https://orcid.org/0000-0002-3651-1836
Figures
Abstract
In this study, the high-entropy alloy AlCoCrFeNiCu0.5 − xNbx (x = 0, 0.1, 0.2, 0.3, 0.4, 0.5) is selected as the research subject. Coatings with varying Cu/Nb ratios are fabricated by laser cladding. The phase composition, microstructure, microhardness, tribological behavior, and corrosion resistance are systematically investigated. The results reveal that when x = 0.2, the coating exhibits a dual-phase structure consisting of FCC and Laves phases. A large cellular structure forms and maintains good microstructural continuity. The coating shows a low friction coefficient of 0.327, and no significant spalling is observed on the worn surface. In both 3.5% NaCl and 0.5 M H₂SO₄ solutions, the coating exhibits the most positive corrosion potential, the lowest corrosion current density, and the highest polarization resistance, indicating excellent electrochemical stability.
Citation: Zhang M, Sun S, Wang J, Zhang F, Wang Y (2026) Study on the microstructure and properties of laser-cladded AlCoCrFeNiCu0.5-xNbₓ high-entropy alloy coatings. PLoS One 21(3): e0342446. https://doi.org/10.1371/journal.pone.0342446
Editor: Wislei Riuper Osório, UNICAMP, University of Campinas, BRAZIL
Received: August 28, 2025; Accepted: January 22, 2026; Published: March 16, 2026
Copyright: © 2026 Zhang et al. This is an open access article distributed under the terms of the Creative Commons Attribution License, which permits unrestricted use, distribution, and reproduction in any medium, provided the original author and source are credited.
Data Availability: We confirm that the submission contains the minimal dataset required to replicate all findings reported in the paper. All raw data underlying the results (including the values used to generate all figures/graphs and the values behind reported summary statistics) are provided in the Supporting Information files.
Funding: This study was supported by the 111 Project of China in the form of a grant awarded to SS (D21017); the Shandong Provincial Key Research and Development Program (International Science and Technology Cooperation) in the form of a grant awarded to SS (2024KJHZ002); the Shandong Provincial Key Research and Development Program (Competitive Innovation Platform) in the form of a grant awarded to SS (2025CXPT081); the Natural Science Foundation of Shandong Province in the form of a grant awarded to SS (ZR2023ME156); the Double Hundred Plan Talent Program of Shandong Province in the form of a grant awarded to SS (WSR2023055) and SS (WSR2025056); the National Natural Science Foundation of China in the form of a grant awarded to JW (52405492); and the Qingdao Natural Science Foundation in the form of a grant awarded to SS (24-4-4-zrjj-68-jch) and FZ (25-1-1-207-zyyd-jch). The funders had no role in study design, data collection and analysis, decision to publish, or preparation of the manuscript.
Competing interests: The authors have declared that no competing interests exist.
1. Introduction
With the development of the marine economy [1], ships have become a crucial mode of marine transportation. However, this has also led to significant emissions of exhaust gases. Desulfurization towers are widely employed as exhaust gas treatment systems on ships. Components such as valves, pistons, and impellers within these towers operate under harsh and complex conditions for prolonged periods, often suffering from wear and corrosion failures [2–5]. Additionally, these components are difficult to replace and involve high maintenance costs. Therefore, it is imperative to develop high-performance coatings to protect key components and enhance their operational longevity [6].
High-entropy alloys (HEAs) are metallic materials composed of five or more principal elements in near-equiatomic ratios [7–9], and have attracted significant attention since their initial development [10]. HEAs are characterized by four core effects: the high-entropy effect, lattice distortion effect, sluggish diffusion effect, and cocktail effect [11–12]. As a result, HEAs demonstrate distinct advantages over conventional alloys in terms of hardness, wear resistance, corrosion resistance, and oxidation resistance [13]. Researchers have employed cladding technologies to design and fabricate high-performance HEAs coatings for surface modification and protection of conventional alloy components. Current cladding techniques primarily include thermal spraying, surfacing welding, electroplating, plasma cladding, laser cladding, and related methods [14]. Although traditional cladding methods such as thermal spraying, surfacing welding, and electroplating are widely used and relatively mature, they each exhibit specific limitations. For example, coatings produced via thermal spraying typically exhibit poor adhesion to the substrate. Surfacing welding tends to result in a high dilution rate. Electroplated metal coatings are generally limited in thickness. As an emerging surface engineering technology [15–17], laser cladding offers several advantages, including high energy density, low dilution rate, strong interfacial bonding, and broad material compatibility. Consequently, laser cladding has been extensively applied in the fabrication of high-performance HEAs coatings for engineering applications.
AlCoCrFeNi is one of the most representative systems among high-entropy alloys (HEAs) and is considered a promising material for laser cladding coatings [18]. With technological advancements, researchers have applied various strengthening treatments to AlCoCrFeNi HEAs coatings during their fabrication process [19]. One approach involves optimizing the processing parameters to enhance coating properties. For example, A. Gunen et al. [20] fabricated a laser-cladded AlCrFeCoNi HEAs coating on 316L stainless steel and subjected it to high-temperature filling and expansion. The surface of the coating formed (CoFe) B2, (CrFe) B2, and Cr₂Ni₃B₆ phases, increasing the coating hardness to20.15 GPa and exhibiting excellent wear resistance at 650 °C. Wang et al. [21] applied ultrasonic surface rolling extrusion (USRE) to AlCoCrFeNi HEAs coatings and found that the magnitude of static load significantly affected the coatings’ mechanical and tribological properties. Under a 300 N static load, the coating underwent martensitic transformation, grain refinement, and dislocation strengthening, resulting in a microhardness of 753 HV. Wang et al. [22] also investigated a hybrid method combining spark plasma sintering (SPS) and laser cladding (LC) to prepare AlCoCrFeNi HEAs with low porosity, high hardness, and excellent wear resistance, achieving a maximum hardness of 644.1 HV.
In this study, to enhance the performance of high-entropy alloy (HEA) coatings, researchers incorporated specific alloying elements into the HEA system. Jaturapronperm et al. [23] systematically investigated the effects of Sn and Ti additions on the microstructure and properties of CoCrFeNi-based alloys. The CoCrFeNiSn coating exhibited a BCC + FCC dual-phase structure, while the CoCrFeNiTi coating showed an FCC + Laves dual-phase structure. Compared with Sn, Ti provided a more pronounced improvement in the microhardness of the coating.
Numerous studies have demonstrated that the addition of copper (Cu) can significantly improve the mechanical properties and optimize the microstructure of alloys, primarily due to its excellent thermal conductivity and high diffusion efficiency. As a result, Cu is considered to play an irreplaceable role in enhancing the performance of high-entropy alloys (HEAs). To address the corrosion challenges of 5083 aluminum alloy in marine environments, Wu et al. [24] fabricated Al0.7FeCoCrNiCux (x = 0.30, 0.60, 0.80, 1.00) HEA coatings on the alloy surface via laser cladding. The influence of Cu addition on the phase evolution from a single BCC phase to a dual-phase structure consisting of BCC and FCC was investigated. With increasing Cu content, the corrosion current density decreased, and the Al0.7FeCoCrNiCu1.00coating exhibited the best corrosion resistance. Yang et al. [25] prepared FeCoCrNiCux (x = 0, 0.5, 1.0) coatings to evaluate the effect of Cu on frictional behavior. The results indicated that the friction coefficient decreased progressively with increasing Cu content. When x = 1.0, the coating displayed a single FCC phase, with Cu/Ni segregation observed at grain boundaries, resulting in optimal strength and wear resistance. Wu et al. [26] incorporated Cu into FeCoNiAl coatings to investigate its effect on wear and corrosion resistance. Their findings revealed that when the Cu atomic content was below 4.76 at%, the phase structure remained unchanged. At a Cu content of 0.4, the coating achieved the lowest friction coefficient (0.381) and exhibited enhanced resistance to cracking. Furthermore, Cu-containing coatings significantly reduced the passivation current density, thereby improving corrosion resistance.
As a large-sized, high-melting-point metallic element, Nb exhibits extremely strong oxygen affinity, which plays a crucial role in forming an oxide film to enhance the corrosion resistance of the alloy. Consequently, many scholars have incorporated Nb as one of the alloying elements in high-entropy alloy coatings. To prevent wear and corrosion of Ti6Al4V material during application, Li et al. [27] fabricated CoCrFeNiNbx (x = 0, 1.0) coatings. The frictional corrosion and corrosion resistance of the coatings in environments with pH = 7 and pH = 3 were investigated. The results indicate that Nb facilitates the formation of a dense oxide film on the coating surface, thereby reducing the wear rate and enhancing the corrosion resistance of the coating. Feng et al. [28] systematically studied the microstructure and properties of CoCrFeMnNiNbx (x = 0, 0.25, 0.5, 0.75, 1.0) coatings by controlling the Nb content. Their findings revealed that with the incorporation of Nb, the coating transitions from an FCC phase to an FCC+Laves dual-phase structure. When x = 1, the hardness of the coating reaches its optimum value (416.54 HV0.5). When x = 0.75, the passivation film of the coating is most stable, exhibiting the best corrosion resistance. Wu et al. [29] investigated the electrochemical corrosion and friction wear properties of CoCrFeNiNbx (x = 0, 0.1, 0.2, and 0.3) coatings in sulfuric acid solution. The incorporation of Nb achieved solid-solution strengthening and grain refinement in high-entropy alloy coatings. The Nb0.3 HEACs exhibited an Nb-rich Laves phase and demonstrated the best acid corrosion resistance in sulfuric acid solution.
To explore the synergistic regulation effects of multiple elements, numerous researchers have introduced two or more metallic elements into existing alloy systems. These six- or seven-component alloy systems further refine the microstructure and enhance the stability of the resulting coatings. Moreover, their practical applicability has been validated in various engineering scenarios. Ding et al. [30] fabricated a two-phase gradient CoCrFeNiAlₓMn(1-x) coating with a layered structure and found that with increasing Al content, the coating transformed from a single FCC phase to a dual-phase structure comprising FCC and BCC phases, achieving a microhardness of 530 HV. Both double-layer and four-layer gradient coatings exhibited mixed brittle and ductile fracture modes. Ma et al. [31] incorporated Ti and Zr into the AlCoCrFeNi system to prepare AlCoCrFeNiTi1-xZrₓ (x = 0, 0.25, 0.5, 0.75, and 1.0) coatings via laser cladding. The results indicated that with increasing x, the phase structure evolved from BCC to a BCC + FCC dual-phase structure. Notably, the coating with x = 0.5 exhibited the highest microhardness and wear resistance. Huang et al. [32] fabricated (FeCrCoNi)₈₄AlₓNbᵧ (x = 12, 8, 4; y = 4, 8, 12 at%) coatings via laser cladding and investigated the effect of different Al/Nb ratios on the micro-wear performance. When the Al/Nb ratio was 1:3, the coating transitioned from a BCC phase to an FCC + Laves dual-phase structure with a refined grain size of 4.6 μm. At 285 °C, the wear volume decreased to 0.39 × 10⁵ μm³.
Copper (Cu) is frequently employed as a key alloying element for synergistically regulating alloy properties and enhancing performance in combination with other metallic elements.
In this study, AlCoCrFeNiCu₀.₅ ₋ ₓNbₓ (x = 0, 0.1, 0.2, 0.3, 0.4, 0.5) high-entropy alloy (HEA) coatings with varied Cu and Nb contents were fabricated by laser cladding. The synergistic effects of Cu and Nb on the coatings’ microstructure, phase constitution, hardness, friction and wear behavior, and corrosion resistance were systematically evaluated. Guided by a “structure–interface–environment” co-optimization framework, Cu was employed to stabilize the solid-solution matrix, densify the microstructure, and improve heat/mass-transport pathways, while Nb was utilized to regulate phase architecture and reinforce the robustness of the surface passive film, thereby balancing the dual requirements of wear resistance and corrosion protection. This strategy differs from conventional single-element approaches and highlights a key novelty: by tuning the Cu/Nb ratio, a simultaneous compromise between “toughness–hardness” and “densification–passivation” is achieved. The results elucidate how elemental interactions govern performance optimization, providing theoretical guidance and a practical reference for the deployment of AlCoCrFeNiCu₀.₅ ₋ ₓNbₓ coatings in applications demanding enhanced anti-wear and anti-corrosion capabilities.
2. Materials and methods
2.1 Material preparation
In this study, 316L stainless steel is selected as the substrate material. The substrate hardness is approximately 200 HV (typical value for 316L; measurement uncertainty is within the microhardness tester specification), and the thickness is 10.0 ± 0.05 mm (measured by a digital caliper). The detailed chemical composition is provided in Table 1. Prior to the experiments, the substrate surface is sequentially polished using 80-, 320-, and 800-grit SiC sandpapers, followed by ultrasonic cleaning in anhydrous ethanol for 15.0 min to remove surface contaminants. Laser cladding is performed using a laser power of 1400 ± 10 W, a nominal spot diameter of 3.0 ± 0.1 mm, and a scanning speed of 3.00 ± 0.05 mm/s. High-purity argon is used as the shielding gas.
2.2 Coating materials
To minimize defects such as pores and cracks during cladding, the physical/chemical compatibility between the cladding material and the 316L substrate, as well as the microstructural stability and intrinsic properties of the feedstock, are considered. Accordingly, high-entropy alloy (HEA) powder blends containing seven metallic elements (Al, Co, Cr, Fe, Ni, Cu, and Nb) are selected as the cladding materials. Previous studies [33–35] show that AlCoCrFeNi-based alloys form single-phase or multiphase structures with high hardness and good wear resistance, supporting their use in surface engineering. In the present alloy design, Cu and Nb serve as compositional variables, and the nominal compositions of the cladded coatings are listed in Table 2.
The laser system is an FL020 fiber laser (ROFIN, Germany) with a rated maximum output power of 2.0 kW. The purity of the elemental powders (Al, Co, Cr, Fe, Ni, Cu, and Nb) is ≥ 99.9% (supplier specification). The AlCoCrFeNi premixed powder consists of spherical particles with a particle-size range of 0–53 μm (as specified by the supplier/sieving), and the Cu and Nb powders show a similar nominal size range (0–53 μm), as illustrated in Fig 1. The elemental molar ratio is controlled as Al:Co:Cr:Fe:Ni:(Cu + Nb) = 1:1:1:1:1:0.5, with an overall weighing uncertainty within ±0.01 g for each component, corresponding to a composition uncertainty on the order of the balance accuracy.
The powder blends are homogenized in a ball mill for 4.0 h. Subsequently, the mixtures are dried at 120 ± 2 °C for 6.0 h in a drying chamber prior to laser cladding.
2.3 Experimental methods
2.3.1 Microhardness.
The hardness test was conducted using a digital microhardness tester (TMVP-1, Beijing Times Peak Technology Co, Ltd.). Measurements were taken along a vertical line extending from the surface of the cladding layer to the substrate, with one test point selected at intervals of 100 μm. The test was performed under a load of 0.981N, with a loading time of 10 s, a dwell time of 15 s, and an unloading time of 10 s. Subsequently, additional vertical lines spaced 100 μm apart were selected in both horizontal directions (left and right) for supplementary hardness measurements. The average hardness value along each horizontal line was calculated to reduce experimental error.
2.3.2 Tribological test.
Friction testing of the cladding layer was carried out using a Bruker UMT-TriboLab tribometer (USA). Prior to testing, the surface of the cladding layer was ground, leveled, and polished to ensure uniform flatness and consistency. A 9.525 mm Al₂O₃ ceramic ball was used as the counterbody, with an applied normal load of 20 N and a reciprocating frequency of 2 Hz. The test was conducted in reciprocating mode for 1800 s, with a stroke length of 4 mm. Following the test, the friction coefficient curve was recorded, and the wear morphology was analyzed using scanning electron microscopy (SEM).
2.3.3 Corrosion-friction test.
Corrosion–friction tests are carried out using a reciprocating tribo-corrosion system (MFT4000, Lanzhou, China) in 3.5 wt% NaCl solution (prepared by weighing; concentration uncertainty follows the balance and volumetric glassware accuracy). Tests are conducted at 25 ± 2 °C.Prior to testing, the sample is immersed for 3600 ± 10 s to stabilize the electrochemical response. During testing, a normal load of 5.00 N, a sliding frequency of 0.10 Hz, and a reciprocating stroke length of 5.00 ± 0.05 mm are applied. The procedure consists of three stages: (i) open-circuit potential (OCP) monitoring for 10.0 ± 0.2 min, (ii) friction under load for 30.0 ± 0.2 min, and (iii) OCP monitoring for 10.0 ± 0.2 min after load removal. The time-dependent OCP and coefficient of friction (COF) are recorded continuously.
2.4 Electrochemical corrosion test
The corrosion resistance of the cladded coatings is evaluated using an electrochemical workstation (CHI760E, China) at room temperature (25 ± 2 °C). The cladded layer is sectioned, ground, and polished, and then embedded in epoxy resin to prepare cylindrical specimens with a diameter of 20.0 ± 0.1 mm and a height of 10.0 ± 0.1 mm (measured by a digital caliper). The exposed working area is controlled to 1.00 ± 0.02 cm² using a masking method, and the area uncertainty is determined from dimensional measurements.
A 3.5 wt% NaCl solution is used for neutral corrosion tests and a 0.5 M H₂SO₄ solution is used for acidic tests. Electrolytes are prepared by weighing and volumetric dilution; the concentration uncertainty follows the analytical balance and volumetric glassware specifications. A conventional three-electrode configuration is employed with the coating as the working electrode, a saturated calomel electrode (SCE) as the reference electrode, and a platinum sheet as the counter electrode. Prior to each test, the specimen is immersed for 3600 ± 10 s to stabilize the open-circuit potential (OCP).
Potentiodynamic polarization is recorded within ±0.5 ± 0.01 V relative to the stabilized OCP at a scan rate of 1.00 ± 0.02 mV s ⁻ ¹. The potentiodynamic scan rate can influence the apparent Tafel slopes and the corrosion current densities obtained by Tafel extrapolation, and an appropriate scan rate is therefore required to minimize possible distortions [36–38]. In the present work, a scan rate of 1 mV s ⁻ ¹ is adopted following these recommendations, and the resulting polarization curves show well-defined Tafel regions with consistent fitting behavior across repeated measurements, supporting reliable extraction of corrosion parameters. Electrochemical impedance spectroscopy (EIS) is performed over a frequency range of 10 ⁻ ²–10⁶ Hz; the frequency accuracy follows the instrument specification. EIS fitting parameters are reported together with fitting errors provided by the fitting software. All electrochemical tests are repeated at least three times, and the results are expressed as mean ± standard deviation (SD) unless otherwise stated.
2.5 Microstructure and phase characterization
The microstructure and elemental distribution of the cladded coatings are characterized using scanning electron microscopy (SEM) equipped with energy-dispersive spectroscopy (EDS). Prior to observation, the cladded layer is sectioned into specimens of 10.0 ± 0.1 mm × 10.0 ± 0.1 mm × 5.0 ± 0.1 mm using wire electrical discharge machining (WEDM), and the dimensions are verified with a digital caliper. The specimens are then ground and polished, followed by etching in freshly prepared aqua regia with a nominal volume ratio of HCl:HNO₃ = 3:1 to enhance microstructural contrast.
Phase identification is performed by X-ray diffraction (XRD) over a 2θ range of 20.0°–90.0° (angular accuracy within ±0.02° per instrument specification) at a scan rate of 5.0° min ⁻ ¹.
3. Results and discussion
3.1 XRD analysis
Fig 2a presents the XRD patterns of AlCoCrFeNiCu0.5-xNbₓ coatings (x = 0, 0.1, 0.2, 0.3, 0.4, 0.5). Phase identification is carried out by matching the diffraction peaks with standard ICDD PDF/JCPDS cards. When x = 0, the coating C1 (AlCoCrFeNiCu0.5) exhibits a single FCC-phase (ICDD PDF#47–1417) solid solution, with diffraction peaks corresponding to the (111), (200), and (220) planes. This indicates that in the absence of Nb, the coating forms a typical single-phase solid solution, in which Cu is uniformly dissolved into the high-entropy AlCoCrFeNi matrix to form a stable FCC structure. According to the Hume-Rothery rules, elements with similar atomic radii tend to form solid solutions with high solubility. The average atomic radius of the AlCoCrFeNi system is approximately 1.29 Å, while that of Cu is 1.28 Å, facilitating the formation of a stable FCC phase. With the introduction of Nb, the phase composition of the coatings changes significantly. In coating C2 (AlCoCrFeNiCu0.4Nb0.1), the C15-Laves phase (ICDD PDF#15–0499), associated with of Co2Nb, appears with a (311) diffraction peak. This suggests that Nb promotes the precipitation of new phases, resulting in a dual-phase structure comprising FCC and C15-Laves phases.
Phase identification is performed using ICDD PDF/JCPDS reference cards: FCC (PDF#47-1417), C15-Laves/Co₂Nb (PDF#15-0499), C14-Laves/Fe₂Nb (PDF#17-0908), and FCC Cu–Ni solid solution (Cu0.81Ni0.19, PDF#47-1406).
As the Nb content increases (x ≥ 0.2) and Cu content decreases, the diffraction peaks of the Laves phase become more intense. In coatings C3–C6, the C14-Laves phase (ICDD PDF#17–0908) is detected, with a characteristic (112) plane, and is mainly composed of Fe2Nb. This evolution indicates that increasing Nb and decreasing Cu content facilitates the substitution of Nb atoms for Cr and Co in the lattice, which increases lattice distortion and promotes the formation of Laves phases. Given that Nb has a significantly larger atomic radius (1.43 Å) than the average atomic radius of the base alloy, excessive lattice distortion may exceed the solubility limit of the FCC solid solution, thereby favoring the formation of the more ordered Laves phase. Additionally, the enthalpy of mixing is a key factor in predicting the formation of intermetallic compounds. The mixing enthalpies of Co–Nb and Fe–Nb are –25 kJ/mol and –16 kJ/mol, respectively, both negative, indicating a strong tendency to form stable compounds such as Co2Nb and Fe2Nb, which further promotes Laves phase formation. These findings are consistent with the evolution of diffraction peaks observed in the XRD patterns.
As shown in the magnified XRD pattern in Fig 2b, the peak of C1 within the selected 2θ range is consistent with an FCC Cu–Ni solid solution reference (ICDD PDF#47–1406, Cu0.81Ni0.19), which primarily consists of a Cu0.81Ni0.19 solid solution, with the (111) plane as the dominant reflection. Upon the addition of Nb, the C15-Laves phase (ICDD PDF#15–0499) appears within the same angular range in coatings C2–C6.
Notably, the diffraction peak positions for coatings C2 and C3 show significant shifts. In particular, the peak in coating C3 exhibits a pronounced shift toward lower angles, indicating an increase in the interplanar spacing (d-value) compared to the other coatings. This observation is consistent with Bragg’s law (Equation 1):
Among them, n is the order of reflection, λ is the wavelength of the incident X-rays, d is the interplanar spacing, and θ is the diffraction angle. According to this relationship, as the diffraction angle θ decreases, the interplanar spacing d increases. Furthermore, based on the interplanar spacing equation (as shown in Equation 2), the lattice constant a can be calculated from d. It follows that the lattice constant a varies with θ, exhibiting a negative correlation.
Here, h, k, and l are the Miller indices, and a is the lattice constant. As the lattice constant a represents the edge length of the unit cell, coating C3 (AlCoCrFeNiCu0.3Nb0.2) exhibits the largest interplanar spacing and, correspondingly, the largest lattice constant, indicating that its unit cell has the largest volume. The lattice constant of C2 (AlCoCrFeNiCu0.4Nb0.1) is slightly smaller than that of C3. As the Cu content decreases and the Nb content increases, the diffraction angles (θ) of coatings C4, C5, and C6 gradually shift back to higher values, reflecting a notable reduction in the lattice constant. This suggests significant lattice distortion within the system, which corresponds to the extensive formation of the Laves phase.
3.2 Microstructure analysis
Fig 3 shows the microstructures of the AlCoCrFeNiCu0.5-xNbx HEA coatings. As observed in Fig 3a, coating C1 (AlCoCrFeNiCu0.5) exhibits a uniform equiaxed crystal structure, with no obvious phase precipitation or elemental segregation. The microstructural morphology indicates that in coatings with high Cu content and without Nb, the microstructure formation is primarily governed by solidification dynamics. Grain growth is complete, thermal conduction is relatively efficient, and the grain size tends to stabilize. With the gradual incorporation of Nb, significant changes occur in the microstructure. As shown in Fig 3b and 3c, coatings C2 (AlCoCrFeNiCu0.4Nb0.1) and C3 (AlCoCrFeNiCu0.3Nb0.2) exhibit similar morphologies. The lattice structure expands significantly, evolving into large cellular crystals. The grains appear irregular in shape, and the grain boundaries become less distinct. Compared to C1, the increased lattice size suggests that the addition of Nb alters the solidification behavior of the molten pool. Due to Nb’s high melting point and large atomic radius, it tends to segregate during rapid solidification, particularly at grain boundaries.
In contrast, Cu possesses excellent thermal conductivity and low viscosity, which helps maintain a more uniform temperature in the molten pool, reduces melt viscosity, improves fluidity, and mitigates the segregation tendency of Nb. The rapid diffusion of Cu atoms in coatings C1, C2, and C3 allows them to preferentially migrate toward grain boundaries, where they form short-range ordered structures. This “nail effect” restricts the enrichment of Nb, stabilizes grain growth, and promotes the formation of large cellular crystals.
When the Nb content is further increased (x ≥ 0.3), the microstructures of coatings C4, C5, and C6 (Fig 3d–3f) undergo additional evolution. The previously observed large cellular crystal structures are replaced by finer, more intricate petal-like cellular grains. Grain boundary width increases, and the overall microstructural density decreases.This transformation is primarily attributed to the continuous reduction in Cu content, which weakens its regulatory effect on Nb diffusion. Simultaneously, the incorporation of excessive Nb surpasses the solid solubility limit of the base alloy, resulting in severe Nb segregation along grain boundaries. The local precipitation of Nb disrupts microstructural continuity, thereby compromising the integrity of the coating.
Fig 4 presents the EDS elemental distribution map of the C3 coating, illustrating the spatial distribution of key elements including Fe, Ni, Co, Cu, Nb, Cr, and Al. It is evident that all elements are uniformly distributed across the observed region, with no apparent elemental enrichment or phase segregation. This uniformity suggests that the microstructure of the C3 coating is stable, and that a sufficient solid solution has formed. In particular, the homogeneous distribution of Nb further confirms the regulatory role of Cu in stabilizing the grain structure via enhanced diffusion.
Fig 5 shows the EDS elemental map of the C5 coating (AlCoCrFeNiCu0.1Nb0.4), revealing significant Nb enrichment along the grain boundaries, indicating pronounced segregation. As a result, grain refinement occurs and the microstructural morphology changes markedly, transitioning from large cellular crystals to petal-like fine cellular structures. This transformation trend is consistent with the XRD results: as Cu content decreases and Nb content increases, the single-phase FCC structure progressively evolves into a dual-phase system comprising FCC and Laves phases, accompanied by a variation in the lattice constant.
3.3 Microhardness
Fig 6 illustrates the microhardness distribution of each coating along the depth direction from the surface. As shown in the figure, the region within 0–1 mm corresponds to the cladding layer, where all coatings exhibit relatively high microhardness values, ranging from 600 to 690 HV. With increasing depth, the microhardness gradually decreases. In the molten zone (1.0–1.3 mm), the microhardness drops to 310–460 HV, and further decreases to approximately 200 HV near the substrate, matching the base material hardness.
The hardness data in Table 3 are presented as mean ± SD (n = 3), and “HV” denotes Vickers hardness. According to Table 3, the microhardness of coating C1 is 643.89 HV, with a corresponding molten zone hardness of 310.2 HV. Upon increasing the Nb content to 0.1 (C2), the coating microhardness rises slightly to 644.91 HV, and the molten zone hardness increases to410.97 HV. With further Nb addition in coating C3 (Nb = 0.2), the microhardness reaches 664.55 HV, and 469.2 HV in the molten zone. These results indicate that Nb addition significantly enhances the hardness of both the coating and the molten zone, with the coating microhardness of C3 increasing by over 100 HV compared to C1. As confirmed by the XRD analysis, the increased Nb content promotes the formation of Laves phases, which are known to improve the mechanical strength of the coating, thereby contributing to the observed hardness enhancement.
When the Nb content exceeds 0.2, evident Nb segregation and enrichment are observed in coatings C4, C5, and C6. Grain refinement occurs, and dislocation slip within the grains is simultaneously suppressed. The microhardness of these coatings ranges from 640 to 690 HV. Among them, coating C5 exhibits the highest microhardness, reaching 696.11 HV—representing a 284.8% increase relative to the substrate and a 31.2% improvement compared to C1, which contains no Nb.
Once the Nb content reaches a certain level, the influence of further Nb addition on hardness gradually diminishes. Coatings C3, C4, C5, and C6 all exhibit relatively high microhardness values, but the increasing trend becomes less pronounced. Notably, the microhardness of C5 (696.11 HV) is higher than that of C6 (615.87 HV), indicating that both Nb and Cu contribute positively to enhancing the hardness of the coating.
3.4 Wear resistance
Fig 7 presents the friction coefficient curves of the AlCoCrFeNiCu0.5-xNbₓ HEA coatings tested under ambient air conditions, while Table 4 summarizes the corresponding steady-state friction coefficients. As observed in Fig 7 and Table 4, the time-dependent evolution of the friction coefficients for all coatings follows a similar trend and can be divided into two distinct stages: the initial running-in stage and the subsequent stable wear stage. During the running-in stage, the friction coefficient rapidly increases from zero to its peak value and then transitions into a relatively stable state. This stage corresponds to the initial contact between the counterface and the coating surface, where surface asperities interact and generate high localized pressures due to the limited contact area. These localized stresses cause significant microstructural damage, resulting in pronounced fluctuations in the friction coefficient and an initially high wear rate. Among all samples, the C1 coating exhibits a friction coefficient of 0.404 with minimal curve fluctuation, indicating relatively stable sliding behavior. The C2 coating shows a slightly lower coefficient, reduced by 0.009 compared to C1, and displays less fluctuation during the initial stage before reaching stability, suggesting improved frictional stability. Notably, the C3 coating demonstrates the lowest friction coefficient, approximately 0.327, reflecting superior tribological performance.
In contrast, while the C4 coating shows relatively minor fluctuations during the stable stage, its average friction coefficient is 0.096 higher than that of C3, representing the highest among all coatings. This indicates inferior frictional performance and a stronger tendency toward wear. The C5 coating also exhibits a relatively high friction coefficient accompanied by larger fluctuations, and its curve shows a gradual increase in the later stage of sliding, implying a deterioration in wear resistance. The C6 coating exhibits pronounced fluctuations throughout the test, with an average friction coefficient of approximately 0.39, indicating unstable frictional behavior under the given conditions.
Fig 8 shows the three-dimensional wear track morphologies of coatings C1–C6 obtained by laser confocal microscopy, together with the corresponding cross-sectional profiles. The quantitative measurements further substantiate the above trends. For C1, the wear track width and depth are about 635.2 μm and 10.83 μm, respectively, indicating a relatively large material removal volume, which is consistent with its higher specific wear rate (9.72 × 10 ⁻ ⁵ mm³·N ⁻ ¹·m ⁻ ¹). The C2 coating exhibits a narrower and slightly shallower wear scar (471.9 μm, 9.45 μm), corresponding to a reduced wear rate of 6.25 × 10 ⁻ ⁵ mm³·N ⁻ ¹·m ⁻ ¹. The C3 coating shows the narrowest track width (296.5 μm) with a moderate depth (13.23 μm), resulting in the smallest effective cross-sectional area and thus the lowest wear rate (2.70 × 10 ⁻ ⁵ mm³·N ⁻ ¹·m ⁻ ¹) among all coatings. This observation is in good agreement with its lowest COF, confirming that the optimized Cu/Nb ratio in C3 effectively suppresses both friction and wear.
In contrast, the C4 coating presents the widest and deepest wear track (707.9 μm, 16.27 μm), implying a significantly larger worn volume, which corresponds well to its highest wear rate (1.60 × 10 ⁻ ⁴ mm³·N ⁻ ¹·m ⁻ ¹) and highest COF. The C5 coating exhibits an intermediate wear scar (368.1 μm, 8.96 μm) and a moderate wear rate (4.58 × 10 ⁻ ⁵ mm³·N ⁻ ¹·m ⁻ ¹), while the C6 coating shows a relatively wide and deep track (469.7 μm, 10.41 μm) with an increased wear rate of 6.77 × 10 ⁻ ⁵ mm³·N ⁻ ¹·m ⁻ ¹. As illustrated in Fig 9, the variation of specific wear rate follows a similar tendency to that of the steady-state COF: C3 exhibits the lowest COF and wear rate, C4 shows the highest values for both, and the other compositions lie in between. This consistent correlation between friction behavior, wear rate, and 3D wear morphology further verifies the reliability of the tribological evaluation.
Fig 10 presents the surface morphologies of coatings C1 to C6 after the tribological tests conducted in air. SEM observations reveal that the predominant wear mechanisms across all coatings are adhesive wear and abrasive wear. For coatings C1 and C2, large areas of metal delamination and spalling are observed on the worn surfaces, indicating severe plastic deformation during the friction process. This is primarily attributed to the relatively low hardness of these coatings. With the incorporation of Nb, the hardness of the coatings increases, resulting in a noticeable reduction in metal spalling; however, plastic deformation and local fracture are still evident on the worn surfaces of coatings C4, C5, and C6. This is due to Nb segregation in the microstructure, which promotes the formation of brittle Co₂Nb and Fe₂Nb phases. These intermetallic compounds reduce the coating’s toughness and increase its susceptibility to brittle fracture.
In addition, deep grooves are observed on the worn surfaces of coatings C4 and C5. During the adhesive wear process, delaminated metal debris forms fine particles that, under vertical loading, become embedded in the coating surface. These embedded particles act as abrasive agents during reciprocating motion, further damaging the surface and forming grooves. By comparison, coating C3 exhibits only mild abrasive wear, characterized by minor scratches and no evident plastic deformation. This is consistent with its narrow and shallow 3D wear track and the lowest specific wear rate, indicating that C3 possesses the most favorable combination of hardness, toughness, and interfacial stability, and thus exhibits the best wear resistance among all investigated coatings.
3.5 Corrosion friction performance
Fig 11 presents the coefficient of friction (COF) and open-circuit potential (OCP) curves obtained from corrosion-friction tests for coatings C1 through C6. Specifically, Fig 11a shows the COF variation after load application, while Fig 11b depicts the evolution of OCP throughout the testing process. Table 5 summarizes the fitted COF values as well as the open-circuit voltages (OCP₁ and OCP₂) measured before and after the friction test.
During load application, C1 exhibited the highest COF (0.133) with pronounced fluctuations, indicating severe adhesive and abrasive wear. In contrast, C2 and C5 demonstrated lower COFs (0.089 and 0.094, respectively), suggesting improved tribological performance. C6 showed a slightly higher COF than C3, but significantly lower than C1, indicating that Cu can effectively regulate friction behavior within an optimal compositional range, whereas excessive Cu may be detrimental. Notably, C3 displayed the lowest average COF (0.082) and stabilized rapidly after initial loading, reflecting excellent frictional stability.
Analysis of the OCP before and after loading revealed an overall decline in potential during the friction period (10–40 min), suggesting disruption of the surface passive film and subsequent exposure of the underlying metal to the corrosive environment. For C1, the OCP dropped sharply at the onset of loading, indicating poor oxide film integrity and the lowest corrosion and wear resistance among the samples. C2 and C5 exhibited marginally better OCP values but significant fluctuations, implying unstable surface electrochemical activity, potentially due to alternating dominance of localized corrosion and mechanical wear. In contrast, the OCP of C4 and C6 decreased only slightly and stabilized with time. Remarkably, C3 showed the highest OCP values before and after the test (OCP₁ = −0.134 V, OCP₂ = −0.110 V), indicating that the removal and regeneration of its passive film reached dynamic equilibrium, demonstrating superior electrochemical stability and re-passivation capability.
Fig 12 illustrates the corrosion-friction mechanism. Upon load application after OCP stabilization, the surface passive film is mechanically disrupted, causing a rapid drop in potential and the onset of anodic dissolution. This allows corrosive media to penetrate wear tracks, inducing pitting corrosion. As friction continues, the newly formed passive film is repeatedly worn and regenerated, establishing a dynamic equilibrium characterized by alternating “wear–corrosion–passivation” cycles.
The strong oxidation tendency of Nb facilitates the formation of Nb₂O₅, which integrates with Al₂O₃ and Cr₂O₃ to produce a dense and protective passive layer. This composite film effectively suppresses corrosion and enhances the coating’s wear and electrochemical stability. In addition, the large cellular structure observed in the C3 coating improves ductility and toughness, thereby reducing shear stress and minimizing friction. The presence of Nb further supports the formation of a compact, stable oxide layer, leading to superior corrosion resistance and wear performance. Overall, the optimal Cu and Nb content in C3 promotes the formation of a structurally stable, dense passive film that achieves an excellent balance between tribological and electrochemical protection.
3.6 Corrosion resistance in 3.5%NaCl solutions
Fig 13 displays the electrochemical impedance spectroscopy (EIS) results of the HEA coatings in a 3.5 wt% NaCl solution. The EIS spectra are fitted by complex non-linear least squares (CNLS) using the equivalent circuit Rs–(CPEf ‖ Rf)–(CPEdl ‖ Rct), where Rs is the solution resistance. The high-frequency branch (CPEf ‖ Rf) represents the non-ideal capacitive response and resistance of the surface film (passive/corrosion product layer), while the low-frequency branch (CPEdl ‖ Rct) corresponds to the double-layer behavior and charge-transfer process at the coating/electrolyte interface. Constant phase elements (CPEs) are used instead of ideal capacitors to account for time-constant dispersion caused by surface heterogeneity (n < 1) [39,40]. The CNLS fitted curves (solid lines) are overlaid with the experimental data (symbols) in both Nyquist plots to demonstrate the fitting quality and to support the reliability of the extracted parameters [41]. As shown in the Nyquist plots (Fig 13a), all coatings exhibit incomplete semicircular capacitive arcs, with the arc radius increasing as the Nb content rises. Among them, coating C3 exhibits the largest arc radius, indicating superior corrosion resistance in the neutral solution. The corresponding Bode plots (Fig 13b and 13c) further reveal differences in phase angle responses across various frequency regions. Specifically, the phase angles of C1 and C2 reach peak values of approximately 75° in the mid-to-high frequency range (100–1000 Hz). In comparison, C4 and C5 coatings show enhanced performance, with peak phase angles around 78.5° at ~100 Hz.
These results indicate that decreasing Cu content while increasing Nb content improves the coating’s impedance behavior. Notably, C3 demonstrates the highest phase angle—approximately 80°—within a broad frequency range of 0.1–100 Hz, suggesting excellent electrochemical stability and impedance performance throughout the entire measurement process. Moreover, as Nb content increases, the peak phase angle shifts upward, particularly in the low-frequency region, indicating enhanced corrosion resistance. This improvement is attributed to Nb-induced densification of the coating structure, which reduces the penetration of corrosive species and impedes charge transfer, thereby enhancing the protective and electrochemical stability of the coating.
The corrosion potential (Ecorr) and corrosion current density (icorr) are obtained from the potentiodynamic polarization curves by Tafel extrapolation of the linear portions of the anodic and cathodic branches. It should be noted that icorr (A·cm−2) is an electrochemical kinetic parameter derived from the Tafel model; its accuracy depends on the selection of valid Tafel regions and the applicability of the Tafel relationship. Therefore, the extracted icorr values are primarily used for comparative evaluation among the coatings under identical test conditions. In this work, icorr is reported rather than a corrosion rate, since such conversion requires additional assumptions (e.g., equivalent weight and density). Based on the polarization curve fitting shown in Fig 13d, the polarization resistance (Rp), corrosion potential (Ecorr), and corrosion current density (icorr) for each coating were obtained and are summarized in Table 6. As seen in Fig 13d and Table 6, the coatings show minimal differences in the cathodic region; however, notable variations are observed in the anodic region, ultimately leading to the formation of a distinct passivation plateau. This behavior indicates the formation of a protective passivation film on all coatings during electrochemical polarization. Among the samples, coating C3 exhibits the most positive corrosion potential (–0.3577 V), the lowest corrosion current density (3.68 × 10 ⁻ 8 A·cm-2), and the highest polarization resistance (5.13 × 105 Ω·cm²), indicating superior corrosion resistance.
In contrast, coating C1, which does not contain Nb, shows a more negative corrosion potential (–0.3598V), a higher corrosion current density (4.08 × 10 ⁻ ⁷ A·cm-2), and a significantly lower polarization resistance (7.71 × 10⁴ Ω·cm²), suggesting inferior corrosion resistance compared to C3. Coating C6 which excludes Cu, exhibits a corrosion potential of –0.417 V, a corrosion current density of 4.84 × 10 ⁻ 8 A·cm-2, and a polarization resistance of 3.83 × 10⁵Ω·cm². These values indicate that while the absence of Cu slightly reduces the corrosion resistance compared to C3, the coating still demonstrates excellent electrochemical stability.
3.7 Corrosion resistance in 0.5M H2SO4 solution
Fig 14 displays the EIS results of the HEA coatings in 0.5 M H₂SO₄ solution. The spectra are analyzed by CNLS fitting using the same equivalent circuit model as in Fig 13 (Fig 14e), and the fitted curves (solid lines) are plotted together with the experimental data (symbols) to validate the extracted parameters. As shown in the Nyquist plot (Fig 14a), all coatings exhibit distinct capacitive arc structures. With increasing Nb content, the arc radius shows a clear upward trend. Among them, coating C3 presents the largest arc radius, indicating the best corrosion resistance in acidic conditions. Coating C4 also demonstrates favorable acid corrosion resistance. In contrast, coating C1 (with high Cu content but no Nb) and coating C6 (with high Nb content but no Cu) show significantly smaller arc radii, suggesting inferior corrosion resistance. This trend is consistent with the polarization resistance (Rp) values listed in Table 7, which follow the order: Rp_C3 > Rp_C4 > Rp_C6>> Rp_C1. These results confirm that the C3 coating provides the most effective protection against acid corrosion among all tested compositions.
Analysis of the Bode modulus plot in Fig 14b reveals that the impedance modulus and phase angle values of the coatings differ significantly across various frequency regions. In the low-frequency range (10 ⁻ ² to 10² Hz), the impedance modulus values of coatings C3 and C4 are markedly higher than those of the other samples, suggesting that these coatings provide better resistance to the ingress of corrosive species during the critical low-frequency stage of the corrosion process. Fig 14c shows the Bode phase angle plot. All coatings reach their maximum phase angles within the 10–100 Hz frequency range. Among them, coating C3 exhibits the highest peak phase angle, exceeding 70°, while the peak values for C2, C4, C5, and C6 fall between 60° and 70°. In contrast, coating C1 shows a maximum phase angle below 55°, indicating a thinner or less stable surface passivation film and poor electrochemical stability. These results suggest that reducing Cu content and appropriately increasing Nb content enhances the formation and stability of the passivation layer, thereby improving the corrosion resistance of the coating in acidic environments.
By fitting the polarization curves of each coating in Fig 14d, key electrochemical parameters—including polarization resistance (Rp), corrosion potential (Ecorr), and corrosion current density (icorr)—were obtained and are summarized in Table 7. In the anodic region (–0.3 V to +1.0 V), the polarization curves exhibit significant divergence, reflecting differences in the passivation behavior of the coatings. Coating C3 shows the most positive corrosion potential (−0.2478 V) among all samples and the lowest corrosion current density (1.09 × 10 ⁻ ⁶ A·cm-2), indicating the lowest corrosion rate. In addition, it has the highest polarization resistance, confirming its excellent resistance to acid-induced electrochemical degradation. In contrast, coating C1 exhibits a more negative corrosion potential (−0. 2937 V), the highest corrosion current density (1.82 × 10 ⁻ ⁶ A·cm-2), and the lowest polarization resistance (1.60 × 10⁴Ω·cm²), indicating the poorest corrosion resistance in acidic environments. Although coatings C5 and C6 show corrosion potentials close to that of C3, their icorr values are significantly higher. This suggests that although high Nb content increases hardness, the precipitation of brittle intermetallic phases such as Co₂Nb and Fe₂Nb compromises the stability of the passive film, thereby reducing corrosion resistance.
3.8 Corrosion mechanism of AlCoCrFeNiCu0.5-xNbₓ coatings
In this study, the incorporation of Cu and Nb into high-entropy alloy (HEA) coatings plays a crucial role in enhancing corrosion resistance. The addition of Cu stabilizes the FCC solid solution phase, while Nb promotes the formation of a dual-phase FCC + Laves structure, leading to grain refinement and the development of corrosion-resistant intermetallic compounds. Furthermore, Cu and Nb significantly affect the formation and stability of the passive film. In AlCoCrFeNi alloys, Cr is the primary element responsible for passive film formation. However, excessive Cu can partially substitute for Cr in the oxide layer, resulting in the formation of unstable Cu₂O and compromising the integrity of the film. In contrast, a moderate amount of Cu improves the structural cohesion of the passive film, enhancing its stability. Nb, due to its strong affinity for oxygen, contributes to the formation of a stable oxide film composed of Al₂O₃, Cr₂O₃, and Nb₂O₅. Additionally, Nb integrates into the passive film, reducing defect sites and improving the re-passivation capability of the coating.
As schematically summarized in Fig 15(a), moderate Cu works in concert with Cr, Nb and Al to build a dense composite film (Cr₂O₃ backbone, Nb₂O₅ defect suppression/re-passivation, and Al₂O₃ stability), which is reflected by the upward arrows indicating Ecorr ↑ , icorr ↓ , and Rp ↑ . In contrast, Fig 15(b) illustrates two non-optimal cases: (i) excess Cu favors a Cu₂O-containing, less cohesive film; and (ii) excess Nb yields continuous grain-boundary (GB) Laves (purple band), which facilitates intergranular attack and undermines passivation.
Since Cu and Nb influence both the microstructure and passive film formation, they exert complementary effects: Nb facilitates the rapid formation of a dense passive layer, while Cu elevates the steady-state potential of the alloy in its passivated state. Their synergistic interaction significantly reduces the corrosion current density (icorr), decreases the corrosion rate, and enhances polarization resistance (Rp), thereby improving the electrochemical stability of the coating. In this study, as Nb content increased from 0 to 0.2 (accompanied by a decrease in Cu from 0.5 to 0.3), the corrosion potential (Ecorr) shifted positively, icorr continuously decreased, and Rp increased markedly. These findings indicate that the improvement in corrosion resistance is intrinsically linked to the structural and passivation changes induced by the co-alloying of Cu and Nb. This behavior is consistent with Fig 15, where the balanced Cu–Nb condition produces a structurally coherent, self-healing film (a), whereas Cu-rich or Nb-rich conditions deteriorate film integrity and electrochemical metrics (b).
Based on the above analysis, it is concluded that optimal corrosion resistance in AlCoCrFeNiCu0.5-xNbx coatings is achieved only when microstructure and composition are well balanced. The C3 coating exemplifies this optimal state. Its microstructure consists of an FCC matrix with a minor amount of dispersed Laves phase, exhibiting uniform elemental distribution without significant Cu enrichment or continuous Nb segregation. The moderate reduction in Cu mitigates micro-galvanic corrosion effects associated with Cu-rich regions, while the appropriate Nb content prevents excessive Laves phase precipitation at grain boundaries and reduces the lattice distortion caused by Nb segregation. As a result, the C3 coating combines the homogeneous passivation behavior of a single-phase solid solution with the barrier effect provided by finely dispersed secondary phases. The small Laves phase particles act as “pinning points,” impeding the ingress of corrosive media along grain boundaries. Additionally, the sufficient retention of Cr and Ni supports the formation of a Cr₂O₃-dominated passive layer with strong self-healing capabilities. The presence of Nb further improves the thermal stability and re-passivation efficiency of the oxide film. This structural synergy enables the C3 coating to achieve the highest corrosion potential, the lowest corrosion current, and the greatest polarization resistance among the tested compositions.
Accordingly, Fig 15(a) corresponds to the C3 condition: an FCC matrix + finely dispersed Laves “pinning” particles beneath a Cr₂O₃–Nb₂O₅–Al₂O₃ film, delivering the best combination of highest Ecorr, lowest icorr, and largest Rp among the investigated coatings; the right-hand panels depict the Cu-rich and Nb-rich departures from this optimal balance.
4. Conclusion
- 1). With increasing Nb content and decreasing Cu content, the AlCoCrFeNiCu0.5-xNbx (x = 0–0.5) high-entropy alloy (HEA) coatings evolve from a single FCC phase to an FCC + Laves dual-phase structure. The incorporation of an optimal Nb content (x = 0.2) increases the lattice constant and stabilizes the solid solution phase. However, excessive Nb promotes extensive Laves phase formation and induces significant lattice distortion.
- 2). The microstructure of the C3 coating (AlCoCrFeNiCu₀.₃Nb₀.₂) is predominantly composed of large cellular grains with blurred boundaries. Cu diffusion alleviates Nb segregation and suppresses brittle phase formation. When Nb content exceeds 0.2 (x ≥ 0.3), pronounced Nb segregation at grain boundaries leads to microstructural refinement into petal-like cellular grains, reduced density, and compromised interfacial continuity.
- 3). The coating achieves its highest hardness (696.11 HV) at x = 0.4, representing a 228.4% increase over the substrate. Both Cu and Nb contribute to the enhancement of mechanical properties: Cu improves plasticity and lubricity, while Nb enhances strength through lattice distortion and Laves phase strengthening.
- 4). At x = 0.2, the C3 coating exhibits the lowest coefficient of friction and the mildest wear features in both air and NaCl environments, primarily showing slight abrasive wear. When Nb content is low (x ≤ 0.1), a combination of adhesive and abrasive wear is observed. At higher Nb contents (x ≥ 0.3), despite increased hardness, the enrichment of brittle phases reduces toughness, resulting in deep grooves and brittle fracture.
- 5). The C3 coating demonstrates the most positive corrosion potential, the lowest corrosion current density, and the highest polarization resistance in both 3.5 wt% NaCl and 0.5 M H₂SO₄ solutions, indicating excellent electrochemical stability and re-passivation capability. These findings highlight the critical synergistic role of Cu and Nb in promoting the formation of a dense and stable passivation film.
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