Abrasive Wear Behaviour of WC–CoCr and Cr3C2–20(NiCr) Deposited By

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Surface & Coatings Technolog

Abrasive wear behaviour of WC–CoCr and Cr3C2–20(NiCr) deposited by

HVOF and detonation spray processes

J.K.N. Murthy, B. Venkataraman*

Defence Metallurgical Research Laboratory, P.O. Kanchanbagh, Hyderabad-500048, India

Received 18 June 2004; accepted in revised form 28 October 2004

Available online 24 December 2004

Abstract

Thermally sprayed tungsten carbide-based and chromium carbide-based coatings are being widely used for a variety of wear resistance

applications. These coatings deposited by high velocity processes like high velocity oxy-fuel (HVOF) and detonation gun spray (DS)

techniques are known to provide improved wear performance. In the present study, WC–10Co–4Cr and Cr3C2–20(NiCr) coatings are

deposited by HVOF and pulsed DS processes, and low stress abrasion wear resistance of these coatings are compared. The abrasion tests

were done using a three-body solid particle rubber wheel test rig using silica grits as the abrasive medium. The results show that the DS

coating performs slightly better than the HVOF coating possibly due to the higher residual compressive stresses induced by the former

process and WC-based coating has higher wear resistance in comparison to Cr3C2-based coating. Also, the thermally sprayed carbide-based

coatings have excellent wear resistance with respect to the hard chrome coatings.

D 2004 Elsevier B.V. All rights reserved.

Keywords: WC–CoCr; Cr3C2–20(NiCr); HVOF and detonation spray processes

1. Introduction

Thermally sprayed cermet coatings have emerged as a

viable solution for a wide range of wear resistance

applications to improve the service life of machine

components. Tungsten carbide and chromium carbide-based

coatings are frequently used for many of the applications in

gas turbine, steam turbine and aero-engine to improve the

resistance to sliding, abrasive and erosive wear [1,2]. The

former is used up to 500 8C and the latter up to 800 8C[3,4]. Also, for sliding wear and abrasive wear resistance,

the carbide coatings are considered to be a viable

alternative to hard chrome platings due to the strict

environmental regulations and cost concerns with regard

to the electroplating process [5,6]. These cermet coatings

0257-8972/$ - see front matter D 2004 Elsevier B.V. All rights reserved.

doi:10.1016/j.surfcoat.2004.10.136

* Corresponding author. Tel.: +91 40 24586476; fax: +91 40 24340683/

24341439.

E-mail addresses: [email protected],

[email protected] (B. Venkataraman).

are deposited by plasma spray and high velocity processes

namely high velocity oxy-fuel (HVOF) and detonation gun

spray (DS) processes. The high velocity processes namely

the HVOF and DS are usually employed for depositing

these coatings to avoid significant amount of reduction of

carbides to brittle carbides and oxy-carbides due to the

much lower temperature of the powder particles in the

exhaust gas stream and less in-flight time as compared to

that in plasma [7,8]. Also, the higher particle velocities in

the high velocity processes lead to better coating properties

like higher bond strength, density and lower oxide content.

It has been reported that carbide containing coatings

deposited by high velocity processes have good wear

resistance [9] compared to plasma-sprayed coatings due to

the better coating properties achievable in case of high

velocity processes as mentioned earlier. WC and Cr3C2

with different metallic binders like Co, Ni and Fe have been

studied using different amounts of binder contents with Co

and Ni most commonly used. Addition of Cr to the matrix

has been found to improve the wear and oxidation

resistance of these cermets [8,10]. The wear behaviour of

y 200 (2006) 2642–2652

Table 2

Spraying conditions adopted for HVOF and DS processes

Coating process Coating

SM 5847 D-3007

1. HVOF process

Spray gun DJ2600 DJ2600

O2 pressure (MPa) 1.14, 1.8 1.17, 1.92

O2 Flow rate (m3/h)

H2 pressure (MPa) 0.93, 3.6 0.96, 3.84

H2 Flow rate (m3/h)

Air pressure (MPa) 0.69, 1.2 0.55, 1.68

Air Flow rate (m3/h)

Spray rate (kg/h) 3.6 3.6

Carrier gas Nitrogen Nitrogen

Spraying distance (m) 0.2 0.2

2. DS process

O2/acetylene volume ratio 1:1.23 1:1.21

Carrier gas flow rate (m3/h) 1.6 1.2

Spray distance (m) 0.165 0.17

Frequency of shots (shot/s) 3 3

J.K.N. Murthy, B. Venkataraman / Surface & Coatings Technology 200 (2006) 2642–2652 2643

WC with varying amounts of Co content and Cr3C2–

25(NiCr) coatings deposited by different thermal spray

processes has been studied by various researchers [7,9,11].

It has been reported that the abrasive wear rate for the

cermet coatings is controlled by several factors like the

morphology of the starting powder, the size and distribution

of the carbide particles, hardness of the carbide particles

relative to the abrasive, properties of the matrix and its

volume fraction and the coating process, which determines

the coating characteristics like the phases, density, macro-

hardness and the residual stresses [11–13]. It has been

found that the Cr3C2–25(NiCr) coating has less wear

resistance compared to WC–Co system [14–16]. Several

mechanisms of material removal have been proposed—(i)

extrusion of the binder phase and removal by plastic

deformation and fatigue, (ii) undermining of the particles

and subsequent particle pull-out, (iii) microcutting, (iv)

carbide grain fracture and (v) delamination of the coating

[12,13,17]. In the present study, the abrasive wear

behaviour of WC–10Co–4Cr and Cr3C2–20(NiCr) coatings

deposited by HVOF and pulsed DS processes have been

compared. The wear performance of these coatings is also

compared with the hard chrome coating.

2. Experimental

2.1. Material

The WC–10Co–4Cr (Metco 5847) and Cr3C2–20(NiCr)

(Diamalloy 3007) powders were coated on mild steel

substrates. The coating powder characteristics as supplied

by the manufacturer are given in Table 1. Prior to the

coating, the mild steel substrates of dimensions 70�25�10

mm thick were ultrasonically cleaned with acetone and grit

blasted using Al2O3 grits on the 70 by 25 mm coating face

and again cleaned ultrasonically with acetone and dried.

The grit blasted substrate was held suitably in a fixture and

the coating deposition was carried out with the samples in

the stationary condition and the gun traversing to and from

to obtain the desired coating thickness. The above powders

were deposited by the high velocity processes namely

HVOF and DS. The spraying conditions adopted for the

two processes are given in Table 2. The coating

Table 1

Coating powder characteristics

Characteristic Coating powder

SM 5847 D-3007

Composition 10 wt.% Co 80 wt.% Cr3C2

4 wt.% Cr 20 wt.% (Ni20Cr)

86 wt.% WC

Particle size �53+11 Am �45+5.5 AmShape Mostly spherical Irregular

Manufacturing route Agglomerated/sintered Clad

thicknesses for the two coatings were between 250 and

350 Am.

2.2. Characterization

The following characterization tests were carried out:

1. X-ray diffraction (XRD) analysis of the powders and

the coating was done using a Phillips PW 1320

diffractometer with Cu-Ka radiation operated at 40

kV and 25 mA.

2. Surface roughness measurements of the coated surface

using a surface roughness tester (Make: Mitutoyo,

Model: Surftest 211). The cut-off length was 0.8 mm.

An average of five readings is reported.

3. Scanning electron microscopy (SEM) of the sectioned

and polished surface of the coating and also the worn

surface was obtained using a Phillips make SEM.

4. Porosity measurements were done using a Leitz micro-

scope fitted with a Biovis image analyzer, on the

sectioned and polished surface of the coating. Ten

readings were obtained and the average is reported.

5. Microhardness measurements were done on the sectioned

and polished surface of the coating with a Vickers

indenter at a load of 300 g using a Leica Microhardness

Tester (Model: VMHT Auto). An average of five

readings is reported.

6. Scratch test was also performed using a pin-on-disc

tribometer. The tribometer was modified to hold a

Vickers microindenter to perform the scratch test. The

test was done on a polished surface with a surface

roughness of Rab0.15 Am. Before each test, the coated

surface and the indenter tip were cleaned with acetone.

The axis of the indenter was normal to the coating

surface and a constant load of 10 N was applied during

Table 3

Test conditions

Normal load (N) 45

Wheel (rpm) 201

Total sliding distance (m) 8657

Wheel surface speed (m s�1) 2.4

Abrasive material Silica

Particle size range (Am) 150–300

Feed rate (kg/h) 19.32

J.K.N. Murthy, B. Venkataraman / Surface & Coatings Technology 200 (2006) 2642–26522644

the scratch. The scratch length was approximately 15

mm. The scratch test was done to simulate the behaviour

of the coating material to a single asperity contact.

2.3. Abrasion wear test

The coated samples of 70�25�10 mm dimensions were

tested using the well known dry abrasive rubber wheel

tester. The coated sample was mounted firmly in the sample

holder and was allowed to press against the rim of the

rubber wheel with the desired normal force by applying a

known dead weight using a lever arrangement. The dry

silica sand was then allowed to fall freely between the

wheel and the coated surface while the rubber wheel was

rubbing against the coated surface. The abrasive particles

used were not re-cycled. Fig. 1 shows the SEM micrograph

of the abrasive particles used in the present study. Prior to

the test, the coated sample was ultrasonically cleaned with

acetone, dried and weighed using an electronic weighing

balance (Make: Sartorius) with an accuracy of 0.01 mg.

The coating mass loss was measured at every 10-min

interval. The total duration of the test was 60 min. The

mass loss obtained was normalized with the coating density

to obtain the volume wear loss. The test conditions

followed are given in Table 3. The coating density values

were estimated by the coating weight gain method in which

the mass of the coating deposited was normalized with the

coating volume. The coating density values so obtained

ranged from 11,000 to 12,500 kg m�3 for WC–10Co–4Cr

and 6100 to 7000 kg m�3 for Cr3C2–20(NiCr) coating. The

average coating density values of 12,000 and 6400 kg m�3

for the tungsten carbide-based coating and chromium

carbide-based coating were taken, respectively. For com-

parison, abrasive wear test was also done on hard chrome

plating deposited on mild steel substrate. The hard chrome

plating was done by an electroplating unit using 1%

sulphuric acid bath. Standard procedures normally followed

Fig. 1. The SEM micrograph of the silica abrasive particles.

for commercial applications were adopted. The plating

thickness was approximately 175 Am.

3. Results and discussion

3.1. X-ray diffraction analyses

The X-ray diffraction pattern for the coating powders are

shown in Fig. 2a and b, respectively. In the case of WC–

CoCr, WC and Co were detected, whereas for Cr3C2–

20(NiCr), Cr3C2 and the binder NiCr were the major phases

identified. Fig. 2c shows the XRD pattern for as-sprayed

WC–CoCr coating by HVOF process. It shows partial

decarburisation of tungsten carbide to di-tungsten carbide

(W2C). Similar partial decarburisation occurs during DS

process as well shown in Fig. 2e. Decarburisation of WC to

W2C during the deposition process has been observed [8,18].

Distinct Co peak was not present, however, broadening of

the peak was observed. Probably the binder material is

present in amorphous/nanocrystalline form, presumably due

to the high cooling rates (typically 106–107 K/s) occurring in

such deposition processes upon impact of the molten

particles on the target material/sample. Also, the occurrence

of W and C in the binder phase due to the dissolution of WC

in molten Co during the deposition process has been reported

[18]. The XRD pattern for the as-sprayed Cr3C2–20(NiCr)

coating by HVOF and DS processes are shown in Fig. 2d and

f, respectively. Both show diffused X-ray diffraction patterns

with a number of overlapping diffraction lines of carbides—

Cr3C2, Cr7C3 (formed by decarburisation of Cr3C2) and

binder NiCr as reported earlier [9,19,20] for the composition

Cr3C2–25(NiCr). The only notable difference is the amount

of retained Cr3C2 phase, which is slightly higher for the

HVOF-sprayed coating. It has been reported that decarbur-

isation of Cr3C2 to Cr7C3 or Cr23C6 does not have a

detrimental effect on the wear resistance of the coating [9].

However, in case of WC-based coatings, the decomposition

may deteriorate the wear properties of the coating due to the

formation of brittle carbides and oxy-carbides [8].

3.2. Characterisation of the coatings

3.2.1. Microstructural characterisation

The SEM micrographs of the coating powders used are

shown in Fig. 3. The WC–CoCr powder, Fig. 3a, has mostly

Fig. 2. X-ray diffraction patterns of WC–CoCr and Cr3C2–20(NiCr) (a) and (b) coating powders, (c) and (d) HVOF coating, and (e) and (f) DS coating.

Fig. 3. SEM micrographs of the coating powders. (a) WC–10Co–4Cr and (b) Cr3C2–20(NiCr).


Table 4

Coating characteristics

Characteristics Coating

WC–CoCr Cr3C2NiCr

HVOF DS HVOF DS

Surface Roughness,

Ra (Am)

3.66F0.19 4.5F0.28 2.86F0.18 4.38F0.33

Porosity (%) 2.1F1.1 1.38F0.3 1.3F0.6 0.65F0.3

Microhardness

(HV0.3)

836F30 1096F50 880F30 894F35

Macrohardness

(HV10)

524F25 1007F35 715F20 810F25

Indentation fracture

toughness

(MPa m1/2)

3.1F0.4 4.12F0.4 2.77F0.3 3.4F1


spherical particles, whereas, Cr3C2–20(NiCr) powder par-

ticles (Fig. 3b) have irregular shape. The surface roughness

values of the as-sprayed samples are given in Table 4. The

DS-sprayed coatings in general resulted in slightly higher

surface roughness possibly due to the slightly higher particle

velocity, which causes more particle deformation after impact

[21]. The porosity measurements for the tungsten carbide-

based and chromium carbide-based coatings deposited by the

high velocity processes are given in Table 4. It shows that the

DS coating results in a slightly lower porosity than that of the

HVOF coating. This is evident from the SEM micrograph of

Fig. 4. SEM micrographs of the transverse section of the coatings. (a) and (b) are

are for Cr3C2(NiCr) coated by HVOF and DS processes, respectively.

the transverse section of the coatings deposited by HVOF and

DS processes shown in Fig. 4. The DS-coated samples appear

to be denser compared to the HVOF-coated samples.

3.2.2. Hardness measurements

The microhardness measurements show that DS-coated

samples result in slightly higher hardness values (Table 4).

The WC–CoCr coating deposited by HVOF process resulted

in the lowest microhardness, possibly due to the higher

percentage of porosity. The macrohardness measurements on

the coating (top surface) using Vickers indenter with a normal

load of 10 kg was also carried out as it gives an indication of

the denseness of the coating. The hardness values are given in

Table 4. The macrohardness values were less than the

microhardness values. Such differences have been reported

earlier in case of thermally sprayed coatings [22]. This has

been attributed to the planar pores (pores parallel to the

coating–substrate interface) and the microcracks within the

coating. During microindentation the deformation is highly

localized, whereas in indentations at higher loads, the

influence of planar pores and cracks are more pronounced,

thus resulting in lower macrohardness values. This is

reflected in the hardness values for HVOF deposited WC–

CoCr coating which has higher porosity content that causes

substantial reduction in macrohardness value. The SEM

micrographs of the indentations produced by macroindenta-

for WC–CoCr coated by HVOF and DS processes, respectively; (c) and (d)

Fig. 5. SEM micrographs of Vickers indentations on the coated surface at a load of 98.4 N. (a) and (b) are for WC–CoCr coated by HVOF and DS processes,

respectively; (c) and (d) are for Cr3C2(NiCr) coated by HVOF and DS processes, respectively.


tion are shown in Fig. 5. It can be seen that edge cracks are

formed in all the coatings. Such cracks are normally observed

in thermally sprayed coatings as the cracks at the edges are

formed by the coalescence of microcracks originating at the

pores. In the case of HVOF WC–CoCr (Fig. 5a) coating, the

edge crack intensity surrounding the indentation was higher

probably due to the higher porosity content in the coating.

3.2.3. Indentation fracture toughness measurements

The indentation technique was used to obtain the fracture

toughness of the coatings using a Vickers indenter. The

Fig. 6. The indentation cracks induced in WC–CoCr

indentation was carried out on the transverse section of the

coating in the mid-plane region to minimize the edge and

interface effects [23]. The indenter was loaded such that one

of the horizontal diagonals was parallel to the coating–

substrate interface. A load of 2 kg was applied for a dwell

time of 25 s at a rate of 25 Am/s. Figs. 6 and 7 show the typical

indentations on the transverse section with in-plane cracks for

WC–CoCr and Cr3C2–20(NiCr) coatings respectively. In the

thermally sprayed coatings, the cracks parallel to the coating–

substrate interface are more easily formed in comparison to

the perpendicular direction [24,25]. This has been attributed

coating: (a) HVOF coating and (b) DS coating.

Fig. 7. The indentation cracks induced in Cr3C2–20(NiCr) coating: (a) HVOF coating and (b) DS coating.

Fig. 8. The indentation crack propagation path in (a) DS-coated WC–CoCr

and (b) HVOF-coated Cr3C2–20(NiCr).


to the characteristics of the thermally sprayed coatings [24].

The coatings are built to the desired thickness by the

deposition of molten or semi-molten particles in the form of

plate-like structure called splats. The weak bonding between

the splats results in such an anisotropic crack formation. In

some indentation tests, the cracks were not formed. As can be

seen (Figs. 6 and 7) in most of the cases, the cracks were not

initiated at the corners of the indentation where the stress is

highest but are seen originating from the sides of the

indentation, this has been attributed to the non-uniform

microstructure in the as-sprayed coatings [23]. The crack

length from the center of the indent, c, was used for

determining the fracture toughness of the coatings. The

fracture toughness values of the coating were calculated

according to the method suggested by Evans and Wilshaw

[26]. The length of the cracks was measured from the SEM

images. The Table 4 shows fracture toughness values

(average of 10 readings) obtained for the coatings. The

fracture toughness values showed some scatter for the two

coatings indicating non-homogeneous coating microstruc-

tural features. Such a variation in the fracture toughness

values in thermally sprayed coatings has been observed [23].

In case of WC–CoCr HVOF coating, the crack lengths could

not be measured accurately due to the interference of the

porosity. WC–CoCr coating had marginally higher fracture

toughness compared to Cr3C2–20(NiCr) coatings. Further,

examination of the crack features show that the indentation

crack propagates along a region between the carbide particles

and the binder phase in case of both WC–CoCr and Cr3C2–

20(NiCr) coatings as shown in Fig. 8a and b, respectively.

Similar observation on the indentation crack propagation has

been reported earlier in case of WC–CoCr coating [25].

Further investigation is currently being carried out in our

studies to understand the metallurgical characteristics influ-

encing the crack propagation path.

3.2.4. Scratch behaviour

To understand the wear behaviour when a single abrasive

particle rubs against the coating surface, a pin-on-disc

tribometer was used. It was modified to hold a Vickers

microindenter to perform the scratch test at a constant load.

This may be an over-simplification considering the fact that

the indenter used is diamond whose hardness is significantly

higher compared to the coating materials and acts as a hard

abrasive resulting in severe wear conditions. However, this

test may give a qualitative understanding of the response of

the coating materials under identical test conditions

employed. Fig. 9 shows the groove morphology formed

Fig. 9. SEM micrographs of the surface scratched by Vickers microindenter at 10 N load. (a) and (b) are for WC–CoCr coated by HVOF and DS processes,

respectively; (c) and (d) are for Cr3C2(NiCr) coated by HVOF and DS processes, respectively.

Fig. 10. Volume wear loss as a function of time. The filled symbols

represent HVOF process and the unfilled symbols DS process.


along with the wear debris scattered on either side during the

scratch test. There is a good correlation between the scratch

width and the hardness of the coating. As the indenter

scratches the surface, material is displaced to the sides and

detached. Some of the displaced material will be accom-

modated below the indenter by the porous volume and also

possibly by the surrounding elastic strain field. The displaced

volume of material by the indenter due to ploughing forms

side ridges, which may consist of the binder phase with few

carbide particles. Due to the near zero hydrostatic compres-

sive forces on the side of the indenter at the free surface,

plastic strain in the material flowing on the sides reaches a

high value [22,27]. Beyond a certain critical strain, the

material detaches to form debris. Here, the material detach-

ment is observed to occur in a single pass for both the coated

samples tested. The morphology of the debris also shows the

evidence of microcutting when the indenter (hard abrasive)

scratches the surface. The displacement of the material to the

sides is slightly more in the case of DS-coated samples

possibly due to the slightly higher density of the DS coatings.

3.3. Abrasive wear behaviour

The incremental volume wear loss of the coating as a

function of time and the cumulative wear loss of the coatings

are shown in Figs. 10 and 11, respectively. It shows that WC–

CoCr coating has a lower volume wear loss compared to

Cr3C2(NiCr) coating. The results show that the abrasive wear

volume was slightly lower in the case of DS coating

compared to that of HVOF coating. The SEM micrographs

of the worn surface of the coatings are shown in Fig. 12. In the

case of WC–CoCr coating (Fig. 12a and b), mechanism of

wear was by selective removal of the binder caused probably

by plastic deformation and fatigue due to the repeated action

Fig. 11. The cumulative volume wear loss of the thermally sprayed coatings in comparison with the hard chrome plating.


of the abrasive particles followed by the undermining of the

carbide particles resulting in their eventual pullout. Some

evidence of microcutting may also be noticed indicating the

removal of the binder phase by this mechanism. Also, very

Fig. 12. SEM micrographs of the worn surfaces. (a) and (b) are WC–CoCr coating

coating by HVOF and DS processes, respectively. The direction of abrasion is fr

little carbide grain fracture was observed. In the case of

Cr3C2(NiCr) coating (Fig. 12c and d) in addition to the above

mechanism of wear, material removal by delamination in

both HVOF- and DS-coated samples are observed. The

by HVOF and DS processes, respectively; (c) and (d) are Cr3C2–20(NiCr)

om top to bottom.


transverse section of the worn region was examined to

observe the sub-surface damage in both the coating systems.

The transverse section of the central area of the worn region

along the direction of sliding is shown in Figs. 13 and 14. In

the case ofWC–CoCr (Fig. 13), there is little evidence of sub-

surface damage. However, in very few regions, some

microcracks were observed below the surface running

parallel to the surface. With Cr3C2–NiCr coating (Fig. 14),

material removal by sub-surface crack formation and its

propagation to the free surface to form large chunks of debris

is evident. Such sub-surface cracking was more extensive for

Cr3C2–20(NiCr) coating. Such wear behaviour was common

to both the high velocity coatings as seen in Fig. 14.

The higher wear resistance of WC–CoCr coating over

Cr3C2–20(NiCr) may probably be due to the higher fracture

strength and better adhesive strength of the CoCr matrix

with the carbides compared to the NiCr matrix. The slightly

better indentation fracture toughness of the WC–COCr

coating observed may support this behaviour. The delami-

nating-type cracking has been observed earlier in case of

Cr3C2–25(NiCr) coating, indicating that the matrix phase is

not acting as an efficient toughening phase [14,28]. Such

Fig. 13. SEM micrographs showing SE images of the transverse section of

the worn region for WC–CoCr coating by (a) HVOF process and (b) DS

process.

Fig. 14. SEM micrographs showing SE images of the transverse section of

the worn region for Cr2C3–20(NiCr) coating by (a) HVOF process and (b)

DS process.

type of wear results in the removal of much higher material

thus leading to higher wear rates. This may explain the

better wear resistance offered by HVOF WC–CoCr coating

even in the presence of slightly higher porosity. Also, the

higher hardness of WC grains compared to the Cr3C2 grains

will enable the former to withstand higher loads which

results in fracture of less number of WC grains.

The residual stress measurements by X-ray diffraction

technique was also carried on the WC–CoCr coating. The

Bragg reflection from the (256) planes of WC was used for

determining the residual stresses. More detailed description

on the residual stress measurements is given in our earlier

paper [29]. The results show that compressive residual

stresses are induced in the coating deposited by both the

processes. In fact, higher compressive residual stresses were

induced in the DS coating (�104F20 MPa) compared to the

HVOF coating (�25F10 MPa) probably due to the higher

particle velocities in the former process. This may also

explain the lower wear rate observed in the DS coating

where the compressive stresses impedes or delays the crack

initiation and propagation, thus generating less wear debris.


Abrasive wear test of hard chrome plating was also carried

out for comparison with the thermally sprayed coatings. The

hardness of the hard chrome plating was 871HV0.3 (an

average of five readings). The results (Fig. 11) show that the

carbide-based thermally sprayed coatings have better abra-

sive wear resistance compared to the hard chrome plating. It

is also important to mention that the wear test on the hard

chrome plating had to be discontinued after 20 min as the

plating got completely peeled off the substrate at the test

region under identical test conditions, indicating the poor

adhesion of the plating to the substrate. The tensile stresses

induced during the plating process and the microcracks

present may also lead to its rapid wear suggesting the superior

performance of both the carbide-based coatings against

abrasion.

4. Conclusions

1. WC–10Co–4Cr has better abrasive wear resistance

compared to Cr3C2–20(NiCr) coating possibly due to

the higher hardness of WC particles and better matrix

properties of the CoCr binder material. The indentation

fracture toughness of WC–CoCr coating was slightly

better than that for Cr3C2–20(NiCr), which may also

cause the wear resistance to improve in the former.

2. The coatings deposited by the DS process had slightly

improved wear resistance compared to the HVOF process

possibly due to the higher density and compressive

residual stresses induced in the DS coating.

3. Both the thermally sprayed coatings had superior wear

performance in comparison to the hard chrome coating.

Acknowledgements

The authors wish to thank the Director, DMRL, for

granting permission to publish the paper. Also, the support

provided by the Electron Microscopy group and the X-ray

group of DMRL is acknowledged.

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