ti (D0 22 ) intermetallic phase at 1108 • C. This phase changes to Al 3 Ti (LT) at 760 • C and represent the 7 mol % at equilibrium. The remaining phases at RT are the ternary Al 18 Mg 3 Ti 2 (17 mol %), and the previously defined Mg 2 Sn (4 mol %) and MgZn 2 (5 mol %). In this case, T-phase precipitates at 416 • C and dissolve at 265 • C in Al 18 Mg 3 Ti 2 phase.

Microstructural Characterization
The overall composition of the alloys was estimated using EDS over large areas. The average values of at least three random measurements are presented in Table 2. The overall values of all alloys deviate slightly from the target compositions. At least one other residual components element was measured by EDS in each alloy. There is little difference between the target and experimental alloy composition, because some elemental vaporization and/or oxidation may have occurred during melting.  Figure 2a,b show the SEM micrograph of as-cast Al 80 Mg 5 Sn 5 Zn 5 Ni 5 alloy at low and high magnification respectively. In Figure 2a, a coarse dendritic structure is observed, where the matrix has a black contrast and the interdendritic space has the brightest contrast. A platelet-like region with grey contrast is also distinguished in the same figure. In Figure 2b, the dendritic matrix was labelled as D and the platelet like morphology as IC. In the interdendritic space, two different regions are distinguished, a white contrast region (ID1) and a grey contrast region (ID2).

Microstructural Characterization
The overall composition of the alloys was estimated using EDS over large areas. The average values of at least three random measurements are presented in Table 2. The overall values of all alloys deviate slightly from the target compositions. At least one other residual components element was measured by EDS in each alloy. There is little difference between the target and experimental alloy composition, because some elemental vaporization and/or oxidation may have occurred during melting.  Figure 2a,b show the SEM micrograph of as-cast Al80Mg5Sn5Zn5Ni5 alloy at low and high magnification respectively. In Figure 2a, a coarse dendritic structure is observed, where the matrix has a black contrast and the interdendritic space has the brightest contrast. A platelet-like region with grey contrast is also distinguished in the same figure. In Figure 2b, the dendritic matrix was labelled as D and the platelet like morphology as IC. In the interdendritic space, two different regions are distinguished, a white contrast region (ID1) and a grey contrast region (ID2).  Figure 2c,d show the SEM micrograph of as-cast Al80Mg5Sn5Zn5Mn5 alloy at low and high magnification respectively. It is evident that the substitution of Ni by Mn lead to a significant microstructural alteration. Thereby, in Figure 2c, the microstructure shows fine net-like globular dendrites, even though the interdendritic space also displays the brightest contrast. In this case, spherical, rhomboidal, and plate-like morphologies with the same contrast are distinguished. These different morphologies were identified as Al6Mn phase independent of the particle morphology [35], and attributed to different dislocation substructures in dependence on the matrix orientation. In Figure 2d, the dendritic matrix was labelled as D, the spherical, rhomboidal, and plate-like morphologies as IC. In this   Figure 2c, the microstructure shows fine net-like globular dendrites, even though the interdendritic space also displays the brightest contrast. In this case, spherical, rhomboidal, and plate-like morphologies with the same contrast are distinguished. These different morphologies were identified as Al 6 Mn phase independent of the particle morphology [35], and attributed to different dislocation substructures in dependence on the matrix orientation. In Figure 2d, the dendritic matrix was labelled as D, the spherical, rhomboidal, and plate-like morphologies as IC. In this alloy, the interdendritic space also was composed of a white contrast region (ID1) and a grey contrast region (ID2).
Figure 2e,f show the SEM micrograph of as-cast Al 80 Mg 5 Sn 5 Zn 5 Ti 5 alloy at low and high magnification respectively. In Figure 2e, the dendrites with a black contrast show pillar-like morphology, and two regions are distinguished in the interdendritic space at low magnification. Finally, a rounded region inside the dendrites with a gray contrast is also distinguished in the same figure. In Figure 2f, the dendrites are labelled as D, and the interdendritic space was composed of a major white contrast region (ID1) and a minor grey contrast region (ID2). Figure 3 shows the SEM + EDS micrograph and the elemental mapping of as-cast Al 80 Mg 5 Sn 5 Zn 5 Ni 5 alloy. The major dendritic phase (D) is mainly composed of Al with a small amount of Mg and Zn. The platelet-like particles (IC) are composed of Al and Ni. The bright white region in the interdendritic space (ID1) is composed of Sn and Mg, and the grey contrast region (ID2) is mainly composed of Zn. alloy, the interdendritic space also was composed of a white contrast region (ID1) and a grey contrast region (ID2). Figure 2e,f show the SEM micrograph of as-cast Al80Mg5Sn5Zn5Ti5 alloy at low and high magnification respectively. In Figure 2e, the dendrites with a black contrast show pillar-like morphology, and two regions are distinguished in the interdendritic space at low magnification. Finally, a rounded region inside the dendrites with a gray contrast is also distinguished in the same figure. In Figure 2f, the dendrites are labelled as D, and the interdendritic space was composed of a major white contrast region (ID1) and a minor grey contrast region (ID2). Figure 3 shows the SEM + EDS micrograph and the elemental mapping of as-cast Al80Mg5Sn5Zn5Ni5 alloy. The major dendritic phase (D) is mainly composed of Al with a small amount of Mg and Zn. The platelet-like particles (IC) are composed of Al and Ni. The bright white region in the interdendritic space (ID1) is composed of Sn and Mg, and the grey contrast region (ID2) is mainly composed of Zn.      alloy, the interdendritic space also was composed of a white contrast region (ID1) and a grey contrast region (ID2). Figure 2e,f show the SEM micrograph of as-cast Al80Mg5Sn5Zn5Ti5 alloy at low and high magnification respectively. In Figure 2e, the dendrites with a black contrast show pillar-like morphology, and two regions are distinguished in the interdendritic space at low magnification. Finally, a rounded region inside the dendrites with a gray contrast is also distinguished in the same figure. In Figure 2f, the dendrites are labelled as D, and the interdendritic space was composed of a major white contrast region (ID1) and a minor grey contrast region (ID2). Figure 3 shows the SEM + EDS micrograph and the elemental mapping of as-cast Al80Mg5Sn5Zn5Ni5 alloy. The major dendritic phase (D) is mainly composed of Al with a small amount of Mg and Zn. The platelet-like particles (IC) are composed of Al and Ni. The bright white region in the interdendritic space (ID1) is composed of Sn and Mg, and the grey contrast region (ID2) is mainly composed of Zn.       are not stabilized in this alloy. The EDS elemental mapping shows similar elemental composition of the dendrites. The interdendritic space also shows the segregation in Sn and Mg-rich (ID1) and Zn-rich (ID2) regions. Finally, Ti was precipitated in the dendritic region in combination with Al. This region labelled as IC presents a globular morphology.  Figure 6 shows the XRD patterns of the experimental alloys in the as-cast condition, the Al80Mg5Sn5Zn5X5 alloys were confirmed to be multiphase microstructures. According to the XRD data, at least the diffraction patterns of α-Al (cubic, Fm-3m), Zn (hexagonal, P63/mmc) and Mg2Sn (cubic, Fm-3m) phases were indexed in all the alloys. In the XRD diagram of Al80Mg5Sn5Zn5Ni5 alloy, the expectedAl3Ni (orthorhombic, Pnma) phase also was indexed, which agrees with the result obtained by SEM + EDS. In the XRD diagram of Al80Mg5Sn5Zn5Mn5 alloy, the peaks of the orthorhombic Al6Mn phase were found. From equilibrium thermodynamic calculations, a phase transformation from Al6Mn to Al12Mn phase was expected at 141°C, but instead of the cubic Al12Mn phase the orthorhombic Al6Mn phase was indexed in the XRD diagram. Thus, the results obtained by XRD patterns are consistent with the thermodynamic equilibrium above 141°C, which is related to the non-equilibrium solidification of the permanent mold casting process. Finally, in the XRD diagram of Al80Mg5Sn5Zn5Ti5 alloy, the predicted Al3Ti (tetragonal, I39/mmm) phase also was indexed. The phases indexed in the diagram showed a better agreement with the equilibrium phases predicted above 421 °C instead of the phases at RT. This is consistent with the non-equilibrium solidification of the casting process.
XRD patterns in Figure 6 lead to some discrepancies between CALPHAD predictions and experimental procedures. The Zn phase was not expected to be in thermodynamic equilibrium of the studied systems. However, the different microstructural characterization techniques employed in this paper, including diffraction peaks of Zn phase (hexagonal, P63/mmc), confirmed the formation of Zn-rich region. Thus, in Figure 7 the semi-quantitative chemical composition of the interdendritic regions defined as ID2 in all the alloys is summarized.  Figure 6 shows the XRD patterns of the experimental alloys in the as-cast condition, the Al 80 Mg 5 Sn 5 Zn 5 X 5 alloys were confirmed to be multiphase microstructures. According to the XRD data, at least the diffraction patterns of α-Al (cubic, Fm-3m), Zn (hexagonal, P6 3 /mmc) and Mg 2 Sn (cubic, Fm-3m) phases were indexed in all the alloys. In the XRD diagram of Al 80 Mg 5 Sn 5 Zn 5 Ni 5 alloy, the expectedAl 3 Ni (orthorhombic, Pnma) phase also was indexed, which agrees with the result obtained by SEM + EDS. In the XRD diagram of Al 80 Mg 5 Sn 5 Zn 5 Mn 5 alloy, the peaks of the orthorhombic Al 6 Mn phase were found. From equilibrium thermodynamic calculations, a phase transformation from Al 6 Mn to Al 12 Mn phase was expected at 141 • C, but instead of the cubic Al 12 Mn phase the orthorhombic Al 6 Mn phase was indexed in the XRD diagram. Thus, the results obtained by XRD patterns are consistent with the thermodynamic equilibrium above 141 • C, which is related to the non-equilibrium solidification of the permanent mold casting process. Finally, in the XRD diagram of Al 80 Mg 5 Sn 5 Zn 5 Ti 5 alloy, the predicted Al 3 Ti (tetragonal, I39/mmm) phase also was indexed. The phases indexed in the diagram showed a better agreement with the equilibrium phases predicted above 421 • C instead of the phases at RT. This is consistent with the non-equilibrium solidification of the casting process.  Figure 6 shows the XRD patterns of the experimental alloys in the as-cast condition, the Al80Mg5Sn5Zn5X5 alloys were confirmed to be multiphase microstructures. According to the XRD data, at least the diffraction patterns of α-Al (cubic, Fm-3m), Zn (hexagonal, P63/mmc) and Mg2Sn (cubic, Fm-3m) phases were indexed in all the alloys. In the XRD diagram of Al80Mg5Sn5Zn5Ni5 alloy, the expectedAl3Ni (orthorhombic, Pnma) phase also was indexed, which agrees with the result obtained by SEM + EDS. In the XRD diagram of Al80Mg5Sn5Zn5Mn5 alloy, the peaks of the orthorhombic Al6Mn phase were found. From equilibrium thermodynamic calculations, a phase transformation from Al6Mn to Al12Mn phase was expected at 141°C, but instead of the cubic Al12Mn phase the orthorhombic Al6Mn phase was indexed in the XRD diagram. Thus, the results obtained by XRD patterns are consistent with the thermodynamic equilibrium above 141°C, which is related to the non-equilibrium solidification of the permanent mold casting process. Finally, in the XRD diagram of Al80Mg5Sn5Zn5Ti5 alloy, the predicted Al3Ti (tetragonal, I39/mmm) phase also was indexed. The phases indexed in the diagram showed a better agreement with the equilibrium phases predicted above 421 °C instead of the phases at RT. This is consistent with the non-equilibrium solidification of the casting process.
XRD patterns in Figure 6 lead to some discrepancies between CALPHAD predictions and experimental procedures. The Zn phase was not expected to be in thermodynamic equilibrium of the studied systems. However, the different microstructural characterization techniques employed in this paper, including diffraction peaks of Zn phase (hexagonal, P63/mmc), confirmed the formation of Zn-rich region. Thus, in Figure 7 the semi-quantitative chemical composition of the interdendritic regions defined as ID2 in all the alloys is summarized.  XRD patterns in Figure 6 lead to some discrepancies between CALPHAD predictions and experimental procedures. The Zn phase was not expected to be in thermodynamic equilibrium of the studied systems. However, the different microstructural characterization techniques employed in this paper, including diffraction peaks of Zn phase (hexagonal, P6 3 /mmc), confirmed the formation of Zn-rich region. Thus, in Figure 7 the semiquantitative chemical composition of the interdendritic regions defined as ID2 in all the alloys is summarized.
(c) As can be seen in Figure 7, ID2 is mainly composed of Zn with a small amount of O and Al. Therefore, the discrepancies between CALPHAD and experimental results may be due to the high melting temperatures reached during the casting process to dissolve Ni, Mn, and Ti. This promoted a partial oxidation of Zn at high temperatures stabilizing Zn-O phase instead of the thermodynamically stable phase (T-phase) predicted by Thermo-Calc.

Mechanical Properties
To evaluate the mechanical properties of the as-cast alloys, Vickers hardness and compressive test were conducted at RT. The compressive engineering stress-strain curves of the alloys are plotted in Figure 8. The highest strength is obtained by Ti addition, while the alloys containing Ni and Mn showed similar compressive mechanical properties. The curve corresponding to Al80Mg5Sn5Zn5Ti5 alloy (the blue line denoted as c), shows the highest hardening rate.
In general terms of mechanical properties, the as-cast alloys showed a good balance of mechanical properties, combining high strength and hardness with moderate deformation. The compressive mechanical properties showed that the Al80Mg5Sn5Zn5Ti5 alloy exhibited the highest strength and hardening rate. This is correlated with the formation of globular Al3Ti phase, instead the formation of the acicular phases in the alloys containing Ni and Mn. There is also a good correlation between hardness and elongation values in the alloys studied, with a decrease in elongation values when the hardness of the alloy is increased, as is usual in metallic alloys.
The experimental density, hardness, and compressive mechanical properties with the standard deviations of the developed alloys are summarized in Table 3. The calculated densities in Table 1 showed good agreement with the experimental results. The test results revealed that all the alloys exhibited high compressive offset yield strength (σC0.2) and ultimate compressive strength (σCU) exceeding 317 MPa and 368 MPa, respectively. The As can be seen in Figure 7, ID2 is mainly composed of Zn with a small amount of O and Al. Therefore, the discrepancies between CALPHAD and experimental results may be due to the high melting temperatures reached during the casting process to dissolve Ni, Mn, and Ti. This promoted a partial oxidation of Zn at high temperatures stabilizing Zn-O phase instead of the thermodynamically stable phase (T-phase) predicted by Thermo-Calc.

Mechanical Properties
To evaluate the mechanical properties of the as-cast alloys, Vickers hardness and compressive test were conducted at RT. The compressive engineering stress-strain curves of the alloys are plotted in Figure 8. The highest strength is obtained by Ti addition, while the alloys containing Ni and Mn showed similar compressive mechanical properties. The curve corresponding to Al 80 Mg 5 Sn 5 Zn 5 Ti 5 alloy (the blue line denoted as c), shows the highest hardening rate.    In general terms of mechanical properties, the as-cast alloys showed a good balance of mechanical properties, combining high strength and hardness with moderate deformation.
The compressive mechanical properties showed that the Al 80 Mg 5 Sn 5 Zn 5 Ti 5 alloy exhibited the highest strength and hardening rate. This is correlated with the formation of globular Al 3 Ti phase, instead the formation of the acicular phases in the alloys containing Ni and Mn. There is also a good correlation between hardness and elongation values in the alloys studied, with a decrease in elongation values when the hardness of the alloy is increased, as is usual in metallic alloys.
The experimental density, hardness, and compressive mechanical properties with the standard deviations of the developed alloys are summarized in Table 3. The calculated densities in Table 1 showed good agreement with the experimental results. The test results revealed that all the alloys exhibited high compressive offset yield strength (σ C0.2 ) and ultimate compressive strength (σ CU ) exceeding 317 MPa and 368 MPa, respectively. The hardness of the alloys demonstrated high values exceeding those of Al-based alloys, and the Al 80 Mg 5 Sn 5 Zn 5 Ti 5 alloy showed the highest hardness averaging 159 HV.

Conclusions
In this work, three new lightweight CCAs were successfully developed with a density below 3.33 g/cm 3 and manufactured by a standard permanent mold gravity die casting process. From the equilibrium phase diagrams and thermo-physical parameters for phase formation in HEAs, all the alloys were expected to be composed of a mixture of SS + ICs microstructures. This is consistent with the experimental results, since all the as-cast microstructures are composed at least of a major α-Al SS enriched with Zn and Mg, reinforced with a cubic Mg 2 Sn intermetallic phase, and the unexpected hexagonal Zn phase. The discrepancies between CALPHAD and experimental results in the formation of this phase were attributed to partial oxidation of Zn during the manufacturing process. The microstructural evolution as a function of the alloying element revealed that, depending on the variable element, Al 3 Ni, Al 6 Mn, or Al 3 Ti phases were also stabilized.
The variable element also affected the grain size of the alloys, with the finest microstructure being that of Mn-containing alloy, followed by Ni-containing alloy, and finally, Ti-containing alloy. This can be related to the precipitation temperatures of the Al 3 Ni, Al 6 Mn and Al 3 Ti phases. The equilibrium phase diagrams of the alloys revealed that Al 3 Ti (precipitates at 1108 • C) is the compound with the highest precipitation temperature. The formation and growth of this phase at high temperatures has not allowed it to act as grain refiner, as is common in commercial aluminum alloys. In contrast, Al 3 Ni (precipitates at 672 • C) and Al 6 Mn (precipitates at 749 • C), with a precipitation temperature closer to the formation of α-Al matrix (precipitation temperature around 610 • C), are more efficient as grain refiners. This can be clearly seen in the equilibrium phase diagrams. When the α-Al matrix starts the solidification in Al 80 Mg 5 Sn 5 Zn 5 Ti 5 alloy, the phase fraction of Al 3 Ti phase reached the maximum volume, so there is an insufficient quantity of Ti in the liquid to promote the nucleation of Al-Ti compounds. On the other hand, when the α-Al matrix starts the precipitation in Al 80 Mg 5 Sn 5 Zn 5 Ni 5 and Al 80 Mg 5 Sn 5 Zn 5 Mn 5 alloys, the amount of phase of Al-Ni and Al-Mn phases are growing. This means that some Ni and Mn remain in the liquid phase, promoting the grain refinement of the alloy.
The alloys containing Ni and Mn showed the formation of acicular phases. On the other hand, the addition of Ti led to the formation of a globular phase. This is reflected in the better mechanical properties of Al 80 Mg 5 Sn 5 Zn 5 Ti 5 alloy over alloys containing Ni and Mn. Instead, all the as-cast alloys showed a good balance of mechanical properties, combining high strength and hardness with a deformation up to 15%.
In future studies, some adjustment must be made in the manufacturing process in order to ensure that the final composition is closer to the nominal one, and the effect of different heat treatments in the mechanical properties at RT and high temperatures should be considered. Based on the obtained results, it could be interesting to study the tribological properties of the developed alloys or related systems.