Structure and properties of diamondlike carbon films produced by ...

Structure and properties of diamondlike carbon films produced by pulsed laser deposition A. A. Voevodina) and S. D. Walck MLBT, Wright Laboratory, Wright-Patterson Air Force Base, Ohio 45433-7750

J. S. Solomon University of Dayton Research Institute, Dayton, Ohio 45469-0167

P. J. John MLBT, Wright Laboratory, Wright-Patterson Air Force Base, Ohio 45433-7750

D. C. Ingram Department of Physics and Astronomy, Ohio University, Athens, Ohio 45701-2979

M. S. Donley and J. S. Zabinski MLBT, Wright Laboratory, Wright-Patterson Air Force Base, Ohio 45433-7750

~Received 17 October 1995; accepted 4 March 1996! Pulsed laser deposition was used to produce hydrogen-free amorphous diamondlike carbon (a-C) and hydrogenated amorphous diamondlike carbon (a-C:H) from graphite and polycarbonate targets, respectively. Films were grown under identical conditions in high vacuum at low temperatures onto steel and Si substrates. The a-C films were uniform, while a-C:H films contained a great number of particles ejected from the target surface. The a-C films have hydrogen contamination about 0.1 at. %, while a-C:H have about 25 at. % H and 10 at. % O. High percentages of sp 3 bonding were found in both films. Film densities were estimated to be 3.0 g cm23 for a-C films and 2.2 g cm23 for a-C:H films. Chemical and structural characteristics of the films were correlated with their thermal stability and mechanical properties. Temperatures for starting graphitization were about 500 °C for a-C and 350 °C for a-C:H. The presence of hydrogen reduced film hardness from 60 GPa for a-C films to 14 GPa for a-C:H films. Hydrogen was also associated with dependence of a-C:H films friction coefficient on environment and with higher wear rates of a-C:H films in comparison to a-C films. © 1996 American Vacuum Society.

I. INTRODUCTION Diamondlike carbon ~DLC! films typically have two structures: hydrogen-free amorphous carbon (a-C! or hydrogenated amorphous carbon (a-C:H) with domination of tetrahedral bonding. Their deposition is based on ~i! producing high-energy carbon species to force atoms into a diamondlike metastable configuration, and ~ii! low substrate temperatures to suppress diffusion and structure relaxation. The lack of long-range order is the main structural feature of lowtemperature-deposited DLC. In amorphous DLC, there are always some graphite-like s p 2 bonds together with diamondlike s p 3 bonds.1–3 The s p 3 /s p 2 ratio determines film properties, which can vary over a broad range depending on preparation conditions. Although pulsed laser deposition ~PLD! of DLC has been known for some time,4,5 the intense developments of the method was started by different researchers only in the last few years. A recent review6 of these developments helped to establish laser wavelengths and energy densities required for DLC deposition. The review also indicated the uncertainty concerning the influence of hydrogen on low-temperature PLD of DLC. Comparative studies of a-C and a-C:H PLD films were performed in a number of works.7–12 In these works, hydrogen flow into the vacuum chamber was used to a

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prepare a-C:H films, raising the pressure in the chamber. However, a-C films are typically grown in high vacuum environments to maximize deposition energy and minimize contamination from residual water pressure. Increased chamber pressure and collisions of ablated carbon species with gas molecules reduces the energy of adiabatically expanding plumes. The difference in deposition energies prohibits direct comparison of DLC formation with and without hydrogen. In fact, deposition energy can be more critical for influencing on film properties than the incorporation of hydrogen. Additionally, in some works,9,11,12 a hydrogen atmosphere was accompanied by substrate temperature between 300 and 860 °C, which changed film growth mechanism9,11 and did not allow comparisons of the results with a-C film growth at low temperatures. Recently, studies have been initiated to compare a-C and a-C:H films deposited at low substrate temperatures under identical conditions using ultraviolet ~UV! laser ablation of graphite and polycarbonate targets.13–15 In PLD with graphite targets, hydrogen-free a-C films are produced. Ablation of polycarbonate targets @ C13O3H14# n with a UV laser provides carbon, hydrogen, and oxygen atoms by breaking the polymer bonds. A hydrogen atmosphere is not required and the same high vacuum environment can be used to prepare both a-C and a-C:H films. Thus, the difference in the properties of

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Voevodin et al.: DLC films produced by PLD

these films can be directly related to their chemical composition. The present study is the continuation of the previous research on the comparison of a-C and a-C:H films prepared by ablation of graphite and polycarbonate targets. Specifically, in this article, the results of investigations on surface topography, chemical composition, and thermal stability are correlated to earlier structural,13 mechanical, and tribological14 evaluations. II. EXPERIMENTAL PROCEDURES A. Deposition

A Lambda Physik PPX 110i laser charged with KrF generating 248 nm wavelength pulses of 17 ns width with 20 Hz repetition rate and 200 mJ pulse energy was used to deposit films in high vacuum by focusing laser beam on high-purity graphite and polycarbonate targets to achieve ablation powder density of 109 W cm22 . A detailed description of the deposition system is provided in Refs. 13 and 15. Films were deposited onto polished 25-mm-diam 400 C steel disks or silicon wafers. All substrates were ultrasonically cleaned in acetone, then placed into the chamber at about 4 cm from the target. The substrates were then etched in a 1 kV glow discharge in 5 Pa Ar for 15 min to remove surface contamination. The etch procedure increased substrate temperature from room to 80 °C. The chamber was evacuated to a base pressure of 1026 Pa prior to deposition. A quartz-crystal monitor was used for in situ film thickness monitoring. The substrate temperature during deposition was about 60 °C, resulting from the residual heat of the etching process. Films with thicknesses of about 0.5 mm were deposited with growth rates of 0.01 mm min21 for a-C and 0.09 mm min21 for a-C:H. B. Characterization

After deposition, the film thickness was measured by a Dektak profilometer using a masked area on the samples. This measurement also provided film roughness estimations. Film topography was investigated with a high-resolution Hitachi S-900 field emission scanning electron microscope ~SEM! on cleaved Si substrate coated with the films. Rutherford backscattering spectroscopy ~RBS! and elastic recoil spectroscopy ~ERS! analyses were performed on films deposited on Si wafers to determine the film elemental composition and to evaluate the film density. For RBS studies, a beam of alpha particles was accelerated to the energy of 2.452 MeV to avoid non-Rutherford scattering from carbon. Samples were placed at normal incidence to the beam, and the detector was positioned at a backscattering angle of 168°. In ERS analyses, the beam energy was 3.035 MeV. The sample was tilted at 75° with respect to the incident beam and the detector was positioned at 30° to the beam axis. A Mylar (C12O5H16!n foil was placed in the front of detector to stop forward scattered alpha particles. At these conditions the sensitivity to hydrogen was 0.01 at. %. J. Vac. Sci. Technol. A, Vol. 14, No. 3, May/Jun 1996

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Secondary-ion-mass spectrometry ~SIMS! was also used for analyzing the film elemental composition and detect CHn groups in the films. A quadrupole based instrument with 5 keV Cs ions was used for SIMS analyses. Fourier transform infrared ~FTIR! spectroscopy analyses were performed to analyze a-C:H structural bonding. A Perkin–Elmer 1750 instrument was used for FTIR analyses in 1000– 4000 cm21 frequency range. Film thermal stability was investigated by annealing specimens in vacuum with in situ x-ray photoelectron spectroscopy ~XPS!, using a Surface Science Instruments M-Probe instrument with a 50 eV pass energy and monochromatic Al K a x rays. Sample heating caused the increase of the base pressure in XPS chamber from 1028 Pa at room temperature to 1025 Pa at the highest annealing temperature. The sample holder of the XPS instrument was modified to perform resistive heating and thermocouple based control of the sample temperature. Heating current was switched off when XPS data were acquired. A Voight function was used to approximate C 1s binding energy peak and calculate peak position. Structural changes caused by the sample heating were verified by Raman spectroscopy. Raman spectra were recorded in the range from 1000 to 1800 cm21 using a laser wavelength of 514.5 nm. III. RESULTS AND DISCUSSION A. Film morphology and surface roughness

Studies of film morphology with a cross-sectional highresolution SEM13 found that a-C films are homogeneous, while the a-C:H films have a heterogeneous morphology. When ablating graphite with an UV laser, the plume is a mixture of atoms and ions with energies up to 1.5 keV.13 Plume contamination with splashed microparticles is minimized for graphite due to it good thermal conductivity, higher density, and optical absorption coefficient. A few ejected particles were radiatively cooled during time of flight and were incorporated as random spherical particles in a-C films. This affected the a-C film topography, shown in Fig. 1~a! for a sample cleaved and tilted to provide a view of incorporated spherical particles. The maximum of surface roughness was estimated to be about 100–200 nm, which basically reflects dimensions of largest particles @Fig. 1~a!#. However, due to the spatial distribution of incorporated particles, the average roughness (R a ) did not exceed 10 nm. In fact, R a measured between particles was about 2–3 nm. Polycarbonate has poor thermal conductivity, low density, and low absorption coefficient. Plumes produced by ablation of polycarbonate have lower kinetic energy and contain a large number of melted particles.13 This results in the development of surface topography of a-C:H films shown in Fig. 1~b!. The films did not contain large particles as did a-C films, but their heterogeneous morphology13 leads to an increase of R a values from 2 nm of the substrate to 3 nm of the film, and the maximum roughness was estimated to be around 20–30 nm. Reported R a values for both a-C and a-C:H films are two orders of magnitude less than the roughness of films depos-

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FIG. 1. High-resolution SEM images of cross sections of films deposited on Si substrates: ~a! a-C film and ~b! a-C:H film. A 40° tilt in respect to the sample surface plane is used to show development of the surface roughness.

ited with an infrared 1064 nm laser.16 Shorter laser wavelengths are more favorable for the production of smooth carbon films, since they reduce the quantity of ejected particles.6 The results show, that the main contribution to the surface topography development in both a-C and a-C:H films are melted particles ejected from targets. Thus, PLD, in principle, can be used to produce amorphous films with a featureless topography, if methods eliminating particles contamination are applied.17 B. Chemical composition and density

RBS and ERS studies were undertaken to investigate film chemical composition and density. Area densities of carbon atoms are found from RBS analyses were 6.2 3 1018 atom cm22 for a-C film and 5.031018 atom cm22 for a-C:H film. From Dektak profilometer measurements the thicknesses of the a-C and a-C:H films used in RBS studies were 400 and 450 nm, respectively. Based on area density and film thickness measurements, the densities of the films were determined to be approximately 3.0 g cm23 for a-C and 2.2 g cm23 for a-C:H. They reflect the difference in the film compositions. Results of ERS studies showed that a-C films consist mainly of carbon, while a-C:H films, about 25 at. % hydrogen and 10 at. % oxygen were found. Figure 2 shows the JVST A - Vacuum, Surfaces, and Films

FIG. 2. Elastic forward recoil spectra showing hydrogen content in a-C and a-C:H films. RBS analyses was used to normalize the data. Hydrogen contamination on the both sides of a-C films can be seen.

ERS measurements of the hydrogen content as a function of depth. In a-C films, several atomic layers of hydrogen contamination were found on the front and back film surfaces ~Fig. 2!. Inside a-C films the hydrogen content was about 0.1 at. %. This result was used as a criteria for describing a-C films as hydrogen-free in the present study. For a-C:H films, possible surface contamination with hydrogen was convoluted with considerably higher hydrogen content in the bulk of the film ~Fig. 2!. It is interesting to compare between the chemical compositions a-C:H films and polycarbonate target ~43% C, 46% H, and 10% O!. Almost half of the hydrogen was ‘‘lost’’ in transferring ablated material to the film. A possible explanation could be backscattering of hydrogen atoms from film surface. Another possibility is the hydrogen preferential sputtering. This was recently suggested by Robertson18 to describe formation of hydrogenated DLC and is based on the difference in the displacement energies of carbon ~26 eV! and hydrogen ~3 eV! in a-C:H groups.19 Backscattered and sputtered hydrogen can recombine with other low-energy atoms of H, C, or O to form gas molecules that are then pumped out of the chamber.

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FIG. 3. Results of SIMS analyses of a-C and a-C:H films deposited on steel substrates. A 5 keV Cs ion beam was used to produce secondary ions from film surface.

C. Structure and interatomic bonding

In a previous study, a reflected electron energy-loss spectroscopy ~EELS! analysis was performed for valence bands of a-C and a-C:H films.13 Results did not allow determination of the exact s p 3 /s p 2 ratio, but indicated the dominance of s p 3 bonds in both films by the absence of a p-electron associated losses of 7 eV and the presence of plasmon excitation losses at energies corresponding to diamondlike materials.13 For a-C films, plasmon losses were registered at around 33 eV, the same energy as in diamond.20 Incorporation of 25 at. % hydrogen in the a-C:H film caused a shift in plasmon losses to around 29 eV.13 This shift agrees with results published earlier for hydrogenated amorphous carbon.21 Results of EELS investigations were compared with results of Raman spectroscopy.13 This comparison was used to conclude that both films have an amorphous carbon network, and a-C films are more diamondlike and possess higher disorder in their microstructure.13 Thus, hydrogen lead to a less diamondlike structure, when comparing the microstructure of a-C and a-C:H films. An important question is what type of chemical bonding hydrogen has in a-C:H films. SIMS analyses showed the presence of CH and CH2 groups sputtered from the film surface ~Fig. 3!. Note that 13C and 12CH could not be resolved in the instrument used for the analyses. To determine CH groups, a comparison was made between intensities of masses 12 and 13 in SIMS spectra for a-C and a-C:H films. For a-C films, the ratio of 13C/12C intensities was close to the natural isotopic abundance ratio of 0.011. For a-C:H films, this ratio is much higher ~Fig. 3!, indicating the presence of CH groups. The combined results of SIMS and EELS indicated that s p 3 C–H bonds are present in a-C:H films. This was confirmed by analyzing FTIR absorbance spectra for a-C:H films shown in Fig. 4. Characteristic absorption occurred at wavelengths corresponding to sp 3 CH3 J. Vac. Sci. Technol. A, Vol. 14, No. 3, May/Jun 1996

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FIG. 4. FTIR absorption spectra for a-C and a-C:H films recorded in a reflective mode from films deposited on steel substrates.

stretchings22,23; O–CH3 groups were also observed. This stretching indicated the presence of unbroken polymer chains transferred to the film surface with splashed particles. For a-C films, both SIMS and FTIR analyses did not reveal CH2 , CH3 , or O–CH groups ~Figs. 3 and 4!, verifying that 0.1 at. % of hydrogen does not produce detectable changes to the interatomic arrangements in a-C films. The results of chemical analyses and structural investigations show films deposited from graphite to be amorphous hydrogen-free DLC with dominance of sp 3 bonding. In contrast to other researchers, we deliberately do not term these films as ‘‘nanocrystalline tetrahedral carbon,’’24 or ‘‘nanophase diamond.’’25,26 Films produced by PLD at low temperatures are amorphous in nature, i.e., they lack longrange ordering necessary to claim crystallinity. Films deposited from polycarbonate targets can be described as amorphous hydrogenated DLC with mainly sp 3 bonding in CH3 configurations and some amount of polymer O–CH3 groups, which were probably transferred from the target in splashed microparticles. The difference in composition and structure of a-C and a-C:H films was reflected in film thermal stability and mechanical properties discussed in the following sections. D. Thermal stability

Thermal stability of films deposited onto steel substrates was investigated with XPS. Changes in the position of the C 1s peak have been monitored as a function of temperature ~Fig. 5!. The shift of binding energies between a-C and a-C:H films is within instrumental error associated with sample positioning. An assumption was made that shifts in binding energies after samples were installed and data accumulation began are due to structural modifications. Abrupt changes in binding energies occurred after 350 °C for a-C:H films and after 500 °C for a-C films, indicating sp 3 – sp 2

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FIG. 5. Variation of the C 1s binding energy for a-C and a-C:H films with annealing temperature in in situ XPS analyses of film thermal stability. The relative shift of curves for a-C and a-C:H films is insignificant, since it is caused by the sample positioning. The beginning of graphitization was detected from the drops of the binding energy.

transformations. These transformations were confirmed by Raman investigations of the films before and after annealing. Graphitization was indicated by the change of a single maxima spectrum recorded for as-deposited a-C and a-C:H films13 in the range from 1200 to 1600 cm21 to a spectrum with two separated maxima located at around 1350 and 1580 cm21 for annealed films, which is characteristic of graphitelike amorphous carbon.27 From these studies, hydrogenated DLC is less thermally stable than hydrogen-free DLC. The termination of sp 3 carbon bonds with hydrogen reduces stability of the amorphous carbon network. Thermally induced losses of weakly bonded hydrogen most probably lead to the collapse of sp 3 CH3 configurations into s p 2 CH2 and CvC configurations.28 Recently Friedmann et al.29 performed in situ Raman studies of thermal stability of a-C films grown on Si with a 248 nm laser. Graphitization was found to begin at around 800 °C and degradation of film transparency started from temperatures below 100 °C.29 The difference in graphitization temperatures reported by Friedmann et al. and in this study for a-C films may be due to different substrate materials, since carbon dissolving phases in annealed 440C steel accelerate heat induced s p 3 – s p 2 transitions of the films.

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hydrogen-free sp 3 carbon bonds in a-C films. This caused a decrease in elastic modules and hardness. Incorporation of polymer bonds in a-C:H films is another reason for their lower mechanical properties. The difference in film chemical composition, structure and mechanical properties correlates with results of tribological studies of a-C and a-C:H films.14 Friction coefficients of a-C films sliding against steel and sapphire were around 0.1 in both cases, and were almost not sensitive to the change of the test environment from dry nitrogen, to humid air, and to 10 Pa vacuum.14 In contrast, friction coefficients of a-C:H films against steel varied from 0.19 in humid air to 0.26 in 10 Pa vacuum. A similar behavior was found in sliding a-C:H films against sapphire, where friction coefficients varied from 0.07 to 0.12.4 There are several reasons for the difference in the friction behavior of a-C and a-C:H films, which can be associated with their composition and structure. One is the hydrogen losses under vacuum conditions in wear tracks of a-C:H films by breaking weakly bonded CH3 groups. This process removes termination of carbon–carbon bonds and leads to an increase in friction similar to discussed, for example, by Zaidi et al.31 Another reason can be 10 at. % oxygen inside the films, which may initiate surface oxidation and affect friction. Since a-C:H film contains both hydrogen and oxygen, the results of friction studies14 reflect their synergetic effect which needs additional investigations. Another result of tribological studies,14 related to structural differences between a-C and a-C:H films, is the difference in their sliding wear mechanisms. For a-C:H films, a constant rate wear from the film surface to the substrate interface was observed. In contrast to hydrogen-free films, a-C:H films did not exhibit brittle failure on high contact loads,14 indicating their greater flexibility due to the incorporation of CH3 groups and ejected polymer particles. Their normalized wear rates of 1026 mm3 N21 m21 is comparable with those reported for hydrogenated DLC produced by other techniques.32–34 For a-C films, a friction-induced sp 3 – sp 2 phase transition was found in wear tracks.35 This resulted in the formation of an sp 2 -rich transfer layer on the counterpart friction surface after 105 cycles. This layer performs as a lubricant in humid environments and was responsible for extremely low wear rates of 1029 mm3 N 21 m21 reported for sliding wear of a-C films in an ambient environment.14 The combination of good mechanical characteristics and thermal stability of a-C films makes them superior to a-C:H films for application as protective coatings in sliding water.

E. Mechanical and tribological properties

The hydrogen termination of s p 3 bonds in an amorphous carbon network influence not only on the film thermal stability, but also their mechanical properties. Elastic modules and hardness of a-C films were found to be about 550 and 60 GPa, respectively.14 For a-C:H films, elastic modulus and hardness had considerably lower values of 150 and 14 GPa, respectively.14 The present study shows that hydrogen participated in formation of s p 3 CH3 bonds in a-C:H films, which reduced the average bond strength in comparison to JVST A - Vacuum, Surfaces, and Films

IV. CONCLUSIONS Films produced from graphite had a uniform morphology and contained impurities, including hydrogen, below 1.0 at. %. They were characterized as hydrogen-free DLC and were found to have an amorphous microstructure with mainly sp 3 interatomic bonds and the density around 3.0 g cm23 . Films produced from polycarbonate targets are best described as hydrogenated DLC. They had heterogeneous morphology created by splashed particles and contained ap-

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proximately 25 at. % H and up to 10 at. % O. Their microstructure was also amorphous with a dominance of the sp 3 -type CH3 configuration, as well as some amount of polymer O–CH3 bonds. Their density was estimated to be around 2.2 g cm23 . The presence of hydrogen and polymer bonds in a-C:H films reduced their graphitization temperature to 350 °C in comparison to 500 °C found for a-C films. This also affected film mechanical properties, reducing elastic modules and resulting in softer and more flexible films. The presence of hydrogen resulted in higher friction sensitivity of a-C:H films to the environment in comparison to a-C films. The difference in structure and mechanical properties of a-C and a-C:H films was reflected in considerably lower wear rates for hydrogen-free films in comparison to hydrogenated films. In general, the study showed that introduction of hydrogen into DLC using low-temperature PLD reduces interatomic strength of the amorphous carbon network and leads to deterioration of film stability and mechanical properties. ACKNOWLEDGMENTS The authors are pleased to thank the staff and are grateful for the use of the facilities of Wright Laboratory Materials Directorate. They especially wish to thank J. E. Bultman for help with film deposition. The authors would also like to acknowledge Dr. J. Liang for recording FTIR spectra. This work was performed while one of the authors ~A.V.! held a National Research Council USAF/WL Research Associateship. The work of another author ~J.S.S.! was sponsored by the Materials Directorate, WL/AFMC, US Air Force, Wright Patterson Air Force Base. M. A. Tamor and C. H. Wu, J. Appl. Phys. 67, 1007 ~1990!. C. Z. Wang and K. M. Ho, Phys. Rev. Lett. 71, 1184 ~1993!. 3 P. C. Keleris, C. H. Lee, and W. R. Lambrecht, J. Non-Cryst. Solids 164, 1131 ~1993!. 4 C. L. Marquardt, R. T. Williams, and D. J. Nagel, Mater. Res. Soc. Symp. Proc. 38, 325 ~1985!. 5 T. Sato, S. Furuno, S. Iguchi, and M. Hanabusa, Jpn. J. Appl. Phys. 26, L1487 ~1987!. 6 A. A. Voevodin and M. S. Donley, Surf. Coat. Technol. ~in press!. 7 A. P. Malshe, S. M. Kanetkar, S. B. Ogale, and S. T. Kshirsagar, J. Appl. Phys. 68, 5648 ~1990!. 1 2

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