Microstructure Evolution and Properties of Laser Cladding CoCrFeNiTiAl_x High-Entropy Alloy Coatings

Abstract

High-entropy alloy (HEA) coatings of CoCrFeNiTiAl_x (x = 0, 0.5, 1, 1.5, 2) were prepared on the surface of AISI1045 steel by laser cladding. The effects of the Al content on the microstructure, composition, phase constitution, and wear and corrosion resistance of the coatings were investigated. The results showed that when increasing the Al element content from 0 to 0.5, the phase constitution of the CoCrFeNiTiAl_x coating changed from a single Face-centered cubic (FCC) phase to Body-centered cubic 1 (BCC1) and Body-centered cubic 2 (BCC2) phases, with a small amount of Laves phase, which obviously improved the friction and corrosion resistance of the coating. With further enhancing of the Al content, the amount of BCC1 phase increased, while the BCC2 phase and the Laves phase decreased. The CoCrFeNiTiAl₂ HEA coating transformed into a single BCC1 phase, with retrogressive wear and corrosion resistance. It was found that the Al_0.5 alloy coating exhibits excellent wear resistance, high hardness, and corrosion resistance in a 3.5 wt.% NaCl solution. Furthermore, the effect of the Al content on the microstructure, phase, and the relating properties of the CoCrFeNiTiAl_x HEA coatings is also discussed.

1. Introduction

Traditional alloy design considers that the presence of multiple elements in an alloy results in the formation of intermetallic compounds (IMCs), which is detrimental to alloy performance. However, Ye JW proposed the concept of high-entropy alloys (HEAs) in 2004 [1]. An HEA is composed of five or more constituent elements, as per the mole ratio or the nearly mole ratio. The design principle prioritizes the high entropy effect, which inhibits the formation of brittle IMCs with complex crystal structures, instead forming a simple type of alloy solid solution phase [2]. The formation of high-entropy solid solution alloys must satisfy certain conditions. Zhang et al. used three main parameters (8.5 ≥ δ ≥ 0, 7 kJ/mol ≥ ΔH_mix ≥ −22 kJ/mol, and 19.5 J/(K·mol) ≥ ΔS_mix ≥ 11 J/(K·mol)) to characterize the common behavior of constituent elements in HEAs [3]. They are atomic size difference δ, mixing enthalpy ΔH_mix, and mixing entropy ΔS_mix, and their respective expressions are:

$$\mathsf{\delta }=100\sqrt{{\displaystyle \sum }_{i=1}^n{c}_i{\left(1-r_i/\overline{r}\right)}^2}$$

(1)

$$\mathsf{\Delta }{H}_{\mathrm{mix}}={\displaystyle \sum }_{i=1,i\ne j}^n{\Omega }_{ij}{c}_i{c}_j$$

(2)

$$\mathsf{\Delta }{S}_{\mathrm{mix}}=-R{\displaystyle \sum }_{i=1}^n{c}_i\ln c_i$$

(3)

A lot of research has reported HEAs with excellent hardness, strength, and corrosion and wear resistances [2,3,4,5,6], a typical example being that of the CoCrFeNi HEA system. For example, CoCrFeNiMn HEA has higher yield strength (518 MPa) and ultimate tensile strength (660 MPa), and the maximum yield strength of Al_0.7CoCrFeNi HEA is 1022 MPa [7,8]. Due to their excellent performance, HEA coatings are also prepared to protect the substrates in different ways. The commonly used processes for preparing alloy coatings are plasma spraying and high-velocity oxygen fuel (HVOF) spraying [9,10,11]. However, recently, the process of preparing alloy coating by laser cladding has become common [12,13]. Laser cladding is characterized by high energy density, fast heating and cooling rates, a lower heat effect on the base material, and a low dilution rate. The coating microstructure is of lower uniform density and micro-defects and forms excellent metallurgical bonds with the matrix material independent of size [14], Matilainen et al. prepared AlCoCrFeNi HEAs by laser additive manufacturing, and it is found that the solidification rate is a very important factor that determines the existence and distribution of phases in AlCoCrFeNi HEA, fast solidification substantially increases the preference for the BCC phase with a corresponding substantial increase in hardness. It should be noted that this is also related to some parameters such as specific powder, laser beam speed, and power density [15].

As reported, the key factor affecting the alloying behavior and properties of HEA is its elemental composition design. Al and Ti have been investigated in many articles involving HEA due to the larger atomic radius and characteristics of these elements having greater impact on the structural property characteristic of HEAs [12,16,17,18,19]. Liu et al. prepared an AlCoCrFeNiTi_x (x = 0, 0.2, 0.4, 0.6, 0.8, 1.0) coating on AISI1045 steel by laser cladding technology and studied the corrosion behavior of the HEA coating in a 3.5 wt.% NaCl solution. The results proved that the introduction of Ti improves the corrosion resistance of the coating [12]. Zhang et al. used laser cladding technology to prepare a FeCoCrAlNi HEA coating on 304 stainless steel, which exhibited good corrosion resistance and microhardness [13]. However, Kao et al. studied the corrosion resistance of Al_xCoCrFeNi HEA and found that the self-corrosion potential (E_corr) and self-corrosion current (I_corr) were independent of the Al content [20], which conflicts with previous studies. Thus, it is important to study the effect of additive elements for better understanding the HEA systems.

In this work, CoCrFeNiTiAl_x HEA coatings were prepared using laser cladding technology. The performance of the as-prepared HEA coatings and the effect of Al content on the microstructure, phase constitution, and wear and corrosion resistance of the coatings were investigated.

2. Materials and Methods

AISI1045 steel with dimensions of 80 by 60 by 10 mm³ was selected as the substrate. It was cleaned with absolute ethanol as the pretreatment in order to guarantee high matrix-coating bonds. Commercial Co, Cr, Fe, Ni, Ti, and Al powders with purities higher than 99.9 wt.% were used as raw materials. The particle sizes of the powders were determined to be 75–100 μm. The powders were weighed according to the nominal composition of CoCrFeNiTiAl_x (x = 0, 0.5, 1, 1.5, 2 in molar ratio) using a precise electronic balance, the quality of each component is shown in

Table 1. The mass of elements of CoCrFeNiTiAl_x high entropy alloy coatings (g).

Element	Co	Cr	Fe	Ni	Ti	Al
CoCrFeNiTi	23.56	20.80	22.40	23.48	19.16	0
CoCrFeNiTiAl0.5	23.56	20.80	22.40	23.48	19.16	5.40
CoCrFeNiTiAl1	23.56	20.80	22.40	23.48	19.16	10.80
CoCrFeNiTiAl1.5	23.56	20.80	22.40	23.48	19.16	16.20
CoCrFeNiTiAl2	23.56	20.80	22.40	23.48	19.16	21.60

. To ensure uniform mixing of the HEA powder during cladding, the ball-to-powder ratio (mass ratio) was set to 2:1 and mixed at a speed of 300 r/min in a QE-04 planetary ball mill for an hour. Then, a single lap multi-layer of laser cladding feeding powder was carried out under an argon atmosphere with an LSF-SI6000A equipment (BLT, Xian, China), as shown in Figure 1. The specific process parameters are listed in

Table 2. Process parameters of laser cladding.

Parameters	Data
Laser power	2100 W
Scan speed	9 mm/s
Laser beam diameter	5 mm
Single layer thickness	0.2 mm
Layers	30
Mass of feed powder	8 g/min
Cladding area	60 × 30 mm2

.

After laser cladding, the prepared coating material was sectioned into samples measuring 8 by 8 by 4 mm³ using a wire electric discharge machine. The samples were ground and polished using a polishing machine. The microstructure of the HEA coatings was observed by the Hitachi S-4800 cold field emission scanning electron microscope (SEM, Hitachi, Tokyo, Japan) applying an energy dispersive spectrometer (EDS, Hitachi, Tokyo, Japan) to determine the element composition and the field emission transmission electron microscope (TEM, FEI, Hillsboro, OR, USA) Talos F200X. The phase constitution was determined using X-ray diffraction (Bruker D8 Advance model), using Cu–Kα radiation scanning from 30° to 90° at a scanning rate of 8° min⁻¹, the volume fraction of the phases was calculated by using Image-Pro Plus software (Image-Pro Plus 6.0). The microhardness of the sample was measured using a microhardness tester (HV-1000A, Shanghai Optical Instrument Factory, Shanghai, China) with a load value of 200 g and a holding time of 10 s. The average value of seven different points was taken as the hardness of the sample. Before the friction and abrasion test, the mass of each sample was weighed by an electronic balance. At room temperature and air relative humidity of 25%–30%, the friction and wear behavior of the coating was evaluated using a pin-disk wear machine (MMQ-02G, KASON, Jinan, China). The GCr15 bearing steel ball with a diameter of 6 mm was selected as the abrasive material, and the prepared sample was fixed on the disk of the worktable. The GCr15 bearing steel ball started the friction and wear test at the center of the sample. The process parameters are as follows: The load is 10 N, the speed is 100 r/min, and the time is 15 min. Considering that the initial stage of the friction and wear test did not reach a stable state, we took the value in the interval of 301 s–900 s to calculate the average friction coefficient. After the friction coefficient measurement was completed, the mass of each sample after the friction test was re-weighed. The corrosion behavior of the HEA coating in a 3.5 wt.% NaCl solution was studied by using an electrochemical workstation. The sample to be tested was used as a working electrode (WE), with the platinum plate electrode used as a counter electrode (CE), and a saturated calomel electrode (Hg/Hg₂Cl₂ saturated KCl) used as a reference electrode (RE). The potential scanning speed was 2 mV/s during the test, at a scanning range of approximately −1.5 to 0.5 V. The corrosion current density (I_coor) and self-corrosion potential (E_coor) were obtained by fitting the test results using the CorShow software (CorShow software 1.1.16).

3. Results and Discussion

3.1. Phase Constitution

Figure 2 shows the XRD diffraction patterns of the as-prepared laser cladding CoCrFeNiTiAl_x HEA coatings with different Al contents. The Al₀ coating was composed of a single FCC solid solution phase. When increasing the molar ratio of Al (x = 0.5–1.5), the phase constitution of the coatings change from a single FCC phase to two BCC (BCC1 and BCC2) phases, and a small amount of Laves phase peaks can also be found. As can be seen with the increased Al content, the intensity of the diffraction peak of the BCC2 phase and the Laves phase begin to weaken, and simultaneously, the BCC1 phase peak increases. The Al₂ alloy coating transforms to a single phase with the BCC1 crystal structure. Using JADE software analysis, the BCC1 and BCC2 phases were identified as the Al–Ni phase and the Fe–Cr phase, and the Laves phase was identified as FeTi type, respectively.

3.2. Microstructure Evolution

The microstructures of the CoCrFeNiTiAl_x high-entropy alloy coatings are shown in Figure 3. The Al₀ coating in Figure 3a has a uniform single-phase structure with tiny voids, which is consistent with the XRD results. The laser cladding technology is characterized by rapid heating/cooling, which results in poor diffusion during the solidification process that prevents a small amount of gas from escaping when the molten pool solidifies, generating a small number of voids. The Al_0.5 coating shows a typical dendritic crystal, consisting of the dendritic region (DR) and interdendritic region (ID). With the increase of Al content, the number of dendrites gradually increases and the distance between dendrites becomes smaller. This is attributed to the introduction of Al with a relatively larger atomic radius, which leads to grain refinement. When the Al content increases to two, the alloy coating exhibits a uniform BCC1 phase structure. Compared to Al_0.5 in Figure 3a, the Al₂ coating shows more defects, such as the voids indicated by the arrows. Figure 3 proves that the Al content significantly influences the microstructure evolution of the CoCrFeNiTiAl_x high-entropy alloy coatings.

Figure 4a illustrates the morphology of a cross-section of the Al_0.5 coating. It can be seen that three zones comprise the cross-section, including the substrate, the heat-affected zone (HAZ), and the HEA coating (HEAC). The dilution rate of the coating can be calculated by Equation (4):

$$\eta =\frac{h}{H+h}$$

(4)

where η is the dilution rate, h is the substrate penetration, and H is the cladding layer width, coating dilution rate is 32.93%, the combination of the coating and substrate is metallurgical and continuous. Due to the dilution effect of the Fe element in the matrix, the content of each alloy element in the composition transition region significantly deviates from the nominal composition proved by the EDS line scan result. This dilution phenomenon generally affects the whole HAZ zone of the coating. As the distance increases from the substrate to the HEAC, the content of each element gradually approaches the nominal composition.

Elemental distribution maps of Al_0.5 HEA coating are illustrated in Figure 4b. It can be found that the DR is rich in Ni and Al elements, while the IR is rich in Fe and Cr elements. It has been reported that among multi-element metal elements, the more negative the mixing enthalpy, the greater the affinity between the two atoms, the more positive the mixing enthalpy, and the smaller the affinity [2,3,21]. According to the molar mixing enthalpy among the six elements shown in

Table 3. Molar mixing enthalpy (∆H_mix KJ·mol⁻¹) of binary alloys [3,21].

Element	Co	Cr	Fe	Ni	Ti	Al
Co	-	−4	−1	0	−28	−19
Cr	−4	-	−1	−7	−1	−10
Fe	−1	−1	-	−2	−17	−11
Ni	0	−7	−2	-	−35	−22
Ti	−28	−1	−17	−35	-	−30
Al	−19	−10	−11	−22	−30	-

, the mixing enthalpy of Fe and Cr with the other five elements presents as the most positive, which signifies relatively small affinity. Consequently, when solidified, Fe and Cr elements are more easily repelled to the interdendritic region.

3.3. Microhardness

The results of the microhardness test of the CoCrFeNiTiAl_x HEA coatings are shown in Figure 5. The hardness distribution curves are divided into three parts in a step shape—the HEA coatings, the HAZ, and the substrate. The thickness of the coating is ~3.5 mm, compared with 2.5 mm of the HAZ. From the HEA coatings to the substrate, the hardness distribution shows a downward trend, and the decreasing interval is concentrated in the heat-affected zone. The Al_0.5 coating shows the highest hardness value of about 880 HV, which is 6.5 times as much as the substrate.

3.4. Friction and Wear

The curves of the friction coefficients versus the sliding time acquired from the substrate and the CoCrFeNiTiAl_x HEA coatings against the steel ball are shown in Figure 6a. All of the curves consist of a running-in stage initially, which is a process of removing the surface roughness and flattening the uneven surface with relatively small contact areas between the worn surface and the counterpart of the steel ball. After the rough initial surface becomes smooth, the wear enters a stable stage, exhibiting a stable friction coefficient. The average friction coefficient of the substrate is the largest, which is 0.54, and the Al₀ coating is also at a high level with an average friction coefficient of 0.42. It is obvious that, along with the addition of Al, the friction coefficient decreases sharply. The Al_0.5 HEA coating exhibits the smallest friction coefficient of about 0.28, compared with the Al₁, Al_1.5, and Al₂ coatings of about 0.31, 0.33, and 0.35, respectively.

The surface morphology of the coatings and substrate after wear are shown in Figure 6b. As can be seen, the substrate has the widest and deepest wear track, indicating its poor resistance to wear. Yet many furrows formed in the worn surface of the Al₀ coating, which may be due to the coating and abrasive ball being slightly adhered in the early stages of wear. The generated abrasive debris, the coating, and the abrasive material formed the three-body wear, indicating the abrasive wear mechanism of the Al₀ coating. The wear result of the Al_0.5 coating shows no furrows existing on the wear surface, and only slight adhesive wear occurred. With increasing Al content, the wear of the Al₁ and Al_1.5 coatings becomes more serious. The number of furrows and the depth increase, and the wear mechanism changes to micro-abrasive wear. In comparison, the wear morphology of the Al₂ coating emerges as obvious spalling, thus the wear mechanism of the Al₂ coating can be deduced to be micro-fracture (flake) wear.

The wear mass loss of the CoCrFeNiTiAl_x high-entropy alloy coating and the substrate is also presented in Figure 6c. The wear amount of the Al₀ coating is 4.4 mg, and that of the Al_0.5 coating suddenly drops to 2.9 mg, reaching its minimum value. As the Al content increases, the wear amount of the Al₁, Al_1.5, and Al₂ coatings become 3.7 mg, 4.0 mg, and 4.2 mg, respectively. However, the wear of the five HEA coatings is still significantly lower than that of the substrate (19.7 mg). The Archard wear equation is an equation derived from a simplified model of abrasive wear, and although it does not take into account factors such as work hardening and frictional heat, it links the wear rate with the engineering variables and mechanical properties of the material. The expression of the Archard wear equation is [22]

$$\frac{W_v}{S}=K\frac{F}{H}$$

(5)

where W_v is the volume wear, S is the sliding distance, F is the normal load, H is the hardness of the wear material, and K is the wear coefficient. According to the previous microhardness tests, the microhardness descends in the order of Al_0.5, Al₁, Al_1.5, Al₂, and A_l0, AISI1045 steel. The relationship between hardness and the amount of wear is inverse: The higher the hardness, the lower the wear. The experimental results are consistent with Archard’s theoretical equations. Under the same wear condition, the wear rate of the as-prepared coatings is proportional to the friction coefficient and inversely proportional to the hardness.

3.5. Corrosion Resistance

The potential dynamic polarization curve of the CoCrFeNiTiAl_x HEA coatings and substrate in a 3.5 wt.% NaCl solution is shown in Figure 7. The corrosion dynamics parameters obtained using the Tafel epitaxy are shown in

Table 4. Corrosion dynamics parameters for the CoCrFeNiTiAl_x HEA coatings and the AISI1045 steel.

Sample No.	ba (mV)	bc (mV)	Ecorr/V	Icorr/(A·cm−2)
Al0	227.05	149.37	−0.7382	5.60 × 10−6
Al0.5	140.14	78.60	−0.5766	9.13 × 10−7
Al1	90.25	102.11	−0.6791	5.43 × 10−6
Al1.5	157.32	146.30	−0.8591	8.03 × 10−6
Al2	48.80	58.86	−0.8802	9.14 × 10−6
AISI1045	40.97	46.244	−0.9701	4.71 × 10−5

. The self-corrosion potential of the substrate is −0.9701 V, which is lower than the self-corrosion potential of the coating, and the self-corrosion current density is the largest, reaching 4.71 × 10⁻⁵ A·cm⁻². When increasing the Al content from 0 to 2, the corrosion resistance of the CoCrFeNiTiAl_x HEA coating gradually increases, and then decreases. The Al_0.5 coating shows the best corrosion resistance; the maximum self-corrosion potential is −0.5766 V, and the minimum self-corrosion current density is 9.13 × 10⁻⁷ A·cm⁻².

3.6. Discussion

It has been reported that the addition of Al can change the phase constitution and the microstructure of HEA [16,17,18]. Figure 8a displays the TEM bright-field image of the Al_0.5 coating. The EDS mapping in Figure 8d proves that the white area is rich in Fe and Cr elements, while the gray area is rich in Al and Ni elements. Being consistent with the XRD result, the gray and white areas are confirmed to be the BCC1 and BCC2 phases, BCC1 is an ordered phase, respectively. According to our XRD and SEM results, it can be concluded that compared with the XRD data of Al_0.5 coating, the BCC1 phase diffraction peaks of Al_1.0 and Al_1.5 shifted slightly to the left, the BCC2 phase diffraction peaks did not change, and the lattice constant of BCC1 phase of Al_0.5 coating was 0.2881 nm and BCC2 phase lattice constant is 0.29 nm. The Laves phase is a transition region from BCC2 to BCC phase, and its lattice constant is 0.2912 nm, which can be identified as FeTi type. Figure 8c clearly shows that there are some nano-precipitates between the BCC1 and BCC2 phases, with a diameter of about 15–20 nm, and these nano-precipitates rarely occur in the Laves phase. Based on the kinetic theory, due to the relatively slow long-distance diffusion of phase separation in high-entropy alloys with multiple major elements, when laser cladding, the cooling rate is very fast, substitutional diffusion of the elements becomes difficult, resulting in the generation of nano-precipitates. Figure 8e displays the line scan results and

Table 5. EDS composition of different points corresponding to the three cross marks in Figure 8c.

Element	Co	Cr	Fe	Ni	Ti	Al
BCC1	20.58	3.99	9.55	25.14	25.58	15.16
BCC2	15.79	28.05	29.94	10.04	14.47	1.71
Laves	18.02	12.80	17.25	20.24	22.01	9.69

lists the composition analysis. It is obvious that the elements of the Laves phase are more uniformly distributed compared to the enriched BCC1 and BCC2 phases, which proves that the Laves phase is a transition region from the BCC2 phase to the BCC1 phase. Simultaneously, this phenomenon was proved by XRD, seen in Figure 2, where the BCC1 phase peak shifts to the left as the Al element increases. The scan angles corresponding to the main diffraction peaks of the BCC1 phase of the Al_0.5, Al_1.0, and Al_1.5 coatings were approximately 43.513°, 43.525°, and 43.548°, respectively. The left shift of the BCC1 phase peak can be explained by the Bragg diffraction law, which is expressed as:

$$2dsin\theta =n\lambda \left(n=1,2\dots \right)$$

(6)

where d is the interplanar distance, θ is the angle between the incident X-ray and the corresponding crystal plane, n is the diffraction order, and λ is the wavelength of the X-ray. It can be seen from Equation (6) that as the interplanar spacing increases, the diffraction angle decreases; because the atomic radius of the Al atom is large, the introduction of the Al atom will lead to an increase in the lattice parameter of the BCC1 phase, and when Al is added to CoCrFeNiTi alloys with FCC structure, the elements solid-solve with each other to form a higher lattice distortion energy, which results in a high energy metastable state of the alloy system. The structural changes of high-entropy alloys caused by Al addition can be explained by the atomic packing efficiency of FCC and BCC [23]. As the FCC structure has the highest density of atomic arrangements (74%), in order to release the increasing lattice distortion energy, the structure tends to transform into a BCC structure with a lower atomic density (68%), forming a stable structure with lower energy; this trend becomes more apparent as the Al content increases [24]. In addition, the FCC phase in HEAs transforming to the BCC phase through the addition of Al element has also been proven by many research works [25].

In order to further explore the influence of the coating’s microstructure change on the performance, the crystal size D calculation was performed by the Scherrer equation:

$$D=\frac{k\lambda }{\beta cos\theta }$$

(7)

where D is mean crystalline size in the nm, k is the Scherrer constant, which is usually considered to be 0.89, λ is 0.154056 nm, θ is identified as the position of the diffraction angle of the XRD peak, and β is the Full width at half maximum (FWHM) of the peak at θ (rad). The crystallite size of the Al₀ FCC phase is 36.31 nm. With the increase of the Al element, the crystallite sizes of the BCC1 phase of Al_0.5, Al₁ and Al_1.5 are 39.23 nm, 41.54 nm, and 44.68 nm, and the crystallite size of the BCC2 phase are gradual decline is 34.10 nm, 29.98 nm, and 28.74 nm. The crystallite size of Al₂ coating of a single BCC1 phase is 32.90 nm. It can be found that rapid heating and rapid cooling are the grain refinement of the coating’s microstructure, which increases the solubility, produces a lattice distortion effect and a solid solution effect, and improves the hardness and corrosion resistance of the coating. Figure 9 shows the relationship between the phase volume fraction and the average microhardness of the CoCrFeNiTiAl_x HEA alloy coatings. It can be found that the hardness of the alloy is closely related to the addition amount of Al element, and the intrinsic factor is the final phase constitution and microstructure of the alloy. The single-solid solution FCC minimizes the hardness of the Al₀ alloy, due to the fact that the slip system of the FCC structure is much larger than that of the BCC structure. It has been reported that the addition of Al can cause lattice distortion and solid solution effect [22,23]. Referring to the Al_0.5 HEA coating, the phase transformed from a single FCC into a stable BCC structure, leading to a significantly improved abrasion resistance and hardness of the coating. That is also the reason why only slight adhesive wear occurred during the wear process, which indicates that the BCC2 and Laves phases play a strengthening role in the alloy. When the addition of Al increased, solid solution saturation occurred and a low hardness phase of BCC1 became the majority. However, the Al₂ coating experienced large flaky flaking in the wear resistance experiment, possessing a micro-fracture (peeling) wear mechanism, which may be attributed to the brittleness of the BCC1 phase, the result is consistent with the required change in crystallite size. This also confirms the results of previous studies, because laser cladding has a higher cooling rate, the higher solidification rate is beneficial to the formation of BCC phase and improves the hardness.

The corrosion test proves that the addition of Al also has a significant effect on the corrosion resistance of the CoCrFeNiTiAl_x HEA alloy coatings. Using the Al₀ coating with a single FCC structure containing a large amount of Cr and Ti elements, it is easy to form Cr₂O₃, TiO₂, and Ti₂O₃ passivation films, which enhances the corrosion resistance of the coating. As a passive element, the addition of a small amount of Al will form an Al₂O₃ passivation film, which can further improve the corrosion resistance of the Al_0.5 coating. With the increase of Al content to one or higher, the structure of the coating changed significantly. The dendrite content gradually increased and the interdendritic content decreased. Consequently, the increased grain boundaries exhibiting a higher dislocation density and corrosion activity made them more susceptible to corrosion by the medium. Furthermore, the secondary precipitation effect produced by the higher content accelerated the pitting corrosion of Al in the coating. Thus, the Al₂ coating shows the lowest corrosion resistance among the five coatings. At the same time, according to the previous calculation of the crystallite size, it can be proved that when x = 0.5, the crystallite size difference between the BCC1 and BCC2 phases is the smallest, so its corrosion resistance performance is the best.

Figure 10 shows the comparison of the self-corrosion potential E_corr and the self-corrosion current density I_corr of the six HEAs: CoCrFeNiTiAl_x, CoCrFeNi, Al_xCoCrFeNi, Al_xCoCrFeMnNi, CoCrFeNiAiTi_x, and Al_xCoCrFe_2.7NiMo [6,12,26,27]. The introduction of Ti element can promote the passivation behavior of the coating due to the formation of TiO₂ and Ti₂O₃ phases. Adding Mo to the CoCrFeNiAl alloy can result in a significant improvement in the corrosion resistance of the alloy, which is ascribed to the fact that Mo can easily produce a dense passivation film. It is generally believed that Al₂O₃ with a dense structure can effectively resist Cl¯ ion corrosion. While a further increase in Al content leads to an increase in the volume fraction of (Ni,Al)-rich BCC1 phases, which is susceptible to Cl¯ ion corrosion and reduces the local corrosion resistance of the passivation film. Based on this study, the reason for the decrease in corrosion resistance may be attributed to the microstructure evolution. In addition, the changed metal oxide content in the passivation film may play a certain role. Consequently, there should be an interaction between the passivation behavior and the effect of the metal oxide content on the corrosion resistance of HEA alloys and coatings.

4. Conclusion

In this work, the CoCrFeNiTiAl_x (x = 0, 0.5, 1, 1.5, 2) coatings were prepared on the surface of AISI1045 steel by laser cladding technology, and its microstructure, hardness, wear resistance, and corrosion resistance were analyzed.

• Under the conditions of high laser power and fast solidification rate, the increase in Al content promoted the transition from a single FCC phase to the BCC1 and BCC2 phases, and a small amount of Laves phase. Due to the larger atomic radius of Al, as the Al element increases, the lattice distortion energy increases, and the BCC2 phase gradually decreases to form a relatively stable BCC1 phase.
• As the Al content increases, the hardness and abrasion resistance of the coating first increase and then decrease. The crystallite size of each phase of the coating also changed, and the grains gradually refined. The grain size difference between the two phases of the Al_0.5 coating was the smallest. The grain refinement resulted in the highest hardness of the Al_0.5 coating, reaching 880 HV, 6.5 times that of the substrate. The abrasion mechanism was the transition from abrasive wear to slight adhesive wear, and finally to micro-fracture (flake) wear.
• The corrosion resistance of the CoCrFeNiTiAl_x HEA coatings in a 3.5% NaCl solution increased first, then decreased with increasing Al content. The addition of a small amount of Al will form an Al₂O₃ passivation film, the self-corrosion potential of Al_0.5 was the most positive, at −0.5766 V, while the self-corrosion current density was the smallest, at 9.13 × 10⁻⁷ A·cm⁻². The Al_0.5 coating showed good corrosion resistance. This also proves that when there are multiple phases, the performance of the alloy may be better when the crystallite size difference of each phase is small.
• The results show that adding a small amount of aluminum can refine the grains of the CoCrFeNiTiAl_x coating and improve the corrosion resistance and mechanical properties of the coating. CoCrFeNiTiAl_0.5 exhibits the best performance among the as-prepared coatings, can significantly improve the surface properties of AISI1045 steel.