Laser-cladded coatings on titanium alloys: processing-microstructure relationships, phase evolution, and interfacial mechanisms

Process principle and non-equilibrium solidification behavior

In laser cladding, a focused laser beam locally melts pre-placed or continuously delivered powder together with a shallow region of the Ti substrate, producing a transient molten pool. Representative thermal-field evolution and solidification characteristics of Ti6Al4V under laser-based deposition are shown in Figure 2, including the temperature distribution, grain-growth morphology, cooling rate, and temperature gradient. Under such processing conditions, rapid melting and solidification promote the formation of non-equilibrium microstructures, which are typically characterized by refined grains, microsegregation, supersaturated solid-solution matrices, metastable phases, and in situ reaction products^[39]. Laser cladding differs from many conventional surface treatments in its localized heat input, adjustable dilution, short processing time, limited heat-affected zone (HAZ), and ability to form metallurgical bonding with the Ti substrate. These characteristics make it a particularly suitable method for Ti-alloy surface modification^[22,39].

Figure 2. Thermal-field evolution, grain-growth behavior, and solidification characteristics of Ti6Al4V during laser-based deposition. (A) Simulated temperature field and molten-pool characteristics during single-track deposition; (B) transverse-section grain morphology showing columnar and equiaxed grains; (C) longitudinal-section grain-growth morphology along the scanning direction; (D) cooling-rate distribution within the molten pool; and (E) temperature-gradient distribution within the molten pool. This figure is adapted from Ref.^[37]. Licensed under CC-BY 4.0.

For Ti alloy substrates, this non-equilibrium solidification behavior has a dual effect. On the one hand, rapid solidification promotes grain refinement, metastable phase formation, and fine dispersion of reinforcing phases, which are beneficial for hardness and wear resistance. On the other hand, it may intensify elemental segregation, constitutional supercooling, dendritic heterogeneity, and local enrichment of brittle reaction products. These features can generate spatially heterogeneous mechanical properties and unstable crack-propagation paths. Therefore, the key issue is not merely to exploit rapid solidification for strengthening, but to control the solidification pathway so that refined microstructures are obtained without forming continuous brittle networks or highly segregated interfacial regions.

In Ti-alloy claddings, solidification often starts epitaxially from the substrate, leading to columnar grains adjacent to the bonding interface and more equiaxed grains toward the upper region of the coating^[40,41]. As shown in Figure 2, the resulting grain morphology is closely related to the spatial variation in cooling rate and temperature gradient within the molten pool. During this process, constitutional supercooling may develop in front of the advancing solid/liquid interface, promoting columnar-to-equiaxed transition (CET) when newly formed equiaxed nuclei progressively restrict the continued growth of columnar grains^[24,27]. The final growth mode is mainly controlled by the balance between the thermal gradient (G) and the solidification rate (R), which governs the transition from planar growth to cellular and dendritic morphologies^[42].

Microstructural evolution, Ti-induced interfacial reactions, and cracking susceptibility

The solidification-driven features not only define the microstructural morphology but also dictate subsequent phase evolution, element segregation, and strengthening behavior across various coating systems. Because atomic diffusion is strongly restricted during rapid solidification, solutes may be retained unevenly within the solidifying coating, resulting in microsegregation and metastable solid-solution formation^[43-46]. These microstructural characteristics significantly influence the hardness, toughness, and wear resistance of the cladding layer.

Notwithstanding these advantages, Ti alloys exhibit intrinsic chemical and thermo-mechanical characteristics that make coating formation particularly challenging. Ti’s high chemical reactivity promotes interfacial chemical ordering and intermetallic compound formation through reactions with transition metals, ceramic-forming elements, and oxygen, while the large mismatch in thermal expansion generates high residual stresses during rapid heating and cooling^[47-49]. During the short molten-pool lifetime, dilution from the Ti substrate may trigger non-equilibrium interfacial reactions, leading to the formation of Ti-Ni, Ti-Fe, Ti-Co, Ti-Si, TiC, TiB, TiN, or oxide-containing reaction products depending on the coating composition^[22,50,51]. These products are not inherently beneficial or detrimental. Finely dispersed or discontinuous reaction phases may enhance metallurgical bonding and local hardening, whereas continuous or coarse brittle reaction layers reduce interfacial deformability and become preferential sites for crack initiation and delamination. Therefore, interfacial reaction products should be evaluated according to their morphology, continuity, thickness, and compatibility with the adjacent matrix, rather than simply being regarded as strengthening phases.

Residual stress represents another critical challenge specific to laser cladding on Ti alloys. The combined influence of intermetallic compound formation and residual stress greatly increases the cracking susceptibility of laser-cladded coatings^[52,53]. Severe thermal gradients and rapid solidification during laser processing generate tensile stresses in the cladding layer, which are intensified by the mismatch in thermophysical and mechanical properties at the interface. When brittle intermetallic compounds or continuous ceramic networks are present, residual stress tends to concentrate at grain boundaries (GBs), phase interfaces, and dilution-affected regions, thereby promoting hot cracking, cold cracking, or interfacial spallation. This explains why coatings with higher hardness do not necessarily exhibit better wear resistance if their stress-accommodation capacity and interfacial toughness are insufficient. Accordingly, controlling interfacial reactions, mitigating residual stresses, and optimizing elemental partitioning remain key challenges for coating design.

Design strategies and relevance to coating systems

To overcome persistent metallurgical and thermo-mechanical challenges, particularly Ti-induced interfacial reactions, brittle reaction-layer formation, residual stress accumulation, and cracking, recent research has progressively shifted from passive process optimization toward active material system design. This transition reflects the recognition that adjusting processing parameters alone cannot fundamentally resolve the inherent incompatibilities between coating materials and Ti substrates.

The five coating systems discussed in the following sections, namely self-fluxing alloys, ceramic particle-reinforced composites, rare-earth-modified coatings, self-lubricating coatings, and HEAs, embody distinct materials design pathways developed to mitigate these inherent conflicts. However, each strategy inevitably introduces trade-offs, including hardness-toughness imbalance, interfacial chemical reactivity, and lubrication instability under complex service conditions. Consequently, coating reliability is governed not by a single strengthening mechanism but by the coordinated regulation of compositional design, phase architecture, and interfacial stability.

Collectively, these approaches establish a unified conceptual framework linking processing thermodynamics, microstructural evolution, and multi-mechanism strengthening behavior. This framework provides the scientific basis for the comparative analyses presented in the subsequent sections and underscores the necessity of synergistic design beyond isolated performance enhancement.

Nickel-based coatings

Nickel-based self-fluxing alloy coatings are among the most extensively investigated systems in wear- and corrosion-resistant surface engineering due to their mature engineering foundation and favorable comprehensive performance^[66,67]. Compared with Fe-based counterparts, Ni-based alloys exhibit superior toughness and corrosion resistance while retaining relatively high hardness. In contrast to Co-based systems, they offer lower material cost and a wider processing window^[49,68]. Consequently, Ni-based self-fluxing coatings are widely considered as versatile materials capable of balancing mechanical performance and economic efficiency across diverse service conditions.

In terms of compositional design, typical Ni-based self-fluxing alloys are built around a γ-Ni solid-solution matrix. B, together with Si, is introduced to lower the melting temperature and improve melt fluidity, while Cr and C favor in situ precipitation of carbides and borides^[33,61]. Rapid cooling during laser cladding promotes the formation of a solid solution, borides, carbides, and minor Ni-Si intermetallic compounds. This multiphase constitution reflects a typical non-equilibrium phase-selection process in laser cladding, in which rapid melting, dilution with the Ti substrate, and subsequent solidification jointly determine the formation of borides, carbides, and intermetallic products. In this sense, the wear behavior of Ni-based self-fluxing coatings should also be interpreted in relation to phase evolution and interfacial reaction pathways. This multiphase architecture, characterized by a continuous tough matrix reinforced by dispersed hard phases, enables an effective balance among hardness, toughness, and crack resistance. Representative studies further indicate that the effect of compositional optimization in Ni-based self-fluxing coatings is not monotonic. Fan et al.^[33] showed that increasing the Cr and B contents from Ni40 to Ni60 promoted the formation of carbide- and boride-type hard phases, thereby increasing coating hardness and reducing the COF. However, the wear resistance did not improve monotonically, and the Ni50 coating exhibited the best overall tribological performance^[33]. This behavior suggests that the optimum reinforcement level is governed not simply by hardness, but by the balance between hard-phase strengthening and microstructural integrity. At moderate hard-phase fractions, dispersed carbides and borides can effectively share the contact load and suppress matrix plowing, while the γ-Ni matrix still retains sufficient deformability to alleviate local stress concentration. By contrast, when the hard-phase fraction increases further, the reinforcing phases tend to become coarser and less mechanically compatible with the surrounding matrix, which promotes particle fracture, interfacial decohesion, and abrasive debris generation during sliding^[33,69]. In this case, the hard phases no longer act only as load-bearing units, but also as potential crack-initiation and micro-cutting sources. Therefore, the better overall behavior of the Ni50 coating can be understood as the result of a more favorable balance between reinforcement efficiency and damage tolerance, whereas the inferior wear behavior of Ni60 reflects the transition from beneficial hard-phase dispersion to excessive brittle reinforcement.

Mechanistically, Ni-based self-fluxing coatings resist wear through the combined contribution of matrix hardening and load support by dispersed hard phases. The γ-Ni matrix accommodates part of the applied stress through plastic deformation, thereby mitigating local stress concentration and delaying crack initiation, while dispersed carbides and borides sustain the majority of the contact load and suppress plowing and micro-cutting wear^[33,70]. Consistent with this mechanism, Wang et al.^[61] fabricated a laser-cladded Ni60 coating on TC4 Ti alloy, where in situ-formed TiC and Cr₃C₂ increased the hardness to 837 HV₁ (~2.3 times that of the substrate). The synergistic strengthening reduced material loss from 115 mg to 21 mg and shifted the dominant wear mechanism from severe adhesive wear to mild abrasive wear, highlighting the effectiveness of multiphase strengthening.

Cobalt-based coatings

Cobalt-based self-fluxing alloy coatings, typically represented by the Stellite series, are widely applied in demanding service environments, including aircraft-engine parts, gas-turbine components, and high-temperature valve seats due to their superior high-temperature mechanical stability and wear reliability^[71]. Although their material cost exceeds that of Fe- and Ni-based counterparts, Co-based alloys provide distinct advantages under severe operating conditions, particularly in retaining hardness at elevated temperatures, resisting adhesive wear, and tolerating thermal fatigue^[71,72].

Microstructurally, Co-based self-fluxing alloys generally consist of a face-centered cubic (FCC) Co solid solution matrix reinforced by dispersed Cr-rich carbides, predominantly Cr₇C₃ and M₂₃C₆^[64,65]. This phase constitution is also a result of non-equilibrium solidification and interfacial dilution during laser cladding, in which rapid cooling and local compositional redistribution govern the selection, dispersion, and stability of carbide phases within the FCC Co matrix. The FCC Co matrix exhibits good plastic deformability and strain accommodation at elevated temperatures, thereby alleviating stress concentration during thermal cycling. Simultaneously, high-melting-point carbides form a rigid, load-bearing skeleton that preserves structural integrity under frictional loading^[63,65]. Representative studies have further clarified these strengthening mechanisms. Liu et al.^[65] fabricated a Stellite 6 coating on TC4 Ti alloy via laser cladding and reported a microstructure composed of an FCC γ-Co matrix with dispersed M₂₃C₆, Cr₇C₃, TiC, and CoC_x phases. The incorporation of Fe, Ni, and C stabilized the γ-Co structure during thermal cycling, enhancing strain accommodation through lattice distortion under residual tensile stress. From a crystallographic perspective, this stabilization is important because it implies that the matrix can retain a deformation-tolerant crystal structure under thermal and mechanical loading, thereby reducing the tendency for premature crack initiation at stress-concentrated regions. This synergistic matrix-carbide architecture resulted in an average microhardness of 542 HV, significantly reducing the COF relative to TC4.

From a tribological perspective, Co-based self-fluxing coatings exhibit strong resistance to adhesive and fatigue wear, particularly at elevated temperatures. During high-temperature sliding, the Co matrix exhibits pronounced work-hardening, enabling continuous renewal of load-bearing regions via surface plastic deformation. In parallel, Cr-rich and Co oxides formed at the friction interface form a dense, stable tribo-oxide layer that mitigates adhesive failure. Shi et al.^[63] reported that Co12 coatings undergo a transition from abrasive wear at room temperature (RT) to oxidation-dominated wear at elevated temperatures. At 500 °C, the formation of a dense glaze layer reduced the COF to approximately 0.35 via third-body rolling, resulting in wear resistance superior to that at RT.

Iron-based coatings

Fe-based self-fluxing systems are mainly valued for their cost advantage, since their material cost is usually much lower than that of Ni- or Co-based coatings, while they can still offer high hardness and good abrasive-wear resistance at RT^[73].

From a microstructural design perspective, boron (B) functions not only as a fluxing element but also as a key regulator of reinforcing phase evolution. Because B has low solubility in the metallic matrix and diffuses slowly in hard phases, local B enrichment can reduce the energy barrier for phase formation and promote phase transformation. Conventional Fe-based hardfacing materials tend to form brittle carbides such as M₂C, MC, or M₆C. With suitable B addition, however, the activity of carbon in the melt increases, which promotes M₃C-type eutectic precipitation, such as Fe₃C, with relatively better fracture toughness^[74].

In Fe-based hardfacing coatings, Feng et al.^[74] reported that increasing B content induces a transition of reinforcing phases from NbC or Cr₂₃C₆ carbides toward Fe₃C-type carbides. When the B content reached 0.75%, a Fe₃C + γ-Fe eutectic microstructure was formed, enabling the microhardness to exceed 900 HV_0.5. Under rapid solidification, a skeletal architecture composed of blocky or plate-like M₇C₃ and Fe₂B phases developed, supported by a martensitic or bainitic composite matrix. This multiphase framework exhibits strong room-temperature wear resistance^[75].

Despite these advantages, Fe-based self-fluxing coatings are seldom applied to Ti alloys due to severe metallurgical incompatibility and thermophysical mismatch at elevated temperatures. Owing to the strong Ti-Fe affinity, laser cladding induces the rapid formation of brittle interfacial Ti-Fe intermetallics, while significant thermal expansion mismatch generates tensile residual stresses upon cooling, frequently resulting in cracking or even coating failure^[76]. Even when crack-free coatings are achieved, a fundamental performance conflict persists: Ti alloys are commonly employed in medium- to high-temperature environments, whereas Fe-based coatings undergo rapid matrix softening and loss of wear resistance above approximately 400 °C, limiting their practical applicability for Ti substrates.

Nevertheless, Fe can serve as a reactive alloying element to modify in situ phase evolution in Ti-based systems. Zhao et al.^[51] fabricated a Ti-40SiC transition layer on TC4, followed by a Ti-xFe-80SiC working layer with varying Fe contents (0-15 vol.%). Their results showed that Fe addition altered the reaction pathway of the Ti-Si-C system. At higher Fe contents (10 and 15 vol.%), an ultra-hard TiFeSi₂ phase formed in situ, significantly enhancing coating performance^[51]. Furthermore, Fe-based alloys improved wettability and self-fluxing behavior, effectively reducing pores and cracks. As the Fe content increased, defect density within the coating decreased markedly [see Figure 3]. These findings indicate that controlled Fe incorporation can tailor phase evolution and construct ultra-hard wear-resistant architectures, offering an alternative design strategy for Ti-based composite coatings.

Figure 3. SEM images of the cross-sectional microstructure of the coatings: (A) WL-F0; (B) WL-F5; (C) WL-F10; (D) WL-F15; (A₁-D₁) Higher-magnification views of the regions in (A-D), respectively; (A₂-D₂) Further enlargements of (A₁-D₁), respectively. This figure is adapted from Ref.^[51]. Licensed under CC-BY 4.0. SEM: Scanning electron microscopy.

Overall, self-fluxing systems are typically characterized by a ductile metallic matrix, such as γ-Ni- or γ-Co-based phases, reinforced by in situ formed carbides and/or borides generated during rapid solidification. Their wear-resistant performance mainly originates from the synergistic effects of matrix-supported stress accommodation, hard-phase load bearing, and strong metallurgical bonding.

Although self-fluxing systems are important in surface engineering because they can provide good resistance to both wear and corrosion, their inherent limitations become pronounced when applied to highly reactive substrates such as Ti alloys^[48,49]. The drawbacks of traditional self-fluxing coatings on Ti substrates mainly arise from three mutually reinforcing factors.

First, the high-volume fraction of carbides and borides required to improve wear resistance inevitably increases microstructural brittleness and reduces impact toughness, making self-fluxing alloy coatings susceptible to cracking^[63]. As shown in Figure 4A, cracking can be observed in a representative Ni-based self-fluxing alloy coating^[69]. For Fe- and Co-based coatings, worn surface morphologies further reveal pitting, spalling, peeling, and cracking, indicating that crack formation is closely coupled with tribological damage during sliding contact [Figure 4B and C]^[77]. Second, reactions between Ti and transition metals promote the formation of brittle Ti-Ni, Ti-Co, and Ti-Fe intermetallic compounds near the coating/substrate interface, which favor crack nucleation and interfacial debonding^[78-80]. Third, mismatches in thermal expansion and elastic properties generate high tensile residual stresses under steep thermal gradients. As schematically illustrated in Figure 4D-F, crack formation can be associated with thermal stress, organizational stress, and restraint stress. These stresses accelerate interfacial damage and crack propagation. Although optimized powder feeding and process control can partially suppress cracking, process-parameter optimization alone cannot fully resolve the intrinsic incompatibility between conventional self-fluxing alloys and Ti substrates.

Figure 4. Cracking, spalling, and stress-induced failure mechanisms of laser-cladded self-fluxing alloy coatings. (A) Crack morphology in a Ni-based self-fluxing alloy coating; (A) is adapted from Ref.^[69]. Licensed under CC-BY 4.0; (B) worn surface morphology of a Fe-based coating showing pitting, spalling, and cracking; (C) worn surface morphology of a Co-based coating showing slight spalling and cracking; (B and C) are adapted from Ref.^[77]. Licensed under CC-BY 4.0; (D-F) schematic illustrations of crack formation induced by thermal stress, organizational stress, and restraint stress, respectively. (D-F) are adapted from Ref.^[81]. Licensed under CC-BY 4.0.

Taken together, these limitations have driven the development of composite and multiphase coatings that integrate ceramic load-bearing frameworks, rare-earth oxide grain refiners, solid-lubricant phases, and high-entropy matrices. These strategies aim to balance hardness, toughness, lubrication, and interfacial stability through synergistic microstructural design, forming the basis for the following sections on ceramic reinforcement, rare-earth modification, self-lubricating coatings, and HEA systems.

Carbide-reinforced systems

Carbide reinforcements are primarily introduced to suppress severe abrasive and erosive wear by constructing a high-hardness load-bearing skeleton within the cladding layer. This strategy is particularly effective under wear regimes dominated by micro-cutting and plowing, provided that the service temperature remains within a range where the metallic matrix retains sufficient strength^[102,103].

Carbide-reinforced coatings can be produced either by ex situ incorporation of stable carbides (e.g., TiC, WC, and SiC) or via in situ formation during reaction-assisted cladding. In situ TiC typically forms because of the strong chemical affinity between Ti and C, for example, via reactions involving graphite, B₄C, or carbide-containing powders^[85-87]. WC may partially dissolve in the molten pool and reprecipitate as (Ti, W)C-type carbides, while SiC can react with Ti to generate TiC, accompanied by Ti-Si intermetallic or silicide phases, depending on local composition and thermal history^[94,104,105]. Representative TiC-, SiC-, and WC-reinforced coating microstructures and their hard-phase-assisted wear-resistance mechanisms are summarized in Figure 5, highlighting the important role of carbide ceramics in microstructural strengthening and load-bearing protection. These phase reactions show that carbide reinforcement in laser cladding is closely associated with non-equilibrium solidification and local solute redistribution in the molten pool. The final reinforcing architecture is governed by the competition among dissolution, interfacial reaction, and carbide reprecipitation during rapid cooling.

Figure 5. Representative microstructures and wear-resistance mechanism of carbide-reinforced laser-cladded composite coatings. (A and B) Cross-sectional microstructures of TiC-reinforced coatings; (A and B) are adapted from Ref.^[69]. Licensed under CC-BY 4.0; (C) cross-sectional morphology and basket-weave microstructure of a SiC-containing laser-cladded coating; (D) high-magnification SEM image showing α-laths, residual SiC particles, and Ti₃Al phase in the SiC-containing coating; (C and D) are adapted from Ref.^[36]. Licensed under CC-BY 4.0; (E) schematic illustration of the WC-assisted wear-resistance mechanism during sliding contact. (E) is adapted from Ref.^[94]. Licensed under CC-BY 4.0. SEM: Scanning electron microscopy; WC: tungsten carbide.

The morphology, size, and spatial distribution of carbide phases critically govern coating behavior and damage tolerance. Wei et al.^[106] synthesized in situ TiC/Ti composite coatings on Ti6Al4V using graphene as a carbon precursor. With increasing carbon content, the TiC morphology evolved from short, fibrous structures to near-spherical particles, accompanied by increases in volume fraction and hardness. This microstructural transition improved wear resistance by enhancing load sharing and stabilizing the matrix. Such observations indicate that carbide efficiency is controlled not simply by phase fraction, but by whether the carbide phase remains finely dispersed or develops into a more continuous load-bearing framework.

Consistently, increasing the TiC fraction together with grain refinement leads to higher microhardness and a reduced COF, as reflected by the systematic variation in hardness and COF with SiC content. Fine carbide particles promote Orowan strengthening and increase dislocation density through thermal-misfit effects, whereas coarser carbides enhance the effective composite modulus and share contact loads^[107-109]. During sliding, a well-developed carbide skeleton stabilizes the near-surface deformation layer and can shift the wear mode from severe adhesive or oxidative wear to controlled abrasive wear, thereby reducing wear-track depth and counter-body damage^[110-113]. Wear-track profiles and wear-rate measurements further indicate the existence of an optimal carbide fraction: moderate reinforcement maximizes load sharing and minimizes material removal, while excessive ceramic content deteriorates structural integrity and promotes crack initiation. Wang et al.^[87] further confirmed these strengthening mechanisms in SiC-reinforced Ti-based composite coatings fabricated on Ti6Al4V via laser cladding. Rapid SiC dissolution and strong Ti-C affinity promoted in situ TiC formation as the dominant reinforcing phase, while residual SiC and reaction products provided heterogeneous nucleation sites. Enhanced carbon activity and melt-pool convection promoted TiC growth into dendritic and flower-like morphologies, resulting in refined microstructures and significantly improved hardness. The resulting improvement in tribological performance therefore arose not only from increased hardness, but also from grain refinement, load transfer by TiC, and stabilization of the near-surface deformation layer.

Despite the substantial hardness enhancement provided by carbide reinforcements, wear resistance does not increase monotonically with reinforcement content. The key factor is not the carbide fraction itself, but whether the reinforcing phase remains finely dispersed and mechanically compatible with the surrounding matrix. At moderate reinforcement levels, TiC particles can effectively share contact stresses, suppress matrix plowing, and stabilize the near-surface deformation layer during sliding, thereby improving wear resistance^[87,100,114]. However, once the carbide content exceeds an optimal range, the reinforcing architecture gradually loses its damage tolerance. Particle agglomeration, incomplete melting, and pore formation become more pronounced, while the growing continuity and coarsening of the hard phase increase local stress concentration^[87,114]. Under such conditions, the carbide phase no longer acts only as a load-bearing constituent, but also becomes a preferred site for crack initiation and brittle fracture. Cai et al.^[114] reported that increasing the TiC content in TiC-reinforced Ni60 composite coatings reduced both wear resistance and flexural strength, indicating that excessive, coarse TiC particles can promote brittle failure rather than further strengthening. Similarly, Wang et al.^[87] showed that excessive SiC addition during in situ TiC formation led to crack development because agglomerated particles and associated pores created stress-concentration zones under rapid solidification. Therefore, the deterioration in wear performance at excessive carbide contents originates fundamentally from the microstructural transition of the coating from a dispersed strengthening architecture to a defect-assisted brittle network.

These findings indicate that the effectiveness of carbide reinforcement is controlled by the balance between hard-phase strengthening and structural integrity. Accordingly, the design of carbide-reinforced composite coatings should focus on controlling phase fraction, particle morphology, and distribution uniformity, for example, through compositional gradient design, controlled particle-size distribution, or the incorporation of ductile binder phases, to retain high load-bearing capacity while mitigating brittleness. More fundamentally, carbide-reinforced coatings are effective only when phase constitution, reinforcing morphology, and defect distribution are jointly optimized to prevent the hard ceramic framework from evolving into a continuous brittle crack-initiation path.

Boride-reinforced systems

Boride reinforcements are introduced to enhance hardness while retaining higher fracture tolerance compared with carbide-dominated architectures^[115,116]. For Ti alloys, the TiB/TiB₂ system is particularly attractive because the in situ formation of TiB whiskers or needles can generate a semi-continuous reinforcing framework, enabling efficient load transfer while mitigating crack propagation under impact-sensitive wear and thermal cycling^{[84,115,117,118]}.

Boride-reinforced coatings are fabricated either through ex situ addition of TiB₂ particulates or via in situ reactions from B-containing precursors such as B₄C, BN, or elemental B during laser cladding. Typical microstructures exhibit a multiscale architecture containing primary TiB, eutectic TiB whiskers, and residual TiB₂ embedded within an α-Ti matrix. As shown in Figure 6, scanning electron microscopy (SEM) and transmission electron microscopy (TEM) observations reveal TiB/TiB₂-reinforced microstructures and the interfacial characteristics between α-Ti and TiB, while the corresponding selected-area electron diffraction (SAED) patterns further confirm the crystallographic characteristics of the α-Ti matrix and TiB phase. These interfacial features are closely related to load transfer and crack-propagation behavior in boride-reinforced coatings. Representative studies have demonstrated the feasibility of gradient and composite designs; for example, Lin et al.^[99] produced a gradient coating containing 30% TiB₂/Ti on Ti alloy substrates, while Li et al.^[84] prepared multiple TiB-reinforced titanium matrix composite coatings via in situ laser cladding to improve wear performance.

Figure 6. Representative microstructures and interfacial crystallography of TiB/TiB₂-reinforced laser-cladded composite coatings. (A) Morphology of TiB₂ particles; (B) microstructure showing TiB and residual TiB₂ phases in the coating; (C and D) TEM images showing the α-Ti/TiB interfacial characteristics; (E and F) SAED patterns corresponding to the α-Ti matrix and TiB phase, respectively. This figure is adapted from Ref.^[99]. Licensed under CC-BY 4.0. TEM: Transmission electron microscopy; SAED: selected-area electron diffraction.

The hallmark of TiB-reinforced claddings lies in the development of a multiscale whisker network inside a rapidly solidified matrix. High thermal gradients near the substrate promote epitaxial growth and columnar alignment of TiB whiskers, whereas reduced gradients and solute redistribution favor more equiaxed morphologies in upper regions^[83]. Clean interfaces and favorable wettability between TiB whiskers and Ti-rich matrices facilitate effective load transfer and suppress interfacial decohesion during wear^[119,120]. With increasing TiB₂ addition, the microstructure evolves from fine eutectic whiskers to a mixed architecture containing primary TiB and residual TiB₂ particles. Moderate reinforcement levels generate an interlocked whisker-like TiB framework that enhances load transfer and effectively impedes crack propagation. In this condition, the TiB phase acts as a crack-deflection and crack-bridging unit, forcing propagating cracks to change direction or become blunted, thereby improving fracture toughness and wear resistance. However, with further increases in boride content, the morphology and spatial distribution of TiB change significantly. Coarsening and local connectivity of whiskers may gradually develop into a quasi-continuous brittle network. Under such conditions, the role of TiB shifts from inhibiting crack propagation to facilitating it, as cracks can propagate preferentially along or through the interconnected boride structure. This transition from crack-deflection-dominated behavior to crack-propagation-assisted failure is responsible for the observed reduction in fracture toughness and wear resistance at excessive reinforcement levels^[84].

Beyond simple load transfer, TiB networks contribute to crack deflection, bridging, and pull-out resistance, as the interlocking skeleton redistributes local stresses^[121,122]. Wear cross-sectional analyses further reveal a gradual transition in wear mechanism: from matrix-dominated plastic deformation to skeleton-controlled abrasion with increasing reinforcement content. The whisker architecture also induces back-stress hardening, as deformation incompatibility between whiskers and the matrix generates geometrically necessary dislocations that improve the strength-ductility synergy^[123,124]. Compared with TiC, TiB exhibits a closer thermal-expansion match with Ti and higher intrinsic fracture tolerance, thereby reducing susceptibility to thermal cracking during processing and service^[125,126].

However, excessive boron activity or overly high TiB₂ addition can lead to whisker coarsening, increased melt viscosity, clustering, and lack-of-fusion defects, thereby elevating crack susceptibility^[127]. Schematic analyses suggest that while moderate TiB content forms an interlocked network capable of effectively distributing contact stresses, excessive reinforcement results in brittle phase continuity and particle fracture. Consistent with this trend, Wang et al.^[128] reported that although increasing TiB₂ content improved microhardness in laser melting deposition coatings on TiAl alloys, fracture toughness decreased markedly at high addition levels.

An effective design window for boride-reinforced coatings requires controlled whisker aspect ratio and spacing, avoidance of continuous brittle networks, and regulation of thermal gradients through optimized energy density, track overlap, and preheating when necessary. These factors govern not only hardness development but also the stability of the whisker framework and the coating’s capacity to relieve stress before premature cracking occurs. The goal is to obtain a better balance among hardness, crack resistance, and structural reliability.

Nitride-reinforced systems

Nitride reinforcements are introduced to enhance wear resistance under oxidative and high-temperature conditions while reducing friction through chemically stable, low-shear interfaces^[129,130]. Among them, TiN is particularly suitable for Ti substrates because it combines high hardness and thermal stability with relatively good fracture tolerance and wettability, making it favorable for constructing stable hard-phase architectures in Ti-based cladding layers^[131,132].

TiN-containing cladding layers are typically fabricated through three principal routes: (i) direct ex situ incorporation of TiN into the cladding powder^[133]; (ii) laser-assisted in situ nitriding under N₂-rich shielding gas or reactive atmospheres^[131]; (iii) introduction of nitrogen-containing precursors, such as boron nitride (BN) or Si₃N₄, which react with molten Ti to form TiN during cladding^[134]. In precursor-assisted systems, BN can simultaneously introduce boron into the melt pool, resulting in the concurrent formation of TiB as a secondary reinforcement. Under such conditions, TiN, Ti(C,N), sub-stoichiometric titanium nitrides, and nano-scale TiB may coexist, producing a composite nitride-boride architecture that contributes to microstructural refinement and improved tribological performance^[134].

Rapid solidification during laser cladding promotes the formation of fine TiN particles within the Ti matrix. Optimal performance is generally associated with a fine and discontinuous distribution of TiN-based phases, which enhances second phase strengthening and nitrogen-induced solid-solution strengthening while avoiding continuous brittle nitride bands^[83]. In BN-derived systems, TiB can interlock with TiN particles and help deflect or arrest cracks, thereby reducing the crack sensitivity of otherwise brittle TiN-rich regions^[135,136].

Fine TiN dispersions increase dislocation density, promote grain refinement, and improve load-bearing capacity. In addition, chemically stable nitride surfaces can facilitate the formation of oxide- or oxynitride-rich tribofilms, which reduce interfacial shear stress during sliding^[83,137]. However, excessive nitride or nitride-boride content may cause particle coarsening, brittle connectivity, porosity, and localized corrosion or cracking, especially under strong compositional gradients or repeated cladding passes. Therefore, the effectiveness of nitride reinforcement depends not only on phase fraction, but also on the size, continuity, and spatial distribution of nitride-containing regions within the coating.

The primary risks associated with nitride systems include nitrogen-related porosity and residual-stress-driven cracking under steep thermal gradients and thermomechanical mismatch^[138,139]. Accordingly, an effective design window requires balancing nitrogen activity to ensure adequate TiN formation while maintaining melt-pool stability to prevent gas entrapment. Control of particle fraction, spatial distribution, and thermal gradients through powder design, shielding strategy, and energy-density optimization is essential for achieving structural integrity and long-term reliability. From a microstructural standpoint, the key challenge is to retain the strengthening benefits of nitride phases while preventing defect accumulation, brittle connectivity, and interfacial instability during processing and service.

Overall, ceramic-particle-reinforced coatings are characterized by metallic matrices containing dispersed or interconnected carbide, boride, and/or nitride phases that provide substantial improvements in hardness and load-bearing capacity. Their tribological performance nevertheless depends strongly on reinforcement fraction, spatial distribution, and matrix/ceramic interface stability, because excessive phase connectivity may reduce damage tolerance and compromise coating integrity. These limitations have encouraged the incorporation of complementary regulatory mechanisms, including rare-earth grain refinement, self-lubricating phases, and high-entropy matrices, to achieve more balanced and stable service performance.

CeO₂-modified coatings

As a typical active rare-earth oxide, CeO₂ is often introduced into Ti-alloy coating systems not as a primary load-bearing phase, but to regulate microstructural evolution. Its design objective is to optimize melt-pool behavior, solidification dynamics, and interfacial stability during laser cladding. CeO₂ possesses good thermal stability and sufficient chemical activity, which allows it to participate in melt purification and interfacial reactions^[149]. In addition, the reversible Ce³⁺/Ce⁴⁺ redox capability mitigates the detrimental effects of oxygen on coating integrity, thereby indirectly enhancing wear resistance and long-term service stability^[150,151].

During laser cladding, CeO₂ is typically introduced via ex situ powder addition. Owing to its high melting point, CeO₂ particles remain largely solid or partially dissolved in the molten pool, where they may undergo local reduction or interfacial reactions^[142]. A portion of Ce segregates toward GBs and melt-pool edges, reacting with impurity elements such as O and S to reduce their effective concentration^[152]. Simultaneously, CeO₂ particles and reaction products promote grain refinement and modify solidification pathways^[149].

Microstructural modification following CeO₂ addition commonly manifests as reduced columnar grain fractions and increased equiaxed grains or fine dendrites, thereby improving structural uniformity. Wang et al.^[142] reported that introducing 3 wt.% nano-CeO₂ into TiCN/SiO₂ composite coatings produced a refined and compact microstructure, increasing microhardness to 1,230 HV_0.2 and reducing wear volume loss by 14% in simulated body fluid. In the absence of CeO₂, coarse dendritic grains dominated the coating thickness, whereas CeO₂ addition significantly refined dendrites and improved microstructural homogeneity. GB Ce segregation suppressed continuous brittle phase formation and reduced defect density, thereby enhancing coating stability^[153]. Additionally, Gong et al.^[154] observed that 3 wt.% CeO₂ addition to Ni60A coatings on TC4 yielded crack-free, pore-free layers with a microhardness of 1,164 HV.

The regulatory effect of CeO₂ exhibits a distinct concentration window. At excessive addition levels, its ability to effectively regulate melt-pool dynamics and microstructural refinement becomes progressively limited. Instead, particle agglomeration, localized embrittlement, and increased substrate dilution may occur, replacing the beneficial refinement effect observed at moderate additions^[50,155]. Liu et al.^[50] demonstrated that moderate CeO₂ addition eliminated characteristic herringbone cracks and promoted TiC spheroidization, transforming dendritic TiC morphologies into more homogeneous granular structures. However, with further increases in CeO₂ content, microstructural coarsening was observed^[149]. Elemental analyses revealed Ce enrichment along TiC dendrites, consistent with Ti-Ce mutual solubility at elevated temperatures under non-equilibrium cladding conditions. This indicates that excessive Ce alters the solidification pathway, shifting it from refinement-dominated behavior toward dendritic growth, thereby weakening the beneficial effects observed at intermediate additions.

Additional studies indicate that excessive CeO₂ may increase the dilution rate by enhancing melt-pool fluidity, which in turn leads to reduced hardness at high addition levels^[144]. Consequently, the effectiveness of CeO₂ modification is intrinsically limited to a controlled compositional window. Since CeO₂ does not directly participate in primary load bearing, its primary role lies in regulating microstructural evolution and defect formation, rather than providing direct mechanical reinforcement. Therefore, its beneficial effect depends on maintaining a balance between refinement efficiency and structural stability through synergistic interaction with both hard phases and the metallic matrix.

Y₂O₃-modified coatings

The primary objective of Y₂O₃-modified coatings is to enhance microstructural stability, interfacial bonding quality, and reliability at high temperatures. Compared with CeO₂, Y₂O₃ exhibits higher thermal stability and greater chemical inertness, making it particularly suitable for elevated-temperature and complex service environments^[35,156]. Rather than acting as an active melt regulator, Y₂O₃ primarily serves as a structural stabilizer, strengthening the microstructural foundation of wear-resistant coatings through grain-boundary regulation and interfacial control.

During laser cladding, Y₂O₃ particles are typically introduced as dispersed solids. Owing to their high melting point and limited solubility, they largely remain undissolved within the molten pool, while a small fraction of Y undergoes limited diffusion and interfacial reactions. Y₂O₃ particles preferentially segregate at GBs and exert a pinning effect that inhibits grain growth. Although heterogeneous nucleation has been proposed in earlier interpretations, lattice-misfit analyses indicate that the dominant mechanism is grain-boundary pinning and dragging rather than nucleation-rate enhancement^[35]. This restriction on solid-liquid interface migration promotes significant grain refinement, leading to transformation from coarse columnar structures to fine equiaxed grains or dendrites and resulting in a denser, more homogeneous microstructure^[157].

Experimental evidence supports this mechanism. Zhang et al.^[145] reported that Y₂O₃ addition to laser-cladded Ti6Al4V coatings completely suppressed pore formation and improved overall forming quality. The refined microstructure, predominantly composed of acicular α’-Ti, exhibited disrupted β-Ti GB continuity due to adsorption and pinning effects. Elemental mapping and lattice mismatch evaluation revealed Y and O enrichment along original β-Ti grain boundaries, while the lattice mismatch between Y₂O₃ and β-Ti exceeded 12%, excluding effective heterogeneous nucleation. These findings confirm that refinement originates primarily from grain-boundary pinning rather than increased nucleation.

Beyond microstructural refinement, Y₂O₃ improves interfacial bonding and crack resistance. By regulating interfacial reactions, it reduces the thickness of brittle interfacial layers and alleviates residual stress concentration, thereby enhancing structural stability during service. Zhang et al.^[157] demonstrated that Y₂O₃ addition in TiC-Ti6Al4V composite coatings induced the formation of numerous tertiary dendrites and improved metallurgical bonding. The semi-coherent epitaxial relationship between TiC particles and the β-Ti matrix further strengthened interfacial adhesion, thereby improving resistance to dislocation motion and delamination under frictional loading. Likewise, surface and cross-sectional observations revealed that moderate Y₂O₃ additions suppressed crack formation and porosity, significantly enhancing coating integrity^[146].

As with CeO₂, the strengthening effect of Y₂O₃ exhibits a concentration-dependent window. Moderate additions refine the microstructure and, by enhancing structural stability, indirectly improve wear resistance. However, excessive Y₂O₃ content leads to particle agglomeration, localized embrittlement, and surface instability. At high concentrations, crack reappearance and localized surface collapse have been observed, indicating that over-addition disrupts the beneficial pinning mechanism and compromises mechanical integrity. Hence, Y₂O₃ functions most effectively as a microstructural regulator when combined with hard reinforcing phases, enabling a better balance among hardness, toughness, and high-temperature reliability.

In summary, rare-earth-oxide-modified coatings are characterized by refined, dense, and more homogeneous microstructures resulting from the regulation of melt-pool behavior, solidification, and interfacial reactions. Their principal contribution lies in improving microstructural quality and interfacial stability rather than providing direct load-bearing reinforcement. Consequently, although rare-earth oxides can enhance the structural basis of wear resistance, they have limited ability to directly regulate friction or suppress severe wear under high loads, dry sliding, or elevated temperatures.

For this reason, coating design strategies have gradually extended beyond structural stabilization toward more active control over friction and wear. One promising direction is to design self-lubricating coating systems by incorporating solid lubricants into wear-resistant matrices, enabling stable transfer films to form during sliding to reduce COF and suppress adhesive wear.

At the same time, increasingly demanding service environments require coating systems with more comprehensive performance. Because a single strengthening mechanism is often insufficient to simultaneously satisfy the demands for wear resistance, thermal stability, and environmental adaptability, HEA coatings have attracted growing attention. Through multi-principal-element alloying, high mixing entropy effects, and enhanced structural stability, these systems offer unique potential for synergistically improving wear resistance, thermal stability, and environmental tolerance under complex service conditions, and therefore represent an important emerging direction in Ti alloy surface modification.

Solid lubricant synergistic composite coatings

Currently, solid lubricants applied to Ti alloy surfaces mainly include layered materials such as graphite, h-BN, MoS₂, and WS₂; soft metals including Ag, Au, and Cu; and alkaline-earth fluorides such as CaF₂ and BaF₂.

Layered materials

Layered materials represent the earliest and most extensively studied class of solid lubricants in self-lubricating coatings for Ti alloys. Their lubrication mechanism originates from weak van der Waals bonding between adjacent atomic layers within their lamellar crystal structures. During sliding, these materials readily undergo interlaminar shear along basal planes, which helps develop continuous or partially continuous transfer films at the contact interface. Such films significantly reduce interfacial shear resistance and mitigate adhesive wear and micro-cutting, which are common in Ti alloys under dry sliding conditions. Therefore, the effectiveness of layered lubricants depends not only on their intrinsic low-shear crystal structure, but also on their ability to form and maintain a stable tribofilm during service.

Liu et al.^[159] prepared a series of TiN/WS₂ + h-BN/NiCrBSi self-lubricating coatings with different WS₂/h-BN ratios to evaluate the effect of solid-lubricant composition on wear behavior. As shown in Figure 9, the optimized coating exhibited a shallower and flatter wear scar, indicating that an appropriate WS₂/h-BN balance can reduce material removal during sliding.

Figure 9. Wear-scar morphology and wear volume of TiN/WS₂ + h-BN/NiCrBSi self-lubricating coatings with different solid-lubricant compositions. (A-E) Wear-scar morphologies and wear volumes of coatings 1-5 obtained by 3D microscopy, respectively. Coatings 1-5 were prepared with different WS₂/h-BN ratios while maintaining NiCrBSi and TiN as the matrix and hard reinforcing phases. This figure is adapted from Ref.^[159]. Licensed under CC-BY 4.0. h-BN: Hexagonal boron nitride.

The temperature-dependent evolution of tribofilms in graphene-containing coatings provides a representative example. Zhang et al.^[165] reported that tribological performance is governed by the composition and architecture of the interfacial film. At room temperature, the tribofilm mainly consists of graphene nanosheets with a small amount of oxides. SEM, elemental mapping, and Raman analyses confirmed the formation of a dense carbon-rich transfer film that preserves the lamellar structure of graphene within the wear track. This low-shear barrier minimizes direct metal contact and suppresses adhesive wear.

At elevated temperatures, however, the tribofilm composition undergoes a pronounced transition. X-ray photoelectron spectroscopy (XPS) analysis at 600 °C revealed the dominance of oxide phases, including CuO, Cr₂O₃, CoWO₄, TiO₂, and SiO₂. Although graphene retains partial lubricating functionality, oxide products become the primary load-bearing and protective constituents. The resulting graphene-oxide composite tribofilm enhances oxidation resistance and stabilizes the sliding interface. Therefore, lubrication is governed mainly by a graphene-dominated low-shear mechanism at room temperature but shifts toward an oxide-dominated protective mechanism at high temperature. This transition reflects an intrinsic trade-off between shear accommodation and chemical stability in layered lubricants and indicates that lubrication in such systems is controlled by the temperature-dependent evolution of tribofilm chemistry and architecture.

Similar synergistic behavior has been reported for MoS₂-containing systems. Qu et al.^[163] observed that when MoS₂ content exceeded 6 wt.%, sliding-induced reactions generated a dense TiS₂ surface layer, which facilitated the formation of a stable oxide film. The TiS₂ layer reduced friction, while the oxide layer provided structural protection, cumulatively suppressing micro-cutting and reducing wear loss. This result further suggests that the lubricating function of layered phases is often coupled with interfacial reaction products rather than being solely determined by the original lubricant phase.

Despite their effectiveness at moderate temperatures, layered materials exhibit limited thermal and chemical stability. Graphite and MoS₂ are susceptible to oxidation or decomposition in oxygen-rich environments at elevated temperatures, leading to diminished lubrication performance. Consistent with this limitation, Zhang et al.^[165] reported higher COF and wear rates at elevated temperatures than at RT, indicating that the long-term effectiveness of layered lubricants depends not only on their low-shear crystal structure but also on the stability and continuity of the tribofilm formed under service conditions.

MAX phase ceramic self-lubricating coatings

Beyond conventional self-lubricating systems based on externally introduced lubricant phases, a further developmental trend in Ti alloy surface engineering involves shifting toward coatings that intrinsically combine wear resistance with self-lubricating capability. This strategy promotes the design of synergistic composite coatings that combine rigid ceramic frameworks with solid lubricating constituents. Within such architectures, hard ceramic phases provide structural load-bearing support, while high-temperature-resistant lubricants such as MoS₂, WS₂, h-BN, or CaF₂ regulate interfacial shear. The ultimate objective is to construct surfaces capable of resisting abrasive wear while adaptively forming lubricating films during sliding.

In this context, MAX phases represent a unique class of nanolaminated ceramics that intrinsically integrate hardness and lubricity. These materials follow the general formula M_n+1AX_n, where M is an early transition metal, A is a group A element, and X represents carbon or nitrogen [see Figure 10]^[177,178]. Their crystal structure consists of alternating M-X octahedral layers and A-element atomic layers stacked along the c-axis. Strong M-X covalent bonds confer high hardness and load-bearing capability, whereas weaker metallic M-A bonds enable interlayer shear. This layered architecture allows MAX phases to act simultaneously as structural reinforcements and internal lubricating planes, helping adaptive transfer films form during sliding without requiring external lubricants^[177,179]. As a result, MAX phases offer a rare combination of mechanical robustness and frictional adaptability^[180]. Unlike conventional lubricant-containing coatings, whose lubricating function is directly rooted in their crystallographic architecture, these coatings are particularly relevant to a microstructure-centered interpretation of self-lubricating behavior.

Figure 10. Elements of the Periodic Table found in MAX phases and unit cell for n = 1. This figure is adapted from Ref.^[177]. Licensed under CC-BY 4.0. MAX: MAX phases, a family of layered ternary carbides and nitrides generally represented by the formula Mₙ₊₁AXₙ.

Among them, Ti₃SiC₂ and Ti₂AlC are the most extensively investigated for Ti alloy coatings. Ti₃SiC₂ exhibits dual metallic-ceramic characteristics. Its weakly bonded Si layers readily undergo shear and selective delamination, forming graphite-like lubricating films with low shear strength. In laser-cladded coatings, Ti₃SiC₂ participates in load sharing while simultaneously reducing friction through self-adaptive delamination^[181]. It maintains relatively stable tribological performance from RT to 600 °C^[171]. Wang et al.^[171] demonstrated that Ti₃SiC₂-containing Ni60-based coatings exhibited significantly improved tribological behavior across this temperature range, as the lamellar structure reduced shear stress while hard phases inhibited plastic deformation. The key advantage of Ti₃SiC₂, therefore, lies in the coexistence of a load-bearing ceramic framework and a shear-accommodating lamellar substructure within the same phase architecture.

Ti₂AlC also possesses a layered structure but relies on Al-layer plasticity and oxidation for lubrication. During sliding, Al layers migrate or oxidize, promoting the formation of Al-O-rich transfer films with low shear strength^[172,182]. Compared with Ti₃SiC₂, Ti₂AlC has a lower density and superior thermal shock resistance, making it attractive for lightweight high-temperature applications. Hua et al.^[173] showed that laser-cladded TC4-Ti₂AlC composites exhibited improved wear resistance through the formation of continuous Ti₂AlC-derived transfer films at the sliding interface. However, Ti₂AlC generally provides weaker lubrication than Ti₃SiC₂ due to stronger interlayer bonding, and Al loss during cladding may reduce phase stability^[172]. It is therefore often used as an auxiliary lubricating phase in combination with other soft metals or high-temperature lubricants. This indicates that the tribological role of Ti₂AlC is strongly dependent on phase stability during cladding and on the retention of its layered structural characteristics during service.

The in situ formation of MAX phases further enhances their structural integration. Chu et al.^[178] revealed that Ti₂AlC formation occurs through both direct precipitation from Ti-Al-C melts and interfacial transformation of TiC via Al diffusion. Moderate Al additions improved melt fluidity and promoted uniform MAX phase distribution. Similarly, Sun et al.^[182] demonstrated that coatings with optimized MAX phase fractions exhibited balanced phase distribution and enhanced wear resistance at both RT and 400 °C. These findings indicate that an optimal MAX phase fraction is critical because its effectiveness depends on the retention and distribution of the lamellar structure within the coating. At appropriate fractions, MAX phases can simultaneously contribute to load support and shear accommodation through their layered architecture, enabling a balance between strengthening and adaptive lubrication. However, when the phase fraction deviates from this optimal range, either due to insufficient presence or structural degradation, the coordinated load-sharing and shear-relief mechanisms are disrupted, leading to diminished tribological performance.

Overall, MAX-phase-based self-lubricating coatings are typically characterized by a laminated ceramic architecture that intrinsically combines load-bearing capability with interlayer shear accommodation. Their tribological benefit mainly arises from the coexistence of structural reinforcement and adaptive lubrication, enabling transfer-film formation while preserving mechanical support. However, their effectiveness still depends on phase stability during cladding, the retention of lamellar structural features, and the control of MAX phase fraction and distribution within the coating. These characteristics make MAX phases an important bridge between conventional lubricant-containing coatings and more intrinsically coupled multifunctional surface systems.

Conventional wear-resistant coatings predominantly rely on hardness enhancement to mitigate material removal. However, under severe operating conditions involving high contact stress, elevated temperature, or dry sliding, excessive frictional heating and stress concentration can undermine structural stability and accelerate failure. The synergistic coating systems discussed above address these limitations by integrating structural reinforcement with active friction regulation. By integrating load-bearing frameworks and lubricating phases, these coatings simultaneously suppress abrasive and adhesive wear, thereby extending the applicability of laser cladding technology to complex service environments characterized by high sliding speeds and unlubricated contact.

Building upon this concept of multifunctional synergy, HEA coatings represent a further evolution in surface engineering strategies. By leveraging multi-principal-element alloying and tunable phase stability, these systems offer expanded opportunities for in situ phase regulation, structural adaptability, and potentially intrinsic friction modulation. Consequently, they provide a promising pathway toward next-generation multifunctional coatings that can sustain performance under increasingly demanding service conditions.

Phase-structure-dominated HEA coatings

BCC-Laves/B2 synergistically strengthened HEA coatings

Building on BCC solid-solution-dominated systems, the introduction of intermetallic phases to construct BCC-Laves or BCC-B2 synergistic architectures represents an effective strategy to further enhance the load-bearing capacity of HEA coatings on Ti alloys. These coatings typically consist of a BCC matrix combined with a controlled volume fraction of Laves or ordered B2 phases, which commonly precipitate in interdendritic regions or form semi-continuous network morphologies. Such architectures differ from single-phase BCC coatings in that strengthening is no longer derived solely from lattice distortion within the matrix, but from microstructural partitioning between a continuous solid-solution framework and harder intermetallic domains.

A representative BCC-B2 phase-partitioned architecture is shown in Figure 12. In laser-deposited AlCoCrFeNiTi_0.5 HEA coatings on Ti6Al4V, the coating consists of disordered A2 and ordered B2 BCC phases, with Fe/Cr enrichment in the A2 regions and Al/Ni/Ti enrichment in the B2 regions. This phase partitioning demonstrates how elemental segregation during rapid solidification can construct a dual-BCC architecture that differs from single-phase BCC solid-solution coatings.

Figure 12. Representative BCC-B2 phase architecture and interfacial microstructure of a laser-deposited AlCoCrFeNiTi_0.5 HEA coating on Ti6Al4V. (A and B) XRD patterns of the Ti6Al4V substrate and as-deposited HEA coating, respectively; (C) coating/substrate interfacial morphology; (D) EDS line-scan profile across the HEA/Ti6Al4V interface; (E and F) SEM microstructures showing A2/B2 phase partitioning in the HEA coating. This figure is adapted from Ref.^[220]. Licensed under CC-BY 4.0. BCC: Body-centered cubic; HEA: high-entropy alloy; EDS: energy-dispersive X-ray spectroscopy; XRD: X-ray diffraction; SEM: scanning electron microscopy.

For example, Meng et al.^[202] laser-cladded a TiZrNbCrCo coating on Ti6Al4V and observed that while the coating was predominantly BCC, increasing cladding layer number and reducing substrate dilution promoted the formation of Zr/Co-enriched C14-type Laves phases in interdendritic regions, resulting in a characteristic BCC-Laves dual-phase microstructure. This result shows that the phase balance in such coatings is highly sensitive to solidification conditions and dilution history, both of which directly affect interdendritic segregation and intermetallic precipitation.

The improved wear resistance of such dual-phase systems arises from synergistic load sharing between the BCC matrix and intermetallic phases. The BCC matrix provides a continuous load-transfer pathway, whereas high-hardness phases such as Laves or B2 preferentially sustain contact stresses, thereby suppressing plastic plowing and micro-cutting. Consequently, the dominant wear mechanism transitions toward a mixed abrasive-oxidative or an adhesive-oxidative regime. Hao et al.^[203] reported that an AlTiVCrNb coating composed of disordered BCC, Laves phases, and dispersed Ti₂AlNb nanoscale precipitates exhibited increased hardness and approximately 1.58-fold improved wear resistance compared with the substrate. This improvement can be attributed to the formation of a hierarchical load-bearing architecture across multiple length scales. The key feature here is not simply phase coexistence, but the establishment of a hierarchical stress-bearing network in which nanoscale precipitates, interdendritic intermetallics, and the BCC matrix operate at different structural levels during sliding.

From a performance standpoint, BCC-Laves/B2 systems often outperform single-phase BCC coatings in hardness and, under properly controlled phase fractions and morphologies, can also exhibit superior wear resistance. Meng et al.^[202] reported an average hardness of approximately 748 HV for TiZrNbCrCo multiphase coatings, with a wear-volume reduction of 68.6% relative to the substrate. In addition, Zhang et al.^[221] prepared FeCrAlMoSi_x coatings containing B2 and μ phases, achieving hardness values up to approximately 805 HV at x = 0.10. This improvement is closely related to the introduction of ordered or topologically close-packed strengthening units that reduce local deformation compatibility but markedly increase resistance to penetration and plowing.

The advantages of BCC-intermetallic systems become more pronounced at elevated temperatures, especially when intermetallic-assisted oxidation promotes the formation of dense and adherent protective films that reduce interface damage and suppress oxidation-assisted cracking. The incorporation of Cr-rich or Al-rich intermetallic phases facilitates the formation of dense, adherent oxide films during sliding or thermal exposure, thereby reducing oxidation wear and enhancing service stability. Deng et al.^[198] demonstrated that a CoCrFeNiMo_0.2 coating composed of BCC and σ phases exhibited hardness of approximately 900 HV and superior wear resistance at 600 °C due to the formation of a protective Cr-rich oxide layer. In addition, Yang et al.^[208] reported that increasing Ni content in (AlCrFeTiV)_100-xNi_x coatings transformed the phase constitution from BCC/Laves to BCC+NiAl-type B2+Laves, resulting in improved oxidation resistance at 800 °C due to the formation of a compact spinel oxide film. In other words, intermetallic phases influence high-temperature tribological behavior not only through direct hardening, but also through their effect on oxidation pathways and the chemistry of protective surface scales.

Despite these benefits, excessive volume fractions or continuous networks of intermetallic phases may introduce brittleness and increase susceptibility to crack initiation, particularly under thermal cycling or impact loading. For instance, an overly high μ-phase content in FeCrAlMoSi_x coatings was shown to promote stress concentration and microcrack formation^[221]. Once the intermetallic constituent evolves from a dispersed strengthening unit into a continuous brittle path, the same microstructural feature that improves hardness begins to accelerate crack propagation and structural instability. As a result, while BCC-Laves/B2 synergistic strengthening enhances load-bearing capability and wear resistance, careful control of phase fraction, morphology, and distribution is essential to balance mechanical strength with structural reliability. The design problem in this class of coatings is therefore not how to maximize intermetallic content, but how to confine hard phases to morphologies and length scales that strengthen the matrix without converting the coating into a brittle interconnected framework. This inherent trade-off has motivated further exploration of coatings with enhanced thermal stability and multifunctional composite architectures capable of sustaining performance under complex and extreme service conditions.

Tribological performance of RHEA coatings

With the expanding application of Ti alloys in aerospace, energy systems, and high-temperature tribological components, surface-modification requirements have progressively shifted from room- and intermediate-temperature wear resistance toward service conditions involving elevated temperatures, high contact stresses, and coupled oxidation-wear environments. Under such conditions, conventional HEA coatings dominated by FCC or BCC phases may experience microstructural softening, phase instability, or accelerated oxidation. In response, RHEA coatings, mainly composed of refractory elements with high melting points, have emerged as a promising solution. The central design philosophy of RHEAs is to exploit strong atomic bonding, sluggish diffusion, and exceptional thermal stability to maintain wear resistance at elevated temperatures. What sets this class apart is that coating design is no longer centered mainly on ambient-strength enhancement, but on preserving phase constitution and near-surface microstructural stability during prolonged thermal-mechanical exposure.

From a phase-structure standpoint, RHEA coatings are typically characterized by a BCC solid-solution matrix, occasionally accompanied by minor secondary phases or localized segregation. Although this configuration resembles conventional BCC-type HEAs, the design emphasis is fundamentally different. Whereas BCC-type HEAs are often optimized for room- or intermediate-temperature load-bearing enhancement, RHEAs prioritize phase stability and mechanical property retention under high-temperature exposure. The distinction therefore lies less in the nominal crystal structure itself than in the extent to which that structure can resist thermally induced softening, diffusion-assisted coarsening, and oxidation-related degradation. Representative microstructures and tribological properties of laser-cladded WMoTaNbV RHEA coatings on Ti6Al4V are shown in Figure 13, illustrating the relationship between BCC-based refractory microstructures, coating hardness, friction stability, and wear resistance.

Figure 13. Cross-sectional microstructure and tribological properties of laser-cladded WMoTaNbV refractory high-entropy alloy coatings on TC4 substrate. (A-D) Cross-sectional SEM images of the coating fabricated at 3.7 kW, showing (A) the overall coating morphology and the microstructures of the (B) top, (C) middle, and (D) bottom regions; (E-H) Tribological results of the TC4 substrate and coatings fabricated at different laser powers, including (E) microhardness profiles, (F) friction coefficient curves, (G) wear-volume loss, and (H) wear rate and wear resistance. This figure is adapted from Ref.^[204]. Licensed under CC-BY 4.0. BCC: Body-centered cubic; SEM: scanning electron microscopy.

BCC structures enriched with refractory elements such as Nb, Mo, Ta, W, and V possess high melting points and elastic moduli, rendering them less susceptible to softening or microstructural degradation during high-temperature sliding. This intrinsic stability provides a robust load-bearing framework for elevated-temperature wear resistance. For example, Hu et al.^[225] reported that a lightweight AlNbTiVZr RHEA coating fabricated on Ti6Al4V exhibited a predominantly BCC microstructure with a minor fraction of C14-type Laves phase, achieving enhanced hardness and wear resistance while maintaining reduced density. Here, the presence of only a limited amount of secondary phase is important, because it preserves the thermal stability of the BCC framework without introducing excessive brittle connectivity. In the same way, Du et al.^[211] demonstrated that in TiZrHfNbTaMo_x RHEA coatings, Mo addition stabilized the BCC structure and improved hardness retention at elevated temperatures, establishing a structural basis for reliable service in complex environments.

In terms of tribological behavior, the advantages of RHEA coatings become particularly evident at medium to high temperatures. During elevated temperature sliding or fretting, these coatings generally exhibit lower wear rates and more stable COF. Their performance derives not only from high hardness but also from the suppression of microstructural degradation. Strong atomic bonding and relatively low diffusion rates among refractory elements effectively inhibit dislocation recovery and grain coarsening, thereby delaying near-surface softening and preventing the wear mechanism from transitioning to severe adhesive or spallation modes. Liu et al.^[200] reported that at 600 °C, the wear rate of Al₂₀V₂₀Cr₂₀Nb_40-xMo_x RHEA coatings decreased by nearly two orders of magnitude, primarily due to the rapid formation of a dense oxide film and the mechanical anchoring effect of Mo-rich oxides at the friction interface. In this temperature regime, wear resistance is governed by the combined persistence of the refractory matrix and the development of a compact oxide-bearing surface layer, rather than by matrix hardness alone; in particular, the mechanical integrity, adhesion, and continuity of the oxide layer are critical in determining whether oxidation becomes protective or damage-accelerating.

Because high-temperature wear is often strongly coupled with oxidation, the compositional design of RHEAs frequently incorporates elements such as Cr and Al that form dense, adherent oxide scales. These oxide layers serve dual functions: they retard further matrix oxidation and participate in load sharing and friction modulation at the sliding interface. However, performance does not improve monotonically with increasing alloying content or temperature. Zhao et al.^[212] observed that in AlNbTaZr_x RHEA coatings, increasing Zr content induced a transition from a single BCC phase to a BCC + hexagonal close-packed (HCP) dual-phase structure, accompanied by a shift in wear mechanism from micro-cutting to oxidation-assisted brittle spallation. Similarly, Liu et al.^[226] reported that TiZrNbAl_x lightweight RHEA coatings exhibited inferior wear resistance at 500 °C compared with RT, which was associated with enhanced thermal vibration and strain accumulation at elevated temperatures. Once phase stability is disrupted or the oxide layer loses mechanical integrity, the same high-temperature environment that favors protective oxidation can instead accelerate embrittlement and surface damage.

Although RHEA coatings exhibit pronounced advantages in high-temperature wear resistance and oxidation resistance, their design is often associated with relatively high density, increased material cost, and potentially insufficient room-temperature plasticity. These constraints mean that RHEAs are better viewed as environment-specific coating solutions than as universally optimal systems for Ti-alloy surfaces. Consequently, RHEAs are better suited to specific service conditions that demand exceptionally high-temperature stability and long-term durability, rather than serving as general-purpose surface modification solutions for Ti alloys.

Overall, RHEA coatings represent a significant extension of HEA design for Ti alloy surface engineering, marking a shift from phase-centered strengthening strategies toward service-environment-driven optimization. Together with BCC-type, BCC-intermetallic, FCC-BCC dual-phase, and in situ hard-phase-reinforced systems, RHEAs complete a hierarchical material design framework capable of addressing diverse combinations of load, temperature, and environmental conditions. Within this framework, their defining contribution lies in coupling refractory phase stability with oxidation-assisted interfacial protection under severe thermal tribological conditions.

HEA-based composite coatings with ceramic reinforcements, rare-earth oxides, and lubricating phases

While monolithic HEA coatings provide enhanced structural stability and load-bearing capacity, their performance under extreme wear, corrosion, or high-temperature environments may still be limited by insufficient hardness, interfacial instability, or frictional heat accumulation. To overcome these limitations, recent research has increasingly focused on HEA-based composite coatings that integrate ceramic reinforcements, rare-earth oxides, and lubricating phases within a high-entropy matrix, thereby constructing multifunctional microstructural architectures with coupled strengthening, stabilization, and friction-regulation mechanisms.