Effect of Multi-Phase Composite Structure on the Mechanical Properties of AlxFe1.5CoNiC0.12 High Entropy Alloys

Abstract

Single-phase high entropy alloys (HEAs) exhibit limited mechanical properties while dual-phase and multi-phase HEAs offer better strength, toughness, and stability. In this paper, the as-cast Al_xFe_1.5CoNiC_0.12 HEAs with triple-phase dendritic composite structure is studied, and the influence of the composite structure on the mechanical properties is discussed. The interdendrite (ID) of this structure is composed of a uniformly distributed high-density ordered face-centered cubic structure (L1₂) precipitate phase and face-centered cubic (FCC) matrix, while the dendrite (DR) consists of an ordered body-centered cubic (B2) single-phase. The high density L1₂ precipitate phase leads to a higher hardness in the FCC+L1₂ dual-phase region compared to the B2 single-phase region. The decrease in Al content can greatly improve mechanical performance. The improvement was attributed to the higher volume fraction of the ID and the smaller particle size of the precipitates. The L1₂ phase nano-precipitates exhibit minimal lattice mismatch with the FCC matrix, thereby significantly enhancing the stability of the alloy at the nanoscale. This stability is reflected in the fracture morphology. Modulating the triple-phase dendritic composite structure effectively improves the mechanical properties of the alloy.

1. Introduction

The exploration of high entropy alloys (HEAs) with four or more principal elements has become hotspots in metal materials research, driven by their unique structures and application potentials [1,2,3,4,5]. Interestingly, despite containing various major elements, high configurational entropy in HEAs facilitates the formation of stable a single-phase structure. These structures are typically simple, optimizing atomic; for example, face-centered cubic (FCC), body-centered cubic (BCC) and their ordered derivatives [6,7,8,9,10,11]. Studies have shown that single-phase HEAs often cannot achieve both excellent strength and plasticity simultaneously. The FCC phase has high symmetry and low shear resistance, which grants alloys good ductility and plasticity, but results in lower strength [12,13]. In contrast, HEAs with single-phase BCC possess high tensile strength but poor toughness and are prone to brittle fracture [14,15]. Both single-phase FCC and BCC HEAs have distinct advantages and limitations. Therefore, dual-phase or multi-phase HEAs, which exhibit superior mechanical properties, are often preferred. A common approach to optimizing alloy properties is designing alloys by doping different elements to obtain dual-phase or multi-phase structures [16]. The Ni₂₀Al₂₀Co₂₀Fe₂₀Nb₂₀ alloy [17], which consists of ordered body-centered cubic (B2) and Laves phase, exhibits an ultra-high hardness of 818.1 HV. Through varying Al content in the AlNiCoFeCr HEAs, Ma et al. [18] achieved a triple-phase composite structure with BCC + B2 matrix with a minor FCC phase. As a result of the presence of this three-phase structure, the alloy achieves a compressive yield strength (σ_y) of 1394 MPa. Research indicates that adding Al to HEAs with an FCC phase can induce the formation of a BCC/B2 phase. The volume fraction of the precipitation phase increases with increasing Al content, eventually leading to a complete transformation into a BCC/B2 phase [19,20]. The addition of some non-metallic elements significantly influences the mechanical behavior and microstructural features of HEAs [21,22]. Carbon (C) has garnered increasing interest within the realm of HEAs, as a result of its strong alloying ability and potential to enhance mechanical properties and phase stability [23,24]. Jae et al. [25] found that adding C to the (NiCo)₇₅Cr₁₇Fe₈ alloy significantly improves its mechanical properties. Controlling the carbon content is crucial, as excessive carbide precipitation can reduce plasticity and toughness [26]. The trade-off between toughness and strength in HEAs still needs further exploration.

Among traditional strengthening mechanisms, grain refinement and solid solution strengthening contribute to the improvement of alloy mechanical properties. However, these mechanisms provide limited enhancement to the overall performance of the material [27]. The synergistic interactions between distinct phases in multi-phase materials can greatly enhance their overall mechanical properties [28]. Alloys exhibiting dendritic structures are representative multi-phase materials, with more complex microstructures [29]. Dendrite (DR) and interdendrite (ID) typically exhibit distinct phase compositions, and the resulting differences in their mechanical properties contribute to a complementary effect that enhances the overall performance of the alloy [30]. Santiago et al. [31] modified the microstructure of CoCrFeMoNi HEAs by doping with Ti and found that the resulting dendritic composite structure enhanced the alloy hardness by 22.2%. Multi-phase dendritic composite structures with appropriate DR and ID volume fractions may enhance mechanical properties, which warrants further study.

This study explores the microstructural evolution and phase composition of the dendritic structure, where the volume fractions of the dendritic region (DR) and interdendritic region (ID) were systematically varied by adjusting the aluminum (Al) content. The mechanisms leading to the observed mechanical properties are discussed by analyzing the dendrite morphology, precipitate phase size, and fracture characteristics.

2. Materials and Methods

In this paper, Al_xFe_1.5CoNiC_0.12 (x = 0.5, 1.0) HEAs have been studied. The constituent elements (purity > 99.95 wt.%) were mixed in a vacuum arc-melting furnace with a water-cooled Cu crucible under an argon atmosphere. Each ingot was remelted a minimum of six times to achieve uniformity. The alloys studied are denoted as Al_0.5 and Al_1.0, respectively. The alloys’ crystal and microstructural features were examined through X-ray diffraction (XRD, D/Max-2500/PCX, Rigaku, Tokyo, Japan) employing Cu-Kα radiation, field-emission scanning electron microscopy (SEM, Hitachi S4800, Hitachi High-Tech, Tokyo, Japan) with backscattered electron imaging, and transmission electron microscopy (TEM, JEM-2010, JEOL, Tokyo, Japan) combined with selected area diffraction pattern (SADP). Energy-dispersive X-ray spectroscopy (EDS), integrated with the SEM setup, was employed for analyzing the chemical composition. The SEM samples were first mechanically ground and polished and were then corroded with diluted corrosive solution (HCl: HNO₃ = 3:1). Samples prepared for TEM analysis were mechanically ground to approximately 50 μM, punched into a 3 mm diameter round plate, and then thinned using twin-jet electro-polishing in a 10 vol% HClO₄ + 90 vol% CH₃OH solution at −20 °C with a 20 V potential.

Cylindrical samples with dimensions of Φ3 mm × 6 mm were prepared using wire electro spark machining for compression testing, following the ASTM E9-19 standard [32] for testing the compressive properties of materials. The sample ends were polished and lubricated with graphite sheets alongside high-pressure grease prior to testing, aimed at reducing friction and improving result accuracy. An Instron 3382 electronic testing machine (Thermo Fisher Scientific, Norwood, MA, USA) was employed for compression tests at ambient temperature, applying a strain rate of 5 × 10⁻⁴ s⁻¹, in accordance with the ISO 6892-1 standard [33] for tensile and compressive testing of metallic materials. An extensometer was utilized to monitor the deformation strain throughout the compression testing process. Three samples were tested for each alloy composition to verify the reliability of the experimental data.

The hardness measurements were conducted using a TriboIndenter Ti-900 nanoindenter (Hysitron, Minneapolis, MN, US), with a load range of 1 μN to 10 mN. The measurements followed the ISO 14577-1 standard [34] for instrumented indentation testing of metallic materials. To ensure accurate results, the thermal drift was maintained below 0.05 nm/s, and the maximum displacement resolution was set to 0.0002 nm. The constant load rate mode was used for indentation, with both loading and unloading rates set at 1000 μN/s, and the loading duration was 8 s. For each phase, six points were measured, and the average value was used as the hardness for that phase. The distance between adjacent nanoindents was maintained at 5 μM, ensuring that each indentation was performed in a non-overlapping region to avoid interference. The nanoindenter tip used in this study was a Berkovich indenter, with a three-sided pyramidal geometry.

3. Results

3.1. Microstructure Characteristics

Figure 1 shows the XRD patterns of the Al_xFe_1.5CoNiC_0.12 (x = 0.5, 1.0) HEAs. The diffraction patterns reveal that the Al_1.0 alloy predominantly exhibits an ordered body-centered cubic structure (B2), with an additional peak corresponding to the (200) reflection of the FCC structure. The Al_1.0 alloy is primarily characterized by the B2 phase, with a small proportion of the FCC phase. In contrast, the Al_0.5 alloy is dominated by the FCC phase, and the B2 phase exhibits a markedly lower peak strength. This indicates that a phase transition occurs with decreasing Al content in Al_xFe_1.5CoNiC_0.12 HEAs, promoting the conversion from B2 to FCC. Notably, the Al_0.5 alloy shows a weak (100) peak around 23.9° in its diffraction pattern, characteristic of the ordered face-centered cubic structure (L1₂), indicating the presence of a certain amount of the L1₂ phase. In summary, the diffraction pattern analysis reveals that the Al_1.0 alloy primarily consists of a dominant B2 phase and a minor FCC phase. The Al_0.5 alloy exhibits a multi-phase structure comprising a predominant number of FCC and B2 phases, along with a minor fraction of L1₂ phases.

To investigate the microstructure of the HEAs, SEM test was performed. Figure 2 presents SEM images of Al_xFe_1.5CoNiC_0.12 (x = 0.5, 1.0) HEAs with varying Al contents in the as-cast state. Figure 2a,c display low-magnification SEM images of the Al_1.0 and Al_0.5 alloys, respectively. The two alloys both display typical dendritic morphology, which can be categorized into two regions: DR and ID. With decreasing Al content, the proportion of the DR in the Al_xFe_1.5CoNiC_0.12 HEAs significantly decreases, while the volume fraction of the ID correspondingly increases. Figure 2b,d display detailed images of the ID for Al_1.0 and Al_0.5 alloys, respectively. A high density of cubic precipitation is observed within the ID in both compositions. This microstructure is similar to that observed in the (FeCoNi)₈₆Al₇Ti₇ alloy, which is strengthened by multicomponent intermetallic nanoparticles (MCINPs) [35]. The primary difference between the ID structures of the two alloys lies in the size of the cubic precipitation.

Statistical analysis of the SEM images for each alloy composition was performed to determine the volume fraction of individual microstructural region and the average size of the precipitation phases in the ID, as summarized in Table 1. The volume fraction of the ID in Al_xFe_1.5CoNiC_0.12 (x = 0.5, 1.0) HEAs microstructure increases from 16% (Al_1.0) to 93% (Al_0.5). The average particle sizes of the cubic precipitation in the ID of Al_1.0 and Al_0.5 alloys are approximately 123 nm and 38 nm, respectively. The particle size distribution histograms for the precipitation in the ID of Al_1.0 (Figure 3a) and Al_0.5 (Figure 3b) alloys are shown clearly, with decreasing Al content, the cubic precipitation size transition from ultrafine (>100 nm) to nanoscale (≤100 nm). This demonstrates that changes in Al content significantly affect the size of the precipitation phase.

In addition, the element distribution in different regions of Al_0.5 alloy was analyzed using EDS. Figure 4 is the element distribution map, which clearly illustrates the compositional segregation within the alloy, particularly the distribution of Al and Fe. The element distribution map reveals that the Al element is predominantly enriched in the ID and notably depleted in the DR, suggesting that alterations in Al content markedly affect the makeup of the dendritic structure.

The microstructures of the two alloys were further characterized using the TEM. Figure 5 shows the TEM image of the Al_1.0 alloy. The low-magnification bright-field TEM image reveals that the alloy structure consists of both DR and ID. The SADPs of the DR along the [001] crystalline zone axis is shown in the inset at the lower right corner of Figure 5a. The superlattice diffraction spots, indicated by blue circles, confirm that the DR region possesses a B2 single-phase structure. Figure 5b presents high-magnification TEM image of the ID, which clearly shows many ultra-fine precipitates, with particle sizes ranging from 90 to 140 nm, distributed throughout the region. The SADPs of both the matrix (red circle c) and precipitate phases (red circle d) between the ID, along the [011] crystal axis, are shown in Figure 5c,d. These images reveal that the precipitate phase exhibits L1₂ structure, while the matrix adopts FCC structure. Therefore, the ID is composed of a uniformly distributed high-density L1₂ cubic precipitate phase and FCC matrix.

The TEM images of the ID of the Al_0.5 alloy are shown in Figure 6. The image (Figure 6a) presents a low-magnification bright-field TEM image of the ID, along with the SADPs along the [011] crystalline zone axis. The ID of the Al_0.5 alloy consists of an FCC matrix phase with a uniformly distributed L1₂ precipitates phase, and it resembles the microstructure in the Al_1.0 alloy. The difference lies in the size of the precipitate phase, which has been reduced to 20–65 nm, in agreement with the SEM analysis. Figure 6b exhibits a representative high-resolution transmission electron microscopy (HRTEM) image of the interface between the precipitate phase and the matrix. The precipitate phase exhibits an ordered structure, whereas the matrix is disordered. The corresponding fast Fourier transform (FFT) images of the precipitate phase and matrix are shown in Figure 6c,d, respectively. These images further confirm that the nano-precipitate phase exhibits L1₂ structure, while the matrix shows an FCC structure. It is noteworthy that the precipitate phase exhibits a coherent interface relationship with the matrix. This interface interaction ensures an exceptionally low lattice mismatch, which significantly enhances the overall mechanical properties of the alloy [35].

3.2. Mechanical Property

To investigate the effect of Al content on the mechanical properties of Al_xFe_1.5CoNiC_0.12 HEAs, compression tests were conducted. Figure 7a presents the compressive stress–strain curves of the as-cast Al_1.0 and Al_0.5 alloys at room temperature. The σ_y of the Al_1.0 and Al_0.5 alloys is 1064 MPa and 1260 MPa, the ultimate compressive strength (σ_u) is 2523 MPa and 3321 MPa, and the fracture strain (ε) is 20.1% and 37.1%, respectively. Compared with the Al_1.0 alloy, the Al_0.5 alloy exhibits superior overall compressive properties, with increases of 18.4%, 31.6%, and 84.6% in σ_y, σ_u and ε, respectively. This enhancement is primarily attributed to the increase in the volume fraction of ID and the reduction in the size of the L1₂ precipitate phase to the nanoscale, which results in the formation of a stable composite structure in the Al_0.5 alloy. To more accurately assess the micromechanical property of each region in the alloy, the hardness (Hc) of the DR and ID was measured. The test results are shown in Figure 7b,c. The detailed information is presented in Table 2. The typical nanoindentation loading-displacement curves of the as-cast Al_xFe_1.5CoNiC_0.12 HEAs at room temperature in the DR and ID, reveal that the hardness of the DR in both Al_0.5 and Al_1.0 alloys is similar. The hardness of the ID in the Al_0.5 alloy is higher than that in the Al_1.0 alloy, with values of 7.2 GPa and 6.5 GPa, respectively. The difference is attributed to the finer precipitate phase in the Al_0.5 alloy. It is important to note that the hardness of the ID in both alloys is higher than that of the DR, primarily due to the presence of numerous ultrafine/nano precipitate particles in the ID. These precipitates effectively impede dislocation motion through the coherent interface effect, thereby enhancing hardness.

4. Discussion

Based on the above results, for the Al_1.0 alloy, reducing the Al content contributes to enhancing its mechanical properties. Notably, both the σ_y and ε exhibit a significant increase, rising from 2523 MPa and 20.1% for the Al_1.0 alloy to 3321 MPa and 37.1% for the Al_0.5 alloy. The lower Al content imparts excellent strength and ductility to the alloy. This outstanding synergy of strength and ductility results from two primary factors. On one hand, the Al_0.5 alloy exhibits a more stable DR-to-ID volume fraction ratio, and the cubic precipitate phase in the ID refined to the nanoscale. On the other hand, the coherent nano precipitate exhibits very low lattice misfit with the matrix, which effectively enhances the stability of the Al_0.5 alloy at the nanoscale [28,35].

The strengthening mechanism of the Al_xFe_1.5CoNiC_0.12 HEAs is discussed below. In general, the σ_y of HEAs is the result of the combined contributions from intrinsic strength and the four strengthening mechanisms (precipitation strengthening, solid solution strengthening, grain boundary strengthening, and dislocation strengthening). As-cast alloys typically exhibit large grain sizes and low dislocation densities. Consequently, the effects of grain boundary strengthening and dislocation strengthening are minimal and are not considered significant in this study. The theoretical yield strength (σ_ty) of the alloys studied can be calculated using the following equation [36]:

(1)${\sigma }_{\mathrm{ty}}={\sigma }_0+{\sigma }_{\mathrm{s}}+{\sigma }_{\mathrm{p}}$

where σ₀ (123 MPa) is the intrinsic strength [37,38]. σ_s represents the strength provided by solid solution strengthening and σ_p represents the strength provided by precipitation strengthening. σ_s is calculated by taking into account the modulus (${\eta }_{\mathrm{i}}^{’}$) and lattice mismatches (${\delta }_{\mathrm{i}}$) using the following equation [39,40]:

(2)${\sigma }_{\mathrm{s}}={{\displaystyle \frac{\mathrm{M}\mu }{\mathrm{Z}}}({\eta }_{\mathrm{i}}^2+{\alpha }^2{\delta }_{\mathrm{i}}^2)}^{2/3}{\mathrm{X}}_{\mathrm{i}}^{2/3}$

(3)${\eta }_{\mathrm{i}}^{’}={\displaystyle \frac{{\eta }_{\mathrm{i}}}{1+0.5\left|{\eta }_{\mathrm{i}}\right|}},{\eta }_{\mathrm{i}}={\displaystyle \frac{\mathrm{d}\mu }{\mathrm{d}{\mathrm{X}}_{\mathrm{i}}}}{\displaystyle \frac{1}{\mu }},{\delta }_{\mathrm{i}}={\displaystyle \frac{\mathrm{da}}{\mathrm{d}{\mathrm{X}}_{\mathrm{i}}}}{\displaystyle \frac{1}{\mathrm{a}}}$

where M is the Taylor factor, which is 2.73 for BCC phase and 3.06 for FCC phase [41]. The Z value is 820. μ is the shear modulus of the element. α is a constant with a value of 16 [42]. X_i denotes the molar fraction of solute species i, and a represents the lattice constant. The contribution of solid solution strengthening in the DR and ID regions for the Al_0.5 alloy is 36 MPa and 55 MPa, respectively. In the case of the Al_1.0 alloy, the contributions are 33 MPa and 52 MPa for the DR and ID regions, respectively. The total solid solution strengthening of the alloy can be calculated using a weighted average, where the weighting coefficient corresponds to the volume fraction of each phase. Therefore, the strength resulting from solid solution strengthening in the Al_0.5 and Al_1.0 alloys is 54 MPa and 36 MPa, respectively.

The precipitation strengthening effect of precipitates in the HEAs can be categorized into two mechanisms: Orowan dislocation bypass and dislocation cutting, based on the interaction between dislocations and the precipitated phase [43]. The size of the precipitate particles and their coherence with the matrix determine which mechanism predominates during alloy deformation. The Orowan dislocation bypass mechanism is dominant when the particle radius exceeds a critical threshold or when the precipitates are incoherent with the matrix. However, when the precipitates are small enough and coherent with the matrix, the dislocation cutting mechanism becomes predominant [44].

For the Al_0.5 alloy, the L1₂ precipitates at the nanoscale in the FCC matrix have a coherent interface, and the dislocation cutting mechanism is dominant. To assess the impact of the dislocation cutting mechanism on precipitates, three factors that contribute to the enhancement of σ_y include precipitate-matrix coherency (Δσ_CS), modulus mismatch (Δσ_MS) and atom ordering (Δσ_OS). The calculation equation is expressed as follows [45]:

(4)$\Delta {\sigma }_{\mathrm{CS}}=\mathrm{M}{\alpha }_{\epsilon }{\left(\mathrm{G}{\epsilon }_{\mathrm{c}}\right)}^{{\displaystyle \frac{3}{2}}}{\left({\displaystyle \frac{\mathrm{rf}}{0.5\mathrm{Gb}}}\right)}^{{\displaystyle \frac{1}{2}}}$

(5)$\Delta {\sigma }_{\mathrm{MS}}=0.0055 \mathrm{M}{\left(\Delta \mathrm{G}\right)}^{{\displaystyle \frac{3}{2}}}{\left({\displaystyle \frac{2\mathrm{f}}{\mathrm{G}}}\right)}^{{\displaystyle \frac{1}{2}}}{\left({\displaystyle \frac{\mathrm{r}}{\mathrm{b}}}\right)}^{{\displaystyle \frac{3\mathrm{m}}{2}}-1}$

(6)$\Delta {\sigma }_{\mathrm{OS}}=0.81 \mathrm{M}\left({\displaystyle \frac{{\gamma }_{\mathrm{apb}}}{2\mathrm{b}}}\right){\left({\displaystyle \frac{3\pi \mathrm{f}}{8}}\right)}^{{\displaystyle \frac{1}{2}}}$

where m = 0.85 is the constant for the edge dislocation [41]. α_ε = 2.6 [43] is a constant. G is the shear modulus of CoFeNi alloy (68 GPa) [46]. ε_c is the constrained lattice parameter mismatch [41]. b = ${\displaystyle \frac{\sqrt{2}}{2}}\alpha$ is the Burgers vector [47]. r is the average particle size of the nanoscale L1₂ precipitate phase. f is the volume fraction of the nano-precipitate. ΔG (7.97 GPa) is the modulus mismatch between the FCC matrix phase and the L1₂ phase. γ_apb is the APB energy of the L1₂ phase [48]. Based on Equations (4)–(6), the precipitation strengthening contributions from precipitate-matrix coherency (Δσ_CS), modulus mismatch (Δσ_MS) and atom ordering (Δσ_OS), are 478, 147 and 1025 MPa, respectively. Since Δσ_CS + Δσ_MS is less than Δσ_OS, atomic ordering largely determines the precipitation strengthening of L1₂ phase in Al_0.5 alloy. The contribution of precipitation strengthening caused by L1₂ phase to yield strength is 1025 MPa.

For the Al_0.5 alloy, the size of the B2 precipitates is larger, and the Orowan dispersion mechanism predominates. The strength increment of the Orowan-type mechanism (Δσ_Orowan) can be calculated using Orowan-Ashby equation [49]:

(7)$\Delta {\sigma }_{\mathrm{Orowan}}=0.4 \mathrm{MGb}{\displaystyle \frac{\ln (2\sqrt{\frac{2\mathrm{r}}{3}/\mathrm{b}})}{{\lambda }_{\mathrm{P}}\pi \sqrt{1-\nu }}}$

where ν = 0.3 is the Poisson’s ratio of the alloy [50], r = 5.6 μM and f = 5.1% are the average particle size and volume fraction of the precipitated phase in the alloy, respectively. λ_P is the precipitation spacing of the B2 phase, defined as follows:

(8)${\lambda }_{\mathrm{P}}=2\sqrt{{\displaystyle \frac{2}{3}}\mathrm{r}}\left(\sqrt{{\displaystyle \frac{\pi }{4\mathrm{f}}}-1}\right)$

By incorporating these parameters into the above equation, the contribution of precipitation strengthening caused by B2 phase to yield strength is 45 MPa. In summary, the σ_ty of Al_0.5 alloy is 1247 MPa.

For Al_1.0 alloy, due to the large size of L1₂ phase and B2 phase, the precipitation strengthening of these two phases is calculated by Orowan dispersion mechanism. r = 123 nm and f = 14.3% are the average particle size and volume fraction of the L1₂ phase. r = 29 μM and f = 63.3% are the average particle size and volume fraction of the B2 phase. According to the equation, the contribution of L1₂ and B2 phase precipitation strengthening to the yield strength is 586 MPa and 305 MPa, respectively. The σ_ty of Al_1.0 alloy is finally obtained as 1050 MPa.

As shown in Figure 8, σ_ty agrees well with σ_y. The difference in σ_y between the two alloys is primarily due to the variation in strength provided by precipitation strengthening. This confirms that the difference in σ_y results from the increase in the volume fraction of the ID and the reduction in the precipitate particle size.

To gain a deeper insight into how the mechanical properties, microstructure, and fracture mechanism of Al_xFe_1.5CoNiC_0.12 HEAs are interconnected, we investigated the fracture morphology of Al_1.0 and Al_0.5 alloys. The compression property of Al_1.0 alloy is poor, with ε of only 20.1%. As seen in Figure 9a, the DR and ID can be clearly distinguished in the fracture morphology. Figure 9b shows a high-magnification fracture morphology image of the Al_1.0 alloy. The fracture surface in the image reveals distinct regions exhibiting different fracture behaviors. The region corresponding to DR displays typical brittle fracture characteristics, while the region corresponding to ID exhibits ductile fracture features, characterized by a dimpled morphology. An obvious tearing edge exists between the two regions, indicating that the alloy exhibits intergranular fracture during compression, resulting in poor compression plasticity. The fracture morphology of the Al_0.5 alloy are shown in Figure 9c,d. Its compression fracture exhibits a typical dimple shear flow pattern [51,52], suggesting that the alloy possesses excellent compression plasticity. The fracture morphology of the Al_0.5 alloy exhibits a higher dimple density, reflecting excellent compressive performance.

5. Conclusions

In this study, by changing the Al content, the Al_xFe_1.5CoNiC_0.12 (x = 0.5, 1.0) HEAs with different precipitate phase sizes and proportions of dendrite and interdendrite were obtained. The microstructure, mechanical properties, and strengthening mechanisms of the alloy were investigated. The following conclusions can be drawn:

1. The Al_xFe_1.5CoNiC_0.12 HEAs exhibit triple-phase composite structure comprising B2, FCC, and L1₂. The decrease in Al content leads to an increase in the interdendritic volume fraction and a shift in the size of the L1₂ precipitation phase from ultrafine to nanoscale.

2. The results of the mechanical properties test shows that hardness of the interdendrite of Al_xFe_1.5CoNiC_0.12 HEAs is higher than that of the dendrite. The interdendritic hardness, yield strength, and strain increased from 6.5 GPa, 1064 MPa and 20.1% for Al_1.0 alloy to 7.2 GPa, 1260 MPa and 37.1% for Al_0.5 alloy, respectively.

3. For the Al_0.5 alloy, the high-density nanoscale L1₂ precipitate phase in the interdendrite exhibits minimal lattice mismatch with the FCC matrix, which effectively enhances the stability at the nanoscale. This composite structure contributes to its excellent performance. These properties suggest that the alloy is promising for structural applications in industries requiring high-performance materials, such as aerospace or automotive.

Author Contributions

Conceptualization, Y.J. and A.L.; methodology, Y.J. and A.L.; software, B.Z., K.K. and D.H.; validation, J.Z., C.C. and S.L.; formal analysis, Y.J. and J.Z.; investigation, D.H. and M.X.; resources, G.L.; data curation, Y.J., S.L. and M.X.; writing—original draft preparation, Y.J.; writing—review and editing, G.L.; visualization, C.C. and Y.L.; supervision, Y.L., B.Z. and K.K.; project administration, G.L.; funding acquisition, G.L. All authors have read and agreed to the published version of the manuscript.

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding authors.

Conflicts of Interest

The authors declare no conflicts of interest.

Footnotes

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Figures and Tables

Figure 1. XRD patterns of Al0.5 and Al1.0 alloys.

Figure 2. Microstructure of Al1.0 and Al0.5 alloys observed through SEM: (a,c) low magnification images, (b,d) detailed views of the ID. (a,b) are Al1.0 alloy, while (c,d) are Al0.5 alloy.

Figure 3. Statistical chart of particle size distribution of ID precipitation, (a) Al1.0 alloy and (b) Al0.5 alloy.

Figure 4. The SEM-EDS element distribution map of the Al0.5 alloy. (a) SEM image of the alloy. (b) EDS mapping of the Al element. (c) EDS mapping of the C element. (d) EDS mapping of the Fe element. (e) EDS mapping of the Co element. (f) EDS mapping of the Ni element.

Figure 5. TEM image of the Al1.0 alloy. (a) represents triple-phase composite structure. The blue arrow indicates DR, and the red arrow indicates ID. The inset in (a) is the SADPs in the DR. The detailed view of the ID (b–d) are the SADPs of the precipitate phase and the matrix in the ID, respectively.

Figure 6. The TEM micrograph of the ID in the Al0.5 alloy, accompanied by the corresponding SADPs (a), HRTEM image of the interface between the precipitate phase and the matrix (b) and the corresponding Fourier transform image of the precipitate phase (c) and the matrix (d). The blue box represents the L12 precipitate phase, and the red box represents the FCC matrix.

Figure 7. Room-temperature compressive curves (a) of the as-cast AlxFe1.5CoNiC0.12 (x = 0.5, 1.0) HEAs tested. Nanoindentation load–displacement curves of alloys DR (b) and ID (c).

Figure 8. The results of the derivation of the strengthening mechanism and the σy, (a) and (b) are Al0.5 and Al1.0 alloys, respectively. Blue, green, yellow, and red represent intrinsic strength, solid solution strengthening, precipitation strengthening and experimental yield strength, respectively.

Figure 9. Fracture morphologies of the Al1.0 alloy (a,b), fracture morphologies of the Al0.5 alloy (c,d). (b,d) are high-magnification TEM images of the square areas marked in (a) and (c), respectively.

Table 1

Volume fractions of the different regions in the Al_0.5 and Al_1.0 alloys, and the average size of the precipitation phases in the ID.

Alloys	Volume Fraction of Region (%)		Average Size of the Precipitation Phases in the ID (nm)
	DR	ID
Al1.0	84	16	123 ± 2
Al0.5	7	93	38 ± 3

Table 2

Nanoindentation test results and standard deviation of the different regions of as-cast Al_xFe_1.5CoNiC_0.12 (x = 0.5, 1.0) HEAs.

HEAs	Hc (GPa)
	DR	ID
Al1.0	6.1 ± 0.1	6.5 ± 0.1
Al0.5	6.2 ± 0.1	7.2 ± 0.2

References

1. Miracle, D.B.; Senkov, O.N. A critical review of high entropy alloys and related concepts. Acta Mater.; 2017; 122, pp. 448-511. [DOI: https://dx.doi.org/10.1016/j.actamat.2016.08.081]

2. Krishna, S.A.; Noble, N.; Radhika, N.; Saleh, B. A comprehensive review on advances in high entropy alloys: Fabrication and surface modification methods, properties, applications, and future prospects. J. Manuf. Process.; 2024; 109, pp. 583-606. [DOI: https://dx.doi.org/10.1016/j.jmapro.2023.12.039]

3. Ma, Y.; Wang, W.; He, J.; Zhu, Y.; Wu, X.; Yuan, F. Medium/high entropy alloys with heterogeneous structures for superior properties: A review. Mater. Today Adv.; 2025; 25, 100553. [DOI: https://dx.doi.org/10.1016/j.mtadv.2024.100553]

4. Wang, M.; Lu, Y.; Wang, T.; Zhang, C.; Cao, Z.; Li, T.; Liaw, P.K. A novel bulk eutectic high-entropy alloy with outstanding as-cast specific yield strengths at elevated temperatures. Scr. Mater.; 2021; 204, 114132. [DOI: https://dx.doi.org/10.1016/j.scriptamat.2021.114132]

5. Wang, M.; Lu, Y.; Lan, J.; Wang, T.; Zhang, C.; Cao, Z.; Li, T.; Liaw, P.K. Lightweight, ultrastrong and high thermal-stable eutectic high-entropy alloys for elevated-temperature applications. Acta Mater.; 2023; 248, 118806. [DOI: https://dx.doi.org/10.1016/j.actamat.2023.118806]

6. Senkov, O.N.; Scott, J.M.; Senkova, S.V.; Miracle, D.B.; Woodward, C.F. Microstructure and room temperature properties of a high-entropy TaNbHfZrTi alloy. J. Alloys Compd.; 2011; 509, pp. 6043-6048. [DOI: https://dx.doi.org/10.1016/j.jallcom.2011.02.171]

7. Ng, C.; Guo, S.; Luan, J.; Shi, S.; Liu, C.T. Entropy-driven phase stability and slow diffusion kinetics in an Al0.5CoCrCuFeNi high entropy alloy. Intermetallics; 2012; 31, pp. 165-172. [DOI: https://dx.doi.org/10.1016/j.intermet.2012.07.001]

8. Yuan, Y.; Wu, Y.; Tong, X.; Zhang, H.; Wang, H.; Liu, X.J.; Ma, L.; Suo, H.L.; Lu, Z.P. Rare-earth high-entropy alloys with giant magnetocaloric effect. Acta Mater.; 2017; 125, pp. 481-489. [DOI: https://dx.doi.org/10.1016/j.actamat.2016.12.021]

9. Li, Z.; Zhao, S.; Ritchie, R.O.; Meyers, M.A. Mechanical properties of high-entropy alloys with emphasis on face-centered cubic alloys. Prog. Mater. Sci.; 2019; 102, pp. 296-345. [DOI: https://dx.doi.org/10.1016/j.pmatsci.2018.12.003]

10. Ostovari Moghaddam, A.; Fereidonnejad, R.; Cabot, A. Semi-ordered high entropy materials: The case of high entropy intermetallic compounds. J. Alloys Compd.; 2023; 960, 170802. [DOI: https://dx.doi.org/10.1016/j.jallcom.2023.170802]

11. Wu, Y.; Liaw, P.K.; Li, R.; Zhang, W.; Geng, G.; Yan, X.; Liu, G.; Zhang, Y. Relationship between the unique microstructures and behaviors of high-entropy alloys. Int. J. Miner. Metall. Mater.; 2024; 31, pp. 1350-1363. [DOI: https://dx.doi.org/10.1007/s12613-023-2777-4]

12. Sohn, S.; Liu, Y.; Liu, J.; Gong, P.; Prades-Rodel, S.; Blatter, A.; Scanley, B.E.; Broadbridge, C.C.; Schroers, J. Noble metal high entropy alloys. Scr. Mater.; 2017; 126, pp. 29-32. [DOI: https://dx.doi.org/10.1016/j.scriptamat.2016.08.017]

13. Otto, F.; Dlouhý, A.; Somsen, C.; Bei, H.; Eggeler, G.; George, E.P. The influences of temperature and microstructure on the tensile properties of a CoCrFeMnNi high-entropy alloy. Acta Mater.; 2013; 61, pp. 5743-5755. [DOI: https://dx.doi.org/10.1016/j.actamat.2013.06.018]

14. Zhang, Y.; Zuo, T.T.; Tang, Z.; Gao, M.C.; Dahmen, K.A.; Liaw, P.K.; Lu, Z.P. Microstructures and properties of high-entropy alloys. Prog. Mater. Sci.; 2014; 61, pp. 1-93. [DOI: https://dx.doi.org/10.1016/j.pmatsci.2013.10.001]

15. Senkov, O.N.; Wilks, G.B.; Miracle, D.B.; Chuang, C.P.; Liaw, P.K. Refractory high-entropy alloys. Intermetallics; 2010; 18, pp. 1758-1765. [DOI: https://dx.doi.org/10.1016/j.intermet.2010.05.014]

16. Filipoiu, N.; Nemnes, G.A. Prediction of Equilibrium Phase, Stability and Stress-Strain Properties in Co-Cr-Fe-Ni-Al High Entropy Alloys Using Artificial Neural Networks. Metals; 2020; 10, 1569. [DOI: https://dx.doi.org/10.3390/met10121569]

17. Gu, L.; Liang, N.; Liu, Y.; Chen, Y.; Liu, J.; Sun, Y.; Zhao, Y. Novel as-cast NiAlCoFeNb dual-phase high-entropy alloys with high hardness. Mater. Lett.; 2022; 324, 132676. [DOI: https://dx.doi.org/10.1016/j.matlet.2022.132676]

18. Ma, Y.; Jiang, B.; Li, C.; Wang, Q.; Dong, C.; Liaw, P.; Xu, F.; Sun, L. The BCC/B2 Morphologies in Al_xNiCoFeCr High-Entropy Alloys. Metals; 2017; 7, 57. [DOI: https://dx.doi.org/10.3390/met7020057]

19. Joseph, J.; Stanford, N.; Hodgson, P.; Fabijanic, D.M. Understanding the mechanical behaviour and the large strength/ductility differences between FCC and BCC Al_xCoCrFeNi high entropy alloys. J. Alloys Compd.; 2017; 726, pp. 885-895. [DOI: https://dx.doi.org/10.1016/j.jallcom.2017.08.067]

20. Zuo, T.T.; Li, R.B.; Ren, X.J.; Zhang, Y. Effects of Al and Si addition on the structure and properties of CoFeNi equal atomic ratio alloy. J. Magn. Magn. Mater.; 2014; 371, pp. 60-68. [DOI: https://dx.doi.org/10.1016/j.jmmm.2014.07.023]

21. Wei, D.; Gong, W.; Tsuru, T.; Lobzenko, I.; Li, X.; Harjo, S.; Kawasaki, T.; Do, H.-S.; Bae, J.W.; Wagner, C. et al. Si-addition contributes to overcoming the strength-ductility trade-off in high-entropy alloys. Int. J. Plast.; 2022; 159, 103443. [DOI: https://dx.doi.org/10.1016/j.ijplas.2022.103443]

22. Wu, P.; Zhang, Y.; Han, L.; Gan, K.; Yan, D.; Wu, W.; He, L.; Fu, Z.; Li, Z. Unexpected sluggish martensitic transformation in a strong and super-ductile high-entropy alloy of ultralow stacking fault energy. Acta Mater.; 2023; 261, 119389. [DOI: https://dx.doi.org/10.1016/j.actamat.2023.119389]

23. Baker, I. Interstitials in f.c.c. High Entropy Alloys. Metals; 2020; 10, 695. [DOI: https://dx.doi.org/10.3390/met10050695]

24. Lukianova, O.; Rao, Z.; Kulitckii, V.; Li, Z.; Wilde, G.; Divinski, S.V. Impact of interstitial carbon on self-diffusion in CoCrFeMnNi high entropy alloys. Scr. Mater.; 2020; 188, pp. 264-268. [DOI: https://dx.doi.org/10.1016/j.scriptamat.2020.07.044]

25. Song, J.S.; Lee, B.J.; Moon, W.J.; Hong, S.I. Mechanical Performance and Microstructural Evolution of (NiCo)₇₅Cr₁₇Fe₈C_x (x = 0~0.83) Medium Entropy Alloys at Room and Cryogenic Temperatures. Metals; 2020; 10, 1646. [DOI: https://dx.doi.org/10.3390/met10121646]

26. Ali, N.; Zhang, L.; Dongming, L.; Zhou, H.; Sanaullah, K.; Zhang, C.; Nian, Y.; Cheng, J. Enhancing the strength-ductility synergy in hot-forged Fe50Mn30Co10Cr10 high entropy alloy (HEA) through carbon additions. J. Mater. Res. Technol.; 2024; 29, pp. 5646-5655. [DOI: https://dx.doi.org/10.1016/j.jmrt.2024.02.232]

27. Ren, H.; Chen, R.R.; Gao, X.F.; Liu, T.; Qin, G.; Wu, S.P.; Guo, J.J. Phase formation and mechanical features in (AlCoCrFeNi)100-Hf high-entropy alloys: The role of Hf. Mater. Sci. Eng. A; 2022; 858, 144156. [DOI: https://dx.doi.org/10.1016/j.msea.2022.144156]

28. Jiang, S.; Wang, H.; Wu, Y.; Liu, X.; Chen, H.; Yao, M.; Gault, B.; Ponge, D.; Raabe, D.; Hirata, A. et al. Ultrastrong steel via minimal lattice misfit and high-density nanoprecipitation. Nature; 2017; 544, pp. 460-464. [DOI: https://dx.doi.org/10.1038/nature22032]

29. Ye, X.; Diao, Z.; Lei, H.; Wang, L.; Li, Z.; Li, B.; Feng, J.; Chen, J.; Liu, X.; Fang, D. Multi-phase FCC-based composite eutectic high entropy alloy with multi-scale microstructure. Mater. Sci. Eng. A; 2024; 889, 145815. [DOI: https://dx.doi.org/10.1016/j.msea.2023.145815]

30. Shang, G.; Zheng, W.; Wang, J.; Lu, X.-G. Microstructural evolution and local mechanical properties of dendrites in Al0.6CoCrFeNi high entropy alloy. Mater. Sci. Eng. A; 2022; 846, 143294. [DOI: https://dx.doi.org/10.1016/j.msea.2022.143294]

31. Brito-Garcia, S.J.; Mirza-Rosca, J.C.; Jimenez-Marcos, C.; Voiculescu, I. Impact of Ti Doping on the Microstructure and Mechanical Properties of CoCrFeMoNi High-Entropy Alloy. Metals; 2023; 13, 854. [DOI: https://dx.doi.org/10.3390/met13050854]

32. Padilla-González, A.; González, G.; Alfonso, I.; Vidilli, A.L.; Otani, L.B.; Figueroa, I.A. Development and Mechanical Characterization of a CoCr-Based Multiple-Principal-Element Alloy. Metallogr. Microstruct. Anal.; 2024; 13, pp. 730-740. [DOI: https://dx.doi.org/10.1007/s13632-024-01111-z]

33. Piątkowski, J.; Hejne, M.; Wieszała, R. Influence of manganese content on the microstructure and properties of AlSi10MnMg(Fe) alloy for die castings. Arch. Mater. Sci. Eng.; 2023; 123, pp. 5-12. [DOI: https://dx.doi.org/10.5604/01.3001.0053.9750]

34. Genta, G.; Maculotti, G.; Barbato, G.; Levi, R.; Galetto, M. Effect of contact stiffness and machine calibration in nano-indentation testing. Procedia Cirp.; 2018; 78, pp. 208-212. [DOI: https://dx.doi.org/10.1016/j.procir.2018.08.313]

35. Yang, T.; Zhao, Y.; Tong, Y.; Jiao, Z.; Wei, J.; Cai, J.; Han, X.; Chen, D.; Hu, A.; Kai, J. et al. Multicomponent intermetallic nanoparticles and superb mechanical behaviors of complex alloys. Science; 2018; 362, pp. 933-937. [DOI: https://dx.doi.org/10.1126/science.aas8815]

36. Kamikawa, N.; Sato, K.; Miyamoto, G.; Murayama, M.; Sekido, N.; Tsuzaki, K.; Furuhara, T. Stress–strain behavior of ferrite and bainite with nano-precipitation in low carbon steels. Acta Mater.; 2015; 83, pp. 383-396. [DOI: https://dx.doi.org/10.1016/j.actamat.2014.10.010]

37. Wu, Z.; Gao, Y.; Bei, H. Thermal activation mechanisms and Labusch-type strengthening analysis for a family of high-entropy and equiatomic solid-solution alloys. Acta Mater.; 2016; 120, pp. 108-119. [DOI: https://dx.doi.org/10.1016/j.actamat.2016.08.047]

38. Yoshida, S.; Ikeuchi, T.; Bhattacharjee, T.; Bai, Y.; Shibata, A.; Tsuji, N. Effect of elemental combination on friction stress and Hall-Petch relationship in face-centered cubic high/medium entropy alloys. Acta Mater.; 2019; 171, pp. 201-215. [DOI: https://dx.doi.org/10.1016/j.actamat.2019.04.017]

39. Toda-Caraballo, I.; Rivera-Díaz-del-Castillo, P.E.J. Modelling solid solution hardening in high entropy alloys. Acta Mater.; 2015; 85, pp. 14-23. [DOI: https://dx.doi.org/10.1016/j.actamat.2014.11.014]

40. Walbrühl, M.; Linder, D.; Ågren, J.; Borgenstam, A. Modelling of solid solution strengthening in multicomponent alloys. Mater. Sci. Eng. A; 2017; 700, pp. 301-311. [DOI: https://dx.doi.org/10.1016/j.msea.2017.06.001]

41. He, J.Y.; Wang, H.; Huang, H.L.; Xu, X.D.; Chen, M.W.; Wu, Y.; Liu, X.J.; Nieh, T.G.; An, K.; Lu, Z.P. A precipitation-hardened high-entropy alloy with outstanding tensile properties. Acta Mater.; 2016; 102, pp. 187-196. [DOI: https://dx.doi.org/10.1016/j.actamat.2015.08.076]

42. Moravcik, I.; Hornik, V.; Minárik, P.; Li, L.; Dlouhy, I.; Janovska, M.; Raabe, D.; Li, Z. Interstitial doping enhances the strength-ductility synergy in a CoCrNi medium entropy alloy. Mater. Sci. Eng. A; 2020; 781, 139242. [DOI: https://dx.doi.org/10.1016/j.msea.2020.139242]

43. Seidman, D.N.; Marquis, E.A.; Dunand, D.C. Precipitation strengthening at ambient and elevated temperatures of heat-treatable Al (Sc) alloys. Acta Mater.; 2002; 50, pp. 4021-4035. [DOI: https://dx.doi.org/10.1016/S1359-6454(02)00201-X]

44. Ma, Y.; Wang, Q.; Jiang, B.B.; Li, C.L.; Hao, J.M.; Li, X.N.; Dong, C.; Nieh, T.G. Controlled formation of coherent cuboidal nanoprecipitates in body-centered cubic high-entropy alloys based on Al₂(Ni,Co,Fe,Cr)₁₄ compositions. Acta Mater.; 2018; 147, pp. 213-225. [DOI: https://dx.doi.org/10.1016/j.actamat.2018.01.050]

45. Nembach, E. Precipitation hardening caused by a difference in shear modulus between particle and matrix. Phys. Status Solidi A; 1983; 78, pp. 571-581. [DOI: https://dx.doi.org/10.1002/pssa.2210780223]

46. Laplanche, G.; Gadaud, P.; Bärsch, C.; Demtröder, K.; Reinhart, C.; Schreuer, J.; George, E.P. Elastic moduli and thermal expansion coefficients of medium-entropy subsystems of the CrMnFeCoNi high-entropy alloy. J. Alloys Compd.; 2018; 746, pp. 244-255. [DOI: https://dx.doi.org/10.1016/j.jallcom.2018.02.251]

47. Lu, T.; He, T.; Andreoli, A.F.; Yao, N.; Wan, B.; Scudino, S. The origin of good mechanical and soft magnetic properties in a CoFeNi-based high-entropy alloy with hierarchical structure. Mater. Charact.; 2024; 215, 114237. [DOI: https://dx.doi.org/10.1016/j.matchar.2024.114237]

48. Zhou, X.; Chen, J.; Ding, R.; Wu, H.; Du, J.; He, J.; Wang, W.; Sun, W.; Liu, Y.; Sha, G. et al. Carbon-driven coherent nanoprecipitates enable ultrahigh yield strength in a high-entropy alloy. Mater. Today Nano; 2023; 22, 100331. [DOI: https://dx.doi.org/10.1016/j.mtnano.2023.100331]

49. Martin, J.W. Particle strengthening of metals and alloys. Mater. Sci. Technol.; 1997; 13, 705.

50. Zhang, L.J.; Zhang, M.D.; Zhou, Z.; Fan, J.T.; Cui, P.; Yu, P.F.; Jing, Q.; Ma, M.Z.; Liaw, P.K.; Li, G. et al. Effects of rare-earth element, Y, additions on the microstructure and mechanical properties of CoCrFeNi high entropy alloy. Mater. Sci. Eng. A; 2018; 725, pp. 437-446. [DOI: https://dx.doi.org/10.1016/j.msea.2018.04.058]

51. Louzguine-Luzgin, D.V.; Louzguina-Luzgina, L.V.; Kato, H.; Inoue, A. Investigation of Ti–Fe–Co bulk alloys with high strength and enhanced ductility. Acta Mater.; 2005; 53, pp. 2009-2017. [DOI: https://dx.doi.org/10.1016/j.actamat.2005.01.012]

52. He, F.; Wang, Z.; Cheng, P.; Wang, Q.; Li, J.; Dang, Y.; Wang, J.; Liu, C. Designing eutectic high entropy alloys of CoCrFeNiNbx. J. Alloys Compd.; 2016; 656, pp. 284-289. [DOI: https://dx.doi.org/10.1016/j.jallcom.2015.09.153]