Tribological study of hydrogenated amorphous carbon films with tailored microstructure and composition produced by bias-enhanced plasma chemical vapour deposition

Diamond & Related Materials 19 (2010) 1093–1102

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Diamond & Related Materials j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / d i a m o n d

Tribological study of hydrogenated amorphous carbon films with tailored microstructure and composition produced by bias-enhanced plasma chemical vapour deposition J.G. Buijnsters a,b,⁎, M. Camero a, L. Vázquez a, F. Agulló-Rueda a, R. Gago a, I. Jiménez a, C. Gómez-Aleixandre a, J.M. Albella a a b

Instituto de Ciencia de Materiales de Madrid (CSIC), c/ Sor Juana Inés de la Cruz 3, 28049 Madrid, Spain Institute for Molecules and Materials, Radboud University Nijmegen, Heijendaalseweg 135, 6525 AJ Nijmegen, The Netherlands

a r t i c l e

i n f o

Article history: Received 16 September 2009 Received in revised form 10 March 2010 Accepted 23 March 2010 Available online 30 March 2010 Keywords: Hydrogenated amorphous carbon (a-C:H) Tribology Microstructure Graphitization Fullerene Friction Wear Plasma-enhanced CVD

a b s t r a c t We investigated the mechanical and tribological properties of hydrogenated amorphous carbon (a-C:H) films on silicon substrates by nanoindentation, ball-on-disc tribotesting and scratch testing. The a-C:H films were deposited from an argon/methane gas mixture by bias-enhanced electron cyclotron resonance chemical vapour deposition (ECR-CVD). We found that substrate biasing directly influences the hardness, friction and wear resistance of the a-C:H films. An abrupt change in these properties is observed at a substrate bias of about −100 V, which is attributed to the bias-controlled transition from polymer- to fullerenelike carbon coatings. Friction coefficients in the range of 0.28–0.39 and wear rates of about 7 × 10−5 mm3/Nm are derived for the polymeric films when tested against WC–Co balls at atmospheric test conditions. On the other hand, the fullerenelike hydrogenated carbon films produced at ion energies N 100 eV display a nanohardness of about 17 GPa, a strong reduction in the friction coefficient (∼ 0.10) and a severe increase in the wear resistance (∼ 1 × 10−7 mm3/Nm). For these films, relative humidity has a detrimental effect on friction but no correlation with the wear rate was found. © 2010 Elsevier B.V. All rights reserved.

1. Introduction Hard carbon coatings display a number of highly interesting material properties that have led to their use in a wide range of tribological applications. Their high hardness, high wear resistance, low friction and strong corrosion resistance allow using them as protective layers on bearings, gears, engine components and even beverage bottles [1–3]. The amorphous carbon bonding structure is characterized by the presence of tetrahedral sp3-diamondlike and trigonal sp2-graphitic bonds. Based on the relative percentage of these sp3–C and sp2–C bonds and the amount of bonded hydrogen within the film microstructure, four types of a-C:H materials have been distinguished [4], i.e. polymerlike (PLCH), graphitelike (GLCH), diamondlike (DLCH or DLC) and tetrahedral (TACH) hydrogenated amorphous carbon. PLCH displays the highest H content and its polymerlike structure governed by the H terminated sp3 bonds results in a low density and soft matrix. GLCH combines a low H content with a high sp2 content. TACH films have a relatively low H content (25–30 at.%) and a large sp3–C content, ⁎ Corresponding author. Present address: Katholieke Universiteit Leuven, Department of Metallurgy and Materials Engineering, Kasteelpark Arenberg 44, 3001 Leuven, Belgium. Tel.: + 31243653012; fax: + 31243653311. E-mail address: [email protected] (J.G. Buijnsters). 0925-9635/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.diamond.2010.03.017

whereas the content of sp3–C in DLC films can vary significantly as a result of the varying contents of atomic hydrogen present in the film matrix. Young's modulus and density of DLC films are generally much lower than those of TACH films. Within the class of DLC coatings, the bonding structure and composition can vary largely depending on the applied source gases as well as on the deposition techniques and growth conditions. As a result, the mechanical and tribological properties of DLC films differ substantially from one to another [5,6]. The ion energy and ion density are crucial factors in the plasma growth of a-C:H films. The ion–(sub)surface interactions such as the dangling bond formation, sputtering, atomic displacement and hydrogen depletion are strongly depending on the ion energy [7]. The application of a radio frequency or (pulsed) direct current substrate bias voltage allows the control of the energy of the ions bombarding the surface. Thus, the substrate bias can be regarded as a decisive parameter controlling the composition, structure, properties and even surface morphology [8–10] of the deposited a-C:H films. Recent advances in the development of alloyed and nanostructured carbon resulted in the production of hard carbon films with even superior friction and wear properties. Doping carbon films with elemental nitrogen is interesting due to its direct effects on the surface roughness and changes in film microstructure [11–13]. The so-called fullerenelike carbon nitride coatings with three-dimensional, cross-

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linked graphite sheets proved to supply an extreme elasticity and excellent tribological performance [14]. In this respect, we recently demonstrated that also hydrogenated carbon films with a fullerenelike (FL) structure and self-formed C60 inclusions can be produced by bias-enhanced electron cyclotron resonance chemical vapour deposition (ECR-CVD) [15]. Due to the low surface energy, spherical shape, high load bearing capacity and high chemical stability of the embedded C60 molecules, combined with the structural and chemical nature of the bulk, these fullerenelike hydrogenated carbon films are very promising for low-friction and tribological applications. Please note that these fullerenelike carbon films differ from the ordered onionlike carbon materials typically obtained by other methods, e.g. high-temperature annealing of diamond nanoparticles, electron irradiation of carbon soot, carbon ion implantation, or low-pressure fuel-rich flames. In contrast to our dense fullerenelike carbon films which are synthesized by plasma CVD, the onionlike quasi-spherical nanoparticles do not contain significant quantities of bonded hydrogen. In this study the tribological behaviour of a-C:H films with tailored microstructure and hydrogen content is discussed. More specifically, the effect of the application of substrate biases in the range of −300 to +100 V on the mechanical and tribological properties of a-C:H films deposited by ECR-CVD is investigated. We have chosen a gas mixture of argon/methane because (ultra)smooth carbon films with relatively high H content are readily obtained in the ECR plasma for ion energies larger than 100 eV [10]. In previous works [15,16] we already showed that there is a strong variation in structure and composition of the deposited carbon films over the substrate bias ranging from +100 V down to −300 V. Therefore, we will denote the a-C:H films obtained at substrate biases above −100 V as polymerlike hydrogenated carbon (PLCH) films. The coatings produced at more negative substrate bias (b−100 V) are referred to as fullerenelike hydrogenated carbon (FLCH) films. In a previous study [15] we showed that the FLCH films contain self-formed C60 inclusions and that high-resolution transmission electron microscopy analysis of these films indicates the existence of curved graphitic planes. The presence of C60 inclusions and the possible existence of a network of curved graphitic planes for a-C:H films grown under similar plasma conditions were reported in earlier works [17,18] as well. Note that the FLCH films described here are produced under growth conditions comparable to those applied by others for the production of the hard hydrogenated carbon coatings generally referred to as DLC; not to be confused with the strongly graphitic GLCH films that are generally produced at much higher bias and that contain significantly less hydrogen. We would like to stress that the present work is aimed towards completing our studies on the transition from PLCH to FLCH material by supplying a detailed analysis of the differences in tribological behaviour between the different types of a-C:H films rather than being a conclusive study on any effect of incorporated C60 and fullerenelike arrangements in the tribological process. Since the C60 inclusions are even present in small quantities only [15], it is also extremely difficult to ascribe the elastic behaviour of our FLCH films to the excellent elastic properties of the C60 inclusions. 2. Experimental details a-C:H films have been grown by ECR-CVD (ASTEX, mod. AX4500) in a two-zone vacuum chamber operating with a 2.45 GHz microwave plasma source at 208–210 W input power. Gas mixtures of methane/ argon are used with a total flow rate of 50 sccm keeping the operating pressure at 1.1 · 10−2 Torr. The argon gas (35 sccm) is fed to the upper discharge chamber while methane (15 sccm) is introduced into the process chamber about 10 cm from the substrate holder. 1-Hour deposition runs were carried out on 15 × 15 mm2 p-type silicon (100) substrates (280 µm thickness, double-side polished) without film delamination. Prior to deposition, the substrates were cleaned ex-situ

with trichloroethylene, acetone and ethanol ultrasonic baths and insitu by Ar ion bombardment at a direct current (DC) bias of −50 V for 6 min. During film deposition a DC bias varying from −300 to +100 V was applied to the silicon substrates while no intentional heating was employed. The temperature reached at the end of each deposition run never exceeded 120 °C and the thickness of the grown a-C:H films varies between about 630 and 1150 nm (i.e., the deposition rate varies between 0.18 and 0.32 nm s−1) depending on the applied substrate DC bias. Nanoindentations using a Nanotest system (MicroMaterials, Ltd.) applying a Berkovitch diamond indenter were performed to obtain the hardness and elastic modulus of the a-C:H films. Young's modulus and nanohardness values are derived from the slope of the unloading curves at the maximum loads and the projected contact areas [19]. The analysis is based on analytical solutions accounting for the curvature of the unloading data and the indenter shape function to establish the contact area at peak load. Load–displacement curves were recorded at 27.4 ± 0.1 °C never exceeding a maximum displacement of about 10% of the total film thickness (to eliminate substrate effects) and maintaining a dwell time of 5 s at maximum load. Tribotesting of the coatings was performed by ball-on-disc experiments using a TPD/10 Microtest® system with mirror-polished, 3-mm diameter WC/Co balls. Previous to each test, a new WC/Co spherical tip was selected. The ball-on-disc test parameters include the normal load, sample rotation speed, relative humidity and temperature. The apparatus was covered by an isolation box to control the test atmosphere, including the relative humidity. Unless mentioned otherwise, normal loads of 3 and 5 N at rotation speeds of 375 rev/min, a relative humidity (RH) of about 23% and a fixed temperature of 22 ± 1 °C were applied. The radius of the wear tracks was 1.0 mm and the relative sliding speed between contacting surfaces was about 40 mm s−1. The Hertzian contact pressures for the applied normal loads are 1.25 and 2.35 GPa, respectively. The coefficient of friction (CoF) was measured continuously during each test. Explorative experiments were programmed to stop once the monitored friction value reached that of a Si–WC/Co sliding contact (∼0.55), i.e. at complete penetration of the coating. Wear rates were obtained from the volume loss of the a-C:H film material during new ball-on-disc tribotests, which were stopped prior to coating failure. The wear rates were then calculated from the width and depth of the wear track measured with a profilometer, normalized to the applied load and the total distance covered by the tip. The width and depth of the wear tracks were measured with a Dektak 3030 profilometer with a resolution of ∼5 nm. Scratch tests were performed on a Revetest setup (CSM Instruments) with a diamond Rockwell indenter (200 µm in diameter). Scratching was carried out with normal loads in the range of 0.9–20 N in atmospheric conditions (22 °C and about 55% RH) using loading rates of 1.5 and 4.0 N min−1 and a fixed scratch length of 3.0 mm. The normal load at which chipping at the edges of the scratch initiates, is mostly defined as a representative, quantitative measure for the adhesion value of the film [20] and was taken as the critical load. Though the acoustic emission, penetration depth and tangential force were measured during each scratch test as well, the critical loads were derived from careful optical examination of the scratch tracks only. Scanning electron microscopy (SEM; Hitachi S-2700) was used to study the wear tracks produced by the ball-on-disc tribotests. Atomic force microscopy (AFM) measurements (using Nanoscope IIIa (Veeco) equipment operating in tapping mode with silicon cantilevers (nominal radius of 10 nm)) were performed to derive the root mean square (rms) surface roughness of the as-grown a-C:H films using typically 4 × 4 µm2 scan areas. The rms roughness values are averaged numbers over three measurements made on different locations being about 5 mm apart. Differences in local surface roughness were always less than 10%, indicating a very uniform film roughness for each sample [10]. Also, AFM was employed to analyze

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the centre of the wear tracks. The bonding structure of the asdeposited a-C:H films and the microstructural changes within the wear track and wear debris after ball-on-disc tests were analyzed by visible light (514.5 nm wavelength) micro-Raman spectroscopy using a Renishaw Ramascope 2000 micro-Raman spectrometer. Unpolarized Raman spectra were collected with a 100× objective and a spectral resolution of about 4 cm−1. The relative content of H, Ar, and O elements within the a-C:H films has been derived from ion beam analysis techniques. In particular, simultaneous elastic recoil detection analysis (ERDA) and Rutherford backscattering spectrometry (RBS) were performed (see Ref. [16] for details).

3. Results 3.1. a-C:H film characteristics In Table 1, the film thickness, rms surface roughness, hydrogen content and critical load (adhesion value) of a selection of a-C:H films are listed as a function of the applied DC substrate bias, Vb. Like we reported in detail elsewhere [10,15,16], two growth regimes separated by a relatively sharp transition at about −100 V can be distinguished. In the cauliflowerlike surface morphology regime (regime A), i.e. from +100 V ≥ Vb ≥ −80 V, the film roughness varies from 4.7 to 5.7 nm, whereas rms roughness values of 0.13–0.20 nm are derived for the films grown with −120 V ≥ Vb ≥ −300 V (regime B) [10]. The hydrogen content of the cauliflowerlike PLCH films is relatively high (N40 at.%), whereas it is close to 30 at.% for the ultrasmooth FLCH films obtained at Vb b −100 V [15]. A clear distinction in adhesion values between regimes A and B is evident as well. Values close to 20 N are observed for regime B, whereas much lower values are obtained for regime A. Fig. 1. Sets of ten load–displacement curves of the PLCH and FLCH films grown at (a) 0 V and (b) −300 V, respectively. Indicated are the loading and unloading curves, the maximum load, Pmax, and the maximum displacement, hmax.

3.2. Nanoindentation Nanoindentation measurements demonstrate specific differences between growth regimes A and B. In Fig. 1, a series of ten load– displacement curves of the PLCH and FLCH films grown at 0 V (a) and −300 V (b), respectively, are shown. Each individual indentation has been performed on a fresh surface spot. For the PLCH film grown at floating bias, only a load of 0.3 mN is sufficient to result in indenter displacements of about 100 nm, whereas a ten-fold load of 3 mN is needed to indent the FLCH coating to the same extent. Next to the difference in maximum load, also the elastic recovery, which is defined as the difference between the displacements at maximum load and permanent displacement after unloading, varies significantly between the two bias regimes. The poorer reproducibility (less overlap between different curves) for the PLCH films (Fig. 1a) might be ascribed to the relatively high surface roughness and pile-up of film material around the indent [19]. For the FLCH films of regime B the elastic recovery is in the order of 80%, significantly higher than the value of about 68% derived for the PLCH films of regime A.

Moreover, the nanoindentation measurements clearly show that a drastic change in nanohardness (H) and Young's modulus (E) between both regimes is observed (Fig. 2). In regime A, the values of H and E vary from 1.4 to 2.4 GPa and from 22 to 34 GPa, respectively. The H/E ratio varies between 0.05 and 0.08. A significant increase in both nanohardness and Young's modulus is observed for Vb ≤ −120 V. In regime B, the nanohardness reaches 18 GPa in

Table 1 Thickness, rms surface roughness, H content and critical load as a function of the applied DC substrate bias for a selection of a-C:H films. DC bias (V)

Thickness (nm)

Roughness (nm)

H content (at.%)

Critical load (N)

+ 100 + 50 0 −50 −80 −120 −200 −300

599 687 834 864 1003 995 1120 951

4.8 5.7 5.3 4.9 4.7 0.20 0.13 0.17

47 46 45 41 40 32 32 29

0.9 1.1 1.1 2.8 4.3 18 20 18

Fig. 2. Values for the nanohardness (■) and Young's modulus (○) of the a-C:H films as a function of the applied DC bias. Indicated are growth regimes A (PLCH) and B (FLCH). The dashed line indicates the transition between both regimes.

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combination with a Young's modulus of about 155 GPa. The H/E ratio for these FLCH films is about 0.12. The sample grown at −120 V shows a relatively low nanohardness and a value for H/E of only 0.06, which can be attributed to being at the microstructural transition from PLCH to FLCH films (dashed line). 3.3. Ball-on-disc tribotests To show the difference in tribomechanical behaviour of the a-C:H films between both bias regimes, the CoFs recorded during explorative tribotests on the films deposited at + 100 and −200 V are displayed in Fig. 3(a) and (b), respectively. For both curves three different regions can be distinguished. A very short initial stage (up to about 50 rev) acts as the running-in period [21]. The rotating disc is accelerating and the CoF reaches a steady-state value. Due to the relatively small surface roughness of all a-C:H films with respect to the large ball contact, steady-state tribotest conditions are rapidly attained. Then, the second stage is made up by a steady-state situation for which the CoF stays constant for the major part of the test. Averaged values of the CoF obtained for this second stage are considered as the equilibrium CoF. Finally, a third stage is observed for which a sudden minor change in CoF is detected before the CoF abruptly jumps to a value of 0.55. This point, prior to full penetration of the a-C:H coatings, can be considered as the moment of coating failure. At 3 N normal load, the film grown at +100 V is already fully penetrated after about 400 rev (Fig. 3a), whereas the film deposited at −200 V only fails under more severe test conditions, i.e. at 650 rev/ min and 20 N normal load (Fig. 3b). In this case, coating failure takes place after about 23,000 rev, corresponding to a distance of about 150 m. Also, note the large difference in absolute value of the CoF

between both types of film. The ultrasmooth and hard FLCH film grown at −200 V displays a steady-state CoF of about 0.095, whereas a 4-fold value is obtained for the PLCH film grown at +100 V. In Fig. 3(c) and (d) the wear profiles corresponding to the tribotests of Fig. 3(a) and (b) are shown. These profiles are recorded immediately after coating failure and, thus, material removal of the soft silicon substrate is kept to a minimum. The wear profile for the sample grown at +100 V (Fig. 3c) is characterized by rather steep sidewalls and a flat bottom. Film material is pushed sideways and is piled up along the edges of the wear track. Together with the very low number of cycles needed for coating failure we can conclude that the soft PLCH film is unable to withstand the relatively high Hertzian contact pressure of 1.25 GPa, being close to the derived value of its nanohardness. On the other hand, the wear track for the hard and ultrasmooth FLCH film grown at −200 V (Fig. 3d) exhibits more shallow sidewalls. Here, piled-up film material and wear debris can be found over a rather large distance aside the wear track. The two lateral minima within the wear track are probably produced by the debris that accumulates at both sides of the contact point between ball and film, once the soft silicon substrate is reached [21]. Next, ball-on-disc tests with loads of 3 and 5 N were then repeated several times up to a pre-selected number of revolutions lower than the failure limit. Each test was performed on a fresh area of the coating and the wear track was measured with the profilometer. A selection of tested samples was also analyzed by SEM, AFM and Raman spectroscopy. The CoF during the steady-state stage is taken as the representative CoF for each test, yielding an error bar of about 0.005 for the films of growth regime B and 0.01 for regime A. Averaged values of the CoF and wear rate derived from 2 up to 4 tests are presented in Fig. 4 together with those from several reference

Fig. 3. Plots of the coefficient of friction versus number of revolutions during the ball-on-disc tribotests for the a-C:H films grown at + 100 V (a) and −200 V (b), respectively. Note that under the applied test conditions (normal loads and rotation speeds of 3 N/20 N and 375/650 rev per minute, respectively) coating failure (i.e. film penetration) leading to the programmed stopping takes place in both tests. (c, d) Wear track profiles of the PLCH and FLCH films corresponding to tests (a) and (b), respectively.

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Fig. 4. Wear rate (upper panel) and coefficient of friction (lower panel) as a function of the applied DC substrate bias. Growth regimes A and B and the transition (dashed line) at about −100 V are indicated together with the values for silicon, BC0.3N, evaporated carbon and ta-C references [22].

Fig. 5. Wear rate (upper panel) and coefficient of friction (lower panel) of the FLCH film deposited at −200 V as a function of the relative humidity. Tests were carried out at 5 N normal load. The dashed lines are guides to the eye.

coatings tested under similar conditions in our laboratory. Again, the outcomes for growth regimes A and B differ remarkably. The relatively soft PLCH films from regime A display CoF values of 0.30–0.40, slightly increasing for more positive bias. These films also exhibit a poor wear resistance which is only slightly better than the silicon substrate material. On the contrary, the higher hardness and Young's modulus of the FLCH films of regime B are expressed by the drastically reduced wear rates (∼ 1 × 10−7 mm3/Nm). Moreover, the CoF derived for these films is much lower than those obtained for regime A and reaches values similar to hydrogen-free tetrahedral amorphous carbon (ta-C) [22].

(b) display SEM images of the wear tracks corresponding to the a-C:H films grown at +100 and −200 V during coating failure and prior to it, respectively. Different spots are indicated by numbers and are treated below. The wear track for the soft PLCH film (Fig. 6(a)) already reveals grinding lines by the WC–Co ball in contact with the soft silicon substrate, but it does not show large amounts of wear debris within the wear track (1). Only, coating material piled up around the edges of the wear track is observed. The soft a-C:H film material is easily pushed outwards and the coating is penetrated after a very short number of cycles. On the other hand, for the FLCH film deposited at −200 V (Fig. 6(b)), the presence of wear debris particles inside (2) and around (3) the wear track is manifest. Particularly, close to the edges of the wear track huge amounts of wear debris (3) are collected and these particles vary in size from about 0.5 to 8 µm. Within the track itself, debris particles are observed that have been trapped at the sliding contact interfaces and have undergone severe grinding action. Moreover, in few localized areas within the wear track we can observe the formation of microholes (Fig. 6(c)), indicating that delamination wear might play a significant role in the tribotest, although the number of microholes is rather small. Apart from these microholes and trapped debris particles, the wear track itself is extremely smooth, like its film surface roughness prior to testing. In order to get more detailed information on this ultrasmooth wear track we employed AFM analysis. In Fig. 7(a) and (b) typical AFM images of the FLCH film are displayed before and after the wear process, respectively. The as-deposited film morphology is very flat (rms roughness in the 0.1–0.2 nm range) and featureless. In contrast, the morphology of the worn film surface presents parallel shallow grooves, 1–4 nm deep, with a wavelength in

3.4. Tribology of fullerenelike hydrogenated carbon films As mentioned earlier, the FLCH films obtained in growth regime B are ultrasmooth and hard. This makes these layers very interesting for tool industry applications, with the possibility of performing high speed machining in air without the aid of lubricants. Therefore, we have studied the effect of the relative humidity on the CoF and wear rate for these types of films as well (Fig. 5). The effect of the relative humidity on the CoF is profound (lower panel). At nearly saturated moist conditions the CoF is about 0.15, but for decreasing humidity the CoF drops significantly reaching values as low as 0.02 in dry air. Although the relative humidity has a strong effect on the friction of the FLCH films, the influence on the wear rate is less clear (upper panel). In order to get more information on the underlying phenomena during tribotesting, the wear tracks were analyzed in detail. Fig. 6(a) and

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Fig. 6. SEM images of the wear tracks corresponding to a-C:H films grown at + 100 V (a) and −200 V (b, c) during coating failure and prior to it, respectively. The sliding direction is indicated by the white arrows. The numeric labels correspond to representative spots on the wear tracks (see text).

Fig. 7. 4 × 4 µm2 AFM images taken on a FLCH film grown at −200 V before (a) and after (b) undergoing a ball-on-disc wear experiment. (c) Logarithmic plot of PSD(k,t) functions versus wave number k, corresponding to the FLCH film before (solid line) and after (-○-) the wear test. The PSD curves were calculated from AFM measurements made at 3 different locations on the pristine film surface and on the wear track produced on the sample, respectively. Image (b) and PSD curve (-○-) were measured on the wear track produced on the sample.

the 1.5–2 µm range. The rms roughness increases up to 0.9 nm. Evidently, the wear process changes the film surface morphology over distances equal or larger than the groove wavelength, i.e. 1.5–2 µm. Further insight on the way in which the wear process alters the film

morphology inside the track can be obtained by analyzing the Power Spectral Density (PSD) functions [10] of both images (Fig. 7c). The PSD function is defined as PSD(k,t) = 〈 H(k,t)H(−k,t) 〉, with the function H(k,t) being the Fourier transform of the surface height

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function h(r,t), which represents the local surface height at position r and time t. Clearly, both PSD curves overlap for k N 0.003 nm−1, i.e. for length scales smaller than 350 nm. For larger distances (smaller k values) the PSD function corresponding to the film subjected to wear is higher than that of the as-deposited film. This increase is clearly related to the groove structure produced during the wear test. Eventually, we attempted to indent the as-deposited FLCH film with the silicon tip, with a spring constant close to 40 N/m, employed for the AFM measurements. Thus, we realized single force curve experiments where the maximum deflexion was 83 nm, i.e. a maximum load of 3.2 µN. Afterwards, we imaged the surface morphology and we did not detect any change on the morphology. This is consistent with its high value of the Young's modulus (∼155 GPa) derived from the separate nanoindentation measurements described above. In Fig. 8, SEM images of different areas on the ball contact for the tested FLCH film produced at −200 V are shown. Fig. 8(a) displays the overall picture of the material that is left on the WC–Co ball after the ball-on-disc test. A transfer layer is formed at the contact area (4) that extends along the sliding direction to a much thicker piled-up film at the back part (5) of it. Particles varying in size from about 1 to 20 µm are spread around the contact area, mainly distributed as side bands (6) along the sliding direction. The very thin transfer layer at the contact area (Fig. 8(b)) reveals the presence of grooves running in the sliding direction and some light and dark spots related to trapped wear debris particles. A detailed SEM image of the wear particles present in the side bands (6) is given in Fig. 8(c). Most of these wear particles are curved and are relatively thick (up to 0.7 µm). Thus, entrapped wear debris is crushed and pressed at the back part of the contact area and accumulates as a paste (5). Then, as a result of the backward movement the compacted debris is pushed away and forms individual wear particles at both sides of the ball. To obtain a more complete view of the tribological process, we used micro-Raman spectroscopy to study the changes in the bonding structure for the FLCH film. In Fig. 9 the Raman spectra taken at the spots indicated by the numbers in Figs. 6 and 8 are displayed. The Raman spectrum of the as-grown FLCH film (f) in the range of 800 to 2000 cm−1 consists of two broad bands centred at 1545 cm−1 and about 1360 cm−1 , which correspond to the G and D bands characteristic for hard a-C:H films [23]. The Raman spectrum taken at the centre of the wear track (2) shows a similar profile but a slight shifting of the G peak towards higher wave numbers can be observed. This is even more pronounced for the spectra taken from the wear debris particles present on the film (3) and ball (6). Here, the G band shifts towards 1560 and 1583 cm−1, respectively, and the D band can be distinguished as a broad shoulder or as a local maximum. Also, the transfer film on the ball contact (in-between (4) and (5)) shows a clear shift of the G band despite the very low intensity of the Raman signal, probably as a result of the small layer thickness. In Fig. 10, the G peak position together with the ratio between the D and G band intensities (ID/IG) are shown. These data are obtained from the Raman spectra displayed in Fig. 9 by fitting D and G bands with two Gaussians. There is a clear correlation between both parameters as they both shift to higher values in the order: film (f) → centre of wear track (2) → wear debris on film (3) → transfer film on ball (5) → wear debris on ball (6). Both the G peak shift and increasing ID/IG are indicative for the graphitization of the FLCH film material, as will be discussed below.

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Fig. 8. SEM images of the contact area on the surface of the WC–Co ball used for the testing of the FLCH film grown at −200 V. The numeric labels (see text) indicate representative spots on the ball contact. (a) Overall view of ball contact, (b) transfer layer at spot (4) and (c) debris particles at (6).

4. Discussion 4.1. Effect of DC substrate bias In previous works [10,15,16,24] we demonstrated that the application of an externally applied DC substrate bias during the ECR-CVD process has a strong effect on the morphology, roughness and bonding

structure of the produced a-C:H films. In the present work, we have demonstrated that the two bias regimes leading to a-C:H films with highly different compositions and microstructures and separated around −100 V can be distinguished by clear differences in the mechanical and tribological properties as well. The bulk properties of

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So, whereas in previous works [28,29] a smooth transition between soft PLCH and hard DLCH films as a function of radio frequent (RF) substrate bias or RF plasma power is reported, we find a threshold between PLCH and FLCH films at about −100 V when applying a DC self-bias to the substrate directly. The ion energy distribution (IED) of argon and methane ECR plasmas at pressures above 0.15 Pa can be very narrow with a full width at half-maximum in the order of several eV [25]. A DC self-bias can also be generated at the substrate by the application of a RF power to it [30,31]. However, much broader IEDs are produced when applying such a RF substrate bias [30,31] or RF plasmas [32]. Therefore, it is not surprising that sharp transitions or threshold behaviour are generally not observed. 4.2. Mechanical properties

Fig. 9. Micro-Raman spectra taken at the spots indicated by the numeric labels (see text) in Figs. 6(b) and 8(a). Due to the extremely weak Raman signal from the transfer layer on the ball contact, the corresponding data are multiplied by a factor of 20 for representation. The dashed line indicates the relative shift in G peak positions.

the a-C:H material are determined by the operating film growth and (sub)surface processes initiated by the incoming hydrocarbon and argon ions. The a-C:H films are generally dielectric materials and the application of a decreasing DC bias divides the deposition conditions into roughly 2 growth regimes: A and B. For regime A (+100 V ≥ Vb ≥ −80 V), the a-C:H films are grown at rather low (or even floating) potential. The bulk of arriving ions will have an ion energy in the order of 15–20 eV [25], which implies that the impacting ions will only have direct consequences at the very surface, as the energy is insufficient to penetrate into subsurface layers. For this situation, a-C:H film growth is mainly determined by the physisorption, chemisorption, and incorporation of the hydrocarbon growth precursors. The ioninduced surface mobility and sticking of radical growth precursors seem to increase for decreasing bias, since continuously enhanced film growth rates are observed for more negative Vb. A sudden change of growth mode (regime B) occurs at a certain limit for Vb (i.e., −100 V), where the growing a-C:H film is subjected to ion bombardment controlled by the DC bias. Here, the ions will impact with relatively high ion energy and also initiate more subsurface processes such as implantation, atomic displacement, hydrogen depletion, and sputtering of bonded film atoms [26,27].

Fig. 10. Raman G peak position (■) and ID/IG ratio (○) for different spots on the wear track and ball contact of the FLCH film grown at −200 V after ball-on-disc tribotesting. Data are obtained from the micro-Raman spectra (Fig. 9) by fitting G and D peaks to two Gaussians.

The values of the hardness and Young's modulus of the FLCH films are similar to the numbers reported for DLC films grown from linear inductively coupled Ar/CH4 plasmas [33] and ECR methane plasmas [34] at RF and DC negative substrate biases larger than 100 V. The H/E ratio appears in the so-called “plasticity index” and is considered a valuable measure in evaluating the limit of elastic behaviour in a surface contact [35,36], essential for accurate wear resistance. The high H/E values for our FLCH films are also comparable to those reported for DLCH films produced by plasma-enhanced CVD [33,37] and support the high wear resistance of these types of films. The elastic recovery (∼80%) for our FLCH films is higher than the maximum value (68%) measured for hydrogen-free DLC films grown by DC magnetron sputtering [38], but comparable to the values (70– 90%) for DLC films produced by a CH4/Ar RF plasma at negative pulsed high-voltage-bias [39]. It is, therefore, extremely difficult to ascribe the elastic behaviour of our FLCH films to the excellent elastic properties of the C60 inclusions, which are even present in small quantities only [15]. It is likely that the combination of a dense film matrix consisting of a mixture of sp2 and sp3 bonded carbon together with the self-formed C60 particles gives rise to the observed mechanical and tribological characteristics. 4.3. Wear process It is well known that the wear mechanism and wear rate of films tested under different conditions can vary significantly [5,40,41]. Although the underlying friction and tribological mechanisms operating at the sliding contact faces are not fully understood yet, the common theory of the wear mechanism of a-C:H films includes the contribution of three components [6,40]: i) hydrogen passivates dangling bonds on the film surface and thus permits only weak interactions at the sliding interface, ii) friction-induced shear forces facilitate the breaking of C–H and C–C bonds and the transformation into graphitic structures and iii) lubrication effects are present due to the formation of a transfer film on the counter wear surface. The absence of catastrophic mechanisms like film cracking or spallation, the nearly constant CoF, the extremely smooth wear track and reproducibility in performance during ball-on-disc tribotesting of our FLCH films indicate a consistent wear mechanism. The volumetric wear rates for our FLCH films are comparable to those reported for DLC in literature [42,43]. The steady-state condition can be ascribed to the formation of protective transfer layers, i.e. graphitized tribolayers, on the wear surfaces of both the FLCH film and the WC–Co ball. Our Raman analysis (Figs. 9 and 10) evidences graphitization within the wear track and for the transfer layer and wear particles on the ball contact. The film surface can gradually reconstruct and transform to an energetically more stable graphitelike structure as a result of thermal and strain effects from the repeated friction [44]. This graphitization process is accompanied by hydrogen desorption and under sliding contact the phase transition can take place at much lower temperatures than needed for thermal annealing alone [45].

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Therefore, we believe that the wear behaviour for our FLCH coatings is based on a tribochemical reaction of the coating material into a graphitic transfer layer with lower elastic constants. The increasing degree of graphitization going from the original FLCH film, via the centre of the wear track and wear debris, towards transfer layer and wear particles on the ball contact, as measured by Raman spectroscopy, supports this hypothesis. Wear debris that is piled up at the back part of the contact area and that after severe accumulation makes up a paste before breaking up into smaller fragments, experiences repeated strain and heat effects, which in turn explains the higher degree of graphitization. The AFM analysis of a worn FLCH film (Fig. 7) shows that shallow grooves that run parallel to the sliding direction and that are typical for sliding contacts, are produced on the micron scale. These are likely generated by the action of hard debris particles which play a role as abrasives. However, for submicrometric scales the wear process does not alter the surface morphology. This finding would imply a sort of layer by layer wear process.

therefore, little effect of irregular surface contacts is expected. Noticeably, the friction and wear processes do not induce any roughness at submicrometric scales. Indeed, the running-in period is extremely short and the steady-state is rapidly attained. Intrinsically, the hydrogen content of the a-C:H films undoubtedly determines the frictional properties to a large extent. Passivation of dangling bonds by the formation of C–H bonds reduces the attractive forces between the two surfaces in sliding contact. Interaction between C–H bonds at both the film surface and counterface occurs via weak Van der Waals forces, resulting in low friction [49]. However, in the case of direct interaction of C atoms without passivation by hydrogen, the π–π* interaction that corresponds to higher binding energies would cause high friction, as is the case for H-free carbon films [6]. The FLCH films contain about 30 at.% H (see Table 1) and, therefore, under dry nitrogen the weak C–H interaction guarantees very low friction. For increasing humidity, water molecules adsorbing at the sliding surfaces displace the C–H bonds and, thus, increase the attractive forces at the interface (see Fig. 5).

4.4. Friction behaviour

5. Conclusions

As reviewed recently [6], the factors influencing the frictional behaviour of a-C:H films are plenty. Intrinsic factors such as the ratio of sp2-to-sp3 bonded carbon and the amount of bonded hydrogen as well as extrinsic factors like the surface roughness and test conditions have a strong effect on friction [6]. This complicates the direct comparison between results obtained in different studies. Nevertheless, the relations between these factors and the frictional behaviour of our FLCH films can be explained as follows. The CoFs (0.3–0.4) for the PLCH films are very similar to the numbers (0.27–0.41) collected for high molecular weight polyethylene sliding against steel [46]. On the other hand, the CoFs for the FLCH films are close to 0.1 at 23% RH. These values resemble those in earlier works on DLC films sliding against WC–Co [42] and FR601 steel [47] in air but are in contrast to those found for plasma grown FLCH films tested under similar conditions [48]. Fig. 5 shows that ultralow friction is combined with good wear resistance under dry nitrogen conditions, whereas the CoF increases for increasing humidity. This can be attributed to the physical separation of the FLCH interface by weakly adsorbed gas molecules [41]. The surfaces of the sliding contacts interact with oxygen and water molecules. For example, the highly polar water molecules can physically interact with the surface carbon atoms of the FLCH films and those of the transfer layer on the counter ball to form a layer of physisorbed species. In fact, the friction of a-C:H films generally increases with humidity. When oxygen and/ or moisture are introduced, the friction coefficients may increase substantially and in highly moist air, condensed water molecules can give rise to capillary forces that increase friction [6]. The counterface as well as the test environment might give rise to tribochemical interactions and influence the kinetics of the formation and composition of the transfer films. Hence, they highly determine the frictional and wear behaviour of the FLCH coatings. The uncoated WC–Co that was applied as counterface in our ball-on-disc measurements generally acts as a strong carbide former. The generation of transfer films is then fast and with a high degree of coverage and strong bonding [6]. The relative humidity also has a strong effect on the generation of transfer films and the graphitization process of third-body particles. Increased humidity leads to more efficient cooling by water molecules on flash heating at the contact spots and tends to lower the nominal contact temperature [6], thus, resulting in reduced graphitization and thinner transfer layers. The surface roughness of the as-deposited a-C:H films plays a significant role in the friction and wear behaviour as well. Mechanical interlocking between surface asperities can lead to high frictional losses that are most profound during the early stages of ball-on-disc tribotests. The FLCH films are ultrasmooth as measured by AFM and,

We have grown a-C:H films from an argon/methane gas mixture by ECR-CVD applying a direct current bias to the substrate. In the bias range going from +100 to −300 V two growth regimes with an abrupt transition at about −100 V are distinguished. Hydrogen-rich, polymeric a-C:H films are obtained at relatively minor ion energies. These films are soft and display poor tribological behaviour (CoF of 0.28–0.39 and wear rate ∼7 × 10−5 mm3/Nm) in ball-on-disc testing against WC–Co at atmospheric conditions. On the contrary, fullerenelike hydrogenated carbon films are produced at ion energies starting from about 100 eV. These highly elastic and hard coatings display an ultrasmooth surface finish, present a strong reduction in the CoF (∼0.10) and wear loss (∼1 × 10-7 mm3/Nm) and exhibit very low friction (CoF ∼ 0.02) under dry nitrogen. Increasing humidity has a detrimental effect on the friction behaviour of these films, but no direct correlation with the wear loss was found. Detailed investigation of the wear track and ball counterface shows that graphitization and formation of a transfer layer are the most important phenomena related to the excellent tribological performance of the fullerenelike hydrogenated carbon coatings. Acknowledgements Financial support is acknowledged from the European project, FOREMOST, No. NMP3-CT-2005-515840 and the Spanish MEC Project No. FIS2009-12964-C05-04. J.G.B. acknowledges support from MECD through the Juan de La Cierva program, the Dutch Veni grant, and the FWO (Fund for Scientific Research, Flanders, Belgium). The authors would like to thank J. Ortiz for his help with the nanoindentation measurements. References [1] J. Robertson, Mater. Sci. Eng. 37 (2002) 129. [2] R. Hauert, U. Müller, Diamond Relat. Mater. 12 (2003) 171. [3] A. Shirakura, M. Nakaya, Y. Koga, H. Kodama, T. Hasebe, T. Suzuki, Thin Solid Films 494 (2006) 84. [4] C. Casiraghi, A.C. Ferrari, J. Robertson, Phys. Rev. B 72 (2005) 085401. [5] A. Erdemir, G.R. Fenske, J. Terry, P. Wilbur, Surf. Coat. Technol. 94–95 (1997) 525. [6] A. Erdemir, C. Donnet, J. Phys. D Appl. Phys. 39 (2006) R311. [7] I. Ahmad, S.S. Roy, P.D. Maguire, P. Papakonstantinou, J.A. McLaughlin, Thin Solid Films 482 (2005) 45. [8] C. Corbella, E. Pascual, M.A. Gómez, M.C. Polo, J. García-Céspedes, J.L. Andújar, E. Bertran, Diamond Relat. Mater. 14 (2005) 1062. [9] M. Moseler, P. Gumbsch, C. Casiraghi, A.C. Ferrari, J. Robertson, Science 309 (2005) 1545. [10] J.G. Buijnsters, M. Camero, L. Vázquez, Phys. Rev. B 74 (2006) 155417. [11] M. Camero, J.G. Buijnsters, C. Gómez-Aleixandre, R. Gago, I. Caretti, I. Jiménez, J. Appl. Phys. 101 (2007) 063515. [12] J.G. Buijnsters, L. Vázquez, J. Phys. D Appl. Phys. 41 (2008) 012006.

1102

J.G. Buijnsters et al. / Diamond & Related Materials 19 (2010) 1093–1102

[13] L. Vázquez, J.G. Buijnsters, J. Appl. Phys. 106 (2009) 033504. [14] J. Neidhardt, L. Hultman, E. Broitman, T.W. Scharf, I.L. Singer, Diamond Relat. Mater. 13 (2004) 1882. [15] J.G. Buijnsters, M. Camero, R. Gago, A.R. Landa-Canovas, C. Gómez-Aleixandre, I. Jiménez, Appl. Phys. Lett. 92 (2008) 141920. [16] J.G. Buijnsters, R. Gago, I. Jiménez, M. Camero, F. Agulló-Rueda, C. GómezAleixandre, J. Appl. Phys. 105 (2009) 093510. [17] J.Q. Chen, D.L. Meeker, N.N. Barashkov, Appl. Phys. Lett. 71 (1997) 216. [18] Q. Wang, C. Wang, Z. Wang, J. Zhang, D. He, Appl. Phys. Lett. 91 (2007) 141902. [19] S.J. Bull, J. Phys. D Appl. Phys. 38 (2005) R393. [20] K. Vercammen, K. Van Acker, A. Vanhulsel, J. Barriga, A. Arnsek, M. Kalin, J. Meneve, Tribol. Int. 37 (2004) 983. [21] I. Caretti, J.M. Albella, I. Jiménez, Diamond Relat. Mater. 16 (2007) 63. [22] I. Caretti, I. Jiménez, R. Gago, D. Cáceres, B. Abendroth, J.M. Albella, Diamond Relat. Mater. 13 (2004) 1532. [23] A.C. Ferrari, J. Robertson, Phys. Rev. B 61 (2000) 14095. [24] J.G. Buijnsters, L. Vázquez, Surf. Coat. Technol. 201 (2007) 8950. [25] W. Jacob, P. Reinke, W. Möller, Diamond Relat. Mater. 2 (1993) 378. [26] W. Jacob, Thin Solid Films 326 (1998) 1. [27] Y. Lifshitz, S.R. Kasi, J.W. Rabalais, W. Eckstein, Phys. Rev. B 41 (1990) 10468. [28] E. Braca, J.M. Kenny, D. Korzec, J. Engemann, Thin Solid Films 394 (2001) 30. [29] O. Durand-Drouhin, M. Lejeune, M. Benlahsen, J. Appl. Phys. 91 (2002) 867. [30] Q. Zhang, S.F. Yoon, Rusli, J. Ahn, H. Yang, D. Bahr, Appl. Phys. 86 (1999) 289. [31] R.M. Dey, S.B. Singh, A. Biswas, R.B. Tokas, N. Chand, S. Venkateshwaran, D. Bhattacharya, N.K. Sahoo, S.W. Gosavi, S.K. Kulkarni, D.S. Patil, Curr. Appl. Phys. 8 (2008) 6.

[32] E. Kawamura, V. Vahedi, M.A. Lieberman, C.K. Birdsall, Plasma Sources Sci. Technol. 8 (1999) R45. [33] M. Rybachuk, J.M. Bell, Thin Solid Films 515 (2007) 7855. [34] K.H. Lai, C.Y. Chan, M.K. Fung, I. Bello, C.S. Lee, S.T. Lee, Diamond Relat. Mater. 10 (2001) 1862. [35] A. Leyland, A. Matthews, Wear 246 (2000) 1. [36] R.C. Powles, D.R. McKenzie, S.J. Meure, M.V. Swain, N.L. James, Surf. Coat. Technol. 201 (2007) 7961. [37] P. Gupta, Synthesis, structure and properties of nanolayered DLC/DLC films, M.S. Thesis, Louisiana State University, USA, 2003. [38] E. Broitman, N. Hellgren, Zs. Czigány, R.D. Twesten, J. Luning, I. Petrov, L. Hultman, B.C. Holloway, J. Vac. Sci. Technol. A 21 (2003) 851. [39] Y. Ohtsu, H. Noda, C. Nakamura, T. Misawa, H. Fakiyama, C. Dislasu, Surf. Coat. Technol. 201 (2007) 6674. [40] J. Goldsmith, E. Sutter, J.J. Moore, B. Mishra, M. Crowder, Surf. Coat. Technol. 200 (2005) 2386. [41] H.I. Kim, J.R. Lince, O.L. Eryilmaz, A. Erdemir, Tribol. Lett. 21 (2006) 53. [42] C. Tsotsos, A.L. Yerokhin, A.D. Wilson, A. Leyland, A. Matthews, Wear 253 (2002) 986. [43] A. Tanaka, M.W. Ko, S.Y. Kim, S.H. Lee, T. Kumagai, Diamond Films Technol. 8 (1998) 51. [44] Y. Liu, A. Erdemir, E.I. Meletis, Surf. Coat. Technol. 82 (1996) 48. [45] T. LeHuu, H. Zaidi, D. Paulmier, P. Voumard, Thin Solid Films 291 (1996) 126. [46] H. Unal, U. Sen, A. Mimaroglu, Tribol. Int. 37 (2004) 727. [47] F. Rabbani, Surf. Coat. Technol. 184 (2004) 194. [48] C.B. Wang, S.R. Yang, Q. Wang, Z. Wang, J.Y. Zhang, Nanotechnology 19 (2008) 225709. [49] R. Hauert, Tribol. Int. 37 (2004) 991.