24
Tribology of Diamond,
Diamond-Like Carbon,
and Related Films*
24.1
24.2
Introduction
Diamond Films
Microcrystalline Diamond Films • Nanocrystalline
Diamond Films • Tribology of Diamond Coatings •
Tribological Applications
24.3
Argonne National Laboratory
Christophe Donnet
École Centrale de Lyon
Diamond-like Carbon (DLC) Films
Background on DLC Films • Tribology of DLC
Films • Applications of DLC Films
Ali Erdemir
24.4
Other Related Films
Cubic 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 boron
nitride [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 sliding
bearings, machining, mechanical seals, biomedical implants, microelectromechanical systems (MEMS),
etc. The dry sliding friction and wear coefficients of these materials are among the lowest recorded to
date (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 choice
for a wide range of applications. Besides their exceptional mechanical and tribological properties, most
of these superhard materials offer broad optical transparency, high refractive index, wide bandgap, low
or negative electron affinity, transparency to light from deep UV through visible to far infrared, excellent
thermal conductivity, and extremely low thermal expansion. Briefly, these exceptional qualities make
diamond, DLC, and other related materials ideal for numerous industrial applications in addition to
tribology.
* The submitted manuscript has been created by the University of Chicago as Operator of Argonne National
Laboratory (“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 license
in said article to reproduce, prepare derivative works, distribute copies to the public, and perform publicly and
display 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.
© 2001 by CRC Press LLC
The purpose of this chapter is to provide an overview of the tribology of diamond and related coatings
that have attracted overwhelming interest in recent years. Emphasis is placed on the current state-of-theart of our understanding of the friction and wear mechanisms of these films as well as on their uses for
demanding tribological applications. Referring to the structural and fundamental tribological knowledge
gained during the past decade, this chapter emphasizes the importance of surface physical and chemical
effects on friction and wear of these materials and describes new deposition procedures that can lead to the
production 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 recent
tribological studies, one observes that the surface roughness of these materials plays a major role in their
friction and wear performance. Smooth diamond films obtained by polishing and by controlling grain size
hold 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 third
is 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 industrial
applications. Except for a few cases (e.g., diamond-studded rotary drill bits, dressers, diamond-tipped
glass cutters, and fine powders as superabrasives in grinding wheels), natural diamond is rarely used for
tribological purposes. However, synthetic polycrystalline diamonds (PCDs) have been available for more
than 30 years and they are widely used in many industrial applications. PCDs are produced in large
quantities by a high-pressure/high-temperature (HPHT) method in which diamond is crystallized from
a 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 polishing
media. The phase diagram in Figure 24.1 indicates the stability ranges of diamond, graphite, and other
FIGURE 24.1
indicated.
Carbon diagram in which the stability ranges of diamond, graphite, and other forms of carbon are
© 2001 by CRC Press LLC
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 metalcutting or machining operations.
24.2.1 Microcrystalline Diamond Films
The prospects for inexpensive and large-scale production of diamond increased tremendously in the
1970s, 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). Early
attempts 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 diamond
from hydrocarbon gas mixtures by Deryagin and Fedoseev (1975) and Fedoseev (1994) can be regarded
as the earliest attempts to produce synthetic diamond by CVD. In fact, Eversole of Union Carbide was
able to produce diamond phases by a thermal pyrolysis method in 1962 (Eversole, 1962). Progress slowed
thereafter until Angus of Case Western University found a more efficient method in which graphite is
etched simultaneously with atomic H in the deposition system (Angus, 1994). His work motivated
Deryagin and co-workers to focus on the uses of H/hydrocarbon mixtures for low-pressure diamond
growth (Deryagin and Fedoseev, 1975). Since these studies, there has been an explosion in research on
diamond and related materials, with the expectation that low-pressure methods will allow faster, less
expensive, and easier production of synthetic diamonds.
During the past 2 decades, great strides have been made in both the production and industrial
applications of diamond films. At present, several techniques can be used to produce high-quality
diamond films with micro- and nanocrystalline structures at fairly high deposition rates over reasonably
large 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; Butler
and Windischmann, 1998). These methods have greatly reduced the unit production cost of diamond
and 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-quality
diamond coatings produced by CVD methods exhibit most of the desired mechanical and tribological
properties 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) gas
mixtures. Typical deposition pressures in these syntheses are 102 to 103 Pa, and the gas mixture primarily
consists of H and very little CH4. In the case of microwave CVD, 0.5 to 5% CH4 is mixed with H gas and
introduced into the deposition reactor. Small amounts of O and N can also be blended. In the case of
microwave CVD, a 2.45-GHz power source is used to initiate and maintain a gas discharge plasma in
which the substrate material is immersed. At very high plasma temperatures, electrons intensely interact
with the H-rich gas mixture, and the ionic species that are created during this interaction lead to the
formation of diamond nuclei on the surface of a suitable substrate (i.e., Si, SiC, W, Mo). Plasma conditions
in deposition reactors are tailored to control the film microstructure and chemistry, as well as growth
rates and orientation. Typical substrate temperatures for high-quality diamond deposition can range
from 700 to 1000°C. Deposition of diamond at temperatures below 700°C has also been demonstrated
but growth rates are reduced significantly and the amount of non-diamond precursors in these coatings
is increased (Hatta and Hiraki, 1998).
Hydrogen is extremely important for the synthesis of high-quality diamond coatings in most CVD
processes. It plays a critical role in the stabilization of the sp3 bond character of the growing diamond
surfaces. It also controls the size of the initial nuclei, dissolution of C and generation of condensable C
radicals in the gas phase, abstraction of H from hydrocarbons attached to the surfaces, and production
of vacant surface sites where sp3-bonded C precursors can be inserted. It etches most of the double or
sp2-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
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continuous renucleation. Only those grains that attain a certain critical ratio of bulk-to-surface atoms
survive. Consequently, these diamond coatings consist primarily of large grains with highly faceted surface
finishes 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, independent
nuclei form. As these nuclei grow larger, they close the gaps between them and merge to form a continuous
film. Thereafter, the growth process and growth rates are dominated by nearby neighbors and growth
orientation. Grains with more favorable growth orientation grow most quickly and overshadow the grains
with less favorable growth orientation. Grains with less favorable growth orientation will be buried
between the large grains. Figure 24.2 shows initial, intermediate, and final growth stages of a diamond
film on an Si substrate.
High-quality microcrystalline diamond coatings are rough and often nonuniform in thickness or
slightly bowed because of internal stresses and differences in deposition rate from the edges to the center
of the samples. The generally rough surface finish of these coatings precludes their immediate use for
most tribological applications. As discussed in detail later, when used in sliding-wear applications, such
rough coatings cause high friction and very high wear losses on mating surfaces. The rough diamond
coatings can be polished by laser beams, mechanical lapping with fine diamond powders, ion-beam or
plasma etching, and thermomechanical polishing with hot Fe or Ni plates (Zaitsev et al., 1998; Ramesham
and 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 of
large areas. Depending on the polishing method, one can achieve mirror finish surfaces with an rms
surface roughness of 10 nm or less. As demonstrated by numerous investigators, the polished diamond
coatings 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 are
tedious, time-consuming, and, in the case of complex geometries, highly impractical. Depending on the
desired film thickness or the degree of original roughness, it may be necessary to remove large amounts
of material before a smooth surface is obtained. Figure 24.3 shows the surface morphology of a rough
microcrystalline 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 for
most tribological applications. Recently, new methods have been developed for the deposition of finegrained 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 methane
fraction in the source gas, substrate temperature, and substrate pretreatment were also shown to exert a
strong effect on the crystallinity and grain size of these films (Gilbert et al., 1998; Erz et al., 1993; Hogmark
et al., 1996). Although growth rates are somewhat reduced and the amount of non-diamond phase is
somewhat 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. The
source 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 a
modified version of a conventional microwave CVD reactor, details of which are schematically depicted
in Figure 24.4. To introduce C60 into the reactor, a quartz transpirator, also shown in Figure 24.4 is
attached. Fullerene-rich soots that contain ≈10% C60 are placed in the transpirator. The soot is heated
to 200°C under vacuum for 2 h to remove residual gases and hydrocarbons. The tube furnace and
transport tube are heated to between 550 and 600°C to sublime C60 into the gas phase. Argon gas is
passed through the transpirator to carry the C60 vapor into the plasma. To ensure that C60 is transported
into 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 deposited
on the wafer in 1 h. The film displayed only strong C60 infrared absorption features. The measured Raman
© 2001 by CRC Press LLC
FIGURE 24.2 (a) Initial, (b) intermediate, and (c) final growth stages of a microcrystalline diamond film produced
in a microwave CVD reactor.
spectrum was also attributable to C60. Based on these measurements, the transpirator is considered an
effective source of C60 for diamond growth. The NCD coatings can also be grown under similar low-Hcontent 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. Film
growth is monitored in situ by laser reflectance interferometry to determine growth rate and to stop
growth at the desired film thickness. The substrate temperature can vary between 700 and 950°C, and
© 2001 by CRC Press LLC
FIGURE 24.3
Surface morphology of a rough diamond film (a) before and (b) after laser polishing.
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 24.5
shows atomic force microscopy (AFM) and scanning electron microscopy (SEM) images of the surface
of 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 into
the diamond surface has been proposed for the growth of these coatings (Gruen, 1998). The resultant
coatings are phase-pure NCD, with an average grain size of ≈15 nm. The C dimer growth mechanism is
unique 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 lowest
friction coefficients to sliding tribological interfaces when tested in open air. This combination of extreme
hardness (and hence wear resistance) and ultralow friction in one material is very rare in the field of
tribology and it renders diamond ideal for a wide range of tribological applications. However, recent
fundamental studies have confirmed that there are several factors that can adversely affect the friction
and wear performance of bulk diamond and thin diamond films (Bowden and Tabor, 1950; Bowden and
Young, 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
© 2001 by CRC Press LLC
FIGURE 24.4 Schematic illustration of a modified microwave CVD reactor used for deposition of nanocrystalline
diamond 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, Thin
Solid Films, 270, 154-159. With permission.)
of bulk diamond sliding against itself in open air is 0.02 to 0.05. However, when tested in ultrahigh
vacuum (UHV) or at high temperatures, the friction coefficient of diamond increases by more than
1 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 diamond
are contact pressure, surface roughness, crystallographic orientation, and the presence or formation of
gaseous, 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 fundamental
tribological studies at atomic-to-nano scales have greatly increased our understanding of the influence
of each of the above-mentioned factors on friction and wear of diamond. In particular, with the advent
of new experimental tools (e.g., AFM) and computer simulation methods (e.g., molecular dynamic
simulation), it is now possible to visualize and better understand the extent of chemical, physical, and
mechanical events that occur at atomic or molecular levels and hence gain a better understanding of the
mechanisms that govern the friction and wear behavior of diamonds.
24.2.3.1 Early Studies
Early tribological studies have led to the conclusion that the low-friction property of diamond is largely
due 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 was
associated with the lack of adhesive forces between sliding diamond surfaces. Gaseous adsorbates in the
surrounding atmosphere, such as H, O, or water molecules, can attach to and effectively passivate the
dangling σ-bonds of diamond surfaces, and hence lead to reduced adhesion and/or friction. This hypothesis 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 and
Bhushan (1992) has confirmed that the presence or absence of surface contaminants in the test chamber
© 2001 by CRC Press LLC
FIGURE 24.5 (a) AFM and (b) SEM images of the surface morphology of nanocrystalline diamond film produced
by microwave CVD in Ar-C60 plasma.
makes a huge difference in the friction and wear performance of bulk diamond and/or thin diamond
films. When tested in ultraclean and ultradry test environments (i.e., UHV) or at high ambient temperatures, sliding diamond surfaces always exhibit high friction and wear, mainly because surface contaminants are desorbed or mechanically removed from the sliding surfaces and hence are not available to
passify 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 bond
configuration. However, C atoms on diamond surfaces can only establish three covalent bonds with their
near neighbors, while the fourth bond is left open and dangling out of the surface. Gaseous species in
the surrounding air such as water molecules, O, and H, can chemisorb and effectively passivate these
dangling 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 between
sliding diamond surfaces is drastically reduced. The extent of interaction between passive diamond
surfaces 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 effective
means to passivate them with, for example, F atoms, and have thus achieved extremely low friction and
wear 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 the
sliding surfaces of diamond by thermal desorption at high temperatures, the friction coefficient increases
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Dangling bond orbital
Free electron
Shear
plane
C
C
C
-C-C-C-C-C-C-C-C-C-CI I I I
I I I I I I
H H H H H H H H H H
H H H H H H H H H H
I I I I
I I I I I I
-C-C-C-C-C-C-C-C-C-C-
C
C
C
C
C
C
0.8
µ
0.6
UHV
air, 55% RH
1.0
air, 52% RH
FIGURE 24.6 Schematic illustration of dangling surface bonds of diamond and weak shear plane between hydrogenterminated diamond surfaces. (From Holmberg, K. (1998), Tribology of diamond and diamond coatings — A review,
Tribologia, 12, 33-62. With permission.)
0.4
0.2
0
100
200
300
400
500
Cycles
FIGURE 24.7 Friction coefficient vs. number of sliding passes for a CVD-coated SiC pin sliding against a CVD
diamond-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 Surface
Science 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.)
rapidly; presumably, the dangling surface bonds are reactivated and are available to form strong adhesive
bonds 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 coefficient
drops precipitously, presumably because of the repassivation of the dangling surface bonds, as shown in
Figure 24.7 (Dugger et al., 1992).
Most of the vacuum tribological studies on diamond were performed by Gardos and co-workers. Using
an SEM tribometry test device, they have explored the effects of surface chemistry-induced changes on
friction and wear of self-mated sliding diamond surfaces as a function of ambient temperature. As
summarized in a recent review article by Gardos (1999), their studies provide circumstantial evidence
for the effect of surface dangling bonds on friction (in the absence of shear-induced phase transformation
and/or graphitization). Specifically, their studies show that the lowest friction is achieved under conditions
where complete passivation of dangling bonds of diamond surfaces is feasible (i.e., by atomic or molecular
adsorbates such as H2O, H, or O). However, reconstructed diamond surfaces provide relatively reduced
friction and, finally, the incipient linking of the sliding diamond surfaces by totally unsaturated dangling
bonds (created by thermal desorption) causes the highest friction. The repassivation of dangling bonds
by atomic or molecular adsorbates during cooling or exposure to gaseous species causes reduced adhesion
and hence low friction. Figure 24.8 illustrates the frictional behavior of diamond films in the passive,
© 2001 by CRC Press LLC
FIGURE 24.8 Variation of friction coefficient of CVD diamond films sliding against themselves in vacuum at
temperatures to 850°C. (From Gardos, M.N. (1999), Tribological fundamentals of polycrystalline diamond films,
Surf. Coat. Technol., 113, 183-200. With permission.)
active, and reconstructed regimes of sliding. Apparently, as the temperature increases, surface adsorbates
begin to desorb and leave some of the dangling surface bonds exposed and available for strong adhesive
interactions. As the extent of adhesive interactions increases, the friction coefficient increases sharply and
reaches 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 normal
sliding 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 high
temperatures 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 temperature
reaches 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 tribotested at elevated temperatures, bulk diamond and thin diamond films tend to gradually oxidize and
transform 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 not
undergo oxidation or graphitization. Briefly, diamond films are not suitable for use at high temperatures
in open air.
In addition to the effects of surface chemistry and/or environment, which largely control the adhesion
component of friction, the effects of crystal orientation and surface roughness on friction and wear of
diamond have also been studied rather extensively in previous years. It has been shown that surface
roughness profoundly influences the friction and wear performance of thin diamond films. As mentioned
above, the surface rms roughness of synthetic diamond films can vary from 10 to >1000 nm, depending
on deposition conditions, growth orientation, grain size, and film thickness. As demonstrated by Casey
and Wilks (1973), and Samuels and Wilks (1988), the extent of interaction between sliding-surface
asperities of diamond may contribute to the amount of total energy being dissipated in the form(s) of
phonons 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
© 2001 by CRC Press LLC
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 of
polished 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 produced
by 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 of
smooth 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 diamondlike
carbon films, in Protective Coatings and Thin Films: Synthesis, Characterization and Applications, Pauleau, Y. and
Barna, B. (Eds.), NATO ASI-High Technology Series, Vol. 21, Kluwer Academic, London, 169-184. 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 coefficients
are attained (Erdemir et al., 1997a; Erdemir et al., 1999; Miyoshi et al., 1993) (Figure 24.9). Reported
friction 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, but
decreases substantially as sliding continues in open air. The trend appears to be the opposite during tests
in UHV environments where, regardless of the counterface materials, the friction coefficients are initially
low 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 the
measured friction coefficients of bulk diamond and/or diamond films. Usually, diamond-against-diamond or diamond films give lower friction than diamond-against-nondiamond counterfaces (Hayward
et al., 1992; Miyoshi et al., 1993). When bulk diamond or diamond films are slid against relatively softer
metals, alloys, or ceramic counterfaces, the steady-state friction and wear are largely controlled by the
extent of surface roughness of the sliding diamond side and by the tendency of the counterface material
to form a transfer layer (Hayward et al., 1992; Gangopadhyay et al., 1993; Erdemir et al., 1997b). Rougher
diamond surfaces usually cause greater plowing and abrasive wear on the softer metallic or ceramic
counterfaces and also facilitate the transfer of worn debris particles to its sliding surface. Fine debris
particles can easily fill in the valleys between sharp asperities of diamond films and, eventually, sliding
contact largely occurs between the transferred and original material of the non-diamond counterfaces.
Previous research has demonstrated that growth orientation of microcrystalline diamond films can
significantly affect friction and wear. Films with large grains and highly faceted surface finish usually
cause 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 orientation
is very desirable and can be achieved by controlling deposition conditions and by pretreatment of the
substrate surfaces (Yee et al., 1997; Wolden et al., 1997; Wolter et al., 1996). Diamond films with coarse
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FIGURE 24.10 Friction coefficients of Si3N4 balls sliding against diamond films with ⟨111⟩ and ⟨100⟩ type growth
orientations. (From Erdemir, A., Bindal, C., Fenske, G.R., and Wilbur, P. (1996a), Tribological properties of hard
carbon films on zirconia ceramics, Tribol. Trans., 39, 735-741. 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 of
both 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 micropyramidal
asperities 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 to
the abrasive cutting and ploughing effects of sharp surface asperities on the softer counterface pins. If a
favorable ⟨100⟩ growth orientation is present, such coatings can also afford low friction coefficients to
sliding surfaces, despite relatively higher measured surface roughness. In general, previous studies confirmed that regardless of the grain size, diamond coatings with a smooth surface finish provide very low
friction to sliding counterfaces.
Apart from physical roughness and chemical passivation effects, phase transformation or structural
changes can also play a major role in the friction and wear performance of diamond coatings. The extent
of 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 natural
diamond (see Gogotsi et al., 1998) when extreme contact pressures and/or high frictional heating are
present at local asperity levels. Real contact occurs first between these asperities, and their tips can either
fracture 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 also
known 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 interface
and then begin to dominate the long-term sliding friction and wear performance of these coatings.
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Previous research by Erdemir et al. (1997b) and Jahanmir et al. (1989) revealed that most of the debris
particles derived from sliding diamond surfaces exhibited a graphitic microstructure. Raman spectroscopy, electron diffraction, EELS, and transmission electron microscopy have concurrently confirmed the
presence 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., friction
force 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; Miyake
et al., 1995). These studies have greatly increased our fundamental knowledge base of the tribology of
diamond.
Superhigh-speed computers with high memory and processing capability have allowed simulation and
real-time visualization of atomic-scale phenomena that occur at sliding diamond surfaces. For the most
part, results from fundamental nanoscale experimental studies are consistent with those from earlier
micro- to mesoscale studies, in which standard tribology test machines were used. The nanoscale studies
have revealed that surface adsorbates, ambient temperature, and gaseous, liquid, or solid debris particles
that are present or generated at the sliding diamond interfaces play critical roles in the friction and wear
of diamond. For example, it was shown that when adsorbed gases are removed from sliding diamond
surfaces, friction increases rapidly, because the dangling surface bonds are free and highly activated to
form 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 provided
significant insight into the extent of physical, chemical, and mechanical interactions that occur at sliding
diamond interfaces on the atomic scale.
24.2.4 Tribological Applications
Over the years, great strides have been made in the deposition, characterization, and industrial utilization
of diamond films (Seal, 1995; Ravi, 1994; Lux and Haubner, 1993; Murakawa, 1997). Our fundamental
understanding of the structure, properties, and performance of these coatings has increased tremendously
in recent years and this understanding has been used to optimize, design, and customize new diamond
coatings that can meet the increasingly stringent needs of advanced tribological applications. While some
uses are still in the exploratory stage and require further development, others have been fully established
and 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 laboratory
and field tests confirm that these coated tools provide much improved performance during metal-cutting
operations by allowing high-speed machining at increased feed rates. Some of the advantages that
diamond provides in these applications include extreme hardness and wear resistance; good fatigue
strength; high chemical inertness; excellent resistance to abrasion, erosion, and corrosion; high thermal
conductivity; low friction; and excellent environmental compatibility.
The current market share of diamond-coated cutting tools is relatively small, and the widespread
utilization of diamond coatings in other tribological fields has not yet been fully realized. Some of the
major reasons for the slow progress in large-scale utilization of diamond coatings are high fabrication
cost, 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 poor
adhesion, high deposition temperatures, and slow growth rates. Diamond-coated tool inserts cannot be
used to machine pure Fe and Fe-base materials and the alloys of groups IVa, Va, VIa, VIIa, and VIIIa of
the Periodic Table, which represent the largest segment of industrial machining. Diamond can chemically
react with and/or dissolve in these materials at the high temperatures generated during machining, and
thus it wears out rather quickly. Also, diamond coatings cannot be produced on the surfaces of tools that
are made of materials such as high-speed steels, which are widely used in the tooling industry. Thus far,
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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 for
commercial machining. Although prototypes of other tribological parts (e.g., mechanical seals; Hollman
et al., 1998) have also been demonstrated, their large-scale utilization has not yet been realized because
of high fabrication costs and insufficient reliability in actual operation. Other potential applications for
diamond coatings include sliding bearings, MEMS, wire-drawing dies, and various wear parts that are
used 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 of
synthetic diamond powders produced by HPHT methods. Specifically, fine diamond powders are first
sintered 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 an
Si wafer from which small bits are laser-cut into a desired cutting-edge shape. Subsequently, the Si
substrate is etched out and the cut diamond tips are recovered. As a last step, the freestanding diamond
tips are bonded or brazed onto the cutting edges of the tool inserts or end mills. The third form is the
direct deposition of diamond onto cutting tools by various CVD methods.
There are some fundamental differences between brazed and coated diamond tool bits. First, diamondcoated tools can be tailor-made to meet the specific machining needs of an application, whereas the
brazed diamond tips come in a strict geometric form and cover only the sections of the sharp tips of
tool inserts. Second, the entire functional surface of a tool insert can be coated with diamond, and it can
be prepared in any style or geometry to provide better efficiency and overcome limitations on depth of
cuts — always a problem with brazed tool bits. Finally, the shape and area of a coated carbide insert can
be prepared in such a way that it will allow better chip handling and breaking capability than the very
small brazeable diamond tips. CVD diamond coatings perform equally well or even better than PCD
diamonds in cutting and turning operations. However, edge chipping or deterioration due to easy cleavage
of highly oriented needle-like diamond grains of CVD diamond films can occasionally occur and limit
the lifetime of the coated tools. Figure 24.11 shows the morphology of the cutting edge of a diamondcoated 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.
© 2001 by CRC Press LLC
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 their
light 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-Si
alloys include pistons, engine blocks and heads, wheels and chassis, and some transmission parts. The
increased use of lightweight materials in the automotive industry is likely to continue and, hence, the
need for diamond-coated tools will increase. Other engineering materials suitable for machining by
diamond-coated tools include alloys of Mg, Cu, Pb, and Mn, as well as various composites (e.g., fiberreinforced 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 operations is strong bonding between diamond coating and substrate insert (i.e., WC-Co). If bonding is not
intimate and strong, diamond coatings can easily be removed from the cutting edges, which thus lose
their cutting performance and effectiveness. In fact, premature delamination of the diamond film in early
trials represented the biggest challenge and most common mode of tool failure. Residual stresses are
inherent in all CVD-diamond coatings, and these stresses can significantly affect the film/substrate
adhesion (Olson and Dawes, 1996; Gunnars and Alahelisten, 1996). Diamond coatings are often deposited
at high temperatures, and a large mismatch in thermal-mechanical properties between the diamond
coating and the underlying substrate leads to large compressive stresses. The presence of a high stress in
the 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 be
considered along with the residual stresses whenever the quality of a diamond coating is evaluated. Two
other problems can occur when diamond films are used in machining operations. One is oxidation; the
other 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 a
binder and controls the toughness of these materials. In addition to strict process control during deposition, the selection and pretreatment of the substrate materials are extremely important for achieving
strong adhesion between the diamond coating and the substrate material. In particular, the selection of
the 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 can
adversely affect the adhesion of diamond coatings to carbide inserts. Specifically, during deposition at
high temperatures, Co can play a catalytic role in the formation of non-diamond phases and complex
Co carbides that are highly undesirable. It also slows the nucleation process and reduces nucleation
density. Interfaces with a high proportion of non-diamond phases and low nucleation density suffer from
poor adhesion. Too much Co can also increase the difference in thermal expansion coefficient between
diamond and substrate; hence, during cooling, very high compressive stresses can build up at the interface
region. When these stresses are combined with the stresses associated with cutting action, premature
adhesive failures can occur at or near the cutting edges. Briefly, it is important to use carbide substrates
with low Co content or to etch the Co out completely from the near surface; otherwise, the coatings will
not 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. Most
of the methods involve selective etching of Co by chemical means. Some of the chemicals that have proved
to be effective in etching Co are HNO3, HCl, and CH3COOH. In one case, tool inserts are dipped in a
nitric acid and water solution (mixed 1:1 by volume) and ultrasonically agitated for 15 min. They are
then rinsed and ultrasonically cleaned in ultrapure water for 5 min. In another case, researchers have
heat treated WC-Co substrates at very high temperatures for an extended period of time in a protective
© 2001 by CRC Press LLC
FIGURE 24.12 Cutting performance of an uncoated and CVD diamond-coated insert while machining a cast
aluminum alloy containing 7% Si (cutting conditions: speed, 150 to 2500 m/min; depth of cut, 0.5 to 3 mm; feed
rate, 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.)
environment. Heat treatment causes grain growth in WC and hence roughens the surface; it also evaporates Co and thus reduces the amount of Co at or near the surface. Increased surface roughness provides
better mechanical interlocking between diamond and WC grains, whereas the absence of Co ensures
higher nucleation density and better adhesive bonding.
Because Co is the binder that holds the tool’s WC grains together, Co removal must be done with
great care. If too much Co is removed from the surface, the tool will become weakened and its integrity
near the surface will be impaired. Formation of stable Co borides and silicides that can endure the high
deposition 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 been
tried on tool inserts to achieve strong bonding. The effectiveness of bond layers is very much dependent
on the exact composition and thickness of the bond layer material. Few other procedures have been
developed 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 coated
with CVD diamond can exhibit tool life improvements that range from 25% to more than a factor of
40. Greater improvements in tool life are reported when cutting Al-Si alloys, green ceramics, and fiberreinforced composites. Figure 24.12 shows the performance of uncoated and CVD-diamond-coated
carbide 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 in
preventing edge wear and loss of sharpness after machining of a high-Si-content Al alloy. CVD diamondcoated tools were proven to be as good as or better than the PCD tool inserts when used during turning
of Al-Si alloys and graphite.
The market share of coated tools is expected to increase steadily because prices are declining while
performance and reliability are improving. Better quality and reliability will certainly result in diamondcoated tools that last longer, cost less, and increase productivity. Several companies are engaged in the
production of CVD diamond-coated tool inserts. DeBeers, Mitsubishi, Nachi-Fujikoshi, Norton, Kennametal, Sandvik, and Sumitomo are the major firms. Kennametal and Sandvik provide diamond-coated
tools commercially. As discussed above, CVD diamond-coated inserts are primarily targeted for machining applications that involve high-Si-Al alloys. The projected growth of Al use in automobiles indicates
© 2001 by CRC Press LLC
FIGURE 24.13 Extent of wear damage on (a) uncoated and (b) nanocrystalline diamond-coated tool inserts after
machining of a hypereutectic aluminum-silicon alloy.
that the growth of the diamond-coated tool insert market will be significant. When the performance and
price of CVD tools compare favorably with those of PCD tools, CVD tools may displace PCD in some
applications. Woodworking tools represent another major market segment for diamond coatings. Tungsten carbide blades already comprise a large part of the woodworking tool market. CVD diamond-coated
tools 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 diamond
films. The mechanical-seals market is rather large, ranging from the very simple seals used in automotive
water 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 resistance
to corrosion, erosion, and abrasion is needed. Because of its low friction, a diamond-coated seal can also
be very useful in applications where a liquid lubricant is not permissible, for example, in pharmaceutical
and food processing plants.
Over 3 million higher-end pumps rely on the tribological performance of SiC seals, and SiC is an ideal
substrate for diamond film growth. Improper use of or leakage in mechanical pump systems can be very
costly and may lead to environmental disasters. Total energy losses due to high friction and wear can be
very high. Because properly applied diamond coatings afford very low friction to the sliding surfaces of
seal 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 in
substantial 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 made
of soft graphite or C composite. When rubbed against a rough and superhard surface, these inserts wear
out rather quickly. Hard SiC inserts can also be used, but they are much more expensive and, when
rubbed against rough diamond, they too suffer major wear losses. Polished or smooth nanocrystalline
diamond 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 be
used for a wide range of applications, including sensors, actuators, high-precision positioning devices,
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microelectronics, etc. MEMS devices such as gear assemblies and micromotors have been largely fabricated by Si micromachining technology; however, Si exhibits very poor mechanical and tribological
properties. When these microdevices are used in very high-speed applications, they suffer from unacceptably high friction and wear losses and thus are unsuitable for applications that involve high speeds
and realistic loads. Researchers have developed alternative fabrication methods that allow production of
MEMS devices from SiO2, Si3N4, or SiC-type materials. However, the tribological performance of these
materials 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 diamond
coatings 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 in
fabricating diamond coatings and/or components with a morphology suitable for MEMS. The NCD
films discussed previously appear to be better suited because of their very small grain size (i.e., 10 to
30 nm). With these films, it will be possible to attain film coverage at very small thicknesses (as thin as
50 nm) and hence preserve the fine structural features of MEMS devices. Conventional CVD films with
thicknesses 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, push
pins, high-precision microdrills, surgical blades, wire-drawing dies, and various wear parts that are used
against 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 wide
range of structures and compositions. The tribological behavior of these films strongly depends on the
deposition technique. Consequently, the first part of this discussion is dedicated to a background on
DLC, highlighting the relationships between the deposition process, the film structure, and some film
properties. The second part summarizes the tribological behavior of DLC films, through a description
of 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 general
term “DLC” describes hydrogenated and hydrogen-free metastable amorphous carbon materials, prepared 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 film
structure 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 for
obtaining a wide optical gap and a high electrical resistivity, removing midgap defect states, stabilizing
the 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 phase
diagram in Figure 24.14. The sp2-bonded amorphous C lies in the lower left corner. Phases with an H
content that is too high, lying at the far right of the diagram, cannot form an interconnected network
and 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), with
a predominance of the sp3 hybridization of up to 85%. The most common type of DLC is the hydrogenated
a-CH film, with only moderate sp3 content and a rather high H content of up to ≈50 at%. A lightly
hydrogenated 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, and
various metals.
© 2001 by CRC Press LLC
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., Robertson, J., Milne, W.I., and Amaratunga, G.A.J. (Eds.), World Scientific Publishing, Singapore, 32-45. With permission.)
Such a wide range of film structures and composition and the diversity of methods used for DLC film
deposition 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 practical
applications have apparently matured. The study of hydrogen-free DLCs has not yet reached this state
and 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 for
effective process control, primarily the tailoring of adhesion properties. Indeed, the adhesion of DLC
films varies widely, depending on the nature of the substrate. To be used as tribological coatings, DLC
films must adhere well to the substrate material, and the adhesive forces must overcome the high internal
stresses that would otherwise cause film delamination (Holmberg and Matthews, 1994). Adhesion can
be affected by the deposition method, in combination with the nature of the substrate. Good adhesion
of DLC films is observed on carbide- and silicide-forming substrates. The adhesion of DLC coatings to
silicide-forming metals can be improved by depositing a 2- to 4-nm-thick interfacial layer of amorphous
Si 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 intercalated
between the substrate and the DLC top layer has also been studied (Voevodin et al., 1997). The deposition
of 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 superimposed 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 a
nanometer and are therefore referred to as nanosmooth. The roughness thus conforms to the underlying
topography. All methods for the deposition of DLC are nonequilibrium processes characterized by the
interaction of energetic species with the surface of the growing film. The growth and properties of DLC
films are controlled by the substrate temperature and the energy of the species impinging on the surface
during the deposition process, the latter exerting dominant control. A wide literature survey on the
deposition procedures has been assembled by Grill and Meyerson (1994). The following discussion is
limited to the techniques that are most commonly used to deposit DLC films, with emphasis on the
relationships between deposition conditions and film structure. The deposition of ta-C films is summarized 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 of
sp3 bonding is maximized for ion-dominated processes with ion energies ≈100 eV. The specific properties
of ta-C films are achieved by the high energies of the impinging particles, with film growth governed by
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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 transition temperature of 250°C, a noticeable decrease in the high sp3 content and density is generally
observed. This transition temperature, which is much lower than the thermal stability temperature of
ta-C films (which can reach relatively high values approaching 600°C), decreases with increasing ion
energy.
Magnetron sputtering is probably the most commonly used industrial method for depositing DLC
films. 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 massselected ion-beam technique, it is possible to obtain ta-C films that contain up to 80% sp3-hybridized C
by varying the acceleration, mass, and energy selection of a C ion beam that is condensing onto the
substrate. ta-C films can also be obtained by a cathodic vacuum process on both the laboratory and
industrial scales. An arc is struck on a graphite cathode, creating an intense, 30%-ionized plasma that
expands through the deposition chamber and condenses onto the substrate. The filtered cathodic vacuum
arc configuration avoids degradation of the growing film by particulates produced in the arc along with
the 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 produced
by UV laser ablation of a graphite target.
Plasma-enhanced CVD, in which any gaseous hydrocarbon species is used as precursor, is probably
the most popular means to deposit a-CH films. The chamber is configured with cathode and anode
plates, with the substrate attached to the cathode to maximize ion bombardment. Parallel plate reactors
are preferred because they allow the deposition of uniform films over large areas. RF power can be
capacitively coupled to retain the ability to deposit the films on a wide variety of substrates, including
insulators. Because the film structure is a strong function of the energy of the impacting species, this
ability depends on several plasma parameters — primarily the bias and the gas pressure. The bias is
indirectly controlled by the RF power and the gas pressure in the RF configuration. The sp3 and H content
are strongly controlled by the average impact energy. More precisely, at the lowest impact energies, the
gaseous precursor is not sufficiently decomposed, and polymer-like C films with a predominance of
CH2 groups are generally obtained. At intermediate impact energies, the H content has declined
sufficiently so the C–C sp3 bonding has reached a maximum, thus leading to the so-called diamond-like
qualities. However, if the impact energies become too high, graphite-like C structures are generally
deposited 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 carbonaceous
network. Free H fraction increases when the impact energy, and hence the dissociation of precursor, is
increased. This also leads to higher C network crosslinking. Consequently, the nature of the precursor
plays a direct role in the fraction of H bonded to C. However, based on recent comparisons between
FTIR 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 remains
ambiguous.
As with ta-C films, temperature during deposition is another parameter that generally affects the DLC
structure. Deposition should be carried out below 250°C to stabilize a significant sp3 content. The increase
of diamond-like qualities of a-CH films has been achieved using a plasma beam source with acetylene
to reduce the amount of incorporated H, and by operating at lower gas pressure to increase plasma
ionization (Weiler et al., 1996). The resultant films are called ta-CH, by analogy to ta-C. The sp3 content
can reach a maximum near 80% and H content in the range of 25 to 30%, with an effect of deposition
temperature 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 regular
films, by introducing into the reactor species or precursors that contain the modifying elements.
© 2001 by CRC Press LLC
TABLE 24.1
Range of DLC Structure, Composition, and Properties
Variable
Hydrogen content, at%
sp3, %
Density, g/cm3
Thermal stability, °C
Optical gap, eV
Electrical resistivity, Ω/cm
Index of refraction
Compressive stress, GPa
Hardness, GPa
Young’s modulus, GPa
a
b
a-Ca
a-CHb
<5
5–90
1.9–3.0
<600
0.4–1.5
20–60
20–65
0.9–2.2
<400
0.8–4.0
102–1016
1.8–2.4
0.5–5
<80
<900
<60
<300
a–C = amorphous carbon.
aCH = amorphous carbon, hydrogenated.
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 H
bonded to C, C hybridizations, nature of bonds, alloying elements) are strongly influenced by the average
impact 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 deposition
processes and conditions. Table 24.1 summarizes the range of various typical properties that depend on
the presence or lack of H in the film. Nonhydrogenated films exhibit generally higher hardness, Young’s
modulus, density, thermal stability, and compressive stress; and lower index of refraction, resistivity, and
bandgap, when compared with hydrogenated films. For each category of film, the property variation is
directly related to the structure, as controlled by the deposition process. The relationships between the
structure and the properties of DLC films have been more extensively studied on a-CH than on ta-C
films. In hydrogenated DLC films, an increase in H content is generally associated with a decrease in
hardness, Young’s modulus, stress, thermal stability, density, and index of refraction, and an increase in
the gap and electrical resistivity. A more accurate discussion remains impossible because most of these
properties also depend on the fraction of H that is bonded to C, which remains difficult to investigate
for the reasons mentioned previously.
The difficulties in generalizing the structure/property relationships are due to a lack of systematic and
unambiguous standardized characterization of DLC films, the wide range of structures and compositions,
and the experimental difficulties encountered when thin metastable amorphous films are accurately
probed. Differing models, such as the graphitic cluster, the random covalent network, or the defective
graphite network reviewed by Grill and Meyerson (1994), constitute interesting guidelines to improve
our 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, and
various metals. The introduction of dopants generally brings about a decrease in compressive stress, due
to fewer interconnections in the random carbonaceous network. Grischke et al. (1995) have shown that
the specific chemical composition of the modified films strongly influences the surface energy, depending
on the nature of the dopant, and may modify various physical properties, making some doped DLCs
suitable for various applications — for example, to improve field emission characteristics. As highlighted
in Section 24.3.2, the introduction of alloying elements into DLC affects many properties, including
tribological behavior.
24.3.2 Tribology of DLC Films
Much effort has been invested in the characterization of the tribological behavior of DLC films, mainly
because of the very low friction coefficients and high wear resistance of these materials. Thus, the tribology
of DLC coatings has been discussed extensively in the literature, and has been summarized in recent
© 2001 by CRC Press LLC
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 by
the 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), kinematic parameters (nature of motion, speed), physical parameters (temperature during friction), and
chemical parameters (nature of the environment). An exhaustive presentation of the tribological behavior
of DLC is thus impossible. A more interesting and useful method of presentation will be used here, in
which a top-down approach to DLC tribology is used to focus on the various phenomena occurring at
three different scales. The macroscopic scale allows the establishment of relationships between the nature
of the films and friction levels and wear rates under various testing conditions. The microscopic scale
focuses 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 reactivity
of the top surface in relationship to the environment that surrounds the contact spot, and the surface
chemistry 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 to
0.6, which probably represents the widest range of friction behavior in any category of coatings. Wear
rates are less systematically investigated and are often recorded in nonstandard units, thus increasing the
difficulty of comparing results. DLC films that are easily scratchable with no wear resistance at all, as
well 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 nature
of the films. Hydrogen-free films generally exhibit lower friction (<0.15) when compared with a-CH
coatings. This finding is consistent with the expected role of H in stabilizing the random covalent network
and preventing its collapse into a graphitic network. In ta-C films, friction easily causes a local shearinduced graphitization at the microscopic scale, as described in Section 24.3.2.2. Consequently, it is not
surprising 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 over
each 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 in
Sections 24.3.2.2 and 24.3.2.3.
Unlike friction, wear resistance performance is not so easy to discuss by simply discriminating between
the two previous categories of DLC films. Extremely low wear rates can be achieved for both ta-C and
a-CH coatings. In the case of H-containing films deposited by PECVD, Grill (1997b) has correlated wear
resistance (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 low
DC bias) are less wear resistant than diamond-like films. Thus, based on the accumulated knowledge
described in the previous sections on deposition/structure/property relationships, this finding shows that
the optimization of the wear resistance of DLC films requires strong attention to the control of the
deposition procedure.
In inert environments, including dry N and vacuum, friction coefficients can reach either ultralow
values (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 be
controlled by the H content of the films. Hydrogen concentrations lower than 34 at% systematically lead
to 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 friction
level is highlighted in the chapter section dedicated to the tribology of DLC at the nanometric scale
(Section 24.3.2.3).
© 2001 by CRC Press LLC
A compilation of the tribology of doped DLC and C alloy coatings has been recently prepared by
Donnet (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 appears
to be significantly reduced (<0.1), where compared with that of conventional undoped DLC in ambient
humid air, with comparably high wear resistance. However, this tribological behavior appears to be
observed when the contact pressure remains below 1 GPa. At higher contact pressures, conventional ta-C
and a-CH films cannot be surpassed. Consequently, ta-CHSi films can be used in applications that require
both low friction (<0.1) and high wear resistance (<10–7 mm3/N-m) under moderate mechanical conditions, for the protection of low-stress aerospace or automotive components, precision ball bearings
and gears, sliding bearings, and magnetic recording media.
The incorporation of Si and F into the DLC structure affects the surface properties. The reduction of
stress, when compared with conventional DLC, is in the same range as that in a-CHSi. However, the
reduction 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 by
deposition conditions to obtain films with comparable wear resistance and friction levels of conventional
a-CH films, but with lower stress and surface energy. Less research has been performed on the tribological
investigation of N-containing DLC films because of their recent discovery when compared with conventional DLC and other doped DLC. A review of the tribological behavior of C nitride films is presented
elsewhere in this chapter (Section 24.4.2).
The range of composition and structures attainable with metal-containing DLC coatings appears to
be enormous. One should keep in mind that the optimization of the material combination and deposition
parameters is a challenging subject for each element or combination of elements. When optimization is
achieved, metal-containing DLC films can exhibit promising tribological properties in terms of steadystate friction levels and wear rates for various applications. Dimigen and Klages (1991) observed friction
coefficients in the range of 0.02 to 0.04 in ambient air when a significant amount of Ta or W was
incorporated 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, a
transfer film buildup that is followed by an interfilm sliding mechanism is the most frequently observed
velocity accommodation mode (Figure 24.15). The kinetics of transfer film formation, together with its
thickness, strongly depend on the nature of the counterfaces and the environmental conditions. The
interfilm sliding mechanism occurs when the DLC film adheres strongly to both contacting surfaces, but
separates into two distinct layers that slide across one another. The resultant shear strength is that of the
two layers sliding against one another. Neither of the two original surfaces is in contact, except if the film
is completely worn. At this scale, the behavior of ta-C and a-CH is rather similar, and DLC films are
generally considered to slide according to this accommodation mode. One should keep in mind that the
initial formation of a transfer film is associated with significant wear of the initial coated part. Consequently, 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
coatings.
Interfilm sliding mechanism consecutive to transfer film formation, frequently observed with DLC
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The nature of the counterfaces can significantly influence the size and nature of the transfer layer (both
of which sometimes differ from the initial composition of the film). The transfer layers may become a
mixture of the original carbon film and elements from the counterface, indicating a strong triboreactivity
with the initial pin counterface. Such a mixture has already been observed by secondary ion mass
spectroscopy (SIMS) analysis of inside wear tracks of DLC-coated steels (Ronkainen et al., 1993). The
analysis revealed a transfer film composed of a mixture of C, Fe, Cr, H, and O. A lower wear rate of the
pin, associated with a modification of the transfer composition, was observed by SIMS analysis at high
loads and high speeds. The authors also investigated, with Al2O3 pins, the coatings and experimental
setup that had been studied earlier with metal counterfaces. The thickness of the transfer increases with
higher loads and sliding velocities; this is associated with a drastic decrease in friction coefficient from
0.13 to 0.02. Alumina was transferred to the surface of the disc wear track. These results were observed
both on H-free and hydrogenated DLC films. Thus, similar mechanisms are generally observed when
ceramic materials are used. Transfer layers are also formed on ceramic counterfaces such as Si3N4, with
composition 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 observed
in 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 water
molecules is reduced. A decreased graphitization rate is also observed at lower temperatures and higher
humidity 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), strongly
affects the transfer film thickness, which appears to influence the friction level. This effect has been
observed in hydrogenated DLC film deposited on an Si substrate and tested against a steel pin counterface
in 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, as
shown by in situ Auger electron spectroscopy (AES) performed inside the transfer film. Higher humidity
values are associated with a thinner transfer film, so that AES detects the substrate elements that underlie
this 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 of
the 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 transfer
film and the counterface can induce the formation of complex structures and compositions on the top
surface of the film. Moreover, the nature and thickness of contaminant or passivation layers will probably
play a role, but this kind of assumption can be checked only by tribological tests under UHV conditions
and sliding in situ on fresh surfaces.
An exception to this conventional transfer film formation mechanism can be found when DLC-coated
steel slides slowly against PTFE (Yang et al., 1991). Without the DLC film, very little PTFE is directly
transferred 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 between
PTFE 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 tribochemistry of DLC coatings.
© 2001 by CRC Press LLC
TABLE 24.2
Friction Mechanism of Carbonaceous Compounds at the Molecular Scale
Friction range
Nature of the interaction
Energy (eV/bond)
Schematic of the
interaction
(environment)
< 0.02
van der Waals
0.08
a-C:H
(UHV)
0.1–0.2
Hydrogen
0.2
a-C:H or a-C
(RH > 4%)
>0.5
σ or π
0.4–0.8
Diamond, graphite,
DLC with low H content (UHV)
Note: Energy values are indicated by Gardos (1994), and experimental friction ranges have been compiled by
Donnet (1997; 1998b) and Gardos (1994). UHV means ultrahigh vacuum.
The role of the H content and sp2/sp3 hybridizations in a-CH films is discussed first. The discussions
are based on Table 24.2. With this approach, it should be possible to explain the role of the high H content
in lowering the friction of DLC in UHV from high values that are in the same range as values related to
diamond or graphite under the same environmental conditions. When transfer film buildup during the
running-in period is taken into account, steady-state sliding with DLC films is accommodated between
the two counterfaces by interfilm sliding of C-covered smooth surfaces. Under UHV conditions, the
contaminant topcoats are removed and sliding occurs between the two carbonaceous fresh surfaces. In
the case of diamond/diamond or graphite/graphite, as with poorly hydrogenated a-CH films, the UHV
friction 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. If
these free bonds do not reconstruct or are not activated by any adsorbates, they interact with high energies
and 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 of
these 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 band
parameters influence the net inter- and intraplanar attraction of a family of imperfect graphitic planes
of particles interacting in a solid lubricant layer. Gardos (1994) indicates that the 0.4 to 0.8 eV binding
energy between the basal planes is enough to render pristine highly oriented graphite, a high friction
and wear material.
The discussion related to a-CH films is more complex in that it considers the wide distribution of
H content and C hybridizations, and determines various degrees of C crosslinking. The ultralow friction
of highly hydrogenated DLC films in UHV is consistent with the predominance of hydrocarbon polymerlike topcoats mutually interacting through weak van der Waals interactions (Table 24.2). As observed by
Donnet 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. The
exact origin of the threshold between ultralow friction and high friction observed with a difference limited
to 6 at% is not completely understood. Others have already observed that aliphatic-type hydrocarbon
chains with high flexibility undergo easy friction-induced orientation along the sliding direction, as
shown by polarized infrared microscopy performed on inside wear tracks of hydrogenated DLC after a
UHV friction test (Sugimoto and Miyaki, 1990).
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The tenfold increase in friction from <0.02 to >0.1 caused by humidity and oxidation indicates that
the increase in bond strength from 0.08 eV/bond to 0.2 eV/bond (H bonding at CO sites by water
molecules) associated with a progressive friction-induced graphitization should be the main cause of the
change in friction (Table 24.2). However, such a description fails to account for the modification of the
rheology of the contact by triboreactivity between the transfer layer and the counterface, and by the
formation of friction-induced wear particles. This modification of the rheology has been shown by Kim
et al. (1991) by performing FTIR microprobe tests to identify the debris and film structures observed on
various interfacial films formed by a Si3N4 ball rubbed against DLC-coated Si in dry Ar, dry air, and
moist air conditions. Humidity caused a tribochemical reaction and led to a transfer film composed of
oxidized hydrocarbon and hydrated silica (from the ball). Consequently, friction was rather low, but wear
was significantly high. The opposite tribological behavior was observed in dry air, where the transfer
layer and debris were composed of a carbonyl compound formed by the tribo-oxidation of hydrocarbons
in the DLC films. The triboreactivity was reduced in dry Ar, thus lowering both friction and wear, with
a transfer film and debris whose composition was similar to that of the initial film.
24.3.3 Applications of DLC Films
The unique properties of DLC films and their modifications, together with the possibility of adjusting
the 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, chemical
inertness, infrared transparency, high electrical resistivity, and, potentially, field emission and low dielectric constant. Although ta-C has properties similar to or (in some instances) better than those of
conventional DLC (i.e., a-CH), thus far, only conventional DLC films have been used for practical
applications (see extensive reviews by Grill and Meyerson, 1994; Grill, 1997a). Because DLC is IRtransparent, 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 temperature
of DLC allows its use as a wear-protective layer on products made of plastic; therefore, it is used to protect
polycarbonate 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 protective
coatings for both magnetic disks and magnetic heads (Bhushan, 1999). Video recording or magnetic data
storage tapes, in which ferromagnetic metal is a recording medium, and the metallic capstans in contact
with the tapes are also being protected with DLC coatings to reduce wear and friction, thus extending
both the life of the tapes and their reliability. The announcements of the latest MACH3 razor blades by
Gillette 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. Its
potential 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 biological
environments. Because they are chemically inert and impermeable to liquids, DLC coatings could protect
biomedical implants against corrosion and serve as diffusion barriers. DLC is being considered for coating
metallic 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, which
are components of artificial heart valves, has been found capable of satisfying both the mechanical and
biological requirements and improving the performance of these components. The same properties may
make DLC useful as a protective coating for hip joint implants. Improvement of carbon/carbon composite
prosthetics by DLC coatings has also been demonstrated. Currently, DLC and its modifications are being
considered as low-dielectric materials for the interconnect structures of ultra-large-scale integrated circuits (ULSI). A better understanding of the means to control the thermal stability and other integration
problems 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 promising
application appears to be for cathodes in field-emission-based flat-panel displays or as pixel elements in
large outdoor displays.
© 2001 by CRC Press LLC
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 nitrides
and carbides used by industry today. It is the second hardest material known. Unlike diamond, it does
not dissolve in or interact with Fe-based alloys, and thus can be used to machine ferrous materials (which
account 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 excellent
resistance to wear and oxidation at elevated temperatures. CBN can also work very well under dry
machining 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 pand n-type elements; hence, like diamond, CBN can also be used in various optical and electrical
applications.
Like diamond, CBN can be synthesized by the HPHT method in the powder form and can be deposited
in thin films on various metallic and ceramic substrates by both PVD and CVD methods. The most widely
used deposition methods include magnetron sputtering, plasma-enhanced CVD, ion plating, pulse laser
deposition, ion-beam-assisted deposition, etc. (Murakawa and Watanabe, 1990; Chiang et al., 1997; Smeets
et al., 1997; Konyashin et al., 1997). The deposition temperatures needed for the formation of the CBN
phase are substantially lower than those required for diamond films. Also, the surface finish of CBN films
is 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, achieving
and maintaining a 1:1 stoichiometry between B and N atoms tends to be rather difficult. Most films are
slightly N-deficient and, hence, are not perfectly crystalline. Increases in B above stoichiometry and
concomitant formation of vacancies on the N sublattice lead to distortions and significant changes in
crystal structure. These changes can adversely affect the mechanical, thermal, and other important
properties of CBN. Depending on deposition temperature, ion energy, and current density, resultant
films may contain appreciable amounts of amorphous or hexagonal BN phases. For example, the formation of an sp3-hybridized CBN phase is largely controlled by the substrate temperature and the energy
of N and Ar ions in sputtering and ion-beam deposition. The presence or absence of H in the plasma
can also affect film quality. For higher quality CBN films, moderate ion energies are needed (Schaffnit
et al., 1996).
Growing CBN films tend to accumulate increasingly high compressive stresses in their microstructures
and this represents a major problem in attaining films that are thick enough (at least 3 to 5 µm) for
machining operations. The lack of suitable tool substrates with a good lattice and thermal expansion match
is a major part of this problem. During deposition, compressive stress within the film tends to increase
with increasing thickness and eventually becomes so high that CBN film delaminates prematurely or
fractures severely (Murakawa and Watanabe, 1990). Currently, safe film thickness for crystalline CBN films
is in the range of 0.2 to 0.5 µm. For most machining and other tribological applications, films thicker than
this 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 various
substrate 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 relieve
some of the residual stresses, and hence reduce the amount of stress buildup at the coating-substrate
interface. The use of a Ti layer or first layer, followed by a second layer of graded BN, was shown to be
quite effective in achieving strong adhesion between CBN and several tool materials (i.e., WC-Co alloy
and high-speed steel). Furthermore, it was shown that CBN can be directly grown on diamond films
without 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 improvements in achieving strong adhesion between CBN and various substrate materials, high residual stresses
in these films still hinder the deposition of very thick films with good adhesion.
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24.4.1.1 Tribology of CBN Films
Because of its excellent mechanical properties, chemical inertness, and high-temperature durability, CBN
holds 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 been
explored. Tribological studies have indicated that the friction and wear performance of CBN is highly
dependent on the test environment and counterface materials.
In a search of self-lubricating materials for space applications, Miyoshi (1999) performed comprehensive tests with diamond, DLC, and CBN in UHV environments. Surprisingly, it was found that CVD
diamond films against CBN exhibited the lowest coefficients of friction. Such impressive results led
Miyoshi to the conclusion that a combination of these two hardest materials may serve as an effective
self-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 TiN
or TiC films. When CBN is slid against steel in air or under water, wear is low; but in high vacuum, it
is high. Sliding CBN against Al led to high friction and wear in air and vacuum; but in water, wear was
almost zero.
Studies by Miyake et al. (1992) and Watanabe et al. (1991) demonstrated that the crystallinity of CBN
coatings strongly affects the friction and wear properties of ion-plated CBN films. In their tests, highly
crystalline CBN films provided the best wear resistance and lubrication, whereas films with an amorphous
structure exhibited good frictional properties but very short wear lives. Poor adhesion of hexagonal BN
led to premature debonding and hence poor wear performance. In the end, these authors concluded that
perhaps 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 CBN
and amorphous BN. More specifically, an initial layer of CBN was deposited onto an Si substrate by ion
plating, followed by a second layer of amorphous BN on top of the CBN layer with a gradient interlayer
formed between the two BN layers. The results of tribological tests showed that this composite film
exhibited 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 coefficients
during sliding in air. Comprehensive studies by Murakawa (1997) and Watanabe et al. (1991, 1995) have
demonstrated 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 their
excellent tribological properties, CBN films were deposited on WC-Co drills and tool inserts and subjected
to cutting operations. Except for partial delamination in a few cases, CBN coatings reduced torque while
drilling 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 resistance
to corrosion and high-temperature oxidation, and low-to-moderate deposition temperatures), CBN
suffers from poor adhesion, limited film thickness, high internal stresses, and poor reproducibility of
structure and properties. Currently, researchers are concentrating on overcoming these problems. Except
for a few case studies where superior machining and lower torque were demonstrated while cutting or
drilling 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 new
form of carbon nitride (i.e., cubic beta carbonitride β-C3N4) that exhibits an extremely low compressibility and superhigh hardness that exceed those of natural diamond. The unique crystal structure of
cubic β-C3N4 is based on the β-Si3N4 structure, in which C is substituted for Si to achieve an atomic
bond configuration that favors extreme hardness and elastic modulus. Because hardness is inversely
proportional to the bond length and hence to the radii of the atoms that form that bond, Cohen predicted
that such compounds based on C and N, which have very small atomic radii and the strongest bonds,
© 2001 by CRC Press LLC
will exhibit the highest hardness. Since these initial theoretical studies, many researchers have attempted
to synthesize β-C3N4 in both the bulk and thin-film forms by various methods, with the expectation that
this will lead to new possibilities in the fields of superhard and wear-resistant materials. In addition to
its extreme mechanical properties, researchers have predicted that β-C3N4 will provide excellent thermal
conductivity and wide bandgap (Cohen, 1995; Wang, 1997).
During the past decade, several new methods have been developed and used to synthesize crystalline
β-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 DC
magnetron sputtering, ion-beam deposition, ion implantation, plasma arc deposition, UV-assisted chemical synthesis, and hot-filament CVD) have produced amorphous carbon nitride films with a relatively
low N content (≈20 to 30%), mainly because of an extremely high bond dissociation energy for N in gas
discharge plasmas or ion fluxes. Occasionally, within the predominantly amorphous structure of these
films, some researchers have reported the presence of small crystallites with electron diffraction patterns
that 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 researchers
have tried to stabilize the crystalline phase in a pseudomorphic state using appropriate structural templates 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) and
ZrN 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. Although
these 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:N
ratio), amorphous C nitride films produced by various research groups exhibit large variations in mechanical properties and tribological performance. Reported friction values range from 0.05 to >0.5 in air. In
dry environments or high vacuum, much higher friction coefficients were observed. Such a large scatter
in friction has been mainly attributed to the differences in N content, extent of sp3 and/or sp2 hybridization, deposition methods and conditions, and, ultimately, the mechanical properties of the films. Kusano
et al. (1998) have found that CNx films deposited with >70% N in the sputtering gas exhibit higher
friction 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 between
friction and wear properties of CNx films, C:N ratio, and deposition conditions (i.e., temperature, ion
energy, ion current density, etc.). Li et al. (1994) have observed that in dry sliding contacts, CNx gives
an initial friction coefficient of 0.1 against a 52100 steel ball, but this value increases to ≈0.5 in steady
states. Zou et al. (1999) reported steady-state friction coefficients of 0.08 to 0.14 for CNx films produced
by a reactively ionized cluster beam technique.
As far as tribological applications are concerned, amorphous C nitride films are currently used as
protective coatings in magnetic hard disks. Khurshudov and Kato (1996) reported that the wear rates of
the carbon nitride-coated disks were 10 times lower than those of disks coated with a commercial DLC
film of the same thickness. These C nitride coatings also provided ≈3 to 30-times longer wear life than
commercial C coatings. Carbon nitride coatings can also be used in other tribological fields where
conventional DLC films have already been exploited, but the progress in implementation has been rather
slow.
24.5 Summary and Future Direction
Over the last 2 decades, great strides have been made in the deposition, characterization, and tribological
utilization 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
© 2001 by CRC Press LLC
properties that can be very useful for some demanding tribological applications. The current state-ofthe-art in diamond, DLC, and other related films allows many tribocomponents to be coated with these
films and offered for practical applications. Although some applications are still in the exploratory stage
and require further refinement, prototypes of others have been successfully produced and are currently
being evaluated for endurance and reliability.
As a true reflection of the growing interest in these films, the number of publications that deal with
them has steadily increased over the years. It is not possible to refer to all of them in this chapter, but
even a simple literature survey will yield several key publications for those who are interested in more
in-depth information about these films. In addition to some archival journals, several technical/scientific
books and book chapters, as well as conference proceedings, are now available. These publications have
led to a better understanding of the microstructure and chemistry of these films and, hence, have led to
their greater utilization by industry.
The relatively high cost of diamond and poor adhesion and reliability of CBN are delaying their wider
applications in industry. Until now, only tool inserts and drills have been coated successfully with diamond
and CBN and been made available for limited commercial use. However, these superhard coatings have
much to offer future tribological applications. It is anticipated that their use in the cutting-tool industry
will 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 also
been 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 coatings
are 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 actual
application. DLC and C nitride films are more cost-effective and much easier to produce than diamond
and CBN; hence, they are now used in various tribological applications.
The combination of high wear resistance (due to high mechanical hardness) and low friction makes
diamond, 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 and
control their friction and wear performance. Existing data suggest that, with proper control of their
microstructure, chemistry, and surface topography, these films may live up to their promise. The high
surface roughness of microcrystalline diamond films presents serious problems in many tribological
applications. These microcrystalline films can be polished by various methods or, alternatively, nanocrystalline diamond films with a very smooth surface finish can be used to achieve low friction and wear
in sliding tribological applications. These smooth films are particularly suitable for mechanical shaft seals
and MEMS applications. As with other types of hard coatings, strong bonding between these films and
their 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 life
improvement to cutting tool inserts when they are used to machine Al-Si alloys and graphite, whereas CBN
is more suitable for machining ferrous and nonferrous alloys. When compared to diamond, CBN has not
yet found large-scale applications in the industrial world. Poor adhesion, insufficient thickness, high
internal stresses, and difficulties in attaining and maintaining 1:1 stoichiometry between B and N atoms
in a cubic crystalline structure during deposition are some of the major problems that hinder further use
of CBN in tribological fields. With recent advances in deposition technology, the quality and reliability of
these coatings are expected to improve in the near future and, hence, increase their chances for large-scale
applications. Briefly, the future of diamond, DLC, and related coatings in tribological applications looks
promising 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 Contract
W-31-109-Eng-38.
© 2001 by CRC Press LLC
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