ra

ag 3

ondWC(Eve wtere osmnd

hermalus studoice ofcoating

ort the plasticity during grit-blasting prior to ther

properties and abrasive wear resistance [8]. A microstructural and ana-

g and fracture [9].alling in thermal spraye coating and the sub-ay be either benecial

Int. Journal of Refractory Metals and Hard Materials 44 (2014) 6876

Contents lists available at ScienceDirect

Int. Journal of Refractory M

l seity by most binder metals. Cobalt is the most commonly-used binder ifabrasion resistance is required due to its excellent carbide wetting andadhesion properties [4].

or detrimental, depending upon the sign and distribution of stresseswith respect to some external factors (e.g. load).Montay et al. [11] iden-tied that the residual stress gradient in the depth of the top coat, bondCermet thermal spray coatings are widely used in industry due totheir characteristics such as resistance to abrasion, erosion, high tem-perature and corrosive atmospheres [3]. The choice of WC over othercarbides is prefer due to the good adhesive characteristic and wettabil-

adhesive and cohesive strength leading to spallinAmajor factor inuencing the fracture and sp

coating is the residual stress prole between thstrate material [10]. The residual stress effects m higher coefcients of thermal expansion than the coatingmaterial togenerate a limited degree of compressive residual stress in the coat-ing. An important step to know the particular substrate for a particu-lar engineering application is to understand more precisely how themicrostructure of the coating affects the coating properties.

lytical study of thermally sprayed WCCo coatings deposited unto aus-tenitic stainless steel substrate was done [9]. The results showed thatthe microstructure of the coating cross section comprised islands, elon-gated in directions parallel to the substrate. Despite these improve-ments, microstructure defects within the coating sometimes lower theRecently [2,3], many thermal spraying techspraying (APS), high velocity oxyl fuel (HVO

Corresponding author at: School of Chemical & Metalof the Witwatersrand, Private Bag 3, WITS, 2050, South A

E-mail address: seyiphilip@gmail.com (O.P. Oladijo).

0263-4368/$ see front matter 2014 Elsevier Ltd. All rihttp://dx.doi.org/10.1016/j.ijrmhm.2014.01.009mal spraying to attain ahanical interlock as well

and 17% cobalt WC/Co coatings deposited by VPS, APS, HVOF and DG(detonation gun), results had shown that HVOF and DG gave the bestcritical surface roughness, to provide mec

as increased contact area.1. Introduction

The choice of a metal substrate for tattention in the industrial world. Previoacterisation [1,2] has shown that the chbological applications of thermal spray

ability of substrate material to suppspraying needs adequateies based on global char-substratematerial for tri-s relies upon:

coating.

plasma spraying (VPS) have been applied to deposit WCCo coatings,although the coatings properties depend on spraying parameters.HVOF spraying is one of the best methods for depositing conventionalWCCo cermet powder due to the high velocities and lower tempera-tures experienced by the powder, which results in less decompositionof the WC during spraying [5,6]. In addition, it offers an effective andeconomic method of enhancingwear resistance without compromisingother attributes of the component. Mantyla et al. [7] have evaluated 12%Abrasive wearCorrelation between residual stress and abWC17Co coatings

O.P. Oladijo a,b,, A.M. Venter b,c, L.A. Cornish a,b

a School of Chemical and Metallurgical Engineering, University of the Witwatersrand, Private Bb DST/NRF Centre of Excellence in Strong Materials, South Africac Research & Development Division, NECSA Limited, Pretoria, South Africa

a b s t r a c ta r t i c l e i n f o

Article history:Received 2 September 2013Accepted 26 January 2014Available online 30 January 2014

Keywords:Residual stressHVOFWCCo coatingX-ray diffraction

This investigation had been ctance of HVOF thermal sprayelectric discharge machiningASTM-G65 three body abrasiby means of X-ray diffractionpressive residual stresses weation results after cutting intostresses in the surface layer a

j ourna l homepage: www.eniques such as air plasmaF) spraying and vacuum

lurgical Engineering, Universityfrica.

ghts reserved.sive wear of

, WITS, 2050, South Africa

ucted to determine the inuence of residual stresses on the abrasive wear resis-17 wt.% Co coatings, as well as to derive stress relaxation after cutting by wireDM). The abrasive wear properties of the coatings were characterised using anear machine with silica sand as the abrasive. The residual stress was measured

chniques, on the coated samples before and after the abrasive wear tests. Com-bserved in the surface layer of the large coated samples. However, stress relax-all sizeswere distinctly different. Therewas strong correlation between residualabrasive wear resistance, as well as yield strength of a material.

2014 Elsevier Ltd. All rights reserved.

etals and Hard Materials

v ie r .com/ locate / IJRMHMcoat and substrate came from three sources. The rst is due to thethermal and kinetic energy of individual sprayed particles that are inu-enced by the process itself. The second comes from the thermal expan-sion mismatch between the coating and the substrate. The third is thesubstrate preparation, which inuences the residual stress gradient.Venter et al. [12] investigated the residual strain associated with WC

17%Co coatings thermally sprayed onto metal substrates by means ofhigh energy synchrotronX-rays. He found that the large strainmist be-tween the coating and substrates emanated from the grit-blast surfacepreparation of the surface, extending down to 0.5 mm, whereas theresidual strain over the gauge volumes employed were low, owing tolocalised relaxation due to the microcracking.

X-ray diffraction was used to determine the residual stress in thiswork. This techniquewas chosen because of the following reasons [13]:

It is non destructive, and the same specimens can be used for otherinvestigations, or repeated measurements can be performed.

It is phase distinctive, and capable of stress investigation in eachphase.

It has a moderate restriction on specimen dimensions and shape,therefore, it can be used for measurements on different specimen

planes are parallel to the surface. The compressive stress observedwould not affect this lattice spacing, as they are acting parallel to thediffracting lattice planes. The sample has been rotated through aknown angle in Fig. 1b). The presence of compressive stress causesthe lattice spacing to be smaller than in the non-stressed state, whichcan be measured through determination of the shift in the diffractedintensity peaks. Once the shift is measured for at least two orientationsof y (the rotation of axis), then the lattice spacing, and hence residualstresses, can be resolved.

Using the Bragg equation:

n 2d sin 1

where = wavelength (nm), n = constant, = Diffraction angle(2 Theta in ).

otat

69O.P. Oladijo et al. / Int. Journal of Refractory Metals and Hard Materials 44 (2014) 6876sizes and components.

The residual stress investigation of this research work had been re-ported elsewhere [14], but inconclusive. This is due to the lattice spacing{WC 202}, as well as singled spot considered for residual stresses deter-mination on the large coated sample. Also, the correlation between theresidual stresses and abrasion resistance could not be nalised [16]since the abrasive wheel irradiated area was larger than initial singledspot. Thus, the objective of this work was to further explore the use ofX-ray diffraction to determine the residual stress associated with ther-mally sprayed WC17Co coating on metal substrates. The work wasalso done to investigate the inuence of residual strain on the abrasivewear resistance, as well as coating properties (yield strength, surfaceroughness).

2. Principles of X-rays diffraction

The principles governing the 2D detectors had been intensivelydiscussed elsewhere [14], whilst the principles concerning the sin2shall be discussed in details.

Residual stress magnitudes are determined through measurementof changes in thematerial lattice spacing, d-spacing, due to the presenceof a stress. Based on the knowledge of the non-stressed lattice spacing,any stresses present in the sample or material can be calculated usingthe established sin2 equation from Noyan [15]. In this method, acolliminated X-ray beam of wavelength similarly to the interplanarspacing is focussed onto a specimen and the number of X-rays diffractedis counted as the angle between the X-ray detector and X-ray tube isvaried [16]. This allows a plot of diffracted intensity versus 2Theta tobe achieved. From these peaks, the lattice spacing, which varies fromstressed to non-stressed material can be determined using the Braggequation. Fig. 1 indicates the impinging and diffracted X-ray beam ona magnied level. The angle is the angle between the surface normaland the bisector of the incident and diffracted X-ray beam. It is alsothe normal between the diffracted lattice planes and the sample's sur-face. Fig. 1a) shows the sample orientated so that the diffracting lattice

Fig. 1. Sample and laboratory coordinate systems [16]: a) = 0 and b) = i (sample is r

Surface normal.The strain is obtained as the change in d-spacings between crystallo-graphic planes and is given by:

ddodo

1 vE

sin2 v

E11 22 2

This predicts a linear variation of strain or interplanar spacing varia-tionwith sin2 so that stress can be obtained from the slope of a plotof strain vs Sin2. The geometry for the biaxial stress is shown in Fig. 1.where is the stress component along the SS direction in the plane,and is given by:

11 cos2 12 sin2 22 sin2d spacing in the direction defined by and m ;v Poisson ratio and E Youngs modulus Pa ; Surface stress defined by the angle Pa

3

A graph of d vs sin2 is plotted with the stress value determinedfrom the slope [15]. This approach is called the sin2 technique.

The slope m, is given by:

m 1 n =E d 4

The parametersv/E and (1 + v)/E are generally known as S1 and1/2S2 respectively, and are referred to as the X-ray elastic constants. Alinear behaviour of d vs sin2 is predicted by Eq. (2) [15].

3. Experimental procedure

A 2xxx Aluminium alloy, 304 L stainless steel, brass and super-invarof 75 25 10 mm3 in size were used as substrates. Wire EDM wasused to cut the larger coated samples into 30 25 10 mm3 pieces(i.e. small sample). The small samples were used for stress relaxationinvestigations. The residual stress analysis was taken at three differentpositions on the surface of the coated large samples (typical coated sam-ple used for abrasive wear tests discussed in previous publication [17]).

ed through some known angle i), where D= X-ray detector, S = X-ray source, and N=

and rotational wheel speed of 140 revs/min were employed on thelarge coated samples. Each coated sample was abraded for 30 min. A

in a

70 O.P. Oladijo et al. / Int. Journal of Refractory Metals and Hard Materials 44 (2014) 6876The sample sizes and measuring positions are shown in Fig. 2. Micro-structures investigation were done on 25 25 mm2 sample cut fromthe remaining procured substrate after the HVOF process, in order toprevent stress induced in the material needed for residual stressinvestigation.

Commercially available WC17 wt.%Co was used as a feedstock.The powder was received with a nominal size distribution of (45 +15 m) and was sprayed in this form. The surface of the substrate wasgrit-blasted with alumna before deposition. Coatings of about 200 mthickness were deposited by High Velocity Oxyl-fuel (HVOF) on 25 75 10 mm3 top surfaces for four different substrates namely brass,304 L stainless steel, super-invar and 2xxx aluminium alloy samples.These substrates were selected because they have coefcients of ther-mal expansion different from the coated material, and to each other toassess the residual stress associated with the each coefcient of expan-sion. Deposition parameters were the same on all coated samples, andwere as follows: 4 in. gun barrel; 380 mm spray distance; 0.0227 m3/hfuel (kerosene) ow rate; and 56.6 m3/h oxygen ow rate, thus similarto the parameters used elsewhere [14].

Fig. 2. Sample geometry and measurement position for the straResidual stressmeasurements taken on coating surfacewere done onan X-ray analyser with Co-K radiation, using the sin2 method. Theanode settings were 40 kV and 40 mA. For the WC coating, the {112}peak (reection at 123.62) that presented the highest diffracted intensi-ties were used. All other peaks had signicant overlap with uorescencepeaks from the high energy of the X-ray beam. The strainwere convertedto stress using bulk elastic constants S1=3.247 107MPa, 1/2S2=1.948 106 MPa and v = 0.20 [14]. Diffraction collection was doneusing a two-dimensional High Star (Bruker AXS) detector. Data wereanalysed using Leptos software version 6, as part of the Bruker AXSsuite of software for residual stress analysis. The full techniquesemployed to determine the residual stress measurement is describedelsewhere [14]. Similarly, coating characterisation had been reported

Table 1Material characteristics [6,17].

Substrate CTE [106/C] WC grain size (m)

Aluminium 2xxx series 23 0.150 0.01Brass 19 0.127 0.01304 L SS 17.3 0.185 0.02Super-invar 1.2 0.133 0.02

CTE = Coefcient of thermal expansion.detailed description of the abrasive wear test and analysis is given else-where [17] with the results listed in Table 2 for ease of reference.

The wear damage of the coating surface was further studied byatomic force microscopy (AFM) using a Veeco Dimension 3100 andXRD for qualitative analysis. While the state of abrasive silica sandafter post wear test was subsequently investigated using sieves analysismethod.

4. Results

The starting powder morphology and its cross-sectional part areshown in Fig. 3(a and b). The typical SEM/EDX image of the thermalelsewhere [6,17], which the results relevant to this research work listedin Table 1.

The wear analysis was done using three abrasive wear tests on aASTM-G65 dry sand rubber wheel apparatus. An applied load of 25 N

nalyses. Strain measurement was taken at positions 1, 2 and 5.spray WC17Co coatings showed that the light phase consisted of WCgrains and dark phase was the cobalt binder while the black portion isthe pore (Fig. 3(c)).

To understand the inuences of residual stress on the abrasive wearresistance, the residual stress was determined at three different posi-tions on large coated sample (typical sample used for wear test analysiselsewhere [17]). These positionswere chosen due to the abrasive wheelirradiated area being larger than initial spot employed in the previouspublished paper. The results are shown in Fig. 4, and the total averagevalues are given in Table 2. The residual stresses in WC phase over theentire sample were compressive, but quite different in the differentpositions, which implied that additional factors determined the stressresponse on the large coated substrates.

XRD results on coating Coating HV (GPa) Substrate HV5 (GPa)

WC, W2C, & Co 10.22 0.02 0.09 0.01WC 10.04 0.01 1.40 0.01WC 9.41 0.01 2.71 0.01W3C, Co 7.91 0.01 1.53 0.01

Figs. 5 and 6 show the correlation that exists between wear loss,yield strength and average residual stress determined on the large coat-ed samples, which was remarkable.

The residual stress measured on the worn surfaces of the abrasivecoated samples reported elsewhere [17], and their results were com-pared to the average residual stress on the large coated sample shownin Fig. 7. The compressive residual stress valuesmeasured on the coated

Fig. 10 shows the particle size distribution of the silica sand beforeand after wear test. After post wear test, the sand showed slightly little

The experimental results of the residual stresses found on the largeas-sprayed coated samples (Fig. 4) were compressive stress but quitedifferent from each other. The compressive stress in coated 2xxx alu-minium alloy and 304 L stainless steel increased, despite change thesample measuring position. However, the compressive stress on thesurface of coated super-invar decreases with respect to the measuringposition. The coated brass showed a distinct difference between the re-

Table 2Yield strength, abrasive wear results and average residual stress (i.e. total) of the large coated samples.

Substrate Aluminium 2xxx series Brass 304 L stainless steel Super-invar

Coating yield strength (GPa) [17] 2.41 2.37 2.22 1.87Wear mass loss (g) [17] 0.10 0.01 0.093 0.01 0.085 0.01 0.092 0.01Average residual stress (MPa) 224.23 42.3 149.3 42.7 93.06 43.8 78.86 41.2Surface roughness (nm) [6] 35 32 10 24

71O.P. Oladijo et al. / Int. Journal of Refractory Metals and Hard Materials 44 (2014) 6876changes distribution.In order to investigate the effect of stress relaxation in the coating.

The large coated sample cut by EDMwere re-examined by X-ray diffrac-tion. The residual stress determined under the assumption of planarstress conditions after cutting bywire EDM is shown in Fig. 11. After cut-ting, the small coated samples size showed a general trend of decreasingresidual stress. The coated 2xxx aluminium alloy and super-invar sub-strates consistently showed small changes. The coated 304 L stainlesssteel changed from compressive stress to tensile stress, while the stressof the coated brass sample was about a third of the large coated samplebefore cutting (but still compressive).

5. Discussion

The morphology of the dried starting powder (Fig. 3a) and its crosssectional images (Fig. 3b) was relatively simple. The powder showed adense structure with WC clusters cemented by cobalt, which wereboth identied fromXRD and EDX.Microstructural analyses of the coat-ing cross sections indicates that the typical thermal sprayed WCCocoating, comprised equiaxed carbide particles distributed in the Comatrix (Fig. 3c).

a bsample increased after wear test.XRD pattern of the worn coating surface are shown in Fig. 8. The

worn coating surfaces were largely dominated byWC peak as expectedwith a minor W3O peak, except worn coated brass which had SiO2 inadditions.

The AFM three dimensional views of the worn surfaces of WC17Cocoatings were shown in Fig. 9. The images showed similar imperfectionthat occurred on theworn surface after the abrasion. However, thewearfeatures were different among the substrates.Fig. 3. SEM/BSE images showing (a & b) WCCo starting powder and its cross sectional view, ssubstrates, showing WC grains (light), cobalt binder (grey), and black pores.sults measured on its surface, which might have been due to the tex-turing effect observed on its surface during measurement. Therefore,the non-uniform measurements observed on the coatings with re-spect to the measuring position might be due to the followingmechanism:

(1) porosity;(2) irregular distribution of powder resulting from spray kinetic

[18];(4) different grain size distribution;(5) shape of the coated samples;

In addition, the rate of cooling can signicantly inuence the residu-al stress or strain prole in the coating, which was not considered here,because the cooling rates were assumed to be same.

The average residual stress (i.e. total) on the large coated substratesamples are given in Table 2. The coated 2xxx aluminium sample hadthe highest compressive stress, and the residual stresses of the coatedbrass and 304 L stainless steel were in the middle, with the coatedsuper-invar having the lowest compressive stress value. The differencein residual stress results measured on all the coated samples, despitethe same WC17Co powder being used as a feedstock was deduced tobe due to their differences in the coefcient of thermal expansion, inter-action between coating and the substrate, as well as decarburisationprocess, which had been discussed by the current author elsewhere[6,14].

Many factors such as thermal history and deposition parametersmake it increasingly difcult to compare results of coating from othersources [14]. Stokes and Looney [19] determined the residual stressesin WCCo coatings by an analytical method similar to that of Clyneet al. [20], and reported tensile stresses of 82 MPa and 15 MPa for coat-ings of thicknesses of 200 and 600 m respectively. Ahmed et al. [10]investigated the inuence of vacuum heat treatment on the residualstress of thermal spray cermet coatings using neutron diffraction. They

chowing WC clusters embedded in a cobalt matrix (c) coatings on the 304 L stainless steel

observed the average values of the stress in the as-sprayed and heat-treated coating layers to be 553 MPa and 492 MPa respectively.The changes in the stress gradient after heat treatment, was related to

the compositional changes, caused by diffusion zones at the coatingsubstrate interface. The residual stresses of the samples investigatedwere much lower than those of Ahmed et al. [10], which could be dueto the differences in feedstock powder, process, substrate sample andsurface roughness. The aws in the coating [21], localisedmicrocrackingwithin the coatings, surface roughness, as well as surface stress relaxa-tion might also have effects.

It is thought that differences in the employed techniquesmight con-tribute to the lowvalues. Clyne et al. [20] highlighted some difculties inusing X-ray measurement to determine the coating stress, including:

(1) Limitation of the range of , since large values would require theincidence X-ray beam to penetrate appreciable thicknesses ofcoating and/or substrate.

(2) The penetration depth of X-ray diffraction is proportional to thesurface roughness of as-sprayed coating only, as well as

(3) Error from variation in stress levels.

However, the greater penetration of neutrons offers severaladvantages:

(1) Surface effects, such as deformation due to grinding or polishing,or oxidation are avoided:

(2) Complete rotation of the sample is achieved, giving more thor-ough sampling.

In general, the difference between the residual stress values ob-served and those in the literature depend not only on the techniqueand deposition parameters, but also on the method employed for

Fig. 4. Residual stresses taken at the three different positions on the large coatedsubstrates.

ba

72 O.P. Oladijo et al. / Int. Journal of Refractory Metals and Hard Materials 44 (2014) 6876c

Fig. 5. Correlation between compressive the residual stress and (a) yield strength, (b) wearmass loss, (c) surface roughness of the WC17Co coatings on the different substrates.

and (

73O.P. Oladijo et al. / Int. Journal of Refractory Metals and Hard Materials 44 (2014) 6876deposition, which can have new phases, resulting in residual stresses asthe volume changes due to different coefcients of thermal expansion[22].

There was a concern that the stresses were too small to have realeffect on the abrasion resistance and coating properties, althoughFig. 5 and Table 2 demonstrated strong correlations between yieldstrength, residual stress, surface roughness and wear resistance of theWC17Co coatings. The higher yield strength samples had higher resid-ual stresses, showing that the residual stress was proportional to theyield strength. Surface topography plays an important part in under-standing the nature of coatings [23]. It is thought that the surfaces onwhich thermal spraying is done have to be chemically andmechanicallywell-prepared for good adhesion. With an increased surface roughness,the residual stresses increased as expected, thus the residual stress wasproportional to the surface roughness.

In order to investigate the inuence of residual stresses on abrasivewear resistance of the WCCo coating, it is better to understand thesource of residual stresses in the abrasive wear measurement. It isthought that the residual stresses in the wear process comes from twomain sources [24]: (1) residual stresses, resulting from friction andwear process, that aid (tensile stresses) or retard (compressive stresses)the wear process; (2) induced residual stresses of the required sign(tensile or compressive) before the wear process or friction. Higherwear resistance is accompanied by increased compressive residual

Fig. 6. Correlation between wear mass loss and (a) coating yield strength, astresses, demonstrating that the abrasive wear resistance was propor-tional to the compressive residual stress, similar to the work of Garba[24]. However, the discrepancy of the as-coated super-invar was due

Fig. 7. Residual stresses of the coated WC17Co samples, in the as-coated condition, andafter wear testing.to partial stress relaxation, caused by the microcracking in the coating.The wear resistance of all coating increased with an increasing yieldstrength and surface roughness (Fig. 6).

In order to fully understand the stress state condition of the worncoating surface after abrasive wear test, it is better to investigate thestate of residual stresses after thewearmeasurement (i.e. most abradedpart of the sample). Fig. 7 and Table 2 shows the residual stresses of as-coated samples compared to after wear testing. The residual stressesthat were determined after wear test were compressive. It is thoughtthat the grinding of the coating by the abrasive sand under action offorce could have given rise to the plastic deformation and the introduc-tion of compressive stresses. In addition, the compressive residualstresses after the wear tests were higher than the as-coated samples(i.e. before the wear test). This is expected as the surface of all theworn samples had been altered by silica sand [17], as well as the inu-ence of material shakedown. Stoica et al. [25] investigated the residualstress and strain results of WC12Co coating in pre- and post tribologi-cal test. He found that the residual stress before and afterwear testwerecompressive, with the residual stress value slightly higher at the end ofwear test due to inuence of material shakedown.

The XRD analyses of the worn coating surfaces (Fig. 8) showed newphases as expected. All theworn coating surfaces comprises ofWC peakand W3O peak (minor), except worn coated brass which had SiO2 inadditions. It is thought that W O peaks could be formed as a result of

b

b) surface roughness of the WC17Co coatings on the different substrates.3

the interaction that existed between W (from the coating) and oxideresulting from their wear and friction. While SiO2 found on the worncoated brass is in agreement with the SEM/EDX result. Comparison ofthe XRD spectra of the as-sprayed coating [17] (with result listed inTable 1) and the worn coating surface showed greater reduction in theXRD intensity. This reduction might occur as a result of the residualstresses generated from the interaction of abrasive wheel, silica sandand the coating surface. Surprisingly, the initial secondary phases ob-served on the as-sprayed coating (Table 1) could not be found on theworn coating surface. It is thought that the secondary phases mighthave been abraded or eroded during the interaction of the silica sandon the coating. Although, the actual amount of the secondary phaseswas not determined, but their peak were all minor which can be easilyget abraded by the sand particles.

Atomic force microscopy (AFM) was used to obtained more detailsabout the wear mechanisms, as its give qualitative information on thebinder and carbide elimination. Inspection of these images (see Fig. 9)revealed that similar wear features were observed on the entire worncoating surface. However, the degree of damage or evolution wasseemed different to one another. The worn coated 2xxx aluminiumsample suffered the greatest damage due to extensive matrix and

30 40 50 60 70 80 90 100 110 120 130 1400

10000

20000

30000

40000

50000

60000

SiO2

WCWC WCWCW3O WCWC

WC

WCWC

WCWCWC

WC

WC

304L SS

Super-Invar

Brass

2xxx Al

Inte

nsity

Position (2 Theta)

Fig. 8. XRD pattern of the worn coating surfaces.

a b

c d

Fig. 9. AFM Three dimensional view of the worn surfaces of WC17Co coatings on (a) 2xxx aluminium (b) brass (c) super-invar (d) 304 L stainless steel.

74 O.P. Oladijo et al. / Int. Journal of Refractory Metals and Hard Materials 44 (2014) 6876

Fig. 10. Particle size distribution of silica sand before and after wear tests.

75O.P. Oladijo et al. / Int. Journal of Refractory Metals and Hard Materials 44 (2014) 6876carbide removal. Fewer binder removal with strong standing carbide inmost area unaltered were observed on the worn coated brass andsuper-invar, with extremely few found on the worn coated 304 L stain-less steel. The worn coated 304 L stainless steel showed limited binderremoval with high carbide grains (upper area) unaltered, while 2xxxaluminium reected the highest level of carbide grain and binderremoval. Thus in agreement with the SEM/EDX results discussed else-where [17]. Guilemany et al. [26] reported that the excellent wettingand adhesive characteristics of cobalt being binder make it difcult forthe carbide to be eliminated. Thus, a large quantity of binder materialelimination is necessary to begin removal of carbide. In addition, the dif-ferences in the degree of damage on all fourworn coating surfaces couldbe due to their differences in substrate hardness, surface roughness, andtheir physical properties that played a crucial role in cooling and solidi-cation of the coating after HVOF process.

The particle size distribution of the abrasive silica sand used wasfurther studies by Mabotja [27]. This was done to verify its conditionafter crushing the silica sand particle during the wear test. The results(Fig. 10) showed no signicant changes of the particle sizes distributionin comparison to before wear test. Thus, suggested that the silica sandcan still be used for more measurement, provided of the same proper-ties to avoid any contamination.

Some understanding of the stress relaxation may be gained bycutting the material into small sizes. After cutting the samples to pro-mote stress relaxation by EDM(Fig. 11), therewas little stress relaxationin the coated 2xxx aluminium alloy, brass, and super-invar sampleswhich are still relevant during service (i.e. still benecial for an engi-neering application). However, high stress relaxation was found in thecoated 304 L stainless steels, and it had changed to tensile residualstress, which can leads to cracking and promote fatigue failure if themagnitude exceeds the tensile strength of the coating [18]. A possibleFig. 11. Residual stress comparison in large sample before and after cutting to mediumsize.explanation for this behaviour is that melting was observed to be theprimary material removal mechanism during EDM machining of thesample. Thus, upon solidication, considerable thermal residual stresseson the surface layer occurredwhichmight inuence the performance ofthe service.

It is thought that thermal treatmentmay affect the stress state of thecoating [12]. Thus a better knowledge of the material behaviour at hightemperature (i.e. heat treatment) is therefore needed in our futurestudy to improve the quality of the coating during service.

6. Conclusions

The following conclusions were drawn based from the determina-tion of residual stress of WC7Co thermally sprayed onto differentmetal substrates. The surface residual stresses determined by X-ray dif-fraction are compressive and could be due to differences in coefcient ofthermal expansion and cooling rate. After cutting into small specimensvia wire EDM, the stress relaxation effect was high on coated brass, butmuch less on the other samples. Therewas correlation between residualstresses, abrasive wear and coating properties. While AFM gives moredetails information of the damage produced on worn coating surface.

Acknowledgements

The authorswish to acknowledge thenancial and technical supportreceived from the Department of Science and Technology in SouthAfrica, National Research Foundation, the Nuclear Energy Corporationof South Africa (NECSA), and the University of Witwatersrand. Theauthors also wish to thank Tshepo Ntsoane and Ryno van der Merwefrom Necsa for assistance with XRD and SEM.

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76 O.P. Oladijo et al. / Int. Journal of Refractory Metals and Hard Materials 44 (2014) 6876

Correlation between residual stress and abrasive wear of WC17Co coatings1. Introduction2. Principles of X-rays diffraction3. Experimental procedure4. Results5. Discussion6. ConclusionsAcknowledgementsReferences