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HVOF sprayed WC-CoCr coatings on aluminum: Tensile and tribological
properties
A Koutsomichalis1, M Vardavoulias2, P Psyllaki3 and N Vaxevanidis4
1 Ass. Professor, Hellenic Air-Force Academy, Athens E-mail: akouts@eeogroup.gr
2 Managing Director, Pyrogenesis S.A, Lavrio, GR, E-mail: mvardavoulias@pyrogenesis-sa.gr
3 Assoc. Professor, Piraeus University of Applied Sciences, Department of Mechanical Engineering, Athens
GR E-mail: psyllaki@teipir.gr
4 Professor, School of Pedagogical & Technological Education, Department of Mechanical Engineering
Educators, Athens, GR, E-mail: vaxev@aspete.gr
Abstract. This paper provides a comprehensive assessment of the tensile and sliding wear behaviour of WC–
10Co4Cr agglomerated and sintered powder deposited on aluminum by HVOF process. The microstructure
of the HVOF sprayed WC–CoCr coatings were characterized by scanning electron microscopy and optical
microscopy. Microstructural analysis identified grains of tungsten carbide (WC) in the metal matrix of the
cobalt-chromium (Co-Cr), which was confirmed by the elemental analysis of the electron microscope. A
transformation of the WC in the W2C phase was observed through a carbonization process and
decomposition of WC in the metal matrix. The HVOF WC-Co-Cr coating was found to decrease tensile
behaviour of the aluminum substrate. It was found that initial crack initiation and transverse crack
propagation occurred in the top coat layer. The number of transverse cracks increased rapidly and then
saturated with an increase in tensile strain. After the saturation of the multiple cracks, the cracks stopped on
the interface between the top coat and substrate. Finally, decohesion occurred at the bond coating-substrate
interface when the transverse cracks reached the interface. Tribological properties were examined using the
pin-on-disk method under various loads. The friction coefficient rose abruptly at the start-up phase and
stabilized at almost the same sliding distance independently of the applied load. The friction coefficient found
to increase with increasing applied load. Similarly, the wear volume increased linearly with increasing applied
load. Study of the wear mechanisms revealed surface micro-cracking and fragmentation of flattened coating
layers with subsequent gradual pull out of the carbide particles. Micro-cracks along the sidewalls of the crater
wear groove found to be perpendicular to the layer of the coating (brittle cracking) and sometimes parallel to
the coating layers. The latter leads to delamination of the adjacent layers interface.
Keywords:
High Velocity Oxygen-Fuel (HVOF) spraying, WC–10Co4Cr, Aluminum alloy, Pin-on-disk wear tests,
tensile behaviour
1. Introduction
Tungsten carbide coatings have been studied for a few decades because of their superior wear properties.
WC–Co cermet surface coatings are commonly used to enhance the wear resistance of many types of
engineering components, deposited via air plasma (APS) and HVOF spraying. These coatings accordingly
find a large variety of industrial applications for the protection of mechanical components against sliding
and abrasive wear at various temperatures [1] and in different environments. They are also listed among
the most promising alternatives to hard chromium electroplating [2], due to their technical advantages [3].
Tungsten carbide coatings are used for numerous industrial applications like aircraft, automotive,
petrochemical etc in solving severe abrasion, erosion and sliding wear problems. In addition to the coated
form, there are also used in sintered form for structural applications, for making components like cutting
tools, plungers, bearings, gears, steam turbine, hydro turbines etc. [4]. The sprayed WC–Co coatings can
exhibit complex, multi-phase microstructures, with a significantly lower volume fraction of primary
carbide than that of their starting powders [5]. Decomposition of WC–Co powders occurs due to high
temperatures and low velocities and the use of small carbide grain sizes within the powder particles, all of
which promote carbide dissolution in the molten matrix and subsequent decarburization. The degree of
decomposition of WC–Co powders during spraying, resulting in complex microstructures, depends
primarily upon the time–temperature history of the particle and the particle characteristics such as size,
porosity and WC grain size within the particle. However, while decomposition has been reported to affect
wear behaviour, low temperature spraying has, in some cases, shown to produce coatings with poorer
adhesion between the splats, again resulting in high rates of wear [6].
WC-CoCr can be successful applied to protect metal surface against pure wear. To achieve the desired
properties of wear it is important to optimize the powder composition, structure and the spraying
parameters. The properties of these coatings depend on the powder size and volume fraction of WC
powder feedstock and matrix composition [7]. WC-CoCr coatings derive its wear resistance properties
from the presence of high volume fraction of hard, wear resistant WC grains in a CoCr based metallic
binder phase. The presence of the metallic binder provides some toughness in the coating in comparison
with pure ceramic coatings. However, the relevantly high roughness of thermal-sprayed cermet coatings
often imposes their post-deposition conventional machining, which is difficult due to the inherent
hardness of the carbide. For this purpose, recent studies are focused on the tribological behaviour of
composite cermet/metal coatings, in which the metallic phase participates not as binder metal, but as a
distinct matrix of the product in percentages higher than 20 % [8].
HVOF is particularly suitable for the deposition of WC-based cermet HVOF coatings, since is
designed to retain a larger fraction of WC in the coating. The hypersonic velocity of the flame shortens
the powder-flame interaction time and together with lower flame temperature limit WC decomposition
[9]. However, its decomposition to W2C and metallic W and the WC reaction with the metal binder both
deteriorating the coatings hardness values and consequently their antiwear performance, have been
frequently reported [10, 11-16]. Recent studies, addressing this issue are focused on the optimization of
the deposition process parameters (e.g. spray distance, oxygen and flue flow rates) via Taguchi analysis
using as a criterion the coating hardness values achieved [15, 16]
The purpose of this work is to study the effect of WC-CoCr coatings applied by HVOF on the tensile
strength and wear behaviour of aluminum alloy. In specific, the objective of this work was to investigate
the friction and wear behaviour of HVOF WC-CoCr coating at different loads and identify the effect of
load on the wear mechanism of the carbide coating. This paper also attempts to provide an in situ
microscopic observation on the crack initiation, propagation and fracture behavior of the coated system
under uni-axial tensile loading.
2. Experimental part
A commercially available WC-Co-Cr agglomerated and sintered powder (WOKA 3653) was deposited on
the surface of Al-4%Cu alloy. The powder characteristics are given in Table 1.SEM micrograph of
agglomerated micro-structured WC–10Co–4Cr powder with EDX analysis is shown in fig. 1.
Table 1 Coating powder characteristic
Composition WC: 86%, Co:10%, Cr: 4%, Fe<0.3%
particle size -45+11 m
shape spherical
density 14.7gr/cm3
Two different types of specimens were prepared by the aluninum substrate: tensile specimens with
dimensions 130x90x2.1mm3and flat cylindrical test pieces for wear tests, 38.1 mm in diameter and 10 mm
thick. All specimens were grit blasted, prior to spraying, with aluminum oxide mesh 90.
The cermet powder was deposited by HVOF equipment. This step was accomplished by means of a
Diamond Jet 2700 Sulzer Metco gun, employing a mixture of propane as fuel and oxygen. The spraying
conditions are listed in Table 2.
WC-CoCr coating was deposited on both sides of the tensile specimens in equal thicknesses in order to
enable balanced loading during tensile testing. Cylindrical wear test pieces were sprayed to an average
coating thickness of 400m. The microstructure of the coated samples was observed under a
metallographic microscope (Leica DMR) and SEM (Jeol); see Fig.1, with image analysis software (Image
Pro).
For the evaluation of friction and wear characteristics sliding friction tests were performed (according
to ASTM G 99-05 standard) on a state-of-the-art CSEM pin-on-disc apparatus using, as counterbody, a
cutting insert with a diamond-coated tip. Four normal loads were selected, namely 1 N, 2 N, 5 N and
10 N. All tests were conducted at 16 mm track radius and with constant linear speed 0.6 m/s.
Table 2 Spraying parameters
Pressure Oxygen: 220 psi, Propane: 30 psi
Air 320 psi, Argon (carrier): 140 psi
Flow rate Oxygen: 240 lt/min, Propane: 74 lt/min
Air: 800 lt/min, Argon (carrier): 30 lt/min
Coolant: 24 lt/min
Powder feed 2,5 rpm
Spraying distance 200 mm
Revolutions 53 rpm
Gun speed 0,044 m/s
Flame temperature 2500
Figure 1 SEM micrograph of WoKA 3653 micron-sized feedstock powder showing spherical geometry
3. Results and discussion
Fig. 2 shows the cross section of the coating which is exhibiting dense structure, low porosity and is
macro-crack free. Coating-substrate interface shows no gaps or cracks indicating good adhesion between
the coatings and the substrate. The coating is built up by a deposit of lenticular splats, one over the other,
in a uniform manner, throughout the coating. Pores appear black in the micrographs. The coating
microstructure is observed to be composed of WC particles, which are uniformly distributed within the
coating.
(a) (b)
Figure 2 (a) Optical and (b) SEM migrographs of the WC-Co-Cr coating cross section
The coating microstructure typically consists of adjacent layers of dark and light areas (Fig. 3). The
microstructure of the dark areas was revealed to be mostly Co particles with aggregate of WC, while the
microstructure of the light areas was mostly consisted of sintered aggregate WC particles. Close
examination of light areas of coatings showed very small WC particles forming larger particles. Tungsten
carbide grains with sharp edges are attributed to their incomplete melting, while rounded carbide grains
result from the dissolution of WC, due to local overheating, in the Co matrix.
Figure 3 SEM/EDS analysis of HVOF coatings
It has been reported [17-21] that during spraying of WC–Co powder, particles are subjected to
decarburization leading to a decrease of the total carbon content. The extent of the WC transformation
has been associated with the starting powder type (size, morphology, and carbide size), the type of spray
process, the amount of oxygen in the environment, and the spray parameters [22, 23].
During decarburization tungsten carbide (WC) is transformed to tungsten semicarbide (W2C)
according to the following reactions [17]:
2WC W2C + C
W2C 2W + C
The XRD pattern of the coating, shown in Fig. 6, indicates the presence of both WC and the
decarburized brittle W2C phase formed during deposition. These brittle phases have been correlated with
reduced toughness and overall wear resistance [24] and are normally distributed in the peripheries of
splats, leading to a weak interface between the splats [25]. The formation of hard crystalline phases and
tough crystalline Co, without the consumption of a great deal of the WC phase, results in higher wear
resistance coatings [26].
Figure 4 XRD diffractogram of the WC-Co-Cr coating
Figure 5 shows the engineering stress – strain curves for the WC-CoCr coated Al substrate with
various coating thickness. The presence of the WC-CoCr coating results in a lower tensile strength of the
coated aluminum in comparison with the tensile strength of the uncoated aluminum. The reduction of the
tensile properties may be attributed to the brittle characteristics of the coating, while fracture analysis
revealed that brittle cracks, formed in the coating, are the primary cause of the tensile failure of the coated
specimens. Tensile tests were conducted for a number of different coating thicknesses. The coated steel
substrate with equal WC-CoCr coating thickness on both sides can be treated as a laminated composite
stressed parallel to the lamellas with the coating and the substrate being the constituents, which determine
the elastic and elastic-plastic behaviour.
Since the applied strain in the loading direction is same for all the components of the specimen, an iso-
strain condition in the substrate and coating layers is assumed. According to the mixture law for
composite materials, following relationship can be given for the present coating system:
Ecs(hAl + hWC) = EWChWC + EAlhAl (1)
where Ecs, EAl and EWC are the Young’s modulus of the coated system, the substrate (Al) and coating
layer (WC-Co-Cr), respectively, and hAl, hWC are the thicknesses of the substrate and the coating layers.
From the measured uniaxial stress–strain curves, the average Young’s modulus of three coated specimens,
Ecs, is 372 GPa. The Young modulus of the substrate is 70GPa, as measured, and the reported Young modulus
of the coating [27] is 310 GPa, the estimated Young’s modulus of the coated system, according to Eq. (1), is
400 GPa, which is in good approximation with the experimental values.
Figure 5 Tensile stress–strain curves of HVOF WC-Co-Cr coated Al substrate with various coating thicknesses
During tensile testing the coating exhibited cracking in the transverse direction of the tensile load.
After the saturation of the multiple cracks, the cracks stopped on the interface between the top coat and
substrate. Cracks in the coating surface were observed prior to the transition in the plastic region and are
attributed to its brittle nature. While tensile testing was deployed, cracks multiplied (fig. 6a). Multiplication
of transverse cracks in the coating proceeds until the applied strain reaches ~1.6 % and then the crack-to-
crack spacing rapidly decreases. As tensile strain increased, the initiated cracks propagate transversely
towards the coating-substrate in tree like manner. When the initially nucleated cracks approach the
interface, they stop and then new cracks form and propagate (fig. 6b). The cracks upon reaching the
interface deflect into parallel microcracks and propagate causing delamination of the coating especially if
the coating thickness is increased. After the cracks propagate and pass through the coating-substrate interface,
decohesion occurs just after the transverse crack tips arrive at this interface (Fig. 6c). The interfacial
delamination cracks have been studied in coated systems and are attributed to the linkage of previously
formed cracks and are relative to the residual stresses and are associated with the fracture toughness of the
coating [28, 29]. Another factor that strongly influences the tensile performance of the HVOF WC-CoCr
coatings is adhesion of the coating to the substrate. This is largely influenced by the pre-spraying surface
preparation which is grit blasted to roughen the surface to allow mechanical interlocking of the splats with
the blasted surface. The strength between the coating and substrate can be compromised by embedment
of grit into the substrate during blasting and by the presence of grit remnants prior to deposition. These
grit particles act as stress raisers and are agents for interfacial crack propagation and delamination of the
coating when stress is induced.
(a) (b) (c)
Figure 6 (a) Plan view of the fractured coated system (b) transverse cracks in the cermet coating passing through the
interface (c) enlargement of crack tip region in (b)
Figure 7 shows the friction coefficient of WC-CoCr coating for the four different loads 1N, 2N, 5N
and 10N as a function of the sliding distance (or equivalently, time). Initially and for the first 250m
approximately the friction coefficient rapidly increases. After this initial running-in period the friction
coefficient tends to a constant steady-state value, for all the four loads applied. [30]. This constant value
was found to be increasing with increasing the normal load applied. More precisely, the steady-state
friction coefficient value increases from 0.36 to 0.44 and to 0.55, as the normal load increases from 2 to 5
and to 10 N, respectively. During the application of the 10N load an increase of temperature was
observed affecting the ductility of the coating thus leading in higher friction coefficient since as it has been
elsewhere reported [31] even a relatively modest temperature increment can have a large effect on the
tribological properties. The friction coefficient is in the same range with most literature values [32-34].
The wear volume loss after 10000 m sliding versus the applied load is shown in Fig. 8. It is evident that
wear volume (equivalently wear mass losses) increases with increasing applied normal load following the
Lancaster model [35]:
V = ksN (2)
where V is the volume wear, s is the total sliding distance, N is the normal load and k is the wear
coefficient or specific wear rate.
For relatively low normal load applied (2 N), the pressure at the contact area does not exceed the
fracture stress of the ceramic coating; during sliding, friction obeys to the classical mechanical mechanisms
leading to higher friction coefficient values due to the high roughness of these coatings. At the level of the
contact area, sliding follows the topography of the coating’s protrusions and wear takes place through
micro-cutting. For high normal load applied (10 N), the pressure at the contact area exceeds the fracture
stress of the ceramic leading to lower friction coefficient values, since extensive micro-fragmentation at
the level of the coating’s protrusions results in an intermediate layer of debris that facilitates sliding. For
intermediate normal load applied (5 N), the pressure at the contact area exceeds only locally the fracture
stress of the ceramic, e.g. at the level the higher protrusions leading to a mixed wear mechanism.
Figure 7 Variation of the friction coefficient with the sliding distance for various loads
Figure 8 Wear volume as a function of applied load
Stereoscopic analysis showed that the wear grooves are wider with increasing normal load and during
sliding material is removed and deposited on the groove edges (fig. 9). The cermet coating combines
extremely hard carbide grains with softer matrix. These are indicative of typical abrasion wear,
encountered when a ceramic counterbody (diamond tip) is sliding against an elastoplastic metallic surface
(WC-CoCr coating) [36].
Figure 9 Plan view of the wear grooves (a) 1 , (b) 2 , (c) 5 , (d) 10
In the case of WC-Co-Cr coating layers consisting of hard particles and a tough matrix, the wear
behaviour is related to the general wear mechanism but also to the exclusion of the hard particles of WC,
and the wear loss can be explained by the delamination from the interface [37]. The wear mechanism is
associated with an initial removal of the binder phase followed by a fragmentation or uprooting of carbide
particles. As a consequence of that the wear debris are trapped between the coating surface and the
counterbody. While the wear debris is small, damage along the wear groove is low. Polishing lines along
the sliding direction, together with micro-ploughing traces are clearly seen in Fig. 10a. The decohesion of
the metallic matrix led to pull-out of the tungsten carbides, as also observed previously [4], that remained
on the worn surface during testing (Fig. 10b). Further observation of the wear tracks revealed that the
coating was worn out through the normally expected material removal via debris creation.
Figure 10 SEM top-view micrographs of worn surface
SEM observations on composite coatings cross sections revealed subsurface coating damage as well.
Cracks were developed at the wear track boundaries on the contact surface and propagated parallel to the
coating surface (Fig. 11a and b), practically splitting them and resulting in partial decohesion. Such cracks
have been reported elsewhere [34, 38] but are found to propagate perpendicular leading to eventual splat
exfoliation and consequently to relatively increased wear rates. In this case cracks are localized to surface
and were not found to reach the coating-substrate interface. EDX microanalysis showed that the debris
layer consists of a mixture of metallic and carbide components occurring from the simultaneous wear of
both (fig. 11c, d). The carbides particles in the debris act as a third abrasive body, destroying locally the
tribo-layer that acts as an in-situ lubricant.
Figure 11 SEM cross section of micrographs of the WC-Co/Cr composite coating: (a) Typical view of the entire
wear track width; (b) Cracks at the boundaries of the wear track (c) Magnification of relevant marked area in (b) and
(d) EDX analysis of marked area in (c)
4. Conclusions
1. WC-10Co-4Cr micron sized powder was deposited by HVOF on the surface of aluminum alloy. A
typical microstructure characterized by solidified flattened droplets and the coexistence of semi-molten
particles with a certain degree of porosity was identified.
2. Tensile tests were conducted for a number of different coating thicknesses. In all cases examined the
tensile strength of the coated aluminum was lower than that of the uncoated one indicating that the
presence of the WC-CoCr coating reduces the tensile strength. During tensile loading, cracks first
initiated in the outer layer of the coating and transversely propagated towards the substrate, and then
cracks multiplied and saturated in the coating; finally interfacial decohesion occurred at the coating -
substrate interface.
3. The variation of the friction coefficient with the sliding distance followed similar patterns for four
loads applied (1, 2, 5, 10N). After an initial running-in period, the friction coefficient tends to a
constant steady-state value, for all the cases examined. This steady-state value was found to be
increasing with the applied normal load.
4. Wear volume increases with increasing applied normal load. The wear mechanism was micro-cracking
of the metal binder leading to tungsten carbide particles pull-out. Cracks were developed at the wear
track boundaries on the contact surface and propagated parallel to the coating surface causing partial
decohesion.
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