- Email: [email protected]
Microstructural changes in DLC films due to tribological contact
Surface and Coatings Technology 108–109 (1998) 385–390
Microstructural changes in DLC films due to tribological contact a, b a a a b J. Koskinen *, D. Schneider , H. Ronkainen , T. Muukkonen , S. Varjus , P. Burck , K. Holmberg a , H.-J. Scheibe b a
b
VTT Manufacturing Technology, P.O. Box 1703, FIN-1703, 02044 VTT, Finland ¨ Werkstoff- und Strahltechnik, Helmholtzstraße 20, 01069 Dresden, Germany Fraunhofer-Institut f ur
Abstract The hydrogenated (a-C:H) and hydrogen free (ta-C) films were tested in sliding wear tests. The tests were carried out in air at room temperature (RH 1565%). The coefficient of friction in the tests was in the range 0.03–0.12 for the a-C:H films and 0.17–0.28 for the ta-C films. The wear volume of the DLC films was sufficiently low to enable the direct comparison of the film properties of the worn film to the original film. The microstructure of the films after wear tests has been characterized by micro RAMAN, the mechanical properties were measured by nano indentation and the elastic modulus was measured with the laser-acoustic method. No change of the micro structure of the DLC films was observed. The observed reduction of the elastic modulus of the ta-C film is explained by the evolution of mechanical defects such as micro-cracks and cracking at the nodular defects typical in non filtered vacuum arc depositions 1998 Elsevier Science S.A. All rights reserved. Keywords: Diamond-like carbon; Tribology; Surface acoustic wave SAW
1. Introduction
2. Experimental
Diamond-like carbon (DLC) films have successfully been applied as wear resistant coatings in a growing number of applications. In the literature the tribological performance of the DLC films has been widely reported and for example the formation of graphitic tribolayers on the contact surfaces of DLC films and counter parts have been reported [1,2]. However, the microstructural characterization of the DLC films subjected to repeated tribological contact has been seldom reported. In vacuum arc deposition macroparticles interfere the growth of a smooth film causing accumulation of defects visible as spots in the grown film. These defects can be described as growing nodules within the film [3]. When high loads are applied on the film–substrate system, the failure of DLC films is due to cracking and delamination of the film leading to catastrophic failure of the system in particular at the initially defected locations of the film. Since the mechanical properties have an important effect on the tribological performance of the DLC films, the attempt in this paper was to evaluate the changes in the mechanical properties of the film microstructure due to repeated sliding contact.
2.1. Deposition techniques
*Corresponding author. Tel.: 1358-9-4565413; fax: 1358-9-463118; e-mail: [email protected]
Hydrogenated and non-hydrogenated DLC films were deposited on high speed steel (M2) discs, which were hardened to 760 HV. The discs were polished mechanically until a surface roughness R a of 0.02 mm. Prior to deposition the surfaces were cleaned by solvents in an ultrasonic bath. The non-hydrogenated diamond-like carbon films, often referred as tetrahedral amorphous carbon (ta-C) [4,5] films were grown by using pulsed vacuum arc. The pulsed vacuum arc-discharge source is mounted to a 150-l vacuum chamber evacuated to a base pressure of about 33 10 24 Pa by using an oil diffusion pump. The cathode is a cylindrical graphite electrode (diameter, 30 mm) and the anode is a copper cylinder (diameter, 150 mm). The arc is ignited with an electrical spark between ignition electrodes at a frequency of 3 Hz. The capacitor bank (2600 mF) is discharged yielding a current pulse with a maximum current of about 3 kA and a half width of 150 ms. The distance from the sample to the cathode was about 300 mm. The nominal deposition rate was about 0.3 nm / s corresponding to an accumulation of about 1.6310 15 at / cm 2 during each pulse. The substrate temperature during deposition was under 708C. The film thickness was 0.5 mm. The amorphous hydrogenated carbon films (a-C:H) were
0257-8972 / 98 / $ – see front matter 1998 Elsevier Science S.A. All rights reserved. PII: S0257-8972( 98 )00656-2
386
J. Koskinen et al. / Surface and Coatings Technology 108 – 109 (1998) 385 – 390
deposited in a capacitively coupled rf (13.56 MHz) plasma. The substrates were placed directly on the powered cathode. The bias voltage was 2670 V and the deposition temperature was in the region of 2508C. Methane was used as the source gas and the pressure was kept constant at 2 Pa during deposition. Prior to deposition the substrates were sputter cleaned in an argon atmosphere for 5 min (bias 2400 V, pressure 0.7 Pa). A thin (,100 nm) titanium intermediate layer was deposited on the substrates by evaporating titanium into an argon glow discharge to improve the adhesion of the a-C:H coating. The film thickness was 1.1 mm.
2.2. Tribological tests The tribological performance was assessed with pin-ondisc equipment in normal atmosphere (22628C, 1565% RH). The tests were carried out using polished alumina (a-Al 2 O 3 ) balls with a diameter of 10 mm. The surface roughness of the balls was R a 50.04 mm. The normal force was 10 N and the sliding speed 1 m / s. In order to obtain a 5-mm wide worn area required for the elastic modulus measurements, a set of 12 adjoining wear tracks were generated next to each other with the wear track diameter changing from 18 to 28 mm. The wear tests were performed for 25 000 revolutions in a random order to avoid the gradual changes in the worn area. Since the ta-C films caused severe wear of the counter part, the contact area of the alumina pin was changed for the ta-C films three times during the tests. After the change of the alumina ball the run-in wear of the ball was carried out outside the actual test area. Thus the contact area of the ball varied in the range f 50.6–1 mm during the tests, causing contact pressure in the range 35–12 MPa, respectively. In order to reach the same level of contact pressure for the a-C:H film, the run-in of the alumina ball was carried out against the ta-C film. In the case of a-C:H films the contact area of the alumina ball changed from f 50.6 to 0.65 mm causing contact pressure in the range 35–30 MPa, respectively.
frequency. Surface waves with higher frequencies are more influenced by the film. The dependence of the SAW velocity on the frequency represents the so-called dispersion curve. The shape of the curve depends on the film parameters, Young’s modulus, density, and film thickness. The basic principle of the method is to measure the dispersion curve and fit a theoretical curve for deducing film parameters such as Young’s modulus. The theoretical dispersion curve can be derived from the equation of elastic wave motion and the boundary conditions for traction forces and displacement. Details of the method are described elsewhere [6]. Fig. 1 shows the schematic representation of the test equipment. Short laser pulses (pulse duration, 0.5 ns; energy, 0.4 mJ) of a nitrogen laser are generated. A cylindrical lens focuses the laser beam on the specimen surface. The short-time local heating within the laser focus line launches the acoustic impulse on the surface of the test piece. The surface wave impulse is detected by a wideband piezoelectric transducer (bandwidth, 200 MHz). Specimen and transducer are fixed to a translation stage that moves perpendicular to the laser beam position to vary the distance between laser focus line and transducer. At different distances between laser focus line and transducer the surface acoustic waveform is detected. Fourier transforming the waveform yields the phase spectrum that enables the phase velocity to be determined in dependence on frequency (dispersion curve). The measuring results are fitted with a theoretical dispersion curve to obtain the film parameters. Young’s modulus of films with thickness less than 100 nm can be determined. In order to determine the effect of the wear to the elastic modulus the SAW measurements were performed in the 5-mm wide wear track (three measurements) and at a virgin untouched surface at the immediate vicinity of the wear track (three measurements). RAMAN spectroscopy represents an effective method of characterizing the bonding structure of DLC films [7,8].
2.3. Characterization The wear tracks were measured by using an optical profilometer (UBM). Laser-induced surface acoustic waves (SAW) were used to determine Young’s modulus of the film material. These waves propagate along the surface of the materials. The concentration of the wave motion near the surface makes the SAW technique very sensitive for characterization of films, even for films much thinner than the penetration depth of the surface wave. In a homogeneous isotropic material the SAW velocity is constant. However, in coated materials it depends on frequency, termed dispersion. The reason for this is that the SAW penetration depth decreases with increasing
Fig. 1. Schematic representation of the surface acoustic wave (SAW) equipment.
J. Koskinen et al. / Surface and Coatings Technology 108 – 109 (1998) 385 – 390
The monochromatic light of an incident laser beam is scattered by interaction with lattice vibrations (phonons). The frequency of the scattered light is partially shifted by the frequency of the scattering phonons allowing to characterize the microstructure. The investigations were performed with a Renishaw Imaging Microscope 2000 (with Ar ion laser, l5514 nm; resolution of the shift of the wave number, 1 cm 21 ). Hardness and elastic modulus measurements were carried out on the samples by using an instrumented nanohardness tester (NanoTest 550). A Berkovich indenter was used. The calibration of the diamond tip was done deriving the contact area from the indentations in a quartz sample with known mechanical properties [9]. Indentations with varying plastic penetration depths from 42 to 271 nm were done. Measurements were made at loads from 5 to 40 mN using a loading rate of 0.91 mN / s. The indentation curves were analyzed according to the method proposed by Oliver and Pharr [9].
3. Results and discussion The surface profilometer measurements indicate mild wear of the DLC films. The thickness reduction of the ta-C
387
film was below the detection limit of the surface profilometer and for the a-C:H film it was about 100 nm. For the ta-C film the change of local film thickness is reflected in the interference color of the film, which would indicate an average reduction of the film thickness due to wear of about 30 nm. These values are used in the SAW results. Inspection of the worn surface by optical microscope shows merely smoothening of the DLC layer. Pin holes and local isolated delaminated spots varying in diameter from less than 1 mm to over 10 mm can be observed. In scanning electron microscopic (SEM) investigations cracking or pullout of the nodular defects on the ta-C film were revealed. This produces porosity in the worn film. The friction coefficient varied in the range 0.17–0.28 for the ta-C films and in the range 0.03–0.12 for the a-C:H films. The ta-C films typically exhibited a high peak value in friction (0.4–0.6) during the first revolutions of the test. The wear of the Al 2 O 3 pin was significantly higher for the ta-C film. The wear area of the ball typically changed from 0.6 to 1 mm in diameter for the ta-C film and from 0.6 to 0.65 mm in diameter for the a-C:H film during the tests. In the Raman measurements of the ta-C film no change in the Raman spectra before and after wear was observed (Fig. 2). This indicates that no accumulation of graphitic
Fig. 2. Micro-Raman spectra of the ta-C film measured at three different locations: virgin surface, wear track, and a detached spot in the wear track.
388
J. Koskinen et al. / Surface and Coatings Technology 108 – 109 (1998) 385 – 390
material had occurred in the bulk of the film neither on the surface of the film. However, in the spectra measured in the vicinity of the delaminated spots graphitic material is evident. This might suggest that the tribolayer on the counter surface, the formation of which has earlier been reported, contains graphitic spices [10]. The SAW measurement results of the elastic modulus of the films are shown in Table 1. The results have been obtained by using the measured average reduction of film thickness due to wear. The value of the elastic modulus varied at different locations of the disk which may result from the uncertainty of the film’s precise film thickness (non-uniform film thickness on the sample). On the other hand there is uncertainty in the local reduction of the film thickness caused by wear. Nevertheless, a small reduction in the elastic modulus of the DLC films can be observed. The error bars of the measured elastic modulus values reflect the repeatability of the SAW measurements at the particular measuring locations provided that the film parameters such as film thickness and density are known exactly. The effect of the uncertainty of these parameters is minimized by using relative measurements (DE at adjacent locations). Since the Raman results indicated no change of the microstructure (bonding) of ta-C caused by the wear experiment, the reduction of the film modulus may be attributed to the defects in the worn film. It may be assumed that the mechanical defects are micro-cracks and local detachments of the defects in the film (nodules). It is well-known that micro-cracks, porosity and delamination reduce the stiffness of the film that is quantified by the Young’s modulus. Studies at plasma-sprayed ZrO 2 films which contained long-stretched pores (porosity, about 5%) show a reduction of elastic modulus by about 80% compared to the defect-free bulk material [11]. Similarly TiN films with reduced adhesion to steel substrate were studied with the laser-acoustic technique. It was found that the elasticity was reduced by up to 8%, which was attributed to the higher defect density in the interface [12]. The effect of defects on measured elasticity can be estimated by the theory of Kreher and Pompe [13]. Fig. 3 shows a theoretical curve of Young’s modulus of DLC depending on the volume fraction of defects. The defects
were assumed to be disc-like pores with an aspect ratio of 100. The Figure illustrates the magnitude of the modulus reduction caused by the mechanical defects. The model enables to estimate the volume fraction of the defects in the film after the wear treatment as deduced from the measured film modulus. Of course, the theory does not consider all aspects of mechanical degradation of the film, but it demonstrates a means for non-destructive evaluation of the wear state. The nano-hardness measurements indicate a small reduction of the film hardness after the wear experiment as can be seen from Fig. 4. The difference of the hardness values is in the order of 10% with a plastic depth of about 50 nm. The reduction of hardness is of the same order as observed with the SAW measurements. However, identical values with less than 3% difference (maximum 300 GPa) were measured for the elastic modulus in the ta-C film both before and after wear. The elastic modulus values obtained from nano indentation of the hard DLC films are clearly lower than by using SAW [11]. This is attributed to the fact that the elastic response of the coated sample is dominated by the substrate in the nano-indentation test. Thus the small variations of the film elastic modulus were not detectable in indentation measurements.
4. Conclusions It can be concluded that based on RAMAN measurements no microstructural change in the bond structure of the DLC is caused by a dry sliding of a ceramic ball in such conditions which yet caused a noticeable wear of the DLC and the counter body. The elastic modulus of the ta-C films was found to be reduced slightly by the wear in dry sliding. Since Raman spectroscopy did not show variation in the DLC bond structure, the degradation of the mechanical properties (elastic modulus and hardness) was attributed to mechanical defects in form of cracking at the nodular defects and possible micro cracks not easily observed by microscopy. In the case of a-C:H films, such a reduction of the elastic modulus induced by wear could not be detected. This effect is within the error limit of the applied measuring method. Improving the accuracy of the
Table 1 The effect of pin-on-disk wear to the elastic modulus values measured by SAW for ta-C and a-C:H films Film, measurement no.
ta-C, 1 ta-C, 2 ta-C, 3 ta-C, average a-C:H, 1 a-C:H, 2 a-C:H, 3 a-C:H, average
DE (reduction of E) (GPa)
E (GPa) Original
After wear
59865 64865 67565 640 7962.5 8962.5 8162.5 83
61065 61565 59765 610 7462.5 8362.5 7762.5 78
212610 33610 78610 33 565 665 465 5
J. Koskinen et al. / Surface and Coatings Technology 108 – 109 (1998) 385 – 390
389
Fig. 3. Theoretical dependence of DLC film elastic modulus to the film on the defect concentration.
Fig. 4. Nano-hardness measurement of the ta-C film. The lines are least square fits of the measurement points. The circles and the solid line refer to the virgin ta-C, and the stars and the dashed line refer to the worn ta-C.
390
J. Koskinen et al. / Surface and Coatings Technology 108 – 109 (1998) 385 – 390
measurement of the local film thickness would enable this effect to be detected by the SAW method. In the case of the ta-C films mechanical defects such as micro cracks are responsible for the apparent reduction of the elastic modulus (and hardness) values. More work is required in order to quantify the magnitude and quality of the defects but the SAW measurements provided an initial estimate for the defect content of about 0.01% of the volume of the ta-C film.
Acknowledgements The authors are indebted to Tom E. Gustafsson for the SEM analysis. The financial support of the Technical Development Centre in Finland (TEKES) and Finnish Academy are gratefully acknowledged.
References [1] A. Erdemir, C. Bindal, J. Pagan, P. Wilbur, Surf. Coat. Technol. 76–77 (1995) 559–563.
[2] Y. Liu, A. Erdemir, E.J. Meletis, Surf. Coat. Technol. 82 (1996) 48–56. [3] D. Drescher, J. Koskinen, H.-J. Scheibe, A. Mensch, Diamond Rel. Mater. in press. [4] D. McKenzie, Y. Tin, N. Marks, C. Davis, B. Pailthrope, G. Amaratunga, V. Veerasamy, Diamond Rel. Mater. 3 (1994) 353–360. [5] J.-P. Hirvonen, J. Koskinen, M. Kaukonen, R. Nieminen, H-J. Scheibe, J. Appl. Phys., 81, 1997, 7248–7254; J. Koskinen, J.-P. ¨ Hirvonen, J. Keranen, J. Appl. Phys., in press. [6] D. Schneider, Th. Schwarz, Surf. Coat. Technol. 91 (1997) 136. [7] S. Prawer, K.W. Nugnet, Y. Lifshitz, G.D. Lempert, E. Grossman, J. Kulik, I. Avigal, R. Kalish, Diamond Rel. Mater. 5 (1996) 433. [8] M.A. Tamor, W.C. Vassell, J. Appl. Phys. 76 (1994) 3823. [9] D. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564–1583. [10] H. Ronkainen, J. Koskinen, J. Likonen, S. Varjus, J. Vihersalo, Diamond Rel. Mater. (and Refs. therein) 3 (1994) 1329–1336. ¨ [11] D. Schneider, Th. Schwarz, H.-P. Buchkremer, D. Stover, Thin Solid Films 224 (1993) 177–183. [12] D. Schneider, H. Ollendorf, Th. Schwarz, Appl. Phys. A61 (1995) 2779. [13] W. Kreher, W. Pompe, Internal Stresses in Heterogeneous Solids, Physical Research, vol. 9, ch. 4.3, Akadamie-Verlag, Berlin, 1989.
Microstructural changes in DLC films due to tribological contact a, b a a a b J. Koskinen *, D. Schneider , H. Ronkainen , T. Muukkonen , S. Varjus , P. Burck , K. Holmberg a , H.-J. Scheibe b a
b
VTT Manufacturing Technology, P.O. Box 1703, FIN-1703, 02044 VTT, Finland ¨ Werkstoff- und Strahltechnik, Helmholtzstraße 20, 01069 Dresden, Germany Fraunhofer-Institut f ur
Abstract The hydrogenated (a-C:H) and hydrogen free (ta-C) films were tested in sliding wear tests. The tests were carried out in air at room temperature (RH 1565%). The coefficient of friction in the tests was in the range 0.03–0.12 for the a-C:H films and 0.17–0.28 for the ta-C films. The wear volume of the DLC films was sufficiently low to enable the direct comparison of the film properties of the worn film to the original film. The microstructure of the films after wear tests has been characterized by micro RAMAN, the mechanical properties were measured by nano indentation and the elastic modulus was measured with the laser-acoustic method. No change of the micro structure of the DLC films was observed. The observed reduction of the elastic modulus of the ta-C film is explained by the evolution of mechanical defects such as micro-cracks and cracking at the nodular defects typical in non filtered vacuum arc depositions 1998 Elsevier Science S.A. All rights reserved. Keywords: Diamond-like carbon; Tribology; Surface acoustic wave SAW
1. Introduction
2. Experimental
Diamond-like carbon (DLC) films have successfully been applied as wear resistant coatings in a growing number of applications. In the literature the tribological performance of the DLC films has been widely reported and for example the formation of graphitic tribolayers on the contact surfaces of DLC films and counter parts have been reported [1,2]. However, the microstructural characterization of the DLC films subjected to repeated tribological contact has been seldom reported. In vacuum arc deposition macroparticles interfere the growth of a smooth film causing accumulation of defects visible as spots in the grown film. These defects can be described as growing nodules within the film [3]. When high loads are applied on the film–substrate system, the failure of DLC films is due to cracking and delamination of the film leading to catastrophic failure of the system in particular at the initially defected locations of the film. Since the mechanical properties have an important effect on the tribological performance of the DLC films, the attempt in this paper was to evaluate the changes in the mechanical properties of the film microstructure due to repeated sliding contact.
2.1. Deposition techniques
*Corresponding author. Tel.: 1358-9-4565413; fax: 1358-9-463118; e-mail: [email protected]
Hydrogenated and non-hydrogenated DLC films were deposited on high speed steel (M2) discs, which were hardened to 760 HV. The discs were polished mechanically until a surface roughness R a of 0.02 mm. Prior to deposition the surfaces were cleaned by solvents in an ultrasonic bath. The non-hydrogenated diamond-like carbon films, often referred as tetrahedral amorphous carbon (ta-C) [4,5] films were grown by using pulsed vacuum arc. The pulsed vacuum arc-discharge source is mounted to a 150-l vacuum chamber evacuated to a base pressure of about 33 10 24 Pa by using an oil diffusion pump. The cathode is a cylindrical graphite electrode (diameter, 30 mm) and the anode is a copper cylinder (diameter, 150 mm). The arc is ignited with an electrical spark between ignition electrodes at a frequency of 3 Hz. The capacitor bank (2600 mF) is discharged yielding a current pulse with a maximum current of about 3 kA and a half width of 150 ms. The distance from the sample to the cathode was about 300 mm. The nominal deposition rate was about 0.3 nm / s corresponding to an accumulation of about 1.6310 15 at / cm 2 during each pulse. The substrate temperature during deposition was under 708C. The film thickness was 0.5 mm. The amorphous hydrogenated carbon films (a-C:H) were
0257-8972 / 98 / $ – see front matter 1998 Elsevier Science S.A. All rights reserved. PII: S0257-8972( 98 )00656-2
386
J. Koskinen et al. / Surface and Coatings Technology 108 – 109 (1998) 385 – 390
deposited in a capacitively coupled rf (13.56 MHz) plasma. The substrates were placed directly on the powered cathode. The bias voltage was 2670 V and the deposition temperature was in the region of 2508C. Methane was used as the source gas and the pressure was kept constant at 2 Pa during deposition. Prior to deposition the substrates were sputter cleaned in an argon atmosphere for 5 min (bias 2400 V, pressure 0.7 Pa). A thin (,100 nm) titanium intermediate layer was deposited on the substrates by evaporating titanium into an argon glow discharge to improve the adhesion of the a-C:H coating. The film thickness was 1.1 mm.
2.2. Tribological tests The tribological performance was assessed with pin-ondisc equipment in normal atmosphere (22628C, 1565% RH). The tests were carried out using polished alumina (a-Al 2 O 3 ) balls with a diameter of 10 mm. The surface roughness of the balls was R a 50.04 mm. The normal force was 10 N and the sliding speed 1 m / s. In order to obtain a 5-mm wide worn area required for the elastic modulus measurements, a set of 12 adjoining wear tracks were generated next to each other with the wear track diameter changing from 18 to 28 mm. The wear tests were performed for 25 000 revolutions in a random order to avoid the gradual changes in the worn area. Since the ta-C films caused severe wear of the counter part, the contact area of the alumina pin was changed for the ta-C films three times during the tests. After the change of the alumina ball the run-in wear of the ball was carried out outside the actual test area. Thus the contact area of the ball varied in the range f 50.6–1 mm during the tests, causing contact pressure in the range 35–12 MPa, respectively. In order to reach the same level of contact pressure for the a-C:H film, the run-in of the alumina ball was carried out against the ta-C film. In the case of a-C:H films the contact area of the alumina ball changed from f 50.6 to 0.65 mm causing contact pressure in the range 35–30 MPa, respectively.
frequency. Surface waves with higher frequencies are more influenced by the film. The dependence of the SAW velocity on the frequency represents the so-called dispersion curve. The shape of the curve depends on the film parameters, Young’s modulus, density, and film thickness. The basic principle of the method is to measure the dispersion curve and fit a theoretical curve for deducing film parameters such as Young’s modulus. The theoretical dispersion curve can be derived from the equation of elastic wave motion and the boundary conditions for traction forces and displacement. Details of the method are described elsewhere [6]. Fig. 1 shows the schematic representation of the test equipment. Short laser pulses (pulse duration, 0.5 ns; energy, 0.4 mJ) of a nitrogen laser are generated. A cylindrical lens focuses the laser beam on the specimen surface. The short-time local heating within the laser focus line launches the acoustic impulse on the surface of the test piece. The surface wave impulse is detected by a wideband piezoelectric transducer (bandwidth, 200 MHz). Specimen and transducer are fixed to a translation stage that moves perpendicular to the laser beam position to vary the distance between laser focus line and transducer. At different distances between laser focus line and transducer the surface acoustic waveform is detected. Fourier transforming the waveform yields the phase spectrum that enables the phase velocity to be determined in dependence on frequency (dispersion curve). The measuring results are fitted with a theoretical dispersion curve to obtain the film parameters. Young’s modulus of films with thickness less than 100 nm can be determined. In order to determine the effect of the wear to the elastic modulus the SAW measurements were performed in the 5-mm wide wear track (three measurements) and at a virgin untouched surface at the immediate vicinity of the wear track (three measurements). RAMAN spectroscopy represents an effective method of characterizing the bonding structure of DLC films [7,8].
2.3. Characterization The wear tracks were measured by using an optical profilometer (UBM). Laser-induced surface acoustic waves (SAW) were used to determine Young’s modulus of the film material. These waves propagate along the surface of the materials. The concentration of the wave motion near the surface makes the SAW technique very sensitive for characterization of films, even for films much thinner than the penetration depth of the surface wave. In a homogeneous isotropic material the SAW velocity is constant. However, in coated materials it depends on frequency, termed dispersion. The reason for this is that the SAW penetration depth decreases with increasing
Fig. 1. Schematic representation of the surface acoustic wave (SAW) equipment.
J. Koskinen et al. / Surface and Coatings Technology 108 – 109 (1998) 385 – 390
The monochromatic light of an incident laser beam is scattered by interaction with lattice vibrations (phonons). The frequency of the scattered light is partially shifted by the frequency of the scattering phonons allowing to characterize the microstructure. The investigations were performed with a Renishaw Imaging Microscope 2000 (with Ar ion laser, l5514 nm; resolution of the shift of the wave number, 1 cm 21 ). Hardness and elastic modulus measurements were carried out on the samples by using an instrumented nanohardness tester (NanoTest 550). A Berkovich indenter was used. The calibration of the diamond tip was done deriving the contact area from the indentations in a quartz sample with known mechanical properties [9]. Indentations with varying plastic penetration depths from 42 to 271 nm were done. Measurements were made at loads from 5 to 40 mN using a loading rate of 0.91 mN / s. The indentation curves were analyzed according to the method proposed by Oliver and Pharr [9].
3. Results and discussion The surface profilometer measurements indicate mild wear of the DLC films. The thickness reduction of the ta-C
387
film was below the detection limit of the surface profilometer and for the a-C:H film it was about 100 nm. For the ta-C film the change of local film thickness is reflected in the interference color of the film, which would indicate an average reduction of the film thickness due to wear of about 30 nm. These values are used in the SAW results. Inspection of the worn surface by optical microscope shows merely smoothening of the DLC layer. Pin holes and local isolated delaminated spots varying in diameter from less than 1 mm to over 10 mm can be observed. In scanning electron microscopic (SEM) investigations cracking or pullout of the nodular defects on the ta-C film were revealed. This produces porosity in the worn film. The friction coefficient varied in the range 0.17–0.28 for the ta-C films and in the range 0.03–0.12 for the a-C:H films. The ta-C films typically exhibited a high peak value in friction (0.4–0.6) during the first revolutions of the test. The wear of the Al 2 O 3 pin was significantly higher for the ta-C film. The wear area of the ball typically changed from 0.6 to 1 mm in diameter for the ta-C film and from 0.6 to 0.65 mm in diameter for the a-C:H film during the tests. In the Raman measurements of the ta-C film no change in the Raman spectra before and after wear was observed (Fig. 2). This indicates that no accumulation of graphitic
Fig. 2. Micro-Raman spectra of the ta-C film measured at three different locations: virgin surface, wear track, and a detached spot in the wear track.
388
J. Koskinen et al. / Surface and Coatings Technology 108 – 109 (1998) 385 – 390
material had occurred in the bulk of the film neither on the surface of the film. However, in the spectra measured in the vicinity of the delaminated spots graphitic material is evident. This might suggest that the tribolayer on the counter surface, the formation of which has earlier been reported, contains graphitic spices [10]. The SAW measurement results of the elastic modulus of the films are shown in Table 1. The results have been obtained by using the measured average reduction of film thickness due to wear. The value of the elastic modulus varied at different locations of the disk which may result from the uncertainty of the film’s precise film thickness (non-uniform film thickness on the sample). On the other hand there is uncertainty in the local reduction of the film thickness caused by wear. Nevertheless, a small reduction in the elastic modulus of the DLC films can be observed. The error bars of the measured elastic modulus values reflect the repeatability of the SAW measurements at the particular measuring locations provided that the film parameters such as film thickness and density are known exactly. The effect of the uncertainty of these parameters is minimized by using relative measurements (DE at adjacent locations). Since the Raman results indicated no change of the microstructure (bonding) of ta-C caused by the wear experiment, the reduction of the film modulus may be attributed to the defects in the worn film. It may be assumed that the mechanical defects are micro-cracks and local detachments of the defects in the film (nodules). It is well-known that micro-cracks, porosity and delamination reduce the stiffness of the film that is quantified by the Young’s modulus. Studies at plasma-sprayed ZrO 2 films which contained long-stretched pores (porosity, about 5%) show a reduction of elastic modulus by about 80% compared to the defect-free bulk material [11]. Similarly TiN films with reduced adhesion to steel substrate were studied with the laser-acoustic technique. It was found that the elasticity was reduced by up to 8%, which was attributed to the higher defect density in the interface [12]. The effect of defects on measured elasticity can be estimated by the theory of Kreher and Pompe [13]. Fig. 3 shows a theoretical curve of Young’s modulus of DLC depending on the volume fraction of defects. The defects
were assumed to be disc-like pores with an aspect ratio of 100. The Figure illustrates the magnitude of the modulus reduction caused by the mechanical defects. The model enables to estimate the volume fraction of the defects in the film after the wear treatment as deduced from the measured film modulus. Of course, the theory does not consider all aspects of mechanical degradation of the film, but it demonstrates a means for non-destructive evaluation of the wear state. The nano-hardness measurements indicate a small reduction of the film hardness after the wear experiment as can be seen from Fig. 4. The difference of the hardness values is in the order of 10% with a plastic depth of about 50 nm. The reduction of hardness is of the same order as observed with the SAW measurements. However, identical values with less than 3% difference (maximum 300 GPa) were measured for the elastic modulus in the ta-C film both before and after wear. The elastic modulus values obtained from nano indentation of the hard DLC films are clearly lower than by using SAW [11]. This is attributed to the fact that the elastic response of the coated sample is dominated by the substrate in the nano-indentation test. Thus the small variations of the film elastic modulus were not detectable in indentation measurements.
4. Conclusions It can be concluded that based on RAMAN measurements no microstructural change in the bond structure of the DLC is caused by a dry sliding of a ceramic ball in such conditions which yet caused a noticeable wear of the DLC and the counter body. The elastic modulus of the ta-C films was found to be reduced slightly by the wear in dry sliding. Since Raman spectroscopy did not show variation in the DLC bond structure, the degradation of the mechanical properties (elastic modulus and hardness) was attributed to mechanical defects in form of cracking at the nodular defects and possible micro cracks not easily observed by microscopy. In the case of a-C:H films, such a reduction of the elastic modulus induced by wear could not be detected. This effect is within the error limit of the applied measuring method. Improving the accuracy of the
Table 1 The effect of pin-on-disk wear to the elastic modulus values measured by SAW for ta-C and a-C:H films Film, measurement no.
ta-C, 1 ta-C, 2 ta-C, 3 ta-C, average a-C:H, 1 a-C:H, 2 a-C:H, 3 a-C:H, average
DE (reduction of E) (GPa)
E (GPa) Original
After wear
59865 64865 67565 640 7962.5 8962.5 8162.5 83
61065 61565 59765 610 7462.5 8362.5 7762.5 78
212610 33610 78610 33 565 665 465 5
J. Koskinen et al. / Surface and Coatings Technology 108 – 109 (1998) 385 – 390
389
Fig. 3. Theoretical dependence of DLC film elastic modulus to the film on the defect concentration.
Fig. 4. Nano-hardness measurement of the ta-C film. The lines are least square fits of the measurement points. The circles and the solid line refer to the virgin ta-C, and the stars and the dashed line refer to the worn ta-C.
390
J. Koskinen et al. / Surface and Coatings Technology 108 – 109 (1998) 385 – 390
measurement of the local film thickness would enable this effect to be detected by the SAW method. In the case of the ta-C films mechanical defects such as micro cracks are responsible for the apparent reduction of the elastic modulus (and hardness) values. More work is required in order to quantify the magnitude and quality of the defects but the SAW measurements provided an initial estimate for the defect content of about 0.01% of the volume of the ta-C film.
Acknowledgements The authors are indebted to Tom E. Gustafsson for the SEM analysis. The financial support of the Technical Development Centre in Finland (TEKES) and Finnish Academy are gratefully acknowledged.
References [1] A. Erdemir, C. Bindal, J. Pagan, P. Wilbur, Surf. Coat. Technol. 76–77 (1995) 559–563.
[2] Y. Liu, A. Erdemir, E.J. Meletis, Surf. Coat. Technol. 82 (1996) 48–56. [3] D. Drescher, J. Koskinen, H.-J. Scheibe, A. Mensch, Diamond Rel. Mater. in press. [4] D. McKenzie, Y. Tin, N. Marks, C. Davis, B. Pailthrope, G. Amaratunga, V. Veerasamy, Diamond Rel. Mater. 3 (1994) 353–360. [5] J.-P. Hirvonen, J. Koskinen, M. Kaukonen, R. Nieminen, H-J. Scheibe, J. Appl. Phys., 81, 1997, 7248–7254; J. Koskinen, J.-P. ¨ Hirvonen, J. Keranen, J. Appl. Phys., in press. [6] D. Schneider, Th. Schwarz, Surf. Coat. Technol. 91 (1997) 136. [7] S. Prawer, K.W. Nugnet, Y. Lifshitz, G.D. Lempert, E. Grossman, J. Kulik, I. Avigal, R. Kalish, Diamond Rel. Mater. 5 (1996) 433. [8] M.A. Tamor, W.C. Vassell, J. Appl. Phys. 76 (1994) 3823. [9] D. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564–1583. [10] H. Ronkainen, J. Koskinen, J. Likonen, S. Varjus, J. Vihersalo, Diamond Rel. Mater. (and Refs. therein) 3 (1994) 1329–1336. ¨ [11] D. Schneider, Th. Schwarz, H.-P. Buchkremer, D. Stover, Thin Solid Films 224 (1993) 177–183. [12] D. Schneider, H. Ollendorf, Th. Schwarz, Appl. Phys. A61 (1995) 2779. [13] W. Kreher, W. Pompe, Internal Stresses in Heterogeneous Solids, Physical Research, vol. 9, ch. 4.3, Akadamie-Verlag, Berlin, 1989.