Sliding Wear of Conventional and Suspension Sprayed Nanocomposite WC-Co Coatings: An Invited Review

The global thermal spray coatings market was valued at USD 10.1 billion in 2019 and is expected to grow at a compound annual growth rate of 3.9% from 2020 to 2027. Carbide coatings form an essential segment of this market and provide cost-effective and environmental friendly tribological solutions for applications in aerospace, industrial gas turbine, automotive, printing, oil and gas, steel, and pulp and paper industries. Almost 23% of the world’s total energy consumption originates from tribological contacts. Thermal spray WC-Co coatings provide excellent wear resistance for industrial applications in sliding and rolling contacts. Some of these applications in abrasive, sliding and erosive conditions include sink rolls in zinc pots, conveyor screws, pump housings, impeller shafts, aircraft flap tracks, cam followers and expansion joints. These coatings are considered as a replacement of the hazardous chrome plating for tribological applications. The microstructure of thermal spray coatings is however complex, and the wear mechanisms and wear rates vary significantly when compared to cemented WC-Co carbides or vapour deposition WC coatings. This paper provides an expert review of the tribological considerations that dictate the sliding wear performance of thermal spray WC-Co coatings. Structure–property relationships and failure modes are discussed to grasp the design aspects of WC-Co coatings for tribological applications. Recent developments of suspension sprayed nanocomposite coatings are compared with conventional coatings in terms of performance and failure mechanisms. The dependency of coating microstructure, binder material, carbide size, fracture toughness, post-treatment and hardness on sliding wear performance and test methodology is discussed. Semiempirical mathematical models of wear rate related to the influence of tribological test conditions and coating characteristics are analysed for sliding contacts. Finally, advances for numerical modelling of sliding wear rate are discussed.

Almost 23% of the world's total energy consumption originates from tribological contacts ( Ref 1). The global thermal spray coatings market size was valued at USD 10.1 billion in 2019 and is expected to grow at a compound annual growth rate (CAGR) of 3.9% from 2020 to 2027 (Ref 2). Carbide coatings form an essential segment of this market and provide a cost-effective and environmentally friendly tribological solution to many applications ( Ref 2). The microstructure (Ref [3][4][5] of thermal spray coatings is, however, complex and the wear mechanisms and wear rates vary significantly when compared to cemented WC-Co carbides (Ref [6][7][8][9][10] or vapour deposition WC coatings ( Ref 11). Thermal spray coatings improve the wear and frictional performance of components (Ref 3). These coatings are deposited through a line of sight process where the coating material is heated and propelled on to the substrate to form a lamella coating (Ref 4). Thermal spray methods allow for the deposition of hard metals, or their composites, on large surfaces where parts cannot be produced using other techniques due to technical or economic reasons ( Ref 3,4). These coatings are used to provide a combination of coating and substrate properties to improve the service life, restore dimensions of machined parts and repair expensive worn components ( Ref 12,13). Thermal spray carbide coatings are used for a range of tribological applications in abrasive, sliding, corrosive and erosive conditions for the aforementioned industrial sectors. In particular, WC-Co coatings are used for applications such as steel rolls, sink rolls in zinc pots, corrugating rollers, agricultural rasp bars, conveyor screws, pump housings, impeller shafts, conveyor screws, compressor stators, impeller shafts, aircraft flap tracks, cam followers, and expansion joints ( Ref 12). Thermal spray coatings can also replace electrolytic hard chrome (EHC) plating, which is considered environmentally hazardous, for industrial applications ( Ref 13).
Thermal spray coatings are deposited through various generations of high velocity oxygen fuel (HVOF), vacuum plasma spray (VPS), air plasma spray (APS), warm spray (WS), cold spray (CS), and high velocity air fuel (HVAF) systems (Ref 4). Although conventional HVOF and APS systems have been used to deposit nanocomposite WC-Co coatings, recent progress has introduced suspension sprayed HVOF (S-HVOF) coating systems, which can deposit nanocomposite coatings using liquid suspensions (Ref [14][15][16]. Coating particle temperature and velocity in thermal spray systems dictate the coating microstructure and mechanical strength. Figure 1 shows the range of temperature and velocity of spray particles for thermal and cold spray systems (Ref 4). Thermal sprayed coatings are formed by the continuous layering of splats, and the process creates an anisotropic coating microstructure, where the mechanical properties such as fracture toughness show relatively low values parallel to the coating surface (Ref 4,17,18).
Ceramic carbides are often used to form wear-resistant coatings for tribological applications; whereas other ceramic oxides and nitrides are used for applications such as thermal or electrical resistance ( Ref 3,12,13). Table 1 summarises commonly used carbides such as WC, Cr 3 C 2 and TiC for tribological applications ( Ref 10). The selection of these carbides is based on their solubility and wettability in metal binders during thermal spraying. WC generally exhibits good wettability with metallic liquids, e.g., Ni and Fe, but its wettability with Co in particular is excellent (Ref 10). Similarly, the dissolution of carbides harden the binder ( Ref 3,10). WC-Co coatings are generally used for wear resistance applications up to an operating temperature of 500°C, whereas Cr 3 C 2 and WC-Co-Cr is preferred for corrosion resistance applications ( Ref 10,[19][20][21]. WC-Co has long been used as a cemented carbide for producing wear-resistant parts, which has then been brought forward to thermal sprayed coatings ( Ref 3,10). Figure 2 illustrates applications for WC-Co cemented carbides along with their design features such as the WC carbide size, binder content and hardness ( Ref 10). Thermal spray WC-Co coatings share applications in areas such as those for wear resistance ( Ref 3,4,12,13). Figure 3 shows a comparison of the mechanical properties of cemented WC-Co with other engineering materials that are either used as a (i) wear-resistant material, e.g., SiC, alumina and Stellites TM (typically of composition; Balance-Co, 28-32%Cr, 4-17.5%W, 1-3.3%C, in weight percentages), (ii) a counter body or (iii) substrate material for coatings in sliding contacts (Ref 10). The properties presented in Fig. 3 show the high bound and low bound values of hardness, toughness, elastic modulus and compressive strength achievable for these materials. The properties of thermal spray coatings are generally towards the lower values due to their complex microstructure as summarised in Table 2 (Ref . For example, hardness values of thermal spray WC-Co coatings range from * 795 to 1568 HV in comparison with * 750 to 2250 HV for cemented WC-Co carbides (Fig. 3, Table 2). Similarly, the fracture toughness of thermal sprayed WC-Co varies between * 1.2 and 25 MPa m 1/2 in comparison with * 8 and 30 MPa m 1/2 for cemented WC-Co carbides (Fig. 3, Table 2). The elastic modulus of as-sprayed thermal spray WC-Co deposits is around one-third of the bulk material, typically around * 250 ± 50 GPa ( Ref 14,15) Fig. 1 Spectrum of thermal spray and cold spray processes across particle velocity and flame temperature attainable during coating deposition (Ref 4) in comparison with 400-650 GPa for cemented carbides ( Ref 10).

WC-Co Materials and Deposits
Cobalt (* 4-17 wt.%) is generally used with WC as a metal binder due to its high fracture toughness, good wettability and dry lubricating effects ( Ref 3,9,10). The cobalt also provides a good combination of hardness and toughness for thermal spray coatings ( Ref 37,[41][42][43]45,46). Other variations in binder composition with Ni and Cr allow for improved performance under high temperature and corrosive environments, respectively ( Ref 47).
The sliding wear performance (Table 2) of WC-Co coatings depends on the coating process and the process parameters that dictate the microstructure and the tribological test conditions leading to differences in the wear mechanisms. Figure 4 summarises the factors that influence the sliding wear performance of these coatings in terms of the structure-property relationship and the underlying wear mechanism. The upper half of the figure shows the parameters influencing the coating microstructure and its mechanical properties, whereas the lower half summarises the tribological parameters influencing coating performance. The outer ring in this figure depicts design factors that need to be selected for a given application. The middle ring summarises the interdependency of these design factors on coating performance and highlights a complex behaviour of microstructure and tribological conditions. The central ring shows the relationship between the structure-property and the sliding wear process.

Organisation of This Review
Corrosive wear and tribo-corrosion conditions of thermal sprayed carbide coatings have been reviewed (Ref 8,37,45) as has the microstructural evolution of cemented carbides (Ref 10). This contribution reviews the structureproperty relationships and wear mechanisms in sliding wear of conventional and nanocomposite suspension sprayed WC-Co coatings and suggests future trends to further improve their sliding wear performance.
This review is arranged in the eight sections. ''Introducion'' section provides an introduction of the coating processes, whereas 00 Microstructure of WC-Co Coatings 00 section concentrates on the evolution of the WC-Co coating microstructure. 00 Sliding Wear Mechanisms 00 section relates the complexity of this microstructure to the wear mechanisms and wear rates. Comparison of failure modes of WC-Co coatings with cemented WC-Co carbides is also discussed in 00 Sliding Wear Mechanisms 00 section. The dependency of coating microstructure and mechanical properties on the prediction of wear rates is discussed in 00 Structure-Property Relationship in Sliding Wear'' section. The influence of post-treatment on the structureproperty relationship is discussed in this Table 1. Numerical modelling of wear rates is discussed in ''Computational Wear Modelling''. Conclusions and future recommendations are included in ''Conclusion'' and ''Microstructural Quality and Sliding'' sections, respectively. In this paper, the term wear is specifically applied to sliding wear unless otherwise stated.

Microstructure of WC-Co Coatings Distinguishing Features of Thermal Sprayed Microstructures
The microstructure of WC-Co coatings is significantly different from cemented carbides ( Ref 3,9,10,16). Thermal sprayed WC-Co coatings reveal a lamella structure where the bonding mechanism is predominantly mechanical interlocks. Despite some claims of metallurgical bonding in thermal spraying process, for example, using patented technology (Ref 48) and/or using a spray and fuse technique (Ref 49), there is no evidence of metallurgical bonding for WC-Co coatings in the published literature. For the case of cold-or warm-sprayed materials in general, although there is a growing discussion of metallurgical bonding in the deposits due to adiabatic shear instabilities, which causes metallurgical bonding due to explosive cladding (Ref 50), no such metallurgical bonding in cold sprayed WC-Co has been reported ( Ref 51). Thermal spray coating microstructure contains pores and micro-cracks and exhibit anisotropy due to layering of splats during coating deposition ( 17,26,45). Due to this complex microstructure, the mechanical properties such as hardness, modulus and fracture toughness are relatively lower than the cemented carbides (Table 2, Fig. 3). Through-thickness residual stress is caused by differences in the coefficients of thermal expansion of the coating substrate system as well as quenching of lamella particles also influence coating properties (Ref [52][53][54][55]. The residual stress profile is superimposed on the mechanical and/or thermal service loading of the coated component and need to be considered when evaluating coating performance ( Ref 50,[54][55][56][57][58][59]. Despite relatively lower mechanical strength than the cemented carbides (Tables 1, 2 and Fig. 3), these coatings are frequently used in tribological applications because of their ability to provide a cost-effective and environmentalfriendly solution to the design of components requiring surface engineering. Carbide size, carbon loss leading to complex microstructural phases, and microstructural features such as porosity and carbide distribution, therefore, impact the tribo-mechanical performance ( Table 2) (Ref [60][61][62][63]. Figure 5 shows the typical x-ray diffraction (XRD) pattern of WC-Co coatings deposited by conventional HVOF (JP5000 and Diamond Jet (DJ)), S-HVOF, VPS and APS coatings ( Ref 15,16,21,22). This figure also shows the typical XRD pattern for cemented WC-Co carbides (Ref 8). Figure 6 shows the SEM and TEM images of the deposited coatings (Ref [14][15][16]21). This figure shows that in conventional HVOF spraying with liquid fuel HVOF systems such as JP5000, most of the WC phase is retained as blocky carbides from the starting powder ( Fig. 6c and d). A comparison of the typical microstructure for cemented WC-Co carbides is shown in Figure 6(i) and (j) ( Ref 8,10).

Impact of Decarburisation of WC During Thermal Spray Processing
As shown in Fig. 1, particle velocity and temperature in thermal processes vary for different spray systems (Ref 4). The oxygen atmosphere in HVOF systems can, however, promote oxidation reactions ( Ref 64). HVOF systems are, therefore, preferred over the air plasma spray systems (APS) due to their ability to deposit coating material at a relatively lower temperature and higher velocity (Ref 4,60,64). Nevertheless, decarburisation of WC leading to W 2 C, W and eta-phases ('g-phases') such as Co 3 W 3 C, Co 2 W 4 C (M 6 C) and Co 6 W 6 C (M 12 C) can occur, depending upon the coating process and the process parameters (Ref The decarburised W 2 C phase formed from reaction shown in Eq 1 is precipitated along the WC grain boundary, together with a ring of metallic W between the WC grain and W 2 C (Ref 19) formed by reactions shown in Eq 2 and 3. These reactions take place mainly in WC grains that interact with oxygen (Eq 1,2, and 3). In addition, WC grains can be degraded in an oxygen-free atmosphere (i.e., those that are in the core of the particles), according to Eq 4 (Ref 33  shows that in conventional HVOF spraying there is limited decarburisation of WC leading to the peak broadening in the 2-theta range of 35°-48°containing predominantly an amorphous/nanocrystalline phase and the formation of W 2 C and Co 3 W 3 C ( Ref 14). The extent of carbon loss depends on the spray system and conditions, for instance, when conventional coatings are sprayed with excess oxygen in the case of detonation gun spraying (DGS) (Ref 63) or coatings prepared from strongly carbon-deficient powders ( Ref 66). The increase in carbon loss also leads to peak broadening in the 2-theta range of 35°-48°( Ref 14,15). In the case of the APS process, the higher particle temperature and lower velocity leads to higher dwell time that results in (i) higher decarburisation due to thermal loading of powder particles and (ii) higher porosity due to lower velocity. These effects are observed in Figure 5(d), where both conventional and nanostructured powders show decarburisation when sprayed with hydrogen or helium ( Ref 21). APS processing (Ref 21) with hydrogen leads to relatively higher decarburisation for both conventional and nanostructured coatings, which is caused by a higher enthalpy. APS coatings also show high porosity of up to * 10% ( Ref 21,60).
The lower velocity in some HVOF systems can also lead to porosity, as shown in Figure 6(d), due to either an earlier generation HVOF system or HVOF spraying under nonoptimised spray conditions ( Ref 14,60). Higher velocity in the later generation HVOF spraying not only reduces the  dwell time but also increases the beneficial peening effect that leads to compressive residual stress in the deposit and increases coating density ( Ref 52,53,56,57). VPS coatings should show a lower tendency of decarburisation due to the absence of oxygen in the system. However, as shown in Figure 5c, when nanostructured powders are VPS processed there is formation of eta-phases (M 2 C, M 6 C) associated with the decarburisation of the WC particles (Ref 22). This suggests that small amounts of oxygen during VPS spraying, in conjunction with the high temperature of the plasma system, can also lead to decarburisation (Eq 1 to 4). Zhu et al. (Ref 22) concluded that the extent of decarburisation was relatively low when compared to studies reported by Stewart et al. that used a HVOF top gun system (Ref 62). However, even with the VPS systems, the decarburisation of HVOF nanostructured coatings cannot be avoided. This stipulates that the higher particle temperature has a more dominant effect on carbon loss in WC-Co coatings instead of the available oxygen content around the powder particles. Some studies (Ref 33,65,(67)(68)(69) have reported that the carbon loss leads to the formation of CO 2 during spraying; however, these reactions require careful interpretation in oxygen-deficient environments such as for the VPS process. The reactions in the W-C-Co phase system have been studied intensely by Matthews et al. (Ref 65).
The production of nanostructured layers is also possible with nanostructured powder feedstock ( Ref 32,[34][35][36]70). Suspension HVOF (S-HVOF) systems deposit nanostructured coatings using nanopowder suspensions ( Ref 14,16,20,71,72). Nanoparticles cannot be injected directly in HVOF systems; instead, they are suspended in a liquid suspension such as water or ethanol and a dispersant such as polyethyleneimine is added for HVOF spraying (Ref 14-16, 71, 73). A stable suspension of nanoparticles is needed to deposit reproducible coatings (Ref 14). Due to the high surface area-to-volume ratio of the nanosized particles, thermal loading of WC-Co particles in S-HVOF spraying can lead to carbon loss. WC-Co suspensions pose additional challenges; for example, the surface of WC is acidic and that of Co (CoO) is basic in nature ( Ref 14,73,74), which influences both the stability and pH value of the suspension due to Co dissolution. This difference in the acidic and alkaline nature can, however, be avoided by employing a preformed WC-Co composite powder ( Ref 14,16). Berghaus et al. (Ref 73) studied the development of WC-12Co nanocomposite coatings by suspension spraying using a commercial APS (AXIAL III, Northwest Mettech Corp., Canada) process with an internal injection module that allowed in-flight particle velocities up to 800 m/s at particle temperatures up to * 2400°C. Coatings were prepared mostly from an agglomerated powder (60-250 nm nominal carbide grain size and particle sizes\ 20 lm) in a 20 wt. % ethanol suspension. Coatings with low porosity were produced with a maximum hardness of about 700 HV0.3. It was concluded that the relatively lower coating hardness was a function of both the coating porosity and carbide degradation caused by carbon loss (Ref 73). These coatings showed a pronounced amorphous ''hump'' in the XRD pattern due to decarburisation (Ref 73).
In summary, thermal loading of WC particles tends to increase carbon loss leading to M 2 C, M 6 C, M 12 C phases along with metallic W ( Ref 14,17,64). Thermal loading of particles is influenced by the temperature and velocity of the particles that depend on the coating process and the process parameters (Fig. 1), along with the surface area to volume ratio of carbides. In the case of suspension spraying, the carbon loss can be reduced by (i) changing the coating process from APS to HVOF, (ii) optimising the suspension parameters by using water instead of ethanol as suspension media to decrease the process temperatures during spraying, (iii) optimising the concentration of powder particles in the suspension, and (iv) optimising the spray parameters ( Ref 14,16,73). Heat treatment of WC-Co coatings can lead to the transformation of the amorphous phases and also improve the mechanical strength as seen later in 00 Structure-Property Relationship in Sliding Wear'' section ( Ref 15).

Eta-Phase Generation
Studies by ) employed a HVOF system to reduce the thermal loading on particles and increase particle velocity to deposit dense S-HVOF coatings ( Fig. 6a and b). The lamella structure in the S-HVOF coatings is finer, Fig. 6. TEM images in Figure 6 (e) and (f) show a uniform distribution of nanosized particles within the S-HVOF microstructure. Figure 5(a) shows that decarburisation still occurs in the S-HVOF spraying as indicated by the hump in the 2-theta range of 35°-48°( Ref 16). The amorphous and the nanocrystalline phases can be crystallised through heat treatment at 920°C into M 6 C and M 12 C (Ref 15). Figure 7 shows the phase composition and amounts of equilibrium phases along the path of carbon loss in the phase diagram (Ref 15). This figure stipulates that as the carbon loss increases during S-HVOF spraying, there is an increase in eta-phases from M 6 C to M 12 C after heat treatment. M 12 C appears twice in this equilibrium diagram, once with the relatively lower carbon loss and then with higher carbon loss. Metallic W also appears as a stable phase at higher carbon loss as an equilibrium phase ( Ref 15). W 2 C is not a stable microstructural phase at this temperature and hence also does not appear in the XRD pattern (Fig. 5a, 7). However, there is crystallisation of amorphous/nanostructured phases and better inter-splat bonding as indicated by the indentation modulus analysis after the heat treatment (Ref 15).
Eta-phases ('g-phases') can also occur in cemented carbides, and steps are taken to prevent such phase transformation (Ref 10). Garcia et al. ( Ref 10) described that gettering of oxygen by inhibitors prevent decarburisation of WC and the formation of pure W during the consolidation and densification stages of sintering WC-Co-Cr/V cemented carbides. The inhibitors are added as carbides and may thus act as a carbon source. Eta-phases, however, are sometimes also intentionally integrated within the cemented carbide microstructure. For example, decarburized cemented carbides containing eta-phase in the microstructure are used to moderate property gradients in cemented carbides (Ref 10).
It is informative to contrast the microstructures of cemented WC-Co carbides with those formed by thermal spray methods. The eta-phases of cemented WC-Co carbides are controlled during the manufacturing process to provide WC and Co as the only crystalline phases ( Fig. 5e and 6i, j) ( Ref 8,10). These cemented carbides are produced using degassing and HIPing (hot isostatic pressing) cycles around 1400-1500°C. Figure 8(a) reveals the microstructural evolution during HIPing of cemented WC-Co carbides (Ref 10). Grain growth and densification occur between 800 and 1190°C. This provides critical understanding concerning the microstructural evolution during the HIPing post-treatment of thermal spray coatings; for example (Ref 15) at a HIPing temperature of 920°C.
The carbon loss that can occur during thermal spraying restricts grain growth due to existing eta-phases as indicated in Fig. 7 (Ref 15). Additional factors dictate the microstructure for heat-treated thermal spray WC-Co coatings as shown in Figure 8(b) ( Ref 75). In this instance, Co diffuses across the coating substrate interface from the coating, whereas Fe diffuses across this interface from the steel substrate ( Ref 75). The mechanism of diffusion is illustrated in Fig. 8(c) and discussed by Stoica (Ref 75). A similar diffusion mechanism of elemental species during HIP treatment of WC-NiCrBSi coatings is discussed by Stewart (Ref 76). These microstructural changes also influence the hardness, modulus and residual stress profile in heat-treated WC-Co coatings ( Ref 75,76).
During sintering of cemented carbides at temperatures higher than 1190°C, viscous flow densification, liquid phase sintering and particle rearrangement occurs up to 1400°C (Fig. 8a) (Ref 10). Eta-phases can also form during the manufacturing of cemented WC-Co carbides and result in reduced toughness (Ref 10). The control of eta-phases is not a function of temperature in cemented carbides because the temperature is tightly controlled between 1400 and 1500°C. This is unlike thermal spray coatings where the thermal loading of particles (Fig. 1

Mitigating Decarburisation During Thermal Spraying
There are various methods for limiting the decarburisation of WC-Co coatings to improve wear performance. One method is to control the spray parameters, but this requires measurements of particle velocity and temperature and design of experiments that account for carbide sizes, spray particle morphology, fuel/plasma gas chemistry, environmental factors, etc. [see (Ref 43,44)]. Developments in this area include automated feedback control systems to alter coating process parameters and can further improve the reproducibility of coated components. In suspension spraying, liquid suspensions can reduce the thermal loading of WC particles ( Ref 14,16). Another method of limiting decarburisation is to engineer the powder particles to reduce the thermal loading of carbides through cladding; for example by having the carbides concentrated in the core and the binder/protective coating on the outer shell (Ref 77,78).
One method used by Jafari et al. milled powder of conventional carbide grain size (CGS) to nano-dimensions, which was then electroplated with nickel ( Ref 38). The powder was then spray-dried and agglomerated ( Ref 38). Three coatings were prepared with the conventional, nanostructured and engineering feedstock for wear testing. The same spray parameters were maintained for all the feedstocks, whereas it is common practise that each  Table 3 feedstock would need to be optimised. Nevertheless, the microstructural observations indicated that the conventional coating exhibited some carbide dissolution, whereas the uncoated nanostructured coating revealed extensive decarburisation that led to dissolution into the binder (Ref 38). On the other hand, the nickel-coated nanostructured coating showed no evidence of dissolution into the binder. The best wear performance was achieved for the nickelcoated nanostructured coating, which was followed by the conventional and then the nanostructured coating (Ref 38). The wear mechanisms of the conventional coating incorporated the formation of tribo-films, subsurface cracking of splats and then delamination. The wear of nickel-coated nanostructured coating was mainly comprised of binder extrusion followed by carbide pull out.
Engineered particles also improve wear as a by-product of enhanced mechanical properties at inter-splat boundaries. Yuan   and cracks. These microstructural features preferentially originate and propagate along the interfaces between brittle and softer zones in the coating structure. The WC particles reinforced the splat boundaries; therefore, mitigating crack propagation under splats and reducing splat delamination. This reinforcement also increased the sliding wear resistance of the coating ( Ref 79).
The cladding of particles to mitigate carbon loss and improve coating strength has improved wear performance when compared to conventionally produced particles (Ref 79). However, the additional particle technology and cost of modifying the particles has restricted widespread use. Several patents have considered the cladding of coating particles to improve the coating microstructure for several applications ( Ref 77,78). These patents exploit techniques such as plating, physical vapour deposition (PVD) and chemical vapour deposition (CVD) to protect the coating particle during thermal spraying. In the case of CVD and PVD coatings, several cladding materials are possible; for example, coating of hydroxyapatite with titanium ( Ref 78). Cladding with alloys such as Ti6Al4V using these processes is difficult since the deposition of elemental species cannot be controlled precisely to form the specific alloy through CVD and PVD processes.
Another approach to mitigate carbon loss in WC-Co coatings was proposed by Matthews et al. (Ref 65). A combination of powder design, spray parameters and subsequent heat treatment of coatings was employed to minimise the carbon loss in WC-Co coatings. This approach is based on the tendency for carbides to dissolve in-flight and for precipitation of submicron carbide particles on subsequent heat treatment (Ref 65). Similar concepts have been applied to Cr 3 C 2 -NiCr coatings to take advantage of the tendency for Cr 3 C 2 carbides to dissolve into the molten NiCr binder and maximise carbide dissolution (Ref 80). Post-coating heat treatment is then employed to precipitate the dissolved elements out of solid solution to re-form the initial carbide particles. The time and temperature is used to control the extent of grain growth to form unique microstructures that range from high concentrations of nanostructured particles, through to complex interconnected carbide networks (Ref 65).

Sliding Wear Mechanisms
The complex interdependencies of wear mechanisms must be understood to predict the sliding wear performance of coatings (Fig. 4). The cost of coating development can be significantly reduced using standard tests that evaluate the failure modes and sliding wear rate of coatings. There are three principal tests that assess sliding wear, i.e., ASTM G133 . Wear volume is measured in all of these tests, either by the mass loss or surface topography that is used to calculate the wear rate, i.e., wear rate = wear volume / (normal load 9 total sliding distance) typically in the units of mm 3 /N m (Ref [81][82][83]. In situ and post-test evaluations help identify wear mechanisms. Table 2 summaries the sliding wear test evaluations for the coating and counter body properties along with the sliding wear test method and wear rates. Some of the commonly used counter body materials in these investigations include stainless steel, WC-Co, Si 3 N 4 and Al 2 O 3 ( Table 2).
Several studies have proposed wear mechanisms in sliding wear of coated surfaces ( Ref 16,37,[39][40][41]79,84). The wear mechanism of WC-Co coatings depends not only on the coating microstructure but also on the properties of the counter body and the tribological test conditions of load and velocity during sliding wear (Fig. 4). In terms of coating microstructure, the size and distribution of carbides, degree of carburisation leading to eta-phases and mechanical strength of the coating material are the most important factors affecting sliding wear ( Ref 16,37,[39][40][41]79,84). The counter body can also significantly alter the wear mechanism based on its tribo-mechanical and tribo-chemical properties (Ref 14-16, 21, 41). Similarly, the changes in the contact pressure, contact geometry and wear debris in the contact region during sliding wear alter the wear mechanism ( Table 2). The wear mechanisms in the sliding wear of WC-Co coatings can be broadly classified as (i) abrasive wear or macro-mechanical wear, (ii) fracture and fatigue and (iii) tribo-chemical. Figure 10 summarises some of the main factors contributing to the sliding wear of WC-Co coatings ( Ref 16,37,79,85). SEM observations indicating the underpinning wear mechanism are also shown in Fig. 10.

Abrasive or Macro-mechanical Wear
Abrasive wear during sliding contact in WC-Co coatings starts as a two-body wear mechanism due to asperity contact ( Ref 35,37,62,75,86,87). The contact pressure is highest at this stage of the sliding wear process (Ref 87). For a given coating material, this process depends on the hardness, fracture toughness and tribological conditions of the test couples (Table 2). In the case of a ball-on-flat test  pressure is shared between the carbides and the matrix. The surface asperities interact within the contact region. Thus, there can be preferential wear of the matrix (Fig. 10f) that can expose carbides ( Ref 16,40). The extent of this matrix erosion depends on the surface roughness, hardness and toughness of the counter body.
For a given surface roughness, a harder counter body can increase matrix erosion because the harder surface asperity of the counter body and/or entrapped wear debris within the contact region can plough through the softer matrix (Ref 16,40,79). The carbides dislodge due to matrix erosion as the sliding wear process continues, hence leading to three-body abrasion of the test couple ( Ref 14,37,40,79). This is schematically represented in Figure 11(a), where matrix erosion leads to the removal of carbides ( Ref 75). As the wear process continues, the conformity of the contact region increases leading to an increase in the apparent area of contact (Fig. 10a), which reduces the contact pressure ( Ref 14,16,87). However, the wear debris within the contact region (Fig. 10j) increases the contact pressure as the real area of contact is now dictated by three-body abrasion of the (i) coating material, (ii) counter body and (iii) wear debris (Ref 14-16, 75, 87). The wear debris can roll and slide within the contact region, leading to a change in friction and wear behaviour. The debris is composed of both the coating and counter body material. Even for a relatively lower hardness counter body such as steel sliding against a WC-Co coating, the preferential wear of the metal matrix and carbide pull out is accelerated by the presence of detached carbides in the

WC-17%Co
Knoop ( Table 2.   [90][91][92][93][94][95], which can influence matrix erosion. Such transformations play an important role in the mechanical properties and wear mechanism of Co-based Stellite TM alloys (Fig. 3). However, as the weight percentage of Co in WC-Co coatings is relatively low, there has been no reported observation of a strain induced FCC to HCP transformation in the WC-Co coatings during sliding contact. The presence of tungsten in the cobalt matrix due to carbide dissolution and formation of etaphases due to carbon loss can, however, alter its properties, especially hardness and fracture toughness ( Ref 30,36,38,42,44).
The binder mean free path decreases as the carbide size decreases, which decreases the extent of preferential matrix erosion ( Ref 41). Nanocomposite coatings are expected to have better resistance to abrasive wear processes due to reduction in carbide size, since smaller carbides resist fracture, and also leads to a reduction in the binder mean free path that reduces matrix erosion ( Ref 25,30,36). However, the carbon loss and formation of eta-phases during thermal spraying can make the matrix relatively brittle, which can increase matrix erosion. Usmani et al. (Ref 41) indicated that the average diameter of WC grains (D WC ) and the mean free path of the binder (L) can be related as: where N L is the number of non-contiguous grains intersected on a metallographic plane by a line of unit length, N S is the number of non-contiguous grains per unit area, and f is the volume fraction of the dispersed phase. They indicated that the binder mean free path (L) is nonlinearly related to the fine (WC size = 1.2 lm), medium (WC size = 3.8 lm) and coarse (WC size = 7.9 lm) carbide size distributions ( Figure 12a One aspect of three-body abrasion is the increase in the wear rate of the coating materials caused by the relatively higher wear of the counter body (Ref 14-16, 21, 89). Therefore, the higher fraction of wear debris in the contact region increases the coating wear rate. Guilemany et al. In summary, apart from the coating characteristics such as its hardness, carbide size and strength, the hardness and surface chemistry of the counter body also plays an important role in the sliding wear of WC-Co coatings due to its effect on the three-body wear and stability of the tribo-film ( Ref 15,16,40,75,79,89).

Fracture and Fatigue
The compressive strength of WC is * 1.4 GPa and its tensile strength 0.3 GPa (Ref 96). As discussed above, a typical contact pressure and the tensile stress of 1. This stress can, therefore, reach the strength of WC at asperity interactions, which can lead to fracture of carbides within the contact region. However, as the load is shared between several carbides and the matrix, the stress within the carbide only reaches these values if the carbide size increases, leading to carbide fracture as shown in Fig. 10(e) and (i) ( Ref 37,75,79). In the case of finer carbides or nanocomposite S-HVOF coatings, the carbide fracture is not observed as the stress on the individual carbides reduces due to the increase in the number of carbides within the contact region. Repeated sliding within the contact region, however, leads to fatigue crack initiation and propagation as shown in Fig. 10(i), which can lead to coating delamination ( Ref 37,75). The critical crack length for fast fracture in WC-Co is larger than the carbide size in WC-Co coatings; hence, crack propagation occurs through the metal matrix binder or along boundary regions with brittle phases for coating delamination ( Ref 75,79).
As the carbide size decreases, the binder mean free path also decreases ( Figure 12a) which can reduce the probability of fatigue crack initiation and propagation during sliding wear ( Ref 40,41). However, this behaviour is not always observed in the fracture toughness measurements as shown in Figure 12(b) ( Ref 40,41). The microstructure of the binder and type of carbides therefore play a more dominant role in this process. As the carbon loss increases within the WC-Co coating, the coating becomes more brittle due to the formation of M 6 C and M 12 C carbides, increasing coating wear ( Ref 15,16). Fracture toughness of the WC-Co coatings thus plays an important role in resisting coating fracture and delamination. The literature adapts different approaches to calculate the fracture toughness of thermal spray coatings (Ref 4). One approach for measuring fracture toughness of WC-Co coatings uses the radial crack length and has been adopted by Yang   where H is the Vickers micro-hardness, D is the diagonal of Vickers indentation (lm), C L is the total crack length (lm) emanating from the four corners of the Vickers indentation.
where a is the average indent half-diagonal size, P is the indentation load, and K L is an empirical constant that can be determined for any given indenter/specimen/indentation system combination.  41) also measured values perpendicular to the coating surface and indicated that the coatings are relatively tough and resist crack propagation when indented perpendicular to the coating surface. Their work is summarised in Figure 12(b) and shows a coating anisotropy of * 0.34-0.5 based on K IC values measured parallel and perpendicular to the coating surface ( Ref 40,41). The value of anisotropy is, however, sensitive to the measurement technique, for example the fracture toughness value measured on the coating surface can be influenced by the substrate which supports the coating during the indentation fracture, whereas the cross section measurements are not influenced by the substrate properties. Several studies have discussed the influence of complexity of lamella structure on the anisotropy (hardness, elastic modulus and fracture toughness) of thermal spray coatings ( 20) . Faisal et al. (Ref 97) also showed that the measured K IC values are also sensitive to the measurement technique since not all WC-Co coatings present classical radial cracks emanating from the indent edges for fracture toughness measurements. Fracture toughness of WC-Co coatings is, therefore, not an easily measurable quantity and literature on the sliding wear of WC-Co coatings that include K IC values is sparse ( Table 2). The fracture behaviour of WC-Co coatings during sliding wear is mostly intergranular as the evidence of carbide fracture is only observed for relatively large carbide size coatings (  concluded that submicron-sized WC particles with an even distribution at splat interfaces could cause deflection or impede cracks. This improvement in the fracture behaviour from transgranular to intergranular by reducing the carbide size to nanocomposite is consistent with the approach of S-HVOF coatings (Ref [14][15][16] for improving sliding wear performance. However, the carbon loss needs to be carefully controlled before a significant improvement in the sliding wear performance of nanocomposite coatings under a range of sliding wear conditions can be observed (Ref 14-16, 22, 25, 30, 44).
The fracture and fatigue behaviour of thermal sprayed WC-Co coatings can also be determined through rolling contact fatigue (RCF) investigations ( Ref 88). In RCF, the maximum shear stress occurs in subsurface and, hence, crack initiation and propagation are always subsurface. Asperity contact is possible under certain lubrication conditions, which leads to surface initiation and propagation of cracks ( Ref 68,69,88,98). Figure 13 . This investigation indicated that as the mechanical properties of the WC Co improved after heat treatment, then the RCF and sliding wear performance also increases. The improved behaviour was attributed to improved inter-splat bonding after the heat treatment ( Ref 98). Generally, as the fracture toughness of the coating increases, then the sliding wear and RCF attributes also increase. Higher metallic Co binder content contributes to higher ductility and fracture toughness of the WC-Co coatings and also results in less cracking during rolling contact test (Ref 79). However, no direct mathematical relationship between the RCF and sliding wear performance has been reported.

Tribo-Chemical Behaviour
The oxidation of elemental species within the test couples leads to the formation of tribo-films, Fig. 10(g) [see (Ref 14,89)]. In some cases, the tribo-film is also formed by the transfer of material in its original or oxidised form, for example there is evidence of Fe in the oxidised film forming during the sliding wear of WC-Co coatings against steel ( Ref 14,16). Oxidation occurs due to the flash temperature during asperity interaction in sliding contact (Ref 87). In the case of WC-Co coatings in sliding contacts, regardless of the counter body material, tribo-films of WO 3 and CoO caused by the isothermal oxidation of WC-Co are frequently reported ( Ref 16,40,43,89). This oxidised tribo-film covers the surface of the wear track and also fills the delaminated areas and pores within the coating ( Fig. 10  and 11). In addition to these oxides, when sliding wear tests are conducted at higher temperatures (* 200-300°C), an additional oxide of WCoO 4 has been reported ( Ref 99). A tribo-film reduces the friction and wear of the test couples by providing a low shear strength film during sliding contact. The oxide WO 3 , in particular, has been reported to exhibit good lubricating properties (Ref 89). Counter body materials in sliding wear such as Si 3 N 4 and steel also indicate tribological reactions leading to the formation of a tribo-film (Ref 100).
The formation and removal of the tribo-film during sliding wear of WC-Co coatings is a continuous process where the film thickness grows until it becomes unstable and is then removed due to sliding action to make way

Influence of Counter Body on the Tribo-Film
The W-rich tribo-film generally covers the surface of the wear track when sliding wear tests on WC-Co coatings are conducted against hard metals such as WC-Co and Al 2 O 3 (Ref 15, 75, 101). In conventional WC-Co coatings, for tests against steel and Si 3 N 4 surfaces, the oxidised layer is

Influence of Decarburisation on the Tribo-Film
The formation of a W-rich tribo-film in WC-Co coatings is not influenced by the microstructural phase of tungsten, which can consist of eta-phases, WC or W (Ref 14-16, 43, 89). In the case of decarburisation of WC-Co coatings during coating deposition, leading to eta-phases, there is still formation of a W-rich tribo-film during sliding wear (Ref 14-16, 78, 89). Hence, it is not the phase of tungsten but instead its presence on the surface of the wear track that helps WO x tribo-film formation. This understanding is consistent with other engineering alloys such as Co-based Stellite TM alloys that contain tungsten (Ref 104  [91][92][93][94][95]. This observation concerning the nature of the tribo-film further confirms that decarburisation in the WC-Co coatings leading to etaphases or W has no significant effect on the formation of a protective W-rich tribo-film (Ref 14-16, 43, 89).
In summary, the tribo-film can significantly influence the sliding wear and friction of WC-Co coatings (Ref 14-16, 43, 89). The formation and removal of the beneficial WO x tribo-film in sliding wear of WC-Co coatings is influenced significantly by the surface chemistry of the counter body and the test load (Ref 100, 101). It is, however, not influenced by the microstructural phase of tungsten within the coating material. The wear debris within the contact region also influences considerably the stability of the tribo-film (Ref [14][15][16]75).

Comparison of Failure Modes of WC-Co Coatings with Cemented WC-Co Carbides
Although the microstructure of cemented WC-Co carbides is different from WC-Co coatings ( Fig. 5 and 6), there are similarities in their wear mechanisms ( Ref 6,7,9). The sliding wear mechanism in cemented WC-Co carbides is categorised into six steps as (i) polishing of grains due to micro-abrasion, (ii) wear debris formation due to microabrasion, (iii) adhesion of wear debris to form agglomerates, (iv) removal of binder, (v) grain cracking and/or grain Binder removal also leads to cracking and/or pullout of cemented carbides (Ref 6, 7). Another factor influencing the role of the binder in sliding wear was reported by Saito et al. (Ref 9); where the increase in carbide mean free path led to increased plastic deformation and, thus, increased sliding wear rate.
The above-mentioned six wear mechanisms in cemented WC-Co are also observed in the case of WC-Co coatings ( Fig. 10) but not necessarily in the same order of these steps ( Fig. 10 and 11). For example, in the case of thermal sprayed WC-Co coatings, synergism among these mechanisms can occur, and the dictating wear mechanisms may change over time. The wear can start with loosely bonded carbides, which can then influence both three-body abrasion and stability of the tribo-film due to carbide pullout ( Ref 75). Alternatively, the wear can start with carbide fracture or subsurface delamination of the coating material ( Fig. 10 and 11). Similarly, subsurface fracture or delamination is not a wear mechanism observed in cemented WC-Co carbides, which can be attributed to the strength of material due to sintering manufacturing route (Ref 10), whereas microstructural defects lead to subsurface delamination of thermal sprayed coatings that accelerate the sliding wear process (Fig. 10).
Although tribo-film formation is not indicated within the wear mechanism of WC-Co cemented carbides by Bonny et al. (Ref 7), the presence of elemental W, Co, C and O within the wear track was observed, that is consistent with the formation of a tribo-film in thermal sprayed WC-Co coatings ( Fig. 10)  considered the sliding wear of WC against W and also discussed that the wear particles revealed the crystalline structure of W, WC, W 2 C, WO x and Co. They also conducted a molecular dynamics simulation of the test couples. It was concluded that for rough-on-rough tribocouples, the WC asperities ploughed through the W asperities. The exposed carbon layer of the WC gradually mixes with the disordered W atoms, which accommodates the sliding motion. Hence, the beneficial effects of the tribo-film, which can include W in both the elemental and oxidised form, reduces the wear and friction of sliding test couples for cemented WC-Co carbides and thermal sprayed WC-Co coatings ( Ref 11). The role of a beneficial WO x tribo-film is very similar in the sliding wear of WC-Co coatings and cemented WC-Co carbides (Ref 2, 6-9, 11, 14-16, 43, 89, 101). The authors concluded that the wear of cemented carbide in this case was controlled by the transfer of counter body (steel) material on to WC-Co during sliding wear (Ref 9). This is consistent with the observation in thermal sprayed WC-Co coatings tested against steel balls (Ref [14][15][16]75). This investigation revealed that removal of the binder phase can play a dominant role during the wear of cemented WC-Co carbides (Ref 9). Despite this difference in the sliding wear mechanism due to the steel counter body, the authors concluded that the wear rate for WC-16wt.% Co varied between 0.1 910 -6 and 0.65 910 -6 mm 3 /Nm for the carbide grain size range of 0.6-6 lm (Ref 9). Although the wear rate in their investigation was proportional to the Vickers hardness, which is consistent with other investigations (Ref 6, 7), it was discussed that the hardness and carbide mean free path cannot completely clarify the wear mechanism. This was observed when the manufacturing route of cemented carbides was adjusted to a tighter impurity composition and mean carbide sizes from those specifications of commercially available (Ref 9) feedstocks. The authors highlighted that changes in feedstock manufacture need to be considered to understand the wear mechanism since these influence the microstructure of cemented carbides and affect properties such as fracture toughness. They also concluded that the wear rate should consider the shape and size distribution of WC grains (Ref 9).
The investigation by Saito et al. (Ref 9) concerned cemented WC-Co carbides and related wear mechanisms to the microstructure. These considerations are as equally important concerning the wear rate and wear mechanism of thermal sprayed WC-Co coatings due to their complex coating microstructure, as discussed in 00 Structure-Property Relationship in Sliding Wear'' and 00 Semiempirical Mathematical Relationship in Sliding Wear of WC-Co Coatings'' sctions.

Structure-Property Relationship in Sliding Wear
This section develops wear performance relationships between the sliding wear test parameters and material properties of the test couples. These include design parameters such as normal loads, sliding distance, sliding velocity for service loading and material properties such as hardness, elastic modulus and fracture toughness of the test J Therm Spray Tech (2021) 30:800-861 829 couples (Fig. 4). This objective requires a rigorous assessment of the literature, where there are often gaps in the provision of explicit materials characterisation data. Care has been exercised to select high pedigree data for analysis, and in some cases, values have been estimated from graphs or converted from wear volume loss to wear rate. These estimated and 1. The features of the coating microstructure that include the phases present, the extent of decarburisation, the mean free path of the binder, porosity, inter-splat bond strength, cracks and impurities. 2. The bulk properties of the coating itself, e.g., hardness, fracture toughness and elasticity. 3. The wear test parameters such as test type, material couple and test conditions. 4. Tribological incidents in the wear environment, e.g., three-body abrasion, tribo-film, binder and carbide interactions.
The following sections analyse these design and performance parameters to develop structure-property relationships.

Influence of Carbide Grain Size (CGS) on Hardness and Sliding Wear Resistance
The usual methods of manufacturing feedstocks for thermal spray coatings are by agglomeration and fuse/sintering processes. Commercial powders are manufactured with different carbide grain size (CGS) and binder composition that influence the wear rate (Table 2). In cemented WC-6 wt.%Co carbides, the resistance to abrasive wear depends upon CGS and can be approximated by the relationship K IC 3/8 H 1/2 /D wc , where K IC is the fracture toughness, H is the hardness, and D wc average WC grain size diameter (Ref 41). Although this relationship does not hold for thermal sprayed WC-Co coatings, it does show the dependency of wear rate on CGS. A reduction in CGS is expected to reduce the binder mean free path (Fig. 12a) that forces the crack propagation behaviour from transgranular to intergranular ( Fig. 10 and 11), hence encouraging cracks to propagate through the tougher binder and thereby increasing the toughness of the material. As discussed in 00 Sliding Wear Mechanisms 00 section, a reduction in the binder mean free path also reduces preferential matrix erosion during sliding wear ( Ref 40,41). Reduction in CGS is also expected to reduce the variation in hardness due to evenly distributed carbides within the microstructure, as described by Shipway et al. (Ref 37, 105).
Conventional WC-Co coatings exhibit a CGS of roughly 2-7 lm (Table 2). Several investigations have considered reducing the CGS to sub-micrometer-sized and nanosized carbides for wear-resistant applications (Ref 14,22,25,38,42). However, there is a compromise since investigations on WC-12Co coatings showed that finer carbides are associated with a lower thermal load capacity, which leads to decarburisation and dissolution of WC into the binder that increases binder brittleness (Ref 14-16, 35, 65, 73). The increase of thermal loading on smaller CGS affects the carbides on the surface of the powder particles, thereby affecting the inter-splat features. The resulting brittle phases become a path for crack propagation around and under the splats that leads to splat delamination (Ref 37). These earlier studies, therefore, reported that the wear performance of reduced CGS coatings was inferior to conventional coatings (Ref 37). An important retrospect comment, however, is that the spray parameters were not optimised in these studies for the reduced CGS and optimum performance would not be expected for these coatings that were manufactured using the same spray parameters.
Further studies showed that the reduction in CGS produced more wear-resistant WC-Co coatings, especially where nanosized carbides are used in HVOF ( ) indicated that as long as the carbon loss was below 60 wt.% then the hardness and sliding wear resistance of HVOF WC-Co coatings was not affected significantly. In plasma-sprayed coatings, the plasma spray torch and gas mixture heavily influence the microstructure and properties of the coating. Zhu et al. (Ref 22) produced coatings using VPS, which limited the decarburisation due to the high temperatures inherent in such plasma spray technologies (Fig. 5). The nanostructured VPS coating performed significantly better than the conventional counterpart at loads of 20-80 N when using the ball-on-flat methodology against an alumina ball ( Ref 22). The improved wear resistance of the nanostructured coating was attributed to its higher hardness and toughness. This allowed the nanostructured coating to perform well compared to the conventional coating. Sanchez et al. (Ref 21) studied the wear performance of coatings produced using Ar-H 2 and Ar-He. The H 2 gas mixture provides a higher enthalpy and produced a coating with higher decarburisation and more amorphous phase that resulted in relatively poor sliding wear performance (Ref 21).
Ahmed et al. (Ref 14) produced a nanostructured coating using S-HVOF where the total wear rate of the test couples was better than a conventional coating. When comparing the influence of the counterbody on the total wear rate, S-HVOF coatings performed relatively better when tested against a WC-Co ball (Ref 16, 107). The S-HVOF coatings exhibited a higher hardness but also a higher carbon loss, as inferred by the presence of W that was detected by XRD of the S-HVOF coating, Fig. 5  A further point is that the S-HVOF process has the capability of manufacturing thin coatings that are unachievable by conventional thermal spray technologies. Similarly, improved sliding wear performance of nanostructured coatings depends on an even distribution of nanosized particles (Ref 79), which is not a limitation in the S-HVOF coatings as all WC particles are nanosized (Ref [14][15][16]73). There is, accordingly, an optimisation trade-off concerning the design specifications for these nanocomposite coatings. Figure 14(a) summarises the relationship between CGS and Vickers hardness and indicates that the reduction in carbide size can lead to increased hardness (Ref 21-23, 25, 28-31, 33, 34, 38-40, 43, 44, 105). There are several outlier data points that are attributed to non-optimised spray parameters or eta-phase-transformations ( Ref 21,22). Increased retention of WC in coatings tends to produce coatings with higher hardness. Although the hardness of W 2 C is higher (HV = 3000 kgf/mm 2 ) than WC (HV = 2400 kgf/mm 2 ) (Ref 41), the greater proportion of retained WC still produces a coating with higher average hardness and lower wear rate. This is consistent with the findings of Ang et al. (Ref 60) where HVOF coatings produced a higher normalised hardness (^8-16 GPa) in comparison of APS (^6-12 GPa) and cold-sprayed (^7-12 GPa) WC-Co coatings. The materials science reasoning is that decarburisation of WC not only produces W 2 C but also W ( Fig. 5 and 6), which significantly reduces hardness (Ref 41). In addition, the eta-phases of M 6 C and M 12 C reduce fracture toughness; thereby reducing wear performance. Figure 14(b) indicates that, although the CGS has no direct correlation to the sliding wear performance, the wear performance is related to the extent of decarburisation. The paucity of data in some publications restricted their use in Fig. 14; but these investigations related to the influence of CGS have been included in Table 2.
In summary, the wear performance tends to improve when CGS is reduced, if decarburisation and other mechanical properties are controlled, which demands optimised spray parameters when depositing WC-12Co coatings. Furthermore, HIPing can increase the hardness and elastic modulus and influence the wear performance ( Ref 15,75,98,102,108) in conventional and S-HVOF coatings ( Table 2).

Influence of Binder Composition, Content and Test Temperature On Sliding Wear Performance
The binder composition plays an important role in WC cermet coatings by allowing the coating to perform under high temperature and corrosive environments (Ref    [120][121][122][123][124][125][126][127] presents the wear performance of WC cermet coatings with respect to binder composition. The cermet coating composed of WC-NiCrBSi is not well-documented in terms of its sliding wear in comparison to WC-12Co coatings ( Table 4). The NiCrBSi alloy has attractive properties as a binder material that is attributed to (i) the wettability of Ni, (ii) an increase in the resistance to oxidation and corrosion by Cr, and (iii) a reduced melting temperature conferred by boron and silicon (Ref 122). Coatings of WC-NiCrBSi, when compared to WC-12Co coatings, demonstrate more consistent performance in high-temperature environments, as observed by Bolelli et al. (Ref 128). These authors also reported that the sliding wear performance of a WC-12Co blend was better than a WC-NiCrBSi coating. . This investigation was conducted from room temperature to 700°C. The Vickers micro-hardness values of these coatings were similar despite the difference in microstructures between these HVOF and HVAF coatings. The reinforcing effect of WC-Co decreases the wear rate of composite coatings (&1910 -6 mm 3 /Nm) by more than an order of magnitude, compared to unreinforced ones (&1-2910 -5 mm 3 /Nm). The increase in test temperature led to an increase in the wear rate, which arose from the oxidation of WC-Co. Variations of the NiCrBSi alloy composition with the addition of WC have been investigated in other sliding wear investigations (Ref 56, 102, 108).
WC-12Co, under sliding wear conditions, is considered as an ideal coating material amongst commercially available WC cermet coatings (Tables 2 and 4). Table 4 and Fig. 15 specify that adjustments in binder composition are justified when the environmental factors are changed to high temperatures or corrosive environments. Figure 15(d) shows an analysis of the variation in sliding wear rates for APS, HVOF, HVAF, LPPS coatings over 115,116,121,128). The data in this analysis are presented in the log scale to visualise the large variations in sliding wear rate with temperature and argon      The bold text identifies the coating composition under the column''Coating parameter'' and the parameters underneath signify any other variation in the coating parameter for the same composition *Converted (9 10 -6 mm 3 /NÁm), * Estimated from illustrations environment. The lower part of the figure shows a comparison of wear rate up to * 50 9 10 -6 mm 3 /Nm to highlight the wear rate differences in oxygen environment up to a temperature of 600°C. A comparison of Fig.15(d) with Fig. 14(b) shows that the wear rate for higher temperature tests can increase by orders of magnitude (* 0.1-400 9 10 -6 mm 3 /Nm) in comparison with room temperature tests (* 0.1-6 9 10 -6 mm 3 /Nm). A further inspection of the data shows that in most cases the wear rate over the temperature range of 25-650°C is below 50 9 10 -6 mm 3 /Nm, as the higher wear rates, shown in Fig. 15 115) indicated that the volume loss of WC-Co coatings is significantly higher than that in air for sliding wear tests conducted in the range of * 25-600°C. They concluded that the WC-Co coatings should not be used for applications in oxygen-deficient environments.
Two factors play an important role as the temperature increases from room temperature up to * 600°C. The first factor is the temperature-dependent changes in the mechanical properties of the coatings. As the temperature increases, the elastic modulus and hardness decreases, whereas fracture toughness is expected to show an increase due to thermal softening. However, due to microstructural phase changes leading to eta-phases at a higher temperature, the toughness also decreases as indicated by Geng et al. (Ref 115). These temperature-dependent variations in the mechanical strength are well documented for cemented WC-Co carbides (Ref 129,130); however, they are more complex in WC-Co coatings due to their microstructure ( Ref 115). In particular, the changes in the cracking and delamination behaviour of WC-Co coatings at higher temperatures depend on the coating microstructure and, thus, the coating process and parameters. This was discussed by Geng et al. (Ref 31) where APS coatings showed significant delamination at * 250°C. Other studies using HVOF, HVAF and LPPS coating processes did not show such an aggressive delamination behaviour until a much higher temperature of * 650°C, as shown in Fig. 15(d) (Ref 31, 115, 116, 121, 128).
The second factor depends on the tribological test conditions and surface chemistry of the coating material. Geng et al. (Ref 115) investigated the microstructural phase changes in HVOF-sprayed WC-Co coatings using XRD performed under air and argon environments. They indicated that in air, WC peaks remain dominant in the microstructure up to 600°C; whereas WO 3 and CoWO 4 peaks are also observed above * 450°C. As discussed earlier in ''Tribo-Chemical Behaviour'' section, WO 3 and CoWO 4 are important tribo-films that reduce the sliding friction and wear behaviour of WC-Co coatings. There were no significant changes below 550°C for the XRD measurements carried out in an oxygen-deficient, argon environment, whereas M 12 C (Co 6 W 6 C) forms at 600°C ( Ref 115). Some oxygen is present even in a presumed argon environment, so the formation of some CoWO 4 at 650°C was reported.
Note that even for sliding wear tests conducted at room temperature, the flash temperature caused by the asperity interaction still exceeds * 500°C, leading to the formation of a beneficial tribo-film WO 3 (Ref 16,40,43,89,101). Hence the role of oxidation leading to the formation of tribo-film at higher temperature sliding wear tests is both beneficial and counterproductive. It is beneficial for protective tribo-film formation and stability, but counterproductive due to detrimental phase changes that lead to loss of carbides due to oxidation and reduced mechanical strength. In some respects, the WC-Co coating becomes sacrificial at higher temperatures due to material loss through oxidation during the formation of the protective oxide film. Other wear mechanisms of preferential matrix erosion and cracking are also observed at higher temperature sliding wear tests; however, the wear mechanism is dominated by the formation and stability of the oxidised tribo-film.
Structure-property relationship and tribological test conditions dictate the performance during high-temperature sliding wear tests. The test load and temperature influences the stability of the tribo-film. This effect was J Therm Spray Tech (2021) 30:800-861 839 indicated by Geng et al. (Ref 31) where back-to-back tests conducted at 8 and 28 N load using Si 3 N 4 counter body ball material revealed a significantly lower friction coefficient and sliding wear rate due to the stability of the tribofilm at the lower load (Fig. 15d). The figure shows that a change in binder Co content from 12 to 17 wt.% increases the wear resistance due to the relatively lower oxidation of WC and beneficial influence of Co in improving both the mechanical strength and formation of CoWO 4 tribo-film at higher temperatures. This outcome is consistent with the investigation of Xie et al. (Ref 116) where the sliding wear rate at 500°C was lowest for WC-17Co by * 50%, followed by Cr 3 C 2 -25NiCr, WC-12Co, WC-10Co-4Cr, respectively (Fig. 15d).

Influence of Spray Parameters on Sliding Wear Performance
HVOF and APS processes are the most commonly used methods for the deposition of WC-Co coatings ( Table 2). APS coatings are inferior due to their higher porosity, low adhesion strength and adverse phase reactions to those produced using HVOF (Table 2). Hence, there are less recent studies of WC-Co deposition by APS processes. HVOF systems use liquid or gas fuels and plasma systems use plasma gases, both of which need to be optimised to keep the feedstock from overheating in-flight and decomposing into eta-phases. Qiao et al. (Ref 44) studied the effect of changing the fuel to oxygen ratio (F/O) on the wear performance of coatings produced with different CGS feedstocks. The fuel ratios were classified as neutral (complete combustion), oxidising (excess oxygen) and reducing (excess propylene) and are indicated in Table 2. In-flight temperature and velocity of the powders were measured using a Technar DPV 2000 system. In the neutral condition, the feedstock powders achieved the highest temperature, * 2200°C, and the highest velocity, * 700 m/s. Since the combustion in the reducing and oxidising flame is not complete, powder temperatures (1550-1800°C) and velocities (450-550 m/s) were lower than in the neutral flame. The neutral ratio produced the highest temperatures and also exhibited more decarburisation in the deposited coatings, followed by the oxidising ratio in most cases. The importance of the in-flight temperatures was highlighted in this study since the coatings with high degrees of decarburisation (neutral ratio) performed better in sliding wear than their counterparts. The lower temperatures affect the melting of particles in-flight and, therefore, weakening the adhesion strength of the coating which affected the wear performance. The dominating wear mechanism for coatings formed from the neutral ratio was binder extrusion followed by carbide pull out with a small degree of three-body abrasion. The carbides were well-embedded in the binder with low porosity. On the other hand, the coatings produced under reducing or oxidising ratios were dominated mainly by three-body abrasion since the porosity was higher and bonding between splats was weak, which lead to carbide pull out (Ref 44) Table 5 (Ref 131) shows the changes in bulk mechanical properties of the coating and the wear track width for changing fuel/ oxygen ratios. The coating with the highest wear resistance (96/1020 fuel to oxygen ratio) showed good carbide retention, whereas the coating with the highest degree of decarburisation (136/1020 fuel to oxygen ratio) displayed brittle fracture in the worn surface. The brittle fracture of the coating would contribute strongly to three-body abrasion (Ref 131).
The preferred wear mechanism for WC-Co-based coatings was binder extrusion followed by carbide pull out with small degrees of three-body abrasion (Ref 131). The coating wear rate becomes significantly higher when threebody abrasion becomes the main wear mechanism. Improvement of the sliding wear resistance carbide bonding is achieved by the partial dissolution of the carbide into the binder. Girolamo et al. investigated different compositions of gases to produce WC-12Co wear-resistant coatings (Ref 132). Table 6 summarises the parameter and bulk mechanical properties and their findings of wear performance (Ref 132, 133). Decarburisation was extensive for the coating produced using Ar-H gases. Decomposition of WC to b-WC, which occurs at high temperatures, led to brittle phases and promoted fracture-based wear mechanisms followed by three-body abrasion. The Ar-H-He and Ar-He coatings exhibited heterogeneous microstructures with embedded carbides in both un-dissolved and dissolved states within the binder ( Ref 132). The authors reported that wear mechanics for these (Ar-H-He and Ar-He) coatings was composed mainly of binder extrusion and carbide pull out or fragmentation. The coating produced using He showed the highest porosity of these coatings and revealed the second-highest degree of decarburisation. The porosity was linked to the low impact velocity of the particles due to the low mass flow rate of He. It was suggested that the porosity promoted fracturing since it presented a defective microstructure that incorporated the presence of brittle phases (Ref 132).
A similar study by Girolamo et al. (Ref 133) investigated nanostructured coatings manufactured under different gas compositions. Reported parameters, properties and results are summarised in Table 6. The Ar-H gas coating had lower wear resistance than the Ar-He coatings (Ref 133). The Ar-He coating wear performance was similar (* 4.4-4.8 910 -6 mm 3 /Nm) in both studies (Ref 132,133).
In summary, spray parameter selection is critical for producing a wear resistance coating with high wear resistance. The main aspects that are affected by these parameters are the in-flight temperature of the particles and inflight velocity (Fig. 1). Optimisation of these parameters, therefore, contribute to the bulk mechanical properties and tribo-mechanical performance. A controlled degree of carbide dissolution into the binder is important to reduce carbide pullout during sliding wear of WC-Co coatings ( Ref 43,44,131).

Influence of Coating Post-treatment on Wear Performance
Post-treatment of WC-Co coatings improves their wear performance by reducing porosity, and enhancing adhesion between splats and to the substrate ( Ref 75,76). Heat treatment, laser remelting and hot isostatic pressing (HIPing) are the main post-treatment methods used for thermal spray coatings ( Ref 15,75,102,[134][135][136]. Each of these methods has its respective benefits, as listed below. Heat treatment increases hardness and is cost-effective but does not reduce porosity. Laser remelting reduces porosity to almost zero but is not recommended in a lubrication environment due to the lack of cavities that act as a reservoir for lubricant. Furthermore, the high energy density of the laser creates unwanted heat affected zones and undesirable phase reactions on the surface. HIPing can deliver the benefits of heat treatment and laser remelting, while avoiding their deficiencies. A WC-NiCrSiBFeC coating was produced (Stoica et al. (Ref 102)) in the form of two layers, an initial layer of WC-40 NiCrSiBFeC that was followed with a top layer of WC10 NiCrSiBFeC. These coatings were HIPed at 850 and 1200°C and allowed the elemental species across the distinct boundaries of the as-sprayed coatings to diffuse. The degree of diffusion of the coating with the substrate at the interface increases with higher HIPing temperatures as shown in Figure 16 (Ref 56). A similar influence on the diffusion zone in WC-Co coatings was observed in another study (Ref 75) (Fig. 8b and c). Their work on HIPing of WC-NiCrBSi coatings showed that post-treatment improved the sliding wear resistance of these coatings by *40-50% for steel and Si 3 N 4 counter body materials, respectively, at a HIP treatment of 1200°C (Ref 102). Similar studies also showed improved performance after HIPing for WC-Co coatings in the range of *15-28% for a  127 HV0.3). HIPing increased the hardness of the nanostructured coatings to 1130 ± 140 and 1201 ± 74 HV0.3. The conventional coating performed better than both types of nanostructured coatings under counter bodies of Si 3 N 4 and AISI 440C stainless steel, whereas the S-HVOF coating performed better against a WC-Co counter body (Ref 15, 16). The HIPed coatings exhibited better performance (*12-22%) than the as-sprayed ones when a ceramic counter body was used.
Laser remelting was used by Mateos et al. to treat a plasma sprayed WC-17Co coating ( Ref 135). The microstructure of the porous as-sprayed coating showed that carbides were partially dissolved into the binder. The laser-treated coatings indicated that carbides regained their polygonal structure as well as revealing dendritic features in the binder, as shown in Figure 17 (Ref 135). The hardness of the laser-treated coating doubled and the sliding wear performance was enhanced under a variety of loads and sliding speeds. The as-sprayed coating showed that the tribo-film was composed mainly of counter body material. There was crack propagation through the binder that resulted in coating fragmentation, leading to threebody abrasion. The post-treated coating also showed similar mechanisms but wear rate improved (Ref 135) as shown in Figure 18.
Sohi et al. reported the results of heat-treated plasma sprayed WC-12Co coatings in argon, at 650, 900 and 1150°C (Ref 136). The coatings treated at 900 and 1150°C exhibited better average wear resistances of *61910 -2 mg/m and *72910 -2 mg/m, respectively, than the as-sprayed coating (*83910 -2 mg/m), which has been linked to the formation of eta-phases in the coatings.
The coating treated at 900°C demonstrated the best wear performance (*26.5% reduction in the wear rate when compared to the as-sprayed coating) of all the coatings, whereas the one treated at 650°C (average wear rate In summary, the post-treatment of thermal sprayed coatings generally improves the wear performance. But there is a technical limit to this improvement, as can be seen from the investigation by Sohi et al. (Ref 136). Most heat treatment methods also enable diffusion of the substrate elements into the coating ( Ref 136). Careful selection of post-treatment parameters is critical and must take into account the substrate material and its effect on residual stress (Ref 56, 57, 137).

Comparison of Wear Rates of WC-Co Coatings with Cemented WC-Co Carbides
Cemented WC-Co carbides have been extensively investigated for their sliding wear performance and wear mechanisms (Ref [6][7][8][9]11). Several investigations, summarised in Table 2, also use a cemented WC-Co counter body against the WC-Co thermal sprayed coatings, which provides a direct comparison of sliding wear performance between these two materials. The hardness of cemented WC-Co varies between 1200 and 1913 HV for cobalt compositions of 6-15wt.% (Ref 6, 7). Typically for WC-12Co, cemented carbides exhibit Vickers hardness of 1240-1253 HV, which is comparable to the hardness values of WC-12Co coatings indicated in Table 2. The fracture toughness of cemented WC-Co carbides varies between 8.8 and 20.6 MPa.m 1/2 for cobalt compositions of 12-6wt.%, respectively (Ref 7). The anisotropy in thermal sprayed WC-Co coatings presents lower values of fracture toughness parallel to the coating surface (Fig. 12b), and the fracture toughness values in Table 2 are relatively lower than those for cemented carbides (Fig. 3). Despite this difference, the sliding wear rates of cemented WC-Co carbides shown in Figure 19a indicate values of 0.3910 -6 -0.8910 -6 mm 3 /N m for cobalt contents of 6-20 wt.%. This is a narrow range of wear rate when compared to the wear rate of WC-Co coatings that vary from 0.04910 -6 to 19.2910 -6 mm 3 /N m for cobalt contents of 9-17 wt.%, respectively, as summarised in Table 2.
Similar to thermal spray coatings, the wear rates in cemented WC-Co carbides are sensitive to the tribological test conditions. Bonny et al. (Ref 7) investigated the influence of test parameters by using ASTM G133 method (Ref 81) for the sliding wear performance of cemented WC-Co carbides. The load was varied between 15 and 100 N, and the reciprocating speed from 10 to 30 Hz (equivalent to 0.3-0.9 ms -1 ) up to a sliding distance of 10 km. Tests were conducted against WC-6Co pins of 4.08 mm radius. This study indicated that the increase in the normal load increases the wear rate of cemented WC-Co carbides, whereas the friction displayed an opposite trend, whereas both wear and friction increased with the increase in the sliding velocity. Regardless of these changes in the test conditions, the wear rates for cemented WC-Co carbides stayed in a small range of 0.016910 -6 -0.022910 -6 mm 3 /N m, which is comparable to the lowest values of wear rate in Table 2 for thermal sprayed WC-Co coatings.
In summary, the wear rate of cemented WC-Co carbides is relatively low (Ref [6][7][8][9]11). The change in the wear rate with tribological test conditions is also within a narrow range when compared to thermal sprayed WC-Co coatings (Table 2). Similarly, the wear resistance of cemented WC-Co carbides is the highest when compared with other carbides such as Cr 3 C 2 -Ni and TiC-NiMo (Ref 6) (Fig. 19). The investigation of these cemented carbides also confirms that the sliding wear volume is almost directly proportional to the Vickers hardness of cemented WC-Co and TiC-NiMo, as shown in Figure 19b Figure 4 shows the complex interdependency of sliding wear test conditions on the microstructure and tribo-mechanical properties of coatings. As reported by the authors ( Ref 15,16), correlations of the influential factors dictating the sliding wear performance have not been successful. Wood (Ref 45) indicated that in the case of thermal spray coatings, the dependencies of wear rate on the binder metal, porosity, hardness and Young's modulus were not found. Single variable relationships with the wear rate with all other factors constant have, however, been reported. Mateos et al. (Ref 135) studied the effect of test parameters such as normal load and velocity in a block-on-ring test method with steel as a counter body. In Figure 18, the sliding wear rate is presented against the velocity at a constant load of 377 N. A second-order polynomial relationship exists between these parameters ( Ref 135).

Semiempirical Mathematical Relationship in Sliding Wear of WC-Co Coatings
Generic relationships that can be applied to a combination of test conditions and tribo-mechanical properties of the test couples during sliding wear of coatings are not available in published literature. Authors have, therefore, extracted these relationships from the available studies, where Q is the wear volume depth (mm 3 ), l is the sliding distance (m), N is the normal force (N), H is the hardness of the worn surface, and K is a dimensionless wear coefficient. The term (K/H) was replaced by a dimensional wear rate (k D ) in the units of mm 3 /N m. However, in the current analysis of WC-Co coatings, since the relative ball hardness is lower for steel and higher for the WC-Co and Si 3 N 4 , the ball to the coating hardness ratio (H ball /H coating ) is considered. 3 The ball or pin diameter and material can vary between the investigations. Therefore, a variable based on the Hertzian contact pressure (C p ) is introduced to cater for the differences in ball geometries and elastic properties of test couples. 4 Fracture of surfaces inevitably leads to the formation of wear debris so the K IC of the coatings is considered as another variable. The fracture toughness of the ball or pin material (K IC ball or K IC pin ) did not show a strong dependency on the wear rate and is not included in this analysis. 5 The relationship between the coefficient of friction (CoF) is expected to be linear, where an increase in CoF can lead to an increase in the wear rate. However, this relationship is not always valid for the WC-Co test couples since the role of tribo-film changes for different test couples. As well, wear debris in the wear track can present a complex interdependency. Hence the CoF is introduced as a variable in the current analysis. 6 The coating porosity can influence the wear performance since dense coatings are more desirable. Hence coating porosity was also included as a variable in the current analysis. Figure 20(a) shows the relationship between the contact pressure (C p ) and CoF against the wear rate for a variety of test couples. The wear rates lie within the C p 9CoF envelope of 0.5-2.5. Differences between plasma spray and HVOF coatings can be separated using the 'C p 9CoF' parameter; however, there is no correlation with respect to the wear rate. An alternative graphical representation of this combined data set relates the wear rate with another parameter that includes the hardness of the ball and coating materials; i.e., '(C p 9 H coating 9 CoF 0.1 )/(H ball 9 v 0.5 )', Figure 20(b). Although there is no straightforward relationship between the wear rate and (C p 9 H coating 9 -CoF 0.1 )/(H ball 9 v 0.5 ), this figure facilitates comparisons of WC-Co coatings performance with other wear-resistant coatings or cemented carbides. Figure 21 shows the values of wear rate against coating hardness and indicates that there is no direct relationship between the coating hardness and sliding wear rate. This behaviour is significantly different to cemented WC-Co carbides, where sliding wear rate is proportional to hardness (Fig. 16). However, if C p and the ratio of 'ball to coating hardness' are considered, as shown in Figure 22, there is a trend for an increase in the wear rate with C p 9(H ball /H coating ). The analyses presented in Fig. 20, 21, and 22 are based on physical quantities of CoF, hardness and contact pressure that are generally reported.
Relationships between wear rates and extrinsic materials properties would be clearer with the introduction of the coating fracture toughness, K IC , which is often not available within published literature. Figure 23(a) presents a compilation of data in cases where the K IC values for the coating material can be inferred from published literature. The wear rate is plotted against (H ball /H coating ) 3/4 9 (C p / H ball ) 2/3 /(K IC coating ) 1/2 /(H ball /E ball ) 1/7 , which was previously applied to S-HVOF coatings (Ref 15) with success. A linear trend is more distinct for the plasma sprayed coatings when using this relationship.
The correlation between the wear rate and tribo-mechanical properties of test couples and test conditions is not clear, primarily due to the way in which the wear rate is defined in the published literature. The force term in the wear rate equation depends on the geometry and material. For example, a load of 25 N on a steel ball of 10 mm diameter is not the same as a load of 25 N on a WC or ceramic ball in terms of the contact pressure. Even for the same ball material, as the load changes the contact pressure changes nonlinearly.
The analysis is more complex when comparing the results with ASTM G77 (Ref 83) block-on-ring tests because the loads are relatively higher, but the contact geometry is more conforming, so the contact pressure is relatively low. Hence, the influence of geometry in the wear rate calculation can be included by replacing the force term with the Hertzian contact pressure (C p ). This approach enables a comprehensive analysis of wear rate and permits differentiation between coating processes. Therefore, the wear rate has been expressed in terms of contact pressure instead of load, i.e., in units of 'mm 3 /GPa m' instead of 'mm 3 /N m', and it is possible to differentiate the HVOF and plasma sprayed coating behaviours. The change of wear rate from load to contact pressure is more rigorous because it includes the dimensions and elastic properties of both the coating and ball or pin materials. This is represented in Figure 23(b) where the oval-shaped envelope distinguishes the HVOF and APS coatings, much in the same fashion depicted by Ashby schematic diagrams for materials selection. The x-axis is presented as a log scale to identify differences among the coating types that display wide ranges of material properties. Figure 23(c) shows a modification to this relationship where coating porosity is also incorporated into the mathematical construct. Here the distinction between the HVOF and APS coatings is more subtle, as APS coatings generally exhibit higher porosity and relatively lower fracture toughness due to carbon loss.  mathematical modelling of wear rates. In this section, the numerical modelling of sliding wear rate is considered. Finite element numerical models that consider the evolution of wear volume during sliding contacts, and hence predict the wear rate, can provide useful predictions of the wear rate. Two approaches can model sliding wear by a finite element method (FEM), i.e., mesh deletion and node adjustment (Ref [139][140][141][142][143][144]. The mesh deletion method is based on the failure of elements that are then removed from the model to estimate the wear rate. This method is illustrated in Figure 24 Once the failure criteria is reached within the element, the element is deleted from the model as wear debris and new surfaces are created. This iterative approach has also modelled machining processes such as milling or the interaction of the cutting tool on workpiece leading to the formation of chips and a newly machined surface (Ref [139][140][141]. This approach can be applied to the sliding wear process using an algorithm that redefines the contact interactions when new surfaces are formed after mesh deletion for continuous modelling of the wear process (Ref 139,140). The minimum amount of wear volume that can be removed depends on the mesh density, or the size of elements, which need to be considered in terms of the computing cost and required dimensional precision. The Fig. 21 Wear rate of coatings from literature against their respective Vickers hardness, the counter bodies used are mentioned as well (*value estimated from graphs, *value converted to 910 -6 mm 3 /Nm, pl plasma-sprayed coating, hv HVOF-sprayed coating, shv S-HVOF-sprayed coating) Fig. 22 Wear rate from literature against the product of maximum contact pressure (Cp) in GPa with the ratio of the hardness of the ball (H ball ) to the coating (H coating ) (*value estimated from graphs, *value converted to 910 -6 mm 3 /Nm, pl plasma-sprayed coating, hv HVOF-sprayed coating, shv S-HVOF-sprayed coating) second approach adjusts the nodes in the contact zone based on the contact pressure (Fig. 24b). As the body wears, the nodes are adjusted to account for the worn material ( Ref 85,93,[142][143][144]. The dependency of mesh density indicated above for the mesh deletion technique can, therefore, be reduced using the node adjustment technique.

Computational Wear Modelling
Although the finite element technique has not yet modelled the prediction of sliding wear rate in thermal spray coatings, Bolelli  . These numerical models require experimental data that is often difficult to attain, and, in principle, this approach can be applied to thermal spray coatings. An SEM image of the cast Stellite TM alloy is shown (Fig. 25b) since it should always be recalled that FEM models must always be related to the microstructure.
Wear modelling using computational methods are based on the Archard's wear rate which is further defined in ''Influence of Spray Parameters on Sliding Wear Performance'' section ( Ref 138). The FEM approach requires that the wear rate needs to be in a form where it can be applied to the nodes in the model. Archard's wear rate is used in FEM to simplify the computational model. Nevertheless, the FEM approach depends on experimentally measured values of wear coefficient. Williams (Ref 147) describes that wear can be generally categorised as ductile or brittle. However, brittle wear can transition into ductile wear under high compressive stress, thereby adding to the complexities of a predictive wear model. Figure 26  Lim and Ashby (Ref 149) developed a wear mechanism map for steel-on-steel contacts using pin-on-disc experiments. This was based on Archard's wear law using velocity and contact pressure as the sliding wear experimental parameters. Williams also suggested using such a wear map for accurate computational modelling, but it was essential to select the appropriate wear coefficient based on experimental parameters ( Ref 147). It is, however, challenging to identify transitions in wear mechanisms due to changes in the experimental parameters. Therefore, developing a model that simulates wear based on wear mechanism is not currently feasible by a FEM approach. Numerical models based on the Archard's wear rate (k D ) are therefore employed in FEM ( Ref 65,143,144).
Investigators have proposed models for predicting wear for materials, and the outcome is a compromise by exploiting Archard's wear law. Meng  Expanding wear mechanism map methodology to other materials is challenging. There are no wear maps, especially for composite materials such as thermal sprayed cermets, where coating microstructure depends not only on the coating process but also on process parameters and powder characteristics (Fig. 4). Wear maps are subject to change based on the type of sliding wear test, such as pinon-disc, ball-on-flat or block-on-ring. The proposed semiempirical mathematical approach in ''Semiempirical Mathematical Relationship in Sliding Wear of WC-Co Coatings'' section may develop into a computational wear model. Such a model, for instance, may be applied to thermal sprayed WC-12Co coatings as a predictive tool for the Archard's wear rate (Fig. 25). However, currently there are no such models available in the literature for thermal spray coatings.

Archard's Wear Rate in FEM
In FEM modelling, the elements are defined by nodes. Archard's wear rate (k D ) introduced in Eq 12 (Ref 85,138,150) for the sliding wear rate can be expressed in terms of the displacement of nodes in FEM. The surface nodes need to be displaced appropriately based on the wear rate. The following equation allows for the calculation of the nodal displacement.
where h is the linear wear and P n is the contact pressure at the node. The term h is further discretized concerning the sliding distance.
Equation 14 provides the change in the height of individual nodes based on the nodal contact pressure. The adjustment of the node has the magnitude of dh, in the direction of the inward surface normal, which is calculated with regard to the local contact pressure and local sliding distance. The Euler integration scheme integrates the wear law over the sliding distance ( Ref 138,151).
where h j is the total wear up to the jth increment, h j-1 is the total wear up to the (j -1)th increment, and Dh j is the wear depth of the current increment. This approach calculates the adjustment of the surface nodes that undergo wear with variations due to different wear test modelling ( Ref 143,[152][153][154][155][156][157][158][159][160][161]. The contact region undergoes several wear mechanisms concurrently under service loading. The experimentally measured Archard's wear rate is assumed to cover this aspect of wear (Ref 85,138). Another limitation of computational wear modelling is the transition from severe wear to mild wear after the running-in period. Thus, a comparison between simulation and experimental data is difficult because FEM simulation uses a constant global wear rate, whereas the wear rate changes with test duration (Ref 85) and tribological test conditions (Fig. 25, 26) Fig. 23 (a) Wear rate from literature against the logarithmic scale of the algorithm, (b) adjusted wear rate parameter against the algorithm and (c) adjusted wear rate parameter against the algorithm values further corrected with non-porosity factor for WC-12Co coatings. (*value estimated from graphs, *value converted to 910 -6 mm 3 / Nm, pl plasma-sprayed coating, hv HVOF-sprayed coating, shv S-HVOF-sprayed coating) calculated based on the global contact pressure rather than the local contact pressure on the nodes (Ref 162). This reduces the computational time, but it comes at the cost of accuracy.
Another approach is to use the ABAQUS/standard user subroutine UMESHMOTION (based on FORTRAN script) ( Ref 155,158). In this approach, for every wear increment, the counter body is displaced on the basis of the wear coefficient, and the contact node is, respectively, displaced ( Ref 153,[155][156][157]. This method allows the translation of surface nodes that is followed by adaptive meshing and advection. The adaptive meshing avoids severe deformation of elements that causes errors where the model fails to converge. The mesh is re-located without creating new elements, therefore keeping the number of elements constant. Advection is the process of transferring the nodal and elemental data after the adaptive meshing has taken place, allowing for the re-mapping of the previous solution onto the new mesh. This procedure is termed as arbitrary Lagrangian-Eulerian (ALE) adaptive meshing within ABAQUS software. Figure 27 depicts the method used in wear simulations that are based on Archard's wear rate in Eq 14. Variations to this method depend on the type of sliding wear experiment. Similar methods can be used in other software such as ANSYS (Ref 143).

Computation Time Management
The wear modelling described in Fig. 27  wear step, the accuracy will be high and computation time longer, but if the wear step was too big then the model will lose accuracy. One such wear step is defined as the execution of the loop described in Fig. 27.

Validation of Wear Modelling
An  As discussed above, although it is assumed that the global wear coefficient would cover the effects of wear mechanisms, it cannot account for the transitions in wear mechanisms as the number of cycles increase. One possible solution is to use different wear coefficients at different stages of the model, but that would negate the purpose of a computational model since the wear coefficient is derived from empirical data. Another limitation of these models is their inability to account for phase transformations and the formation or removal of a tribo-layer during the sliding wear process (Fig. 4 and 10

Role of Residual Stress in Modelling Wear
The semiempirical mathematical relationship of the wear rate in ''Semiempirical Mathematical Relationship in Sliding Wear of WC-Co Coatings'' section and the FEM modelling in ''Archard's Wear Rate in FEM'' to ''Validation of Wear Modelling'' sections does not consider the role of residual stress. For a given coating substrate system, the measured through-thickness residual stress profile in thermal spray coatings depends not only on the coating process and process parameters, but also on the stress measurement technique ( Ref 58,[163][164][165][166]. The throughthickness magnitude of residual stress can be significant, as shown for the WC-Co coatings in Figure 30 Compressive residual stress prohibits mode one crack initiation and propagation, whereas tensile residual stress can be detrimental to the coating performance. Higher velocity in HVOF spraying not only reduces the dwell time but also increases the beneficial peening effect leading to compressive residual stress in the deposit and increasing coating density ( Ref 52,53,56,57). Figure 30(a) shows compressive residual stress of a few hundred megapascals (MPa) in the as-sprayed and HIPed conditions, which can be advantageous in resisting fracture and delamination.
Measurement of residual stress is sensitive to the measurement method as shown in Figure 30(b) (Ref 57). This figure provides a comparison of residual stress measured by the non-destructive neutron diffraction and the destructive hole drilling techniques. The difference between the two results lies in the type of stress being measured, specifically, macro stress contrasted to stress based on lattice deformation. The choice of the residual stress measurement technique for wear analysis should be based on the accuracy of the measurement. For example, if the as-received residual stress profile is intended to be compared to the profile generated after the wear test, then a non-destructive technique such as x-ray, synchrotron or neutron diffraction is preferred (Ref 134). However, the surface roughness of the specimen created by the wear test method increases the standard deviation of the measured stress. Thus, the penetration depth of the diffraction method should be selected so that the measurement is representative of the material just beneath the surface. Consequently, care is necessary  The role of residual stress in the sliding wear process has been investigated ( Ref 54,134,166); however, it has not been included in the semiempirical mathematical models of sliding wear rate or numerical modelling. Note though that the numerical approach discussed in ''Computational Wear Modelling'' section can be adapted to include the throughthickness residual stress profile incorporated within the FEM of sliding wear. This approach can further reduce the error between the experimentally measured and numerically predicted values of sliding wear ( Ref 52,[55][56][57][58]163,166). Furthermore, the resulting microstructure associated with the numerical models of residual stress during spray coatings (Ref 50) can serve as the ab initio point to integrate sliding wear models and investigate both the influence of residual stress and coating architecture.

Influence of Test Temperature on Modelling the Sliding Wear Resistance of WC-Co Coatings
As shown in Figure 15(c) and (d) and discussed in ''Influence of Binder Composition, Content and Test Temperature On Sliding Wear Performance'' section, two aspects influence wear mechanisms as the test temperature increases for sliding. The first one relates changes in the mechanical properties such as elastic modulus, hardness and fracture toughness: factors that can be captured in the FEM models using temperature-dependent mechanical properties. The second factor is related to phase transformations occurring with respect to temperature during the sliding wear test. As it is not possible to model temperature-dependent phase transformations using FEM, the result of these phase changes on material properties may also be approximated using temperature-dependent material properties. The significant influence caused by the stability of the beneficial tribo-film can only be captured in the global wear rate (k D ), which contributes to complex mechanisms since the tribo-film properties change with temperature and load. Further work in this area will benefit both room temperature and higher temperature predictions of sliding wear rate using FEM.

Conclusion
The sliding wear performance of WC cermet coatings depends on parameters that link structure-property relationships to the tribological test conditions (Fig. 4). This review focuses on the effects of coating microstructure, carbide size, binder content, binder composition, spray manufacturing method, post-treatment, wear testing environments and counter body on the sliding wear of WC-12Co thermal-sprayed coatings. Sliding wear mechanisms and wear rates are also compared with cemented WC-Co carbides. The specific conclusions are as follows.
1. Coating process parameter optimisation is critical in reducing the thermal loading of WC particles, which can result in carbon loss leading to M 2 C, M 6 C, M 12 C phase transformations ( Fig. 5 and 6). A controlled amount of carbide dissolution is, however, desirable in improving the wear resistance of WC-Co coatings because it reduces carbide pullout by strengthening the matrix and improving the interface bonding between the matrix and carbide constituents. Strategies of reducing carbon loss due to thermal loading are discussed, which include cladding of WC particles and heat treatment of WC-Co coatings leading to precipitation of dissolved carbides. Heat treatment of WC-Co coatings can also lead to crystallisation of amorphous phases and improved mechanical properties, which enhances the sliding wear resistance. There is a relatively higher carbon loss in S-HVOF coatings. However, S-HVOF coatings provide a dense microstructure with hardness and elastic modulus similar to or better than the conventional HVOF coatings. 2. The sliding wear mechanism of WC-Co coatings ( Fig. 10 and 11) include carbide pullout, subsurface delamination, carbide cracking, preferential wear of the matrix and three-body abrasion caused by wear debris in the contact region. Formation and stability of tribo-films (WO x ) on the surface of the wear track is critical in reducing wear and friction during the sliding wear of WC-Co coatings. The stability of this tribofilm depends not only on the tribological test conditions but also on the chemical composition of the counter body material in both WC-Co coatings and cemented WC-Co carbides. Some mechanisms of sliding wear in WC-Co coatings such as carbide fracture/pullout, matrix erosion and three-body abrasion are consistent with the sliding wear of cemented WC-Co carbides, whereas subsurface delamination and wear related to carbon loss and the formation of etaphases is unique to thermal spray WC-Co coatings. 3. The sliding wear rate of WC-Co coatings varies significantly based not only on the structure-property relationship but also the tribological test methodology and test parameters. This leads to large variations in sliding wear rates. The sliding wear rate of WC-Co coatings varies between 0.04910 -6 and 19.2910 -6 mm 3 /N m ( Table 2). On the other hand, the wear rates of cemented WC-Co carbides vary in a narrow range of 0.03910 -6 -0.8910 -6 mm 3 /N m regardless of the tribological test conditions (Fig. 19a). The hardness and wear resistance increase in WC-Co coatings as the carbide size is reduced and decarburisation is controlled. The sliding wear rate of S-HVOF coating is similar or slightly better than the conventional HVOF coatings under specific tribological conditions. 4. The contact stress related to the load term in sliding wear rate calculation is geometry dependent. Thus, the wear rate is re-defined in terms of contact pressure (mm 3 /GPa m). Semiempirical mathematical models based on this modified wear rate are developed to link sliding wear with the structure-property relationship using (i) coating fracture toughness, (ii) the hardness of the coating and counter body material and (iii) elastic modulus of the counter body material (Fig. 23). This interdependency is substantially different to the sliding wear rate of cemented WC-Co carbides that show an inversely proportional relationship between sliding wear rate and hardness (Fig. 19b). 5. In sliding wear under non-corrosive environments up to a temperature of * 500°C, WC-Co coatings provide the best wear performance. Other binder compositions of Ni and Cr are more justified for corrosive environments (Fig. 15). In terms of the wear mechanism, the oxidation of WC becomes dominant as the sliding test temperature increases. This behaviour is both beneficial and counterproductive: i.e., oxidation at higher temperatures stabilises the tribo-film, thereby reducing friction and wear rate. But this also reduces the mechanical strength of the coating, which increases the wear rate due to oxidation. A stable tribo-film at higher temperature has a more dominant effect in controlling both the wear mechanism and wear rate than the reduction in mechanical strength. A higher Co binder content of 17% instead of 12%, increases the sliding wear resistance by almost 50% at * 500°C ( Fig. 15d) at high temperatures. The sliding wear resistance decreases according to the progression ''WC-17Co (best) ? Cr3C2-25NiCr ? WC-12Co ? WC-10Co-4Cr (worst)'' at * 500°C.
In terms of coating processes, cracking and delamination has been observed for APS coatings at * 250°C. This behaviour was not observed for the HVOF, HVAF and LPPS coatings due to their improved microstructure until a higher temperature of * 650°C. The corrosion resistance, however, shows an opposite trend to the high-temperature sliding wear rate, where the corrosion resistance increases as WC-17Co (worst) ? WC-12Co ? WC-10Co-4Cr ? WC-Ni (best), Fig. 15(a). 6. The numerical modelling of sliding wear rate using FEM can be implemented using strategies of mesh deletion and node adjustment, which have been applied to other materials. These models are in their infancy and are based on experimentally evaluated wear rate (k D ). They have the potential to superimpose the residual stress profile and microstructural phases during sliding wear. Some aspects of sliding wear at higher temperatures, such as the variation in the temperature-dependent material properties, can also be captured by the FEM analysis. However, challenges remain when modelling the stability of the tribo-film at higher temperatures and loads.

Future Projections
Several developments can further improve the microstructural quality and sliding wear of WC coatings. These developments include the following five areas of science and technology.
(1) Closed-loop monitoring and feedback control of the thermal spray process can link the coating process parameters with particle temperatures, their velocity distributions and the generation of thermal residual stresses. This knowledge will lead to optimisation of coating powders and process parameters to reduce carbon loss, especially in the emerging technological area of nanocomposite coatings. (2) New material developments are needed to design coatings that can perform as well as WC-Co coatings in sliding wear and, in addition, include corrosion and high-temperature tribological environments. (3) Numerical modelling of the sliding wear process is in its infancy. Multiscale modelling in this area, along with the optimisation of these models to reduce computational cost and improve the accuracy of results, will enable the design of improved wearresistant coatings for existing and new markets. (4) A better understanding of residual stress is needed to understand its role during the service loading of WC-Co coatings in sliding contacts. The influence of residual stress also impacts heavily on fatiguerelated failure mechanisms of components. The intention is to design a through-thickness residual stress profile in the coating substrate system by controlling the coating deposition process and parameters and post-spray heat treatment of coatings, for improved wear resistance. (5) Further development of thermal spray processes such as 'warm spray' of WC-Co coatings can be coupled with an in situ heat treatment to improve the tribomechanical properties of WC-Co coatings.
Acknowledgements CCB has the support of the Australian Research Council under the Industrial Transformation Training Centre Project IC180100005 that is titled ''Surface Engineering for Advanced Materials'', SEAM.
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