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Enhanced oxygen-induced properties of bulk oxygenated amorphous carbonfilms deposited with an anode layer ion source

Abdul Mateen Qasima, Qingdong Ruana, Ricky K.Y. Fua, Farhat Alib, Babak Mehrjoua, Hao Wud,Liangliang Liua, Zhongzhen Wua,c, Paul K. Chua,∗

a Department of Physics, Department of Materials Science and Engineering, and Department of Biomedical Engineering, City University of Hong Kong, Tat Chee Avenue,Kowloon, Hong Kong, ChinabDepartment of Materials Engineering, Pakistan Institute of Engineering & Applied Sciences, Nilore, Islamabad, Pakistanc School of Advance Materials, Peking University Shenzhen Graduate School, Shenzhen, 518055, Chinad School of Mechanical and Automotive Engineering, South China University of Technology, Guangzhou, Guangdong, 510641, China

A R T I C L E I N F O

Keywords:Anode layer ion sourceOxygenHydrogenated amorphous carbon filmMechanical propertiesTribology

A B S T R A C T

Bulk oxygenated amorphous carbon films (a-C:H:O) with different oxygen concentrations are deposited on high-speed steel with a chromium carbide (CrCx) interlayer using an anode layer ion source (ALIS) and their struc-tural, mechanical, and tribological properties are determined. Films with 7 at%, 9 at%, 13 at%, and 27 at %oxygen are prepared by changing the O2 to C2H2 ratio in the gas mixture during deposition. The film depositedwithout intentionally introducing oxygen has an oxygen concentration of 5% and serves as the control sample.The surface of the control film contains nanodots which can be reduced with 10% O2 in the gas mixture. The bestproperties are obtained from the film containing 7 at% oxygen. The hardness increases from 19.5 GPa for thecontrol film to 29.0 GPa of the a-C:H:O film with 7 at% oxygen. At the same time, the wear rate decreases by 2.5times to 2.0× 10−7 mm3/m-N and the friction coefficient by 10 times to 0.09. If the oxygen concentration in thefilm exceeds 7%, the hardness begins to decrease and wear rate increases while the friction coefficient is reducedslightly. Raman scattering reveals a smaller ID/IG ratio and more disorder in the 7% sample and XPS shows thatthe improvement stems from increased sp3 hybridization and C–C bonding. However, if the oxygen con-centration in the gas mixture is increased further, oxygen ions begin to replace carbon atoms in the sp3 structureand more O–C–O, C]O, and O–C]O bonds are created. The carbonyl and carboxylic groups soften the matrixand degrade the mechanical properties. Our results clarify the effects and mechanism of oxygen on the me-chanical and tribological properties of bulk amorphous carbon films.

1. Introduction

Amorphous carbon (a-C) or diamond-like carbon (DLC) coatingswhich have many outstanding properties such as high hardness, smallfriction coefficients, low wear rates, high chemical stability, high op-tical transparency, wide bandgap, and good biocompatibility [1] arewidely used in the tool, automobile, aerospace, and electronics in-dustries. DLC is classified according to the hydrogen content as hy-drogenated (a-C:H) and non-hydrogenated (a-C) amorphous carbon.They have different properties due to the different sp3 and sp2 fractions.A small sp2/sp3 ratio introduces diamond characteristics and a largeratio produces graphite properties. The physical and mechanicalproperties are mainly determined by the sp3 fraction [1] and a largersp2 content can enhance the tribological properties such as lower

friction [2]. Hence the characteristics of amorphous carbon coatingscan be tailored to a wide range. This capability imparts great flexibilityto satisfy a wide range of industrial demands.

The sp2/sp3 ratio in DLC films can be adjusted by physical andchemical means [3,5]. There are number of physical techniques todeposit the hydrogenated amorphous carbon (a-C:H) films, for example,physical vapor deposition (PVD) techniques that includes magnetronsputtering and cathodic arc evaporation (CVA) and chemical vapordeposition techniques such as plasma-enhanced -CVD (PECVD) relyingon the generation of a gaseous discharge in the presence of a hydro-genated gas [3,4]. In these processes, the sp2/sp3 ratio is controlled bythe energy and flux of the ions. Another way to control the properties ofa-C:H films is to modify the structure by chemical means and metalslike Cr, Ni, Ti, Cu, Al, W, and Co [5–12] and non-metals such as N, Si, F,

https://doi.org/10.1016/j.vacuum.2019.108915Received 22 July 2019; Received in revised form 27 August 2019; Accepted 28 August 2019

∗ Corresponding author.E-mail address: paul.chu@cityu.edu.hk (P.K. Chu).

Vacuum 169 (2019) 108915

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and O [13–15] have been shown to tailor the properties of a-C:H films.Among the various types of metal/non-metal doped a-C films,

oxygen-doped or oxygen plasma-treated a-C (a-C:O/a-C:H:O) films haverecently attracted interest on account of their unique properties [16].For example, Santini et al. [17] showed that a-C:O films deposited bymagnetron sputtering had potential in resistive memory applications byreplacing graphitic oxide in carbon-based electronics. Palomaki et al.[18] demonstrated good electrochemical performance in various redoxsystems and the optical characteristics of a-C:H:O films were better[19,20]. Safaie et al. [21] used the mixture of CH4 and O2 in a radiofrequency -PECVD reactor to study the structure and mechanicalproperties of oxygen-doped films. They used a maximum of 10 vol%oxygen that corresponded to a concentration of 11 at% in the film andthe oxygen plasma-treated a-C:H films had improved scratch resistanceand mechanical properties [5,16,22–24]. The common fabricationmethods involve either in situ oxygen doping during magnetron sput-tering [17,18,25] and RF-PECVD [21] or the use of an oxygen plasmasustained by predominantly RF or ECR sources to modify the surfaceproperties [22–24]. In the case of plasma treatment, the modified sur-face layer is normally quite thin and can be abraded easily to expose theunderlying a-C:H layer. Therefore, owing to the increasing interest in a-C:H:O films and concomitant limitations, it is important to study themechanical and tribological properties of a-C:H:O films with differentoxygen contents.

In this work, 1750 ± 150 nm thick a-C:H:O films are depositedusing an anode layer ion source (ALIS) on a magnetron-sputtered gra-dient CrCx interlayer on high-speed steel or silicon. ALIS is a widelyused industrial source capable of depositing large-area DLC coatingswith good coverage at high deposition rates [26]. Another advantage ofALIS is its ability to provide a large ion flux to form dense coatings [27].In our experiments, the oxygen concentration is adjusted by varying theratio of oxygen to acetylene in the gas mixture and the hardness (H),elastic modulus (E), friction coefficients, and wear rates are in-vestigated systematically. Our results indicate that an optimal range ofoxygen concentration is crucial to the mechanical and tribologicalproperties and the enhancement mechanisms involve structural mod-ification and formation of new bonds.

2. Experimental details

4241 high speed steel (4241-HSS) and p-type silicon (100-Si) waferswere used as the substrates. The HSS samples were cut into25×25×1mm pieces, ground with SiC abrasive paper, and me-chanically polished with diamond paste. The HSS and Si substrateswere ultrasonically cleaned with acetone and ethanol for 15min eachfollowed by drying with nitrogen. The samples were placed on a100×250mm plate and mounted onto the substrate holder positionedcoaxially in the vacuum chamber at the center of the plasma immersionion implantation and deposition instrument in the Plasma Laboratory ofCity University of Hong Kong [28]. The system was equipped with twomiddle-frequency magnetron sputtering targets and a gridless anodelayer ion source (ALIS) and both units were connected to current-con-trol power units operated at a frequency of 40.0 kHz and duty cycle of80%.

After the chamber was evacuated to a base pressure of

3.0× 10−3 Pa, an argon glow discharge (Ar flow rate= 100 sccm,chamber pressure= 1.0 Pa, and bias voltage=−1.2 kV) was sustainedfor 15min to clean the sample. An interlayer of Cr and gradient Cr–Cwas first deposited using a chromium (99.9% pure) magnetron sput-tering target (7.5 x 34.5 cm2). The distance between the substrate andthe target was 10 cm. The interlayer was deposited in a pure Ar dis-charge at a discharge current of 6.0 A and C2H2 was introduced toproduce the gradient CrCx layer, i.e., from Cr-rich to C-rich. The biaswas −4.0 kV during the process and the average thickness of the in-terlayer was controlled to be 1.7 μm. Finally, the DLC layer was de-posited with the ALIS. The discharge current was 1.0 A, power was700W, and Ar to C2H2 ratio was 1:5 and no bias was applied. Thesample stage was stationary during deposition of the interlayer and DLCfilm. The oxygenated amorphous carbon film (a-C:H:O) was producedusing the same ALIS parameters but with different ratios of O2:C2H2 inthe gas mixture to obtain films with different oxygen concentrations.Here, the ratio is defined as O2/(O2 + C2H2) and the sample designa-tions and characteristics of the amorphous carbon coatings are listed inTables 1 and 2.

The film thickness, composition, structure, and morphology wereexamined on a field-emission scanning electron microscope (FE-SEM,ZEISS SUPRA® 55) equipped with an energy-dispersive x-ray spectro-meter (EDS). Atomic force microscopy (AFM, Auto Probe CP, ParkScientific Instruments, USA) was conducted to observe the surfacemorphology and roughness in the tapping mode in an area of5× 5 μm2. The bonding information was obtained at room temperatureby Raman scattering (LabRamHR, Horiba) using a 514.5 nm laser. X-rayphotoelectron spectroscopy (Physical electronics PHI5802) was utilizedin this study to determine the chemical states and sp3 content of carbonin the amorphous carbon films. The monochromatic Al Kα line withexcitation energy of 15 keV was used to acquire the spectra. Theequipment was calibrated with the Cu 2p3/2 and Au 4f7/2 standardsfor BE scale alignment. The samples were stored at a pressure of 6.0 x10−5 Pa prior loading into the spectrometer. A thin gold layer wasdeposited on half of the samples and all the BEs were referenced to Au4f7/2 at 84.0 eV. The XPS spectrum of the surface gold layer is shown inFig. 7f. Ar ion sputtering (2 kV and sputtering rate of about 9 nm/minbased on a SiO2/Si reference standard) was used to clean the surfaceprior to acquiring the spectra from an area of 0.5mm2 [29]. Thestructure and bonding were also characterized by FTIR (Fourier trans-form infrared spectroscopy, PerkinElmer) in the total transmissionmode in the range between 500 cm−1 and 4000 cm−1.

The mechanical properties such as hardness (H) and elastic modulus

Table 1Sample designations, mechanical, and Raman characteristics of the amorphous carbon coatings.

Sample name % O2 in the gas mixture % O2 in the film (at %) Hardness (GPa) Elastic Modulus (GPa) ID/IG FWHM (G) (cm−1)

Native 0 ≤5 19.52 ± 0.35 125.53 ± 0.89 0.71 152.3O-1 10 7.43 27.88 ± 2.33 172.77 ± 7.28 0.55 170.46O-2 20 9.6 26.06 ± 1.30 158.36 ± 6.98 0.66 162.11O-3 30 12.7 22.24 ± 1.60 138.50 ± 3.19 0.64 156.2O-4 40 27.7 10.66 ± 1.20 114.32 ± 4.8 1.03 116.95O-5 50 – 8.6 ± 0.58 93.43 ± 2.739 1.19 135.35

Table 2Sample designations and XPS characteristics of the amorphous carbon coatings.

Samplename

% O2 in thegas mixture

% O2 inthe film(at %)

XPS

sp3 sp2 O–C–O C = O O–C]O

Native 0 ≤5 55.1 29.39 3.83 11.68 –O-1 10 7.43 58.12 22.76 6.69 2.13 10.3O-2 20 9.6 49.76 30.84 9.24 1.83 8.33O-3 30 12.7 44.78 32.49 8.43 2.94 11.36O-4 40 27.7 20.99 58.63 3.08 2.96 14.34

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(E) were determined on a microhardness tester (Fisheroscope,HM2000XYp; Fischer Technology, Inc. Windsor) with a Vickers in-denter tip in the constant load mode. The instrument was calibratedwith standard glass (BK7) which represented 0.06% error and thestandard calibration had a maximum uncertainty up to 4%. A load of10mN was utilized to limit the maximum depth to 10% of the depositedfilm and avoid the effects of the interlayer and substrate. The data re-presented the average of six measurements with the average distancebetween each indentation being ≥200 μm. The wear rate and coeffi-cient of friction were measured on a ball-on-disk instrument (Rtec MFT-5000) using a silicon nitride (Si3N4) ball 3 mm in diameter at roomtemperature and ambient environment. A normal force of 4 N was ap-plied at a speed of 200 rpm and the track diameter was 5mm. The trackwas examined by optical microscopy, whereas the wear volume wasdetermined by 3D optical profilometry (3D profiler). The wear rate wascalculated based on the corresponding wear volume by the followingequation:

=K υsF (1)

where K is the wear rate, ʋ is the wear volume in m3, s is the totalsliding distance in meters, and F is the applied normal force in Newton[30]. The adhesion strength of the film was determined on a scratchtester (WS-2005, Zhongke Kaihua Technology, China). A preload of 5 Nwas set and the reading was adjusted to 0 N. Afterwards, the load wasincreased progressively from 0N to 100 N at a rate of 50 N/min for ascratch length of 3mm and the scratches were examined by SEM andoptical microscopy.

3. Results and discussion

The effects of the oxygen content in the gas mixture on the de-position rate and oxygen concentration in the films are presented inFig. 1. The deposition rate decreases from 700 nm/h for the control film(without the intentional introduction of oxygen) to almost zero as theoxygen content in the gas mixture reaches 50%. Oxygen is used to etchthe a-C:H layers [31] and the etching rate depends on the density ofoxygen radicals in the plasma. McKindra et al. [25] and Adamopouloset al. [19] prepared oxygenated a-C and a-C:H films by sputtering andobserved diminishing deposition rates with increasing oxygen partialpressure. It is believed that oxygen chemically etches the carbon surfaceby forming CO or CO2 and oxygen radicals formed by the ALIS alsoproduce larger physical and chemical effects to accelerate etching [23].Eventually, for a larger oxygen ratio, the etching rate dominates thedeposition rates resulting in a condition in which no net depositionoccurs. The presence of chromium oxide on the surface as indicated bythe Raman spectrum in Fig. 6 and shows the dominant effects of

etching. The oxygen contents determined by EDS (Fig. 1) increase from5 at% to 27 at%.

The cross-sectional FE-SEM images of the control and O-3 films aredepicted in Figs. 2a and 3a, respectively. The film consists of threedifferent layers. The interlayer has a columnar structure and the columnwidth increases from bottom to top with a tightly packed morphologydue to 4 kV ion bombardment. The first layer consists of Cr followed bya gradient CrCx layer deposited by progressively increasing the flowrate of C2H2. The second layer is very thin and consists of a largerproportion of carbon with a graphite-like hydrocarbon structure. Thisinterlayer architecture has been shown to improve the adhesionstrength between the a-C:H film and underlying HSS substrate [32,33].For example, Liu et al. fabricated an a-C:H layer with a thickness of over50 μm with excellent adhesion with the gradient interlayer of CrC. Thethird layer at the top is the a-C:H layer. The carbon layer is amorphousand dense without visible pores. ALIS is capable of producing a highlyionized plasma that forms a dense coating with good coverage andcharacteristics [34]. Fig. 2c–d and 3c-d show the EDS line profiles andelemental maps of the control and O-3 films, respectively, of oxygen.Owing to the non-UHV conditions in the instrument, a small amount ofoxygen in the film comes from the residual vacuum and upon exposureto air after deposition. Comparison of Fig. 2c–d and 3c-d show thatmore oxygen can be incorporated into films by increasing the oxygenpercent in the gas mixture and oxygen is uniformly distributed in thefilms.

Fig. 4a and b display the FE-SEM images of the control sample atdifferent magnifications. The surface of the control film is smooth ingeneral and there are many tiny particles about 5 nm in size on thesurface. During sputtering of graphite in Ar and C2H2, Wasy et al. andDe Vriendt et al. [35,36] observed generation of nanoparticles andnanoscale powders were produced by the C2H2 plasma [37]. De vriendtet al. also showed that sputtering in the pulsed mode produced particleswith different shape and size depending on the pulse-off time. In thisstudy, ALIS is operated in the pulsed mode with a 5 μs pulse-off timewhich is less than the minimum duration reported before, i.e., 10 μs.According to the model, during the pulse-off period, fragments of C2H2

with high kinetic energy overcome Coulomb repulsion and coagulate inthe form of nanoparticles, whereas a shorter pulse-off time hindersparticle growth and agglomeration. These particles or nanodots arereduced to a significant extent in the oxygenated films as shown inFig. 4c and d and the convex network become prominent with the ad-dition of oxygen due to etching artifacts. This effect can be explained asfollows. When O2 and C2H2 are introduced to the ALIS that generatesthe gas plasma and transports it to the substrate, O2 radicals in theplasma are highly reactive and hinder the formation of carbon particlesdue to etching of the graphite component. Ono et al. observed that anO2 plasma with additional ICP source in PLD of graphite was capable ofrestricting the formation of carbon particles compared to oxygen mo-lecules without the ICP plasma [38]. The size of the convex network isabout 500 nm in diameter and decreases with larger oxygen contents inthe plasma further confirming the etching effects. Our results indicatethat 10 vol% of O2 is sufficient to impede particle formation and growthin the deposited film. The sample deposited under O-5 conditions, i.e.,50 vol%, oxygen does not have a measurable a-C:H:O layer (Fig. 4f)indicating dominant etching effects.

The AFM images in Fig. 5a–e reveal features similar to those ob-served by FE-SEM. Ra of the control sample is 13 nm and small particlesare visible. Ra of the oxygenated films decreases to 3.7 nm and theconvex surface features visible in the scanned images become morepronounced with increasing oxygen content in the gas mixture. A de-crease in the surface roughness with increasing oxygen content is at-tributed to the uniform etching effects produced by oxygen radicals andremoval of nanodots. Increase in the roughness for sample depositedunder the O-5 condition is due to the simultaneous effect of etching andformation of Cr–O compounds on the surface which acts as a maskhindering uniform etching.

Fig. 1. The deposition rate and oxygen concentration in the a-C:H coatings as afunction of oxygen volume percent in the gas mixture.

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Raman scattering is a useful tool to determine the structure andbonding nature of carbon in the films. The typical Raman spectrum ofamorphous carbon can be deconvoluted into the D-peak at1358 ± 2 cm−1 and G-peak at 1575 ± 10 cm−1. In amorphouscarbon, the D-peak originates from the breathing modes of A1g sym-metry phonons arising from sp2 carbon atoms in the rings, whereas theG-band originates from the E2g mode due to bond stretching of all pairsof sp2 atoms in both rings and chains [39–41]. According to the pro-posed extended three-stage model for amorphous carbon films [39,40],the ID/IG ratio is a measure of the size of the sp2 phase organized inrings and FWHM (G) is a measure of the structural disorder arising fromthe bond angle and bond length disorder. If ID/IG is reduced, the sp2

phase is mainly organized in chains or if rings are present, the bonds are

not fully delocalized. Hence, a smaller ID/IG ratio suggests a larger sp3

content. The FWHM (G) is linked to the structural disorder and asmaller FWHM (G) corresponds to defect-free sp2 clusters in theamorphous matrix. On the other hand, a larger value represents dis-order stemming from the larger bond length and angle [34,39].

Fig. 6a presents the Raman scattering spectra of the a-C:H films withdifferent O2 concentrations. The control film shows an ID/IG ratio of0.71 with the G-peak at 1560 ± 1.05 cm−1 and FWHM (G) of156 cm−1. According to the three-stage model [39,40], the film iscomposed of amorphous carbon. For the film with 7.0 at% oxygen (O-1), the ID/IG ratio and G-peak position decrease while FWHM (G) in-creases, indicating that oxygen increases the sp3 content, C–C bonding,disorder and reduces the sp2 cluster size [39]. If the oxygen

Fig. 2. (a) Cross-sectional FE-SEM image showing the three layers on Cr-CrCx/C and CrCx/a-C:H (Control); (b) Enlarged view of the deposited film; (c) EDS lineprofile as shown in (e); (d) EDS oxygen concentration map from the image shown in (b). The data are from the a-C:H control.

Fig. 3. (a) Cross-sectional FE-SEM imageshowing the three layers on Cr-CrCx/C andCrCx/a-C:H:O (O-3); (b) Enlarged view of thedeposited a-C:H:O film only; (c) EDS line profilefrom the region as shown in (e); (d) EDS oxygenconcentration map from the image shown in (b).The data shown are from sample O-3 and a-C:H:O film.

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concentration is higher, for instance, O-2, O-3, and O-4, the ID/IG ratioand G-peak position increase again and FWHM (G) decrease indicatingan ordering of sp2 clusters. Mckindra et al. [25] and Palomaki et al.[18] prepared oxygen-doped a-C films by magnetron sputtering andobserved that the ID/IG ratio decreased when the oxygen concentrationwas 6–8 at%. However, the ID/IG ratio observed by the former groupdecreased continuously but increased again as observed by the latter

group, but no information about FWHM (G) was provided. It was at-tributed to the different deposition parameters. Dwivede et al. [22]etched a-C:H films with an oxygen plasma and found that the ID/IG ratiodecreased in addition to decreased G-peak position indicating higher

Fig. 4. Plain-view FE-SEM images: (a, b) Control a-C:H film at low and highmagnifications, respectively and small nanodots can be observed from themagnified image; (c, d) O-1 a-C:H:O at low and high magnifications, respec-tively. The nanodots are not present on the surface and convex surface featurescan be seen from the magnified image; (e, f) Images of samples deposited underO-3 and O-5 conditions.

Fig. 5. AFM surface scan from a 5×5 μm area: (a)Native a-C:H film, (b) Sample O-1, (c) Sample O-3,(d) Sample O-4, and (e) Sample O-5 a-C:H:O films.The insets show the phase-contrast mode. Smallnanodots can be observed from the control film butare significantly reduced in the oxygenated films; (f)Evolution of surface roughness (Ra) with increasingoxygen contents in the gas mixture.

Fig. 6. (a) Raman scattering spectra of the deposited films with different con-centrations of oxygen; (b) Evolution of the ID/IG ratio and FWHM (G) peakalong with G-Peak position with increasing oxygen percentage in the gas mix-ture.

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disorder and reduced sp2 clustering. The Raman scattering spectrum ofthe film prepared with 50 vol% oxygen shows the extra peaks at674.24 cm−1 and 998.68 cm−1 which can be assigned to different Cr–Orelated compounds on the surface of the film [42–44]. The presence ofthese compounds on the surface reveals that there is no carbon filmdeposition but instead the CrCx adhesion layer is etched and newcompounds of Cr–O has been formed.

The high-resolution C1s spectra of the control and oxygenated filmsare shown in Fig. 7 a- e. The spectra are deconvoluted by the XPSpeak4.0 software [45,46]. In brief, the Shirley background was subtractedfrom the peaks at 280.0 eV–295.0 eV and the Voight function was ap-plied to fit the peaks to reveal the sp3 and sp2 phases. The C1s spectrumof the control film is deconvoluted into four sub-peaks. The peaks at283.79 eV and 284.29 eV correspond to carbon bonding in sp2 and sp3,respectively, and the other two at 285.59 eV and 286.79 eV are relatedto C–O and C]O. The deconvoluted spectrum of the O-1 sample inFig. 7b shows sp2, sp3, C–O, C]O, and O–C]O at 284.23 eV,284.73 eV, 286.03 eV, 287.23 eV, and 288.63 eV, respectively. Thepoints of interest are the larger sp3 and smaller sp2 fractions resulting inincreasing sp3/sp2 ratios from 1.87 to 2.85. Compared to the controlsample, oxygen has two effects. The resulting larger sp3 content is

consistent with the Raman data and a larger degree of C–O bondingtogether with the presence of carbonyl (C]O) and carboxyl (O–C]O)groups is observed. If the oxygen content is further increased, the sp3/sp2 ratio decreases and the oxygen-related components increase as inthe case of sample O-3 and O-4. Similar results have been obtained fromnon-hydrogenated a-C films deposited by sputtering in the Ar and O2

plasma or ex-situ oxygen plasma etching of a-C:H films [14,15,17,23].The oxygen plasma serves several purposes: preferential etching of thesp2 phase, breakup of C]C and C–H bonds, promotion of C–C bonding,and creation of carbonyl and carboxyl groups. According to our find-ings, an optimal amount of oxygen in the gas mixture and consequentlyin the films is necessary to enhance sp3 bonding in the films.

The FTIR spectra of the control, O-1, and O-3 samples confirm theincrease in oxygen-induced compounds after addition of oxygen asshown in Fig. 8. The presence of intense peaks from the O-1 and O-3films in the region below 1200 cm−1, around 1500 cm-1, 2350 cm−1,and 2850 to 2950 cm-1 and a broad peak after 3000 cm−1 indicates out-of-plane vibration of C–O–H, aliphatic and benzene rings, stretchingvibrations due to C]O, -C-O, vibrations due to sp2 and sp3 bonding,and stretching vibrations due to hydroxyl and carboxyl groups, re-spectively [3,47–50]. According to the FTIR and XPS spectra, it can be

Fig. 7. H-res XPS spectra; (a) Native film; (b) O-1 film; (c) O-2 film; (d) O-3 film; (e) O-4 film; (f) Au4f spectrum for reference.

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concluded that oxygen in the films tends to be retained in the form ofdifferent bonding structures with increasing content in form of C–O–H,carbonyl, and carboxyl.

Fig. 9 shows the hardness (H) and elastic modulus (E) of the films.The control sample has a hardness of 20 GPa and elastic modulus of125 GPa. The former is less than that observed by Kahn et al. [34] butmore than that determined by Liu et al. and Zhiqiang et al. [32]. Thedifference in hardness is due to the processing parameters and ionsource. Kahn et al. operated the ion source at a higher discharge voltageand the energy (eV) per deposited carbon ion was higher, whereasZhiqiang et al. used a greater discharge current to promote local gra-phitization. Liu et al. applied high-voltage pulses for a short period oftime to enhance adhesion of the a-C:H films and introduced stress re-lieving. The hardness increased initially from 20 GPa to 28 GPa whenthe oxygen volume content in the gas mixture was 10%, which corre-sponds to the 7 at% oxygen in the film. These hardness values are on

higher sides compared to those in Ref. [21] and the difference arisesfrom the difference in the deposition technique. They used RF-PECVDto deposit the a-C:H:O films with CH4 as a precursor gas while ALIS isused to deposit the films via C2H2 gas in our case. Further increase inthe oxygen content decreases the H and E values. As shown in theRaman scattering and XPS spectra, the sp3 content increases for the filmdeposited with 10 vol% oxygen, i.e., 7 at% oxygen in film. In a-C:Hfilms, a larger sp3 fraction does not necessarily translate into higherhardness as the mechanical properties depend more on C–C bonding.The hardness increases only if C–C bonding increases instead of C–Hbonding. Here, C–C bonding increases due to the incorporation ofoxygen in two ways. Firstly, oxygen preferentially etches graphite-likesp2 clusters more quickly than sp3 and destroys the sp2 carbon networkmore efficiently thus leaving behind a network rich in sp3 hybridization[17,18,32]. Secondly, during exposure to the oxygen plasma, C]Cbreaks first leading to the formation of new bonds between carbon andcarbon and oxygen as well as bonds between oxygen and hydrogen.Adamopoulos et al. [19] observed reduced hydrogen contents afteraddition of oxygen. However, if the concentration of oxygen in the gasmixture is too large, C–C bonds begin to break and new carbonyl andcarboxyl bonds form leading to softening of the diamond-like matrix asindicated by XPS and FTIR [22]. Hence, the hardness starts to decreaseafterward and for sample O-4, i.e., 40 vol% oxygen in the gas mixtureand about 27% in the film, the hardness drops to 10.66 GPa due tohigher sp2 ordering and carboxyl groups.

Fig. 10a shows the friction characteristics of the HSS substrate anddeposited films. The friction coefficient of HSS starts at a small value of0.15 and increases abruptly after 5,000 cycles to 0.75. In comparison,the friction coefficient of the a-C:H coated HSS is smaller and settles at0.1. The friction behavior of the control and oxygenated a-C:H filmsexhibit the typical DLC coating characteristics with a fast transitionfrom boundary lubrication to the hydrodynamic zone. The frictioncoefficients of the oxygenated films are slightly less than that of thenative film. It is well known that the free carbon on the surface ofhydrogenated a-C films forms bonds with hydrogen resulting in lessdrag and adhesion between the free carbon and environment [51]. In

Fig. 8. FTIR spectra of the native, O-1 (10% oxygen by volume in the gas mixture), and O-3 samples (30% oxygen by volume in the gas mixture).

Fig. 9. Hardness (H) and elastic modulus (E) as a function of oxygen content inthe gas mixture.

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cases when oxygen is added in the films, the surface of the film containsthe hydroxyl group and the surface carbon atom contains bonds withoxygen and hydrogen causing surface passivation and less friction.Jongwannasiri et al. also observed the least friction offered by theoxygen plasma-treated a-C:H films in dry and humid environments andrelated it to surface passivation of carbon with oxygen and hydrogen[16]. Fig. 10b shows the wear rate of the control and oxygenated a-C:Hfilms as determined by Eq. (1). The wear rate of the control film againstthe SiN3 ball is 5× 10−1 mm3/m-N and decreases to 2.0 mm3/m-N forsample O-1. The wear rate increases with a further addition of oxygen(samples O-2 and O-3) but remains less than that of the control sample.The insets show the wear tracks on the samples revealing smootheninginside the tracks and graphitization at the near-surface of the trackunder a tribological load is considered the reason for this behavior. Thegraphitic sp2 clusters can slide upon each other acting as a self-lubricant[33]. Fig. 11d shows the Raman data of the tracks indicating larger ID/IG and more extensive graphitization. The trend is similar to that in-ferred from the hardness and wear rates. Fig. 11a shows wear tracks ofthe films after the wear test and it can be noted that by adding oxygen,sp3 C–C bonds and disorder are enhanced as indicated by Ramanscattering and XPS and responsible for the improved tribologicalproperties indicated by lower width and depth of wear track. However,too much oxygen leads to a decrease in disorder, increase in carbonyland carboxylic groups, and larger wear rates due to softening of thematrix. Fig. 11b and c show the 3-D optical micrographs of the weartracks of the control and O-1 film and the smoothening of the track andreduction in wear volume are observed. Based on these findings, it canbe inferred that addition of oxygen with the appropriate concentrationenhances the properties of the a-C:H films.

Another important property of a-C:H films is the adhesion strengthwith the underlying HSS substrate and a gradient CrC interlayer hasbeen reported to improve adhesion [32,33]. In this work, magnetronsputtering is employed to form the interlayer and ALIS is employed todeposit the overlying a-C:H layer. The backscattered electron image ofthe scratch on the control is shown in Fig. 12c. According to ASTMC1624-05, the critical load LC2 is defined as that when complete filmdelamination is observed. An average critical load (LC2) of 39 N is re-quired to completely detach the 1.7 μm thick film. In contrast, the a-C:H/CrCx film does not delaminate from the HSS substrate but instead,the a-C:H film detaches from the interlayer. The chromium EDS map inFig. 12a indicates that Cr is present at the track edge where delami-nation is initiated. The Raman results in Fig. 12b resemble those of theinterlayer indicating weak adhesion between the film and CrCx inter-layer. After addition of a small amount of oxygen (O-1), the scratchtrack is quite different as shown in Fig. 12d. The average critical load

Fig. 10. (a) Evolution of the friction coefficients of the different samples andHSS substrate; (b) Wear rate of the deposited films for different oxygen contentin the plasma with the insets showing the optical images of the wear tracks.

Fig. 11. (a) Wear tracks on samples with different oxygen contents; (b and c) 3-D optical images of the wear tracks on the control and sample deposited under O-1conditions; (d) Raman scattering spectra of the wear tracks indicating graphitization.

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(LC2) increases to more than 55 N thereby reflecting a drastic increase inthe interfacial strength between the a-C:H film and CrCx interlayer. TheEDS Cr maps at the same distance on the control sample for the first andsecond frames in Fig. 12e and f reveal the absence of CrCx and presenceof the CrCx layer in the respective frames. In this case, delamination isnot in the form of large-area spallation, but rather conformal crackinginside the track and buckling spallation. The observation discloses re-structuring of the amorphous carbon film after addition of oxygen. Thenew bonds with oxygen and modified bonds in the CrCx layer increasethe interfacial strength between the interlayer and amorphous carboncoating [52] leading to a strong interface between the CrCx layer and a-C:H:O film.

4. Conclusion

The effects of oxygen incorporation into a-C:H films deposited withan anode layer ion source on the structural, mechanical, and tribolo-gical properties are investigated. Films with 7 at%, 9 at%, 13 at%, and27% oxygen are fabricated by varying the oxygen concentrations in thegas mixture. The control film deposited without the intentional additionof oxygen contains 5% oxygen due to the non-UHV conditions in theinstrument and shows nanodots. The film containing 7 at% oxygenshows the maximum hardness of 29.0 GPa compared to 19.5 GPa of thecontrol as well as a smooth morphology with a significant reduction innanodots. The wear rate is reduced by 2.5 times to 2.0× 10−7 mm3/m-N and the friction coefficient by 10% to 0.09. However, increasing theoxygen concentration beyond 7 at% decreases the hardness and wearrate while the friction coefficient decreases slightly. Raman scatteringand XPS show that by adding a small amount of oxygen, the sp3 contentand C–C network increase, disorder increases, and the aromatic ringsdecrease. However, films with more than 7% oxygen become orderedand the carbonyl and carboxylic groups in the amorphous networksoften the matrix consequently degrading the mechanical and tribolo-gical performance. Our results reveal that an optimal amount of oxygenis important to the mechanical and tribological properties and provideinsights into the underlying mechanism.

Acknowledgments

The work was financially supported by Hong Kong Research GrantsCouncil (RGC) General Research Funds (GRF) No. CityU 11205617 aswell as Hong Kong Innovation and Technology Fund (ITF) No. ITS/452/17FP.

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