Wear Behavior Analysis of Al 2 O 3 Coatings Manufactured by APS and HVOF Spraying Processes Using Powder and Suspension Feedstocks

: Thermally sprayed ceramic coatings are applied for the protection of surfaces that are exposed mainly to wear, high temperatures, and corrosion. In recent years, great interest has been garnered by spray processes with submicrometric and nanometric feedstock materials, due to the reﬁnement of the structure and improved coating properties. This paper compares the microstructure and tribological properties of alumina coatings sprayed using conventional atmospheric plasma spraying (APS), and various methods that use ﬁnely grained suspension feedstocks, namely, suspension plasma spraying (SPS) and suspension high-velocity oxy-fuel spraying (S-HVOF). Furthermore, the suspension plasma-sprayed Al 2 O 3 coatings have been deposited with radial (SPS) and axial (A-SPS) feedstock injection. The results showed that all suspension-based coatings demonstrated much better wear resistance than the powder-sprayed ones. S-HVOF and axial suspension plasma spraying (A-SPS) allowed for the deposition of the most dense and homogeneous coatings. Dense-structured coatings with low porosity (4 vol.%) and good cohesion to the metallic substrate, containing a high content of α –Al 2 O 3 phase (56 vol.%) and a very low wear rate (0.2 ± 0.04 mm 3 × 10 − 6 /(N · m)), were produced with the S-HVOF method. The wear mechanism of ceramic coatings included the adhesive wear mode supported by the fatigue-induced material delamination. Moreover, the presence of wear debris and triboﬁlm was conﬁrmed. Finally, the coefﬁcient of friction for the coatings was in the range between 0.44 and 0.68, with the highest values being recorded for APS sprayed coatings.


Introduction
The development of highly advanced spray torches and a better understanding of coating deposition mechanisms have contributed to rapid progress in surface engineering [1]. Requirements in many fields of machine operation and production techniques led to applications for various alumina-based coatings, for example, as electrically insulating or wear-protection layers [2]. Such coatings are primarily found in the machinery, marine, chemical, food, textile, printing, and energy industries [3][4][5].
One of the primary methods of ceramic coating manufacturing is atmospheric plasma spraying (APS) [6] and high-velocity oxy-fuel spraying (HVOF) [7]. Plasma spraying involves the use of a high-temperature plasma jet to produce wear-and erosion-resistant ceramic [8,9], composite [10,11], or metallic coatings [12]. The APS process uses a powder Table 1. Particle size distribution of Al 2 O 3 powders used for deposition of the coatings; d v -particle size by volume [µm]; APS-atmospheric plasma spraying, SPS-suspension plasma spraying with radial injection, A-SPS-suspension plasma spraying with axial injection. The distribution of powder particles sizes was confirmed by FEG-SEM investigation (Leo 1530, Zeiss, Zürich, Switzerland) and SEM TESCAN VEGA 3 SBH (Tescan, Brno, Czech Republic), as shown in Figure 1.

Spraying Processes
To avoid clogging the injector nozzle during APS spraying, the micrometer-sized powder was previously dried for 2 h at 120 • C. The suspensions were redispersed by continuous mechanical stirring. The stainless-steel coupons, of 3 mm thickness and a diameter equal to 25 mm, were grit-blasted and cleaned with ethanol prior to spraying.

Spraying Processes
To avoid clogging the injector nozzle during APS spraying, the micrometer-sized powder was previously dried for 2 h at 120 °C. The suspensions were redispersed by continuous mechanical stirring. The stainless-steel coupons, of 3 mm thickness and a diameter equal to 25 mm, were grit-blasted and cleaned with ethanol prior to spraying.
The main spraying parameters are shown in Table 2. APS spraying, performed at Wroclaw University, was carried out with one anode and one cathode SG-100 torch (Praxair, Indianapolis, IN, USA). SPS spraying with radial injection was performed at Fraunhofer IWS, using a cascaded plasma torch KK (AMT AG, Kleindöttingen, Switzerland), while A-SPS spraying using a Mettech Axial III (Northwest Mettech Corp., Vancouver, Canada) was performed at University West. In both SPS processes, the standoff distances were similar to the powder-based APS process. S-HVOF spraying was performed at Fraunhofer IWS using an HVOF Top Gun (GTV Verschleißschutz GmbH, Luckenbach, Germany) with a modified combustion chamber that allowed internal and axial feeding of the suspension, as previously described by Toma et al. [22,25,42].

Coating Characterization
Microstructural characterization of the coatings' surfaces and cross-sections was carried out using the scanning electron microscopes PHENOM G2 PRO (Phenom World BV, Eindhoven, The Netherlands) and TESCAN VEGA 3 SBH (Tescan, Brno, Czech Republic). Prior to cross-sectional investigations, the metallographic preparation included precise cutting, cold mounting in low viscosity epoxy resin, grinding, end polishing, and The main spraying parameters are shown in Table 2. APS spraying, performed at Wroclaw University, was carried out with one anode and one cathode SG-100 torch (Praxair, Indianapolis, IN, USA). SPS spraying with radial injection was performed at Fraunhofer IWS, using a cascaded plasma torch KK (AMT AG, Kleindöttingen, Switzerland), while A-SPS spraying using a Mettech Axial III (Northwest Mettech Corp., Vancouver, Canada) was performed at University West. In both SPS processes, the stand-off distances were similar to the powder-based APS process. S-HVOF spraying was performed at Fraunhofer IWS using an HVOF Top Gun (GTV Verschleißschutz GmbH, Luckenbach, Germany) with a modified combustion chamber that allowed internal and axial feeding of the suspension, as previously described by Toma et al. [22,25,42].

Coating Characterization
Microstructural characterization of the coatings' surfaces and cross-sections was carried out using the scanning electron microscopes PHENOM G2 PRO (Phenom World BV, Eindhoven, The Netherlands) and TESCAN VEGA 3 SBH (Tescan, Brno, Czech Republic). Prior to cross-sectional investigations, the metallographic preparation included precise cutting, cold mounting in low viscosity epoxy resin, grinding, end polishing, and gold sputtering. To determine the porosity of the coatings, ImageJ 1.50i software was used, and the standard procedure described by standard ASTM E2109-01(2014) was performed [43,44]. For this purpose, the cross-section micrographs were taken in at least 20 random areas.
The crystalline phases in the feedstocks and as-sprayed coatings were investigated by X-ray diffraction (XRD) using a D8 Discover diffractometer (Bruker AXS, Karlsruhe, Germany) with CuK α radiation. The range of 2θ acquisition was 10 • to 80 • . The crystalline phases were identified using the JCPDS standard cards: 00-046-1212 (α-Al 2 O 3 ) and 00-010-0425 (γ-Al 2 O 3 ). The volume percentage of the α-Al 2 O 3 and γ-Al 2 O 3 phases was determined using Equation (1) (following [42]): where C is the content of γ-Al 2 O 3 phase content, and I γ , I α is the intensity of the peak diffraction for the corresponding plane of γ-Al 2 O 3 and α-Al 2 O 3 , accordingly.
The microhardness was measured by Vickers indentations with 0.2 N loading, using a Sinowon HV-1000 apparatus (Sinowon Innovation Metrology, DongGuan, China). For each coating sample, 10 indentations were made.
The intensity of wear behavior and the friction coefficients of the coatings were investigated using a ball-on-disk (BOD) tribometer CSM (CSM Instruments SA, Peseux, Switzerland) in dry sliding contact conditions, at room temperature. The WC counterpart ball diameter was 6 mm with a hardness of 16 GPa, which exceeds the approximate 15 GPa of the hardness reported for alumina [45]. The test conditions were: 10 N normal load, sliding speed of 0.1 m/s, 1000 m sliding distance, and 3 mm radius of the wear track. The wear rate was calculated according to the formula given elsewhere [46,47], taking into account the material volume loss, load, and sliding distance. Prior to testing, the coatings' surfaces were ground and polished to assure the surface roughness (Ra) was below 0.2 µm, as specified by the ASTM G99 standard [48]. The morphology of the wear track was examined using a scanning electron microscope PHENOM G2 PRO (Phenom World BV, Eindhoven, The Netherlands).

Morphology and Microstructure
Both as-sprayed conventional APS and suspension-sprayed coatings exhibited a characteristic topography that is typical for thermally sprayed deposits ( Figure 2). APS coating ( Figure 2a) had some microcracks and partially melted powder particles but, generally, was built from micrometer-sized splats. Cracks emerged in the coatings due to the high energy input during spraying. Unmelted powder particles remaining in the structure were due to the residence of Al 2 O 3 particles in the plasma within less than 1/1000 s [49], or from particles traveling in the external and colder plasma jet region. It was already reported [50] that in the case of agglomerated powders, such a period is not sufficient to completely melt the agglomerated powder sized 20-50 µm. It results in lower thermal conductivity when compared to the fused powder. Additionally, the radial injection meant that some powders traveled in hotter and some in colder plasma jet regions, so this causes some differences in the melting of the feedstock during spraying. In the case of suspension-based processes, the melting process of fine alumina powders is different. During injection into the plasma, they melt fully, so, basically, the structure consists of well-melted and fine splats. Some of the particles, which stay in the colder plasma regions, transform into a liquid state, but only partially. Therefore, some partially melted or re-solidified powder particles are also observed in suspension-sprayed coatings (SPS, A-SPS, and S-HVOF). However, for the same suspension feedstock, the effect was more visible when radial injection was used (SPS, Figure 2b) than in the S-HVOF process ( Figure 2d). The presence of very fine particles, observed in A-SPS coatings (Figure 2c), was mainly because of the wider powder particle size distribution in the suspension, where the starting powder contained 10 vol.% of particles smaller than 0.5 µm but also because of the greater stand-off distance, which could lead to the re-solidification of the finest powder particles. where the starting powder contained 10 vol.% of particles smaller than 0.5 µm but also because of the greater stand-off distance, which could lead to the re-solidification of the finest powder particles.  Figure 3 presents cross-section SEM images of the deposited coatings. As already discussed, the splat size in conventional APS coatings was larger than in fine-grained, suspension-based ones. It was observed that the axially sprayed coatings A-SPS and S-HVOF (see Figure 3c,d) were characterized by a dense, uniform microstructure with the lowest porosity of 6 vol.% and 4 vol.%, respectively. Pores were observed mainly in the APS and SPS deposits. The main difference was in the size of the pores-although the coatings were of similar porosity (12 vol.%-APS coating; 15 vol.%-SPS coating), the pores in APS coatings were of micrometer size, while SPS coating had refined porosity.   Figure 3 presents cross-section SEM images of the deposited coatings. As already discussed, the splat size in conventional APS coatings was larger than in fine-grained, suspension-based ones. It was observed that the axially sprayed coatings A-SPS and S-HVOF (see Figure 3c,d) were characterized by a dense, uniform microstructure with the lowest porosity of 6 vol.% and 4 vol.%, respectively. Pores were observed mainly in the APS and SPS deposits. The main difference was in the size of the pores-although the coatings were of similar porosity (12 vol.%-APS coating; 15 vol.%-SPS coating), the pores in APS coatings were of micrometer size, while SPS coating had refined porosity.
where the starting powder contained 10 vol.% of particles smaller than 0.5 µm but als because of the greater stand-off distance, which could lead to the re-solidification of th finest powder particles.  Figure 3 presents cross-section SEM images of the deposited coatings. As alread discussed, the splat size in conventional APS coatings was larger than in fine-grained suspension-based ones. It was observed that the axially sprayed coatings A-SPS and S HVOF (see Figure 3c,d) were characterized by a dense, uniform microstructure with th lowest porosity of 6 vol.% and 4 vol.%, respectively. Pores were observed mainly in th APS and SPS deposits. The main difference was in the size of the pores-although th coatings were of similar porosity (12 vol.%-APS coating; 15 vol.%-SPS coating), th pores in APS coatings were of micrometer size, while SPS coating had refined porosity.

Phase Composition of Alumina Coatings
The phase analysis of all feedstock powders confirmed that in all cases, 100 vol.% rhombohedral α-Al 2 O 3 phase was found (see example in Figure 4a). During the coating deposition, the γ-Al 2 O 3 crystalline phase has also been formed; the ratios of α and γ alumina were different, depending on the spray method (Table 3). Such differences resulted from heat input and the mechanism of coating deposition, as well as particle heat capacity. An example XRD diffractogram obtained for S-HVOF coating is given in Figure 4b.

Phase Composition of Alumina Coatings
The phase analysis of all feedstock powders confirmed that in all cases, 100 v rhombohedral α-Al2O3 phase was found (see example in Figure 4a). During the co deposition, the γ-Al2O3 crystalline phase has also been formed; the ratios of α a alumina were different, depending on the spray method (Table 3). Such differe resulted from heat input and the mechanism of coating deposition, as well as particle capacity. An example XRD diffractogram obtained for S-HVOF coating is given in Fi 4b. In thermal spray coatings, due to the rapid heating and cooling of powder part phase transformation occurs from α to γ-Al2O3. This transformation is a result of the temperature and the preferential nucleation of γ-Al2O3 at high cooling rates [51]. For coating, the quantitative results (see Table 3) are similar to the ones in the literature [25 On the other hand, for SPS coating, the content of γ-Al2O3 was the lowest. This probably connected with a relatively high fraction of not fully melted particles, which caused by radial injection. In such an injection mode, there is a problem with achie good penetration of the plasma plume by submicrometer-powder particles. Sim  In thermal spray coatings, due to the rapid heating and cooling of powder particles, phase transformation occurs from α to γ-Al 2 O 3 . This transformation is a result of the high temperature and the preferential nucleation of γ-Al 2 O 3 at high cooling rates [51]. For APS coating, the quantitative results (see Table 3) are similar to the ones in the literature [25,52]. On the other hand, for SPS coating, the content of γ-Al 2 O 3 was the lowest. This was probably connected with a relatively high fraction of not fully melted particles, which was caused by radial injection. In such an injection mode, there is a problem with achieving good penetration of the plasma plume by submicrometer-powder particles. Similar results could be found, for instance, in [44]. This could be explained by the two-zone microstructure of SPS coatings [53]. A higher content of α-Al 2 O 3 phase in SPS coating, in comparison to the APS one, resulted also from the fact that APS particles are in direct contact with the plasma, whereas in SPS some energy is consumed by the evaporation of the liquid phase [13]. In the case of A-SPS coating, the injection mode was axial. This change resulted in better penetration of the plasma flame, reaching its hot core. Consequently, the powder particles have been exposed to a higher heat load, which resulted in a higher content of α-Al 2 O 3 phase in the coating. Nevertheless, in this case, some energy was also consumed by the evaporation of the liquid phase from suspension. For S-HVOF coating, the heat treatment by flame is much lower than in the case of plasma. On the other hand, the axial injection mode promoted better particle heat treatment.

Microhardness
The microhardness of the coatings was dependent on the phase composition, the porosity, the presence of not fully melted particles, and cohesion ( Figure 5). In general, a coating that contains non-melted powders may exhibit either lower microhardness (because of poor interlamellar bonding), or higher microhardness, resulting from the retained α phase [49]. APS coatings exhibit relatively high values of microhardness (see Figure 5). This could be related to a rather dense structure and good interlamellar cohesion. Similar results could be found, e.g., in [54]. For SPS coating, the non-homogeneous microstructure plays an important role in the microhardness value. The porosity is the highest among all tested samples (15 vol.%). There is a high number of non-melted and re-solidified particles, which weaken the coating. However, taking into account such defects, the microhardness value is still relatively high. Similar results are mentioned by [55]. In the case of A-SPS coating, in this structure, many re-solidified particles could be observed. Moreover, the much smaller size of splats and pores means that the indent spans several splats. This phenomenon [56] could decrease the real microhardness value (see Figure 5), but it is still in accordance with the literature [32]. The highest values of microhardness were obtained for S-HVOF coating. This could be explained by coating giving the most compact microstructure, the lowest porosity level (only 4 vol.%), and very strong interlamellar cohesion. Moreover, small pores are homogeneously distributed within the coatings, which do not greatly affect the microhardness, as is in line with other reported findings [57].

Wear Behavior
The sliding wear results given in Figure 6 present the wear rate and coefficient of friction (COF) of the suspension-sprayed and conventionally deposited alumina coatings. Comparative analysis indicates the superior wear resistance of S-HVOF coating, which presents a negligible wear rate, equal to 0.2 ± 0.04 mm 3 × 10 −6 /(N•m). This was one order of magnitude lower than A-SPS coating, and about two orders of magnitude lower than the wear rate of APS coating. It is known that hardness is a major factor influencing the wear resistance of solid materials [58], therefore the superior anti-wear properties of S-HVOF coatings derive mainly from the highest hardness of S-HVOF-deposited alumina (average value of 1127 HV0.2 (see Figure 5)), which exceeds the hardness of other samples. Moreover, in the case of thermally deposited ceramic materials, their characteristic microstructural features, likewise their porosity, lamellar microstructure, interlamellar bonding, or various discontinuities, affect the wear behavior. In the case of A-SPS coating, the axial injection of the suspension resulted in a compact microstructure with high wear resistance, despite a lower hardness value.
The dominant wear mechanism of APS coating was fatigue-induced delamination of the splat material and subsequent adhesive smearing debris. The material loss was caused by delamination, initiated at fatigue cracks, pores, and the splat border. The coefficients of friction (COF) of the coatings are comparable to the values reported in the literature [32,54]. In addition, wear investigation confirms the higher COF of APS and A-SPS

Wear Behavior
The sliding wear results given in Figure 6 present the wear rate and coefficient of friction (COF) of the suspension-sprayed and conventionally deposited alumina coatings. Comparative analysis indicates the superior wear resistance of S-HVOF coating, which presents a negligible wear rate, equal to 0.2 ± 0.04 mm 3 × 10 −6 /(N·m). This was one order of magnitude lower than A-SPS coating, and about two orders of magnitude lower than the wear rate of APS coating.

Wear Behavior
The sliding wear results given in Figure 6 present the wear rate and coefficie friction (COF) of the suspension-sprayed and conventionally deposited alumina coat Comparative analysis indicates the superior wear resistance of S-HVOF coating, w presents a negligible wear rate, equal to 0.2 ± 0.04 mm 3 × 10 −6 /(N•m). This was one o of magnitude lower than A-SPS coating, and about two orders of magnitude lower the wear rate of APS coating. It is known that hardness is a major factor influencing the wear resistance of materials [58], therefore the superior anti-wear properties of S-HVOF coatings d mainly from the highest hardness of S-HVOF-deposited alumina (average value of HV0.2 (see Figure 5)), which exceeds the hardness of other samples. Moreover, in the of thermally deposited ceramic materials, their characteristic microstructural feat likewise their porosity, lamellar microstructure, interlamellar bonding, or va discontinuities, affect the wear behavior. In the case of A-SPS coating, the axial inje of the suspension resulted in a compact microstructure with high wear resistance, de a lower hardness value.
The dominant wear mechanism of APS coating was fatigue-induced delaminati the splat material and subsequent adhesive smearing debris. The material loss was ca by delamination, initiated at fatigue cracks, pores, and the splat border. The coeffic of friction (COF) of the coatings are comparable to the values reported in the liter [32,54]. In addition, wear investigation confirms the higher COF of APS and A It is known that hardness is a major factor influencing the wear resistance of solid materials [58], therefore the superior anti-wear properties of S-HVOF coatings derive mainly from the highest hardness of S-HVOF-deposited alumina (average value of 1127 HV0.2 (see Figure 5)), which exceeds the hardness of other samples. Moreover, in the case of thermally deposited ceramic materials, their characteristic microstructural features, likewise their porosity, lamellar microstructure, interlamellar bonding, or various discontinuities, affect the wear behavior. In the case of A-SPS coating, the axial injection of the suspension resulted in a compact microstructure with high wear resistance, despite a lower hardness value.
The dominant wear mechanism of APS coating was fatigue-induced delamination of the splat material and subsequent adhesive smearing debris. The material loss was caused by delamination, initiated at fatigue cracks, pores, and the splat border. The coefficients of friction (COF) of the coatings are comparable to the values reported in the literature [32,54]. In addition, wear investigation confirms the higher COF of APS and A-SPS coatings than in those deposited by the SPS and S-HVOF processes. The α-Al 2 O 3 is characterized by higher toughness than that reported for the γ-Al 2 O 3 phase [59], therefore the S-HVOF coating, which is α-Al 2 O 3 -enriched, presents superior hardness; this explains the beneficial frictional behavior. Undoubtedly, the microstructure, cohesion, splat size, and occurrence of defects such as cracks and pores affect the wear mechanisms of ceramic materials [60]. Therefore, the lowest average COF of SPS coating can be explained by the reduction of friction, due to the smearing action of the fine and loosely worn debris. Hence, even SPS coating has a relatively high mass loss; however, it exhibits a low coefficient of friction.
Although all ceramic coatings presented the adhesive wear mode supported by the fatigue-induced material delamination, their overall wear mechanism was different, see Figure 7. In all coating wear tracks, the presence of wear debris and tribofilm was identified. Similar findings were already mentioned in our previous study [39], in which both the wear of the WC counterpart and debris material transfer of the alumina coatings/counterpart ball were identified. This was also proved in this work by the SEM-EDS investigation of the tribofilm. Transfer of the counter ball material into the wear debris was confirmed by the presence of tungsten in the wear track (25.3 ± 2.8 wt.% W). Although, the identified tungsten content in tribofilm was at a comparable level, with a relatively low standard deviation (the mean tribofilm's chemical composition in wt.% is: O = 35.8 ± 2.5; Al = 27.7 ± 2.2; W = 25.3 ± 2.8; and N = 11.2 ± 0.7). Coating wear behavior was affected by the deposition technique and process parameters themselves. APS coating had severe adhesive wear, followed by grooving wear and delamination of the material, which relies on the brittle mode (the presence of cracks). APS coating's microstructural non-uniformities, such as initial cracks, the lamellar microstructure, and porosity, facilitating material loss were confirmed in the previous work [39]. On the other hand, SPS coating presented the adhesive transfer of coating material, strongly influenced by poor coating cohesion (see Figure 7c). SPS deposited coatings had better wear performance than APS ones, because of the transfer of fine wear debris along the wear track. It mitigated the COF and wear rate when compared to APS coatings. Furthermore, the A-SPS coating wear mechanism presented in Figure 7e,f had the adhesive behavior, the tribofilm was transferred through the wear track, and cyclic fatigue resulted in coating material delamination. Wear debris were uniformly smashed at the edges of the wear track where underwent moderate grooving wear. This dense tribofilm was beneficial until it was removed due to fatigue-induced cracking. In the case of S-HVOF coating (Figure 7g,h), the wear was greatly reduced when compared to other investigated coatings. It derives from the superior properties of S-HVOF coating, likewise the high adhesion [1], dense morphology and microstructure, lack of cracks, pores or other discontinuities, which support the superior antifatigue performance of S-HVOF alumina coating. After testing, the surface material loss was negligible and decreased the surface roughness in the wear track area. The obtained wear debris was preserved in the roughness pits and scratches. The adhesive smashing of debris, produced by the cutting of the surface heights and fatigue-delaminated coating material, was observed. The compact tribofilm was formed. Following the findings of Bolelli et al. [52], it can be claimed that the formation of a compact film visible at the edges of the wear track of A-SPS and S-HVOF supported the anti-wear alumina properties. Finally, it should be pointed out that the dense, tough, and uniform structure likewise manufactured by the A-SPS and S-HVOF processes, followed by the dense tribofilm formation, mitigated the adhesive wear mechanism of alumina coatings, and was beneficial for their tribological performance. The obtained findings confirmed that, among the four investigated alumina coatings, the S-HVOF sprayed ones were characterized by the most favorable wear behavior.
The results clearly showed that the wear behavior of alumina coatings is strongly determined by the size of the feedstock material and the spray process itself. The results clearly showed that the wear behavior of alumina coatings is strongly determined by the size of the feedstock material and the spray process itself.

Conclusions
In this work, the microstructure and tribological behavior of conventional (APS) and fine-grained suspension sprayed (SPS, A-SPS, and S-HVOF) Al 2 O 3 coatings have been investigated. The following conclusions can be drawn.
All as-sprayed alumina coatings were fairly uniformly built and showed a typical lamellar morphology, even for suspension-sprayed coatings obtained at rather long spray distances (90 or 100 mm). The typical microstructural features (i.e., the size of lamellas, the number of non-melted or resolidified powder particles embedded in the structure, the size and volume of pores and microcracks) were mainly influenced by the type of feedstock used for spraying (coarse vs. fine powders), the injection mode used (radial vs axial), and the spray process characteristic (plasma spraying vs. HVOF). The use of suspension feedstock, together with axial injection (S-HVOF and A-SPS), resulted in homogeneous, fine-grained, and dense coatings.
The phase composition was strongly affected by the deposition process (heat input for plasma and flame) as well as spray parameters. The axial injection mode promoted a higher amount of α-Al 2 O 3 phase.
In the case of microhardness values, the microstructure played an important role. The highest value was exhibited by S-HVOF coating (1127 HV0.2). This could be related to the low porosity, high interlamellar cohesion, and limited presence of non-melted and re-solidified particles in the coating.
The wear rate of finely grained coatings sprayed from suspensions (SPS, A-SPS, and S-HVOF) is lower than that reported for APS powder deposits. The results indicate the superior wear resistance of S-HVOF coating, which presents a negligible wear rate, equal to 0.2 ± 0.04 mm 3 × 10 −6 /(N·m). This wear rate was one order of magnitude lower than that for A-SPS coating, and almost two orders of magnitude lower than the wear rate for APS coating.
The wear mechanism of ceramic coatings includes the adhesive wear mode supported by the fatigue-induced material delamination. Moreover, the presence of wear debris and tribofilm was confirmed. S-HVOF and axial suspension plasma spraying (A-SPS) allow the deposition of dense and coherent coatings, which mitigates the adhesive action and reduces the coating material mass loss.
The lowest average COF values of 0.44 and 0.49 were obtained for SPS and S-HVOF sprayed coatings, respectively. In the case of SPS coating, this can be explained by the reduction of friction due to the smearing action of the fine and loosely worn debris, while for S-HVOF, this was an effect of the high coating hardness mitigating the adhesive wear mechanism.