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Structural and optoelectronic characteristics of nanocrystalline silicon oxide film as absorber layer for thin film solar cells
Accepted Manuscript Structural and optoelectronic characteristics of nanocrystalline silicon oxide film as absorber layer for thin film solar cells Haixu Liu, Yanbin Yang, Jianping Liu, Zhaoyi Jiang, Yun Li, Wei Yu, Wenge Ding, Guangsheng Fu PII:
S0925-8388(16)30311-5
DOI:
10.1016/j.jallcom.2016.02.042
Reference:
JALCOM 36653
To appear in:
Journal of Alloys and Compounds
Received Date: 14 December 2015 Revised Date:
4 February 2016
Accepted Date: 5 February 2016
Please cite this article as: H. Liu, Y. Yang, J. Liu, Z. Jiang, Y. Li, W. Yu, W. Ding, G. Fu, Structural and optoelectronic characteristics of nanocrystalline silicon oxide film as absorber layer for thin film solar cells, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.02.042. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
ACCEPTED MANUSCRIPT Structural and optoelectronic characteristics of nanocrystalline silicon oxide film as absorber layer for thin film solar cells Haixu Liu, Yanbin Yang, Jianping Liu, Zhaoyi Jiang, Yun Li, Wei Yu*, Wenge Ding
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and Guangsheng Fu Hebei Key Laboratory of Optic-Electronic Information Material, College of Physics Science and Technology, Hebei University, Baoding 071002, China
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The mixed phase nanocrystalline silicon oxide (nc-SiOx:H) thin films have been
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prepared by very high frequency plasma enhanced chemical vapor deposition (VHF-PECVD) system from SiH4, H2 and CO2 mixture. The effect of oxygen incorporation on the structural and opto-electrical properties of nc-Si:H films was investigated. As the CO2 flow rate increased, the crystalline volume fraction in the
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films decreased, while the optical band gap increased slowly. During the phase transition, both oxygen and hydrogen content increased drastically with the increasing CO2 flow rate, while the Urbatch energy and microstructure factor show obvious
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reduction, which indicates an enhancement of the medium range order of the films
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and good passivation of the inner voids. Device quality nc-SiOx:H film was fabricated with low microstructure factor of 0.32 and high crystalline volume fraction around 50%, along with wide optical band gap of ~1.75 eV, low dark-conductivity σd ~7.6×10-6 S/cm and high photosensitivity of about 2×102. On flat substrate a solar cell was prepared with high Voc of 0.62V and Jsc of 14.57mA/cm2. These results reveal the
*
Corresponding author. Tel.: +86 0312 5079560; fax: +86 0312 5079560.
E-mail address: [email protected].
ACCEPTED MANUSCRIPT potential of nc-SiOx:H film as an absorber layer for thin film solar cells. Key Words: nanocrystalline silicon oxide; phase transition; absorber layer; thin film
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solar cells.
ACCEPTED MANUSCRIPT 1. Introduction Micromorph tandem solar cells with hydrogenated amorphous silicon (a-Si:H) and hydrogenated microcrystalline silicon (µc-Si:H) absorber layers have been intensively studied for its potential of making high efficiency thin film silicon solar
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cells. Despite significant progresses in micromorph solar cells [1], the short-circuit current density (Jsc) of the entire cell is limited by the a-Si:H top cell. Single junction a-Si solar cell has Jsc over 18mA/cm2, however in tandem devices top cell thickness is
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reduced to match current, which makes current smaller [2]. To increase the current of tandem device for higher conversion efficiency, thick top cell is needed. Using a-Si:H
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as top cell can not reach this goal and thick a-Si:H top cell has more Staebler-Wronski effect [3].
In recent years, nanocrystalline silicon oxide (nc-SiOx:H) films have proven to be a novel material in photovoltaic applications for its tunable optical and electrical
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properties [4-7]. The optical band gap of nc-SiOx:H can be adjusted by changing the oxygen concentration, while the contained nanocrystalline silicon phase will improve the carrier mobility in the film [8]. We believe nc-SiOx:H is a promising candidate for
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the absorber layer of the top cell, considering its wide optical band gap and high
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crystalline volume fraction. Moreover, nc-SiOx:H film is more stable under incident light than a-Si:H due to the improved medium range order and increased carrier lifetime [9]. However, the oxygen incorporation may result in defects and disordered structure in the film hence the oxygen concentration should be well controlled [10]. Till now, there are hardly any reports on the nc-SiOx:H material as the absorber layer for solar cells, compared with the intensively studied a-SiOx:H thin film solar cells [11-13].
ACCEPTED MANUSCRIPT This work aims to discuss the potential use of nc-SiOx:H film as the absorber layer in thin film solar cells, which may facilitate the fabrication of novel tandem solar cells. Towards this goal, the influences of oxygen incorporation on the structural and optoelectronic properties of nc-Si:H films were thoroughly investigated. Device
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quality nc-SiOx:H material with high conductivity and wide band gap has been fabricated with optimal oxygen incorporation near the phase transition. With low σd of 7.6×10-6 S/cm and high photosensitivity of 2×102, nc-SiOx:H material demonstrates
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its importance for high efficiency silicon-based thin film photovoltaics.
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2. Experiment
Hydrogenated silicon oxide thin films were deposited by VHF-PECVD (multi-chamber system) from a gas mixture of silane (SiH4), hydrogen (H2) and carbon dioxide (CO2) at 60MHz, 200 , 60 Pa. The plasma excitation power density
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was 45mW/cm2. The gas flow rates of SiH4 and H2 were kept constant while the H2 diluted CO2 was oxygen source. Here we defined the parameter Ro as the CO2/SiH4
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gas flow ratio. Material characterization was performed on 400nm thick SiOx layers deposited on glass substrates and silicon wafers (100).
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Solar cells with p-i-n configuration were prepared on AZO (ZnO:Al) coated glass substrates. The p-type nc-SiOx doped layer was deposited from H2, SiH4，CO2 and Trimethylboron (TMB) mixture reported in our previous work [14]. The nc-SiOx:H solar cell fabricated in this work has the following structure: AZO glass /p-nc-SiOx:H (20nm)/buffer layer (4nm)/i-nc-SiOx:H (1µm)/n-a-Si:H (30nm)/ITO/Ag. The thickness of the film was measured by a step profiler (Dektak 150). The structural properties were characterized by Raman spectroscopy (LabRAM-HR,
ACCEPTED MANUSCRIPT 532nm). High resolution transmission electron microscopy (HRTEM) observation of the sample was performed with FEIF20 transmission electron microscope. The TEM sample was scraped from the substrate, triturated into powder, and then sonically
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dispersed in ethanol. A copper TEM microgrid was immersed into the ethanol to collect the fragment of the sample. FTIR spectra were measured to investigate the bonding configuration of the as-deposited films (Bruker Tensor27). Absorption
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spectra of the layers were recorded by UV-VIS-NIR spectrometer (Hitachi UV-4100).
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The electrical conductivity was characterized by current density versus voltage (J-V) measurements. The photoconductivity and the solar cell performance were measured by B1500A semiconductor device parameter analyzer under AM1.5 solar simulator. 3. Results
To characterize the structural properties of the materials, normalized Raman
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spectra of nc-SiOx:H thin films deposited with different Ro are shown in figure 1(a). Firstly, an oxygen free nc-Si:H sample was deposited as a starting point, whose
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Raman peak locates mainly at 520 cm-1, indicating nanocrystalline phase was dominant in the film. With the increase of Ro from 0% to 14%, a broad peak around
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480 cm-1 related to amorphous phase became dominant. Moreover, an obvious blue shift of peak position could be observed as the miorostructure changed. In order to describe the evolution of the microstructure quantitatively, crystalline
volume fraction XC can be calculated from the following formula: X C = (I 510 + I 520 ) (I 480 + I 510 + I 520 )
(1)
where Ii denotes the integrated intensity of the individual Gaussian peak centered at i cm-1 [15]. The XC as a function of Ro is shown in figure 1(b). The crystalline volume
ACCEPTED MANUSCRIPT fraction XC decreases from ~70% to 8% as Ro increases from 0 to 14%. These results suggest that the incorporation suppresses the formation of crystallites resulting in the phase transition from crystalline to amorphous. It is obvious that there is a transition
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to amorphous growth regime at Ro of ~9-12%. HRTEM characterization was carried out to get detailed structural information of the deposited material. Figure 2 shows a HRTEM image of a nc-SiOx:H sample
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deposited with Ro = 8%, which has a typical mixed-phase structure with the silicon
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nanocrystals (highlighted in circles) embedded in a amorphous matrix. It can be observed that the diameters of the silicon nanocrystals cluster range from 3 nm to 7 nm and the average size of silicon nanocrystals is about 5.3 nm. The inset is a zoom in image of a silicon nanocrystal cluster in the film. It is clear that in the cluster there is
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order organize structure.
FTIR spectrum of the nc-SiOx:H layers are shown in figure 3. From figure 3, one can see that, with the increase of Ro, the peak of Si-O stretching mode (900
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cm-1~1300 cm-1) slightly shifts to higher wavenumber (as denoted by the dashed line).
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This shift of the peak of Si-O stretching mode can be attributed to the oxygen back-bonding or the variation of the Si-O-Si bond angles [16]. The peak around 630cm-1 related to Si-H has no visible variation as Ro increases, because it is silicon rich film.
The bonded hydrogen content (CH) and oxygen content (CO) can be derived from the Si-H peak around 630cm-1 and the Si-O peak (900-1300cm-1), respectively. The calculations are performed with a numerical integration formula:
ACCEPTED MANUSCRIPT C (at %) = ( Aω N Si )∫ [α (ω ) / ω ]dω
(2)
where α(ω) is the absorption coefficient, ω is the wavenumber, NSi = 5×1022 cm-2 is the atomic density of the crystalline silicon, and Aω is the proportionality constant for
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hydrogen (1.6×1019 cm-2) and oxygen (9.5×1018 cm-2) [17]. Figure 4(a) and (b) show the calculated results of CO and CH of nc-SiOx:H films versus Ro. With the increase of Ro from 0 to 9%, the oxygen content CO increases
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gradually from 0% to 4.8%,which indicates an enhancement of incorporation of
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oxygen in the films. Oxygen concentration then declines slightly from 4.8% to 4.3% with further increase of Ro to 14%. The tendency of CH is quite similar to that of CO. With increasing Ro, the CH once increases significantly from 10.7% to 17.9%, and then decreases slightly from 17.9% to 15%. This drastic changes of CH can also be
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found in the transition from nc-Si:H to a-Si:H [18].
The Si-H stretching mode (1900 cm-1~2200 cm-1) includes two vibration peaks related to SiH and SiH2 absorption centered at 2000 cm-1 and 2100 cm-1, respectively.
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The SiH absorption peak corresponds to the amorphous Si phase of compact
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microstructure while SiH2 absorption peak can be attributed to the Si-H bonds located on the surface of inner voids or grain boundaries [19]. To evaluate the quality of the deposited films, the microstructure factor is calculated by the formula: R* = I 2100 /( I 2000 + I 2100 )
(3)
where I2000 and I2100 are the integration intensity of the absorption peaks corresponding to SiH and SiH2, respectively. Here I2000 and I2100 are obtained by Gaussian fitting [20]. The R* is used to evaluate the quality of the deposited films and
ACCEPTED MANUSCRIPT lower indicates better film quality [19]. As shown in figure 4(c), as Ro inceases, R* deceases first and then inceases, and at Ro =8% the R* is lowest. The quality of the film is thus improved with oxygen
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concentration up to 3.2%. Higher oxygen concentration will deteriorate the film quality by introducing more recombination sites as well as amorphous phase becomes dominant. It is well known that the Si-H stretching mode (2000 cm-1) may be shifted
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to higher energy by oxygen back-bonding, which is the so-called induction effect [21].
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It means that the peak at 2100 cm-1 can be attributed to both SiH2 at surfaces of inner voids and oxygen back-bonding of SiH in a high quality film. Considering this fact, the microstructure factor may be even smaller than the calculated results. The optical properties are analyzed by UV-VIS-NIR absorption spectra. The
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optical band gap Eg of the nc-SiOx:H layer can be estimated by the equation:
(αhν ) 12 = B(hν − E g )
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known as Tauc formulation, where α is the absorption coefficient, hν is the energy of
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[22,23].
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incident light, Eg is the optical band gap, and B is a constant related to disorder degree
To characterize the width of band-tail states and the disorder degree in the
amorphous network, the Urbach energy EU can be described as: hν EU
α ≈ exp
(5)
which can be extrapolated from the semilogarithm plot of α versus hν [23]. Figure 5(a) shows the plot of (αhν)1/2 (Ro=8% and 14%) versus photon energy, from which Eg can be extrapolated from a straight line to hν = 0. The solid lines show
ACCEPTED MANUSCRIPT the linear fit of the straight part. The calculated Eg and EU are plotted in figure 5(b). With the increase of Ro, the values of Eg show a general increasing trend with a fluctuation. It is also notable that the band gap is approximately 1.75 eV with only 3.2
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at% oxygen concentration (Ro = 8%). Such wide band gap is comparable to the Eg of a-Si:H and will be favourable for the application as i-layer instead of the a-Si:H layer in the top cells. With the increase of Ro, the Urbach energy EU first decreases and then
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increases. The minimum of EU is also obtained with Ro = 8%, which indicates the
ordered microstructure in the film.
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reduction of the localized band-tail states (particularly the valence band-tail) and more
Figure 6 (a) and (b) show the σd and σph, respectively. The photosensitivity (σph/σd) versus Ro is shown in figure 6 (c). It is obviously that both σd and σph
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decrease with the increasing Ro. The incorporation of oxygen reduces the crystallinity of the nc-Si:H film, which may account for the degradation of σd and σph. However, since σd decreases more rapidly than σph, the photosensitivity increases with Ro from
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10 to 103. Low σd of ~7.6×10-6 S/cm and high photosensitivity of 2×102 can be
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obtained at Ro = 8%, which is near the phase boundary. Similar reports about device-grade µc-Si:H has also been obtained near the phase transition [24]. To reveal the potential of intrinsic nc-SiOx:H films as the absorber layers, a thin
film solar cell was fabricated with a configuration of glass substrate/flat ZnO/p-nc-SiOx:H/buffer layer/i-nc-SiOx:H/n-a-Si:H/ITO/Ag. Figure 7 shows the J-V curves of single junction solar cell containing the nc-SiOx:H i-layer deposited with Ro = 8%. An initial efficiency of about 5% is obtained with Voc = 0.62V, Jsc =
ACCEPTED MANUSCRIPT 14.57mA/cm-1 and FF = 0.533. It should be noted that no textured substrate is adopted here, and the p/i interface needs further optimization. Higher conversion efficiency will be achieved with improvement of the process parameters and the
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structure of the cell.
4. Discussion
To discuss the structural evolution caused by oxygen incorporation, the Si-O
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stretching vibration mode in the FTIR spectra is fitted by several Gaussian peaks.
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Figure 8(a) shows the fitted result of a nc-SiOx:H layer deposited with Ro = 14%. The peaks centered at 980, 1012, 1034, and 1076 cm-1 correspond to the stretching vibration modes of HSi(Si2O), HSi(SiO2), HSi(O3) and Si(O4), respectively [25]. The absorption peak at 1150 cm-1 is attributed to the vibration of the oxygen in the Si(O4)
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network [21,25].
Here we define the fraction of silicon-rich vibration modes and the fraction of oxygen-rich vibration modes as:
(6)
X O − rich = (I1034 + I1076 ) (I 980 + I1012 + I 1034 + I 1076 )
(7)
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X Si −rich = (I 980 + I1012 ) (I 980 + I1012 + I1034 + I1076 )
where Ii denotes the integrated absorption intensity at the i cm-1 [26]. The nc-SiOx:H film actually contains different phases, including the nanocrystalline silicon (nc-Si), hydrogenated amorphous silicon (a-Si:H) and oxygen-rich hydrogenated amorphous silicon oxide (a-SiOx:H) [27]. It is thus reasonable that XSi-rich and XO-rich are related to the a-Si:H phase and the oxygen-rich a-SiOx:H phase in the deposited films, respectively.
ACCEPTED MANUSCRIPT Figure 8(b) shows the XSi-rich and XO-rich versus Ro, which may facilitate the description of the structural evolution. With the increase of Ro from 0 to 8%, the XSi-rich decreases while XO-rich becomes larger. At Ro = 8%, XSi-rich shows a minimum
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while XO-rich reaches its maximum, which is an indication of the Si/SiOx phase separation. As shown in figure 4(c) and 5(b), the microstructure factor R* and the Urbach energy EU are also greatly reduced at Ro = 8%. The phase separation may give
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rise to the enhancement of the medium range order of the film and good surface
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passivation of the grain boundaries.
The broadening of optical band gap Eg can be mainly attributed to oxygen incorporation into the samples. According to the results of mixed-phase nc-Si:H material reported by Yan et al, quantum confinement effect occurs most likely in
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isolated silicon nanocrystals with small grain sizes [28]. In most of the samples deposited here (with XC ≥ 30%), the XC is high enough to form conduction path for carriers and the quantum confinement effect may not have the dominant effect on the
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optical band gap. The broadening of Eg is mostly associated with the increasing
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a-SiOx:H phase in the deposited film. The oxygen induced band gap broadening can be related to the fact that Si-O bond energy (8.4 eV) is stronger than the Si-Si bond (2.4 eV) and the Si-H bond (3.0 eV) [29]. In addition, the oxygen back bonded to Si-Si bond increases the Si-Si bond energy, which results in the increase in the energy difference between the bonding and anti-bonding splitting of the Si-Si bonds. As Eg is a weighted average of the conduction and valence band states (EC and EV), the oxygen back-bonding will lead to the widening of Eg [30].
ACCEPTED MANUSCRIPT To explain the increase of the photosensitivity near the phase transition region, the dark-conductivity σd, photo-conductivity σph and photosensitivity are described as following formulas:
σ ph = e(N d + Gτ )µ ∝ µτ
(8) (9)
τ
E − EF exp − C K BT
(10)
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σ ph e( N d + Gτ ) µ = ∝ σd eN d µ
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EC − EF K BT
σ d = eN d µ ∝ τ exp −
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where Nd is the carrier concentration under thermal equilibrium condition, µ is the carrier mobility, G is the carrier generation rate, τ is the carrier lifetime, EC is the conduction band edge, EF is the Fermi level position, KB is Boltzmann constant and T is the temperature [11]. As oxygen concentration increases, the carrier lifetime τ
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decreases as a consequence of the reduction of the crystalline volume fraction XC and larger scattering probability in the amorphous network, which results in the decreases
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of σd and σph. The increase of photosensitivity thus can be attributed to the rising value of (EC－EF), which can be related to the increase of the optical band gap shown
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in figure 5(b). In general, the Fermi level EF remains almost unchanged in the a-SiOx:H films [10]. However, in this work, the drastic increase of photosensitivity may suggest that the EF is lowered near the phase transition. As known from the former discussion, the localized electronic states on the grain boundaries and valence band-tail states are significantly reduced in the films, which have been demonstrated by the former results of R* and EU in figure 4(c) and figure 5(b). It is thus reasonable that the shift of EF is a consequence of the reduction of R* and EU. .
ACCEPTED MANUSCRIPT 5. Conclusions In summary, the influence of oxygen incorporation on the structural, optical and electrical properties of nc-Si:H layer deposited by VHF-PECVD is investigated. The
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incorporation of oxygen gives rise to band gap broadening and inhibition of the crystalline Si growth, as the nc-Si:H phase is gradually substituted by the a-SiOx:H phase in the films. It is found that the nc-SiOx:H material with optimal optoelectronic
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properties can be obtained near the phase transition. Device quality nc-SiOx:H layer
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has been prepared with wide optical band gap of ~1.75 eV, high crystalline volume fraction of 50% and low oxygen concentration of ~3.2at%, along with low microstructure factor of 0.32, low σd ~7.6×10-6 S/cm and high photosensitivity of 2×102. This material deserves enormous promise for applications in nc-SiOx:H solar
solar cells.
Acknowledgments
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cells, especially for preparing high efficiency and stable top cell in thin film tandem
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This work is financially supported by the Specialized Research Fund for the
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Doctoral Program of Higher Education, PRC (Grant No. 20131301120003), the Science and Technology Planning Project of Hebei Province, PRC (Grant No. 13214315) and the Colleges and universities in Hebei province science and technology research youth fund, PRC (Grant No. QN20131115).
References [1] J. S. Cashmore et al, Sol. Energy Mater. Solar Cells 144 (2016) 84. [2] K Tabuchi, W. Wenas, A.Yamada, M Konagai, K. Takahashi, Jpn. J. Appl. Phys.
ACCEPTED MANUSCRIPT 32 (1993) 3764. [3] B. J. Yan, G. Z. Yue, L. Sivec, J. Yang, S. Guha, C. S. Jiang, Appl. Phys. Lett. 99 (2011) 113512.
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[4] P. Cuony, D. T. Alexander, I. Perez-Wurfl, M. Despeisse, G. Bugnon, M. Boccard, T. Soderstrom , A. Hessler-Wyser, C. Hebert, C. Ballif, Adv. Mater. 24 (2012) 1182.
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[5] P. D. Veneri, L. V. Mercaldo, I. Usatii, Prog. Photovolt: Res. Appl. 21(2013) 148.
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[6] P. Cuony, M. Marending, D. T. L. Alexander, M. Boccard, G. Bugnon, M. Despeisse, C. Ballif, Appl. Phys. Lett. 97 (2010) 213502.
[7] A. Lambertz, V. Smirnov, T. Merdzhanova, K. Ding, S. Haas, G. Jost, R. E. I. Schropp, F. Finger, U. Rau, Sol. Energy Mater. Solar Cells 119 (2013) 134.
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[8] V. Smirnov, W. Böttler, A. Lambertz, H. Y. Wang, R. Carius, F. Finger, phys. status solidi C 7 (2010) 1053.
[9] S. B. Zhang, X. B. Liao, L. An, F. H. Yang, G. L. Kong, Y. Q. Wang, Y. Y. Xu, C.
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Y. Chen, H. W. Diao, Acta Physica Sinica 51 (2002) 1811.
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[10] J. E. Lee, J. H. Park, J. Yoo, K. H. Yoon, D. Kim, J. S. Cho, Solid State Sci. 20 (2013) 70.
[11] S. Inthisang, K. Sriprapha, S. Miyajima, A. Yamada, M. Konagai, Jpn. J. Appl. Phys. 48 (2009) 122402.
[12] D. Y. Kim, E. Guijt, R. van Swaaij, M. Zeman, Prog. Photovolt.: Res. Appl. 23 (2015) 671. [13] D. Y. Kim, E. Guijt and F. T. Si, R. Santbergen, J. Holovský, O. Isabella, R. van
ACCEPTED MANUSCRIPT Swaaij, M. Zeman, Sol. Energy Mater. Solar Cells 141 (2015) 148. [14] Z. L. Shi, Y. Ji, W. Yu , Y. B. Yang, R. D. Cong, Y. J. Chen, X. W. Li, S. F. Guang,
Chin. Phys. B 24 (2015) 078105.
Phys. Lett. 90 (2007) 263509.
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[15] S. Y. Myong, K. Sriprapha, S. Miyajima, M. Konagai and A. Yamada, Appl.
[16] F. Einsele , W. Beyer, U. Rau, J. Appl. Phys. 112 (2012) 054905.
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[17] P. H. Gaskell, D. J. Wallis, Phys. Rev. Lett. 76 (1996) 66.
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[18] A. Shah, J. Meier, E. Vallat-Sauvain, N. Wyrsch, U. Kroll, S. Droz, U. Graf, Sol. Energy Mater. Solar Cells 78 (2003) 469.
[19] Y. Adjallah, C. Anderson, U. Kortshagen, J. Kakalios, J. Appl. Phys. 107 (2010) 043704.
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[20] D. X. Han, K. D. Wang, J. M. Owens, L. N. Gedvilas, H. Habuchi and M. Tanaka, J. Appl. Phys. 93 (2003) 3776. [21] D. V. Tsu, G. Lucovsky, B. N. Davidson, Phys. Rev. B 40 (1989) 1795.
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[22] Q. J. Cheng, S. Y. Xu, K. Ostrikov, Nanotechnology 20 (2009) 215606.
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[23] T. Li , J. Kanicki, W. Kong, F. L. Terry, J. Appl. Phys. 88 (2000) 5764. [24] S. Y. Myong, K. Sriprapha, Y. Yashiki, S. Miyajima, A. Yamada, M. Konagai, Sol. Energy Mater. Solar Cells 92 (2008) 639.
[25] L. He, T. Inokuma , Y. Kurata, S. Hasegawa, J. Non-Crystal. Solids 185 (1995) 249. [26] H. M. Li, W. Yu, Y. M. Xu, Y. Ji, Z. Y. Jiang, X. Z. Wang, X. W. Li, G. S. Fu, Chin. Phys. B 24 (2015) 028102.
ACCEPTED MANUSCRIPT [27] A. Richter, L. Zhao, F. Finger, K. N. Ding, Surf. Coat. Technol. (2015) In Press. [28] B. J. Yan, G. Z. Yue, J. Yang, S. Guha, Sol. Energy Mater. Solar Cells 111 (2013) 90.
Barua, Mater. Chem. Phys. 157 (2015) 130.
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[29] S. Mandal, G. Das, S. Dhar, R. M. Tomy, S. Mukhopadhyay, C. Banerjee, A. K.
[30] I. Umezu, K. Miyamoto, N. Sakamoto, K. Maeda, Jpn. J. Appl. Phys. 34 (1995)
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Figure captions: Fig.1 (a) Raman spectra of nc-SiOx:H layers grown with different Ro. (b) The crystalline volume fraction and growth rate of nc-SiOx:H layers versus Ro.
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Fig.2 HRTEM image of nc-SiOx:H layer deposited with Ro = 8%, and the inset shows
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a magnified image of a typical silicon nanocrystal embedded in silicon oxide matrix. Fig.3 FTIR spectra of nc-SiOx:H layers deposited with different Ro. Fig.4 (a) The oxygen content, (b) the hydrogen content and (c) the microstructure factor plotted as functions of Ro.
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Fig.5 (a) (αhν)1/2 (Ro=8% and 14%) versus the photon energy. (b) The calculated optical band gap Eg and Urbach energy EU versus Ro. Fig.6 (a) The dark-conductivity, (b) photo-conductivity and (c) photosensitivity of
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nc-SiOx:H films plotted as a function of Ro.
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Fig.7 J-V characteristics of the nc-SiOx:H solar cells with the intrinsic layer deposited with Ro = 8%.
Fig.8 (a) Gaussian fitting results of the Si-O stretching mode of a nc-SiOx:H film (Ro = 14%). (b) Plots of XSi-rich and XO-rich in the nc-SiOx:H films versus Ro.
(a )
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3
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1 4 %
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1 2 % 1 0 % 8 %
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2
5 %
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N o r m a liz e d In te n s ity (a .u .)
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4 0 0
4 5 0
R 5 0 0
R a m a n s h ift (c m
5 5 0 -1
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O
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ACCEPTED MANUSCRIPT Reaserch Highlights
► The structural and opto-electrical properties of nc-SiOx:H layer is investigated. ► Device quality nc-SiOx:H material is obtained near the phase transition region.
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► A nontextured solar cell was prepared with Voc of 0.62V and Jsc of
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14.57mA/cm2.
S0925-8388(16)30311-5
DOI:
10.1016/j.jallcom.2016.02.042
Reference:
JALCOM 36653
To appear in:
Journal of Alloys and Compounds
Received Date: 14 December 2015 Revised Date:
4 February 2016
Accepted Date: 5 February 2016
Please cite this article as: H. Liu, Y. Yang, J. Liu, Z. Jiang, Y. Li, W. Yu, W. Ding, G. Fu, Structural and optoelectronic characteristics of nanocrystalline silicon oxide film as absorber layer for thin film solar cells, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.02.042. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
ACCEPTED MANUSCRIPT Structural and optoelectronic characteristics of nanocrystalline silicon oxide film as absorber layer for thin film solar cells Haixu Liu, Yanbin Yang, Jianping Liu, Zhaoyi Jiang, Yun Li, Wei Yu*, Wenge Ding
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and Guangsheng Fu Hebei Key Laboratory of Optic-Electronic Information Material, College of Physics Science and Technology, Hebei University, Baoding 071002, China
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The mixed phase nanocrystalline silicon oxide (nc-SiOx:H) thin films have been
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prepared by very high frequency plasma enhanced chemical vapor deposition (VHF-PECVD) system from SiH4, H2 and CO2 mixture. The effect of oxygen incorporation on the structural and opto-electrical properties of nc-Si:H films was investigated. As the CO2 flow rate increased, the crystalline volume fraction in the
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films decreased, while the optical band gap increased slowly. During the phase transition, both oxygen and hydrogen content increased drastically with the increasing CO2 flow rate, while the Urbatch energy and microstructure factor show obvious
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reduction, which indicates an enhancement of the medium range order of the films
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and good passivation of the inner voids. Device quality nc-SiOx:H film was fabricated with low microstructure factor of 0.32 and high crystalline volume fraction around 50%, along with wide optical band gap of ~1.75 eV, low dark-conductivity σd ~7.6×10-6 S/cm and high photosensitivity of about 2×102. On flat substrate a solar cell was prepared with high Voc of 0.62V and Jsc of 14.57mA/cm2. These results reveal the
*
Corresponding author. Tel.: +86 0312 5079560; fax: +86 0312 5079560.
E-mail address: [email protected].
ACCEPTED MANUSCRIPT potential of nc-SiOx:H film as an absorber layer for thin film solar cells. Key Words: nanocrystalline silicon oxide; phase transition; absorber layer; thin film
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solar cells.
ACCEPTED MANUSCRIPT 1. Introduction Micromorph tandem solar cells with hydrogenated amorphous silicon (a-Si:H) and hydrogenated microcrystalline silicon (µc-Si:H) absorber layers have been intensively studied for its potential of making high efficiency thin film silicon solar
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cells. Despite significant progresses in micromorph solar cells [1], the short-circuit current density (Jsc) of the entire cell is limited by the a-Si:H top cell. Single junction a-Si solar cell has Jsc over 18mA/cm2, however in tandem devices top cell thickness is
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reduced to match current, which makes current smaller [2]. To increase the current of tandem device for higher conversion efficiency, thick top cell is needed. Using a-Si:H
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as top cell can not reach this goal and thick a-Si:H top cell has more Staebler-Wronski effect [3].
In recent years, nanocrystalline silicon oxide (nc-SiOx:H) films have proven to be a novel material in photovoltaic applications for its tunable optical and electrical
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properties [4-7]. The optical band gap of nc-SiOx:H can be adjusted by changing the oxygen concentration, while the contained nanocrystalline silicon phase will improve the carrier mobility in the film [8]. We believe nc-SiOx:H is a promising candidate for
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the absorber layer of the top cell, considering its wide optical band gap and high
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crystalline volume fraction. Moreover, nc-SiOx:H film is more stable under incident light than a-Si:H due to the improved medium range order and increased carrier lifetime [9]. However, the oxygen incorporation may result in defects and disordered structure in the film hence the oxygen concentration should be well controlled [10]. Till now, there are hardly any reports on the nc-SiOx:H material as the absorber layer for solar cells, compared with the intensively studied a-SiOx:H thin film solar cells [11-13].
ACCEPTED MANUSCRIPT This work aims to discuss the potential use of nc-SiOx:H film as the absorber layer in thin film solar cells, which may facilitate the fabrication of novel tandem solar cells. Towards this goal, the influences of oxygen incorporation on the structural and optoelectronic properties of nc-Si:H films were thoroughly investigated. Device
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quality nc-SiOx:H material with high conductivity and wide band gap has been fabricated with optimal oxygen incorporation near the phase transition. With low σd of 7.6×10-6 S/cm and high photosensitivity of 2×102, nc-SiOx:H material demonstrates
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its importance for high efficiency silicon-based thin film photovoltaics.
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2. Experiment
Hydrogenated silicon oxide thin films were deposited by VHF-PECVD (multi-chamber system) from a gas mixture of silane (SiH4), hydrogen (H2) and carbon dioxide (CO2) at 60MHz, 200 , 60 Pa. The plasma excitation power density
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was 45mW/cm2. The gas flow rates of SiH4 and H2 were kept constant while the H2 diluted CO2 was oxygen source. Here we defined the parameter Ro as the CO2/SiH4
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gas flow ratio. Material characterization was performed on 400nm thick SiOx layers deposited on glass substrates and silicon wafers (100).
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Solar cells with p-i-n configuration were prepared on AZO (ZnO:Al) coated glass substrates. The p-type nc-SiOx doped layer was deposited from H2, SiH4，CO2 and Trimethylboron (TMB) mixture reported in our previous work [14]. The nc-SiOx:H solar cell fabricated in this work has the following structure: AZO glass /p-nc-SiOx:H (20nm)/buffer layer (4nm)/i-nc-SiOx:H (1µm)/n-a-Si:H (30nm)/ITO/Ag. The thickness of the film was measured by a step profiler (Dektak 150). The structural properties were characterized by Raman spectroscopy (LabRAM-HR,
ACCEPTED MANUSCRIPT 532nm). High resolution transmission electron microscopy (HRTEM) observation of the sample was performed with FEIF20 transmission electron microscope. The TEM sample was scraped from the substrate, triturated into powder, and then sonically
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dispersed in ethanol. A copper TEM microgrid was immersed into the ethanol to collect the fragment of the sample. FTIR spectra were measured to investigate the bonding configuration of the as-deposited films (Bruker Tensor27). Absorption
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spectra of the layers were recorded by UV-VIS-NIR spectrometer (Hitachi UV-4100).
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The electrical conductivity was characterized by current density versus voltage (J-V) measurements. The photoconductivity and the solar cell performance were measured by B1500A semiconductor device parameter analyzer under AM1.5 solar simulator. 3. Results
To characterize the structural properties of the materials, normalized Raman
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spectra of nc-SiOx:H thin films deposited with different Ro are shown in figure 1(a). Firstly, an oxygen free nc-Si:H sample was deposited as a starting point, whose
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Raman peak locates mainly at 520 cm-1, indicating nanocrystalline phase was dominant in the film. With the increase of Ro from 0% to 14%, a broad peak around
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480 cm-1 related to amorphous phase became dominant. Moreover, an obvious blue shift of peak position could be observed as the miorostructure changed. In order to describe the evolution of the microstructure quantitatively, crystalline
volume fraction XC can be calculated from the following formula: X C = (I 510 + I 520 ) (I 480 + I 510 + I 520 )
(1)
where Ii denotes the integrated intensity of the individual Gaussian peak centered at i cm-1 [15]. The XC as a function of Ro is shown in figure 1(b). The crystalline volume
ACCEPTED MANUSCRIPT fraction XC decreases from ~70% to 8% as Ro increases from 0 to 14%. These results suggest that the incorporation suppresses the formation of crystallites resulting in the phase transition from crystalline to amorphous. It is obvious that there is a transition
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to amorphous growth regime at Ro of ~9-12%. HRTEM characterization was carried out to get detailed structural information of the deposited material. Figure 2 shows a HRTEM image of a nc-SiOx:H sample
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deposited with Ro = 8%, which has a typical mixed-phase structure with the silicon
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nanocrystals (highlighted in circles) embedded in a amorphous matrix. It can be observed that the diameters of the silicon nanocrystals cluster range from 3 nm to 7 nm and the average size of silicon nanocrystals is about 5.3 nm. The inset is a zoom in image of a silicon nanocrystal cluster in the film. It is clear that in the cluster there is
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order organize structure.
FTIR spectrum of the nc-SiOx:H layers are shown in figure 3. From figure 3, one can see that, with the increase of Ro, the peak of Si-O stretching mode (900
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cm-1~1300 cm-1) slightly shifts to higher wavenumber (as denoted by the dashed line).
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This shift of the peak of Si-O stretching mode can be attributed to the oxygen back-bonding or the variation of the Si-O-Si bond angles [16]. The peak around 630cm-1 related to Si-H has no visible variation as Ro increases, because it is silicon rich film.
The bonded hydrogen content (CH) and oxygen content (CO) can be derived from the Si-H peak around 630cm-1 and the Si-O peak (900-1300cm-1), respectively. The calculations are performed with a numerical integration formula:
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(2)
where α(ω) is the absorption coefficient, ω is the wavenumber, NSi = 5×1022 cm-2 is the atomic density of the crystalline silicon, and Aω is the proportionality constant for
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hydrogen (1.6×1019 cm-2) and oxygen (9.5×1018 cm-2) [17]. Figure 4(a) and (b) show the calculated results of CO and CH of nc-SiOx:H films versus Ro. With the increase of Ro from 0 to 9%, the oxygen content CO increases
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gradually from 0% to 4.8%,which indicates an enhancement of incorporation of
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oxygen in the films. Oxygen concentration then declines slightly from 4.8% to 4.3% with further increase of Ro to 14%. The tendency of CH is quite similar to that of CO. With increasing Ro, the CH once increases significantly from 10.7% to 17.9%, and then decreases slightly from 17.9% to 15%. This drastic changes of CH can also be
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found in the transition from nc-Si:H to a-Si:H [18].
The Si-H stretching mode (1900 cm-1~2200 cm-1) includes two vibration peaks related to SiH and SiH2 absorption centered at 2000 cm-1 and 2100 cm-1, respectively.
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The SiH absorption peak corresponds to the amorphous Si phase of compact
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microstructure while SiH2 absorption peak can be attributed to the Si-H bonds located on the surface of inner voids or grain boundaries [19]. To evaluate the quality of the deposited films, the microstructure factor is calculated by the formula: R* = I 2100 /( I 2000 + I 2100 )
(3)
where I2000 and I2100 are the integration intensity of the absorption peaks corresponding to SiH and SiH2, respectively. Here I2000 and I2100 are obtained by Gaussian fitting [20]. The R* is used to evaluate the quality of the deposited films and
ACCEPTED MANUSCRIPT lower indicates better film quality [19]. As shown in figure 4(c), as Ro inceases, R* deceases first and then inceases, and at Ro =8% the R* is lowest. The quality of the film is thus improved with oxygen
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concentration up to 3.2%. Higher oxygen concentration will deteriorate the film quality by introducing more recombination sites as well as amorphous phase becomes dominant. It is well known that the Si-H stretching mode (2000 cm-1) may be shifted
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to higher energy by oxygen back-bonding, which is the so-called induction effect [21].
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It means that the peak at 2100 cm-1 can be attributed to both SiH2 at surfaces of inner voids and oxygen back-bonding of SiH in a high quality film. Considering this fact, the microstructure factor may be even smaller than the calculated results. The optical properties are analyzed by UV-VIS-NIR absorption spectra. The
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optical band gap Eg of the nc-SiOx:H layer can be estimated by the equation:
(αhν ) 12 = B(hν − E g )
(4)
known as Tauc formulation, where α is the absorption coefficient, hν is the energy of
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[22,23].
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incident light, Eg is the optical band gap, and B is a constant related to disorder degree
To characterize the width of band-tail states and the disorder degree in the
amorphous network, the Urbach energy EU can be described as: hν EU
α ≈ exp
(5)
which can be extrapolated from the semilogarithm plot of α versus hν [23]. Figure 5(a) shows the plot of (αhν)1/2 (Ro=8% and 14%) versus photon energy, from which Eg can be extrapolated from a straight line to hν = 0. The solid lines show
ACCEPTED MANUSCRIPT the linear fit of the straight part. The calculated Eg and EU are plotted in figure 5(b). With the increase of Ro, the values of Eg show a general increasing trend with a fluctuation. It is also notable that the band gap is approximately 1.75 eV with only 3.2
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at% oxygen concentration (Ro = 8%). Such wide band gap is comparable to the Eg of a-Si:H and will be favourable for the application as i-layer instead of the a-Si:H layer in the top cells. With the increase of Ro, the Urbach energy EU first decreases and then
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increases. The minimum of EU is also obtained with Ro = 8%, which indicates the
ordered microstructure in the film.
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reduction of the localized band-tail states (particularly the valence band-tail) and more
Figure 6 (a) and (b) show the σd and σph, respectively. The photosensitivity (σph/σd) versus Ro is shown in figure 6 (c). It is obviously that both σd and σph
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decrease with the increasing Ro. The incorporation of oxygen reduces the crystallinity of the nc-Si:H film, which may account for the degradation of σd and σph. However, since σd decreases more rapidly than σph, the photosensitivity increases with Ro from
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10 to 103. Low σd of ~7.6×10-6 S/cm and high photosensitivity of 2×102 can be
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obtained at Ro = 8%, which is near the phase boundary. Similar reports about device-grade µc-Si:H has also been obtained near the phase transition [24]. To reveal the potential of intrinsic nc-SiOx:H films as the absorber layers, a thin
film solar cell was fabricated with a configuration of glass substrate/flat ZnO/p-nc-SiOx:H/buffer layer/i-nc-SiOx:H/n-a-Si:H/ITO/Ag. Figure 7 shows the J-V curves of single junction solar cell containing the nc-SiOx:H i-layer deposited with Ro = 8%. An initial efficiency of about 5% is obtained with Voc = 0.62V, Jsc =
ACCEPTED MANUSCRIPT 14.57mA/cm-1 and FF = 0.533. It should be noted that no textured substrate is adopted here, and the p/i interface needs further optimization. Higher conversion efficiency will be achieved with improvement of the process parameters and the
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structure of the cell.
4. Discussion
To discuss the structural evolution caused by oxygen incorporation, the Si-O
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stretching vibration mode in the FTIR spectra is fitted by several Gaussian peaks.
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Figure 8(a) shows the fitted result of a nc-SiOx:H layer deposited with Ro = 14%. The peaks centered at 980, 1012, 1034, and 1076 cm-1 correspond to the stretching vibration modes of HSi(Si2O), HSi(SiO2), HSi(O3) and Si(O4), respectively [25]. The absorption peak at 1150 cm-1 is attributed to the vibration of the oxygen in the Si(O4)
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network [21,25].
Here we define the fraction of silicon-rich vibration modes and the fraction of oxygen-rich vibration modes as:
(6)
X O − rich = (I1034 + I1076 ) (I 980 + I1012 + I 1034 + I 1076 )
(7)
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X Si −rich = (I 980 + I1012 ) (I 980 + I1012 + I1034 + I1076 )
where Ii denotes the integrated absorption intensity at the i cm-1 [26]. The nc-SiOx:H film actually contains different phases, including the nanocrystalline silicon (nc-Si), hydrogenated amorphous silicon (a-Si:H) and oxygen-rich hydrogenated amorphous silicon oxide (a-SiOx:H) [27]. It is thus reasonable that XSi-rich and XO-rich are related to the a-Si:H phase and the oxygen-rich a-SiOx:H phase in the deposited films, respectively.
ACCEPTED MANUSCRIPT Figure 8(b) shows the XSi-rich and XO-rich versus Ro, which may facilitate the description of the structural evolution. With the increase of Ro from 0 to 8%, the XSi-rich decreases while XO-rich becomes larger. At Ro = 8%, XSi-rich shows a minimum
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while XO-rich reaches its maximum, which is an indication of the Si/SiOx phase separation. As shown in figure 4(c) and 5(b), the microstructure factor R* and the Urbach energy EU are also greatly reduced at Ro = 8%. The phase separation may give
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rise to the enhancement of the medium range order of the film and good surface
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passivation of the grain boundaries.
The broadening of optical band gap Eg can be mainly attributed to oxygen incorporation into the samples. According to the results of mixed-phase nc-Si:H material reported by Yan et al, quantum confinement effect occurs most likely in
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isolated silicon nanocrystals with small grain sizes [28]. In most of the samples deposited here (with XC ≥ 30%), the XC is high enough to form conduction path for carriers and the quantum confinement effect may not have the dominant effect on the
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optical band gap. The broadening of Eg is mostly associated with the increasing
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a-SiOx:H phase in the deposited film. The oxygen induced band gap broadening can be related to the fact that Si-O bond energy (8.4 eV) is stronger than the Si-Si bond (2.4 eV) and the Si-H bond (3.0 eV) [29]. In addition, the oxygen back bonded to Si-Si bond increases the Si-Si bond energy, which results in the increase in the energy difference between the bonding and anti-bonding splitting of the Si-Si bonds. As Eg is a weighted average of the conduction and valence band states (EC and EV), the oxygen back-bonding will lead to the widening of Eg [30].
ACCEPTED MANUSCRIPT To explain the increase of the photosensitivity near the phase transition region, the dark-conductivity σd, photo-conductivity σph and photosensitivity are described as following formulas:
σ ph = e(N d + Gτ )µ ∝ µτ
(8) (9)
τ
E − EF exp − C K BT
(10)
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σ ph e( N d + Gτ ) µ = ∝ σd eN d µ
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EC − EF K BT
σ d = eN d µ ∝ τ exp −
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where Nd is the carrier concentration under thermal equilibrium condition, µ is the carrier mobility, G is the carrier generation rate, τ is the carrier lifetime, EC is the conduction band edge, EF is the Fermi level position, KB is Boltzmann constant and T is the temperature [11]. As oxygen concentration increases, the carrier lifetime τ
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decreases as a consequence of the reduction of the crystalline volume fraction XC and larger scattering probability in the amorphous network, which results in the decreases
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of σd and σph. The increase of photosensitivity thus can be attributed to the rising value of (EC－EF), which can be related to the increase of the optical band gap shown
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in figure 5(b). In general, the Fermi level EF remains almost unchanged in the a-SiOx:H films [10]. However, in this work, the drastic increase of photosensitivity may suggest that the EF is lowered near the phase transition. As known from the former discussion, the localized electronic states on the grain boundaries and valence band-tail states are significantly reduced in the films, which have been demonstrated by the former results of R* and EU in figure 4(c) and figure 5(b). It is thus reasonable that the shift of EF is a consequence of the reduction of R* and EU. .
ACCEPTED MANUSCRIPT 5. Conclusions In summary, the influence of oxygen incorporation on the structural, optical and electrical properties of nc-Si:H layer deposited by VHF-PECVD is investigated. The
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incorporation of oxygen gives rise to band gap broadening and inhibition of the crystalline Si growth, as the nc-Si:H phase is gradually substituted by the a-SiOx:H phase in the films. It is found that the nc-SiOx:H material with optimal optoelectronic
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properties can be obtained near the phase transition. Device quality nc-SiOx:H layer
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has been prepared with wide optical band gap of ~1.75 eV, high crystalline volume fraction of 50% and low oxygen concentration of ~3.2at%, along with low microstructure factor of 0.32, low σd ~7.6×10-6 S/cm and high photosensitivity of 2×102. This material deserves enormous promise for applications in nc-SiOx:H solar
solar cells.
Acknowledgments
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cells, especially for preparing high efficiency and stable top cell in thin film tandem
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This work is financially supported by the Specialized Research Fund for the
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Doctoral Program of Higher Education, PRC (Grant No. 20131301120003), the Science and Technology Planning Project of Hebei Province, PRC (Grant No. 13214315) and the Colleges and universities in Hebei province science and technology research youth fund, PRC (Grant No. QN20131115).
References [1] J. S. Cashmore et al, Sol. Energy Mater. Solar Cells 144 (2016) 84. [2] K Tabuchi, W. Wenas, A.Yamada, M Konagai, K. Takahashi, Jpn. J. Appl. Phys.
ACCEPTED MANUSCRIPT 32 (1993) 3764. [3] B. J. Yan, G. Z. Yue, L. Sivec, J. Yang, S. Guha, C. S. Jiang, Appl. Phys. Lett. 99 (2011) 113512.
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[4] P. Cuony, D. T. Alexander, I. Perez-Wurfl, M. Despeisse, G. Bugnon, M. Boccard, T. Soderstrom , A. Hessler-Wyser, C. Hebert, C. Ballif, Adv. Mater. 24 (2012) 1182.
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[5] P. D. Veneri, L. V. Mercaldo, I. Usatii, Prog. Photovolt: Res. Appl. 21(2013) 148.
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[6] P. Cuony, M. Marending, D. T. L. Alexander, M. Boccard, G. Bugnon, M. Despeisse, C. Ballif, Appl. Phys. Lett. 97 (2010) 213502.
[7] A. Lambertz, V. Smirnov, T. Merdzhanova, K. Ding, S. Haas, G. Jost, R. E. I. Schropp, F. Finger, U. Rau, Sol. Energy Mater. Solar Cells 119 (2013) 134.
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[8] V. Smirnov, W. Böttler, A. Lambertz, H. Y. Wang, R. Carius, F. Finger, phys. status solidi C 7 (2010) 1053.
[9] S. B. Zhang, X. B. Liao, L. An, F. H. Yang, G. L. Kong, Y. Q. Wang, Y. Y. Xu, C.
EP
Y. Chen, H. W. Diao, Acta Physica Sinica 51 (2002) 1811.
AC C
[10] J. E. Lee, J. H. Park, J. Yoo, K. H. Yoon, D. Kim, J. S. Cho, Solid State Sci. 20 (2013) 70.
[11] S. Inthisang, K. Sriprapha, S. Miyajima, A. Yamada, M. Konagai, Jpn. J. Appl. Phys. 48 (2009) 122402.
[12] D. Y. Kim, E. Guijt, R. van Swaaij, M. Zeman, Prog. Photovolt.: Res. Appl. 23 (2015) 671. [13] D. Y. Kim, E. Guijt and F. T. Si, R. Santbergen, J. Holovský, O. Isabella, R. van
ACCEPTED MANUSCRIPT Swaaij, M. Zeman, Sol. Energy Mater. Solar Cells 141 (2015) 148. [14] Z. L. Shi, Y. Ji, W. Yu , Y. B. Yang, R. D. Cong, Y. J. Chen, X. W. Li, S. F. Guang,
Chin. Phys. B 24 (2015) 078105.
Phys. Lett. 90 (2007) 263509.
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[15] S. Y. Myong, K. Sriprapha, S. Miyajima, M. Konagai and A. Yamada, Appl.
[16] F. Einsele , W. Beyer, U. Rau, J. Appl. Phys. 112 (2012) 054905.
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[17] P. H. Gaskell, D. J. Wallis, Phys. Rev. Lett. 76 (1996) 66.
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[18] A. Shah, J. Meier, E. Vallat-Sauvain, N. Wyrsch, U. Kroll, S. Droz, U. Graf, Sol. Energy Mater. Solar Cells 78 (2003) 469.
[19] Y. Adjallah, C. Anderson, U. Kortshagen, J. Kakalios, J. Appl. Phys. 107 (2010) 043704.
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[20] D. X. Han, K. D. Wang, J. M. Owens, L. N. Gedvilas, H. Habuchi and M. Tanaka, J. Appl. Phys. 93 (2003) 3776. [21] D. V. Tsu, G. Lucovsky, B. N. Davidson, Phys. Rev. B 40 (1989) 1795.
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[22] Q. J. Cheng, S. Y. Xu, K. Ostrikov, Nanotechnology 20 (2009) 215606.
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[23] T. Li , J. Kanicki, W. Kong, F. L. Terry, J. Appl. Phys. 88 (2000) 5764. [24] S. Y. Myong, K. Sriprapha, Y. Yashiki, S. Miyajima, A. Yamada, M. Konagai, Sol. Energy Mater. Solar Cells 92 (2008) 639.
[25] L. He, T. Inokuma , Y. Kurata, S. Hasegawa, J. Non-Crystal. Solids 185 (1995) 249. [26] H. M. Li, W. Yu, Y. M. Xu, Y. Ji, Z. Y. Jiang, X. Z. Wang, X. W. Li, G. S. Fu, Chin. Phys. B 24 (2015) 028102.
ACCEPTED MANUSCRIPT [27] A. Richter, L. Zhao, F. Finger, K. N. Ding, Surf. Coat. Technol. (2015) In Press. [28] B. J. Yan, G. Z. Yue, J. Yang, S. Guha, Sol. Energy Mater. Solar Cells 111 (2013) 90.
Barua, Mater. Chem. Phys. 157 (2015) 130.
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[29] S. Mandal, G. Das, S. Dhar, R. M. Tomy, S. Mukhopadhyay, C. Banerjee, A. K.
[30] I. Umezu, K. Miyamoto, N. Sakamoto, K. Maeda, Jpn. J. Appl. Phys. 34 (1995)
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1753.
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Figure captions: Fig.1 (a) Raman spectra of nc-SiOx:H layers grown with different Ro. (b) The crystalline volume fraction and growth rate of nc-SiOx:H layers versus Ro.
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Fig.2 HRTEM image of nc-SiOx:H layer deposited with Ro = 8%, and the inset shows
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a magnified image of a typical silicon nanocrystal embedded in silicon oxide matrix. Fig.3 FTIR spectra of nc-SiOx:H layers deposited with different Ro. Fig.4 (a) The oxygen content, (b) the hydrogen content and (c) the microstructure factor plotted as functions of Ro.
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Fig.5 (a) (αhν)1/2 (Ro=8% and 14%) versus the photon energy. (b) The calculated optical band gap Eg and Urbach energy EU versus Ro. Fig.6 (a) The dark-conductivity, (b) photo-conductivity and (c) photosensitivity of
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nc-SiOx:H films plotted as a function of Ro.
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Fig.7 J-V characteristics of the nc-SiOx:H solar cells with the intrinsic layer deposited with Ro = 8%.
Fig.8 (a) Gaussian fitting results of the Si-O stretching mode of a nc-SiOx:H film (Ro = 14%). (b) Plots of XSi-rich and XO-rich in the nc-SiOx:H films versus Ro.
(a )
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3
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1 4 %
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1 2 % 1 0 % 8 %
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2
5 %
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1
0
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N o r m a liz e d In te n s ity (a .u .)
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4 0 0
4 5 0
R 5 0 0
R a m a n s h ift (c m
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O
3 % = 0 % 6 0 0
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400
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300
TE D
200
RO=14%
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100
0
RO=8%
EP
1/2
(h) (eV/cm)
1/2
(a)
1.7
1.8
1.9
2.0
2.1
2.2
2.3
2.4
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1 6
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1 4
V
o c
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6
J 4
F F = 0 .5 3 3 C E = 5 %
= 1 4 .5 m A /c m
2
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s c
TE D
8
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1 0
J (m A /c m
2
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SC
1 2
0
2 0
1 0 0
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0 .6 5
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0 .4 5
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0 .4 0
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6
R
8 O
1 0 (% )
1 2
1 4
F r a c tio n o f O X y g e n -r ic h
F r a c tio n o f S ilic o n -r ic h
(b )
ACCEPTED MANUSCRIPT Reaserch Highlights
► The structural and opto-electrical properties of nc-SiOx:H layer is investigated. ► Device quality nc-SiOx:H material is obtained near the phase transition region.
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► A nontextured solar cell was prepared with Voc of 0.62V and Jsc of
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14.57mA/cm2.