Met. Mater. Int., Vol. 18, No. 2 (2012), pp. 231~236doi: 10.1007/s12540-012-2005-7 Published 27 April 2012

Effect of Metallic Interfacial Layers on the Properties of Diamond-LikeCarbon Thin Films

Neeraj Dwivedi1,2, Sushil Kumar1,*, Sreekumar C.1, Saurabh Dayal1,C. M. S. Rauthan1, and O.S. Panwar1

1Physics of Energy Harvesting Division, National Physical Laboratory (CSIR), K.S. Krishnan Road,New Delhi - 110012, India

2Department of Physics, Indian Institute of Technology Delhi, New Delhi-110016, India

(received date: 13 April 2011 / accepted date: 26 May 2011)

Diamond like carbon (DLC) thin films with metallic interfacial layers of aluminum and nickel-chromium(Al and Ni-Cr) were grown using a low cost hybrid technique involving a resistive heating thermal evapo-rator and radio frequency plasma enhanced chemical vapor deposition techniques. Stress, hardness, elasticmodulus, bonding, phase, and electrical conductivity of these films were investigated. Introduction ofinterfacial Al and Ni-Cr layers in DLC led to drastic improvement of its conductivity along with a signifi-cant reduction in residual stress but with some reduction of hardness and the elastic modulus. The struc-tural and surface properties of thin films were studied using X-ray diffraction, X-ray photoelectronspectroscopy, and scanning electron microscopy techniques.

Key words: carbon and graphite, plasma deposition, electrical properties, mechanical properties, hardness test

1. INTRODUCTION

Recently, diamond like carbon (DLC) thin films haveattracted considerable interest in both scientific and indus-trial communities, not only due to their versatile mechanicaland tribological properties such as high hardness, high elas-tic modulus, and low coefficient of friction but also owing totheir outstanding electronic properties such as field emissionand low dielectric constant (k) electronics [1-8]. Currently,low k materials are strongly demanded in very large scaleand ultra large scale integrated circuit (IC) designs wheremillions of transistors are used in a small single chip. DLChas huge potential as a low k material for IC designs. Fur-ther, many studies have been carried out on chemical vapordeposition (CVD) grown diamond for electronic applicationsdue to its surface stability under high temperature (~500-600 °C in air and 1400-1700 °C in inert atmosphere), highelectron saturation velocity (~2.8×107 cm/s), high thermalconductivity (~20 W/cm), high break down voltage that is20-40 times greater than that of Si, high electron and holemobility (~2000 cm2V−1s−1). This leads to its application forhigh frequency and high power operated transistors. More-over, negative electron affinity, i.e. the energy of electrons inthe conduction band is above the minimum energy of elec-

trons in the vacuum level, thereby allowing its application incold cathode emission devices [9,10]. DLC films also showthese characteristics. However, one of the major drawbacksof DLC films is their poor adhesion with the substrates dueto high residual stress [11]. The introduction of a thin layerof metal prior to DLC deposition was found to be an appro-priate approach to resolve this problem. Moreover, a thininterfacial metal layer may also provide a platform to enhancethe conductivity of DLC films. Presently, metal/DLC bilayerstructures are grown using a hybrid system combining radiofrequency plasma enhanced chemical vapor deposition (RF-PECVD) and RF- sputtering units. However, this hybrid systemis found to be entail complex operation and is very expen-sive. In addition, during growth of the DLC layer carbonspecies adhere to and contaminate the target. This is referredto as “target poisoning” [12]. Thus, techniques combiningthermal evaporation and RF-PECVD were found to be anexcellent alternative for the deposition of metal/DLC struc-tures. Furthermore, among various metals, Ti, Cu, and W havealready been either incorporated in DLC films or used as anadhesive interfacial layer for successive DLC layer. The intro-duction of newer interfacial layers such as Al and compositeNi-Cr in DLC films is thus quite interesting.

In the present study, (Al, Ni-Cr)/DLC films have beengrown using combined thermal evaporation and RF-PECVDtechniques. Interfacial Al and Ni-Cr layers act as not only anadhesive layer for the successive DLC layer but also improve

*Corresponding author: skumar@nplindia.org©KIM and Springer

232 Neeraj Dwivedi et al.

its conductivity with some reduction of hardness. In addi-tion, bonding, composition, and phase analyses have beenperformed to support the obtained results.

2. EXPERIMENTAL PROCEDURES

Three sets of DLC films were deposited on a well-cleanedSi <100> wafer, Corning 7059 glass, and indium tin oxidecoated glass substrates. Among them, two sets of films weredeposited with a configuration of substrate/Al/DLC and sub-strate/Ni-Cr/DLC. The interfacial Al and Ni-Cr metal layerswere deposited using a thermal evaporation techniquewhereas the top DLC layers were deposited using the RF-PECVD technique. A third set of films, having a configu-ration of substrate/DLC, was grown using the RF-PECVDtechnique alone. A Al and Ni-Cr composite with purity of99.99 % and RF frequency of 13.56 MHz was used in thepresent study. Thin layers of Al and Ni-Cr of ~90 nm weregrown at base pressure exceeding 10−5 Torr with evaporationof Al and Ni-Cr while keeping the substrate to filament dis-tance of ~15 cm. DLC films were grown using an asym-metric capacitive coupled RF-PECVD system. The chamberof the system was initially evacuated using a rotary-rootsvacuum pump combination to a base pressure of ~10−3

Torr. The working gas pressure was kept constant at 30mTorr, which was achieved by feeding Ar and C2H2 gasesin the process chamber.

The mechanical properties of these films were estimatedat an indentation load of 10 mN by IBIS-Nanoindentation(M/S Fisher-Cripps Laboratories Pvt. Limited, Australia),equipped with a Berkovich indenter. The residual stress(S) in these films deposited on a silicon wafer was deter-mined from the change in curvature method, using a500TC temperature controlled film stress measurementsystem (M/s FSM Frontier Semiconductor, USA). Thesubstrate curvature method relies on the Stoney formula,which is given in equation (1).

S= (1)

where Es, νs, and ds are the Young’s modulus, Poisson ratio,and thickness of the substrate, respectively, and R0 and Rf

are the radii of the substrate curvature before and after filmdeposition. X-ray diffraction (XRD) patterns of Al/DLC,Ni-Cr/DLC and DLC films were obtained by a RigakuMiniflex II instrument. The X-ray photoelectron spectros-copy (XPS) spectrum of Ni-Cr/DLC was obtained fromPerkin Elmer 1257 instrument. Scanning electron micros-copy (SEM) micrographs were taken to study the morphol-ogy of the film. Conductivity measurements were performedusing a Keithley 610C Solid State electrometer from roomtemperature to 100 °C.

3. RESULTS AND DISCUSSION

Bonding, phase, and microstructural studies of these filmswere carried out using XPS, XRD, and SEM instruments.XPS measurement was performed at room temperature onNi-Cr/DLC films grown at a negative self bias of 100 V. TheXPS core level spectrum of C 1s is shown in Fig. 1(a). Fromthe figure it is evident that the peak corresponding to C wasobtained at 284.2 eV. However, no peak ascribed to the exist-ence of Cr or Ni was obtained in the Cr 2p and Ni 2p corelevel spectra, respectively. Since the Ni-Cr layer was grownbelow the 130 nm thicker DLC layer, it was difficult to visu-alize the presence of both Cr and Ni under this configuration.Thus, initial sputtering for 7 minutes was performed on theNi-Cr/DLC film and the various core level spectra wererecorded. The Cr 2p core level spectrum for the 7 minutesputtered Ni-Cr/DLC film is shown in Fig. 1(c). The peakscentered at ~576 and 583.4 eV correspond to the atomictransition of Cr from 2p3/2 to 2p1/2. However, as compared tostandard data [13], some shifting in the peak positions of Crwas observed, which may be due to a sputtering effect. Onthe other hand, no evidence for the presence of Ni in the Ni-Cr/DLC film was observed. Therefore, sputtering for 15minutes on the Ni-Cr/DLC film was further performed. TheNi 2p core level spectrum for the 22 minute sputtered Ni-Cr/DLC film is shown in Fig. 1(d). The Ni 2p core level spectrareveal the presence of two peaks centered at 852.4 eV and869.8 eV, assigned to the atomic transition of Ni from 2p3/2 to2p1/2. The observed peaks in the Ni 2p core level spectra weregood in agreement with standard results [13]. In addition, C

Esds2

6 1 νs–( )df------------------------ 1

Rf---- 1

R0-----–⎝ ⎠

⎛ ⎞

Fig. 1. Core level spectra of, (a) C 1s before sputtering, (b) C 1s aftersputtering of 22 minute, (c) Cr (2p), and (d) Ni (2p) for Ni-Cr/DLCfilms.

Effect of Metallic Interfacial Layers on the Properties of Diamond-Like Carbon Thin Films 233

1s core level spectra of the 22 minute sputtered Ni-Cr / DLCfilm were also recorded, as shown in Fig. 1(b). C 1s corelevel spectra clearly show the strong peak at 283 eV. It isinteresting to note that sputtering of 22 minutes changes theposition of the C 1s peak toward the lower binding energyside. The shifting of the C 1s peak was found to be ~1.2 eV(284.2 eV to 283.0 eV). Thus, the XPS results reveal the for-mation of an intrinsic multilayer in the Ni-Cr/DLC film.Effort has also been made to prove the formation of anintrinsic multilayer thermodynamically. It is well known thatthe melting point and temperature at vapor pressure of 10−2

Torr of Ni are lower than those of Cr (as given in Table 1)[14]. Therefore, it may be possible that during evaporation,Ni evaporates first, and beyond the critical temperature ofNi, the evaporation of Cr starts. When the DLC layer isdeposited over the Ni-Cr layer, and then the complete struc-ture may be referred to as a multilayer structure. These findingsmade through thermodynamic studies were also proventhrough XPS measurement; after sputtering of 22 minutes,evidence for the presence of Ni was observed, whereas thepresence of Cr was visualized after 7 minutes of sputtering.Furthermore, no evidence for the formation of oxide in Niwas observed, but the Cr peak exhibited some oxide forma-tion. Generally, the peaks corresponding to the atomic transi-tion of 2p3/2 to 2p1/2 in pure Cr are observed at 574.1 and583.4 eV, respectively, whereas peaks assigned to the sametransition in chromium oxide (Cr2O3) are observed at 576.6and 586.3 eV, respectively. However, in the present study,the peaks in the Cr 2p core level spectra corresponding to thesame atomic transition were observed at ~576 and 583.4 eV,respectively. Thus, it has been found that the peak corre-sponding to 2p3/2 lies in the region between Cr and Cr2O3.The reason for oxide formation in Cr only rather than in Nimay be that during evaporation of Ni-Cr, the Cr layer wasgrown on top (as confirmed by XPS as well as thermody-namic studies). Furthermore, when the Ni-Cr coated sub-strates were placed in the RF-PECVD chamber for successiveDLC layer growth they were exposed in ambient air, and atthat time oxygen reacted only with the top Cr layer. Thus,the oxide in the Cr layer may be due to surface contamina-tion only.

XRD patterns of DLC, Al/DLC, and Ni-Cr/DLC filmsgrown at a negative self bias of 100 V are shown in Fig. 2.The pure DLC film was found to be amorphous in nature,which is confirmed by the observation of a broad hump,

comprising several small peaks, in the 2θ range of 21° to36°. The XRD spectrum of the Al/DLC film shows a sharppeak at 2θ of 38.4°, which corresponds to Al <111> [15]with a broad hump in the range of 21° to 36° assigned to theamorphous nature of carbon. Similarly, the Ni-Cr/DLC filmshows a broad peak at 2θ of 44° and a broad hump in therange of 21° to 36°. The cubic Ni <111> peak is obtained at44.5° [16] whereas the cubic Cr <111> peak is derived at43.5° [17], as confirmed by the JCPDS card. Therefore, thepeak at 44° may be ascribed to either Ni or Cr alone or maybe due to the contribution of both Ni and Cr. However, nooxidation state was observed with metals, because the growthof a metal (Ni-Cr and Al) layer was carried out at base pres-sure exceeding 10−5 Torr.

Figure 3 shows SEM micrographs of DLC and Ni-Cr/DLC films grown at a negative self bias of 100 V. Reports inthe literature describe agglomeration in DLC films with anamorphous structure [18]. The SEM micrograph of the DLCfilm clearly reveals a bulk amorphous structure. However,the Ni-Cr/DLC film showed clustering and agglomeration ofthe particles. These agglomerates in the Ni-Cr/DLC filmmay be of sp2 clusters, which enhance its adhesion and con-ductivity.

Adhesion of DLC films deposited on Al and Ni-Cr coatedSi substrates was investigated in term of stress and the obtainedresults were compared with the pure DLC film. The varia-tion of the residual stress for various types of DLC films isshown in Fig. 4. From the figure it is evident that the maxi-mum residual stress (2.8 GPa) was observed in the pureDLC film, which decreases to 1.45 and 1.2 GPa in the Al/DLC and Ni-Cr/DLC films, respectively. Interfacial Al and

Table 1. Melting point and temperature vapor pressure at 10−2 mbar of Ni, Cr and Ni-Cr composite

Metal Melting Point(°C)

Temperature (°C) Vapor Pressure at 10−2 Mbar

Ni 1453 1130Cr 1875 1360

Ni-Cr 1500 1600

Fig. 2. XRD patterns of (a) Ni-Cr/DLC, (b) Al/DLC, and (c) DLCfilms.

234 Neeraj Dwivedi et al.

Ni-Cr layers act as adhesive layers for the successive DLClayer. Notably, there is competition between adhesive forces(which exist at the interface) and cohesive forces (whichexist within the material). If adhesive forces dominant overcohesive forces, then the film will properly adhere to thesubstrate; otherwise it is peeled off. It is realized that theintroduction of interfacial Al and Ni-Cr layers enhances theadhesive forces at the interface, which leads to greater adhe-sion and lower residual stress in the Al/DLC and Ni-Cr/DLCfilms. Fallon et al. [19] systematically studied the stress incarbon films while S. Paul [20] suggested the mechanismregarding stress and adhesion of DLC films on metals. SinceAl and Ni-Cr layers are soft in nature, the soft/hard layerconcept may also be valid to provide the needed relaxationof residual stress in the Al/DLC and Ni-Cr/DLC films. Giotiet al. also reported similar work previously [21]. Recently,we also observed very low residual stress in DLC films byemploying a Cu/a-C:H multilayer structure [22]. Anotherpossible explanation regarding stress relaxation in the Al/DLC and Ni-Cr/DLC films involves a bilayer and multilayerstructure concept [23]. Recently, Han et al. suggested a mul-tilayer structure concept to obtain lower stress in tetrahedralamorphous carbon films [24]. It was noted that compara-

tively higher residual stress in a pure DLC film may also bedue to interdiffusion of carbon into the substrate Si; thisleads to the formation of a mixed a-SiC or a-SiC:H layer atthe Si/DLC interface, which results in interfacial mismatchand therefore interfacial stress [24]. In contrast, the introduc-tion of interfacial Al and Ni-Cr layers prevents the interdiffu-sion of carbon into Si and circumvents interfacial mismatch.This may also explain why Al/DLC and Ni-Cr/DLC filmsexhibited lower residual stress than the pure DLC film. AmongAl/DLC and Ni-Cr/DLC films, the observed comparativelylower residual stress in the Ni-Cr/DLC film may be due tothe formation of an intrinsic multilayer in its structure (whereasthe Al/DLC film shows a bilayer structure), as revealed byXPS and thermodynamics studies. It is also noted that theresidual stress in DLC films strongly depends on graphite-like sp2 bonding and varies inversely with it; i.e. higher sp2

bonding lowers the residual stress. It is found in the literaturethat the addition of metal in the DLC film can enhance thesp2 bonding [25]. Thus, the obtained lower stress in the Al/DLC and Ni-Cr/DLC films may also be due to increased sp2

fraction in its structure. Further, the observed lower residualstress in Al/DLC and Ni-Cr/DLC may also accompany inco-herent stress relaxation as well as a percolation structuraltransition from a-C:H rich structure to a carbide rich struc-ture [26]. Taking into consideration these findings, a lowerstress DLC film can be grown by employing low cost evap-oration and RF-PECVD techniques rather than complex andcostly RF-sputtering and RF-PECVD techniques.

Nano indentation is an advanced technique to study themechanical properties of thin films. Figure 5(a) and 5(b)show the variation of hardness (H) and elastic modulus (E)for DLC, Al/DLC, and Ni-Cr/DLC films. It is evident fromthe figure that the DLC films attained maximum values of Hand E relative to the Al/DLC and Ni-Cr/DLC films. The val-ues of H and E for the DLC film were found to be 12 and136 GPa, respectively. The Al/DLC film exhibited H and Evalues of 7.3 GPa and 66 GPa, respectively, versus 8.5 GPaand 74 GPa, respectively, for the Ni-Cr/DLC film. Observedlower H value in the Al/DLC and Ni-Cr/DLC films may bedue to enhancement of sp2 bonding or clustering [25]. Inaddition, the mechanical properties were also stronglydependent on the substrate effects. If the film is thin and theindenter penetrates one-third of the total film thickness, thesubstrate effect cannot be ignored [27]. Zhang and Komvopou-los extensively studied the effects of the substrate on themechanical properties of very thin tetrahedral amorphouscarbon films [28,29]. In light of these considerations, the lowH and E values observed in these films may also be due tosubstrate effects. Furthermore, interfacial metal layers alsoinfluence H and E, because Al and Ni-Cr are soft and toughin nature and deformed under applied load. Thus, the addi-tion of these interfacial layers to the DLC film lowers its Hand E values. Plastic deformation energy (Ur) was found to

Fig. 3. SEM micrographs of (a) DLC, and (b) Ni-Cr/DLC films.

Fig. 4. Variation of stress for DLC, Al/DLC and Ni-Cr/DLC films.

Effect of Metallic Interfacial Layers on the Properties of Diamond-Like Carbon Thin Films 235

be an important factor that explores the elastic-plastic prop-erties of thin films in the form of energy. The plastic defor-mation energy for the pyramid indenter was estimated usingthe relation [27] given as Eq. (2)

(2)

where wo is the geometry constant and takes a value of 1.3for the pyramid indenter, P is the load, and ψ is the halfangle of the Berkovich indenter and has a value 65.3°. Thevariation of Ur for various films is shown in Fig. 5(c). Ur

varies inversely with H. It can be seen from the figure thatthe DLC film exhibited the minimum value of Ur amongthe three kinds of films. The value of Ur for DLC wasfound to be 1.23×10−9 joule, and increased to 1.57×10−9

joule for the Al/DLC film. However, a small decrease in Ur

was observed for the Ni-Cr/DLC film, which showed avalue of 1.46×10−9 joule.

Figure 6 shows the variation of conductivity as a functionof temperature for as grown DLC, Al/DLC, and Ni-Cr/DLCfilms. The conductivity of these films was examined at threedifferent temperatures, 27, 50, and 100 °C. The figure showsthat the pure DLC film showed a value of conductivity2×10−12 Ω−1cm-1 at 27°C, which increased to 9×10−12 Ω−1cm−1

at 50 °C. At 100 °C, the conductivity further increased andattained a maximum value 5×10−10 Ω−1cm−1. The Al/DLCfilm exhibited very high conductivity of 1.7×10−1, 2.4×10−1,and 2.7×10−1 Ω−1cm−1 at temperatures of 27, 50, and 100 °C,respectively. However, the Ni-Cr/DLC film showed conduc-

tivity values of 1.0, 1.7, and 1.9 Ω−1cm−1 at 27, 50, and 100 °C,respectively. The effect of temperature on the conductivity ofthe Al/DLC and Ni-Cr/DLC films was found to be negligi-ble, as these films exhibited almost the same values of con-ductivity between 27 and 100 °C. However, the conductivityof the Al/DLC and Ni-Cr/DLC films at room temperature(27 °C) was found to be very high. It was thought that thiscould be attributed to diffusion of metal during the fabricationof the metal contact for the coplanar structure. However,through repeated experiments we confirmed that metal wasnot diffused into the films. Meanwhile, significant variationin the values of conductivity with temperature from 27 to100 °C was observed for the pure DLC film. As in pure met-als (bulk), the conductivity decreased with increasing tem-perature due to collisions and scattering of charge carriers.However, in the present case, Al and Ni-Cr were used in filmform and the DLC layer was always located at the top of theconfiguration. Therefore, the effect of temperature on theconductivity of Al/DLC and Ni-Cr/DLC films was found tobe negligible.

4. CONCLUSIONS

Good quality DLC films on a metallic layer (depositedusing thermal evaporator) were deposited in a RF PECVDsystem (having high base pressure, 10−3 Torr), thus potentiallyaffording a low cost process. It was found that the metallicinterface layers of Al and Ni-Cr played an important role inreducing the residual compressive stress of the DLC films.The metal interface improves the adhesion of the DLC filmwith the substrate with some fall of H and E. However, thereduction of H and E values is lower in comparison to the

Ur13--- 1

ω0tan2Ψ-------------------- 1

H--------P3/2=

Fig. 5. Variation of (a) H, (b) E and (c) Ur for DLC, Al/DLC and Ni-Cr/DLC films.

Fig. 6. Variation of conductivity for DLC, Al/DLC and Ni-Cr/DLCfilms. The figure given in the inset shows the variation of conductivityversus temperature for DLC film.

236 Neeraj Dwivedi et al.

reduction of stresses in these films. Room temperature con-ductivity was improved significantly due to Al and Ni-Crinterface layers. The observed drastic improvement of con-ductivity in the Al/DLC and Ni-Cr/DLC films at room tem-perature can be useful for room temperature hard electronicdevices. XPS and XRD results confirmed the presence of Al,Ni, and Cr in the Al/DLC and Ni-Cr/DLC films. SEM micro-graphs showed agglomeration and a metallic network in thebackground.

ACKNOWLEDGMENTS

The authors are grateful to the Director, National PhysicalLaboratory, New Delhi (India) for his kind support. Theauthors are very thankful to Dr. Govind for providing a XPSmeasurement facility. One of the authors, ND, acknowledgesCSIR, Govt. of India for providing a SRF fellowship. Weacknowledge CSIR, Govt. of India for sponsoring NetworkProject NWP-0027 and for their financial support.

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