24Tribology of Diamond,Diamond-Like Carbon,

and Related Films*

24.1 Introduction24.2 Diamond Films

Microcrystalline Diamond Films • Nanocrystalline Diamond Films • Tribology of Diamond Coatings • Tribological Applications

24.3 Diamond-like Carbon (DLC) FilmsBackground on DLC Films • Tribology of DLC Films • Applications of DLC Films

24.4 Other Related FilmsCubic Boron Nitride (CBN) • Carbon Nitride Films

24.5 Summary and Future Direction

24.1 Introduction

Diamond, diamond-like carbon (DLC), and other related materials (i.e., carbon nitride and cubic boronnitride [CBN]) are some of the hardest materials known and offer several other outstanding properties,such as high mechanical strength, chemical inertness, and very attractive friction and wear properties,that make them good prospects for a wide range of tribological applications, including rolling and slidingbearings, machining, mechanical seals, biomedical implants, microelectromechanical systems (MEMS),etc. The dry sliding friction and wear coefficients of these materials are among the lowest recorded todate (Brookes and Brookes, 1991; Feng and Field, 1991; Field, 1992; Miyoshi, 1995; Erdemir, 2001a,b).In fact, if they were inexpensive and readily available, they would undoubtedly be the materials of choicefor a wide range of applications. Besides their exceptional mechanical and tribological properties, mostof these superhard materials offer broad optical transparency, high refractive index, wide bandgap, lowor negative electron affinity, transparency to light from deep UV through visible to far infrared, excellentthermal conductivity, and extremely low thermal expansion. Briefly, these exceptional qualities makediamond, DLC, and other related materials ideal for numerous industrial applications in addition totribology.

Ali ErdemirArgonne National Laboratory

Christophe DonnetÉcole Centrale de Lyon

* The submitted manuscript has been created by the University of Chicago as Operator of Argonne NationalLaboratory (“Argonne”) under Contract No. W-31-109-ENG-38 with the U.S. Department of Energy. The U.S.Government retains for itself, and others acting on its behalf, a paid-up, nonexclusive, irrevocable worldwide licensein said article to reproduce, prepare derivative works, distribute copies to the public, and perform publicly anddisplay publicly, by or on behalf of the Government.

Work supported by U.S. Department of Energy, Office of Energy Research, under Contract W-31-109-Eng-38.

The purpose of this chapter is to provide an overview of the tribology of diamond and related coatingsthat have attracted overwhelming interest in recent years. Emphasis is placed on the current state-of-the-art of our understanding of the friction and wear mechanisms of these films as well as on their uses fordemanding tribological applications. Referring to the structural and fundamental tribological knowledgegained during the past decade, this chapter emphasizes the importance of surface physical and chemicaleffects on friction and wear of these materials and describes new deposition procedures that can lead to theproduction of novel films that afford ultralow friction and very long wear life to sliding tribological interfaces.Tribological issues associated with metal cutting and contact sliding are also addressed. By referring to recenttribological studies, one observes that the surface roughness of these materials plays a major role in theirfriction and wear performance. Smooth diamond films obtained by polishing and by controlling grain sizehold promise for future tribological applications, including mechanical seals and MEMS.

The chapter is divided into three main subject areas. The first area is devoted to the synthesis, tribology,and applications of diamond films. The second deals with DLC films and their tribology; and the thirdis devoted to the tribology and applications of other related films such as CBN and carbon nitride.

24.2 Diamond Films

Natural diamonds are expensive and difficult to machine into useful wear parts for large-scale industrialapplications. Except for a few cases (e.g., diamond-studded rotary drill bits, dressers, diamond-tippedglass cutters, and fine powders as superabrasives in grinding wheels), natural diamond is rarely used fortribological purposes. However, synthetic polycrystalline diamonds (PCDs) have been available for morethan 30 years and they are widely used in many industrial applications. PCDs are produced in largequantities by a high-pressure/high-temperature (HPHT) method in which diamond is crystallized froma graphitic or carbonaceous precursor at pressures of 50 to 100 kbar and temperatures of 1600 to 2000°C.The end product is a fine diamond powder that is used mostly as a superabrasive in grinding and polishingmedia. The phase diagram in Figure 24.1 indicates the stability ranges of diamond, graphite, and other

FIGURE 24.1 Carbon diagram in which the stability ranges of diamond, graphite, and other forms of carbon areindicated.

forms of C. It is possible to sinter these diamond powders with metallic or ceramic binders (e.g., Co, Ni,TiC, etc.) at high pressures and temperatures and then use them as tool bits or sharp blades in metal-cutting or machining operations.

24.2.1 Microcrystalline Diamond Films

The prospects for inexpensive and large-scale production of diamond increased tremendously in the1970s, when researchers discovered that diamond can be grown as a thin film at low deposition pressures(102 to 103 Pa) from a hydrocarbon/hydrogen mixture by chemical vapor deposition (CVD). Earlyattempts to produce diamond by CVD date to when HPHT synthesis of diamond was being explored.The tube oven experiments of Eversole (1962) and Angus (1994) and low-pressure synthesis of diamondfrom hydrocarbon gas mixtures by Deryagin and Fedoseev (1975) and Fedoseev (1994) can be regardedas the earliest attempts to produce synthetic diamond by CVD. In fact, Eversole of Union Carbide wasable to produce diamond phases by a thermal pyrolysis method in 1962 (Eversole, 1962). Progress slowedthereafter until Angus of Case Western University found a more efficient method in which graphite isetched simultaneously with atomic H in the deposition system (Angus, 1994). His work motivatedDeryagin and co-workers to focus on the uses of H/hydrocarbon mixtures for low-pressure diamondgrowth (Deryagin and Fedoseev, 1975). Since these studies, there has been an explosion in research ondiamond and related materials, with the expectation that low-pressure methods will allow faster, lessexpensive, and easier production of synthetic diamonds.

During the past 2 decades, great strides have been made in both the production and industrialapplications of diamond films. At present, several techniques can be used to produce high-qualitydiamond films with micro- and nanocrystalline structures at fairly high deposition rates over reasonablylarge surface areas (Banholzer, 1992; Gruen et al., 1994; Hatta and Hiraki, 1998; Komplin and Hauge,1998). The most widely used methods are plasma-enhanced CVD, hot-filament CVD, microwave CVD,DC-arc jet, combustion flame, and laser-assisted CVD (Haubner and Lup, 1992; Anthony, 1992; Butlerand Windischmann, 1998). These methods have greatly reduced the unit production cost of diamondand increased the prospects for large-scale industrial applications in tribological and other fields. In fact,several commercial companies now offer diamond-coated products at reasonable cost. The high-qualitydiamond coatings produced by CVD methods exhibit most of the desired mechanical and tribologicalproperties of natural diamonds (Field, 1992; Miyoshi, 1995; Lux et al., 1997).

Low-pressure synthesis of diamond coatings involves the use of hydrogen (H) and methane (CH4) gasmixtures. Typical deposition pressures in these syntheses are 102 to 103 Pa, and the gas mixture primarilyconsists of H and very little CH4. In the case of microwave CVD, 0.5 to 5% CH4 is mixed with H gas andintroduced into the deposition reactor. Small amounts of O and N can also be blended. In the case ofmicrowave CVD, a 2.45-GHz power source is used to initiate and maintain a gas discharge plasma inwhich the substrate material is immersed. At very high plasma temperatures, electrons intensely interactwith the H-rich gas mixture, and the ionic species that are created during this interaction lead to theformation of diamond nuclei on the surface of a suitable substrate (i.e., Si, SiC, W, Mo). Plasma conditionsin deposition reactors are tailored to control the film microstructure and chemistry, as well as growthrates and orientation. Typical substrate temperatures for high-quality diamond deposition can rangefrom 700 to 1000°C. Deposition of diamond at temperatures below 700°C has also been demonstratedbut growth rates are reduced significantly and the amount of non-diamond precursors in these coatingsis increased (Hatta and Hiraki, 1998).

Hydrogen is extremely important for the synthesis of high-quality diamond coatings in most CVDprocesses. It plays a critical role in the stabilization of the sp3 bond character of the growing diamondsurfaces. It also controls the size of the initial nuclei, dissolution of C and generation of condensable Cradicals in the gas phase, abstraction of H from hydrocarbons attached to the surfaces, and productionof vacant surface sites where sp3-bonded C precursors can be inserted. It etches most of the double orsp2-bonded C from the surface, and thus hinders the formation of graphitic and/or amorphous C(Banholzer, 1992). Hydrogen also continuously etches out the smaller diamond grains and suppresses

continuous renucleation. Only those grains that attain a certain critical ratio of bulk-to-surface atomssurvive. Consequently, these diamond coatings consist primarily of large grains with highly faceted surfacefinishes that exhibit a surface roughness value of ≈10% of the film thickness.

During the initial stage of film growth, nucleation starts at preferred sites; eventually, independentnuclei form. As these nuclei grow larger, they close the gaps between them and merge to form a continuousfilm. Thereafter, the growth process and growth rates are dominated by nearby neighbors and growthorientation. Grains with more favorable growth orientation grow most quickly and overshadow the grainswith less favorable growth orientation. Grains with less favorable growth orientation will be buriedbetween the large grains. Figure 24.2 shows initial, intermediate, and final growth stages of a diamondfilm on an Si substrate.

High-quality microcrystalline diamond coatings are rough and often nonuniform in thickness orslightly bowed because of internal stresses and differences in deposition rate from the edges to the centerof the samples. The generally rough surface finish of these coatings precludes their immediate use formost tribological applications. As discussed in detail later, when used in sliding-wear applications, suchrough coatings cause high friction and very high wear losses on mating surfaces. The rough diamondcoatings can be polished by laser beams, mechanical lapping with fine diamond powders, ion-beam orplasma etching, and thermomechanical polishing with hot Fe or Ni plates (Zaitsev et al., 1998; Rameshamand Rose, 1998; Erdemir et al., 1997a; Bhushan et al., 1994; Gupta et al., 1994; Pimenov et al., 1996).Mechanical lapping and polishing with hot Fe plates, the most widely used methods, allow polishing oflarge areas. Depending on the polishing method, one can achieve mirror finish surfaces with an rmssurface roughness of 10 nm or less. As demonstrated by numerous investigators, the polished diamondcoatings can provide friction coefficients comparable to that of natural diamond (Erdemir et al., 1997a;Bhushan et al., 1993; Gupta et al., 1994; Pimenov et al., 1996). However, the polishing processes aretedious, time-consuming, and, in the case of complex geometries, highly impractical. Depending on thedesired film thickness or the degree of original roughness, it may be necessary to remove large amountsof material before a smooth surface is obtained. Figure 24.3 shows the surface morphology of a roughmicrocrystalline diamond film before and after laser polishing.

24.2.2 Nanocrystalline Diamond Films

The very rough nature of the above-described microcrystalline diamond films renders them useless formost tribological applications. Recently, new methods have been developed for the deposition of fine-grained or nanocrystalline diamond (NCD) films with a very smooth surface finish. In these methods,a higher than normal C:H ratio is used in a microwave plasma, or a DC bias is applied to the substrates(Hollman et al., 1998; Hogmark et al., 1996; Cappelli et al., 1998; Bhusari et al., 1998). The methanefraction in the source gas, substrate temperature, and substrate pretreatment were also shown to exert astrong effect on the crystallinity and grain size of these films (Gilbert et al., 1998; Erz et al., 1993; Hogmarket al., 1996). Although growth rates are somewhat reduced and the amount of non-diamond phase issomewhat increased, the surface finish of the resultant coatings is very smooth (i.e., 25 nm, rms).

Deposition of NCD coatings has also been achieved by microwave CVD in the near absence of H. Thesource gas consists of mostly Ar, with either C60 or CH4 as the C precursor (Gruen et al., 1994; Gruen,1998; Zuiker et al., 1995; Erdemir et al., 1996b,c). The CVD reactor used in this process is essentially amodified version of a conventional microwave CVD reactor, details of which are schematically depictedin Figure 24.4. To introduce C60 into the reactor, a quartz transpirator, also shown in Figure 24.4 isattached. Fullerene-rich soots that contain ≈10% C60 are placed in the transpirator. The soot is heatedto 200°C under vacuum for 2 h to remove residual gases and hydrocarbons. The tube furnace andtransport tube are heated to between 550 and 600°C to sublime C60 into the gas phase. Argon gas ispassed through the transpirator to carry the C60 vapor into the plasma. To ensure that C60 is transportedinto the chamber, an Si wafer is placed in front of the transport tube while maintaining a 14-sccm Ar flow.With the reactor pressure at 100 torr and the transpirator at 600°C, a 1.7-mg brown film was depositedon the wafer in 1 h. The film displayed only strong C60 infrared absorption features. The measured Raman

spectrum was also attributable to C60. Based on these measurements, the transpirator is considered aneffective source of C60 for diamond growth. The NCD coatings can also be grown under similar low-H-content conditions, with CH4 instead of C60 as the C source.

The NCD coatings can be grown on various substrates, including single-crystal Si wafers, sintered SiC,W, WC, Si3N4, etc. Initially, a bias of –150 V is applied to enhance diamond nucleation density. Filmgrowth is monitored in situ by laser reflectance interferometry to determine growth rate and to stopgrowth at the desired film thickness. The substrate temperature can vary between 700 and 950°C, and

FIGURE 24.2 (a) Initial, (b) intermediate, and (c) final growth stages of a microcrystalline diamond film producedin a microwave CVD reactor.

the total gas flow rate is ≈100 sccm. A typical gas composition for NCD diamond film can be 97% Ar,2% H2, and 1% C60 or CH4 at a total pressure of 1.33 × 104 Pa and a microwave power of 800 W. Figure 24.5shows atomic force microscopy (AFM) and scanning electron microscopy (SEM) images of the surfaceof an NCD film produced by the method depicted in Figure 24.4.

The growth mechanism of NCD differs radically from that of a microcrystalline diamond film.Specifically, extraction of a C dimer (C2) from C60 and CH4 molecules and its subsequent insertion intothe diamond surface has been proposed for the growth of these coatings (Gruen, 1998). The resultantcoatings are phase-pure NCD, with an average grain size of ≈15 nm. The C dimer growth mechanism isunique in that it is capable of producing a continuous diamond coating that can be as thin as 30 to 60 nm.

24.2.3 Tribology of Diamond Coatings

In addition to being the hardest material known to mankind, diamond provides some of the lowestfriction coefficients to sliding tribological interfaces when tested in open air. This combination of extremehardness (and hence wear resistance) and ultralow friction in one material is very rare in the field oftribology and it renders diamond ideal for a wide range of tribological applications. However, recentfundamental studies have confirmed that there are several factors that can adversely affect the frictionand wear performance of bulk diamond and thin diamond films (Bowden and Tabor, 1950; Bowden andYoung, 1951; 1964; Bowden and Hanwell, 1966; Tabor, 1979; Enomoto and Tabor, 1981; Feng and Field,1992; Hayward and Field, 1987; Gardos 1994; Chandrasekar and Bhushan, 1992; Miyoshi et al., 1993;Kohzaki and Noda, 1994). Most previous studies have confirmed the finding that the friction coefficient

FIGURE 24.3 Surface morphology of a rough diamond film (a) before and (b) after laser polishing.

of bulk diamond sliding against itself in open air is 0.02 to 0.05. However, when tested in ultrahighvacuum (UHV) or at high temperatures, the friction coefficient of diamond increases by more than1 order of magnitude (Hayward and Field, 1987; Chandrasekar and Bhushan, 1992; Miyoshi et al., 1993;Dugger et al., 1992). Some of the other factors that influence the friction and wear behavior of diamondare contact pressure, surface roughness, crystallographic orientation, and the presence or formation ofgaseous, liquid, or solid third-body and/or transfer film on the sliding surface.

In addition to the early studies that were performed at micro-to-meso scales, recent fundamentaltribological studies at atomic-to-nano scales have greatly increased our understanding of the influenceof each of the above-mentioned factors on friction and wear of diamond. In particular, with the adventof new experimental tools (e.g., AFM) and computer simulation methods (e.g., molecular dynamicsimulation), it is now possible to visualize and better understand the extent of chemical, physical, andmechanical events that occur at atomic or molecular levels and hence gain a better understanding of themechanisms that govern the friction and wear behavior of diamonds.

24.2.3.1 Early StudiesEarly tribological studies have led to the conclusion that the low-friction property of diamond is largelydue to the highly inert or very passive nature of its sliding contact surfaces (Bowden and Young, 1951;Bowden and Hanwell, 1966). Specifically, it has been proposed that the low friction of diamond wasassociated with the lack of adhesive forces between sliding diamond surfaces. Gaseous adsorbates in thesurrounding atmosphere, such as H, O, or water molecules, can attach to and effectively passivate thedangling σ-bonds of diamond surfaces, and hence lead to reduced adhesion and/or friction. This hypoth-esis for the low-friction mechanism of diamond has existed for a long time. More recent work by Gardos(1999), Gardos and Ravi (1994), Miyoshi et al. (1993), Dugger et al. (1992), and Chandrasekar andBhushan (1992) has confirmed that the presence or absence of surface contaminants in the test chamber

FIGURE 24.4 Schematic illustration of a modified microwave CVD reactor used for deposition of nanocrystallinediamond films. (From Zuiker, C., Krauss, A.R., Gruen, D.M., Pan, X., Li, J.C., Csencsits, R., Erdemir, A., Bindal, C.,and Fenske, G. (1995), Physical and tribological properties of diamond films grown in argon-carbon plasmas, ThinSolid Films, 270, 154-159. With permission.)

makes a huge difference in the friction and wear performance of bulk diamond and/or thin diamondfilms. When tested in ultraclean and ultradry test environments (i.e., UHV) or at high ambient temper-atures, sliding diamond surfaces always exhibit high friction and wear, mainly because surface contam-inants are desorbed or mechanically removed from the sliding surfaces and hence are not available topassify the dangling σ-bonds of the sliding diamond surfaces.

24.2.3.2 Friction and Wear Mechanisms

Carbon atoms in bulk diamond crystals are held together by strong covalent bonds in a tetrahedral bondconfiguration. However, C atoms on diamond surfaces can only establish three covalent bonds with theirnear neighbors, while the fourth bond is left open and dangling out of the surface. Gaseous species inthe surrounding air such as water molecules, O, and H, can chemisorb and effectively passivate thesedangling surface bonds (Figure 24.6) (Holmberg, 1998; Pate et al., 1982; Derry et al., 1983; Wei et al.,1995). When dangling bonds are tied up and passivated, the extent of adhesion and hence friction betweensliding diamond surfaces is drastically reduced. The extent of interaction between passive diamondsurfaces that are sliding against each other drops to the level of very weak van der Waals attractions.Knowing the critical role that dangling bonds play on friction, researchers have developed more effectivemeans to passivate them with, for example, F atoms, and have thus achieved extremely low friction andwear coefficients for bulk diamond and diamond films (Miyake et al., 1994; Smentkowski et al., 1996,1997; Molian et al., 1993).

Recent fundamental studies have shown very clearly that, when adsorbed gases are removed from thesliding surfaces of diamond by thermal desorption at high temperatures, the friction coefficient increases

FIGURE 24.5 (a) AFM and (b) SEM images of the surface morphology of nanocrystalline diamond film producedby microwave CVD in Ar-C60 plasma.

rapidly; presumably, the dangling surface bonds are reactivated and are available to form strong adhesivebonds with the surface atoms of the counterface materials (Gardos, 1994; Dugger et al., 1992). Conversely,if the surface of diamond is exposed to gaseous contaminants or open air again, the friction coefficientdrops precipitously, presumably because of the repassivation of the dangling surface bonds, as shown inFigure 24.7 (Dugger et al., 1992).

Most of the vacuum tribological studies on diamond were performed by Gardos and co-workers. Usingan SEM tribometry test device, they have explored the effects of surface chemistry-induced changes onfriction and wear of self-mated sliding diamond surfaces as a function of ambient temperature. Assummarized in a recent review article by Gardos (1999), their studies provide circumstantial evidencefor the effect of surface dangling bonds on friction (in the absence of shear-induced phase transformationand/or graphitization). Specifically, their studies show that the lowest friction is achieved under conditionswhere complete passivation of dangling bonds of diamond surfaces is feasible (i.e., by atomic or molecularadsorbates such as H2O, H, or O). However, reconstructed diamond surfaces provide relatively reducedfriction and, finally, the incipient linking of the sliding diamond surfaces by totally unsaturated danglingbonds (created by thermal desorption) causes the highest friction. The repassivation of dangling bondsby atomic or molecular adsorbates during cooling or exposure to gaseous species causes reduced adhesionand hence low friction. Figure 24.8 illustrates the frictional behavior of diamond films in the passive,

FIGURE 24.6 Schematic illustration of dangling surface bonds of diamond and weak shear plane between hydrogen-terminated diamond surfaces. (From Holmberg, K. (1998), Tribology of diamond and diamond coatings — A review,Tribologia, 12, 33-62. With permission.)

FIGURE 24.7 Friction coefficient vs. number of sliding passes for a CVD-coated SiC pin sliding against a CVDdiamond-coated Si wafer in air and UHV environments. (From Dugger, D., Peebles, E., and Pope, L.E. (1992),Counterface material and ambient atmosphere: role in the tribological performance of diamond films, in SurfaceScience Investigations in Tribology, Experimental Approaches, Chung, Y.-W., Homolo, A.M., and Street, G.B. (Eds.),ACS Symposium Series 485, American Chemical Society, Washington, D.C., 72-102. With permission.)

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active, and reconstructed regimes of sliding. Apparently, as the temperature increases, surface adsorbatesbegin to desorb and leave some of the dangling surface bonds exposed and available for strong adhesiveinteractions. As the extent of adhesive interactions increases, the friction coefficient increases sharply andreaches values ≈0.7 until some of them reconstruct and reduce the extent of adhesion and hence friction.Among all of the surface adsorbates, H provided the lowest friction (Gardos and Gabelish, 1999).

Because of its extreme hardness, bulk diamond and thin diamond film hardly wear under normalsliding conditions (i.e., open air, room temperature, low velocity). However, under certain test conditions,diamond may experience some wear by shear or ambient-induced phase transformation. For example,a few studies have indicated that micrographitization of sliding diamond surfaces can occur at hightemperatures or at high sliding velocities, and the surfaces of diamond films begin to act like graphite;that is, they give low friction in moist air, but relatively high friction in dry, inert, or vacuum environments(Gardos et al., 1997; Gardos and Ravi, 1994; Erdemir et al., 1997b).

Natural diamond is chemically very stable and does not appreciably oxidize until the temperaturereaches 600°C. The oxidation behavior of diamond films is also similar to that of natural diamond but,because they contain some non-diamond phases, high levels of structural defects, and grain boundaries,they tend to oxidize at relatively lower temperatures (Tankala, 1990; Johnson et al., 1990). When tri-botested at elevated temperatures, bulk diamond and thin diamond films tend to gradually oxidize andtransform to graphite in open air (Erdemir et al., 1996d; Erdemir et al., 1997c; Kohzaki et al., 1992).However, in inert test environments or under high vacuum, diamond films are more stable and do notundergo oxidation or graphitization. Briefly, diamond films are not suitable for use at high temperaturesin open air.

In addition to the effects of surface chemistry and/or environment, which largely control the adhesioncomponent of friction, the effects of crystal orientation and surface roughness on friction and wear ofdiamond have also been studied rather extensively in previous years. It has been shown that surfaceroughness profoundly influences the friction and wear performance of thin diamond films. As mentionedabove, the surface rms roughness of synthetic diamond films can vary from 10 to >1000 nm, dependingon deposition conditions, growth orientation, grain size, and film thickness. As demonstrated by Caseyand Wilks (1973), and Samuels and Wilks (1988), the extent of interaction between sliding-surfaceasperities of diamond may contribute to the amount of total energy being dissipated in the form(s) ofphonons and/or vibrations. In general, higher surface roughness leads to higher levels of vibration and,hence, high friction.

Recent systematic studies by Hayward (1991), and Hayward et al. (1992), Bhushan et al. (1993),Miyoshi et al. (1993), and Erdemir et al. (1996b) further confirm the results of above-mentioned studies

FIGURE 24.8 Variation of friction coefficient of CVD diamond films sliding against themselves in vacuum attemperatures to 850°C. (From Gardos, M.N. (1999), Tribological fundamentals of polycrystalline diamond films,Surf. Coat. Technol., 113, 183-200. With permission.)

in that the diamond films with high surface roughness always lead to high friction and wear. Conversely,when polished or NCD coatings are used in sliding-contact experiments, very low friction coefficientsare attained (Erdemir et al., 1997a; Erdemir et al., 1999; Miyoshi et al., 1993) (Figure 24.9). Reportedfriction values typically range from 0.03 (for ultra-smooth nanocrystalline or polished diamond films)to >0.5 (for very rough microcrystalline films). In all cases, friction tends to be higher initially, butdecreases substantially as sliding continues in open air. The trend appears to be the opposite during testsin UHV environments where, regardless of the counterface materials, the friction coefficients are initiallylow but go up significantly as sliding continues (Miyoshi et al., 1993).

The type of counterface material used during friction tests also makes a substantial difference in themeasured friction coefficients of bulk diamond and/or diamond films. Usually, diamond-against-dia-mond or diamond films give lower friction than diamond-against-nondiamond counterfaces (Haywardet al., 1992; Miyoshi et al., 1993). When bulk diamond or diamond films are slid against relatively softermetals, alloys, or ceramic counterfaces, the steady-state friction and wear are largely controlled by theextent of surface roughness of the sliding diamond side and by the tendency of the counterface materialto form a transfer layer (Hayward et al., 1992; Gangopadhyay et al., 1993; Erdemir et al., 1997b). Rougherdiamond surfaces usually cause greater plowing and abrasive wear on the softer metallic or ceramiccounterfaces and also facilitate the transfer of worn debris particles to its sliding surface. Fine debrisparticles can easily fill in the valleys between sharp asperities of diamond films and, eventually, slidingcontact largely occurs between the transferred and original material of the non-diamond counterfaces.

Previous research has demonstrated that growth orientation of microcrystalline diamond films cansignificantly affect friction and wear. Films with large grains and highly faceted surface finish usuallycause high friction and wear (Hayward, 1991; Blau et al., 1990; Yee et al., 1995; Erdemir et al., 1996b).In recent years, researchers have developed special procedures to control growth orientation and, hence,the surface roughness of diamond films. From a tribological standpoint, ⟨100⟩ type growth orientationis very desirable and can be achieved by controlling deposition conditions and by pretreatment of thesubstrate surfaces (Yee et al., 1997; Wolden et al., 1997; Wolter et al., 1996). Diamond films with coarse

FIGURE 24.9 Relationship between rms surface roughness of CVD diamond films and their friction coefficients.(Data from Bhushan, B., Subramanian, V.V., Malshe, A., Gupta, B.K., and Ruan, J. (1993), Tribological properties ofpolished diamond films, J. Appl. Phys., 74, 4178-4185; Erdemir, A., Halter, M., Fenske, G.R., Krauss, A., Gruen, D.M.,Pimenov, S.M., and Konov, V.I. (1997a), Durability and tribological performance of smooth diamond films producedby Ar-C60 microwave plasmas and by laser polishing, Surf. Coat. Technol., 94/96, 537-541; Erdemir, A., Halter, M.,Fenske, G.R., Zuiker, C., Csencsits, R., Krauss, A.R., and Gruen, D.M. (1997b), Friction and wear mechanisms ofsmooth diamond films during sliding in air and dry nitrogen, Tribol. Trans., 40, 667-673; Erdemir, A., Fenske, G.,and Wilbur, P. (1997c), High-temperature durability and tribological performance of diamond and diamondlikecarbon films, in Protective Coatings and Thin Films: Synthesis, Characterization and Applications, Pauleau, Y. andBarna, B. (Eds.), NATO ASI-High Technology Series, Vol. 21, Kluwer Academic, London, 169-184. With permission.)

grains and ⟨111⟩ type growth orientation tend to be very rough and can cause severe abrasive wear(Erdemir et al., 1996b).

Figure 24.10 shows the frictional performance of two films with ⟨111⟩ and ⟨100⟩ growth orientation.These films were grown in an H2/CH4 plasma. After an initial run-in period, the friction coefficients ofboth coatings decreased. The friction coefficients of the film with ⟨100⟩ growth orientation stabilized at≈0.1, whereas the friction coefficient of the film with ⟨111⟩ growth orientation and micropyramidalasperities remained high and unsteady but continued to decrease steadily during successive sliding passes.

The high friction coefficients of rough diamond coatings with ⟨111⟩ orientation can be attributed tothe abrasive cutting and ploughing effects of sharp surface asperities on the softer counterface pins. If afavorable ⟨100⟩ growth orientation is present, such coatings can also afford low friction coefficients tosliding surfaces, despite relatively higher measured surface roughness. In general, previous studies con-firmed that regardless of the grain size, diamond coatings with a smooth surface finish provide very lowfriction to sliding counterfaces.

Apart from physical roughness and chemical passivation effects, phase transformation or structuralchanges can also play a major role in the friction and wear performance of diamond coatings. The extentof such changes can be dominated by environmental species or by ambient temperature (Jahanmir et al.,1989; Gardos and Ravi, 1994; Erdemir et al., 1997). Phase transformation can readily occur even in naturaldiamond (see Gogotsi et al., 1998) when extreme contact pressures and/or high frictional heating arepresent at local asperity levels. Real contact occurs first between these asperities, and their tips can eitherfracture or undergo phase transformation because of the extreme pressures and high frictional heating.Thermodynamically, graphite is the most stable form of C, whereas diamond is metastable. It is alsoknown that when excited thermally or by ion bombardment, diamond can transform to a graphitic form(Lee et al.,1993). The graphitic debris particles can gradually accumulate at the sliding contact interfaceand then begin to dominate the long-term sliding friction and wear performance of these coatings.

FIGURE 24.10 Friction coefficients of Si3N4 balls sliding against diamond films with ⟨111⟩ and ⟨100⟩ type growthorientations. (From Erdemir, A., Bindal, C., Fenske, G.R., and Wilbur, P. (1996a), Tribological properties of hardcarbon films on zirconia ceramics, Tribol. Trans., 39, 735-741. With permission.)

Previous research by Erdemir et al. (1997b) and Jahanmir et al. (1989) revealed that most of the debrisparticles derived from sliding diamond surfaces exhibited a graphitic microstructure. Raman spectros-copy, electron diffraction, EELS, and transmission electron microscopy have concurrently confirmed thepresence of highly disordered graphitic debris particles at sliding contact interfaces (Erdemir et al., 1997b).

24.2.3.3 Tribology of Diamond at Nanoscale

Recent fundamental studies on nanometer and molecular scales with new experimental tools (e.g., frictionforce microscopy, surface force apparatus, and quartz crystal microbalance) and computational methods(e.g., molecular dynamics simulation) have allowed friction studies on very fine scales (Mate, 1993;Bhushan, 1998; Germann et al., 1993; Harrison et al., 1992a,b; 1993; Harrison and Brenner, 1994; Miyakeet al., 1995). These studies have greatly increased our fundamental knowledge base of the tribology ofdiamond.

Superhigh-speed computers with high memory and processing capability have allowed simulation andreal-time visualization of atomic-scale phenomena that occur at sliding diamond surfaces. For the mostpart, results from fundamental nanoscale experimental studies are consistent with those from earliermicro- to mesoscale studies, in which standard tribology test machines were used. The nanoscale studieshave revealed that surface adsorbates, ambient temperature, and gaseous, liquid, or solid debris particlesthat are present or generated at the sliding diamond interfaces play critical roles in the friction and wearof diamond. For example, it was shown that when adsorbed gases are removed from sliding diamondsurfaces, friction increases rapidly, because the dangling surface bonds are free and highly activated toform strong bonds across the sliding contact interface (Harrison et al., 1992a; Harrison and Brenner,1994). Despite a large discrepancy in time and length scales, molecular dynamics simulation has providedsignificant insight into the extent of physical, chemical, and mechanical interactions that occur at slidingdiamond interfaces on the atomic scale.

24.2.4 Tribological Applications

Over the years, great strides have been made in the deposition, characterization, and industrial utilizationof diamond films (Seal, 1995; Ravi, 1994; Lux and Haubner, 1993; Murakawa, 1997). Our fundamentalunderstanding of the structure, properties, and performance of these coatings has increased tremendouslyin recent years and this understanding has been used to optimize, design, and customize new diamondcoatings that can meet the increasingly stringent needs of advanced tribological applications. While someuses are still in the exploratory stage and require further development, others have been fully establishedand are offered on a commercial scale. In particular, cutting tools (e.g., inserts, end mills, microdrills,push pins, tab tools, etc.) are offered commercially by several industrial manufacturers. Recent laboratoryand field tests confirm that these coated tools provide much improved performance during metal-cuttingoperations by allowing high-speed machining at increased feed rates. Some of the advantages thatdiamond provides in these applications include extreme hardness and wear resistance; good fatiguestrength; high chemical inertness; excellent resistance to abrasion, erosion, and corrosion; high thermalconductivity; low friction; and excellent environmental compatibility.

The current market share of diamond-coated cutting tools is relatively small, and the widespreadutilization of diamond coatings in other tribological fields has not yet been fully realized. Some of themajor reasons for the slow progress in large-scale utilization of diamond coatings are high fabricationcost, limitations on the type and size of substrate materials, oxidation when used at elevated temperatures,insufficient reliability and reproducibility, high surface roughness of finished products, somewhat pooradhesion, high deposition temperatures, and slow growth rates. Diamond-coated tool inserts cannot beused to machine pure Fe and Fe-base materials and the alloys of groups IVa, Va, VIa, VIIa, and VIIIa ofthe Periodic Table, which represent the largest segment of industrial machining. Diamond can chemicallyreact with and/or dissolve in these materials at the high temperatures generated during machining, andthus it wears out rather quickly. Also, diamond coatings cannot be produced on the surfaces of tools thatare made of materials such as high-speed steels, which are widely used in the tooling industry. Thus far,

diamond films can only be deposited on certain refractory metals (e.g., W and Mo) and a few ceramics(e.g., WC, Si, SiC, and Si3N4) that can endure the high deposition temperatures of the CVD process.Among others, WC-based tools have been coated successfully with diamond and made available forcommercial machining. Although prototypes of other tribological parts (e.g., mechanical seals; Hollmanet al., 1998) have also been demonstrated, their large-scale utilization has not yet been realized becauseof high fabrication costs and insufficient reliability in actual operation. Other potential applications fordiamond coatings include sliding bearings, MEMS, wire-drawing dies, and various wear parts that areused against erosion and abrasion (e.g., jet nozzles).

24.2.4.1 Industrial Machining

Diamond has been used for machining purposes in three distinct forms. One is based on the use ofsynthetic diamond powders produced by HPHT methods. Specifically, fine diamond powders are firstsintered into the shape of a tool tip and then brazed directly onto the cutting edges of the tool inserts.The second form is based on the CVD of a relatively thick (i.e., 0.5 to 1 mm thick) diamond film on anSi wafer from which small bits are laser-cut into a desired cutting-edge shape. Subsequently, the Sisubstrate is etched out and the cut diamond tips are recovered. As a last step, the freestanding diamondtips are bonded or brazed onto the cutting edges of the tool inserts or end mills. The third form is thedirect deposition of diamond onto cutting tools by various CVD methods.

There are some fundamental differences between brazed and coated diamond tool bits. First, diamond-coated tools can be tailor-made to meet the specific machining needs of an application, whereas thebrazed diamond tips come in a strict geometric form and cover only the sections of the sharp tips oftool inserts. Second, the entire functional surface of a tool insert can be coated with diamond, and it canbe prepared in any style or geometry to provide better efficiency and overcome limitations on depth ofcuts — always a problem with brazed tool bits. Finally, the shape and area of a coated carbide insert canbe prepared in such a way that it will allow better chip handling and breaking capability than the verysmall brazeable diamond tips. CVD diamond coatings perform equally well or even better than PCDdiamonds in cutting and turning operations. However, edge chipping or deterioration due to easy cleavageof highly oriented needle-like diamond grains of CVD diamond films can occasionally occur and limitthe lifetime of the coated tools. Figure 24.11 shows the morphology of the cutting edge of a diamond-coated tool insert.

As mentioned above, diamond-coated tools are primarily used in the machining of nonferrous metals,alloys, and composite materials that are inherently very difficult to cut or machine. In particular, diamond

FIGURE 24.11 Cutting-edge morphology of a CVD diamond-coated tool insert.

has been the material of choice for machining of high-Si-content Al alloys for automotive applications.Hypereutectic Al-Si alloys are increasingly being used in this industrial sector, mainly because of theirlight weight, relatively high strength, and high resistance to wear (Deuerler et al., 1996; Keipke et al.,1998; Prengel et al., 1998; Shen, 1996, 1998; Trava-Airoldi et al., 1998). The major applications for Al-Sialloys include pistons, engine blocks and heads, wheels and chassis, and some transmission parts. Theincreased use of lightweight materials in the automotive industry is likely to continue and, hence, theneed for diamond-coated tools will increase. Other engineering materials suitable for machining bydiamond-coated tools include alloys of Mg, Cu, Pb, and Mn, as well as various composites (e.g., fiber-reinforced plastics, fiberglass composites, C-C composites, and metal-matrix composites), bulk graphite,plastics, epoxy resins, green ceramic, concrete, and various mining and rock-drilling operations (Kubel,1998; Lux and Haubner, 1998; Komanduri and Nandyal, 1993; Inspektor et al., 1997; Karner et al., 1996).Another key application for diamond coatings is in woodworking tools.

The most important factor that governs the success of a diamond-coated tool insert in cutting oper-ations is strong bonding between diamond coating and substrate insert (i.e., WC-Co). If bonding is notintimate and strong, diamond coatings can easily be removed from the cutting edges, which thus losetheir cutting performance and effectiveness. In fact, premature delamination of the diamond film in earlytrials represented the biggest challenge and most common mode of tool failure. Residual stresses areinherent in all CVD-diamond coatings, and these stresses can significantly affect the film/substrateadhesion (Olson and Dawes, 1996; Gunnars and Alahelisten, 1996). Diamond coatings are often depositedat high temperatures, and a large mismatch in thermal-mechanical properties between the diamondcoating and the underlying substrate leads to large compressive stresses. The presence of a high stress inthe coating can lead to various undesirable failure modes that can cause the film to delaminate, crack,or blister and thus destroy the entire structure. Other factors, such as film thickness, film roughness,substrate preparation, and substrate grain size, also affect film/substrate adhesion, and they must beconsidered along with the residual stresses whenever the quality of a diamond coating is evaluated. Twoother problems can occur when diamond films are used in machining operations. One is oxidation; theother is fracture or chipping. Despite its extreme hardness, diamond is brittle and will readily fracture,especially when forced in the direction of relatively weak columnar grains.

Currently, diamond coatings are primarily produced on WC-Co-based cemented carbide tool inserts,mainly because of their excellent toughness, hardness, and high-temperature durability. Co serves as abinder and controls the toughness of these materials. In addition to strict process control during depo-sition, the selection and pretreatment of the substrate materials are extremely important for achievingstrong adhesion between the diamond coating and the substrate material. In particular, the selection ofthe correct types of WC-Co material is a must for achieving strong bonding and, hence, long wear life.In principle, a low Co content (<5%) in cemented carbide is highly desired. Higher Co contents canadversely affect the adhesion of diamond coatings to carbide inserts. Specifically, during deposition athigh temperatures, Co can play a catalytic role in the formation of non-diamond phases and complexCo carbides that are highly undesirable. It also slows the nucleation process and reduces nucleationdensity. Interfaces with a high proportion of non-diamond phases and low nucleation density suffer frompoor adhesion. Too much Co can also increase the difference in thermal expansion coefficient betweendiamond and substrate; hence, during cooling, very high compressive stresses can build up at the interfaceregion. When these stresses are combined with the stresses associated with cutting action, prematureadhesive failures can occur at or near the cutting edges. Briefly, it is important to use carbide substrateswith low Co content or to etch the Co out completely from the near surface; otherwise, the coatings willnot stick to the tools (Menningen et al., 1994; Cappelli et al., 1998; Toenshoff et al., 1998).

Several methods are used by industry to remove surface Co from cemented carbide tool inserts. Mostof the methods involve selective etching of Co by chemical means. Some of the chemicals that have provedto be effective in etching Co are HNO3, HCl, and CH3COOH. In one case, tool inserts are dipped in anitric acid and water solution (mixed 1:1 by volume) and ultrasonically agitated for 15 min. They arethen rinsed and ultrasonically cleaned in ultrapure water for 5 min. In another case, researchers haveheat treated WC-Co substrates at very high temperatures for an extended period of time in a protective

environment. Heat treatment causes grain growth in WC and hence roughens the surface; it also evap-orates Co and thus reduces the amount of Co at or near the surface. Increased surface roughness providesbetter mechanical interlocking between diamond and WC grains, whereas the absence of Co ensureshigher nucleation density and better adhesive bonding.

Because Co is the binder that holds the tool’s WC grains together, Co removal must be done withgreat care. If too much Co is removed from the surface, the tool will become weakened and its integritynear the surface will be impaired. Formation of stable Co borides and silicides that can endure the highdeposition temperatures of diamond has also been used in the past to prevent Co-associated problems(Kupp et al., 1994; Kubelka et al., 1994). Various bond layers (i.e., W, Si, SiC, and Si3N4) have also beentried on tool inserts to achieve strong bonding. The effectiveness of bond layers is very much dependenton the exact composition and thickness of the bond layer material. Few other procedures have beendeveloped by industry to achieve strong adhesion, but they are considered highly proprietary; hence,very little is known about them.

Published case studies indicate that, depending on the material cut and cutting conditions, tools coatedwith CVD diamond can exhibit tool life improvements that range from 25% to more than a factor of40. Greater improvements in tool life are reported when cutting Al-Si alloys, green ceramics, and fiber-reinforced composites. Figure 24.12 shows the performance of uncoated and CVD-diamond-coatedcarbide inserts during machining of an Al-Si alloy for a truck wheel application.

Figure 24.13 demonstrates the effectiveness of a nanocrystalline diamond-coated carbide insert inpreventing edge wear and loss of sharpness after machining of a high-Si-content Al alloy. CVD diamond-coated tools were proven to be as good as or better than the PCD tool inserts when used during turningof Al-Si alloys and graphite.

The market share of coated tools is expected to increase steadily because prices are declining whileperformance and reliability are improving. Better quality and reliability will certainly result in diamond-coated tools that last longer, cost less, and increase productivity. Several companies are engaged in theproduction of CVD diamond-coated tool inserts. DeBeers, Mitsubishi, Nachi-Fujikoshi, Norton, Kenna-metal, Sandvik, and Sumitomo are the major firms. Kennametal and Sandvik provide diamond-coatedtools commercially. As discussed above, CVD diamond-coated inserts are primarily targeted for machin-ing applications that involve high-Si-Al alloys. The projected growth of Al use in automobiles indicates

FIGURE 24.12 Cutting performance of an uncoated and CVD diamond-coated insert while machining a castaluminum alloy containing 7% Si (cutting conditions: speed, 150 to 2500 m/min; depth of cut, 0.5 to 3 mm; feedrate, 0.25 to 0.8 mm/rev). (From Karner, J., Pedrazzini, M., Reineck, I., Sjostrand, M.E., and Bergmann, E. (1996),CVD diamond-coated cemented carbide cutting tools, Mater. Sci. Eng., A209, 405-413. With permission.)

that the growth of the diamond-coated tool insert market will be significant. When the performance andprice of CVD tools compare favorably with those of PCD tools, CVD tools may displace PCD in someapplications. Woodworking tools represent another major market segment for diamond coatings. Tung-sten carbide blades already comprise a large part of the woodworking tool market. CVD diamond-coatedtools have yet to be used in woodworking and, thus, they offer an excellent business opportunity.

24.2.4.2 Mechanical Seals

After cutting tools, mechanical seals represent the next best application possibility for CVD diamondfilms. The mechanical-seals market is rather large, ranging from the very simple seals used in automotivewater pumps to higher-end seals that are used at high temperatures and in aggressive environments.Coated seals would be especially attractive in applications where high chemical stability and resistanceto corrosion, erosion, and abrasion is needed. Because of its low friction, a diamond-coated seal can alsobe very useful in applications where a liquid lubricant is not permissible, for example, in pharmaceuticaland food processing plants.

Over 3 million higher-end pumps rely on the tribological performance of SiC seals, and SiC is an idealsubstrate for diamond film growth. Improper use of or leakage in mechanical pump systems can be verycostly and may lead to environmental disasters. Total energy losses due to high friction and wear can bevery high. Because properly applied diamond coatings afford very low friction to the sliding surfaces ofseal components, they can extend the useful life of these components and also reduce energy losses.Recent studies by Hollman et al. (1998a,b) have demonstrated that smooth diamond films can result insubstantial reductions in frictional torque and wear rates of shaft seals.

Rough diamond coatings cannot be used in seal applications. One of the sealing faces (inserts) is madeof soft graphite or C composite. When rubbed against a rough and superhard surface, these inserts wearout rather quickly. Hard SiC inserts can also be used, but they are much more expensive and, whenrubbed against rough diamond, they too suffer major wear losses. Polished or smooth nanocrystallinediamond coatings will be more desirable for seal applications.

24.2.4.3 Other Applications

Microelectromechanical systems represent a new class of moving mechanical assemblies that can beused for a wide range of applications, including sensors, actuators, high-precision positioning devices,

FIGURE 24.13 Extent of wear damage on (a) uncoated and (b) nanocrystalline diamond-coated tool inserts aftermachining of a hypereutectic aluminum-silicon alloy.

microelectronics, etc. MEMS devices such as gear assemblies and micromotors have been largely fabri-cated by Si micromachining technology; however, Si exhibits very poor mechanical and tribologicalproperties. When these microdevices are used in very high-speed applications, they suffer from unac-ceptably high friction and wear losses and thus are unsuitable for applications that involve high speedsand realistic loads. Researchers have developed alternative fabrication methods that allow production ofMEMS devices from SiO2, Si3N4, or SiC-type materials. However, the tribological performance of thesematerials is also very poor; hence, they may not function well in a dynamic MEMS application.

Because of its excellent properties, researchers have recently been exploring possible uses of diamondcoatings for MEMS applications (Bjorkman et al., 1998; Aslam and Schulz, 1995; Mao et al., 1995;Ramesham, 1999; Gardos, 1998; Cagin et al., 1999). However, there have been some difficulties infabricating diamond coatings and/or components with a morphology suitable for MEMS. The NCDfilms discussed previously appear to be better suited because of their very small grain size (i.e., 10 to30 nm). With these films, it will be possible to attain film coverage at very small thicknesses (as thin as50 nm) and hence preserve the fine structural features of MEMS devices. Conventional CVD films withthicknesses of 0.5 to 10 µm may not achieve complete coverage on the surface of Si-based MEMS devices.

Other tribological applications for CVD diamond coatings include woodworking tools, tab tools, pushpins, high-precision microdrills, surgical blades, wire-drawing dies, and various wear parts that are usedagainst erosion and abrasion (e.g., jet nozzles).

24.3 Diamond-like Carbon (DLC) Films

DLC films represent a noteworthy example of thin films whose properties can be varied over a widerange of structures and compositions. The tribological behavior of these films strongly depends on thedeposition technique. Consequently, the first part of this discussion is dedicated to a background onDLC, highlighting the relationships between the deposition process, the film structure, and some filmproperties. The second part summarizes the tribological behavior of DLC films, through a descriptionof the friction and wear mechanisms, successively at the macroscopic, microscopic, and nanometer scales.

24.3.1 Background on DLC Films

24.3.1.1 Classification of DLC Films

Diamond-like carbon coatings have been the subject of intensive studies for the last 20 years. The generalterm “DLC” describes hydrogenated and hydrogen-free metastable amorphous carbon materials, pre-pared by a wide variety of PVD and CVD techniques. The films exhibit a wide range of structure,composition, and attractive mechanical, optical, electrical, chemical, and tribological properties. The filmstructure and properties are determined by the H content and the relative ratio of the two sp2 and sp3

carbon hybridizations, the sp1 C hybridization being negligible. Hydrogen in DLC is important forobtaining a wide optical gap and a high electrical resistivity, removing midgap defect states, stabilizingthe random network, and preventing its collapse into a graphitic phase.

As proposed by Robertson (1998), the wide range of DLCs is conveniently displayed on a ternary phasediagram in Figure 24.14. The sp2-bonded amorphous C lies in the lower left corner. Phases with an Hcontent that is too high, lying at the far right of the diagram, cannot form an interconnected networkand form gas or liquid molecules. The nonhydrogenated films are classified either as sputtered amorphous C(Sputtered C) with a predominance of the sp2 hybridization, or as tetragonal amorphous C (ta-C), witha predominance of the sp3 hybridization of up to 85%. The most common type of DLC is the hydrogenateda-CH film, with only moderate sp3 content and a rather high H content of up to ≈50 at%. A lightlyhydrogenated analogue of ta-C, termed ta-CH (Weiler et al., 1996), is also shown. Donnet (1998)attempted to modify some DLC properties by the addition of alloying elements, mainly Si, F, N, andvarious metals.

Such a wide range of film structures and composition and the diversity of methods used for DLC filmdeposition provide the flexibility to tailor their properties according to specific needs and applications.The study of hydrogenated DLC and the state of our knowledge concerning its properties and practicalapplications have apparently matured. The study of hydrogen-free DLCs has not yet reached this stateand practical applications for it have yet to be found.

24.3.1.2 Influence of Deposition Methods on Film Microstructure and Composition

From the point of view of optimizing the routine production of DLC films, the primary need is foreffective process control, primarily the tailoring of adhesion properties. Indeed, the adhesion of DLCfilms varies widely, depending on the nature of the substrate. To be used as tribological coatings, DLCfilms must adhere well to the substrate material, and the adhesive forces must overcome the high internalstresses that would otherwise cause film delamination (Holmberg and Matthews, 1994). Adhesion canbe affected by the deposition method, in combination with the nature of the substrate. Good adhesionof DLC films is observed on carbide- and silicide-forming substrates. The adhesion of DLC coatings tosilicide-forming metals can be improved by depositing a 2- to 4-nm-thick interfacial layer of amorphousSi between the metal and the C film, thus forming an interfacial silicide layer promoted by the plasma,even at relatively low substrate temperatures. The use of Ti-based, functionally gradient films intercalatedbetween the substrate and the DLC top layer has also been studied (Voevodin et al., 1997). The depositionof adhesion and interface layers is ideally achieved in the same reactor as the deposition of the DLC films,by multiplex processes that minimize the introduction of defects and impurities between the superim-posed layers and allow precise control of the entire deposition methodology.

Whatever the deposition conditions, DLC films are generally smooth to the level of tenths of ananometer and are therefore referred to as nanosmooth. The roughness thus conforms to the underlyingtopography. All methods for the deposition of DLC are nonequilibrium processes characterized by theinteraction of energetic species with the surface of the growing film. The growth and properties of DLCfilms are controlled by the substrate temperature and the energy of the species impinging on the surfaceduring the deposition process, the latter exerting dominant control. A wide literature survey on thedeposition procedures has been assembled by Grill and Meyerson (1994). The following discussion islimited to the techniques that are most commonly used to deposit DLC films, with emphasis on therelationships between deposition conditions and film structure. The deposition of ta-C films is summa-rized prior to a discussion of the deposition of a-CH films.

Today, the main processes used to deposit nonhydrogenated DLC films are magnetron sputtering,mass-selected ion beam, cathodic vacuum arc, and laser plasma. Whatever the technique, the fraction ofsp3 bonding is maximized for ion-dominated processes with ion energies ≈100 eV. The specific propertiesof ta-C films are achieved by the high energies of the impinging particles, with film growth governed by

FIGURE 24.14 Ternary “phase diagram” showing the various forms of diamond-like carbon. (From Robertson, J.(1998), Deposition mechanism of diamond-like carbon, in Amorphous Carbon: State of the Art, Silva, S.R.P., Robert-son, J., Milne, W.I., and Amaratunga, G.A.J. (Eds.), World Scientific Publishing, Singapore, 32-45. With permission.)

a subplantation process instead of conventional condensation as in a-CH films (Robertson, 1998).Consequently, the temperature during the deposition process is also a crucial parameter. Above a tran-sition temperature of 250°C, a noticeable decrease in the high sp3 content and density is generallyobserved. This transition temperature, which is much lower than the thermal stability temperature ofta-C films (which can reach relatively high values approaching 600°C), decreases with increasing ionenergy.

Magnetron sputtering is probably the most commonly used industrial method for depositing DLCfilms. In the unbalanced configuration, the substrate is bombarded with Ar ions, thus increasing the sp3

content of the films. Ar/H mixtures allow the production of hydrogenated DLC films. With the mass-selected ion-beam technique, it is possible to obtain ta-C films that contain up to 80% sp3-hybridized Cby varying the acceleration, mass, and energy selection of a C ion beam that is condensing onto thesubstrate. ta-C films can also be obtained by a cathodic vacuum process on both the laboratory andindustrial scales. An arc is struck on a graphite cathode, creating an intense, 30%-ionized plasma thatexpands through the deposition chamber and condenses onto the substrate. The filtered cathodic vacuumarc configuration avoids degradation of the growing film by particulates produced in the arc along withthe desired reactive species. Laser plasma deposition is a versatile method used to deposit many materials;it is widely used for DLC film deposition by creating a plasma plume similar to the cathodic arc producedby UV laser ablation of a graphite target.

Plasma-enhanced CVD, in which any gaseous hydrocarbon species is used as precursor, is probablythe most popular means to deposit a-CH films. The chamber is configured with cathode and anodeplates, with the substrate attached to the cathode to maximize ion bombardment. Parallel plate reactorsare preferred because they allow the deposition of uniform films over large areas. RF power can becapacitively coupled to retain the ability to deposit the films on a wide variety of substrates, includinginsulators. Because the film structure is a strong function of the energy of the impacting species, thisability depends on several plasma parameters — primarily the bias and the gas pressure. The bias isindirectly controlled by the RF power and the gas pressure in the RF configuration. The sp3 and H contentare strongly controlled by the average impact energy. More precisely, at the lowest impact energies, thegaseous precursor is not sufficiently decomposed, and polymer-like C films with a predominance ofCH2 groups are generally obtained. At intermediate impact energies, the H content has declinedsufficiently so the C–C sp3 bonding has reached a maximum, thus leading to the so-called diamond-likequalities. However, if the impact energies become too high, graphite-like C structures are generallydeposited because of an increase of disordered sp2-like bonding, at the expense of a decline of C–C sp3

bonding.From characterization by Fourier transform infrared (FTIR) combined with forward recoil elastic

scattering experiments, a significant amount of H in a-CH films has been unbound to C (Grill and Patel,1992). This free H, optically inactive by FTIR, should be trapped in some voids of the carbonaceousnetwork. Free H fraction increases when the impact energy, and hence the dissociation of precursor, isincreased. This also leads to higher C network crosslinking. Consequently, the nature of the precursorplays a direct role in the fraction of H bonded to C. However, based on recent comparisons betweenFTIR and 13C/1H NMR experiments performed on typical DLC films obtained under various conditions(Donnet et al., 1999), the absolute quantification of the fraction of H that is bonded to C remainsambiguous.

As with ta-C films, temperature during deposition is another parameter that generally affects the DLCstructure. Deposition should be carried out below 250°C to stabilize a significant sp3 content. The increaseof diamond-like qualities of a-CH films has been achieved using a plasma beam source with acetyleneto reduce the amount of incorporated H, and by operating at lower gas pressure to increase plasmaionization (Weiler et al., 1996). The resultant films are called ta-CH, by analogy to ta-C. The sp3 contentcan reach a maximum near 80% and H content in the range of 25 to 30%, with an effect of depositiontemperature that is more complex than that for ta-C films.

Metal- and/or nonmetal-doped diamond-like films are deposited by the same techniques as the regularfilms, by introducing into the reactor species or precursors that contain the modifying elements.

24.3.1.3 Influence of Film Structure and Composition on Some Typical Properties

As described in Section 24.3.1.2, the structure and composition of the films (H content, fraction of Hbonded to C, C hybridizations, nature of bonds, alloying elements) are strongly influenced by the averageimpact energy of the particles that impinge on the growing surface and by the deposition temperature.Thus, most film properties depend on key structural parameters that are related to specific depositionprocesses and conditions. Table 24.1 summarizes the range of various typical properties that depend onthe presence or lack of H in the film. Nonhydrogenated films exhibit generally higher hardness, Young’smodulus, density, thermal stability, and compressive stress; and lower index of refraction, resistivity, andbandgap, when compared with hydrogenated films. For each category of film, the property variation isdirectly related to the structure, as controlled by the deposition process. The relationships between thestructure and the properties of DLC films have been more extensively studied on a-CH than on ta-Cfilms. In hydrogenated DLC films, an increase in H content is generally associated with a decrease inhardness, Young’s modulus, stress, thermal stability, density, and index of refraction, and an increase inthe gap and electrical resistivity. A more accurate discussion remains impossible because most of theseproperties also depend on the fraction of H that is bonded to C, which remains difficult to investigatefor the reasons mentioned previously.

The difficulties in generalizing the structure/property relationships are due to a lack of systematic andunambiguous standardized characterization of DLC films, the wide range of structures and compositions,and the experimental difficulties encountered when thin metastable amorphous films are accuratelyprobed. Differing models, such as the graphitic cluster, the random covalent network, or the defectivegraphite network reviewed by Grill and Meyerson (1994), constitute interesting guidelines to improveour current knowledge about structure/property relationships.

Various C-based materials, similar in structure to a-CH or ta-C, include elements such as Si, F, N, andvarious metals. The introduction of dopants generally brings about a decrease in compressive stress, dueto fewer interconnections in the random carbonaceous network. Grischke et al. (1995) have shown thatthe specific chemical composition of the modified films strongly influences the surface energy, dependingon the nature of the dopant, and may modify various physical properties, making some doped DLCssuitable for various applications — for example, to improve field emission characteristics. As highlightedin Section 24.3.2, the introduction of alloying elements into DLC affects many properties, includingtribological behavior.

24.3.2 Tribology of DLC Films

Much effort has been invested in the characterization of the tribological behavior of DLC films, mainlybecause of the very low friction coefficients and high wear resistance of these materials. Thus, the tribologyof DLC coatings has been discussed extensively in the literature, and has been summarized in recent

TABLE 24.1 Range of DLC Structure, Composition, and Properties

Variable a-Ca a-CHb

Hydrogen content, at% <5 20–60sp3, % 5–90 20–65Density, g/cm3 1.9–3.0 0.9–2.2Thermal stability, °C <600 <400Optical gap, eV 0.4–1.5 0.8–4.0Electrical resistivity, Ω/cm 102–1016

Index of refraction 1.8–2.4Compressive stress, GPa 0.5–5Hardness, GPa <80 <60Young’s modulus, GPa <900 <300

a a–C = amorphous carbon.b aCH = amorphous carbon, hydrogenated.

reviews such as those of Holmberg and Matthews (1994), Grill and Meyerson (1994), Grill (1997a),Donnet (1995, 1998), and Erdemir (1998). Friction and wear of DLC coatings are strongly affected bythe nature of the films, as controlled by the deposition process, and by the tribotesting conditions,including material parameters (nature of the substrate), mechanical parameters (contact pressure), kine-matic parameters (nature of motion, speed), physical parameters (temperature during friction), andchemical parameters (nature of the environment). An exhaustive presentation of the tribological behaviorof DLC is thus impossible. A more interesting and useful method of presentation will be used here, inwhich a top-down approach to DLC tribology is used to focus on the various phenomena occurring atthree different scales. The macroscopic scale allows the establishment of relationships between the natureof the films and friction levels and wear rates under various testing conditions. The microscopic scalefocuses on the accommodation modes, in terms of interfacial shearing mechanisms, transfer film buildup,and friction-induced wear-particle formation. The nanometric scale focuses on the chemical reactivityof the top surface in relationship to the environment that surrounds the contact spot, and the surfacechemistry of the films deposited under differing conditions.

Below, the current knowledge at these three complementary scales is summarized.

24.3.2.1 Tribology of DLC at the Macroscopic Scale: Friction and Wear Behavior

The friction data on DLC films show that their sliding friction coefficients, µ, span the range of 0.003 to0.6, which probably represents the widest range of friction behavior in any category of coatings. Wearrates are less systematically investigated and are often recorded in nonstandard units, thus increasing thedifficulty of comparing results. DLC films that are easily scratchable with no wear resistance at all, aswell as films with normalized wear rates as low as 10–8 mm3/N-m, can be produced.

Friction and wear control can be achieved, primarily by considering the nature of the DLC films,together with the environmental conditions. In ambient humid air (typically 20% < Relative Humidity <60%), friction generally ranges between 0.05 and 0.3, with wear rates strongly depending on the natureof the films. Hydrogen-free films generally exhibit lower friction (<0.15) when compared with a-CHcoatings. This finding is consistent with the expected role of H in stabilizing the random covalent networkand preventing its collapse into a graphitic network. In ta-C films, friction easily causes a local shear-induced graphitization at the microscopic scale, as described in Section 24.3.2.2. Consequently, it is notsurprising to observe a decrease in the friction of ta-C films with increasing humidity (Voevodin et al.,1996), because water molecules are known to intercalate between graphite layers and ease their slip overeach other. On the other hand, friction of a-CH films generally increases with humidity (Franks et al.,1990). The role of water vapor at both the microscopic and nanometric scales is highlighted inSections 24.3.2.2 and 24.3.2.3.

Unlike friction, wear resistance performance is not so easy to discuss by simply discriminating betweenthe two previous categories of DLC films. Extremely low wear rates can be achieved for both ta-C anda-CH coatings. In the case of H-containing films deposited by PECVD, Grill (1997b) has correlated wearresistance (down to 0.1 nm for 1000 rotations in ambient air) with deposition conditions (precursor,bias, and pressure). Polymer-like films obtained at low average impact energy (controlled through lowDC bias) are less wear resistant than diamond-like films. Thus, based on the accumulated knowledgedescribed in the previous sections on deposition/structure/property relationships, this finding shows thatthe optimization of the wear resistance of DLC films requires strong attention to the control of thedeposition procedure.

In inert environments, including dry N and vacuum, friction coefficients can reach either ultralowvalues (down to 0.01 or less [Erdemir et al., 2000a,b]) or high values (>0.5). Donnet and Grill (1997)have shown that this binary friction behavior, under UHV conditions during tribological tests, can becontrolled by the H content of the films. Hydrogen concentrations lower than 34 at% systematically leadto high friction, similar to the friction level of diamond or graphite in the same UHV environment.Ultralow friction is achieved with films that contain at least 40 at% H. The effect of H on the frictionlevel is highlighted in the chapter section dedicated to the tribology of DLC at the nanometric scale(Section 24.3.2.3).

A compilation of the tribology of doped DLC and C alloy coatings has been recently prepared byDonnet (1998). Silicon incorporation into the DLC structure is shown to affect most film properties(including a decrease in surface energy and internal stress) and its tribological behavior. Friction appearsto be significantly reduced (<0.1), where compared with that of conventional undoped DLC in ambienthumid air, with comparably high wear resistance. However, this tribological behavior appears to beobserved when the contact pressure remains below 1 GPa. At higher contact pressures, conventional ta-Cand a-CH films cannot be surpassed. Consequently, ta-CHSi films can be used in applications that requireboth low friction (<0.1) and high wear resistance (<10–7 mm3/N-m) under moderate mechanical con-ditions, for the protection of low-stress aerospace or automotive components, precision ball bearingsand gears, sliding bearings, and magnetic recording media.

The incorporation of Si and F into the DLC structure affects the surface properties. The reduction ofstress, when compared with conventional DLC, is in the same range as that in a-CHSi. However, thereduction of the surface energy is higher with F than with Si. Highly fluorinated DLC (F/[F+C] > 0.4)appears to be soft, with no wear resistance. Moderate fluorination (F/[F+C] < 0.2) can be controlled bydeposition conditions to obtain films with comparable wear resistance and friction levels of conventionala-CH films, but with lower stress and surface energy. Less research has been performed on the tribologicalinvestigation of N-containing DLC films because of their recent discovery when compared with conven-tional DLC and other doped DLC. A review of the tribological behavior of C nitride films is presentedelsewhere in this chapter (Section 24.4.2).

The range of composition and structures attainable with metal-containing DLC coatings appears tobe enormous. One should keep in mind that the optimization of the material combination and depositionparameters is a challenging subject for each element or combination of elements. When optimization isachieved, metal-containing DLC films can exhibit promising tribological properties in terms of steady-state friction levels and wear rates for various applications. Dimigen and Klages (1991) observed frictioncoefficients in the range of 0.02 to 0.04 in ambient air when a significant amount of Ta or W wasincorporated in a-CH films.

24.3.2.2 Tribology of DLC at the Microscopic Scale: Accommodation Modes

Most experimental observations indicate that for the low friction and reasonably long life of DLC, atransfer film buildup that is followed by an interfilm sliding mechanism is the most frequently observedvelocity accommodation mode (Figure 24.15). The kinetics of transfer film formation, together with itsthickness, strongly depend on the nature of the counterfaces and the environmental conditions. Theinterfilm sliding mechanism occurs when the DLC film adheres strongly to both contacting surfaces, butseparates into two distinct layers that slide across one another. The resultant shear strength is that of thetwo layers sliding against one another. Neither of the two original surfaces is in contact, except if the filmis completely worn. At this scale, the behavior of ta-C and a-CH is rather similar, and DLC films aregenerally considered to slide according to this accommodation mode. One should keep in mind that theinitial formation of a transfer film is associated with significant wear of the initial coated part. Conse-quently, this initial running-in wear may be crucial to ensure low friction and wear over long periods,and wear does not systematically vary linearly with time.

FIGURE 24.15 Interfilm sliding mechanism consecutive to transfer film formation, frequently observed with DLCcoatings.

The nature of the counterfaces can significantly influence the size and nature of the transfer layer (bothof which sometimes differ from the initial composition of the film). The transfer layers may become amixture of the original carbon film and elements from the counterface, indicating a strong triboreactivitywith the initial pin counterface. Such a mixture has already been observed by secondary ion massspectroscopy (SIMS) analysis of inside wear tracks of DLC-coated steels (Ronkainen et al., 1993). Theanalysis revealed a transfer film composed of a mixture of C, Fe, Cr, H, and O. A lower wear rate of thepin, associated with a modification of the transfer composition, was observed by SIMS analysis at highloads and high speeds. The authors also investigated, with Al2O3 pins, the coatings and experimentalsetup that had been studied earlier with metal counterfaces. The thickness of the transfer increases withhigher loads and sliding velocities; this is associated with a drastic decrease in friction coefficient from0.13 to 0.02. Alumina was transferred to the surface of the disc wear track. These results were observedboth on H-free and hydrogenated DLC films. Thus, similar mechanisms are generally observed whenceramic materials are used. Transfer layers are also formed on ceramic counterfaces such as Si3N4, withcomposition strongly depending on the environment during friction (Kim et al., 1991; Miyoshi et al.,1992). This can be explained in terms of tribochemical reactions, as described in Section 24.3.2.3.

The structure of the transfer layer can also be modified by the friction process, as frequently observedin a wear-induced graphitization mechanism (Erdemir et al., 1991, 1993, 1995, 1996; Liu et al., 1996).Lower humidity generally increases the graphitization rate, more than likely because the effect of watermolecules is reduced. A decreased graphitization rate is also observed at lower temperatures and higherhumidity during friction, and can be attributed to the suppression of temperature rises at hot spot.Moreover, the environment, controlled by the partial pressure of pure water vapor (pH2O), stronglyaffects the transfer film thickness, which appears to influence the friction level. This effect has beenobserved in hydrogenated DLC film deposited on an Si substrate and tested against a steel pin counterfacein UHV at various pH2O values (Donnet et al., 1998). A transition between ultralow friction (10–2 range)and higher friction (10–1 range) is observed within a water vapor pressure range between 0.1 hPa (RH =0.4% at 23°C) and 1 hPa (RH = 4% at 23°C).

Ultralow friction at the lowest pH2O values is associated with a homogeneous C-based transfer, asshown by in situ Auger electron spectroscopy (AES) performed inside the transfer film. Higher humidityvalues are associated with a thinner transfer film, so that AES detects the substrate elements that underliethis transfer, and friction rises to higher values. The transfer is thus impeded at high humidity values.Consequently, the nature of the counterface and the tribotesting parameters, together with the nature ofthe environment, play a crucial role in the kinetics of formation and composition of the transfer film,and thus strongly influence the friction levels and wear rates. Tribochemical reactions between the transferfilm and the counterface can induce the formation of complex structures and compositions on the topsurface of the film. Moreover, the nature and thickness of contaminant or passivation layers will probablyplay a role, but this kind of assumption can be checked only by tribological tests under UHV conditionsand sliding in situ on fresh surfaces.

An exception to this conventional transfer film formation mechanism can be found when DLC-coatedsteel slides slowly against PTFE (Yang et al., 1991). Without the DLC film, very little PTFE is directlytransferred onto the steel counterface and friction remains between 0.15 and 0.20. With the DLC film,friction decreases to 0.04, without any transfer film formation, because of the low adhesion betweenPTFE and the DLC coating.

24.3.2.3 Tribology of DLC at the Nanometric Scale: Tribochemistry and Molecular Interactions

During a sliding process, interfacial rheology is associated with exposure of the top surfaces and interfaces(third body) to one another and, except for UHV tribotests, to the surrounding atmosphere. Therefore,the chemical composition of the first layers is of paramount importance in understanding the tribochem-istry of DLC coatings.

The role of the H content and sp2/sp3 hybridizations in a-CH films is discussed first. The discussionsare based on Table 24.2. With this approach, it should be possible to explain the role of the high H contentin lowering the friction of DLC in UHV from high values that are in the same range as values related todiamond or graphite under the same environmental conditions. When transfer film buildup during therunning-in period is taken into account, steady-state sliding with DLC films is accommodated betweenthe two counterfaces by interfilm sliding of C-covered smooth surfaces. Under UHV conditions, thecontaminant topcoats are removed and sliding occurs between the two carbonaceous fresh surfaces. Inthe case of diamond/diamond or graphite/graphite, as with poorly hydrogenated a-CH films, the UHVfriction coefficient is initially in the 0.1 to 0.2 range, as long as the surface sites are saturated by H, O,or H2O molecules. Steady-state friction of diamond and un-intercalated graphite is generally very high(>0.5), as reviewed by Gardos (1994).

Desorption of adsorbates upon rubbing in vacuum creates σ dangling bonds on diamond surfaces. Ifthese free bonds do not reconstruct or are not activated by any adsorbates, they interact with high energiesand thus contribute to the high friction. In the case of un-intercalated graphite, the π-π* orbitals (i.e.,sub-bands of the graphite band structure) overlap at selected sites of the Brillouin zone. The overlap ofthese sub-bands can lead to attractive interactions that range from 0.4 eV (37 kJ/mol) to 0.8 eV(66 kJ/mol). The actual interaction force is somewhere within this range, depending on how the bandparameters influence the net inter- and intraplanar attraction of a family of imperfect graphitic planesof particles interacting in a solid lubricant layer. Gardos (1994) indicates that the 0.4 to 0.8 eV bindingenergy between the basal planes is enough to render pristine highly oriented graphite, a high frictionand wear material.

The discussion related to a-CH films is more complex in that it considers the wide distribution ofH content and C hybridizations, and determines various degrees of C crosslinking. The ultralow frictionof highly hydrogenated DLC films in UHV is consistent with the predominance of hydrocarbon polymer-like topcoats mutually interacting through weak van der Waals interactions (Table 24.2). As observed byDonnet and Grill (1997), this mechanism appears to be predominant if the H content is higher than≈40 at%, whereas DLC with H content lower than 34 at% behaves in UHV as graphite or diamond. Theexact origin of the threshold between ultralow friction and high friction observed with a difference limitedto 6 at% is not completely understood. Others have already observed that aliphatic-type hydrocarbonchains with high flexibility undergo easy friction-induced orientation along the sliding direction, asshown by polarized infrared microscopy performed on inside wear tracks of hydrogenated DLC after aUHV friction test (Sugimoto and Miyaki, 1990).

TABLE 24.2 Friction Mechanism of Carbonaceous Compounds at the Molecular Scale

Friction range < 0.02 0.1–0.2 >0.5Nature of the interaction van der Waals Hydrogen σ or πEnergy (eV/bond) 0.08 0.2 0.4–0.8Schematic of the

interaction (environment)

a-C:H a-C:H or a-C Diamond, graphite,(UHV) (RH > 4%) DLC with low H content (UHV)

Note: Energy values are indicated by Gardos (1994), and experimental friction ranges have been compiled byDonnet (1997; 1998b) and Gardos (1994). UHV means ultrahigh vacuum.

The tenfold increase in friction from <0.02 to >0.1 caused by humidity and oxidation indicates thatthe increase in bond strength from 0.08 eV/bond to 0.2 eV/bond (H bonding at CO sites by watermolecules) associated with a progressive friction-induced graphitization should be the main cause of thechange in friction (Table 24.2). However, such a description fails to account for the modification of therheology of the contact by triboreactivity between the transfer layer and the counterface, and by theformation of friction-induced wear particles. This modification of the rheology has been shown by Kimet al. (1991) by performing FTIR microprobe tests to identify the debris and film structures observed onvarious interfacial films formed by a Si3N4 ball rubbed against DLC-coated Si in dry Ar, dry air, andmoist air conditions. Humidity caused a tribochemical reaction and led to a transfer film composed ofoxidized hydrocarbon and hydrated silica (from the ball). Consequently, friction was rather low, but wearwas significantly high. The opposite tribological behavior was observed in dry air, where the transferlayer and debris were composed of a carbonyl compound formed by the tribo-oxidation of hydrocarbonsin the DLC films. The triboreactivity was reduced in dry Ar, thus lowering both friction and wear, witha transfer film and debris whose composition was similar to that of the initial film.

24.3.3 Applications of DLC FilmsThe unique properties of DLC films and their modifications, together with the possibility of adjustingthe properties by choosing the right deposition parameters, make them suitable for various applications.The exploited properties thus far include high wear resistance and low friction coefficients, chemicalinertness, infrared transparency, high electrical resistivity, and, potentially, field emission and low dielec-tric constant. Although ta-C has properties similar to or (in some instances) better than those ofconventional DLC (i.e., a-CH), thus far, only conventional DLC films have been used for practicalapplications (see extensive reviews by Grill and Meyerson, 1994; Grill, 1997a). Because DLC is IR-transparent, it can be used as an antireflective and scratch-resistant, wear-protective coating for IR optics(at wavelengths of 8 to 13 µm) made of Ge, ZnS, and ZnSe (Grill, 1999). The low deposition temperatureof DLC allows its use as a wear-protective layer on products made of plastic; therefore, it is used to protectpolycarbonate sunglass lenses from abrasion.

The most widespread use of DLC films is in wear and corrosion protection of magnetic storage media.Nanosmooth and very thin (even <5 nm) DLC films are now used as corrosion and wear protectivecoatings for both magnetic disks and magnetic heads (Bhushan, 1999). Video recording or magnetic datastorage tapes, in which ferromagnetic metal is a recording medium, and the metallic capstans in contactwith the tapes are also being protected with DLC coatings to reduce wear and friction, thus extendingboth the life of the tapes and their reliability. The announcements of the latest MACH3 razor blades byGillette underscore the use of DLC as a coating to improve the quality and performance of the blades.DLC seems also to have found its uses in tribological coatings for metal bearings, gears, and seals. Itspotential use for phase-shift masks for DUV lithography has also been demonstrated.

DLC films appear to be biocompatible, and applications are being developed for their use in biologicalenvironments. Because they are chemically inert and impermeable to liquids, DLC coatings could protectbiomedical implants against corrosion and serve as diffusion barriers. DLC is being considered for coatingmetallic and polymeric-substrate biocomponents, such as polyurethane, polycarbonate, and polyethylene,to improve their compatibility with body tissues. DLC deposited on stainless steel and Ti alloys, whichare components of artificial heart valves, has been found capable of satisfying both the mechanical andbiological requirements and improving the performance of these components. The same properties maymake DLC useful as a protective coating for hip joint implants. Improvement of carbon/carbon compositeprosthetics by DLC coatings has also been demonstrated. Currently, DLC and its modifications are beingconsidered as low-dielectric materials for the interconnect structures of ultra-large-scale integrated cir-cuits (ULSI). A better understanding of the means to control the thermal stability and other integrationproblems of DLC and its modifications will potentially expand their use in ULSI chips.

Nonhydrogenated ta-C films have yet to find such widespread application. The most promisingapplication appears to be for cathodes in field-emission-based flat-panel displays or as pixel elements inlarge outdoor displays.

24.4 Other Related Films

24.4.1 Cubic Boron Nitride (CBN)

Like diamond, CBN exhibits many exceptional properties that render it far superior to other hard nitridesand carbides used by industry today. It is the second hardest material known. Unlike diamond, it doesnot dissolve in or interact with Fe-based alloys, and thus can be used to machine ferrous materials (whichaccount for 80% of all cutting and machining operations). Because CBN has a very high melting point,is chemically very stable, and exhibits excellent hot hardness and strength, it can provide excellentresistance to wear and oxidation at elevated temperatures. CBN can also work very well under drymachining conditions. In addition to its excellent thermal and chemical stability and high wear resistance,CBN offers high thermal conductivity and a very large band gap (≈6 eV). It can be doped with both p-and n-type elements; hence, like diamond, CBN can also be used in various optical and electricalapplications.

Like diamond, CBN can be synthesized by the HPHT method in the powder form and can be depositedin thin films on various metallic and ceramic substrates by both PVD and CVD methods. The most widelyused deposition methods include magnetron sputtering, plasma-enhanced CVD, ion plating, pulse laserdeposition, ion-beam-assisted deposition, etc. (Murakawa and Watanabe, 1990; Chiang et al., 1997; Smeetset al., 1997; Konyashin et al., 1997). The deposition temperatures needed for the formation of the CBNphase are substantially lower than those required for diamond films. Also, the surface finish of CBN filmsis much smoother; hence, there is no need for post-deposition polishing (Watanabe et al., 1999).

Several factors can affect the synthesis of high-quality CBN films during deposition. First, achievingand maintaining a 1:1 stoichiometry between B and N atoms tends to be rather difficult. Most films areslightly N-deficient and, hence, are not perfectly crystalline. Increases in B above stoichiometry andconcomitant formation of vacancies on the N sublattice lead to distortions and significant changes incrystal structure. These changes can adversely affect the mechanical, thermal, and other importantproperties of CBN. Depending on deposition temperature, ion energy, and current density, resultantfilms may contain appreciable amounts of amorphous or hexagonal BN phases. For example, the for-mation of an sp3-hybridized CBN phase is largely controlled by the substrate temperature and the energyof N and Ar ions in sputtering and ion-beam deposition. The presence or absence of H in the plasmacan also affect film quality. For higher quality CBN films, moderate ion energies are needed (Schaffnitet al., 1996).

Growing CBN films tend to accumulate increasingly high compressive stresses in their microstructuresand this represents a major problem in attaining films that are thick enough (at least 3 to 5 µm) formachining operations. The lack of suitable tool substrates with a good lattice and thermal expansion matchis a major part of this problem. During deposition, compressive stress within the film tends to increasewith increasing thickness and eventually becomes so high that CBN film delaminates prematurely orfractures severely (Murakawa and Watanabe, 1990). Currently, safe film thickness for crystalline CBN filmsis in the range of 0.2 to 0.5 µm. For most machining and other tribological applications, films thicker thanthis are needed. Deposition of such thick films is quite possible with diamond, but not with CBN.

In recent years, approaches have been tried to achieve strong bonding between CBN and varioussubstrate materials. In one approach, CBN was produced over a graded Ti interlayer. In another approach,annealing of the CBN films both during and after deposition was tried. Annealing is thought to relievesome of the residual stresses, and hence reduce the amount of stress buildup at the coating-substrateinterface. The use of a Ti layer or first layer, followed by a second layer of graded BN, was shown to bequite effective in achieving strong adhesion between CBN and several tool materials (i.e., WC-Co alloyand high-speed steel). Furthermore, it was shown that CBN can be directly grown on diamond filmswithout using a bond or a graded BN layer (Murakawa and Watanabe, 1990; Murakawa et al., 1991);hence, a duplex coating of the two hardest materials can be obtained. Despite some incremental improve-ments in achieving strong adhesion between CBN and various substrate materials, high residual stressesin these films still hinder the deposition of very thick films with good adhesion.

24.4.1.1 Tribology of CBN Films

Because of its excellent mechanical properties, chemical inertness, and high-temperature durability, CBNholds significant promise for various tribological applications, including wear parts, cutting tools, dies,seals, bearings, etc. Accordingly, it has been the subject of several tribological studies in recent years.Specifically, the friction, wear, and usefulness of CBN for extreme machining operations have beenexplored. Tribological studies have indicated that the friction and wear performance of CBN is highlydependent on the test environment and counterface materials.

In a search of self-lubricating materials for space applications, Miyoshi (1999) performed comprehen-sive tests with diamond, DLC, and CBN in UHV environments. Surprisingly, it was found that CVDdiamond films against CBN exhibited the lowest coefficients of friction. Such impressive results ledMiyoshi to the conclusion that a combination of these two hardest materials may serve as an effectiveself-lubricating, wear-resistant couple for UHV or space applications. Tribological tests by Watanabe et al.(1999) have revealed that CBN exhibited much better wear performance than amorphous BN and TiNor TiC films. When CBN is slid against steel in air or under water, wear is low; but in high vacuum, itis high. Sliding CBN against Al led to high friction and wear in air and vacuum; but in water, wear wasalmost zero.

Studies by Miyake et al. (1992) and Watanabe et al. (1991) demonstrated that the crystallinity of CBNcoatings strongly affects the friction and wear properties of ion-plated CBN films. In their tests, highlycrystalline CBN films provided the best wear resistance and lubrication, whereas films with an amorphousstructure exhibited good frictional properties but very short wear lives. Poor adhesion of hexagonal BNled to premature debonding and hence poor wear performance. In the end, these authors concluded thatperhaps an ideal film should contain a mixture of CBN (for high wear resistance) and amorphous BN(for good lubricity) phases. To test this idea, they prepared a duplex film that consisted of both CBNand amorphous BN. More specifically, an initial layer of CBN was deposited onto an Si substrate by ionplating, followed by a second layer of amorphous BN on top of the CBN layer with a gradient interlayerformed between the two BN layers. The results of tribological tests showed that this composite filmexhibited significantly lower friction coefficients and much longer wear life than the single-layer CBN film.

Compared to other hard nitride and carbide coatings, CBN provides very low friction coefficientsduring sliding in air. Comprehensive studies by Murakawa (1997) and Watanabe et al. (1991, 1995) havedemonstrated that the friction coefficients of high-quality CBN films are in the range of 0.1 to 0.2,whereas those of TiC, TiN, and TiCN films are 0.3 to 0.5 under the same test conditions. Because of theirexcellent tribological properties, CBN films were deposited on WC-Co drills and tool inserts and subjectedto cutting operations. Except for partial delamination in a few cases, CBN coatings reduced torque whiledrilling 440C steel work pieces.

Large-scale applications of CBN films have not yet been realized. Despite its many attractive properties(i.e., high thermal stability, low reactivity with ferrous alloys, smooth surface finish, excellent resistanceto corrosion and high-temperature oxidation, and low-to-moderate deposition temperatures), CBNsuffers from poor adhesion, limited film thickness, high internal stresses, and poor reproducibility ofstructure and properties. Currently, researchers are concentrating on overcoming these problems. Exceptfor a few case studies where superior machining and lower torque were demonstrated while cutting ordrilling steels, CBN films have not yet been offered for large-scale applications.

24.4.2 Carbon Nitride Films

First-principles calculations by Cohen (1985) and Liu and Cohen (1989, 1990) have predicted a newform of carbon nitride (i.e., cubic beta carbonitride β-C3N4) that exhibits an extremely low compress-ibility and superhigh hardness that exceed those of natural diamond. The unique crystal structure ofcubic β-C3N4 is based on the β-Si3N4 structure, in which C is substituted for Si to achieve an atomicbond configuration that favors extreme hardness and elastic modulus. Because hardness is inverselyproportional to the bond length and hence to the radii of the atoms that form that bond, Cohen predictedthat such compounds based on C and N, which have very small atomic radii and the strongest bonds,

will exhibit the highest hardness. Since these initial theoretical studies, many researchers have attemptedto synthesize β-C3N4 in both the bulk and thin-film forms by various methods, with the expectation thatthis will lead to new possibilities in the fields of superhard and wear-resistant materials. In addition toits extreme mechanical properties, researchers have predicted that β-C3N4 will provide excellent thermalconductivity and wide bandgap (Cohen, 1995; Wang, 1997).

During the past decade, several new methods have been developed and used to synthesize crystallineβ-C3N4 films. Despite intense research efforts around the world, successful synthesis of crystalline β-C3N4

has not yet been fully realized. Most PVD and CVD methods (such as laser ablation, RF and DCmagnetron sputtering, ion-beam deposition, ion implantation, plasma arc deposition, UV-assisted chem-ical synthesis, and hot-filament CVD) have produced amorphous carbon nitride films with a relativelylow N content (≈20 to 30%), mainly because of an extremely high bond dissociation energy for N in gasdischarge plasmas or ion fluxes. Occasionally, within the predominantly amorphous structure of thesefilms, some researchers have reported the presence of small crystallites with electron diffraction patternsthat match the pattern of β-C3N4 (Chen et al., 1997; Yu et al., 1997; Yu et al., 1994).

Frustrated by slow progress in producing β-C3N4 in the bulk or thin-film forms, a few researchershave tried to stabilize the crystalline phase in a pseudomorphic state using appropriate structural tem-plates such as TiN or ZrN(111), which can provide the correct unit cell geometry and lattice match.Specifically, using crystalline ZrN(111) as the structural template, Li et al. (1995) and Wu et al. (1997,1998) have recently produced a multilayer film that consists of very thin β-C3N4 (1 to 2 nm thick) andZrN films with a hardness value above 40 GPa and an elastic modulus of 400 GPa.

Most research groups have been able to successfully produce amorphous C nitride coatings. Althoughthese coatings are much softer than diamond, they still provide excellent tribological performance,especially when used as a protective overcoat for computer hard drive systems (Cutiongco et al., 1996;Khurshudov and Kato, 1996).

Primarily because of a large variation in their microstructure and chemical stoichiometry (i.e., C:Nratio), amorphous C nitride films produced by various research groups exhibit large variations in mechan-ical properties and tribological performance. Reported friction values range from 0.05 to >0.5 in air. Indry environments or high vacuum, much higher friction coefficients were observed. Such a large scatterin friction has been mainly attributed to the differences in N content, extent of sp3 and/or sp2 hybridiza-tion, deposition methods and conditions, and, ultimately, the mechanical properties of the films. Kusanoet al. (1998) have found that CNx films deposited with >70% N in the sputtering gas exhibit higherfriction and wear coefficients than films grown in the presence of lower N content. Czyzniewski et al.(1998), Hajek et al. (1997), and Fendrych et al. (1998) have also found a very close relationship betweenfriction and wear properties of CNx films, C:N ratio, and deposition conditions (i.e., temperature, ionenergy, ion current density, etc.). Li et al. (1994) have observed that in dry sliding contacts, CNx givesan initial friction coefficient of 0.1 against a 52100 steel ball, but this value increases to ≈0.5 in steadystates. Zou et al. (1999) reported steady-state friction coefficients of 0.08 to 0.14 for CNx films producedby a reactively ionized cluster beam technique.

As far as tribological applications are concerned, amorphous C nitride films are currently used asprotective coatings in magnetic hard disks. Khurshudov and Kato (1996) reported that the wear rates ofthe carbon nitride-coated disks were 10 times lower than those of disks coated with a commercial DLCfilm of the same thickness. These C nitride coatings also provided ≈3 to 30-times longer wear life thancommercial C coatings. Carbon nitride coatings can also be used in other tribological fields whereconventional DLC films have already been exploited, but the progress in implementation has been ratherslow.

24.5 Summary and Future Direction

Over the last 2 decades, great strides have been made in the deposition, characterization, and tribologicalutilization of diamond, DLC, and related films, which represent some of the hardest materials known.In addition to their exceptional mechanical properties, these coatings incorporate several other attractive

properties that can be very useful for some demanding tribological applications. The current state-of-the-art in diamond, DLC, and other related films allows many tribocomponents to be coated with thesefilms and offered for practical applications. Although some applications are still in the exploratory stageand require further refinement, prototypes of others have been successfully produced and are currentlybeing evaluated for endurance and reliability.

As a true reflection of the growing interest in these films, the number of publications that deal withthem has steadily increased over the years. It is not possible to refer to all of them in this chapter, buteven a simple literature survey will yield several key publications for those who are interested in morein-depth information about these films. In addition to some archival journals, several technical/scientificbooks and book chapters, as well as conference proceedings, are now available. These publications haveled to a better understanding of the microstructure and chemistry of these films and, hence, have led totheir greater utilization by industry.

The relatively high cost of diamond and poor adhesion and reliability of CBN are delaying their widerapplications in industry. Until now, only tool inserts and drills have been coated successfully with diamondand CBN and been made available for limited commercial use. However, these superhard coatings havemuch to offer future tribological applications. It is anticipated that their use in the cutting-tool industrywill further increase if and when their reliability is improved and their unit cost is further reduced.Although prototypes of other tribological parts (e.g., mechanical seals) with diamond coatings have alsobeen prepared, their large-scale utilization has not yet been realized.

Some of the reasons for the slow progress in the commercialization of diamond and CBN coatingsare rough surface finish (in the case of diamond), problems with adhesion (CBN), high fabrication cost(both diamond and CBN), and most important, insufficient reliability and reproducibility in actualapplication. DLC and C nitride films are more cost-effective and much easier to produce than diamondand CBN; hence, they are now used in various tribological applications.

The combination of high wear resistance (due to high mechanical hardness) and low friction makesdiamond, DLC, and related films very unique and ideal for demanding tribological applications. However,the results of previous studies suggest that, depending on the tribological and environmental constraints,tribochemical and thermomechanical interactions can occur at the sliding interfaces of these films andcontrol their friction and wear performance. Existing data suggest that, with proper control of theirmicrostructure, chemistry, and surface topography, these films may live up to their promise. The highsurface roughness of microcrystalline diamond films presents serious problems in many tribologicalapplications. These microcrystalline films can be polished by various methods or, alternatively, nano-crystalline diamond films with a very smooth surface finish can be used to achieve low friction and wearin sliding tribological applications. These smooth films are particularly suitable for mechanical shaft sealsand MEMS applications. As with other types of hard coatings, strong bonding between these films andtheir substrates is an important prerequisite for long service life in most tribological applications.

Field and laboratory test results confirm the notion that diamond coatings can afford excellent lifeimprovement to cutting tool inserts when they are used to machine Al-Si alloys and graphite, whereas CBNis more suitable for machining ferrous and nonferrous alloys. When compared to diamond, CBN has notyet found large-scale applications in the industrial world. Poor adhesion, insufficient thickness, highinternal stresses, and difficulties in attaining and maintaining 1:1 stoichiometry between B and N atomsin a cubic crystalline structure during deposition are some of the major problems that hinder further useof CBN in tribological fields. With recent advances in deposition technology, the quality and reliability ofthese coatings are expected to improve in the near future and, hence, increase their chances for large-scaleapplications. Briefly, the future of diamond, DLC, and related coatings in tribological applications lookspromising because the industrial need for materials with unusual properties is constantly increasing.

Acknowledgment

This work is supported by the U.S. Department of Energy, Office of Energy Research, under ContractW-31-109-Eng-38.

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