Cermet coatings with Fe-based matrix as alternative to WC–CoCr: Mechanical and tribological behaviours

Surface & Coatings Technology 206 (2012) 4079–4094

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Cermet coatings with Fe-based matrix as alternative to WC–CoCr: Mechanical and tribological behaviours Giovanni Bolelli a,⁎, Tim Börner a, Francesco Bozza a, Valeria Cannillo a, Gennaro Cirillo b, Luca Lusvarghi a a b

Department of Materials and Environmental Engineering, University of Modena and Reggio Emilia, Via Vignolese 905, I-41125 Modena (MO), Italy Parma Spray Italia S.r.l., Via Giovanni XXIII, I-43040 Varano de'Melegari (PR), Italy

a r t i c l e

i n f o

Article history: Received 20 February 2012 Accepted in revised form 31 March 2012 Available online 10 April 2012 Keywords: WC–FeCrAl cermet High Velocity Oxygen-Fuel (HVOF) spraying Sliding wear Abrasive wear Residual stress Cyclic impact behaviour

a b s t r a c t Recently, cermet coatings with Fe-based metal matrix have emerged as a less hazardous and more environmentally friendly alternative to WC–Co-based ones, which have known inhalation toxicity problems. This study therefore aimed to validate WC-based cermet coatings with Fe-based matrix, obtained using a commercially available feedstock powder, as an alternative to WC–CoCr. HVOF-sprayed WC–15 wt.%FeCrAl layers were therefore obtained using different oxygen and fuel (kerosene) flow rates and powder feed rates; their mechanical and tribological properties were compared to HVOFsprayed WC–10 wt.% Co–4 wt.%Cr. The WC–FeCrAl coatings always exhibited equi-biaxial compressive residual stress state and possessed dense microstructures, with homogeneous metal matrix, but they contained more oxide inclusions than WC–CoCr. Their characteristics were significantly affected by the normalised oxygen-fuel ratio (λ). Small but meaningful differences existed between the ball-on-disc sliding wear rates of the various WC– FeCrAl coatings, the best sample being that with the most favourable combination of compressive residual stress, low oxidation and high hardness/modulus (H/E) ratio. Its sliding wear resistance was comparable to that of WC–CoCr. The cyclic ball impact resistance of WC–FeCrAl layers was also comparable to that of WC–CoCr, but the dry particle abrasion resistance was inferior, because of the brittleness induced by the oxide inclusions. © 2012 Elsevier B.V. All rights reserved.

1. Introduction Cermets consisting of WC particles embedded in a Co- or Ni-based metal matrix or of CrxCy in a Ni-based matrix are frequently employed as coating materials for the protection against various forms of wear (sliding, abrasive, erosive, etc.) in industrial machinery [1–8]. Thermal spray processes, especially those involving high kinetic energy, such as the High Velocity Oxygen-Fuel (HVOF) or High Velocity AirFuel (HVAF) techniques, are usually the preferred deposition methods, on account of the rather high productivity and of the high quality of the resulting layers, which possess low porosity, high cohesive strength and satisfactory adhesion to the substrate [9,10]. These coatings, however, pose some safety and health issues. First of all, handling of the feedstock powder in thermal spray shops implies some hazards. Ni-based alloys are allergenic and are labelled as suspect carcinogenic agents, as seen e.g. in [11]: for instance, Nibased powders come under the H351 hazard statement according to European Commission regulation EC 790/2009 [12]. WC–Co materials are also toxic by inhalation [13] and are currently under consideration

⁎ Corresponding author. Tel.: + 39 0592056233; fax: + 39 0592056243. E-mail address: [email protected] (G. Bolelli). 0257-8972/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2012.03.094

for inclusion in the list of suspect human carcinogens as well [14]. On the other hand, Ni- and Co-containing coatings sometimes fail to conform to the requirements of some very specific but economically significant application fields, most notably the food and beverage processing industry, where usable materials must meet stringent requirements. Specifically, strict limitations exist against the potential contamination of the products with toxic or harmful elements, released in the form of fine wear debris particles or of ions, leached in liquid media [15,16]. Searching for replacements to Ni- and Co-based coating materials, capable of overcoming the above-mentioned shortcomings, the authors recently investigated the properties of HVOF-sprayed Fe-based metal alloy coatings [17]. Although they were shown to be viable alternatives for electroplated hard chromium and for other thermallysprayed metal alloy coatings, their wear resistance could not compare with that of cermet materials. These coatings, therefore, would not be suitable in applications where the tribological characteristics of conventional cermets, such as WC–Co-based ones, would be required. A better solution for this purpose might be offered by thermal spray-grade cermet powders free of Ni and Co, consisting of WC particles in a Fe alloy matrix, which have recently been introduced in the market [11]. The properties of the thermally-sprayed coatings obtained from those powders were considered in a few publications

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[13,18], but a complete and systematic investigation has not been performed yet. The purpose of the present paper, therefore, was to provide a comprehensive mechanical and tribological characterisation of WC– FeCrAl coatings deposited by liquid-fuelled HVOF-spraying technique. Their microstructural features, depth-sensing indentation response, residual stresses, sliding and abrasive wear behaviour and cyclic impact resistance were investigated as a function of some key spraying parameters and were compared to those of a standard WC–CoCr layer, employed as a term of comparison.

Table 2 HVOF process parameters. Parameter set

O2 flow rate (SL/min)

Kerosene flow rate (L/h)

λa

Powder feed rate (g/min)

Stand-off distance (mm)

#1 #2 #3 #4 #5

950 850 950 850 900

23 18 18 23 21

1.27 1.46 1.63 1.14 1.32

100 100 60 60 80

400

a λ = normalised oxygen fuel ratio, computed based on a simplified kerosene combustion reaction C12H26(l) + 37/2 O2(g) → 12 CO2(g) + 13 H2O(g) with ρ = 0.8 g/cm3 as density of kerosene and ρ = 1.4290 kg/m3 as standard density of O2.

2. Materials and methods size analyser with Hydro-S wet dispersion unit, Malvern Instruments Ltd., Malvern, UK).

2.1. Coating deposition Commercially available thermal spray-grade WC–15 wt.% FeCrAl and WC–10 wt.% Co–4 wt.% Cr feedstock powders, manufactured by agglomeration and sintering, were employed (Table 1). The substrates were C40 plain carbon steel plates (100 × 100 × 5 mm) and AA 6082T6 aluminium alloy plates (80 × 80 × 8 mm); the substrates were preliminarily degreased with acetone, grit-blasted using 240 μm alumina grits and pre-heated at 60 °C for 1 h in a drying oven to remove adsorbed moisture. The powers were sprayed onto the substrate plates using a kerosene-fuelled Praxair-Tafa JP5000 HVOF torch, equipped with a 152.4 mm-long nozzle. As specified in Table 1, the WC–FeCrAl powder was deposited onto the carbon steel plates using 5 distinct parameter sets (labelled from #1 to #5), which differed in oxygen and kerosene flow rates and powder feed rates (Table 2), in order to study their influence on coating properties. An additional WC–FeCrAl coating was deposited onto the aluminium alloy plate using parameter set #1, in order to check possible influences of substrate properties on the characteristics of the coating. Parameter set #1, representing standard HVOF processing conditions for WC–CoCr cermets, was also employed in order to spray the WC–CoCr feedstock powder onto the carbon steel plates, in order to obtain reference coatings, which serve as a term of comparison to evaluate the characteristics of the WC–FeCrAl layers. 2.2. Characterisation of the WC–FeCrAl powder The phase composition of the WC–FeCrAl powder was analysed by X-ray diffractometry (XRD: X'Pert PRO diffractometer, PANAlytical, Almelo, The Netherlands), using Cu–Kα radiation emitted from an X-ray tube operated at 40 kV energy and 40 mA current, and employing a gas-proportional X-ray detector equipped with single-crystal monochromator. A 2θ range from 20° to 100° was scanned with a step size of 0.025° and an acquisition time of 0.90 s/step. Cross-sections of the powder particles, obtained by cold-mounting in polyester resin followed by grinding using SiC papers (up to 2500 mesh) and polishing with diamond slurries (up to 3 μm size), were observed by scanning electron microscopy (SEM: XL-30, FEI, Eindhoven, The Netherlands); qualitative and semi-quantitative chemical analyses were also obtained by energy-dispersive X-ray microanalysis (EDX: INCA, Oxford Instruments Analytical, Abingdon, UK). The particle size distribution was assessed by laser diffraction technique using wet dispersion method (Mastersizer 2000 particle Table 1 Characteristics of feedstock powders (manufacturer: H.C. Starck, Laufenburg, Germany). Composition (wt.%) Powder trade name Particle size range (μm) Process parameter sets (see Table 2)

WC–15(Fe–20Cr–7Al) Amperit 618.074 − 45 + 15 #1–#5

WC–10Co–4Cr Amperit 556.074 − 45 + 15 #1

2.3. Characterisation of coatings Cross-sections of the WC–FeCrAl and WC–CoCr coatings were prepared by metallographic cutting, hot-mounting in phenol resin, grinding with diamond papers (up to 5 μm particle size) and polishing using diamond slurries (up to 3 μm particle size). Polished cross-sections were employed for microstructural observations through SEM and for qualitative and semi-quantitative chemical analysis through EDX. Specifically, semi-quantitative assessments of chemical compositions were obtained by performing large-area EDX scans on two cross-sectional views at 400× magnification for each coating, with electron beam energy of 15 keV. The amount of oxide inclusions in each coating was assessed by image analysis technique (software: ImageJ version 1.43u, NIH, Bethesda, USA) onto 5 backscattered electron (BSE)-SEM micrographs at 1000× magnification. Coating thickness was also assessed by image analysis on low magnification micrographs. Depth-sensing Berkovich micro-indentation tests were also performed onto the polished cross-sections, with maximum load of 3 N, loading/unloading rate of 2.4 N/min, holding time at maximum load of 15 s (Micro-Combi Tester, CSM Instruments, Peseux, Switzerland). Instrumented Vickers hardness and elastic modulus were assessed using the Oliver–Pharr procedure [19], in accordance with the ISO 14577 standards. Square samples of 23 × 23 mm were cut from the coated plates and polished using diamond papers and diamond slurries as described above, achieving a final roughness Ra ≈ 0.02 μm (assessed through optical confocal profilometry: Conscan profilometer, CSM Instruments). The thickness reduction during grinding and polishing was assessed using a digital outside micrometre. Phase composition of the coatings was assessed onto polished surfaces by XRD, using the same experimental set-up described in Section 2.1. Residual stress was also measured on the polished surfaces by Xray diffractometry. The classical sin 2ψ technique was employed [20]: 11 different ψ-tilt values symmetrically distributed in the [−45°; + 45°] range, corresponding to sin 2ψ values of 0, 0.1, … , 0.5, were adopted in order to follow the displacement of the (112) peak of WC, located at 2θ ≈ 98.7° (according to JCPDF 51-939). Using the same diffractometer mentioned above in a point-focus configuration, a 2θ scan from 97.0° to 100.5°, with a step size of 0.017° and an acquisition time of 151.5 s/step, was performed at each ψ-tilt position. The complete ψ-scan was repeated along three different φ orientations (0°, 45°, 90°). The normal stress along each φ orientation was computed using the equation for plane-stress condition (1) [20]:

ε φψ ¼

dφψ −d0 1 ðhklÞ ðhklÞ 2 ¼ S1 ðσ 1 þ σ 2 Þ− S2 σ φ sin ψ: 2 d0

ð1Þ

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Where: dφψ d0 σφ ðhklÞ S1

interplanar spacing of the selected diffraction peak ((112) peak of WC), measured at a given φ,ψ position; interplanar spacing of the selected diffraction peak ((112) peak of WC) in the unstressed material; normal stress along the φ direction;

ðhklÞ ¼ Ev ðhklÞ ; 12 S2 ¼ 1þv X-ray elastic constants along the E ðhklÞ (hkl) direction.

Similar to the measurement procedure described in [21], the coating's elastic modulus E assessed by depth-sensing indentation was employed in Eq. (1), while the Poisson's ratio was assumed to be υ = 0.3: although this may imply a certain degree of approximation, the results are certainly useful in order to rank the residual stress magnitude in the various coated samples. The unstressed interplanar spacing d0 was taken from the XRD pattern of the WC–FeCrAl feedstock powder. In any case, the error on σφ due to uncertainties on d0 is generally of ≈0.1% [22]. Accordingly, it was ascertained that, if this d0 value was replaced with the dψ = 0° value in the analysis of the present results, the change in the value of σφ was practically negligible. Dry sliding wear tests were performed on the polished, 23× 23 mmsized square samples under ball-on-disc configuration (according to ASTM G99-05), using a pin-on-disc tribometer (High-Temperature Tribometer, CSM Instruments) equipped with 6 mm-diameter sintered α-Al2O3 spheres (nominal Vickers hardness = 1900 HVN). Tests were performed at a temperature of 25 ± 2 °C with a relative humidity of 58± 2%. The normal load was 10 N, the linear relative sliding speed of the ball on the sample was 0.10 m/s, the radius of the wear track was 7 mm and the overall sliding distance was 5000 m. Each sample was tested at least twice to ensure repeatability of the results.

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The volume wear loss of the sample was assessed by optical confocal profilometry (Conscan profilometer, CSM Instruments) and was converted to wear rate, i.e. volume loss per unit normal load and unit sliding distance, according to Archard equation [23]. The worn surface of the sample was also inspected by SEM. Dry particle abrasion tests were performed on as-deposited coating surfaces (i.e. without polishing) using a dry sand-steel wheel apparatus (a modified version of the ASTM B611-85(2005) test method for abrasive wear resistance of cemented carbides), previously employed by the authors in [24,25] as well. The sample was pressed with a 40.2 N load against a 200 mm-diameter Fe360 steel wheel rotating at 75 rpm, with a tangential flow (75 g/min) of angular alumina particles (180 μm average size). A total of 90 disc revolutions were performed (corresponding to a wear distance of 56.5 m). All samples were tested at least three times. The volume wear loss of the samples was computed based on the measured wear scar length and was converted to wear rate as described previously. The wear scars were also observed by SEM (Quanta 200, FEI). Cyclic ball drop impact tests were performed on the polished 23 × 23 mm square samples: a 39 mm-diameter X200Cr11 steel ball, attached to an overall mass of 1.35 kg, was dropped 200 times (3 Hz impact frequency) onto a fixed location on the sample surface. The schematic of the test equipment is shown in [26]. The surface morphology of the tested coatings was inspected by optical microscopy and the volume of the impact crater was assessed by optical confocal profilometry (Conscan profilometer). Cross-sectional morphologies were further inspected by SEM (Quanta-200): the samples were embedded in epoxy resin, dissected by metallographic cutting, ground and polished with diamond papers and diamond slurries as described previously.

Fig. 1. Cross-sectional backscattered electron (BSE)-SEM micrographs of the WC–FeCrAl feedstock powder (A: overview; B: magnified view of one particle; C: detail of WC particles) and EDX microanalyses (D) acquired on the regions marked as “1” and “2” in panel B. In panel B: label 1 = WC-rich area; label 2 = Cr-rich area without WC particles.

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Fig. 2. XRD patterns of the WC–FeCrAl feedstock powder and of corresponding HVOFsprayed coatings. Legend: = WC (JCPDF 51-939); = W2C (JCPDF 35-776); = W (JCPDF 4-806) or (W,Cr)2C [32]; α = Fe — b.c.c. (JCPDF 6-696); η = M6C (JCPDF 271125).

3. Results and discussion 3.1. Microstructural characterisation of powders and coatings The narrow, monomodal particle size distribution of the WC– FeCrAl powder, characterised by d0.1 = 19.0 μm, d0.5 = 30.7 μm, and d0.9 = 47.0 μm, matches remarkably well with the nominal values indicated in Table 1. Its particles exhibit a rounded, porous morphology (Fig. 1A), typical of agglomerated and sintered powders, as seen e.g. in [27,28]. Backscattered electron views highlight a clear separation (better seen in the higher magnification micrograph of Fig. 1B) between darker regions, rich in Cr but devoid of WC (EDX spectra in Fig. 1D), and brighter regions, rich in WC particles embedded in a Fe–Al matrix. Specifically, micrometric and sub-micrometric WC particles, having quite angular morphology, co-exist in these regions (Fig. 1C). This separation is reflected by the XRD pattern of the powder (Fig. 2), where, apart from the obvious diffraction peaks of WC, peaks ascribable to α-Fe (b.c.c. structure) and to M6C (η-phase) are detectable: they presumably correspond to the Fe–Al matrix surrounding the WC particles and to the dark Cr-rich regions, respectively. This means the η-phase, in this case, consists of a (Cr,Fe)6C structure. After HVOF spraying, dense and homogeneous coatings, with thicknesses lying between 130 μm and 220 μm (Table 3), free of any major defect and of visible porosity, and macroscopically similar to the WC–CoCr reference, are obtained (Fig. 3). By comparing the thickness data in Table 3 and the process parameters in Table 2, it can be inferred that thickness is primarily affected by powder feed rate. The overall chemical composition of the coatings, determined by semi-quantitative EDX analysis (Table 3: the values are the average

± standard deviation of the two large-area EDX scans acquired on each coating, as specified in Section 2.3), matches reasonably well with the nominal one listed in Table 1; however, significant amounts of oxygen seem to have been incorporated, particularly in the WC– FeCrAl layers. It should be remarked that oxygen concentrations in Table 3 are to be considered as indicative values only, since the reliability of a quantitative assessment of oxygen by the EDX technique is limited. Some approximation in the chemical analysis in Table 3 (where the sum of all weight percentages was normalised to 100) is due to the inability of the EDX technique to quantify C. The clear distinction between bright WC-rich areas and dark Crrich areas, which characterised the powder particles, is no more present in all of the WC–FeCrAl coatings (Figs. 4A–E and 5). By contrast, the WC–FeCrAl coatings exhibit a homogeneous matrix (Fig. 5) simultaneously containing Fe, Cr, Al and, presumably, some dissolved W (as shown by the EDX spectrum in Fig. 6A). Consistently, the X-ray diffraction peaks of α-Fe and η-M6C are not detectable in the XRD patterns of the WC–FeCrAl coatings (Fig. 2). The coatings also contain a significant amount of oxide inclusions (easily recognisable from their dark contrast in BSE micrographs, Fig. 4A–E), consistent with the EDX results in Table 3. These inclusions appear both as elongated stringers at lamellar boundaries (Fig. 5, label 1) and as small spherical intralamellar particles (Fig. 5, label 2). The former are somewhat enriched in Al compared to the base alloy composition, but contain significant amounts of W, Cr and Fe as well, as shown by the EDX spectra in Fig. 6B; the latter, by contrast, are extremely rich in Al (Fig. 6B): probably, Al2O3 is their main constituent. All of these microstructural features clearly indicate that, during spraying, irrespective of the processing conditions, the matrix phases (α-Fe and η-M6C) of the original WC–FeCrAl powder particles were completely melted. During flight, the molten phase experienced interdiffusion, homogenisation, and some degree of oxidation; then, upon impact onto the substrate, it solidified into an amorphous and/or nanocrystalline structure, as a consequence of the high quenching rate. This phenomenon is analogous to the amorphisation of the metal matrix in thermally-sprayed WC–Co-based cermets [29]. In-flight interaction of the molten matrix with the O2 left in the HVOF gas jet after the combustion of kerosene (all of the process parameters indeed feature oxygen-rich mixtures with λ > 1, as shown in Table 2, since the kerosene-fuelled JP5000 torch can only operate with excess oxygen) and/or with surrounding air entrained in the jet by turbulent mixing causes the formation of oxides. The oxides are originally developed on the surface of the droplet and are then dragged towards the centre because of turbulent flow within the molten metal, giving rise to the spherical, Al-rich inclusions seen inside the splats in Fig. 5A and B. Their morphology is indeed identical to that of oxide inclusions developed in flight inside thermally-sprayed metal droplets [30,31]. After impact quenching onto the substrate, the splats still remain at high temperature for a short time, and experience subsequent brief re-heating during successive passes of the torch. During these periods, oxidation of the surface of the deposited lamella occurs,

Table 3 Chemical composition of the coatings from EDX analysis, index of carbide retention (I) computed from Eq. (2), oxide content and thickness from image analysis on BSE-SEM micrographs. Sample

W (wt.%)

Fe (wt.%)

Cr (wt.%)

Al (wt.%)

O (wt.%)

I

Oxide content (vol.%)

Thickness (μm)

#1 #2 #3 #4 #5 WC–CoCr

80.0 ± 0.6 78.8 ± 0.5 78.4 ± 0.1 79.9 ± 0.3 80.8 ± 0.9 81.7 ± 0.6

10.8 ± 0.1 11.7 ± 0.1 11.0 ± 0.1 10.9 ± 0.2 10.9 ± 0.5 Co: 11.3 ± 0.2

2.6 ± 0.4 2.9 ± 0.2 2.7 ± 0.1 2.8 ± 0.1 2.7 ± 0.3 4.4 ± 0.2

0.6 ± 0.1 0.8 ± 0.1 0.9 ± 0.1 0.7 ± 0.1 0.6 ± 0.1 –

6.0 ± 0.4 5.9 ± 0.4 6.9 ± 0.1 5.7 ± 0.3 5.0 ± 0.1 2.7 ± 0.1

0.78 0.89 0.83 0.75 0.81 0.88

7.3 ± 2.0 9.6 ± 0.6 7.9 ± 1.0 6.1 ± 0.6 5.5 ± 1.1 1.8 ± 0.7

184 ± 6 217 ± 8 151 ± 9 133 ± 10 164 ± 8 158 ± 12

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Fig. 3. Cross-sectional BSE-SEM micrographs (overviews) of the HVOF-sprayed coatings deposited onto C40 steel plates: WC–FeCrAl #1 (A), #2 (B), #3 (C), #4 (D), #5 (E) and WC– CoCr reference (F).

according to the mechanisms illustrated in [31], developing the splatboundary oxide stringers also seen in Fig. 5A and B. This interpretation of the oxidation mechanisms, based on extensive previous literature studies [30,31], accounts for the clear difference between the morphology and chemical composition of the two oxide types. Such difference is indeed a clear indication that the oxides were developed by two distinct and different processes (in-flight oxidation and post-deposition oxidation, respectively). In both cases, Al is oxidised preferentially; indeed, according to Ellingham's diagram, Al has greater high-temperature affinity for oxygen than Cr, W and Fe. During in-flight oxidation, however, Al is almost the only oxidised element, whereas, during post-deposition oxidation, the other metallic elements are also affected. Presumably, WC particles were also affected during spraying: they experienced some decarburisation and dissolution in the molten matrix. Microstructural evidence of these phenomena includes the rounded edges of the carbide particles in the coatings (compare Fig. 5A, B to the original morphology of WC in the feedstock powder, Fig. 1C), the disappearance of some of the original sub-micrometric WC particles (again, Fig. 5A, B can be compared to Fig. 1C), and the formation of a few bright, W-rich matrix areas with no recognisable

WC (Fig. 5A, label 3: the corresponding EDX spectrum in Fig. 6A confirms that these areas are richer in W than ordinary matrix areas). Further evidence comes from the XRD patterns of the coatings, which exhibit diffraction peaks of W2C witnessing some degree of decarburisation of WC. A weak secondary peak at 2θ ≈ 41° also appears in the diffraction patterns of the coatings. Its precise identification is hindered by its very low intensity; however, it is hypothesised that it might belong to W (although its position does not match exactly with that reported in the JCPDF 4-806 file) or to (W,Cr)2C (according to [32]). The formation of W would further corroborate to the occurrence of some decarburisation of WC during spraying, whereas, if (W,Cr)2C were developed, this might suggest some interaction of dissolved WC with the metal matrix. Presumably, the fraction of W and C which remained dissolved within the matrix as a consequence of carbide dissolution (without entering the mixed carbide phase (W,Cr)2C which might have appeared in the coatings) further promoted its amorphisation during impact quenching (as discussed above), by opposing further hindrance to crystallisation [29]. The two key microstructural features mentioned above, namely decarburisation and oxidation, were quantified by peak fitting on XRD patterns and by image analysis on SEM micrographs, respectively.

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Fig. 4. Cross-sectional BSE-SEM micrographs (1000 × magnification) of the HVOF-sprayed coatings deposited onto C40 steel plates: WC–FeCrAl #1 (A), #2 (B), #3 (C), #4 (D), #5 (E) and WC–CoCr reference (F).

Specifically, decarburisation was quantified through the index of carbide retention I, defined in accordance with [33] (Eq. (2)): I¼

IWC : I WC þ I W2C þ IW

ð2Þ

Where: IWC IW2C IW

integral intensity of the main WC diffraction peak at 2θ = 35.6°; integral intensity of the main W2C diffraction peak at 2θ = 39.8°; integral intensity of the main W diffraction peak at 2θ = 40.3°.

Integral intensities were obtained by fitting the diffraction peaks shown in Fig. 2 using pseudo-Voigt peak functions. Since the peak at 2θ ≈ 41° cannot be undoubtedly ascribed to W, it was not employed for the computation of I. This might cause I to be somewhat overestimated, since the amount of W possibly lost in the formation of W or of (W,Cr)2C is neglected. The ranking between the various WC–FeCrAl

coatings, however, can be considered meaningful, particularly because the intensity of the peak at 2θ ≈ 41° varies between samples in the same way as that of the W2C peaks, i.e. the coatings with lower amount of W2C also contain less W or less (W,Cr)2C. The results (Table 3) suggest that the degree of decarburisation is primarily controlled by the oxygen/fuel ratio (λ, Table 2) during spraying; accordingly, highest decarburisation (lowest I value, samples #1 and #4) is associated to the lowest values of λ, whereas the less decarburised samples (highest I) are those sprayed with highest λ values (samples #2 and #3). Quantification of oxide contents by image analysis (Table 3) shows that, first of all, the WC–FeCrAl coatings are definitely more oxidised than the WC–CoCr reference, consistent with the EDX results. This can also be qualitatively seen by comparing Fig. 4A–E to F and Fig. 5A–B to C. Specifically, since the WC–FeCrAl #1 coating and the WC–CoCr reference were deposited under identical process parameters using feedstock powders with similar particle size distributions (Table 1), the different oxide contents must have been caused by higher oxidation tendency of the FeCrAl matrix compared to the CoCr one. This is probably due to the large reactivity between Al and oxygen at high temperature, as discussed previously.

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on the key role of the oxygen-fuel ratio to determine properties such as oxide content in HVOF-sprayed coatings were drawn in [34]. 3.2. Microindentation testing Indentation testing (Fig. 8) shows that the hardness of the WC– FeCrAl coatings (roughly comprised between 1000 HV3N and 1100 HV3N) and their modulus (around 200 GPa) is lower than those of the WC–CoCr reference (approximately 1200 HV3N and 350 GPa, respectively). However, the mechanical properties of the WC–FeCrAl coatings can be modified, to some extent, by the selection of process parameters, just as it happened for the carbide retention and oxide content discussed in Section 3.1. Specifically, Fig. 8 suggests that hardness and modulus of the WC–FeCrAl coatings follow roughly inverse trends, i.e. samples exhibiting the highest hardness (samples #4 and #5) possess lower modulus and vice versa. Specifically, the differences between the modulus values of the various WC–FeCrAl coatings are significantly large, compared to the associated standard deviations (plotted in Fig. 8 as error bars), whereas the differences between hardness values are quite small. It should preliminarily be remarked that the standard deviations of the hardness measurements, of about 70–90 Vickers units in all cases, are definitely consistent with the typical standard deviations associated to the Vickers hardness values of HVOF-sprayed cermet coatings, as seen e.g. in [35,36]. Due to the presence of scattered defects and inhomogeneities (such as oxide inclusions, weak interlamellar boundaries, etc.), indeed, hardness values measured on thermally sprayed coatings usually exhibit significant scatter [36]. The actual statistical

Fig. 5. Cross-sectional BSE-SEM micrographs (details) of the WC–FeCrAl coatings #1 (A) and #5 (B) and of the WC–CoCr reference (C). In panel A: 1 = splat-boundary oxide stringer; 2 = spherical intralamellar oxide inclusion; 3 = decarburised area with no recognizable WC particles.

Another important feature highlighted by image analysis is the possibility to vary, within certain limits, the oxide content of WC– FeCrAl coatings by adjusting the process parameters (Table 3). The lowest volume fraction of oxide inclusions is exhibited by samples #4 and #5, which are also those with the lower oxygen content, as determined by EDX micro-analysis (although such quantitative assessment has important limitations, as mentioned previously). These coatings were deposited using oxygen/fuel ratios closer to stoichiometry (Table 2), which implies less residual O2 in the flame to promote in-flight and post-impact oxidation. Parameter set #1 did not produce equally low oxidation in spite of its similarly low oxygen/fuel ratio because, since the flow rates of oxygen and kerosene are the highest, the thermal power in the HVOF gas jet and the resulting heat flux into the substrate are maximised: this presumably promoted post-deposition oxidation at lamellar boundaries, according to the mechanisms outlined previously. Fig. 7A effectively summarises the fact that the oxygen/fuel ratio (λ) is the primary, though not the only, factor controlling microstructural features such as decarburisation and oxide content. Similar conclusions

Fig. 6. (A) EDX spectra of the metal matrix between WC particles and of the decarburised area marked by label “3” in Fig. 5A, and (B) EDX spectra of oxides (splatboundary oxide marked by label “1” and intralamellar oxide marked by label “2” in Fig. 5A) compared to the overall spectrum of the complete coating.

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G. Bolelli et al. / Surface & Coatings Technology 206 (2012) 4079–4094 Table 4 Two-sample Student's t-test on hardness values. The null hypothesis corresponds to two populations having statistically identical average values.

Fig. 7. (A) Microstructural features (carbide retention, oxide inclusions) and (B) micromechanical properties (micro-hardness, elastic modulus and H/E ratio assessed by depth-sensing micro-indentation) as a function of the normalised oxygen-fuel ratio λ.

significance of the differences between average hardness values was checked by applying the two-sample Student's t-test to all pairs of hardness measurements (Table 4). The two-sample tests comparing the hardness values of coatings #1, #2 and #3 to those of coatings #4 and #5 return low probabilities (almost always around or below 0.05) that the “null hypothesis” (i.e. the hypothesis that the two

Fig. 8. Vickers microhardness and elastic modulus values assessed by depth-sensing micro-indentation on all HVOF-sprayed coatings deposited onto carbon steel substrate.

Sample pair

t Value

Degrees of freedom

Prob. > |t|

#1 #1 #1 #1 #1 #2 #2 #2 #2 #3 #3 #3 #4 #4 #5

− 0.41077 − 0.78002 − 4.64379 − 2.74808 − 8.79491 − 0.11684 − 2.63533 − 1.57368 − 5.42737 − 3.63208 − 1.97875 − 7.57162 1.43618 − 3.82645 − 5.05585

33 31 28 30 29 36 33 35 34 31 33 32 30 29 31

0.68389 0.44129 0.00007 0.01005 b 0.00001 0.90764 0.01271 0.12456 b 0.00001 0.00100 0.05624 b 0.00001 0.1613 0.00006 0.00002

vs. #2 vs. #3 vs. #4 vs. #5 vs. WC–CoCr vs. #3 vs. #4 vs. #5 vs. WC–CoCr vs. #4 vs. #5 vs. WC–CoCr vs. #5 vs. WC–CoCr vs. WC–CoCr

tested populations of hardness values have identical average) is verified. This corroborates to the previous assumption that samples #4 and #5 are significantly harder than the other WC–FeCrAl samples. Obviously, the probability of the “null hypothesis” also becomes negligibly small when comparing the hardness of all WC–FeCrAl coatings to that of the WC–CoCr reference. The behaviour of hardness and modulus data of WC–FeCrAl layers is better clarified by plotting the results of micro-indentation tests against the oxygen/fuel ratio λ (Fig. 7B). Although the statistical significance of the differences between the average hardness values is not guaranteed in all cases (as seen in Table 4), the graph shows that hardness and modulus (just like the oxide content and carbide retention index discussed in Section 3.1) primarily depend on λ: they exhibit opposite linear dependence from λ, namely hardness decreases while elastic modulus increases with increasing λ. When λ is higher, indeed, coatings are generally less decarburised but more oxidised. During the loading stage of depth-sensing indentation experiments, oxides, especially those located at splat boundaries, are fractured (cracks propagating along splat boundaries, following oxide stringers, are clearly seen in optical micrographs of Berkovich micro-indentations, Fig. 9A), allowing inelastic interlamellar sliding. This results in larger indenter penetration and, consequently, lower hardness. On the other hand, during the unloading stage, when the elastic portion of deformation is recovered, the extremely large stiffness of WC (modulus values > 500 GPa are usually found in the literature [37,38]) increases the modulus of the whole coating. Conversely, coatings sprayed with low λ values are generally less oxidised, which means better interlamellar cohesion and less irreversible inelastic deformation, but they are more decarburised, which means that the elastic stiffness is decreased. The reason why the differences between hardness values are less marked than those between modulus values (compared to the respective standard deviations) may lie in the fact that scattered defects (such as oxide inclusions at lamellar boundaries) affect hardness more than modulus. Depending on the presence or absence of such defects in the indented area, inelastic deformation during indenter penetration can occur to significantly different extents, causing larger scatter in hardness values, whereas the elastic recovery during unloading is comparatively less affected, resulting in lower data scatter. As a result of the trends outlined above, the H/E ratio between hardness (converted in GPa) and elastic modulus also decreases linearly with λ (Fig. 7B). The parameter H/E indicates the ability of the material to accommodate strains within elastic deformation regime, without reaching the yield or failure limit [39,40]; it has important implications for the mechanical and tribological behaviours of coatings and will be referred to in the forthcoming Sections 3.3 and 3.4

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Fig. 10. Plots of εφψ as a function of sin2ψ acquired on WC–FeCrAl samples #1 and #5 at φ = 0°, with linear fits to the experimental data points.

propagating along the weakest and most oxidised splat boundaries (Fig. 9C).

3.3. Residual stress

Fig. 9. Optical micrographs of Berkovich micro-indentations on the polished crosssections of WC–FeCrAl #3 (A) and WC–CoCr (B) coatings and of a high-load (10 N) Vickers micro-indentation on the polished cross-section of sample WC–FeCrAl #3 (C). The arrows indicate recognizable microcracks.

to assist in the interpretation of tribological behaviour and impact behaviour. The WC–CoCr reference seems comparatively tougher than the WC–FeCrAl coatings, since Berkovich micro-indentations performed on the former are definitely less microcracked (compare Fig. 9B to A). Quantification of the indentation fracture toughness of WC– FeCrAl coatings, however, is unfeasible: attempts to produce highload, microcracked Vickers indentations usable for this purpose [41] were frustrated by the irregular cracking behaviour of the material. Instead of developing radial cracks propagating from the indentation corners, which are required for the assessment of indentation fracture toughness, indeed, the WC–FeCrAl coatings exhibit multiple cracks

The plane stress assumption underlying the use of Eq. (1) for Xray residual stress analysis is confirmed by the remarkably good linearity of the εφψ vs. sin 2ψ plots (two of which are shown in Fig. 10). This indicates that no sub-surface shear stress gradients exist. As such gradients are typically induced by surface machining (e.g. grinding) [42], it may be assumed that the grinding and polishing operations performed on the present samples (as specified in Section 2.2) did not alter significantly the stress state of the material. Manual grinding with fine diamond papers and final polishing with diamond slurries should indeed ensure that no significant near-surface deformation is induced in the material. Accordingly, surface polishing of thermally-sprayed coatings was also performed in [43,44] before Xray residual stress measurement. Such polishing operation is important because the surface roughness of the as-deposited surface, which is comparable to the penetration depth of X-rays in the material, would significantly affect and alter the stress measurement [43]. All of the coatings (Table 5) exhibit compressive near-surface residual stresses. The values measured along the three φ directions do not differ much from one another, which means that the nearsurface stress state is substantially equi-biaxial. The values are quite consistent with those reported for other cermet coatings deposited by liquid-fuelled HVOF torches [9,45–48]. These torches indeed confer large kinetic energy to sprayed particles: compressive stresses are therefore caused by peening of previously deposited layers by new particles impinging at high velocity [9,49,50]. The fact that deposition-related phenomena (i.e. peening effects), not the thermal expansion mismatch between coating and substrate at the end of the

Table 5 Residual stresses in HVOF-sprayed cermet coatings. The elastic modulus values employed for the computation according to Eq. (1) are also shown in the table. Coating/substrate

E (GPa)

σφ = 0° (MPa)

σφ = 45° (MPa)

σφ = 90° (MPa)

#1/C40 #2/C40 #3/C40 #4/C40 #5/C40 #1/AA6082 WC–CoCr/C40

194 ± 22 226 ± 16 260 ± 15 176 ± 11 203 ± 18 194 ± 22 352 ± 23

− 260 ± 42 − 219 ± 29 − 165 ± 22 − 241 ± 38 − 331 ± 43 − 270 ± 47 − 415 ± 66

− 287 ± 44 − 242 ± 24 − 160 ± 16 − 255 ± 36 − 360 ± 45 − 291 ± 41 − 420 ± 49

− 288 ± 42 − 273 ± 28 − 148 ± 20 − 305 ± 29 − 406 ± 48 − 287 ± 43 − 453 ± 58

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deposition process, are responsible for the compressive stress state in WC–FeCrAl coatings is confirmed by the substantial identity of the residual stress states measured on coating #1 deposited onto two substrates with largely different thermal expansion coefficients (carbon steel and aluminium alloy: Table 5). As mentioned in Section 2.1, indeed, the deposition process was periodically interrupted in order to avoid system overheating, which keeps stresses due to thermal expansion mismatch down to a minimum. It is important to remark that coating material subjected to peening was exposed onto the sample surface after the outermost layer was removed by grinding and polishing operations: the very first surface layer, indeed, often has slightly tensile residual stress because of the absence of peening, as shown e.g. in [45,48–50]. Specifically, micrometre measurements indicated that approximately 60 μm of material were removed from the coating surface after grinding + polishing. Unlike the thickness values assessed by image analysis (listed in Table 3), which are referred to the mean surface line, micrometre measurements are referred to the highest surface peaks. Since all of the coatings possess Rmax ≈ 45–50 μm (assessed through optical confocal profilometry), grinding and polishing took the original surface roughness away entirely and further removed approximately 10–15 μm of material from the coating surface, which roughly corresponds to the layer deposited during one torch cycle. The overall thickness reduction of each coating was therefore of ≈35 μm. Measurements performed under these experimental conditions are certainly more significant than stress assessments on asdeposited surfaces (where no peening has occurred), not only because surface roughness may alter the stress measurement, as mentioned previously, but also because, in many practical applications (especially those involving sliding wear contacts), HVOF-sprayed cermet coatings are ground and polished in order to obtain smooth surface finish. More specifically, sample #5 exhibits the largest compressive stress (which is also qualitatively seen through the large slope of its εφψ vs. sin 2ψ graphs, Fig. 10): this is probably due to the fact that the intensity of peening is maximised by the combined effects of rather large particle impact velocity (quite large oxygen and kerosene flow rates were employed, Table 2, which accelerates the particles at large speeds [50,51]) and low oxide content (Section 3.1 and Table 3). The latter makes the deposited material more prone to be plastically deformed and, therefore, peened by new impinging particles. Differently from the properties examined in Sections 3.1 and 3.2, therefore, residual stresses do not depend only (or primarily) on λ, but on the combined effect of λ and of the total gas flow rate; the latter, indeed, has huge influence on in-flight particle velocity [50,51]. Accordingly, sample #4, sprayed not only with low λ value but also with low oxygen flow rate, exhibits lower compressive stress than sample #5. On the other hand, when oxidation becomes excessively large because of high λ values, the sensitivity of the material to peening is reduced; sample #3 indeed exhibits the lowest compressive residual stress of all WC–FeCrAl coatings. The thickness of the as-deposited coatings seems to have no obvious relation with the residual stress state. In general, the thickness of a thermally-sprayed coating can affect the residual stress state in two ways. On the one hand, as a coating grows thicker, the bending moment, which arises during the final cooling because of the thermal expansion mismatch with the substrate, becomes increasingly large. This bending moment causes a stress gradient across the coating, giving a tensile stress contribution near the top surface [52]. On the other hand, higher coating thickness implies (for a fixed number of torch scans) larger thickness deposited per torch pass: this is usually believed to increase the tensile quenching stress contribution [53]. Both of these factors, however, seem not to have significant effects in the present case. For instance, based on the above considerations, the thinnest coating (sample #3) should exhibit large compressive residual stress, since tensile stress contributions due to bending

moments and to quenching should be minimum. This sample, by contrast, is the one exhibiting the lowest compressive residual stress (Table 5). This confirms that the thermal expansion mismatch between coating and substrate does not have a major influence on the residual stress state of the present coatings, as observed previously, and that peening effects are largely preponderant over quenching stress contributions.

3.3. Dry sliding wear behaviour The sliding wear rates measured on the various WC–FeCrAl coatings after ball-on-disc tests (Table 6), roughly comprised between 5 ∗ 10 − 8 mm 3/(Nm) and 1 ∗ 10 − 7 mm 3/(Nm), exhibit limited, yet significant differences. The best wear resistance is shown by samples #4 and (particularly) #5: based on Sections 3.1–3.3, these coatings possess lower oxide content, higher micro-hardness and higher (H/E) ratio than the other WC–FeCrAl samples; sample #5 also exhibits the largest compressive residual stress. All of these factors probably concur to promote highest wear resistance. Accordingly, sample #3, which features low hardness, the lowest (H/E) ratio (Fig. 7B), high oxide content and the lowest compressive stress, also has the lowest wear resistance. This confirms that sliding wear resistance depends on a combination of these factors. The substrate material, on the other hand, has little or no influence on the sliding wear behaviour, as the wear rates of coating #1 deposited onto steel and aluminium substrates are very similar (Table 6). In order to provide an explanation for the dependence of sliding wear on factors such as oxidation, hardness, H/E ratio and residual stress, the wear mechanisms were investigated. All of the WC–FeCrAl coatings experience identical wear mechanisms (Fig. 11). On the one hand, single carbide particles are pulled out of the coating surface (Fig. 11B, label 1), since abrasion of the surrounding metal matrix leaves these particles progressively unsupported (as seen e.g. in Fig. 11C). On the other hand, the formation and the propagation of brittle cracks on the coating surface (as seen in Fig. 11D) lead to the removal of near-surface portions of material (Fig. 11B, label 2). Both carbide pull-out and brittle cracking are well known sliding wear processes for thermally-sprayed cermets (see [54–56] and [56–58], respectively), but the latter is probably the one responsible for most of the recorded wear loss. Brittle cracking is likely caused by the combination of hertzian surface stresses (under ball-on-plane configuration, tensile radial stresses are produced around the edge of the contact area) with frictional stresses: this combination indeed produces a maximum in the tensile normal stress at the trailing edge of the contact area [59], and the cyclic application of this stress at every disc revolution, causing fatigue loading, increases the severity of the contact condition [60]. Large compressive residual stresses can clearly hinder the formation and propagation of cracks by reducing the magnitude of the overall surface tensile stresses. Moreover, when the H/E ratio is

Table 6 Results of ball-on-disc dry sliding wear tests (sample wear rate), dry sand-steel wheel abrasion test (sample wear rate) and cyclic ball drop test (impact crater volume). Sample

Dry sliding Impact volume Abrasive wear rate [mm3] wear rate −8 3 [∗10 [∗10− 2 mm3/(Nm)] mm /(Nm)]

#1 C40 substrate 10.57 ± 0.97 AA6082 substrate 11.84 ± 2.83 #2 7.12 ± 0.49 #3 12.47 ± 2.43 #4 6.19 ± 1.75 #5 4.94 ± 1.29 WC–CoCr 3.94 ± 1.42

0.244 0.849 0.249 0.203 0.175 0.174 0.192

1.53 ± 0.07 1.46 ± 0.21 1.57 ± 0.03 1.63 ± 0.10 1.55 ± 0.06 1.20 ± 0.08

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Fig. 11. BSE-SEM micrographs of the wear scar produced after ball-on-disc dry sliding wear testing on the WC–FeCrAl #3 coating. A: overview; B: higher magnification (label 1 = single carbide pull-out; label 2 = near-surface delamination of material); C: detail of matrix abrasion between WC particles; D: detail of microcracked area (the arrows indicate a microcrack); and E: detail of a highly oxidised area with numerous microcracks (marked by arrows).

higher, the coating can accommodate larger near-surface strain before the elastic limit is reached [39,40]: since WC–FeCrAl coatings are quite brittle (as shown by indentation testing in Section 3.2), cracking, rather than plastic deformation, occurs when the elastic limit is reached, so that higher H/E ratio eventually means lower tendency to near-surface brittle fracture. Finally, brittle cracking is seen to be more frequent in areas where oxidation of the material is more pronounced (Fig. 11E: in the area shown in this micrograph, which exhibits many microcracks highlighted by arrows, a large number of dark oxide inclusions can be seen), so that coatings with low oxidation are less prone to cracking. As a result, the sliding wear resistance of sample #5, which exhibits the most favourable combination between microstructural features, micromechanical properties and compressive residual stresses, approaches that of the WC–CoCr reference (Table 6), although the wear mechanisms of the latter are somewhat different (Fig. 12). Matrix abrasion (Fig. 12C) with single carbide pull-out occur on the WC– CoCr layer as well; however, on a larger scale, the WC–CoCr coating does not undergo cracking and near-surface removal of material. By

contrast, “wavy” regions, clearly resulting from near-surface plastic deformation, appear (Fig. 12B). The orientation of these features (forming an angle with the plane of the original surface) indicates that deformation is induced by shear stress acting on the coating surface. It is inferred that shear stresses caused by the tangential friction force [59] deform the ductile CoCr alloy matrix, which drags the fine WC particles, giving rise to the observed undulations. An important difference therefore emerges between the tribological behaviour of the WC–FeCrAl layers and that of the WC–CoCr reference. The WC–FeCrAl coatings fail like brittle materials, with surface cracks induced by the maximum tensile stress, as seen in Fig. 11B, D, and E, whereas the WC–CoCr coating follows the failure criterion for ductile materials, with surface damage caused by shear stresses (the well known Tresca failure criterion for ductile materials). This difference is explained by considering that, on the one hand, the value of the H/E ratio is not larger in the WC–CoCr coating (H/ E = 3.8 ∗ 10 − 2, based on the data in Fig. 8) than in the WC–FeCrAl layers (H/E comprised between 4.2 ∗ 10 − 2 and 6.9 ∗ 10 − 2, as shown in Fig. 7B). The ability of the WC–CoCr coating to accommodate

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3.4. Dry particle abrasion resistance The wear rates produced in this test are much larger than those found in sliding wear tests (Table 6), which indicates greater severity of the particle abrasion tests over the sliding wear test. In this case, the abrasive wear rates experienced by the various WC–FeCrAl coatings do not differ significantly within experimental

Fig. 12. BSE-SEM micrographs of the wear scar produced after ball-on-disc dry sliding wear testing on the WC–CoCr coating. A: Overview; B: higher magnification of an area with numerous “wavy” features; and C: detail of matrix abrasion between WC particles.

surface strains without reaching the elastic limit is, therefore, not superior to that of the WC–FeCrAl coatings, but, on the other hand, when the elastic limit is reached, WC–FeCrAl layers crack almost immediately, as discussed previously, whereas the WC–CoCr coating can deform plastically without cracking, on account of its lower oxide content and better toughness (qualitatively discussed in Section 3.2). Its highly compressive residual stress state also helps in mitigating the maximum surface tensile stress. Though it is not central to the purpose of the present paper, it is interesting to compare the present observations on the sliding wear behaviour of the WC–CoCr coatings with previous studies by the same authors on similar coatings. The near-surface plastic deformability during dry sliding contacts indeed turns out to be a general characteristic of WC–CoCr coatings deposited by liquid-fuelled HVOF spraying. Identical “wavy” features related to surface plastic shearing, without brittle cracking, were accordingly reported on coatings deposited with GTV-K2 and Thermico-CJS torches [61,62]. This might be related to common features of coatings deposited by this kind of torches, including low oxide content, low decarburisation and compressive residual stress state [9,10].

Fig. 13. Secondary electron (SE)-SEM micrographs of the wear scar produced by dry particle abrasion testing on the WC–FeCrAl #4 coating. A: Overview (label 1 = abrasive groove; label 2 = indentation; label 3 = surface crack); B: higher magnification view (label 1 = abrasive groove; label 3 = delaminated area); and C: detail of the area circled in panel B (arrows: chipping of the material along the side of the abrasive groove).

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error (Table 6). All of the WC–FeCrAl samples experience a combination of abrasive grooving (ploughing/cutting: label 1 in Fig. 13A, B), particle indentation (Fig. 13A, label 2) and brittle cracking. Some cracks are clearly seen in the overview micrograph of Fig. 13A (label 3): they probably originate from the accumulation of elastoplastic deformation in the material subjected to repeated abrasive grooving and particle indentation. Cracking and chipping of the lateral ridges of an abrasive groove are indeed seen in Fig. 13C (arrows), which

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shows a detail of the circled area of the groove labelled “1” in Fig. 13B. As deformation accumulates because of repeated abrasion and indentation events, cracks grow on the surface and immediately below it, until an entire portion of material is delaminated. Accordingly, a region where delamination of near-surface material occurred as a consequence of surface and sub-surface crack propagation is recognisable in Fig. 13B (label “3”). The morphology seen in Fig. 13B is indeed remarkably similar to that produced by model experiments on brittle coatings subjected to multi-cycle scratch tests (as shown e.g. in [63,64]). The WC–CoCr coating, by contrast, possesses somewhat better abrasive wear resistance (Table 6). This sample indeed experiences no obvious microcracking and no brittle removal of material (Fig. 14A–C). The dominant wear mechanisms are abrasive grooving (Fig. 14A, B, label “1”) and indentation of loose particles (Fig. 14A, B, label “2”). Specifically, particle indentation seems much more frequent, which is probably a consequence of the higher toughness and ductility of this material, allowing penetration of the indenter with limited or no cracking, in accordance with the results of indentation testing discussed in Section 3.2. It would therefore seem that the resistance to cracking plays a more important role in determining particle abrasion resistance than it does in dry sliding wear conditions, where other factors (e.g. the H/E ratio, the residual stress) are equally influential (as discussed in Section 3.3). Toughness has accordingly been identified in the literature as an important parameter in controlling particle abrasion resistance of brittle materials [65], including cermet coatings [56,66]. The lower oxide content, better ductility and better resistance to interlamellar cracking of the WC–CoCr sample are therefore decisive in determining its better abrasive wear performance, whereas the slight differences existing between the properties of the WC–FeCrAl samples are presumably too small to be appreciated by this testing technique, which is probably not sensitive enough. 3.5. Cyclic impact behaviour

Fig. 14. particle groove; groove;

Secondary electron (SE)-SEM micrographs of the wear scar produced by dry abrasion testing on the WC–CoCr coating. A: overview (label 1 = abrasive label 2 = indentation); B: higher magnification view (label 1 = abrasive label 2 = indentation); and C: detail of abrasive grooves.

The impact volumes of the various samples deposited onto the steel substrate (Table 6) do not differ much from one another. The largest difference occurs between the impact volumes of coating #1 deposited onto steel and onto aluminium: under the present impact test conditions, deformation occurs in the coating as well as in the substrate, so that the overall system deformation depends much on the substrate properties. This means that all of the coatings deposited onto the same substrate material are subjected to similar deformations. Quite remarkably, the WC–FeCrAl coatings and the WC–CoCr reference exhibit very similar cracking behaviour, dominated by the formation of circumferential ring cracks along the periphery of the contact region (Fig. 15). These ring cracks propagate from the coating surface and extend across most of the thickness (Figs. 16A and 17A), although they do not reach the substrate interface. No cracks appear inside the coatings in the middle of the impact region (Figs. 16C and 17C), and no delamination from the substrate is visible (Figs. 16A, C and 17A, C). This cracking behaviour can be interpreted with reference to the spherical indentation models given by Chai and Lawn for hard coatings deposited onto ductile substrates [67]. According to Chai and Lawn, throughthickness cracking in the coating begins in the middle of the contact area (radial cracks) when the radius of the contact region (a) is comparable to the coating thickness (d), whereas it occurs preferentially along the periphery of the contact area (circumferential ring cracks) when d/a >0.1. In the present case, from optical profilometry measurements on the impact mark, a ≈ 1.7 mm→ d/a ≈ 0.1, i.e. the present test conditions are roughly at the boundary between indentation on intermediate thickness coatings and indentation on thin coatings. Accordingly, circumferential cracks occur preferentially, but they do not extend across the entire coating thickness, as it would be expected in the Chai and Lawn model for thin layers [67].

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The critical condition for the onset of circumferential ring cracking given by Chai and Lawn for thin coatings is therefore applicable to the present systems only with a certain degree of approximation; nonetheless, it can provide very useful indications in order to interpret the results shown in Figs. 15–17. According to this model [67], the critical load Pcrit for the onset of circumferential ring cracking is (Eq. (3)):

P crit ¼ A

Fig. 15. Optical micrographs of the surfaces of the WC–FeCrAl #1 (A) and WC–CoCr (B) coatings deposited onto steel substrate and of the WC–FeCrAl #1 coating deposited onto AA6082 substrate (C), after cyclic ball drop testing.

ES 2 S a : EC C

ð3Þ

Fig. 16. Cross-sectional BSE-SEM micrographs of the WC–FeCrAl sample #1 deposited onto C40 steel substrate, after cyclic ball drop test. A: Edge of the impact area, with cross-sectional view of circumferential ring cracks; B: detail of a circumferential ring crack; and C: middle of the impact area, with no visible cracks.

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This model is based on classical mechanic equations, not on fracture mechanics; therefore it neglects the effects of defects, pores, inclusions, etc. However, it is apparent from Eq. (3) that circumferential ring cracking depends on the matching between the elastic modulus of the coating and that of the substrate, and on the SC/EC ratio of the coating. Since the mechanical strength SC of a material is related to its hardness [68], the SC/EC ratio depends on the H/E ratio of the coating. Although the WC–FeCrAl coatings are more brittle and contain more inclusions (e.g. oxide inclusions) than the WC–CoCr one, which may favour the onset of cracking, their lower elastic modulus (Fig. 8) is better matched to that of the substrate materials (ES ≈ 210 GPa for steel, ES ≈ 70 GPa for aluminium alloys [69]) and their H/E ratio is larger. Consequently, the circumferential cracking behaviour of all cermet coatings (all subjected to analogous deformations, as noted previously) is similar. It is also observed that the deformation of the coated systems during ball drop induces normal stresses which are mainly parallel to the coating / substrate interface [67], i.e. these stresses lay in the plane of the lamellae and are directed parallel, not perpendicular, to the interlamellar oxide stringers. Consequently, these interlamellar oxides are not much affected by the deformation-induced stresses. Accordingly, circumferential cracks have rather straight propagation paths across the thickness of WC–FeCrAl coatings, without branching along interlamellar oxides, as seen in Fig. 16B. Interestingly, higher degree of crack branching occurs in the WC–CoCr coating (Fig. 17B). The circumferential cracks in this coating indeed develop numerous secondary branches which follow quite tortuous propagation paths: in low magnification micrographs, this appears as an apparent increase in porosity at the boundary of the impact region (Fig. 17A). This might be ascribed to the greater stiffness of this coating, which tends to develop larger stresses for a given imposed deformation. 4. Conclusions This research focused on the characterisation of WC–FeCrAl cermet coatings deposited by liquid-fuelled HVOF-spraying onto steel and aluminium alloy substrates. Their microstructural features, micromechanical properties, sliding and abrasive wear behaviour and cyclic impact resistance were studied as a function of the process parameters and were compared to the characteristics of a reference WC–CoCr layer, manufactured using the same technique. The main conclusions of this study are as follows:

Fig. 17. Cross-sectional BSE-SEM micrographs of the WC–CoCr coating deposited onto C40 steel substrate, after cyclic ball drop test. A: edge of the impact area, with crosssectional view of circumferential ring cracks; B: detail of branching in a circumferential ring crack; and C: middle of the impact area, with no visible cracks.

Where: ES, EC SC A

elastic modulus of the substrate and of the coating (respectively); mechanical strength of the coating; proportionality coefficient (in hertzian contacts, A = 2π / (1 − 2υS) with υS = Poisson's ratio of the substrate, but this formulation is not applicable to the present case, where significant plastic deformation of the substrate occurs).

• Deposition mechanisms of WC–FeCrAl powder particles involve inflight melting and homogenisation of the metal matrix, in-flight and post-deposition interaction with oxygen, as well as some decarburisation and dissolution of WC. Decarburisation decreases and oxidation increases as the oxygen/fuel ratio (λ) increases. • Micro-hardness, elastic modulus and H/E ratio also vary with λ (low λ = higher hardness, lower modulus, higher H/E ratio), because these properties depend on oxidation and decarburisation. • Compared to the WC–CoCr reference, the WC–FeCrAl layers are more oxidised; they possess lower hardness, lower modulus, but higher H/E ratio. Higher tendency towards interlamellar cracking is also qualitatively noted in WC–FeCrAl coatings. • The WC–FeCrAl coatings possess compressive residual stresses, whose magnitude seems to depend both on λ and on the total gas flow rate during spraying. WC–CoCr is also under compression. • Small but significant differences exist between the sliding wear rates of the various WC–FeCrAl coatings. Sliding wear resistance is enhanced by a combination of high compressive residual stresses, low oxide content, high H/E ratio. The sliding wear performance of the best WC–FeCrAl coating is comparable to that of the WC– CoCr reference. WC–FeCrAl layers are therefore a good alternative to WC–CoCr when sliding wear is involved (e.g. protection of shafts, bearings, seals, piston sleeves, etc.).

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• In dry particle abrasion, by contrast, no WC–FeCrAl coating approaches the performance of WC–CoCr, as their abrasive wear resistance is impaired by excessive oxidation-induced brittleness. WC–FeCrAl coatings are less suitable for operation under particle abrasion conditions (e.g. mills, crushers, conveyors for particulate materials, etc.). • The cyclic impact behaviour of the WC–FeCrAl coatings does not differ from that of the WC–CoCr reference. All coatings exhibit circumferential ring cracking with no interface delamination. Higher H/E ratio and lower elastic modulus compensate the greater brittleness of WC–FeCrAl by reducing stress concentration in the coating during impact events. Further developments of this research will include an assessment of the corrosion resistance of WC–FeCrAl coatings deposited onto different substrate materials and on their ionic release in liquid media of interest for possible applications to the food processing industry (as mentioned in the Introduction). Acknowledgements The authors are grateful to Mr. Ferri (Parma Spray Italia S.r.l.) for his precious assistance during coating deposition. Many thanks to Dr. Ing. Benedetta Bonferroni for her help with the experimental activities and to Dr. Miriam Hanuskova (University of Modena and Reggio Emilia) for the particle size distribution measurement. The research was funded by Centro Interdipartimentale INTERMECH MO.RE. located at the Faculty of Engineering “Enzo Ferrari”, University of Modena and Reggio Emilia and by Regione Emilia Romagna, Italy. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11]

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