Structure, mechanical, and tribological properties of ...

Structure, mechanical, and tribological properties of titanium implanted alumina S. M. M. Ramos, B. Canut, L. Gea, and P. Thevenard Universite Claude Bernard Lyon I, Departement Physique des Materiaux, URA CNRS 172, Campus de la Doua, 43 boulevard du 11 Novembre 1918, 69622 Vdleurbanne Cedex, France

M. Bauer, Y. Maheo, Ph. Kapsa, and J. L. Loubet Ecole Centrale de Lyon, Laboratoire de Technologie des Surfaces, URA CNRS 855, 36 avenue Guy de Collongue, B.P. 163, 69131 Ecully Cedex, France (Received 1 February 1991; accepted 3 September 1991)

A study of the effects of titanium ion implantation on the structural, mechanical, and tribological properties of single crystal and polycrystalline a-alumina has been carried out. Rutherford Backscattering Spectrometry (RBS) in channeling geometry shows that a great proportion of implanted titanium ions are substitutional at low fluence. This fraction falls to near zero as an amorphous layer is formed. The chemical states for implanted titanium are determined by X-ray Photoelectron Spectroscopy (XPS). Titanium is present in the Ti3+ state near the surface and as metallic Ti° and as Ti3+ at depths corresponding to higher local concentration of titanium. The same behavior is observed for polycrystalline and single crystal a-alumina. Nanoindentation experiments show that low fluence implantation of titanium results in an increase of mechanical properties whereas high fluence implanted samples exhibit reduced hardness and Young's modulus compared to unimplanted samples. The friction coefficient is not changed by titanium ion implantation. Likewise, the wear characteristics were not changed by low fluence implantation, but amorphization at high fluence leads to a greater disk wear rate.

I. INTRODUCTION Structural changes induced by radiation damage and chemical effects strongly affect the surface mechanical properties of implanted ceramics.1"9 Some of the relationships between the microstructure of the implanted layers and the surface mechanical properties of single crystal a-alumina have been shown by McHargue et al.2 However, a comparison of the effects of ion implantation on polycrystalline materials and single crystals has not yet been studied in detail. The misorientation of grains and grain boundary effects should have a strong influence on the mechanical properties of polycrystalline samples. It has been observed, for example, that the critical fluence for amorphization of a-alumina sapphire is dependent on the bombarded surface orientation.10 The aim of this paper is to investigate the effects of titanium ion implantation on the structure and the mechanical and tribological properties of alumina. A comparison between the tribological behavior of implanted single crystal and polycrystalline a-alumina is made. For this purpose titanium ions have been implanted in both single crystal and polycrystalline a-alumina. Rutherford Backscattering Spectrometry (RBS) in channeling geometry and X-ray Photoelectron Spectroscopy (XPS) were used for structural and chemical characterization. The mechanical properties are described by measurements of the Young's modulus and hardness of 178

J. Mater. Res., Vol. 7, No. 1, Jan 1992

sapphire. Tribological properties (wear and friction coefficient) were studied on both single crystal and polycrystalline samples using a pin on disk apparatus. II. EXPERIMENTAL A. Materials Two types of alumina samples have been used: (1) commercial sintered polycrystalline a-alumina, T195 from Cice Company, with a porosity of about 6% and a grain size of about 10 /xm, and (2) single crystals of a-alumina (sapphire) from Criceram Company, the sample surfaces used for the tests being parallel to the basal plane (0001). The specimens were polished to a mirror-like surface finish with a fine diamond paste, and before tests they were cleaned in an ultrasonic bath with pure acetone and rinsed with pure propanol. The total surface roughness measured was less than 100 nm for sapphire, while for polycrystalline alumina it was less than 100 nm between surface voids and 4000 nm for the entire surface. B. Implantation Ion implantation was performed using a 200 kV Balzer's implantor with titanium ions of 150 KeV and fluences from 1016 up to 1017 ions • cm"2. The target temperature during implantation was maintained at © 1992 Materials Research Society

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S.M.M. Ramos et al.: Structure, mechanical, and tribological properties of titanium implanted alumina

300 K. In this study, no effect of irradiation on the size of surface voids has been observed. C. Physico-chemical characterization The implanted samples were characterized by RBS in random and channeling geometry and by XPS. RBS measurements were performed using a 2 MeV 4 He + beam generated by a van de Graaff accelerator, with a beam current around 10 nA. The sample was mounted on a three axis goniometer-head whose angle resolution is better than 0.05°. A 15 KeV resolution surface barrier detector was used with a scattering angle of 150°. In these experiments, the cross section of the beam impinging the sample never exceeded 1 mm 2 . XPS measurements were performed using Riber's spectrometer and Al KQ irradiation. The titanium depth profile was observed by successive sputterings. A depth calibration performed by RBS measurements on stripped samples allowed us to determine the sputtering velocity that was assumed to be constant during all the process. Moreover, we have also supposed that the sputtering does not induce chemical reactions with different atoms of implanted samples. D. Mechanical properties In order to measure the mechanical properties, Young's modulus (E) and hardness (//) of the implanted sapphire surfaces, nanoindentation experiments with a trigonal indenter were carried out. A comparison of the results with those of unimplanted sapphire is given. Unfortunately, no mechanical properties measurements could be done on polycrystalline a-alumina: voids prevent one from giving reliable results for such surfaces. Static and dynamic forces were determined as a function of the indentation depth h (0-300 nm). The apparatus and the method used are described elsewhere.11 The addition of a vibrating component to the linear displacement of the indenter allows the plasticity and the elasticity to be continuously and simultaneously characterized along the indentation depth in a nanoindentation experiment. Two types of measurements are made: static measurements of force (P) and displacement (h), and dynamic measurement of the mechanical transfer function. An experimental transfer function of the apparatus is taken into account in order to calculate the transfer function of the contact from the measurements. Its real and imaginary parts represent the stiffness of the contact, K, and the damping coefficient of the plane/indenter contact, respectively. In the case of a diamond tip-sapphire plane contact, the imaginary part is negligible compared to the real part.

The hardness and Young's modulus can be calculated from the two records (P) and (K). The hardness is defined by H = P/A where (A) is the projected area of the residual indent, geometrically related to an ideal indentation depth, (6), by:

A = 82/a with a = 0.0364 for a 115.2° trigonal indenter. The composite modulus of contact Ec, obtained by modeling the indenter/surface contact,11 is: -l

Em

Ed

where u is the Poisson's ratio and E the elastic modulus; subscripts m and d relate to the tested material and to the diamond, respectively. We assume that the Poisson's ratio of sapphire is vm = 0.3 and that for the diamond (1 - v\)IEd = 10 12 Pa"1. Then: 8

\(hT - h'R)8]

with 0 = 0.1782 for the 115.2° trigonal indenter. For this type of indenter, 8 = (h'R + h0) where h'R = h — P(h)/K(h) and h0 is an experimentally determined length, characteristic of the particular indenter. It has the same height as the distance over which the tip of the indenter deviates from the true pyramid. h0 is calculated using the constant modulus hypothesis, taking unimplanted sapphire as a reference.11 In our experiments the tip defect equivalent depth, h0, is 16 nm and so the indentation depth (hj) where the contact can be considered as an ideal pyramid on a plane is hj = 25 nm.11 E. Tribological testing Friction tests were performed using a rotating pin on a flat tribometer. The sphere and the disk are of the same type of alumina as is being tested: single crystal sapphire or polycrystalline a-alumina. Tribological properties were studied on high fluence, implanted, poly- and single a-alumina crystals (1017 Ti+ • cm"2) and on low fluence, implanted, polycrystalline a-alumina. Only the disk was implanted. The tribological test conditions were: sphere diameter : 10 mm, sliding speed : 24 mm s~\ normal applied load : 5 N, test duration : 1000 cycles, and ambient temperature and atmosphere (relative humidity between 40 and 60%). Under these conditions, the contact is elastic, the Hertzian radius of the contact, a, is 50 /zm, and the mean


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S. M. M. Ramos et al.: Structure, mechanical, and tribological properties of titanium implanted alumina

pressure at the beginning of a test is Pm = 7.108 Pa. No cracks due to tensile stress were created during testing. The increase in temperature due to sliding is negligible.12 The tangential force is measured during the whole experiment and the friction coefficient is then calculated as the tangential force divided by the normal force. After the test, the wear scars are observed with an optical microscope and with a scanning electron microscope, in order to obtain qualitative information on wear phenomena. Surface profiles perpendicular to the wear track on the disks give quantitative information regarding the amount of material removed.

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A. Microstructural evolution Figure 1 shows the random (a) and aligned (b) spectra obtained on sapphire implanted at 150 keV with 1016 Ti+ • cm"2 [Fig. l(a)] and 1017 Ti+ • cnr 2 [Fig. l(b)]. The high energy peak, «1400 keV, of the random spectrum is related to the titanium distribution, and the humps appearing in the substrate region of the channeling curve (b) give information about the disorder induced by titanium implantation in aluminum (1000-1200 keV) or oxygen (600-800 keV) sublattices. In Fig. l(a), these humps do not coincide with the random spectrum, which means that the amorphization of the substrate did not occur. In this case, a significant decrease in the titanium backscattering yield is observed in the channeling position, caused by the presence of implanted ions in substitutional sites. On the other hand in Fig. l(b), the amorphization of the substrate is shown by superposition of random and aligned spectra. The fraction of substitutional titanium was calculated by using the following expression: a = where XTi is the ratio between titanium peak areas in aligned and random configurations and Xmin is the dechanneling yield determined by the ratio between aligned and random yields in the aluminum sublattice. In a-alumina implanted with 1016 Ti+ • cm""2 the fraction of substitutional titanium is 53%. For 1017 Ti+ • cm"2, an amorphous layer was formed for which a has no physical meaning. Similar behavior has been shown by Romana et a/.13 for single crystal a—alumina implanted with niobium. However, for a fluence of 1016 Nb+ • cm"2 the niobium substitutional fraction was 28%, about one-half the titanium substitutional fraction measured in this study. The depth distribution profile of implanted titanium does not show any change in shape up to a high fluence and the measured parameters are in good agreement with theoretical predictions (TRIM). 14 180

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(b) FIG. 1. RBS channeling spectra for sapphire implanted with two different fiuences of titanium at an energy of 150 keV: (a) 1016 Ti+ • cm"2, random spectrum (implanted sample); (b) 10 17 Ti+ • cm"2, channeled spectrum (implanted sample).

The same behavior is observed for polycrystalline samples. The random RBS spectra show a depth distribution profile of implanted particles, in good agreement with TRIM predictions. The different chemical states of implanted titanium were determined by XPS analysis along the titanium depth profile. For a sample implanted with 1017 Ti+ • cm"2 the oxidation state of titanium near the surface corresponds to Ti3+ with the characteristics of Ti 2 O 3 . The relative amount of Ti in this oxidation state decreases as the local concentration of titanium increases. Near the maximum of the implantation profile the oxidation state corresponds to a mixture of metallic Ti and Ti3+ (Ti2O3) (Fig. 2). Complementary experiments by x-ray diffraction at glancing incidence have not shown the presence of a large size metallic Ti cluster. So, if some are present, their size is probably less than 1 nm because for Nb implantation, this last technique has allowed the observation of 1 nm size metallic Nb clusters.13




50 ,

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DEPTH (nm) FIG. 2. Depth distribution of the oxidation state of titanium implanted in sapphire (implantation fiuence: 10 17 Ti+ • cm" 2 ).

In a recent work15 we have shown that the amorphization phenomenon in some refractory oxides (sapphire, TiO2, MgO) implanted with niobium can be correlated with different final charge states of the implanted particles. In that work it was shown that for the same energy and fluence of niobium implantation, titanium oxide and sapphire are amorphized and the niobium charge state varies from Nb5+ (at the surface) up to Nb° (at the region of maximal local concentration). For magnesium oxide implanted in the same conditions, amorphization was not observed and the niobium charge state varies from Nb5+ up to Nb 4+ with characteristics of NbO2. So, the charge state of niobium always remains at a high value independent of the local concentration. A similar process, observed for sapphire and titanium oxide implanted with niobium, seems to occur for titanium implantation on sapphire. The interpretation of these results is that a local charge modification can produce a deformation of octahedral sites, which can also be correlated with the amorphization process. B. Mechanical properties Figure 3 shows the change in the hardness (H) and Young's modulus (£) with depth and fluence for implanted sapphire. The broken lines of the curves correspond to the layer where the values of E and H are not reliable.11 For a low dose, 1016 Ti+ • cm"2, where radiation damage does not produce amorphization, the hardness and Young's modulus increase. For an indentation depth of 50 nm (point M in Fig. 3) there was an increase of about 16% for the hardness and 17% for the Young's modulus compared to the values of the unimplanted sapphire. For the higher dose, 1017 Ti+ • cm"2, where the implantation produced the amorphization, the values of these mechanical properties of the implanted sapphire

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(b) FIG. 3. Evolution of the mechanical properties (E) and (H) with depth for unimplanted low fluence implanted sapphire (10 16 Ti+ • cm"2) and high fluence implanted sapphire (10 17 Ti+ • cm" 2 ): (a) hardness (//) versus depth and (b) Young's modulus (E) versus depth.

are lower than those of the unimplanted samples. For an indentation depth of 50 nm (point M), decreases of about 57% for the hardness and 30% for the Young's modulus are observed with respect to the unimplanted sapphire. This behavior is in agreement with results for other implanted ions16'17; amorphization due to high dose implantations of Ti lead to a decrease in the hardness and elastic modulus of sapphire. For the higher dose 1017 Ti+ • cm"2, the decrease in hardness is observed to a depth of about 250 nm. The decrease in Young's modulus E with respect to the unimplanted sapphire extends to a 150 nm indentation depth. This depth corresponds to the thickness of the implanted layer. This difference can be explained on the basis of the relative volumes of materials affected by the plastic and elastic stress fields during indentation. The plastically deformed volume is hemispherical


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S. M. M. Ramos et si.: Structure, mechanical, and tribological properties of titanium implanted alumina

with a radius of about 2.5a, where a is the radius of the sapphire-indenter contact circle.18 The elastically deformed volume is hemispherical, too, but its radius is about 10a.19 Thus, the influence of the substrate is observed for smaller indentation depths in the case of Young's modulus E than in the case of hardness H. C. Tribological behavior 1. Friction coefficient The change in the friction coefficient versus the number of cycles for implanted and unimplanted a-alumina is shown in Fig. 4, where the curve in Fig. 4(a) represents the friction coefficient of

POLYCRYSTALLINE «-AI2O3

0.5

polycrystalline a-alumina while the curve in Fig. 4(b) concerns sapphire. The initial value is 0.1. This low value is related to the adsorption of water molecules and physical-chemical state of the surface created by the cleaning process.12 After the beginning of sliding, an increase in the friction coefficient is observed. The steady state values are 0.32, 0.35, and 0.38 for unimplanted polycrystalline a-alumina, low fluence implanted polycrystalline a-alumina (1016 Ti+ • cm 4 ), and high fluence implanted polycrystalline a-alumina (1017 Ti+ • cm"2), respectively. Similar values were obtained for sapphire: 0.3 for unimplanted sapphire and 0.29 for high fluence implanted sapphire. The difference between the friction coefficients measured is not significant when fluctuations in their values during sliding ( A / = 0.07) and the test reproducibility are taken into account. A friction coefficient of 0.3 to 0.4 is classically obtained for a-alumina sliding on a-alumina in mild tribological conditions.20"22 2. Wear

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NUMBER OF CYCLES (b) FIG. 4. Friction coefficient (mean value) evolution versus the number of cycles. The vertical bar represents the amplitude of variations during sliding, (a) Polycrystalline a-alumina disks implanted and unimplanted, and unimplanted a-alumina polycrystalline spheres 10 mm diameter, (b) Single crystal a-alumina disks and unimplanted sapphire spheres (10 mm diameter): unimplanted sapphire disk and implanted sapphire disk (10 17 Ti+ • cm" 2 ).

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Unimplanted a-alumina. In our conditions, after the tests, the surface degradation is slight. For disks (single crystal or polycrystalline a-alumina), the wear tracks are not perceptible to the eye. The wear scars on the a-alumina spheres are circular, with a diameter of about 130 //m. Some wear debris accumulates in front of the contact area. These wear particles were found to be amorphous alumina by examination by transmission electron microscopy. Their diameters are less than 10 nm. As-implanted polycrystalline alumina. For the low fluence implanted polycrystalline a-alumina (1016 Ti+ • cm"2), the mechanical properties of the surface are not very different from the unimplanted case, although the nature and the structure of the material are different. Observations on the worn surface are similar. The wear scar diameter on the sphere is about 130 /xm. The wear scar depth of the disk was estimated to be about 20 nm. The wear of sliding bodies corresponds to the presence of a few particles that are compacted at the front of the contact area and that adhere to the sphere. Some of them are visible on both sides of the disk wear track, and a very thin surface layer is created (Fig. 5). The behavior of the high fluence implanted a-alumina (1017 Ti+ • cm"2) is different. The surface has a reddish appearance in optical microscopy examination and the disk wear track is more apparent. The wear scar of the sphere is still about 130 /im in diameter, but the wear of the disk is greater than in the other cases. This is probably related to the relatively low mechanical properties of the amorphous surface layer created under high dose implantation. The depth of the wear scar is estimated at 120 nm [Fig. 6(a)]; this value is similar to the thickness of the amorphous



S.M.M. Ramos et a!.: Structure, mechanical, and tribological properties of titanium implanted alumina

FIG. 5. Observation of low dose implanted polycrystalline a-alumina disk wear scars: (a) SEM micrograph inside the disk wear track. A very thin transfer layer is observable in some areas (arrows), (b) Optical micrograph. A few small wear particles have been pushed outside the sliding contact and are visible on both sides of the wear track (labels 1 and 2). The arrow indicates the direction of the disk motion.

layer. The wear of the amorphous layer is very fast and probably corresponds to the first variations of the friction coefficient. More wear particles are formed, leading to a greater amount of accumulated matter adherent to the sphere at the front of the contact and to the formation of a transfer film in the disk wear track. The latter is due to compaction of elemental wear particles, and its thickness can be estimated at 100 nm, as in the case of lubricated conditions.23 The transfer film thickness is uneven, and its appearance indicates a highly deformed material [Fig. 6(b)]. Our results have not shown any important relationship between the presence of this surface layer and the friction coefficient value. As-implanted sapphire. The wear scar of the sphere is still about 130 /xm, but the wear of the disk is very

FIG. 6. Observations of high dose implanted polycrystainne aalumina disk wear scars: (a) optical micrograph with a cross surface profile showing the depth of the wear track. The left part of the wear scar is covered by a transfer layer, (b) SEM micrograph of this transfer layer showing its heterogeneous structure.

low (less than 20 nm). No transfer film was formed in the disk wear track and only a few wear particles were present on both sides of the track (Fig. 7). This result indicates that there are different wear behaviors for single crystal and polycrystalline a-alumina. For the same high fluence of implantation, disk wear volumes are different whereas both surfaces are found to be of similar structure. IV. CONCLUSION In this work, we have shown that ion implantation of titanium into sapphire and polycrystalline (1-AI2O3


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FIG. 7. Optical observation of high dose implanted sapphire disk wear scar. Network of lines visible on the surface are very smooth scratches created during the polishing process of samples.

results in substantial changes in surface mechanical properties relative to the unimplanted state. For a low fluence implant of sapphire (150 keV, 1016 Ti+ • cm' 2 ), the lattice remains crystalline and the substitutional fraction of titanium is high («53%). From nanoindentation measurements an increase in hardness and Young's modulus of about 16% with respect to the values of unimplanted sapphire occurs near the maximum of the implantation depth profile. For a high fluence implant (150 keV, 1017 Ti+ • 2 cm" ), amorphization of the implanted layer occurred, and significant decreases in the values of the hardness («57%) and Young's modulus (30%) are observed. The change in the friction coefficient versus the number of cycles for both sapphire and polycrystalline samples is not very sensitive to the microstructure of the implanted layer. However, the wear behavior strongly depends on the implanted fluence and on the target structures. For polycrystalline alumina implanted at the low fluence, the wear test results are similar to those for the unimplanted sample, but at high fluence the wear is greater and the depth of the wear scar corresponds to the thickness of the amorphous layer. In the case of heavily implanted sapphire, the wear is very low compared to polycrystalline samples while the microstructural characteristics of the outer implanted surfaces appear to be similar. Further investigations are necessary to explain these differences in wear behavior.

1. C. J. McHargue, Structure Property Relationships in Surface Modified Ceramics, NATO ASI Series, edited by C. J. McHargue, R. Kossowsky, and W. O. Hofer (Kluwer, London, 1989), p. 253. 2. C.J. McHargue, Nucl. Instrum. Methods in Phys. Res. B 19/20, 787 (1987). 3. C.J. McHargue, C.G. Farlow, C.W. White, B.R. Appleton, J. M. Williams, P. S. Sklad, and P. Angelini, Application of Ion Plating and Ion Implantation to Materials, edited by R. Hochman (American Society for Metals, Metals Park, OH, 1986), p. 255. 4. C.J. McHargue, C.G. Farlow, C.W. White, J.M. Williams, P. Angelini, and G. M. Bergun, Mater. Sci. Eng. 69, 123 (1985). 5. T. Hioki, A. Itoh, S. Noda, H. Doi, J. Kawamoto, and O. Kamiagaito, Nucl. Instrum. Methods in Phys. Res. B 7/8, 521 (1985). 6. T. Hioki, A. Itoh, S. Noda, H. Doi, J. Kawamoto, and O. Kamiagaito, J. Mater. Sci. 21, 1321 (1986). 7. T. Hioki, A. Itoh, M. Ohkubo, S. Noda, H. Doi, and J. Kawamoto, Structure Property Relationships in Surface Modified Ceramics, NATO ASI Series, edited by C. J. McHargue, R. Kossowsky, and W. O. Hofer (Kluwer, London, 1989), p. 303. 8. P.J. Burnett and T.F. Page, J. Mater. Sci. 19, 845 (1984). 9. P. J. Burnett and T. F. Page, Plastic Deformation of Ceramic Materials, edited by R. C. Bradt and R. E. Tressler (Plenum Press, New York, 1984), p. 669. 10. S.J. Bull, "The Mechanical and Tribological Properties of Ion Implanted Ceramics", Thesis, University of Cambridge, 1987. 11. J. L. Loubet, M. Bauer, A. Tonck, and Ph. Kapsa, to be published. 12. Ph. Kapsa, M. Bauer, and D. Mazuyer, Proc. Int. Tribology Conference, edited by the Japanese Society of Tribologists, Nagoya, Japan, November 1990. 13. L. Romana, P. Thevenard, B. Canut, G. Massouras, R. Brenier, and M. Brunei, Nucl. Instrum. Methods B46, 94 (1990). 14. J. F. Ziegler, J. P. Biersack, and U. Littmark, Stopping Power and Ranges of Ions in Matter (Pergamon Press, Oxford, 1985), Vol. 1. 15. S. M. M. Ramos, B. Canut, L. Gea, L. Romana, J. C. Brusq, P. Thevenard, and M. Brunei, Nucl. Instrum. Methods in Phys. Res. (in press). 16. C. W. White, C. J. McHargue, P. S. Sklad, L. A. Boatner, and G. C. Farlow, Mater. Sci. Rep. 4, 41 (1989). 17. W.C. Oliver, C.J. McHargue, G.C. Farlow, and C.W. White, Rad. Eff. 25, 1-12 (1975). 18. S. S. Chiang, "Response of Solids to Elastic-Plastic Indentation and the Application to Adhesion Measurements", Ph.D. Thesis, Lawrence Berkeley, Lab., CA (1981). 19. K.L. Johnson, J. Mech. Phys. Solids 18, 115-126 (1970). 20. H. Czichos, S. Becker, and J. Lexow, Wear 135, 171-191 (1989). 21. S. Sasaki, Journal of J.S.L.E. 10, 2 1 - 2 6 (1989). 22. N. Wallbridge, D. Dowson, and E. W. Roberts, Proc. Int. Conf. on Wear of Materials, edited by K. C. Ludema, Reston, VA, April 1983. 23. A. Tonck, Ph. Kapsa, and J. Sabot, ASME Trans., Journal of Tribology 108, 117-122 (1986).

ACKNOWLEDGMENTS The authors thank SEPR Company for financial support. S. M. M. Ramos further acknowledges support given by CNPq, Brazilian Agency.

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