Post-print of: Surface and Coatings Technology Volume 206, Issue 7, 25 December
2011, Pages 1913–1920
Identification of the wear mechanism on WC/C nanostructured coatings
S. El Mrabet, M.D. Abad, J.C. Sánchez-López
Instituto de Ciencia de Materiales de Sevilla (CSIC-US), Avda. Americo Vespucio 49,
41092 Sevilla, Spain
Abstract
A series of WC/C nanostructured films with carbon contents ranging from 30 to 70 at.%
was deposited on M2 steel substrates by magnetron sputtering of WC and graphite
targets in argon. Depending on the amorphous carbon (a-C) incorporated in the
coatings, nanocrystalline coating (formed mainly by WC1 − x and W2C phases) or
nanocomposite (WC1 − x/a-C) were obtained with tunable mechanical and tribological
properties. Ultrahardness values of 36–40 GPa were measured for the nanocrystalline
samples whilst values between 16 and 23 GPa were obtained in the nanocomposite ones
depending on the a-C content. The tribological properties were studied using a pin-ondisk tester versus steel (100Cr6) balls and 5 N of applied load in dry sliding conditions
and the failure modes by scratch adhesion tests. Three different zones were identified
according to the observed tribological behavior: I (μ > 0.8; adhesive wear), II (μ: 0.3–
0.6; abrasive wear) and III (μ ~ 0.2; self-lubricated). The wear tracks and the ball scars
were observed by scanning electron microscopy (SEM) and Raman spectroscopy in
order to elucidate the tribochemical reactions appearing at the contact and to determine
the wear mechanism present in each type. A correlation among structure, crystalline
phases, a-C content and tribomechanical properties could be established for the series of
WC/C coatings and extended to understand the trends observed in the literature for
similar coatings.
Keywords
WC; Coatings; Amorphous carbon; Structure; Mechanical properties; Tribology
1. Introduction
Tungsten carbide (WC), a well-known refractory material is widely used in the
industrial applications because of its high hardness, high elastic modulus, wear
resistance and chemical inertness [1], [2], [3], [4], [5], [6] and [7]. One of the main
peculiarities of the WC material is the high number of compositional and structural
forms that can exist according to the W–C phase diagram [8] and [9]. The deposition of
tungsten carbide films have been done by many chemical vapor deposition (CVD) [10],
[11], [12] and [13] and physical vapor deposition (PVD) methods [1], [2], [3], [4], [5],
1
[6], [7], [14], [15], [16], [17], [18], [19], [20], [21], [22], [23], [24], [25], [26], [27],
[28], [29], [30], [31], [32], [33], [34], [35] and [36]. Gouy-Pailler and Pauleau[3] in
1993 found by X-ray diffraction analysis (XRD) that coatings with carbon contents
below 25 at.% contained α-W phase, named W(C) films, whereas cubic WC1 − x
phases were detected when carbon content surpassed 30 at.%. Palmquist et al. [14]
demonstrated by XRD, transmission electron microscopy (TEM) and electron
diffraction (ED) that a hexagonal W2C phase could be formed between the solid
solution of C in α-W and the cubic WC1 − x phase for C contents between 20 and 35
at.%. The appearance of amorphous carbon (a-C) surrounding the W2C and WC1 − x
crystalline phases was reported for the ranges superior to 30 or 40 at.% of C depending
on the author [15], [16], [17], [18], [19] and [20]. The properties of such nanocomposite
coatings depend critically on the ability to co-deposit both the nanocrystalline (nc-WC)
and amorphous phases (a-C, a-C:H, DLC) controlling its relative amount and
distribution inside the nanocomposite.
Focusing on the tribological applications, many groups have synthesized WC/C
composites with the aim of combining the benefits of a hard nanocrystalline material
with a soft solid lubricant [15], [16], [17], [18], [20], [21], [22], [23] and [24]. Hardness
(H), stress and friction coefficient (μ) are usually the most studied parameters to
determine the quality of the coatings for this functionality. In Fig. 1 it is shown a data
review on hardness, stress and friction coefficient values vs. carbon content of different
nanocomposites composed of WC nanocrystals (whichever stoichiometry) embedded in
an amorphous carbon matrix (hydrogenated or hydrogen-free) found in the literature. A
wide dispersion of the results and a lack of correlation can apparently be inferred. One
can observe the existing differences in hardness and friction coefficient values between
samples containing similar carbon content but obtained by different groups. This could
be partially attributed to the employment of different deposition techniques and
synthesis conditions. For instance, Voevodin et al. [15], [16] and [17] prepared
WC/DLC nanocomposite coatings using laser ablation of graphite, meanwhile,
Czyzniewski et al. [22], [23] and [24] prepared WC/a-C:H with acetylene as a precursor
gas. Another property less studied that can provide complementary information on the
mechanical behavior is the scratch adhesion [37], [38] and [39]. Hardness and scratch
adhesion testing share many features and therefore that scratch testing constitutes a
valuable tool to gain understanding of the deformation processes induced by the
cumulative action of indentation and friction.
The main motivation of this paper is to investigate the friction and wear mechanisms of
nanocomposite WC/a-C coatings using a set of well characterized samples as base
material [20]. The determination of the fraction of carbon atoms bonded to tungsten or
bonded to carbon (as a-C matrix phase) by X-ray photoelectron spectroscopy (XPS)
resulted crucial for understanding the changes in structure and mechanical and
tribological properties. In this paper, the acquired knowledge on the phase composition
and microstructure together with a detailed analysis by microscopy and Raman
techniques on the wear tracks originated by scratch and friction tests are used to explain
2
the observed tribomechanical properties and correlate it with the reported literature data
for similar WC/C or WC/C:H compounds.
2. Experimental details
WC/a-C coatings were prepared by Ar sputtering of WC (Kurt J. Lesker, 99.5% purity)
and graphite (Goodfellow, 99.5% purity) targets connected to radio frequency (r.f.) and
direct current (d.c.) power sources respectively. A series of samples has been prepared
by changing the sputtering power ratio, defined as R = PC/PWC, from 0 to 3. The
typical power values (PWC) applied to the WC target were 150 and 250 W while those
applied to the graphite target (PC) were varied from 0 to 450 W. The obtained films are
labeled as R0, R0.1, R0.3, R0.5, R1, R2 and R3. Further experimental details
concerning the synthesis conditions can be found in reference [20].
Nanoindentation experiments were performed with a Nanoindenter II (Nano
Instruments, Inc., Knoxville, TN) microprobe. All tests were carried out at room
temperature with a diamond Berkovich (three-sided pyramid) indenter tip. The load–
displacement data obtained were analyzed using the method of Oliver and Pharr [40] to
determine the hardness and the elastic modulus as a function of the displacement of the
indenter. The maximum load was selected in such a way that the maximum indentation
depth did not exceed 10–15% of the coating thickness in order to avoid the influence of
the substrate. The film stress was measured by measuring the substrate bending using a
profilometer and the Stoney's equation.
The scratch tests were carried out using a TRIBOtechnic Millenium 200 scratch-tester.
A Rockwell C diamond tip (200 μm radius) was used as an indenter. During the test, the
indenter was drawn over the coated surface for 10 mm-length as the applied normal load
increased continuously up to 100 N (loading rate of 50 N/min and a scratching speed of
10 mm/min). The diamond tip was cleaned after each scratch. Acoustic emissions were
recorded during the scratching, but optical microscopy was applied to assess the critical
normal load values and to indicate the coating fracture and delamination modes. Two
critical loads were determined, the lower (LC1) being the start of the cohesive failure
within the coating, and the upper (LC2) the onset of an adhesive failure of the coating
[39]. Each sample was submitted at least to three scratch experiments to determine the
general trends in performance of the coatings. From the collected results, mean values
for the critical loads corresponding to specific damage of the films, and experimental
variations were determined.
Tribological tests were carried out using 100Cr6 6 mm-diameter steel balls in a pin-ondisk CSM tribometer with a sliding speed of 10 cm/s and 5 N of applied load
(maximum initial Herztian contact pressure of 1.12 GPa) in ambient air (30–60% of
relative humidity). The sliding distance was 1000 m with typical track radius between 6
and 10 mm. Normalized wear rates (mm3/Nm) were evaluated from cross-sectional
profiles taken across the disk-wear track after testing by means of stylus profilometry.
Scanning electron microscopy (SEM) data were recorded in a FEG Hitachi S5200
microscope operating at 5 kV. Micro-Raman measurements were performed using a
3
LabRAM Jobin Yvon spectrometer equipped with a microscope. Laser radiation (λ =
532 nm) was used as excitation source at 5 mW. All measurements were recorded under
the same conditions (10 s of integration time and 10 accumulations) using a 100×
magnification objective and a 100 μm pinhole.
3. Results
3.1. Hardness, critical loads and tribological properties
In previous works a series of WC/a-C coatings was prepared by magnetron sputtering
and exhaustively characterized by XRD, TEM, ED, Raman and XPS techniques [20]. In
summary, the results of the chemical and microstructural analysis revealed that
nanocrystalline hexagonal W2C phase is the main phase by single sputtering of WC
target. Then, the subsequent incorporation of carbon leads to a progressive reduction of
the crystalline domain size and the nucleation of the cubic WC1 − x phase. From a total
C content of approximately 50 at.% the formation of composite films containing
nanocrystallites of cubic WC1 − x phase dispersed in an amorphous carbon matrix is
clearly manifested. Further increase of the power applied to the graphite target leads to a
progressive increment of the free amorphous carbon content becoming comparable to
the crystalline fraction from overall 70 at.% of C.
Table 1 summarizes the mechanical and tribological properties of the set of WC/C
coatings prepared varying the R parameter as a function of the total and amorphous free
carbon contents. In Fig. 2a the hardness and friction coefficient of the samples are
represented as a function of the a-C (at.%). The maximum hardness are obtained for the
R0 and R0.1 samples (36 and 40 GPa respectively) with very low of a-C contents (< 10
at.%). The remaining samples exhibited a progressive diminution in hardness and
friction coefficient values (μ) by increasing the a-C content to 16–20 GPa and μ ~ 0.2
respectively for the richest carbon samples. In Fig. 2b the values of the ball and film
wear rates (K) are represented as a function of the a-C content. It should be mentioned
that the Kfilm values for samples R0 and R0.1 are not provided due to the transfer of
mating material (steel) to the surface making impossible the estimation of the wear track
as it will be explained in the next sub-section. In Fig. 2c both critical load values have
been plotted for each coating. It should be mentioned that the first deposition step
consisted in switching on only the WC target (for 1 h) and then combined with the
graphite one in order to have a graded transition from WC to WC/C layered system. As
a result of this procedure a first layer of 100–200 nm of nanocrystalline W2C is located
at the substrate interface. The LC1, the starting point of crack formation, was found in a
very close range between 12 and 15 N for the coatings containing up to 16 at.% of a-C.
For a-C contents higher than 26 at.%, the LC1 value was not observed as they deform
mainly plastically as described in the next section. The found values for LC2 exhibited
larger differences depending on the amount of the a-C phase. The samples with the
lowest a-C showed very high values around between 63 and 69 N, followed by a
significant decrease to 30 N when a-C is between 10 and 16 at.%. In the last group, the
samples with a-C > 25 at.% the trend is reversed, reaching almost 50 N for the sample
4
with the highest a-C content. In these coatings the critical load corresponded to the point
where coating is scrapped off exposing the substrate.
In summary, the study of the dependence of the tribomechanical properties with the
phase composition allows to establish three different zones using the amount of atomic
carbon present in the a-C phase as main criteria: I (< 10 at.%); II (10–30 at.%) and III (>
30 at.%). This classification will be used in the following sections for the discussion of
deformation modes and wear mechanisms.
3.2. Scratch tests and failure modes
A deeper study of the failure mode by means of optical microscopy was carried out in
order to understand the differences in the critical load values. In Fig. 3 it is depicted the
picture taken at 12.3 N for the coating R0.1 as starting point of the crack formation
(conformal cracks). As the indenter moves along, several partial ring cracks are formed
along the track and these rings may intercept with each other causing a network of
cracks (LC2 = 63.0 N). Conformal cracking occurs by the compressive stress ahead of
the moving indenter and this driving force is similar to bucking spallation. The reason
for coating cracking without delamination indicates a sufficient interface substrate-film
adhesion. This kind of failure response was also observed by scratching nanocrystalline
coatings of WC by Czyzniewski [22].
Fig. 3 also shows the failure mode of the coating R0.3 representing the transition from
the region I to II. Some diagonal lines are formed when the coatings is deformed by the
indenter in its moving direction (LC1 = 13.6 N). These marks are called as “chevron
tensile cracks” because they point to the crack origin. This mode of failure also occurs
at the trailing end, but the fracture initiates near the two edges of the contact groove,
forming a slanted angle to the sliding direction. As can be seen in the Fig. 3 (LC2 = 31.5
N) with further increase of the load, both the density and the length of the cracks
increase, and spallation occurs due to the compressive stress generated by the indenter
[41]. In Fig. 3 it is likewise shown the typical failure mode of the coating R0.5 (LC1 =
14.1 N). In this case the scratch failure mode corresponds to the formation of “tensile
trailing cracks”. The tensile cracks are formed under the dominant effect of frictional
traction stress. Besides the stresses caused by frictional pulling, coating bending due to
the contact action also contributes to the tensile stress that opens up the crack
perpendicular to the sliding direction behind the moving stylus. When the contact load
is large severe grooving in the substrate will also cause cracking in the coating
alongside the groove edge (LC2 = 27.9 N). The crack lines do not overlap with each
other. This results are in good agreement with Czyzniewsky [22] where he observed
that increasing carbon content in nanocomposite WC/a-C:H from 48.1 to 56.5 at.%
there is an increase of the number of cracks and their propagation in this matrix.
Czyzniewski proposed that together with the increasing carbon content, the thickness of
the amorphous a-C matrix enclosing WC nanograins is also growing and in
consequence number of cracks forms and propagates in this matrix.
5
In Fig. 3 we can see the wear scars of the R2 and R3 coating taking at LC2 = 23.0 and
45.6 N, respectively. The aspects of the wear tracks reveal that the coatings exhibited
surface deformation, which visually appear to be “plastic”. Until above 20 N, it seems
that the indenter produces a plastic deformation and just it is possible to observe some
marks or longitudinal cracks at the border parallel to the indenter movement indicating
material pile-up along the track borders. This behavior can be interpreted as the coatings
have sufficient ductility to accommodate plastic deformation. These samples
corresponded to the third region described in the previous section whose a-C contents
were maximum (around 30 at.%). Voevodin and Zabinski [17] in their early work on
WC/DLC coatings explained that this was not true plasticity since dislocation sources
were prohibited, but rather the result of WC grain boundary sliding in the a-C matrix.
The increase of the applied load led an increased extent of the “plastic” deformation
until the point that the coating is scrapped off exposing the substrate.
3.3. Analysis post-tribo by SEM/EDX
In order to have a better insight about the chemical phenomena occurring at the contact
area, analysis of the ball and film surfaces after friction tests is carried out by
SEM/EDX and Raman on selected samples. R0.1 (3 at.% of a-C), R0.5 (16 at.% of a-C)
and R2 (30 at.% of a-C) coatings have been selected as representative of high, medium
and low friction regimes respectively. Ball and disk worn surfaces of R0.1 sample are
shown in Fig. 4 as representative example of the high friction coefficient region (I)
formed by the samples with less a-C content (R0 and R0.1). When sliding against steel,
a friction coefficient of about 0.8 was found for these two coatings. This can be
understood considering the great difference in hardness properties between film (36–40
GPa) and counterface (5–6 GPa) materials. Under high contact pressure (> 1 GPa), the
softer steel is rapidly worn away leading to an increased friction. The more ductile steel
is also transferred to the film surface appearing smeared onto the film wear track as can
be seen in Fig. 4 in the track disk at higher magnifications. A SEM/EDX line scan of the
wear track (depicted as a white mark) revealed that the composition of this adhered
layer was mainly iron oxide. The analysis of the debris deposited on the ball scar by
EDX was not concluding and so was lately assessed by Raman.
In the second region (R0.3 to R1), friction coefficients varied between 0.6 and 0.35 and
the values of wear rates were relatively high for film (~ 10−6) and ball (~ 10−5–10−6
mm3/Nm) respectively. Ball and wear track for R0.5 film with 16 at.% a-C are shown
in Fig. 4. The wear track is characterized by surface deformation in the form of
longitudinal grooves produced by hard debris particles entrapped within the track. In
this case, both high friction and wear values registered could be originated from a high
concentration of hard phases in this zone. Therefore the production of the abrasive wear
particles during frictional contact contributed to accelerate the coating damage by
cracking and spallation achieving high friction and high wear rate.
For the third region corresponding to highest content of a-C (R2 and R3), low friction
values in combination with low wear rates were obtained (μ < 0.2 and K ~ 10−7–10−8
6
mm3/Nm). The SEM pictures of both counterfaces obtained for the sample R2 are
depicted in Fig. 4. The wear marks are less visible indicative of the improved
tribological behavior in this region. The film track is characterized by shallow grooves
and no loose debris appears covering the ball surface. This is in good agreement with
the nanocomposite phase composition and the results observed by the scratch test where
the high content of amorphous phase allowed to accommodate the shear induced by the
tip deforming pseudoplastically.
3.4. Raman analysis of the friction contact regions
Raman spectra from the as-deposited coatings, the wear track surface and the transfer
film formed on the ball counterfaces for the selected representative (R0.1, R0.5 and R2)
samples are given in Fig. 5. The Raman spectrum of the transfer layer formed on the
steel counterpart of sample R0.1 (Fig. 5a) showed the presence of two bands at 720 and
945 cm−1 which can be assigned to a mixture of iron and tungsten oxides or
ferritungstite [42]. The strong broadening of these bands suggests that the formed
compounds present a high structural disorder, indicating low crystallinity. The peak at
945 cm−1 is also possible to be assigned to the stretching mode of W=O bonds that
appear on the boundaries of amorphous or nanostructured tungsten oxides [43]. These
results correlate with the strong interaction of the steel counterpart with the film surface
for these samples highlighted by the SEM/EDX analysis. In the track, very weak peaks
in the range of 1300–1600 cm−1 corresponding to the D and G bands typical of
disordered amorphous carbon were also observed with an additional peak at 880 cm−1
from crystalline WO3[16] and [44]. The Raman spectrum of the initial film does not
display any features in the region of D and G bands indicating that this carbon Raman
signal is only present in the contact after friction tests. The oxidation of the WC
nanocrystals in air generates tungsten oxides and free carbon although insufficient or
ineffective to lubricate the contact.
Fig. 5b shows the spectra obtained for the sample R0.5 (16 at.% a-C). The Raman
spectrum of material adhered on the steel counterpart showed broad peaks within the
range 220–650 cm−1 characteristic of iron oxides phases [45]. The two intense peaks at
220 and 283 cm−1 can be assigned to hematite (α-Fe2O3) and the broad band situated
at 350–390 cm−1 to the maghemite (γ-Fe2O3) [46]. The two broad Raman peaks at
about 1360 cm−1 and 1600 cm−1 indicated that some a-C phase, initially present in the
film (cf. bottom spectrum) or induced by friction, was transferred to the ball. For higher
amorphous carbon contents (Fig. 5c) the Raman spectrum of the as-deposited film
exhibits a broad band at 1300–1500 cm−1 indicative of the presence of a disordered
network of sp2-bonded carbon matrix. In this case, we found the typical carbon features
both in the track and the ball surfaces with only minor signs of oxidation so the adhered
material on the ball scars must come from the nanocomposite WC/a-C film. Besides, we
could observe a significant sharpening in the signal of D and G peaks and shift of the Gposition towards higher frequencies. These changes are consistent with an increase
ordering of the initial disordered sp2-C phase in agreement with the behavior
demonstrated by DLC and other carbon based coatings [47].
7
3.5. Wear modes
Attending to the film microstructure, crystalline and chemical composition previously
presented [20] and the results obtained in this paper after the scratch and friction tests
we can try to explain the observed tribo-mechanical performance. In the Fig. 6 we have
summarized the hardness and friction coefficients measured for the coatings of this
work plotted versus the total carbon content and other structural and chemical aspects.
Although we were using the a-C content as key-parameter to discriminate the different
type of behavior we have plotted here again the total C content in order to enable the
comparison with the literature data (mostly in overall C content). Thus, the observed
performance can be broadly grouped into three categories:
I. Group nanocrystalline W2C and/or WC1 − x coatings (approx. 30–40 at.% C). The
main characteristic of these coatings is the major presence of crystalline W2C and WC1
− x phases with grain sizes ranging 5 to 10 nm with scarce a-C contents (< 7 at.%).
They possess high hardness (35–40 GPa) but poor lubricant properties. Although the
formation of cracks begins early due to their brittle ceramic character, they
demonstrated a maximum fracture toughness and resilience (resistance to plastic
deformation). These coatings interact severely with the steel counterface producing
wearing of the ball and iron transfer to the wear track. The formation of tungsten oxide
and mixed iron–tungsten-oxide (ferritungstite) are the evidences of the tribochemical
reactions occurred in air.
II. Group hard nanocomposite WC1 − x/a-C coatings (approx. 40–65 at.% C). This class
can be defined as hard nanocomposite coatings (hardness about 21–23 GPa) composed
of small WC1 − x crystals surrounded by an amorphous carbon matrix (about 10 to 25
at.% of a-C). For the lowest of this a-C content range, where there is accordingly a
higher concentration of the hard WC1 − x phase, the production of abrasive debris
particles contributes to accelerate the degradation of the coating achieving high frictions
and high wear rates. In many cases these reaction products adhere strongly to the
surface. An abrasive/adhesive (predominantly abrasive) wear mode can be identified for
these coatings. When the a-C concentration is increased, it can be seen how there is a
diminution of friction concomitant with the wear rate. The reasons behind this
tribological improvement must be found both in the reduction of hard WC1 − x debris
particles promoting abrasive wear of the coatings and, in less extent, to a selflubrication promoted by the a-C phase. This latter becomes the dominant factor in the
third group of WC/C samples.
III. Group lubricant nanocomposite WC1 − x/a-C coatings. These nanocomposites are
formed by a poor crystallized WC1 − x phase embedded in a major amorphous carbon
matrix. The hardness of these coatings are lower (16–20 GPa) although are best
qualified in terms of tribological properties. This good tribological behavior compared
with the hard-moderate nanocomposites can be explained by a change in the wear
mechanism from mixed abrasive/adhesive to pure sliding controlled by the a-C phase
8
supply to the contact. When the a-C phase is over 30 at.% a solid lubricant mechanism
is found. The formation of a carbonaceous third body material in the contact preserves
the counterfaces from degradation and promotes an easy shear of the sliding surfaces.
Under these conditions, the specific wear rates reach values in the range of 10−8
mm3/Nm, which are much below typical wear rates of hard metal carbides and nitrides
and comparable to that of metal doped or pure DLC coatings. Concerning friction,
values as low as 0.2 for coatings with ~ 30 at.% of a-C are appropriated for operating in
dry lubrication conditions similarly to many DLC and carbon-based compounds used in
tribological applications [47]. In comparison with the well studied load-adaptative
TiC/C self-lubricant nanocomposites, the WC/C represents the additional advantage of
requiring less amount of a-C to reduce friction below 0.2, exhibiting higher hardness
and fracture toughness. This has been explained by considering a friction induced
decomposition of the non-stoichiometric carbides releasing free carbon and forming
structures that accommodate easily carbon vacancies [21].
4. Concluding remarks
The investigation of the mechanical and tribological properties of a set of well
characterized WC/a-C nanostructured coatings in connection with a deep analysis of the
structural and chemical features responsible of this behavior has been very useful for
clarifying the trends observed. At first glance, if we except the low carbon region (< 20
at.%) that can be rather defined as a solid solution of W(C) than a nanocomposite, we
can find that our set of coatings reproduces the same trend observed in the literature.
According to phase composition and wear mode identified within these coatings the
following conclusions can be inferred:
a) For carbon contents ranging between 30 to 40 at.% C, the highest values of hardness
and friction are found. This region corresponds to films made of nanocrystalline W2C
and/or WC1 − x, with a minimum quantity of free amorphous carbon. An adhesive wear
mode was identified and Fe–W–O phases were observed by Raman either in the ball
and track counterfaces. The critical loads for these coatings were very high showing
high adhesion to the steel substrate and the failure mode was coincident with the
formation of conformal cracks.
b) For carbon contents ranging between 40 and 65 at.% C a continuous decrease of the
hardness and friction coefficient with the increment of the carbon content is noticed as
the main particularity of this region. The structure of these coatings is composed of
nanocrystalline WC1 − x embedded in a matrix of a-C but insufficient to lubricate the
contact (10–30 at.%). Therefore moderate high values of hardness (20–25 GPa),
together with intermediate friction coefficients (0.3–0.6) and high film wear rates are
found. The wear mechanism found for these coatings is mainly influenced by its
abrasive component. Iron oxides were produced together with hard debris of WC1 − x
which became abrasive and high friction coefficients are obtained. The scratch adhesion
9
strength in this region is limited. Forward chevron and tensile cracks were identified as
failure modes.
c) At high carbon contents (> 65 at.% C), the amount of free a-C (over 30 at.%) inside
the nanocomposite is enough to make reducing the friction below 0.2, with very low
wear rates (in the range of 10−8 mm3/Nm). The formation of a carbonaceous third body
material in the contact of these coatings preserved the counterfaces from degradation
and promotes an easy shear of the sliding surfaces. If the concentration of carbon is not
increased in excess a good tribological behavior is possible to be obtained with a
reasonable moderate hardness (10–15 GPa). These coatings were able to withstand
higher deformation without fracture as the enrichment in the amorphous matrix enabled
to deform pseudoplastically.
The obtained information can be very useful for tailored design of this group of
nanocomposite coatings depending on the requirements needed for specific applications
as gear, bearings, tools and mechanical components in general submitted to wear.
Acknowledgments
The authors are grateful to the Spanish Ministry of Science and Innovation (project nos.
MAT2007-66881-C02-01, MAT2010-21597-C02-01 and Consolider FUNCOAT
CSD2008-00023), Junta de Andalucía (TEP217) and I3P program of CSIC for financial
support.
10
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14
Figure captions
Figure 1. Comparative analysis of hardness (a), stress (b) and friction coefficient (c)
values of WC/C films found in the literature.
Figure 2. Dependence of tribo-mechanical properties as a function of the a-C content:
hardness and friction coefficient (a); film and ball wear rates (b) and scratch critical
loads (c).
Figure 3. Optical micrographs of the scratch tracks taken for coatings R0.1, R0.3, R0.5
at different critical loads: LC1 (left column) and LC2 (right column). The last two
pictures correspond to LC2 values of R2 and R3. The direction of the indentor
displacement is from left to right.
Figure 4. SEM micrographs of ball wear scar and wear track associated to the sample
R0.1, R05 and R2. Also a zoom of the zone marked in the wear track of R0.1 showing
the material transfer from steel to the film wear track and EDX line scan profile taken
across the wear track are shown.
Figure 5. Raman spectra obtained from the ball scars and the wear tracks after the
tribo-testing in comparison with the initial film spectra for films R0.1 [7 at.% a-C] (a),
R0.5 [16 at.% a-C] (b) and R2 [30 at.% a-C] (c).
Figure 6. Summary of hardness and friction coefficient values found for the WC/a-C
nanocomposite coatings deposited in this work along with the defined regions according
to the phase composition, a-C (at.%) and identified wear mechanism.
15
Table 1
Table 1. Mechanical and tribological properties for the WC/a-C samples.
Tribo-mechanical properties
Ctotal a-C
(at.%) (at.%) H
(GPa)
R0 33
7
36 ± 4
R0.1 37
3
40 ± 10
R0.3 50
10
23 ± 2
R0.5 55
16
22 ± 4
R1 64
26
21 ± 1
R2 69
30
20 ± 2
R3 71
31
16 ± 3
Film
μ
0.84 ± 0.01
0.81 ± 0.01
0.59 ± 0.13
0.49 ± 0.03
0.35 ± 0.01
0.20 ± 0.01
0.19 ± 0.04
Film wear rate, Kfilm
(mm3/Nm)
–
–
9.9 × 10−6 ± 3.3 × 10−6
4.4 × 10−6 ± 2.1 × 10−6
2.0 × 10−6 ± 1.7 × 10−6
7.2 × 10−8 ± 8.6 × 10−9
4.4 × 10−8 ± 1.0 × 10−8
Ball wear rate, Kball
(mm3/Nm)
6.5 × 10−7 ± 3.5 × 10−8
6.1 × 10−7 ± 8.6 × 10−8
8.8 × 10−6 ± 9.5 × 10−6
1.6 × 10−5 ± 1.3 × 10−6
1.7 × 10−6 ± 2.2 × 10−6
2.1 × 10−8 ± 1.4 × 10−8
3.2 × 10−8 ± 3.4 × 10−9
LC1
(N)
14.8 ± 2.5
12.3 ± 0.5
13.6 ± 2.7
14.1 ± 4.0
–
–
–
16
LC2
(N)
68.3 ± 6.4
63.0 ± 2.1
31.5 ± 4.5
27.9 ± 7.0
17.4 ± 2.4
23.0 ± 1.8
45.6 ± 1.
Figure 1
17
Figure 2
18
Figure 3
19
Figure 4
20
Figure 5
21
Figure 6
22