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Materials Chemistry and Physics 130 (2011) 775– 785

Contents lists available at ScienceDirect

Materials Chemistry and Physics

j ourna l ho me pag e: www.elsev ier .com/ locate /matchemphys

nfluence of bonding environment on nano-mechanical properties of nitrogenontaining hydrogenated amorphous carbon thin films

eeraj Dwivedia,b, Sushil Kumara,∗, H.K. Malikb, C.M.S. Rauthana, O.S. Panwara

Physics of Energy Harvesting Division, National Physical Laboratory (CSIR), K.S. Krishnan Road, New Delhi 110 012, IndiaDepartment of Physics, Indian Institute of Technology Delhi, New Delhi 110 016, India

r t i c l e i n f o

rticle history:eceived 28 March 2011eceived in revised form 28 June 2011ccepted 24 July 2011

eywords:hin films

a b s t r a c t

High quality nitrogen incorporated hydrogenated amorphous carbon (a-C:N:H) thin films have beengrown, using radio frequency-plasma enhanced chemical vapor deposition technique, under varied selfbiases from −25 to −400 V. Presence of nitrogen in carbon–nitrogen network and their bondings withcarbon and hydrogen is confirmed by time of flight-secondary ion mass spectroscopy, energy dispersiveX-ray analysis, and Fourier transform infrared spectroscopy measurements, respectively. By deconvolut-ing FTIR spectra, the various sp3CHn and sp2CHn bonding as well as % sp3 and sp2 fractions are estimated,

VDTIRechanical properties

which further correlated with nano-mechanical properties of a-C:N:H films. Observed maximum hard-ness and elastic modulus as 42 and 430 GPa, respectively in a-C:N:H film deposited at −100 V revealsuperhard behaviour of coating. Residual stress in a-C:H:N films is found to be lower due to nitrogenincorporation that further decreases with increasing self bias. In addition, load versus displacementcurves have also been employed to calculate various other nano-mechanical parameters, which further

s we 3 2

correlated with self bias a

. Introduction

After theoretical prediction of existence of super hard �-C3N4ompound by Liu and Cohen [1], the researchers from world-ide have devoted their keen interest towards carbon nitride. It

s reported in literature that bulk modulus and hardness of �-C3N4ompound can be comparable to that of diamond [2]. However,he creation of predicted crystalline phase in carbon nitride in theaboratory is still a complex task. Nitrogen incorporated hydro-enated amorphous carbon (a-C:N:H) is an another important classf diamond-like carbon (DLC), which have potential to exhibitremendous mechanical, electrical and optical properties [3–7].onetheless, DLC films have main drawback of high residual stress

hat restrict its potential technological and industrial applications.ddition of nitrogen in a-C:H has been found an appropriate way

o minimize its residual stress [3,5,7]. Self bias (which is a linearunction of ion energy) is a very important growth parameter thatnfluence the bonding environment and various properties of a-C:Hnd a-C:N:H films [8].

Further, in a-C:H films the two sources of diamond-like sp3 and

raphite-like sp2 bonding are due to ion and due to C–H bondings.he ion induced sp3 and sp2 bonding have been studied extensivelynd correlated with its electrical, optical and mechanical proper-

∗ Corresponding author. Tel.: +91 114 560 8650; fax: +91 114 560 9310.E-mail address: skumar@nplindia.org (S. Kumar).

254-0584/$ – see front matter © 2011 Elsevier B.V. All rights reserved.oi:10.1016/j.matchemphys.2011.07.060

ll as estimated % sp and sp fractions.© 2011 Elsevier B.V. All rights reserved.

ties. However, influence of C–H bonding (sp3 C–Hn and sp2 C–Hn

bonding) on various properties of a-C:H films is not explored yet,despite definite amount of hydrogen in a-C:H films is quite essen-tial to obtain diamond-like character and stabilizing the diamondstructure by passivating dangling bonds through maintaining thesp3 hybridization configuration. Titantah et al. [9] suggested thathydrogen stabilizes the diamond-like structure by maintaining sp3

carbon content, thereby enabling its usage for the coating on cut-ting tools, high shear stress devices, solar cells and hard disk. This iswhat said for a-C:H is also true for a-C:N:H films. Thus, sp3 and sp2

hybridization of C–H bonding may also influence electrical, opticaland mechanical properties of a-C:N:H films significantly.

Presently, development of hard, superhard and ultrahard coat-ing is found to be an interesting research topic due to their directindustrial applications such as cutting tools, automobile parts, mag-netic storage media and glasses. The materials can be classified ashard, superhard and ultrahard if they posses hardness better than20, 40 and 80 GPa, respectively [10]. Diamond with hardness andelastic modulus ∼100 and ∼1000 GPa, respectively is well knownultrahard material. �-C3N4 has also great potential to exhibit prop-erties similar to that of diamond. However, the formation of thesetwo materials are found to be quite complex. In addition, variousother nitrides and carbides such as Si3N4, TiN, TiC, and SiC have also

potential to exhibit hard and superhard coating [11,12]. But thereis great interest towards a-C:N:H as a hard and superhard coat-ing because a-C:N:H films also exhibit excellent biocompatibility,chemical inertness, high wear resistance capability [13,14].

776 N. Dwivedi et al. / Materials Chemistry and Physics 130 (2011) 775– 785

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te (n

m/m

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Typical ToF-SIMS depth profiles for C, H, N, and Si from a-C:N:H

ig. 1. Schematic representation of deposition unit used for the growth of a-C:N:Hlms.

Further, in order to obtain accurate nano-mechanical proper-ies from thin films, the selection of appropriate characterizationool is very important. If indenter penetrates one third of total filmhickness then substrate also influences the mechanical proper-ies and result shows composite substrate/film effect. Generallyimple indentation and microindentation, which are not depthensitive, are used to measure mechanical properties of a-C:N:Hlms. However, simple indentation is useful to characterize bulkamples whereas microindentation is useful to characterize veryhick films (micron order). In order to estimate accurate nano-

echanical properties of thin films then depth sensitive and highesolution indentation is required. Nanoindentation has been foundn important tool to characterize nano-mechanical properties ofhin a-C:N:H films due to its depth sensitive ability as well asigh resolution characteristic. Recently, we have also estimatedhe nano-mechanical properties of various types of DLC films byanoindentation [3,6,15–17].

Therefore in the present paper, the influence of bonding envi-onment on nano-mechanical properties of a-C:N:H films has beenxplored. The effort has been made to deposit hard and superhard-C:N:H coating to enhance its possible applications. In addition,rocess optimization has also been done to minimize excessiveesidual stress through nitrogen incorporation.

. Experimental details

The a-C:N:H films have been grown using asymmetric capacitive coupled radiorequency (13.56 MHz)-plasma enhance chemical vapor deposition RF-PECVD tech-ique, on well cleaned Si and corning 7059 glass substrates, at a base pressure

× 10−5 Torr. The schematic representation of deposition unit used for growth of a-:N:H films is depicted in Fig. 1. The deposition unit consist of four segments namely1) vacuum pumps & components, (2) gas manifold, (3) electrical power supply and4) process chamber and electrode geometry. The rotary and turbo molecular pumpsombination have been used to achieve the base pressure better than 1 × 10−5 Torr.he gas section contains acetylene (C2H2), nitrogen (N2) and argon (Ar) gases lines.he power section contains RF power generator (Comdel CPS-2000), matching net-ork and the self bias indicator. The process chamber is a very important part of the

ystem which also acts as ground electrode where power electrode is fixed inside.he power electrode (cathode, diameter 8 in.) on which the substrates are placed isept much smaller than the ground electrode (anode), which included whole bodyf the deposition chamber. The plasma is sustained between these two electrodes.rior to deposition, the substrates are also cleaned in Ar plasma for 10 min to avoidoisture and contamination from surface. The growth of a-C:N:H films have been

arried out under varied negative self biases from −25 to −400 V with maintainingonstant acetylene (C2H2) and nitrogen (N2) gas pressures and therefore constantorking pressure in all the depositions. The constant working pressure is kept at

0 × 10−3 Torr that is achieved by feeding first C2H2 gas, which changed the pressure

Self Bias (V)

Fig. 2. Variation of deposition rate versus self bias for different a-C:N:H films.

from 1 × 10−5 Torr to 20 × 10−3 Torr then by N2 gas that changed the pressure from20 × 10−3 Torr to 40 × 10−3 Torr.

Thicknesses of these a-C:N:H films have been measured by Taylor–HobsonTalystep instrument, which are found in the range from 250 to 412 nm. The depthprofile and mass spectra of a-C:N:H film is recorded by time of flight-secondaryion mass spectroscopy (ToF-SIMS), which is equipped with sputter oxygen gun.Plasma emission of nitrogen containing hydrocarbon plasma, in the range from 300to 700 nm, is recorded by M/s Avantes multichannel spectrometer connected withthe optical fiber cable. Energy dispersive X-ray analysis (EDAX) measurement isperformed to confirm the presence of carbon and nitrogen in a-C:N:H film. Themicrostructure of a-C:N:H films has been studied by scanning electron microscope(SEM, model: JEOL, JSM-35 instrument). FTIR measurements are carried out, toinvestigate the bonding environment of different a-C:N:H films, using M/s Varian660 IR spectrophotometer. The residual stress (S) in these films is measured using500TC temperature controlled film stress measurement system (M/s FSM FrontierSemiconductor, USA). The Stoney formula, which is based on curvature method, wasused to measure S, which is given in Eq. (1)

S = Esd2s

6(1 − �s)df

(1Rf

− 1R0

)(1)

where Es , �s , df and ds are Young’s modulus, Poisson ratio and thickness of the filmand substrate, respectively and R0 and Rf are the radii of substrate curvature beforeand after film deposition. Fully automatic software controlled IBIS nano indenta-tion (M/s Fisher-Cripps Laboratories Pvt. Ltd., Australia) having diamond Berkovichindenter is employed to measure the nano-mechanical properties of a-C:N:H films.The maximum indentation load of 10 mN is applied in nanoindentation testing.

3. Results and discussion

3.1. Deposition rate

The variation of deposition rate versus self bias for a-C:N:Hfilms is shown in Fig. 2. The deposition rate increases with increas-ing the self bias. Initially, deposition rate increases considerablywith increasing self bias from −25 to −100 but beyond −100 V itincreases steadily. The value of deposition rate at self biases −25,−100, −200, −300 and −400 V are found to be 5.6, 51.7, 56.7, 60.8and 68.7 nm min−1, respectively. Observed increase in depositionrate with increase in self bias is due to increase in dissociation rateof hydrocarbon gas precursor.

3.2. ToF-SIMS analysis

film deposited at self bias −100 V is shown in Fig. 3(a). From ToF-SIMS depth profile, it is clearly confirmed that the species C, H,N constitute the a-C:N:H films. As expected, among C, H, and N, C

N. Dwivedi et al. / Materials Chemistry and Physics 130 (2011) 775– 785 777

10 3

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(a.u

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Inte

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Time (S)

504030201000

900

1800

2700

3600 -100 V C3H3+C2H3

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C+

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Inte

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Mass

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Inte

nsity

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Wavelength (nm)

from bulk. EDAX uses high energy X-ray (typically 10 KeV) and it ispossible to figure out information even from bulk of the material.However, XPS (X-ray photoelectron spectroscopy) measurementsmay provide better information about surface contamination. Ear-

ig. 3. (a) Typical SIMS depth profile of a-C:H:N film and (b) mass spectra of a-C:N:Hlm grown at self bias −100 V.

ith its high intensity is found to be a dominant element. However,s compared to C and H, the intensity of N in the depth profile isound to be lower but sufficient to influence the properties of a-:N:H films. FTIR results as discussed later also reveal sufficient N

n a-C:N:H film. In addition, depth profile spectrum contains a veryigh intensity peak, which corresponds to the substrate Si. TypicaloF-SIMS mass spectra of a-C:N:H film deposited at self bias −100 Vs shown in Fig. 3(b). The presence of key ions such as H+, C+ and

+ can be clearly seen in the mass spectra of film. Besides H+, C+

nd N+, other abundance positive ions are found to be C2H3+ and

3H3+.

.3. Optical emission spectroscopy (OES) analysis

Typical OES spectra of nitrogen containing hydrocarbon plasmased for the deposition of a-C:N:H film at self bias −100 V is shown

n Fig. 4. Figure clearly reveals the existence of various emissioneaks, which corresponds to various CH and CN transitions. Theeak found near 337 nm exhibits two distinct transitions. Intensemission found at 336.7 nm (as shown by No. 2 in the spectra) revealhe transition from N(1+) band whereas weaker emission obtainedt 333.1 nm (as shown by No. 1) correspond to NH transition. Aery intense peak centered at 381.6 nm shows emission throughontribution of both CH and CN bond. It may be possible that themission peak centered at 594.8 nm (as given by No. 4) represents

he A2A1 → X2B1 rovibronic transition of NH2. The peak obtained at70.4 nm (represented by No. 5) may correspond to 4d4P → 3p4P0

tomic N transition. OES spectra also reveal the existence of N2 sec-nd band (N2

(2+)) beyond 575 nm. H� transition also appeared in

Fig. 4. Typical OES spectra of a-C:N:H film deposited at self bias of −100 V.

the N2(2+) band as shown in the spectra. In comparison to reported

literature [18], observed little peak shifting may be due to the dif-ference in the processing parameters used in the present study. Thisis to be noted that these plasma emission peaks are very useful forreproducibility of the results.

3.4. Energy dispersive X-ray analysis (EDAX)

EDAX measurement has been performed to confirm the pres-ence of carbon and nitrogen in a-C:N:H films. Fig. 5 shows theTypical EDAX spectrum of a-C:N:H film deposited at self bias of−100 V. EDAX spectrum clearly reveal the presence of carbon aswell as nitrogen in a-C:N:H structure. In addition to carbon andnitrogen, some oxygen is also obtained in spectra which may bedue to surface contamination and could also have contribution

Fig. 5. Typical EDAX spectrum of a-C:N:H film deposited at self bias of −100 V.

778 N. Dwivedi et al. / Materials Chemistry and Physics 130 (2011) 775– 785

Fig. 6. (a)–(f) SEM micrographs, 3D view of SEM micrographs and roughness of a-C:H:N films grown at self bias −100 and −400 V.

istry and Physics 130 (2011) 775– 785 779

lNmoCt

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Inte

nsity

(a.u

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Wavenumber (cm-1 )

1 -100 V2 -200 V3 -300 V4 -400 V

N. Dwivedi et al. / Materials Chem

ier, we have also performed the XPS measurements on DLC and-DLC films and discussed the role of contamination on nano-echanical properties of these films [3,17]. Thus, presence of

xygen may improve the nano-mechanical properties of these a-:N:H films as oxygen enhance the diamond-like sp3 bonding dueo etching of graphite-like sp2 bonding.

.5. Surface morphology studies

Scanning electron microscope (SEM) has been used to investi-ate the surface morphology of a-C:N:H films. The SEM pictures of-C:N:H films deposited at self bias −100 and −400 V are shownn Fig. 6(a) and (b), respectively. From figure, it is confirmed thathese a-C:N:H films exhibit nano-structured morphology. The a-:N:H film deposited at −100 V shows highly dense network alongith uniform distribution of spherical shaped nano-particles over

ntire surface. The size of these spherical nano-particles is foundn between ∼120 and 130 nm. This can be seen from figure thatt few places the clustering of these spherical nano-particles (asarked by circles) have also been observed but this clustering do

ot influence the mechanical properties of films. Somani et al. [19]lso observed spherical shaped nano-structured morphology in a-:N:H films. Further, it is interesting to note that with increase inelf bias for the growth of a-C:N:H film to −400 V, drastic change ints surface morphology has been obtained. The film exhibits columnike nano-structures on entire surface. However, the size of theseano-structured columns is found to be uneven throughout the sur-

ace. The uneven sized column-like nano-structured morphologyn carbon nitride and related materials have also been reported by

ang [20]. Besides isolated nano-structured columns, the bundlingf these nano-structured columns have also been observed, whichs marked by white dot circles. Fig. 6(c) and (d), show the 3D versionf SEM micrographs of a-C:N:H films grown at −100 and −400 V,espectively. The generated 3D micrograph of a-C:N:H films revealhe uniform surface topography. These 3D micrographs are good ingreement with 2D micrographs. The surface line profile has alsoeen performed to evaluate the surface roughness of a-C:N:H films,hich as shown in Fig. 6(e) and (f), respectively. The roughness data

s taken at three different places on the surface of films. The aver-ge roughness (Ra) in these films is found to be significantly lower.owever, increase in self bias from −100 V to −400 V increases thea from 0.46 nm to 1.2 nm. Observed increase in Ra with the increase

n self bias is due to increase in energy of bombarding ions on grow-ng film surfaces. Recently Ishpal et al. [21] have also proposed this

ay to produce 3D SEM images and the roughnesses data. Thesenteresting surface features are correlated with nano-mechanicalroperties of a-C:N:H films.

.6. FTIR analysis

FTIR spectra of a-C:N:H films deposited at different self biasesf −100, −200, −300 and −400 V is shown in Fig. 7. Spectra clearlyeveal the presence of nitrogen in the carbon–nitrogen network.ands appear in the range between 1200 and 1800 cm−1 and250–2400 cm−1 may ascribe to bonding states of C with N alongith water vibrations and CO2 signal due to moisture (impurity)

n surface of films. The peak near1590 cm−1 shows the existencef imino C N −group as well as sp2 C C vibrations. The peakbtained at ∼1110 cm−1 may assign to Si–O bonding, which fur-her reduces with increasing the self bias. Various peaks observedn the range between 500 and 1000 cm−1 correspond to bond-ng states of Si with H. Peaks obtained in the range from 2800

o 3090 cm−1 assigned to C–H stretching mode. These results areound to be good in agreements with reported literature [22–25].ince C–H stretching vibrations are very dominant bondings in a-:H and a-C:N:H films, therefore deconvolution of these vibrating

Fig. 7. FTIR spectra of a-C:N:H film deposited at self biases of −100, −200, −300 and−400 V.

modes has also been performed to figure out more informationabout the structures. The deconvoluted FTIR spectra in the range2700–3200 cm−1 of a-C:N:H films deposited at different self biasesof −100, −200, −300 and −400 V are shown in Fig. 8(a)–(d), respec-tively. The 8 peaks have been fitted with considering Gaussiancomponent in each a-C:N:H films. The film deposited at −100 Vexhibit peaks at 2860.2, 2869.1, 2917.9, 2954.7, 2972.8, 3006.1,3036.0 and 3059.1 cm−1 (represented by digits 1, 2, 3, 4, 5, 6, 7 and 8,respectively), which correspond to sp3CH2, sp3CH3, sp3CH, sp3CH3,sp2CH2, sp2CH, sp2CH and sp2CH2 bonding, respectively. Similarly,in other self biases (−200, −300 and −400 V) the numeric digits1, 2, 3, 4, 5, 6, 7 and 8 represents sp3CH2, sp3CH3, sp3CH, sp3CH3,sp2CH2, sp2CH, sp2CH and sp2CH2 bondings, respectively but withsome changes in their peak positions due to change in depositionparameters. In addition, the estimation of % sp3 and sp2 fraction ofC–H bonding has also been performed by considering peaks arearatio method. In all films, the peaks represented by 1, 2, 3 and 4showed sp3 CHn bonding whereas peaks represented by 5, 6, 7 and8 exhibited sp2CHn bonding. Hence % sp3 and % sp2 bonding havebeen calculated using Eqs. (2) and (3)

% sp3 =∑4

i=1Ai∑8i=1Ai

× 100 (2)

% sp2 =∑8

i=5Ai∑8i=1Ai

× 100 (3)

where i = 1, 2, 3, 4, 5, 6, 7 and 8 and A1, A2, A3, A4 are area under thepeaks for sp3CHn bonding (peaks 1–4) whereas A5, A6, A7 and A8 arearea under the peaks for sp2CHn bonding (peaks 5–8), respectively.The variation of % sp3 and % sp2 fractions versus self bias is shownin Fig. 9. % sp3 fraction has been found to continuously decreasewith increasing self bias from −100 to −400 V. This can be seenthat up to −300 V the value of % sp3 fraction decreases graduallybut beyond −300 it decreases drastically. The values of sp3 fractionat self biases −100, −200, −300 and −400 V are found to be 82.1,78.1, 76.7 and 60.9%, respectively. On the other hand % sp2 fractionhas been found to continuously increase with increasing self bias

from −100 to −400 V. The values of sp2 fraction at self biases −100,−200, −300 and −400 V are found to be 17.9, 21.9, 23.3 and 39.1%,respectively. Observed decrease in sp3 and increase in sp2 fractionswith increase in self bias from −100 to −400 V is due to initiation

780 N. Dwivedi et al. / Materials Chemistry and Physics 130 (2011) 775– 785

320031003000290028002700

a

(8)

(7)

(6)

(5)

(4)

(2)

(3)

(1)

-100 VIn

tens

ity (a

.u.)

Wave Number (cm-1)

320031003000290028002700

b

(8)

(7)

(6)(5)(4)

(3)

(2)

(1)

-200 V

Inte

nsity

(a.u

.)

Wave Number (cm-1)

320031003000290028002700

c

(8)

(7)

(6)(5)(4)

(3)

(2)

(1)

-300 V

Inte

nsity

(a.u

.)

-1

320031003000290028002700

d

87

6

5

4

3

2

1

-400 VIn

tens

ity (a

.u.)

-1

eposi

oPFtbHtacethsevbF(

B

Wave Number (cm )

Fig. 8. Deconvoluted FTIR spectra of a-C:N:H film d

f graphite-like nature of the a-C:N:H films. Sharon et al. [26] andaul et al. [27] have also calculated the sp3 and sp2 fractions byTIR measurements. Further, the estimation of hydrogen concen-ration (HC) in a-C:H and a-C:N:H films have also been performedy FTIR analysis. It is important to mention that the estimation ofC by FTIR is also a subject of debate, because stretching mode in

he range 2700–3150 cm−1 where C atoms properly bond with Htoms is used for the estimation. Thus, estimated HC by FTIR mayall as bound hydrogen concentration (BHC). On the other hand,lectron recoil detection analysis (ERDA) and Rutherford back scat-ering (RBS) techniques estimate the total HC (THC) that is boundydrogen as well as trapped hydrogen [28]. The bonded hydrogentabilize the diamond-like bonding, and therefore directly influ-nce the nano-mechanical properties. Thus, estimation of BHC isery important in present study. The variation of BHC versus selfias for different a-C:N:H films, estimated by FTIR is depicted inig. 10. The BHC in these films was evaluated by employing an Eq.

4) [29]

HC = As(CHx)

∫˛(k)

kdk (4)

Wave Number (cm )

ted at (a) −100, (b) −200, (c) −300 and (d) −400 V.

where ˛(k) is the absorption coefficient at wavenumber k, As(CHx)is the proportionality constant. We used As(CHx) = 1 × 1021 cm−2 forestimation of hydrogen concentration [29]. The BHC in these filmsis found to be lower but sufficient to stabilize the diamond-likesp3 bonding. Moreover, the variation of BHC with self bias is alsofound to be very small, this is the reason why a-C:N:H films grownat and above −100 V showed better nano-mechanical properties.The calculated % sp3 and sp2 fractions have been further correlatedwith nano-mechanical properties in the next section.

3.7. Nano-mechanical properties

Due to its high resolution characteristics nanoindentationis found to be an important technique to measure the nano-mechanical properties of thin films. The load versus displacementcurves of fused silica and a-C:N:H films deposited at negative self

biases of −25 −100, −200, −300 and −400 V are depicted in Fig. 11.The mechanical property of fused silica is also examined and con-sidered as a reference. The values of hardness and elastic modulus infused silica are found to be ∼9 GPa and ∼78 GPa, respectively. Fur-

N. Dwivedi et al. / Materials Chemistry and Physics 130 (2011) 775– 785 781

15

20

25

30

35

40

-400-300-200-1000

60

65

70

75

80

85

sp2 (%

)Self Bias (V)

sp3 (%

)

tfiibtpaItsetatiit−T

0.30.20.10.0

0

2

4

6

8

10

Load

(mN

)

Displacement (μm)

25 V 100 V 200 V 300 V 400 V Fused Silica

namely polymer-like, diamond-like and graphite-like carbon films.Erdemir and Donnet [31] reported that at low ion energy, thehydrocarbon precursor is not sufficiently decomposed and contain

Fig. 9. Variation of sp3 and sp2 fractions of a-C:N:H films with self bias.

her, the penetration depth (h) is found to be minimum in a-C:N:Hlm deposited at −100 V and beyond this self bias (up to −400 V)

t is increased due to initiation of graphite-like sp2 bonding andelow −100 V (at −25 V) observed higher penetration depth is dueo presence of high amount of unbound hydrogen that promotes theolymer-like structure. The penetration depth at 10 mN in all the-C:N:H films is found to be in the range between 100 and 190 nm.t is important to mention that substrate significantly influencehe hardness (H) of thin films, therefore H of Si wafer is also mea-ured before the depositions and then H of substrate/film has beenstimated by composite hardness model [30]. Observed least hys-eresis between loading–unloading curve in a-C:N:H film depositedt −100 V revealed their superhard and highly elastic behaviour andherefore better quality of film. The variation of H versus self biass plotted in Fig. 12(a). It is evident from figure that the value of Hncreases drastically from 8 to 42 with increase in self bias from −25

o −100 V GPa. However, further increase in self bias from −100 to400 V results in continuously decreases of H from 42 to 23 GPa.he variation of elastic modulus (E) versus self bias is shown in

-400-300-200-1000

9.6x1 021

1.0x1 022

1.1x1 022

1.2x1 022

1.3x1 022

1.4x1 022

BH

C (a

tm. /

cm

2 )

Self Bias (V)

Fig. 10. Variation of BHC versus self bias for different a-C:N:H films.

Fig. 11. Load versus displacement curves of fused silica and a-C:N:H films depositedat self biases −25, 100, −200, −300 and −400 V.

Fig. 12(b). Experimentally observed E values also followed similartrend and first abrupt increases from 120 to 430 GPa with increasein self bias from −25 to −100 V, beyond this voltage it is continu-ously decrease to 302 GPa with increasing self bias to −400 V. Themechanical properties (H and E) of a-C:N:H films is controlled bycontrolling their bonding network. It is to be noted that nitrogenpartial pressure and self bias both influence the bonding network ofa-C:N:H film. But in the present study only self bias is varied (nitro-gen and acetylene gas pressures have been kept constant) to seethe changes in bonding environment of a-C:H:N films. Ion energy,which is a linear function of self bias and RF power is used to explorebonding structure of a-C:N:H film. On the basis of ion energy, thea-C:H and a-C:N:H thin films can be divided into three categories,

-400-300-200-1000

100

200

300

400

10

20

30

40

(b)

E (G

Pa)

Self Bias (V)

(a)

H (G

Pa)

Fig. 12. Variation of self bias versus (a) H and (b) E for a-C:N:H films.

7 istry and Physics 130 (2011) 775– 785

htafild[hvsntabomttgCaoowihohop−mt2ieH[pStHpcpiiiiCtdssehiHsIchadea

0.80.70.60.50.40.3

10

20

30

40

50a

2001801601401201005

10

15

20

25

30

35

40

45

45

2

3

1

H (G

Pa)

Penetration Depth (nm)

54

3

2

1

H (G

Pa)

Penetration depth / film thickness

1: -25 V2: -100 V3: -200 V4: -300 V5: -400 V

40302010

100

200

300

400

b a-C:N:H films

E (G

Pa)

H (GPa)

Experimental value E/H = 10 Thumb rule

82 N. Dwivedi et al. / Materials Chem

igh amount of hydrogen, which resulted into soft polymer-likehin films. At intermediate ion energy, hydrogen content reducesnd hydrocarbon precursor is decomposed sufficiently, depositedlms possess high density and therefore, revealing their diamond-

ike character. At very high ion energy, graphite-like sp2 bondingominant due to increase in sp2 induced disorder. Fallon et al.32] obtained maximum sp3 bonding at about 100 eV in tetra-edral amorphous carbon films deposited using filtered cathodicacuum arc technique, beyond which initiation of graphite-likep2 bondings has been observed. Ito et al. [33] reported carbonitride film deposited at 40 V exhibit H ∼ 3.5 GPa, which increaseso ∼36 GPa at 100 V. Singh et al. [34] observed diamond-like filmt −150 V. We have also observed diamond-like properties at selfias in between −100 and −130 V [3,6]. Thus, in the present studybserved lower value of H at −25 V and higher value of H at −100 Vay correspond to their polymer-like and diamond-like charac-

ers, respectively. However, beyond −100 V observed decreasingrend of H with increase in self bias accompanied by initiation ofraphite-like sp2 bondings. The observed excellent H and E in a-:N:H films deposited at −100 V and above self biases have alsopproved by its highly dense nano-structured surface morphol-gy. According to Hall–Petch theory, H varies inversely with sizef the particles [35]. Therefore, observed trend of decrease in Hith increase in self bias from −100 to −400 V may due to increase

n sizes of the particles and the clusters. As sufficient amount ofydrogen is required to stabilize sp3 network and then used tobtained diamond-like properties. Therefore, H of a-C:N:H filmsas also been expressed in term of sp3 and sp2 CHn bonding. The Hf a-C:H:N film varied proportional with sp3 fractions and inverselyroportional with sp2 fractions. The H of a-C:N:H film deposited at100 V posses maximum sp3 fraction 82.1% and therefore maxi-um H 42 GPa. However, with decrease in sp3 fraction from 82.1

o 60.9% the H has been found to continuously decrease from 42 to3 GPa. The same trend with sp3 fractions has also been observed

n E. It is worth noting that the value of H is considerably highven at −400 V. Recently, we have demonstrated the variation of

of a-C:H films with sp3 fractions as well as with sp3/sp2 ratios17]. Xu et al. [36] improved the mechanical properties of amor-hous carbon nitride films by additional titanium doping. Similarlyhi [37] studied systematically the mechanical properties of sput-ered amorphous carbon nitride films. However, obtained higher

values in a-C:H:N films may also be due to cross linking of basallanes (graphite induced hardening). Generally the basal planes areurved, resulting in buckling in flat graphite-like sp2 hybridizedlanes. Existence of this buckling in basal planes leads to intersect-

ng planes. Thus, it is possible that basal planes in this network maynterlock with covalent bonds of shorter bond length and resultingn high H [38]. In contrast doping of foreign elements as suggestedn literature, we obtained better nano-mechanical properties in a-:N:H films just by optimizing the process parameters. The H ofhese a-C:N:H films is also examined as a function of penetrationepth/film thickness (h/t) and penetration depth only, which ashown in Fig. 13(a). Saha and Nix [39] and Chen and Vlassak [40]ystematically studied the effect of substrate on mechanical prop-rties of thin films. They considered two cases first, soft film onard substrate and second, hard film on soft substrate. They real-

zed that soft film deposited on hard substrate showed increase in with increase in penetration depth whereas hard film on soft sub-trate exhibited decrease in H with increase in penetration depth.n the present study the case hard film on soft substrate may appli-able for a-C:N:H films deposited at and above −100 V. On the otherand the case soft film on hard substrate in a-C:N:H film deposited

t −25 V may considered. It is evident from figure that a-C:N:H filmeposited at −25 V showed highest h/t ∼ 0.77 and therefore, low-st H (very close to substrate Si) due to not only bonding but maylso be due to dominance of substrate. On the other hand a-C:N:H

Fig. 13. (a) Plot of H versus penetration depth/film thickness and (b) variation of Hversus E for different a-C:N:H films.

film deposited at −100 V exhibit lowest h/t ∼ 0.33 (one third of totalfilm thickness) and therefore, highest H with negligible substrateeffect. The H values in a-C:N:H films deposited at −200 V, −300 Vand −400 V are found to be quite close to each other due to veryslight variation in h/t as represented by dashed circle in Fig. 13(a).This can be seen that film deposited at −200 V exhibited compara-tively higher H despite of more h/t than films deposited at −300 Vand −400 V. This is due to the fact that in this film although h/t isslightly higher but graphite-like sp2 bonding is significantly lowerthan films deposited at −300 V and −400 V. Similarly due to moresp3 bonding, a-C:N:H film deposited at −300 V exhibited compar-atively higher H than film deposited at −400 V in spite of slightlymore h/t. The inset of Fig. 13(a) shows the variation of H versus pen-etration depth only. Obtained result of H as a function of penetrationdepth is also good in agreement with the result of H as a functionof h/t. Generally H varies linearly with E and follows the relationH ∼ (1/10)E, which is well proven rule of thumb for a-C:H and a-C:N:H films. The variation of E versus H for experimental measureda-C:N:H films and their comparison with standard rule of thumb[H ∼ (1/10)E] is plotted in Fig. 13(b). It is evident from figure that

in the beginning experimental values are found to be higher thanthat of the values of rule of thumb (H ∼ (1/10)E). However, beyondH of 29 GPa the experimental values of E are bend and expected tomeet with standard values (H ∼ (1/10)E) at H of 42 GPa or above.

N. Dwivedi et al. / Materials Chemistry and Physics 130 (2011) 775– 785 783

-400-300-200-100050

60

70

80

90

0.07

0.08

0.09

0.10

(b)

ER (%

)

Self Bias (V)

(a)H

/E

TatH

iotcHdbibdmfiiau

%

wrtEabriitrittahd

-400-300-200-1000

6.0x10-10

9.0x10-10

1.2x10-9

1.5x10-9

0.1

0.2

0.3

0.4

0.5

(b)

U (J

)r

(a)

h res/h

max

Fig. 14. Variation of self bias versus (a) H/E and (b) ER for a-C:N:H films.

hus, it has been further confirmed that a-C:N:H film depositedt self bias of −100 V with H and E of 42 and 430 GPa, respec-ively exhibited highly diamond-like character (diamond possesses

∼ 100 GPa and E ∼ 1000–1100 GPa).Plastic resistance parameter (H/E), which included both H and E,

s found to be an important parameter to describe elastic propertiesf thin films. The high (low) values of H/E correspond to highly elas-ic (plastic) behaviour of thin films. Also for high wear resistanceoating the parameter H/E must be very high [3]. The variation of/E versus self bias is shown in Fig. 14(a). The value of H/E increasesrastically with increasing self bias from −25 to −100 V. However,eyond −100 V H/E decreases gradually. The value of H/E is varied

n the range between 0.067 and 0.098. Observed decrease in H/Eeyond −100 V shows more fraction of work consumed in plasticeformation and large plastic strain is expected when contactingaterial. However, observed considerable higher H/E in a-C:N:H

lms made it a material of great utility in wear resistance coat-ng on magnetic storage media. In addition, elastic properties of-C:N:H films is also estimated by calculating elastic recovery (ER)sing Eq. (5)

ER = (hmax − hres)hmax

× 100 (5)

here hmax and hres are the displacement at maximum load andesidual displacement after load removal, respectively. The varia-ion of ER versus self bias for a-C:N:H films is depicted in Fig. 14(b).R values also followed similar trend as observed in H, E and H/End initially rapidly increases from 54.7 to 93% with increasing selfias from −25 to −100 V. Further increase in self bias up to −400 V,esults in gradual decrease in ER. Observed decrease in ER withncrease of self bias beyond −100 V is due to increase in % sp2 bond-ngs as confirmed by FTIR. Whereas lower ER at −25 V may be dueo less dense and polymer-like character of the films. Observed EResults are good in agreement with H, E and H/E results. The moremportant fact regarding ER is that ER always strong if the penetra-ion depth is small. This can be seen from loading–unloading curves

hat minimum penetration depth is obtained for film depositedt −100 V, revealing less or negligible fracture of film and thus,igher ER whereas maximum penetration depth is observed for filmeposited at −25 V, which results in lower ER in later film. It is to be

Self Bias (V)

Fig. 15. Variation of self bias versus (a) hres/hmax and (b) Ur for a-C:N:H films.

noted that ratio of residual displacement after load removal (hres)with displacement at maximum load (hmax) i.e. (hres/hmax) providesinformation similar to that of ER with different domain of validity.Typically, the domain of validity of hres/hmax is found to in between0 and 1. The lower limit corresponds to fully elastic whereas upperlimit corresponds to the elastic–plastic behaviour, respectively. Thevariation of hres/hmax versus self bias for a-C:N:H films is shown inFig. 15(a). Values of hres/hmax are found to be in the range between0.07 and 0.45. The maximum value of hres/hmax is observed at −25 Vwhereas minimum at −100 V. Thus, from this parameter, it is againconfirmed that a-C:N:H film deposited at −100 V exhibit excellentelastic properties. The deformation behaviour in a-C:N:H films isconfirmed by calculating plastic deformation energy (Ur) [3,41].Generally Ur of the film is estimated by area surrounded by theloading–unloading curve in the load–displacement profile. Ur pos-sesses a relation with H, which is given by Eq. (5)

Ur =[

13

√1

ω0 tan2 �

]1√H

P3/2 (6)

where ω0 is the geometry constant and attains the value of 1.3for pyramid indenter, P is the load, and � is the half angle ofBerkovich indenter and have the value 65.3◦. The variation of Ur

against self bias for a-C:N:H films is shown in Fig. 15(b). MinimumUr of 6.6 × 10−10 J is observed in a-C:N:H film deposited at −100 V.The values of Ur are found to in the range between 6.6 × 10−10

and 1.5 × 10−9 J. The H and Ur strongly depends on ER. If ER in theunloading process is strong then H rises whereas Ur falls. Our Ur

results are found to be good in agreements with H and ER results.As high level of residual stress (S) is a big drawback of a-C:H films

that restrict its potential industrial applications. Several fundamen-tal and experimental efforts pertained to solution for said problemhave already been made and most commonly used experimentalway to reduce S in a-C:H is found to be the incorporation of nitro-gen in its structure. Nevertheless, nitrogen incorporation althoughminimize the S but it also reduce the H of a-C:H films. Thus, there isgreat interest to reduce S without affecting its H values. Variationof S versus self bias for a-C:N:H films is depicted in Fig. 16(a). S in a-

C:N:H films decreases with increasing self bias and found in rangebetween 1.1 and 2.1 GPa. It is evident from figure that initially withincrease in self bias from −25 to −100 V the value of S is increasedbut beyond this self bias S decreases continuously. Observed maxi-

784 N. Dwivedi et al. / Materials Chemistry and Physics 130 (2011) 775– 785

-400-300-200-10000.002

0.004

0.006

0.008

1.2

1.6

2.0

(b)

S / t

(GPa

/ nm

)

Self Bias (V)

(a)

S (G

Pa)

mbachplsfitSttsDcttntbftia

-400-300-200-1000

0.04

0.05

0.06

0.074

8

12

16

20

(b)

(H/E

)/S (G

Pa)-1

(a)

H/S

[(H/E)]/S] at higher self biases (−200, −300 and −400 V), made it

TH

Fig. 16. Variation of self bias versus (a) S and (b) S/t for a-C:N:H films.

um value of S at −100 V is due to dominance of diamond-like sp3

onding and beyond −100 V observed decrease in S is due to initi-tion of sp2 bonding, which as confirmed by FTIR results. However,omparatively higher S at −25 V may be due to presence of unboundydrogen in a-C:N:H structure because at −25 V hydrocarbon gasrecursor may not sufficiently decomposed. Observed lower S at

ow self bias may also be due to its polymer-like structure (H resultupport the formation of polymer-like character). The S of a-CN:Hlms strongly depends on the thickness of the films. Up to certainhickness, which may also called as critical thickness, the value of

in DLC and a-C:N:H films decreased. However, beyond criticalhickness the films get peel off. Although in the present study, thehicknesses of a-C:N:H films are lower than 0.5 �m, then these filmseems to below the critical thickness (that may be few microns inLC films). But the contribution of thickness in S of a-C:N:H filmsannot be ignored. Thus, in order to visualize the effect of bias andherefore, bonding on S of a-C:N:H films, we have also normalizedhe S results by thickness and variation of stress per unit thick-ess (S/t) versus self bias is depicted in Fig. 16(b). This can be seenhat the variation of S/t beyond −100 V was similar as observed in S,ut was significantly different below −100 V. Considering thicknessactors, the maximum S/t was observed at −25 V. This may be due

o fact that the fraction of unbound hydrogen in per unit thicknessn a-C:N:H film deposited at −25 V was more than film depositedt −100 V.

able 1, E, H/E, ER, hres/hmax, Ur , S, H/S and (H/E)/S values of various a-C:N:H films deposited at s

Parameters Self bias (V)

−25 −100

H(GPa) 8 42

E (GPa) 120 430

H/E 0.067 0.098

ER (%) 54.7 93

hres/hmax 0.45 0.07

Ur (J) 1.5 × 10−9 6.6 × 10−10

S (GPa) 1.8 2.1

H/S 4.4 20

(H/E)/S (GPa)−1 0.037 0.047

Self Bias (V)

Fig. 17. Variation of self bias versus (a) H/S and (b) (H/E)/S for various a-C:N:H films.

High S always accompanied by high H and both S and H variesproportional with diamond-like sp3 bonding. Observed results ofS and H follow the said statement. Thus due to combined effect,estimation of hardness per unit stress (H/S) is also found to be animportant parameter to figure out the actual nano-mechanical per-formance of a-C:N:H films. It is to be noted that higher the H/S,better the film quality is, thus may be treated as figure of merit.Variation of H/S against self bias for a-C:N:H films is plotted inFig. 17(a). It is evident from figure that initially with increase inself bias from −25 to −100 V, H/S drastically increases but beyond−100 V saturation in its values is observed. Observed higher H/Sat −100 or above −100 V, reveal better quality of films. Further,plastic resistance parameter (H/E) has already been explored asan important parameter for wear resistance coatings. Thus, esti-mation of plastic resistance parameter per unit stress [(H/E)/S]may also be very useful parameter to understand stress inducedwear resistance mechanism due to combined effect. Higher the(H/E) and lower the S makes higher the [(H/E)/S] is, thus bet-ter the wear resistance coating. The variation of [(H/E)/S] againstself bias is depicted in Fig. 17(b). From figure, it is evident that[(H/E)/S] varies linearly with self bias and found to be maximumat −400 V. Thus, despite of comparatively lower H, observed higher

an effective wear resistance coating material. On the other handfilm deposited at −25 V possesses lower H, lower H/E and lower[(H/E)/S] seems to be poor for its mechanical application. The esti-

elf biases of −25, −100, −200, −300 and −400 V.

−200 −300 −400

29 26 23350 330 3020.083 0.079 0.07678.4 74.3 72.30.22 0.26 0.287.9 × 10−10 8.3 × 10−10 8.9 × 10−10

1.6 1.4 1.118.1 18.6 20.90.052 0.056 0.069

istry

mT

4

i−stSfste4ivrmpnatvtc

A

otfmocC

R

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[[[38] H. Sjostrom, S. Stafstrom, M. Boman, J.E. Sundgren, Phys. Rev. Lett. 75 (1995)

1336.

N. Dwivedi et al. / Materials Chem

ated various different mechanical parameters are summarized inable 1.

. Conclusions

a-C:N:H thin films have been deposited by asymmetric capac-tive coupled RF-PECVD technique under varied self biases from25 to −400 V. Optical emission spectroscopy is used to diagno-

is various species in the plasma. The presence of N as well asheir bondings with H and C in these films is confirmed by ToF-IMS, EDAX and FTIR measurements, respectively. sp3 and sp2

ractions and BHC have been estimated by deconvoluting FTIRpectra and then correlated with their nano-mechanical proper-ies. Nano-mechanical properties are discussed in frame of bondingnvironment of a-C:N:H films. Superhard a-C:N:H coating with H2 GPa is obtained at −100 V. Thus, −100 V is optimized for obtain-

ng better quality film with highly diamond-like character. Loadersus displacements curves are also employed to estimate plasticesistance parameter, elastic recovery, ratio of residual displace-ent after load removal with displacement at maximum load,

lastic deformation energy, stress, stress per unit thickness, hard-ess per unit stress and wear resistance parameter per unit stressnd observed results are correlated each other. Thus, due to versa-ile mechanical properties these a-C:N:H films seems to be ideal forarious applications including hard and protective coating on cut-ing tools, automobile parts, solar cells and high wear resistanceoating on hard disk.

cknowledgments

The authors are grateful to the Director, National Physical Lab-ratory, New Delhi (India) for his kind support. The authors wisho thank Dr. Ranjana Mehrotra for providing FTIR measurementsacility. Authors also wish to thank Dr. M.K. Dalai for providing SIMS

easurement facility and Mr. Ishpal for helps in SEM analysis. Onef author N.D. acknowledges CSIR, Govt. of India for providing finan-ial support through SRF fellowship. This research is sponsored bySIR, Govt. of India, through the Network Project NWP-0027.

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