Bhushan B. Handbook of Micro/Nano Tribology, Second Edition

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were detected occasionally. However, no AE signals were detected during the distinct load–displacement

discontinuity on unloading which is seen in Figure 10.42a at 60 mN. Since no emission was detected,

the source of the discontinuity is more likely to be a sluggish transformation (Hu et al., 1986; Clarke

et al., 1988; Pharr, 1992) than a rapid crack growth.

Gupta et al. (1993, 1994) and Gupta and Bhushan (1994) reported that hysteresis in cyclic indentation

and discontinuity kinks in the unloading curve are considerably reduced by ion implantation of com-

pound forming species 0

and N

into single-crystal silicon. They further reported that amorphous silicon

ﬁlms did not exhibit either hysteresis in the cyclic indentation or a discontinuity kink in the indentation

loads ranging from 1 to 90 mN. This suggests that for a perfect amorphous structure, the hysteresis does

FIGURE 10.40 Load vs. displacement plots for (110) silicon (a) at a peak load of 120 mN and (b) at a peak load

of 4.5 mN. (From Pharr, G.M., 1992, in Thin Film: Stresses and Mechanical Properties III, W.D. Nix et al., eds., Vol. 239,

pp. 301–312, Materials Research Society, Pittsburgh. With permission.)

not occur because of absence of crystallographic pressure-induced phase transition during cyclic loading

and unloading. In addition, the disordered structure does not allow the nucleation and propagation of

the lateral cracks beneath the indentation. They have also reported that ion implantation has an insig-

niﬁcant effect on hardness and elastic modulus. Improvement in the hardness is about 10% as a result

FIGURE 10.41 A schematic illustration of the coupling of the phase transformation and cracking during indenta-

tion of single-crystal silicon. (From Pharr, G.M., 1992, in Thin Film: Stresses and Mechanical Properties III, W.D. Nix

et al., eds., Vol. 239, pp. 301–312, Materials Research Society, Pittsburgh. With permission.)

FIGURE 10.42 (a) Applied force as a function of tip displacement for Si(100) and (b) AE amplitude as a function

of time during loading. (From Weihs, T.P. et al., 1992, in Thin Films: Stresses and Mechanical Properties III, Symp.

Proc., W.E. Nix, et al., eds., Vol. 239, pp. 361–370, Materials Research Society, Pittsburgh. With permission.)

of implantation with carbon and boron ions. Gupta et al. (1993, 1994) reported that ion-bombarded

silicon exhibits a very low coefﬁcient of friction (~0.05) and a low wear factor (~10

–6

–1

) when

slid against alumina and 52100 steel balls.

Weihs et al. (1992) indented Ni ﬁlms on a glass substrate. The Ni ﬁlms debonded from their substrates

at forces ranging between 130 and 250 mN. The debonding events were marked by the indenter tip

jumping downward as “chunks” of the Ni ﬁlm buckled away from the underneath. The indenter tip

jumped a distance equal to the ﬁlm thickness as it initially debonded. Figure 10.43 also shows the

corresponding AE trace with a rise time of 1.8 µs. In this particular test, debonding continued at higher

forces and a second AE event was recorded. After each test, optical microscopy conﬁrmed the delamina-

tion of the ﬁlm from underneath the indenter.

Gupta and Bhushan (1995) indented 20-nm-thick amorphous carbon (a–C, also known as diamond-

like carbon or DLC) ﬁlms deposited by four different deposition techniques, sputtered Al

and sput-

tered SiC ﬁlms on Al

–TiC substrate, Figure 10.44. The indentation depths are larger than the ﬁlm

thickness, which means that the tip had penetrated through the ﬁlms into the substrate material during

the indentation process. Most of the ﬁlms exhibit a discontinuity or pop-in marks in the loading curve,

shown by arrows in Figure 10.44, which indicate a sudden penetration of the tip into the sample. A

nonuniform penetration of the tip into a thin ﬁlm possibly results due to formation of cracks in the ﬁlm,

formation of cracks at the ﬁlm–substrate interface, or debonding or delamination of the ﬁlm from the

substrate. The Al

–TiC substrate exhibits a smooth loading curve free from any pop-in marks, which

indicates that a 0.5-mN peak indentation load is insufﬁcient to damage the surface. All of the ﬁlms

exhibited indentation depths at the peak load higher than that of the Al

–TiC substrate.

Oliver and Pharr (1992) have reported that load–displacement curves during unloading are not linear,

rather each unloading curve is slightly concave over its entire span. They showed that the unloading

FIGURE 10.43 (a) Applied force as a function of tip displacement for Ni ﬁlm on glass substrate and (b) AE

amplitude as a function of time during loading. (From Weihs, T.P. et al., 1992, in Thin Films: Stresses and Mechanical

Properties III, Symp. Proc., W.E. Nix, et al., eds., Vol. 239, pp. 361–370, Materials Research Society, Pittsburgh. With

permission.)

curves are well described by a power law relation like that of Equation 10.7 presented earlier. For the six

materials tested, the exponents varied from 1.25 to 1.50 and were not distinctly greater than 1 which

implies that none of the data is consistent with ﬂat punch behavior.

10.4.2 Continuous Stiffness Measurements

Continuous stiffness measurements are used for continuous measurement of sample stiffness or com-

pliance (related to elastic modulus) during indentation. Oliver and Pharr (1992) measured continuous

stiffness of electropolished single-crystal tungsten at low peak loads of 0.5 mN or less, for purely elastic

contact. Continuous stiffness data for the indentation data in Figure 10.37b are presented in Figure 10.45a.

Comparisons of these data with the load–time history, shown in Figure 10.35, show that the measured

contact stiffness and thus the contact area does increase and decrease in the way that would be expected

based on the loading history. The continuous stiffness data for the indentation at 1.5-mN load (plastic

contact, Figure 10.37c) is shown in Figure 10.45b. It is seen that for each of the four unloadings, the

contact stiffness changes immediately and continuously as the specimen is unloaded. Thus, the contact

area, which varies in that same way as the contact stiffness, is not constant during the unloading of the

plastic hardness impression, even during the initial stages of unloading.

10.4.3 Hardness and Elastic Modulus Measurements

Nanoindenter is commonly used to measure surface mechanical properties of bulk materials (Pharr et al.,

1990; Lucas et al., 1991; Oliver and Pharr, 1992; Bhushan and Gupta, 1995; Bhushan et al., 1996).

FIGURE 10.44 Load vs. displacement plots for various 20-nm-thick amorphous carbon ﬁlms and ion beam–sput-

tered Al

and RF-sputtered SiC coatings on Al

–TiC substrate and bare Al

–TiC substrate. (From Gupta, B.K.

and Bhushan, B., 1995, Wear 190, 110–122. With permission.)

Hardness and elastic moduli for six bulk materials and three materials used as a substrate for the

construction of magnetic rigid disks and single-crystal silicon are presented in Figure 10.46. The data

show that there is a very little indentation size effect in several materials on the hardness values. In the

case of aluminum and tungsten in Figure 10.46a, there is a modest increase in hardness at low loads.

According to Oliver and Pharr (1992), this could be due to surface-localized cold-work resulting from

polishing. The modulus data also show that there is very little evidence for an indentation size effect;

i.e., the moduli remains more or less constant over the entire range of load. Oliver and Pharr (1992)

reported that the hardness and modulus values at the two highest loads are comparable with the literature

values. Comparison of published values and measured values of magnetic disk substrates in Figure 10.46b

is shown in Table 10.4. (Fracture toughness data will be discussed later.)

Pharr et al. (1990) reported a nanoindentation hardness value of (110) single-crystal silicon to be

about 11.4 GPa at a peak normal load of 120 mN and corresponding peak indentation depth of 750 nm.

This value was slightly higher than the macrohardness value of silicon of 9 to 10 GPa (Anonymous, 1988).

Using an atomic force microscope, Bhushan and Koinkar (1994) measured hardness of single-crystal

FIGURE 10.45 Contact stiffness vs. time. (a) for a fully elastic contact and (b) for a fully plastic contact on an

electropolished single-crystal tungsten measured with the continuous stiffness technique. (From Oliver, W.C. and

Pharr, G.M., 1992, J. Mater. Res. 7, 1564–1583. With permission.)

silicon at an ultrasmall normal load of 65 µN and corresponding indentation depth of 2.5 nm and

reported a value of 16.6 GPa. It is clear that the hardness of silicon increases with a decrease in the load

and corresponding indentation depth. An increase in hardness at lower indentation depths may result

from contributions of the surface ﬁlms. At smaller volumes, there is a high probability that indentation

would be made into a region that was initially dislocation free. Furthermore, at small volumes, there is

an increase in the stress necessary to operate dislocation sources (Gane and Cox, 1970; Sargent, 1986).

These are some of the plausible explanations for the increase in hardness at smaller volumes.

A number of investigators have used the nanoindenter to study the effect of ion implantation on the

hardness and elastic modulus of metals (Nastasi et al., 1988; Lee and Mansur, 1989; Bourcier et al., 1990,

1991; Was, 1990), silicon (Gupta et al., 1993, 1994; Gupta and Bhushan, 1994), ceramics (McHargue,

1989; McHargue et al., 1990; O’Hern et al., 1990; Was and DeKoven, 1991) and polymers (Lee et al.,

1992, 1993; Rao et al., 1993). Generally, the ion-implanted zone is shallow and the nanoindenter is capable

of measuring the changes in mechanical properties.

The nanoindenter is ideal for measurement of mechanical properties of thin ﬁlms and composite

structures. A number of investigators have reported hardness of composite structures with thin ﬁlms on

a substrate (Pethica et al., 1985; Burnett and Rickerby, 1987c; Stone et al., 1988; Wu et al., 1988, 1989,

1990b; O’Hern et al., 1989; Bhushan and Doerner, 1989; Jiang et al., 1989; Joslin et al., 1989; Baker et al.,

FIGURE 10.46

FIGURE 10.46 Hardness and elastic modulus (a) as a

function of load for six bulk materials — mechanically

polished single-crystal aluminum, electropolished single-

crystal tungsten, soda lime glass, fused silica, (001) single-

crystal sapphire, and (001) single-crystal quartz (Oliver

and Pharr, 1992), and (b) as a function of indentation

depth (load ranging from 0.1 to 2.5 mN) for mechanically

polished 10-µm-thick electroless Ni–P ﬁlm on Al–Mg

alloy 5086, chemically strengthened alkali-aluminosilicate

glass, chain silicate glass–ceramic (polycrystalline) (Can-

asite by Corning), and single-crystal silicon (111). (From

Bhushan, B. and Gupta, B.K., 1995, Adv. Info. Storage Syst.

6, 193–208. With permission.)

TABLE 10.4 rms Roughness Values Measured by an Atomic Force Microscope, Hardness and Elastic

Modulus Values Measured by Using a Nanoindenter (at an indentation depth of about 20 nm), and Fracture

Toughness Measured by Using a Vickers Microindenter

Material

rms

Roughness

(nm)

Hardness

(GPa)

Elastic

Modulus

(GPa)

Crack Length,

c (µm)

Fracture

Toughness

(MPa·m

1/2

)

Ni–P/Al–Mg 3.6 6.0 (5.5) 130 (200) No cracks —

Chemically strengthened glass 1.1 6.0 (5.8) 85 (73) Signiﬁcant cracking — (0.9)

Glass–ceramic 6.1 8.5 (5.5) 100 (83) 59.4 0.65 (4.0)

Single-crystal silicon (111) 0.95 11.0 (9-10) 200 (180) — —

Values in parentheses are the reported values measured by conventional Vickers indentation method.

Values in parentheses are the reported values measured by conventional tensile pull test method.

Values in parentheses are the reported values measured by the chevron notched short bar method.

1990; Bull and Rickerby, 1990; Cammarata et al., 1990; Cho et al., 1990; Cornett et al., 1990; Rubin et al.,

1990; Chou et al., 1991; Cooper, 1991; Cooper and Beetz, 1991; Gissler et al., 1991; Schlesinger et al.,

1991; Stone et al., 1991; Bhushan et al., 1992, 1995, 1997; Fabes et al., 1992; Knight et al., 1992; Lucas

and Oliver, 1992; Savvides and Bell, 1992; Wang et al., 1992; Whitehead and Page, 1992; Li et al., 1993;

Nastasi et al., 1993; Vancoille et al., 1993; Gupta and Bhushan, 1995a,b; Patton and Bhushan, 1996;

Bhushan and Li, 1997; Li and Bhushan, 1998b, c).

True hardness of the ﬁlms can be obtained if the indentation depth does not exceed about 30% of the

ﬁlm thickness. At higher indentation depths, the composite hardness changes with the indentation depth.

Measured hardness values of soft Ti ﬁlms on a hard sapphire substrate are presented in Figure 10.47. We

note that hardness increases with a decrease in the ﬁlm thickness or increase in the indentation depth,

as expected. The ﬁlm hardness is the steady-state hardness independent of the indentation depth.

Hardness and elastic modulus proﬁles as a function of indentation depth at the peak load for 400-nm-

thick amorphous carbon ﬁlms deposited by various deposition techniques, SiC ﬁlms, and single-crystal

silicon substrates are presented in Figure 10.48. We note that the hardness and elastic modulus of cathodic

arc carbon and radio frequency (RF)-sputtered SiC ﬁlms tend to decrease with the indentation depth,

which is attributed to increased contributions of the silicon substrate with a lower hardness of 11 GPa,

at larger depths. Cathodic arc carbon ﬁlm exhibits the highest hardness of 38 GPa and elastic modulus

of 300 GPa as compared with that of other coatings. The high hardness and elastic modulus of cathodic

arc ﬁlm are followed by RF-sputtered SiC, ion beam carbon, and plasma enhanced-chemical vapor

deposition/direct current (PECVD/DC) sputtered carbon ﬁlms. High hardness and elastic modulus of

cathodic arc carbon ﬁlm are attributed to the high kinetic energy of the carbon species involved in the

cathodic arc deposition. The differences in the hardness and elastic modulus of carbon ﬁlms is attributed

to their varying sp

- to sp

-bonding ratio and the amount of hydrogen (Gupta and Bhushan, 1995a,b).

The hardness and elastic modulus data are summarized in Table 10.5 (Gupta and Bhushan, 1995a,b).

FIGURE 10.47 Indentation hardness as a function of indenter displacement for Ti coatings on sapphire substrates.

Numbers next to each set of data correspond to the coating thickness. (From Fabes, B.D. et al., 1992, J. Mater. Res.

7, 3056–3064. With permission.)

10.5 Microscratch Resistance Measurement of Bulk

Materials Using Micro/Nanoscratch Technique

Microscratch technique is commonly used for screening of bulk materials for wear resistance. As stated

earlier, normal load applied to the scratch tip is gradually increased during scratching until the material

is damaged. Friction force is sometimes measured during the scratch test (Bhushan and Gupta, 1995;

Bhushan et al., 1996). After the scratch test, the morphology of the scratch region including debris is

FIGURE 10.48 Hardness and elastic modulus as a function of indentation depth at the peak load of 400-nm-thick

carbon ﬁlms deposited by cathodic arc, ion beam, PECVD, and DC-sputtered and 400-nm-thick and RF-sputtered

SiC ﬁlm on single-crystal silicon substrates. Hardness and elastic modulus proﬁles for single-crystal (100) silicon are

also included for comparison. (From Gupta, B.K. and Bhushan, B., 1995, Thin Solid Films 270, 391–398. With

permission.)

TABLE 10.5 Comparison of Hardness and Elastic Modulus of Various

400-nm-Thick Coatings Measured by Nanoindentation and Bulk Samples

Data Reported by Other Researchers

Coating Hardness (GPa) Elastic Modulus (GPa)

Cathodic arc carbon (a–C) 38 300

Ion beam carbon (a–C:H) 19 150

PECVD carbon (a–C:H) 17 140

DC sputtered carbon (a–C:H) 15 140

Bulk graphite (for comparison) Very soft 9–15

Diamond (for comparison) 80–104 900–1050

Ion beam sputtered Al

8.3 140

RF sputtered SiC 27 255

Single-crystal silicon (bulk) 11 220

observed in an SEM. Based on the combination of the changes in the friction force as a function of

normal load and SEM observations, the critical load is determined and the deformation mode is inden-

tiﬁed. Any damage to the material surface as a result of scratching at a critical ramp-up load results in

an abrupt or gradual increase in friction. The material may deform either by plastic deformation or

fracture. Ductile materials (all metals) deform primarily by plastic deformation, resulting in signiﬁcant

plowing during scratching. Tracks are produced whose width and depth increase with an increase in the

normal load. Plowing results in a continuous increase in the coefﬁcient of friction with an increase in

the normal load during scratching. Debris is generally ribbonlike or curly, whereas brittle materials

deform primarily by brittle fracture with some plastic deformation. In the brittle fracture mode, the

coefﬁcient of friction increases very little until a critical load is reached at which the material fails

catastrophically and produces ﬁne debris, which is rounded, and the coefﬁcient of friction increases

rapidly above the critical load.

The coefﬁcient of friction proﬁles as a function of normal load for scratches made on various ceramic

substrates and corresponding magnetic disks (substrates coated with 75-nm-thick sputtered Co–Pt–Ni

FIGURE 10.49 Coefﬁcient of friction proﬁles as a function of normal load for 500-µm-long scratches made using

a diamond tip with 1-µm tip radius, at a normal load ranging from 1 to 12 mN on various ceramic substrates,

corresponding magnetic disks, and single-crystal silicon. (From Bhushan, B. and Gupta, B.K., 1995, Adv. Info. Storage

Syst. 6, 193–208. With permission.)