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Optimization of Sb2S3 Nanocrystal Concentrations in P3HT: PCBM Layers to Improve the Performance of Polymer Solar Cells

In this study, polymer solar cells were synthesized by adding Sb2S3 nanocrystals (NCs) to thin blended films with polymer poly(3-hexylthiophene)(P3HT) and [6,6]-phenyl-C61-butyric-acid-methyl-ester (PCBM) as the p-type material prepared via the spin-coating method. The purpose of this study is to investigate the dependence of polymer solar cells’ performance on the concentration of Sb2S3 nanocrystals. The effect of the Sb2S3 nanocrystal concentrations (0.01, 0.02, 0.03, and 0.04 mg/mL) in the polymer’s active layer was determined using different characterization techniques. X-ray diffraction (XRD) displayed doped ratio dependences of P3HT crystallite orientations of P3HT crystallites inside a block polymer film. Introducing Sb2S3 NCs increased the light harvesting and regulated the energy levels, improving the electronic parameters. Considerable photoluminescence quenching was observed due to additional excited electron pathways through the Sb2S3 NCs. A UV–visible absorption spectra measurement showed the relationship between the optoelectronic properties and improved surface morphology, and this enhancement was detected by a red shift in the absorption spectrum. The absorber layer’s doping concentration played a definitive role in improving the device’s performance. Using a 0.04 mg/mL doping concentration, a solar cell device with a glass /ITO/PEDOT:PSS/P3HT-PCBM: Sb2S3:NC/MoO3/Ag structure achieved a maximum power conversion efficiency of 2.72%. These Sb2S3 NCs obtained by solvothermal fabrication blended with a P3HT: PCBM polymer, would pave the way for a more effective design of organic photovoltaic devices.


Introduction
Solar cell-based organic semiconductors have many advantages, such as low cost, lightweight, flexibility, low material consumption, easy fabrication, and large area production [1,2].The organic photovoltaic (PV) solar cells continue to find a widespread application, in particularly in the following areas: solar farms, remote locations, powering stand-alone devices, powering earth-orbiting satellites and space stations, building structures such as windows and roof tiles, martial utilizations, and in aviation to power aircrafts at high altitudes. Polymers are often used as adjustment additives in dye-sensitized solar cells to provide a desirably resilient substrate, a frame structure of the semi solid state electrolytes, in addition to the pore/film formation in photoanode films. Moreover, polymers are added to enhance a solar cell device's performance, such as reinforcing the processes of crystallization and nucleation in the perovskite solar cell films. Polymers are applied as buffer layers or donor layers to improve a device's efficiency. Polymers are also applied as electron transmitters, hole transmission materials, as well as interfacial layers, which improve the carrier separation efficacy and minimize the recombination.
Poly (3-hexylthiophene) and [6,6]-phenyl-C61-butyric acid methyl ester, P 3 HT-PCBM blends are promising organic polymers that have been used as photovoltaic materials [3]. They are considered to be a promising fullerene derivative-based donor-acceptor electron material for organic solar cells. PCBM polymers are fullerene derivatives as electron acceptors for organic photovoltaics because of their high electron mobility. P 3 HT is a member of the polythiophene-conducting polymer family, in which excitation of the π-orbit electron in P 3 HT produces photovoltaic effects in the blend [4]. The blend's energy gap is approximately 1.8 eV and should exhibit a high absorption wavelength around 650 nm.
The lifetimes of organic solar cells remain short due to the degradation mechanisms that occur in organic compounds and oxidation of the electrode materials. However, strictly controlling this morphology is severely limited, which will lead to charge carrier recombinations because of incomplete pathways for both types of charge carriers if the films are thicker than 150 nm [5]. The donor-acceptor morphology can be controlled using the metal sulfide geometry, which can create good percolation pathways. Using ordered nanostructures such as nanocrystals or nanoparticles is promising for controlling the final structure, as electrons can be transported along one-dimensional structures over many micrometers, which reduces carrier recombinations [5,6]. P 3 HT forms ordered microcrystalline structures in the solid state. The presence of ordered crystalline structures in solid thin films helps to obtain high device performance because of the improved hole mobility from stacking thiophene rings and forms an enhanced light absorption with ordered structures in longer wavelength regions. Thus, incorporating inorganic and nanostructure semiconductors has several advantages, such as high electron mobility and physical and chemical stability [7,8]. However, some disadvantages are present, such as the formation of large aggregates of nanostructures, which may be effective in active layer morphology charge mobility.
Many nanostructure materials, such as PbS, ZnO, CdS, TiO 2 , and Sb 2 S 3 , have been used in polymer solar cells due to their high electron mobility. Sb 2 S 3 is a semiconductor material with a unique one-dimensional crystal structure, higher stability in air, and diverse Sb-S bond lengths. Sb 2 S 3 is a non-toxic, abundant material with an indirect band gap of 1.7-1.8 eV, an absorption coefficient higher than 104 cm −1 , and high electric conductivity, which make it a suitable material for use as light harvesters in photovoltaic applications. Sb 2 S 3 nanocrystals may provide a significant contribution to the absorption in polymer/Sb 2 S 3 nanocrystal solar cells [9][10][11]. In a study incorporating semiconductor P 3 HT:PCBM nanoparticles, Zhao et al. [12] showed that adding PbS quantum dots (QDs) to P 3 HT changed the chemical structure, which improved the active layer via an optimized phase separation and increased carrier transfers. Kim et al. [13] reported a PCE of approximately 2.98% by combining various concentrations of ZnO nanoparticles grown using the hydrothermal method with a P 3 HT: PCBM blend. They suggested that the improved charge balance and performance may have been due to reducing the charge recombinations and oxygen vacancies from the cathodes.
In this study, we report the fabrication of Sb 2 S 3 NCs for their application in one of the most efficient organic solar cells consisting of PTB7: PCBM organic photovoltaic blends. To the best of our knowledge, the effect of the Sb 2 S 3 NC additive on the properties of P 3 HT:PCBM blend solar cells that employ solid hole conductors has not been investigated. High-quality Sb 2 S 3 NPs were synthesized using solvothermal techniques at a hydrothermal temperature of 180 • C. The effect of different Sb 2 S 3 NC concentrations incorporated with PTB7: PCBM was investigated in detail using different characterization methods. The experimental results indicated that the self-assembly of the PTB7: PCBM Sb 2 S 3 NC polymer and intermolecular orientation in the P 3 HT crystallite was greatly influenced by Sb 2 S 3 nanocrystal doping. By adding Sb 2 S 3 NCs, the electrons and holes transfer in the polymer active layer enhanced our results, which led to an improved carrier separation efficiency and a reduced recombination. The absorption spectra of the sample, after adding Sb 2 S 3 NCs, was better attributed to the π-π* transitions. The Sb 2 S 3 NPs help to disperse the P3HT chains in a solution and promote transformation and, thus, improve the crystallization of P3HT (during film forming process) as well as the device's efficiency. The solar cell device in a structure glass/ITO/PEDOT:PSS/P 3 HT: PCBM:Sb 2 S 3 :NCs/MoO 3 /Ag achieved a maximum power conversion efficiency of 2.72% using an Sb 2 S 3 NC concentration of 0.04 mg/mL. Hence, Sb 2 S 3 NC-doped P3HT-PCBM thin films have great potential applications as active layers in solar cell devices and can be an efficient method to improve thin film properties.

Sb 2 S 3 Nanocrystals Fabrication
For the typical solvothermal fabrication of Sb 2 S 3 nanocrystals, 0.2 mmol of SbCl 3 and 4 mmol of I-cystine were dissolved in 10mL of oleylamine. Then, 4 mmol of thiourea was dissolved in 10 mL of oleylamine in a separate beaker under stirring for 2 h. The antimony solution was dropped slowly into the latter solution under vigorous magnetic stirring. The mixed solutions were then transferred into a 50-milliliter Teflon-lined, stainless steel autoclave and placed in a tubular furnace at 180 • C for 24 h. The resulting precipitate was first centrifuged, and the black product was washed with ethyl alcohol and deionized water and dried at 70 • C for 5 h.

Device Fabrication
Indium tin oxide (ITO)-coated glass substrates, approximately 120 nm thick with a sheet resistance of around 15 Ω/sq and a 1.5 × 1.5 cm 2 device area were used as an anode contact in the organic solar cell device. Acetone or isopropanol were soaked in an ultrasonic bath for 5 min, followed by drying in an N 2 -filled glove box. A PEDOT:PSS polymer was deposited at 3500 rpm for 50 s, using spin coating followed by annealing at 120 • C for 15 min in a furnace in air, resulting in a thickness of ∼60 nm. The PE-DOT:PSS/ITO/glass substrates were moved to a nitrogen-supplied glove box and annealed again at 130 • C for 10 min to avoid humidity. The blended P 3 HT and PCBM solution was prepared by dissolving the polymers in 1,2-dichlorobenzene with a ratio of 1:1 and 50 mg of each polymer. The mixed solutions were stirred at 60 • C overnight. The Sb 2 S 3 nanocrystals were blended in the P 3 HT: PCBM in a concentration range of 0.01, 0.02, 0.03, and 0.04 mg/mL. The P 3 HT:PCBM:Sb 2 S 3 :NC nanocomposites were spin-coated on top of PEDOT:PSS layers at 1500 rpm for 40 s inside a glove box. The active layer's average thickness was 100-120 nm, and the spin-coated layer was annealed at 140 • C for 10 min in a nitrogen-filled glove box. A buffer layer, with approximately 20 nm of MoO 3 , was deposited using RF spurting. The thick Ag electrode (70 nm) was thermally evaporated, and the active area was approximately 0.1 cm 2 . The final device had the following structure: ITO/PEDOT:PSS/P 3 HT:PCBM:Sb 2 S 3 :NCs/MoO 3 /Ag. A schematic illustration of a typical device structure is shown in Figure 1.

Characterization
The structural properties of the Sb2S3 nanocrystals and P3HT:PCBM:Sb2S3 NCs were investigated using an X-ray diffractometer (MiniFlex, Rigaku, Tokyo, Japan) with monochromatic Cu Kα radiation (λ = 1.5405 Å). The UV-visible spectroscopy investigation was performed using a Varian Cary 100 spectrophotometer (Agilent Technologies, Santa Clara, California, USA). A transmission electron microscopy (TEM) analysis was obtained using a JEOL 2010F (JEOL 2010F field emission high resolution scanning/transmission electron microscope, Akishima, Tokyo, Japan) to examine the Sb2S3 nanocrystal's properties. Atomic force microscopy (AFM) measurements were performed on a standard Keysight 5500 scanning probe microscope (Keysight Technologies Fountain grove Parkway, Santa Rosa, CA 95403) in the intermittent contact mode in air. The photoluminescence spectra were collected using a Hitachi F-7000 spectrometer (Hitachi F-7000 fluorescence spectrophotometer, Toranomo, Minato-ku, Tokyo, Japan) equipped with a red-sensitive detector. The Fourier-transform infrared spectroscopy (FTIR) of the polymer/NC blended films was investigated using a Nicolet 8700 (Thermo Fisher Scientific, Madison, WI, USA) spectrometer. Raman spectroscopy of the active layer was performed using a Renishaw inVia Raman microscope (Renishaw inVia Raman microscope, Gloucestershire, United Kingdom) (λ = 514 nm). The current density-voltage (J-V) characteristics of the polymer device were measured under AM 1.5 G and 100 mW cm 2 illumination using a Keithley 2400 (Keithley, Tektronix, Solon, Ohio, USA) for source measurement.

Results and Discussion
The crystal sizes and phase information on the Sb2S3 were confirmed using X-ray diffraction (XRD) patterns from the pure Sb2S3 nanocrystals prepared using the solvothermal method as shown in Figure 2a. All of the XRD patterns in Figure 2a

Characterization
The structural properties of the Sb 2 S 3 nanocrystals and P 3 HT:PCBM:Sb 2 S 3 NCs were investigated using an X-ray diffractometer (MiniFlex, Rigaku, Tokyo, Japan) with monochromatic Cu Kα radiation (λ = 1.5405 Å). The UV-visible spectroscopy investigation was performed using a Varian Cary 100 spectrophotometer (Agilent Technologies, Santa Clara, California, USA). A transmission electron microscopy (TEM) analysis was obtained using a JEOL 2010F (JEOL 2010F field emission high resolution scanning/transmission electron microscope, Akishima, Tokyo, Japan) to examine the Sb 2 S 3 nanocrystal's properties. Atomic force microscopy (AFM) measurements were performed on a standard Keysight 5500 scanning probe microscope (Keysight Technologies Fountain grove Parkway, Santa Rosa, CA 95403) in the intermittent contact mode in air. The photoluminescence spectra were collected using a Hitachi F-7000 spectrometer (Hitachi F-7000 fluorescence spectrophotometer, Toranomo, Minato-ku, Tokyo, Japan) equipped with a red-sensitive detector. The Fourier-transform infrared spectroscopy (FTIR) of the polymer/NC blended films was investigated using a Nicolet 8700 (Thermo Fisher Scientific, Madison, WI, USA) spectrometer. Raman spectroscopy of the active layer was performed using a Renishaw inVia Raman microscope (Renishaw inVia Raman microscope, Gloucestershire, United Kingdom) (λ = 514 nm). The current density-voltage (J-V) characteristics of the polymer device were measured under AM 1.5 G and 100 mW cm 2 illumination using a Keithley 2400 (Keithley, Tektronix, Solon, Ohio, USA) for source measurement.

Results and Discussion
The crystal sizes and phase information on the Sb 2 S 3 were confirmed using X-ray diffraction (XRD) patterns from the pure Sb 2 S 3 nanocrystals prepared using the solvothermal method as shown in Figure 2a (511) orientations, respectively, with a preferred orientation along the (130) plane. The sharp and high-intensity peaks indicated that the product had high crystallinity. No peaks were associated with other phases, which indicated the sample's high purity. The lattice parameter identical to the orthorhombic type was a = 11.23 Å, b = 11.28 Å, and c = 3.83 Å. These results agreed with previously reported studies of Sb 2 S 3 materials [14,15]. The crystallite size of the Sb 2 S 3 nanocrystals was 41 nm, as determined using the Scherrer formula [16]. peaks were associated with other phases, which indicated the sample's high purity. The lattice parameter identical to the orthorhombic type was a = 11.23 Å, b = 11.28 Å, and c = 3.83 Å. These results agreed with previously reported studies of Sb2S3 materials [14,15]. The crystallite size of the Sb2S3 nanocrystals was 41 nm, as determined using the Scherrer formula [16]. To assess the crystalline development of the P3HT:PCBM: Sb2S3 NCs' active layer, the sample's XRD patterns were recorded. Figure 2b shows the XRD profiles of P3HT:PCBM: Sb2S3 NC blended films spin-coated at 1500 rpm for 40 s. The XRD analysis was recorded in a narrow range (2θ = 3-10°). The increase in the (100) peak intensity corresponding to P3HT was observed at 2θ = 5.4°, which agreed with prior studies [17,18]. The diffraction peak located at 2θ = 19° corresponded to crystalline PCBM [19]. The XRD (100) peaks shifted to lower angles, from 5.58 to 5.39°, as the dispersion degree increased. This shift (change in D spacing) indicated the improved diffusion of the PCBM into the P3HT, decreasing the distance between them [20]. The diffraction peaks became narrower and sharper as the Sb2S3NCs' concentration increased, indicating an increase in the crystallinity, suggesting orderliness, and increasing the intermolecular π plane. The P3HT:PCBM: Sb2S3 NCs' blend displayed a high crystallinity that improved regardless of the presence of fullerene. The growth of the P3HT aggregate improved the chains' crystallinity and hole mobility, which resulted in positive J-V curves [21].
Raman spectroscopy in a range of 250 to 2500 cm −1 was used to investigate the molecules' irrational mode in the P3HT:PCBM:Sb2S3 NC blended films. The Raman spectra of the P3HT:PCBM: Sb2S3 NC blended films shown in Figure 3 features the vibration modes, as reported in previous studies, including the vibration spectra, and the stretching and bending modes with different relative intensities varying from high to low values. The amount of Sb2S3 NCs contributed to the changes in the peak intensity. The peaks at ∼578 to ∼782 cm −1 coincided with C-H out-of-phase bending [22]. As most studies of polymers suggested, the peak located at 728cm −1 was related to rocking vibrations in the C-S-C thiophene ring of the P3HT [22,23]. The peak found at ∼1086 cm −1 was not from the P3HT molecules; therefore, this vibration peak may have corresponded to interactions of the P3HT molecules and the Sb2S3 NCs. The peak located at ∼1378 cm −1 was associated with the asymmetric vibrations of C=C skeletal stretching deformation, and the peak appearing at∼1518 cm −1 corresponded to stretching vibrations from the P3HT,which agrees with a previous study [24]. A higher intensity peak at ∼1442 cm −1 was related to the high P3HT structural order [25]. At this peak, the combined skeletal stretching of the complete chain, or at least a large part of it, was found. This mode was dominated by the inter-ring C-C stretching vibration mode and was the origin of the vi- To assess the crystalline development of the P 3 HT:PCBM: Sb 2 S 3 NCs' active layer, the sample's XRD patterns were recorded. Figure 2b shows the XRD profiles of P 3 HT:PCBM: Sb 2 S 3 NC blended films spin-coated at 1500 rpm for 40 s. The XRD analysis was recorded in a narrow range (2θ = 3-10 • ). The increase in the (100) peak intensity corresponding to P 3 HT was observed at 2θ = 5.4 • , which agreed with prior studies [17,18]. The diffraction peak located at 2θ = 19 • corresponded to crystalline PCBM [19]. The XRD (100) peaks shifted to lower angles, from 5.58 to 5.39 • , as the dispersion degree increased. This shift (change in D spacing) indicated the improved diffusion of the PCBM into the P 3 HT, decreasing the distance between them [20]. The diffraction peaks became narrower and sharper as the Sb 2 S 3 NCs' concentration increased, indicating an increase in the crystallinity, suggesting orderliness, and increasing the intermolecular π plane. The P 3 HT:PCBM: Sb 2 S 3 NCs' blend displayed a high crystallinity that improved regardless of the presence of fullerene. The growth of the P 3 HT aggregate improved the chains' crystallinity and hole mobility, which resulted in positive J-V curves [21].
Raman spectroscopy in a range of 250 to 2500 cm −1 was used to investigate the molecules' irrational mode in the P 3 HT:PCBM:Sb 2 S 3 NC blended films. The Raman spectra of the P 3 HT:PCBM: Sb 2 S 3 NC blended films shown in Figure 3 features the vibration modes, as reported in previous studies, including the vibration spectra, and the stretching and bending modes with different relative intensities varying from high to low values. The amount of Sb 2 S 3 NCs contributed to the changes in the peak intensity. The peaks at ∼578 to ∼782 cm −1 coincided with C-H out-of-phase bending [22]. As most studies of polymers suggested, the peak located at 728 cm −1 was related to rocking vibrations in the C-S-C thiophene ring of the P 3 HT [22,23]. The peak found at ∼1086 cm −1 was not from the P 3 HT molecules; therefore, this vibration peak may have corresponded to interactions of the P 3 HT molecules and the Sb 2 S 3 NCs. The peak located at ∼1378 cm −1 was associated with the asymmetric vibrations of C=C skeletal stretching deformation, and the peak appearing at∼1518 cm −1 corresponded to stretching vibrations from the P 3 HT, which agrees with a previous study [24]. A higher intensity peak at ∼1442 cm −1 was related to the high P 3 HT structural order [25]. At this peak, the combined skeletal stretching of the complete chain, or at least a large part of it, was found. This mode was dominated by the inter-ring C-C stretching vibration mode and was the origin of the vibronic structure of the absorption and emission spectra [25,26]. The intensity of the ∼1442 cm −1 peak related to the phonon features was reduced due to the presence of the Sb 2 S 3 NCs, suggesting an increased order in the blended system and enhanced conjugation lengths. bronic structure of the absorption and emission spectra [25,26]. The intensity of the ∼1442 cm −1 peak related to the phonon features was reduced due to the presence of the Sb2S3 NCs, suggesting an increased order in the blended system and enhanced conjugation lengths. To obtain accurate morphological and size information, TEM images and diffraction patterns were obtained for the sample. The optical absorption spectra of the Sb2S3 nanocrystals recorded in a range of 200-800 nm is shown in Figure 5a. The Sb2S3 nanocrystals had high absorption coefficients above 5 × 10 4 cm −1 in a wavelength range of 350-750 nm. The strong absorption intensity was attributed to the good quality of the Sb2S3 nanocrystals. A large absorption coefficient is important for solar cell applications, which implies a high short-circuit current density. The energy gaps of the Sb2S3 NCs were calculated by plotting the αhv 1/2 Figure 3. Raman spectra of theP 3 HT:PCBM:Sb 2 S 3 NC blended thin films prepared using spin coating.
To obtain accurate morphological and size information, TEM images and diffraction patterns were obtained for the sample. Figure 4a-c show TEM images of the Sb 2 S 3 nanocrystals prepared using the solvothermal method at 230 • C. The typical TEM images shown in Figure 4a confirm that the morphology of the Sb 2 S 3 was a nanocrystalline along the (001) direction with a diameter of approximately 100 nm. The selected area electron diffraction SAED images in Figure 4b show the sharp spots of the Sb 2 S 3 , where the diffraction spots indicate the fully crystalline nature of the Sb 2 S 3 NCs and the main diffraction plane matched with the standard XRD patterns shown in Figure 2. The high-resolution TEM (HRTEM) images are shown in Figure 4c and the lattice spacing of 0.32 nm agrees well with the (130) diffraction of the Sb 2 S 3 . The SAED images and corresponding HRTEM images demonstrated that the crystalline nature of the Sb 2 S 3 NCs agreed with the XRD and Raman results.
The optical absorption spectra of the Sb 2 S 3 nanocrystals recorded in a range of 200-800 nm is shown in Figure 5a. The Sb 2 S 3 nanocrystals had high absorption coefficients above 5 × 10 4 cm −1 in a wavelength range of 350-750 nm. The strong absorption intensity was attributed to the good quality of the Sb 2 S 3 nanocrystals. A large absorption coefficient is important for solar cell applications, which implies a high short-circuit current density. The energy gaps of the Sb 2 S 3 NCs were calculated by plotting the αhv 1/2 vs. hv and extrapolating the linear portion of the curve to αhv = 0, as shown in Figure 5a (insert). The optical bandgap (Eg) determined using the Tauc equation [27] was 1.79 eV. vs. hv and extrapolating the linear portion of the curve to αhv = 0, as shown in Figure 5a (insert).The optical bandgap (Eg) determined using the Tauc equation [27] was 1.79 eV. To investigate the photon-gathering ability of the P3HT:PCBM:Sb2S3 NCs' active layer, UV-visible absorption spectra of the sample were recorded in a wavelength range of 300-800 nm. Figure 5b shows the UV-visible absorption spectra of the P3HT:PCBM:Sb2S3 NC blended films prepared using different concentrations of Sb2S3 NCs (0.01, 0.02, 0.03, and 0.04 mg/mL). The absorption spectra in a range of 450-600 nm were attributed to the main P3HT polymer. The absorption spectra of the P3HT:PCBM:Sb2S3NCs' active layer showed an absorption peak of π-π* aggregate formation at a wavelength of 512 nm, with two small shoulders at 550 and 604 nm. The P3HT polymer displayed an atypical absorption band at approximately 512 nm that was attributed to the π-π* transitions, this result is in good agreement with previous studies [28,29]. The π-π* band of the active layer showed a 3-nanometer red shift corresponding to the P3HT polymer, suggesting more efficient π stacking. Sb2S3 NCs had absorption spectra below 550 nm in previous studies. Sb2S3 NC doping may have strongly contributed to the buffer absorbance within a 300-550 nm range. Doping resulted in a clear red shift of the optical absorption, mainly in the 512 nm band, which shifted from 500 to 512 nm. The red shift in the absorption band could be attributed to increasing the π electron delocalization, lowering the energy band of the π and π*, and improving the optical π-π* transitions [30,31]. The increase in the light-harvesting properties of the P3HT:PCBM:Sb2S3 NC blend improved the photo-generated carriers and enhanced the charge transport due to the π-π interactions between the Sb2S3 NCs and the P3HT molecules. The absorption strength was increased, along with the different concentrations, due to the improved polymer crystallinity. However, the improved absorption may have been due to the decreased film roughness, which affected the light scattering in the blended films, resulting in increased absorption in the active layer. Figure 6 shows a schematic energy level diagram of the energy and charge transfer effects of the P3HT:PCBM:Sb2S3 NC blend. There are the following three possible reasons for the electron movement in the P3HT:PCBM:Sb2S3 NC system: electron transport from To investigate the photon-gathering ability of the P 3 HT:PCBM:Sb 2 S 3 NCs' active layer, UV-visible absorption spectra of the sample were recorded in a wavelength range of 300-800 nm. Figure 5b shows the UV-visible absorption spectra of the P 3 HT:PCBM:Sb 2 S 3 NC blended films prepared using different concentrations of Sb 2 S 3 NCs (0.01, 0.02, 0.03, and 0.04 mg/mL). The absorption spectra in a range of 450-600 nm were attributed to the main P 3 HT polymer. The absorption spectra of the P 3 HT:PCBM:Sb 2 S 3 NCs' active layer showed an absorption peak of π-π* aggregate formation at a wavelength of 512 nm, with two small shoulders at 550 and 604 nm. The P 3 HT polymer displayed an atypical absorption band at approximately 512 nm that was attributed to the π-π* transitions, this result is in good agreement with previous studies [28,29]. The π-π* band of the active layer showed a 3-nanometer red shift corresponding to the P 3 HT polymer, suggesting more efficient π stacking. Sb 2 S 3 NCs had absorption spectra below 550 nm in previous studies. Sb 2 S 3 NC doping may have strongly contributed to the buffer absorbance within a 300-550 nm range. Doping resulted in a clear red shift of the optical absorption, mainly in the 512 nm band, which shifted from 500 to 512 nm. The red shift in the absorption band could be attributed to increasing the π electron delocalization, lowering the energy band of the π and π*, and improving the optical π-π* transitions [30,31]. The increase in the light-harvesting properties of the P 3 HT:PCBM:Sb 2 S 3 NC blend improved the photogenerated carriers and enhanced the charge transport due to the π-π interactions between the Sb 2 S 3 NCs and the P 3 HT molecules. The absorption strength was increased, along with the different concentrations, due to the improved polymer crystallinity. However, the improved absorption may have been due to the decreased film roughness, which affected the light scattering in the blended films, resulting in increased absorption in the active layer. Figure 6 shows a schematic energy level diagram of the energy and charge transfer effects of the P 3 HT:PCBM:Sb 2 S 3 NC blend. There are the following three possible reasons for the electron movement in the P 3 HT:PCBM:Sb 2 S 3 NC system: electron transport from the P 3 HT to the PCBM; electron transfer from the P 3 HT to the Sb 2 S 3 NCs; and electrons moving from the P 3 HT to the Sb 2 S 3 NCs, and then to the PCBM. In these cases, the electrons transformed into Ag, while the holes transformed out of the P 3 HT to the ITO layer, which reduced the chances of carrier recombination. Thus, incorporating the Sb 2 S 3 NCs with the polymer contributed to more photo-induced charge carrier separations/transfers that increased the photo-generated exciton dissociation [8,32]. The processes can occur depending on the excitation energy. When excitons are formed upon light absorption in the Sb 2 S 3 NCs, it is expected that the electrons will be transferred to PCBM and holes to P3HT. When excitations are generated in P3HT, we can predict from the energy levels that an electron transfer will occur toward the Sb 2 S 3 NCs and/or PCBM [33]. To learn more about the processes occurring between the blend components, we investigated steady state and time-resolved PL. the P3HT to the PCBM; electron transfer from the P3HT to the Sb2S3 NCs; and electrons moving from the P3HT to the Sb2S3 NCs, and then to the PCBM. In these cases, the electrons transformed into Ag, while the holes transformed out of the P3HT to the ITO layer, which reduced the chances of carrier recombination. Thus, incorporating the Sb2S3 NCs with the polymer contributed to more photo-induced charge carrier separations/transfers that increased the photo-generated exciton dissociation [8,32].The processes can occur depending on the excitation energy. When excitons are formed upon light absorption in the Sb2S3 NCs, it is expected that the electrons will be transferred to PCBM and holes to P3HT. When excitations are generated in P3HT, we can predict from the energy levels that an electron transfer will occur toward the Sb2S3 NCs and/or PCBM [33]. To learn more about the processes occurring between the blend components, we investigated steady state and time-resolved PL.   the P3HT to the PCBM; electron transfer from the P3HT to the Sb2S3 NCs; and electrons moving from the P3HT to the Sb2S3 NCs, and then to the PCBM. In these cases, the elec trons transformed into Ag, while the holes transformed out of the P3HT to the ITO layer which reduced the chances of carrier recombination. Thus, incorporating the Sb2S3 NCs with the polymer contributed to more photo-induced charge carrier separations/transfers that increased the photo-generated exciton dissociation [8,32].The processes can occu depending on the excitation energy. When excitons are formed upon light absorption in the Sb2S3 NCs, it is expected that the electrons will be transferred to PCBM and holes to P3HT. When excitations are generated in P3HT, we can predict from the energy levels that an electron transfer will occur toward the Sb2S3 NCs and/or PCBM [33]. To learn more about the processes occurring between the blend components, we investigated steady state and time-resolved PL.   To study the effect of the Sb 2 S 3 NCs on the exciton dissociation, the PL spectra of the P 3 HT:PCBM:Sb 2 S 3 NC thin films with different concentrations of Sb 2 S 3 NCs were investigated using a PL spectral system, as shown in Figure 7. The PL peak of the P 3 HT:PCBM:Sb 2 S 3 NCs was observed at 630 nm, in good agreement with previous studies for P 3 HT:PCBM [34,35]. At higher Sb 2 S 3 NC concentrations, the peaks become more intense and emission peaks at 635nm red-shifted to 639 nm, which may have been related to band-to-band emission in the P 3 HT, since it had a narrow band gap of approximately 1.9 eV.
The increased PL intensity indicated the improved phase separation size, which benefited the charge transport and collection. The size range of the P 3 HT and PCBM exceeded the exciton diffusion length because of the increased phase separation between the PCBM and P 3 HT after doping. A large increase in the PL intensity was observed, consistent with the increased diffusion of the PCBM in the P 3 HT matrix, leading to increased carrier transfers [36]. Consequently, the exciton-dissociation efficiency decreased, while the PL efficiency increased. These results indicated that the energy and charge transfer occurred between the P 3 HT, PCBM, and Sb 2 S 3 NC material.
To study the effect of the Sb2S3 NCs on the exciton dissociation, the PL spectra of the P3HT:PCBM:Sb2S3 NC thin films with different concentrations of Sb2S3 NCs were investigated using a PL spectral system, as shown in Figure 7. The PL peak of the P3HT:PCBM:Sb2S3 NCs was observed at 630 nm, in good agreement with previous studies for P3HT:PCBM [34,35]. At higher Sb2S3 NC concentrations, the peaks become more intense and emission peaks at 635nm red-shifted to 639 nm, which may have been related to band-to-band emission in the P3HT, since it had a narrow band gap of approximately 1.9 eV.
The increased PL intensity indicated the improved phase separation size, which benefited the charge transport and collection. The size range of the P3HT and PCBM exceeded the exciton diffusion length because of the increased phase separation between the PCBM and P3HT after doping. A large increase in the PL intensity was observed, consistent with the increased diffusion of the PCBM in the P3HT matrix, leading to increased carrier transfers [36]. Consequently, the exciton-dissociation efficiency decreased, while the PL efficiency increased. These results indicated that the energy and charge transfer occurred between the P3HT, PCBM, and Sb2S3 NC material. Atomic force microscopy (AFM) was used to investigate the tapping mode operation and phase distribution of the blends. Figure 8a-d shows AFM height images of the surface roughness and grain size of the P3HT:PCBM:Sb2S3 NC thin films prepared with different Sb2S3 NC concentrations. The images show the network structures of self-organized P3HT chains similar to the seed-like polymer chains reported in previous studies [37,38]. The bright areas (higher phases) in the phase images can be identified as P3HT-rich regions, while the dark areas (lower phases) can be identified as PCBM-rich regions. The phase images of the active layer display a considerable variation in the phase segregation length scale. The size of the P3HT area and PCBM domains, in a range of 24.3 to 59.27 nm, depended on the NC concentration. This indicated that two sequential polymer chains were separated within the exciton diffusion length, with PCBM nanoclusters established between the chains.
Evidently, the Sb2S3 NC acted as a compatibilizer and modified the average domain size of the PCBM by the intermolecular hydrogen bonds generated from the C-H-Os Atomic force microscopy (AFM) was used to investigate the tapping mode operation and phase distribution of the blends. Figure 8a-d shows AFM height images of the surface roughness and grain size of the P 3 HT:PCBM:Sb 2 S 3 NC thin films prepared with different Sb 2 S 3 NC concentrations. The images show the network structures of self-organized P 3 HT chains similar to the seed-like polymer chains reported in previous studies [37,38]. The bright areas (higher phases) in the phase images can be identified as P 3 HT-rich regions, while the dark areas (lower phases) can be identified as PCBM-rich regions. The phase images of the active layer display a considerable variation in the phase segregation length scale. The size of the P 3 HT area and PCBM domains, in a range of 24.3 to 59.27 nm, depended on the NC concentration. This indicated that two sequential polymer chains were separated within the exciton diffusion length, with PCBM nanoclusters established between the chains.
Evidently, the Sb 2 S 3 NC acted as a compatibilizer and modified the average domain size of the PCBM by the intermolecular hydrogen bonds generated from the C-H-Os bonds; therefore, the active layer became smoother. The increase in the surface roughness could also have been due to the improved crystallinity of the P 3 HT in the film. This occurred because of the decreased P 3 HT aggregation size and increased number of P 3 HT single chains. The P 3 HT crystallization degree and effective phase blending had a major influence on the electron and hole mobility in the photovoltaic blend. The P 3 HT crystallization degree and effective phase blending had a major influence on the electron and hole mobility, which enhanced the carrier recombination and reduced the recombination layer of the solar cell device. As reported in prior studies, a high-efficiency photovoltage has a high surface roughness in P 3 HT:PCBM blended thin films [39].
bonds; therefore, the active layer became smoother. The increase in the surface roughness could also have been due to the improved crystallinity of the P3HT in the film. This occurred because of the decreased P3HT aggregation size and increased number of P3HT single chains. The P3HT crystallization degree and effective phase blending had a major influence on the electron and hole mobility in the photovoltaic blend. The P3HT crystallization degree and effective phase blending had a major influence on the electron and hole mobility, which enhanced the carrier recombination and reduced the recombination layer of the solar cell device. As reported in prior studies, a high-efficiency photovoltage has a high surface roughness in P3HT:PCBM blended thin films [39].  Figure 9 displays the Fourier transform infrared (FT-IR) spectra recorded in the 250-4000 cm −1 spectral regions in the P3HT:PCBM:Sb2S3 NC blended films. In the 2700-3100 cm −1 spectral region, the methylene and methyl stretching vibration bands from the hexyl side chains of the thiophene rings were notable. The P3HT polymer exhibited vibration bands at 2954 and 3054 cm −1 corresponding to the stretching and asymmetric vibrations of =C-H and C=C [40,41]. The band at 1260 cm −1 corresponded to the dipole-derivative vector perpendicular to the ring plane, and the band near 1048 cm −1 was due to C-H in-plane bending [42]. The vibration mode at 1454 cm −1 was due to the deformation vibrations of the CH and CH3 [41]. The vibration modes in the 1600-1800 cm −1 region, such as 1715 cm −1 ,were due to photo-degradation products, including C=O groups [43].The vibration mode at 1107 cm −1 corresponded to O=S stretching due to sulfonic esters [44]. The band near 820 cm −1 was related to the asymmetric deformation of  Figure 9 displays the Fourier transform infrared (FT-IR) spectra recorded in the 250-4000 cm −1 spectral regions in the P 3 HT:PCBM:Sb 2 S 3 NC blended films. In the 2700-3100 cm −1 spectral region, the methylene and methyl stretching vibration bands from the hexyl side chains of the thiophene rings were notable. The P 3 HT polymer exhibited vibration bands at 2954 and 3054 cm −1 corresponding to the stretching and asymmetric vibrations of =C-H and C=C [40,41]. The band at 1260 cm −1 corresponded to the dipolederivative vector perpendicular to the ring plane, and the band near 1048 cm −1 was due to C-H in-plane bending [42]. The vibration mode at 1454 cm −1 was due to the deformation vibrations of the CH and CH 3 [41]. The vibration modes in the 1600-1800 cm −1 region, such as 1715 cm −1 , were due to photo-degradation products, including C=O groups [43]. The vibration mode at 1107 cm −1 corresponded to O=S stretching due to sulfonic esters [44]. The band near 820 cm −1 was related to the asymmetric deformation of CH3 vibrations (aromatic out-of-plain vibrations) [45]. Finally, the band at 1591cm −1 was due to C=C bands [46]. CH3 vibrations (aromatic out-of-plain vibrations) [45]. Finally, the band at 1591cm −1 was due to C=C bands [46]. Our results were in good agreement with previous studies of FT-IR vibrations of pristine P3HT:PCBM. Very considerable changes suggested chemical degradation, as demonstrated after doping, leading to an essential increase in the peaks' intensity in the vibration bands at 820, 2954, 2363, and 1260 cm −1 , features of the hexyl side chains. This indicated that the hexyl side chains began to separate from the thiophene rings and eventually volatilized. As a result, as the doping concentration increased, the thiophene rings strengthened, affecting the increase in the P3HT band's intensity, with a small shift related to the improved polymer chains leading to the phase separation of the P3HT:PCBM blend. The FT-IR results described the phase separation processes in the blend and vibration modes of the polymer chemical groups.
To evaluate the influence of the Sb2S3 NCs on the performance of theP3HT:PCBM:Sb2S3 NCs'photovoltage, four polymer devices with different concentrations of Sb2S3 NC-doped active layers (displayed in Figure 10) were fabricated. Figure 10 shows the current-voltage (IV) curves of the devices under AM 1.5G simulation and 100 mW/cm 2 illumination. Table 1 summarizes the results of the open-circuit voltage (Voc), short-circuit current (Isc), fill factor (FF), and solar energy conversion efficiency (η) of the devices obtained with differentSb2S3 NC concentrations. The IV curves display the increases in the photocurrent that corresponded to the presence of theSb2S3 NCs in the active layer. In reference to the doping concentration of the Sb2S3 NCs shown in Table 1, 0.04 mg/mL was the optimum weight concentration that produced the highest efficiency η = 2.72% related to Isc = 10.04 mA/cm 2 , Voc = 412 mV, and FF = 66%. Our results were in good agreement with previous studies of FT-IR vibrations of pristine P 3 HT:PCBM. Very considerable changes suggested chemical degradation, as demonstrated after doping, leading to an essential increase in the peaks' intensity in the vibration bands at 820, 2954, 2363, and 1260 cm −1 , features of the hexyl side chains. This indicated that the hexyl side chains began to separate from the thiophene rings and eventually volatilized. As a result, as the doping concentration increased, the thiophene rings strengthened, affecting the increase in the P 3 HT band's intensity, with a small shift related to the improved polymer chains leading to the phase separation of the P 3 HT:PCBM blend. The FT-IR results described the phase separation processes in the blend and vibration modes of the polymer chemical groups.
To evaluate the influence of the Sb 2 S 3 NCs on the performance of theP 3 HT:PCBM:Sb 2 S 3 NCs'photovoltage, four polymer devices with different concentrations of Sb 2 S 3 NC-doped active layers (displayed in Figure 10) were fabricated. Figure 10 shows the current-voltage (IV) curves of the devices under AM 1.5G simulation and 100 mW/cm 2 illumination. Table 1 summarizes the results of the open-circuit voltage (V oc ), short-circuit current (I sc ), fill factor (FF), and solar energy conversion efficiency (η) of the devices obtained with differentSb 2 S 3 NC concentrations. The IV curves display the increases in the photocurrent that corresponded to the presence of theSb 2 S 3 NCs in the active layer. In reference to the doping concentration of the Sb 2 S 3 NCs shown in Table 1, 0.04 mg/mL was the optimum weight concentration that produced the highest efficiency η = 2.72% related to I sc = 10.04 mA/cm 2 , V oc = 412 mV, and FF = 66%. mination with light power intensity of 100 mW/cm 2 .  Generally, the increase in the FF and Voc was due to the increase in the Rsh and decrease in the Rs. The device-resistant reduction was primarily attributed to the improved morphology in the active layer. The significantly increased Jsc and FF may have been due to a new network of Sb2S3 NCs that facilitated electron transport in the polymer's active layer and a small range of PCBM aggregation that decreased the carrier recombination losses and increased the current density. In addition, the improved photocurrent corresponded to enhanced entrapment and light absorption, which were demonstrated by an energetic disorder or the improved crystallinity of the P3HT. The increase in the Jsc may also have been due to the increased incorporation of the Sb2S3 NCs in the P3HT, which demonstrated that the Sb2S3 NCs could work as electron acceptors comparable to the PCBM, which helped separate the bound photo-generated excitons. The decreases in Voc with a lower concentration of Sb2S3 may be due to the weak interaction that occurs between the polymer and the nanocrystals' interface, which leads to the incomplete charge transfer between the polymer and the nanocrystals, that could decrease the transport of charge-carriers in the hybrid solar cell. This mechanism could have caused many free electrical charges, increased the short-circuit current density, and thus improved the power conversion efficiency. Thus, Sb2S3 NCs blended in a polymer matrix provide a large interfacial area for fast charge dissociation at the interface to ensure that a maximum number of charge carriers contribute to improve the short current density. The performances of fabricated devices can be improved further by incorporating a thin layer of MoO3 as a hole transport layer between the hybrid blend and the back electrodes' Ag.  Generally, the increase in the FF and V oc was due to the increase in the R sh and decrease in the R s . The device-resistant reduction was primarily attributed to the improved morphology in the active layer. The significantly increased J sc and FF may have been due to a new network of Sb 2 S 3 NCs that facilitated electron transport in the polymer's active layer and a small range of PCBM aggregation that decreased the carrier recombination losses and increased the current density. In addition, the improved photocurrent corresponded to enhanced entrapment and light absorption, which were demonstrated by an energetic disorder or the improved crystallinity of the P 3 HT. The increase in the J sc may also have been due to the increased incorporation of the Sb 2 S 3 NCs in the P 3 HT, which demonstrated that the Sb 2 S 3 NCs could work as electron acceptors comparable to the PCBM, which helped separate the bound photo-generated excitons. The decreases in V oc with a lower concentration of Sb 2 S 3 may be due to the weak interaction that occurs between the polymer and the nanocrystals' interface, which leads to the incomplete charge transfer between the polymer and the nanocrystals, that could decrease the transport of charge-carriers in the hybrid solar cell. This mechanism could have caused many free electrical charges, increased the short-circuit current density, and thus improved the power conversion efficiency. Thus, Sb 2 S 3 NCs blended in a polymer matrix provide a large interfacial area for fast charge dissociation at the interface to ensure that a maximum number of charge carriers contribute to improve the short current density. The performances of fabricated devices can be improved further by incorporating a thin layer of MoO 3 as a hole transport layer between the hybrid blend and the back electrodes' Ag.

Conclusions
In conclusion, we successfully synthesized highly crystalline Sb 2 S 3 NCs using the solvothermal method. We synthesized bulk heterojunction organic solar cells based on P 3 HT:PCBM:Sb 2 S 3 NCs by adding Sb 2 S 3 NCs to the active layer. The device's electrical, morphological, and optical properties were significantly affected by the Sb 2 S 3 NC concentration in the P 3 HT:PCBM. The doping concentration improved the surface roughness of the active layer and tapping mode operation. The phase distribution of the blends was also investigated. The high crystalline polymer enhanced the red shift of the optical absorption, increased the photoluminescence intensity, and narrowed the full width at half the maximum of the Raman peaks. The results indicated that the Sb 2 S 3 NCs strongly affected the flat-on orientation, which increased the charge carrier transport assisted by the π-π interactions. The increased number of P 3 HT single chains and the phase separation increased the free electrons, which effected the absorption and mobility and reduced the charge recombination in the active layer blend prepared using a concentration of 0.04 mg/mL. Our studies suggest that the charge separation and current generation in P 3 HT:PCBM:Sb 2 S 3 :NC-based devices result mainly from Sb 2 S 3 NCs' light absorption and subsequent hole-transfer from the inorganic semiconductor to the organic hole transporting material. The best power conversion efficiency (PCE) of the polymer solar cells was 2.72%, using a glass/ITO/PEDOT:PSS/P 3 HT:PCBM:Sb 2 S 3 :NC/MoO 3 /Ag device structure, under AM 1.5 sun, and global irradiation of 1000 W m −2 by blending the Sb 2 S 3 NCs with a concentration of 0.04 mg/mL in the active layer. An expected mechanism is also proposed to explain the superior performance of Sb 2 S 3 NC-doped P 3 HT:PCBM at optimal content.

Data Availability Statement:
The data presented in this study are available on request from the corresponding author.