Phenomena in microwear experiments on metal-free and metal-containing diamond-like carbon coatings: friction, wear, fatigue and plastic deformation

Surface and Coatings Technology 177 – 178 (2004) 453–458

Phenomena in microwear experiments on metal-free and metal-containing diamond-like carbon coatings: friction, wear, fatigue and plastic deformation Kirsten Ingolf Schiffmann* ¨ Schicht- und Oberflachentechnik, ¨ Fraunhofer Institute fur Bienroder Weg 54E, Braunschweig 38108, Germany

Abstract Atomic force microscopy based microwear experiments allow deeper understanding of mechanisms of wear on the microscopic scale due to a well defined single asperity contact, elimination of roughness effects, and the high lateral and vertical resolution in the micro- and nanometer range. In this paper, results of linear oscillating microwear experiments on metal containing diamond like carbon (DLC) coatings, thin pure DLC-coatings, Si- and Si:O-doped DLC-coatings will be presented. The experiments have been performed using either a diamond tipped cantilever-based system or an electrostatic transducer system (Hysitron Inc.) in conjunction with a standard AFM. Diamond tip radii of less than or equal to 1 mm and loads in the range of some millinewtons lead to contact areas of only 0.1–0.2 mm2 and contact pressures in the range 2–20 GPa. Under these conditions, for metal-DLC coatings (e.g. W-DLC) material fatigue on a nanometer scale can directly be observed and identified as an important wear mechanism. Furthermore, the columnar growth structure of the film and percolation of the metallic nano-particles inside the film, i.e. metal content and metal type, strongly influence the fatigue and wear resistance of the coatings. Oscillating microwear experiments on thin DLC, Si-DLC and Si:O-DLC coatings on glass substrates are analysed with regard to friction, wear and plastic deformation. The friction coefficient m can be understood in terms of a combination of Hertzian elastic contact and an additional ploughing term msaLy1y3 qbLn (Lsload) where n strongly changes during the first wear cycles. Comparison of residual wear depth and residual indentation depth shows that the wear volume may completely be due to plastic deformation and not to a real material loss. In other cases, material loss and plastic deformation both contribute to the observed wear volume. Therefore, evaluating only residual wear marks may lead to a misinterpretation of nanowear results. 䊚 2003 Elsevier Science B.V. All rights reserved. Keywords: Diamond like carbon; Microwear; Friction; Fatigue; Plastic deformation; Atomic force microscopy

1. Introduction In the past years, microtribology has become an increasing field in science w1x. The main driving force for this evolution was the development of ultra thin protective coatings for the hard disk industry, which needed to be characterized concerning their friction and wear properties by new tribological techniques w2x. Meanwhile, microtribological tests have been found useful in a much wider field of applications w3x, due to the new possibilities and benefits of this technique. In the following context microtribological test will mean: make use of a well defined single asperity contact (area -1 mm2) and low loads (-10 mN) to perform *Tel.: q49-531-2155577; fax: q49-531-2155906. E-mail address: [email protected] (K.I. Schiffmann).

tribological tests and the analysis of these tests by insitu or ex-situ high-resolution topographic imaging. The benefits of these type of tests are: (a) The possibility to characterise ultra thin coatings due to the high vertical resolution and low loads; (b) characterisation on small areas, as micro components, small material phases or single grains, due to the high lateral resolution of the method; (c) contributions to the basic understanding of the mechanisms of friction and wear due to the controlled single asperity contact geometry; and (d) better understanding of the influence of the material microstructure on friction and wear due to the high lateral resolution. In the following some examples will be shown where microtribological tests have given new insight into wear phenomena and wear mechanisms of DLC-based coatings.

0257-8972/04/$ - see front matter 䊚 2003 Elsevier Science B.V. All rights reserved. doi:10.1016/j.surfcoat.2003.08.064

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2. Experiment Metal containing DLC coatings (Me-DLC) have been prepared on silicon substrates in a combined CVD-PVD process by sputtering from a metallic target in an argon– hydrocarbon plasma. The coating thickness was in the range of 1 mm, and tungsten (W-DLC) and gold (AuDLC) have been used as metallic components. The metal content has been varied between 0 and 50 at.%. The coatings consist of nanocrystalline metal clusters, which are embedded in the amorphous diamond like carbon matrix. More details about the deposition process and the coating structure can be found in Ref. w4x. Pure DLC-films, Si-DLC and Si:O-DLC films have been prepared in a RF plasma process at a pressure of 1.5 Pa using a gas mixture of argon and either 10% acetylene, 40% tetramethylsilane (TMS; Si(CH3)4 ) or 40% hexamethyldisiloxane (HMDSO, ((CH3)3Si)2O), respectively. A substrate bias of 550–700 V has been used and a power between 420 and 580 W has been applied. The coatings have been prepared on glass substrates with coating thickness of only 50 nm (DLC), 320 nm (Si-DLC) and 660 nm (Si:O-DLC). Besides this, also the pure glass substrate has been investigated for reference. The Si content of Si-DLC coating was 33%. The Si:O-DLC coating contains 28% Si and 18% O. There is no evidence for the existence of nanoclusters as in the case of Me-DLC. Si and Si:O seem to be distributed homogeneously in the DLC matrix. The linear oscillating wear experiments on the MeDLC coatings have been performed using a standard AFM with a stainless steel cantilever and pyramidal diamond tip (1 mm tip radius) w1x. Up to 256 linear reciprocating wear cycles have been used to produce wear scars which subsequently are imaged by the same tip under minimum load giving the wear volume as a function of test parameters (load, number of cycles). Also, the relative changes in depth during the wear process (i.e. tip depth under load) were measured by imaging during the wear process, giving information about time evolution of the wear process. The friction could not be measured with this system. The experiments on DLC, Si-DLC and Si:O-DLC have been performed using an electrostatic transducer system (Hysitron Inc.) in conjunction with a standard AFM. A conical diamond tip was used with a tip radius of 600 nm. Each experiment consists of (a) a prescan of the surface topography under minimum load, (b) 25 wear cycles under a fixed higher load and (c) a postscan under minimum load to determine the depth profile of the wear track. During the wear experiment the friction forces and the depth under load are measured, giving more information about time evolution of the wear process. The loads in all experiments were between 0.2 and 7 mN, leading to contact areas of only 0.1–0.2 mm2 and contact pressures in the range of 2–20 GPa.

Fig. 1. Evolution of the depth of wear tracks (under load) during 256 wear cycles for W-DLC coatings with (a) 7% W, (b) 9% W and (c) 12% W, showing periodic break-off of material.

3. Results and discussion 3.1. Me-DLC Fig. 1 shows the depth of the tip in the wear track during the 256 linear oscillating wear cycles (track length 5 mm, load 1.9 mN) for three W-DLC coatings with increasing W content. Beside the continuous removal of material by the tip, a periodic break-off of material can be observed whose frequency increases with increasing W content. This can be interpreted as a direct observation of material fatigue: The periodic loading of the surface by the back and forth moving tip leads to an accumulation of plastic deformation andyor densification of the material, increasing the subsurface stress in the coating. Reaching the critical shear stress

K.I. Schiffmann / Surface and Coatings Technology 177 – 178 (2004) 453–458

Fig. 2. Critical number of wear cycles for first break-off of material (a) as a function of load and (b) as a function of metal content at a load of 1.9 mN for selected W-DLC and Au-DLC coatings.

of the coating, microcracks are formed, finally leading to a sudden removal of material. Fig. 2 shows the critical number of wear cycles Nc for the first break-off of material as a function of load and as a function of metal content. As expected, Nc decreases with increasing load, since higher contact pressure reduces the number of cycles needed to reach the critical internal stress for crack formation. However, increasing the volume fraction of metallic nanoparticles embedded into the DLC matrix leads to a decrease of Nc, too. This indicates that at a load of 1.9 mN and low metal content the volume fraction of highly cross-linked DLC matrix is the decisive factor for the amount of material fatigue. But, for W-DLC this is only true below 15 at.% W. At higher metal content Nc becomes infinite (at 1.9 mN load), i.e. there occurs no material fatigue at all. This can be explained if the microstructure of the coating is taken into account. 15 at.% W corresponds to approximately 30 vol.% WC which is the percolation threshold for the tungsten carbide nanoparticles w5x, i.e. the nanoparticles begin to touch each other and grow into one another. This leads to a strong additional cross-linking within the coating due to the coalescence of carbidic particles and, therefore, leads to a reinforcement of the film structure. At higher loads of 3.2 mN material fatigue can be observed also above the percolation threshold, i.e. higher pressures

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are able to break the polycristalline particle network. On Au-DLC films no reinforcement of the film structure due to particle percolation could be detected. Nc approaches zero at approximately 14 at.% Au and severe wear is observed at higher metal contents. This is probably due to the low hardness and high plasticity of gold. Another microstructural factor has an influence on the wear behaviour of the coatings: Me-DLC coatings often have a well-known columnar growth structure. We clearly could observe this columnar structure at our WDLC films by cross-sectional AFM-imaging, indicating growth columns of approximately 200 to 300 nm in diameter. Using the technique of AFM-imaging during the wear process w7x results in the cycle-by-cycle depth profiling of the wear track shown in Fig. 3. The top end of growth columns can be seen as small hillock along the wear track and can be followed in time as vertical stripes in the image. After approximately 27 cycles, one first growth column breaks-off, visible as dark vertical stripe of approximately 300-nm width, while the neighbouring columns still persist. After 40 cycles other single columns break off, until after 88 cycles all columns finally have broken off. Thus, it can be concluded that the growth column interfaces are the weakest parts of the film structure where the material breaks first under cyclic loading. To finish with Me-DLC, Fig. 4 shows the wear volume of W-DLC (1.9 mN, 256 cycles) and Au-DLC (0.74 mN, 50 cycles) as a function metal content. One notes the different test conditions that have been chosen due to the strongly different wear resistance of W-DLC and Au-DLC. These curves can now be understood easily: For W-DLC up to 15% W, an increase of wear due to the decrease of volume fraction of highly crosslinked DLC-matrix is observed. Above 15% W the percolation of carbidic nanoparticles leads to a reinforcement of the film structure resulting in a strong reduction

Fig. 3. One dimensional depth profile of the wear track as a function of wear cycles for 7 at.% W-DLC coating, showing preferential breakoff of complete growth columns.

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modelled by Hertz theory of elastic contact w8x. Using Hertz contact area and the definition of the shear stress SsFyA where F is the friction force and A is the contact area leads to a friction coefficient of the form msFy LsaLy1y3, where L is the load. This equation represents well the lower load part of the curves. For the higher loads an additional ploughing term has to be added w9x, which describes the plastic deformation and wear above a critical load: msbLn. The whole curve can be well described by an expression of the form: msaLy1y3qbLn

Fig. 4. Volume of the wear track normalized by the length of the track (a) for W-DLC at 1.9 mN and 256 cycles and (b) for Au-DLC at 0.74 mN and 50 cycles.

of wear. Approaching 50% W, i.e. 100% tungsten carbide wear increases again due to the omission of the lubricating DLC matrix resulting in increased friction coefficients, which has been verified by macroscopic oscillating pin-on-disk tests. Likewise, for Au-DLC an increase in wear rate with increasing metal content can be observed which is attributed to the replacement of the wear resistant DLC by the soft gold. The percolation threshold for Au is at approximately 50 at.% w6x. No samples above the percolation threshold have been investigated, but it is believed that coalescence of Au-particles would not reinforce the film structure due to the low hardness and high plasticity of gold. For a more detailed and quantitative evaluation of microwear experiments on Me-DLC please refer to Ref. w4x.

(1)

where a and b are constants and the ploughing exponent n is in the range of 0.7–1.2. Switching to the 50th half cycle as expected the lower load regime does not change very much, but the ploughing exponent n has reduced significantly to 0.1–0.6. This can be understood by the fact that at the end of the experiment the tip runs in a predefined groove and it has not to push as much material away as in the first cycle. This is confirmed by the time evolution of friction coefficient which in the case of lower loads shows a constant friction coefficient from the first up to the last cycle, while for higher loads

3.2. DLC, Si-DLC and Si:O-DLC Fifty nanometres of DLC, 320 nm Si-DLC and 660 nm Si:O-DLC on glass and the pure glass substrate have been investigated by 25 back and forth cycles under different constant loads. First it will be focused on the friction coefficient as a function of load which is shown in Fig. 5 for the first half cycle and for the 50th half cycle of the wear test. During the first half cycle, i.e. when the tip runs on a virgin surface, two load regimes can be distinguished: in the lower load regime a decrease in friction with increasing load can be observed, while in the higher load regime the friction coefficient increases again with load. This can be understood in terms of elastic and plastic deformation: in the lower load regime the surface is only elastically deformed. Elastic deformation of a flat surface by the spherical apex of the tip can be

Fig. 5. Friction coefficient of three coatings on glass and the uncoated glass as a function of load during (a) the first half cycle and (b) the 50th half cycle. The full lines correspond to the fit with Eq. (1).

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mation during the wear process can be measured independently. The corresponding residual indentation depths are indicated by open symbols. For the two thick coatings (Si-DLC and Si:O-DLC) at loads below 2–3 mN only a small material loss is found which does not increase with load. The increase of wear depth can completely be attributed to an increase in plastic deformation. Above 3 mN material loss starts to increase even though plastic deformation is still the dominant fraction. The thin DLC coating and the uncoated glass sample show completely different behaviour. At loads below 1 mN material removal dominates the residual wear depth and the amount of removed material increases with load. At higher loads of 2 to 3 mN, again, plastic deformation becomes more important and finally surpasses the material loss which nevertheless still increases. We can conclude that there are two components contributing to the wear depth in microwear experiments: one is plastic deformation, i.e. the material is only displaced but not removed from the surface. The second is real material removal, i.e. stripping off material from the surface. An overview of the different contributions for a fixed load of 3 mN is given in Fig. 7 showing the fraction of elastic deformation, plastic deformation and removed material. Beside the high fraction of elastic deformation, which recovers after unloading, there remains a fraction of plastic deformation of approximately 50% for DLC and uncoated glass and 67% for Si-DLC and Si:O-DLC. i.e. more than half of the measured wear depth can be attributed to material displacement and not to material removal. We found no evidence for material fatigue as in the case of Me-DLC coatings shown before. The critical number of wear Fig. 6. Residual depth of wear track (closed symbols) and residual depth of corresponding indentation experiments (open symbols) as a function of load for (a) 320 nm Si-DLC and 50 nm DLC and (b) 620 nm Si:O-DLC and uncoated glass indicating real material loss (shaded grey) and plastic deformation (white area below the curves).

it starts with a high friction coefficient in the first cycles which decreases during the first 1 to 3 cycles and then remains more or less constant. In the following section the wear depth created during the oscillating sliding experiment will be analysed. Fig. 6 shows the residual wear depth of the four respective materials as a function of load (closed symbols). To get insight into the significance of plastic deformation during the wear process all experiments have been performed with and without lateral movement of the tip. The latter is essentially equal to making nanoindentation with the sphero–conical tip, using a certain hold time at maximum load. In nanoindentation, the shear forces due to the lateral movement of the tip are avoided. Therefore, no material loss, i.e. no wear will occur and thus a lower limit for the contribution of plastic defor-

Fig. 7. Fraction of elastic deformation, plastic deformation and material loss at a load of 3 mN.

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cycles is either higher than 25 cycles or no material fatigue but only continuous material removal occurs in the case of these DLC derivates. 4. Conclusion Microtribological experiments may give new insight into mechanisms of wear processes. For Me-DLC coatings besides continuous removal of material, low cycle fatigue could be observed as an important wear mechanism. The fatigue can be suppressed by particle percolation in W-DLC coatings and the growth column interfaces could be identified as preferential fracture planes during the fatigue wear. Beside these effects the wear resistance of Me-DLC seems to be mainly determined by the volume fraction of the highly cross-linked DLC matrix. Experiments on DLC, Si-DLC and Si:ODLC showed that the apparent wear depth in microwear experiments can be attributed to more than 50% to plastic deformation, i.e. material displacement, and only a smaller fraction is due to real material loss. The plastic deformation mainly occurs during the first one to three cycles, leading to higher friction coefficients during this phase, while afterwards the tip runs in a predefined wear track and only minor ploughing occurs.

Acknowledgments The author would like to acknowledge A. Hieke for preparation of Si-, Si:O- and pure DLC coatings, H. ¨ Hubsch and R. Thyen for the preparation of Me-DLC coatings. References w1x B. Bhushan, Wear 225–229 (1999) 465–492. w2x B. Bhushan, Diamond Relat. Mater. 8 (1999) 1985–2015. w3x B. Bhushan (Ed.), Tribology Issues and Opportunities in MEMS, Kluwer Academic Publishers, London, 1998. w4x K.I. Schiffmann, M. Fryda, G. Goerigk, R. Lauer, P. Hinze, A. Bulack, Thin Solid Films 347 (1999) 60. w5x M. Wang, Research Center Julich, ¨ Report No. Jul-2595, ISSN 0366–0885, 1991, p. 45. w6x H. Koberle, ¨ ¨ Struktur und elektrische Leitfahigkeit von HFPlasma-erzeugten metallhaltigen Kohlenwasserstoffschichten, Thesis, University of Hamburg, 1989, p. 124. w7x J.L. Loubet, M. Belin, R. Durand, H. Pascal, Thin Solid Films 253 (1994) 194. w8x K.L. Johnson, Contact Mechanics, Cambridge University Press, Cambridge, 1987, p. 427. w9x F.P. Bowden, D. Tabor, The Friction and Lubrication of Solids, Oxford University Press, London, 1958, p. 90.