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Nano- and micro-scale wear of fluorinated carbon films
Surface and Coatings Technology 182 (2004) 335–341
Nano- and micro-scale wear of fluorinated carbon films P. Ayala1, M.E.H. Maia da Costa, R. Prioli, F.L. Freire Jr.* ´ ´ ´ Departamento de Fısica, Pontifıcia Universidade Catolica do Rio de Janeiro, Caixa Postal 38071, 22452-970 Rio de Janeiro, RJ, Brazil Received 19 May 2003; accepted in revised form 19 August 2003
Abstract Wear of fluorinated carbon films deposited by plasma assisted chemical vapor deposition has been studied by atomic force microscopy and the sliding sphere test. The films microstructure and chemical composition were also investigated. The films became polymeric upon fluorine incorporation, as revealed by density measurements, Raman and X-ray photoelectron spectroscopies. It was shown that the wear marks produced with a microscope diamond tip are strongly related to the film hardness and microstructure. The wear depth measured on a polymeric film is greater compared with the ones obtained from hard carbon and carbon fluorinated films. At nanometer scale, the minimum force necessary to scratch the films increases with the film hardness, while the wear rate decreases. Similar behaviors were observed in both nano- and micro-scale wear experiments. 䊚 2003 Elsevier B.V. All rights reserved. Keywords: Abrasive wheel test; Amorphous; Atomic force microscopy; Photoelectron spectroscopy; Plasma assisted chemical vapor deposition; Carbon
1. Introduction The tribological behavior of moving micro- and nanodevices is a challenging subject that impacts their reliabilities and operational lifetimes. The fundamental understanding of the friction and wear in nanometer scale is important to achieve such goals w1x. Amorphous hydrogenated carbon films (a-C:H) have been under investigation for a long time due to their properties such as high hardness, chemical inertness and low friction coefficients w2,3x. These films have found their way into industrial applications as tribological protective coatings. One of the main drawbacks of this application is the need of applying a fluorinated lubricant layer onto the surface of the a-C:H protective coatings in order to reduce the film wear w4x. Good candidates to substitute this lubricant layer are selflubricant films deposited by a plasma technique. Fluor´ *Corresponding author. Departamento de Fısica, PUC-Rio, Rua ˆ de Sao ˜ Vicente, 225, Caixa Postal 38071, 22453-970 Rio Marques de Janeiro, RJ, Brazil. Tel.: q55-21-3114-1272; fax: q55-21-31141040. E-mail address: [email protected] (F.L. Freire Jr.). 1 Present address: Departamento de Fisica, Escuela Politecnica Nacional, Quito, Ecuador.
inated a-C:H films (a-C:H:F) appear as natural candidates to be investigated w5x. Another important application of a-C:H:F films in ultra large-scale integration (ULSI) technology is as low dielectric constant interlayer. Low dielectric interlayers are expected to decrease the interconnection parasitic capacitance and to improve switching performance of ULSI circuits w6x. Finally, a-C:H:F films were proposed as corrosion- and abrasion-resistant antireflection coatings for infrared optics w7x since the incorporation of fluorine, which essentially replaces hydrogen in the film structure, leads to the suppression of CHn absorption bands in the region of 3.0–3.5 mm, broadening the transparency window up to 9 mm, where the CFx absorption bands appear. Amorphous fluorinated carbon films deposited by plasma assisted chemical vapor deposition (PACVD) are able to combine the mechanical properties of hard a-C:H films and the tribological properties of fluorinebased lubricants w8x. They can be grown in either C2H2 – CF4 or CH4 –CF4 atmospheres w9,10x. Depending on the deposition parameters, like the self-bias voltage and the fluorocarbon gas partial pressure in the precursor atmosphere, fluorine may replace either hydrogen or carbon atoms in the amorphous film network. Deposition parameters also determine whether the film is a diamond-
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or polymer-like film controlling, thus, its mechanical properties w8,11x. Also, the fluorine content of the films controls its surface energy and friction. It has been shown that the friction coefficient measured by lateral force microscopy decreases upon the increase of fluorine in a-C:F:H films w12x. It was attributed to the reduction of the capillary meniscus formed in the asperities existing at the interface between the microscope tip and the film surface, confirming that friction properties at nanometer scale are strongly dependent on the surface energy w13,14x. In this work, we have studied the nano-scale wear of carbon and fluorinated carbon films using an atomic force microscope equipped with a stainless steel cantilever and a diamond tip, and the micro-scale wear using a sphere abrasion apparatus. Similar behaviors were observed in micro- and nano-scale wear. In both cases, wear is strongly related to the film hardness, that can be correlated with the film microstructure. The carbon chemical environment in a-C:F:H films deposited by PACVD were studied by X-ray photoelectron spectroscopy (XPS), while structural modifications of the carbon matrix were probed by Raman spectroscopy. The chemical composition and the atomic density of the films were determined by ion beam techniques. 2. Experimental procedures The a-C:H:F films were deposited by PACVD employing an asymmetrical capacitively coupled deposition system. In order to deposit a-C:H:F films, C2H2 – CF4 (1:1) and CH4 –CF4 (1:2) gas mixtures were used as precursor atmosphere. The a-C:H films were deposited in a pure acetylene atmosphere. Si substrates were mounted on a water-cooled copper cathode fed by an r.f. (13.56 MHz) power supply. In all cases, the total pressure was 10 Pa and the total gas influx was 3 sccm. The self-bias voltage was y350 V and the r.f. input power was approximately 20 W for all gas mixtures employed. In order to improve the adhesion of the aC:H:F films to the substrate, a thin (10-nm thick) buffer layer of a-C:H was deposited from a pure acetylene atmosphere. The film thickness measured by a stylus profilometer was 1 mm. The chemical composition of the films was determined by ion beam analysis (IBA), while the film density was obtained by combining IBA results and the film thickness. A detailed description of the IBA measurements can be found in Ref. w8x. The film microstructure was probed by XPS and Raman spectroscopy. XPS was employed to investigate the chemical state of the carbon atoms at the film nearsurface layers. It was performed using an Omicron UHV station based on an EA125 hemispherical analyzer. The photoelectron spectra of C 1s and F 1s core levels were monitored, excited by Mg Ka (1253.6 eV) radiation,
with an overall resolution of 1 eV for 10-eV pass energy. The angle between the surface normal and the electron energy analyzer axis was 458. No surface cleaning procedure was adopted. Raman spectroscopy was carried out in air and at room temperature by employing an Arq-ion laser (ls488 nm, powers100 mW), a double monochromator (Jobin Yvon Ramanor HG2-S) equipped with holographic gratings (2000 linesymm) and a photon counting system. The hardness of the films was measured employing a Hysitron TriboIndenter with loads up to 6 mN. The film hardness was obtained according to the Oliver and Pharr method w15x. The final hardness values were taken as the average of 10 indentations carried out in different spots for penetration depths of approximately 40 nm that are shallower than 10% of the sample thickness. In order to measure wear of the carbon films a Multimode AFM (Nanoscope IIIa, Digital Instruments) equipped with a diamond tip mounted on a stainless steel cantilever (Ks165 Nym) was used. The microscope was operated in tapping mode for both checking the surface before scratching and measuring the wear marks after the scratch. To produce the wear marks, the tip was first pushed against the sample surface until a desired normal force was reached and then, the sample was scanned perpendicularly to the cantilevers main axis at a constant velocity of 1 mmys. The measured wear depth was taken as the vertical distance between the average groove depth and the average height of the image in a region without the wear mark. We did not take into account tipysample convolution effects on the evaluation of the image features. The micro-scale wear measurements were performed by the sliding sphere test using a Calowear娃 equipment (CSEM Instruments). In our experiment, a hardened 25.4-mm diameter steel ball slides over the samples, with a normal load of 180"10 mN and sliding velocity of 80 mmys. Abrasive slurry composed of 100 ml of water and 75 g of silicon carbide (SiC) particles, 4 mm in diameter, was continuously dropped onto the ball during the tests. The rotation of the sphere with the slurry creates a spherical crater on the film and on the Si (1 0 0) substrate that were measured with an optical microscope. The morphology of the craters surface was observed with the AFM in order to gain information on the wear mechanisms. 3. Results and discussion 3.1. Structural characterization The chemical composition and atomic density of the films obtained by IBA measurements, together with the film hardness, are presented in Table 1 for the different precursor atmospheres. The errors in the atomic concentrations and film densities are of the order of 10%, while
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Table 1 Film composition, atomic density and hardness for films deposited in different precursor atmospheres: pure C2H2, C2H2 –CF4 (1:1) and CH4 –CF4 (1:2) Plasma atmosphere C2H2 C2H2qCF4 CH4qCF4
Composition (%) C
H
F
O
77 73 78
22 4 8
0 22 14
1 1 0
Atomic density (=1023 at.ycm3)
Hardness (GPa)
1.3 0.9 1.2
13.0 4.0 7.5
the errors in the film hardness are 1.5 GPa. A small contamination with oxygen was observed for films deposited in C2H2-containing atmospheres. The films deposited in pure acetylene atmosphere have a diamond-like character, since they are dense and hard while the films deposited in C2H2 –CF4 atmosphere
Fig. 2. Raman spectra obtained from films deposited in different precursor atmospheres. The arrows indicate the positions of the D- and G-bands.
Fig. 1. XPS spectra taken from films deposited in different precursor atmospheres. The main features of each spectrum are indicated in the figure, as well as, the simulated spectra composed by the different bands and the Shirley background.
have the higher fluorine content, the lowest atomic density and hardness, suggesting a polymeric character. These results indicate a transition from diamond-like to polymer-like arrangements upon fluorine incorporation that was confirmed by the XPS results shown in Fig. 1. It has to be noted that only the carbon peaks are shown in Fig. 1 because the fluorine peaks at ;690 eV do not provide any information about the different chemical bonds. According to the binding energy values taken from Ref. w16x, higher fluorine contents lead to a relative increase of the intensity of fluorinated carbon peaks, the band at 289.5 eV (C–F) corresponds to the C 1s binding energy of a carbon atom directed bonded to a fluorine atom, while the band at 287.1 eV (C–CF) is related to carbon atoms bonded to another carbon in the neighborhood of a fluorine atom. In addition to the carbon chemical environment analysis performed by XPS, Raman spectroscopy was employed to probe the C–C structural modifications. In Fig. 2 we show the Raman spectra obtained from different samples. The Raman spectra present two partially overlapping bands typical of all amorphous carbon materials. They are usually referred as the D (disorder) and the G (graphite) bands, since they recall more or
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Table 2 The peak position, width and IDyIG peak intensity ratio, obtained by Raman spectroscopy are presented Films
VG
GG
VD
GD
IDyIG
C2H2 C2H2qCF4 CH4qCF4
1550 1556 1548
138 140 137
1371 1380 1370
298 303 290
1.10 1.05 1.03
less close the Raman spectrum of polycrystalline graphite w17x as well as other forms of non-crystalline graphitic materials w18x. The first band, centered at approximately 1370ycm, corresponds to the so-called D-band and is associated with disorder-allowed zone edge modes of graphite that become Raman active due to the lack of long-range order w17x. The second one, peaked approximately 1550ycm, is known as the Gband and is attributed to E2g-symmetry optical modes occurring at the Brillouin center zone of crystalline graphite w17x. This band is observed in the Raman spectra of the a-C:H films as well as in tetrahedral amorphous carbon films but it may appears at slightly different energies in the two forms of amorphous carbon, depending on their sp3 fractional content w19x and on the size of the sp2 domains w20x. The experimental spectra, after background subtraction, can be fitted using two Gaussian lines. The peaks position and width as well as their intensity ratios are presented in Table 2. It is clear that the differences between these parameters are small. However, it should be pointed out that Raman spectroscopy is extremely sensitive to sp2-hybridized carbon domains, so that the analysis of the D- and Gbands cannot completely describe the structural modifications induced by changing the deposition conditions. New information can be obtained perhaps from the evolution of the luminescence background, underlying the Raman spectra. In fact, while the luminescence intensity is negligible in the spectra for films deposited in pure C2H2 and CH4 –CF4 atmospheres, as typically occurs in diamond-like films, it becomes important in the spectra obtained from films deposited in C2H2 –CF4 atmosphere. The increase of the luminescence background has been discussed in terms of a structural rearrangement towards a polymer-like structure w11,21x, and this interpretation is reinforced by the XPS identification of C–Fn bands, characteristic of fluorinated polymers, in the spectra obtained from films deposited in C2H2 –CF4 atmosphere. These polymeric films are also the softest ones. Summarizing this section, IBA, XPS and Raman results show that films deposited in pure C2H2 atmosphere have a diamond-like character, while there is a structural rearrangement towards a less dense and soft polymer-like structure upon fluorine incorporation into the amorphous skeleton. Films deposited in CH4 –CF4 (1:2) atmosphere also have diamond-like properties.
3.2. Wear results AFM images of typical scratches produced after one scanning cycle are presented in Fig. 3. The applied force was 50 mN. Note that almost no debris was observed on the images. Meanwhile, a shallow groove can be seen even in the case of the films deposited in C2H2q CF4 atmosphere. The dependence of the scratch depth with the normal force applied between the microscope tip and the film is presented in Fig. 4. The depth is observed to increase linearly with the normal load. An analysis of the results presented in Fig. 4 shows that the lower wear resistance is obtained for films deposited in C2H2qCF4 atmosphere, while the other two series of films present higher wear resistance. Despite that, all films provide good wear resistance compared to bare Si. In fact, when a Si (1 0 0) wafer was scratched with a normal force of 50 mN, the scratch depth is 40 nm.
Fig. 3. AFM images of typical sctraches (normal load of 50 mN). Note the different vertical scales.
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the sharpness of the asperity tips is, besides surface hardness, another film property that controls wear w22,23x. The elasticity may as well influence the surface wear. When the applied load between the tip and the film surface increases there is an increase in the tipsurface contact area, leading to a decrease on the pressure at the contact. Note that, this increase in the contact area is higher for films with higher elasticity. The wear depth as a function of the number of scanning cycles was also investigated. In order to minimize any possible damage or wear of the tip, we adopt the line-scratch mode and low sliding speed, instead of the area-scratch mode w24x. The results for a constant normal load of 50 mN are shown in Fig. 5. In all of the cases, the wear depth follows an approximately linear behavior. This linear behavior might be an indication that while scanning, the tip is removing a constant
Fig. 4. Scratch depth as a function of the normal force. As a shake of simplicity, in (a) we present the results obtained from films deposited in C2H2 and C2H2 –CF4 atmospheres and in (b) we present the results obtained from films deposited in CH4 –CF4 and C2H2 –CF4 atmospheres. The lines are only an eye-guide.
The minimum force necessary to scratch the surface of the C2H2qCF4 deposited film, that can be obtained by extrapolating the wear depth vs. force curve to zero, is less than 1.0 mN while to scratch the hardest films normal forces of the order of 10 mN are necessary. The scratch depths are higher for the softer film since a higher indentation depth will be achieved when the microscope tip is pushed against the sample surface, leading to a higher scratch depth. However, when we compare the results obtained for the films deposited in C2H2 and CH4 –CF4 atmospheres, quite similar scratch depths were obtained for the same normal load. These results indicate that the film hardness cannot be the only parameter used to explain the results of wear at nanometer scale; surface asperities and the film elasticity, for example, may also play a role. It is known that, when surfaces are brought into contact, it occurs predominantly at the surface asperities. When they are sliding against each other, the dissipated energy might be sufficient to modify the surface topography. It was shown that the nano-scale film morphology, particularly,
Fig. 5. Wear depth as a function of the wear cycles. In (a) we present the results obtained from films deposited in C2H2 and C2H2 –CF4 atmospheres and in (b) we present the results obtained from films deposited in CH4 –CF4 and C2H2 –CF4 atmospheres. The lines are only an eye-guide.
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Fig. 6. Wear rate and critical force to scratch the films as functions of the hardness. As a shake of comparison, the wear rate obtained for the Si substrate is 20=10y3 mm3yNm. The lines are only an eyeguide.
number of atomic layers per scratch. It has to be noted that, to produce the wear marks, the microscope tip was first pushed against the sample surface until a desired load was reached and then, the sample was scanned. This procedure explains why the extrapolation to zero cycles of the depth vs. number of cycles curve is not the origin. In Fig. 5a, one can see that the incorporation of fluorine in a-C:H films deposited in acetylene atmospheres induce an increase not only in the wear depth but in the wear rate as well. The rate goes from 0.17"0.07 for the a-C:H film to 0.29"0.04 nmycycle for the polymeric a-C:H:F film. From the data presented in Fig. 5b, the direct comparison of the wear of a-C:H:F films deposited in C2H2qCF4 and CH4qCF4 atmospheres can be made. It can be observed that the wear rate for the soft and polymeric a-C:H:F film is also higher than the rate for the harder a-C:H:F film. As the results presented in Figs. 4 and 5 show, both the wear rate and the scratch depth measured in the same experimental conditions for the a-C:H film and for the hard a-C:H:F films are quite similar, despite the fact that the fluorinated film has a hardness value that is two-thirds of the hardness of the a-C:H film. These results suggest that at nanometer scale, other film characteristics may be considered. In order to evaluate the effectiveness of each of the studied film as a protective coating at micrometer scale, we have used the sliding sphere test. The wear volume of the films was measured as a function of the ball sliding distance. A method of analysis allows the simultaneous evaluation of the intrinsic wear for both the substrate and the coating in a coated sample for their combined wear data w25x. From Fig. 6, it is seen that the wear rate of the polymeric a-C:H:F film is higher than the values
obtained for the harder a-C:H:F films deposited in CH4qCF4 gas mixture and for the a-C:H film. The results show that when the film hardness decreases, an increase in the wear rate is observed. At the micrometer scale, the hardness of the films appears to be the dominant factor in determining the wear resistance. In this figure, we also plot the critical force to scratch the film surface using the microscope tip as function of the film hardness. In this last set of data, the hardness also appears as a good parameter to present the data. In Fig. 7, a typical AFM image of the wear crater on an a-C:H:F film is presented and it is representative of all studied samples. The image shows the absence of long grooves in the direction of the ball sliding in both film and substrate. The surface morphology suggested a three-body abrasive wear mechanism involving plastic deformation w25x. This is an indication that the SiC particles might have rolled over the surface during the tests leading to small indentations and ploughing on the worn surface. 4. Summary and conclusions Nano- and micro-scale wear behavior of a-C:H and a-C:H:F films deposited by PACVD have been studied with the use of the atomic force microscopy and the sliding sphere test. At nanometer scale, the wear depth is observed to increase linearly with the increase of the tip-surface normal force. An analysis of the minimum force necessary to scratch the surface, obtained by extrapolating the wear vs. force curve, as well as the wear rates and the scratch depth test results, show that the wear is strongly related to the film hardness and microstructure. This fact has been observed either in a very low forces regime, in the range of micronewton using the AFM, or with forces in the range of millinewton with the sphere abrasion equipment, since similar
Fig. 7. An AFM image of the wear crater produced on an a-C:H:F film surface by the sphere abrasion test.
P. Ayala et al. / Surface and Coatings Technology 182 (2004) 335–341
trend is observed in both scales. Even though the range of forces and the wear process in both tests are quite different, we are lead to the conclusion that the hardness of the films plays the dominant role in determine its wear performance in both nano- and micro-scales. However, the results at nanometer scale also suggested that other film characteristics, such as roughness and elasticity, must be taken into consideration in order to evaluate a candidate to be used as protective coating. Acknowledgments This work was partially supported by the Brazilian Agencies: Conselho Nacional de Desenvolvimento ´ ´ ¸ ˜ de AperCientıfico e Tecnologico (CNPq), Coordenacao ´ Superior (CAPES) and ¸ feicoamento de Pessoal de Nıvel ¸ ˜ de Amparo a` Pesquisa do Estado do Rio de Fundacao Janeiro (FAPERJ). We would like to thank Prof. G. Mariotto for the Raman experiments, Dr C. Radke for the XPS analysis and Dr L.G. Jacobsohn for the hardness measurements. References w1x B. Bushan (Ed.), Handbook of MicroyNanoTribology, second ed., CRC Press, Boca Raton, 1999. w2x F.L. Freire Jr., R. Prioli, in: H.S. Nalwa (Ed.), Encyclopedia of Nanoscience and Nanotechnology, American Scientific Publishers, California, Stevenson Ranch, 2003. w3x A. Grill, Diamond Relat. Mater. 8 (1999) 428. w4x T.E. Karis, G.W. Tyndall, D. Fenzel-Alexander, M.S. Crowder, J. Appl. Phys. 81 (1997) 5378.
341
w5x C. Donnet, J. Fontaine, A. Grill, V. Patel, C. Jahnes, M. Belin, Surf. Coat. Technol. 94 (1997) 531. w6x J. Theil, J. Vac. Sci. Technol. B 17 (1999) 2397. w7x R.E. Sah, B. Dischler, A. Bubenzer, P. Koidl, Appl. Phys. Lett. 46 (1985) 739. w8x L.G. Jacobsohn, D.F. Franceschini, M.E.H. Maia da Costa, F.L. Freire Jr., J. Vac. Sci. Technol. A 18 (2000) 2230. w9x F.L. Freire Jr., M.E.H. Maia da Costa, L.G. Jacobsohn, D.F. Franceschini, Diamond Relat. Mater. 10 (2001) 125. w10x F.L. Freire Jr., M.E.H. Maia da Costa, V.J. Trava-Airoldi, L.G. Jacobsohn, Diamond Relat. Mater., in press. w11x L.G. Jacobsohn, S.S. Camargo Jr., F.L. Freire Jr., Diamond Relat. Mater. 11 (2002) 1831. w12x R. Prioli, L.G. Jacobsohn, M.E.H. Maia da Costa, F.L. Freire Jr., Tribol. Lett. 15 (2003) 177. w13x E. Riedo, F. Levy, ´ H. Brune, Phys. Rev. Lett. 88 (2002) 185505. w14x B. Bushan, S. Sundararajan, Acta Mater. 46 (1998) 3793. w15x W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564. w16x S. Agraharam, D.H. Hess, P.A. Kohl, S. Bistrup Allen, J. Vac. Sci. Technol. A 17 (1999) 3265. w17x F. Tuinstra, J.L. Koenig, J. Chem. Phys. 53 (1970) 1126. w18x R.O. Dillon, J.A. Woollam, V. Karkanant, Phys. Rev. B 29 (1984) 3482. w19x R. Kalish, Y. Lifshitz, K. Nugent, S. Prawer, Appl. Phys. Lett. 74 (1999) 2936. w20x G. Mariotto, C.A. Achete, F.L. Freire Jr., Thin Solid Films 241 (1994) 255. w21x R. d’Agostino, R. Lamendola, P. Favia, A. Giquel, J. Vac. Sci. Technol. A 12 (1994) 308. w22x S.J. Harris, A.M. Weiner, C.H. Olk, M. Crischke, Wear 219 (1998) 98. w23x R. Prioli, D.C. Reigada, F.L. Freire Jr., J. Appl. Phys. 87 (2000) 1118. w24x M. Bai, K. Kato, N. Umehara, Y. Miyake, J. Xu, H. Tokisue, Surf. Coat. Technol. 126 (2000) 181. w25x K.L. Rutherford, I.M. Huychings, Surf. Coat. Technol. 79 (1996) 231.
Nano- and micro-scale wear of fluorinated carbon films P. Ayala1, M.E.H. Maia da Costa, R. Prioli, F.L. Freire Jr.* ´ ´ ´ Departamento de Fısica, Pontifıcia Universidade Catolica do Rio de Janeiro, Caixa Postal 38071, 22452-970 Rio de Janeiro, RJ, Brazil Received 19 May 2003; accepted in revised form 19 August 2003
Abstract Wear of fluorinated carbon films deposited by plasma assisted chemical vapor deposition has been studied by atomic force microscopy and the sliding sphere test. The films microstructure and chemical composition were also investigated. The films became polymeric upon fluorine incorporation, as revealed by density measurements, Raman and X-ray photoelectron spectroscopies. It was shown that the wear marks produced with a microscope diamond tip are strongly related to the film hardness and microstructure. The wear depth measured on a polymeric film is greater compared with the ones obtained from hard carbon and carbon fluorinated films. At nanometer scale, the minimum force necessary to scratch the films increases with the film hardness, while the wear rate decreases. Similar behaviors were observed in both nano- and micro-scale wear experiments. 䊚 2003 Elsevier B.V. All rights reserved. Keywords: Abrasive wheel test; Amorphous; Atomic force microscopy; Photoelectron spectroscopy; Plasma assisted chemical vapor deposition; Carbon
1. Introduction The tribological behavior of moving micro- and nanodevices is a challenging subject that impacts their reliabilities and operational lifetimes. The fundamental understanding of the friction and wear in nanometer scale is important to achieve such goals w1x. Amorphous hydrogenated carbon films (a-C:H) have been under investigation for a long time due to their properties such as high hardness, chemical inertness and low friction coefficients w2,3x. These films have found their way into industrial applications as tribological protective coatings. One of the main drawbacks of this application is the need of applying a fluorinated lubricant layer onto the surface of the a-C:H protective coatings in order to reduce the film wear w4x. Good candidates to substitute this lubricant layer are selflubricant films deposited by a plasma technique. Fluor´ *Corresponding author. Departamento de Fısica, PUC-Rio, Rua ˆ de Sao ˜ Vicente, 225, Caixa Postal 38071, 22453-970 Rio Marques de Janeiro, RJ, Brazil. Tel.: q55-21-3114-1272; fax: q55-21-31141040. E-mail address: [email protected] (F.L. Freire Jr.). 1 Present address: Departamento de Fisica, Escuela Politecnica Nacional, Quito, Ecuador.
inated a-C:H films (a-C:H:F) appear as natural candidates to be investigated w5x. Another important application of a-C:H:F films in ultra large-scale integration (ULSI) technology is as low dielectric constant interlayer. Low dielectric interlayers are expected to decrease the interconnection parasitic capacitance and to improve switching performance of ULSI circuits w6x. Finally, a-C:H:F films were proposed as corrosion- and abrasion-resistant antireflection coatings for infrared optics w7x since the incorporation of fluorine, which essentially replaces hydrogen in the film structure, leads to the suppression of CHn absorption bands in the region of 3.0–3.5 mm, broadening the transparency window up to 9 mm, where the CFx absorption bands appear. Amorphous fluorinated carbon films deposited by plasma assisted chemical vapor deposition (PACVD) are able to combine the mechanical properties of hard a-C:H films and the tribological properties of fluorinebased lubricants w8x. They can be grown in either C2H2 – CF4 or CH4 –CF4 atmospheres w9,10x. Depending on the deposition parameters, like the self-bias voltage and the fluorocarbon gas partial pressure in the precursor atmosphere, fluorine may replace either hydrogen or carbon atoms in the amorphous film network. Deposition parameters also determine whether the film is a diamond-
0257-8972/04/$ - see front matter 䊚 2003 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2003.08.075
336
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or polymer-like film controlling, thus, its mechanical properties w8,11x. Also, the fluorine content of the films controls its surface energy and friction. It has been shown that the friction coefficient measured by lateral force microscopy decreases upon the increase of fluorine in a-C:F:H films w12x. It was attributed to the reduction of the capillary meniscus formed in the asperities existing at the interface between the microscope tip and the film surface, confirming that friction properties at nanometer scale are strongly dependent on the surface energy w13,14x. In this work, we have studied the nano-scale wear of carbon and fluorinated carbon films using an atomic force microscope equipped with a stainless steel cantilever and a diamond tip, and the micro-scale wear using a sphere abrasion apparatus. Similar behaviors were observed in micro- and nano-scale wear. In both cases, wear is strongly related to the film hardness, that can be correlated with the film microstructure. The carbon chemical environment in a-C:F:H films deposited by PACVD were studied by X-ray photoelectron spectroscopy (XPS), while structural modifications of the carbon matrix were probed by Raman spectroscopy. The chemical composition and the atomic density of the films were determined by ion beam techniques. 2. Experimental procedures The a-C:H:F films were deposited by PACVD employing an asymmetrical capacitively coupled deposition system. In order to deposit a-C:H:F films, C2H2 – CF4 (1:1) and CH4 –CF4 (1:2) gas mixtures were used as precursor atmosphere. The a-C:H films were deposited in a pure acetylene atmosphere. Si substrates were mounted on a water-cooled copper cathode fed by an r.f. (13.56 MHz) power supply. In all cases, the total pressure was 10 Pa and the total gas influx was 3 sccm. The self-bias voltage was y350 V and the r.f. input power was approximately 20 W for all gas mixtures employed. In order to improve the adhesion of the aC:H:F films to the substrate, a thin (10-nm thick) buffer layer of a-C:H was deposited from a pure acetylene atmosphere. The film thickness measured by a stylus profilometer was 1 mm. The chemical composition of the films was determined by ion beam analysis (IBA), while the film density was obtained by combining IBA results and the film thickness. A detailed description of the IBA measurements can be found in Ref. w8x. The film microstructure was probed by XPS and Raman spectroscopy. XPS was employed to investigate the chemical state of the carbon atoms at the film nearsurface layers. It was performed using an Omicron UHV station based on an EA125 hemispherical analyzer. The photoelectron spectra of C 1s and F 1s core levels were monitored, excited by Mg Ka (1253.6 eV) radiation,
with an overall resolution of 1 eV for 10-eV pass energy. The angle between the surface normal and the electron energy analyzer axis was 458. No surface cleaning procedure was adopted. Raman spectroscopy was carried out in air and at room temperature by employing an Arq-ion laser (ls488 nm, powers100 mW), a double monochromator (Jobin Yvon Ramanor HG2-S) equipped with holographic gratings (2000 linesymm) and a photon counting system. The hardness of the films was measured employing a Hysitron TriboIndenter with loads up to 6 mN. The film hardness was obtained according to the Oliver and Pharr method w15x. The final hardness values were taken as the average of 10 indentations carried out in different spots for penetration depths of approximately 40 nm that are shallower than 10% of the sample thickness. In order to measure wear of the carbon films a Multimode AFM (Nanoscope IIIa, Digital Instruments) equipped with a diamond tip mounted on a stainless steel cantilever (Ks165 Nym) was used. The microscope was operated in tapping mode for both checking the surface before scratching and measuring the wear marks after the scratch. To produce the wear marks, the tip was first pushed against the sample surface until a desired normal force was reached and then, the sample was scanned perpendicularly to the cantilevers main axis at a constant velocity of 1 mmys. The measured wear depth was taken as the vertical distance between the average groove depth and the average height of the image in a region without the wear mark. We did not take into account tipysample convolution effects on the evaluation of the image features. The micro-scale wear measurements were performed by the sliding sphere test using a Calowear娃 equipment (CSEM Instruments). In our experiment, a hardened 25.4-mm diameter steel ball slides over the samples, with a normal load of 180"10 mN and sliding velocity of 80 mmys. Abrasive slurry composed of 100 ml of water and 75 g of silicon carbide (SiC) particles, 4 mm in diameter, was continuously dropped onto the ball during the tests. The rotation of the sphere with the slurry creates a spherical crater on the film and on the Si (1 0 0) substrate that were measured with an optical microscope. The morphology of the craters surface was observed with the AFM in order to gain information on the wear mechanisms. 3. Results and discussion 3.1. Structural characterization The chemical composition and atomic density of the films obtained by IBA measurements, together with the film hardness, are presented in Table 1 for the different precursor atmospheres. The errors in the atomic concentrations and film densities are of the order of 10%, while
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Table 1 Film composition, atomic density and hardness for films deposited in different precursor atmospheres: pure C2H2, C2H2 –CF4 (1:1) and CH4 –CF4 (1:2) Plasma atmosphere C2H2 C2H2qCF4 CH4qCF4
Composition (%) C
H
F
O
77 73 78
22 4 8
0 22 14
1 1 0
Atomic density (=1023 at.ycm3)
Hardness (GPa)
1.3 0.9 1.2
13.0 4.0 7.5
the errors in the film hardness are 1.5 GPa. A small contamination with oxygen was observed for films deposited in C2H2-containing atmospheres. The films deposited in pure acetylene atmosphere have a diamond-like character, since they are dense and hard while the films deposited in C2H2 –CF4 atmosphere
Fig. 2. Raman spectra obtained from films deposited in different precursor atmospheres. The arrows indicate the positions of the D- and G-bands.
Fig. 1. XPS spectra taken from films deposited in different precursor atmospheres. The main features of each spectrum are indicated in the figure, as well as, the simulated spectra composed by the different bands and the Shirley background.
have the higher fluorine content, the lowest atomic density and hardness, suggesting a polymeric character. These results indicate a transition from diamond-like to polymer-like arrangements upon fluorine incorporation that was confirmed by the XPS results shown in Fig. 1. It has to be noted that only the carbon peaks are shown in Fig. 1 because the fluorine peaks at ;690 eV do not provide any information about the different chemical bonds. According to the binding energy values taken from Ref. w16x, higher fluorine contents lead to a relative increase of the intensity of fluorinated carbon peaks, the band at 289.5 eV (C–F) corresponds to the C 1s binding energy of a carbon atom directed bonded to a fluorine atom, while the band at 287.1 eV (C–CF) is related to carbon atoms bonded to another carbon in the neighborhood of a fluorine atom. In addition to the carbon chemical environment analysis performed by XPS, Raman spectroscopy was employed to probe the C–C structural modifications. In Fig. 2 we show the Raman spectra obtained from different samples. The Raman spectra present two partially overlapping bands typical of all amorphous carbon materials. They are usually referred as the D (disorder) and the G (graphite) bands, since they recall more or
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Table 2 The peak position, width and IDyIG peak intensity ratio, obtained by Raman spectroscopy are presented Films
VG
GG
VD
GD
IDyIG
C2H2 C2H2qCF4 CH4qCF4
1550 1556 1548
138 140 137
1371 1380 1370
298 303 290
1.10 1.05 1.03
less close the Raman spectrum of polycrystalline graphite w17x as well as other forms of non-crystalline graphitic materials w18x. The first band, centered at approximately 1370ycm, corresponds to the so-called D-band and is associated with disorder-allowed zone edge modes of graphite that become Raman active due to the lack of long-range order w17x. The second one, peaked approximately 1550ycm, is known as the Gband and is attributed to E2g-symmetry optical modes occurring at the Brillouin center zone of crystalline graphite w17x. This band is observed in the Raman spectra of the a-C:H films as well as in tetrahedral amorphous carbon films but it may appears at slightly different energies in the two forms of amorphous carbon, depending on their sp3 fractional content w19x and on the size of the sp2 domains w20x. The experimental spectra, after background subtraction, can be fitted using two Gaussian lines. The peaks position and width as well as their intensity ratios are presented in Table 2. It is clear that the differences between these parameters are small. However, it should be pointed out that Raman spectroscopy is extremely sensitive to sp2-hybridized carbon domains, so that the analysis of the D- and Gbands cannot completely describe the structural modifications induced by changing the deposition conditions. New information can be obtained perhaps from the evolution of the luminescence background, underlying the Raman spectra. In fact, while the luminescence intensity is negligible in the spectra for films deposited in pure C2H2 and CH4 –CF4 atmospheres, as typically occurs in diamond-like films, it becomes important in the spectra obtained from films deposited in C2H2 –CF4 atmosphere. The increase of the luminescence background has been discussed in terms of a structural rearrangement towards a polymer-like structure w11,21x, and this interpretation is reinforced by the XPS identification of C–Fn bands, characteristic of fluorinated polymers, in the spectra obtained from films deposited in C2H2 –CF4 atmosphere. These polymeric films are also the softest ones. Summarizing this section, IBA, XPS and Raman results show that films deposited in pure C2H2 atmosphere have a diamond-like character, while there is a structural rearrangement towards a less dense and soft polymer-like structure upon fluorine incorporation into the amorphous skeleton. Films deposited in CH4 –CF4 (1:2) atmosphere also have diamond-like properties.
3.2. Wear results AFM images of typical scratches produced after one scanning cycle are presented in Fig. 3. The applied force was 50 mN. Note that almost no debris was observed on the images. Meanwhile, a shallow groove can be seen even in the case of the films deposited in C2H2q CF4 atmosphere. The dependence of the scratch depth with the normal force applied between the microscope tip and the film is presented in Fig. 4. The depth is observed to increase linearly with the normal load. An analysis of the results presented in Fig. 4 shows that the lower wear resistance is obtained for films deposited in C2H2qCF4 atmosphere, while the other two series of films present higher wear resistance. Despite that, all films provide good wear resistance compared to bare Si. In fact, when a Si (1 0 0) wafer was scratched with a normal force of 50 mN, the scratch depth is 40 nm.
Fig. 3. AFM images of typical sctraches (normal load of 50 mN). Note the different vertical scales.
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the sharpness of the asperity tips is, besides surface hardness, another film property that controls wear w22,23x. The elasticity may as well influence the surface wear. When the applied load between the tip and the film surface increases there is an increase in the tipsurface contact area, leading to a decrease on the pressure at the contact. Note that, this increase in the contact area is higher for films with higher elasticity. The wear depth as a function of the number of scanning cycles was also investigated. In order to minimize any possible damage or wear of the tip, we adopt the line-scratch mode and low sliding speed, instead of the area-scratch mode w24x. The results for a constant normal load of 50 mN are shown in Fig. 5. In all of the cases, the wear depth follows an approximately linear behavior. This linear behavior might be an indication that while scanning, the tip is removing a constant
Fig. 4. Scratch depth as a function of the normal force. As a shake of simplicity, in (a) we present the results obtained from films deposited in C2H2 and C2H2 –CF4 atmospheres and in (b) we present the results obtained from films deposited in CH4 –CF4 and C2H2 –CF4 atmospheres. The lines are only an eye-guide.
The minimum force necessary to scratch the surface of the C2H2qCF4 deposited film, that can be obtained by extrapolating the wear depth vs. force curve to zero, is less than 1.0 mN while to scratch the hardest films normal forces of the order of 10 mN are necessary. The scratch depths are higher for the softer film since a higher indentation depth will be achieved when the microscope tip is pushed against the sample surface, leading to a higher scratch depth. However, when we compare the results obtained for the films deposited in C2H2 and CH4 –CF4 atmospheres, quite similar scratch depths were obtained for the same normal load. These results indicate that the film hardness cannot be the only parameter used to explain the results of wear at nanometer scale; surface asperities and the film elasticity, for example, may also play a role. It is known that, when surfaces are brought into contact, it occurs predominantly at the surface asperities. When they are sliding against each other, the dissipated energy might be sufficient to modify the surface topography. It was shown that the nano-scale film morphology, particularly,
Fig. 5. Wear depth as a function of the wear cycles. In (a) we present the results obtained from films deposited in C2H2 and C2H2 –CF4 atmospheres and in (b) we present the results obtained from films deposited in CH4 –CF4 and C2H2 –CF4 atmospheres. The lines are only an eye-guide.
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Fig. 6. Wear rate and critical force to scratch the films as functions of the hardness. As a shake of comparison, the wear rate obtained for the Si substrate is 20=10y3 mm3yNm. The lines are only an eyeguide.
number of atomic layers per scratch. It has to be noted that, to produce the wear marks, the microscope tip was first pushed against the sample surface until a desired load was reached and then, the sample was scanned. This procedure explains why the extrapolation to zero cycles of the depth vs. number of cycles curve is not the origin. In Fig. 5a, one can see that the incorporation of fluorine in a-C:H films deposited in acetylene atmospheres induce an increase not only in the wear depth but in the wear rate as well. The rate goes from 0.17"0.07 for the a-C:H film to 0.29"0.04 nmycycle for the polymeric a-C:H:F film. From the data presented in Fig. 5b, the direct comparison of the wear of a-C:H:F films deposited in C2H2qCF4 and CH4qCF4 atmospheres can be made. It can be observed that the wear rate for the soft and polymeric a-C:H:F film is also higher than the rate for the harder a-C:H:F film. As the results presented in Figs. 4 and 5 show, both the wear rate and the scratch depth measured in the same experimental conditions for the a-C:H film and for the hard a-C:H:F films are quite similar, despite the fact that the fluorinated film has a hardness value that is two-thirds of the hardness of the a-C:H film. These results suggest that at nanometer scale, other film characteristics may be considered. In order to evaluate the effectiveness of each of the studied film as a protective coating at micrometer scale, we have used the sliding sphere test. The wear volume of the films was measured as a function of the ball sliding distance. A method of analysis allows the simultaneous evaluation of the intrinsic wear for both the substrate and the coating in a coated sample for their combined wear data w25x. From Fig. 6, it is seen that the wear rate of the polymeric a-C:H:F film is higher than the values
obtained for the harder a-C:H:F films deposited in CH4qCF4 gas mixture and for the a-C:H film. The results show that when the film hardness decreases, an increase in the wear rate is observed. At the micrometer scale, the hardness of the films appears to be the dominant factor in determining the wear resistance. In this figure, we also plot the critical force to scratch the film surface using the microscope tip as function of the film hardness. In this last set of data, the hardness also appears as a good parameter to present the data. In Fig. 7, a typical AFM image of the wear crater on an a-C:H:F film is presented and it is representative of all studied samples. The image shows the absence of long grooves in the direction of the ball sliding in both film and substrate. The surface morphology suggested a three-body abrasive wear mechanism involving plastic deformation w25x. This is an indication that the SiC particles might have rolled over the surface during the tests leading to small indentations and ploughing on the worn surface. 4. Summary and conclusions Nano- and micro-scale wear behavior of a-C:H and a-C:H:F films deposited by PACVD have been studied with the use of the atomic force microscopy and the sliding sphere test. At nanometer scale, the wear depth is observed to increase linearly with the increase of the tip-surface normal force. An analysis of the minimum force necessary to scratch the surface, obtained by extrapolating the wear vs. force curve, as well as the wear rates and the scratch depth test results, show that the wear is strongly related to the film hardness and microstructure. This fact has been observed either in a very low forces regime, in the range of micronewton using the AFM, or with forces in the range of millinewton with the sphere abrasion equipment, since similar
Fig. 7. An AFM image of the wear crater produced on an a-C:H:F film surface by the sphere abrasion test.
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trend is observed in both scales. Even though the range of forces and the wear process in both tests are quite different, we are lead to the conclusion that the hardness of the films plays the dominant role in determine its wear performance in both nano- and micro-scales. However, the results at nanometer scale also suggested that other film characteristics, such as roughness and elasticity, must be taken into consideration in order to evaluate a candidate to be used as protective coating. Acknowledgments This work was partially supported by the Brazilian Agencies: Conselho Nacional de Desenvolvimento ´ ´ ¸ ˜ de AperCientıfico e Tecnologico (CNPq), Coordenacao ´ Superior (CAPES) and ¸ feicoamento de Pessoal de Nıvel ¸ ˜ de Amparo a` Pesquisa do Estado do Rio de Fundacao Janeiro (FAPERJ). We would like to thank Prof. G. Mariotto for the Raman experiments, Dr C. Radke for the XPS analysis and Dr L.G. Jacobsohn for the hardness measurements. References w1x B. Bushan (Ed.), Handbook of MicroyNanoTribology, second ed., CRC Press, Boca Raton, 1999. w2x F.L. Freire Jr., R. Prioli, in: H.S. Nalwa (Ed.), Encyclopedia of Nanoscience and Nanotechnology, American Scientific Publishers, California, Stevenson Ranch, 2003. w3x A. Grill, Diamond Relat. Mater. 8 (1999) 428. w4x T.E. Karis, G.W. Tyndall, D. Fenzel-Alexander, M.S. Crowder, J. Appl. Phys. 81 (1997) 5378.
341
w5x C. Donnet, J. Fontaine, A. Grill, V. Patel, C. Jahnes, M. Belin, Surf. Coat. Technol. 94 (1997) 531. w6x J. Theil, J. Vac. Sci. Technol. B 17 (1999) 2397. w7x R.E. Sah, B. Dischler, A. Bubenzer, P. Koidl, Appl. Phys. Lett. 46 (1985) 739. w8x L.G. Jacobsohn, D.F. Franceschini, M.E.H. Maia da Costa, F.L. Freire Jr., J. Vac. Sci. Technol. A 18 (2000) 2230. w9x F.L. Freire Jr., M.E.H. Maia da Costa, L.G. Jacobsohn, D.F. Franceschini, Diamond Relat. Mater. 10 (2001) 125. w10x F.L. Freire Jr., M.E.H. Maia da Costa, V.J. Trava-Airoldi, L.G. Jacobsohn, Diamond Relat. Mater., in press. w11x L.G. Jacobsohn, S.S. Camargo Jr., F.L. Freire Jr., Diamond Relat. Mater. 11 (2002) 1831. w12x R. Prioli, L.G. Jacobsohn, M.E.H. Maia da Costa, F.L. Freire Jr., Tribol. Lett. 15 (2003) 177. w13x E. Riedo, F. Levy, ´ H. Brune, Phys. Rev. Lett. 88 (2002) 185505. w14x B. Bushan, S. Sundararajan, Acta Mater. 46 (1998) 3793. w15x W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564. w16x S. Agraharam, D.H. Hess, P.A. Kohl, S. Bistrup Allen, J. Vac. Sci. Technol. A 17 (1999) 3265. w17x F. Tuinstra, J.L. Koenig, J. Chem. Phys. 53 (1970) 1126. w18x R.O. Dillon, J.A. Woollam, V. Karkanant, Phys. Rev. B 29 (1984) 3482. w19x R. Kalish, Y. Lifshitz, K. Nugent, S. Prawer, Appl. Phys. Lett. 74 (1999) 2936. w20x G. Mariotto, C.A. Achete, F.L. Freire Jr., Thin Solid Films 241 (1994) 255. w21x R. d’Agostino, R. Lamendola, P. Favia, A. Giquel, J. Vac. Sci. Technol. A 12 (1994) 308. w22x S.J. Harris, A.M. Weiner, C.H. Olk, M. Crischke, Wear 219 (1998) 98. w23x R. Prioli, D.C. Reigada, F.L. Freire Jr., J. Appl. Phys. 87 (2000) 1118. w24x M. Bai, K. Kato, N. Umehara, Y. Miyake, J. Xu, H. Tokisue, Surf. Coat. Technol. 126 (2000) 181. w25x K.L. Rutherford, I.M. Huychings, Surf. Coat. Technol. 79 (1996) 231.