Physical properties of aC: N films produced by ion ...

Physical properties of a-C: N films produced by ion beam assisted deposition Francois Rossi Commission of the European Communities, Advanced Coating Centre, P.O. Box 2, 1755 ZG Petten, The Netherlands

Bernard Andre CEREM, Centre d'Etudes Nucleates de Grenoble, BP85X 38041 Grenoble Cedex, France

A. van Veen, P. E. Mijnarends, H. Schut, and F. Labohm IRI, Delft University of Technology, Mekelweg 15, 2629 JB Delft, The Netherlands

Hugh Dunlop Pechiney CRV, BP 27 38040 Voreppe, France

Marie Paule Delplancke Service de Metallurgie et Electrochimie, Universite Libre de Bruxelles, 50 Avenue Franklin Roosevelt, CP 165 1050 Bruxelles, Belgium

Kevin Hubbard Center for Materials Science, MS K765, Los Alamos National Laboratory, Los Alamos, New Mexico 87454 (Received 7 July 1993; accepted 28 April 1994)

Carbon films with up to 32 at. % of nitrogen have been prepared with ion beam assisted magnetron, using a N2 + /N + beam at energies between 50 and 300 eV. The composition and density of the films vary strongly with the deposition parameters. EELS, SXS, XPS, and IR studies show that these a-C: N films are mostly graphitic and have up to 20% sp3 bonding. Nitrogen is mostly combined with carbon in nitrile ( C = N ) and imine ( C = N ) groups. It is shown by RBS and NDP that density goes through a maximum as the average damage energy per incoming ion increases. Positron annihilation spectroscopy shows that the void concentration in the films goes through a minimum with average damage energy. These results are consistent with a densification induced by the collisions at low average damage energy values and induced graphitization at higher damage energy values. These results are similar to what is observed for Ar ion assisted deposition of a-C films. The mechanical properties of these films have been studied with a nanoindenter, and it was found that the hardness and Young's modulus go through a maximum as the average damage energy is increased. The maximum of mechanical properties corresponds to the minimum in the void concentration in the film. Tribological studies of the a-C: N show that the friction coefficient obtained against diamond under dynamic loading decreases strongly as the nitrogen composition increases, this effect being more pronounced at low loads.

I. INTRODUCTION Diamond and diamond-like carbon films are materials of considerable interest because of their special properties: they are hard, with high thermal conductivity, and with high electron and hole mobilities.1 In particular, diamond is the hardest material known, with a hardness value about 100 GPa. In comparison, the second hardest material known is c-BN which lies far behind diamond with a hardness value of 50 GPa. Recently, some reports have been published about the interest of doping carbon films with nitrogen. The main 2440

J. Mater. Res., Vol. 9, No. 9, Sep 1994

motivation for studying C: N films was a paper by Liu and Cohen2 calculating the properties of a hypothetical carbon-nitrogen compound. The model predicts the bulk modulus values of diamond and /3-Si 3 N 4 with good accuracy. When applied to a hypothetical compound /3-C 3 N 4 , where C was replaced by Si in the /3-Si 3 N 4 structure, the results of the calculations indicated that this material should have mechanical and thermal properties similar to those of diamond. However, C3N4 is probably metastable and has not been synthesized to date. Several attempts have been made to synthesize C: N compounds, © 1994 Materials Research Society

F. Rossi et al: Physical properties of a-C: N films

using Physical Vapor Deposition (PVD) such as reactive dc magnetron sputtering,3 RF sputtering,4 and Chemical Vapor Deposition (CVD).5"8 In the PVD techniques, carbon film was deposited under nitrogen partial pressure. In the case of CVD, the carbon films were deposited from a carbon containing gas in the presence of hydrogen and nitrogen. In all cases, the maximum nitrogen concentration obtained in the films was ca. 10 at. % for CVD techniques and between 10 and 30 at. % in the case of magnetron sputtering,3 i.e., much lower than the theoretical stoichiometry of C 3 N 4 . Moreover, these films were amorphous or nanocrystalline. The samples prepared by CVD had a high hydrogen concentration and a density of about 1.2 g/cm 3 , a value similar to those obtained for a-C: H.5>6 Deposition by reactive magnetron sputtering3 leads to a density of 2.2 g/cm 3 , similar to that obtained for non-doped a-C. In this case, the electron density calculated from Electron Energy Loss Spectroscopy (EELS), as well as the diamond character, was found to increase with the nitrogen content. However, no systematic study of the nitrogen incorporation and film densification was reported. In fact, the different techniques mentioned above do not allow an independent control of nitrogen and carbon flux to the film surface. This might be a reason why the nitrogen incorporation in the films and the density of the C: N coatings were relatively low. In this paper, we present the characterization of C: N samples, from here on called a-C: N in reference to a-C: H, produced by ion beam assisted deposition. This technique allows a high flexibility in adjusting the relative N and C fluxes, as well as the energy of the nitrogen ions. We present the evolution of film density, composition, microstructure, and mechanical properties of the films as a function of the deposition parameters, i.e., nitrogen flux and the energy of the nitrogen ions. II. EXPERIMENTAL

Substrate k\\\\\\\\\\\\\\\\\\J»---r a rhnn

FIG. 1. IB AM system composed of a dc magnetron sputtering system and a Kaufman source using nitrogen gas.

bombardment of the substrate during film growth. Neutralization of the space charge was performed with a tungsten filament located in the beam. The energy of the ions varied between 50 and 300 eV, and the beam current varied between 5 and 35 mA. The beam is composed of N + and N 2 + , the ratio between the two types of ions varying with the source parameters. Hubler et al.9 reported a N + concentration varying between 2.5 and 25%, depending on the discharge voltage for conditions similar to ours. We could not measure the ratio between N + and N 2 + , but ion flux/atom flux (I/A) values were estimated for the different conditions tested by measuring the deposition rates of non-assisted samples and taking the nitrogen flux values reported in Ref. 9. B. Film characterization

A. Film deposition Samples of a-C:N were deposited on (100) Si wafers by ion beam assisted magnetron sputtering (Fig. 1). The sputtering of the carbon target was obtained with a dc circular magnetron system. This magnetron worked with a 700 W argon discharge at low pressure, typically a few 10" 2 Pa, and all the sputtering parameters remained constant in the experiments described below. The limit pressure in the vacuum chamber prior to deposition was 6 X 10~5 Pa, and the working pressure was 6 X 10~2 Pa for all cases. The deposition rate obtained during the experiments under these conditions was 2000 A/h. The temperature of the samples during deposition was always kept below 100 °C. Ion assistance was performed using a Kaufman source with 3 cm dual graphite extraction grids fed with N 2 gas. The nitrogen beam was defocused to ensure a homogeneous

Characterization of the films was made by Neutron Depth Profiling (NDP) using the following nuclear reaction between thermal neutrons and nitrogen: 14*

14

Since H and C are emitted at a very well defined energy, their energy loss can be used to determine the depth profile of nitrogen. Neutron depth profiling was used to calculate the composition of the films, which was also evaluated with Rutherford Backscattering Spectrometry (RBS) using a 1 MeV He + beam. Precise evaluations of the densities could not be obtained from mass measurements, owing to the low values of both the atomic mass of carbon and the film thickness (from 400 to 500 nm). The density of the films was thus calculated by dividing the areal densities given by RBS and NDP,


2441

F. Rossi ef al: Physical properties of a-C: N films

by the thickness of the layers. Thicknesses were measured using scanning electron micrographs of crosssectioned films. The defect concentration in the a-C: N films was studied with the aid of the positron annihilation 5-parameter method, using a beam of slow positrons (e+).w'n This method is based on the measurement of the shape of the 511 keV annihilation line from thermalized positrons annihilating in the sample under investigation. The 511 keV line is Doppler broadened as a result of the motion of electrons in the sample, and hence its shape is a measure of the momentum distribution of the electrons. The line shape is characterized by a parameter S, defined as the ratio between the area under a suitably defined central part of the peak and the total area under the spectrum line. The S-parameter is particularly sensitive to the presence of open volume defects in the sample such as vacancies or voids, which constitute a negative potential and thus can trap positrons. In a vacancy, the density of high momentum core electrons is lower than in the bulk, while the density of low momentum valence electrons is comparable to the bulk value. As a result, the S-parameter increases when open volume defects are present. The structural properties of the films were studied with X-ray Photoelectron Spectroscopy (XPS), Infra Red spectroscopy (IR), and TEM. XPS measurements have been carried out on a CAMECA Nanoscan 50 system using an incident x-ray beam (MgK a line). The sp3 hybridization content was measured by EELS. EELS was performed by measuring the energy loss of reflected electrons from the sample, using a fixed primary beam energy of 500 eV. The EELS experiments were performed under a working pressure of 2 X 10~8 Pa. Auger Electron Spectroscopy (AES) was used to examine the surface composition of the samples. In case of oxygen contamination, sputtering by a 2 keV Ar + beam was performed before spectrum acquisition. We checked that sputtering did not induce changes in the EELS spectra.

Raman spectroscopy is sensitive to changes in translational symmetry. The Raman signature is strongly dependent on the short-range structure, typically on the scale of a few A and thus is useful for the study of disorder in carbon films. The deposits were analyzed using a micro Raman spectrometer equipped with a DILOR optical multichannel analyzer working at 514.5 nm. The size of the laser spot was 50 fim and the power was limited to 60 mW. Mechanical properties were evaluated with a Nanoindenter II from Nano Instruments, and the tribological properties were evaluated with a Scratch Tester (Revetest) under dynamic loading between 10 and 30 N over a length of 10 mm. Pin-on-disk tests were performed on films deposited on stainless steel (thickness 500 nm approximately). The tests were performed using an alumina pin under a constant load of 1 N. Humidity during the tests varied between 45 and 75%. III. RESULTS A. Composition The composition of the films was studied with RBS and NDP. The experimental results obtained are presented in Table I. It can be seen that the composition of the films varies strongly with the deposition parameters. Since we varied both the ion energy and the ratio between the ion flux and atom flux, the effect of ion bombardment per deposited atom changed during our experiments. This effect can be taken into account by the average damage energy per deposited atom. / p N

p

I

(Ecoii + Eph)

I

where Ecoll is the nuclear collision energy, Eph is the phonon energy, R the ion range, and I/A the ratio of the ion and atom fluxes. Both Ed and R have been calculated by TRIM9212 using a displacement energy of 35 eV and a surface binding energy of 7.4 eV, the values generally accepted now for this type of calculation.13

TABLE I. Processing parameters and corresponding a-C: N film characteristics. / is the beam current of the ion assistance source, E the ion energy, e the coating thickness, and N the nitrogen concentration. Sample no.

/(mA)

£(eV)

e(A)

N (at. %) N R A

N (at. %) RBS

2 3 4 5 6 7 8 9 10 11 12

5 5 5 5 10 15 23 35 5 5 20

200 300 100 50 100 100 100 100 150 250 100

4500 6000 5600 6000 5000 5000 5700 6000 5500 6000 5700

21 n/a 17 22 26 26 26 28 30 28 27

18 21 14 23 23 29 32 25 25 25

2442


14

Density (g/cm 3 ) 2.23 1.94 2.18 1.85 2.30 2.31 2.09 1.97 2.09 1.92 2.14


35.00-

C. Transmission electron microscopy

30.00--

25.00-—

>-

20.00-15.00--

•

•

li NDP - • — RBS

10.00 0

1

0.5

1.5

2

2.5

E/I/A (eV/A) FIG. 2. Evolution of the N content in at. % in the film with the average damage energy during ion beam assistance.

Figure 2 shows that the N concentration of the films increases steadily with the average damage energy. Concentrations up to 30-32 at. % of nitrogen were obtained for {Ed) of 3 eV/A, which corresponds to a source current of 35 mA for a beam energy of 100 eV. The hydrogen concentration was not measured for this set of experiments, but is expected to be of the same order as the one obtained for a-C films prepared in similar conditions under Ar and Xe bombardment,10 i.e., about 5% to a maximum of 10%. B. Density of the films Figure 3 shows the evolution of density of the films with the average damage energy. The film density reaches a maximum value around {Ed) ~ 1 eV/A. Compared to Fig. 2, these results show that, within the range of experimental values tested, no relationship exists between the nitrogen content and the density of the films.

TEM observation of as-deposited samples showed a strong central background with two very diffuse rings at 1.13 and 3.1 A, corresponding to (112)graphite and (104)graphite. We did not observe a ring at 2.08 A corresponding to (lll)diamond a s m Tomg et al.4 However, we could not observe indications of microcrystallinity in the TEM pictures, even at the highest magnification (X 200 000). D. Electron energy loss spectroscopy Figure 4 shows the EELS spectra of film numbers 7, 8, and 9 of Table I. For comparison, spectra obtained with Highly Oriented Pyrolitic Graphite (HOPG) and a PACVD diamond film are reported. For HOPG, the two peaks at 6 eV and 28 eV can be attributed to the (TT - IT*) and (TT + a) electron plasma resonance, respectively.14 In the case of diamond, two plasmon peaks can be observed around 24 eV and 35 eV. In the case of our a-C: N samples, the (TT — TT*) peak can hardly be detected in the direct signal, and its position was found using the derivative of the spectrum. The position of the (TT + a) plasmon peak changes between 24 and 28 eV, indicating a low concentration of sp3 sites present. These results are similar to those reported by Torng et al. ,4 although the nitrogen content of our sample is twice as large. A calculation of the sp3 concentration in non-doped a-C: H samples was made by Wang et a/.14 The sp^/sp2 ratios can be calculated following the formula 0.23

CO2P(TT + a)

4-n-e2

-

m

CH

~ 1

where CO(TT) and co (TT + a) are the (TT — 77*) and (TT + a) plasmon energies, C H is the hydrogen con-

2.40. Diamond

2.30-|-

Sample 9

2.20-i —

g 13

2.10- -

-

2.00-1

-

Sample 8 Sample 7

_

1.90-f : a 1.80-1 • I

ffl •

1.700

0.5

1

1.5

2

NDP RBS

40

FIG. 3. Variation of the film density measured by NDP and RBS as a function of average damage energy.

50

Energy (eV)

2.5

Ed*I/A (eV/A)

_It HOPG ; _

FIG. 4. EELS spectra of a-C: N films obtained with 10 mA, 23 mA, and 35 mA nitrogen beam currents at 100 eV. The corresponding average nitrogen content of these films was 24 at. %, 27 at. %, and 30 at. %. The EELS spectra of HOPG and CVD diamond films are also presented.


2443


centration, and m and e are the mass and charge of an electron. In the present case, the above formula can be used only as an indication since TT and a electron densities are probably strongly affected by the presence of nitrogen. However, by neglecting this interaction and the possible contribution of hydrogen (limited in our case to between 5 and 10 at. %), an order of magnitude can be estimated. Values of sp3 concentrations of 13, 15, and 23% could be calculated in the three tested a-C: N samples. These values are of the same order of magnitude as the ones obtained under similar conditions for Ar and Xe assisted a-C films.10

Kinetic Energy (eV)

E. XPS and AES Figures 5 and 6 show the surface XPS carbon and nitrogen peaks of the same sample. A comparison with XPS signals obtained on pure graphite and polyacrylonitrile fibers could be made. 15 From this comparison, the nitrogen signal in our samples can be deconvoluted into two peaks at 398.2 and 400.2 which can be attributed to N = C (nitrile) and N = C (imine), respectively. The carbon signal in a-C: N is very similar to pure graphite except for the presence of a new peak at 288.4 which can be attributed to N 2 C. These results could be confirmed by IR analysis after dissolution of the film in KBr. The spectrum obtained on one sample (density 2.1 g/cm 3 , nitrogen concentration of 25 at. %) shows three isolated peaks at 1220, 1390, and 1625 cm" 1 , which could be attributed to C—N, sp3 C—C, and C = C , C = N , respectively.3-4 Figure 7 shows the comparison of the Auger spectra of HOPG, diamond, and a-C: N sample no. 7. In comparison with HOPG, we observe the appearance of a new peak in the a-C: N sample similar to the one observed in diamond. This was already observed

500 • 400--

SB

I

I

284.5 eV

f~

( 286.3 eV -vf

H h

294 292 290 288 286 284 282 280 278

Kinetic Energy (eV) FIG. 5. XPS carbon peak of a C: N sample prepared with a beam energy of 100 eV and a beam current of 15 mA. The main peaks at 284.5 and 286.3 are similar to the peaks obtained in pure graphite. The peak at 288.4 could be attributed to N2C. 2444

408 406 404 402 400 398 396 394 392

FIG. 6. XPS nitrogen peak of the C: N sample presented in Fig. 7.

in Ref. 4 and was interpreted in terms of the presence of two phases (graphitic and diamond-like). Moreover, when considering the difference D between the position of the maximum of the positive excursion and the negative excursion in these spectra, we find that D is equal to 19.7 eV in HOPG and 14.1 eV in a-C:N. Comparison of these values with those for diamond16 for which D ~ 14.2 indicates that our a-C: N samples have a well-defined sp3 bonding character. This is consistent with the previous XPS and IR results where sp2 and sp3 bonding in the carbon phase could be observed. F. Raman spectroscopy Raman spectroscopy is often used to characterize the crystallinity of diamond and graphite thin films.17 This technique has been applied to a-C: N sample nos. 5, 9, and 10. Figure 8 shows the Raman spectra of the three samples. These Raman spectra are representative of disordered graphite18"20 and very different from the results of Kaufman et al.21 The experimental results are best fit by four Gaussians and a constant background. Two D and G peaks can be observed about 1350 and 1580 cm" 1 and their intensities vary with the film preparation conditions. These results can be interpreted by a comparison with graphite and diamond spectra. The Raman spectrum of crystalline diamond consists of a single sharp peak at 1332 cm" 1 . The Raman spectrum of crystalline graphite consists of two peaks: the G peak centered on 1550 cm" 1 is the zone center E2g mode of the perfect graphite crystal and the D peak centered on 1350 cm" 1 is a zone edge Aig mode activated by disorder in the graphite crystal.22'23 The D mode is a common feature of disordered graphitic carbon. Its intensity relative to the G mode (as measured by the ratio ID/IG) changes with the disorder.24 Tuinstra and Koenig23 have shown that the ratio ID/IG is inversely proportional to the graphitic crystallite size La as measured by XRD. In mixed sp2sp3 bonded carbon



0.0 2000

1800

1600

1400

1200

1000

Raman shift (cm'1) —i

/i A

3

a-C:N:

MM

M

.

c o

—i

/

V

100.0

^

—i-

0.0 2000

1800

1600

1400

1200

1000

Raman shift (cm'1)

—i 100.0

200

220

240

260

280 300

Kinetic Energy (eV)

0.0 2000

1800

1600

1400

1200

1000

Raman shift (cm 1 )

FIG. 7. Auger electron spectroscopy of HOPG, a-C: N (sample no. 7), and diamond.

FIG. 8. Raman spectra of a-C: N samples no. 5, 9, and 10.

layers the overall Raman spectrum is dominated by the G component, because the cross section of the graphite stretching mode is much higher than that of the 1332 diamond mode. From the fitting parameters obtained from the spectra of Fig. 8, the peak position, peak width, and integrated intensity ratio of the two G and D peaks can be obtained. Taking the relationship proposed by Tuinstra and Koenig23 between La and ID/IG leads to small

values of La which are plotted versus density in Fig. 9. The highest density corresponds to the lowest crystallite size. Moreover, the ratios T of the halfwidths of the G and D peaks are indicative of strong disorder within graphitic bonding and a large variation in bond angle.18 The highest degree of disorder is found for sample no. 10 with a value of T equal to 0.46 to be compared with 0.89 and 0.48 for sample nos. 5 and 9, respectively. This result is consistent with our previous results obtained


2445

F. Rossi et at: Physical properties of a-C: N films

80.0-

10.0 1.80 1.85 1.90 1.95 2.00 2.05 2.10 2.15 2.20 0

Density (g/cm3) FIG. 9. Graphitic crystallite size of samples 5, 9, and 10 calculated from the Raman spectra as a function of density.

with Ar assisted a-C growth where the increase of density was accompanied by a decrease in the graphitic cluster size and an increase of disorder.10 G. Defect microstructure Positron annihilation spectroscopy has been used to evaluate the void density in the samples. The ^-parameter measured is an increasing function of the void concentration. Figure 10 shows the S-parameter value as a function of the average damage energy. Although the scattering of the results is large, it can be seen that the S-parameter reaches a minimum for a value of 1 to 1.5 eV/A. This is approximately the value for which density reaches a maximum. For comparison, the value of the relative S parameter for pyrolitic graphite is 0.925. In the experimental range tested, the a-C:N layers have a concentration of voids similar to graphite, as well as a similar density.

0.5

1.5

E *I/A (eV/A) d FIG. 11. Young's modulus of the a-C:N films measured by nanoindentation as a function of the average damage energy.

H. Mechanical properties The mechanical properties of the a-C:N samples were studied with a Nanoindenter II from Nano Instruments. Figure 11 shows Young's modulus E of our a-C:N films as a function of average damage energy. Here again, E reaches a maximum about 1 to 1.5 eV/A, the range of values for which the density is maximum, and the defects concentration minimum. Figure 12 shows the relationship between Young's modulus and hardness for a-C: N films. For comparison, values obtained for a-C films deposited under Ar and Xe bombardment and similar conditions10 are displayed. The hardness values obtained for a-C: N are lower than those obtained for diamond-like carbon films deposited with Ar assistance under similar conditions. A linear dependence between E and H is found in both cases with a slope dH/dE ~ 0.13 for a-C: N. This value is slightly higher than the one found by other authors for a-C: H and a-C films (H/E = 0.11).25'26 From semi-empirical

0.940

40- a-CN

35--

o

-a-C

." - . * ! : .

30--

H

15150

E *I/A (eV/A) d FIG. 10. S-parameter value, measured by positron annihilation spectroscopy as a function of the average damage energy. The S-parameter is a direct indication of the concentration of voids in the film. 2446

200

250

E (GPa) FIG. 12. Relationship between hardness and elastic modulus of the a-C: N films measured by nanoindentation. For comparison, results obtained on a-C films prepared by ion beam assisted deposition10 under Ar and Xe assistance are also presented.


F. Rossi ef a/: Physical properties of a-C: N films

0.14

15

20

25

30

35

Nat% FIG. 13. Average friction coefficient of a-C: N against a diamond tip under dynamic loading, as a function of nitrogen concentration.

considerations, a value of H/E ~ 1/6 was estimated by Robertson.1 High H/E values are an indication of a good wear resistance of the material,27'28 and from this point of view, a-C: N films should have a better friction behavior than non-doped films. I. Tribological properties Tribological properties of the C: N films were studied with a scratch tester equipped with a friction measurement outfit. The friction force during the scratch test is measured continuously against a Vickers diamond pyramid (radius 200 yu-m), as the load is increased progressively up to 30 N. The friction coefficient was measured continuously between 10 and 30 N load. In all cases, the friction coefficient measured increased linearly with the load. The values of the friction coefficient against diamond (at 10, 20, and 30 N load under dynamic loading) are plotted in Fig. 13 as a function of nitrogen concentration. For clarity, only the RBS measurements are displayed in the plot. It can be seen that in all cases, an increase in nitrogen content produces a decrease in the friction coefficient. No obvious relationship between the friction coefficient and hardness or elastic modulus can be found. The tribological properties of these films were also studied with a pin-on-disk test against an alumina ball of 5 mm diameter, under a load of 1 N at a velocity of 4 m/min. The friction coefficients obtained were between 0.15 and 0.20. For comparison, the friction coefficient under similar conditions for a-C coatings obtained by DIBS with Ar assistance was between 0.06 and 0.15. 29 IV. DISCUSSION The effect of nitrogen ion beam assistance on the properties and composition of the a-C: N films can be interpreted by a ballistic effect due to ion bombardment. These results are similar to our previous results of

ion beam assisted deposition of diamond-like carbon films.10 The incorporation of nitrogen in the films shown in Fig. 2 is much easier than in the case of Ar or Xe bombardment, since values of 20 to 30 at. % nitrogen can be easily obtained, to be compared to 1 to 5% for Xe and Ar.10 This can be explained by the fact that at low energies the sputtering rate, as calculated by TRIM12 for Ar or Xe incorporated in the film, is two to three times larger than for nitrogen. It is thus much easier to enrich the growing layer in nitrogen than in Ar or Xe. It seems also that by using larger fluxes, a much greater concentration of nitrogen could be reached. The changes in density illustrated by Fig. 3 show the similarity between the ballistic effect in these experiments and densification under Ar + bombardment in similar conditions. Comparison between Figs. 2 and 3 indicates that nitrogen content and density are not related. The densification mechanism at low ion energies can be attributed in large extent to a decrease in the void concentration. This effect can be clearly seen in the PAS results (Fig. 10) where the changes of the S parameter correspond to a variation of the void concentration in the film. This is an ion pinning effect already reported in the literature.31 Without ion bombardment, the carbon phase is mostly composed of graphitic domains embedded in a disordered phase. The ion bombardment produces displacement of the atoms in the growing film and increases the disorder as described, for instance, by Kaukonen and Nieminen31 or Tamor and Wu.32 The disorder induces cross linking between graphitic planes by sp3 bonds which can explain the observed increase in density, hardness, and the corresponding sp3 concentration. This strongly disordered phase grows at the expense of the sp2 bonded domains and reduces their size, as can be seen from the Raman results (Fig. 9). At larger damage energies, the film density decreases (Fig. 3) and the void concentration increases (Fig. 10). A possible explanation could be an effect similar to what was observed in the case of Ar assisted a-C, where the decrease of density could be attributed to a damage induced graphitization of the films. This idea is supported by the fact that the maximum of density occur§ for N + , Ar + , and Xe + at energies where vacancy production, as calculated by TRIM 12 , begins to be significant, i.e., a few percent.10 Because carbon has so few electrons, the energy associated with disorder and in particular bond angle distortion is very large. It is possible that excessive disorder provides the energy necessary to break the sp3 bonds and graphitize the structure during the defect recombination phase. The decrease in densification efficiency at higher energies could thus be explained by damage-induced graphitization of the films under ion bombardment. Such a mechanism has also


2447


been invoked to explain conductivity changes in ion irradiated a-C:H. 33 After incorporation, nitrogen is bonded to carbon by single, double, and triple bonds as shown by XPS and IR. This result is similar to the observations of Ricci et al.6 or Kaufman et al.21 This incorporation influences principally the friction coefficient, as shown in Fig. 13. It would be interesting to know whether nitrogen doping increases the hardness or Young's modulus. Our results show that the mechanical properties of the films are mostly influenced by the densification changes induced by the bombardment, and our experiments do not allow a discrimination between this effect and the influence of nitrogen. It can be said, however, that this chemical effect, if it exists, is probably of the second order compared to ballistic effects of ion bombardment. We cannot infer from our results an increase of the diamond-like character of the films related to the incorporation of nitrogen, as was proposed by Torng et al.4 As in Ref. 4, our EELS results show a disappearance of the (TT-TT*) plasmon at 6 eV, but this result was already observed in a-C samples prepared under Ar or Xe bombardment where no nitrogen was incorporated. In the same way, the AES spectra of a-C: N films show similarities with those of diamond, indicating the possible presence of sp3 bonding, also suggested by IR and EELS. In fact, sp3 bonding formation and nitrogen incorporation seem to be related to each other only through the ion bombardment effects, and it is not possible to conclude from our results that nitrogen stabilizes the sp3 bonding, as was suggested by Torng et al. In fact, it is expected that nitrogen adds w electrons and has an interaction with the a bonding, but a stabilization mechanism of the sp3 bonding similar to the one observed with atomic hydrogen in diamond formation is difficult to imagine in our case. The beneficial effect of nitrogen on the friction coefficient against diamond is illustrated in Fig. 13. This effect is more pronounced under low loads owing to the fact that the thickness of the films is small and the influence of the substrate is more important at high loads. The friction coefficient obtained between a-C: N and alumina is greater than the one reported between a-C and alumina29 under similar conditions owing to the fact that the hardness of the latter was much larger than in the case of the present samples (ca. 30 against 20 GPa). Moreover, under low loads, surface effects related to water adsorption are more important, and adsorption mechanisms for these two materials could be very different. V. CONCLUSION We have prepared carbon films with up to 30% nitrogen concentration with IBAM, by using a N 2 + / N + beam at energies between 50 and 300 eV. From the 2448

XPS studies, we have found that nitrogen was bonded in imine ( C = N ) and nitrile ( C = N ) groups. It was found that the film properties were governed by damage energy deposition during irradiation. These results can be interpreted as a combination of ballistic effects leading to densification for low ion energies, and graphitization at higher ion energies. The changes in the mechanical properties are related to a variation in the void concentration, but no direct relationship between the nitrogen content and the materials properties (diamondlike character, sp3 content, . . . ) could be found, except for the friction coefficient which decreases as the N concentration increases. Further work is needed to reach larger nitrogen concentrations with the purpose of preparing a pure carbonitride compound which has been predicted to exhibit properties similar to diamond. ACKNOWLEDGMENTS We would like to thank Martine Ricci (University of Bordeaux) for IR measurement, Philippe Renaux of CENG for pin-on-disk tests, and Eric Anger of LIMHP (University Paris Nord) for Raman investigations on our samples. Scientific discussions with H. del Puppo, W. Gissler, and J. Haupt, of the Joint Research Centre of the European Communities, as well as with M. Nastasi of Los Alamos National Laboratory, are gratefully acknowledged. This work was carried out within the Research Programme of the European Communities.

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