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Structural and optical properties of silicon thin-films deposited by hot-wire chemical vapor deposition: The effects of silane concentrations
Thin Solid Films 542 (2013) 139–143
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Thin Solid Films journal homepage: www.elsevier.com/locate/tsf
Structural and optical properties of silicon thin-films deposited by hot-wire chemical vapor deposition: The effects of silane concentrations A.K. Panchal a, Vivek Beladiya b, Vipul Kheraj b,⁎ a b
Electrical Engineering Department, S.V. National Institute of Technology, Ichchhanath, Surat 395007, India Department of Applied Physics, S.V. National Institute of Technology, Ichchhanath, Surat 395007, India
a r t i c l e
i n f o
Article history: Received 20 September 2012 Received in revised form 24 June 2013 Accepted 27 June 2013 Available online 12 July 2013 Keywords: Silicon Thin-film Raman spectra Microcrystalline Grain
a b s t r a c t In this paper, the structural and optical properties of a series of silicon (Si) thin-films deposited using hot-wire chemical vapor deposition with different silane concentrations (SCs) are presented. All the films are characterized by Raman spectroscopy, scanning electron microscopy (SEM) and photoluminescence (PL). In the Raman analysis, the first order and specifically the second order Raman spectra indicate increase in crystalline grain size as well as crystalline volume fraction in the films with a reduction in SC, which is also confirmed by the SEM analysis. At the higher SC, the Si microcrystalline grains get embedded in the nanocrystalline Si network. The Gaussian fitted peaks in the PL analysis reveal the emission due to either band to band tail-state transitions or tail-state to mid-gap defect-state transitions due to Si-dangling bonds present in the films. © 2013 Elsevier B.V. All rights reserved.
1. Introduction The solar photovoltaics (PV) has been considered to be one of the most viable renewable energy technologies. It could be an important energy source to deal with the current scenario of the energy crisis as well as for the sustainable development [1]. From the first generation crystalline silicon (Si) cells to the third generation thin-film PV technologies, the efficiency and the cost-effectiveness have improved significantly. Especially, the developments of materials and improved growth techniques have caused a remarkable progress in the field of PV. However, Si is still one of the most important materials for the PV technologies with its total share of more than 87% in the current market [2]. With increasing penetration of amorphous and microcrystalline Si materials in the thin-film solar photovoltaics (PV) now-a-days, it is extremely important to understand the thin-film growth mechanism for these materials for further developments of the technology. The crystalline Si, due to its indirect bandgap, has relatively poor light-absorption properties. Moreover, the Si bandgap (~1.12 eV) is considerably lower from the optimum value of ~1.4 eV for the solar PV materials [3]. With the developments of thin-film amorphous and microcrystalline Si technology, it has become possible to engineer the Si bandgap. Amorphous and microcrystalline Si also exhibit better light-absorption properties than crystalline Si [4]. However, these materials have their own limitation in the form of high defect-densities, ⁎ Corresponding author. Tel.: +91 9904334220. E-mail address: [email protected] (V. Kheraj). 0040-6090/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.tsf.2013.06.098
which lead to the hindered carrier-transport in the material. Further, the light-induced degradation in these materials raise a great concern related to the device lifetime as well as performance reliability of the solar cells [5]. This trade off requires the combination of amorphous and crystalline Si to exploit the solar energy in the optimum manner. Quite a few schemes have been taken up to combine the two phases together in a PV cell. The recent advances in micro-morph tandem solar cells have promised the great deal of improvements in the maximum solar spectrum utilization [6]. Nearly 6% of the 27.2 GW power generated by solar PV in 2010 was contributed by thin film amorphous and microcrystalline Si PV and it is likely to be taken over in the near future by the tandem amorphous/microcrystalline Si modules [6]. Recently, Dalal et al. [7] have also reported the improved performance of Si-based cells with multilayered structure of alternating amorphous and microcrystalline layers. Thus, the study of Si as an absorber material and its optical properties with the variation in the crystal structure is of great interest. The optical and electrical behaviors of Si thin-film are very much dependent on the structural properties, which in turn rely up to a great extent on the growth conditions. It is very important to understand the growth mechanism of Si thin-films as the crystalline fraction of the material can be tuned with the deposition parameters. Various deposition techniques have been employed to grow the Si thin-films with varying crystalline structure [8]. Although plasma enhanced chemical vapor deposition (PECVD) is one of the widely used techniques for the deposition of amorphous, nanocrystalline and microcrystalline Si thin-films, the growth rate in PECVD is relatively low. Moreover, the lower precursor gas utilization efficiency in PECVD
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A.K. Panchal et al. / Thin Solid Films 542 (2013) 139–143
Table 1 Details of sample numbers and corresponding silane concentration (SC). Proportional flow SiH4:H2
Sample number
SC (%) (SiH4/(SiH4 + H2))
1:00 1:05 1:10 1:20 1:40
S100 S105 S110 S120 S140
100 16.7 9.1 4.8 2.4
results in higher manufacturing costs. In light of these limitations, the hot-wire chemical vapor deposition (HWCVD) is a promising technique for the deposition of Si thin-films with significantly high growth rate (up to 10 nm/s) and efficient precursor-gas utilization [9,10]. It is also observed that the open circuit voltage and the efficiency of the solar cells grown by HWCVD are higher than those grown by the PECVD [11]. Quite a few groups have employed HWCVD for the growth of amorphous as well as microcrystalline Si [10–12]. The transition from amorphous to microcrystalline phase in Si has also been investigated [13]. The feasibility of HWCVD for the deposition of Si on moving glass substrate for the roll to roll mass production has also been reported [14]. Moreover, Dusane has shown that the crystalline volume fraction of the film, grain orientation, grain size and the deposition rate are very
sensitive to the working gas concentrations in HWCVD [15]. Thus, it is possible to fine-tune the structural properties and consequently the optical and electrical properties of the films by controlling the precursor gas flow. Here, we report the study of growth of Si thin-films using HWCVD. The present work is an effort to understand the effects of precursors' concentration, mainly the silane concentration (SC) on the structural and optical properties of the films. We present a comprehensive analysis using first and second order Raman spectra in tandem with photoluminescence (PL) and scanning electron microscopy (SEM) analysis in order to study the properties of Si films and to investigate the influence of SC. The systematic data are also reported to understand the HWCVD growth parameters for the fine-tuning of properties of Si films in the perspective of solar PV applications. 2. Experimental details All films were deposited on RCA cleaned p-type crystalline Si wafer (one sided polished, 0.01–0.02 Ω-cm resistivity, 2 inch diameter, b100N orientation) pieces in HWCVD chamber. The films were deposited with a substrate temperature 200 °C. The silane (SiH4) and hydrogen (H2) gas mixture was cracked at the hot tungsten filament kept at 1900 °C. The distance between the filament and the substrate was maintained 5 cm during the process. Films with nearly 300 nm thickness were prepared with fixed SiH4 flow (1 sccm) and varying H2
Fig. 1. SEM micrographs of films: (a) S100, (b) S105, (c) S110, (d) S120 and (e) S140. All micrographs have a bar scale of 100 nm.
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Fig. 2. (a) Raman spectrum of the a-Si film S100, (b) first order Raman spectra of the films with decreasing SC, (c) Gaussian peaks fit in the first order Raman spectra of film S105 and (d) second order Raman spectra of the films with decreasing SC.
flow (0, 5, 10, 20 and 40 sccm) and the precursor gases were cracked for 30 min. The base pressure in the chamber was ~10−5 Pa before the deposition. The films are identified as film (letter-S) followed by fixed SiH4 flow (numeral-1) followed by H2 flow (numerals-00, 05, 10, 20 and 40) and are listed in Table 1. To investigate the structural morphology of all films, SEM and Raman spectroscopy characterizations were performed. SEM was done using Reith 150 model with beam energy 10 keV. Raman analysis was conducted using a Confocal Jobin-Yvon Raman Spectrometer with a 514.5 nm argon laser beam with a power intensity 18 mW and an aperture 100 μm. In order to study the light emission mechanisms from the amorphous and microcrystalline Si films, PL measurement Table 2 Details of Raman peaks and the crystalline fraction obtained from the Gaussian fit for all films. Films Peak 1
Peak 2
Raman FWHM Integrated Raman FWHM Integrated shift (cm−1) intensity (cm−1) intensity shift (cm−1) (cm−1) S100 S105 S110 S120 S140
480.0 519.9 519.9 519.9 519.9
– 3.6 3.9 3.7 4.4
– 793.4 1669.7 3053.6 8569.5
– 517.1 517.2 518.0 516.2
– 15.5 16.8 15.0 20.5
– 1101.2 1747.4 3189.6 8638.4
Crystalline fraction (%)
– 41.87 48.86 48.91 49.80
was performed with a focused 532 nm, 25 mW argon laser excitation source. The emitted light was collected by a lens and detected by an array detector mounted on a triple grating spectrometer. All the measurements were performed at room temperature. 3. Results and discussions Fig. 1(a)–(e) shows the surface morphology of all films as SC decreases. No Si grain growth is seen in Fig. 1(a) of film S100 as it is completely amorphous in nature. The grain growth starts appearing in film S105 (Fig. 1(b)) and the horizontal grain growth increases gradually in the films S110 to S140 as observed in Fig. 1(c)–(e). The measured average grain size is 44, 79 and 235 nm for the films S110, S120 and S140, respectively. SEM analysis reveals that the grain size increases as SC decreases in the films, which is in agreement with the literature [16]. To further understand this observation, films were subjected to the first and second order Raman spectroscopy. Fig. 2(a) shows a broad spectrum centered at the Raman stokes line at 480 cm−1, which corresponds to the transverse optical (TO) phonon mode of a-Si in film S100 [17]. This indicates that S100 has Si film which is completely amorphous in nature. Film S100 is deposited directly on a pure crystalline Si substrate, but no peak is detected at Raman line 520 cm−1 (corresponding to crystalline Si) in Fig. 2(a). Moreover, the estimated integrated intensity at 520 cm−1 in all other films increases with decrease in SC as seen in Table 2. These facts
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Fig. 3. Room temperature PL spectra with Gaussian fitted peaks for the films: (a) S100, (b) S105, (c) S110, (d) S120 and (e) S140.
clearly decline the possibility of detection of Raman signal from the crystalline Si substrate and confirm that the signals are detected only from the films. Fig. 2(b) illustrates the first order Raman spectra centered at the Raman stokes line 520 cm−1 for different films. The spectra have asymmetric shoulder in the lower Raman frequency range and are broader than that of the monocrystalline Si (nearly 2–3 cm−1 FWHM). It is also observed that the intensity increases with decrease in SC. Increasing intensity suggests the increased crystalline fraction in the films. To investigate the presence of different Si phases in the films, the Raman spectra are best fitted with Gaussian distribution function. Fig. 2(c) represents the Gaussian fitted Raman spectrum of film S105. A narrow peak centered at 519.9 cm−1 (with FWHM 3.6 cm−1) is corresponding to the crystalline phase of Si and a broad peak at 517.1 cm−1 (with FWHM 15.5 cm−1) is corresponding to the nanocrystalline phase of Si. Similarly, these two major peaks are also observed in all the films as seen in Table 2. The first peak is at 519.9 cm−1 (with FWHM ranging from 3.6 to 4.4 cm−1) and the second is in the range 516 to 518 cm−1 (with FWHM ranging from 15 to 20.5 cm−1). Thus, Raman study of the films confirms the presence of two phases of Si in all the films. This analysis of the first order Raman spectra helps to conclude that the crystalline Si grains
Table 3 Gaussian fitted PL peaks and corresponding FWHM for all films. Films Peak 1
S100 S105 S110 S120 S140
Peak 2
Peak 3
Wavelength (λ1)
FWHM
Wavelength (λ2)
FWHM Wavelength (λ3)
FWHM
956.53 989.82 990.37 990.05 951.57
172.716 104.00 101.18 95.02 66.63
995.92 1008.41 1009.54 1012.24 1004.07
88.50 57.81 56.06 52.59 74.09
50.538 68.00 65.22 65.03 58.23
1089.02 1080.54 1081.93 1076.38 1077.84
are embedded in the nanocrystalline Si networks in the films grown with the lower SC. The second order transverse optical (2TO) mode of Si changes with decrease in SC as depicted in Fig. 2(d). The 2TO mode appears in the range 900 to 1000 cm−1 for all films [18]. Film S100 has no signal in this range. The intensity increases slowly for the 2TO mode as SC decreases in the films and ultimately the spectrum takes the flat top shape at very low SC. This analysis is in the similar line with that of the increase in the intensity of the first order Raman spectra of the films as SC decreases. Observing the last column of Table 2, it is seen that the crystalline fraction in the film increases from 41.87% to 49.8% as SC decreases. This confirms that the crystalline Si grain size increases with the decrease SC and can be correlated to the results obtained from the SEM images of all films. However, a small hump at about 490 cm−1 appears in film S140 as shown in Fig. 2(b), which probably indicates the presence of amorphous Si components in the film grown with excess H2 flow. The increment in the silicon grain size with decrement in the SC as confirmed by both, SEM and Raman analysis can be explained by means of the hydrogen induced annealing effect as the hydrogenated amorphous Si film being deposited is exposed to the excess hydrogen. The hydrogen induced annealing is dominant in HWCVD and assists the grain growth as the hydrogen content increases, unlike PECVD where this effect is countered by the ion-bombardment effect after certain amount of excess hydrogen [19]. Thus, as the SC decreases more hydrogen becomes available during the growth to induce the crystallization of amorphous silicon films even at lower substrate temperature and facilitates the grains to grow. Fig. 3 shows the PL measured as a function of emission wavelength from all films. As seen from this figure, a broad luminescence centered about 1000 nm is observed in all films. The broad asymmetric PL band indicates that the recombination takes place between multiple emission bands at room temperature. Since it is quite difficult to understand the origin of PL emission from the room temperature PL measurements, it is important to analyze the RTPL results in the context of the Raman and SEM observations. It is apparent from the Raman and SEM analysis that the film grown with the silane to hydrogen ratio 1:00 (film S100) is amorphous in nature. As the SC decreases, the films exhibit microcrystalline behavior. Thus, in order to investigate the roots of PL emission, the PL emission bands are de-convoluted in different Gaussian peaks for all the films as shown in Fig. 3. The peak wavelengths and FWHM obtained by Gaussian fits are listed in Table 3 for all films. As shown in Fig. 3(a), the film S100 emits a broad peak centered at 956 nm (~1.3 eV). As film S100 exhibits amorphous nature, the origin of this peak can be correlated to the structural disorder of the a-Si phase [20]. A strong peak with maximum intensity at 995 nm (~1.24 eV) observed in S100 could be originated from the radiative transitions between the band-tail states of the conduction band and valence band [21]. A feeble shoulder line is also observed at 1089 nm (~1.14 eV). This line is associated with the radiative transitions between tail states of the conduction and the mid-gap defect states due to Si dangling bonds [22]. Fig. 3(b) shows the PL emission from the film S105. As evident from the SEM and Raman analysis, the formation of crystalline grains begins in the film S105 and the crystalline fraction increases. The presence of crystalline grains in the films causes the PL peak to shift to the lower energy, which can clearly be observed in Fig. 3(b) in the form of peaks in the range from about 989 nm (~1.25 eV) to 1008 nm (~ 1.23 eV). These peaks, which are present in all films grown in the presence of hydrogen dilution, result from the shifted a-Si:H tail to tail transition in the presence of crystalline grains [23]. The PL peak at about 1080 nm (1.15 eV) originating from the radiative transitions between the tail states and the defect states of the Si dangling bonds is also clearly observed in all films as evident from Fig. 3 and the FWHM of this peak decreases with increasing size of the grains as a consequence of reduction in SC.
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Moreover, the presence of amorphous Si component indicated by the Raman analysis is confirmed by a distinct peak at about 951 nm (~ 1.3 eV), which appears in the S140. This blue shift of the peak in S140 grown with excess hydrogen flow arises due to the formation of amorphous Si tissues around the crystalline grains which is in agreement with the literature [24]. 4. Conclusions Nearly 300 nm Si thin-films were deposited with varying SC using HWCVD. Planar views of all films in SEM micrographs indicate an increase in the grain size with a reduction in SC. The first order Raman spectra of all films are in agreement with the SEM results. Interestingly, the second order Raman spectra in the range 900–1000 cm−1 also confirm the same results by the way of exhibiting an increased Raman intensity with a reduction in SC. PL analysis of all films except S100 shows the emission peak existence due to tail to tail transition between the conduction and valence bands as well as the emission due to the tail-state to the mid-gap defect state transition which clearly indicates the presence of microcrystalline phase in all the films. The experiments and the results presented here convey that the crystalline volume fraction and the grain size of the silicon thin films can be effectively controlled by varying the SC during the growth in HWCVD. The tuning of crystalline volume fraction and the grain size can be very important for the development of tandem and Si-multilayer based micromorph solar cells for the future PV technology since the deposition of multilayered high quality Si films with varying crystallinity can be accomplished in a single step. Acknowledgment Authors thank the Center for Nanoelectronics, IIT Bombay for the support to deposit Si thin-films using HWCVD and to characterize
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the films by SEM and PL under the Indian Nanoelectronics Users Programme (INUP). Authors are also grateful to Sophisticated Analytical Instruments Facility (SAIF), IIT Bombay for the Raman characterization of all films. References [1] V. Devabhaktuni, M. Alam, S.S.S.R. Depuru, R.C. Green II, D. Nims, C. Near, Renew. Sustain. Energy Rev. 19 (2013) 555. [2] L.M. Peter, Phil. Trans. R. Soc. A 369 (2011) 1840. [3] A. Goetzberger, C. Hebling, H.W. Schock, Mater. Sci. Eng. R 40 (2003) 1. [4] K. Shimakawa, J. Non-Cryst. Solids 266–269 (2000) 223. [5] D.L. Staebler, C.R. Wronski, Appl. Phys. Lett. 31 (1977) 292. [6] M. Zeman, R.E.I. Schorpp, Compr. Renew. Energy 1 (2012) 389. [7] A. Madhavan, D. Ghosh, M. Noack, V. Dalal, in: V. Chu, S. Miyazaki, A. Nathan, J. Yang, H.W. Zan (Eds.), Amorphous and Polycrystalline Thin-film Silicon Science and Technology—2007, San Francisco, CA, USA, 9–13 April 2007, Material Research Society Symposium Proceedings, 989, 2007, (A18-02). [8] K.L. Chopra, S.R. Das, Thin Film Solar Cells, Springer, New York, 1983. [9] Jean-Eric Bourée, in: Arvind Shah (Ed.), Thin Film Silicon Solar Cells, EPFL Press, Lausanne, Switzerland, 2010, p. 307. [10] Qi Wang, Thin Solid Films 517 (2009) 3570. [11] S. Klein, T. Repmann, T. Brammer, Sol. Energy 77 (2004) 893. [12] F. Villar, A. Antony, J. Escarré, D. Ibarz, R. Roldán, M. Stella, D. Muñoz, J.M. Asensi, J. Bertomeu, Thin Solid Films 517 (2009) 3575. [13] Guozhen Yue, J.D. Lorentzen, Jing Lin, Daxing Han, Qi Wang, Appl. Phys. Lett. 75 (1999) 492. [14] A. Blink, M. Brinza, J.P.H. Jongen, R.E.I. Schropp, Thin Solid Films 517 (2009) 3588. [15] R.O. Dusane, Thin Solid Films 519 (2011) 4555. [16] Friedhelm Finger, in: Arvind Shah (Ed.), Thin Film Silicon Solar Cells, EPFL Press, Lausanne, Switzerland, 2010, p. 121. [17] A.K. Panchal, C.S. Solanki, J. Cryst. Growth 311 (2009) 2659. [18] P. Mishra, K.P. Jain, Phys. Rev. B 64 (2001) 073304. [19] M.H. Gullanar, H. Chen, W.S. Wei, R.Q. Cui, W.Z. Shen, J. Appl. Phys. 95 (2004) 3961. [20] C.W. Tsang, R.S. Street, Phys. Rev. B 19 (1979) 3027. [21] P.C. Taylor, Braz. J. Phys. 23 (1993) 132. [22] T.K. Chini, D.P. Datta, S. Facsko, A. Mücklich, Appl. Phys. Lett. 92 (2008) 101919. [23] E.V. Johnson, S. Hoogland, E. Klem, N. Kherani, S. Zukotynski, J. Mater. Sci. Mater. Electron. 17 (2006) 789. [24] G.G. Qin, G.L. Kong, Philos. Mag. Lett. 57 (1988) 117.
Contents lists available at SciVerse ScienceDirect
Thin Solid Films journal homepage: www.elsevier.com/locate/tsf
Structural and optical properties of silicon thin-films deposited by hot-wire chemical vapor deposition: The effects of silane concentrations A.K. Panchal a, Vivek Beladiya b, Vipul Kheraj b,⁎ a b
Electrical Engineering Department, S.V. National Institute of Technology, Ichchhanath, Surat 395007, India Department of Applied Physics, S.V. National Institute of Technology, Ichchhanath, Surat 395007, India
a r t i c l e
i n f o
Article history: Received 20 September 2012 Received in revised form 24 June 2013 Accepted 27 June 2013 Available online 12 July 2013 Keywords: Silicon Thin-film Raman spectra Microcrystalline Grain
a b s t r a c t In this paper, the structural and optical properties of a series of silicon (Si) thin-films deposited using hot-wire chemical vapor deposition with different silane concentrations (SCs) are presented. All the films are characterized by Raman spectroscopy, scanning electron microscopy (SEM) and photoluminescence (PL). In the Raman analysis, the first order and specifically the second order Raman spectra indicate increase in crystalline grain size as well as crystalline volume fraction in the films with a reduction in SC, which is also confirmed by the SEM analysis. At the higher SC, the Si microcrystalline grains get embedded in the nanocrystalline Si network. The Gaussian fitted peaks in the PL analysis reveal the emission due to either band to band tail-state transitions or tail-state to mid-gap defect-state transitions due to Si-dangling bonds present in the films. © 2013 Elsevier B.V. All rights reserved.
1. Introduction The solar photovoltaics (PV) has been considered to be one of the most viable renewable energy technologies. It could be an important energy source to deal with the current scenario of the energy crisis as well as for the sustainable development [1]. From the first generation crystalline silicon (Si) cells to the third generation thin-film PV technologies, the efficiency and the cost-effectiveness have improved significantly. Especially, the developments of materials and improved growth techniques have caused a remarkable progress in the field of PV. However, Si is still one of the most important materials for the PV technologies with its total share of more than 87% in the current market [2]. With increasing penetration of amorphous and microcrystalline Si materials in the thin-film solar photovoltaics (PV) now-a-days, it is extremely important to understand the thin-film growth mechanism for these materials for further developments of the technology. The crystalline Si, due to its indirect bandgap, has relatively poor light-absorption properties. Moreover, the Si bandgap (~1.12 eV) is considerably lower from the optimum value of ~1.4 eV for the solar PV materials [3]. With the developments of thin-film amorphous and microcrystalline Si technology, it has become possible to engineer the Si bandgap. Amorphous and microcrystalline Si also exhibit better light-absorption properties than crystalline Si [4]. However, these materials have their own limitation in the form of high defect-densities, ⁎ Corresponding author. Tel.: +91 9904334220. E-mail address: [email protected] (V. Kheraj). 0040-6090/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.tsf.2013.06.098
which lead to the hindered carrier-transport in the material. Further, the light-induced degradation in these materials raise a great concern related to the device lifetime as well as performance reliability of the solar cells [5]. This trade off requires the combination of amorphous and crystalline Si to exploit the solar energy in the optimum manner. Quite a few schemes have been taken up to combine the two phases together in a PV cell. The recent advances in micro-morph tandem solar cells have promised the great deal of improvements in the maximum solar spectrum utilization [6]. Nearly 6% of the 27.2 GW power generated by solar PV in 2010 was contributed by thin film amorphous and microcrystalline Si PV and it is likely to be taken over in the near future by the tandem amorphous/microcrystalline Si modules [6]. Recently, Dalal et al. [7] have also reported the improved performance of Si-based cells with multilayered structure of alternating amorphous and microcrystalline layers. Thus, the study of Si as an absorber material and its optical properties with the variation in the crystal structure is of great interest. The optical and electrical behaviors of Si thin-film are very much dependent on the structural properties, which in turn rely up to a great extent on the growth conditions. It is very important to understand the growth mechanism of Si thin-films as the crystalline fraction of the material can be tuned with the deposition parameters. Various deposition techniques have been employed to grow the Si thin-films with varying crystalline structure [8]. Although plasma enhanced chemical vapor deposition (PECVD) is one of the widely used techniques for the deposition of amorphous, nanocrystalline and microcrystalline Si thin-films, the growth rate in PECVD is relatively low. Moreover, the lower precursor gas utilization efficiency in PECVD
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Table 1 Details of sample numbers and corresponding silane concentration (SC). Proportional flow SiH4:H2
Sample number
SC (%) (SiH4/(SiH4 + H2))
1:00 1:05 1:10 1:20 1:40
S100 S105 S110 S120 S140
100 16.7 9.1 4.8 2.4
results in higher manufacturing costs. In light of these limitations, the hot-wire chemical vapor deposition (HWCVD) is a promising technique for the deposition of Si thin-films with significantly high growth rate (up to 10 nm/s) and efficient precursor-gas utilization [9,10]. It is also observed that the open circuit voltage and the efficiency of the solar cells grown by HWCVD are higher than those grown by the PECVD [11]. Quite a few groups have employed HWCVD for the growth of amorphous as well as microcrystalline Si [10–12]. The transition from amorphous to microcrystalline phase in Si has also been investigated [13]. The feasibility of HWCVD for the deposition of Si on moving glass substrate for the roll to roll mass production has also been reported [14]. Moreover, Dusane has shown that the crystalline volume fraction of the film, grain orientation, grain size and the deposition rate are very
sensitive to the working gas concentrations in HWCVD [15]. Thus, it is possible to fine-tune the structural properties and consequently the optical and electrical properties of the films by controlling the precursor gas flow. Here, we report the study of growth of Si thin-films using HWCVD. The present work is an effort to understand the effects of precursors' concentration, mainly the silane concentration (SC) on the structural and optical properties of the films. We present a comprehensive analysis using first and second order Raman spectra in tandem with photoluminescence (PL) and scanning electron microscopy (SEM) analysis in order to study the properties of Si films and to investigate the influence of SC. The systematic data are also reported to understand the HWCVD growth parameters for the fine-tuning of properties of Si films in the perspective of solar PV applications. 2. Experimental details All films were deposited on RCA cleaned p-type crystalline Si wafer (one sided polished, 0.01–0.02 Ω-cm resistivity, 2 inch diameter, b100N orientation) pieces in HWCVD chamber. The films were deposited with a substrate temperature 200 °C. The silane (SiH4) and hydrogen (H2) gas mixture was cracked at the hot tungsten filament kept at 1900 °C. The distance between the filament and the substrate was maintained 5 cm during the process. Films with nearly 300 nm thickness were prepared with fixed SiH4 flow (1 sccm) and varying H2
Fig. 1. SEM micrographs of films: (a) S100, (b) S105, (c) S110, (d) S120 and (e) S140. All micrographs have a bar scale of 100 nm.
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Fig. 2. (a) Raman spectrum of the a-Si film S100, (b) first order Raman spectra of the films with decreasing SC, (c) Gaussian peaks fit in the first order Raman spectra of film S105 and (d) second order Raman spectra of the films with decreasing SC.
flow (0, 5, 10, 20 and 40 sccm) and the precursor gases were cracked for 30 min. The base pressure in the chamber was ~10−5 Pa before the deposition. The films are identified as film (letter-S) followed by fixed SiH4 flow (numeral-1) followed by H2 flow (numerals-00, 05, 10, 20 and 40) and are listed in Table 1. To investigate the structural morphology of all films, SEM and Raman spectroscopy characterizations were performed. SEM was done using Reith 150 model with beam energy 10 keV. Raman analysis was conducted using a Confocal Jobin-Yvon Raman Spectrometer with a 514.5 nm argon laser beam with a power intensity 18 mW and an aperture 100 μm. In order to study the light emission mechanisms from the amorphous and microcrystalline Si films, PL measurement Table 2 Details of Raman peaks and the crystalline fraction obtained from the Gaussian fit for all films. Films Peak 1
Peak 2
Raman FWHM Integrated Raman FWHM Integrated shift (cm−1) intensity (cm−1) intensity shift (cm−1) (cm−1) S100 S105 S110 S120 S140
480.0 519.9 519.9 519.9 519.9
– 3.6 3.9 3.7 4.4
– 793.4 1669.7 3053.6 8569.5
– 517.1 517.2 518.0 516.2
– 15.5 16.8 15.0 20.5
– 1101.2 1747.4 3189.6 8638.4
Crystalline fraction (%)
– 41.87 48.86 48.91 49.80
was performed with a focused 532 nm, 25 mW argon laser excitation source. The emitted light was collected by a lens and detected by an array detector mounted on a triple grating spectrometer. All the measurements were performed at room temperature. 3. Results and discussions Fig. 1(a)–(e) shows the surface morphology of all films as SC decreases. No Si grain growth is seen in Fig. 1(a) of film S100 as it is completely amorphous in nature. The grain growth starts appearing in film S105 (Fig. 1(b)) and the horizontal grain growth increases gradually in the films S110 to S140 as observed in Fig. 1(c)–(e). The measured average grain size is 44, 79 and 235 nm for the films S110, S120 and S140, respectively. SEM analysis reveals that the grain size increases as SC decreases in the films, which is in agreement with the literature [16]. To further understand this observation, films were subjected to the first and second order Raman spectroscopy. Fig. 2(a) shows a broad spectrum centered at the Raman stokes line at 480 cm−1, which corresponds to the transverse optical (TO) phonon mode of a-Si in film S100 [17]. This indicates that S100 has Si film which is completely amorphous in nature. Film S100 is deposited directly on a pure crystalline Si substrate, but no peak is detected at Raman line 520 cm−1 (corresponding to crystalline Si) in Fig. 2(a). Moreover, the estimated integrated intensity at 520 cm−1 in all other films increases with decrease in SC as seen in Table 2. These facts
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Fig. 3. Room temperature PL spectra with Gaussian fitted peaks for the films: (a) S100, (b) S105, (c) S110, (d) S120 and (e) S140.
clearly decline the possibility of detection of Raman signal from the crystalline Si substrate and confirm that the signals are detected only from the films. Fig. 2(b) illustrates the first order Raman spectra centered at the Raman stokes line 520 cm−1 for different films. The spectra have asymmetric shoulder in the lower Raman frequency range and are broader than that of the monocrystalline Si (nearly 2–3 cm−1 FWHM). It is also observed that the intensity increases with decrease in SC. Increasing intensity suggests the increased crystalline fraction in the films. To investigate the presence of different Si phases in the films, the Raman spectra are best fitted with Gaussian distribution function. Fig. 2(c) represents the Gaussian fitted Raman spectrum of film S105. A narrow peak centered at 519.9 cm−1 (with FWHM 3.6 cm−1) is corresponding to the crystalline phase of Si and a broad peak at 517.1 cm−1 (with FWHM 15.5 cm−1) is corresponding to the nanocrystalline phase of Si. Similarly, these two major peaks are also observed in all the films as seen in Table 2. The first peak is at 519.9 cm−1 (with FWHM ranging from 3.6 to 4.4 cm−1) and the second is in the range 516 to 518 cm−1 (with FWHM ranging from 15 to 20.5 cm−1). Thus, Raman study of the films confirms the presence of two phases of Si in all the films. This analysis of the first order Raman spectra helps to conclude that the crystalline Si grains
Table 3 Gaussian fitted PL peaks and corresponding FWHM for all films. Films Peak 1
S100 S105 S110 S120 S140
Peak 2
Peak 3
Wavelength (λ1)
FWHM
Wavelength (λ2)
FWHM Wavelength (λ3)
FWHM
956.53 989.82 990.37 990.05 951.57
172.716 104.00 101.18 95.02 66.63
995.92 1008.41 1009.54 1012.24 1004.07
88.50 57.81 56.06 52.59 74.09
50.538 68.00 65.22 65.03 58.23
1089.02 1080.54 1081.93 1076.38 1077.84
are embedded in the nanocrystalline Si networks in the films grown with the lower SC. The second order transverse optical (2TO) mode of Si changes with decrease in SC as depicted in Fig. 2(d). The 2TO mode appears in the range 900 to 1000 cm−1 for all films [18]. Film S100 has no signal in this range. The intensity increases slowly for the 2TO mode as SC decreases in the films and ultimately the spectrum takes the flat top shape at very low SC. This analysis is in the similar line with that of the increase in the intensity of the first order Raman spectra of the films as SC decreases. Observing the last column of Table 2, it is seen that the crystalline fraction in the film increases from 41.87% to 49.8% as SC decreases. This confirms that the crystalline Si grain size increases with the decrease SC and can be correlated to the results obtained from the SEM images of all films. However, a small hump at about 490 cm−1 appears in film S140 as shown in Fig. 2(b), which probably indicates the presence of amorphous Si components in the film grown with excess H2 flow. The increment in the silicon grain size with decrement in the SC as confirmed by both, SEM and Raman analysis can be explained by means of the hydrogen induced annealing effect as the hydrogenated amorphous Si film being deposited is exposed to the excess hydrogen. The hydrogen induced annealing is dominant in HWCVD and assists the grain growth as the hydrogen content increases, unlike PECVD where this effect is countered by the ion-bombardment effect after certain amount of excess hydrogen [19]. Thus, as the SC decreases more hydrogen becomes available during the growth to induce the crystallization of amorphous silicon films even at lower substrate temperature and facilitates the grains to grow. Fig. 3 shows the PL measured as a function of emission wavelength from all films. As seen from this figure, a broad luminescence centered about 1000 nm is observed in all films. The broad asymmetric PL band indicates that the recombination takes place between multiple emission bands at room temperature. Since it is quite difficult to understand the origin of PL emission from the room temperature PL measurements, it is important to analyze the RTPL results in the context of the Raman and SEM observations. It is apparent from the Raman and SEM analysis that the film grown with the silane to hydrogen ratio 1:00 (film S100) is amorphous in nature. As the SC decreases, the films exhibit microcrystalline behavior. Thus, in order to investigate the roots of PL emission, the PL emission bands are de-convoluted in different Gaussian peaks for all the films as shown in Fig. 3. The peak wavelengths and FWHM obtained by Gaussian fits are listed in Table 3 for all films. As shown in Fig. 3(a), the film S100 emits a broad peak centered at 956 nm (~1.3 eV). As film S100 exhibits amorphous nature, the origin of this peak can be correlated to the structural disorder of the a-Si phase [20]. A strong peak with maximum intensity at 995 nm (~1.24 eV) observed in S100 could be originated from the radiative transitions between the band-tail states of the conduction band and valence band [21]. A feeble shoulder line is also observed at 1089 nm (~1.14 eV). This line is associated with the radiative transitions between tail states of the conduction and the mid-gap defect states due to Si dangling bonds [22]. Fig. 3(b) shows the PL emission from the film S105. As evident from the SEM and Raman analysis, the formation of crystalline grains begins in the film S105 and the crystalline fraction increases. The presence of crystalline grains in the films causes the PL peak to shift to the lower energy, which can clearly be observed in Fig. 3(b) in the form of peaks in the range from about 989 nm (~1.25 eV) to 1008 nm (~ 1.23 eV). These peaks, which are present in all films grown in the presence of hydrogen dilution, result from the shifted a-Si:H tail to tail transition in the presence of crystalline grains [23]. The PL peak at about 1080 nm (1.15 eV) originating from the radiative transitions between the tail states and the defect states of the Si dangling bonds is also clearly observed in all films as evident from Fig. 3 and the FWHM of this peak decreases with increasing size of the grains as a consequence of reduction in SC.
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Moreover, the presence of amorphous Si component indicated by the Raman analysis is confirmed by a distinct peak at about 951 nm (~ 1.3 eV), which appears in the S140. This blue shift of the peak in S140 grown with excess hydrogen flow arises due to the formation of amorphous Si tissues around the crystalline grains which is in agreement with the literature [24]. 4. Conclusions Nearly 300 nm Si thin-films were deposited with varying SC using HWCVD. Planar views of all films in SEM micrographs indicate an increase in the grain size with a reduction in SC. The first order Raman spectra of all films are in agreement with the SEM results. Interestingly, the second order Raman spectra in the range 900–1000 cm−1 also confirm the same results by the way of exhibiting an increased Raman intensity with a reduction in SC. PL analysis of all films except S100 shows the emission peak existence due to tail to tail transition between the conduction and valence bands as well as the emission due to the tail-state to the mid-gap defect state transition which clearly indicates the presence of microcrystalline phase in all the films. The experiments and the results presented here convey that the crystalline volume fraction and the grain size of the silicon thin films can be effectively controlled by varying the SC during the growth in HWCVD. The tuning of crystalline volume fraction and the grain size can be very important for the development of tandem and Si-multilayer based micromorph solar cells for the future PV technology since the deposition of multilayered high quality Si films with varying crystallinity can be accomplished in a single step. Acknowledgment Authors thank the Center for Nanoelectronics, IIT Bombay for the support to deposit Si thin-films using HWCVD and to characterize
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the films by SEM and PL under the Indian Nanoelectronics Users Programme (INUP). Authors are also grateful to Sophisticated Analytical Instruments Facility (SAIF), IIT Bombay for the Raman characterization of all films. References [1] V. Devabhaktuni, M. Alam, S.S.S.R. Depuru, R.C. Green II, D. Nims, C. Near, Renew. Sustain. Energy Rev. 19 (2013) 555. [2] L.M. Peter, Phil. Trans. R. Soc. A 369 (2011) 1840. [3] A. Goetzberger, C. Hebling, H.W. Schock, Mater. Sci. Eng. R 40 (2003) 1. [4] K. Shimakawa, J. Non-Cryst. Solids 266–269 (2000) 223. [5] D.L. Staebler, C.R. Wronski, Appl. Phys. Lett. 31 (1977) 292. [6] M. Zeman, R.E.I. Schorpp, Compr. Renew. Energy 1 (2012) 389. [7] A. Madhavan, D. Ghosh, M. Noack, V. Dalal, in: V. Chu, S. Miyazaki, A. Nathan, J. Yang, H.W. Zan (Eds.), Amorphous and Polycrystalline Thin-film Silicon Science and Technology—2007, San Francisco, CA, USA, 9–13 April 2007, Material Research Society Symposium Proceedings, 989, 2007, (A18-02). [8] K.L. Chopra, S.R. Das, Thin Film Solar Cells, Springer, New York, 1983. [9] Jean-Eric Bourée, in: Arvind Shah (Ed.), Thin Film Silicon Solar Cells, EPFL Press, Lausanne, Switzerland, 2010, p. 307. [10] Qi Wang, Thin Solid Films 517 (2009) 3570. [11] S. Klein, T. Repmann, T. Brammer, Sol. Energy 77 (2004) 893. [12] F. Villar, A. Antony, J. Escarré, D. Ibarz, R. Roldán, M. Stella, D. Muñoz, J.M. Asensi, J. Bertomeu, Thin Solid Films 517 (2009) 3575. [13] Guozhen Yue, J.D. Lorentzen, Jing Lin, Daxing Han, Qi Wang, Appl. Phys. Lett. 75 (1999) 492. [14] A. Blink, M. Brinza, J.P.H. Jongen, R.E.I. Schropp, Thin Solid Films 517 (2009) 3588. [15] R.O. Dusane, Thin Solid Films 519 (2011) 4555. [16] Friedhelm Finger, in: Arvind Shah (Ed.), Thin Film Silicon Solar Cells, EPFL Press, Lausanne, Switzerland, 2010, p. 121. [17] A.K. Panchal, C.S. Solanki, J. Cryst. Growth 311 (2009) 2659. [18] P. Mishra, K.P. Jain, Phys. Rev. B 64 (2001) 073304. [19] M.H. Gullanar, H. Chen, W.S. Wei, R.Q. Cui, W.Z. Shen, J. Appl. Phys. 95 (2004) 3961. [20] C.W. Tsang, R.S. Street, Phys. Rev. B 19 (1979) 3027. [21] P.C. Taylor, Braz. J. Phys. 23 (1993) 132. [22] T.K. Chini, D.P. Datta, S. Facsko, A. Mücklich, Appl. Phys. Lett. 92 (2008) 101919. [23] E.V. Johnson, S. Hoogland, E. Klem, N. Kherani, S. Zukotynski, J. Mater. Sci. Mater. Electron. 17 (2006) 789. [24] G.G. Qin, G.L. Kong, Philos. Mag. Lett. 57 (1988) 117.