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Wear 268 (2010) 1309–1319

Contents lists available at ScienceDirect

Wear

journa l homepage: www.e lsev ier .com/ locate /wear

ffect of microstructure on abrasive wear behavior of thermally sprayedC–10Co–4Cr coatings

anchan Kumaria,∗, K. Ananda, Michelangelo Bellaccib, Massimo Giannozzib

Materials Research Laboratory, GE India Technology Centre, EPIP-II, Whitefield, Bangalore 560066, IndiaGE Oil & Gas, Nuovo Pignone SPA, Via Felice Matteucci, 2 - 50127 Florence, Italy

r t i c l e i n f o

rticle history:eceived 12 December 2008eceived in revised form 27 January 2010ccepted 2 February 2010vailable online 10 February 2010

eywords:hermal sprayingC–10Co–4Cr

brasionear mechanismicrostructure

a b s t r a c t

Thermally sprayed WC–Co coatings have been widely used in the coatings industry for its superior slid-ing, abrasive and erosive wear properties. In applications where corrosion resistance is also required inaddition to wear resistance, WC–10Co–4Cr is the preferred coating composition. The coatings producedby different thermal spray processes exhibit a broad range of coating hardness, porosity and microstruc-tural features like grain size and volume fraction of individual phases. In this study, we have evaluated thecoating microstructures of various WC–10Co–4Cr coatings produced from different spraying processes,such as high velocity oxy-fuel (HVOF) and pulsed combustion. The objective of our study is to explorethe abrasive wear mechanism of WC–10Co–4Cr coatings in great detail, and determine how these mech-anisms are influenced by the coating microstructure. Dry sand rubber wheel abrasion test rig (based onASTM G65) is used for evaluating the three-body abrasive wear properties of the coatings, using aluminaas the abrasive material. The coating microstructural parameters including WC grain size, volume frac-

tion, binder mean free path have been quantitatively measured, and correlation between abrasive wearbehavior of coatings and its microstructural parameters is sought. This study shows that binder mean freepath of carbides (which is a function of WC grain size and volume fraction) is a very important parameteraffecting the abrasion resistance of good quality coatings (containing low porosity). The lower the bindermean free path, the higher is the abrasive wear resistance offered by the coating. This is because theabrasive wear mechanism of these coatings is dominated by preferential removal of the binder phase,

C gra

followed by pullout of W

. Introduction

WC–Co coatings have been extensively studied for the last twoo three decades because of their superior wear properties, in slid-ng [1–4], abrasion and erosive wear conditions [5–15]. They haveeen widely used for numerous industrial applications like air-raft, oil and gas, mining etc. in solving severe abrasion and erosionroblems. In addition to the coated form, they are also used in sin-ered form for structural applications, for making components likeutting tools, dies, plungers, gears, bearings etc., and their abra-ion behavior have also been widely reported in the literature10,16–18]. Wayne and Sampath [10] have looked at comparison oftructure/property relationships in sintered and thermally sprayed

C–Co materials, and have shown that the same set of equationsor abrasion and erosion resistance hold good for both sintered andhermally sprayed WC–Co materials, relating them to material’sardness, fracture toughness and the Co binder content.

∗ Corresponding author. Tel.: +91 80 4012 2566; fax: +91 80 2841 2111.E-mail address: kanchan.kumari@ge.com (K. Kumari).

043-1648/$ – see front matter © 2010 Elsevier B.V. All rights reserved.oi:10.1016/j.wear.2010.02.001

ins.© 2010 Elsevier B.V. All rights reserved.

Most of the thermally sprayed WC–Co coatings reported in theliterature are produced by either air plasma spray process (APS)or various HVOF processes such as JP5000, DJ, Top Gun etc., andthe coating microstructure and wear behavior are described. How-ever, there are very few papers that compare and contrast themicrostructure and wear behavior of coatings [5,6,8], which isaddressed in this paper. The coating microstructure is a function ofthe starting feedstock powder, thermal spray gun used for deposit-ing the coating, and the spraying parameters followed [5,13,19].The choice of the thermal spray gun largely governs (a) tempera-ture exposure of particles, (b) dwell time in the gun and (c) theirdeposition velocity. The effect of these parameters can vary theporosity levels in the coating, the extent of decarburization, thevolume fraction, grain size and distribution of carbides.

One of the important attributes of WC–Co coating microstruc-ture is the extent of decomposition of WC grains during spraying,

which is a strong function of the particle temperature in the flame,and can result in undesirable W2C and W phases, in addition to pri-mary WC phase in the coating. This results in further dissolution ofW and C in the binder, which is largely amorphous/nanocrystallinein nature, because of the very high cooling rates experienced by the

1 ar 268 (2010) 1309–1319

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Table 1Details of various coatings used in this study.

Coating Coating process Thickness (�m) Roughness (�m)

W1 HVOF (JP5000) 240 0.157

310 K. Kumari et al. / We

owder particles. This aspect of thermally sprayed WC–Co coatingas been extensively studied in the literature [6,8,9,20]. Hence, theesultant coating microstructure can have a much lower volumeraction of the wear resistant primary WC phase and a much higherolume fraction of the binder phase compared to the starting pow-er microstructure [1,6,8]. In another version of thermal spraying,alled cold spraying, where the flame temperatures are extremelyow (<500 ◦C), the starting powder characteristics like crystalline Coinder and fully retained WC phase can be preserved in the coatingery well [21].

WC–Co coatings derive its wear resistant properties from theresence of high volume fraction of hard, wear resistant WC grains

n a Co-based metallic binder phase. The presence of the metallicinder provides some toughness in the coating compared to pureeramic coatings, however, the binder can exhibit some brittlenessf high W and C are dissolved in the binder during spraying. Weave extensive literature on the abrasive wear behavior of WC–Cooatings [6,7,14,22], while very few of them discuss the abrasiveear behavior of WC–10Co–4Cr coating composition [5,13], which

s used in applications that demand some corrosion resistance inddition to wear resistance. In this study, we are interested invaluating this coating microstructure, produced by off the shelfommercially available thermal spray processes, from reputed ven-ors, for applications in equipments that pump or compress gases.e also studied a few coatings produced by the air plasma process

APS), but they were highly porous, and hence, we chose not toescribe them in this paper. Coatings produced by HVOF and APSrocesses have been studied extensively in the literature.

Often the end user does not have control over the coating pro-ess, which is largely controlled by the coating supplier. However,t may be possible to predict the wear performance of coat-ngs by evaluating the coating microstructures. The objective ofhis work is to explore the mechanisms of abrasive wear dam-ge of WC–10Co–4Cr coating in great detail and determine howhese mechanisms are influenced by the coating microstructure.n particular, it would be interesting to explore the scale of the

icrostructural damage relative to low dosage of abrasive par-icles to understand the wear mechanisms better. The scale ofhe damage zone relative to the size of the microstructural fea-ures, such as WC grain size, binder mean free path is worthnvestigating.

There have been several investigations on the abrasive wearehavior of WC–Co materials using micro-scale abrasion tests7,17,23], tests based on dry sand rubber wheel abrasion (ASTM65) [5,12–14,17,24], tests based on ASTM B611 rotating steelheel with curved vanes [6,18,25], pin/block-on-diamond disk

brasion tests [10,16], in either dry/wet conditions. Micro-scalebrasion tests uses an abrasive slurry, containing very fine abra-ives (alumina, silica, diamond etc.) in the size range of 1–10 �mn de-ionized water, to study the abrasion resistance of coatingsressed against a rotating steel ball. These tests represent field con-itions like in the mining industry where the coating is exposed tone abrasive contaminants, while the abrasive wear tests basedn ASTM G65 and ASTM B611 look at the abrasive wear behaviorf coatings using coarser abrasives like alumina, silica, typically inhe size range of 100–600 �m, which can be encountered for e.g.n the oil and gas industry. ASTM G65 based tests are used in thery condition, while the tests based on ASTM B611 uses coarserbrasives in a slurry medium. In the abrasion tests with diamondisk, the pin/block is coated with the desired WC–Co coating, andhe disk is typically covered with a resin bonded 30-�m diamond

late. Bozzi and de Mello [22] have looked at the abrasion wearehavior of WC–12% Co coating as a function of abrasive materialardness (using three different abrasives of silica, alumina and sil-

con carbide), while Stewart et al. [14] have examined the effectf particle size and nature of abrasive (alumina vs. silica) on abra-

W2 HVOF (JP5000) 220 0.118W3 Pulsed combustion 200 0.151W4 Pulsed combustion 270 0.093

sive wear behavior of conventional WC–17Co and nanostructuredWC–15Co coatings.

The coatings studied here had to meet two specific requirementsfor application in dry gas reciprocating compressors: (1) producelow sliding wear of the polymer piston rings and (2) resist abrasionby fine hard particles like alumina, silica, iron oxides etc., that canenter into the compressor, along with the gases. In this work, wereport the abrasion resistance of the various WC–10Co–4Cr coat-ings against fine alumina particles in the size range of 20–70 �m(average size 50 �m), since they have the highest hardness valueamong the various abrading foreign particles like alumina, silica,iron oxides etc.

In our study, we have used modified dry sand rubber wheelabrasion test rig (based on ASTM G65) to study the three-bodyabrasive wear behavior of WC–10Co–4Cr coatings using aluminaparticles. The main reason we have chosen alumina as the abrasiveis because alumina can be present in the gases to be compressed asa contaminant, for e.g. in propylene or it can enter during pressureswing adsorption of hydrogen separation/purification. Though theactual size of the foreign particles in the gases would be smaller(∼10 �m), we were forced to use 50 �m average size alumina par-ticles, since smaller particle sizes (alumina with average particlesize of 10 �m and 25 �m respectively that were also tested) hadthe tendency to clog the feeder and the nozzle of the machine.Hence, we have used the smallest size of the alumina particlesthat could flow continuously through the feeder of the dry sandrubber wheel equipment. We wanted to simulate the three-bodyabrasion condition between the coating, polymer piston ring andthe alumina particles in the dry condition, and hence, dry sandrubber wheel test rig was a good option to use. The material rubberis much closer to simulating the polymer material compared tosteel wheel or ball, as used in micro-scale abrasion tests based onASTM B611. In addition, both of these tests use abrasive in a slurrymedium, while our requirement was to test the coatings in thedry condition. Also, the preference for the rubber wheel abrasiontests over micro-scale abrasion tests comes from the fact that ourtests are more aggressive and produces significant wear depth,and hence, can differentiate various coatings much better.

2. Experimental procedure

2.1. Coating preparation

Thermally sprayed coatings of composition WC–10Co–4Cr,were obtained on 4140 steel substrates of different sizes:3 in. × 1 in. × 0.25 in. (abrasion test), 1 in. × 1 in. × 0.25 in. (X-raydiffraction, metallography), from reputed external vendors in theindustry, using HVOF and pulsed combustion processes. We haveneglected the coatings produced by the APS process in this studysince they had very high porosity. The coatings were all ground andpolished to obtain a surface roughness in the range of 0.1–0.2 �mRa, since the coating application required the coatings to have a

smooth surface finish. Higher coating roughness can result in muchhigher wear of the polymeric piston ring materials, and hence, allthe coatings were ground and polished. They had a coating thick-ness in the range of 200 �m after the finishing operation. Table 1shows a list of the coatings that are investigated in this study. Coat-

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ngs W1 and W2 were deposited by HVOF process using a JP5000un, while coatings W3 and W4 were deposited using pulsed com-ustion process, that are not that commonly used in the coatings

ndustry. The starting feedstock powder was broadly in the sizeange of −45 + 15 �m, and was prepared by agglomeration andintering process. However, the exact details of starting feedstockowder preparation, WC grain size in the powder, and sprayingarameters followed during the spraying process was proprietaryo the coating vendor.

.2. Coating characterization

Samples for metallography were prepared by cutting the smallerest samples in cross-section, mounting them in epoxy with vac-um impregnation, followed by emery paper grinding, diamondolishing and finishing with 0.03 �m colloidal silica suspension.he coating cross-sections were examined using Leica opticalicroscope, equipped with clemex image analysis software, and

EI Quanta 400 scanning electron microscope. From optical micro-raphs, coating thickness of various coatings were measured (asisted in Table 1), and porosity was determined using image anal-sis software from 500× images. A total of 10 micrographs weresed to calculate the average porosity for each coating. Coatingarameters like WC grain size, volume fraction and mean free pathere calculated from high magnification SEM images of coatings

ross-section in the SEI mode (7000× for coatings W1, W3 and W4,nd 20,000–25,000× for coating W2 as it contained much finer WCrains). The high magnifications were necessary to better revealhe WC grain boundaries. SE imaging in preference to BS imag-ng was used for getting the topographical contrast between the

C grains and the metallic binder. Quantitative metallography haseen obtained using the linear intercept method, as proposed byee and Gurland [26], and the equations used are described below:

1) Tungsten carbide grain size:

dWC = 2VWC

2N�� + N��(1)

2) Contiguity of tungsten carbide:

C = 2N��

2N�� + N��(2)

3) Binder mean free path:

� = dWC(1 − VWC)VWC(1 − C)

(3)

here VWC is the volume fraction of WC phase, (1 − VWC) is theolume fraction of the binder phase, N�� is the average number ofntercepts per unit length of test line with carbide–carbide inter-ace, and N�� is the average number of intercepts per unit lengthf test line with carbide–binder interface. SEM micrographs wererinted, and a large test grid of 18 vertical and 15 horizontal lineswith a grid size of the order of WC grain) printed on a trans-arency sheet was placed on top of the micrograph, to count theumber of intersections with WC grains. VWC is the ratio of theumber of intersections with WC grains to the total number of

ntersections. Similarly, N�� and N�� were calculated for each ofhe micrographs, and an average value of WC grain size, volumeraction and binder mean free path were calculated from a setf five representative micrographs for each of the coatings. For

oatings W1, W3 and W4 at 7000× magnification, a total area of2 �m × 41 �m was covered in each micrograph, while a muchmaller area of 6 �m × 6 �m for coating W2 at 25,000× was coveredn each micrograph. Though this represents a rather small portionf the total coated area, but various important observations can

Fig. 1. Schematic diagram of dry sand rubber wheel abrasion test apparatus (ASTMG65) [24].

be drawn using these microstructural parameters, and its effect onabrasive wear behavior of the coatings. We understand that theaccuracy of microstructural parameters can be improved by usinga higher number of micrographs, however, we limited ourselves to5 micrographs per coating.

Microhardness tests on coating cross-sections were performedusing a CLARK, CM-400 AT machine with a 300 gf load and a dwelltime of 15 s. Hardness values reported represent the average of10 individual indentations made on a coating cross-section. Forphase analysis of various coatings, a Philips X-Pert X-ray diffrac-tor with Cu-K� radiation was used. 2� scans were run between 20◦

and 90◦ using a step size of 0.002◦ and a dwell time of 10 s. Thesurface roughness of the various coatings was determined using aZeiss-TSK Surfcom 1800D surface profilometer, using a diamondtip of radius 2.5 �m. Both roughness and contour measurements ofcoatings could be performed using this equipment.

2.3. Abrasive wear testing

The abrasive wear resistance of the various coatings was evalu-ated in three-body abrasive wear condition using a dry sand rubberwheel abrasion test rig based on ASTM G65 [24]. Fig. 1 shows theschematic diagram of the dry sand rubber wheel apparatus, astaken from this ASTM standard. While this standard proposes to use200–300 �m size Ottawa silica sand as the abrasive, we have usedalumina particles of 50 �m average size, as per our requirementdiscussed in Section 1. The 3 in. × 1 in. face of the coated samplewas pressed against the rubber lined wheel, and is abraded by theflowing particles of alumina causing three-body abrasion of thecoating. Based on several trial experiments with various sizes ofalumina abrasive, to get a continuous flow rate of particles, alu-mina of average particle size 50 �m (with 90% of particles fallingin the range of 20–70 �m as per particle size analysis) was used forall the tests. Fig. 2 shows the morphology of the alumina abrasiveparticles, which are highly angular in nature, and can cause severeabrasion of the coatings.

Parameters like load and duration of test in the initial experi-ments were varied to get a measurable weight loss of coatings, anda load of 210 N and test duration of 10 min was finalized for all thetests. The rubber wheel was made of Neoprene rubber, and showeda hardness of Durometer A-60. The flow rate with alumina particles

was measured to be 134 g/min, and a test time of 10 min alloweda dosage of 1.34 kg of alumina to cause the wear of the individualcoatings. This is a very high dosage of abrasive, and is meant toproduce differential wear for the different types of coatings stud-ied. This resulted in a total sliding distance of 1440 m, with the

1312 K. Kumari et al. / Wear 268

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Fig. 2. SEM micrograph showing the morphology of alumina abrasive.

ubber wheel rpm of 200, and rubber wheel diameter of 228.6 mm,n 10 min test duration. No recycling of the abrasive took place andll tests were performed dry. The test samples were ultrasonicallyleaned in acetone before and after conducting the tests, and thebrasive weight loss of the coating was determined. An average of–4 samples were tested for each coating, and their mean valuesere reported as the abrasive wear data.

The worn surfaces of various coatings were examined by SEM

n the top surface, in the middle of the wear scars. Also, surfacerofilometer was used to study the worn surface roughness, andhe maximum depth of coating removal for each of the coatings.n order to understand the wear mechanisms better, a few exper-ments were done later on, for each polished coating, with a very

Fig. 3. Optical micrographs of WC–10Co–4Cr c

(2010) 1309–1319

low dosage (∼7 g) of alumina abrasive, which were quite useful toreveal the abrasive wear mechanism of these coatings.

3. Results and discussion

3.1. Microstructural characterization

The coating thickness for various coatings as measured from thecoating cross-section, are reported in Table 1. The table also showsthe surface roughness values for the various coatings in their as-received conditions, as measured by the surface profilometer. Mostof the coatings obtained had a thickness of ∼200 �m, and a surfaceroughness in the desirable range of 0.1–0.2 �m.

Fig. 3 shows the optical micrographs of the coatings cross-section. All of these coatings have quite low porosity. Coatings W1and W2, produced by JP5000 process, have a very high volume frac-tion of fine, uniformly dispersed WC grains with a low porosity (1%),while coatings W3 and W4, produced by the pulsed combustionprocess, though have low porosity, but have much lower volumefraction of WC grains which are not uniformly distributed in thecoating. It is interesting to note that the coating W4 show closeto zero porosity. So, we can see that the coating microstructure isvery much governed by the type of the thermal spray gun usedfor depositing the coating. In other words, coatings W1 and W2show a high retention of primary WC grains, while coatings W3 andW4 show a much lower retention of WC grains, and an increasedvolume fraction of the binder phase with dissolved W and C in it.Coating W2 shows much finer WC grains than observed in coat-ing W1. The porosity content of these coatings is determined using

clemex image analysis software on optical images, and is presentedin Table 2. Legoux et al. [5] have reported the microstructures of var-ious APS and HVOF WC–10Co–4Cr coatings, with porosities in therange of 5–16% for the APS coatings, and in the range of 0.6–8.5% forvarious HVOF coatings, by varying the spraying parameters. How-

oatings: (a) W1; (b) W2; (c) W3; (d) W4.

K. Kumari et al. / Wear 268 (2010) 1309–1319 1313

Table 2Microstructural parameters measured for various coatings.

Coating Porosity (%) WC grain size (�m) WC volume fraction Mean free path (�m)

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W1 1.1 0.63W2 1.0 0.33W3 1.4 0.75W4 0.2 0.96

ver, Liao et al. [6] have reported porosity in the range of 1% forhe APS coating, and 0.6% for the HVOF coating. Zhao et al. [13]ave reported porosity in the range of 0.2–4% for WC–CoCr coat-

ng deposited by HVOF (DJ 2600) process. The porosity values webtained are in the range of 1%, and hence, are dense, good qualityoatings.

For studying the coating microstructural features, such as grainize, volume fraction, binder mean free path etc., SEM micrographsf the coating cross-sections are also obtained at 7000×, except for2 coating, where 20,000–25,000× magnification was needed to

eveal the finer WC grains (Fig. 4).From this figure, we observe that coating W1 shows a very wide

istribution of WC grain size, which was not evident from the opti-al micrograph. Coatings produced by HVOF process, W1 and W2how high volume fraction of WC grains. It is interesting to see

hat coating W4, produced by the pulsed combustion process, hasbimodal distribution of WC grains (containing two sets of grain

izes; one in the range of a few microns, and another in the range of1 �m), which is resolved only by going to high magnification SEMmages. The finer WC grains could not be resolved from the optical

Fig. 4. SEM SEI images showing details of microstructure of W

0.61 0.270.73 0.080.12 3.670.40 0.96

micrograph. Hence, the total volume fraction of carbide grains incoating W4 is effectively much higher than in coating W3, sincecoating W3 has very low volume fraction of more or less uniformWC grains, in the size range of 1 �m. Microstructural parameters,such as grain size, volume fraction and binder mean free path, forthe various coatings have been obtained from several SEM images(Fig. 4), by using Eqs. (1)–(3), and the obtained values are tabu-lated in Table 2, which also show their porosity values. We see thatcoating W2 has an average WC grain size of 330 nm, which is thesmallest among the four coatings examined, and would fall in thecategory of nanostructured WC–Co–Cr coatings, which has beenwidely reported in the literature [14,19,27]. The measured volumefraction of the WC grains vary a lot among the various coatings:it is lowest for coating W3 (12%), followed by coating W4 (40%),followed by coating W1 (61%), and coating W2 shows the highest

volume fraction of 73%. The calculated WC grain size and volumefraction result in smallest binder mean free path of 0.08 �m forcoating W2, 0.27 �m for coating W1, 0.96 �m for coating W4, anda very large binder mean free path of 3.67 �m for coating W3. It isexpected that the coating mechanical properties like hardness will

C–10Co–4Cr coatings: (a) W1; (b) W2; (c) W3; (d) W4.

1314 K. Kumari et al. / Wear 268 (2010) 1309–1319

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Fig. 5. Microhardness data for various coatings obtained at 300 g load.

e a strong function of microstructural parameter like binder meanree path, since we know that WC phase hardness (∼2400–2500 Hv)s much higher than the hardness of the binder phase (<1000 Hv)ince it is metallic.

One of the other coating properties, fracture toughness, can alsoffect the wear performance of coatings, and it was of our inter-st to measure it. However, our coatings were relatively thinner∼200 �m thick), and the effort to produce indentation cracking onhe coating cross-section, for generating fracture toughness data,as not successful. For example, when we used a 5 kg load on coat-

ng W4 (the thickest of all the coatings observed), the indentationiagonal covered more than 60% of the coating thickness, withoutroducing any cracks. No cracking could be seen with the use of

ower loads. Hence, we were limited by the coating thickness tobtain any fracture toughness data on the coatings.

The microhardness of the various coatings are obtained from theoatings cross-section using a standard load of 300 g. Fig. 5 showshe average hardness for all the coatings, which is obtained from 10

easurement values. The minima and maxima values reported forach of the coating, is indicated by error bars. The coatings obtainedy JP5000 process, W1 and W2, are very much superior in theirardness compared to coatings produced by pulsed combustionrocess. As per the literature, most of the thermally sprayed WC–Coamily of coatings, sprayed by HVOF process have hardness in theange of 900–1330 Hv [5,6,13]. The hardness of most of our coat-ngs is in agreement with the reported literature. The large spreadn the coating hardness of thermal sprayed coating is a function ofhe non-homogenous coating microstructure. Among the coatingsroduced by the pulsed combustion processes, slight improvement

n the microhardness of W4 coating is seen compared to W3 coat-ng, since it has negligible porosity and higher volume fraction of

C grains. However, these coatings are lower in hardness com-ared to coatings produced by JP5000 process, and could performifferently in abrasion tests.

Fig. 6 shows the X-ray diffraction patterns for all the coatingstudied. The most important peak seen in these patterns is that of

C from the starting powders used for coating deposition. How-ver, the peaks for W2C and W are seen because of decarburization.he extent of decarburization is different in coatings obtained byifferent processes like HVOF (JP5000 gun) and pulsed combus-ion. Coatings obtained by JP5000 gun show low decarburization (asndicated by very small intensity W2C peak). However, the coatingsbtained by pulsed combustion process (W3 and W4) have signif-cant amounts of W2C phase, in addition to the primary WC phase.oating W3, also produced by pulsed combustion, shows the max-

mum extent of decarburization (significant presence of elementalpeak, in addition to W2C peaks), and the presence of WC phase is

educed to a minimum. This explains the microstructure observedn our coatings where we saw significant reductions in volume frac-

ion of hard, wear resistant phase for coatings W3 and W4, and aorresponding increase in the volume fraction of the binder phase.he absence of Co peak in these XRD patterns indicate that the Coinder is amorphous or nanocrystalline in nature.

Fig. 6. XRD patterns for various WC–10Co–4Cr coatings examined in this study.

There is vast literature available on the XRD patterns of startingWC–Co family of powders, and their thermally sprayed coat-ings [1,6,9,20], which confirms similar observations. It has beenreported [14,19,27] that WC–Co coatings from nanostructuredpowders have higher extent of decarburization. In these papers,the coatings are produced by Top Gun, DJ etc., but in our study,coatings are produced by JP5000 gun, and we find that coating W2containing 330 nm WC grains show slightly lower intensity of W2Cphase, compared to the coating W1 showing higher WC grain size(630 nm). This illustrates that the coating vendor has a very goodcontrol on the spraying parameters using the nanostructured pow-ders to control decarburization, and confirms to the possibility ofreducing nanocomposite degradation with liquid fuel systems likeJP5000 gun, as pointed out by Stewart et al. [14]. Yandouzi et al. [21]have shown the phase evolution of WC–Co based coatings obtainedby cold spraying process, where the flame temperature is very low(<500 ◦C), and hence the resulting coating has crystalline peaks ofCo binder, in addition to WC peaks. No W2C peaks are seen here,which suggests that decarburization is a result of the higher flametemperatures encountered in the thermal spray guns.

3.2. Abrasive wear behavior of coatings

The weight loss and standard deviation data for the variouscoatings tested in abrasion is presented in Fig. 7. We observe thatcoatings obtained by JP5000 process, W1 and W2, have much lowerweight loss in abrasion testing, compared to the coatings obtainedby pulsed combustion. We obtained a very low standard devia-tion for coatings W2 and W4, while coatings W1 and W3 showedhigh standard deviation, as indicated by error bars. The large dif-ference between the abrasive wear loss of coatings W3 and W4 canbe explained in terms of their volume fraction and mean free pathof carbide grains, as reported in Table 2.

Since the coatings were all ground and polished to a surfaceroughness of ∼0.1 �m, the corresponding coating wear loss val-ues were small, compared to the scenario if the coatings weretested in the as-deposited condition with roughness of 3–5 �m. Weobserved a wear coefficient of ∼2 × 10−5 mg/(N-m) for the coatings

obtained by JP5000 process when abraded by 50 �m alumina par-ticles, while Stewart et al. [14] have observed a wear coefficient of∼4 × 10−4 mg/(N-m) when abraded by 125–150 �m alumina parti-cles in their modified “dry sand rubber wheel technique”, since the

K. Kumari et al. / Wear 268

Fig. 7. Abrasive weight loss data for various WC–10Co–4Cr coatings tested with1.34 kg alumina particles.

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resistance of WC–Co coatings, which could also be attributed to

TA

ig. 8. (a) A typical abrasion tested sample with roughness measured along therrow. (b) Correlation between the abrasive wear loss of coating, and the roughnesseasured in the worn area.

oatings were tested in the as-deposited condition with a surfaceoughness of 5 �m.

The relatively poor performance of coating W3 could also beelated to the high extent of decarburization seen in this coatingrom XRD pattern (Fig. 6), which is a phenomena also reported byeveral other authors [6,14,27]. The insight brought in these paperss well as our study suggest that the key to improving the wearesistant properties of WC–Co based family of coatings is to haveigh levels of retained WC phase, in the well dispersed, fine mor-hology, which is best met with our coating W2. The very largepread in the abrasive wear loss data for coating W3 is because ofhe large variation seen in the coating microstructure, as seen inig. 3, where WC grain size distribution is not homogeneous. Theicrostructural regions, which are devoid of carbides, will offeruch lower abrasive wear resistance than the regions with high

olume fraction of carbide grains. The observed abrasive weightoss trend of these coatings follows the trend in coating hardness.

A typical abrasion tested sample is shown in Fig. 8a, whicheveals the wear scar of WC–10Co–4Cr coating after the rubberheel abrasion test with alumina particles. In the middle section of

he worn area (shown by arrow), surface roughness of the coatings measured using surface profilometry, and a correlation is soughtetween the measured roughness of the abraded area and the cor-esponding abrasive wear loss of the coating. This is shown is Fig. 8b.he data also includes a few other WC–10Co–4Cr coatings evalu-

ted by us that are not described in this paper. In general, coatingshich are smoother in the worn area after the abrasion testing

W1, W2 and W4) show a lower abrasive wear loss compared to theoatings which are rougher after wear. High roughness in the worn

able 3brasive wear loss and worn surface roughness of coatings, along with the parameters af

Coating Porosity (%) Binder mean free path (�m) Hardness

W1 1.1 0.27 1326W2 1.0 0.08 1369W3 1.4 3.67 1021W4 0.2 0.96 1092

(2010) 1309–1319 1315

area is seen for coatings, which either have high porosity, or havelow volume fraction of wear resistant carbides like coating W3, inwhich case, it is easier for the abrasive to preferentially removethe large volume fraction of the binder phase. Table 3 presents themeasured abrasive weight loss data obtained for the various coat-ings, the roughness of the worn area, along with the parametersthat affect abrasive wear loss of coatings.

The SEM micrographs of the worn areas of these coatings, fromthe middle of the wear scar, are shown in Fig. 9. We clearly see thedifferent scale of damage experienced by the various coatings afterabrasion by high dosage of 1.34 kg of alumina particles. Coatings W1and W2 have relatively finer WC grain size and a low binder meanfree path, so the abrasive is not able to sample the WC grains andthe binder separately. The WC grains in coating W2 are extremelyfine, and are not clearly visible at similar magnification. Hence, thedamage caused by alumina particles of average size 50 �m is min-imal. Both of these coatings contain very high volume fraction ofWC grains, and hence the damage to the softer binder phase is sig-nificantly reduced. However, the coatings W3 and W4 showed lowvolume fraction of WC grains, hence there are several pockets inthe coating, which are devoid of hard, wear resistant WC phase,and the binder is preferentially removed from these areas. If weexamine the worn morphology of coatings W3 and W4, we can seethat the binder is gouged much deeper for coating W3, hence, pro-ducing a more rougher surface and higher abrasive wear loss eventhough both of these coatings in this micrograph show decent vol-ume fraction of WC grains. However, the overall abrasive wear lossof coating W3 is much higher than that of coating W4, becausecoating W3 has only 12 vol.% of hard, wear resistant phase to resistabrasion, while coating W4 still has about 40 vol.% of hard phase,resulting in a binder mean free path of 3.67 �m for coating W3,compared to a binder mean free of 0.96 �m for coating W4. Coat-ings W1 and W2 have a mean free path of 0.27 �m and 0.08 �mrespectively, which are significantly lower than that of coatingsW3 and W4, and hence, can significantly resist abrasion. We canclearly see the WC grains standing in relief in the case of coatingsW3 and W4, which shows that they can resist abrasion by aluminaparticles without getting cracked, chipped off, rounded etc. This isbecause the hardness of alumina (1900–2000 Hv) is still lower thanthe hardness of WC grains (2400–2500 Hv) to cause any chippingor rounding of WC grains. Occasional WC grain pull out is noticedin these coatings, which would happen if the surrounding binderphase has been completely removed preferentially, and then it iseasier for the WC grains to get removed. A coating with high volumefraction of fine WC grains, resulting in low binder mean free path,will result in lower damage to the coating. This is well representedfor coating W2, having WC grains in the range of 330 nm, and a vol-ume fraction of WC grains in the range of 73%. Hence, coating W2has performed the best in abrasion tests, with a really low standarddeviation (Fig. 7).

It is well reported in the literature [6,14,27] that higher decar-burization of WC phase results in decreasing the abrasive wear

the poor performance of the coating W3 and W4 in that order, asdiscussed earlier, and supported by XRD results (Fig. 6). Anotherinteresting result is the relatively good performance of coating W4compared to coating W3, and its extremely low standard deviation,

fecting it.

(kgf/mm2) Abrasion loss (mg) Roughness of worn area (�m)

6.24 0.105.40 0.10

36.86 1.0713.59 0.22

1316 K. Kumari et al. / Wear 268 (2010) 1309–1319

ngs af

wtootfit

wtbftwbsfotwasmasowat

Fig. 9. SEM SEI plan view images of wear scars showing WC–10Co–4Cr coati

hich could be attributed to its porosity free coating microstruc-ure. It could have performed even better if we had higher retentionf WC phase in the coating microstructure. Hence, there is lot ofpportunity to control the spraying parameters during sprayingo obtain the desired coating microstructure with high retention ofne WC grains in the coating, with a very low binder mean free path,o improve the abrasion resistance of WC–Co family of coatings.

To summarize the wear mechanism observed by us in this workith 20–70 �m (average 50 �m) alumina abrasive, we can say that

he abrasion of these coatings occur by preferential removal of theinder phase, followed by WC grain pullout. The low binder meanree path (which is a function of the WC grain size and volume frac-ion) ensures that the abrasive is not able to produce significantear loss in the coating. Also, no scoring or ploughing marks has

een noticed for the various coatings since the size of the abra-ive is relatively small (50 �m). No trace of alumina debris could beound on the worn surfaces of these coatings. Particles size analysisf worn alumina particles confirmed that there is no size reduc-ion or wear of alumina particles during abrasion testing, sincee have used a compliant rubber wheel for our tests. Shipway

nd Hogg [17] have reported similar wear mechanism for WC–Cointered materials using alumina and silicon carbide abrasive inicro-scale abrasion testing. Chen et al. [7] have reported similar

brasion wear mechanism for WC–12% Co coatings under micro-

cale abrasion when tested with SiC abrasive. Liao et al. [6] havebserved micro-cutting and micro-ploughing as the main abrasiveear mechanisms, when WC–17% Co coatings were tested in wet

brasion with alumina of 100 �m size, and these disappear whenhe coatings are tested with lower alumina particle sizes, where

ter abrasion testing with alumina particles: (a) W1; (b) W2; (c) W3; (d) W4.

preferential removal of matrix around WC particles was observed.We would like to highlight the fact that even though the averageparticle size of alumina was 50 �m, but the particles showed highlyangular morphology with very sharp edges (Fig. 2), and hence thesmaller dimension of the particle that could cause abrasion of thecoating could still be in the range of 10–20 �m. Thus, the abra-sive wear mechanisms observed by us is similar to that observedin micro-scale abrasion reported by other authors.

In order to understand the abrasion wear mechanisms of thesecoatings better, a few experiments were done under similar testconditions, but with a very low dosage of alumina particles (∼7 g).This can enable us to capture the wear mechanism of our coatingsbetter. The SEM morphology of the four coatings, after low dosagealumina testing, is shown in Fig. 10. The damage zone for the coat-ing W3 is much larger than that of the remaining coatings, with aclear indication of WC grain pull out, since the surrounding binderphase has been completely removed. This represents a region inthe coating where the volume fraction of the WC grains is low,and hence, the preferential removal of the binder phase has causedsignificant damage to the coating. In the remaining coatings, thedamage zone is smaller but the wear mechanism is similar, whichis the preferential removal of the binder phase surrounding the WCparticles, followed by undermining of the WC particles leading toits pull out. Fig. 11 shows the SEM morphology of coating W1 at a

still higher magnification, which clearly shows the individual WCgrains, and way in which the alumina abrasive is able to squeezethe softer binder phase, and preferentially remove it. Hence, thesemicrographs with low dosage alumina particles clearly illustratethe wear mechanism of these coatings, and are quite insightful.

K. Kumari et al. / Wear 268 (2010) 1309–1319 1317

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ig. 10. SEM SEI plan view images of wear scars showing WC–10Co–4Cr coatings a4.

owever, the actual experiments with high dosage of alumina par-icles were still very useful to measure the relative weight loss ofhe various coatings, and their relative ranking in abrasion, whichould not have come from the low dosage experiments.

Based on the abrasive wear mechanism discussed above, weave tried to correlate the observed abrasion wear loss behaviorf coatings with respect to the binder mean free path, as shown in

ig. 12, since it is the binder mean free path which is responsible forhe differential in the hardness of coatings. We can see almost a lin-ar relationship between the abrasive wear loss of various coatingsnd the mean free path of carbide grains.

ig. 11. SEM SEI plan view image of coating W1 after low dosage alumina abrasivexperiment.

ry low dosage of alumina abrasive (∼7 g) experiments: (a) W1; (b) W2; (c) W3; (d)

We would like to bring out the fact that coatings with highporosity, which were also studied by us but not included in thispaper, showed 8–10 times higher wear compared to coatings W1and W2, even though it had much higher retention of WC grains.Porosity most adversely affects the abrasion behavior of coatings,however, if the coatings have low porosity values, then mean freepath of carbides mostly controls the coating abrasion behavior.Hence, one of the important ways of increasing the abrasive wearresistance of WC–Co family of coatings is to continuously decreasethe mean free path of carbides, so that the individual abrasive par-ticle sees a large volume fraction of hard, wear resistant WC grains,and very less area of the relatively softer binder phase, which iseasier to get removed by the harder abrasive like alumina. This canbe achieved by decreasing the size of the WC grains in the start-

ing feedstock powder, reducing decarburization of coating duringspraying by optimizing the spraying parameters, and the choiceof thermal spray gun that will favor minimum dissolution of WCgrains. In this regard, using JP5000 gun among the various HVOFguns available, which uses kerosene as the fuel unlike combustible

Fig. 12. Relationship between abrasive wear loss of coatings and their calculatedbinder mean free path.

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ases like hydrogen or propylene, as used in other HVOF guns,hould be preferred. This is because JP5000 gun offers lowest flameemperature and highest particle velocity among the various HVOFuns, leading to lowest residence time of particles in the flame andence, minimizing dissolution of primary WC grains [5].

Hence, based on our research findings, porosity of the coat-ng is the most important parameter that adversely affects theoating abrasive wear loss, and among good quality coatings withow porosity, it is the binder mean free path that largely governshe abrasion behavior of coating. Binder mean free path of coat-ng is a function of WC grain size and volume fraction, and hencehe way to improve the abrasion wear behavior of WC–Co fam-ly of coatings is to have a high volume fraction of fine carbiderain size. Thus, based on the coating microstructure, it is possi-le to predict the abrasion wear behavior of the coatings. Waynet al. [10,16] and Bolelli et al. [11] have explored the relationshipetween abrasive wear resistance of WC–Co coatings, and coatingarameters like, hardness, fracture toughness, volume fraction ofhe binder phase. It is reported that the higher the coating hard-ess and fracture toughness, the better is the abrasion resistance.he relationship with Co binder content is not that straightfor-ard. In our study, we measured the hardness and volume fraction

f the carbide/binder phase, but we were unable to measure theracture toughness of the coatings because of the limited coat-ng thickness, in the range of ∼200 �m. Our relationship is basedn microstructural parameters, which will ultimately govern theechanical properties of the coating like its hardness and fracture

oughness.In addition to the coating parameters explored in this study for

mproving the abrasion resistance of WC–Co coatings, one morearameter reported by Stewart et al. [28] is post heat-treatmentf the coating in the range of 250–600 ◦C. The post heat-treatmentmproves the coating integrity, reduces tensile residual stresses,mproves the bonding between the binder and the WC parti-les, and hence improves the abrasive wear resistance of coatingsy at least 35%. This appears to be an easy way to increase theear resistance of thermally sprayed WC–Co coatings, post depo-

ition.

. Conclusions

In this work, the microstructure of WC–10Co–4Cr coatingsbtained by HVOF and pulsed combustion process has beenescribed in great detail. All the coatings studied have very loworosity, however the coating microstructure is widely different.icrostructural parameters like WC grain size, volume fraction and

inder mean free path has been calculated for the various coat-ngs using linear intercept method. Low porosity in the coating isundamental requirement to achieving good abrasion resistance ofoatings. We have found that the coatings obtained by JP5000 gunave high retention of finely sized WC grains, leading to low binderean free path, while coatings obtained by the pulsed combus-

ion process show significant decomposition of primary WC grains,eading to presence of W2C and W secondary phases in these coat-ngs. This results in significantly lower volume fraction of primaryard, wear resistant WC grains, and a much enhanced volume frac-ion of the binder phase in the coating, resulting in higher binder

ean free path.The average hardness for coatings obtained by JP5000 process

s in the range of 1300–1370 Hv, while that of coatings obtained by

ulsed combustion process is in the range of 1000–1100 Hv, sincehe volume fraction of hard, wear resistant phase is much highern the case of coatings deposited by JP5000 process. The coatingardness is strongly influenced by the binder mean free path ofarbides.

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In modified dry sand rubber wheel abrasion tests using20–70 �m alumina abrasive particles, coatings obtained by JP5000process, with a low binder mean free path, showed significantimprovement in wear resistance than the coatings obtained bypulsed combustion, which had larger binder mean free path. Astrong linear relationship was observed between the abrasivewear loss of coatings and the binder mean free path. Coating W2,obtained with JP5000 process, showed a very fine WC grain size, inthe range of 330 nm, and a binder mean free path of 0.08 �m, andpresented the best abrasion resistance among the coatings tested,while coating W3, obtained by pulsed combustion process, with thehighest binder mean free path (3.67 �m), showed the worst perfor-mance. Surface roughness of coatings was measured in the wornarea, and strong linear relationship was again observed betweenthe abrasive wear loss of coatings and the measured surface rough-ness in the worn area.

The main abrasive wear mechanism of these coatings with20–70 �m alumina particles is preferential removal of the binderphase followed by pullout of WC grains, since the hardness of thebinder phase is expected to be <1000 Hv, while the hardness of alu-mina abrasive is in the range of 1900–2000 Hv. Hence, the key toimproving the abrasive wear resistance of WC–Co family of coatingsis to obtain a coating microstructure with high retention of fine, uni-form WC grain size, with minimum decarburization, resulting in alow binder mean free path of the coating. This can be achieved dur-ing spraying by using optimized feedstock powder, suitable choiceof the thermal spray gun, and optimizing the spraying parameters.

Acknowledgements

The authors would like to acknowledge the support receivedfrom several people from GE India Technology Centre, Bangalore,in completing this work: Kiran Joseph George for the wear experi-ments, C.H. Sathisha and T. Vishwanath for metallography samplepreparation, T. Shalini for the SEM work, and Srinidhi Sampath andPaul Mathew for various useful technical discussions, and criticallyreviewing this work. This work was supported by GE’s reciprocatingcompressors coatings development program.

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