- Email: [email protected]
A comparative study of Si–C–N films on different substrates grown by RF magnetron sputtering
Journal of Alloys and Compounds 478 (2009) 474–478
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
A comparative study of Si–C–N films on different substrates grown by RF magnetron sputtering A.S. Bhattacharyya a,b,∗ , S.K. Mishra a , S. Mukherjee b , G.C. Das b a b
National Metallurgical Laboratory, Jamshedpur 831007, India Jadavpur University, Kolkata 700032, India
a r t i c l e
i n f o
Article history: Received 23 July 2008 Received in revised form 10 November 2008 Accepted 16 November 2008 Available online 3 December 2008 Keywords: Si–C–N RF sputtering Substrate effect
a b s t r a c t Si–C–N nanocomposite thin films were deposited on industrially important substrates like silicon (1 0 0), borosilicate glass, and stainless steel (304SS) by radio frequency (RF) magnetron sputtering. The microstructural characterization was carried out by transmission electron microscopy (TEM) showing localized -C3 N4 in amorphous Si–C–N matrix, which was confirmed by X-ray photoelectron spectroscopy (XPS) and Raman spectroscopy. The thermal mismatch occurring between the substrate and the coating resulted in variation in deposition rate, roughness and other mechanical properties like hardness and adhesion for the three different substrates. Both microindentation and nanoindentation were performed to estimate the hardness of the coatings. Scratch tests were used for the adhesion studies. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Si–C–N has been a very important nanocomposite material showing promising combination of properties [1] Si–C–N nanocomposite coatings exhibit improved properties compared to conventional coatings in terms of better thermal conductivity, thermal stability, hardness, tunable band gap, chemical inertness, wetting behaviour and wear resistance [2–6]. Due to these unusual combinations of properties, they have a large range of applications, e.g. wear resistant coatings for automotive industry, microelectronic mechanical system (MEMS) device fabrication, high temperature semiconducting and optoelectronic devices [7–10]. Si–C–N also finds application in biological purification step by producing pure gDNA [11]. Nanocomposites consisting of precursor-derived Si–C–N ceramics incorporated with carbon nanotubes (CNTs) were successfully prepared by casting a mixture of CNTs and a liquid precursor polymer followed by cross-linking and thermolysis [12]. The high temperature behaviour of Si–C–N ceramics has been thermodynamically calculated using CALPHAD software [1]. A study of the relationship between the chemical and structural properties with terminological properties of sputtered Si–C–N films have been carried out, where the hardness of the film
∗ Corresponding author at: NML, Jamshedpur, India. Tel.: +91 657 2271709 14; fax: +91 657 2270527. E-mail address: [email protected] (A.S. Bhattacharyya). 0925-8388/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2008.11.105
has been found to vary depending upon the position in the Si–C–N phase diagram [5]. Several methods for the fabrication of amorphous and crystalline Si–C–N films are reported in the literature. Both crystalline and amorphous or nanostructured Si–C–N compounds have been prepared. They are produced in bulk form by pyrolysis of organometallic polysilazane precursors, where polymer to ceramic transformation takes place [11,13,14]. The polymer impregnation and pyrolysis (PIP) process was used to prepare a mullite interphase of C/Si–C–N composites in order to provide an acceptable oxidation protection of these composites [15]. Thin film depositions of Si–C–N have been carried out by plasma and ion assisted deposition, chemical vapour deposition, magnetron sputtering, microwave and electron cyclotron resonance plasma enhanced chemical vapour (ECRPECVD), ion implantation, pulsed laser deposition and Rapid thermal chemical vapour deposition (RTCVD) [2–10]. At substrate temperatures below 1000 ◦ C, amorphous Si–C–N films are reported to be deposited, while higher temperatures produced crystalline composite films of ␣- and Si3 N4 and ␣- and -SiC [6]. Microhardness increase and promising field emission properties were obtained from CGed films in comparison with monolithic SiC and SiNx films deposited by PECVD [16]. Si–C–N thin films with tailored stoichiometries on the tie line SiC–Si3 N4 have been produced by several fold ion implantation and by a combination of RF magnetron sputtering and ion implantation [17]. These high purity thin films have been heat treated at 1250 ◦ C under high vacuum conditions using an electron beam annealing system to enable crystallization and/or phase formation. The formation of an amorphous network of mixed Si(C,N)4 tetrahedrons
A.S. Bhattacharyya et al. / Journal of Alloys and Compounds 478 (2009) 474–478
475
Fig. 1. (a) TEM studies of Si–C–N films, (b) corresponding SAED pattern and (c) C 1s peaks of XPS spectra confirming formation of -C3 N4 .
for the ternary compositions and the crystallization of the binary SiC and Si3 N4 thin films were obtained. Although there is a lot of literature on Si–C–N, a proper investigation of the effect of different substrates in the microstructure and corresponding mechanical properties of the coatings is lacking. In this communication, we have used magnetron sputtering (RF) techniques to deposit coatings on silicon (1 0 0), borosilicate glass and stainless steel (304SS) substrates. A correlation among thickness, hardness, modulus, microstructure and elemental composition with the nature of the substrate is presented. 2. Experimental procedures Si–C–N coatings were deposited on Si (1 0 0), 304SS and borosilicate glass substrates by RF magnetron sputtering (HHV, Bangalore, India) using single target SiC. Silicon and carbon powder was used in 1:1 ratio and was compacted into a 50 mm disc with 3 mm thickness and sintered at graphite furnace at 1900 ◦ C in argon atmosphere. For deposition, the vacuum chamber was first evacuated to 1 × 10−6 mbar pressure. Thin films were deposited on Si (1 0 0) substrate at 1 × 10−2 mbar pressures in Ar/N2 atmosphere in 1:9 ratio. The reason for maintaining this gas ratio was deposition of hardest film as reported in our previous publications. The temperature was kept constant at 500 ◦ C and the power was maintained at 400 W. The deposition was done for 3 h. Thickness of the coatings was measured by surface profilometer (Talysurf, Taylor Hobson, UK). AFM (Seiko Scanning probe, 400, Japan) was used for roughness measurements. Vickers microhardness of the coatings was measured at 15 gf load by Microhardness Tester (The LEICA VMHT AUTO, Austria). Nanoindentation tests were carried out in continuous stiffness mode using XP-Nanoindenter MTS, USA. The load–displacement curves, determined hardness and modulus. Microstructural investigation was done by transmission electron microscope (TEM) of Phillips, EM200, Netherlands. Almega Dispersive Raman Spectroscope having He–Ne laser (532 nm) was used for the Raman studies. It consisted of 25 m pin hole. Olympus microscope was used to select the location and 25% of laser power was used for the spectroscopy. The XPS analysis was carried out using Thermo VG Scientific Multitech 2000 XPS system, UK.
3. Results and discussion TEM studies showed the formation of -C3 N4 crystallites in the amorphous matrix of Si–C–N (Fig. 1a). The ring formation in the SAED pattern indicated the formation of polycrystalline phases (Fig. 1b). The matched d-values and corresponding planes are shown in the figure. The crystallization process was localized. Scanning other area mainly showed amorphous (a-Si–C–N) phase. The formation of -C3 N4 was also confirmed by XPS studies. The C 1s spectra showed peaks at 285.11 eV due to C–C, 290.32 eV due to C–N and 292.57 eV due to C N (Fig. 1(c)) [18–20]. The Raman spectra for all three films are exhibited in Fig. 2(a). In the range 1000–1800 cm−1 the bands were deconvoluted using 3 Gaussian fits as shown in Fig. 2(b) for silicon substrates. The band with peak value at 1529 cm−1 called the G (graphitic) band, arises due to stretching of sp2 , hybridised carbon [19–23]. For purely graphitic phase containing 100% sp2 hybridised carbon the G band peak occurs at 1580 cm−1 . The shift to lower wavenumber in this case is the indication of increase in sp3 carbon percentage in the coating which is due to incorporation of nitrogen in the coating forming mainly amorphous silicon carbonitride (a-Si–C–N) with very small -C3 N4 crystallites as observed in the TEM studies. The band with peak value at 1311 cm−1 is called the D (disorder) band, which arises due to small crystallite sizes, nitrogen incorporation and sp3 hybridised carbon [19]. A band peaked at 1153 cm−1 called the T band was also observed due to stretching modes of sp3 hybridised carbon [7]. The Raman spectra of a bare silicon substrate have been shown for comparison. The first and second order peaks observed at 526 cm−1 and 981 cm−1 , respectively, are typical of pure silicon. The adatoms generated during sputtering when falling on the substrate have to thermally equilibrate themselves and undergo
476
A.S. Bhattacharyya et al. / Journal of Alloys and Compounds 478 (2009) 474–478
coalescence. Compressive stress in the coatings aids in better deposition whereas tensile stress is a detrimental factor. Deposition at a higher substrate temperature increase the compressive stresses, and hence leads to better adherence. The extent of tensile stress is on the other hand is proportional to the difference in CTE values of the coating and the substrate. 304SS has a CTE value of 11 whereas it is 3.0 for Si (1 0 0) and 8.5 for borosilicate glass. Thus the best adhesion is obtained for silicon substrates as Si–C–N which has a CTE value of 0.5. The deposition rate (r) and average roughness (Ra) of the coatings varied due to CTE mismatch; however the variation in ´˚ deposition rate was negligible. Highest deposition rate of 175 A/min ´ ˚ was observed for silicon substrates and the lowest (162 A/min) was
Fig. 2. (a) Raman spectra of Si–C–N films on silicon, glass and steel substrates and (b) their 3 Gaussian fir deconvolutions for silicon substrate and (c) Raman spectra of bare silicon substrate.
for 304SS substrates. Therefore the lowest deposition time was used for coating silicon substrates giving a thickness of 1 m. Roughness was found to be the lowest in case of silicon (20 nm) but interestingly found to be the highest in case of glass (120 nm) instead of 304SS (50 nm), which is due to the presence of an oxide layer usually found in the top surface of a borosilicate glass. Adhesion tests were performed through scratch tests. Fig. 3 shows the variation of coefficient of friction and normal load with stroke length obtained during scratch tests. The change in slope occurring in the C.O.F curve is the indication of coating failure. In the case of silicon substrates (Fig. 3(a) and (b)), cohesive failure was observed initially where the coating starts to penetrate the coating. The adhesive film failure occurred at a critical load of 10 N where brittle failure and chipping of the coating was observed. For glass substrates (Fig. 3(c) and (d)) the cohesive and adhesive failure regions were not separately resolvable. The slope change in the C.O.F curve occurred at a corresponding critical load of 5 N and brittle failure by buckling was observed at higher loads. Lastly in case of steel substrates (Fig. 3(e) and (f)) all the three mechanism of cohesive failure, adhesive failure and complete coating removal through spallation was observed [24]. The critical load in this case was 7 N. Vickers microhardness tests performed at 15 gf showed Vickers hardness number (VHN) of 1434 for Si, 1560 for glass and 2123 for steel substrates. In the nanoindentation tests a sharp indenter (Berkovich) is used to test the properties of a thin coating and to ensure that the contact stresses are high enough to cause yield of the coating before the substrate. The initial contact will always be elastic, dominated by the spherical tip of the indenter. As the load is increased the sloping sides of the indenter will come into contact and the indenter behaves like a truncated cone. At this point the plastic deformation is initiated. This transition is associated with the contact-induced stresses exceeding the yield stress of the material being tested and full plastic deformation occurs beneath the indenter. The nanohardness, at this point is estimated from the ratio of maximum indentation load and projected contact area of the plastically deformed region (Fig. 4) [25,26]. A nanohardness of 17 GPa was found to be the highest in case of 304SS substrates. The lowest hardness of 10 GPa was found in case of silicon substrates and coatings deposited on glass substrates showed a hardness of 14 GPa. Table 1 summarises all the parameters obtained on all the three substrates. The lowest hardness obtained in case of silicon substrates inspite of having the lowest CTE mismatch and the highest deposition rate and the
Table 1 Variation of critical load, microhardness and nanohardness of Si–C–N films deposited on Silicon, glass and steel substrates. Substrate
r (nm/min)
Time (min)
t (m)
Ra (nm)
Lc (N)
VHN (15 gf)
H (GPa)
Silicon (100) Glass 304SS
16.6 12.5 10.8
60 120 240
1.0 1.5 2.6
20 120 50
10 5 7
1434 1560 2123
10 14 17
A.S. Bhattacharyya et al. / Journal of Alloys and Compounds 478 (2009) 474–478
477
Fig. 3. Scratch tests performed on coatings deposited on (a, b) silicon, (c, d) glass and (e, f) steel substrates.
best adhesion is due to the lowest thickness, which also explains the reason for getting the highest hardness in case of 304SS substrates. 4. Conclusions
Fig. 4. Nanohardness of Si–C–N coatings on different substrates.
A systematic comparison of RF magnetron sputtered Si–C–N coatings on Si (1 0 0), 304SS, and borosilicate glass was carried out. Evidence of nucleation of -C3 N4 having both sp2 and sp3 hybridised carbon was obtained from microscopic and spectroscopic studies. The substrate effect was mainly observed in hardness, modulus and adhesion properties of the films. Better deposition and adhesion was seen for silicon substrates due to less difference in CTE values with Si–C–N. Coating thickness was found to have a major influence on hardness, as silicon substrates although having the lowest CTE mismatch, gave the lowest hardness due to lowest thickness. The highest nanohardness of 17 GPa and microhardness of 2123 HV0.015 with a critical load of 7 N was obtained in case of coatings deposited on 304SS substrate.
478
A.S. Bhattacharyya et al. / Journal of Alloys and Compounds 478 (2009) 474–478
Acknowledgements The authors wish to thank Dr. M. Ghosh and Dr. D.D.N. Singh of NML, Jamshedpur, for their cooperation in carrying out TEM and Raman studies. This research work is sponsored by Department of Science and Technology, Government of India. References [1] H.J. Seifert, J. Peng, H.L. Lukas, F. Aldinger, J. Alloys Compd. 320 (2001) 251–261. [2] I.B. Lezak, A.M. Wrobel, M.P.M. Kivitorma, I.J. Vayrynen, A. Tracz, Appl. Surf. Sci. 253 (2007) 7211–7218. [3] I.B. Lezak, A.M. Wrobel, D.M. Bielinski, Diam. Relat. Mater. 15 (2006) 1650–1658. [4] E. Tomasella, F. Rebib, M. Dubois, J. Cellier, M. Jacquet, J. Phys: Conf. Ser. 100 (2008) 1–4. [5] H. Hoche, D. Allebrandt, M. Bruns, R. Riedel, C. Fasel, Surf. Coat. Technol. 202 (2008) 5567–5571. [6] N.J. Wagner, W.W. Gerberich, J.V.R. Heberlein, Surf. Coat. Technol. 201 (2006) 4168–4173. [7] K.B. Sundaram, Z. Alizadeh, R.M. Todi, V.H. Desai, Mater. Sci. Eng. A 368 (2004) 103–108. [8] W. Cheng, J. Jiang, Y. Zhang, H. Zhu, D. Shen, Mater. Chem. Phys. 85 (2004) 370–376. [9] J. Vleck, M. Kormunda, J. Cizek, V. Perina, J. Zemek, Surf. Coat. Technol. 160 (2002) 74–79.
[10] W.R. Chang, Y.K. Fang, S.F. Ting, Y.S. Tsair, C.N. Chang, C.Y. Lin, S.F. Chen, IEEE Electron. Dev. Lett. 24 (2003) 565–567. [11] C. Roh, S.H. Lee, V. Francois, J. Chromatogr. A 1179 (2008) 145–151. [12] Y. Katsuda, P. Gerstel, J. Narayanan, J. Bill, F. Aldinger, J. Eur. Ceram. Soc. 26 (2006) 3399–3405. [13] N. Janakiraman, F. Aldinger, J. Eur. Ceram. Soc. 29 (2009) 163–173. [14] J. Wan, P.R.L. Malenfant, S.T. Taylor, S.M. Loureiro, M. Manoharan, Mater. Sci. Eng. A 463 (2007) 78–88. [15] G. Lu, S. Qiao, C. Zhang, J. Hou, D. Jia, Y. Zhang, Comp. Part A: Appl. Sci. Manuf. 39 (2008) 1467–1470. [16] Y.S. Li, H. Kiyono, S. Shimada, X. Lu, A. Hirose, Diam. Relat. Mater. 15 (2006) 1727–1731. [17] M. Rudolphi, M. Bruns, H. Baumann, U. Geckle, Diam. Relat. Mater. 16 (2007) 1273–1277. [18] M.J. Bojdys, J.O. Muller, M. Antonietti, A. Thomas, Chem. Eur. J. (2008) 8177–8182. [19] A. Champi, F.C. Marques, Braz. J. Phys. 36 (2006) 462–465. [20] G. Peng, X. Jun, P. Yong, D. Wanyu, D. Xinlu, D. Chuang, Plasma Sci. Technol. 8 (2006) 425–428. [21] N.I. Fainer, Y.M. Rumyantsev, M.L. Kosinova, Chem. Sust. Dev. 9 (2001) 31–36. [22] N.I. Fainer, Y.M. Rumyantsev, M.L. Kosinova, Chem. Sust. Dev. 9 (2001) 37–43. [23] E.J. Liang, J.W. Zhang, J. Leme, C. Moura, L. Cunha, Thin Solid Films 469 (2004) 410–415. [24] S.K. Mishra, A.S. Bhattacharyya, Mater. Lett. 62 (2008) 398–402. [25] R.L. Zong, S.P. Wen, F. Zeng, Y. Gao, F. Pan, J. Alloys Compd. 464 (2008) 544– 549. [26] D. Drozdz, T. Kulik, H.J. Fecht, J. Alloys Compd. 441 (2007) 62–65.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
A comparative study of Si–C–N films on different substrates grown by RF magnetron sputtering A.S. Bhattacharyya a,b,∗ , S.K. Mishra a , S. Mukherjee b , G.C. Das b a b
National Metallurgical Laboratory, Jamshedpur 831007, India Jadavpur University, Kolkata 700032, India
a r t i c l e
i n f o
Article history: Received 23 July 2008 Received in revised form 10 November 2008 Accepted 16 November 2008 Available online 3 December 2008 Keywords: Si–C–N RF sputtering Substrate effect
a b s t r a c t Si–C–N nanocomposite thin films were deposited on industrially important substrates like silicon (1 0 0), borosilicate glass, and stainless steel (304SS) by radio frequency (RF) magnetron sputtering. The microstructural characterization was carried out by transmission electron microscopy (TEM) showing localized -C3 N4 in amorphous Si–C–N matrix, which was confirmed by X-ray photoelectron spectroscopy (XPS) and Raman spectroscopy. The thermal mismatch occurring between the substrate and the coating resulted in variation in deposition rate, roughness and other mechanical properties like hardness and adhesion for the three different substrates. Both microindentation and nanoindentation were performed to estimate the hardness of the coatings. Scratch tests were used for the adhesion studies. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Si–C–N has been a very important nanocomposite material showing promising combination of properties [1] Si–C–N nanocomposite coatings exhibit improved properties compared to conventional coatings in terms of better thermal conductivity, thermal stability, hardness, tunable band gap, chemical inertness, wetting behaviour and wear resistance [2–6]. Due to these unusual combinations of properties, they have a large range of applications, e.g. wear resistant coatings for automotive industry, microelectronic mechanical system (MEMS) device fabrication, high temperature semiconducting and optoelectronic devices [7–10]. Si–C–N also finds application in biological purification step by producing pure gDNA [11]. Nanocomposites consisting of precursor-derived Si–C–N ceramics incorporated with carbon nanotubes (CNTs) were successfully prepared by casting a mixture of CNTs and a liquid precursor polymer followed by cross-linking and thermolysis [12]. The high temperature behaviour of Si–C–N ceramics has been thermodynamically calculated using CALPHAD software [1]. A study of the relationship between the chemical and structural properties with terminological properties of sputtered Si–C–N films have been carried out, where the hardness of the film
∗ Corresponding author at: NML, Jamshedpur, India. Tel.: +91 657 2271709 14; fax: +91 657 2270527. E-mail address: [email protected] (A.S. Bhattacharyya). 0925-8388/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2008.11.105
has been found to vary depending upon the position in the Si–C–N phase diagram [5]. Several methods for the fabrication of amorphous and crystalline Si–C–N films are reported in the literature. Both crystalline and amorphous or nanostructured Si–C–N compounds have been prepared. They are produced in bulk form by pyrolysis of organometallic polysilazane precursors, where polymer to ceramic transformation takes place [11,13,14]. The polymer impregnation and pyrolysis (PIP) process was used to prepare a mullite interphase of C/Si–C–N composites in order to provide an acceptable oxidation protection of these composites [15]. Thin film depositions of Si–C–N have been carried out by plasma and ion assisted deposition, chemical vapour deposition, magnetron sputtering, microwave and electron cyclotron resonance plasma enhanced chemical vapour (ECRPECVD), ion implantation, pulsed laser deposition and Rapid thermal chemical vapour deposition (RTCVD) [2–10]. At substrate temperatures below 1000 ◦ C, amorphous Si–C–N films are reported to be deposited, while higher temperatures produced crystalline composite films of ␣- and Si3 N4 and ␣- and -SiC [6]. Microhardness increase and promising field emission properties were obtained from CGed films in comparison with monolithic SiC and SiNx films deposited by PECVD [16]. Si–C–N thin films with tailored stoichiometries on the tie line SiC–Si3 N4 have been produced by several fold ion implantation and by a combination of RF magnetron sputtering and ion implantation [17]. These high purity thin films have been heat treated at 1250 ◦ C under high vacuum conditions using an electron beam annealing system to enable crystallization and/or phase formation. The formation of an amorphous network of mixed Si(C,N)4 tetrahedrons
A.S. Bhattacharyya et al. / Journal of Alloys and Compounds 478 (2009) 474–478
475
Fig. 1. (a) TEM studies of Si–C–N films, (b) corresponding SAED pattern and (c) C 1s peaks of XPS spectra confirming formation of -C3 N4 .
for the ternary compositions and the crystallization of the binary SiC and Si3 N4 thin films were obtained. Although there is a lot of literature on Si–C–N, a proper investigation of the effect of different substrates in the microstructure and corresponding mechanical properties of the coatings is lacking. In this communication, we have used magnetron sputtering (RF) techniques to deposit coatings on silicon (1 0 0), borosilicate glass and stainless steel (304SS) substrates. A correlation among thickness, hardness, modulus, microstructure and elemental composition with the nature of the substrate is presented. 2. Experimental procedures Si–C–N coatings were deposited on Si (1 0 0), 304SS and borosilicate glass substrates by RF magnetron sputtering (HHV, Bangalore, India) using single target SiC. Silicon and carbon powder was used in 1:1 ratio and was compacted into a 50 mm disc with 3 mm thickness and sintered at graphite furnace at 1900 ◦ C in argon atmosphere. For deposition, the vacuum chamber was first evacuated to 1 × 10−6 mbar pressure. Thin films were deposited on Si (1 0 0) substrate at 1 × 10−2 mbar pressures in Ar/N2 atmosphere in 1:9 ratio. The reason for maintaining this gas ratio was deposition of hardest film as reported in our previous publications. The temperature was kept constant at 500 ◦ C and the power was maintained at 400 W. The deposition was done for 3 h. Thickness of the coatings was measured by surface profilometer (Talysurf, Taylor Hobson, UK). AFM (Seiko Scanning probe, 400, Japan) was used for roughness measurements. Vickers microhardness of the coatings was measured at 15 gf load by Microhardness Tester (The LEICA VMHT AUTO, Austria). Nanoindentation tests were carried out in continuous stiffness mode using XP-Nanoindenter MTS, USA. The load–displacement curves, determined hardness and modulus. Microstructural investigation was done by transmission electron microscope (TEM) of Phillips, EM200, Netherlands. Almega Dispersive Raman Spectroscope having He–Ne laser (532 nm) was used for the Raman studies. It consisted of 25 m pin hole. Olympus microscope was used to select the location and 25% of laser power was used for the spectroscopy. The XPS analysis was carried out using Thermo VG Scientific Multitech 2000 XPS system, UK.
3. Results and discussion TEM studies showed the formation of -C3 N4 crystallites in the amorphous matrix of Si–C–N (Fig. 1a). The ring formation in the SAED pattern indicated the formation of polycrystalline phases (Fig. 1b). The matched d-values and corresponding planes are shown in the figure. The crystallization process was localized. Scanning other area mainly showed amorphous (a-Si–C–N) phase. The formation of -C3 N4 was also confirmed by XPS studies. The C 1s spectra showed peaks at 285.11 eV due to C–C, 290.32 eV due to C–N and 292.57 eV due to C N (Fig. 1(c)) [18–20]. The Raman spectra for all three films are exhibited in Fig. 2(a). In the range 1000–1800 cm−1 the bands were deconvoluted using 3 Gaussian fits as shown in Fig. 2(b) for silicon substrates. The band with peak value at 1529 cm−1 called the G (graphitic) band, arises due to stretching of sp2 , hybridised carbon [19–23]. For purely graphitic phase containing 100% sp2 hybridised carbon the G band peak occurs at 1580 cm−1 . The shift to lower wavenumber in this case is the indication of increase in sp3 carbon percentage in the coating which is due to incorporation of nitrogen in the coating forming mainly amorphous silicon carbonitride (a-Si–C–N) with very small -C3 N4 crystallites as observed in the TEM studies. The band with peak value at 1311 cm−1 is called the D (disorder) band, which arises due to small crystallite sizes, nitrogen incorporation and sp3 hybridised carbon [19]. A band peaked at 1153 cm−1 called the T band was also observed due to stretching modes of sp3 hybridised carbon [7]. The Raman spectra of a bare silicon substrate have been shown for comparison. The first and second order peaks observed at 526 cm−1 and 981 cm−1 , respectively, are typical of pure silicon. The adatoms generated during sputtering when falling on the substrate have to thermally equilibrate themselves and undergo
476
A.S. Bhattacharyya et al. / Journal of Alloys and Compounds 478 (2009) 474–478
coalescence. Compressive stress in the coatings aids in better deposition whereas tensile stress is a detrimental factor. Deposition at a higher substrate temperature increase the compressive stresses, and hence leads to better adherence. The extent of tensile stress is on the other hand is proportional to the difference in CTE values of the coating and the substrate. 304SS has a CTE value of 11 whereas it is 3.0 for Si (1 0 0) and 8.5 for borosilicate glass. Thus the best adhesion is obtained for silicon substrates as Si–C–N which has a CTE value of 0.5. The deposition rate (r) and average roughness (Ra) of the coatings varied due to CTE mismatch; however the variation in ´˚ deposition rate was negligible. Highest deposition rate of 175 A/min ´ ˚ was observed for silicon substrates and the lowest (162 A/min) was
Fig. 2. (a) Raman spectra of Si–C–N films on silicon, glass and steel substrates and (b) their 3 Gaussian fir deconvolutions for silicon substrate and (c) Raman spectra of bare silicon substrate.
for 304SS substrates. Therefore the lowest deposition time was used for coating silicon substrates giving a thickness of 1 m. Roughness was found to be the lowest in case of silicon (20 nm) but interestingly found to be the highest in case of glass (120 nm) instead of 304SS (50 nm), which is due to the presence of an oxide layer usually found in the top surface of a borosilicate glass. Adhesion tests were performed through scratch tests. Fig. 3 shows the variation of coefficient of friction and normal load with stroke length obtained during scratch tests. The change in slope occurring in the C.O.F curve is the indication of coating failure. In the case of silicon substrates (Fig. 3(a) and (b)), cohesive failure was observed initially where the coating starts to penetrate the coating. The adhesive film failure occurred at a critical load of 10 N where brittle failure and chipping of the coating was observed. For glass substrates (Fig. 3(c) and (d)) the cohesive and adhesive failure regions were not separately resolvable. The slope change in the C.O.F curve occurred at a corresponding critical load of 5 N and brittle failure by buckling was observed at higher loads. Lastly in case of steel substrates (Fig. 3(e) and (f)) all the three mechanism of cohesive failure, adhesive failure and complete coating removal through spallation was observed [24]. The critical load in this case was 7 N. Vickers microhardness tests performed at 15 gf showed Vickers hardness number (VHN) of 1434 for Si, 1560 for glass and 2123 for steel substrates. In the nanoindentation tests a sharp indenter (Berkovich) is used to test the properties of a thin coating and to ensure that the contact stresses are high enough to cause yield of the coating before the substrate. The initial contact will always be elastic, dominated by the spherical tip of the indenter. As the load is increased the sloping sides of the indenter will come into contact and the indenter behaves like a truncated cone. At this point the plastic deformation is initiated. This transition is associated with the contact-induced stresses exceeding the yield stress of the material being tested and full plastic deformation occurs beneath the indenter. The nanohardness, at this point is estimated from the ratio of maximum indentation load and projected contact area of the plastically deformed region (Fig. 4) [25,26]. A nanohardness of 17 GPa was found to be the highest in case of 304SS substrates. The lowest hardness of 10 GPa was found in case of silicon substrates and coatings deposited on glass substrates showed a hardness of 14 GPa. Table 1 summarises all the parameters obtained on all the three substrates. The lowest hardness obtained in case of silicon substrates inspite of having the lowest CTE mismatch and the highest deposition rate and the
Table 1 Variation of critical load, microhardness and nanohardness of Si–C–N films deposited on Silicon, glass and steel substrates. Substrate
r (nm/min)
Time (min)
t (m)
Ra (nm)
Lc (N)
VHN (15 gf)
H (GPa)
Silicon (100) Glass 304SS
16.6 12.5 10.8
60 120 240
1.0 1.5 2.6
20 120 50
10 5 7
1434 1560 2123
10 14 17
A.S. Bhattacharyya et al. / Journal of Alloys and Compounds 478 (2009) 474–478
477
Fig. 3. Scratch tests performed on coatings deposited on (a, b) silicon, (c, d) glass and (e, f) steel substrates.
best adhesion is due to the lowest thickness, which also explains the reason for getting the highest hardness in case of 304SS substrates. 4. Conclusions
Fig. 4. Nanohardness of Si–C–N coatings on different substrates.
A systematic comparison of RF magnetron sputtered Si–C–N coatings on Si (1 0 0), 304SS, and borosilicate glass was carried out. Evidence of nucleation of -C3 N4 having both sp2 and sp3 hybridised carbon was obtained from microscopic and spectroscopic studies. The substrate effect was mainly observed in hardness, modulus and adhesion properties of the films. Better deposition and adhesion was seen for silicon substrates due to less difference in CTE values with Si–C–N. Coating thickness was found to have a major influence on hardness, as silicon substrates although having the lowest CTE mismatch, gave the lowest hardness due to lowest thickness. The highest nanohardness of 17 GPa and microhardness of 2123 HV0.015 with a critical load of 7 N was obtained in case of coatings deposited on 304SS substrate.
478
A.S. Bhattacharyya et al. / Journal of Alloys and Compounds 478 (2009) 474–478
Acknowledgements The authors wish to thank Dr. M. Ghosh and Dr. D.D.N. Singh of NML, Jamshedpur, for their cooperation in carrying out TEM and Raman studies. This research work is sponsored by Department of Science and Technology, Government of India. References [1] H.J. Seifert, J. Peng, H.L. Lukas, F. Aldinger, J. Alloys Compd. 320 (2001) 251–261. [2] I.B. Lezak, A.M. Wrobel, M.P.M. Kivitorma, I.J. Vayrynen, A. Tracz, Appl. Surf. Sci. 253 (2007) 7211–7218. [3] I.B. Lezak, A.M. Wrobel, D.M. Bielinski, Diam. Relat. Mater. 15 (2006) 1650–1658. [4] E. Tomasella, F. Rebib, M. Dubois, J. Cellier, M. Jacquet, J. Phys: Conf. Ser. 100 (2008) 1–4. [5] H. Hoche, D. Allebrandt, M. Bruns, R. Riedel, C. Fasel, Surf. Coat. Technol. 202 (2008) 5567–5571. [6] N.J. Wagner, W.W. Gerberich, J.V.R. Heberlein, Surf. Coat. Technol. 201 (2006) 4168–4173. [7] K.B. Sundaram, Z. Alizadeh, R.M. Todi, V.H. Desai, Mater. Sci. Eng. A 368 (2004) 103–108. [8] W. Cheng, J. Jiang, Y. Zhang, H. Zhu, D. Shen, Mater. Chem. Phys. 85 (2004) 370–376. [9] J. Vleck, M. Kormunda, J. Cizek, V. Perina, J. Zemek, Surf. Coat. Technol. 160 (2002) 74–79.
[10] W.R. Chang, Y.K. Fang, S.F. Ting, Y.S. Tsair, C.N. Chang, C.Y. Lin, S.F. Chen, IEEE Electron. Dev. Lett. 24 (2003) 565–567. [11] C. Roh, S.H. Lee, V. Francois, J. Chromatogr. A 1179 (2008) 145–151. [12] Y. Katsuda, P. Gerstel, J. Narayanan, J. Bill, F. Aldinger, J. Eur. Ceram. Soc. 26 (2006) 3399–3405. [13] N. Janakiraman, F. Aldinger, J. Eur. Ceram. Soc. 29 (2009) 163–173. [14] J. Wan, P.R.L. Malenfant, S.T. Taylor, S.M. Loureiro, M. Manoharan, Mater. Sci. Eng. A 463 (2007) 78–88. [15] G. Lu, S. Qiao, C. Zhang, J. Hou, D. Jia, Y. Zhang, Comp. Part A: Appl. Sci. Manuf. 39 (2008) 1467–1470. [16] Y.S. Li, H. Kiyono, S. Shimada, X. Lu, A. Hirose, Diam. Relat. Mater. 15 (2006) 1727–1731. [17] M. Rudolphi, M. Bruns, H. Baumann, U. Geckle, Diam. Relat. Mater. 16 (2007) 1273–1277. [18] M.J. Bojdys, J.O. Muller, M. Antonietti, A. Thomas, Chem. Eur. J. (2008) 8177–8182. [19] A. Champi, F.C. Marques, Braz. J. Phys. 36 (2006) 462–465. [20] G. Peng, X. Jun, P. Yong, D. Wanyu, D. Xinlu, D. Chuang, Plasma Sci. Technol. 8 (2006) 425–428. [21] N.I. Fainer, Y.M. Rumyantsev, M.L. Kosinova, Chem. Sust. Dev. 9 (2001) 31–36. [22] N.I. Fainer, Y.M. Rumyantsev, M.L. Kosinova, Chem. Sust. Dev. 9 (2001) 37–43. [23] E.J. Liang, J.W. Zhang, J. Leme, C. Moura, L. Cunha, Thin Solid Films 469 (2004) 410–415. [24] S.K. Mishra, A.S. Bhattacharyya, Mater. Lett. 62 (2008) 398–402. [25] R.L. Zong, S.P. Wen, F. Zeng, Y. Gao, F. Pan, J. Alloys Compd. 464 (2008) 544– 549. [26] D. Drozdz, T. Kulik, H.J. Fecht, J. Alloys Compd. 441 (2007) 62–65.