Effect of Tungsten Content on the Microstructure, Mechanical and Tribological Properties of AlCoCrFeNi High-Entropy Alloys

Abstract

High-entropy alloys (HEAs) have recently attracted considerable attention due to their unique combination of high strength, hardness, and corrosion and wear resistance, making them promising candidates for advanced structural and functional applications. Among these, AlCoCrFeNi-based HEAs are well known for their high hardness and good wear resistance; however, their limited tribological stability under operational conditions restricts their broader application. To address this limitation, tungsten (W) was incorporated into the AlCoCrFeNi system to enhance its mechanical and tribological performance. In this study, the microstructural, mechanical, and tribological properties of AlCoCrFeNiW_x (x = 0, 0.1, 0.25, 0.5 and 1 mol) HEAs were systematically investigated. The alloys were fabricated using the vacuum arc melting method and characterized by XRD, SEM-EDS, elemental mapping, microhardness, and wear tests. The addition of W caused a shift in the 2θ ≈ 44° (110) peak toward lower angles. While the W-free alloy exhibited Body-Centered Cubic (BCC) + B2 phases, W addition led to the formation of a new W-rich phase, and at higher W contents, a pure W phase appeared. The hardness increased from 507.11 HV₁ to 651.81 HV₁ with increasing W content. Furthermore, wear resistance improved and the coefficient of friction decreased with higher W addition. When comparing the W-free alloy to the alloy with the highest W content, the wear rate decreased by approximately 1.85 times under a 2 N load and 1.89 times under a 5 N load. These results demonstrate that W addition significantly enhances the wear resistance of AlCoCrFeNi-based HEAs by nearly twofold.

1. Introduction

High-entropy alloys (HEAs) have attracted increasing attention due to their excellent combination of strength and ductility, outstanding corrosion and wear resistance, and high-temperature oxidation resistance. These superior properties provide significant potential for engineering applications [1,2]. In the literature, hundreds of different HEA systems have been investigated, and AlCoCrFeNi is one of the first HEAs studied in this field [3,4]. AlCoCrFeNi HEAs stand out for their high wear resistance, good corrosion resistance, and high strength-to-weight ratio. These attributes make them promising candidates for components exposed to wear, such as gears, bearings, and cutting tools, as well as in sectors where lightweight materials are crucial, such as aerospace and automotive industries [5].

The interest in high-entropy alloys (HEAs) continues to grow rapidly. As a result, various production methods have been developed for HEA fabrication. HEAs can be synthesized using numerous techniques such as arc melting [6], powder metallurgy [7], solvothermal synthesis [8], mechanical alloying [9], carbothermal shock [10], spray pyrolysis [11], laser cladding [12], magnetron sputtering [13], combustion [14], microwave-assisted synthesis [15], plasma sintering [16], and electrochemical deposition [17,18]. In the arc melting technique, the local temperature generated by arc discharge can exceed 6000 °C, enabling the rapid melting of refractory elements such as W, Nb, Mo, and Ta [19]. The equiatomic AlCoCrFeNi alloy is known for possessing several advantageous mechanical properties that make it suitable for various practical applications. Although the constituent elements of the equiatomic AlCoCrFeNi HEA have different crystal structures, many studies have reported that this alloy forms solely a BCC (body-centered cubic) structure, which significantly influences its mechanical properties. However, the relatively low high-temperature wear resistance of AlCoCrFeNi limits its applicability in certain engineering fields. The addition of alloying elements can enhance mechanical properties by forming precipitation-type intermetallic phases [20].

With the addition of W to HEAs, it has been emphasized that W can form intermetallic phases such as Fe₇W₆ and Co₇W₆ with Fe and Co elements, thereby enhancing the strength of the alloy [21]. The microstructure and mechanical properties of the AlCoCrFeNi alloy vary significantly depending on the alloy fabrication method, production parameters, and the type of heat treatment applied. The AlCoCrFeNi alloy system offers an intriguing structure that exhibits different microstructures and properties depending on the Al content, and it has been reported that the addition of secondary alloying elements can further improve its microstructural and mechanical characteristics [22]. Numerous studies have incorporated secondary elements such as Ti [23,24], V [25], Zr [26], B [27], Nb [22], Mn [28], and Cu [29] into the AlCoCrFeNi high-entropy alloy. These studies have reported that elements such as Ti and Zr strengthen the matrix structure through solid solution and precipitation hardening mechanisms, thereby enhancing wear resistance. Furthermore, the addition of elements like Cu or Mn promotes the formation of FCC phases, contributing to improved ductility and toughness [5].

The addition of Nb to AlCoCrFeNi HEAs alters the original phase structure of the alloy, leading to the formation of an ordered Laves phase in addition to the solid solution phase. This results in increased yield strength, hardness, and coercive force, while simultaneously decreasing plasticity, saturation magnetization, and residual magnetization [22]. In a study involving V addition, it was reported that the alloy maintained a BCC solid solution structure and that the segregation of Al, Fe, and Ni elements was reduced, thereby enhancing microstructural homogeneity. Moreover, increasing V content was found to improve yield strength and hardness, while reducing plastic deformation [25]. In the Al_xCoCrFeB_x alloy system, the addition of B led to enhanced fracture strength, compressive ratio, and hardness values. However, these improvements were only observed when the molar ratio of B was kept at x(B) ≤ 0.02; higher contents resulted in a deterioration of these properties [27]. Due to their exceptional mechanical and physical properties, HEA systems have become the subject of extensive research in the literature. It is generally acknowledged that the addition of a sixth or subsequent element in specific proportions can endow the alloy with superior characteristics. Jiang et al. synthesized CrFeNiV_0.5W_x and CrFeNi₂V_0.5W_x alloys using vacuum arc melting and investigated the influence of W content on microstructure and mechanical properties [30]. Chang et al. reported that the addition of 3% W to a CoCrNi medium entropy alloy increased the yield strength by approximately 33%, maintained high work-hardening capability, and delayed recrystallization behavior with increasing W content [1].

Although several studies have investigated the effects of secondary alloying elements on AlCoCrFeNi HEAs, systematic research on the influence of gradual tungsten (W) addition remains very limited. In particular, the critical W content that governs phase stability, lattice distortion, and tribological behavior under room-temperature conditions has not yet been clearly identified. Therefore, this study aims to fill this gap by systematically evaluating the microstructural, mechanical, and tribological responses of AlCoCrFeNiW_x (x = 0, 0.1, 0.25, 0.5 and 1 mol) alloys synthesized via vacuum arc melting. Considering the solubility behavior of W in the AlCoCrFeNi matrix and the challenges of achieving homogeneous melting at higher W concentrations, the W content range was restricted to 0–1 mol to ensure phase stability and experimental reproducibility. The selected W ratios (x = 0–1 mol) were specifically designed to identify the critical solubility limit of W in the AlCoCrFeNi matrix and to evaluate its influence on the formation of W-rich intermetallic phases such as μ-(FeCoCr)₇W₆.

In this study, varying amounts of W were added to the equimolar AlCoCrFeNi HEA system as an alloying element. W was selected due to its ability to occupy lattice sites and increase lattice distortion energy, contribute to solid solution strengthening, and enhance wear resistance through the formation of hard intermetallic compounds such as μ-(FeCoCr)₇W₆ [31,32,33]. Several studies have clearly stated that W stabilizes the BCC crystal structure. Therefore, maintaining the BCC structure of the AlCoCrFeNi alloy while enhancing its mechanical properties and identifying the critical W content are among the primary objectives of this work. It is also considered that W can form intermetallic compounds such as Fe₇W₆ and Co₇W₆ with Fe and Co, which may further enhance wear resistance. Based on the aforementioned information, this study comprehensively investigates the effects of W addition—still scarcely explored in the literature—on the microstructural characteristics, mechanical performance, and particularly the tribological behavior of the AlCoCrFeNi HEA system. We believe it could be an alternative to lower-density alloys, such as TiV64, used in aircraft and other similar applications. While its high density may seem disadvantageous for aircraft, it has been observed that it will actually be useful in key components requiring corrosion and abrasion resistance. AlCoCrFeNiW_x HEAs were fabricated using the vacuum arc melting method. The effects of W addition on the crystal structure, microstructure, hardness, and wear properties were investigated.

2. Materials and Methods

2.1. The Synthesis of HEAs

This research aimed to determine the critical W content for the equimolar AlCoCrFeNi high-entropy alloy (HEA) system. All HEAs were synthesized using the vacuum arc melting method. By varying the W content, five different HEAs with the compositions AlCoCrFeNiW_x (x = 0, 0.1, 0.25, 0.5 and 1 mol) were produced. The physical properties of the elemental powders used in the synthesis are presented in

Table 1. Physical properties of the powders.

Powder	Supplier	Purity (%)	Size
Al	Roth	99.5	44 µm
Co	Merck	99	44 µm
Cr	Merck	99	44 µm
Fe	Alfa Aesar	99	44 µm
Ni	Merck	99	44 µm
W	Nanokar	99	44 µm

. The effects of W addition on the microstructural, mechanical, and tribological properties of the synthesized alloys were comprehensively investigated. The produced HEAs were designated as HEA-0, HEA-0.1, HEA-0.25, HEA-0.5, and HEA-1, respectively, according to their W content.

The prepared powder mixtures were blended for 30 min using a 3D Turbula mixer (ALPHIE-3, Mixer, manufactured by Hexagon Product Development Pvt Ltd. Gujarat, INDIA) to ensure homogeneous distribution. Laboratory atmospheric conditions were maintained at 23 °C with 20% relative humidity. Under the same conditions, the powders were cold-pressed into circular molds with a diameter of 20 mm under a pressure of 20 MPa. This pre-compaction step was implemented to minimize powder scattering and to improve electrical contact among the particles during the initial arc striking stage. The pressed samples were then melted in a vacuum arc melting furnace (Arc Melter AM, Edmund Bühler GmbH, Bodelshausen, Germany) under an inert argon atmosphere. Each sample was melted at least three times to achieve chemical homogeneity. The codes and chemical compositions (% molar ratio) of the produced HEAs are listed in

Table 2. Chemical compositions (% molar ratio) for AlCoCrFeNiW_x.

Alloy	ID	Al	Co	Cr	Fe	Ni	W
AlCoCrFeNi	HEA-0	20	20	20	20	20	-
AlCoCrFeNiW0.1	HEA-0.1	19.61	19.61	19.61	19.61	19.61	1.95
AlCoCrFeNiW0.25	HEA-0.25	19.05	19.05	19.05	19.05	19.05	4.75
AlCoCrFeNiW0.5	HEA-0.5	18.18	18.18	18.18	18.18	18.18	9.10
AlCoCrFeNiW1	HEA-1	16.67	16.67	16.67	16.67	16.67	16.67

. The flow chart of the synthesized HEA samples is presented in Figure 1.

2.2. Thermodynamic Aspects of AlCoCrFeNiW_x HEA

In high entropy alloys (HEAs), traditional Hume-Rothery rules are insufficient to predict the formation of unlimited solid solutions. Therefore, thermodynamic calculations supported by specific parameters are employed to estimate solid solution phase formation in HEA systems. The key parameters used in these calculations include the entropy of mixing (ΔS_mix), enthalpy of mixing (ΔH_mix), configurational entropy of mixing (ΔSconf), valence electron concentration (VEC), thermodynamic parameter (Ω), electronegativity difference (Δ_x), and atomic size difference (δ%). The equations for calculating ΔS_mix and ΔH_mix for HEA systems are presented in Equation (1) and Equation (2), respectively.

In HEAs, single-phase solid solutions are formed primarily due to the high entropy effect. However, this effect alone is not sufficient to prevent the formation of intermetallic phases. Therefore, criteria based on Hume-Rothery rules, such as the valence electron concentration (VEC), electronegativity difference (Δ_x), and atomic size difference (δ), play a critical role in the design and synthesis of HEAs. Additionally, the thermodynamic parameter Ω, which represents the ratio of mixing enthalpy to mixing entropy, is also considered significant. The relevant thermodynamic parameters Tm (melting temperature), Ω, VEC, Δ_x, and δ—are calculated using Equation (3), Equation (4), Equation (5), Equation (6), and Equation (7), respectively [2].

Here, R represents the gas constant (8.314 J/(K·mol)), n is the number of elements forming the HEA, c_i and c_j denote the atomic fractions of the ith and jth components, respectively, rᵢ is the atomic radius of each alloying element, $\overline{\mathrm{r}}$ is the average atomic radius of the alloy, and (T_m)_i refers to the melting point of the ith element [2,34,35]. The entropy of mixing (ΔS_mix) reaches its maximum value when the HEA is equimolar. The enthalpy of mixing of the solid solution can be determined based on the ordered melt model. In Equation (2), Ω_ij = 4 and ΔHAB are defined as the enthalpy of mixing values of binary alloys, calculated using the tabulated data provided by Takeuchi et al. [36]. The values were calculated using the Miedema model [37] for atomic pairs in AlCoCrFeNi and AlCoCrFeNiW_x HEAs.

$$\Delta {\mathrm{S}}_{\mathrm{m}\mathrm{i}\mathrm{x}}=-\mathrm{R}\sum _{\mathrm{i}}^{\mathrm{n}}{\mathrm{c}}_{\mathrm{i}}\mathrm{l}\mathrm{n}{\mathrm{c}}_{\mathrm{i}}$$

(1)

$$\Delta {\mathrm{H}}_{\mathrm{m}\mathrm{i}\mathrm{x}}=\sum _{\mathrm{i}=1,\mathrm{j}\ne \mathrm{i}}^{\mathrm{n}}{\mathsf{\Omega }}_{\mathrm{ij}}{\mathrm{c}}_{\mathrm{i}}{\mathrm{c}}_{\mathrm{j}}$$

(2)

$${\mathrm{T}}_{\mathrm{m}}=\sum _{\mathrm{i}=0}^{\mathrm{n}}{\mathrm{c}}_{\mathrm{i}}{\left({\mathrm{T}}_{\mathrm{m}}\right)}_{\mathrm{i}}$$

(3)

$$\mathsf{\Omega }={\displaystyle \frac{{\mathrm{T}}_{\mathrm{m}}\Delta {\mathrm{S}}_{\mathrm{mix}}}{\Delta {\mathrm{H}}_{\mathrm{mix}}}}$$

(4)

$$\mathrm{VEC}={\displaystyle \sum _{\mathrm{i}}^{\mathrm{n}}{\mathrm{c}}_{\mathrm{i}}{(\mathrm{VEC})}_{\mathrm{i}}}$$

(5)

$${\Delta }_{\mathrm{x}}=\sqrt{{\displaystyle \sum _{\mathrm{i}=1}^{\mathrm{n}}{\mathrm{c}}_{\mathrm{i}}{({\mathrm{x}}_{\mathrm{i}}-\stackrel{\_}{\mathrm{x}})}^2}}$$

(6)

$$\delta =100\sqrt{{\displaystyle \sum _{\mathrm{i}=1}^{\mathrm{n}}{\mathrm{c}}_{\mathrm{i}}{\left(1-\frac{{\mathrm{r}}_{\mathrm{i}}}{\stackrel{\_}{\mathrm{r}}}\right)}^2}}$$

(7)

Table 3. Chemical mixing enthalpies of the atomic pairs among the alloying elements (unit: kJ/ mol).

Element	Al	Fe	Co	Cr	Ni	W
Al	-	-	-	-	-	-
Fe	−11	-	-	-	-	-
Co	−19	−1	-	-	-	-
Cr	−10	−1	−4	-	-	-
Ni	−22	−2	0	−7	-	-
W	−2	0	−1	1	−3	-

shows the atom pair values between the elements Al, Fe, Co, Cr, Ni and W determined by the Miedema model for the mixing enthalpy calculation.

2.3. Material Characterization Methods

2.3.1. Material Characterization

Prior to the microstructural examinations of the HEAs, grinding, polishing, and etching procedures were applied, respectively. The grinding process was carried out using SiC abrasive papers with grit sizes of #240, #360, #400, #600, #800, #1000, and #1200, respectively. During the polishing stage, diamond suspensions with particle sizes of 6 µm, 3 µm, and 1 µm were used. Etching was performed for 5–10 s using an etchant solution known as aqua regia, consisting of 20 mL HNO₃ and 60 mL HCl. X-ray diffraction (XRD) measurements were performed using a Malvern Panalytical EMPYREAN (Malvern Panalytical Ltd. is a Spectris company, Westborough, MA, USA) equipped with CuKα radiation (λ = 0.15406 nm) operating at 45 kV and 40 mA between 20 and 100 (2θ) with a step size of 0.02°. The obtained data were evaluated using the HighScore Plus XRD analysis (HighScore Plus 4.8 (V4.8.025518), Malvern Panalytical B.V., Almelo, The Netherlands) software. The XRD phase identification and lattice parameter calculations were performed using this software with reference patterns obtained from the Crystallography Open Database (COD), which provides open access crystallographic data compatible with HighScore Plus. Through this software, the lattice constants of the 2θ ≈ 44° (110) BCC solid solution phase of the HEAs, the crystallite sizes using the Scherrer equation (Equation (8)), and the dislocation densities (Equation (9)) were calculated. The microstructure and chemical composition of the alloys were examined using a scanning electron microscope (Thermo Fisher Scientific Apreo S LoVac SEM, Thermo Fisher Scientific Inc., Waltham, MA, USA) in backscattered electron (BSE) mode. Elemental distribution maps were obtained through energy-dispersive spectroscopy (EDS) and mapping analysis.

$$\mathrm{D}={\displaystyle \frac{\mathrm{K}·\mathsf{\lambda }}{\mathsf{\beta }·\cos \mathsf{\theta }}}$$

(8)

$$\mathsf{\rho }={\displaystyle \frac{1}{{\mathrm{D}}^2}}$$

(9)

Here, D represents the crystallite size, K is the Scherrer constant (K = 0.94), λ is the X-ray wavelength, β is the full width at half maximum (FWHM) of the diffraction peak, θ is the Bragg angle, and ρ denotes the dislocation density.

2.3.2. Mechanical and Wear Properties

Vickers hardness measurements of the HEAs were conducted using a universal hardness testing machine (QATM/Qness 750 CS EVO, QATM a Verder Scientific companies, Mammelzen, Germany) under a load of 1000 g and a dwell time of 15 s. For each sample, at least ten different points were measured, and the average hardness values were calculated.

Friction and wear tests were performed under dry sliding conditions at room temperature using a tribometer (CSM Instruments TRB 18-317) (CSM Instrument, Graz, Austria). As the counter-body, a 6 mm diameter alumina (Al₂O₃) ball compliant with wear test standards was used and fixed onto the system. The wear tests were carried out in a reciprocating sliding mode with a full amplitude of 20 mm. The HEAs were tested under two different normal loads, 2 N and 5 N, with a sliding speed of 25 cm/s and a total sliding distance of 500 m.

The specific wear rate (k) was determined using the relation

$$\mathrm{k}={\displaystyle \frac{\mathrm{V}}{\mathrm{N}·\mathrm{L}}}$$

(10)

where V represents the wear volume (mm³), N is the applied load (N), and L is the total sliding distance (m). The wear volume V was quantified by surface profilometry of the reciprocating wear scar. A reference plane was fitted to unworn shoulders adjacent to the scar, and only material removal below this plane was integrated (pile-up excluded). For 3D optical maps, the negative height field was integrated over the stitched scar area; for stylus measurements, we averaged the cross-sectional loss area from M evenly spaced transverse profiles and multiplied by the effective worn length. Each condition was repeated on ≥3 specimens; profilometry was performed at least twice per specimen. Calibration against a step-height standard and re-leveling on unworn areas were used to control baseline drift.

The worn surfaces were examined using a scanning electron microscope (Thermo Fisher Scientific Apreo S LoVac SEM, Thermo Fisher Scientific Inc., Waltham, MA, USA), and the elemental distributions on the surfaces were evaluated through mapping analysis. The wear test parameters of the HEAs are presented in

Table 4. The wear test parameters of the HEAs.

Parameters	Test Conditions
Load (N)	2 and 5
Sliding Speed (cm/s)	25
Ball Diameter (mm)	6
Sliding Distance (m)	500
Full Amplitude (mm)	20
Atmosphere	Room Temperature (23 °C ± 1)

.

3. Results and Discussion

3.1. Structural Properties

The atomic radii, Pauling electronegativity, and VEC values of the constituent elements used in the synthesized HEAs are presented in

Table 5. Properties of the elements used in HEAs.

Element	Atomic Radius (Å)	Pauling Electronegativity	VEC
Al	1.43	1.61	3
Fe	1.24	1.83	8
Co	1.25	1.88	9
Cr	1.25	1.66	6
Ni	1.25	1.91	10
W	1.37	2.36	6

. Additionally, the calculated parameters for the HEAs including ΔH_mix, ΔS_mix, ΔS_conf, VEC, Ω, Δ_x, δ% and Tm are listed in

Table 6. Parameters of ΔS_mix (J.·(K mol)⁻¹), ΔH_mix (kJ·mol⁻¹), ΔS_conf, VEC, Ω, Δ_x, T_m (K) for AlFeCoCrNiW_x.

HEA	ΔSmix	ΔHmix	ΔSconf	VEC	Ω	Δx	δ %	Tm
HEA-0	13.38	−12.32	1.61	7.20	1.82	0.1205	5.78	1673
HEA-0.1	13.92	−11.92	1.67	7.18	2.00	0.1441	5.78	1713
HEA-0.25	14.34	−11.36	1.72	7.14	2.23	0.1709	5.79	1769
HEA-0.5	14.70	−10.51	1.77	7.09	2.59	0.2030	5.78	1856
HEA-1	14.90	−9.11	1.79	7.00	3.28	0.2432	5.74	2008

.

As shown in Table 6, with the increasing W content, the values of ΔS_mix, ΔH_mix, ΔS_conf, Ω, and Δ_x increased, while VEC and δ% decreased. Notably, as the W content increased, the VEC values decreased from 7.20 to 7.00. Numerous studies in the literature have clearly stated that the W element stabilizes the BCC crystal structure [25,30,38]. Tian et al. extensively investigated the correlation between experimental and thermodynamic data of nearly 200 different HEA systems and their corresponding crystal structures. Their findings revealed that when the VEC is in the range of 4.33 ≤ VEC ≤ 7.55, the resulting crystal structure is body-centered cubic (BCC), whereas for 7.80 ≤ VEC ≤ 9.50, a face-centered cubic (FCC) structure is favored. In addition, it was reported that in order to form a stable solid solution, the thermodynamic parameters should fall within the following ranges: −22 ≤ ΔH_mix ≤ 7 kJ/mol, 0 ≤ δ ≤ 8.5%, and 11 ≤ ΔS_mix ≤ 19.5 J/(K·mol). Moreover, it was emphasized that equimolar HEA systems exhibit the maximum possible ΔS_mix values [36].

In another study on the AlCoCrFeNi₂W_x HEA system, it was reported that increasing W content led to an increase in ΔHmix and Ω values, while the VEC value decreased. This observation suggests that the addition of W may promote the formation of BCC phases [38]. Similarly, in a study conducted on the AlCuCrFeMnW_x HEA system, it was found that as the W content increased, the thermodynamic parameters ΔHmix, ΔSmix, ΔSconf, and Ω increased, whereas VEC and δ values decreased. Notably, the increase in ΔSconf with higher W content was interpreted as an indication that W supports the successful formation of solid solutions and enhances the high-entropy effect. Additionally, the decrease in VEC was highlighted as a factor that promotes the formation of BCC crystal structures in the presence of W [39]. Dong et al. reported that in the AlCoCrFeNiV_x HEA system, increasing V content led to a decrease in VEC and δ values, which in turn stabilized the BCC solid solution phase [25]. Similarly, Jiang et al. observed that an increase in W content resulted in a decrease in the VEC value and a corresponding increase in the fraction of the BCC solid solution phase, thereby stabilizing the BCC structure [30]. In light of these findings, the experimental results obtained for the AlCoCrFeNiW_x HEAs in this study are consistent with the literature in terms of solid solution phase formation. Figure 2 presents the XRD patterns of the synthesized AlCoCrFeNiW_x HEAs.

In the HEA-0 alloy, the formation of an ordered BCC-B2 phase at 2θ ≈ 31° (100), and BCC solid solution phases at 2θ ≈ 44° (110), 2θ ≈ 64° (200), and 2θ ≈ 81° (211) was observed. According to numerous studies in the literature, the peak at 2θ ≈ 31° (100) corresponding to the ordered BCC-B2 phase is clearly attributed to the Al-Ni phase [20,27,40,41]. In the HEA-0.1 alloy, an increase in the intensity of the 2θ ≈ 44° (110) peak was observed, while the intensities of the other peaks decreased. With the increase in W content, characteristic W-rich BCC phase peaks appeared at 2θ ≈ 40° (110), 2θ ≈ 58° (200), 2θ ≈ 73° (211), and 2θ ≈ 87° (200). Moreover, the intensities of these peaks increased with increasing W content, becoming more pronounced.

Figure 2b presents a magnified view of the main peak region corresponding to the 2θ ≈ 44° (110) BCC solid solution phase. With the addition of W, the peak position was observed to shift toward lower 2θ angles. This shift can be attributed to the relatively large atomic radius of W, which is the second largest among the alloying elements. Based on the evaluations carried out using HighScore Plus XRD analysis software (HighScore Plus 4.8 (V4.8.025518), Malvern Panalytical B.V. Almelo, The Netherlands), the lattice parameters, crystallite sizes, and dislocation densities of the 2θ ≈ 44° (110) BCC solid solution phase were calculated for the HEAs. The results of these calculations for the HEA-0, HEA-0.1, HEA-0.25, HEA-0.5, and HEA-1 alloys are presented in

Table 7. The average crystallite size (D), dislocation density (ρ) and lattice constant (a) for AlFeCoCrNiW_x alloys.

HEA	D (nm)	ρ (10−3 nm−2)	a (nm)
HEA-0	28.52	1.22	0.2882
HEA-0.1	38.21	0.68	0.2893
HEA-0.25	31.08	1.03	0.2893
HEA-0.5	28.15	1.26	0.2884
HEA-1	22.65	1.94	0.2886

.

As shown in Table 7, when the W content increased from x = 0 to x = 0.1, the crystallite size increased and the dislocation density decreased. However, with further Increase In W content from x = 0.1 to x = 1, the crystallite sizes decreased while the dislocation densities increased. This behavior is attributed to anisotropic variations in residual stress caused by the increasing W content [39]. The lattice constants of the 2θ ≈ 44° (110) BCC solid solution phase for AlCoCrFeNiW_x HEAs were calculated as 0.2882, 0.2893, 0.2893, 0.2884, and 0.2886 nm for x = 0, 0.1, 0.25, 0.5, and 1, respectively. These values were determined using the HighScore Plus software with reference patterns obtained from the Crystallography Open Database (COD); for instance, the lattice constant of HEA-0 corresponds to the COD database code: 9013475. The addition of W was found to increase the lattice constant of the BCC phase, suggesting that lattice distortion occurred due to the increased lattice parameter. The slight decrease in lattice constant observed for HEA-0.5 and HEA-1 may be attributed to the solubility limit of W being exceeded and to the reduced Al content in the BCC solid solution phase [38]. It has been reported that increasing W content induces lattice distortion and increases the lattice constant. However, beyond a certain threshold, W contributes to the formation of a W-rich BCC phase, which limits lattice distortion and thus reduces the lattice constant [30].

Figure 3 presents the SEM images of all HEAs. In Figure 3a–c, the SEM micrographs of the HEA-0 alloy are shown. The HEA-0 alloy exhibits a single-phase BCC solid solution structure with equiaxed grain morphology. Color differences (gray or white) observed in the planar arrangement of grains are attributed to angular misorientations. EDS analysis revealed no significant differences in the chemical composition of the characterized grains.

As shown in Figure 3c, the high-magnification SEM image of the HEA-0 alloy clearly reveals the presence of grain boundaries in its BCC crystal structure. In addition, the alloy exhibits distinct dark and bright regions within the BCC matrix, indicating spinodal decomposition. The chemical compositions of these regions, characterized by EDS, are presented in

Table 8. Elemental Composition of HEA-0 Alloy Determined by EDS Analysis (wt. %).

Alloy	Region	Al	Co	Cr	Fe	Ni
HEA-0	Nominal	10.68	23.34	20.59	22.12	23.24
	Dark Region	10.31	22.93	21.26	22.53	22.97
	Bright Region	9.49	22.44	23.53	23.69	20.85
	Grain Boundary	8.03	20.28	30.41	21.12	20.16
	Pores	33.42	17.45	16.66	16.57	15.90

. The grain boundary region is notably rich in Cr. Moreover, the dark region is enriched in Al and Ni compared to the bright region, whereas the bright region contains higher concentrations of Cr and Fe than the dark region. These findings confirm that the HEA-0 alloy possesses a single-phase BCC structure and that the formation of dark and bright regions is attributed to spinodal decomposition.

As shown in Table 3, the binary mixing enthalpy values between the constituent elements of the HEA-0 alloy reveal that the Al-Ni atomic pair exhibits a significantly higher negative mixing enthalpy compared to other combinations. This suggests that Al and Ni atoms have a strong tendency to associate and segregate as atomic pairs. According to previous studies, the attraction of Al atoms towards Ni and Co promotes the formation of AlNiCo-enriched phases. Conversely, it has been reported that strong mutual interactions also exist between Fe and Cr atoms, leading to the development of regions enriched in these elements [25,42].

Wang et al. reported that the equiatomic AlCoCrFeNi alloy exhibits an equiaxed dendritic microstructure composed of a bright BCC-B2 phase and a dark BCC-A2 phase [41]. In another study, the same alloy was observed to exhibit a microstructure consisting of BCC precipitates distributed within a BCC-B2 matrix. Additionally, it was found that the BCC + B2 microstructures were formed through the spinodal decomposition mechanism. The B2 matrix phase was found to be enriched in Al and Ni, while the BCC phase was enriched in Cr [43]. In the literature, it has been clearly stated that the equiatomic AlCoCrFeNi HEA system consists of a single-phase BCC solid solution with an equiaxed grain structure and is consistent with the spinodal decomposition mechanism [44,45,46,47]. Furthermore, the formation of BCC and B2 phases has been emphasized to occur through the spinodal decomposition mechanism [43,44,48]. In addition, it has been observed that the BCC phase contains bright regions enriched in Fe and Cr, while the B2 phase consists of dark regions enriched in Al and Ni, with both regions interconnected [25,26,27,42,47,49]. Based on the EDS data presented in Table 8, the microstructural features of the HEA-0 alloy are consistent with the findings reported in the literature.

As observed from the microstructural images presented in Figure 3, pores were formed in all the synthesized alloys. According to the EDS results listed in

Table 9. Elemental compositions of HEAs determined by EDS analysis (wt. %).

Alloy	Region	Al	Co	Cr	Fe	Ni	W
HEA-0	Nominal	10.68	23.34	20.59	22.12	23.24	-
	BCC	10.67	22.91	20.52	22.00	23.90	-
HEA-0.1	Nominal	9.96	21.76	19.19	20.62	21.67	6.78
	BCC	8.07	21.57	21.76	21.77	19.97	6.86
HEA-0.25	Nominal	9.04	19.74	17.42	18.71	19.66	15.40
	BCC	8.07	21.16	20.34	21.27	20.16	8.99
	W-rich	5.67	16.25	14.53	14.86	15.27	33.42
HEA-0.5	Nominal	7.83	17.11	15.09	16.21	17.04	26.69
	BCC	5.50	18.87	19.36	18.83	16.67	20.77
	W-rich	2.25	17.54	22.21	19.43	9.87	28.70
	W phase	0.23	0.63	3.35	1.10	0.29	94.40
HEA-1	Nominal	6.18	13.50	11.91	12.80	13.45	42.13
	BCC	4.97	19.15	20.00	19.86	16.41	17.54
	W-rich	2.46	14.96	18.77	16.56	8.21	31.84
	W phase	0.11	0.61	1.90	0.90	0.62	95.85

and the elemental mapping analysis shown in Figure 4, these pores were found to contain a high amount of Al. It was determined that the Al content in these regions was lower than the nominal composition. This phenomenon can be attributed to the evaporation and/or oxidation of Al during the production process due to its low melting point [20].

Figure 3d,g,j,m presents the SEM images of the W-alloyed HEAs. The observed microstructures are consistent with the XRD results, showing a notable reduction in grain boundary visibility with increasing W content.

The microstructure of the HEA-0.1 alloy is similar to that of HEA-0, and the presence of dark and bright regions within the BCC grains is also evident. It can be inferred that all of the W in the HEA-0.1 alloy is dissolved within the BCC solid solution. This is supported by both the XRD pattern shown in Figure 2a and the elemental mapping analysis presented in Figure 4. Moreover, EDS analysis confirms that the bright regions contain a higher amount of W compared to the dark regions. The SEM images of the HEA-0.25 alloy (Figure 3g–i) indicate that the coexistence of dark and bright regions in the BCC matrix continues, as observed in the HEA-0 and HEA-1 alloys. In addition to this structure, the formation of a new bright phase enriched in W has also been observed. EDS and mapping analyses reveal that there is no significant chemical composition difference between the dark and bright regions within the matrix. As shown in Figure 3 j–l, the microstructure of the HEA-0.5 alloy indicates that W has reached its solid solubility limit. While a portion of W is incorporated into the BCC solid solution, the excess W precipitates as plate-shaped W-rich phases due to the solubility limit being exceeded. The microstructure of the HEA-1 alloy is presented in Figure 3m,n,o. In the HEA-0.5 alloy, white plate-shaped W phases are observed, whereas in the HEA-1 alloy, these transform into white irregular and/or suspended spherical forms. Similarly, Wu et al. reported that the microstructure of the AlCoCrFeNiW alloy contains white, suspended spherical or near-spherical W-enriched particles, which were identified as characteristic W phases [20].

Figure 4 presents the elemental distribution (mapping) analyses of all HEAs. In the microstructures of the HEA-0 and HEA-0.1 alloys, all constituent elements were found to be homogeneously distributed. This observation supports that the W element is fully dissolved in the BCC solid solution phase in the HEA-0.1 alloy. In contrast, in the HEA-0.25, HEA-0.5, and HEA-1 alloys, while a portion of the W is dissolved in the BCC solid solution, another portion segregates along grain boundaries and within the bright phase regions, indicating W enrichment in these areas. Additionally, the dark regions corresponding to the BCC matrix contain only a limited amount of W compared to the nominal composition. This behavior can be attributed to the refractory nature of W, which tends to form secondary phases rather than fully dissolving into the matrix [40]. With the addition of W, grain boundaries appear to become finer, and the number of Al-rich pores decreases. This phenomenon can be attributed to the reduction in Al content within the alloy as the W content increases. Furthermore, in the W-enriched regions—such as grain boundaries, bright zones, and areas containing W-rich phases—the concentrations of Co, Cr, and Fe are found to be relatively close to their nominal values, whereas Ni and especially Al are present at significantly lower levels. According to the literature, these elements tend to form W-rich μ phases. Accordingly, the observed bright regions are considered to correspond to μ-(FeCoCr)₇W₆ intermetallic compounds [50]. It is also known that the solubility of Al in the μ phase is extremely limited [51], which explains the very low Al content observed in the W-enriched regions.

3.2. Microhardness of AlCoCrFeNiW_x HEAs

Microhardness measurements of the HEAs were performed on their microstructures. In HEA-0.5 and HEA-1 alloys, hardness measurements were deliberately avoided in regions containing plate-shaped and spherical W-rich phases, as these regions contain more than 95% tungsten and would directly influence the results. The hardness values of the HEAs are presented in Figure 5. The HEA-0 alloy exhibited a hardness of 507.11 HV1, which increased to 651.81 HV1 in the HEA-1 alloy with the highest W content. A noticeable increase in hardness was observed with the rising W molar ratio.

The increase in W content contributes to the formation of W-rich phases, thereby affecting the overall mechanical properties of the alloys. It is well known from the literature that W acts as a solid solution strengthener when dissolved in the matrix [52]. In the HEA-0.25, HEA-0.5, and HEA-1 alloys, the strengthening effect of the W-rich phases significantly contributed to the increase in hardness. In this context, it can be concluded that W enhances the hardness through a dual strengthening mechanism: by dissolving in the BCC solid solution and by forming intermetallic compounds such as μ-(FeCoCr)₇W₆ [50].

Furthermore, it is well known that the high-entropy effect significantly contributes to the overall properties of HEAs [52]. As shown in Table 6, the mixing entropy (ΔSmix) of the HEAs increased with the rising W content. When comparing HEA-0 and HEA-1, an increase of approximately 11% in mixing entropy was observed, which corresponded to an approximate 28.5% increase in hardness. This finding indicates that the high-entropy effect supports solid solution strengthening and enhances mechanical properties. In addition, the incorporation of W atoms which possess a relatively large atomic radius into the solid solution is known to cause lattice distortion. This distortion increases the resistance to dislocation motion, thereby contributing to improved microhardness [27,53,54].

3.3. Wear Tests of AlCoCrFeNiW_x HEAs

Figure 6 presents the friction coefficient (μ) versus sliding distance curves of AlCoCrFeNiW_x HEAs under applied loads of 2 N and 5 N. At the beginning of the tests, all HEAs exhibited a typical “running-in stage”, during which a sudden increase in the friction coefficient occurred within the first 50 m, followed by stabilization as the system reached equilibrium. A noticeable decrease in the friction coefficient was observed with the addition of W.

Under a 2 N load, the friction coefficient (μ) of the HEA-0 alloy was approximately 0.75 μ during the initial 50 m of sliding distance and gradually increased to around 0.82 μ as the test progressed. This alloy exhibited the highest friction coefficient among all compositions. Moreover, its friction curve displayed significant fluctuations, which are believed to be associated with strong adhesive interactions and/or severe plastic deformation on the contact surface. In comparison, the average friction coefficients of HEA-0.1 and HEA-0.25 alloys were ~0.65 μ and 0.55–0.60 μ, respectively, demonstrating more stable frictional behavior. These findings indicate that the minor addition of W moderately improved the tribological performance. The HEA-0.5 alloy showed a consistent and stable friction response with an average μ value ranging between 0.43 and 0.47, likely due to the presence of W-rich intermetallic phases such as (FeCoCr)₇W₆ in the microstructure. The HEA-1 alloy exhibited the lowest and most stable friction coefficient values, with an average in the range of 0.37–0.40 μ. The formation of (FeCoCr)₇W₆-phase promoted by the increased W content is thought to enhance surface hardness and suppress plastic deformation, thereby improving wear resistance [31].

Under a 5 N load, the HEA-0 alloy exhibited the highest average friction coefficient, ranging from 0.78 to 0.82 μ throughout the test duration. Although the HEA-0.1 alloy demonstrated a slightly more stable friction behavior compared to HEA-0, its average friction coefficient remained between 0.65 and 0.68 μ, with irregular fluctuations still present. This indicates that the low W content contributed partially to the wear behavior but was insufficient to form an effective mechanical barrier. The HEA-0.25 alloy showed an average friction coefficient of approximately 0.58–0.60 μ and a more consistent friction behavior compared to HEA-0.1. This improvement is likely associated with the formation of fine intermetallic phases induced by W addition. The HEA-0.5 alloy demonstrated a stable friction curve with an average μ value of 0.44–0.46. HEA-1 exhibited the lowest average friction coefficient, remaining steady at approximately 0.38–0.40 μ, with an exceptionally smooth friction curve. It can be inferred that the presence of hard phases in the matrix limited plastic deformation, reduced surface damage, and consequently enhanced the stability of the contact surface.

The addition of W element functions as a solid lubricant within the HEA matrix, exhibiting a friction-reducing effect on the contact surface. The incorporation of W significantly improves the tribological performance of HEAs by reducing the coefficient of friction, enhancing hardness, and consequently increasing wear resistance. It is well known that alloying elements with high hardness, such as W and Cr, play an active role in tribological applications due to their wear-resistant properties. When introduced as reinforcing phases into the HEA structure, these elements effectively limit plastic deformation on the surface and improve the overall frictional behavior.

Following the wear tests, the worn surfaces were analyzed using elemental mapping, and the wear track widths as well as the distribution of alloying elements were evaluated. Wear rates were calculated using the tribometer software(V.7.2.6) based on the experimental data. Figure 7 presents the wear rates of the fabricated HEAs under 2 N and 5 N applied loads. The increase in W content exhibited a significant influence on wear behavior under both loading conditions. In particular, the evident differences observed between the HEA-0 and HEA-1 alloys clearly demonstrate the contribution of W addition to the tribological performance. The HEA-0 alloy exhibited the highest wear rate among all the tested alloys under both 2 N and 5 N applied loads. Specifically, the wear rate was measured as 1.52 × 10⁻⁵ mm³/N·m at 2 N and 0.821 × 10⁻⁵ mm³/N·m at 5 N. These findings indicate that HEA-0 demonstrated inadequate tribological performance, which can be attributed to its low surface hardness, limited microstructural stability, and a higher tendency for plastic deformation. Additionally, the higher wear rate under the lower load (2 N) suggests an increased susceptibility to adhesive wear mechanisms. In contrast, the HEA-1 alloy showed the lowest wear rates among all compositions, with values of 0.82 × 10⁻⁵ mm³/N·m at 2 N and 0.433 × 10⁻⁵ mm³/N·m at 5 N. This corresponds to an approximately twofold improvement in wear resistance compared to HEA-0. The enhanced performance of HEA-1 is primarily associated with the formation of high-hardness intermetallic phases, such as (FeCoCr)₇W₆, induced by tungsten addition, which contributed to narrower wear tracks and reduced surface deformation. Moreover, the HEA-1 alloy exhibited a consistent and stable wear behavior under both loading conditions, suggesting that its surface hardness was effectively maintained regardless of the applied load, in agreement with experimental observations.

Figure 8 presents the SEM micrographs of the wear tracks of AlCoCrFeNiW_x HEAs obtained under applied loads of 2 N and 5 N. In addition, SEM line analyses were used to determine the wear track widths induced by the alumina ball on the HEAs. Under a 2 N load, the average wear track widths for the HEA-0, HEA-0.1, HEA-0.25, HEA-0.5, and HEA-1 alloys were measured as 755.62 μm, 695.81 μm, 525.13 μm, 520.35 μm, and 411.58 μm, respectively. Under a 5 N load, the wear track widths were measured as 1023.4 μm, 1002.4 μm, 762.36 μm, 648.22 μm, and 549.26 μm, respectively. The increase in applied load leads to a broader contact area at the wear interface, thereby resulting in more extensive surface interactions and enhanced frictional effects. In contrast, wear tracks formed under lower loads tend to be narrower, and the friction-wear relationship exhibits a more linear trend, particularly in matrix structures that are free from particle clusters or elemental segregations. However, features such as grain boundaries, W-rich hard phases, dendritic regions, and spherical morphological accumulations of alloying elements can significantly influence the frictional behavior. For instance, the hard phases present in the microstructure of the HEAs exert an abrasive effect on the contact surface of the counter ball, leading to a gradual transition from its spherical contact geometry to a flatter surface. In such cases, the worn ball material may become partially embedded into the HEA surface, potentially contributing to a reduction in the friction coefficient. This phenomenon may be associated with surface oxidation reactions. Particularly in the HEA-1 alloy, the narrow and shallow wear tracks observed indicate that wear occurs at a significantly low level, which is directly linked to both the microstructural stability and the wear resistance provided by the hard phases. As a result, it has been determined that increasing W content in the fabricated HEAs leads to a decrease in both the coefficient of friction and wear rate, thereby significantly enhancing the overall wear resistance.

One of the key engineering approaches influencing wear rate is the formation of intermetallic phases or the development of an oxidative surface layer by alloying elements, which helps to reduce material loss. For example, the addition of 7 at.% W to the FeCoCrNi alloy enhances hardness due to lattice distortions caused by the large atomic radius of W in the solid solution phase. The additional hardness observed after heat treatment is attributed to the increased volume fraction of (Co–Fe)₇W₆-type intermetallic phases [33]. In the present study, SEM analyses clearly revealed the formation of both W-rich dendritic structures and (Co–Fe)₇W₆ intermetallic phases, particularly in the HEA-0.5 and HEA-1 alloys. Figure 9a,b present the mapping images of wear tracks formed under 2 N and 5 N loads, respectively. Additionally, the widths of the wear tracks were analyzed in detail. According to the mapping images shown in Figure 9a, which correspond to the wear tracks formed under a 2 N load, branched W-rich intermetallics and agglomerated W-based phases were particularly observed within the grains. The HEA-0.5 alloy emerges as a critical threshold in this context, exhibiting granular W clusters that enhance both the microhardness and wear resistance of the alloy [20]. However, dendritic structures were initially observed in the HEA-0.25 alloy, indicating this composition as the morphological threshold for dendritic formation. In the literature, it has been reported that the addition of 0.5 at.% W to the AlCoCrFeNiWx system results in minimum friction coefficient and wear loss. This improvement is attributed not only to the lattice distortion caused by W, but also to the formation of hard W-rich phases. In conclusion, the addition of W significantly enhances high-temperature friction and wear resistance. Moreover, unmolten W-rich particles further improve wear resistance by acting as hard reinforcing phases [20]. In Figure 9b, the mapping images of the HEAs under a 5 N load reveal that the wear tracks of the HEA-0.5 and HEA-1 alloys exhibit distinct nearly spherical W-rich clusters, with no apparent wear grooves or signs of delamination. In the “grey” image of the HEA-0.5 alloy, it is observed that the reciprocating motion of the spherical alumina ball causes oxidation and debris accumulation, particularly at the edges of the wear track. As observed in both Figure 9a,b, the elements O, Al, Cr, Fe, Co, Ni, and W are clearly distributed throughout the wear track regions. These findings confirm the occurrence of oxidation induced by wear and illustrate the distribution of the constituent elements within the affected area.

4. Conclusions

In this study, AlCoCrFeNiW_x high-entropy alloys were successfully fabricated using the vacuum arc melting method. To evaluate the influence of tungsten (W) addition on the crystal structure, microstructure, mechanical, and tribological properties of the alloys, a series of analyses including XRD, SEM-EDS, microhardness, and wear tests were conducted.

The main findings obtained from this investigation are summarized as follows:

1.

With increasing W content, both the entropy of mixing (ΔSmix) and enthalpy of mixing (ΔHmix) values increased, whereas the valence electron concentration (VEC) decreased, which was observed to stabilize the BCC crystal structure.

2.

According to the XRD analysis, the addition of W caused a shift in the peak position of the (110) BCC solid solution phase at approximately 2θ ≈ 44° toward lower angles.

3.

SEM-EDS and elemental mapping analyses revealed that the HEA-0 alloy consists of BCC + B2 phases formed via spinodal decomposition. In the HEA-0.1 alloy, all constituent elements were observed to be homogeneously distributed within the microstructure. In the HEA-0.25, HEA-0.5, and HEA-1 alloys, a portion of the W element was dissolved in the BCC solid solution, while the remaining W segregated at grain boundaries and bright regions to form W-rich phases. Moreover, plate-like W phases were observed in the HEA-0.5 alloy, whereas spherical W phases were formed in the HEA-1 alloy.

4.

A noticeable increase in the microhardness values of the HEAs was observed with increasing W content. The microhardness of the HEA-0 alloy was measured as 507.11 HV1, while the HEA-1 alloy exhibited a higher value of 651.81 HV1.

5.

Wear tests conducted under both 2 N and 5 N loads revealed that increasing W content led to reductions in the friction coefficient, wear rate, and wear track width. When comparing the HEA-0 and HEA-1 alloys, the wear rate decreased by approximately 1.85 times under a 2 N load and by 1.89 times under a 5 N load. These results indicate that the wear resistance improved by nearly twofold with increasing W content.

The present study is limited to room-temperature wear conditions, and it is recommended that further comprehensive investigations be conducted at elevated temperatures and in oxidative environments to fully determine the performance potential of W-modified AlCoCrFeNi high-entropy alloys (HEAs).