Wear, 168 (1993) 143-153 143

Review of the tribology of diamond-like carbon

A. Grill IBM Research llivirion, T..T. Watson Research Center, ~o~t~wn ~e~~ht~, AT 10598 (USA)

(Received July 20, 1992; accepted January 4, 1993)

Abstract

Diamond-like carbon (DE) films are characterized by very low friction coefficients, high wear resistance and high corrosion resistance. Depending upon the testing environment, the coefficient of friction can be as low as 0.01. As-deposited films are wear resistant in vacuum as well as in atmospheric ambient. However, the tribological properties of DLC are strongly affected by the deposition method. This paper reviews the friction and wear properties of DLC, and of similar materials derived from DLC, as a function of the preparation method and testing environment. Mechanisms proposed to explain the tribological properties are presented and discussed.

Diamond-like carbon (DLC) is a term used to describe

hard carbon films which are mostly metastable amor- phous materials but can include a microcrystalline phase. DLC films have been prepared by a variety of methods

and precursors, ~n~~udin~ r.f. or d.c. plasma-assisted chemical vapor deposition (CVD), sputtering, vacuum arc, and ion beam deposition, from a variety of carbon- bearing solid or gaseous source materials [l]. DLC films are characterized by an extreme hardness, which is ‘measured to be in the range 2000-5000 kg mm-’ [2], a generally low friction coefficient and usually very high internal stresses 13-51. As a function of the de- position conditions, the films may contain varying amounts of hydrogen. The films deposited by plasma- assisted CVD (PACVD) usually incorporate up to 60% hydrogen f6], while those deposited by sputtering or

a vacuum are may contain only small amounts of hydrogen or no hydrogen at all.

The high hardness and chemical resistance of the DLC films makes them good candidates as wear-resistant protective coatings for metals, optical, or electronic components. The use of DLC is especially attractive in applications where it is required that the thickness of the protective film be less than 50 nm, as for example in the case of magnetic recording media. One such application is in hard magnetic disks, where the trend towards higher density data storage has led to the requirement of very low flying heights between a disk and a recording head 171. The protective coating must be resistant to wear and corrosion but also thin enough not to impede the achievement of high recording density.

Therefore, special attention has been given to the tribological study of the head-disk interfaces in magnetic recording drives using thin film media, in order to develop reliable, high capacity disk files. During normal operation, the read/write head flies above the disk surface; however, when the disk starts or stops, the slider rubs on the surface of the disk. The friction and wear which can develop between the siider and disk during contact can result in the failure of the recording media. In addition, very little debris can be allowed to form during the start and stop cycles, because the debris can disrupt the flight of the head and lead to the failure of the disk surface.

At the operating conditions of magnetic-recording media and micromechanics in general, microtribology becomes a key technology for interfaces [$I. Microtri- bology addresses the changes taking place at the atomic level in the very superficial layers in contact. Therefore, the characterization of microtribological properties has to be performed at ultralow loads. Microtribological testing tools, such as a contact profilometer [9] or point contact microscope flO] have been developed for this purpose by modi~ing a scanning tunnelling microscope. Measurements have been performed with this tool at loads as small as 1 pg [9].

In microtribology, the required surface material is not a self-sacrifice-type solid lubricant but has to be both wear resistant and lubricating. According to Miyake

and Kaneko [8], to reduce atomic-scale wear, the struc- ture of the tribological material has to satisfy the following requirements.

~43-1#8/93/$6.~ 0 1993 - EIsevier Sequoia. All rights reserved

(1) top layer - shear should always occur at the sliding interface but not within which requires a material of low surface energy as the topmost layer;

(2) bulk of tribological film - fracture formation and defect growth by sliding should be negligible, which requires a high strength material;

(3) interface to substrate - this should have a large adhesive strength, which should prevent delamination of the tribological structure from the substrate.

The third requirement can be fulfilled by a suitable bonding layer, such as a-Si:H, for deposition of DLC on silicide-forming metals [ll], or by grading the in- terfacial layer with a mixture composition of film and substrate, such as Si-C for deposition of DLC on silicon [S]. The other two requirements will be discussed later, showing that DLC, especially the type prepared by PACVD, and materials derived from it can satisfy these requirements.

The properties of DLC films make them useful for a variety of protective applications [12]. Nevertheless, improvement of their tribological properties and ex- panding the use of DLC films for tribological protection in new applications requires a better understanding of their friction and wear properties, and the dependence of those on the deposition parameters and on the operating environment. This is important taking into consideration that the properties of DLC films can be significantly affected by the deposition system (i.e. method, precursor, parameters) [6]. In an effort to improve the tribological properties of DLC films, a variety of related materials have been developed and investigated.

This paper presents a review of tribological studies of hard carbon films and related materials, and discusses the mechanisms proposed to explain these tribological properties.

2. Review of DIE trhology

The structure of hydrogenated DLC can be described as a randbm network of covalently bonded carbon in hybridized tetragonal, (sp’) and trigonal (sp’) local coordination, with some of the bonds terminated by hydrogen [13]. Robertson [14,15] claimed that, in ad- dition to the short-range order defined by carbon hy- bridization and hydrogen content, a substantial degree of medium-range order on about the 10 8, scale also exists in the films. According to Robertson, the DLC can be described as a network of graphitic clusters, linked into islands by sp3 bonds. Sputtered hard carbon films are often mostly graphite-like, while, for certain deposition methods, the hard carbon films are consti- tuted of a mixture of graphite and diamond micro- crystallites. Because of this spectrum of microstructures,

the tribological properties of diftcrent carbon films caii be related to the tribological behavior of graphite 01 diamond. These properties will be first discussed to serve as a reference to the properties of DLC

The friction coefficient of diamond sliding on diamond is relatively low and adhesion plays an important role in determining the friction of diamond on diamond in air. The diamond surface is normally terminated with a layer of chemisorbed hydrogen and oxygen. This layer renders the surface relatively unreactive and prevents strong C-C bonding across the sliding interface. When diamond-diamond sliding experiments are performed in a high vacuum, the sliding action wears off the chemisorbed hydrogen and oxygen from the surface. Strong bonding then occurs across the interface, causing a large increase in the friction coefficient. The friction can be decreased again by bleeding in hydrogen in the chamber containing the sample 1161. During sliding in air, the adsorbates are also removed by sliding but are replenished through adsorption from the environment, and the friction coefficient remains low. It was found that the friction coefficient of diamond on diamond in air can be changed by heating the samples to 100 “C or by exposing it to water [17]. The heating causes the desorption of materials adsorbed on the surface of’ diamond, leading to better interfacial contact and an increased friction coefficient. In contrast, water has the opposite effect, reducing the friction coefficient. Lu- bricating oils do not change the friction coefficient of diamond on diamond [17].

Wear of diamond occurs through chipping of small fragments from the surface, while wear through a surface graphitization mechanism, as a cause for diamond wear, is generally rejected [16]. Even a soft material, such as gold, can cause wear of diamond. This is explained through a fatigue mechanism which causes the formation and growth of small cracks, leading finally to fracture of the diamond particles. There is also evidence that the debris formed during sliding contributes to the low friction in diamond [18].

It has been shown that water acts as a lubricant for graphite and improves its wear and lubricating behavior [19-211. The lubricating action of the water has been explained as a result of a three-body interaction with the adsorbed water, which causes reduced adhesion between the surfaces in contact. Hydrogen has been found to have a similar lubricating effect on graphite. Zaidi et al. [22] investigated the sliding wear of graphite against graphite in an atomic hydrogen atmosphere produced by passing hydrogen over a tungsten ribbon at 2000 K. The authors found that atomic hydrogen is an effective lubricant for graphite and that there is a close analogy between the lubrication of the graphite by water vapor and by atomic hydrogen. It appears that hydrogen passivates dangling bonds at the edges

A. Grill I Tribology of diamond-like carbon 145

of the graphite crystallites, leaving only the possibility of weak interactions with the r bonds, thus resulting in reduced friction.

2.1. Friction coefficients Despite the high chemical inertness of DLC films,

their tribological behavior is controlled to a large extent, as we shall see in the following, by the surface chemistry of the films. The surface chemistry can be affected by the environment. However, it is also dependent on the method used for the preparation of the films, which determines their structure.

In 1980, Enke et al. [3] reported a very low friction coefficient of DLC films deposited by PACVD from acetylene. In nitrogen at a relative humidity (RH) of less than l%, the friction coefficient between a steel ball and a DLC-coated silicon wafer was p = 0.01-0.02. It increased with increasing RH, attaining values of 0.05 at RH= 10% and up to 0.19 at RH= 100%. A similar behavior was reported for films deposited by PACVD from ethylene [23]. The films showed a con- tinuous increase in the friction coefficient with increasing RH and the friction exhibited a hysteresis effect between RH =50% and 100%. Values as low as p= 0.005 were measured in a vacuum [23]. Enke et al. explained the low friction coefficients by assuming that graphitization of the surface layer of the DLC films takes place during testing.

Low friction coefficients were also reported by Weiss- mantel et al. [4] for carbon films deposited by ion beam methods. The films had a friction coefficient of 0.19 against steel, decreasing to 0.04 after 10 000 passes in a pin-on-disk apparatus.

Memming et al. [18] studied the tribological behavior of amorphous hydrogenated carbon (a-C:H) deposited by r.f. PACVD from acetylene, toluene or benzene. Stationary (as referred to by the authors) friction coef- ficients of DLC against steel were found to decrease to very low values, i.e. p cO.02, at RH < 1% and this low value was maintained in ultrahigh vacuum (UHV). The very low friction coefficient was attained after 25 rotations from an initial value of ~~0.2. The friction coefficient remained low in dry nitrogen (p= 0.02); however, it increased drastically to about 0.6 when nitrogen was substituted by dry oxygen. In humid ni- trogen or humid oxygen, the friction coefficient was 0.25. The authors found by Auger electron spectroscopy that, in the conditions of very low friction, carbonaceous material was transferred from the DLC film to the steel rider, while, in the conditions of high friction (in dry oxygen), iron transferred to the DLC surface. The material transfer changed the chemistry of the interface layer and, according to the authors, this ‘tribochemical’ behavior explained the different values obtained for the friction coefficients under various conditions [18].

Memming et al. also found that the loss of hydrogen after annealing the DLC films above 550 “C caused large changes in the friction coefficient, which reached values of ~=0.68 in UHV or dry nitrogen, indicating that the presence of hydrogen in these films was essential for obtaining a low friction in dry nitrogen or IX-IV. However, in a humid atmosphere, the friction coefficient did not change significantly after annealing 1181. Weiss- mantel et al. also found previously [24] an increase of the friction coefficient of DLC upon heating above 400 “C and attributed this increase to a loss of hydrogen rather than to graphitization of the DLC.

Miyoshi and co-workers studied the tribological be- havior of hydrogenated DLC films deposited on silicon nitride by r.f. PACVD of methane or butane in sliding wear against spherical silicon nitride riders [25,26]. The friction coefficient was found to be 0.1 in dry nitrogen and 0.18 in laboratory air. For films deposited at a power density of 0.08 W cme2, the coefficient of friction was found to increase with increasing number of sliding passes in both air and dry nitrogen. In air, however, the film broke through and was removed from the sliding zone after only 1000 passes, probably as a result of poor adhesion. The coefficient of friction of a film deposited at a higher power density of 0.40 W cm-’ decreased, between 10 and 10 000 passes, reaching a value of 0.01 in dry nitrogen. The authors explained the decrease of the friction coefficient above 1000 passes by assuming the formation of a hydrocarbon rich layer at the sliding interface. The differences in the behavior of the films deposited at different power levels were attributed to differences in film densities; however, density measurements were not reported.

According to Miyoshi and co-workers, annealing in vucuo above 500 “C causes hydrogen loss from the surface of the DLC and the formation of a graphite- like layer through a two-step process [25,26], i.e. a carbonization stage in which the film loses hydrogen, followed by a polymerization stage, forming graphitic crystallites or sheets. During heating in z)acuo, the friction coefficient remained low up to 500 “C (~<0.2) but increased strongly at 600 “C (~=0.7), presumably as a result of the assumed graphitization process. The initial friction coefficient in humid air of the films annealed at 700 “C was lower than that in dry nitrogen, the film thus behaving similarly to graphite, which requires water to produce a low friction value [19,20].

It appears from the results presented above that, in hydrogenated DLC films, hydrogen has a similar effect to that in diamond [16] or in graphite [22]. Hydrogen passivates the dangling bonds in the hydrogenated DLC films and permits only weak interactions between DLC and the sliding partner. When hydrogen is lost from hydrogenated DLC by annealing, the dangling bonds created cause strong interactions between the contacting

surfaces, resulting in increased friction in UHV or dry nitrogen, similar to the cases reported for diamond or graphite [16,22]. Humidity reduces the friction coef- ficient to an intermediate value, as it would do for graphite, but also for diamond, although the values are closer to that of graphite.

The tribological properties of DLC films deposited on silicon by r.f. PACVD from a methane and hydrogen mixture have been reported by Kim et al. 1271. The friction was measured against a silicon nitride ball in argon and air at different humidities. The friction coefficient varied from p= 0.2 in 50% humid argon and 100% humid air to ~=0.06 in dry argon. This variation of the friction coefficient was attributed to material transferred between DLC and the ball in the contact area. In dry argon, material was transferred from the DLC to the silicon nitride ball and this reduced the friction coefficient. In dry or humid air, an interfacial layer - identified by microprobe Fourier transform IR as a carbonyl compound (oxidized hydrocarbon) --- formed by a tribochemical reaction covered both the DLC and silicon nitride contact areas and increased the friction coefficient [27].

The effects of the deposition temperature and sub- strate bias on the tribological properties of DLC films deposited by r.f. PACVD from acetylene have been reported by Grill et al. [28]. The DLC films were deposited on silicon wafers at temperatures between 100 and 250 “C and substrate biases of - 80 and - 150 V (d.c.). The films were annealed in z)acuo at tem- peratures up to 590 “C for 3-4 h. The original hydrogen content in the films decreased from 40% in the as- deposited film to 22% after annealing at 590 “C. The friction coefficients were measured in air at RH=40%-70% against a spherical steel rider. The friction coefficients of the as-deposited DLC films were found to decrease from 0.35 f 0.04 for films deposited at 100 “C to 0.20f 0.04 for films deposited at 250 “C, with the surface bias having a negligible effect. Within the experimental error, the friction coefficients of the annealed films were essentially the same as those of the as-deposited films. The insensitivity of the friction coefficients to the annealing temperature was similar to what was observed by Memming et al. [18] and Miyoshi [26] in humid air.

The tribological properties of sputtered DLC films have been reported on by several authors [29-321. Agarwal et al. [29] measured the sliding friction of amorphous carbon films 30 nm thick sputtered on magnetic recording fihns (CoNi and COP) against single- crystal, hemispherical sapphire. The tests were per- formed at a linear speed of l-5 m s- ’ up to 20 000 disk revolutions. After an initial decrease, as a result of surface burnishing, the friction coefficient increased

from an initiar value of 0.3 I .O with increasing number

Marchon nr~d co-worker< tron-sputtered carbon films ceramic sliders of Mn-%n .md (.‘a l’if&. ‘I’he tests were performed under a controlled atmosphere at <I ronstan! drag speed of’ 0.06 m s ‘. In ;t I-cgular atmosphere, the friction coefficient of the sputtered films was found to increase gradually from 0.2 to I.2 during the testing, while. in dry nitrogen, the value of the friction coeficient remained constant at 0.2 up to 5f! h of testing. When oxygen was introduced into the system, the friction coefficient increased to 1.2 but decreased again when oxygen was replaced with nitrogen. 0wing to the gradual increase of the friction coefficient and the drastic effect produced by oxygen - compared with the effect of an inert gas .- the changes in the friction coefficient were attributed to a chemical (corrosive) type of wear,

If the sliding wear is adhesive isear, the frictlonnl force F between the surfaces of the friction couple is given by [-lo]

where A, is the area of contact and S the shear stress at the contact area.

When the surface of the carbon 1s exposed to oxygen, surface oxides can form. In graphite, chemisorbed ox- ygen is desorbed as CO at temperatures above 600 “C [32]. According to Marchon et al. [33], at sliding speeds of 1 m s---’ -- as typically encountered between head and disk during start-stop operations - flash tem- peratures can reach values of several hundreds of degrees centigrade, which are sufficient to cause thermal desorption of CO and CO,. However, at low sliding speeds of 0.06 m SK’ as used in the experimental testing, such high temperatures are not reached. Nevertheless, it is possible that the desorption temperature could be lowered by the catalytic effect of the slider material in contact with the film (as was shown by a 20-fold increase in the reactivity of graphite with oxygen when CaTiO, powder was added to graphite powder). Another explanation for the desorption of the carbon oxides at the lower temperatures could be the direct transfer of energy from mechanical friction to chemical reaction, as suggested by Fischer [35].

The formation and desorption of the carbon oxides is a tribochemical wear which leads to surface smoothing of the sputtered film (in the specific case, from an initial surface roughness of 20 A to about 6 A), a gradual increase in the true contact area A, and an increase in the friction coefficient in air. Tribochemical wear of the carbon surface is thus the cause of the frictional build-up (as observed from 0.2 to 1.2) in an oxygen-containing atmosphere [30]. In contrast, in a clean environment, when the existing surface oxides are removed from the carbon surface through desorption

by the action of the rubbing slider, the clean carbon surface locally reconstructs to saturate dangling bonds and the friction coefficient remains low. The drop in the friction coefficient when oxygen is replaced by nitrogen was explained by Marchon et al. as resulting from a change in the shear strength with the change of atmosphere [33]. The magnitude of the friction drop when oxygen is replaced by nitrogen is not always predictable. A lubricant could reduce the wear of the carbon film by reducing the contact temperature, thus reducing the removal and regeneration rate of oxides, and by serving as a shield, preventing oxygen chemi- sorbtion on the surface.

The tribology of sputtered carbon films can be im- proved when deposited by sputtering in hydrogen-con- taining argon [36], or by exposing the carbon films to an atmosphere of fluorine, which oxidizes the graphite to graphite-fluoride, even at room temperature 1331. In these cases, the hydrogen or fluoride probably block the active sites at the surface which could interact with oxygen.

Similar results were reported by Strom et al. [32] for tests performed on commercial, circumferentially tex- tured unlubricated disks coated with sputtered carbon films 30 nm thick on top of a Co-Pt-Ni magnetic layer. The tests were performed with sliders composed of Al,O,:TiC= 70:30, in contact with the disk, at a slider load of 15 g. The measurements were performed in oxygen, nitrogen, argon and helium. Similar to that reported by Marchon et al. [30], it was found that the friction increased continuously in dry oxygen, while, in the other gases (nitrogen, argon and helium), the friction increased slightly only during the initial three revolutions and remained constant afterwards. At RH = 4%, little differences were observed in the friction in the different gases. The behavior was explained by the tribochemical mechanism described before. Adhesive friction was assumed to be the dominant friction mechanism and, since the frictional force is proportional to the real contact area, the smoothing of the surface by the tribochemical wear resulted in higher friction. During rubbing in the presence of water, the carbon appears to be removed in the same way as occurs during rubbing in oxygen, with the oxidation of the carbon being preceded by the dissociation of water molecules [32].

At low humidity (RH=O%-OS%), the increase in friction with sliding cycles was smaller than that at higher humidities. In contrast to the case with conditions of higher humidity, at the lower humidity debris was observed to form at the sliding interface as a result of mechanical wear. The debris maintained a small contact area, resulting in corresponding lower friction [32]. After a larger number of cycles, an equilibrium was reached between mechanical wear (which caused debris fo~ation) and chemical wear (which removed

the debris and smoothed the surface), and the friction coefficient reached a constant value [32].

Marchon and Khan [34] used scanning tunnelling microscopy (STM) to obtain topographic and spectro- scopic images of 30 nm thick magnetron-sputtered carbon films 30 nm thick. The spectroscopic images, showing the derivative dZ/ds of the tunnelling current I with respect to the tip-to-surface distance s, provided info~ation on the chemical nature of the film in terms of the local value of the work function. Areas of large d&is fluctuation were observed and were attributed to local changes in the structure of the sputtered carbon from graphitic to diamond-like microdomains, according to the model described by Robertson [ 141. No correlation was observed between the topographic and spectroscopic images of the films. Atomic structures have been resolved locally, showing short-range graphitic order, especially in films of poor durability, while harder DLC films showed a totally disordered structure. Sputtered carbon films with better mechanical properties produced spec- troscopic images of greater homogenei~, indicating a more uniform film microstructure [34].

The effect of the surface topography and chemistry on the tribological behavior of sputtered carbon was addressed by Hilden et al. [31]. Carbon films were sputtered in argon on two types of disk: disks make of particulate media (y - Fe,O, or cobalt-doped y - Fe,O,) in organic binders and thin film disks pre- pared by sputtering a cobalt-based magnetic alloy plus a carbon overcoat on polished NiP substrates. Ruth- erford back scattering (RBS) analysis showed that the carbon films contained small amounts of oxygen and argon and 2%-3% hydrogen, although the source of the hydrogen was unclear. The films were characterized by continuous sliding tests and start-stop tests against sliders of A120,--TiC at a load of 35 sf.

It was found that the behavior of the carbon-over- coated particulate disks did not change significantly with humidity. The friction coefficient changed slightly, remaining between 0.2 and 0.4, with a maximum ob- tained at 45% humidity. However, the dynamic friction coeflicient of the carbon sputtered on the thin film disks showed a high sensitivi~ to the presence of water and decreased from p= 1.3 at RH -5% to ~=0.9 at RH =40%, then decreased drastically to ~=0.3 at RN = 60%. The decrease in the friction coefficient with increasing humidity may indicate that these films were more graphite like than diamond like. The authors explained the differences in the behavior of the two types of disk by the differences in their surface to- pography. During sliding of the head against the carbon surface, the roughness owing to processing corrugations can be ironed out, increasing greatly the contact area in the case of thin film disks. The surface roughness of the particulate disks with a wealth of pores and

148 A. Grill I Tribology of diamond-like carho!!

crevices was not significantly affected by the wear tests and the contact area remained constant.

The tribological properties of a different type of DLC film, deposited by condensing carbon from the plasma flux generated by vacuum arc discharges, have also been investigated [37-40]. Films 3-4 pm thick with a Vickers hardness in the range 40-180 GPa were de- posited by Aksenov and Strel’nitskij [37] on both surfaces of a friction couple (hardened steel-45 disk and spherical ball-bearing steel, Cr-15 rider). In air at a sliding speed of 0.8 m s-‘, a friction coefficient in the range 0.04-0.13 was found to be independent of the load up to a critical load of 70 N. In a vacuum of lo-’ Pa and below a critical load of 18-20 N, the friction coefficient was 0.01-0.14. Above the critical load, the friction coefficient in mcuo increased gradually to 0.6 at 60 N.

DLC film, probably had the donrrnant effect on tht’ tribological behavior of the friction couple, determining the value of the friction coefficient and producing iin extremely low wear coefficient for- the DLC film.

To improve the mechanical and triboiogical properties of DLC, different derivatives of hydrogenated DLO films have been investigated. Such derivatives arc silicon-. metal- or nitrogen-containing DLC. as well as fluorinated variations of some of these materials [&-IO, 41-461. The modifications of the DLC films were gen- erally made (even if not always stated explicitly) to improve their qualities according to requirements (l)-(3) described in the Introduction.

During friction testing in air, a fine powdered debris identified as turbostratic graphite was formed on the friction couple and the authors assumed that, in this situation, the frictional force was determined by the sliding of graphite against graphite. The wear debris formed in vucuo was claimed to contain polycrystalline diamond and carbine in addition to graphite (no in- dication was given how this composition was deter- mined). The graphite particles were presumably carried away from the contact zone during the friction testing and the surfaces interacted directly with the hard diamond and carbine particles. The normal load was thus applied mainly to the hard components of the DLC and the true contact area was smaller than that in air. This could explain why the friction in mcuo was less than that in air. The authors tested different combinations of DLC coating with various hardnesses (40-180 GPa) and found the lowest friction coefficient (0.04 in air and 0.01 in vacm) for the combination of DLC coatings with highest hardness [37].

Tochitsky et al. [38] deposited films 0.5-3 pm thick by accelerated flows of plasma from pulsed vacuum arc discharges at substrate temperatures below 373 K and obtained carbon films with fine-grained structures. The films were tested for up to 100 h by reciprocal sliding at RH = 50%-60% at a rate of 3000-5000 min- ’ without lubrication. The films had a hardness of 70-100 GPa and the friction coefficients against steel were 0.1 (static) and 0.07 (dynamic).

Miyake and co-workers [9,10] performed microtri- bological studies of silicon-containing carbon films and fluorinated silicon-containing carbon films. Point contact microscopy (PCM) (a tool developed from STM) was used to investigate the microscopic wear and frictional properties of different carbon films. ‘The measurements were performed at room temperature in ambient air at RH -450/o-50%. Silicon was incorporated into the film to improve the adhesion of the DLC to the silicon substrates, and fluorination was carried out to reduce the surface energy of the DLC film. Electron cyclotron resonance (ECR) was employed to deposit films 1 ,um thick from silane and ethylene, and the films were subsequently fluorinated by exposure to CF, plasma. The fluorination resulted in the incorporation of up to 28% surface fluorine and reduced the friction coef- ficient of the silicon-containing carbon film to less than 0.3. The hardness of the films increased with increasing silicon content, reaching a maximum at 40% Si, and was not changed by fluorination. The surface energy, evaluated by the contact angle with water, decreased as a result of fluorination and the adhesion of the carbon film to the silicon substrate increased with the addition of small quantities of silicon. The life of the coating was measured in the microtribological tester as the time to reach a friction coefficient of 0.3. The lifetime thus defined was maximal for films containing lo%-20% Si. Films containing more than 20% Si had a higher initial friction coefficient. The silicon-containing carbon films had less wear than did the pure DLC films, while the fluorinated films showed even less weal

WI. Hirvonen et al. [40] have deposited DLC films on The tribological properties of silicon-containing car-

steel using a vacuum arc discharge [40]. In dry sliding bon films (Si-DLC) were also reported by Oguri and wear against a hardened steel pin in air at RH = 45%, Arai [41]. The films were deposited by a d.c. plasma a friction coefficient of 0.18 was measured in a pin- of CH4, Sic&, H, and argon and were found to contain on-disk tribotester at a sliding speed of 10.8 cm s- ’ about 40% H. The wear and friction coefficients were and a hertzian stress of 860 MPa. After 100 000 cycles, measured by a pin-on-disk apparatus, using a ball- the friction coefficient decreased to 0.12. Scanning bearing steel ball (AISI 52100) as a rider. According electron microscopy (SEM) analysis showed that a to the authors, the films with compositions of 70% CC/ transfer layer was formed on the pin during sliding. (C + Si) < 85% had a DLC character. These films were This layer, which appeared to be softer than the initial found to have the lowest friction coefficients of about

A. Gtill / T~boio# of diamond-i~ke carbon 149

0.04 against steel in an ambient atmosphere at RH= 50%-70%. The low friction coefficient was ex- plained by the formation of SiO, debris and its inter- action with ambient gases and the DLC. The authors also suggested that the low friction coefficient could be explained by temporary melting of the SiO, at the interface between the two parts of the friction couple

1411. A material similar to amo~hous carbon was studied

by Hioki and co-workers [44,45]. The films of i-siiicone, prepared by ion-beam-assisted deposition (IBAD), were deposited by combining evaporation of a silicone oil (pentaphenyl-trimethyl-trisiloxilane; (Si,O,C,,H,,),) with simultaneous irradiation of the surface of the growing film with energetic ions (1.5 MeV Ar+, 400 KeV N’ or 200 KeV Ti+). RBS analysis showed that the composition of the IBAD carbonaceous film, with the exception of hydrogen, is C:Si:O = 0.8:0.1:0.1, which is similar to that of the original oil, i.e. C:Si:O = 0.87:0.08:0.05. The evaluated density of the film was 2.6 g crrP3, compared with 1.1 g cmP3 for the original oil. Secondary ion mass spectroscopy (SIMS) measurements indicated that the films also contained about 10% H [45]. The authors explained the findings, assuming that the irradiation caused the emission of the hydrogen from the oil, resulting in its solidification.

Some i-silicone films 0.1-0.3 pm thick were tested on a pin-on-disk apparatus against hardened SUJ2 steel pins. The films had low friction coefficients (~=0.05) in air or nitrogen at RH= 20%-70%. In dry nitrogen, i-silicone showed a very low friction coefficient of ~~0.02. For the same steel disk coated with an unir- radiated oil fihn, the friction coefficient was 0.1. The authors found that i-silicone material was transferred to the steel rider during testing and they related the low friction coefficients to that transfer [44].

Metal-containing hydrogenated carbon films (Me- DLC) were studied by Dimigen and co-workers [42,43]. Metal was incorporated into the films to reduce stresses and improve the adhesion of DLC to the metals. The films, deposited by sputtering a metal or metal carbide target in an Ar + hydrocarbon plasma, had an H:C ratio of 0.2-0.5 and a metal mole fraction xMe=cMe/ (chne -i-c=) > 0.02. The films were identified as having a granular structure consisting of nanocrystalline metal or carbide particles in a DLC matrix. The friction coefficients and adhesive wear were measured with a pin-on-disk apparatus, using uncoated steel or cemented carbide spheres.

The friction coefficient of the Me-DLC films has been found to be strongly dependent on the metal incorporated into the films. In an earlier paper 1421, Dimigen et al. reported that, in humid air or humid nitrogen, the coefficient of friction was found to be ~~0.2. At a very low relative humidity, i.e. RH<O.l%,

the value of the friction coefficient was found to scatter and to depend on the metal, its concentration and sliding speed. Also, the friction coefficient increased with increasing metal concentration. A very low friction coefficient of 0.06 was found for an Ru-DLC film with x Ru= 0.07 at RH=O.l% and a sliding speed of 1.6 cm S --I. However, at higher sliding speeds, the friction coefhcient was much higher (~=0.7). In WHV the friction coefficient became high even at very low speeds. ~though the friction ~e~cient of pure DLC (without metal) was 0.02 in UHV, a small amount (5%) of metal led to a considerable increase in the friction coefficient to 0.15 in UHV. In humid air or humid nitrogen at RH > l%, the friction coefficient of the metal-containing films was generally below 0.2 [42]. More recent results obtained in ambient air at RH= 18%-41% showed a minimal friction coefficient of 0.04 at 12 at.% metal in Ta-DLC films 1431. For W-DLC films containing less than 13 at.% W, the friction coefficient against steel decreased from 1.5 at RH=40%-80% to about 0.02 in dry nitrogen (RH<O.l%). For a tungsten concen- tration higher than 30 at.%, the humidity had an opposite effect on the friction coefficient 1431.

Dimigen and Klages [43] explained the tribological behavior of the Me-DLC films based on their mechanical properties and surface energy. According to them, tribological materials are frequently divided into three groups: metals, ceramics and polymers. The Me-DLC films attain a hardness H typical of inorganic ceramics but have Young’s modulus values E significantly lower than those of ceramics. The corresponding H/E ratios are relatively high and their values are typical for polymers. The surface energy S of the Me-DLC fihns with low metal contents are also more similar to that for polymers than to those for metals or ceramics. The combination of the high H/E with low S/H ratios (S/H values being smallest among the other type of ma- terials) could explain the low adhesion to the Me-DLC films, their low friction coefficient and the low wear

1431. The above explanation of the tribological behavior

of metal-containing DLC films may to a large extent be applied to pure hydrogenated DLC or to silicon- containing and/or fluorinated DLC films. The excellent tribological performance of DLC films and their de- rivatives might be attributed to a combination of fa- vorable properties which are ceramic-like, on the one hand (high hardness), and polymer-like, on the other hand (high elasticity, H/E and low surface energy).

2.2. Wear As shown above, a significant number of studies have

investigated the friction of DLC, starting as early as 1981, and have suggested mechanisms to explain the friction coefhcients of DLC films and related materials.

However, the wear behavior of DLC has been only recently addressed in a systematic way. The following studies will be discussed according to the type of DLC: (corresponding to the deposition method) and not in chronological order.

Namura et al. f7] investigated the failure mechanisms of hard magnetic disks coated with sputtered carbon by contact start-stop (CSS) tests against Al,O,-TiC sliders. The carbon films, 20-30 nm thick were sputtered on top of a sputtered magnetic layer of Co-Cr and permalloy. The tests were performed at rise times in the range of 2-24 s and relatively low humidities of 15%-20% to prevent slider stiction. The slider was pressed on the disk with a load of 5 mN before starting the tests. The acceleration time of the disk was identical to the deceleration time. The wear of the carbon layer was found to be minimal during CSS testing under the conditions investigated. Disk failure was found to occur abruptly. Fatigue cracks were observed to develop in the disk with repeated CSS cycIes. The cracks extended in-plane and propagated to the surface of the disks. When the crack reached the surface, flaking occurred abruptly. The authors concluded that the cyclic load is dominant over the forces acting at the start, because of higher static friction coefficient and dynamic load factor.

Agarwal et al. [29] described the wear of the sputtered carbon films by a wear coefficient K, defined as the ratio of the vertical cross-section of the wear track produced in the pin-on-disk tester to the horizontal contact area of the slider. The values of K obtained were 5 X 10M6 for CoNi and 9 x 10e6 for COP with hard carbon over-coating, U.Y. 50x lo-’ for uncoated COP, indicating an increase in wear resistance by a factor of 6 as a result of coating the magnetic metallic film with the carbon layer, despite the high friction coef- ficient. The wear was found to increase with normal load and the number of revolutions but was independent of the linear speed. According to the authors, the principal contribution to the friction coefficient of DLC is inelastic in nature; the mechanical energy is dissipated primarily in the form of heat through inelastic processes rather than the generation of lattice defects or other mechanisms of plastic defo~ation~ Nevertheless the small fraction of the total energy consumed to form lattice defects causes plastic deformation and this, coupled with fragmentation of surface material, results in macroscopic wear.

Wear tests performed by Strom el al. [32] on thin film magnetic disks with sputtered carbon overcoats, using commercial read/write ceramic heads as sliders, indicated that the carbon overcoats wear through a tribochemical mechanism in the presence of oxygen and through a mechanical mechanism in the absence of oxygen. In dry gas without oxygen, debris is produced

almost mnnedrately through mecnatncal wear irr,o trr; debris precludes the development of high friction itn<i reduces the wear. In the prcserloc ,)1‘ cirq oxygen OI small amounts of water, tribochcmical wear takes place. without the formation of debris, and high friction cir:.- velops between the slider and dii;Ei

In systems which are sensitrvc itt the f~~rrn~~t~~)n of particles. eg. magnetic storage hytcms, tribochemical wear without debris formation may be preferred tc) mechanical wearwhich causes the generation ofparticles [32]. A dry inert environment, which allows rapid me- chanical wear, is undesirable in ~;uch devices using sputtered carbon films.

Hilden d (zl. [31] found that the RH value had a significant effect on the lifetime of carbon-overcoated thin film disks. Values of the lifetime to failure from 2 min at RH = 2% to greater than 1300 min at RH = 70% were observed, corresponding to the decrease of the friction coefficient with increasing humidity. More than 20 000 start-stop cycles have been achieved for carbon- coated disks, though the authors did not indicate for which type of disks or condition,

Miyoshi 1201 found that DLC deposited at a low power density failed in air and was removed from the sliding zone after 1000 passes, while films deposited at higher power density did not wear off the substrate, even after 10 OO(I passes. Even with the superficial graphite-like layer assumed to be formed on the surface of annealed DLC. the wear life of the high density fihn was reduced in humid air %3y a factor of two compared with that in dry nitrogen. The differences in the behaviors of the films deposited at different power levels were attributed to thu differences in the densities of the films.

G-ii1 and co-workers measured lhe wear of the DLC films deposited by PACVD from acetylene in a pin- on-disk tribometer against a ball-bearing steel rider at a hertzian stress of 21.5 MPa in nitrogen at RH=5% and in air at RH=40%-70% [47.483. The wear tests were performed by running the apparatus for a preset number of rotations (260 000) and measuring the re- sulting wear tracks with a profilometer. The wear of the DLC films, expressed as the track depth (A) per lo4 rotations in the tribotester, is presented in Fig. I as a function of the annealing temperature. The values over 3000 indicate that the respective DLC films were completely worn through after only a few thousand rotations.

The as-deposited films showed very little wear, in- dependent of the deposition temperature, with the films deposited at a bias of -150 V (de.) having only a negligible amount of wear. After annealing at 390 “C, the films deposited at -80 V (d.c.) at 100 and 180 “C had very low wear resistances and were completely worn through after a few thousand rotations both in

A. Grill 1 Tribology of diamond-lip carbon 151

L As Dee. 390 440 490 590

02 3.000

‘;’ 1.000

E 300

100

30

10

In Air

1 (a)

As Dep. 390 440 490 590

Temperature ( “C )

DeposTemp

BeIs

Fig. 1. Wear of DLC films vs. annealing temperature: (a) in nitrogen; (h) in air.

nitrogen and humid air 1491. The film deposited at -80 V (dc.) at 250 “C and annealed to 440 “C did wear at a higher rate than the as-deposited film and was worn through in nitrogen after annealing at 490 “C, while in air it lost its wear resistance after annealing at 400 “C. DLC films deposited at 180 or 250 “C at a bias of - 150 V showed negligible wear in nitrogen after annealing at 390 “C, wore more rapidly after annealing at higher temperatures but remained wear resistant in nitrogen, even after annealing at 590 “C (see Fig. 1). In humid air, the wear resistance decreased and the film deposited at 180 “C at - 150 V lost its wear resistance after annealing at 490 “C. However, the film deposited at 250 “C at - 150 V remained wear resistant in humid air, even after annealing at 590 “C

[491. These results also showed that the wear of the DLC

films was not directly related to the values of their friction coefficients. In addition, while other physical properties showed only small differences in behavior between the films deposited at various temperatures and substrate biases [48], the wear resistance of the films and their thermal stability were strongly affected by the deposition conditions. A higher wear resistance after annealing was obtained for DLC deposited at higher temperatures and films deposited at the higher biases remained wear resistant, even after annealing at 590 “C.

Em et al. 1271 measured the wear rates of the r.f. PACVD DLC films against a silicon nitride ball in a pin-on-disk tribometer. The wear tests were performed in argon and air at different humidities. The wear rates of the silicon nitride ball and the DLC at RH =50% were lop8 and 10e7 mm3 N-’ m-’ respectively. In dry argon, dry air and 100% humid air, the wear rates of DLC were two orders of magnitude lower and that of the ball was undetectable. An interfacial oxidized layer formed from mechanically scissored particles, which combined to produce macrowear soft debris, was ob- served between the sliding parts. Kim el al. claimed that, in 100% humid air, water molecules produce strong adhesion of the wear debris and that a dense layer covered the track on the DLC and reduced the oxidation rate in the wear track. This oxidized debris layer was thick enough to separate the DLC from the silicon nitride and so reduce the wear. In 50% humid air, the debris was found to be a mixture of oxidized hydrocarbon and hydrated SiO, from the ball. However, the debris appeared to be loose and not covering the wear track. According to the authors, the wear of DLC against silicon nitride is determined by the transfer layer of wear debris. The wear is determined by an adhesive (mechanical) wear mechanism in dry argon but by tribochemistry in air and humid environments [27]. The authors found no correlation between wear and the friction coefficients in the various testing con- ditions.

Microt~bological studies performed by Miyake et al. [9] for Si-DLC films and fluorinated Si-DLC films showed an increase in the lifetime of the films with increasing silicon content, with a maximal lifetime being obtained for films containing lo%-20% Si. Under the conditions of microtribological testing, wear tracks 3 nm in size were obtained at a load of 7 FN on un~uorinated film, while wear of the ~uorinated film was undetectable. The wear increased with load and was lower for the fluorinated films at all loads. The wear of Si-DLC films was also reported on by Oguri and Arai [41]. The wear of a steel rider against Si- DLC was found to be the same as that against DLC coatings; however, the wear of the Si-DLC coating was found to be higher than that of DLC.

In the study of the wear of Me-DLC films, for W- DLC films with 12 at.% W, the authors observed a two-order reduction in the sliding wear of the coated disk against steel compared with that for the bare substrate and a drastic reduction in the wear of the sliding counterparts 1431. When the test was carried out against TiN or T&N films, significant wear of the steel balls was observed, with the transfer of steel to the nitride films. For the Me-DLC films, such a transfer of steel onto the W-DLC was not observed. Dimigen and Klages [43] also measured the abrasive wear rates

152 A. Grill I Triboloa of diamond-l& c‘crrix )I!

of the Me-DLC using a glycerine suspension of Al,O, particles 5 pm in size. The abrasive wear rates showed a minimum at about 13 at.% W in W-DLC films (a composition at which the lowest friction coefficient was obtained). The minimal abrasive wear rates were two and eight times lower than the abrasive wear rates of TiAlN and TiN respectively. However, the abrasive wear rate of Me-DLC was two times larger than that of pure DLC films. The combination of a low energy surface (low S/H ratio) with a relatively strong, cross- linked microstructure, resulting in high hardness and high H/E values, could explain the low adhesion to the Me-DLC films and their low friction coefficient and low wear [43].

The DLC films deposited by vacuum arc discharge, as reported by Aksenov et al. [37], were found to have a high wear resistance both in air and vacuum. The best wear resistance was obtained for the combination of DLC coatings with highest hardness. DLC coatings improved the lifetime of drills made of tool steel by a factor of 1.5-3 during drilling of abrasive glass- reinforced polymers and increased by a factor of four the lifetime of carbide tools for turning titanium. The application of DLC on burnishers made of hardened tool steel made it possible to reduce the surface rough- ness of non-ferrous metals by two or three classes [37]. Tochitsky et al. [38] found that the service life of the steel parts coated with arc-discharge-deposited DLC increased by 3-40 times.

3. Summarizing remarks

The excellent tribological performance of DLC films and their derivatives might be attributed to a combi- nation of favorable properties which are ceramic-like, on the one hand, i.e. high hardness, and polymer-like, on the other hand, i.e. high elasticity and H/E and low surface energy. The results reviewed show that, although the tribological properties of DLC or derived materials deposited by different methods span a wide range, the materials appear to have several common characteristics:

the value of the friction coefficient of hydrogenated DLC deposited by PACVD is low in humid nitrogen or oxygen, extremely low in dry nitrogen or UHV, and very high in dry oxygen;

the loss of hydrogen through annealing at high tem- peratures causes a marked increase in the friction coefficient in UHV but affects only slightly its value in a humid atmosphere. In UHV or an inert atmosphere, these films display a frictional behavior similar to that of diamond. However, in a humid atmosphere, their behavior seems to be closer to that of graphite. The mixed type of frictional

behavior reflects the mixture ot sp and sp-’ carboti

hybridization in the hydrogenated carbon films. The behavior of DLC deposited by other method:,

is less systematic, though very h t‘riction coefftcicnts in UHV were found for vacuum-arc-deposited films and a behavior similar to that <tf ihe PACVD films was observed in dly oxygen anti nitrogen for- some sputtered films. The behavior of the friction coefhcients of the sputtered carbon films indicates that this type of film is more graphite like than diamond like.

The high friction values found in the tests run in dry oxygen and the low friction in dry nitrogen appear to be common features of amorphous carbon with diamond-like properties.

Lowering of the surface energy ot the DLC films, e.g. by fluorination, can reduce their friction coefficients. The friction of all DLC films, especially after a large number of sliding passes, appears to be controlled in many cases by the formation of an interfacial transfer layer, the properties of which depend on the cnviron- ment and on the sliding partner.

The wear resistance of DLC appears to be strongly dependent upon the deposition conditions of individual films. The data reported do not permit us to establish a correlation between the wear resistance and the friction coefficient. However, it seems that the interfacial transfer layer, which affects the friction, also determines the wear behavior of the DLC films. The testing en- vironment can determine whether the wear mechanism is mechanical (adhesive) or tribochemical and can de- termine the degree of wear.

In addition to its friction coefficient, a factor which most probably affects the wear behavior of DLC, is the adhesion of films to the substrates investigated. This factor, which can be controlled by the preparation of the substrate and by the deposition parameters, is not addressed directly in most cases.

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