Solution-processed solar cells based on inorganic bulk heterojunctions with evident hole contribution to photocurrent generation.

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Cite this: Phys. Chem. Chem. Phys., 2015, 17, 12328

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Solution-processed solar cells based on inorganic bulk heterojunctions with evident hole contribution to photocurrent generation† Zeliang Qiu,a Changwen Liu,a Guoxing Pan,b Weili Meng,a Wenjin Yue,a Junwei Chen,a Xun Zhou,a Fapei Zhangb and Mingtai Wang*a To develop solution-processed and novel device structures is of great importance for achieving advanced and low-cost solar cells. In this paper, we report the solution-processed solar cells based on inorganic bulk heterojunctions (BHJs) featuring a bulk crystalline Sb2S3 absorbing layer interdigitated with a TiO2 nanoarray as an electron transporter. A solution-processed amorphous-to-crystalline transformation strategy is used for the preparation of Sb2S3/TiO2-BHJs. Steady-state and dynamic results demonstrate that the crystalline structure in the Sb2S3 absorbing layer is crucial for efficient devices, and a better Sb2S3 crystallization favors a higher device performance by increasing the charge

Received 4th January 2015, Accepted 8th April 2015

collection efficiency for a higher short-circuit current, due to reduced interfacial and bulk charge

DOI: 10.1039/c5cp00030k

the Sb2S3 layer as well. Moreover, an evident contribution to photocurrent generation from the

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photogenerated holes in the Sb2S3 layer is revealed by experimental and simulated dynamic data. These results imply a kind of potential non-excitonic BHJ for energy conversion.

recombinations, and enhancing the open-circuit voltage and fill factor with the reduced defect states in

1. Introduction Global demand for renewable and substantial energy has encouraged research into photovoltaic materials and devices with a good trade-off between performance and fabrication cost.1 Solutionprocessed inorganic solar cells (ISCs) are a possibility due to their high charge mobility, excellent environmental stability and broad absorption spectrum of inorganic materials. ISCs are often fabricated by depositing sequentially preformed n- and p-type nanocrystals into planar heterojunctions (PHJs)2,3 or blending the nanocrystals into bulk heterojunctions (BHJs).4 Due to their nontoxicity, stability to environment, transparency to visible light, high electron mobility and ready availability by solution reactions, vertically aligned nanorod/nanowire arrays (NAs) of

a

Institute of Plasma Physics, Chinese Academy of Sciences, Hefei 230031, P. R. China. E-mail: [email protected]; Fax: +86-551-65593171; Tel: +86-551-65593171 b High Magnetic Field Laboratory, Chinese Academy of Sciences, Hefei 230031, P. R. China † Electronic supplementary information (ESI) available: TiO2-NA characterization, XRD data of Sb2S3/TiO2-BHJs crystallized at different temperatures, band gaps and UPS data of the Sb2S3 layers in crystalline Sb2S3/TiO2-BHJs, absorption spectrum of MEH-PPV, and additional calculated results for Je, Jh and Jph, SEM data of Sb2S3/TiO2-BHJs crystallized at 300 and 450 1C, steady-state and dynamic data for the solar cell based on Sb2S3/TiO2-BHJs crystallized at 300 1C. See DOI: 10.1039/c5cp00030k

12328 | Phys. Chem. Chem. Phys., 2015, 17, 12328--12339

ZnO5 and TiO26 have been widely used as the ideal electron acceptor and transporter in devices such as dye-sensitized solar cells (DSCs) and hybrid polymer-based solar cells (HPSCs). Filling preformed p-type quantum dots into ZnO-NAs and TiO2-NAs from solutions results in ISCs based on aligned inorganic BHJs with straightforward nanochannels for efficient charge transport.7–10 However, the ISCs from preformed nanocrystals often encounter deficiencies related to their small size, such as high exciton binding energy,11 low carrier concentration and high trap density,12 charge conduction level dispersion and grain boundary issues (e.g. capping ligand and electronic contact).13 Hence, a solution-processed nanocrystal film is often sintered into a more continuous layer with increased electronic contact and eliminated capping ligand between nanocrystals for a high charge transport and collection efficiency.10,14 Recently, Lee and Yong15 filled the molecular precursors for CuInS2 into the CdS nanocrystal coated ZnO-NA and prepared the solar cells after the CuInS2 nanocrystal had formed inside the ZnO-NA by thermal decomposition of the precursor molecules. People have also attempted to sinter the nanocrystal film into a bulk layer for harvesting photons in PHJ devices13,16 or directly electrodeposit a bulk absorbing layer into ZnO and TiO2 nanoarrays to prepare BHJ devices.17,18 However, the photocurrent generation features in those ISCs are not clear yet. Crystalline Sb2S3 is a promising absorbing material for solar cells, due to its high absorption coefficient (a = 105 cm1 at 450 nm)

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in the visible spectrum and appropriate band gap (Eg E 1.70 eV).19 There have been several reports on the application of Sb2S3 in solar cells. A power conversion efficiency (Z) of around 5% was achieved in the hybrid solar cells fabricated by infiltrating a polymer into TiO2/Sb2S3 coaxial nanowire arrays.20 Seok and co-workers21–23 have reached efficiencies of 5–6% in the hybrid solar cells based on Sb2S3-sensitized mesoporous TiO2 films with polymers as hole transporting materials. Li and co-workers24 deposited Sb2S3 nanoparticles into TiO2-NAs and prepared the sensitized solar cells with an efficiency of ca. 1.5% in combination with a liquid electrolyte. Moreover, Ito and co-workers25 deposited Sb2S3 extreme thin absorber (ETA) onto a TiO2 nanoparticle film and obtained all-inorganic ETA solar cells with efficiencies of 5.7%. To prepare those solar cells,20–22,24,25 the chemically deposited Sb2S3 is normally annealed at a certain temperature to form Sb2S3 nanocrystals. Herein, we use the solution-processed amorphous-to-crystalline transformation strategy to fabricate solar cells based on inorganic Sb2S3/TiO2-BHJs, which feature a bulk Sb2S3 absorbing layer interdigitated with a TiO2-NA as a direct electron transporter; our results reveal the crucial importance of the crystalline structure in the Sb2S3 absorbing layer for efficient devices and the evident contribution to photocurrent generation from the photogenerated holes in the bulk Sb2S3 layer in those inorganic BHJ solar cells. The formation of Sb2S3/TiO2-BHJs and solar cell architecture are depicted in Fig. 1. In this experiment, TiO2-NA (nanorod length ca. 600 nm, nanorod diameter 40–50 nm) was hydrothermally grown on a fluorine-doped tin oxide (FTO) coated glass sheet, as shown in Fig. S1 in the ESI.† The TiO2-NA is first filled with amorphous Sb2S3 nanoparticles by chemical bath deposition (CBD) and then a thermally induced crystallization at 350–400 1C in N2 atmosphere turns the amorphous Sb2S3 nanoparticles into a bulk crystalline Sb2S3 layer of 800 nm in thickness. The Sb2S3/TiO2-BHJ solar cells are fabricated by spincoating a thin film (10–15 nm thick) of poly(2-methoxy-5-(2ethylhexyloxy)-1,4-phenylene vinylene) (MEH-PPV) as a buffer layer over Sb2S3, followed by a layer (40–50 nm thick) of

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poly(3,4-ethylene dioxythiophene):poly(styrene sulfonate) (PEDOT: PSS) as hole transporter and a thermally evaporated Au electrode (100 nm) over the PEDOT:PSS layer. Note that our solar cells are completely different from the previously reported hybrid,20–23 sensitized24 and all-inorganic25 solar cells, in that the absorption and charge generation properties of Sb2S3 mainly originate from the nanosized Sb2S3 phases in the previous reports, but from the bulk state Sb2S3 in our case. It is also noticeable that our solar cells are quite different from the ZnO/Sb2S3/poly(3-hexylthiophene) (P3HT) hybrid solar cells fabricated by a thermally evaporated crystalline Sb2S3 layer (B100 nm) into ZnO-NAs (length B100 nm, diameter B30 nm) and spin-coating a P3HT film (40 nm) over the Sb2S3 layer,26 in which the Sb2S3 layer acts as the main absorber to harvest about 67% incident power and P3HT provides additional absorption to the left incident photons.

2. Experimental section 2.1.

TiO2-NA (Fig. S1, ESI†) was grown on a FTO (SnO2:F) coated glass sheet (14 O &1, 400 nm FTO in thickness, Nippon Sheet Glass Co.) by a hydrothermal method, as described elsewhere.27 Typically, the FTO substrate was suspended upside down in an mixture of 30 mL deionized water, 30 mL concentrated HCl (36% by weight) and 1 mL titanium(IV) n-butoxide (Alfa Aesar, 99+%) contained in a Teflon-lined stainless steel autoclave of 100 mL capacity, and the hydrothermal growth was conducted at 180 1C for 2.5 h. Deposition of Sb2S3 into TiO2-NA was carried out by a CBD technique as follows: TiO2-NA was suspended upside-down in the aqueous solution of Na2S2O3 (0.28 M) at about 10 1C and the acetone solution of SbCl3 (0.3 M) was added dropwise into the Na2S2O3 under stirring; then, the Sb2S3 deposition was done at about 10 1C for 5–6 h. After the CBD process, a smooth and uniform Sb2S3 layer was deposited over the TiO2-NA, and the sample was rinsed thoroughly with deionized water and dried in vacuum overnight, resulting in the as-deposited Sb2S3/TiO2-BHJ film. Note that in the as-deposited sample, the back substrate side was decorated by some large Sb2S3 particles over a very thin Sb2S3 layer, which were wiped away with cotton swabs soaked in dilute HCl. The as-deposited Sb2S3/TiO2-BHJ film was further annealed for 0.5 h at a given crystallization temperature (Tc) on a digital hot plate (IKAs C-MAG HP 7, Germany) in a glovebox (O2 r 1 ppm, H2O r 1 ppm) under a N2 atmosphere to obtain a crystallized Sb2S3/TiO2-BHJ film. 2.2.

Fig. 1 Schematic illustrations for the formation of crystallized Sb2S3/TiO2BHJs by amorphous-to-crystalline transformation and the solar cell architecture.


Preparation of Sb2S3/TiO2-BHJs

Solar cell fabrication

The Sb2S3/TiO2-BHJs in solar cells are those crystallized at Tc temperatures. The Sb2S3/TiO2-BHJ solar cells (Fig. 1) were fabricated as follows. The thin MEH-PPV (avg. Mn = 40 000–70 000, Sigma-Aldrich) layer was spin-coated on the Sb2S3/TiO2-BHJ film from the solution in chlorobenzene (5 mg mL1) under ambient conditions, followed by annealing at 150 1C for 10 min under a N2 atmosphere. Afterwards, the film of PEDOT:PSS (Clevios PH1000, H. C. Starck) was spin-coated over the MEH-PPV layer, for which

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the PEDOT:PSS solution with isopropanol (50% in volume) was filtered with a 0.80 mm filter prior to use. After deposition of the PEDOT:PSS layer, the sample was subjected to a thermal treatment of 100 1C for 15 min under a N2 atmosphere. Finally, the Au electrode was thermally evaporated over the PEDOT:PSS layer through a shadow mask to form an overlapped area between FTO and Au of 1 4 mm2, which defined the active area of each device. The solar cells were sealed in a glovebox (O2 r 1 ppm, H2O r 1 ppm) under a N2 atmosphere. 2.3.

Instruments and characterization

X-ray diffraction (XRD) patterns were collected on a Philips X’Pert Pro MPD diffractometer with monochromated Cu Ka radiation (l = 1.54056 Å). Scanning electron microscopy (SEM) measurements were performed on a field-emission scanning electron microscope (FESEM, FEI-Sirion 200), and transmission electron microscopy (TEM) and high-resolution TEM (HRTEM) studies were carried out on a JEOL-2010 microscope under an acceleration voltage of 200 kV. Absorption spectra were recorded on a Shimadzu UV-2550 spectrophotometer, and photoluminescence (PL) spectra were obtained by a Hitachi F-7000 spectrofluorophotometer. Ultraviolet photoelectron spectroscopy (UPS) of the samples was carried out on a Thermao ESCALAB 250 photoelectron spectrometer (Thermo Electron Corporation) under an ultrahigh vacuum chamber (9.5 1010 mbar) with a He ultraviolet photon discharge lamp (21.22 eV) with a resolution of 0.1 eV. During the UPS measurements, a bias voltage of 8.0 V with respect to the ground was applied to the samples for the measurement of the binding energy of the secondary electron cut-off. The Sb2S3 samples for UPS measurements were fabricated by the same method as that for solar cell fabrication. Solar cells were characterized with current–voltage ( J–V) measurements, intensity modulated photocurrent spectroscopy (IMPS) and intensity modulated photovoltage spectroscopy (IMVS), and incident photon-to-current efficiency (IPCE) spectrum. Steady-state J–V characteristics were measured under AM 1.5 illumination with an intensity of 100 mW cm2 from a 94023A Oriel Sol3A solar simulator (Newport Stratford, Inc.), and the light intensity from a 450 W xenon lamp was calibrated with a standard crystalline silicon solar cell; the J–V curves were collected with an Oriels I–V test station (PVIV-1A, Keithley 2400 Source Meter, Labview 2009 SP1 GUI Software, Newport). IMPS and IMVS dynamic spectra were measured by controlled intensity modulated photo spectroscopy (CIMPS, Zahner Co., Germany) in ambient conditions with a background intensity of 15.85 mW cm2 and within the frequency range 1 Hz–25 kHz, as described previously.28 During the J–V and dynamic measurements, the illumination was limited to the overlapped area (1 4 mm2) between FTO and Au by a photomask attached to each device. IPCE spectra were measured with a QE/IPCE Measurement Kit (Newport, USA) that was automatically controlled by Oriels Tracq Basic V5.0 software with a light from a 300 W xenon lamp focusing through a monochromator (74125 Oriel Cornerstone 260 1/4 m) onto the solar cells. The light intensity and photocurrent generated were measured with a


2931-C dual channel power/current meter and a 71675 calibrated UV silicon photodetector.

3. Results and discussion 3.1.

Formation of Sb2S3/TiO2-BHJs

Shown in Fig. 2 are the XRD patterns of the samples. The as-deposited Sb2S3/TiO2-BHJ film does not show any diffraction peaks of crystalline materials, indicating that the Sb2S3 deposited onto TiO2-NA during the CBD process was amorphous. After annealing at 350 and 400 1C, however, the Sb2S3/TiO2-BHJ films display the diffraction peaks of orthorhombic Sb2S3 (stibnite) (JCPDS card 42-1393). The XRD data suggest that thermal annealing at 350 and 400 1C induces the crystallization of amorphous Sb2S3. The Sb2S3/TiO2-BHJ film crystallized at Tc = 400 1C exhibits more intensive diffraction peaks than that crystallized at Tc = 350 1C, indicating a higher crystallinity in the Sb2S3 layer. Note that the TiO2-NA diffraction peaks are not observed when Sb2S3 is present, because it is covered by a thicker Sb2S3 layer. Clearly, a weak diffraction peak for Sb2O3 is observed for Tc = 350 1C, but is not evident for Tc = 400 1C (Fig. 2). Similarly, the presence of Sb2O3 was also observed in the Sb2S3 thermally annealed at 330 1C for 30 min in Ar atmosphere.29 Our further results showed that increasing the Tc leads to a great reduction of the Sb2O3 diffraction peak and enhances the Sb2S3 crystallinity in Sb2S3/TiO2-BHJs (Fig. S2 in ESI†). The SEM image in Fig. 3a shows that the interspaces between TiO2 nanorods are completely filled with amorphous Sb2S3 nanoparticles after the CBD process in as-deposited Sb2S3/ TiO2-BHJs. However, when subjected to thermal annealing at 350 and 400 1C, the amorphous Sb2S3 nanoparticles change into a continuous layer with large grains around the TiO2 nanorods and some grains extending the full Sb2S3 layer in the crystallized Sb2S3/TiO2-BHJs (Fig. 3b and c), similar to the cases of the bulk CuO2 layer in ZnO-NAs17 and TiO2-NAs.18 HRTEM images (Fig. 3d and e) clearly show that the TiO2 nanorods are surrounded, without voids between Sb2S3 and TiO2, by randomly orientated but compactly stacked Sb2S3 crystals.

Fig. 2 XRD patterns of as-deposited and differently crystallized Sb2S3/ TiO2-BHJs. (a) As-deposited, (b) Tc = 350 1C and (c) Tc = 400 1C. The K mark on the plot identifies the diffraction peak for Sb2O3 (senarmontite), which is indexed to its (222) plane according to JCPDS card 72-1334.


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Fig. 3 Sectional SEM (a–c) and HRTEM (d, e) images of the as-deposited and differently crystallized Sb2S3/TiO2-BHJs. (a) As-deposited, (b, d, e) Tc = 350 1C, and (c) Tc = 400 1C. Image (e) is the magnification of the marked region on (d).

Fig. 4 shows the absorption and room temperature PL spectra of the samples. The as-deposited Sb2S3/TiO2-BHJs with amorphous Sb2S3 nanoparticles display an absorption edge at ca. 550 nm, but Sb2S3/TiO2-BHJs with crystallized Sb2S3 layers exhibit an absorption edge at ca. 750 nm. The optical band gap (Eg) of amorphous and crystalline Sb2S3 layers was evaluated by

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the direct band gap method,19 that is, plotting the squared absorbance (Abs2) versus energy (eV) and extrapolating to zero (Fig. S3 in ESI†). Our results showed that the amorphous Sb2S3 material has Eg = 2.40 eV, and the Sb2S3 layers crystallized at Tc = 350 and 400 1C have almost the same band gap with Eg = 1.67 eV; these data are in agreement with the previous reports on similarly prepared amorphous Sb2S3 (Eg = 2.20–2.48 eV) and bulk crystalline Sb2S3 (Eg = 1.70 eV).19 Note that, since the Sb2S3 layers in our samples are quite thick (800 nm, Fig. 3), the absorbance of the samples has exceeded the limitation of our spectrophotometer and only a sharper band edge is observed. The absorption spectra of those crystalline Sb2S3/TiO2-BHJs are a bit different from the absorption spectra of Sb2S3 nanocrystals deposited on mesoporous TiO2 films, which are contoured by a broad shoulder with a trail in the long-wavelength direction.21,29 The valence band maximum (EVBM) of the Sb2S3 layers can be determined by UPS and their conduction band minimum (ECBM) values can be further calculated by adding the corresponding optical band-gap to EVBM.30,31 The EVBM of 5.29 eV and ECBM of 3.63 eV were obtained from the UPS and optical absorption data for the Sb2S3 layers in the Sb2S3/TiO2-BHJs crystallized at Tc = 350 and 400 1C (Fig. S4 in ESI†). As shown in Fig. 4b, TiO2-NA shows a near band edge (NBE) emission at 392 nm (3.16 eV) and deep level emissions (DLE) at 452 nm (2.74 eV) and 469 nm (2.64 eV) due to oxygen vacancy defects on TiO2 nanorods.32,33 Because of the matching lattices between Sb2S3 and TiO2,34 the TiO2 surface defects can be passivated by the Sb2S3 crystallization. With respect to the emission of TiO2-NA, the deposition of Sb2S3 significantly reduces the PL emissions of the surface defect states. The intensity ratio between the DLE and NBE emission (IDLE/INBE) is an indication of crystal quality, and a larger ratio of IDLE/INBE means a higher concentration of surface defects.35 The IDLE/ INBE ratio is almost the same in the Sb2S3/TiO2-BHJs crystallized at 350 and 400 1C, indicating that the TiO2 surface defects in them have been passivated to the same concentration.36 Reasonably, the originally amorphous Sb2S3 nanoparticles in TiO2-NA are transformed into a polycrystalline bulk layer after thermal annealing at 350 and 400 1C, resulting in crystalline Sb2S3/TiO2-BHJs. The crystallization at 400 1C produces a higher crystallinity in the Sb2S3 layer than at 350 1C; however, the surface defects on TiO2 nanorods have been passivated to the same concentration even in those differently crystallized Sb2S3/TiO2-BHJs. 3.2.

Fig. 4 (a) Absorption spectra of the as-deposited and differently crystallized Sb2S3/TiO2-BHJs; (1) as-deposited, (2) Tc = 350 1C, and (3) Tc = 400 1C. (b) Room temperature PL spectra of TiO2-NA (1) and the Sb2S3/ TiO2-BHJs crystallized at (2) Tc = 350 1C and (3) Tc = 400 1C. The PL spectra were measured under excitation at 340 nm with illumination from the FTO side.


Solar cells

The crystallized Sb2S3/TiO2-BHJs were used to fabricate solar cells (Fig. 1). With the valence (ca. 5.29 eV) and conduction (ca. 3.63 eV) band edges of bulk crystalline Sb2S3 (Fig. S4, ESI†), the highest occupied molecular orbital (HOMO) (5.25 eV) level of MEH-PPV,37 the conduction band (CB) edge (4.2 eV) of TiO238 and the work-functions of auxiliary materials (5.1 eV for Au and PEDOT:PSS38 and 4.4 eV for FTO39), it is clear that the materials in the Sb2S3/TiO2-BHJ solar cells form type II heterojunctions with staggered band alignments that facilitate charge transfer during the photovoltaic process (Fig. 5a). Using the


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Fig. 5 (a) Band level alignments and the charge transfer in the Sb2S3/ TiO2-BHJ solar cells, in which the electrons injected into the conduction band of TiO2 may be trapped (k1) by surface defect states (SS) and the trapped electrons can be thermally released (k2) into the conduction band. (b) Calculated Sb2S3 layer thickness required to absorb 90% incident photons of a specific wavelength.

reported dependence of absorption coefficient (a) on photon energy for crystalline Sb2S3,19 the Sb2S3 layer thickness required for absorbing 90% of incident photons of a specific wavelength is calculated (Fig. 5b). MEH-PPV absorbs the photons within the wavelength of 400–600 nm (Fig. S5 in ESI†). When passing through the bulk Sb2S3 layer (i.e., x = 800 nm), more than 99% of the incident photons (Iin) of 400–600 nm are absorbed according to the Beer–Lambert law I = Iin exp(ax), in which a is the absorption coefficient of Sb2S3 as a function of photon wavelength. Therefore, the charge carriers in the Sb2S3/TiO2BHJ solar cells are generated from absorption of the Sb2S3 layer, and the MEH-PPV layer only acts as a buffer layer.40 The photogenerated excitons in the bulk Sb2S3 layer dissociate inside the layer due to its quite low binding energy (o10 meV),26 which is much lower than the thermal energy at ambient temperature (kBT B 26 meV, where kB is the Boltzmann constant and T is temperature), resulting in predominantly non-excitonic solar cells. Finally, the free electrons and holes in the Sb2S3 layer are respectively injected into TiO2 and MEH-PPV for photocurrent generation. Fig. 6 shows the J–V curves and IPCE spectra of the solar cells. The J–V data were measured under AM 1.5 illumination (100 mW cm2) from the FTO side. Table 1 summarizes the averaged overall photovoltaic performance of three individual devices for each sample. With increasing Tc for Sb2S3 crystallization, short-circuit current ( Jsc), open-circuit voltage (Voc) and fill factor (FF) are significantly improved, leading to the efficiency Z of 2.11% for Tc = 400 1C. In particular, the Jsc is very sensitive to Tc, and a greater than 1-fold increment in Jsc is observed for increasing Tc from 350 to 400 1C. The Voc of


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Fig. 6 Typical J–V curves (a) and IPCE spectra (b) of the solar cells based on the Sb2S3/TiO2-BHJs crystallized at 350 and 400 1C.

around 0.35 V is comparable to the BHJ solar cells fabricated by electrodepositing the bulk Cu2O layer into TiO2-NAs,18 but much higher than that in the devices with an electrodeposited bulk Cu2O layer in ZnO-NAs.17 Clearly, increasing Tc also leads to a significant improvement in the device FF. The FF values (B29–37%) in our devices are comparable to the data reported in the ISCs based on the BHJs of Cu2O/ZnO-NAs17 and Cu2O/ TiO2-NAs.18 Moreover, our FF values are close to those reported in the TiO2/Sb2S3 (200–250 nm)/P3HT (150 nm)41 and ZnO/ Sb2S3 (50–350 nm)/P3HT (40 nm)42 planar hybrid solar cells, and in the planar ISCs of CdS/Sb2S3 (50–340 nm)/PbS (200 nm) heterojunctions.43 The small FF in the Sb2S3/TiO2-BHJ devices is related to the charge recombination inside the thick Sb2S3 layer.21,26,42 The Tc influence on IPCE is the same as that on Jsc. Increasing Tc from 350 to 400 1C enhances the IPCE in the 300–750 nm range. Since the Sb2S3 layer thickness required for absorbing 90% photons of the wavelength below 750 nm is around 930 nm (Fig. 5b) and the samples with Tc = 350 and 400 1C exhibit a comparable absorption property (Fig. 4a), the enhanced IPCE indicates that the crystallization at 400 1C makes the charge carriers generated within whole Sb2S3 layer easier to be collected. The IPCE spectra display two typical regions at 300–500 and 500–750 nm. The contributions in these two regions are quantitatively compared using the integrated areas A1 and A2 on the IPCE spectra. When increasing Tc from 350 to 400 1C, the total IPCE within 300–750 nm increases 2.2 fold, in which the A1 and A2 contributions are enhanced by 4.7 and 1.6 times, respectively. As 90% of the photons in the


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Table 1 Device performance of the solar cells based on differently crystallized Sb2S3/TiO2-BHJs. Each of the data with standard deviations represents the average of three individual devicesa

Tc (1C)

Voc (V)

Jsc (mA cm2)

FF (%)

Z (%)

Rs (kO)

Rsh (kO)

tPS (ms)

tVS (ms)

350 400

0.27 0.03 0.35 0.02

6.47 0.27 15.87 0.21

28.85 1.94 36.50 2.04

0.51 0.08 2.11 0.06

0.69 0.07 0.22 0.03

1.21 0.25 1.24 0.21

0.07 0.01 0.04 0.01

1.06 0.38 3.77 0.10


a Rs and Rsh are the series resistance and shunt resistance linked with the slope characteristics at Voc and Jsc, respectively. The time constants tVS and tPS detected by IMVS and IMPS are the te and tD of the photogenerated electrons in TiO2 nanorods, respectively.

range below 500 nm can be absorbed by a Sb2S3 layer of less than 300 nm in thickness (Fig. 5b), the more significant increase in A1 suggests that the charge carriers generated in the TiO2-NA region close to the FTO collection electrode become much more efficient for photocurrent generation for Tc = 400 1C.17 Intensity modulated photovoltage spectroscopy (IMVS) and intensity modulated photocurrent spectroscopy (IMPS) have been successfully used to characterize the charge transport dynamics in DSCs44 and HPSCs.45 During IMVS and IMPS measurements, a sinusoidal perturbation with a small depth of d (i.e., Iac = I0deiot) is superimposed on the background light intensity I0, to form a modulated illumination condition [i.e., I = I0(1 + deiot)]. IMVS measures the photovoltage response to the perturbation under open-circuit conditions, while IMPS measures the photocurrent response to perturbation under short-circuit conditions. Since both photovoltage and photocurrent lag behind illumination, IMVS and IMPS responses generally appear in the fourth (IV) quadrant (positive real, negative imaginary) of complex plane.44 To understand the underlying mechanisms of the improved device performance, the dynamic performances of Sb2S3/TiO2-BHJ solar cells were measured with IMVS and IMPS within the frequency range 1 Hz–25 kHz (Fig. 7). The IMVS responses of the solar cells appear in the IV quadrant of the complex plane (Fig. 7a), similar to the observations on DSCs44 and HPSCs.45 It has been well demonstrated that the IMPS responses of DSCs46 and HPSCs28,47 theoretically appear in the IV quadrant of the complex plane and spiral into the origin in high frequency regimes, when the photogenerated hole contribution to photocurrent and the lag time of charge generation behind excitation are ignorable. However, the measured IMPS responses of the Sb2S3/TiO2-BHJ solar cells appear dominantly in the IV quadrant of the complex plane and enter the third (III) quadrant (negative real, negative imaginary) of the complex plane in high frequency regimes, to cross the real axis without spiraling into the origin (Fig. 7b), which is quite different from the previous reports on DSC,44,46 HPSC28,45,47 and ISC devices formed by infiltrating CuInS2 nanoparticles into porous TiO2 films.48 Note that such a change in the shape of the IMPS response has also been observed in the ISCs based on Cu(Inx,Ga1x)Se2 films,49 for which the reason is not clarified yet. 3.2.1. Dynamic time constants sVS and sPS. The frequency ( fmin) of the lowest imaginary component of IMVS or IMPS responses defines the time constant tVS or tPS according to the relation t = (2pfmin)1. In the cases of DSCs44,46 and HPSCs,28,45,47 the tVS provides the evaluation of the lifetime (te) of photogenerated electrons related to charge recombination under


Fig. 7 IMVS (a) and IMPS (b) spectra of the solar cells based on the Sb2S3/ TiO2-BHJs crystallized at 350 and 400 1C as shown. The inset to (b) magnifies the IMPS spectrum around the origin point. The solid symbols on the plots identify the fmin points, and the arrows show the data points of 6.4 kHz where the IMPS responses change. The crossing points of IMPS response with the real axis at high and low frequencies are referred to as the PHF and PLF points, respectively.

open-circuit conditions, while the tPS reflects the transit time (tD) for the photogenerated electrons to reach the device collection electrode under short-circuit conditions. The tVS and tPS in the Sb2S3/TiO2-BHJ solar cells are listed in Table 1. The tVS data are comparable to the te values for the similarly structured HPSCs formed by infiltrating a conjugated polymer into TiO2-NAs50 and ZnO-NAs;45 however, the tPS data are about one order of magnitude lower than the tD values in the similarly structured HPSCs based on ZnO-NAs.45 Note that both te and tD observed with IMVS and IMPS in those HPSCs correlate with the photogenerated electrons injected into the nanoarrays. Spatial separation and effective transport to collection electrodes of the photogenerated charge carriers before their recombination are crucial for the photovoltaic process. As shown in Fig. 5a, the dominant charge transport processes in


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the Sb2S3/TiO2-BHJ solar cells involve the transport and recombination of photogenerated electrons and holes inside the Sb2S3 layer, electron injection from the Sb2S3 layer into TiO2, and the transport and recombination of the electrons in the TiO2 nanorod. Both tVS and tPS values (Table 1) cannot be direct measures for the transport and recombination of charge carriers in the bulk Sb2S3 layer, because they occur on a much faster time scale as follows: (i) the recombination of photogenerated electron–hole pairs in bulk crystalline Sb2S3 occurs on a time scale from picoseconds up to nanoseconds;51,52 (ii) with an electron diffusion coefficient (De) of 0.25 cm2 s1 (qDe = kBTme, where me = 10 cm2 V1 s1 is the electron mobility in crystalline Sb2S3,53 q is the elementary charge, kB is the Boltzmann constant and T is temperature) and a hole diffusion coefficient (Dh)52 of 7 102 cm2 s1 for crystalline Sb2S3, the time t values for the electrons and holes to diffuse across the Sb2S3 layer with a thickness of L = 800– 1000 nm are estimated to be 0.03–0.04 ms and 0.09–0.14 ms by the relation t = L2/De (or Dh), respectively; and (iii) the electron injection from Sb2S3 into TiO2 occurs on a time scale of less than 10 picoseconds.51 It is rationalized, therefore, that the tVS and tPS values detected by IMVS and IMPS reflect the recombination and transport of the photogenerated electrons in the TiO2 nanorods in the Sb2S3/TiO2-BHJ devices. First, formation of type II heterojunctions with a staggered band alignment between Sb2S3 and TiO2 facilitates charge transfer during the photovoltaic process and back electron transfer from TiO2 into the Sb2S3 conduction band is impossible due to an unfavorable band structure (Fig. 5a), and thereby the electrons injected into the TiO2 nanorods under open-circuit conditions must eventually undergo charge recombination at the Sb2S3/TiO2 interface. Moreover, the time for the injected electrons to diffuse across the TiO2 nanorods with a length of 600 nm for photocurrent generation is estimated to be around 0.07 ms from the De of 5 105 cm2 s1 for TiO2,46 which is close to the experimental tPS data. Clearly, the time constants tVS and tPS detected by IMVS and IMPS in the Sb2S3/TiO2-BHJ solar cells are actually the te and tD of the photogenerated electrons in TiO2 nanorods, in which te correlates with the charge recombination at the Sb2S3/TiO2 interface, and the IMPS responses are dominated by the time constant for the electron transport in TiO2 nanorods. Accordingly, the Sb2S3 crystallization at Tc = 400 1C leads to a reduced interfacial charge recombination (i.e., increased te) and an increased electron diffusion rate (i.e., reduced tD) of the photogenerated electrons in TiO2 nanorods. 3.2.2. Hole contribution to photocurrent generation. An IMPS response is actually the comprehensive results of each of the photovoltaic steps (e.g., generation, interfacial transfer, transport and recombination of charge carriers). The shape of the IMPS response correlates tightly with the underlying photovoltaic mechanism. In inorganic semiconductors, the contribution to photocurrent from photogenerated holes is not ignorable due to their high mobility, and the IMPS spectra of their devices consist of combinational contributions from both electron and hole currents.54 As for the Sb2S3/TiO2-BHJ solar cells, the electron current ( Je) is mainly dominated by the


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transport process of the photogenerated electrons injected into TiO2 nanorods. Hence, under modulated illumination, the Je can be expressed by the normal transport mechanism,55 Je ¼

Ge ; 1 þ iotD

(1)

where Ge is the ac component of electron current generation, o is the modulation frequency, and tD is the transit time of electrons. Eqn (1) describes a single semicircle in the IV quadrant of the complex plane. However, the serious trapping of holes by defect states exists in the Sb2S3 layer.51,52 The hole current ( Jh) therein must dominantly correlate with the relaxation of the holes trapped by defect states. It has a frequency response to the modulated illumination, and can be written as56,57 bh ; (2) Jh ¼ Gh 1 1 þ ioth where Gh is the ac component of hole current generation, bh is the fraction of the holes involved in the relaxation, and th is the relaxation time of the holes. Eqn (2) describes a single semicircle in the I or III quadrant of the complex plane, depending on the sign of Gh.57 In order to account for the experimentally observed IMPS shape, we define the photocurrent ( Jph) generated in the Sb2S3/TiO2-BHJ solar cells as the summation of Je and Jh, both of which have opposite signs, Jph = Je + Jh.

(3)

Detailed calculations (Fig. S6 in ESI†) showed that the semicircles for electron transport (i.e., Je) and hole relaxation (i.e., Jh) can be obtained from eqn (3), and whether the semicircle for Jh in IMPS response is discernible depends significantly on the bh and Gh relative to Ge. In practice, of course, the reduction in bh for the degree of hole trapping should also reduce the charge recombination and bring forth the increase in both Ge and Gh. Given the appropriate values for parameter Ge, Gh and bh, one can get the calculated IMPS response in the shape of a single semicircle going from the IV to III quadrant, with a dominant time constant for electron transport and an indiscernible time constant for the relaxation of trapped holes (Fig. S6, ESI†). Therefore, the IMPS responses going from the IV into the III quadrant of the complex plane are a reasonably strong indication of the contribution to photocurrent generation from the photogenerated holes in the Sb2S3 layer. Using the relative values of the PHF and PLF crossing points on the measured IMPS responses for the cells with Tc = 350 and 400 1C (Fig. 7b), we calculate the IMPS responses within the experimental frequency range, as depicted in Fig. 8. On the simulated IMPS responses, the crossing point PHF at high frequency has the value of Gh, meaning a maximal Jh limit in the absence of relaxation loss, while the crossing point PLF at low frequency has the value of Jph (= Ge + Jh), which is the steady-state photocurrent limit determining the efficiency in practical devices (Fig. S6, ESI†). The shapes of the calculated IMPS responses agree well with experimental observations; in particular, the simulated IMPS spectra are dominated by electron transport. In respect to the IMPS response in the high frequency


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steady-state conditions, Bisquert and Vikhrenko60 showed that the apparent (trap controlled) values of De and te are related to D0 and t0 by the expressions


De ¼

Fig. 8 Calculated IMPS responses with tD = 103 s, th = 105 s, o = 1 Hz– 25 kHz. Other calculation parameters are: (J) Ge = 0.21, Gh = 0.02 and bh = 0.1; (n) Ge = 1.01, Gh = 0.08, and bh = 0.05. The inset compares the influence of the bh values on the simulated IMPS responses at high frequency, where the (K) and (m) lines are calculated with the same Ge and Gh values as the (J) and (n) lines, respectively. Note, the arrows identify the fmin points; J and n are with reference to the experimental IMPS shapes for Tc = 350 1C and 400 1C, respectively.

region with the presence of an indiscernible arc for hole relaxation as a criterion, the values of bh = 0.1 and 0.05 do make a remarkable difference (i.e., bh = 0.05 generates a more indistinguishable arc at high frequency than bh = 0.1) for the device with Tc = 400 1C, but almost no difference for the cell with Tc = 350 1C (inset to Fig. 8). From the simulated results, it is known that larger Ge and Gh but smaller bh are produced in the Sb2S3/ TiO2-BHJs solar cells with Tc = 400 1C, in comparison to that in the device with Tc = 350 1C. Note, the calculated time constant (tD = 103 s) is different from the experimental data (tD = 0.04–0.07 ms). This derivation should correlate with the simplification that our model (eqn (3)) does not consider the influence of charge carrier concentration on electron transport and the detailed kinetics of the transport process; hence, our model should only tentatively describe the IMPS shape and does not provide real dynamic characteristics in practical devices. It is also notable that the alternation of the measured IMPS responses at 6.4 kHz is observed for both samples (Fig. 7b), which is not related to the turning point for the process change from electron transport to hole relaxation (Fig. S7 in ESI†), but may be the result of the modulation frequency close to the detrapping rate constant of the electron in surface defects of TiO2 nanorods.28 3.2.3. Device performance. In comparison with Tc = 350 1C, the Sb2S3 crystallization at Tc = 400 1C leads to a reduced interfacial charge recombination (i.e., larger te) and increased electron diffusion rate (i.e., smaller tD) of the photogenerated electrons injected into TiO2 nanorods, and a larger Jh contribution (i.e., larger Gh and smaller bh) from the photogenerated holes inside the Sb2S3 layer. From the relation De = L2/tD, where De is the electron diffusion coefficient in TiO2 nanorods and L is the length of TiO2 nanorods,58 it is known that the De for the device with Tc = 400 1C is about 2 times the value for that with Tc = 350 1C. The intraband surface defects on TiO2 nanorods can strongly influence the electron transport and recombination kinetics of the photogenerated electrons by trapping/detrapping processes (Fig. 5a).59 With a multiple trapping model under non


nc gðEÞ1 D0 kB T

and

te ¼

kB T gðEÞt0 ; nc

(4)

where nc is the free electron density in the conduction band, D0 is the constant diffusion coefficient for free electrons without trapping, t0 is the constant lifetime of electrons in the absence of trapping, g(E) is the distribution of density of states from the traps, and kBT is the thermal energy. Consequently, both De and te depend on the trap occupancy and the electron concentration in conduction band. The TiO2 surface defects in the Sb2S3/TiO2-BHJ devices are actually passivated to the same concentration after the Sb2S3 crystallization at Tc = 350 and 400 1C (Fig. 4b); moreover, the solar cells were illuminated at the same incident intensity. Therefore, the reduced tD (or the increased De) upon increasing Tc clearly correlates with the change in electron concentration in the TiO2 conduction band with Sb2S3 crystallization,60–62 for which the photocurrent may act as a rough measure (Table 1). According to eqn (4), the De and te depend in opposite senses on the electron density nc, which has been clearly demonstrated in typical DSCs under different illumination intensities.62,63 The increased te accompanied by an increase in De (i.e., reduced tD) for photogenerated electrons in the TiO2 nanorod (Table 1), due to the increase in Tc from 350 to 400 1C, indicates that the factor affecting te is different from the electron concentration in the TiO2 conduction band for De in the Sb2S3/TiO2-BHJ solar cells. We think that the factor affecting te in these devices should predominantly correlate with the charge transporting property inside the bulk Sb2S3 layer related to the Sb2S3 crystallization, rather than the defects at the Sb2S3/TiO2 interface and the electron concentration in the TiO2 conduction band. The better crystallized Sb2S3 layer at Tc = 400 1C (Fig. 2) inevitably has fewer defect states to trap charge carriers, and better electronic contacts between crystalline grains, consequently resulting in the improved transporting pathways for both electrons (e) and holes (h) in the bulk Sb2S3 layer.64,65 This is also supported by the significantly reduced series resistance (Rs) (Table 1). With the improved charge transporting properties in the bulk Sb2S3 layer, the spatial e–h separation inside the Sb2S3 layer becomes more efficient for a reduced bulk charge recombination therein and the population of the trapped holes is reduced,51,52 which is further confirmed by the larger Jh contribution to the photocurrent at Tc = 400 1C (Fig. 7b and 8); meanwhile, as result of the easy escape of holes in the Sb2S3 layer away from the Sb2S3/TiO2 interface, the spatial e–h separation between the electrons in TiO2 and the holes left in Sb2S3 takes place more efficiently for a smaller charge recombination at the Sb2S3/TiO2 interface (i.e., a larger te). With the better crystallization in the Sb2S3 layer, on the other hand, the electron concentration in the TiO2 conduction band is inevitably increased for the transportation towards the collection electrode, due to the reduced interfacial and bulk charge recombinations. Also, the increased FF with increasing Tc from 350 to 400 1C is a result of the reduced charge


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recombinations subjected to significantly decreased Rs and slightly increased Rsh (Table 1). Therefore, increasing Tc from 350 to 400 1C leads to improved charge transporting properties in the bulk Sb2S3 layer in the Sb2S3/TiO2-BHJ devices, which eventually results in a much higher electron concentration in the TiO2 conduction band for a reduced tD (i.e., increased De), but a more efficient spatial e–h separation at the Sb2S3/TiO2 interface for an increased te. As indicated from the IPCE data (Fig. 6b), increasing Tc from 350 to 400 1C makes the contribution to photocurrent generation from the charge carriers in the TiO2-NA region close to the FTO collection electrode (i.e., area A1) much more efficient. This is easily understood by the significantly reduced charge recombination at the Sb2S3/TiO2 interface. The higher Jsc and IPCE in the Sb2S3/TiO2-BHJ solar cells crystallized at 400 1C are intrinsically attributed to the improved charge collection efficiency resulting from the reduced charge recombination, both at the Sb2S3/TiO2 interface and in the bulk Sb2S3 layer. Moreover, an increased Voc by ca. 80 mV is observed for increasing Tc from 350 to 400 1C (Table 1). The Voc in solar cells is generated by the splitting of the quasi-Fermi levels for both photogenerated electrons and holes, and is often greatly influenced by the interfacial defects introduced by the electron transporting channels.66 The electrons trapped by these defect states will have lower quasi-Fermi levels than in the conduction band.67–70 However, the PL data (Fig. 4b) indicate that the devices with Tc = 350 and 400 1C in fact have a similar TiO2 surface defect concentration, suggesting that the quasi-Fermi levels of the electrons in TiO2 nanorods are comparable to each other. Moreover, while the EVBM of the Sb2S3 layer is not remarkably altered (Fig. S4, ESI†), the smaller number of defect states due to the better crystallization at Tc = 400 1C will make the quasi-Fermi levels of holes in Sb2S3 closer to the valence band,71 which inevitably increases the difference between the quasi-Fermi levels of photogenerated electrons and holes for an increased Voc. 3.2.4. Remarks. We use MEH-PPV, an intrinsically amorphous conjugated polymer,72 to prepare the buffer layer in order to obtain the compact formation of the polymer/Sb2S3 interface in the Sb2S3/TiO2-BHJ solar cells.73,74 Moreover, several reports have demonstrated that Sb2S3 nanoparticles/ nanofilms can crystallize well at Tc = 300–330 1C for quite good solar cell efficiency.21,22,24,75 However, we observed that Tc = 400 1C is more suitable for efficient Sb2S3/TiO2-BHJ solar cells; the much higher Tc for better Sb2S3/TiO2-BHJ device performance in our cases may be due to the application of much thicker Sb2S3 layers and the required formation of crystalline and continuous bulk Sb2S3 layers. We also studied the Sb2S3/TiO2-BHJs crystallized at Tc temperatures lower or higher than 350–400 1C; for example, those crystallized at Tc = 300 and 450 1C. XRD data showed that the annealing at 300 1C results in a much lower Sb2S3 crystallinity with respect to the sample crystallized at 400 1C (Fig. S2, ESI†). However, the SEM results showed that thermal annealing at Tc = 450 1C removed the major part of Sb2S3 from TiO2-NA, but the Sb2S3/TiO2-BHJ crystallized at Tc = 300 1C still exhibited a


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nanostructured Sb2S3 layer (Fig. S8 in ESI†). Compared to the solar cells with Tc = 350 and 400 1C (Table 1), the solar cells based on the Sb2S3/TiO2-BHJ crystallized at 300 1C have a much lower Jsc (or Jph), much smaller Voc, much shorter electron lifetime te (= 0.48 ms), and a much longer transit time tD (= 0.42 ms) (Fig. S9 in ESI†). In particular, the IMPS spectrum of the solar cell based on the Sb2S3/TiO2-BHJ crystallized at Tc = 300 1C (Fig. S9, ESI†) is located only in the IV quadrant, indicating no photogenerated hole contribution to the photocurrent according to our theoretical model. The measured IMPS shape of the solar cell with Tc = 300 1C is completely different from those for Tc = 350 and 400 1C (Fig. 7b), suggesting that the photovoltaic mechanism in the solar cell with Tc = 300 1C is different from those with Tc = 350 and 400 1C. Our results (Fig. S9, ESI†) indicated that the Sb2S3/TiO2-BHJ solar cell with Tc = 300 1C predominantly involves the excitonic mechanism,28,47 with exciton diffusion in the Sb2S3 layer toward the Sb2S3/TiO2 interface for dissociation and no evident hole contribution to photocurrent generation due to the very serious hole trapping by the defect states in the nanostructured Sb2S3 layer, which is different from the nonexcitonic mechanism76 in the Sb2S3/TiO2-BHJ solar cells with Tc = 350 and 400 1C (Fig. 5).

4. Conclusions The solution-processed non-excitonic BHJ solar cells featuring a bulk Sb2S3 absorbing layer interdigitated with a TiO2 nanoarray as electron transporter are developed. The crystalline Sb2S3 bulk layer was prepared in the TiO2 nanoarray by a thermally induced amorphous-to-crystalline transformation of solutionprocessed amorphous material at 350 and 400 1C. The crystalline structure in the Sb2S3 absorbing layer is important for efficient devices, and the improved crystallization at Tc = 400 1C leads to a much higher electron concentration in the TiO2 conduction band for a reduced tD (i.e., increased De), and more efficient spatial e–h separations both at the Sb2S3/TiO2 interface for an increased te and inside the Sb2S3 for a reduced bulk charge recombination. The higher Jsc and FF in the solar cells with a better crystallized Sb2S3 layer predominantly originate from the improved transporting pathways for enhancing the charge collection efficiency by reducing the charge recombinations at the Sb2S3/TiO2 interface and in the bulk Sb2S3 layer, while the improved Voc mainly originates from the reduced defect state in the Sb2S3 layer for increasing the difference between the quasiFermi levels of photogenerated electrons and holes. The remarkable contribution of photogenerated holes to photocurrent generation is revealed in those solar cells, and a better crystallized Sb2S3 absorber layer produces a bigger contribution to the photocurrent from both the photogenerated holes and electrons. The amorphous-to-crystalline transformation synthesis of inorganic BHJs offers one new approach to build practical nanostructured devices, and the observation of evident hole contribution to photocurrent generation is widely applicable to solar cells with considerable hole mobility.


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Acknowledgements


This work was supported by the ‘‘100-Talent Program’’ of Chinese Academy of Sciences, the National Natural Science Foundation of China (11274307, 11474286, 91333121, 51202002 and 11074256), the Natural Science Foundation of Anhui Province (1308085ME70), and the President Foundation of Hefei Institute of Physical Sciences.

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