Characteristics of sputtered amorphous carbon films prepared by a closed-field unbalanced magnetron sputtering method

Characteristics of sputtered amorphous carbon films prepared by a closed-field unbalanced magnetron sputtering method

Journal of Non-Crystalline Solids 354 (2008) 5504–5508 Contents lists available at ScienceDirect Journal of Non-Crystalline Solids journal homepage:...

334KB Sizes 4 Downloads 97 Views

Journal of Non-Crystalline Solids 354 (2008) 5504–5508

Contents lists available at ScienceDirect

Journal of Non-Crystalline Solids journal homepage: www.elsevier.com/locate/jnoncrysol

Characteristics of sputtered amorphous carbon films prepared by a closed-field unbalanced magnetron sputtering method Yong Seob Park a, Byungyou Hong a,b,* a b

School of Information and Communication Engineering, Sungkyunkwan University, 300 Cheoncheon-dong, Jangan-gu, Suwon 440-746, Republic of Korea Center for Advanced Plasma Surface Technology (CAPST), Sungkyunkwan University, Republic of Korea

a r t i c l e

i n f o

Article history: Received 25 August 2007 Received in revised form 21 July 2008 Available online 18 September 2008 PACS: 87.64.Je 68.49.Uv 81.15.Cd 81.05.Uw 68.37.Og 62.20.Qp Keywords: Raman scattering UPS/XPS Sputtering Carbon STEM/TEM Hardness

a b s t r a c t We discuss the tribological performance of sputtered amorphous carbon (a-C) films deposited by closedfield unbalanced magnetron (CFUBM) sputtering with a graphite target using a mixture of helium (He) and argon (Ar) as sputtering gases. We investigated the effects of the graphite target power density on the micro-structural and physical properties. In the Raman spectra, the G-peak position moved to the higher wavenumbers. The ID/IG ratio increased with the increase of target power density in the fixed DC bias voltage. This was the result of the structural change in the a-C film that resulted with the increase in sp2 bonding fraction. Also, the maximum hardness of the a-C film was 23 GPa, the friction coefficient was 0.1, and the critical load was 25.9 N on the Si wafer. In addition, the compressive residual stress of the film increased a little with increasing target power density. Consequently, the various properties of aC films, with an increase of the target power density, were associated with the increase of cross-linked sp2 bonding fraction and the cluster size. The tribological properties of a-C film showed clear dependence on the energy of ion bombardment with the increase of plasma density during film growth. Ó 2008 Elsevier B.V. All rights reserved.

1. Introduction Amorphous carbon (a-C) and hydrogenated amorphous carbon (a-C:H) films are composed of a network of sp3 (diamond-like) and sp2 (graphite-like) co-ordinations. They possess excellent mechanical and tribological properties such as elevated hardness and excellent wear resistance [1]. In addition, they have low friction coefficients and provide protection for the counter parts [2]. Therefore, these films’ properties indicate that they have good prospects for use in a wide range of applications not only the typical mechanical applications such as in protective coating, wear resistant coating, corrosion resistant coating, and antireflective coating, but also in optical applications such as in photodiodes, light-emitting diodes, and electroluminescence devices [3,4]. We know that the properties of deposited a-C films depend strongly on the deposition conditions and elaboration methods. We can obtain a-C films with several techniques such as sputtering [2,3,11,14], ion beam deposition [10], CVD [13], and PLD. Sputtering methods are preferred for industrial applications because of * Corresponding author. Tel.: +82 31 290 7209; fax: +82 31 290 5669. E-mail address: [email protected] (B. Hong). 0022-3093/$ - see front matter Ó 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jnoncrysol.2008.08.007

these methods’ versatility, widespread use, and ease of scaling up. Also, these kinds of techniques allow for good control of the various deposition parameters such as the deposition rate, temperature, and gas composition during the deposition process. In this paper, we report on the effects of target power density on the tribological and structural properties of a-C films synthesized by closed-field unbalanced magnetron (CFUBM) sputtering. 2. Experimental The a-C films were deposited onto 3 cm  3 cm p-type (1 0 0) silicon and glass substrates by using a closed-field unbalanced magnetron (CFUBM) sputtering system consisting of two targets with 99.99% pure graphite and a diameter of 10.16 cm. The distance from target to substrate was 60 mm. Also, high purity argon (99.99%) and helium (99.99%) were used as the sputtering gases for the growth of the a-C films. The silicon and glass substrates were cleaned successively for 5 min in acetone, methanol, and D.I. water. Then, the silicon substrates were etched in a hydrofluoric acid solution to strip off any native oxide. For the deposition process, the background pressure of the process chamber was evacuated to below 4.3  10 4 Pa using diffusion pumps and then the gases

5505

Y.S. Park, B. Hong / Journal of Non-Crystalline Solids 354 (2008) 5504–5508 Table 1 Deposition conditions of a-C films prepared by CFUBM sputtering Deposition parameters Base pressure (Pa) Conditions (value)

4.3  10

4

Ar/He flow rate (sccm)

Ar/He pressure (PAr/PHe) (Pa)

Total pressure (Pa)

14.2/2.5

0.36/0.04

0.4

of Ar and He (14.2:2.5 sccm) were introduced into the chamber and were mixed. The total working pressure for the entire deposition run was 0.4 Pa. We obtained the a-C films at various target power densities from 1.2 to 2.0 W/cm2 below the fixed negative DC bias voltage. The thicknesses of the deposited all films were approximately 200 nm. We summarize the experimental parameters in Table 1. The thickness of the a-C films was measured using FE-SEM [Jeol, JSM6700F] and the internal structure of the films was characterized by Raman spectrometry [Jasco, MRS-300], X-ray photoelectron spectroscopy [VG MICROTECH, ESCA-2000], and high resolution transmission electron microscopy (HRTEM) [JEOL, JEOL 300 kV]. The surface morphology of the films was observed using atomic force microscopy [Seiko, SPA-400] and the value of the rootmean-square (RMS) roughness was obtained over an image area of 2 lm2. The residual stress of a-C films was measured using stress tester [J&L Tech, JLCST022]. And also, the hardness and elastic modulus were measured using a commercial nano-indentation instrument [Nano-indenter XP] having a continuous stiffness method (CSM) option. Continuous loading–unloading indentations were applied up to a maximum load of 30 mN. Also, the friction

Substrate bias voltage (V) 200

Target power density (kW/ cm2) 1.2, 1.6, 2.0

coefficients of the a-C films with an increase of the target power density were analyzed using a ball-on disk (BOD) tribometer in normal dry ambient air against a polished AISI 52100 steel ball with a diameter of 4.72 mm. The sliding speed of the steel ball was maintained at a constant value of 60 rpm. Also the adhesion value of a-C films was measured using a nanoscratch tester [J&L Tech. JLST022]. A diamond tip of nominal radius was used to scratch the film surfaces and the normal load was increased to 35 N. Also, the scratch speed v and scratch distance were kept to values of 0.2 mm/s and 10 mm. 3. Results Raman spectroscopy is an effective method for the characterization of carbon materials [1,6]. Raman spectra of the a-C films deposited at various target power densities are shown in Fig. 1(a). The ID/IG ratio and the G-peaks position, which were deconvoluted into two Gaussian curve fits, in order to obtain quantitative information on the sp3 content in the film, are shown in Fig. 1(b) for the various target power densities. From these figures, we observed that the G- and D-bands, corresponding to the gra-

a Intensity [arb. units]

Intensity [arb. units]

a

2

2.0 kW/cm

2

1.6 kW/cm 1375 cm

-1

2

2.0 kW/cm 2 1.6 kW/cm

2

1.2 kW/cm 1562 cm 600

900

1200

1500

2

1.2 kW/cm

-1

1800

280

2100

b

b

5.7

1575

4.8 4.5

ID/IG ratio

4.2 3.9

1.6

2.0

3.6

2

Target power density [kW/cm ] Fig. 1. Raman spectra (a) and the variation of the G-peak position and ID/IG ratio (b) of a-C films prepared at various target power densities.

2

G peak position

1.2

292

296

0.56

0.54 0.53 0.52

3

5.1

1565

1555

288

0.55

sp /sp bonding ratio

1570

ID / IG ratio

G peak position [cm-1]

5.4

1560

284

Binding energy [eV]

Wavenumbers [cm-1]

0.51 0.50

1.2

1.6

2.0 2

Target power density [kW/cm ] Fig. 2. Carbon 1s XPS spectra (a) and the variation of sp3/sp2 bonding ratio of a-C films prepared at various target power densities (b).

5506

Y.S. Park, B. Hong / Journal of Non-Crystalline Solids 354 (2008) 5504–5508

phitic band and disordered band, were situated at approximately 1562 cm 1 and 1375 cm 1 [5,7], respectively. Also, we found that the positions of the G-peaks moved to lower wavenumbers and the ID/IG ratio gradually increased with the increase of target power density. Fig. 2 shows the carbon 1s XPS spectra and the variation of sp3/ 2 sp bonding ratio of a-C films prepared at various target power densities. From the Gaussian fitting results of the XPS spectra, which were deconvoluted into two components with binding energies of 284.4 eV and 286.8 eV corresponding to sp2 C–C and/or C–H and sp3 C–C and/or C–H [8,9], the position of the C 1s peak with the increase of target power density moved to a lower binding energy related to the polymerization of the film [8,10]. Also, the sp3/sp2 bonding ratio with various target power densities showed low values below 0.55 and decreased with an increase of the target power density. We performed HRTEM analyses for a detailed investigation of film structure. Fig. 3 shows the HRTEM micrographs of a-C films. From the figures, the microstructure of a-C film prepared at the target power density of 1.2 kW/cm2 was found to be a nanocrystalline graphite structure within an amorphous carbon matrix. And also, the a-C film prepared at 2.0 kW/cm2 exhibited the dispersion of graphite sp2 bonding clusters of approximately 5 nm size and had an interplanar spacing of (1 1 1) and (0 1 0) planes. Also, we could see that with increasing target power density, the number of clusters in the a-C film prepared at 2.0 kW/cm2 was higher than the number of clusters in the film prepared at 1.2 kW/cm2.

Fig. 4 shows the growth rate, surface roughness, and friction coefficient of a-C films prepared with various target power densities at the fixed negative DC bias voltage. From the figure, this result clearly shows that the growth rate increased significantly and the surface roughness decreased with the increase of target power density. The increase of the growth rate can be explained by the enhancement of the reactive ions with the increase of sputtered carbon flux as a result of the increase of target power density [13]. Also, the decrease of the surface roughness is related to the increase of the energetic ions bombardment at the film surface. And, the friction coefficients of a-C films prepared with various target power densities are shown in Fig. 4(c). It was clear that most of the a-C films showed a low friction coefficient below 0.13, which varied very little with target power density. The hardness, elastic modulus, residual stress, and critical load of the a-C films prepared at various target power densities are shown in Fig. 5. We can observe in this figure that the hardness and elastic modulus increased with the increase of target power density and the maximum values exhibited about 23 GPa and

a

b

c Growth rate [nm/min.]

110

Linear fit 100 90 80 70 60

Coefficient of friction

Rms roughness [nm]

50 2.4 2.2 2.0 1.8 1.6 1.4 0.14 0.13 0.12 0.11 0.10 0.09 1.2

1.6

2.0 2

Target power density[ kW/cm ] Fig. 3. HRTEM micrographs of nano-structured a-C films deposited at a target power density of 1.2 kW/cm2 (a) and 2.0 kW/cm2 (b).

Fig. 4. AFM topography images of a-C films deposited at a target power density of 1.2 kW/cm2 (a) and 2.0 kW/cm2 (b) and the growth rate, rms roughness, and the friction coefficient (c) of the a-C films prepared at various target power densities.

Y.S. Park, B. Hong / Journal of Non-Crystalline Solids 354 (2008) 5504–5508

26

220

Hardness [GPa]

Hardness

200

22 180

20

Elastic modulus

18

160

16

140

14 120

12

100

2.5 10

26 2.0 24

Residual stress

1.5

1.0

Critical load

1.2

22

1.6

2.0

Critical load [N]

Residual stress [GPa]

Elastic modulus [GPa]

24

20

2

Target power density [kW/cm ] Fig. 5. The hardness, elastic modulus, residual stress, and critical scratch load of a-C films deposited with various target power densities.

198 GPa. However, the residual stress of a-C films increased a little. Also, we can see that the adhesion between a-C film and substrate increased with the increase of target power density. 4. Discussion The variation of G-peak position and ID/IG ratio depends on the sp2 bonding fraction in the film. From the Raman and XPS results, the increase of target power density is attributed to the increase of the sp2 bonding fraction in the carbon film. These indicate that the increase of target power density leads to the enhancement of the sputtered carbon flux from the graphite target in the plasma. Then, the increased carbon flux attributed to the increase of ion bombardment at the surface by an applied negative DC bias voltage [9,11,12] and led to the rising of the surface temperature by the frequent bombardments with increasing target power density at the film surface. Consequently, this behavior can be explained by the increase of the sp2 bonding fraction in the film as a result of the increase of the growing temperature with increasing target power density, and also the structural variation of a-C films associated closely with the density of the sputtered carbon ion flux and the applied negative DC bias voltage. Also, from the results of HRTEM analysis, we confirmed the existence of the clusters in the carbon. As seen in the pictures shown fig. 3, the increase of target power density contributed to the increase of cross-linked sp2 clusters in the film. These indicate that the enhancement of ion bombardment caused by energetic ions with increasing carbon flux leads to the rising of the surface temperature at the film surface and contributes to the increase of the number of cross-linked sp2 bonding clusters in the carbon networks [10–12]. In the result, the increase of sp2 bonding clusters in a-C film can mainly be explained by the temperature rising at the film surface caused by the increase of energetic carbon flux. This is in agreement with previous observations. As seen in fig. 4, the smooth surface of a-C film at the high target power density is associated with the effect of ion bombardment as a result of the increase of sputtered carbon flux. Generally, we

5507

know that the coefficient of friction depends on the variation of the surface morphology and the film density caused by the chemical bonding of carbon and hydrogen [14]. Indications are that the low value of the friction coefficient results from the improvement of ion bombardments. This behavior is caused by the increase of the energetic ions at the surface as a result of the increase of the sputtered carbon flux with the increase of target power. Consequently, the increase of target power density and the applied negative DC bias voltage can be attributed to the low friction coefficient and the smooth surface due to the formation of graphite sp2 clusters [11,16]. We confirmed the clusters in a-C films generated with increasing target power density in fig. 3. Also, we have reported the physical changes in fig. 5. In the result, the increase of target power density leads to the increase of the hardness, elastic modulus, and critical load [13,17]. However, this caused the increase of the residual stress. These results indicate that the ion bombardment and the rising of surface temperature lead to the promotion of the formation of cross-liked sp2 carbon bond and the disorder degree of carbon network as a result of the improvement of carbon flux with the increase of target power density and the addition of the energy to ions by supplying negative DC bias voltage during the film deposition [11,15,16]. Also, these lead to the increase of the residual stress due to the increase of the dense degree of the films. Consequently, the cross-linked sp2 bonding clusters generated by the target power density are attributed to the improvement of the physical properties of a-C films. In this work, this ‘sp2 clusters’ can occur due to two main causes: the increase of ion bombardment by the sputtered carbon flux and the relaxation process after the implantation of high energetic ions during the film growth. 5. Conclusion We investigated the effect of target power density on tribological and structural properties the a-C films deposited by a magnetron sputtering technique. From the structural results, we confirmed the existence of the sp2 bonding clusters in carbon networks and knew the promotion of the formation of sp2 bonding clusters by the increase of target power density and the applied negative DC bias voltage. Consequently, the increase of the sputtered carbon flux with the increase of target power density leads to the improvement of the tribological properties of a-C film due to the increase of dispersed cross-linked sp2 clusters. Specially, our CFUBM sputtering method can prepare the a-C films exhibited the excellent physical performance. Acknowledgment The authors are grateful for the financial support provided by Grant No. R11-2000-086-0000-0 and No. R01-2008-000-10690-0 from the Center of Excellency program of the Korea Science and Engineering Foundation and by MOST through the Center for Advanced Plasma Surface Technology (CAPST) at Sungkyunkwan University. References [1] J. Robertson, Mater. Sci. Eng. R 37 (2002) 129. [2] V. Rigato, G. Maggioni, D. Boscarino, G. Mariotto, E. Bontempi, A.H.S. Jones, D. Camino, D. Tear, Surf. Coat. Technol. 116–119 (1999) 580. [3] M. Cremona, R. Reyes, C.A. Achete, R. Tarora Britto, S.S. Camargo, Thin Solid Films 447&448 (2004) 74–79. [4] E. Staryga, G.W. Bals, Diamond Relat. Mater. 14 (2005) 23. [5] D.J. Choi, W.H. Koo, S.M. Jeong, D.W. Han, D.Y. Lee, H.K. Baik, S.W. Jang, S.M. Lee, Vacuum 72 (2004) 445. [6] A.C. Ferrari, J. Robertson, Phys. Rev. B 61 (20) (2000) 14095. [7] S. Zhang, X.L. Bui, Y. Fu, Surf. Coat. Technol. 167 (2003) 137.

5508 [8] [9] [10] [11] [12]

Y.S. Park, B. Hong / Journal of Non-Crystalline Solids 354 (2008) 5504–5508

C. Charitidis, S. Logothetidis, P. Douka, Diamond Relat. Mater. 8 (1999) 558. Y.Y. Chang, D.Y. Wang, C.H. Chang, W. Wu, Surf. Coat. Technol. 184 (2004) 349. M. Silinskas, A. Grigonis, Diamond Relat. Mater. 11 (2002) 1026. D.Y. Wang, D.L. Chang, W.Y. Ho, Thin Solid Films 355&356 (1999) 246. B. Yang, Z.H. Huang, C.S. Liu, Z.Y. Zeng, X.J. Fan, D.J. Fu, Surf. Coat. Technol. 200 (2006) 5812.

[13] X.L. Peng, Z.H. Barber, T.W. Clyne, Surf. Coat. Technol. 138 (2001) 23. [14] P.W. Shum, Z.F. Zhon, K.Y. Li, Wear 256 (2004) 362. [15] V. Kulikovsky, P. Bohac, F. Franc, A. Deineka, V. Vorlicek, L. Jastrabik, Diamond Relat. Mater. 10 (2001) 1076. [16] Y.Y. Chang, D.Y. Wang, Surf. Coat. Technol. 200 (2006) 3170. [17] F. Svaln, A. Kassman-Rucolphi, E. Wallen, Wear 254 (2003) 1092.