Microstructural and Mechanical Characterization of Co-Free Al_xTi_xCrFe₂Ni High-Entropy Alloys

Abstract

This study investigates the effect of Al_x and Ti_x content (x = 0.1, 0.2, 0.3, 0.4, 0.5, and 0.6) on the microstructural evolution and mechanical properties of Co-free high-entropy Al_xTi_xCrFe₂Ni alloys in both as-cast and homogenized conditions. The research focused on the characterization of structural features, melting behavior, and mechanical performance. Microstructural characterization was carried out using optical microscopy, scanning electron microscopy (SEM), electron backscatter diffraction (EBSD), X-ray diffraction (XRD), and differential thermal analysis (DTA). Mechanical properties were evaluated through Vickers hardness testing and uniaxial compression tests. Increasing the Al and Ti content induced a transformation from a single-phase FCC structure to a dual-phase BCC structure, with the primary BCC phase strengthened by spherical precipitates rich in Al, Ti, and Ni. Homogenization annealing at 1100 °C led to an overall improvement in the mechanical properties. The Al_0.3Ti_0.3CrFe₂Ni alloy exhibited the most balanced combination of strength and ductility after annealing, achieving a compressive yield strength of 1510 MPa, a compressive strength of 3316 MPa, and a compressive plastic strain of 45%.

1. Introduction

Research on high-entropy alloys (HEAs) has introduced a novel approach to the design of metallic materials composed of five or more elements in equal or near-equiatomic atomic ratios [1,2]. This concept enables the formation of stable solid solutions with simple ordered structures, such as face-centered cubic (FCC), body-centered cubic (BCC), and hexagonal close-packed (HCP) lattices, which exhibit outstanding mechanical properties. As a result, HEAs have attracted considerable interest for applications requiring high tensile and compressive strength and good ductility [3,4,5], high hardness, and reliable performance under extreme conditions [6,7,8,9].

First-generation HEAs were designed as single-phase solid solutions, based either on FCC structures, which provide high ductility, or BCC structures, which offer superior strength, albeit at the expense of reduced plasticity. Subsequent development of HEAs has focused on so-called multiphase alloys that combine the synergistic advantages of both FCC and BCC structures, enabling a balance of high strength and good ductility. Alloying elements such as aluminum [10], which stabilizes the BCC phase and promotes the formation of strengthening precipitates, play a significant role in these systems. A large part of current research is directed toward so-called Co-free HEAs [11,12], aiming to develop cost-efficient HEA systems capable of competing with advanced steels and superalloys [13,14]. In this type of alloy, stabilization of the FCC phase is achieved by substituting Co with other elements such as Cu and Mn [15,16,17] or by increasing the content of Ni [5,18] or Fe [19,20]. Another important alloying element in HEAs is Ti, which stabilizes the BCC phase and supports the formation of strengthening precipitates, thereby contributing to increased material strength, as in the CrFeNiAl₀.₄Ti₀.₂ alloy [21], which exhibits exceptional mechanical properties with a yield strength (${\mathsf{\sigma }}_{\mathrm{Y}}$) of 1600 MPa, a fracture strength (${\mathsf{\sigma }}_{\mathrm{f}}$) of $3819 \mathrm{M}\mathrm{P}\mathrm{a},$ and a fracture strain (${\mathsf{\epsilon }}_{\mathrm{f}}$) of $45\%.$ Similarly, Ni-rich alloys, such as the FCC matrix Ni₄₅(FeCoCr)₄₀(AlTi)₁₅ alloy [22]_, also demonstrate high mechanical properties with ${\mathsf{\sigma }}_{\mathrm{Y}}=1100 \mathrm{M}\mathrm{P}\mathrm{a}$, ${\mathsf{\sigma }}_{\mathrm{f}}=2530 \mathrm{M}\mathrm{P}\mathrm{a}$, and ${\mathsf{\epsilon }}_{\mathrm{f}}=34\%$, as do Fe-rich alloys such as Fe₃₅Co₂₀Cr₁₇Ni₁₂Al₁₂Ti₄ [23], which achieves ${\mathsf{\sigma }}_{\mathrm{Y}}=1420 \mathrm{M}\mathrm{P}\mathrm{a}$, ${\mathsf{\sigma }}_{\mathrm{f}}=2490 \mathrm{M}\mathrm{P}\mathrm{a}$, and ${\mathsf{\epsilon }}_{\mathrm{f}}=34\%$. On the other hand, excessively high Ti contents can lead to the formation of Laves phases, as in AlCoCrFeNiTi₁.₅ [24], where the high strength of these phases increases yield strength but significantly reduces toughness. Eutectic compositions have also been proposed [25,26,27]. In Fe-rich systems, these alloys often show high yield strength, high compressive strength, and relatively high compressive plasticity, as demonstrated by the AlCrFe₂NiTi₀.₂ alloy [25], with mechanical properties in the as-cast state reaching 1499 MPa, 3353 MPa, and 33.9%, respectively.

Based on the results of previous research [28], a new Co-free alloy system Al_xTi_xCrFe₂Ni with increased iron content was designed to develop cost-effective HEAs exhibiting enhanced mechanical properties. The influence of the alloying elements Al and Ti in the Al_xTi_xCrFe₂Ni system (x = 0.1–0.6) was systematically evaluated with respect to their effect on phase formation, crystal structure, microhardness, and compressive strength through a direct comparison between as-cast samples and those subjected to homogenization annealing.

2. Materials and Methods

A series of medium- to high-entropy alloys with the nominal formula Al_xTi_xCrFe₂Ni₂ (x = 0.1, 0.2, 0.3, 0.4, 0.5, and 0.6—denoted as AT1, AT2, AT3, AT4, AT5, and AT6, respectively) were synthesized as small ingots via arc melting in a vacuum arc furnace, the Mini Arc Melter MAM-1 (Edmund Bühler GmbH, Bodelshausen, Germany). Melting was performed under an argon atmosphere on a water-cooled copper hearth. To ensure chemical homogeneity, each ingot was remelted five times, with the ingot flipped between each pass. High-purity elemental feedstock (Fe, Ni, Cr, Al, Ti) with ≥99.9 wt.% purity was used. All alloys were examined in both as-cast and homogenized conditions. Homogenization was carried out at two temperatures, 1100 °C and 1150 °C, for 4 h each, followed by water quenching.

The crystal structures of the alloys were characterized using a Rigaku Rapid II X-ray diffractometer (Rigaku Holdings Corporation, Tokyo, Japan) equipped with a D-Max two-dimensional curved detector. Measurements were performed in reflection geometry, with the omega angle held constant at 15° and the phi axis oscillating at 15°/s.

For microstructural analysis, the samples were prepared in accordance with the ASTM E3 standard [29], following Struers metallographic procedures, including sectioning, mounting, grinding, polishing, and etching. Cutting was performed using a precision saw with a diamond wire to minimize deformation. The samples were then mounted using Poly-Fast conductive resin. Grinding began with Piano 220 discs, followed by fine grinding on Largo discs using a 9 μm diamond suspension. Polishing was carried out in two stages: first with a 3 μm diamond suspension, and then with a 1 μm diamond suspension to achieve a mirror-like surface. To reveal crystallographic features, the polished specimens were briefly etched (1–5 s) in marble etchant solution (20 mL HCl + 4 g CuSO₄ + 20 mL methanol). For EBSD measurements, the sample surfaces were prepared using oxide polishing with a colloidal silica (SiO₂) suspension.

Thermal analysis of the experimental alloys was performed using simultaneous thermal analysis (STA) based on differential thermal analysis (DTA) (STA-449-F3 Jupiter-NETZSCH, Germany) at a heating rate of 10 K/min up to 1450 °C. Thermal analysis was performed on as-cast samples with a mass of 150 mg, placed in alumina crucibles under a high-purity argon atmosphere.

The structural characteristics of the alloys were examined using an optical microscope ZEISS AXIO VERT A1 (Carl Zeiss, Oberkochen, Germany), while microstructural details were observed at high resolution with a scanning electron microscope (SEM), TESCAN MIRA 3 LMU (TESCAN, Brno, Czech Republic). The chemical composition of individual structural components was determined using energy-dispersive spectroscopy (EDS) (Oxford Instruments X-act; Oxford Instruments, Oxford, UK), with measurements taken in at least three different regions per sample. Phase maps of the selected microstructures were obtained via EBSD using an Oxford NordlysMax² detector (Oxford Instruments, Abingdon, UK). The volume fraction of each phase was quantified by digital image analysis performed over more than five areas at different magnifications.

The mechanical behavior was evaluated by conducting room-temperature compression tests using a Tinius Olsen H300KU universal testing machine (Tinius Olsen, Horsham, PA, USA). Each alloy was tested in both as-cast and homogenized conditions with a minimum of two specimens per condition. The tests were conducted using cylindrical samples with dimensions of Φ 4 × 8 mm, and each specimen was subjected to a constant strain rate of 1 × 10⁻³ s⁻¹. The microhardness of each structural component was measured on the etched sample surfaces using a Struers Duramin-5 Vickers tester (Struers, Copenhagen, Denmark) in accordance with ISO 6507-1 [30]. At least 10 indents per phase were made under a 981 mN (HV 0.1) load for 5 s. The distance between adjacent indents was kept at a minimum of three times the diagonal length of the indent to avoid interaction effects. From the measured values, the average microhardness and standard deviation were calculated for each phase.

3. Results

3.1. Characterization of Structural and Mechanical Properties of Al_xTi_xCrFe₂Ni Alloys in As-Cast State

The XRD patterns from the as-cast Al_xTi_xCrFe₂Ni alloys are displayed in Figure 1, where it is evident that the phase composition gradually changes depending on the chemical composition of the melts.

Table 1. Lattice constants of as-cast Al_xTi_xCrFe₂Ni alloys.

Microstructure	Lattice Constant (Å)
Samples	AT1	AT2	AT3	AT4	AT5	AT6
FCC	3.604 ± 2 × 10−3	3.614 ± 4 × 10−3	3.613 ± 2 × 10−3	-	-	-
BCC1	-	2.877 ± 2 × 10−3	2.878 ± 4 × 10−3	2.879 ± 3 × 10−3	2.882 ± 3 × 10−3	2.884 ± 4 × 10−3
BCC2	-	2.942 ± 4 × 10−3	2.937 ± 2 × 10−3	2.939 ± 3 × 10−3	2.943 ± 3 × 10−3	2.941 ± 5 × 10−3

shows the variation in lattice parameters resulting from the changing chemical compositions of alloys AT1–AT6. In the AT1 alloy, which has the lowest Al and Ti content, the structure consists of a single FCC phase, with characteristic diffraction peaks observed in multiple patterns at specific 2θ positions. Specifically, these are diffraction peaks corresponding to (111) (near 19°), (200) (near 22.5°), and (220) (near 32.5°). Additional reflections observed at higher angles such as (222), (400), (331), and (420) further confirm the face-centered cubic lattice. Based on these peaks, the lattice parameter of the FCC phase was calculated to be 3.604 Å. The distinct formation of the BCC1 phase can be seen in the AT2 alloy, where its dominant peaks are identified for the (110) (near 19°), (200) (near 28°), (211) (near 35°), and (220) (near 41°) planes. Additional reflections corresponding to the (310), (222), and (321) planes further support the identification of the body-centered cubic lattice. The lattice parameter of the BCC1 phase, calculated from these reflections, was found to be 2.877 Å. In this alloy, there is even evidence of a second BCC phase (BCC2), which is identified at the same indices (211) and (220), but whose reflections slightly shift toward lower angles due to a larger lattice parameter (see Table 1). Its presence is clearly visible in the AT3 alloy, where it is also evident at (110) and (200). The lattice parameter of the FCC phase exhibits a slight increase as the Al and Ti content increases (see Table 1). In the AT4 alloy, a significant reduction in the FCC phase can be observed, which only exhibits very weak reflections in the diffraction peaks corresponding to (200) and (220). Complete disappearance of the FCC phase is observed in the AT5 alloy, whose structure is composed solely of the BCC1 and BCC2 phases. A similar trend is seen in the AT6 alloy.

Figure 2 presents the as-cast microstructures of Al_xTi_xCrFe₂Ni alloys at two scales: low-magnification images obtained via optical microscopy and high-resolution views captured using SEM.

Table 2. Chemical compositions of different regions in as-cast Al_xTi_xCrFe₂Ni alloys obtained from EDS (at. %).

Alloys	Region		Chemical Composition (at. %)
		Al	Ti	Cr	Fe	Ni
AT1	Nominal	2.4	2.4	23.8	47.6	23.8
	Actual	2.9	2.4	24.1	47.2	23.4
	DR-FCC	2.4	1.5	24.4	50.0	21.7
	ID-BCC1	2.0	12.3	27.6	37.6	20.5
AT2	Nominal	4.6	4.6	22.7	45.5	22.7
	Actual	4.9	4.9	22.3	44.2	23.7
	DR-FCC	5.1	3.8	21.7	45.5	23.9
	ID-BCC1	7.9	6.2	25.8	40.9	19.3
	ID-precipitate	14.2	10.2	19.2	29.7	26.6
AT3	Nominal	6.5	6.5	21.7	43.5	21.7
	Actual	7.2	7.5	19.9	41.4	24.0
	DR-BCC2	6.7	12.2	13.8	35.0	32.2
	ID-BCC1	6.2	5.4	22.7	45.0	20.6
	ID-precipitate	15.8	13.7	12.4	25.1	33.1
AT4	Nominal	8.3	8.3	20.8	41.7	20.8
	Actual	9.6	9.2	20.0	38.5	22.6
	DR-BCC2	23.0	18.1	4.7	16.2	38.0
	ID-BCC1	8.9	7.4	22.3	41.2	20.1
	ID-precipitate	14.5	14.9	11.8	25.4	33.5
AT5	Nominal	10.0	10.0	20.0	40.0	20.0
	Actual	11.3	11.1	19.3	37.4	20.9
	DR-BCC2	22.9	18.4	5.1	17.3	36.3
	ID-BCC1	9.3	7.2	23.8	44.0	15.7
	ID-precipitate	18.2	16.2	10.2	23.6	31.9
AT6	Nominal	11.5	11.5	19.2	38.5	19.2
	Actual	13.3	12.4	18.3	36.4	19.6
	DR-BCC2	23.3	18.4	5.9	19.7	32.8
	ID-BCC1	11.6	9.5	21.9	40.4	16.6
	ID-precipitate	21.2	15.9	9.4	23.0	30.5

summarizes the chemical compositions of the individual structural phases and regions obtained through EDS measurements, while the microhardness values of these regions are shown in Figure 3. The AT1 alloy exhibits a single-phase homogeneous FCC matrix containing finely dispersed precipitates of a second phase (BCC1), precipitated within the columnar interdendritic (ID) regions. The ID area primarily forms due to the presence of Ti (12.3 at.%), while the Al content remains relatively low (2.2 at.%). The overall volume fraction of this phase is less than 3%, and therefore, it was not detected by XRD measurements. The presence of the FCC phase is further supported by its low hardness, approximately 265 HV0.1 (see Figure 3). In the AT2 alloy, the volume fraction of the ID region increases to approximately 20%. According to XRD measurements, this is the BCC1 phase, characterized by small globular precipitates enriched in Al and Ti, with sizes ranging from 160 to 200 nm, as documented in Figure 2d. A similar morphology was previously observed in the Al_0.5CrFe₂Ni₂ alloy, which was identified via selected area electron diffraction (SAED) in transmission electron microscopy (TEM) as a mixture of BCC + B2 phases [31].

A significant difference in microstructural evolution is observed in the AT3 alloy, which crystallizes into an equiaxed dendritic structure composed of two dominant BCC phases: a primary BCC1 phase labeled ID (dark regions in Figure 2) and a secondary BCC2 phase labeled DR (light regions in Figure 2), which precipitates preferentially along the boundaries of ID. Its volume fraction reaches 13%. The presence of both BCC phases was confirmed by XRD analysis (Figure 1). The ID phase is characterized by a high Fe content (45 at.%) and a slightly reduced content of Al and Ti compared to the nominal alloy composition. Within the ID phase, fine globular precipitates measuring 200–300 nm are detected, enriched in Al, Ti, and Ni, and depleted in Cr and Fe. The hardness of the ID phase increases to 425 HV0.1. The DR phase crystallizes as a single-phase structural constituent with a relatively high content of Al (15.8 at.%), Ti (13.7 at.%), and Ni (33.1 at.%) and a lower content of Cr (12.4 at.%) and Fe (25.1 at.%). This phase also exhibits a higher hardness of 448 HV0.1 compared to BCC1. A pronounced change in the chemical composition of the DR phase is identified in the AT4 alloy, where the content of Al (23 at.%), Ti (18.1 at.%), and Ni (38 at.%) further increases, while Cr (4.7 at.%) and Fe (16.2 at.%) decrease, resulting in an increased hardness of 499 HV0.1. The volume fraction of the DR phase is reduced to 7%. The hardness of the ID phase also increases, reaching 486 HV0.1. In alloys AT5 and AT6, the content of Al and Ti in the DR phase remains relatively stable, while a slight increase in Cr and Fe at the expense of Ni likely contributes to a further increase in hardness, reaching 589 HV0.1 and 686 HV0.1, respectively. With increasing Al and Ti content in the alloys, the volume fraction of the DR phase increases to 20% and 38%, respectively. Additionally, a slight enrichment in Al and Ti is observed within the precipitates in the ID phase. A similar morphology has been reported in Al_xCrFeNiTi_0.25 and AlCrFeNi₂Ti_0.5 alloys [32,33] after homogenization annealing at 1000 °C for 6 h followed by furnace cooling.

Figure 4 presents the DTA traces for alloys AT1–AT6 recorded from room temperature up to 1450 °C, with all characteristic temperatures indicated. Each alloy exhibits a clear endothermic peak at the high-temperature end of the curve, corresponding to its melting event. The onset melting temperature for each alloy is determined by the intersection of tangent lines drawn at the first inflection point of the peak. Alloy AT1 has the highest onset melting temperature (solidus) at 1312 °C. Melting occurs within a relatively narrow temperature interval (ΔT = 80 °C), which is characteristic of the melting of a single phase; this observation is in agreement with the XRD measurements. Additionally, a small endothermic peak is observed in this alloy at 1213 °C, representing the melting of the interdendritic (ID) phase, whose presence was confirmed by metallographic analysis, indicating a volume fraction of approximately 3%. The melting temperature of alloy AT2 shifts significantly to a lower value of 1235 °C. This change results from the coexistence of two phases: the original primary phase (FCC), showing a melting peak at 1345 °C, and the secondary BCC1 phase, which melts over a relatively broad temperature range. In alloy AT3, a higher proportion of the BCC1 phase is evident, with the onset of melting further decreasing to 1191 °C, a structural composition also confirmed by metallographic analysis. Distinct melting behavior is observed in alloy AT4, which shows two peaks: one at 1202 °C corresponding to the BCC1 phase and another at 1230 °C, which, according to XRD analysis, corresponds to the BCC2 phase. A similar trend is seen in alloy AT5, where the second peak (1242 °C) is more pronounced and the end of melting shifts towards lower temperatures, indicating the disappearance of the FCC phase, as revealed by XRD. Alloy AT6 exhibits the lowest melting temperature, recorded at 1182 °C, with the BCC2 phase (peak at 1239 °C) becoming predominant.

To investigate the effects of alloying additions, compression tests were conducted on the studied HEAs containing different amounts of Al and Ti. Figure 5 presents the compressive engineering stress–strain profiles for the as-cast alloys, with their key mechanical parameters summarized in

Table 3. Characteristic mechanical properties of Al_xTi_xCrFe₂Ni alloys in as-cast condition.

	AT1	AT2	AT3	AT4	AT5	AT6
${\sigma }_Y\left[\mathrm{M}\mathrm{P}\mathrm{a}\right]$	411	586	1207	1313	1450	1570
${{\sigma }_{max};\sigma }_f\left[\mathrm{M}\mathrm{P}\mathrm{a}\right]$	2044	2598	3240	2895	2291	2011
${\epsilon }_p;{\epsilon }_f\left[\mathbf{\%}\right]$	54	51	39	31	24	9

. For some alloys (AT1–AT4), no failure or loss of specimen integrity was observed during compression due to their high plasticity. Therefore, their strength characteristics are reported as the compressive maximum strength $\left({\mathsf{\sigma }}_{\mathrm{m}\mathrm{a}\mathrm{x}}\right)$ , corresponding to the uniform compressive plastic strain ${(\mathsf{\epsilon }}_{\mathrm{p}})$. This point is defined as the moment when the engineering stress begins to increase again, indicating a reversal in the slope of the stress increment, caused by the onset of non-uniform deformation. The deformation values ${\mathsf{\epsilon }}_p$ and ${\mathsf{\epsilon }}_{\mathrm{f}}$ were determined from the specimen dimensions after testing.

Alloy AT1 exhibits pronounced strain hardening, with plastic deformation starting at ${\mathsf{\sigma }}_{\mathrm{Y}}=411 \mathrm{M}\mathrm{P}\mathrm{a}$ and reaching a strength ${\mathsf{\sigma }}_{\mathrm{m}\mathrm{a}\mathrm{x}}=2044 \mathrm{M}\mathrm{P}\mathrm{a}$ at a plastic strain ${\mathsf{\epsilon }}_{\mathrm{p}}=54\%$. Alloy AT2 shows similar behavior, with a slightly higher yield strength and maximum strength. Alloy AT3 exhibits a substantially higher yield strength (YS = 1208 MPa) and achieves the highest strength (${\mathsf{\sigma }}_{\mathrm{m}\mathrm{a}\mathrm{x}}=3240 \mathrm{M}\mathrm{P}\mathrm{a}$) among all as-cast states. It is also observed that the elastic modulus increases progressively from alloys AT3 to AT6. A further increase in Al and Ti content in alloys AT4, AT5, and AT6 results in a gradual rise in yield strength, reaching 1313 MPa, 1450 MPa, and 1570 MPa, respectively; however, this is accompanied by a significant reduction in overall strength and ductility (see Table 3). In the case of alloy AT5, a ductile fracture occurred along the direction of maximum shear stress, following the diagonal of the compressed cylinder, while alloy AT6 fractured in a brittle manner into multiple fragments.

3.2. Characterization of Structural and Mechanical Properties of Al_xTi_xCrFe₂Ni Alloys After Heat Treatment

Homogenization annealing of HEAs is carried out to eliminate solidification-induced segregations and is most commonly carried out at temperatures ranging from 1100 °C to 1300 °C [34,35,36,37,38], with the specific temperature depending on the chemical composition and the desired grain size. Based on DTA results, two homogenization temperatures were selected: 1100 °C, which is commonly used for HEA alloys, and 1150 °C, which is approximately 30 °C below the lowest solidus temperature. These conditions were designated as HT1 and HT2, respectively. The selected temperatures also allowed for the investigation of the thermal stability of the alloy structures.

The XRD patterns of the homogenized Al_xTi_xCrFe₂Ni alloys are presented in Figure 6, while

Table 4. Lattice parameters of Al_xTi_xCrFe₂Ni alloys after homogenization annealing, HT1 and HT2.

Microstructure			Lattice Constant (Å)
	Samples	AT1	AT2	AT3	AT4	AT5	AT6
HT1	FCC	3.606 ± 4 × 10−3	3.625 ± 3 × 10−3	3.620 ± 3 × 10−3	-	-	-
	BCC1	-	2.879 ± 3 × 10−3	2.877 ± 2 × 10−3	2.878 ± 2 × 10−3	2.882 ± 3 × 10−3	2.882 ± 3 × 10−3
	BCC2	-	2.932 ± 2 × 10−3	2.932 ± 2 × 10−3	2.934 ± 3 × 10−3	2.937 ± 2 × 10−3	2.937 ± 3 × 10−3
HT2	FCC	3.608 ± 2 × 10−3	3.621 ± 4 × 10−3	3.613 ± 2 × 10−3	-	-	-
	BCC1	-	2.880 ± 3 × 10−3	2.876 ± 4 × 10−3	2.878 ± 3 × 10−3	2.882 ± 2 × 10−3	2.886 ± 2 × 10−3
	BCC2	-	2.938 ± 2 × 10−3	2.929 ± 5 × 10−3	2.930 ± 3 × 10−3	2.936 ± 3 × 10−3	2.937 ± 2 × 10−3

summarizes the variation in lattice parameters of the identified structural phases. Compared to the as-cast state, the AT1 alloy retains an FCC structure at both annealing temperatures. The AT2 alloy exhibits the presence of the BCC2 phase on the (200), (211), and (311) planes after HT1. Higher intensities of the FCC phase are observed in the AT3 alloy after HT1, while in the case of HT2, the diffraction pattern is comparable to the as-cast state. A slight indication of the presence of the FCC phase is detected in the AT4 and AT5 alloys after HT2. The phase composition of the AT6 alloy does not significantly change after either heat treatment compared to the as-cast condition.

Figure 7 shows the microstructural evolution after homogenization treatment under HT1 conditions, while the detailed chemical compositions of each identified phase and region are presented in

Table 5. Chemical compositions of different regions of Al_xTi_xCrFe₂Ni alloys after HT1 and HT2 obtained from EDS (at. %).

Alloys	Region	Chemical Composition (at. %)
		Al		Ti		Cr		Fe		Ni
		HT1	HT2	HT1	HT2	HT1	HT2	HT1	HT2	HT1	HT2
AT1	Actual	2.7	2.7	2.5	2.5	23.8	23.7	47.6	47.5	23.4	23.6
AT2	Actual	5.0	5.2	4.6	4.9	22.4	22.2	45.8	44.9	22.3	22.8
	DR-FCC	4.8	5.1	4.6	4.3	22.0	21.7	45.3	45.0	23.4	24.0
	ID-BCC1	5.8	5.7	5.2	5.3	26.8	24.6	44.2	43.7	18.0	20.7
	ID-precipitate	9.8	4.0	6.4	4.0	24.3	27.2	37.9	47.8	21.5	17.0
AT3	Actual	8.0	6.9	7.1	6.8	21.7	21.6	42.5	42.5	20.8	22.2
	DR-BCC2	7.2	5.1	5.8	13.7	18.3	13.8	42.4	36.2	26.3	31.2
	ID-BCC1	7.2	6.7	7.4	4.9	21.8	23.5	41.9	45.3	21.7	19.6
	ID-precipitate	13.6	6.8	9.5	5.6	17.0	22.5	32.7	45.2	27.2	20.0
AT4	Actual	9.5	4.4	8.3	7.3	20.6	20.4	40.2	44.4	21.5	23.6
	DR-BCC2	23.6	1.7	19.0	9.4	3.4	14.7	13.6	43.0	40.5	31.2
	ID-BCC1	8.4	4.1	7.1	5.5	21.6	22.9	42.5	46.9	20.5	20.6
	ID-precipitate	6.5	3.3	5.9	6.2	23.4	22.1	45.7	47.5	18.5	20.8
AT5	Actual	10.8	5.6	9.9	9.4	18.9	19.4	38.6	42.1	21.8	23.6
	DR-BCC2	23.9	1.7	19.7	15.1	3.5	13.8	14.3	41.9	38.5	27.6
	ID-BCC1	7.7	5.6	7.7	6.5	22.6	23.1	44.1	45.3	18.0	19.6
	ID-precipitate	13.3	7.4	10.3	9.3	17.3	17.9	34.2	37.6	24.9	27.8
AT6	Actual	12.2	6.5	12.3	11.5	18.2	18.9	36.3	39.8	21.1	23.3
	DR-BCC2	23.8	1.8	21.0	16.7	3.5	14.5	15.5	44.5	36.3	22.5
	ID-BCC1	7.3	6.8	7.4	7.6	25.1	23.0	47.5	43.6	12.7	19.0
	ID-precipitate	10.3	7.5	9.3	13.3	21.2	15.6	41.0	36.7	18.2	27.0

. The AT1+HT1 alloy exhibits a homogeneous single-phase microstructure in which all interdendritic structural components observed in the as-cast state have been dissolved, while the hardness remains unchanged (see Figure 3). In the AT2+HT1 alloy, spheroidization of the interdendritic regions in the as-cast state was observed, with their size ranging from 15 to 30 μm. The treatment led to a reduction in the hardness of both the matrix (DR) and the secondary phase (ID). The phase composition of this alloy was further confirmed by EBSD analysis, in which the red color represents the FCC matrix and the blue color indicates the BCC phase in Figure 8. The IPF map reveals that the matrix exhibits a preferred crystallographic orientation along <111>, whereas the ID regions are randomly oriented. Pronounced homogenization of the microstructure is evident in the AT3 alloy, where spheroidization of the ID phase occurred along with the formation of polyhedral grains with diameters ranging from 50 to 80 μm. The hardness of these grains increased by 80 HV0.1 compared to the as-cast state. An increase in the number of precipitates and a more uniform size distribution was also observed. The dendritic (DR) phase is located along the boundaries of the ID phases, mostly at triple junctions, appearing as islands that tend to coalesce. Its hardness markedly decreased to 302 HV0.1 compared to the as-cast condition. In the AT4 alloy, partial spheroidization was also observed, along with an increase in the size of ID grains to approximately 48 mm.

The annealed microstructures of the AT5 and AT6 alloys retained a dendritic morphology, with the DR phases becoming more uniformly distributed. Their hardness decreased to 570 HV0.1 and 586 HV0.1, respectively. As a result of homogenization annealing, the chemical compositions of the BCC1 and BCC2 phases in alloys AT4, AT5, and AT6 became more homogeneous, reaching approximately the same average concentrations of the main elements Al, Ti, and Cr. Specifically, the BCC1 phase contains Al (7.8 at.%), Ti (7.4 at.%), and Cr (3.4 at.%), while the BCC2 phase contains Al (23.7 at.%), Ti (19.8 at.%), and Cr (3.5 at.%). At the same time, the Fe content slightly increased, while the Ni content slightly decreased.

The microstructural evolution following HT2 treatment is shown in Figure 9, and Table 5 summarizes the chemical compositions of the various structural phases and regions. After homogenization annealing at 1150 °C (HT2), the alloys AT1-HT2 and AT2-HT2 exhibit a similar structural morphology to that observed after HT1. In the case of AT2-HT2, the ID phase shows a larger grain size, ranging from 30 to 40 μm, compared to AT1-HT1. Structural instability due to the high annealing temperature (1150 °C) was observed in alloys AT3-HT2 to AT6-HT2, where abnormal growth of ID grains occurred, characterized by spherical islands within the grains and the precipitation of additional phases along the grain boundaries. In alloys AT4-HT2 to AT6-HT2, significant depletion of Al was detected, with the Al content decreasing to approximately half of the nominal composition. These alloys also show pronounced redistribution of chemical composition; the ID regions in the as-cast state are rich in Al, Ti, and Ni, the grain interiors are now enriched in Cr, Fe, and Ni, and the precipitates are enriched in Ti, Fe, and Ni. Their hardness gradually increased to 519 HV0.1, 536 HV0.1, and 605 HV0.1, respectively. Along the grain boundaries and at triple junctions (designated as DR), phases with high hardness were formed reaching values of 660 HV0.1, 728 HV0.1, and 750 HV0.1, respectively. In these regions, Al was almost completely depleted, with its content reduced to as low as 1.7 at.%, while Ti, Fe, and Ni were strongly enriched.

The phase composition determined by EBSD analysis (Figure 8) demonstrates that the grains of alloy AT5-HT2 consist of a mixture of multiple phases, including BCC1, BCC2, B2 (NiAl), and L1₂ (Ni₃Al). The identification of these phases in the EBSD analysis was based on their lattice parameters. Along the grain boundaries and at triple junctions, there is an FCC phase (Ni-based superalloy) which was also identified by X-ray analysis. Detailed identification of these phases requires transmission electron diffraction.

Figure 10 compares the compressive engineering stress–strain curves obtained under different heat treatment conditions (HT1 and HT2). For reference, the stress–strain curves of the as-cast condition, taken from Figure 5, are also included to allow for direct comparison of the different processing states (the as-cast curves are shown as a thin dotted line). The corresponding key mechanical properties are summarized in

Table 6. Characteristic mechanical properties of Al_xTi_xCrFe₂Ni alloys after homogenization annealing under HT1 and HT2 conditions.

		AT1	AT2	AT3	AT4	AT5	AT6
HT1	${\mathit{\sigma }}_{\mathit{Y}}\left[\mathrm{M}\mathrm{P}\mathrm{a}\right]$	368	532	1510	1471	1641	1665
	${{\mathit{\sigma }}_{\mathit{m}\mathit{a}\mathit{x}};\mathit{\sigma }}_{\mathit{f}}\left[\mathrm{M}\mathrm{P}\mathrm{a}\right]$	2059	2466	3316	2997	2700	3120
	${\mathit{\epsilon }}_{\mathit{p}};{\mathit{\epsilon }}_{\mathit{f}}\left[\mathit{\%}\right]$	62	53	45	27	30	25
HT2	${\mathit{\sigma }}_{\mathit{Y}}\left[\mathrm{M}\mathrm{P}\mathrm{a}\right]$	385	620	1440	1572
	${{\mathit{\sigma }}_{\mathit{m}\mathit{a}\mathit{x}};\mathit{\sigma }}_{\mathit{f}}\left[\mathrm{M}\mathrm{P}\mathrm{a}\right]$	2170	2500	3148	2079
	${\mathit{\epsilon }}_{\mathit{p}};{\mathit{\epsilon }}_{\mathit{f}}\left[\mathit{\%}\right]$	60	48	36	15

. The AT1 alloy retains high plasticity after both homogenization annealing treatments (HT1 and HT2), with a slight increase in both strength and plasticity compared to the as-cast state. A reduction in uniform compressive plastic strain due to HT2 annealing can be observed starting from the AT2-HT2 alloy. The most beneficial effect of homogenization annealing was observed under HT1 conditions in the AT3-HT1 alloy, where the strength increased compared to the as-cast state, reaching ${\sigma }_{max}=3316\mathrm{M}\mathrm{P}\mathrm{a}$ at a plastic strain ${\epsilon }_p=45\%$. HT1 treatment also had a positive effect on the AT4-HT1, AT5-HT1, and AT6-HT1 alloys, with a gradual increase in both strength and plastic strain. In the case of AT6, the strength increased to ${\sigma }_f=3120\mathrm{M}\mathrm{P}\mathrm{a}$ and the plasticity to ${\epsilon }_p=25\%$. In contrast, HT2 treatment led to a deterioration of the mechanical properties, as the strength of the AT4-HT2 alloy decreased to ${\sigma }_f=2079 \mathrm{M}\mathrm{P}\mathrm{a}$ from the original as-cast value of ${\sigma }_f=2895\mathrm{M}\mathrm{P}\mathrm{a}$. For the AT5-HT2 and AT6-HT2 alloys, due to their high brittleness, it was not possible to produce test specimens for compression testing; therefore, their stress–strain curves are not presented.

4. Discussion

4.1. Effect of Chemical Composition on Microstructural Evolution

The chemical composition and the thermodynamic solidification conditions determine the resulting phase composition of HEAs. Elements with a higher concentration of valence electrons, such as Fe, Ni, Co, and Mn, primarily stabilize the FCC phase [39,40], while elements with a lower concentration of valence electrons, such as Al, Ti, and Cr, promote and stabilize the BCC phase [41,42]. The key thermodynamic parameters used to characterize the investigated HEAs, most commonly applied to predict structural stability and phase formation, are summarized in

Table 7. The thermodynamic parameters of the investigated as-cast alloys.

Alloys	δ %	ΔHmix, kJ·mol−1	ΔSmix, J·mol−1·K−1	Ω, -	VEC -	Tm °C
AT1	3.62	−6.20	1.23 R	3.04	7.75	1312
AT2	4.77	−8.78	1.33 R	2.30	7.59	1235
AT3	5.67	−11.58	1.41 R	1.83	7.42	1191
AT4	6.19	−13.48	1.47 R	1.62	7.20	1191
AT5	6.58	−15.03	1.50 R	1.47	7.02	1187
AT6	6.85	−16.34	1.53 R	1.36	6.86	1182

. These include values calculated using empirical rules for atomic size difference (δ), valence electron concentration (VEC), mixing enthalpy (ΔH_mix), and the Ω parameter, which combines mixing enthalpy, mixing entropy (ΔS_mix), and melting temperature (T_m).

According to Guo et al. [42,43], FCC phases are stable when the VEC > 8, whereas FCC and BCC phases coexist in alloys with VEC values ranging from 6.87 to 8. This prediction is in full agreement with the results for our investigated alloys, which, starting from the lowest concentrations of Al and Ti, exhibit the presence of both FCC and BCC phases. At higher Al and Ti concentrations, as in the case of AT6, two distinct BCC phases are even observed (see Figure 1 and Figure 3). Figure 11 compares the VEC values of the HEAs investigated in this study, Al_xTi_xCrFe₂Ni, which are Fe-based, with a similar alloy system, Al_xTi_xCrFe₂Ni₂, richer in Fe and Ni, studied in our previous work [28]. From the comparison of VEC values and microstructures, it is evident that the Al_xTi_xCrFe₂Ni alloy system contributes more strongly to the formation of BCC phases. For example, the AT3 alloy (Al_0.3Ti_0.3CrFe₂Ni) exhibits an even higher BCC phase fraction than the Al_0.6Ti_0.6CrFe₂Ni₂ alloy [28]. Several studies [44,45,46,47] have demonstrated that this type of Fe-rich alloy consists of a disordered BCC (A2) matrix rich in Fe and Cr, containing uniformly distributed ultrafine BCC (B2) precipitates with high hardness. Further increases in Al and Ti content promote the development of the BCC structure. According to the solid solution stability assessment by the authors [48,49], the AT5 and AT6 alloys exceed the recommended limits for the atomic size difference (1% < δ < 6.6%) and for mixing enthalpy (−15 kJ·mol⁻¹ ≤ ΔH_mix ≤ 5 kJ·mol⁻¹), indicating a strong tendency toward phase segregation or intermetallic formation. This prediction is consistent with the obtained microstructures of the investigated alloys, as confirmed not only by structural analyses but also by X-ray diffraction measurements.

4.2. Thermal Stability of Investigated Alloys

Thermal analysis was used to measure the solidus temperatures of the investigated alloys, as summarized in Table 7. As can be seen, with increasing Al and Ti content, the solidus temperature shifts to lower values due to changes in phase composition, decreasing from 1312 °C to 1182 °C, which represents a reduction of 130 °C. Several authors [50,51] have confirmed that in AlxCoCrFeNi-type alloys, as the Al content increases from 0 at.% to 0.7 at.%, the solidus temperature decreases from 1690 K to 1621 K, but with further increases in Al content from 0.7 at.% to 1.8 at.%, the solidus temperature rises again to 1680 K. Other authors working with Ni₅₁Co₁₈Fe₅Cr₁₀Al_16–xTi_x and AlCrFeNiTi_x alloys [52] have shown that Ti content also has a significant effect on reducing the solidus temperature. In the alloys investigated in this study (AT3–AT6), the solidus temperatures were considerably lower than those of the Ni-based alloys reported in [52]. This reduction is likely caused by the combined effect of Al, Ti, and Fe.

Homogenization annealing in the HT1 regime led to several beneficial microstructural modifications, including spheroidization of interdendritic regions without a significant change in their size, a more uniform distribution of the BCC2 phase along the ID boundaries, an increased number and more homogeneous distribution of precipitates within the ID phase, and the absence of low-melting phase segregation. These changes contributed to improved strength and plastic strain of the alloys. The results also indicate that the studied HEAs after HT1 exhibit outstanding thermal stability during homogenization annealing at temperatures up to 0.9T_m, which significantly exceeds the typical values applied in conventional steels. This thermal resistance is comparable to values characteristic of high-performance superalloys, such as INCONEL 718 [53].

Heat treatment under the HT2 regime had a similar effect on alloys with higher solidus temperatures (AT1 and AT2) as those observed under the HT1 regime, without significant degradation of the microstructure. In contrast, for alloys AT3 through AT6, pronounced grain growth of ID regions and precipitation of hard phases along grain boundaries were observed. These phenomena indicate localized overheating of the material during annealing, leading to a disruption of phase equilibrium and partial degradation of the characteristic HEA microstructure.

4.3. Mechanical Properties of Investigated Alloys

The mechanical properties of the analyzed alloys are strongly influenced by their chemical composition and structural state. Alloys that contain a combination of FCC and BCC phases in the as-cast condition (AT1 and AT2) are characterized by lower hardness, yield strength, and elastic modulus, while retaining high plastic strain. In contrast, alloys with higher Al and Ti contents (AT3–AT6) are predominantly composed of two BCC phases (BCC1 and BCC2), which leads to a gradual increase in yield strength up to 1570 MPa. This effect is most pronounced in the AT3 alloy, which exhibits the highest strength within this group.

Homogenization annealing at 1100 °C (HT1) had a beneficial effect on the mechanical properties of all investigated alloys. For the AT1-HT1 and AT2-HT1 alloys, a reduction in strengthening was observed compared to the as-cast condition, which can be attributed to the chemical homogenization of the FCC phase. This process reduces local segregation of alloying elements, resulting in a more uniform solid solution and, consequently, a lower resistance to dislocation movement.

In the AT3-HT1 alloy, a significant increase in yield strength along with improved plasticity was observed. These enhancements are attributed to the homogenization of interdendritic (ID) regions, which, after heat treatment, exhibit a more uniform microstructure with a higher number of finer precipitates compared to the as-cast state. As a result, the AT3 alloy demonstrated the best overall mechanical performance, achieving a maximum compressive strength of 3316 MPa at a 45% plastic strain.

The most pronounced improvement in mechanical behavior, especially in terms of plasticity, was found in the AT5-HT1 and AT6-HT1 alloys. This improvement is directly linked to the microstructural evolution of the BCC2 phase, which undergoes noticeable softening after homogenization compared to the as-cast condition. Consequently, the material exhibits enhanced plastic deformability.

The negative effect of high-temperature annealing of alloys composed of two BCC phases is attributed to abnormal grain growth and phase segregation. During homogenization, hard phases (with microhardness values of 660–750 HV0.1) tend to segregate along grain boundaries, leading to the formation of brittle, hard zones at interfaces. This microstructural evolution significantly impairs mechanical performance by promoting crack initiation and reducing the alloy’s ability to undergo plastic deformation.

Figure 12 compares the mechanical properties of the HEAs developed in this work with similar alloys alloyed with aluminum and titanium. From the graph, it is evident that with increasing Al and Ti content in the alloys, the yield strength (YS) also increases.

The compared alloys can be categorized into three groups based on the concentration of Al and Ti, enabling a clearer analysis of the relationship between alloying element content and YS (Figure 12a):

• Alloys containing up to 10 at.% (Al + Ti) exhibit a YS below 1000 MPa;
• Alloys with 10–20 at.% show a YS in the range of 1000–1500 MPa;
• For concentrations above 20 at.%, the YS ranges between 1500 and 2000 MPa, and in some cases even higher.

Most of the BCC-structured alloys in this study, following HT1 annealing, exhibit YS values exceeding 1500 MPa, as indicated by the purple triangles. Figure 12b presents the engineering stress–strain curves, showing that alloys with a high YS generally exhibit reduced plasticity, which limits their compressive strength. However, the AT2, AT3, and AT4 alloys represent a very good compromise between strength and plasticity. Moreover, these alloys offer a cost advantage over comparable CrFeNiAl_xTi_x-type HEAs (represented by black crosses), as they contain lower amounts of expensive alloying elements such as Ti, Cr, and Ni.

5. Conclusions

In this study, six HEAs in the Al_xTi_xCrFe₂Ni system, alloyed with aluminum and titanium, were investigated. The research focused on the characterization of their structure, melting temperatures, and evaluation of their mechanical properties through microhardness measurements and compression tests, both in the as-cast state and after homogenization annealing. The main findings can be summarized as follows:

• With increasing Al and Ti content, the crystal structure evolved from a single FCC phase to a mixture of FCC and BCC phases, and at higher concentrations, to a dual-phase BCC structure (BCC1 + BCC2). The BCC1 phase is strengthened by spherical precipitates rich in Al, Ti, and Ni.
• In alloys containing both BCC1 and BCC2 phases, a relatively low solidus temperature was observed, decreasing progressively with increasing Al and Ti content—from 1191 °C for alloy AT3 to 1182 °C for alloy AT6.
• Homogenization annealing at 1100 °C resulted in enhanced mechanical properties, reflected in the increased yield strength, compressive strength, and ductility, as well as a reduction in the microhardness of individual structural constituents.
• Alloy AT3 exhibited the most balanced combination of strength and ductility, both in the as-cast state and after homogenization annealing at 1100 °C. Compared to other Al- and Ti-alloyed HEAs, it also offers lower production costs due to its relatively lower content of expensive alloying elements such as Ti, Cr, and Ni.
• Homogenization annealing at 1150 °C led to significant microstructural degradation in alloys containing both BCC1 and BCC2 phases, resulting in a marked deterioration of mechanical properties.