Tribocorrosion behaviour of thermally sprayed cermet coatings in paper machine environment

Tribology International 142 (2020) 106006

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Tribocorrosion behaviour of thermally sprayed cermet coatings in paper machine environment E. Huttunen-Saarivirta *, V. Heino, E. Isotahdon, L. Kilpi, H. Ronkainen VTT Technical Research Centre of Finland Ltd, P.O. Box 1000, FI-02044, VTT, Finland

A R T I C L E I N F O

A B S T R A C T

Keywords: Tribocorrosion Thermally sprayed coating Cermet Abrasive wear

Tribocorrosion behaviour of thermally sprayed cermet coatings: WC-CoCr deposited by HVOF and HVAF and Cr3C2-WC-NiCrCo applied by HVAF, was examined in paper machine environment in a pin-on-disc tribometer under the load of 20 N in electrolyte containing chlorides and sulphates, pH 4.5. Wear was the dominant degradation mechanism for all coatings, followed by corrosion-induced wear, the importance of which increased with rising potential. The overall material losses were lowest for HVAF WC-CoCr and highest for HVAF Cr3C2WC-NiCrCo, the latter being related to the evident abrasive wear of the Cr3C2 phase and cracking along Cr3C2matrix interface in the near-surface areas. Among the three coatings, contribution by corrosion to the damage was greatest in HVOF WC-CoCr, likely due to high degree of porosity.

1. Introduction Composite coatings with ceramic particles in an alloy matrix (cer­ mets) are used in many sectors where the components are exposed to harsh mechanical and environmental conditions, such as in trans­ portation, energy, chemical, process, marine as well as pulp and paper industries [1–5]. Typically, the ceramic particles in the structure provide the high hardness for wear resistance and thermodynamic stability for corrosion and chemical resistance, whereas the alloy matrix holds the structure together and provides the toughness and ductility. Thus, the composite material eventually meets even partially conflicting property requirements, such as the hardness and ductility. WC-CoCr, also called hardmetal, has been the most commonly used type of cermet in industry for a long time [5], due to the extremely high hardness of the WC par­ ticles and the strong adhesion of WC particles to the Co-rich matrix. Another common coating system involves the combination of the Cr3C2 hard phase and NiCr as the matrix phase [6], although some instances have speculated that the substitution of Co-based matrix cermets will lead to compromises in the material performance [7]. One attempt to overcome such compromise may be the use of combination of Cr3C2 and WC hard phases in the alloy matrix, the composition of which has been matched accordingly, as in Cr3C2–37WC–18NiCoCr [8]. As the application of thermally sprayed hardmetal coatings is justi­ fied by the high wear and/or corrosion resistance in field, their behav­ iour and performance under a range of wear and corrosion conditions

has received plenty of research effort during the course of years [9–16]. With respect to wear behaviour, the typical observation has been that the harder and denser the coatings, the better their wear resistance, the latter being related to the improved support of the carbide grains by the matrix phase. Indeed, under the sliding type of contact, the matrix phase is worn preferentially, followed by a pull-out of the carbide particles [17,18]. Under the high load intensity, cracking may also be detected [17]. Corrosionwise the main degradation mechanism has been the galvanic coupling between the hard phase and the matrix; this may be slightly decreased by the use of (higher) Cr alloying in the matrix to enable passivation and decrease the driving force for microgalvanic corrosion between the carbide particles and the matrix [16].The topic what is not so commonly addressed by earlier studies of thermally sprayed cermets is the tribocorrosion performance, referring to how the coatings behave and succeed under the combined action of wear and corrosion. The few erosion-corrosion results which have been reported demonstrated that the synergy between wear and corrosion may account for the majority of attack in the material [3], and varying synergy results have been collected depending on the test conditions [3,19,20]. Wood [5] has suggested that repassivation capability of the matrix phase plays a key role with regards to wear-corrosion synergy, whereas recently also the role of other synergistic mechanisms, such as the cracking of carbide grains or other type of fracturing and its accelerating role in the corro­ sion process have been arisen [21,22]. Nevertheless, Pilleggi et al. [23] have concluded that more research is needed to understand the

* Corresponding author. E-mail address: [email protected] (E. Huttunen-Saarivirta). https://doi.org/10.1016/j.triboint.2019.106006 Received 14 August 2019; Received in revised form 3 October 2019; Accepted 7 October 2019 Available online 9 October 2019 0301-679X/© 2019 Elsevier Ltd. All rights reserved.

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materials were obtained as commercial powders, which were deposited by an industrial operator using state-of-the-art equipment and spray parameters optimized for each case on ground and grit-blasted disc substrates of the grade 316 stainless steel using two methods: HVOF and HVAF. The coatings included in this research are shown in Table 1. The diameter of the discs was 40 mm and the original thickness before grit blasting was 6 mm. The coatings with the target thickness of 400–500 μm were applied on the discs, after which the top surface of the coating was finished with grinding, in order to fit the specimens into the holder (target specimen thickness: 6 mm) and allow a designed sphereon-flat contact geometry in the tribocorrosion experiments. The thick­ ness of the coating was selected to be within this range in order to avoid any contribution by the substrate to the results. The final grinding of the coated disc specimens was done to 500 grit surface finish, giving the surface roughness of 0.02–0.03 μm (Ra). Prior to the tribocorrosion ex­ periments, the coated discs were treated in an ultrasonic bath (5 min in ethanol, 5 min acetone) and dried. The drying was conducted first by blowing warm air and then by storing the specimens in the desiccator overnight, in order to remove humidity that might have retained in the pores.

Table 1 Thermally sprayed coatings included in this work and the potential values selected (based on polarisation curves) for the potentiostatic experiments. Method

Composition

HVOF HVAF HVAF

WC-CoCr WC-CoCr Cr3C2-WCNiCrCo

Cathodic potential/mV vs. Ag/AgCl 250 200 200

Anodic potentials/mV vs. Ag/AgCl 100, 100, 300 100, 100, 300 100, 100, 300

tribocorrosion mechanism of thermally sprayed cermet coatings. Cermet coatings, such as WC-CoCr or Cr3C2-WC-NiCoCr, are mainly deposited by thermal spraying methods, particularly by high velocity oxygen fuel (HVOF) or high velocity air fuel (HVAF) [5,8–12,14–27]. In the HVOF coating process, oxygen and fuel are mixed and burnt in a combustion chamber at high flow rates in order to generate a high-speed gas jet, in which the powder particles are directed. The accelerated particles impact and adhere tightly on the substrate material [5]. HVAF method employs compressed air as a combustion gas instead of oxygen, giving rise to somewhat lower deposition temperatures than in HVOF (only slightly higher than the melting temperature of the metal, yet temperatures in HVOF may even exceed the metal boiling temperature), which together with relatively higher gas flow rates result in a dense coating structure. Therefore, HVAF coatings have been reported to be harder than the corresponding HVOF coatings and feature somewhat greater wear resistance [28]. However, it has been reported that the deposition efficiency of the HVAF method is lower than in HVOF pro­ cess, yet there has been a significant improvement in this respect during the last 15 years [29]. Paper machines involve numerous machine elements that are simultaneously subjected to wear and corrosion, such as rolls. These may eventually influence the paper quality if not minimized, and the friction is inherently linked with the energy consumption. The highly durable coatings are seen as one key mechanism to reduce friction in paper machines [30]; carbide coatings find use in, for example, paper machine calendar rolls [31]. The rolls in paper machines become in sliding contact with pulp and paper (depending on the location), and the pulp partly determines the prevailing chemical conditions. This paper provides insights into the tribocorrosion behaviour of three thermally sprayed cermet coating types: HVOF WC-CoCr, HVAF WC-CoCr and HVAF Cr3C2-WC-NiCrCo, under the conditions relevant for paper mills with aqueous environment containing chlorides and sulphates at the pH value of 4.5 as well as high contact pressures of solid surfaces moving relative to each other. The selection of the coating types for the research is justified by the industrial importance of WC-CoCr and aim to under­ stand the correlations between coating method, microstructure and tribocorrosion behaviour. Additionally, Cr3C2-WC-NiCrCo is seen as a possible alternative to WC-CoCr [8] and was deposited by HVAF method based on the results obtained for WC-CoCr (HVOF vs. HVAF) coatings. The selected coating types enable the comparison between the spraying methods (HVOF WC-CoCr vs. HVAF WC-CoCr) and cermet compositions (HVAF WC-CoCr vs. HVAF WC-CoCr). The important understanding on the interactions between the wear and corrosion processes for the coatings facilitates the development of thermally sprayed coatings with enhanced protection properties against wear, corrosion and their com­ binations and the selection of the most appropriate deposition method and coating composition for paper mill conditions.

2.2. Methods 2.2.1. Materials characterisation The powders were examined with respect to size distribution using a Malvern Instruments Mastersizer 3000 laser diffraction device in which the powders were inserted in a dry state. The powders and the deposited coatings were subjected to microstructural examination in cross-section using ground and polished metallographic specimens embedded in resin and a Zeiss ULTRAplus field emission (FE) SEM. The elemental com­ positions of phases were determined using a Noran energy-dispersive spectrometer (EDS) system attached to SEM. Image analysis was car­ ried out with ImageJ software using SEM images taken in back-scattered electron (BSE) mode (4 images, the magnification of 1000). For the coatings, investigations were also carried out using a Leica MEF 4 M optical microscope. Additionally, phase structure of the powders and the coatings was characterised by XRD measurements with a PANalytical X’Pert Powder device and CuKα radiation through 2θ range 30–120� . The deposited coatings were further characterised with respect to hardness and Plane Strain modulus. Hardness of the coatings was determined using a DuraScan20 hardness tester and a Vickers indenter (HV0.3), with the results being presented as the average values of ten measurements. Additionally, the coatings were subjected to micro­ indentation experiments, which yielded indentation hardness (HIT) and Plane Strain modulus (E*). The apparatus used for the microindentation studies was an Anton Paar Tritec Micro-Combi Tester. The indentations were performed using a Vickers indenter under load-controlled mode, with the loading up to 3 N at the loading rate of 1.5 N/min, holding at the peak load for 15 s and unloading again at the rate of 1.5 N/min, respectively. Here, the presented values are the average of nine measurements. 2.2.2. Tribocorrosion experiments Sliding type of tribological contacts are frequent in paper machines [30]. Therefore, tribocorrosion experiments were designed from this perspective and, following the conditions in the paper machines, with the possibility of conducting them in the presence of water and chem­ icals. Pin-on-disc (POD) facility that introduces the sliding type of con­ tact between the pin and disc specimens and enables the contact to occur in the electrolyte was chosen for the tests. An Anton Paar Tritec POD tribometer with an electrochemical cell was used in this research, with the coated disc specimens being fastened in the specimen holder at the bottom of the electrochemical cell, as described in [32]. The test set-up was designed so that the rotating shaft controlled the movement of the electrochemical cell, while counterbody was stationary, introducing a sliding type of contact situation. An inert alumina ball (Al2O3, grade 5,

2. Experimental 2.1. Materials In this research, the test materials of two commercial compositions are included: WC-10Co-4Cr and 45Cr3C2–37WC–18NiCrCo (C: 7.7–8.5; Cr 38.5–43.5; Ni 10.0–13.0; Co 2.9–4.1; Fe < 0.5; W balance; wt.%). The 2

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Ra ¼ 0.031 μm) with the diameter of 10 mm was used as the counterbody and brought into contact with the coated disc working electrode (WE, exposed surface area of 11.3 cm2) with the help of a counterbody holder. The cell as well as the disc and counterbody holders were made of polymer in order not to introduce unwanted electrochemical effects. The experiments were performed in three modes by controlling the oxidizing capacity of the test environment through potential and involving/­ removing the counterbody. Pure corrosion tests were done in the absence of the counterbody, pure wear tests were carried out under cathodic conditions (with no corrosion) in the presence of the counter­ body, whereas tribocorrosion tests were conducted under anodic con­ ditions (corrosion) in the presence of the counterbody. In figures, terms corrosion and tribocorrosion are used for the cases without and with the mechanically loaded alumina counterbody. The tests that were per­ formed in the absence of counterbody were otherwise performed in the same way as those with the counterbody, for example, using the same rotation of the electrochemical cell. This approach allows for the determination of the synergy between wear and corrosion processes under the selected potential levels. Polarisation curves were used to guide the potential selection. The selection of the potentials for poten­ tiostatic experiments was carried out so that it covers one cathodic po­ tential with respect to systems without and with the counterbody, one potential that is cathodic in the system without the counterbody but anodic in its presence, and two potentials anodic irrespective of the test mode. Thus, the system with the alumina counterbody was in the core of the potential level selection. In all cases, the electrochemical cell was added with 30 ml of elec­ trolyte that contained 300 mg/L chlorides, Cl , and 600 mg/L sulphates, SO24 , with the pH value of 4.5, thus in wear and tribocorrosion exper­ iments the contact of the disc and the counterbody happened in the corrosive solution. The electrolyte composition was chosen based on the understanding that paper machine environment features slightly acidic pH values and involves chlorides and sulphates in the approximate concentration ratio of 1:2 [33]. The electrolyte was prepared using ion-exchanged water and analytical grades of NaCl, Na2SO4 and H2SO4. All experiments also involved a Ag/AgCl reference electrode (RE, satu­ rated KCl) that was located in the vicinity to the disc-counterbody contact point and a platinum ring counter electrode (CE) positioned at the periphery of the cylindrical cell. The potential values in this paper are thus given against the saturated Ag/AgCl electrode. During the tests, the speed of the rotating shaft was adjusted to 200 rpm, with the linear sliding speed of 0.2 m s 1. In reality, the rotating speeds in paper ma­ chines are clearly higher, approaching as high as 20 m s 1 [30], but here the used rotational speed was chosen because it still enabled the use of electrolyte in the electrochemical cell (at higher rotation speeds, the electrolyte easily escapes the cell due to centrifugal forces). In wear and tribocorrosion experiments, the normal load of 20 N was applied to the counterbody, yielding the Herzian pressure of 1.5 GPa in the beginning of the test. The contact pressure was higher than in the application in order to produce measurable wear in the coatings and to facilitate their ranking with respect to tribocorrosion performance. The diameter of the wear track in wear and tribocorrosion tests was 20 mm and the sliding distance was 1440 m. Independent of the test mode, electrochemical measurements covered open circuit potential (OCP) and potentiodynamic polarisation measurements followed by potentiostatic experiments under selected constant-potential conditions. OCP measurements were performed for 60 min, after which potentiodynamic polarisation measurements were conducted at a scan rate of 0.5 mV s 1 from the cathodic value of 600 mVOCP up to 800 mVOCP. The results from potentiodynamic polarisation measurements enabled the selection of one cathodic potential level and three anodic potential values for potentiostatic tests, the duration of which was 2 h (7200 s). The potentiostatic test conditions for each of the three coatings are shown in Table 1. All electrochemical measurements were run with a Gamry Reference 600TM potentiostat/galvanostat/zero resistance ammeter. In wear and tribocorrosion experiments, friction

coefficient was measured by the tribometer. Each experiment (poten­ tiodynamic polarisation measurement in the absence and presence of alumina counterbody, potentiostatic tests at the four potentials in the absence and presence of alumina counterbody) was repeated once, allowing for the presentation of electrochemical data separately for each individual test (to demonstrate the consistency of the results), yet for friction coefficient records only one example curve is shown due to the complete overlapping of parallel datasets. All experiments were per­ formed in the laboratory with a controlled temperature and humidity (22 � 1 � C, 50 � 5% RH). After the potentiostatic experiments, the homogenised electrolyte was analysed by inductively coupled plasma optical emission spectros­ copy (ICP-OES) with respect to the total concentration (ions released by corrosion þ solid particles released by wear) of cobalt, chromium, iron and, in the case of HVAF Cr3C2-WC-NiCrCo coating, nickel, thus the main constituents in the coating and the substrate. The inaccuracy of the measurement instrument was �10% for all the analysed elements. The coated disc specimens from potentiostatic tests were examined with respect to morphology and the presence of corrosion products using FESEM and EDS. The specimens from potentiostatic wear and tribocorro­ sion experiments were subjected to profilometric analyses using a Mitutoyo Formtracer SV-C3100 2D profilometer (wear track, every 90� position per disc). The profilometric studies yielded the cross-sectional area of the wear track; the wear track volumes were then calculated by multiplying these by the wear track circumferential length. 2.2.3. Wear-corrosion synergy The total material losses that occur during the simultaneous influ­ ence of wear and corrosion (T) cover the share by pure mechanical wear (W), that by pure corrosion (C) and the synergy term (S) describing the mutual degradational reinforcement of the two processes [34–38]: T¼WþCþS

(1)

Synergy term contains two parts: the influence of mechanical wear on corrosion (ΔCW), called in this paper as wear-induced corrosion, and the effect of corrosion on mechanical wear (ΔWC), referred to as corrosion-induced wear, thus: S ¼ ΔCW þ ΔWC

(2)

Corrosion-induced wear ΔWC may be approached following: ΔWC ¼ WC – W

(3)

In Equation (3), W represents volume losses by pure wear (here: under the counterbody yet at the cathodic potential), whereas Wc de­ notes material volume losses by wear under the simultaneous influence of corrosion (under anodic conditions). In turn, wear-induced corrosion is: ΔCW ¼ CW – C

(4)

with C standing for pure corrosion and CW referring to volume losses by corrosion in the presence of counterbody. Furthermore, the contribution by corrosion, C, and the influence of mechanical wear on corrosion, ΔCW, may be approached through Faraday’s law [39]: C¼

ItM nFρ

(5)

where I refers to the current collected in constant-potential experiments (in A), t represents duration of the test (in s) and M stands for the average molar mass of the coating (for WC-CoCr coatings, M was 176.4 g/mol and for Cr3C2-WC-NiCrCo, M was 163.7 g/mol). The number of electrons involved in the corrosion reaction is indicated by n (here assumed to be 2, because Co2þ and Ni2þ are the most likely ionic species for the Co- and Ni-based matrix alloys, respectively, under the test conditions), Fara­ day’s constant by F (96 500 C/mol) and material density by ρ (based on 3

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Fig. 1. Characteristics of the starting powders. a) Particle size distributions. b, c) Cross-sectional SEM image of WC-CoCr powder. d, e) Cross-sectional SEM image of the Cr3C2-WC-NiCrCo powder. 4

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Fig. 2. Microstructure of the studied thermally sprayed coatings, examined by SEM (a, c, e) and OM (b, d, f). a, b) HVOF WC-CoCr. c, d) HVAF WC-CoCr. e, f) HVAF Cr3C2-WC-NiCrCo.

the densities of the constituents: WC 15.77 g/cm3, Cr3C2 6.74 g/cm3, Co 8.85 g/cm3, Cr 7.19 g/cm3 and Ni 8.91 g/cm3) gave 14.7 g/cm3 for a dense WC-CoCr coatings and 10.4 g/cm3 for a dense Cr3C2-WC-NiCrCo). It is further emphasized that I was obtained by multiplying the average current density (I in A/cm2) during the corrosion experiments by the area of the exposed specimen (11.3 cm2). ΔCW could be defined using the same equation (5) and the difference between IW-I (with IW referring to

average current measured under the contribution by wear, while I is the corresponding case for pure corrosion). In tribocorrosion experiments performed in a pin-on-disc tribometer, both pure wear (W) and wear-corrosion synergy (S) take place in the wear track, while pure corrosion occurs in areas outside of the wear track (in the wear track the contribution by pure corrosion is insignifi­ cant, as the surface is continuously subjected to counterbody contact). 5

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Fig. 3. XRD spectra for a) WC-CoCr powder. b) Cr3C2-WC-NiCrCo powder. c) HVOF WC-CoCr coating. d) HVAF WC-CoCr coating. e) HVAF Cr3C2-WC-NiCrCo coating. The following denotations are used: W– – WC, C– –Cr3C2, N ¼ 75Ni–25Cr, O– – Co, ¼Co3W3C, r ¼ Co6W6C, X ¼ W2C.

Therefore, T is composed of W and S, giving T ¼ W þ S. In this work, W was obtained as total material losses T at the lowest potential (C and S are negligible), based on the profilometric analyses (by multiplying the wear track profile by its circumferential length (2π 2d, where d ¼ 20 mm), whereas the material losses obtained in a similar way at the anodic potentials introduced T. Wear-induced corrosion ΔCW at each anodic potential was defined as described above based on Equation (5), allowing corrosion-induced wear ΔWC to determined accordingly following Equations (1) and (2), hence T ¼ W þ ΔCW þ ΔWC.

particles ranging in size from d10 ¼ 26.6 μm to d90 ¼ 50.7 μm. Here, the particle size distribution was closer to that earlier used for HVOF process [41], yet it is known that HVAF method is not as sensitive to the powder particle size distribution as HVOF. SEM studies of the powders revealed that the particles were actually particle agglomerations of near-spherical shape, Fig. 1b–e. Also the greater particle size of the Cr3C2-WC-NiCrCo powder as compared to the WC-CoCr powder was confirmed by SEM studies. Investigation of the internal microstructure of as-polished WC-CoCr powder cross sections in SEM back-scattered electron mode (Fig. 1b and c) disclosed the presence of two main phases: the WC particles, seen in light contrast, and the Co–Cr matrix phase, seen in grey contrast. The WC particles were block-shaped or angular, typically ranged in size from two hundreds nm up to 2 μm, and were relatively evenly distributed within the powder agglomerations. The black areas in SEM images correspond to pores (resin). In the case of Cr3C2-WC-NiCrCo powder, SEM investigations enabled the detection of three phases (Fig. 1d and e): WC particles, seen in light contrast, Cr3C2 phases seen in dark grey contrast (with two different shades) and the matrix Ni–Cr–Co phase, seen in the lightest grey contrast (also with slightly varying shade of grey, indicating some compositional scatter). The WC particles featured a slightly greater variation in size than in WC-CoCr coatings,

3. Results and discussion 3.1. Test materials The size distribution of the feed powders is shown in Fig. 1a. The average particle size of WC-CoCr powder particles was 19.1 μm, with 80% of the particles (by volume) falling in the size range from d10 ¼ 12.0 μm to d90 ¼ 29.5 μm. Such measured powder particle size range has earlier been reported both for powders intended for both HVOF and HVAF processes [40]. In turn, the average size of particles belonging to Cr3C2-WC-NiCrCo powder was 36.8 μm, with 80% of the 6

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Fig. 4. Phase content in the three coating systems, defined by image analysis.

ranging in length scale from dimensions below hundred nm up to 2–3 μm, yet the particles of hundred nm size range were clearly more abundant than those in μm range. These particles were typically concentrated in some parts of the agglomerated powders. The majority of the powder constituted of Cr3C2 particles, which varied in shape from block-shaped towards more rounded and contained some micro-cracks (Fig. 1e). The size of Cr3C2 particles ranged from a few hundred nm up to 2 μm in length scale, with the particles in μm length scale being more abundant than those in nm length scale. Overall, Cr3C2 particles were relatively larger than the WC particles. The matrix phase then held the carbide particles together. Microstructure of the three coatings included in the work, examined by optical microscopy and scanning electron microscopy, is shown in Fig. 2. The target thickness was 300–400 μm, but there was some scatter between the coatings and individual specimens due to the grinding to final thickness in order to fit the specimen holder. In the case of HVOF WC-CoCr coating, the thickness was ~300 μm, with evident porosity detected in the structure(Fig. 2a and b). SEM investigations (Fig. 2b) disclosed that the block-shaped WC particles had somewhat rounded edges towards the matrix phase, indicating that some dissolution may have happened during the spraying. However, WC particle size distri­ bution was wide thus it is expected that no preferential dissolution of the small particles had occurred. The typical inter-particle sections of the matrix phase were clearly thinner than the size of WC particles. HVAF WC-CoCr coatings featured the thickness of ~250–300 μm, with a clearly denser structure than in the case of corresponding HVOF coating (Fig. 2c and d). The WC particles still featured sharp edges, suggesting no dissolution into the matrix phase. Additionally, the inter-particle sections of the matrix phase were thinner than in the case of the HVOF coating and evidently contained small pores. Among the three coating systems of interest, HVAF Cr3C2-WC-NiCrCo was the thickest, ~400–450 μm (Fig. 2e). Here, the size of both carbide phases, Cr3C2 and WC, ranged from some hundreds of nm to several μm (Fig. 2f), being thus larger than in the case of WC-CoCr coatings and consistent with observations on the powder microstructures. Moreover, the sections of the matrix phase between the hard particles (Cr3C2, WC) were somewhat thicker than in the case of WC-CoCr coatings, which is in agreement with the compositional data given in section 2.1. Here, porosity was also present, seemingly at greater amounts than in the case of HVAF WCCoCr coating.

XRD examinations revealed that the WC-CoCr feed powder con­ tained WC as the major phase and the following minor phases: Co, Co3W3C and Co6W6C (Fig. 3a). Such powder composition is in good agreement with that reported by Picas et al. [16], who suggested that CoxWxC phases developed during powder manufacturing (sintering process). In the case of Cr3C2-WC-NiCrCo powder, the major peaks in the XRD spectrum were related to WC phase, whereas also the peaks related to Cr3C2 and 75Ni–25Cr phases were frequent (Fig. 3b); this is the typical phase structure for such powder composition, reported earlier in refs. [8,41] despite minor changes in the matrix alloy composition (also contained Co in addition to Ni and Cr). XRD spectra for HVOF WC-CoCr coating revealed WC as the major phase and Co as the minor phase (Fig. 3c), with tiny peaks suggesting also the traces of W2C. In the case of HVAF WC-CoCr, only the peaks related to WC and Co were detected (Fig. 3d). Indeed, decarburization of WC particles during HVOF process typically leads the evolution of W2C phase [25]. For HVAF Cr3C2-WC-NiCrCo coating, XRD analyses revealed the presence of WC, Cr3C2, Co and 75Ni–25Cr phases (Fig. 3e), thus in addition to the phases present in WC-CoCr, Cr3C2 and 75Ni–25Cr phases were identified. The results from the image analyses conducted for the coating crosssections are presented in Fig. 4. The porosity was the highest in HVOF WC-CoCr coating, corresponding to 6 area-% of the structure, which is in line the observations by Cho et al. [42] and likely related to WC decomposition. Nevertheless, keeping in mind the challenges related to the determination of porosity in thermally sprayed coatings [6], the ranking of materials with respect to porosity rather than presenting absolute porosity values (which do not necessarily match with vol-%) is preferred here. The second highest porosity level was defined for HVAF Cr3C2-WC-NiCrCo coating, with the lowest porosity being detected in the case of HVAF WC-CoCr coating. Otherwise, WC phase covered ma­ jority of the coating structure in both WC-CoCr coatings, whereas the matrix phase dominated in the HVAF Cr3C2-WC-NiCrCo coating, fol­ lowed by the Cr3C2 phase. In HVAF WC-CoCr, the content of the matrix phase was lowest among the studied coatings. The key mechanical properties of the coatings are shown in Table 2. Among the three cermet coatings, for HVOF WC-CoCr, hardness values from Vickers hardness and microindentation measurements gave slightly conflicting results, with the second highest HV0.3 value (~1400) and clearly the lowest HIT value (~9.5 GPa). Also the Plane Strain Modulus from microindentation experiments was lowest among the 7

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of the material, whereas indentation tests disclose both elastic and plastic phenomena [43].

Table 2 Results from hardness measurements and microindentation experiments, given as average values and the standard deviation. HIT is the indentation hardness, whereas E* is plane strain modulus. HV0.3 HIT, MPa E*, GPa

HVOF WC-CoCr

HVAF WC-CoCr

HVAF Cr3C2-WC-NiCrCo

1412 � 97 9443 � 2756 238 � 34

1486 � 44 16977 � 1381 364 � 14

1238 � 69 13778 � 1223 276 � 13

3.2. Tribocorrosion behaviour Polarisation curves for the three coating systems revealed differences between the coatings and the cases of corrosion and tribocorrosion. None of the coatings exhibited true passive behaviour. Instead, the anodic section of the curves was typically divided into three distinct areas: a potential range where current density values showed a sharp increase with increase in potential just above the corrosion potential, then a potential range where the current density values increased only slightly with increase in potential and finally a steep increase in current density values with increase in potential at high oxidation capacities. Earlier, Picas et al. [14] have related the first of the three potential re­ gions to the oxidation of the CoCr metallic matrix phase at the matrix/carbide interface, the intermediate potential range to the pseu­ dopassive behaviour and the last high potential area to the oxidation and removal of WC particles. For HVOF WC-Co-Cr coatings, the curves that were recorded under pure corrosion and tribocorrosion conditions were closest to each other among the three coating systems, with a minor shift by tribological contact towards higher current density values (Fig. 5a). Such results indicated anodic dissolution at the rate proportional to the detected current density values, which were at the highest level among

three coatings (~240 GPa). It is worth mentioning that for HVOF WCCoCr coating, the scatter in all the values was relatively much higher than for the other two coating systems, referring to a more heteroge­ neous coating microstructure. One explanation for this observation may be the highest porosity level detected among the three coating systems. For HVAF WC-CoCr, both hardness values (HV0.3–1500 and HIT 17 GPa) and the Plane Strain modulus (360 GPa) were systematically highest among the three coating systems. The HVAF Cr3C2-WC-NiCrCo coating was characterized by the lowest HV0.3 hardness value (~1240) among the three coating systems, yet microindentation gave intermediate re­ sults between HVAF WC-CoCr (the highest values) and HVOF WC-CoCr (the lowest values), with the values of 13.8 GPa and 280 GPa, respec­ tively. In general, the main differences between the two hardness measuring methods are the phenomena they reveal: conventional hardness measurements only take into account the plastic deformation

Fig. 5. Polarisation curves for the studied thermally sprayed coatings in the absence (corrosion) and presence (tribocorrosion) of the alumina counterbody. a) HVOF WC-Co-Cr. b) HVAF WC-CoCr. c) HVAF Cr3C2-WC-NiCrCo. Corrosion refers to experiments carried out without the alumina counterbody, whereas tribocorrosion refers to experiments involving the alumina counterbody. The curves presented by the same colour (continuous, dotted) are from the parallel repetitions of the test. 8

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Fig. 6. Current density records for the HVOF WC-CoCr coatings at the selected four potentials. a) 250 mV vs. Ag/AgCl. b) 100 mV vs. Ag/AgCl. c) 100 mV vs. Ag/ AgCl. d) 300 mV vs. Ag/AgCl. Corrosion refers to experiments carried out without the alumina counterbody, whereas tribocorrosion refers to experiments involving the alumina counterbody. Curves from the parallel repetitions of the tests are shown.

the three coating systems for the HVOF WC-Co-Cr coatings but did not vary significantly between the corrosion and tribocorrosion cases. For HVAF WC-Co-Cr (Fig. 5b), the curves followed similar shape than in the case of HVOF WC-Co-Cr, also suggesting active dissolution. However, here, the curves collected in the presence of the mechanically loaded counterbody (tribocorrosion) were shifted to slightly higher current density and lower potential values as compared to the pure corrosion cases, and the overall current density levels were 0.5–1 magnitude lower than for the corresponding HVOF coating, likely due to the lower porosity linked with the spraying method. For HVAF Cr3C2-WC-NiCrCo (Fig. 5c), the difference between corrosion and tribocorrosion cases was more pronounced than for the other two coating systems: the current density values of the coating increased by almost one order of magnitude upon using the mechanically loaded counterpart, simultaneous with a clear (>100 mV) potential shift towards active direction. Additionally, the pseudopassive potential range under tribocorrosion conditions was very broad, with no distinct third potential range, in contrast to three-stage Tafel behavior in the case of corrosion. For this coating system, the overall current density values were the lowest among the three coatings included in the research, implying lowest dissolution rate. Altogether, the curves enabled the selection of the most interesting potential areas for more detailed potentiostatic measurements: one cathodic potential in each case to represent a pure wear situation and three anodic cases to evaluate the contribution by wear-corrosion syn­ ergy. For HVOF WC-CoCr, the selected cathodic potential was 250 mV vs. Ag/AgCl while for the other two coating system, the potential of

200 mV vs. Ag/AgCl was chosen (Table 1). The anodic potentials were 100 mV, 100 mV and 300 mV vs. Ag/AgCl for all three coating systems (Table 1), with even the highest potential levels falling within the normal oxidizing capacity of the environment with dissolved oxygen present [44]. Potentiostatic measurements for HVOF WC-CoCr coatings, Fig. 6, disclosed negative current density values at the lowest potential, 250 mV vs. Ag/AgCl, consistent with the expectations based on the polarisation curves, Fig. 6a; these results also confirmed that the ex­ periments were performed under pure wear (with the electrolyte as the lubricant). With increase in potential to 100 mV vs. Ag/AgCl (Fig. 6b), current density values detected in the absence of alumina counterbody (corrosion) were negative, whereas positive current density values were recorded under the contact of alumina counterbody (tribocorrosion); polarisation curves enabled to expect anodic current density values in both cases, yet the heterogeneous microstructure of the materials with evident local variations may explain the observation. Nevertheless, the corrosion-inducing role of the mechanical contact at this potential was clear, although the anodic current density values under the alumina counterbody were low, of the magnitude of 1–4 μA/cm2. With further increase in potential, Fig. 6c and d, the overall current density levels increased. The common trend was that the current density values associated with the systems with tribological counterbody (tribocorro­ sion) were higher than those without the counterbody (corrosion) and that the absolute values increased slightly with increase in test duration, up to ~ 100 μA/cm2, suggesting that active dissolution of the material 9

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Fig. 7. Current density records for the HVAF WC-CoCr coatings at the selected four potentials. a) 200 mV vs. Ag/AgCl. b) 100 mV vs. Ag/AgCl. c) 100 mV vs. Ag/ AgCl. d) 300 mV vs. Ag/AgCl. Corrosion refers to experiments carried out without the alumina counterbody, whereas tribocorrosion refers to experiments involving the alumina counterbody. Curves from the parallel repetitions of the tests are shown.

progressed during the experiments. The decrease in current density values, detected towards later stages of the experiment at the highest potential (Fig. 6d, tribocorrosion), was likely related to either the development of a corrosion product layer or decrease in the extent of surface irregularities, like pores, due to the levelling effect of the tribological contact. In the case of the corresponding coating composition deposited by HVAF method, Fig. 7, the current density values at the lowest potential revealed cathodic behaviour both in the corrosion and tribocorrosion systems (Fig. 7a), similarly to the corresponding HVOF coating. At the potential of 100 mV vs. Ag/AgCl, again, the measurements performed in the absence of counterbody (corrosion) introduced cathodic current density values, whereas those conducted in the presence of the coun­ terbody yielded anodic values of the magnitude of ~1 μA/cm2, Fig. 6b. At the two highest potentials: at 100 mV and 300 mV Ag/AgCl (Fig. 7c and d), differences between the HVAF and HVOF coatings became more apparent, as the overall current density levels recorded for the HVAF coating were at least one magnitude lower than in the case of HVOF coating (the difference for corrosion systems was even greater than that). This may be explained by the greater density of the HVAF coating due to the spraying process. The relatively constant and low current density values over time at both of these potentials referred to the sta­ bilized and slow anodic dissolution process. The stabilized current density values were systematically higher for the coatings under the tribological contact than in its absence, implying the mechanically

loaded counterbody slightly accelerated the corrosion attack. Only at the highest potential in the presence of the tribological counterbody (tribocorrosion), Fig. 7d, the current density level showed a slight increasing trend with progress of the experiment, referring to a slight increase in corrosion rate with progress of the test, possibly due to continuous exposure of fresh surface to the electrolyte by wear. For HVAF Cr3C2-WC-NiCrCo, Fig. 8, the current density trends at the two lowest potentials again followed those detected for the two other coating systems with simply cathodic values at the lowest potential and dissimilar cathodic (corrosion) and anodic (tribocorrosion) current density values at the potential of 100 mV vs. Ag/AgCl (Fig. 8a and b). At the two highest potentials (Fig. 8c and d), the current density values both in the absence (corrosion) and presence (tribocorrosion) of tribo­ logical counterbody were retained at a constant level during most of the test period, with the level of values being somewhat greater in the presence of tribological counterbody (tribocorrosion) than its absence (corrosion). These trends and even the overall levels of current density were equal to those detected for HVAF WC-CoCr coating. Nevertheless, in the presence of the counterbody (tribocorrosion), the current density values that were detected for HVAF Cr3C2-WC-NiCrCo coating featured a noteworthy scatter (Fig. 8c and d). We propose cracking of the coating under the mechanical load and the resulting intermittent exposure of fresh surface for the electrolyte a possible explanation for the current density fluctuations. Evolution of friction coefficient during the potentiostatic 10

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Fig. 8. Current density records for the HVAF Cr3C2-WC-NiCrCo coatings at the selected four potentials. a) 200 mV vs. Ag/AgCl. b) 100 mV vs. Ag/AgCl. c) 100 mV vs. Ag/AgCl. d) 300 mV vs. Ag/AgCl. Corrosion refers to experiments carried out without the alumina counterbody, whereas tribocorrosion refers to experiments involving the alumina counterbody. Curves from the parallel repetitions of the tests are shown.

experiments as a function of applied potential revealed that for all three coating systems, the overall friction coefficient decreased with increase in potential (Fig. 9). In the case of HVOF WC-CoCr, Fig. 9a, the friction coefficient at the lowest potential increased steadily with time, reaching almost 0.5 towards the end of the experiment. With increase in poten­ tial, friction coefficient values stabilized during the experiment, at the value of ~0.35 at the potential of 100 mV vs. Ag/AgCl, at ~0.25 at 100 mV vs. Ag/AgCl and at ~0.2 at the highest potential. These results reflect that either the surface roughness of the coating decreased with increase in potential or that some kind of lubrication was provided either due to dissolution or by the development of corrosion products film (solid lubrication). Nevertheless, given that the current density values increased with increase potential reflecting higher contribution by corrosion processes and the fact that only occassional corrosion products were detected on the surfaces, the middle explanation is most likely and supported by Fig. 10a. In the case of HVAF WC-CoCr, the overall friction coefficients were lower and the differences between the four potentials less pronounced than in the case of the corresponding HVOF coating, Fig. 9b, yet the trend was a decrease in the friction co­ efficient level with increase in potential. At the lowest potential, the friction coefficient peaked at the value of 0.35 but then slowly decreased to the value of 0.3 towards the end of the experiment. With increase in potential, there was only a minor decrease in the overall friction coef­ ficient level, with the lowest value of ~0.25 being obtained at the

highest potential. This finding illustrates minor changes in the surface state of the tribopair (alumina vs. WC-CoCr) with increase in the envi­ ronment oxidizing capacity (potential), consistent with the relatively low current density levels at all studied potentials. In the case of HVAF Cr3C2-WC-NiCrCo coating, the applied potential level had an evident influence on the friction coefficient, Fig. 9c. At the lowest potential, the highest friction coefficient among the studied coating systems was detected, occasionally exceeding the value of 0.5 and altogether involving sharp local fluctuations. Indeed this corresponds to pure wear situation under lubrication of the electrolyte, thus great contribution by wear to material behaviour is expected. With increase in potential to the value of 100 mV vs. Ag/AgCl, which was just within the anodic regime, friction coefficient values lowered to ~0.3–0.35 and showed much less fluctuations than at the lowest potential. Interestingly, at the two highest potentials, the lowest friction coefficient values among the studied coating systems were recorded, with the stabilized values of ~0.17–0.18. Here, however, the fluctuations were again significant, consistent with the observations on current density values. The comparison between electrolyte total metal contents after the corrosion and tribocorrosion experiments at the four potentials, Fig. 10, provides useful information about the degradation mechanism of the coatings. In the case of HVOF WC-CoCr (Fig. 10a), the issue to be noticed was the much higher concentration level of metals, particularly Co, released in the electrolyte than in the case of other two coating systems, 11

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Fig. 9. Evolution of friction coefficient during the experiments performed at the selected four potentials in the presence of alumina counterbody. a) HVOF WC-CoCr, b) HVAF WC-CoCr, c) HVAF Cr3C2-WC-NiCrCo. Example friction curves from only one replicate test are shown due to overlapping of the parallel data.

implying a higher degradation rate of HVOF WC-CoCr coating than of the two coatings deposited by HVAF method. Another important finding was that the metal concentrations in the electrolyte were approximately constant irrespective of the test mode (corrosion vs. tribocorrosion), thus the mechanical contact to the counterbody did not significantly change metal release in the electrolyte. This observation suggests that corrosion was an important degradation mechanism for the HVOF WCCoCr coatings at the two highest potentials, as charged based on the much greater metal release in the electrolyte than at the two lowest potentials. These findings are also supported by the current density re­ cords, Fig. 6. The observation that minor amount of metals were released in the electrolyte at the lowest potential refers to the presence of micro-galvanic corrosion, thus although the overall potential level falls within the cathodic range of the material behaviour, the most active constituent of the multi-phase structure may have undergone some degradation due to potential difference of the phases. In the case of HVAF WC-CoCr coating, Fig. 10b, metal release in the electrolyte increased almost in a linear manner with increase in potential yet the overall concentrations were much lower than for the corresponding WCCoCr coating. Here, the tribological contact to the counterbody also had a clear influence on the metal release in the electrolyte. At all four po­ tentials, Co was at the greatest amounts released in the electrolyte independently of the test mode (corrosion vs. tribocorrosion), with the contribution by Cr increasing under the tribocorrosion conditions with increase in potential. This may be explained by the removal of the Crrich top surface in the tribocorrosion experiments. For the HVAF

Cr3C2-WC-NiCrCo, Fig. 10c, the overall extent of metal release in the electrolyte was low and practically independent of potential. However, at the three lowest potentials, the metal release in the electrolyte was enhanced by the mechanical contact to the counterbody (tribocorro­ sion); this accounted for Ni and Cr. Only at the highest potential, the release of metals in the electrolyte during tribocorrosion experiments was equal to (Ni) or even lower (Co) than during corresponding corro­ sion experiments, with solely Cr being preferentially removed from the surface by the mechanical contact (tribocorrosion). These results do support the key role of cracking of the material in the degradation given it was not possible to define whether the Cr included in the electrolyte was as Cr or as Cr3C2. It is further emphasized that for this coating system, the high metal concentrations in the electrolyte at the lowest potential in the presence of the alumina counterbody was due to pure wear (with the electrolyte acting as a water-based lubricant). Thus pure wear was clearly an important degradation mechanism for HVAF Cr3C2WC-NiCrCo coating system, particularly keeping in mind the high fric­ tion coefficient at the cathodic potential. Only at the potential of 100 mV vs. Ag/AgCl was the metal release by the combined action of wear and corrosion greater than that by pure wear at 200 mV vs. Ag/AgCl. SEM characterisation was employed to provide deeper insights into and confirm the degradation mechanisms of the three types of coating systems. For HVOF WC-CoCr in the absence of the counterbody, SEM studies verified the corrosion of the Co-based matrix phase, Fig. 11a–d, and that the tribological contact slightly influenced the material removal mechanism. In the case of corrosion experiments, Fig. 11a–d, 12

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Fig. 10. Total metal concentration in the electrolyte after the experiments performed at the selected four potentials in the absence and presence of alumina counterbody. a) HVOF WC-CoCr, b) HVAF WC-CoCr, c) HVAF Cr3C2-WC-NiCrCo. Metal concentrations were defined for electrolyte sample from one parallel experiment only.

SEM examinations disclosed that degradation occurred primarily within the matrix phase exposing the WC phase particles; this is an indication that the mechanism may involve galvanic coupling between the WC hard phase and the alloy matrix resulting in micro-galvanic corrosion. Occasionally, pitting corrosion was also evidenced, Fig. 11d, yet we cannot exclude the possibility of galvanic nature of the attack, with the initial removal of the matrix phase, followed by loss of the WC particles within the local area. Under the counterbody contact, surface SEM in­ vestigations revealed some new marks of abrasive wear, such as in Fig. 11e (indicated in figure, label “A”), and striations or waves related to heavily deformed areas, Fig. 11g (indicated in figure, label “S”), earlier reported by Bolelli et al. [25]. Indeed, such striations have been related to near-surface plastic deformation of the material, primarily the matrix phase along which the WC particles are also dragged into wavy extrusions [25], and here these occasionally included corrosion products that contained plenty of W and O and some Cr, Co and C, as analysed by EDS. Besides plastic deformation, there were also signs of brittle cracking, typically in perpendicular direction to the sliding motion, Fig. 11f, and around the most heavily deformed areas of the matrix, Fig. 11g and h. The brittle cracking mechanism was probably respon­ sible for the majority of material wear losses, as it was evidently con­ nected to the detachment of small sections of material from the coating surface following sliding wear, similarly to findings reported earlier [17, 18,45]. The likely reason is the contribution of electrochemical disso­ lution to the local loss of the matrix phase, decreasing the capability of

the coating for plastic deformation in such region. The cracks and the detached material sections were of the size range of a few micrometers, typically in the length range from 2 to 5 μm. In some cases, the material was removed around the WC particles so that these particles were slightly protruding from the surface; this was likely linked with the easier removal of the relatively softer matrix phase, earlier reported also by Wang et al. [46] but may also be related to the micro-galvanic corrosion reported above. Additionally, surface films (tribocorrosion layer) were occasionally detected in the plastically deformed areas; these were enriched with respect to O and Cr as compared to sur­ rounding areas. In the case of HVAF WC-CoCr, Fig. 12, the phenomena that were observed were the same as for HVOF WC-CoCr, but the signs of materials degradation were less pronounced and their number (sites of note­ worthy general corrosion, brittle cracks or locations of detachment of sections of materials by wear) was lower than for the corresponding HVOF coating. Thus dissolution of the matrix phase under the corrosion conditions and removal of the matrix phase by sliding wear under tri­ bocorrosion conditions were the main features. The counterbody clearly intensified the material damage. Additionally, what was more evident for HVAF WC-CoCr coatings was the association of tribocorrosion products in the waves of plastic deformation, Fig. 12g and h (marked with “S”), which were also more frequent than in the HVOF coating. Again, the products were rich in W, O and Cr, but it was unclear whether WC particles were oxidized or if the deformed matrix was spread over 13

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Fig. 11. SEM images, showing the surface details of HVOF WC-CoCr coating after potentiostatic experiments in the absence (a–d) and presence (e–h) of the me­ chanically loaded alumina counterbody at various potentials. a) 100 mV vs. Ag/AgCl, corrosion. b-d) 300 mV vs. Ag/AgCl, corrosion. e-f) 100 mV vs. Ag/AgCl, tribocorrosion. g-h) 300 mV vs. Ag/AgCl, tribocorrosion. The arrows in images (e–g) indicate the sliding direction.

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Fig. 12. SEM images, showing the surface details of HVAF WC-CoCr coating after potentiostatic experiments in the absence (a–d) and presence (e–h) of the me­ chanically loaded alumina counterbody at various potentials. a-b) 100 mV vs. Ag/AgCl, corrosion. c-d) 300 mV vs. Ag/AgCl, corrosion. e-f) 100 mV vs. Ag/AgCl, tribocorrosion. g-h) 300 mV vs. Ag/AgCl, tribocorrosion. The arrows in images (e–g) indicate the sliding direction.

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Fig. 13. SEM images, showing the surface details of HVAF Cr3C2-WC-NiCrCo coating after potentiostatic experiments in the absence (a–b) and presence (c–h) of the mechanically loaded alumina counterbody at various potentials. a) 100 mV vs. Ag/AgCl. b) 300 mV vs. Ag/AgCl. c-d) 200 mV vs. Ag/AgCl. e-f) 100 mV vs. Ag/AgCl. g-h) 300 mV vs. Ag/AgCl.

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Fig. 14. SEM images, showing the tilted cross sections of the near-surface areas of the wear track in HVAF Cr3C2-WC-NiCrCo coatings (potentiostatic measurements at the potential of 300 mV vs. Ag/AgCl. The specimens were prepared by FIB processing.

the WC particles and undergone oxidation. In the case of HVAF Cr3C2-WC-NiCrCo, SEM investigations consoli­ dated earlier findings about the minor role of pure corrosion in the material losses, Fig. 13a and b. Indeed, grinding grooves from the disc preparation for experiments were still visible, with no clear signs of corrosion being detected even at the highest potential after the experi­ ment (Fig. 13b). In the tests involving mechanical load, the behaviour of the material changed. Under pure wear (at the lowest potential under the mechanical load) SEM studies reveled the most prominent feature to be the abrasive wear of the Cr3C2 phase, Figs, 13c-d. Another central finding was the abrasive wear of the matrix phase near or along the particles of WC phase, leading to protrusion of the WC particles from the surroundings surface. These observations may be easily explained by a large particle size of the Cr3C2 phase and the relatively lower hardness of Cr3C2 as compared to WC particles, allowing the predominant plastic deformation and wear loss in the Cr3C2 in contact to the Al2O3 coun­ terbody. The preferential deformation of the matrix phase has been re­ ported earlier [46], yet also the distribution of the matrix phase is of importance for the progress of the wear loss. Additionally, occasional dissolution of the matrix phase was detected particularly at the matrix-WC interfaces towards the higher potentials. Moreover, at the two highest potentials, cracking was evidenced along the Cr3C2-matrix interface. SEM examinations in cross section of FIB-processed speci­ mens, Fig. 14, confirmed the presence of micro-cracks at the Cr3C2-matrix interface in the near-surface areas of the HVAF Cr3C2-WC-NiCrCo coating in all areas of the wear track. Earlier,

Matthews et al. [47] and Hong et al. [48] have reported a similar brittle cracking behaviour in thermally sprayed Cr3C2–NiCr coatings under erosive wear. Indeed, these micro-cracks explain the high scatter detected in current density records (Fig. 8c and d) and friction coeffi­ cient values (Fig. 9c) at the potentials of 100 and 300 mV vs. Ag/AgCl. 3.3. Material losses and wear-corrosion synergy Material losses in the wear track were quantified through profilo­ metric data, Fig. 15. Wear track width values are shown in Fig. 15a. The main findings were that, first, the overall differences in the wear track widths between the four potentials were small and often within the scatter in the results and, second, there were slight differences in the wear track width values between the three coating systems, with a somewhat wider wear tracks being observed in the case of HVAF Cr3C2WC-NiCrCo coatings than in the case of the other two coating systems. In most of the cases for WC-CoCr coating systems deposited by both HVOF and HVAF methods, the wear track width values ranged from 400 to 440 μm, whereas for Cr3C2-WC-NiCrCo, the values fell within the range from 460 to 490 μm. It is also worth mentioning that the width values were systematically much (6–8 times) wider than the original contact diameter. The wear track depth values, Fig. 15b, demonstrated that the wear track cross-sections were actually wide and shallow, but also that the differences in wear track depth values were dependent on both the potential and the coating system of interest. Here the coatings could be divided into two categories. In the case of HVOF WC-CoCr and HVAF 17

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evidenced after the tests, whereas transferred dark grey layer was detected on the alumina ball wear track after the experiment done for HVAF Cr3C2-WC-NiCrCo at the lowest potential (Fig. 16d). Based on the dark colour, the transfer layer probably contains carbides included in the coating (either Cr3C2 or WC or both). Keeping in mind the abrasive wear of the Cr3C2 particles at the lowest potential (Fig. 13), this is the likely carbide phase attached to the alumina ball. Material tribocorrosion volume losses were here defined based on the wear track cross-sectional area from profilometric studies multiplied by the wear track length. These results, presenting total material losses in the wear track, are presented in Fig. 17a. Material losses for both WCCoCr coatings were quite equal, lower than 0.005 mm3 at the lowest potentials and increasing with increase in potential to the maximum losses of 0.01 mm3 at the highest potential. Material losses for HVAF Cr3C2-WC-NiCrCo were consistently higher, ranging from 0.005 to 0.02 mm3 and featuring a greater scatter than in the case of WC-CoCr coatings. Thus, in all cases, the material losses in the cermet coatings were significantly greater than in the counterbody specimens. Given that these were the total material losses, corrosion-induced wear and wearinduced corrosion within the wear track could be defined based on Equations (1)–(5). The contribution by different mechanisms to total volume losses in the wear track is shown in Fig. 17b. In the case of both WC-CoCr coat­ ings, the contribution by pure wear was an important degradation mode at all potentials and explained by the high friction coefficient values, with the significance of corrosion-induced wear mechanism growing with increase in potential. Indeed, at the highest potential, corrosioninduced wear was the dominant damage mechanism in the wear track in both WC-CoCr coatings. The damage by pure wear was slightly more intensive in HVOF WC-CoCr coatings than in corresponding HVAF coatings, likely due to the greater porosity and lower hardness of HVOF coating, although the total material losses in the wear track were roughly equal for the two WC-CoCr coatings. In the case of Cr3C2-WCNiCrCo, the major degradation mechanism in the wear track under all test conditions was pure wear of the material. This was supported by the microstructural data of the exposed surfaces, Figs. 13 and 14, particu­ larly abrasive wear within the Cr3C2 particles. Furthermore, for HVAF Cr2C3-WC-NiCrCo, the contribution of corrosion-induced wear was negative (beneficial) at the second-lowest potential. This is explained by the fact that the wear track dimensions were used to define the volume losses and the wear track at the potential of 100 mV vs. Ag/AgCl was much shallower than at the reference potential of 200 mV vs. Ag/AgCl (pure wear). For example, the lowered friction coefficient is one possible explanation behind the observation. At the two highest potential, the contribution by corrosion-induced wear was approximately of the same magnitude than for WC-CoCr coatings. In all cases, wear-induced corrosion played a negligible role in the wear track material losses, probably because of overall low current density values under the studied conditions and the fact that the surface films (passive film or uniform corrosion product film) did not play a role in the materials corrosion behaviour (the removal of which would have exposed fresh surface to the electrolyte). The contribution by wear-induced corrosion was greatest in the case of HVOF WC-CoCr at the potential of 100 mV vs. Ag/ AgCl, contributing to the volume losses of 1.3⋅10 5 mm3; these were the conditions where the increase in current density values by the tribo­ logical contact was the greatest. Material-dependent wear coefficient KW may be defined for the coatings based on the material losses by pure wear and Archard’s law [50,51]:

Fig. 15. a) Wear track width for the coatings, determined by profilometric measurements. b) Wear track depth for the coatings, determined by profilo­ metric measurements. In the case of HVOF WC-CoCr coating, the lowest po­ tential was 250 mV vs. Ag/AgCl, but the presentation here is simplified. Values defined based on the specimens from replicate experiments.

Cr3C2-WC-NiCrCo, the wear track depth values at the lowest potential (wear) were quite high; starting from the potential of 100 mV vs. Ag/ AgCl upwards within the anodic regime, the depth values increased steadily with increase in potential. For example, in the case of HVAF Cr3C2-WC-NiCrCo coating, only at the highest potential were the wear track depth values greater than those at the lowest potential. In turn, in the case of HVAF WC-CoCr coating, wear track depth values at the lowest potential were low and they increased systematically with in­ crease in potential throughout the studied potential range. Wear of the alumina ball counterbody was also examined in order to obtain understanding on the behavior of the entire tribosystem. Fig. 16a summarises the wear track diameter under various test conditions. The wear track diameter of 400 μm (average value of all experiments) cor­ responds to the wear volume of 0.00025 mm3, which is negligibly small material loss. The overall small wear volume may be justified by the high hardness of alumina, 1800–2000 HV [49]. There was very little variation in the determined counterbody wear track diameter values between the materials and potentials; the only exception was the experiment carried out for HVAF Cr3C2-WC-NiCrCo at the lowest po­ tential (under pure wear), with somewhat greater wear track diameter than under other conditions. The examination of the counterbody wear surfaces after the experiments supports the wear track diameter in­ vestigations (Fig. 16b–g), thus again the alumina ball surface from the experiment conducted for HVAF Cr3C2-WC-NiCrCo at the lowest po­ tential was different in appearance from the other counterbody surfaces. In most cases, no material transfer to the counterbody surface was

W ¼ KW

FN ⋅L H

(6)

where FN is the applied normal force, L is the sliding distance and H is the hardness of the material. Kw is typically expressed by takin the sliding distance account, thus: 18

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Fig. 16. a) Wear track diameter for the counterbody, determined by optical microscopy and the attached measuring tool (average of 4 measurements). The original contact diameter was 158 μm. b-g) Optical microscopy images of the counterbody wear track from the experiments. b) HVOF WC-CoCr, 250 mV vs. Ag/AgCl. c) HVAF WC-CoCr, 200 mV vs HVAF Cr3C2-WC-NiCrCo, 200 mV vs. Ag/AgCl.. Ag/AgCl. d) HVAF Cr3C2-WC-NiCrCo, 200 mV vs. Ag/AgCl. e) HVOF WC-CoCr, 300 mV vs. Ag/AgCl. f) HVAF WC-CoCr, 300 mV vs. Ag/AgCl. g) HVAF Cr3C2-WC-NiCrCo, 300 mV vs. Ag/AgCl. Values defined based on the specimens from replicate experiments.

K¼

W⋅H L⋅FN

coefficients (at the lowest potential, corresponding to pure wear): HVOF WC-CoCr 1.7⋅10 4 mm3/Nm, HVAF WC-CoCr 1.5⋅10 4 mm3/Nm and HVAF Cr3C2-WC-NiCrCo 5.5⋅10 4 mm3/Nm. These values reflect the lower wear resistance of the HVAF Cr3C2-WC-NiCrCo coating than of the

(7)

For the three coating systems, we obtained the following wear 19

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Fig. 17. a) Total material loss in the wear track, defined based on wear track profiles and the length of the wear track. b) Contribution by various degradation mechanisms to total material losses in the wear track. c) Contribution by various degradation mechanisms to total material loss in the HVOF WC-CoCr specimens. In the case of HVOF WC-CoCr coating, the lowest potential was 250 mV vs. Ag/AgCl, but the presentation in a) and b) is simplified. Values defined based on the specimens from replicate experiments.

two WC-CoCr coatings and are supported by the friction coefficient values and the results from SEM studies. Additionally, the great contribution by corrosion-induced wear to the overall material losses may be explained by the concentration of the corrosion attack at the carbide-matrix interface, thus adding surface roughness and exposing these interfaces, or the protruding carbide particles, to enhanced me­ chanical wear Figs. 11–13. Here, an important contributing issue is a likely micro-galvanic coupling between the phases, the WC (and other carbide phases) is typically relatively more noble than the matrix phase, resulting in local dissolution of the matrix phase around the carbide particles [ [52]]. The produced surface with a greater surface roughness and possibly with locally weakened matrix mechanical properties is worn more easily than the corresponding uniform and unattacked surface. In the wear track, the coating system superiority may be formulated as: HVOF WC-CoCr ~ HVAF WC-CoCr ≫ HVAF Cr3C2-WC-NiCrCo. Almost an equal tribocorrosion performance of HVOF and HVAF WCCoCr coatings in the experiments was surprising, keeping in mind the microstructural and hardness differences. Although the contribution by pure wear was slightly higher for HVOF WC-CoCr than for HVAF WCCoCr, the material losses by corrosion-induced wear then compensated the differences. However, also here wear was the predominant way to remove the material and the overall wear resistance of the coatings was relatively good (as compared to the third coating, HVAF Cr3C2-WCNiCrCo). Between the two HVAF coatings: WC-CoCr and Cr3C2-WC-

NiCrCo, the feature that made the most difference was the presence of Cr3C2 carbides. These were softer as compared to WC particles, as charged based on the abrasive wear that took place against the alumina counterbody. Another noteworthy feature was the relatively poorer interaction between the Cr3C2 carbide and the matrix than the WC carbide phase and the matrix phase, reflected by the cracking tendency in the coating in the vicinity of Cr3C2 carbides under the mechanically loaded counterbody. What actually made a difference between the HVOF WC-CoCr and HVAF WC-CoCr coatings during the experiments was the corrosion behaviour. Indeed, surface studies by SEM revealed corrosion in some of the studied coating systems, particularly in the case of HVOF WC-CoCr, outside of the wear track. This means that although wear and wearcorrosion synergy effects were concentrated in the wear track, mate­ rial losses actually occurred all through the coating surface. In the case of HVOF WC-CoCr, the strong contribution by corrosion in materials degradation was manifested by current density values which were one magnitude higher than for the other two coating systems. Therefore, for HVOF WC-CoCr we also defined the contribution by corrosion to the total material losses in the specimens using Equation (5). Fig. 17c shows the contribution of all involved degradation mechanisms to the total tribocorrosion losses of the HVOF WC-CoCr specimens. At the two lowest potentials, the contribution by other mechanisms than wear was negligible. At the two highest potentials, the contribution by corrosion was minor, at maximum 5% at the highest potential. However, we must 20

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remember that corrosion occurred preferentially in the matrix phase near the WC phase interface and was likely driven by the galvanic coupling between the phases. This implies that corrosion was localized and lead to pit-like attack sites, the growth of which is typically very challenging to predict. In the case of the other two coating systems (HVAF WC-CoCr and HVAF Cr3C2-WC-NiCrCo), corrosion played a negligible role in the overall material losses. By taking these results into account, the final ranking of the coatings is as follows: HVAF WCCoCr > HVOF WC-CoCr ≫ HVAF Cr3C2-WC-NiCrCo. The comparison between the WC-CoCr coatings deposited by the two thermal sprayed methods: HVOF and HVAF, enabled to correlate the better corrosion resistance of HVAF WC-CoCr coating to the lower degree of porosity than in the corresponding HVOF coating, indicating that pores act as the preferential initiation sites for the corrosion attack.

Acknowledgements

4. Conclusions

References

Three types of thermally sprayed cermet coatings: HVOF and HVAF WC-CoCr coatings as well as HVAF Cr3C2-WC-NiCrCo coatings, were deposited on stainless steel substrate and subjected to tribocorrosion experiments in order to compare their performance in paper machine environment and define the wear-corrosion synergy mechanisms. The main conclusions based on the findings of this work are:

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The financial support for the SUBTRIB project from the Finnish Funding Agency for Technology and Innovation (Business Finland; de­ cision number 865/31/2016), participating companies and VTT Tech­ nical Research Centre of Finland Ltd is gratefully acknowledged. Mr. Simo Varjus is cordially thanked for conducting the pin-on-disc experi­ ments. The assisting personnel at VTT, particularly Tuomo Kinnunen and Taru Lehtikuusi, is thanked for the help in the laboratory work. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.triboint.2019.106006.

1) HVOF method yielded WC-CoCr coatings with a greater degree of porosity and lower values for indentation hardness and plane strain modulus than HVAF method. Cr3C2-WC-NiCrCo coatings featured wider continuous sections of the matrix phase than WC-CoCr coat­ ings, due to the relatively larger size of the Cr3C2 than WC particles. The mechanical properties of HVAF Cr3C2-WC-NiCrCo coatings were intermediate to those of HVAF WC-CoCr (highest) and HVOF WCCoCr (lowest). 2) None of the three coatings passivated in the paper machine envi­ ronment, yet pseudopassivity was detected. For HVOF WC-CoCr, the overall level of anodic current density values was highest and pure corrosion, as micro-galvanic corrosion of the matrix phase at WC interface, played a role in its degradation. For coatings deposited by HVAF method, the contribution by pure corrosion was negligible. 3) The contribution by pure wear was greatest for HVAF Cr3C2-WCNiCrCo coating. Besides wear of the matrix phase, the embedded Cr3C2 particles experienced evident abrasive wear. Under pure wear, friction coefficient for Cr3C2-WC-NiCrCo coating was highest among the studied coating systems, above 0.5. At anodic potentials, friction coefficient was the lowest among the three coatings but cracking was evidenced at the Cr3C2 particle-matrix interface. 4) Wear-corrosion synergy for all three coating systems occurred mainly as corrosion-induced wear. Wear tracks grew in depth with increase in potential, probably due to microgalvanic corrosion of the matrix phase and easier wear of the roughened coating surface. In the case of WC-CoCr coatings, the contribution by corrosion-induced wear increased linearly with increase in potential and corre­ sponded to majority of the material losses at the highest potential. 5) Among the three coating systems, the primary choice for the paper machine environment is the studied HVAF WC-CoCr coating, due to the lowest material losses. In the case of HVOF WC-CoCr, the ma­ terial loss by pure wear and corrosion-induced wear were equal to those in HVAF WC-CoCr, but for this coating system the contribution by pure corrosion made a difference to HVAF WC-CoCr coating. The highest overall tribocorrosion losses were evidenced for HVAF Cr3C2WC-NiCrCo coatings. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. 21

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