Effects of Mn and Cu Additions on Solidification Microstructure and High-Temperature Strength of Cast Al–Fe Binary Alloy

Abstract

In order to investigate the effects of Mn and Cu additions on solidification microstructure and high-temperature strength of cast Al–Fe alloys, we have fabricated various Al–Fe-based alloys with compositions of Al–1%Fe, Al–1%Fe–1%Mn, Al–1%Fe–1%Cu, and Al–1%Fe–1%Cu–1%Mn (mol%) solidified at different cooling rates (0.3 K·s⁻¹ and 145 K·s⁻¹). In the Al–1%Fe binary alloy, the coarsened θ-Al₁₃Fe₄ phase with a needle-shaped morphology was often observed in the furnace-cooled sample (0.3 K·s⁻¹), whereas the cast sample (145 K·s⁻¹) exhibited several elongated α phases surrounded by fine α/Al₆Fe eutectic microstructure. Such a solidification microstructure was observed in the cast Al–1%Fe–1%Cu alloy, whereas the Al₂₃CuFe₄ phase was locally formed in the finally solidified zone in the furnace-cooled sample. In the Al–1%Fe–1%Mn alloy, the Al₆(Fe, Mn) phase was formed regardless of the cooling rate. Finer α/Al₆(Fe, Mn) two-phase eutectic microstructure was almost entirely occupied in the cast sample. The fine eutectic microstructure was observed in the cast Al–1%Fe–1%Cu–1%Mn alloy as well. Compression tests for cast alloy specimens revealed that the Al–1%Fe–1%Cu–1%Mn alloy exhibited the highest strength level among the studied alloy specimens, indicating the combined addition of Mn and Cu elements could be effective in improving the high-temperature strength of the cast Al–Fe alloys.

1. Introduction

In recent years, engines with turbochargers have been applied to downsize engines in response to the demand for stricter emission regulations and lower fuel consumption of automobiles. The combustion efficiency of engines can be improved by increasing the turbochargers’ compression ratio, whereby it is required to improve the high-temperature strength of materials for the impeller components.^1,2)

The 2618 alloy with a nominal composition of Al–2.3Cu–1.2Mg–1.1Fe–1.0Ni (mass%)^3,4) is generally used for impeller components in inlet-side compressors, but its application temperature is limited because the strength of the 2618 alloy is significantly reduced above 150°C.⁵⁾ The cause of this strength reduction is that the metastable phase (S′ phase), which plays a role in strengthening the 2618 alloy, becomes coarsened through the phase transformation to the stable phase (S-Al₂CuMg phase⁶⁾) in holding at high temperatures. To solve this problem, we have proposed aluminum (Al) eutectic alloys reinforced with thermodynamically stable intermetallic phases (in equilibrium with the α-Al phase) with high volume fractions.^7–9) The proposed alloys exhibited more than twice the strength of conventional alloys at high temperatures above 250°C.^7,8) However, the eutectic alloys containing more than 50% by volume of intermetallic phases, exhibited poor ductility at room temperature.

Many studies have been conducted on heat-resistant Al alloys with added transition metal elements such as Fe.^10–12) Sintered specimens using rapidly solidified alloy powders were reinforced by a fine intermetallic phase and exhibited excellent high-temperature strength,¹³⁾ whereas they often showed poor ductility. In the Al–Fe binary alloys, it is generally known that coarsened θ-Al₁₃Fe₄ phase (monoclinic mC102, Fig. 1(a))¹⁴⁾ is formed in the α(fcc) matrix. The θ-phase has a low fracture toughness,¹⁵⁾ which significantly reduces the ductility of the alloy. To suppress the formation of the brittle θ-phase, it would be effective to rapidly solidify the alloy ingots using various casting processes. The rapid solidification in the Al–Fe binary alloys enhances the formation of a fine metastable Al₆Fe phase (oC28; Fig. 1(b)¹⁶⁾). In fact, Al–Fe binary alloys additive-manufactured by laser powder bed fusion (one of the representative metal additive manufacturing processes) contained numerous fine particles of the metastable Al₆Fe phase.^17–19) They exhibited high strength with an adequate ductility.^20,21) The fine morphologies of the Al₆Fe phase were relatively stable at elevated temperatures as high as 300°C and significantly contributed to high-temperature strength.²²⁾ This study attempted to stabilize the metastable Al₆Fe phase and increase its volume fraction by adding tertiary elements (Mn, Cu) into the Al–Fe binary alloy.

Fig. 1

Schematics showing unit cells of crystal structures of (a) θ-Al₁₃Fe₄, (b) Al₆Fe, (c) Al₆Mn, and (d) T₁-Al₂₃CuFe₄ phases.

In the Al–Mn binary system, Al₆Mn (oC28; Fig. 1(c)),²³⁾ which has the same crystal structure as Al₆Fe, is in equilibrium with the α-phase. In the Al–Fe–Mn ternary system, Fe atoms would be substituted into the Mn sublattice of the Al₆Mn phase, resulting in stabilizing the Al₆(Fe, Mn) phase.^24,25) The Al₂₃CuFe₄ (oC28, Fig. 1(d))²⁶⁾ phase has the same orthorhombic crystal structure as Al₆Fe and shows a similar atomic arrangement. Therefore, added Cu element is expected to preferentially distribute to the metastable Al₆Fe phase, which may suppress the formation of the coarsened θ-phase. However, few reports are on the solidification microstructure and mechanical properties of Al–Fe–Mn or Al–Fe–Cu ternary alloys. Herein, to understand the effects of Mn and/or Cu additions on the formation of (Al,M1)₆M2 (M1:Cu, M2:Fe, Mn) in the solidification microstructures of Al–Fe alloys, we fabricated Al–Fe–Mn or Al–Fe–Cu ternary alloys, and Al–Fe–Cu–Mn quaternary alloys solidified at different cooling rates. The solidification microstructures and high-temperature strength were systematically investigated.

2. Experimental Procedure

The studied alloy compositions were Al–1%Fe (corresponding to a near-eutectic composition in the binary system), Al–1%Fe–1%Mn (Mn-added), Al–1%Fe–1%Cu (Cu-added), and Al–1%Fe–1%Cu–1%Mn (Mn and Cu added). All alloy compositions are given in mole percent (mol%) unless otherwise noted. Figure 2 shows the Al–Fe binary phase diagram and the Al–Fe–Mn and Al–Fe–Cu ternary liquidus projections prepared by thermodynamic calculations using the Al-based multi-component database PanAluminum²⁷⁾ (CompuTherm LLC). In the calculated equilibrium phase diagrams, the primary solidified phase for the Al–1Fe binary and Al–1Fe–1Mn, Al–1Fe–1Cu ternary compositions were predicted to be the θ-Al₁₃Fe₄ phase.

Fig. 2

Studied alloy compositions plotted on (a) Al–Fe binary system and the calculated liquidus projection of (b) Al–Fe–Mn ternary and (c) Al–Fe–Cu ternary systems.

In this study, Al (purity 99.99%), Fe (purity 99.9%), Mn (purity 99.9%), and Cu (purity 99.9%) were used as raw materials. The samples were melted by high-frequency heating in an Ar atmosphere and maintained in the temperature range of 800–900°C for 1800 s. After that, the alloy samples were cooled by switching off the power source to the furnace (cooling rate: 0.3 K·s⁻¹) to produce a solidified alloy sample (hereafter referred to as “furnace-cooled” samples). The rapidly solidified alloy samples (cooling rate: 145 K·s⁻¹) were prepared by casting into an iron mold (hereafter referred to as “cast” samples). The cooling rates were experimentally measured by inserting a W thermocouple into the molten metal.⁷⁾

The fabricated rod-shaped ingots were cut into pieces about 5 mm thick, embedded in resin, and mechanically polished using SiC paper to prepare the sample for observation. A scanning electron microscopy (SEM; JEOL JSM-6610A) was used to observe the microstructure at an acceleration voltage of 15 kV. X-ray diffraction (XRD) measurements were conducted to obtain diffraction peaks, which were compared with diffraction results calculated from the crystal structures for phase identification. Samples for electron-backscatter diffraction (EBSD) analysis were prepared by ion polishing using a cross-sectional polisher at 5 V. A field emission SEM (FE-SEM, JEOL JSM-7001FA) equipped with an EBSD system was used to identify the constituent phases. Elemental analysis was also performed using energy dispersive X-ray spectrometry (EDS).

The samples fabricated by mold-cast were wet-polished and prepared into flakes with a thickness of about 0.1 mm for TEM observation. These specimens were then polished with argon at 6.0 kV using an ion slicer (JEOL Ion Slicer TM, EM-09100IS) and finished for 300 s at 2 kV. Microstructural observations were performed using a TEM (JEOL JEM 2100F/HK) at 200 kV to obtain bright-field images, high-angle scattering dark-field images (HAADF-STEM) using a scanning transmission electron microscope (STEM), and elemental mapping using STEM-EDX.

The microhardness (HV) values of the specimens were measured by pressing a Vickers indenter into the surface of the specimen with a constant load of 1.96 N. The alloys fabricated by mold-cast were cut into the specimens with a dimension of 4 mm × 4 mm × 8 mm for high-temperature compression tests. These specimens were compressed in a temperature range from room temperature to 300°C using a Shimadzu Autograph AG-2000A and a thermostatic chamber (SHIMADZU THERMOSTATIC CHAMBER TCH-382SP). The initial strain rate was 2.0 × 10⁻³ s⁻¹.

3. Results and Discussion

3.1 Solidification microstructure

Figure 3 displays low-magnification back-scattered electron images (BEIs) showing the prepared samples with different compositions. The microstructural morphologies of the furnace-cooled alloys are shown in Fig. 3(a)–(d). In the furnace-cooled Al–1Fe (Fig. 3(a)) and Al–1Fe–1Cu (Fig. 3(c)) alloys, needle-like intermetallic phases were often observed. In the furnace-cooled Al–1Fe–1Mn alloy (Fig. 3(b)), polyhedral intermetallic phases (corresponding to the primary solidification phase) with a size of approximately 100 µm were locally observed. Such polyhedral intermetallic phases were observed more often in the Al–1Fe–1Cu–1Mn alloy (Fig. 3(d)). These microstructural morphologies became finer in the samples solidified at a higher cooling rate of 145 K·s⁻¹. Figure 3(e)–(h) display the macroscopic morphologies of microstructure in the cast alloys. No coarse primary phase with sizes larger than 100 µm was found in any of the alloys.

Fig. 3

Low-magnification BEIs showing solidification microstructures of (a), (e) Al–1Fe, (b), (f) Al–1Fe–1Mn, (c), (g) Al–1Fe–1Cu and (d), (h) Al–1Fe–1Cu–1Mn alloys prepared by (a)–(d) furnace cooling or (e)–(h) mold-casting.

Figure 4 displays high-magnification back-scattered electron images (BEIs) showing the prepared alloy samples. In the furnace-cooled Al–1Fe alloy (Fig. 4(a)), many needle-like intermetallic phases of several 10 µm in length were dispersed in the α-matrix. In the Al–1Fe–1Mn alloy, needle-like and clumped intermetallic phases were observed (Fig. 4(b)). In addition, cracks were observed inside many of the coarse intermetallic phases. These features were observed in the Al–1Fe–1Cu alloy as well (Fig. 4(c)), whereas fine eutectic microstructures (presumably corresponding to the finally solidified zone) were locally observed. Such a final solidification zone was observed in the Al–1Fe–1Cu–1Mn alloy (Fig. 4(d)) as well, and granular intermetallic phases were also observed. The Al–1Fe alloy fabricated by mold-casting (cooling rate: 145 K·s⁻¹) (Fig. 4(e)) has a number of elongated α phases of several 10 µm in size surrounded by fine eutectic structures. A similar microstructure was observed in the Al–1Fe–1Cu alloy (Fig. 4(g)), where slightly coarsened intermetallic phases appeared localized at the interface between the elongated α phase and the eutectic zone. In the Al–1Fe–1Mn alloy (Fig. 4(f)), the fine eutectic structure was almost fully occupied in the solidification microstructure. Similar microstructural morphology was observed in the Al–1Fe–1Cu–1Mn alloy (Fig. 4(h)).

Fig. 4

High-magnification BEIs showing solidification microstructures of (a), (e) Al–1Fe, (b), (f) Al–1Fe–1Mn, (c), (g) Al–1Fe–1Cu and (d), (h) Al–1Fe–1Cu–1Mn alloys prepared by (a)–(d) furnace cooling or (e)–(h) mold-casting.

3.2 Phase identification

Figure 5 shows the XRD profiles of the experimental alloy samples prepared by furnace-cooling (0.3 K·s⁻¹) and mold-cast (145 K·s⁻¹) processes. Diffractions derived from α and θ-Al₁₃Fe₄ phases were detected in the furnace-cooled Al–1Fe alloy. Therefore, the needle-like Al–Fe intermetallic phase dispersed in the α matrix (Fig. 4(a)) could correspond to the α/θ two-phase eutectic microstructure. In the Al–1Fe–1Mn alloys, there were diffraction peaks derived from α and Al₆Mn phases (hereafter referred to as Al₆(Fe, Mn) phase, because the crystal structure is the same as that of Al₆Fe, as shown in Fig. 1). The diffraction peaks around 22∼24° derived from the θ-phase were not observed in the cast Al–1Fe–1Mn alloy, which is indicative of the microstructures composed of α and Al₆(Fe, Mn) phases in the Al–1Fe–1Mn alloy samples (Fig. 4(b), (f)). In the furnace-cooled Al–1Fe–1Cu alloy, diffractions derived from α, θ, and Al₂₃CuFe₄ phases were detected. There was no diffraction peak derived from the θ-phase in the cast sample. These results indicate the presence of α and Al₂₃CuFe₄ phases. The Al–1Fe–1Cu–1Mn alloy samples exhibited similar diffraction profiles to the Al–1Fe–1Mn alloy samples. This result suggests that not only the added Mn elements but also the Cu elements would contribute to the formation of the (Al,M1)₆M2 (M1:Cu, M2:Fe, Mn) phase in the quaternary alloy.

Fig. 5

X-ray diffraction profiles of Al–1Fe, Al–1Fe–1Mn, Al–1Fe–1Cu and Al–1Fe–1Cu–1Mn alloys prepared by (a) furnace cooling process or (b) mold-casting process.

Figure 6 shows EBSD patterns obtained from local areas in furnace-cooled Al–1Fe, Al–1Fe–1Mn and Al–1Fe–1Cu alloy samples. EBSD patterns acquired from the matrix phase (point A) and bright contrast (point B) in the SEM image (Fig. 6(a)) indicate the presence of α and θ phases in the Al–1Fe alloy. The confidence index (CI) values²⁸⁾ of the EBSD analysis for points A and B are 0.320 and 0.143, respectively. These values indicate sufficient reliability of the present analysis. The identified phases were in good agreement with the result of the XRD measurement (Fig. 5(a)). In the furnace-cooled Al–1Fe–1Mn alloy, EBSD patterns captured from the intermetallic phases with different morphologies (points A and B in Fig. 6(f)) were identified as patterns derived from the Al₆Mn phase (oC28; Fig. 1(c)).²³⁾ The furnace-cooled Al–1Fe–1Mn alloy was composed of the Al₆(Fe, Mn) phase in the α matrix, which corresponded well to the result of the XRD profile (Fig. 5(a)). In the SEM image of the furnace-cooled Al–1Fe–1Cu alloy (Fig. 6(k)), point A is located inside the fine eutectic structure, and point B is in the granular phase. The EBSD patterns obtained from points A and B were identified as the Al₂₃CuFe₄ and θ phases, respectively (Fig. 6(l)–(o)). Therefore, it was considered that the furnace-cooled Al–1Fe–1Cu alloy would have the θ-phase as the primary solidification phase and α/Al₂₃CuFe₄ eutectic regions in the α matrix.

Fig. 6

(a), (f), (k) SEI images showing the microstructure of furnace-cooled Al–1Fe, Al–1Fe–1Mn, Al–1Fe–1Cu alloy, (b), (g), (l) EBSD patterns obtained from location A, and (c), (h), (m) analyzed EBSD patterns, (d), (i), (n) EBSD patterns obtained from location, and (e), (j), (o) analyzed EBSD patterns.

Figure 7 shows EDS element maps and EBSD patterns obtained from local areas in the furnace-cooled Al–1Fe–1Cu–1Mn alloy. Point A in Fig. 7(f) is located on the Fe and Mn-rich phase (Fig. 7(c), (e)), and point B corresponds to the Cu-enriched region finally solidified (Fig. 7(d)). The analyzed EBSD pattern obtained from points A and B presented the crystal structures of Al₆Mn phase (oC28) and Al₇CuFe₂ phase (tP40),²⁹⁾ representatively. No clear diffraction peak derived from Al₇CuFe₂ phase was found in the XRD profile (Fig. 5(a)), which may be due to the local presence of finally solidified regions containing the Al₇CuFe₂ phase.

Fig. 7

(a) SEI image and (b)–(e) corresponding EDS element maps of Al–1Fe–1Cu–1Mn furnace-cooled alloy. (f) SEI image showing the microstructure of Al–1Fe–1Cu–1Mn furnace-cooled alloy, (g) EBSD patterns obtained from location A in (f), and (h) analyzed EBSD patterns using the crystal structure of the Al₆Mn phase, (i) EBSD patterns obtained from location B in (f), and (j) analyzed EBSD patterns using the crystal structure of the Al₇CuFe₂ phase.

Figure 8 presents STEM images showing eutectic microstructures in the Al–1Fe, Al–1Fe–1Mn, and Al–1Fe–1Cu alloy prepared by the mold-cast process and the results of EDS composition analysis. The intermetallic phase was identified as the Al₆Fe phase by calculating the atomic ratio of Al and Fe using point analysis in the Fe-rich region (Fig. 8(c)). The eutectic microstructure included a rod-shaped Al₆Fe phase with a diameter of about 0.5 µm (Fig. 8(a)) in the cast Al–1Fe alloy. Similar microstructural morphologies were observed in the Al–1Fe–1Mn alloy (Fig. 8(d)) and Al–1Fe–1Cu alloy (Fig. 8(h)). The added elements (Mn and Cu) were enriched in intermetallic phases (Fig. 8(g), (k)). Point composition analyses for the Fe and Mn-rich region revealed the Al₆(Fe, Mn) phase formed in the cast Al–1Fe–1Mn alloy (Fig. 8(f), (g)). The atomic ratio in the Al₆(Fe, Mn) phase was measured as approximately Al:Fe:Mn = 6:0.8:0.1. In addition, the composition analyses for the Fe and Cu-rich region suggested the Al₂₃CuFe₄ phase formed in the cast Al–1Fe–1Cu alloy (Fig. 8(j), (k)).

Fig. 8

(a) TEM image and (b), (c) corresponding EDS element maps of the mold-cast Al–1Fe alloy. (d) TEM image and (e)–(g) corresponding EDS element maps of the mold-cast Al–1Fe–1Mn alloy. (h) TEM image and (i)–(k) corresponding EDS element maps of the mold-cast Al–1Fe–1Cu alloy.

Figure 9 depicts a STEM image showing the eutectic structure in the cast Al–1Fe–1Cu–1Mn alloy and the results of EDS composition analysis. The quaternary alloy exhibited finely distributed rod-shaped intermetallic phases in the α matrix (Fig. 9(a)) as well as the Al–Fe binary alloy (Fig. 8(a)). The added Cu and Mn elements were enriched in the rod-shape intermetallic phases (Fig. 9(c), (d), (e)), which may contribute to the stabilization of the (Al,M1)₆M2 phase (M1:Cu, M2:Fe, Mn).

Fig. 9

(a) TEM image and (b)–(d) corresponding EDS element maps of the mold-cast Al–1Fe–1Cu–1Mn alloy.

3.3 Mechanical properties

Figure 10 shows the Vickers hardness of each fabricated alloy as a function of the cooling rate. The present hardness measurements were performed inside eutectic microstructures (Fig. 3(a)–(d)) rather than on the primary solidified phase. The Vickers hardness of the furnace-cooled alloys (0.3 K·s⁻¹) was 53 HV for the Al–1Fe–1Cu alloy, 52 HV for the Al–1Fe–1Cu–1Mn alloy, and 34 HV for the Al–1Fe and Al–1Fe–1Mn alloys. The Vickers hardness of the cast alloys (145 K·s⁻¹) increased for all alloys, and the Al–1Fe–1Cu–1Mn quaternary alloy showed the highest hardness (82 HV). This consistent increase in hardness can be due to the fine morphology of the eutectic microstructures.

Fig. 10

Variation of Vickers hardness of Al–1Fe, Al–1Fe–1Mn, Al–1Fe–1Cu and Al–1Fe–1Cu–1Mn alloys as a function of cooling rate controlled by different cast processes.

Figure 11(a) shows the nominal stress-strain curves obtained from compression tests at room temperature for each cast alloy specimen (145 K·s⁻¹). All alloy specimens showed relatively uniform deformation behavior and did not fracture during the tests. The Al–1Fe–1Cu and Al–1Fe–1Cu–1Mn alloys showed a work-hardening behavior, and their strength increased with progressing the compression deformation. The difference in compressive strengths among experimental alloys exhibited the same trend of the variation in Vickers hardness values (Fig. 10).

Fig. 11

Nominal stress-plastic strain curves of mold-cast Al–1Fe alloy, Al–1Fe–1Mn alloy, Al–1Fe–1Cu and Al–1Fe–1Cu–1Mn alloy specimens obtained from compression tests at an initial strain-rate of 2.0 × 10⁻⁴ s⁻¹; (a) at room temperature and (b) at 200°C.

Figure 11(b) shows the nominal stress-strain curves obtained from compression tests at 200°C for each cast alloy. The strength of all alloys decreased at 200°C, but the relative difference in strength was the same as in the room temperature test. Also, at room temperature, the Al–1Fe–1Cu and Al–1Fe–1Cu–1Mn alloys showed a work-hardening behavior, presumably due to the high solute Cu content in the α matrix. EDS composition analyses revealed 0.2 mol% Mn (in solution) in the α matrix in the Al–1Fe–1Mn alloy, while 0.7 mol% Cu was detected in the cast Al–1Fe–1Cu alloy. These results would reasonably agree with a higher solubility limit of Cu element in the α phase than that of Mn element.^30,31)

Figure 12 presents the 0.2% proof stress (measured by the compression tests) of each cast alloy specimen at various temperatures, together with the result of the conventional cast-type Al alloy (with a nominal composition of Al–12%Si–1%Cu–1%Mg–1%Ni (mass%)).³²⁾ The 0.2% proof stress of the cast Al–1Fe alloy at room temperature is about 90 MPa and decreases with increasing temperature. The 0.2% proof stress at 300°C was about 50 MPa. The addition of 1 mol% Mn and Cu increases the 0.2% proof stress to 120∼130 MPa at room temperature. The combined addition of Mn and Cu increased the 0.2% proof stress to approximately 150 MPa, and the reduced strength level by increasing test temperature was less pronounced in the Al–1Fe–1Mn–1Cu quaternary alloy. The 0.2% proof stress of the AC8A alloy was reduced by 80% from room temperature to 300°C, whereas the reduction was approximately 45% in the quaternary alloy. This might be attributed to the relatively high microstructural stability (eutectic microstructure of α and (Al,M1)₆M2 (M1:Cu, M2:Fe, Mn) phases) at high temperatures. Further analyses of the (Al,M1)₆M2 phase substituted with Mn and Cu, and thermal stability of the (Al,M1)₆M2 phase would be needed to understand the improvement in high-temperature strength of Al–Fe alloys with the α/(Al,M1)₆M2 eutectic structure.

Fig. 12

Variation of 0.2% proof stress of cast Al–1Fe, Al–1Fe–1Mn, Al–1Fe–1Cu and Al–1Fe–1Cu–1Mn alloys as a function of test temperature, together with the result of AC8A alloy.³²⁾

4. Summary

The present study was set to understand the effects of Mn and Cu additions on solidification microstructure and the high-temperature strength of cast Al–Fe alloys. We have prepared various Al–Fe-based alloy ingots with compositions of Al–1%Fe, Al–1%Fe–1%Mn, Al–1%Fe–1%Cu, and Al–1%Fe–1%Cu–1%Mn (mol%) solidified at different cooling rates (0.3 K·s⁻¹ and 145 K·s⁻¹). The solidification microstructures and high-temperature strength were systematically investigated. The key results are summarized as follows;

1. (1)    In the furnace-cooled Al–1%Fe alloy (0.3 K·s⁻¹), the coarsened θ-Al₁₃Fe₄ stable phase was formed in the α(fcc) matrix. In the sample prepared by mold-casting (145 K·s⁻¹), fine metastable Al₆Fe phases were included in the eutectic microstructure. In the Al–1%Fe–1%Mn alloy, the Al₆(Fe, Mn) phase appeared regardless of the cooling rate. A higher cooling rate enhances the refinement of the Al₆(Fe, Mn) phase in the eutectic microstructure. In the furnace-cooled Al–1%Fe–1%Cu alloy, the Al₂₃CuFe₄ phase was formed in the finally solidified zone. The cast Al–1%Fe–1%Cu alloy showed the same microstructural morphology as the cast Al–1%Fe binary alloy, indicating that the added Cu element partitioned into the (Al,Cu)₆Fe phase (corresponding to the Al₂₃CuFe₄ phase in the Al–Fe–Cu ternary system²⁶⁾). The Mn and Cu combined addition enhanced the formation of the (Al,M1)₆M2 phase (M1:Cu, M2:Fe, Mn) with an orthorhombic structure.
2. (2)    The compression tests of four alloy specimens fabricated by casting revealed that the Al–1Fe–1Cu–1Mn quaternary alloy exhibited the highest 0.2% proof stress among experimental alloys at elevated temperatures. This might be attributed to relatively high microstructural stability (α/(Al,M1)₆M2 eutectic structure) at high temperatures.

Acknowledgments

The supports of JSPS KAKENHI (Grant Numbers 20H02462), Japan Aluminum Association and the Light Metal Educational Foundation, Inc. are gratefully acknowledged.

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