Surface & Coatings Technology 206 (2012) 3467–3471

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Microstructure and mechanical properties of graphitic a-C:H:Si films

F. Zhao a,b, H.X. Li a, L. Ji a, Y.J. Wang a, Y.F. Mo c, W.L. Quan a, Q.H. Kong a, Y.X. Wang a,J.M. Chen a,⁎, H.D. Zhou a

a State Key Laboratory of Solid Lubrication, Lanzhou Institute of Chemical Physics, Chinese Academy of Sciences, Lanzhou 730000, P. R. Chinab School of Materials Science and Engineering, Henan University of Science and Technology, Luoyang 471003, P. R. Chinac School of Chemistry and Chemical Engineering, Guangxi University, Nanning 530004, P. R. China

⁎ Corresponding author. Tel.: +86 931 4968018.E-mail address: chenjm@licp.cas.cn (J.M. Chen).

0257-8972/$ – see front matter © 2012 Elsevier B.V. Alldoi:10.1016/j.surfcoat.2012.02.013

a b s t r a c t

a r t i c l e i n f o

Article history:Received 8 June 2011Accepted in revised form 8 February 2012Available online 16 February 2012

Keywords:GraphiticCarbon filmSi-dopedXPSRaman

A series of graphitic a-C:H:Si films with different Si content were prepared by altering the sputtering currentin a hybrid RF-PECVD and magnetron sputtering system. Microstructures and mechanical properties of themwere characterized by IR, Raman, XPS, nanoindentation and scratch tests. Results show that although the sp3/sp2 ratio increases with increasing Si content and as high as 8.2 at.% of silicon was doped, the a-C:H:Si filmsremain graphitic in nature and the ID/IG ratio is nearly constant for all. The coupling effects of sputtering-induced heating and strong ion bombarding due to negatively biasing were considered to be responsiblefor the film graphitization. The graphitic nature also accounts for the lower nanohardness of prepared a-C:H:Si films than the diamond-like a-C:H and a-C:H:Si films.

© 2012 Elsevier B.V. All rights reserved.

1. Introduction

Diamond-like carbon (a-C:H) films have attracted particular at-tention from both industry and research communities due to thecombination of outstanding mechanical and tribological properties[1,2]. However, tribological applications of them are usually restricteddue to the high compressive stress produced in energetic depositionprocesses and the humidity-dependent friction behavior [2,3]. Con-siderable efforts have been taken by researchers to solve this prob-lem. Among them, incorporation of silicon into amorphous carbonmatrix were extensively studied [2,4–27] because it not only reducedinternal stress [16] and lowered humidity sensitivity of friction coef-ficient [2,10,28], but also improved other properties of a-C:H films, in-cluding sp3-C fraction [25], thermal stability [12], wear and corrosionresistance [5,17,29], and tribological performance in water [19,21,30].Si-doped a-C:H (a-C:H:Si) films present promising application poten-tial as environmentally adjustable protective and lubricative coatings.

Tribologically, the low friction performance of a-C:H:Si films wasgenerally attributed to the low shear strength of transferred siliconoxide layer [5,29]. It is completely different from the lubricatingmechanism of diamond-like a-C:H films which mainly rely on gra-phitic transfer layer to reduce interfacial shear strength and friction[31–34]. However, the absence of graphitic transfer layer nevermeans the graphitization is useless in reducing friction for a-C:H:Sifilms. On the contrary, hydrocarbon transfer layer generated under

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specific conditions can lead to an ultralow friction behavior [4]. Hence,to investigate the friction behavior by combining the lubricating effectsof graphitic carbon and silicon oxidewill be a fascinatingwork. However,the graphitic transformation of diamond-like a-C:H films due to thefriction-induced annealing and shearing effects [32] would be very diffi-cult to take place for diamond-like a-C:H:Si films because of their higherthermal stability [12]. In general, before temperature rising high enoughto initiate graphitization (H release and subsequent lattice relaxation[32]), oxidization of silicon might have happened. So it is hard to inves-tigate the synergistic lubricating effects of graphitic carbon and siliconoxide by relying on the tribochemical reactions during sliding tests.One feasible method is to prepare graphitic a-C:H:Si films inherentlycontaining graphitic carbon and silicon oxide under special plasma con-ditions and then investigate their tribological performance.

In this paper, we introduce the preparation of graphitic a-C:H:Sifilms with different Si content and then characterize them from com-position, microstructure and mechanical properties. After comparisonwith previous studies, the origin of graphitization is also discussed.Studies on the lubricating effects and mechanism of graphitic a-C:H:Si films will be reported in another paper.

2. Experimental details

a-C:H:Si films were deposited on polished stainless steel(1Cr18Mo8Ni5N) substrates and n-type Si (100) wafers by a hybridradio frequency plasma-enhanced chemical vapor deposition (RF-PECVD) and unbalanced magnetron sputtering (UMS) technique.The system has been described in detail elsewhere [35]. An inductivelycoupled radio frequency (13.56 MHz) system was used to excite the

Fig. 1. Si-doping content versus sputtering current.

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CH4/Ar plasma and control the ion density or ion flux. A mid-frequency(40 KHz) AC power supply, connected to a pair of dual-magnetron tar-gets, was used to dope Si in a current constant mode. A separate pulsepower supply (DC, 40 kHz) was applied to bias the substrate and accel-erate the ions over the sheath. Hence, the ion density (or ion flux) andion-bombarding energy can be controlled independently. A mixture ofCH4 and Arwas used as the source gas of RF-PECVD to synthesize amor-phous carbon film into which Si was doped by sputtering the Si targets.Detailed deposition parameters are described in Table 1. Prior to depo-sition, substrates were firstly etched with Ar+ plasma at a bias voltageof−1000 V for 30minutes to remove the native oxide layers on samplesurfaces. And then an adhesive interlayer of Si, about 200 nm thick, wasdeposited. The deposition duration is 90 min,with an overall film thick-ness of 800–900 nm. Non-doped a-C:H film was also prepared underthe same condition for comparison.

A PHI-5702 X-ray photoelectron spectroscope (XPS) operatingwith monochromated Al-Ka radiation at a pass energy of 29.35 eVwas employed to analyze the chemical states of some typical ele-ments in the film surfaces, with the binding energy of Au (Au4f:84.0 eV) as the reference. The resolution for the measurement of thebinding energy is±0.3 eV. Prior to each XPSmeasurement, the coatingsurface was etched by Ar+ions for more than 1 minute. A LabRAMHR800 (HORIBA Jobin Yvon, France) micro-Raman spectrometer oper-ating with 532 nm Ar+ laser as the excitation source was used to char-acterize the film structure. The acquired Raman spectra were fittedbased on Gaussian curve shapes with a curve-fitting software. FourierTransformed Infrared (FT-IR) spectra were obtained with a Bruker IFS66v/S spectrophotometer in the domain 400–4000 cm−1 with a resolu-tion of below 0.23 cm-1.

Scratch tests were carried out on an UMT-2MT tribometer (CETR,USA) to evaluate the film-to-substrate adhesion strength. A prism-like diamond micro-blade with a radius of 0.4 mm was used as thescratch tool. All tests were carried out at a speed of 1.5 mm/min anda load-increasing rate of 1 N/min. Film failure was determined by anacoustic emission detector and the load at which fracture took placewas defined as the critical load.

The hardness and elastic modulus values of the films were deter-mined using an NanoTest 600 nanomechanical tester (MicroMaterialsLtd., UK), where the maximum indentation depth was controlled tobe about 80 nm (less than 10% of film thickness) so as to minimizethe effect of substrate [36], and five repeated indentations weremade for each sample.

3. Results and discussion

3.1. XPS analysis

Fig. 1 shows the variation of Si content (atomic concentration, Si/(C+Si+O)) as a function of the sputtering current. As seen, the Sicontent is roughly linearly proportional to the sputtering current, in-creasing from~3.2 at.% at 2 A to~8.2 at.% at 8 A. However, for a-C:H:Sifilms, the mechanical and tribological properties are more deter-mined by the chemical states than the content of the doped Si [15].The bonding details can be found from the Si2p XPS spectra in

Table 1Deposition parameters for a-C:H:Si films.

Parameter Value

Working pressure, Pa 0.5Mass flow rate ratio of CH4 to Ar (CH4/Ar) 5/6RF power, W 1500Pulsed bias, V −800Duty cycle, % 50Sputtering current (Si target), A 2; 4; 6; 8

Fig. 2. As reported previously [25,37–40], the peaks nearby 99.6-99.9 eV, 100.3 eV, 101.3 eV, 102.3 eV and 103.2-103.5 eV wereassigned to Si-Si, Si-C, Si-OxCy, Si-Ox (xb2) and Si-O2, respectively.Therefore, it can be induced from Fig. 2 that most of the doped Siatoms are bonded to C atoms and only a small amount of them existin the form of Si-Si or Si-O bonds. The contribution of Si-O bonds isnegligible until the sputtering current reaches a certain value (~ 4 Ain the present case) above which the Si-O shoulder peaks have be-come clearly noticeable.

Chemical states of carbon atoms in a-C:H:Si films can be revealedby the XPS spectra of C1s core level. As seen in Fig. 3(a), the curveshape becomes more asymmetric as the sputtering current increases.In particular, a shoulder peak at~285.2 eV is clearly discerned at 8 A.In addition, the increased sputtering current also leads to the promi-nence of curves at~283 eV. Both phenomena imply the coexistence ofdifferent chemical states of C atoms. Hence, it is necessary to fit thecurves in order to obtain the bonding details of C atoms. InFig. 3(b), the deconvoluted peaks at 282.9±0.1 eV, 284.2±0.1 eV,285.1±0.1 eV, 286.5±0.1 eV and 288.5±0.1 eV are attributed to C-Si, C=C, C-C/C-H, C-O and C=O, respectively [25,40,41]. Fig. 3(c)shows the variation of sp3 / sp2 ratio as a function of the sputteringcurrent. As seen, the increased Si content generally leads to anincrease in the overall sp3-C fraction (sp3-hybridized C-C, C-H and C-Si). This is consistent with previous studies because Si forms onlysingle bonds and could open up the six-fold aromatic rings and furtherreplace the carbon atoms in the sp2-bonded carbon clusters [20,42].However, it is worth noting in Fig. 3(c) that, both the sp3(C-C+C-H)and the overall sp3-C fraction shows an aberrantly abrupt increaseat 8 A. This should not arise from the increase in the amount of dopedSi or the C-Si bonds since no sudden increase in Si content between 6A and 8 A is observed in Fig. 1. It is more likely due to the hybridizationtransformation from C=C to C-H because while Si is doped into the

Fig. 2. XPS Si2p spectra for a-C:H:Si films.

Fig. 3. (a) XPS C1s spectra for a-C:H:Si films; (b) Typical deconvolution of C1s peak;(c) sp3-C/sp2-C as a function of sputtering current.

Fig. 4. FT-IR spectra for a-C:H:Si films.

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carbon matrix a considerable amount of hydrogen is simultaneouslyintroduced and more hydrogenated sp3-C are generated as well [24].Further information could be found from the IR and Raman spectrabelow.

3.2. FT-IR analysis

FT-IR spectra are widely used to determine the local bonding of C,H and Si sites in different a-C:H:Si films. As seen in Fig. 4, the twostrong peaks located at~1108 cm-1 and~740 cm-1 are assigned tothe stretching modes of Si-O-C and Si-C, respectively [23]. The en-hanced intensity of both peaks with increasing sputtering current in-dicates that more Si-O-C and Si-C bonds are generated at a highersputtering current. The peak centered at 2100 cm-1 is attributed to theSi-H stretching mode [15]. It is hardly distinguished until the sputteringcurrent reaches a higher value of 8A. Thismaybe attributed to the higherthermodynamic stability of Si-H bonds than the Si-C bonds, as reportedin [15].

The peaks at~2855 cm-1 and~2920 cm-1 are assigned to the sym-metric and asymmetric stretching mode of sp3 CH2 [43], respectively.And the peak at~1450 cm-1 is attributed to the asymmetric bendingmode of sp3 CH2. These three peaks are apparently stronger than

the peaks at~2960 cm-1, ~ 2870 cm-1 and~1380 cm-1, which corre-spond to the symmetric, asymmetric stretching mode [44] and thesymmetric bending mode of sp3 CH3, respectively. This suggestsmore sp3 CH2 than sp3 CH3 are present, or in other words, carbonchains with certain length are present.

The wide band nearby 1595 cm-1 should be assigned to C=C inaromatic rings. Although the vibration peak of C=C in aromaticrings usually appears at~1580 cm-1 [43], it would shift to a higherfrequency of~1600 cm-1 and be broadened due to the rigidity of thesp3-C matrix connecting the π-bonded clusters as well as the interac-tion between the C=C vibration and the adjoining sp2 C-H vibration[44].

It should bepointed out that, although Si incorporation is consideredto suppress the formation of aromatic structures [15], the peak intensitycomparison among filmswith different Si content is not obvious in pre-sent study and the films remain graphitic even with a Si content of ashigh as 8.2 at.%. It seems as if the carbon film structure is mainly deter-mined by the plasma condition and the effects of doped Si or H are fairlyfinite.

3.3. Raman analysis

Fig. 5(a) shows the Raman spectra of a-C:H:Si films. A non-dopeda-C:H spectrum is also given for comparison. As seen, the line shapeof a-C:H:Si films is obviously different from that of the non-dopeda-C:H film. For a-C:H, only a peak at~1542 cm-1 corresponding to Gpeak could be discerned and the D peak at a lower frequency is almostundetectable. But for a-C:H:Si, both peaks are strongly protuberant.All of them display a D peak at ~1365 cm-1 and a G peak at~1593 cm-1. The increase in D peak intensity and the up-shift of Gpeak position have already been observed during the post-deposition annealing of amorphous carbon films [45–47], and this isusually due to the removal of bond-angle disorder and the growthof graphitic clusters [45,46]. Hence, one may consider the possibilityof the film surface graphitization caused by the heating effect of thelaser beam. Actually, this phenomenon would hardly occur because,on the one hand, the laser power is controlled to be very low (lessthan 0.3 mW), and on the other hand, if graphitization has takenplace for Si-doped films, the same process must have also happenedfor the a-C:H film, since the latter is thermally less stable [26]. Butno graphitization is observed in the a-C:H film. Therefore, the Ramanspectra in Fig. 5(a) not only reveal the real structures of a-C:H:Si filmsbut also reflect their graphitic nature because the stand-out of Dpeaks, the up-shift of G peaks to 1600 cm−1 and the narrowing ofboth peaks are all the typical features of nanocrystalline graphite or gra-phitic carbon films [48,49]. Furthermore, the constant position and thenearly equivalent intensity and FWHMs (full width at half maximum)of G peaks also prove the graphitic nature since the G peak does not dis-perse in graphite or nanocrystalline graphite, although it usually

Fig. 5. (a) Raman spectra for a-C:H:Si films. Spectrum of a-C:H prepared under identicalconditions without sputtering is given for comparison; (b) Typical deconvolution ofRaman spectrum and peak intensity ratio (ID/IG) in each film.

Table 2Deposition parameters and film nature.

Parameter Film nature

Pressure/Pa CH4/Ar RF power/W Bias/V Sputteringcurrent/A

0.5 5/6 1500 −800 2 Graphitic0.5 5/6 1500 −800 0 Diamond-like0.5 1/2-5/2 1500 −600 3 Diamond-like

Fig. 6. Nanohardness and critical load as a function of sputtering current.

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disperses in more disordered carbon and the dispersion is proportionalto the degree of disorder [48].

In addition, the increased Raman background slope for the film at8 A further reveals a higher H concentration in it since hydrogen in-corporation could lead to the saturation of nonradiative recombina-tion sites (e.g., dangling bonds) [20] and result in an increase in thephotoluminescence intensity of the vis-Raman spectra [14,20,50,51].

Fig. 5(b) shows the typical deconvolution of Raman spectrum andthe peak intensity ratio (ID/IG) in each film. As shown, the ID/IG is al-most independent of the sputtering current, around 1.7 for all films.This further reveals the finite effect of Si incorporation on the graphiticnature of a-C:H films produced under specific plasma conditions.

Therefore, it is necessary to discuss the formation of graphitic struc-ture in present plasma deposition process. As mentioned above, the iondensity (or ion flux) and ion-bombarding energy (or kinetic energy ofions) can be controlled independently by RF power and negative bias,respectively. Under present condition (1500 W and 0.6 Pa), the plas-ma was measured to have a density of~1012 cm-3 and an electrontemperature of several eV. However, once entering the sheath zone,the CHx ions would be considerably accelerated by the high negativebias (−600 to −800 V) and obtain an energy of hundreds of eV per Catom when arriving at film growth surface. Hence, major bombardingenergy of incident ions was obtained from biasing acceleration over thesheath and this energy can be controlled easily by adjusting the negativebias.

According to previous studies, the highest sp3 fraction for a-C:Hfilms was generated at around 100 eV per C atom [43]. Once exceedingthis critical value, the formation of sp2 bonding would be promoted.However, in our experiments, graphitization did not always take placeunder a high bias voltage. As shown in Table 2, even with the same RFpower (1500W) and high negative bias voltages (–600 and –800 V),the prepared a-C:H or a-C:H:Si films remained diamond-like.Moreover,all those diamond-like carbon films contain a dominant amount of sp2

bonding rather than sp3 bonding. As a result, it is unsuitable to judge

graphitizationmerely fromestimation ormeasurement of C ions energyor sp2 fraction in films. The graphitization seems to be more controlledby the order, mainly aromatic ordering, rather than the fraction of sp2

sites [43,52].One can see from Table 2 that both sputtering current and bias

voltage are responsible for the formation of graphitic carbon struc-ture. As far as the bias voltage is concerned, it could accelerate the in-cident ions over the sheath and lead to a larger thermal spike volumeand a better relaxation of stress [43,53–55], thus resulting in the for-mation of sp2 bonding which is more stable thermodynamically. Bycomparison, sputtering appears to be impossible to enhance the ener-gy of ions considerably since it could only produce active particleswith relatively low energy. However, sputtering was able to producea large amount of heat and lead to remarkable rise of chamber tem-perature. As observed, the chamber temperature had risen frombelow 65 °C to above 120 °C after the application of sputteringpower. Such a heating effect will increase the diffusion mobility ofsp2 sites and promote their segregation into larger graphitic clusters[46], apart from strongly affecting the rule of highest sp3 fraction at~100 eV per C atom [54–56]. Therefore, the graphitization in presentcase is mainly due to the coupling effect of heating and strong ion bom-barding brought about by sputtering and negatively biasing, respectively.Moreover, our other studies further reveal that all parameters affectingthe ion energy and temperature, such as RF power, affect the nature ofcarbon film (diamond-like or graphite-like) as well.

In spite of their obvious graphitic feature, however, no graphiticinformation could be detected from the XRD spectra, TEM micro-graphs and selected area electron diffraction patterns. It may be dueto the small size of graphite nanocrystalline or the special carbonfilm structure (graphitic clusters are surrounded by amorphous struc-tures). Further investigation on this phenomenon is needed.

3.4. Mechanical properties

Fig. 6 shows the variation of nanohardness and critical load as afunction of the sputtering current. Firstly, compared with non-doped a-C:H films [57] and diamond-like a-C:H:Si films [35], the pre-sent a-C:H:Si films show a considerably lowered hardness. This ismost likely due to the graphitization of the carbon matrix. In addition,it is seen from Fig. 6 that the hardness decreased with increasing

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sputtering current or increasing sp3-C fraction. It seems to be in con-trast with the generally accepted proposition that the sp3 matrix con-trols the mechanical properties of film [43]. Actually, it is not the casebecause the mechanical properties are really controlled by the con-nectivity of the sp3 matrix [36] rather than the total sp3-C fraction.Among the sp3-C components (sp3C-C, sp3C-Si and sp3C-H), onlythose promoting the matrix cross-linking, such as sp3C-C and sp3C-Si, contribute significantly to the hardness [49,58]. On the contrary,the C-H bonds not only make no contribution to the hardness (orelastic modulus) [58], but largely lower it by destroying the matrixconnectivity (since H only forms single bond) and reducing the cohe-sion energy which determines the hardness of the hard homogeneousbulk materials [49]. The XPS, IR and Raman analysis above commonlyreveal the increased C-H and Si-H content with increasing sputteringcurrent. Therefore, it is reasonable for hardness to drop graduallywith sputtering current.

The critical load (or the practical adhesion strength) displays asimilar variation trend to the hardness. Such a consistence was fre-quently observed in both a-C:H:Si [7,35] and other doped a-C:Hfilms [57]. It may result from the relatively low stress of doped a-C:H films. Because the practical adhesion is up to the balance betweenthe intrinsic adhesion (or cohesion) and the total stress items (in-cluding internal stress and interfacial stress) [49], and in caseswhere the stress is low (e.g. the well relaxed graphitic carbon films)the practical adhesion would be mainly dominated by the cohesionof the film.

4. Conclusions

To investigate the synergistic lubricating effects of graphitic car-bon and silicon oxide, we prepared a series of graphitic a-C:H:Sifilms with different Si content by altering the Si-sputtering currentin a hybrid RF-PECVD and unbalanced magnetron sputtering system.Microstructure and mechanical properties of the films were studied.Both the Si content and the sp3-C fraction increase with increasingsputtering current but the mechanical property varies adversely.The stand-out of D peaks, the non-dispersion and up-shiftto~1600 cm−1 of G peaks in Raman spectra, together with the1595 cm-1 peaks corresponding to aromatic C=C vibration in FT-IRspectra, strongly indicate the graphitic nature of prepared a-C:H:Sifilms. The coupling effect of heating and strong ion bombarding in-duced by sputtering and negatively biasing is considered to be re-sponsible for the film graphitization. Further investigation on thetribological properties of a-C:H:Si films will be reported in anotherpaper.

Acknowledgements

The authors are grateful to the National Natural Science Foundation ofChina (Grant No.50705093 and No.50575217), the Innovative GroupFoundation fromNSFC (GrantNo.50421502) and theNational 973 Project(No.2007 CB607601) for financial support.

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