Hot wire chemical vapor deposited multiphase silicon carbide (SiC ...

J Mater Sci: Mater Electron DOI 10.1007/s10854-016-4995-2

Hot wire chemical vapor deposited multiphase silicon carbide (SiC) thin films at various filament temperatures Amit Pawbake1 • Vaishali Waman1 • Ravindra Waykar1 • Ashok Jadhavar1 • Ajinkya Bhorde1 • Rupali Kulkarni1 • Adinath Funde1 • Jayesh Parmar2 • Somnath Bhattacharyya3 • Abhijit Date4 • Rupesh Devan5 • Vidhika Sharma5 Ganesh Lonkar5 • Sandesh Jadkar5

•

Received: 9 March 2016 / Accepted: 14 May 2016 Ó Springer Science+Business Media New York 2016

Abstract Influence of filament temperature (TFil) on the structural, morphology, optical and electrical properties of silicon carbide (SiC) films deposited by using hot wire chemical vapor deposition technique has been investigated. Characterization of these films by low angle XRD, Raman scattering, XPS and TEM revealed the multiphase structure SiC films consisting of 3C–SiC and graphide oxide embedded in amorphous matrix. FTIR spectroscopy analysis show an increase in Si–C, Si–H, and C–H bond densities and decrease in hydrogen content with increase in TFil. The C–H bond density was found higher than the of Si–H and Si–C bond densities suggesting that H preferably get attached to C than Si. AFM investigations show decrease in rms surface roughness and grain size with increase in TFil. SEM studies show that films deposited at low TFil has spherulites-like morphology while at high TFil has cauliflower-like structure. Band gap values ETauc and E04 increases from 1.76 to 2.10 eV and from 1.80 to 2.21 eV respectively, when TFil was increased from 1500 to 2000 °C. These result show increase in band tail width

& Sandesh Jadkar [email protected] 1

School of Energy Studies, Savitribai Phule Pune University, Pune 411 007, India

2

Tata Institute of Fundamental Research, Colaba, Mumbai 400005, India

3

Department of Metallurgical and Materials Engineering, IIT Madras, Chennai 600036, India

4

School of Aerospace Mechanical and Manufacturing Engineering, RMIT University, Plenty Road, Bundoora 3083, Australia

5

Department of Physics, Savitribai Phule Pune University, Pune 411 007, India

(E04–ETauc) of multiphase SiC films. Electrical properties revealed that rDark increases from *7.87 9 10-10 to 1.54 9 10-5 S/cm and Eact decreases from 0.67 to 0.41 eV, which implies possible increase in unintentional doping of oxygen or nitrogen due to improved crystallinity and Si–C bond density with increase in TFil. The deposition rate for the films was found moderately high ˚ /s) over the entire range of TFil studied. (21 \ rdep \ 30 A

1 Introduction Due to unique physical and chemical properties, cubic silicon carbide (3C–SiC) has attracted more attention in recent years for various potential applications in microelectronics and opto-electronics devices such as mechanical protection coating [1], an optical coating for solar cells [2], in light emitting diodes and solar cells [3], X-ray lithography masks [4], and in color sensors [5] etc. Due to its excellent chemically stability, it has been used as hard protective coating for harsh environment applications [6]. The 3C–SiC films have also been used in fabrication of micro-heaters and resistance thermometer device sensors for MEMS [7]. A variety of preparation methods have been used to synthesize SiC thin films to obtain desired physical and electronic properties at low substrate temperature. These include plasma enhanced chemical vapor deposition (PECVD) [8], electron cyclotron resonance CVD [9], magnetron sputtering [10], pulsed laser deposition (PLD) [11], ion implantation [12], molecular beam epitaxy (MBE) [13], hot wire chemical vapor deposition (HW-CVD) [14], photo chemical vapor deposition (Photo-CVD) [15], low pressure chemical vapor deposition (LP-CVD) [16] etc. The PECVD has been the most widely used to obtain 3C–SiC

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films. However, this method is subject to some drawbacks such as low carbon incorporation efficiency and a low deposition rate [17]. The ECR-CVD technique has the problem of non-uniform deposition over the large area and requires very complicated equipment. Deposition of 3C– SiC layers by sputtering technique requires high substrate temperature (750–800 °C) [18]. The LP-CVD technique requires very high substrate temperature of 800–900 °C for the growth of 3C–SiC films. In context of device fabrication, high deposition rate and low substrate temperature are desirable. Thus, search of alternate deposition methods, which allow high deposition rates at low substrate temperature and device quality 3C–SiC films, are enviable. The hot wire chemical vapor deposition (HW-CVD) method has received considerable attention in recent years owing to its capability to synthesize Si-based films such as a-Si:H [19], nc-Si:H [20], a-SiN:H [21], and a-SiC:H [22] at low substrate temperature at higher deposition rates. It has been reported that the structural and optical properties of materials deposited by HW-CVD method are better compared to those prepared by PE-CVD method [23]. Furthermore, this method has several advantages such as large-area deposition, efficient gas utilization, use of low substrate temperature etc. compared with other deposition methods [24]. The two major advantages of employing HW-CVD method for synthesis of Si based alloy coatings are the absence of the deleterious electrons and ions and surface charges which avoid of powder formation and second is high dissociation rate of source gases which leads to higher deposition rate. Recently, this method has been successfully employed for the synthesis of 3C–SiC films [25]. It has been reported that the microstructure and stoichiometric composition of the film critically depends on deposition parameters such as process pressure [26], the filament-to-substrate distance [27], flow rates of precursor gases [28], and filament temperature [29]. However, capabilities of HW-CVD for obtaining device quality 3C– SiC have not been fully established and few reports on filament temperature dependent studies on 3C–SiC films exist in the literature. With this motivation an attempt has been made to prepare 3C–SiC by HW-CVD method by varying the filament temperature. In this paper, we present results of investigation of influence of filament temperature on electrical, structural, morphology and optical properties of 3C–SiC films deposited by HW-CVD method.

2 Experimental 2.1 Film preparation Undoped 3C–SiC films were deposited simultaneously on Corning #7059 glass and c-Si wafers in a locally fabricated

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HW-CVD system, details of which have been described elsewhere [30]. Films were prepared by using pure silane (SiH4) as Si source gas, methane (CH4) as C source gas and hydrogen (H2) as dilution gas. The flow rate of SiH4, CH4 and H2 are kept constant at 3.5, 15 and 25 sccm respectively while filament temperature was varied from 1500 to 2000 °C in the step of 100 °C. The filament temperature was measured by optical pyrometer (IRCON, USA). The pressure during deposition was kept constant at 500 ± 5 mTorr using manual throttle valve. The substrate temperature was held constant (350 °C) using a thermocouple and temperature controller and filament-to-substrate distance (ds-f) was fixed at 5 cm. The glass substrates were cleaned with double distilled water whereas, the c-Si wafers were etched using solution of HF to remove native oxide layer. The substrates were loaded to the substrate holder and then the deposition chamber was evacuated to the base pressure \10-6 Torr. Prior to each deposition, the substrate holder and deposition chamber were baked for 2 h at 100 °C to remove any water vapor absorbed on the substrates and to reduce the oxygen contamination in the film. After that, the substrate temperature was brought to desired value by appropriately setting the inbuilt thermocouple and temperature controller. The deposition was carried out for desired amount of time and films were allowed to cool to room temperature in vacuum. 2.2 Film characterization Room temperature dark conductivity (rdark) was measured by employing a 2400 Keithley source-meter in planar geometry. The activation energy (Eact) was calculated from, Eact r ¼ r0 exp ð1Þ kT where k is the Boltzmann constant and T is the absolute temperature. The carrier concentration, charge carrier mobility and Hall coefficient were measured using Van der Paw method (Ecopia HMS-3000 Hall Measurement System). Fourier transform infrared (FTIR) spectra of the films were recorded by using FTIR spectrophotometer (Shimadzu, Japan). Bonded hydrogen content (CH) was calculated from wagging mode of IR absorption peak using the method given by Brodsky et al. [31]. The optical band gap of 3C–SiC:H films was deduced from transmittance and reflectance spectra of the films deposited on corning glass and were measured using a JASCO, V-670 UV–visible spectrophotometer in the range 250–1100 nm. Raman spectra were recorded with Raman spectroscopy instrument (Jobin–Yvon Horibra LABRAM-HR) in the range 200–1200 cm-1. The spectrometer has backscattering geometry for detection of Raman spectrum with the


resolution of 1 cm-1. The excitation source was 632.8 nm line of He–Ne laser. The power of the Raman laser was kept to\5 mW to avoid laser induced crystallization on the films. Low angle X-ray diffraction pattern were obtained by X-ray diffractometer (Bruker D8 Advance, Germany) ˚ ). The patterns were using Cu Ka line (k = 1.54056 A taken at a grazing angle of 1°. The average crystallite size was estimated using the classical Scherrer’s formula [32], dXray ¼

0:9k B cos hB

ð2Þ

The XPS studies were carried out using VSW ESCA machine, under the vacuum of better than[10-9 Torr, with AlKa (1486.6 eV) radiation with resolution of 0.1 eV. The XPS signal was obtained after several scans in the acquisition process. The spectra were recorded for specific elements. The surface morphology of the films have been investigated using non-contact atomic force microscopy (NC-AFM) (JEOL, JSPM-5200, Japan) and field emission scanning electron microscopy (FE-SEM) (Hitachi, S-4800, Japan). Transmission electron microscopy images and selected area electron diffraction (SAED) pattern were recorded using a transmission electron microscope (TECNAI G2-20-TWIN, FEI, The Netherlands) operating at 200 keV. Thickness of films was determined by profilometer (KLA Tencor, P-16?) and was further confirmed by UV–visible spectroscopy using the method proposed by Swanepoel [33].

3 Results and discussion In HW-CVD, source gases silane (SiH4) and methane (CH4) are broken by the hot filament, forming radicals that bond to the growing surface. The type of radicals formed for each species is related to the bonding energies of each molecule. The radicals that reach the growing surface may present different sticking coefficients. Hence, the concentration of Si and C in the film may be different from that in the gas phase. In the present study, first the formation of multiphase SiC has been confirmed by low angle XRD, Raman scattering, X-ray photoelectron spectroscopy (XPS) and transmission electron microscopy (TEM) analysis and then H-related features, morphology, optical and electrical properties were discussed later. 3.1 Variation in deposition rate Figure 1 shows the variation of deposition rate as function of filament temperature (Tfil) for the SiC films deposited by HW-CVD method. As seen from the figure for low filament temperature (Tfil \ 1700 °C) the deposition rate decreases ˚ /s with increase in filament temperature from 27 to 21 A

Fig. 1 Variation of deposition rate as a function of filament temperature for SiC films deposited by HW-CVD method

whereas for higher filament temperature (Tfil [ 1700 °C) it ˚ /s. These results suggest that the increases from 21 to 30 A change in the deposition rate is filament temperature dependent reaction rate taking place at the filament surface and/or radical–radical reactions occurring in the vicinity of substrate surface. At lower filament temperatures (\1700 °C), the supply of film forming species may be limited by incomplete decomposition of SiH4 and CH4 at the heated filament. As a result the deposition rate decreases with increase in filament temperature. The increase in deposition rate with further increase in filament temperature ([1700 °C) may be due to increase in decomposition rate of SiH4 and CH4 at the heated filament which increases the density of film forming species. The deposition rate obtained in the present study for SiC films is much higher than the PE-CVD [22], low frequency inductively coupled plasma (ICP) CVD [34], DC saddle PE-CVD [35] and recently laser chemical vapor deposition [36] grown SiC films. 3.2 Low angle XRD analysis Figure 2 shows XRD pattern of SiC films deposited at different filament temperatures (TFil). The films deposited over the entire range of filament temperature shows a broad peak centered at 2h * 12.3° corresponding to graphite oxide [37]. The film deposited at TFil = 1700 °C shows onset of formation of 3C–SiC phase. The film shows a peak centered 2h * 35.4° corresponding to (111) diffraction plane of 3C–SiC structure [38, 39]. With increase in filament temperature two more diffraction peaks emerge in the XRD pattern at 2h * 59.39° and 71° which corresponds to (220) and (311) diffraction planes of 3C–SiC structure [38, 39]. The dominant diffraction peak is (111) signifying that the SiC crystallites have preferred orientation in (111) direction. With increase in filament temperature, the

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Fig. 2 Low angle XRD pattern of SiC films deposited at different filament temperatures

intensity of all diffraction peaks increases indicating increase in crystallinity in the film. Also, there is no significant change in line-width of (111) diffraction peak suggesting that the average grain size (dX-ray) is constant (4–5 nm) for the films deposited at higher filament temperatures. In addition, a tiny shoulder centered *33.8° can be seen for the films deposited at TFil [ 1700 °C which corresponds to hexagonal SiC or due to a high staking density of faults and twins [40]. Thus, based on the low angle XRD results, a multiphase structure of the SiC is suggested, which contain graphide oxide and 3C–SiC phases embedded in amorphous matrix. The Raman spectroscopy results also confirm the presence of multiphase structure of SiC thin films.

Fig. 3 Raman spectra of SiC films deposited at different filament temperatures

3.3 Raman scattering analysis Figure 3 shows Raman spectra of SiC films prepared at different filament temperatures (TFil). As seen from the figure, the Raman spectra show three broad peaks, first in the range of 400–700 cm-1, second in the range of 800–1300 cm-1 and third in the range of 1400–1700 cm-1. The Raman scattering efficiency of the Si–C band is much smaller than that of C–C and Si–Si bands and due to overlapping of Si–C, C–C and Si–Si bands it is necessary to de-convolute the Raman spectrum to acquire change in peak position and its line-shape. Therefore, each Raman spectrum was de-convoluted in the range 300–1800 cm-1 using software XPSPEAK (version 4.1). Typical de-convoluted Raman spectra for the film deposited at filament temperature 1900 °C is shown in Fig. 4. From de-convoluted data, it has been observed that the Raman shoulder in the range of 400–700 cm-1 is superposition of two different peaks one centered *510 cm-1 and other *697 cm-1. The peak *510 cm-1 may be due

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Fig. 4 Typical de-convoluted Raman spectra for the film deposited TFil = 1900 °C

to SiC acoustic phonons particularly reported for small SiC crystallites [41]. This peak shift towards lower wave number and its intensity decline with increase in filament temperature. The other tiny broad shoulder *697 cm-1 may corresponds to Si–H (wagging/bending) [42]. Raman shoulder in the range of 800–1300 cm-1 is superposition of two different peaks, one *968 cm-1 and second *1101 cm-1. The peak at 968 cm-1 can be assigned to longitudinal optic (LO) phonon vibrations of 3C–SiC structure of SiC [43] and second order scattering of Si–Si [44]. With increase in the filament temperature the peak


also shift towards lower wave number from its ideal position 973 cm-1. The shifting of peak towards lower wave number is probably due to small crystallite size and strain associated with small SiC crystallites [45]. The relative strengths and shape of these peaks for the film deposited at different filament temperatures are almost identical signifying the similar degree of crystallinity, however exact determination of crystallite size or crystalline volume fraction using Raman spectra is not possible in the present case. The presence of another peak *1099 cm-1 may be due to the Si–O–Si vibrations [46]. The de-convoluted Raman band in the C–C network in the range of 1400–1700 cm-1 is superposition of three peaks. The peak at *1440 cm-1 is attributed to graphite C attached to Si (graphite phase) [47], the peak at *1523 cm-1 can be attributed to G like C [48, 49] and band at *1641 cm-1 corresponds to sp2 bonded C [50]. In our case, no signature of diamond like carbon (DLC) (peak at *1250 cm-1) [49] or nano-crystalline diamond (NCD) (peak at *1170 cm-1) [51, 52] observed in the Raman spectra over the entire range of filament temperature studied which certainly confirm the presence of C–C bonds in the films. These results further support to the formation of multiphase structure of SiC. 3.4 Transmission electron microscopy (TEM) analysis Further to confirm the formation of multiphase structure of SiC transmission electron microscopy (TEM) analysis was performed. For TEM analysis film deposited at 2000 °C was prepared using gatan method. In this method, first 3 mm disk of Si wafer containing the sample was sliced using ultrasonic cutter and then grinding and dimpling were carried out to thin down the sample to a residual thickness of 10 to 15 lm. Finally single-sided Ar? beam milling was performed at small angle (\5°) and at low energies (acceleration voltage 4 kV, beam current\23 lA) to avoid heating of the sample. The TEM image of SiC film deposited at 2000 °C using HW-CVD method is shown in Fig. 5a. The 3C–SiC crystallites of different size and shapes, observable as deep dark spots, are randomly distributed over a relatively bright and homogeneous a-Si matrix. Figure 5b, c are enlarged views of marked area of A and of B in Fig. 5a, respectively. Figure 5d is selected area electron diffraction (SAED) pattern of SiC film deposited at 2000 °C using HW-CVD method. The enlarged view of marked area of A shown in Fig. 5b shows that the distance between the adjacent lattice planes is *0.65 nm which corresponds to graphene oxide [53]. The enlarged view of marked area of B shown in Fig. 5b shows that the distance between the adjacent lattice planes is *0.28 nm, which is consistent with the spacing of (111)

plane of 3C–SiC [JCPDS data card #73-1665]. Therefore, combined with the low angle XRD analysis it can be concluded that the synthesized films are a composite of two phases, namely 3C–SiC and graphene oxide embedded in the amorphous matrix. The bright and spotty diffraction rings in the SAED pattern (Fig. 5d) signify a high degree of crystallinity of the material. 3.5 X-ray photoelectron spectroscopy (XPS) analysis The formation of multiphase SiC films and, their composition and atomic bonding states has been confirmed by X-ray photoelectron spectroscopy (XPS) analysis. Figure 6a shows the XPS wide scan of HW-CVD deposited SiC:H film prepared at filament temperature 1900 °C. The scan shows the silicon (Si 2p), carbon (C 1s), and oxygen (O 1s) peaks. The measured XPS spectrum is a superposition of peaks from the various types of bonds, and therefore it is necessary to deconvolute the wide scan spectrum to determine the individual chemical bonding configuration present in the film. Figure 6b–d show typical deconvoluted XPS spectra of the Si (2p), C (1s), O (1s) electron state, respectively. The Si 2p spectra (97–105 eV) exhibit four different peaks at the binding energies of 99.2, 100.21, 101.8 and 103.2 eV, which were assigned to the Si–H and/or Si–Si, Si–C, Si–Ox and O–Si–C bonds, respectively [54]. The C 1s peak (280–288 eV) was deconvoluted into two peaks at 282.37 and 284.45 eV. The peak at 282.37 eV is attributed to Si–C bonds and the peak at 284.45 eV is attributed to C–C/C–H bonds [55]. The peak located at 284.45 eV can also be assigned to carbon atoms having graphitic bonding [56]. The O 1s peak (526–540 eV) was de-convoluted into four peaks at 531.19, 532.02, 532.84 and 533.85 eV which were assigned to O– Si, O–CH, O=C and O–O bonds respectively [57–61]. The oxygen peak may be due to adsorbed oxygen and surface oxidation of the film prior to XPS measurement. The oxygen was found to be incorporated in SiC films prepared at even lower base pressure than the pressure employed for the XPS measurements [38]. The elemental mapping for the film deposited at filament temperature 1900 °C have been estimated and it is found that the film consist of 24 % of silicon, 61 % of carbon and 12 % of oxygen. 3.6 Fourier transform infra-red (FTIR) spectroscopy analysis To reveal the H bonding configuration (Si–C, Si–H and C– H bonds) and to estimate the bonded hydrogen content (CH) in multiphasic SiC films, FTIR spectroscopy was used. The FTIR transmission spectra (normalized for thickness) of SiC films deposited by HW-CVD method at

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TFil = 2000 0C

(a)

A B

a-Si (b) Enlarged view of A

(c) Enlarged view of B

(d) SAED Pattern

d = 0.28 nm d = 0.65 nm

Fig. 5 a TEM image of SiC film deposited at TFil = 2000 °C using HW-CVD method. b Enlarged view of marked area of A (a). c Enlarged view of marked area of B (a). d SAED pattern

different filament temperatures is shown in Fig. 7. For clarity, the spectra have been broken horizontally into two parts viz; between 400–1200 and 1900–2300 cm-1. As seen from the FTIR spectra, the film deposited at TFil = 1500 °C, show three low intensity modes at 640, 1000 and 2085 cm-1 which are associated with wagging or rocking mode of Si–H [62], rocking and wagging mode of Si–CHn bonds or Si–O–Si [63] and stretching mode of Si– Hn (n = 1, 2) [64] respectively indicating very low hydrogen content in the films [65]. With increase in filament temperature, modes at 640 and 1000 cm-1 are merged and a vibrational band *780 cm-1 appear in the FTIR spectra which is associated with Si–C stretching mode [63, 66] and its intensity increases with increase in filament temperature. This result suggests increase in Si–C bond density and hence carbon content in the film with increase in filament temperature. To determine the bonded hydrogen content (CH) the absorption coefficient (a) can be obtained through the Beer–Lambert law, T ¼ T0 eaðxÞd

ð3Þ

where d is the film thickness, T and T0 are the transmittance of the film substrate and the substrate, respectively,

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and a(x) is the absorption coefficient at frequency (x). The number of Si–C (NSi–C), Si–H (NSi–H), and C–H(NC–H) bonds has been estimated by taking area under the curve for the respective peak using the relation [67], Z aðxÞ N ¼ Ax dx ¼ Ax Ix ð4Þ x where the oscillator strength, Ax, has a value of 2.13 9 1019 [68], 1.4 9 1020 [67], and 1.35 9 1021 cm-2 [68] for Si–C, Si–H, and C–H, respectively. The hydrogen content was then calculated by using the relation [64], CH ¼

NSiH þ NCH NSiC

ð5Þ

The merged Si–H, C–H, and Si–C bands are de-convoluted to separate out the different bonding configurations present in the films. Variation of Si–C, Si–H, and C–H bond densities and CH as a function of filament temperature is displayed in Fig. 8. As seen from the figure, the bond density of Si–C, Si–H, and C–H increases with increase in filament temperature. All together the hydrogen content decreases from 21.53 to 2.42 at.%. Decrease in hydrogen content with increase in filament temperature for SiC films


Fig. 6 Typical XPS spectra for SiC film deposited at TFil = 1900 °C by HW-CVD method. a Wide scan, b de-convoluted XPS spectra of Si (2p) in the range 97–105 eV, c de-convoluted XPS spectra of C 1s

Fig. 7 FTIR transmission spectra of multiphase SiC films prepared at different filament temperatures

peak in the range 280–288 eV, d de-convoluted XPS spectra of O 1s peak in the range 526–540 eV

Fig. 8 Variation of Si–C, Si–H and C–H bond densities and total hydrogen content as a function of filament temperatures for multiphase SiC films prepared by HW-CVD method

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grown by HW-CVD has been reported previously [69, 70]. The bond density of C–H is higher than bond density of Si– H and Si–C suggesting that the hydrogen preferably get attached to C than Si. 3.7 Atomic force microscopy (AFM) and scanning electron microscopy (SEM) analysis Figure 9a1, a2 shows AFM surface profiles of multiphase SiC films deposited at filament temperature 1600 and 2000 °C respectively. The AFM micrograph for the film deposited at TFil = 1600 °C (Fig. 9a1) reveal that agglomerated tiny clusters of nanocrystalline 3C–SiC embedded in the amorphous matrix. The shape of grains on the surface is somewhat hemispherical or spherical. We may attribute hemispherical or spherical grains to an amorphous cluster, since our low angle XRD pattern for the film deposited at TFil = 1600 °C didn’t show any crystalline phase (see Fig. 2). For this film the average cluster/grain size is in the range *200–260 nm and rms surface roughness *7.1 nm. With change in filament temperature, significant differences in the structure can be clearly seen. The film deposited at TFil = 2000 °C (Fig. 9a2) indicate increase in grain density and each grain Fig. 9 Non contact atomic force microscopy (NC-AFM) images of films deposited at a1 TFil = 1600 °C, a2 TFil = 2000 °C and Scanning electron microscopy (SEM) images of the films deposited at b1 TFil = 1600 °C, b2 TFil = 2000 °C

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has an individual identity with its size in the range *140–170 nm and rms surface roughness *6.3 nm. These results indicate that with increase in filament temperature the rms surface roughness and the grain size of multiphase SiC film decreases. The surface SEM images of multiphase SiC films deposited at filament temperature 1600 and 2000 °C are is shown in Fig. 9b1, b2 respectively. The nanocrystalline 3C–SiC embedded in the amorphous matrix film obtained at TFil = 1600 °C show spherulites formed by small-sized crystallite aggregation and rough surface (Fig. 9b1). In contrast, SiC film deposited at TFil = 2000 °C show highly dense cauliflower-like surface structure. However, surface structure somewhat become porous. It is believed that the cauliflower-like surface structure has lots of grain boundaries and voids [71] in which dandling bonds may exist [72] suggesting that with increasing filament temperature grain boundaries and voids in multiphase SiC films increases. 3.8 Variation in optical band gap The optical band gap of multiphase SiC films has been deduced from the optical transmission measurement using Tauc equation [73],


ðaEÞ1=2 ¼ B1=2 ðE ETauc Þ

ð6Þ

where a is the absorption coefficient, B is the combined optical density of state, E is the photon energy is used to extract the optical band gap ETauc from a linear fit. It is agreed that the determination of band gap is not accurate using Tauc’s formulation because of the extent of the valence and conduction band tails in the gap, especially in case of films where nanocrystallites are embedded in the amorphous matrix. However, it is expected to give an approximate estimate of the band gap. Thus, the energy where the absorption coefficient (a) is equal to 10-4 cm-1, designated E04, is also obtained from the transmission measurements. Figure 10 shows estimated values of ETauc and E04 as a function of filament temperature for SiC films deposited by HW-CVD method. Both, ETauc and E04 increase with increasing filament temperature. As reported by Soloman et al. [74] the optical band gap values estimated from E04 method are slightly higher than the ETauc values calculated from Tauc’s plot. The ETauc increases from 1.76 to 2.10 eV and E04 increases from 1.8 to 2.21 eV when filament temperature was increased from 1500 to 1900 °C. As a result, band tail width of material taken as (E04–ETauc) increases with increase in filament temperature. The main factors that affect the band gap of SiC films are the C-to-Si ratio, amount of hydrogen content and crystalline fraction [63]. Our FTIR spectroscopy analysis showed decrease in hydrogen content with increase in filament temperature (see Fig. 8). This is expected to lower the magnitudes of both ETauc and E04 with increase in filament temperature. On the contrary we observed an increase in the band gap values. Therefore, increase in band gap values may be due to increase in carbon content in the film with increase in filament temperature.

Fig. 10 Variation of ETauc and E04 as a function filament temperature for SiC films deposited by HW-CVD method

3.9 Dark conductivity and activation energy measurements Figure 11 shows the variation of dark conductivity (rDark) and charge carrier activation energy (Eact) of multiphase SiC films as a function of filament temperature (TFil). It can be seen from the figure that rDark increases from *7.87 9 10-10 to 1.71 9 10-5 S/cm when filament temperature increased from 1500 to 1900 °C and then it further decreases slightly to 1.54 9 10-5 S/cm for the film deposited at TFil = 2000 °C. The charge carrier activation energy shows an opposite trend. The minimum Eact (*0.41 eV) was observed for the film deposited at TFil = 1900 °C. Although no intentional doping was performed, the rDark of SiC films obtained in the present study was very high and the Eact was low. The observed changes in rDark and Eact can be attributed to the microstructural changes and unintentional doping of nitrogen or oxygen [75, 76] or the presence of a graphitic phase [77] in the films. The low angle XRD, Raman spectroscopy and TEM support this conjecture. Films prepared at low TFil contain low Si–C bond density and are amorphous in nature. The impurity atoms of nitrogen or oxygen remain in its natural bonding state within host amorphous SiC matrix and do not contribute any mobile electron for conduction, hence the conductivity is low [25, 78]. As density of Si–C bond and crystallinity of films improves with increasing TFil, impurity atoms (nitrogen or oxygen) are forced to have the tetrahedral bonding configuration of host SiC matrix and act as a substitutional donor resulting in an increase in dark conductivity and consequent decrease in charge carrier activation energy [25]. The FTIR spectroscopy analysis further supports this (Fig. 8). Furthermore, the photosensitivity gain obtained for the films deposited with increasing TFil was found negligible suggesting that films are highly crystalline because nano/microcrystalline Si films

Fig. 11 Variation of dark conductivity and charge carrier activation energy as a function of function filament temperature

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prepared by different methods show high dark conductivity and less photosensitivity depending on the crystallite size and its volume fraction [79].

4 Conclusions Influence of filament temperature (TFil) on the structural, morphology, optical and electrical properties of silicon carbide (SiC) films was investigated by using various characterization techniques. The low angle XRD, Raman scattering, XPS and TEM reveal the presence of multiphase structure SiC films consisting of 3C–SiC and graphide oxide embedded in amorphous matrix. In FTIR analysis, with increase in filament temperature the density of Si–C, Si–H, and C–H bonds increases. The C–H bond density was found higher than the of Si–H and Si–C bond densities suggesting that H preferably get attached to C than Si. The AFM investigations show decrease in rms surface roughness and the grain size with increase in filament temperature. The SEM analysis shows that films deposited at low filament temperature have spherulites-like surface morphology while deposited at high filament temperature has cauliflower-like surface structure. Band gap values, both, ETauc and E04 increases with increase in filament temperature. The rDark increases with increase in filament temperature while Eact show opposite trend suggesting possible increase in unintentional doping of oxygen or nitrogen due to improved crystallinity and Si–C bond density in the films. These large band gap and highly conducting films are useful for application as window layers in thin-film solar cells. Acknowledgments Authors are thankful to Department of Science and Technology (DST), Ministry of New and Renewable Energy (MNRE), and University Grants Commission (UGC), Government of India, New Delhi for the financial support. Dr. Vidhika Sharma and Dr. Ganesh Lonkar are thankful to UGC, New Delhi for Dr. D. S. Kothari postdoc fellowship. One of the authors SRJ is thankful to University Grants Commission, New Delhi for special financial support under UPE program.

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