Homojunction Sb2Se3 Solar Cell

Homojunction Sb2Se3 Solar Cell | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Article Homojunction Sb 2 Se 3 Solar Cell Tao Chen, Junjie Yang, Jianyu Li, Shuwei Sheng, Zhiyuan Cai, Ke Li, and 3 more This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-7257351/v1 This work is licensed under a CC BY 4.0 License Status: Published Journal Publication published 13 Apr, 2026 Read the published version in Nature Photonics → Version 1 posted You are reading this latest preprint version Abstract Antimony selenide (Sb 2 Se 3 ) has emerged as a promising thin-film photovoltaic absorber due to its ideal bandgap (1.1-1.3 eV), high absorption coefficient (>10⁵ cm - ¹), and environmentally benign composition. However, Sb 2 Se 3 solar cells (SSCs) often suffer from significant open-circuit voltage ( V OC ) losses, attributed to weak built-in fields and severe non-radiative recombination at interfaces and the absorber layer. Here, we demonstrate a composition‑driven carrier polarity control strategy to form an n-type/p-type Sb 2 Se 3 homojunction. By precisely tuning the chemical potentials of Se and Sb, we reversibly modulate the conductivity type, achieving carrier densities exceeding 10 14 cm - ³ for the n- and p-type states, with Hall coefficients ranging from −3.14×10 - ² m³ C - ¹ (n-type) and +9.51×10 - ² m³ C - ¹ (p-type). More importantly, we incorporate a homojunction structure into the planar SSC, which simultaneously enhances the built-in electric field and passivates deep-level defects. These synergistic effects promote carrier separation, reduce non-radiative recombination, and accelerate carrier extraction. As a result, the study demonstrates a record power conversion efficiency of 10.15% for thermally evaporated Sb 2 Se 3 devices, along with the lowest open-circuit voltage deficit (0.459 V) among all reported SSCs. This work not only establishes a new efficiency benchmark for Sb 2 Se 3 solar cells but also offers a universal approach for defect management and junction design in emerging chalcogenide photovoltaics. Physical sciences/Materials science/Materials for energy and catalysis/Solar cells Physical sciences/Materials science/Materials for devices Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 Figure 6 Main Antimony selenide (Sb 2 Se 3 ) material has attracted substantial research interest due to its outstanding optoelectronic properties and exceptional thermal and chemical stability. 1–3 The nearly ideal bandgap (1.1-1.3 eV) and high absorption coefficient (>10 5 cm -1 ) support the theoretical photoelectric conversion efficiency of Sb 2 Se 3 exceeding 30%. 4,5 As an emerging thin-film photovoltaic technology, Sb 2 Se 3 solar cells have demonstrated remarkable progress in material optimization, interface engineering, and device structure design, achieving power conversion efficiency exceeding 10%. 6–9 However, this performance still lags behind mature chalcogenide thin-film technologies, such as CdTe and Cu(In,Ga)Se 2 photovoltaics (PVs). 10–12 The large deficit in open-circuit voltage ( V OC ) is the primary factor limiting the performance of Sb 2 Se 3 solar cells, with experimentally obtained V OC values typically about 500 mV lower than the theoretical radiation limit voltage. This energy loss is attributed to two key factors: (1) insufficient driving force for charge separation at the junction, due to a weak built‑in electric field; and (2) severe non‑radiative recombination via deep traps at both the Sb 2 Se 3 absorber/buffer interfaces and within the bulk material. 13–15 Effective control of semiconductor polarity, combined with judicious device structure design, has proven essential for overcoming efficiency challenges in established photovoltaic systems. 16–18 In conventional PV technologies such as crystalline silicon and III-V semiconductors, the precisely controlled incorporation of donor or acceptor dopants enables the formation of n-type and p-type regions, respectively, which establish strong p-n homojunctions with well-defined built-in electric fields. 19,20 The enhanced internal electric fields effectively suppress carrier recombination, thereby improving device performance. In emerging photovoltaic technologies (such as metal-halide perovskites and organic PVs), intrinsic/extrinsic doping and charge transfer effects have been employed to modulate charge carrier concentrations and construct heterojunction or homojunction architectures, enhancing charge extraction. 21–24 Remarkably, homojunction engineering has achieved substantial efficiency improvements across diverse perovskite solar cells, spanning lead-based, tin-based, and mixed lead-tin systems in both regular (n-i-p) and inverted (p-i-n) architectures. 25–29 These achievements underscore the significant potential of rational carrier-polarity control and internal junction engineering in reducing voltage losses. Applying these theoretical frameworks to Sb 2 Se 3 solar cells offer a viable approach to mitigate their intrinsic voltage limitations. Therefore, we focus on the co-optimization of junction design and defect passivation to improve the photovoltaic performance of Sb 2 Se 3 solar cells. Here, we develop a composition-driven intrinsic doping strategy for polarity control and constructed Sb 2 Se 3 homojunction solar cells. By precisely tuning the chemical potential of Se and Sb during thermal evaporation, we successfully switched the conductivity of Sb 2 Se 3 films from n-type to p-type, enabling a monolithic p-n homojunction within the absorber. We also conducted density functional theory (DFT) calculations to elucidate the mechanistic details of the conductivity-type transition process. This built-in junction provides an extra internal electric field that widens the depletion region, enhancing charge separation and suppressing non-radiative recombination. Comprehensive spectral and electrical characterizations reveal an approximately 0.2 eV internal potential gradient, along with suppressed defects and extended carrier lifetime in the device. Ultrafast spectroscopy and depth-resolved simulations further demonstrate faster carrier transport rate, with recombination losses reduced by over an order of magnitude. Notably, Sb 2 Se 3 solar cells with p-n homojunction demonstrate a record power conversion efficiency of 10.15% for thermally evaporated devices, along with the lowest V OC deficit reported for all kinds of SSCs so far. This work highlights a promising new strategy for improving SSC performance via homojunction engineer to synergistically enhance the interfacial charges extraction and suppress non-radiative recombination. P olarity conversion and intrinsic doping mechanisms Although Sb 2 Se 3 can exhibit ambipolar charge transport and has been reported as either an n‑type or p‑type semiconductor, deliberate strategies to control its conductivity remain undeveloped. 14,30–32 First‑principles calculations suggest that the material’s conductivity type is governed by growth conditions and the relative chemical potentials of Se and Sb, yet experimental validation of this relationship and methods for targeted polarity tuning are absent. 13,33,34 To achieve controlled conductivity switching in Sb 2 Se 3 , we developed an intrinsic doping method based on precursor composition tuning. During ball‑milling synthesis, we adjusted the Se and Sb chemical potentials by introducing a controllable Se excess into the precursor mixture (detailed preparation process is provided in the experimental section). This approach compensates for selenium loss, which is inevitable in the conventional thermal deposition by the Sb 2 Se 3 compound as the precursor material due to selenium’s high vapor pressure during film deposition and annealing. 35 Our method herein enables precise stoichiometric control of the resulting absorber layer, with Se/Sb ratios tunable from 1.41 to 1.59. For both the control Sb 2 Se 3 absorber and the n-type absorber region in the target Sb 2 Se 3 , we employed a precursor Se/Sb ratio of 1.5. For the p-type absorber region in the target Sb 2 Se 3 device, we increased the precursor Se/Sb ratio to 1.6, as first‑principles calculations indicate that selenium‑rich conditions favor p‑type conductivity in Sb 2 Se 3 . 33 In this study, both n-type and p-type absorbers were prepared by thermal evaporation method, which exhibit compact morphologies and pure orthorhombic Sb 2 Se 3 phase, as confirmed by scanning electron microscopy (SEM) and X-ray diffraction (XRD) analysis (Supplementary Fig. S1). For the convenience of discussion, the SSCs without and with homojunction structure is labeled as Control-Sb 2 Se 3 and Target-Sb 2 Se 3 samples. We first examined the surface composition of the films using X‑ray photoelectron spectroscopy (XPS) (Supplementary Fig. S1 and S2). The Se/Sb atomic ratio on the surface of the Target‑Sb 2 Se 3 film is increased from 1.46 (Control‑Sb 2 Se 3 ) to 1.56, indicating successful selenium enrichment in the p‑type layer. Since XPS can only test information within a few nanometers of the surface depth, we performed cross‑sectional energy‑dispersive X‑ray spectroscopy (EDS). Nine measurement points across both the n‑type and p‑type regions show that the Se/Sb ratio in the n‑type region remains approximately 1.46 for both samples, whereas the p‑type region of the Target‑Sb 2 Se 3 reaches 1.56 (Supplementary Fig. S3). For more accurate bulk composition, inductively coupled plasma atomic emission spectroscopy (ICP‑AES) reveals Se/Sb ratios of 1.45 in the Control‑Sb 2 Se 3 and 1.57 in the p‑type region of the Target‑Sb 2 Se 3 (Supplementary Fig. S4), in excellent agreement with XPS and EDS data. Furthermore, secondary ion mass spectrometry (SIMS) and EDS of transmission electron microscopy (TEM) depth profiling both provide a continuous elemental distribution through the absorber thickness (Supplementary Fig. S5 and S18b). The Control‑Sb 2 Se 3 film exhibits essentially constant Sb and Se signals, whereas the Target‑Sb 2 Se 3 film shows a gradual increase in Se/Sb ratio from the CdS interface toward the film surface. Collectively, these multi-dimensional measurements conclusively demonstrate that our precursor-composition strategy effectively drives Se/Sb variations in the film, successfully realizing the targeted composition‑driven homojunction structure. We then studied the electronic structure of Sb 2 Se 3 thin films. Firstly, we characterized the energy level structures of the Control-Sb 2 Se 3 and Target-Sb 2 Se 3 samples using ultraviolet photoelectron spectroscopy (UPS) (Fig. 1a-c). For the Control-Sb 2 Se 3 samples, the Fermi level is determined to be −4.30 eV, close to the conduction band minimum (CBM) of −3.76 eV, indicating the n-type characteristics of the thin film. For the surface of Target-Sb 2 Se 3 samples, the Fermi level shifts toward the valence band maximum (VBM), with offset reducing from 0.67 eV to 0.54 eV (Fig. 1c), indicating a transition in conductivity type from n-type to p-type. The Kelvin probe force microscopy (KPFM) characterization further confirms the above results by testing the surface potential. Figure 1d and 1e show the KPFM potential images of the Control-Sb 2 Se 3 and the Target-Sb 2 Se 3 , respectively, with relatively uniform contact potential distributions in both samples. In addition, it can be seen from Figure 1e that the surface potential of the Target-Sb 2 Se 3 film is lower than that of the Control-Sb 2 Se 3 film (Fig. 1d), because the accumulated holes cause the Fermi level to be located near the valence band. On the contrary, the Control-Sb 2 Se 3 film exhibits a higher surface potential, which is related to the n-type doping of the sample. To guarantee the reliability of the results, we conducted statistical analysis on the surface potential distribution of the samples (Fig. 1f, g). According to the statistical results, the surface potential difference of the target sample reduced by about 0.2 eV compared to the control sample, consistent with the Fermi level shift measured by UPS. As UPS and KPFM are both surface-sensitive techniques, further evidence is significant for analyzing the effects of doping in the bulk regions of the Sb 2 Se 3 samples. Hall effect measurement is one of the most reliable techniques for characterizing the polarity and concentration of carriers in semiconductors. 19,36 The Hall coefficients ( R H ) and majority carrier concentrations of Sb 2 Se 3 samples with different Se/Sb ratio are shown in Fig. 1h and 1i. For the Control-Sb 2 Se 3 sample with a Se/Sb atomic ratio of 1.46, it exhibits n-type characteristics with a negative R H value, corresponding to an electron concentration of 1.99×10 14 cm -3 , obtained from equation of n = 1/( R H q), where n is the carrier concentration and q is the elementary charge. 36 The more Sb-rich composition (with a Se/Sb atomic ratio of 1.42) results in increased electron concentration and enhanced n-type characteristics. However, as the Se/Sb ratio rises to 1.51, transforming to a Se-rich state, R H turns positive, indicating a transition to p-type behavior where holes dominate charge transport. The Target-Sb 2 Se 3 film with more Se-rich composition (Se/Sb of 1.56) demonstrates a more p-type characteristics with hole concentration reaching 6.57×10 14 cm -3 . Clearly, as the composition of the Sb 2 Se 3 thin film deviates from stoichiometric ratio, the measured carrier concentration increases accordingly. It is especially pronounced in p‑type Sb 2 Se 3 samples, which exhibit higher carrier densities—a phenomenon typically attributed to changes in the semiconductor's intrinsic defect properties. In other words, intrinsic doping is the primary factor governing the conductivity type transition in Sb 2 Se 3 . We then employed density functional theory (DFT) to elucidate the mechanism regarding the transition of thin film conductivity types. Supplementary Fig. S6 and S7 depict the computed formation energies of all intrinsic point defects as a function of Fermi level under Sb‑rich and Se‑rich conditions. These values are consistent with early theoretical calculations. 33,37,38 Evidently, anion vacancies (V Se1 , V Se2 , V Se3 ) exhibit deep donor behavior, possessing thermodynamic transition levels far from the CBM. In contrast, cation vacancies (V Sb1 , V Sb2 ) act as shallow acceptors, with transition levels lying above the VBM by several k B T. All intrinsic anti‑site defects, including cation-replace-anion (Sb Se1 , Sb Se2 , Sb Se3 ) and anion-replace-cation (Se Sb1 , Se Sb2 ) defects, display amphoteric character, acting as donors near VBM and acceptors near CBM; notably, only the Se Sb2 defect exhibits shallow acceptor property (about 0.10 eV above the VBM). In addition, cation interstitial defect (Sb i ) is also amphoteric, with a deep mid‑gap transition, while anion interstitial defect (Se i ) remains electrically inert except under extreme p‑type conditions. To identify the dominant defect species, we selected representative intrinsic defects with low formation energy for discussion and presented their transition level diagrams in Fig. 2a, b. In Sb‑rich Sb 2 Se 3 with a bandgap of 1.16 eV, +2 charged V Se2 defect exhibits the minimal formation energy when the Fermi level is below the mid‑gap (~0.58 eV below the CBM), whereas the −1 charged Sb Se2 defect becomes most favorable upon Fermi level approaching to the CBM. The mutual compensation between these donor and acceptor centers pins the Fermi level above mid‑gap (blue dashed line shown in Fig. 2a), resulting in pronounced n‑type conductivity. In contrast, Se‑rich conditions induce a considerably reduction (>0.6 eV) in the formation energies of shallow acceptors V Sb1 , V Sb2 , and Se Sb2 (Fig. 3d-g). The decrease in formation energy significantly increases the equilibrium concentrations of these defects, shifts the Fermi level towards VBM, which fosters p‑type behavior. Meanwhile, the formation energy of V Se2 rises substantially under Se-rich conditions, reducing the compensation effect on hole carriers. Combining these DFT findings with experimental observations (Fig. 1), we propose the following mechanism (Fig. 2h, i). For Sb‑rich Sb 2 Se 3 , the facile formation of deep donors depletes hole populations and pins the Fermi level on the n‑type side; while for Se‑rich Sb 2 Se 3 , shallow acceptors near VBM are easily formed, which increase the hole density and shift the Fermi level towards p‑type side. Such defect‑engineered control of conductivity type provides a novel approach for constructing homojunction Sb 2 Se 3 solar cells. Homojunction construction and photovoltaic performance of Sb 2 Se 3 solar cells Leveraging the precise control over Sb 2 Se 3 conductivity type, we fabricated a monolithic homojunction structure by integrating n‑type and p‑type Sb 2 Se 3 layers into planar solar cells with the architecture: fluorine-doped tin oxide (FTO)/CdS/Sb 2 Se 3 /Spiro-OMeTAD/Au. In this configuration, CdS, Spiro-OMeTAD (2,2',7,7'-tetrakis(N, N-di-p-methoxyphenyl-amine)-9,9'-spirobifluorene), and Au serve as the electron-transporting layer, hole-transporting layer, and back electrode, respectively. Cross-sectional SEM images of the Control- and Target- SSCs (Fig. 3a, b) show that both Sb 2 Se 3 films display densely packed, columnar grains spanning the full absorber thickness, suggesting high crystallinity and compact morphology of the as-synthesized Sb 2 Se 3 films. Remarkably, this junction does not cause discernible interfaces or grain boundaries, which is consistent with the TEM images (Fig. 3g, j, and m). Furthermore, we compared the lattice fringes of n-type Sb 2 Se 3 near the CdS and p-type Sb 2 Se 3 and performed Fourier transforms (FT) of their respective regions (Fig. 3h, i, k, and l). The lattice spacings in both regions are 0.315 nm, which corresponds to the (211) planes of orthorhombic Sb 2 Se 3 . Fourier transform (FT) analysis further confirms that the crystal plane and orientation of n-type and p-type Sb 2 Se 3 are identical. Combined with the XRD results (Supplementary Fig. S1), these findings confirm the formation of a homojunction without lattice mismatch. To further substantiate the formation of the p-n junction in Sb 2 Se 3 absorber, we performed cross-sectional KPFM potential characterizations of the FTO/Sb 2 Se 3 /Au devices (Fig. 3c and 3d), which were collected under dark conditions on freshly cleaved samples to avoid artefacts. In the control device, the Sb 2 Se 3 absorber exhibits a nearly uniform surface potential (about 180-200 mV) throughout the entire thickness (Fig. 3c), confirming its single-type (n-type) conductivity without internal electric field. In contrast, the Sb 2 Se 3 absorber in the target device exhibits a distinct potential gradient (Fig. 3d), decreasing steadily from the FTO interface to the Au contact. This spatial variation visualizes the built-in electric field generated by the p-n homojunction within Sb 2 Se 3 layer, which facilitates efficient charge separation and suppresses bulk recombination in the absorber, beneficial for improving the optoelectronic performance of the device. Consequently, we investigated the photovoltaic performance of SSCs with and without the homojunction under AM 1.5G solar irradiance (100 mW cm - 2 ), with schematic illustrations of the device structures shown in Fig. 4a-c. In fact, we optimized the composition and thickness of the p-type and n-type regions (see details in Supplementary Note 2 and Supplementary Fig. S8-12). As a result, the establishment of p-n homojunction increases the efficiency from 8% for the Control-Sb 2 Se 3 device to 10.15% for the Target-Sb 2 Se 3 device, which is the highest efficiency reported for thermally evaporated Sb 2 Se 3 solar cells to date (Fig. 4d, g and Supplementary Table S3). Specifically, the open circuit voltage ( V OC ), short-circuit current density ( J SC ), and fill factor (FF) are improved from 428 mV to 492 mV, 29.59 to 31.48 mA cm -2 , and 63.21% to 65.56 %, respectively (Table 1), which are 15%, 6% and 4% higher than that of the control device, respectively. Obviously, the homojunction engineering significantly improves photovoltaic parameters, especially V OC (Fig. 4f). To explore the photocurrent generation of the devices, we conducted external quantum efficiency (EQE) spectroscopy characterization. It shows the enhanced wavelength response in the range of 550-1000 nm and demonstrated that the increase in J SC may may stem from the improved film quality (Fig. 4e), as the target device exhibits the same cut-off edge as the control device (approximately 1030 nm). This consistency with the ultraviolet-visible absorption spectra (measured bandgap of 1.205 and 1.202 eV for the control and target devices, respectively (Supplementary Fig. S1d)). As for the V OC , its significant improvement results in a low V OC deficit of 0.459 V, which is the lowest value reported among all Sb 2 Se 3 solar cells to date (Fig. 4h). The increase in V OC arises from enhanced carrier transport and suppressed nonradiative recombination, benefited from the homojunction engineering (Fig. 4c), as evidenced by a reduction in the dark saturation current density J 0 from 1.43×10 -6 in the Control-Sb 2 Se 3 device to 8.70×10 -8 in the Target-Sb 2 Se 3 device (Supplementary Fig. S13). Concurrently, the fill factor improvement stems from a lowered diode ideality factor, which produces a more ideal J-V characteristic and thus enhances FF. These results demonstrate that composition‑driven homojunction engineering in Sb 2 Se 3 constructs a robust internal electric field for improved photovoltaic performance. Table 1. Photovoltaic parameters of the Control- and Target- Sb 2 Se 3 solar cells. Devices V OC (V) J SC (mA/cm 2 ) FF (%) PCE (%) Control 0.428 (0.408±0.009) 29.59 (29.64±0.43) 63.21 (62.25±1.00) 8.00 (7.53±0.23) Target 0.492 (0.486±0.006) 31.48 (31.35±0.14) 65.56 (65.06±0.47) 10.15 (9.91±0.20) Internal homojunction‐driven modulation of carrier generation, transport, and recombination The modulation of built-in electric fields typically alters charge carrier characteristics. In order to elucidate the mechanism through which homojunction construction enhances Sb 2 Se 3 solar cell performance, we performed equilibrium band-structure analysis, numerical simulations, electrostatic profiling, and ultrafast transient absorption spectroscopy (TAS) to investigate charge carrier dynamics. Figure 5a and 5d present schematic energy level diagram and illustrate the working principle of homojunction-SSCs, demonstrating the transport mode of charge carriers between layers. When photons are absorbed by both p-type and n-type Sb 2 Se 3 , charge carriers are generated on both sides of the junction. Subsequently, electrons and holes drift in opposite directions: electrons inject into the CdS layer and transport to the FTO cathode, whereas holes inject into the spiro-OMeTAD layer and transport to the Au anode. In conventional n-i-p structured SSCs, the limited dynamic driving force for carrier transport within the Sb 2 Se 3 film often results in significant carrier recombination. Typically, photo-induced electrons and holes are extracted mostly at the Sb 2 Se 3 /ETL and the Sb 2 Se 3 /HTL interfaces, respectively. Therefore, our designed homojunction introduces an additional built-in electric field, extending the depletion region deeper into the Sb 2 Se 3 absorber layer. This enhanced field promotes the rapid separation of electrons and holes, driving them toward ETL and HTL, respectively (Fig. 5d), while simultaneously reducing carrier recombination. We further applied the SCAPS-1D simulation tool, widely employed for modeling thin-film photovoltaic devices, focusing on the generation and recombination dynamics of photo carriers. As illustrated in Fig. 5b, both the Control and Target architectures exhibit their maximum generation rates at the CdS/n-Sb 2 Se 3 junction. However, the target device maintains a pronounced secondary peak at the internal p-n interface, clearly demonstrating that the built-in field facilitates spatial carrier separation. The corresponding recombination rate profiles (Fig. 5c) reveal that the control device generates a localized recombination maximum of about 10 21 cm -3 ·s -1 within the central region of Sb 2 Se 3 layer, whereas the homojunction reduces this rate by more than an order of magnitude. We attribute this suppression to the unidirectional electric field induced by the internal p-n junction pushing electrons and holes toward opposite contacts, thereby diminishing interfacial non-radiative recombination. Meanwhile, the p-type region of the homojunction may result in lower density of deep-level trap states, contributing to reduced bulk non-radiative losses. On the other hand, Mott-Schottky analysis shows an increase in V bi from 0.46 V to 0.56 V (Fig. 5e). The capacitance-voltage ( C-V ) and drive-level capacitance profiling (DLCP) reveal that the depletion width expands from 131 nm in the control device to 182 nm in the target device without introducing additional interface traps (Fig. 5f). We then conducted ultrafast TAS analysis on FTO/CdS/Sb 2 Se 3 devices to probe the dynamics of charge carriers excited by pulsed photo. As shown in Fig. 5 g-i, the contour maps based on TA spectra of Control- and Target- Sb 2 Se 3 devices both display a photo-induced absorption (PIA) peak at approximately 690 nm which is attributed to the characteristic absorption of Sb 2 Se 3 . Notably, the TAS signal of the target device decays faster, indicating a more efficient transport of charge carriers from the Sb 2 Se 3 film to the adjacent interface layer. By fitting and analyzing the decay kinetics of PIA signals, the average TA decay lifetime of the Target-Sb 2 Se 3 device (1.75 ns) is much smaller than that of the control device (4.15 ns), as shown in Fig. 5i. These results further confirm the critical role of homojunctions in accelerating the extraction and transport of charge carriers. Deep‐level defects critically govern carrier lifetimes via Shockley-Read-Hall (SRH) recombination, particularly the bulk nonradiative losses associated with lattice imperfections. 16,39 To quantify how homojunction engineering reshapes the defect landscape and suppresses SRH recombination, we performed deep-level transient spectroscopy (DLTS) on Control-Sb 2 Se 3 and Target-Sb 2 Se 3 devices. By carefully adjusting the applied pulse voltage (0.2-0.6 V), we mitigated artifact peaks arising from capacitor bridge recovery delays observed at inappropriate pulse voltage levels, while identifying defects at different depths (Fig. 6a and 6d). 40 In the DLTS spectra, positive peaks correspond to majority-carrier traps, while negative peaks indicate minority-carrier traps. Based on the DLTS spectra, Arrhenius plots are analyzed (Supplementary Figure S14a and b), and defect-related parameters, including the trap energy level ( E T ), trap cross-section (σ), and trap density ( N T ), are extracted (Supplementary Tab. S8). In the Control-Sb 2 Se 3 device, three hole-traps (H1, H2, H3) and one electron-trap (E1) are identified. Their respective energy levels are located at 0.364, 0.548, and 0.679 eV above the VBM for the hole traps and 0.602 eV below the CBM for the electron trap (Fig. 6a-c). Surprisingly, there is only one electron-trap (E2) located at the 0.665 eV energy level below CBM is identified in the target device. Based on the DFT calculations (Figure 2), the H1, H2, H3, and E1 defects can be attributed to Sb Se1 , Sb Se3 , Sb Se2 antisite defects, and Se vacancy defect (V Se3 ) when the film is in Sb-rich state. As for the Sb 2 Se 3 in Se-rich state, the E2 defect can be attributed to Se Sb1 defect. Through comparative analysis of defect characteristics between the two devices, we confirm that homojunction engineering effectively passivate the original Sb Se3 , Sb Se1 , Sb Se2 , and V Se defects in Sb 2 Se 3 . In the target device, the newly emerged Se Sb1 defect demonstrates more favorable properties, exhibiting a lower defect density (6.34×10 13 cm -3 ) and a smaller capture cross-section (3.87×10 -16 cm²). According to the defect-assisted SRH recombination model, the carrier lifetime ( τ ) associated with a specific defect can be evaluated by Equation (1): Where ν th is the charge heating rate, which is 10 7 cm s -1 at room temperature. 35 Specifically, we focus on traps with the shortest carrier lifetime because they have the most significant impact on carrier recombination. Our calculations reveal that in Control-Sb 2 Se 3 , the Sb Se2 defect has the shortest lifetime of 543 ns, followed by V Se3 , identifying them as the dominant recombination centers. In contrast, the Se Sb1 defect in Target-Sb 2 Se 3 shows a substantially longer lifetime (4075.68 ns), indicating a slightly negative contribution to carrier recombination. Moreover, according to the SRH recombination theory for multiple recombination centers, the total SRH recombination rate ( U SRH ) is the sum of the individual recombination rates ( U i, SRH ) of each recombination center i , 41 as expressed in equation (2): Based on the calculations, homojunction engineering reduces both the density and variety of traps in the Sb 2 Se 3 bulk phase. According to equation (2), the reduction in recombination centers consequently reduces the SRH recombination of carriers. In the homojunction target device, the n-type Sb 2 Se 3 layer is only about 50 nm thick— comparable to the thickness of CdS electron-transport layer, so p-type Sb 2 Se 3 forms the main bulk of the absorber. DLTS characterization is primarily applied to probe the deep-level defects in the p-type Sb 2 Se 3 region. Analysis of pure p-type Sb 2 Se 3 -based devices (Supplementary Fig. S15) reveals only a single E2 trap, with energy level, density, and capture cross-section values consistent with those observed in the target device. This discovery indicates that the p-type Sb 2 Se 3 intrinsically maintains a more favorable defect profile. Therefore, it is because of the superiority of p-type Sb 2 Se 3 in reducing bulk SRH recombination along with p-n homojunction expanded space charge region that we achieved a momentous improvement in device performance. Conclusion In conclusion, we demonstrate a cation/anion composition-driven homojunction engineering which achieves an efficiency breakthrough in Sb 2 Se 3 photovoltaics. Our study reveals that the modulation of elemental chemical potential during thermal evaporation enables controls over charge carrier polarity, facilitating the fabrication of a monolithic p-n homojunction. Furthermore, we elucidate the mechanism of this strategy in suppressing non-radiative recombination from two aspects: (i) the homojunction increases the depletion width from 131 nm to 182 nm with introducing an additional built-in field to enhance charge separation; and (ii) the Se-rich components passivate deep-level defects, reducing trap density by over 70%. Depth-resolved KPFM and carrier dynamic analysis confirm the accelerated carrier extraction and an order-of-magnitude reduction in recombination losses. The integration of homojunction into planar n-i-p devices enables a record power conversion efficiency of 10.15 % for thermally evaporated Sb 2 Se 3 solar cells, along with a minimal V OC deficit of 0.459 V—the lowest reported value among Sb 2 Se 3 cells. This investigation underscores the synergistic advantages of homojunction engineering in efficient charge separation and recombination suppression, pointing out a new direction for optimizing optoelectronic performance of antimony chalcogenide thin-film solar cells. Declarations Data availability All of the data needed to evaluate the conclusions in the paper are present in the paper and its Supplementary Information. Acknowledgments This work was supported by National Natural Science Foundation of China (22275180), the Fundamental Research Funds for the Central Universities (No. WK2490000002), Major Science and Technology Projects of Anhui Province (AHZDCYCX-LSDT2023-10), and the University Synergy Innovation Program of Anhui Province (GXXT-2023-031). Author contributions Authors and Affiliations Department of Materials Science and Engineering, School of Chemistry and Materials Science, University of Science and Technology of China, Hefei, People’s Republic of China. Junjie Yang, Jianyu Li, Shuwei Sheng, Zhiyuan Cai, Zichen Ruan, Bo Che, Rongfeng Tang, and Tao Chen Institute of Deep Space Sciences, Deep Space Exploration Laboratory, Hefei 230088, P. R. China. Tao Chen Institute of Optoelectronic Materials and Devices, School of Materials Science and Engineering, Jiangxi University of Science and Technology, Ganzhou 341000, P. R. China. Ke Li Contributions J. Y., J. L. and S. S. contributed equally to this work. J.Y. and T.C. conceived of the original concept and designed the experiments. J.Y., J.L. and S.S. fabricated the devices and conducted the photovoltaic and optical characterization and analysis. J. L. did the TAS measurements and performed the SEM, TEM analysis. K.L. Z. R. and B. C. were involved in materials characterization and device simulation. Z. C. conducted the DFT calculations. J.Y., R.T. and T. C. co-wrote the paper. T.C., R.T., S.S. and J. Y. revised the paper with all authors commenting on the paper. Competing interests The authors declare no competing interests. Additional information Publisher’s note Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations. Source data Source Data References Zhou, Y. et al. Thin-film Sb 2 Se 3 photovoltaics with oriented one-dimensional ribbons and benign grain boundaries. Nat . Photon . 9 , 409–415 (2015). Zhao, Q. et al. Grain‐Boundary Elimination via Liquid Medium Annealing toward High‐Efficiency Sb 2 Se 3 Solar Cells. Adv . Mater . 37 , 2414082 (2025). Wen, X. et al. Vapor transport deposition of antimony selenide thin film solar cells with 7.6% efficiency. Nat . Commun . 9 , 2179 (2018). Dong, J. et al. Carrier management through electrode and electron-selective layer engineering for 10.70% efficiency antimony selenosulfide solar cells. Nat . Energy 10 , 857-868 (2025). Tang, R. et al. 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Statistics of the Recombinations of Holes and Electrons. Phys. Rev. 87 , 835–842 (1952). Methods Materials Cadmium sulfate (CdSO 4 ·8/3H 2 O, AR), thiourea (CH 4 N 2 S, AR), ammonium hydroxide (NH 3 ·H 2 O, 25-28%) were purchased from Sinopharm. Antimony chloride (SbCl 3 , 99%) was purchased from Tokyo Chemical Industry Co. Ltd. Antimony powder and selenium powder purchased from Aladdin. Spiro-OMeTAD (99.8%) was purchased from Yingkou Youxuan Trade Co., lithium bis(trifluoromethylsulfonyl) imide (Li-TFSI, 98%) and 4-tert-butylpyridine (96%), were purchased from J&K Scientific Co. All the chemicals and reagents were used without any treatment. Sb 2 Se 3 solar cells Fabrication Preparation of CdS layer: All devices were deposited on cleaned FTO-coated glass. Prior to deposition, the FTO-coated glasses were ultrasonically cleaned sequentially with detergent, deionized (DI) water, isopropanol, acetone, and ethanol for 40 min, followed by cleaning with ultraviolet ozone for 20 minutes. Then, we deposited SnO 2 nanoparticles onto the FTO substrate by spin-coating SnO 2 colloidal water dispersion (15%) diluted with a water volume ratio of 1:5. The spin-coating was performed at 3000 r.p.m. for 30 seconds, then following a further annealing at 250 °C in air for 30 minutes to prepare the SnO 2 interfacial layer. Afterwards, the electron transport layer, CdS, was deposited on FTO/SnO 2 -coated glass via the CBD method. The preparation of CdS is as follows: First, 17.3 mL of ammonium hydroxide (NH 3 ·H 2 O, 25–28%) was added into 13.3 mL of 15 mmol L -1 Cd(NO 3 ) 2 solution with continuously stirring in the air for 1.5 min. Next, 8.8 mL of 1.5 mol L -1 thiourea solution was added into the above solution. After stirring for 30 s, 93.3 mL of deionized water was added, with continuing stirring for 30 s. The FTO-coated glass was immersed into the mixed solution, and reacted at 65 °C for 16 min. After that, took out the substrates and rinsed them with deionized water and ethanol, and dried them with nitrogen gas. Then 30 mg mL -1 of SbCl 3 methanol solution was spin-coated to treat the CdS surface with 3000 r.m.p. for 30 s. The CdS-coated substrates were then heat-treated at a high temperature of 400 °C for 10 min to obtain crystallized CdS in an air atmosphere. After natural cooling, the FTO/SnO 2 /CdS film was transferred into thermal evaporation equipment for the deposition of Sb 2 Se 3 absorber layer. Preparation of Sb 2 Se 3 absorber layer: The Sb 2 Se 3 alloy precursor was synthesized through ball mill and solid phase reaction approach. Due to selenium’s high vapor pressure, some Se inevitably evaporates during substrate heating and post‑annealing of Sb₂Se₃ films. To investigate this, we ball‑milled a series of Sb₂Se₃ precursor powders with systematically varied Se: Sb ratios. After film deposition and annealing, all samples exhibited Se loss, and the extent of depletion increased with higher initial Se content, yet we were still able to produce Sb₂Se₃ thin films covering a broad range of Se: Sb ratios. Firstly, Sb and Se powders were placed in a ball mill tank according to the different Se/Sb ratio from 1.45-1.65 of Sb 2 Se 3 and ball-milled for 8 hours. The precursor was dried in vacuum environment for 30 min to obtain Sb 2 Se 3 precursor powder, then the precursor powder was pressed into a tray and then annealed at 300 °C for 30 min in an N 2 glove box to obtain Sb 2 Se 3 alloy. Sb 2 Se 3 films were deposited using thermal evaporation deposition under a pressure of 1x10 -4 Pa. For the Control-Sb 2 Se 3 , before the deposition, 0.2 g of Sb 2 Se 3 powder with Se/Sb ratio of 1.5 was placed in the middle of a tungsten boat, with an evaporation rate of 5-8 nm per second, and the substrate temperature was set at 315 °C. The thickness of the deposited film was controlled about 400 nm. Then, the Sb 2 Se 3 film was annealed in a glove box with N 2 atmosphere at 370 °C for 8 minutes. For the Target-Sb 2 Se 3 homojunction, we deposited n-type and p-type Sb 2 Se 3 films sequentially using thermal evaporation. For the n-type film, the preparation method is the same as that in Control- Sb 2 Se 3 , except that the thickness is about 50 nm. Next, we deposited about 350 nm of Sb 2 Se 3 using the precursor with a Se/Sb ratio of 1.6, and then annealed the film in a glove filled with nitrogen at 370 °C for 8 minutes. For the fabrication of Sb 2 Se 3 solar cells: After preparing the Sb 2 Se 3 film, the spiro-OMeTAD HTL was spin-coated onto Sb 2 Se 3 film at 3000 r.p.m. for 30 seconds. The spiro-OMeTAD solution was prepared by dissolving 36.6 mg of spiro-OMeTAD in 1 mL of chlorobenzene, along with 14.5 μL of 4-tert-butylpyridine and 9.5 μL of Li-TFSI solution (520 mg mL -1 Li - 1 TFSI in acetonitrile). After spin coating, the film was annealed at 100 °C in air for 10 minutes. Finally, an Au electrode was thermally evaporated onto the film. Film and Device Characterizations: The crystal structure of Sb 2 Se 3 thin films was analyzed on a Bruker Advance D8 diffractometer using X-ray diffraction of Cu K𝛼 radiation (𝜆 = 1.5416 Å). The morphology and thickness of the films were characterized by scanning electron microscopy (SU8220, Hitachi High-Tech), and the element analysis was conducted by Zeiss G500 SEM equipped with an EDS (Bruker) module. The work function and valence band binding energy were measured by UPS, and the excitation source was a He light source (h v = 21.22 eV) with a beam spot of 2 mm. The Hall effect measurement was exerted by Ecopia Hall Effect Tester (HMS-7000) at 300 K. EQE spectra were measured using a single source lighting system (halogen lamp) and a monochromator (model SPIEQ200). The surface atom chemical state was characterized using X-ray photoelectron spectroscopy (Thermo Fisher ESCALAB 250Xi) which possesses Al k𝛼 150 W with beam spot 500 μsm. The change in mass of the sample with temperature or time was performed by Thermogravimeter (TGA Q5000IR). The J-V and VOC-light response of devices was measured by Keithley 2400 apparatus with a standard xenon lamp-based solar simulator (Oriel Sol 3A) under solar-simulated AM 1.5 sunlight (100 mW cm -2 ). The calibration for the light intensity was conducted in advance by reference cell (Oriel P/N 91 150 V, with KG-5 visible color filter), and the monocrystalline silicon was also calibrated by the National Renewable Energy Laboratory (NREL). Secondary ion mass spectrometry (CAMECA IMS 7f-Auto) was adopted to analyze elemental distribution of the films. AFM topography images and KPFM measurements were acquired on a Bruker Dimension Icon instrument. UPS was performed on a Thermo Escalab Xi+ analyzer and HeI lamp (21.22 eV).. Capacitance-voltage measurement and the point defects in the devices were investigated using the Phystech FT-1230 HERA DLTS system to record Deep-Level Transient Spectrum (DLTS). The reverse bias, pulse voltage, period width, and pulse width were set to -0.4, 0.2-0.6 V, 100 ms, and 10 ms, respectively. The CV and DLCP were characterized using a Keysight B1500A. For the CV measurements, we employed a frequency of 10 kHz, an AC bias of 30 mV, and a scanning voltage range of -1 to 1 V. For DLCP, the measurement was conducted at 10 kHz with a DC bias range from -0.2 to 0.2 V, while the sum of VDC + VAC was fixed at 30 mV. Carrier transport kinetics was studied via transient absorption spectroscopy (TAS), which was performed on a pump-probe system (Helios, Ultrafast System LLC). Additional Declarations There is NO Competing Interest. 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Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-7257351","acceptedTermsAndConditions":true,"allowDirectSubmit":false,"archivedVersions":[],"articleType":"Article","associatedPublications":[],"authors":[{"id":502991373,"identity":"3431b7a1-be98-4157-9c4f-de09db06a248","order_by":0,"name":"Tao 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1","display":"","copyAsset":false,"role":"figure","size":897449,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003en-type to p-type transition in Sb\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e2\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003eSe\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e3\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003e films.\u003c/strong\u003e a) UPS spectra of Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e and Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e films. b) Enlarge valence band region of the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e films. c) Energy-level diagrams of the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e samples. d, e) KPFM images of the Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e and Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e films. f, g) Surface potential difference (V\u003csub\u003eCPD\u003c/sub\u003e) of the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e samples. h, i) Hall effect measurement of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e films: h) Hall coefficient and i) majority carrier concentration as a function of Se/Sb ratio in Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e films.\u003c/p\u003e","description":"","filename":"image1.png","url":"https://assets-eu.researchsquare.com/files/rs-7257351/v1/a2f28a3c7bace5de1d5b9d14.png"},{"id":89622347,"identity":"a3dbbecd-fcb0-4e67-87e8-e28f7457e018","added_by":"auto","created_at":"2025-08-22 04:27:31","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":1160804,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eCarrier polarity transition mechanism.\u003c/strong\u003e a, b) Calculated formation energy of intrinsic defects as a function of Fermi level in Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e under the Sb-rich and Se-rich conditions. The vertical lines indicate the positions at which the Fermi level is pinned by the donor and acceptor defects with the lowest formation energies. c) Calculated transition energy levels of intrinsic defects in the bandgap of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e. d-f) Formation energy of shallow acceptor defects under the Sb-rich and Se-rich condition. g) Formation energy of deep donor selenium vacancy (V\u003csub\u003eSe\u003c/sub\u003e) defects. h, i) Schematic illustrations of the lattice structure and energetics of n-type and p-type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e, depicting the intrinsic defect doping mechanism.\u003c/p\u003e","description":"","filename":"image2.png","url":"https://assets-eu.researchsquare.com/files/rs-7257351/v1/0ea7e55fe8530352ae5e8977.png"},{"id":89622521,"identity":"8955fded-a898-4759-ad52-afd3e8f29322","added_by":"auto","created_at":"2025-08-22 04:35:31","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":4148948,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eDepth-resolved evidence of built-in homojunction in Sb\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e2\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003eSe\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e3\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003e absorber.\u003c/strong\u003e a, b) Cross-sectional SEM images of the Control and Target Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells. c, d) Cross-sectional KPFM images of FTO/Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e/Au devices based on the Control and Target Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e. The red dashed line delineates different interfaces. e, f) Corresponding cross-sectional potential profiles of the devices along the horizontal dashed lines in panels c and d. g) Bright-field TEM image of Target Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e. h, i, k, l) High-resolution TEM images and corresponding Fourier transform (FT) images of the p-type and n-type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e. j, m) Se and Sb element mappings by TEM-EDS.\u003c/p\u003e","description":"","filename":"image3.png","url":"https://assets-eu.researchsquare.com/files/rs-7257351/v1/1eecb9d5683d25f740409dae.png"},{"id":89622520,"identity":"efe098a6-3923-4465-91be-0d5d005b9de2","added_by":"auto","created_at":"2025-08-22 04:35:31","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":1783599,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eDevice structure and PV performance of Sb\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e2\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003eSe\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e3\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003e homojunction solar cells\u003c/strong\u003e. a) Standard n-i-p Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells without the homojunction. b) Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells with the homojunction composed by n-type and p-type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e layers. c) Schematic band diagram of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e homojunction solar cell and its promotion effect on charge transport. d) \u003cem\u003eJ-V\u003c/em\u003e curves of the champion devices based on Control- and Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3 \u003c/sub\u003esolar cells. (e) EQE spectra and the corresponding integrated current density of the champion devices. (f) Statistics of the \u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e, FF, \u003cem\u003eJ\u003c/em\u003e\u003csub\u003eSC\u003c/sub\u003e, and PCE of the Control and Target devices. (g) The summarized PCEs of related researches on Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells prepared by thermal evaporation method. (h) The summarized \u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e deficits and PCEs of related researches on the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells.\u003c/p\u003e","description":"","filename":"image4.png","url":"https://assets-eu.researchsquare.com/files/rs-7257351/v1/6402ddd171b986a11afeafdb.png"},{"id":89622349,"identity":"95c82492-513a-4702-a43c-8b0a712b4b92","added_by":"auto","created_at":"2025-08-22 04:27:31","extension":"png","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":1408599,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eInternal-homojunction‐driven carrier generation, transport, and recombination.\u003c/strong\u003e a, d) Energy level of each layer in Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells with homojunction and energy band diagram of the device in equilibrium state. b, c) Simulation of photoinduced carrier generation and recombination rate in Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells under illumination (AM1.5G) and short-circuit conditions. e) Mott-Schottky plots of the Control- and Target- Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells. f) \u003cem\u003eC-V\u003c/em\u003e and DLCP profiling. g, h) 2D color maps of the TA spectra for the devices. i) Normalized dynamic traces at 690 nm and the corresponding curves fitted by a double exponential equation for the devices.\u003c/p\u003e","description":"","filename":"image5.png","url":"https://assets-eu.researchsquare.com/files/rs-7257351/v1/1f96184f6082367ab3b5b711.png"},{"id":89622353,"identity":"4a08c1fa-62e2-49e2-9e6c-edc669f636ea","added_by":"auto","created_at":"2025-08-22 04:27:32","extension":"png","order_by":6,"title":"Figure 6","display":"","copyAsset":false,"role":"figure","size":1046989,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eBulk nonradiative recombination by deep-level defects. \u003c/strong\u003ea, d) DLTS spectra for the Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e and Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e devices. b, d) Magnified DLTS spectra from grey rectangular boxes in a and d. c, f) Schematic diagram of energy band and defect level of Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e and Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e. \u003cem\u003eE\u003c/em\u003e\u003csub\u003eC\u003c/sub\u003e, \u003cem\u003eE\u003c/em\u003e\u003csub\u003eV\u003c/sub\u003e and \u003cem\u003eE\u003c/em\u003e\u003csub\u003eF\u003c/sub\u003e represent CBM, VBM and Fermi level, respectively. g) Perfect crystal structure of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e, and point defects generated in the Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e and Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e films.\u003c/p\u003e","description":"","filename":"image6.png","url":"https://assets-eu.researchsquare.com/files/rs-7257351/v1/4f250a0b2d569eedc5e76b99.png"},{"id":106854645,"identity":"b6b06dce-e4cd-480e-a610-5064553f1bf3","added_by":"auto","created_at":"2026-04-14 07:12:46","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":10787154,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-7257351/v1/535c161e-1154-40a8-9803-aff30823160a.pdf"},{"id":89622351,"identity":"c3b21a6c-74d7-4ee2-a4a2-696a2a941257","added_by":"auto","created_at":"2025-08-22 04:27:31","extension":"docx","order_by":1,"title":"","display":"","copyAsset":false,"role":"supplement","size":4992161,"visible":true,"origin":"","legend":"Supplementary information","description":"","filename":"Supplementaryinformation.docx","url":"https://assets-eu.researchsquare.com/files/rs-7257351/v1/2a1f3f9cc3f24949ab231de8.docx"}],"financialInterests":"There is \u003cb\u003eNO\u003c/b\u003e Competing Interest.","formattedTitle":"\u003cp\u003e\u003cstrong\u003eHomojunction Sb\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e2\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003eSe\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e3\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003e Solar Cell\u003c/strong\u003e\u003c/p\u003e","fulltext":[{"header":"Main","content":"\u003cp\u003eAntimony selenide (Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e) material has attracted substantial research interest due to its outstanding optoelectronic properties and exceptional thermal and chemical stability.\u003csup\u003e1\u0026ndash;3\u003c/sup\u003e The nearly ideal bandgap (1.1-1.3 eV) and high absorption coefficient (\u0026gt;10\u003csup\u003e5\u003c/sup\u003e cm\u003csup\u003e-1\u003c/sup\u003e) support the theoretical photoelectric conversion efficiency of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e exceeding 30%.\u003csup\u003e4,5\u003c/sup\u003e As an emerging thin-film photovoltaic technology, Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells have demonstrated remarkable progress in material optimization, interface engineering, and device structure design, achieving power conversion efficiency exceeding 10%.\u003csup\u003e6\u0026ndash;9\u003c/sup\u003e However, this performance still lags behind mature chalcogenide thin-film technologies, such as CdTe and Cu(In,Ga)Se\u003csub\u003e2\u003c/sub\u003e photovoltaics (PVs).\u003csup\u003e10\u0026ndash;12\u003c/sup\u003e The large deficit in open-circuit voltage (\u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e) is the primary factor limiting the performance of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells, with experimentally obtained \u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e values typically about 500 mV lower than the theoretical radiation limit voltage. This energy loss is attributed to two key factors: (1) insufficient driving force for charge separation at the junction, due to a weak built‑in electric field; and (2) severe non‑radiative recombination via deep traps at both the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e absorber/buffer interfaces and within the bulk material.\u003csup\u003e13\u0026ndash;15\u003c/sup\u003e\u003c/p\u003e\n\u003cp\u003eEffective control of semiconductor polarity, combined with judicious device structure design, has proven essential for overcoming efficiency challenges in established photovoltaic systems.\u003csup\u003e16\u0026ndash;18\u003c/sup\u003e In conventional PV technologies such as crystalline silicon and III-V semiconductors, the precisely controlled incorporation of donor or acceptor dopants enables the formation of n-type and p-type regions, respectively, which establish strong p-n homojunctions with well-defined built-in electric fields.\u003csup\u003e19,20\u003c/sup\u003e The enhanced internal electric fields effectively suppress carrier recombination, thereby improving device performance. In emerging photovoltaic technologies (such as metal-halide perovskites and organic PVs), intrinsic/extrinsic doping and charge transfer effects have been employed to modulate charge carrier concentrations and construct heterojunction or homojunction architectures, enhancing charge extraction.\u003csup\u003e21\u0026ndash;24\u003c/sup\u003e Remarkably, homojunction engineering has achieved substantial efficiency improvements across diverse perovskite solar cells, spanning lead-based, tin-based, and mixed lead-tin systems in both regular (n-i-p) and inverted (p-i-n) architectures.\u003csup\u003e25\u0026ndash;29\u003c/sup\u003e These achievements underscore the significant potential of rational carrier-polarity control and internal junction engineering in reducing voltage losses. Applying these theoretical frameworks to Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells offer a viable approach to mitigate their intrinsic voltage limitations. Therefore, we focus on the co-optimization of junction design and defect passivation to improve the photovoltaic performance of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells.\u003c/p\u003e\n\u003cp\u003eHere,\u0026nbsp;we develop a composition-driven intrinsic doping strategy for polarity control and constructed Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e homojunction solar cells. By precisely tuning the chemical potential of Se and Sb during thermal evaporation, we successfully switched the conductivity of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e films from n-type to p-type, enabling a monolithic p-n homojunction within the absorber. We also conducted density functional theory (DFT) calculations to elucidate the mechanistic details of the conductivity-type transition process. This built-in junction provides an extra internal electric field that widens the depletion region, enhancing charge separation and suppressing non-radiative recombination. Comprehensive spectral and electrical characterizations reveal an approximately 0.2 eV internal potential gradient, along with suppressed defects and extended carrier lifetime in the device. Ultrafast spectroscopy and depth-resolved simulations further demonstrate faster carrier transport rate, with recombination losses reduced by over an order of magnitude. Notably, Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells with p-n homojunction demonstrate a record power conversion efficiency of 10.15% for thermally evaporated devices, along with the lowest \u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e deficit reported for all kinds of SSCs so far. This work highlights a promising new strategy for improving SSC performance via homojunction engineer to synergistically enhance the interfacial charges extraction and suppress non-radiative recombination.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eP\u003c/strong\u003e\u003cstrong\u003eolarity conversion and intrinsic doping mechanisms\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eAlthough\u0026nbsp;Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e can exhibit ambipolar charge transport and has been reported as either an n‑type or p‑type semiconductor, deliberate strategies to control its conductivity remain undeveloped.\u003csup\u003e14,30\u0026ndash;32\u003c/sup\u003e First‑principles calculations suggest that the material\u0026rsquo;s conductivity type is governed by growth conditions and the relative chemical potentials of Se and Sb, yet experimental validation of this relationship and methods for targeted polarity tuning are absent.\u003csup\u003e13,33,34\u003c/sup\u003e To achieve controlled conductivity switching in Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e, we developed an intrinsic doping method based on precursor composition tuning. During ball‑milling synthesis, we adjusted the Se and Sb chemical potentials by introducing a controllable Se excess into the precursor mixture (detailed preparation process is provided in the experimental section). This approach compensates for selenium loss, which is inevitable in the conventional thermal deposition by the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e compound as the precursor material due to selenium\u0026rsquo;s high vapor pressure during film deposition and annealing.\u003csup\u003e35\u003c/sup\u003e Our method herein enables precise stoichiometric control of the resulting absorber layer, with Se/Sb ratios tunable from 1.41 to 1.59. For both the control Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e absorber and the n-type absorber region in the target Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e, we employed a precursor Se/Sb ratio of 1.5. For the p-type absorber region in the target\u0026nbsp;Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e device, we increased the precursor Se/Sb ratio to 1.6, as first‑principles calculations indicate that selenium‑rich conditions favor p‑type conductivity in Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e.\u003csup\u003e33\u003c/sup\u003e In this study, both n-type and p-type absorbers were prepared by thermal evaporation method, which exhibit compact morphologies and pure orthorhombic Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e phase, as confirmed by scanning electron microscopy (SEM) and X-ray diffraction (XRD) analysis (Supplementary Fig. S1). For the convenience of discussion, the SSCs without and with homojunction structure is labeled as Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e and Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e samples.\u003c/p\u003e\n\u003cp\u003eWe first examined the surface composition of the films using X‑ray photoelectron spectroscopy (XPS) (Supplementary Fig. S1 and S2). The Se/Sb atomic ratio on the surface of the Target‑Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e film is increased from 1.46 (Control‑Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e) to 1.56, indicating successful selenium enrichment in the p‑type layer.\u0026nbsp;Since XPS can only test information within a few nanometers of the surface depth, we performed cross‑sectional energy‑dispersive X‑ray spectroscopy (EDS). Nine measurement points across both the n‑type and p‑type regions show that the Se/Sb ratio in the n‑type region remains approximately 1.46 for both samples, whereas the p‑type region of the Target‑Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e reaches 1.56 (Supplementary Fig. S3). For more accurate bulk composition, inductively coupled plasma atomic emission spectroscopy (ICP‑AES) reveals Se/Sb ratios of 1.45 in the Control‑Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e and 1.57 in the p‑type region of the Target‑Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e (Supplementary Fig. S4), in excellent agreement with XPS and EDS data. Furthermore, secondary ion mass spectrometry (SIMS) and EDS of transmission electron microscopy (TEM) depth profiling both provide a continuous elemental distribution through the absorber thickness (Supplementary Fig. S5 and S18b). The Control‑Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e film exhibits essentially constant Sb and Se signals, whereas the Target‑Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e film shows a gradual increase in Se/Sb ratio from the CdS interface toward the film surface. Collectively, these multi-dimensional measurements conclusively demonstrate that our precursor-composition strategy effectively drives Se/Sb variations in the film, successfully realizing the targeted composition‑driven homojunction structure.\u003c/p\u003e\n\u003cp\u003eWe then studied the electronic structure of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e thin films. Firstly, we characterized the energy level structures of the Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e and Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e samples using ultraviolet photoelectron spectroscopy (UPS) (Fig. 1a-c). For the Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e samples, the Fermi level is determined to be \u0026minus;4.30\u0026thinsp;eV, close to the conduction band minimum (CBM) of \u0026minus;3.76\u0026thinsp;eV, indicating the n-type characteristics of the thin film. For the surface of Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e samples, the Fermi level shifts toward the valence band maximum (VBM), with offset reducing from 0.67 eV to 0.54 eV (Fig. 1c), indicating a transition in conductivity type from n-type to p-type. The Kelvin probe force microscopy (KPFM) characterization further confirms the above results by testing the surface potential. Figure 1d and 1e show the KPFM potential images of the Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e and the Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e, respectively, with relatively uniform contact potential distributions in both samples. In addition, it can be seen from Figure 1e that the surface potential of the Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e film is lower than that of the Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e film (Fig. 1d), because the accumulated holes cause the Fermi level to be located near the valence band. On the contrary, the Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e film exhibits a higher surface potential, which is related to the n-type doping of the sample. To guarantee the reliability of the results, we conducted statistical analysis on the surface potential distribution of the samples (Fig. 1f, g). According to the statistical results, the surface potential difference of the target sample reduced by about 0.2 eV compared to the control sample, consistent with the Fermi level shift measured by UPS.\u003c/p\u003e\n\u003cp\u003eAs UPS and KPFM are both surface-sensitive techniques, further evidence is significant for analyzing the effects of doping in the bulk regions of the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e samples. Hall effect measurement is one of the most reliable techniques for characterizing the polarity and concentration of carriers in semiconductors.\u003csup\u003e19,36\u003c/sup\u003e The Hall coefficients (\u003cem\u003eR\u003c/em\u003e\u003csub\u003eH\u003c/sub\u003e) and majority carrier concentrations of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e samples with different Se/Sb ratio are shown in Fig. 1h and 1i. For the Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e sample with a Se/Sb atomic ratio of 1.46, it exhibits n-type characteristics with a negative \u003cem\u003eR\u003c/em\u003e\u003csub\u003eH\u003c/sub\u003e value, corresponding to an electron concentration of 1.99\u0026times;10\u003csup\u003e14\u003c/sup\u003e cm\u003csup\u003e-3\u003c/sup\u003e, obtained from equation of n = 1/(\u003cem\u003eR\u003c/em\u003e\u003csub\u003eH\u003c/sub\u003eq), where n is the carrier concentration and q is the elementary charge.\u003csup\u003e36\u003c/sup\u003e The more Sb-rich composition (with a Se/Sb atomic ratio of 1.42) results in increased electron concentration and enhanced n-type characteristics. However, as the Se/Sb ratio rises to 1.51, transforming to a Se-rich state, \u003cem\u003eR\u003c/em\u003e\u003csub\u003eH\u003c/sub\u003e turns positive, indicating a transition to p-type behavior where holes dominate charge transport. The Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e film with more Se-rich composition\u003csub\u003e\u0026nbsp;\u003c/sub\u003e(Se/Sb of 1.56) demonstrates a more p-type characteristics with hole concentration reaching 6.57\u0026times;10\u003csup\u003e14\u003c/sup\u003e cm\u003csup\u003e-3\u003c/sup\u003e. Clearly, as the composition of the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e thin film deviates from stoichiometric ratio, the measured carrier concentration increases accordingly. It is especially pronounced in p‑type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e samples, which exhibit higher carrier densities\u0026mdash;a phenomenon typically attributed to changes in the semiconductor\u0026apos;s intrinsic defect properties. In other words, intrinsic doping is the primary factor governing the conductivity type transition in Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e.\u003c/p\u003e\n\u003cp\u003eWe then employed density functional theory (DFT) to elucidate the mechanism regarding the transition of thin film conductivity types. Supplementary Fig. S6 and S7 depict the computed formation energies of all intrinsic point defects as a function of Fermi level under Sb‑rich and Se‑rich conditions. These values are consistent with early theoretical calculations.\u003csup\u003e33,37,38\u003c/sup\u003e Evidently, anion vacancies (V\u003csub\u003eSe1\u003c/sub\u003e, V\u003csub\u003eSe2\u003c/sub\u003e, V\u003csub\u003eSe3\u003c/sub\u003e) exhibit deep donor behavior, possessing thermodynamic transition levels far from the CBM. In contrast, cation vacancies (V\u003csub\u003eSb1\u003c/sub\u003e, V\u003csub\u003eSb2\u003c/sub\u003e) act as shallow acceptors, with transition levels lying above the VBM by several \u003cem\u003ek\u003c/em\u003e\u003csub\u003eB\u003c/sub\u003eT. All intrinsic anti‑site defects, including cation-replace-anion (Sb\u003csub\u003eSe1\u003c/sub\u003e, Sb\u003csub\u003eSe2\u003c/sub\u003e, Sb\u003csub\u003eSe3\u003c/sub\u003e) and anion-replace-cation (Se\u003csub\u003eSb1\u003c/sub\u003e, Se\u003csub\u003eSb2\u003c/sub\u003e) defects, display amphoteric character, acting as donors near VBM and acceptors near CBM; notably, only the Se\u003csub\u003eSb2\u003c/sub\u003e defect exhibits shallow acceptor property (about 0.10 eV above the VBM). In addition, cation interstitial defect (Sb\u003csub\u003ei\u003c/sub\u003e) is also amphoteric, with a deep mid‑gap transition, while anion interstitial defect (Se\u003csub\u003ei\u003c/sub\u003e) remains electrically inert except under extreme p‑type conditions.\u003c/p\u003e\n\u003cp\u003eTo identify the dominant defect species, we selected representative intrinsic defects with low formation energy for discussion and presented their transition level diagrams in Fig. 2a, b. In Sb‑rich Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e with a bandgap of 1.16 eV, +2 charged V\u003csub\u003eSe2\u003c/sub\u003e defect exhibits the minimal formation energy when the Fermi level is below the mid‑gap (~0.58 eV below the CBM), whereas the \u0026minus;1 charged Sb\u003csub\u003eSe2\u003c/sub\u003e defect becomes most favorable upon Fermi level approaching to the CBM. The mutual compensation between these donor and acceptor centers pins the Fermi level above mid‑gap (blue dashed line shown in Fig. 2a), resulting in pronounced n‑type conductivity. In contrast, Se‑rich conditions induce a considerably reduction (\u0026gt;0.6 eV) in the formation energies of shallow acceptors V\u003csub\u003eSb1\u003c/sub\u003e, V\u003csub\u003eSb2\u003c/sub\u003e, and Se\u003csub\u003eSb2\u003c/sub\u003e (Fig. 3d-g). The decrease in formation energy significantly increases the equilibrium concentrations of these defects, shifts the Fermi level towards VBM, which fosters p‑type behavior. Meanwhile, the formation energy of V\u003csub\u003eSe2\u003c/sub\u003e rises substantially under Se-rich conditions, reducing the compensation effect on hole carriers. Combining these DFT findings with experimental observations (Fig. 1), we propose the following mechanism (Fig. 2h, i). For Sb‑rich Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e, the facile formation of deep donors depletes hole populations and pins the Fermi level on the n‑type side; while for Se‑rich Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e, shallow acceptors near VBM are easily formed, which increase the hole density and shift the Fermi level towards p‑type side. Such defect‑engineered control of conductivity type provides a novel approach for constructing homojunction Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eHomojunction construction and photovoltaic performance of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eLeveraging the precise control over Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e conductivity type, we fabricated a monolithic homojunction structure by integrating n‑type and p‑type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e layers into planar solar cells with the architecture: fluorine-doped tin oxide (FTO)/CdS/Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e/Spiro-OMeTAD/Au. In this configuration, CdS,\u0026nbsp;Spiro-OMeTAD\u0026nbsp;(2,2\u0026apos;,7,7\u0026apos;-tetrakis(N, N-di-p-methoxyphenyl-amine)-9,9\u0026apos;-spirobifluorene), and\u0026nbsp;Au\u0026nbsp;serve as the electron-transporting layer, hole-transporting layer, and\u0026nbsp;back electrode, respectively.\u0026nbsp;Cross-sectional SEM images of the\u0026nbsp;Control- and Target-\u0026nbsp;SSCs\u0026nbsp;(Fig. 3a, b)\u0026nbsp;show that both Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e films display densely packed, columnar grains spanning the full absorber thickness, suggesting high crystallinity and compact morphology of the as-synthesized Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e films. Remarkably, this junction does not cause discernible interfaces or grain boundaries, which is consistent with the TEM images (Fig. 3g, j, and m). Furthermore, we compared the lattice fringes of n-type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e near the CdS and p-type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e and performed Fourier transforms (FT) of their respective regions (Fig. 3h, i, k, and l). The lattice spacings in both regions are 0.315 nm, which corresponds to the (211) planes of orthorhombic Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e.\u0026nbsp;Fourier transform (FT) analysis further confirms that the crystal plane and orientation of n-type and p-type\u0026nbsp;Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e are identical. Combined with the XRD results (Supplementary Fig. S1), these findings confirm the formation of a homojunction without lattice mismatch.\u003c/p\u003e\n\u003cp\u003eTo further substantiate the formation of the p-n junction in Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e absorber, we performed cross-sectional KPFM potential characterizations of the FTO/Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e/Au devices (Fig. 3c and 3d), which were collected under dark conditions on freshly cleaved samples to avoid artefacts. In the control device, the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e absorber exhibits a nearly uniform surface potential (about 180-200 mV) throughout the entire thickness (Fig. 3c), confirming its single-type (n-type) conductivity without internal electric field. In contrast, the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e absorber in the target device exhibits a distinct potential gradient (Fig. 3d), decreasing steadily from the FTO interface to the Au contact. This spatial variation visualizes the built-in electric field generated by the p-n homojunction within Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e layer, which facilitates efficient charge separation and suppresses bulk recombination in the absorber, beneficial for improving the optoelectronic performance of the device.\u003c/p\u003e\n\u003cp\u003eConsequently, we investigated the photovoltaic performance of SSCs with and without the homojunction under AM 1.5G solar irradiance (100 mW cm\u003csup\u003e-\u003c/sup\u003e\u003csup\u003e2\u003c/sup\u003e), with schematic illustrations of the device structures shown in Fig. 4a-c. In fact, we optimized the composition and thickness of the p-type and n-type regions (see details in Supplementary Note 2 and Supplementary Fig. S8-12). As a result, the establishment of p-n homojunction increases the efficiency from 8% for the Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e device to 10.15% for the Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e device, which is the highest efficiency reported for thermally evaporated Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells to date (Fig. 4d, g and Supplementary Table S3). Specifically, the open circuit voltage (\u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e), short-circuit current density (\u003cem\u003eJ\u003c/em\u003e\u003csub\u003eSC\u003c/sub\u003e), and fill factor (FF) are improved from 428 mV to 492 mV, 29.59 to 31.48 mA cm\u003csup\u003e-2\u003c/sup\u003e, and 63.21% to 65.56 %, respectively (Table 1), which are 15%, 6% and 4% higher than that of the control device, respectively. Obviously, the homojunction engineering significantly improves photovoltaic parameters, especially \u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e (Fig. 4f). To explore the photocurrent generation of the devices, we conducted external quantum efficiency (EQE) spectroscopy characterization. It shows the enhanced wavelength response in the range of 550-1000 nm and demonstrated that the increase in \u003cem\u003eJ\u003c/em\u003e\u003csub\u003eSC\u003c/sub\u003e may may stem from the improved film quality (Fig. 4e), as the target device exhibits the same cut-off edge as the control device (approximately 1030 nm). This consistency with the ultraviolet-visible absorption spectra (measured bandgap of 1.205 and 1.202 eV for the control and target devices, respectively (Supplementary Fig. S1d)). As for the \u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e, its significant improvement results in a low \u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e deficit of 0.459 V, which is the lowest value reported among all Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells to date (Fig. 4h). The increase in \u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e arises from enhanced carrier transport and suppressed nonradiative recombination, benefited from the homojunction engineering (Fig. 4c), as evidenced by a reduction in the dark saturation current density \u003cem\u003eJ\u003c/em\u003e\u003csub\u003e0\u003c/sub\u003e from 1.43\u0026times;10\u003csup\u003e-6\u003c/sup\u003e in the Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e device to 8.70\u0026times;10\u003csup\u003e-8\u0026nbsp;\u003c/sup\u003ein the Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e device (Supplementary Fig. S13). Concurrently, the fill factor improvement stems from a lowered diode ideality factor, which produces a more ideal \u003cem\u003eJ-V\u003c/em\u003e characteristic and thus enhances FF. These results demonstrate that composition‑driven homojunction engineering in Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e constructs a robust internal electric field for improved photovoltaic performance.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eTable 1.\u003c/strong\u003e Photovoltaic parameters of the Control- and Target- Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells.\u003c/p\u003e\n\u003cdiv align=\"center\"\u003e\n \u003ctable border=\"0\" cellspacing=\"0\" cellpadding=\"0\" width=\"100%\"\u003e\n \u003ctbody\u003e\n \u003ctr\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003eDevices\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003e\u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e (V)\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003e\u003cem\u003eJ\u003c/em\u003e\u003csub\u003eSC\u003c/sub\u003e (mA/cm\u003csup\u003e2\u003c/sup\u003e)\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003eFF (%)\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003ePCE (%)\u003c/p\u003e\n \u003c/td\u003e\n \u003c/tr\u003e\n \u003ctr\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003eControl\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003e0.428\u003c/p\u003e\n \u003cp\u003e(0.408\u0026plusmn;0.009)\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003e29.59\u003c/p\u003e\n \u003cp\u003e(29.64\u0026plusmn;0.43)\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003e63.21\u003c/p\u003e\n \u003cp\u003e(62.25\u0026plusmn;1.00)\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003e8.00\u003c/p\u003e\n \u003cp\u003e(7.53\u0026plusmn;0.23)\u003c/p\u003e\n \u003c/td\u003e\n \u003c/tr\u003e\n \u003ctr\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003eTarget\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003e0.492\u003c/p\u003e\n \u003cp\u003e(0.486\u0026plusmn;0.006)\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003e31.48\u003c/p\u003e\n \u003cp\u003e(31.35\u0026plusmn;0.14)\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003e65.56\u003c/p\u003e\n \u003cp\u003e(65.06\u0026plusmn;0.47)\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd style=\"width: 20px;\"\u003e\n \u003cp\u003e10.15\u003c/p\u003e\n \u003cp\u003e(9.91\u0026plusmn;0.20)\u003c/p\u003e\n \u003c/td\u003e\n \u003c/tr\u003e\n \u003c/tbody\u003e\n \u003c/table\u003e\n\u003c/div\u003e\n\u003cp\u003e\u003cstrong\u003eInternal homojunction‐driven modulation of carrier generation, transport, and recombination\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe modulation of built-in electric fields typically alters charge carrier characteristics. In order to elucidate the mechanism through which homojunction construction enhances Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cell performance, we performed equilibrium band-structure analysis, numerical simulations, electrostatic profiling, and ultrafast transient absorption spectroscopy (TAS) to investigate charge carrier dynamics. Figure 5a and 5d present schematic energy level diagram and illustrate the working principle of homojunction-SSCs, demonstrating the transport mode of charge carriers between layers. When photons are absorbed by both p-type and n-type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e, charge carriers are generated on both sides of the junction. Subsequently, electrons and holes drift in opposite directions: electrons inject into the CdS layer and transport to the FTO cathode, whereas holes inject into the spiro-OMeTAD layer and transport to the Au anode. In conventional n-i-p structured SSCs, the limited dynamic driving force for carrier transport within the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e film often results in significant carrier recombination. Typically, photo-induced electrons and holes are extracted mostly at the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e/ETL and the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e/HTL interfaces, respectively. Therefore, our designed homojunction introduces an additional built-in electric field, extending the depletion region deeper into the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e absorber layer. This enhanced field promotes the rapid separation of electrons and holes, driving them toward ETL and HTL, respectively (Fig. 5d), while simultaneously reducing carrier recombination.\u003c/p\u003e\n\u003cp\u003eWe further applied the SCAPS-1D simulation tool, widely employed for modeling thin-film photovoltaic devices, focusing on the generation and recombination dynamics of photo carriers. As illustrated in Fig. 5b, both the Control and Target architectures exhibit their maximum generation rates at the CdS/n-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e junction. However, the target device maintains a pronounced secondary peak at the internal p-n interface, clearly demonstrating that the built-in field facilitates spatial carrier separation. The corresponding recombination rate profiles (Fig. 5c) reveal that the control device generates a localized recombination maximum of about 10\u003csup\u003e21\u003c/sup\u003e cm\u003csup\u003e-3\u003c/sup\u003e\u0026middot;s\u003csup\u003e-1\u003c/sup\u003e within the central region of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e layer, whereas the homojunction reduces this rate by more than an order of magnitude. We attribute this suppression to the unidirectional electric field induced by the internal p-n junction pushing electrons and holes toward opposite contacts, thereby diminishing interfacial non-radiative recombination. Meanwhile, the p-type region of the homojunction may result in lower density of deep-level trap states, contributing to reduced bulk non-radiative losses.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eOn the other hand, Mott-Schottky analysis shows an increase in \u003cem\u003eV\u003c/em\u003e\u003csub\u003ebi\u003c/sub\u003e from 0.46\u0026thinsp;V to 0.56\u0026thinsp;V (Fig. 5e). The capacitance-voltage (\u003cem\u003eC-V\u003c/em\u003e) and drive-level capacitance profiling (DLCP) reveal that the depletion width expands from 131\u0026thinsp;nm in the control device to 182\u0026thinsp;nm in the target device without introducing additional interface traps (Fig. 5f). We then conducted ultrafast TAS analysis on FTO/CdS/Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e devices to probe the dynamics of charge carriers excited by pulsed photo. As shown in Fig. 5 g-i, the contour maps based on TA spectra of Control- and Target- Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e devices both display a photo-induced absorption (PIA) peak at approximately 690 nm which is attributed to the characteristic absorption of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e. Notably, the TAS signal of the target device decays faster, indicating a more efficient transport of charge carriers from the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e film to the adjacent interface layer. By fitting and analyzing the decay kinetics of PIA signals, the average TA decay lifetime of the Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e device (1.75 ns) is much smaller than that of the control device (4.15 ns), as shown in Fig. 5i. These results further confirm the critical role of homojunctions in accelerating the extraction and transport of charge carriers.\u003c/p\u003e\n\u003cp\u003eDeep‐level defects critically govern carrier lifetimes via Shockley-Read-Hall (SRH) recombination, particularly the bulk nonradiative losses associated with lattice imperfections.\u003csup\u003e16,39\u003c/sup\u003e To quantify how homojunction engineering reshapes the defect landscape and suppresses SRH recombination, we performed deep-level transient spectroscopy (DLTS) on Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e and Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e devices. By carefully adjusting the applied pulse voltage (0.2-0.6 V), we mitigated artifact peaks arising from capacitor bridge recovery delays observed at inappropriate pulse voltage levels, while identifying defects at different depths (Fig. 6a and 6d).\u003csup\u003e40\u003c/sup\u003e In the DLTS spectra, positive peaks correspond to majority-carrier traps, while negative peaks indicate minority-carrier traps. Based on the DLTS spectra, Arrhenius plots are analyzed (Supplementary Figure S14a and b), and defect-related parameters, including the trap energy level (\u003cem\u003eE\u003c/em\u003e\u003csub\u003eT\u003c/sub\u003e), trap cross-section (\u0026sigma;), and trap density (\u003cem\u003eN\u003c/em\u003e\u003csub\u003eT\u003c/sub\u003e), are extracted (Supplementary Tab. S8). In the Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u0026nbsp;\u003c/sub\u003edevice, three hole-traps (H1, H2, H3) and one electron-trap (E1) are identified. Their respective energy levels are located at 0.364, 0.548, and 0.679 eV above the VBM for the hole traps and 0.602 eV below the CBM for the electron trap (Fig. 6a-c). Surprisingly, there is only one electron-trap (E2) located at the 0.665 eV energy level below CBM is identified in the target device. Based on the DFT calculations (Figure 2), the H1, H2, H3, and E1 defects can be attributed to Sb\u003csub\u003eSe1\u003c/sub\u003e, Sb\u003csub\u003eSe3\u003c/sub\u003e, Sb\u003csub\u003eSe2\u003c/sub\u003e antisite defects, and Se vacancy defect (V\u003csub\u003eSe3\u003c/sub\u003e) when the film is in Sb-rich state. As for the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u0026nbsp;\u003c/sub\u003ein Se-rich state, the E2 defect can be attributed to Se\u003csub\u003eSb1\u003c/sub\u003e defect.\u003c/p\u003e\n\u003cp\u003eThrough comparative analysis of defect characteristics between the two devices, we confirm that homojunction engineering effectively passivate the original Sb\u003csub\u003eSe3\u003c/sub\u003e, Sb\u003csub\u003eSe1\u003c/sub\u003e, Sb\u003csub\u003eSe2\u003c/sub\u003e, and V\u003csub\u003eSe\u003c/sub\u003e defects in Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e. In the target device, the newly emerged Se\u003csub\u003eSb1\u003c/sub\u003e defect demonstrates more favorable properties, exhibiting a lower defect density (6.34\u0026times;10\u003csup\u003e13\u003c/sup\u003e cm\u003csup\u003e-3\u003c/sup\u003e) and a smaller capture cross-section (3.87\u0026times;10\u003csup\u003e-16\u003c/sup\u003e cm\u0026sup2;). According to the defect-assisted SRH recombination model, the carrier lifetime (\u003cem\u003e\u0026tau;\u003c/em\u003e) associated with a specific defect can be evaluated by Equation (1):\u003c/p\u003e\n\u003cp\u003e\u003cimg src=\"data:image/png;base64,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\" width=\"460\" height=\"51\"\u003e\u003c/p\u003e\n\u003cp\u003eWhere \u003cem\u003e\u0026nu;\u003c/em\u003e\u003csub\u003eth\u003c/sub\u003e is the charge heating rate, which is 10\u003csup\u003e7\u003c/sup\u003e cm s\u003csup\u003e-1\u003c/sup\u003e at room temperature.\u003csup\u003e35\u003c/sup\u003e Specifically, we focus on traps with the shortest carrier lifetime because they have the most significant impact on carrier recombination. Our calculations reveal that in Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e, the Sb\u003csub\u003eSe2\u003c/sub\u003e defect has the shortest lifetime of 543 ns, followed by V\u003csub\u003eSe3\u003c/sub\u003e, identifying them as the dominant recombination centers. In contrast, the Se\u003csub\u003eSb1\u003c/sub\u003e defect in Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e shows a substantially longer lifetime (4075.68 ns), indicating a slightly negative contribution to carrier recombination. Moreover, according to the SRH recombination theory for multiple recombination centers, the total SRH recombination rate (\u003cem\u003eU\u003c/em\u003e\u003csub\u003eSRH\u003c/sub\u003e) is the sum of the individual recombination rates (\u003cem\u003eU\u003c/em\u003e\u003csub\u003ei, SRH\u003c/sub\u003e) of each recombination center \u003cem\u003ei\u003c/em\u003e,\u003csup\u003e41\u003c/sup\u003e as expressed in equation (2):\u003c/p\u003e\n\u003cp\u003e\u003cimg src=\"data:image/png;base64,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\" width=\"442\" height=\"55\"\u003e\u003c/p\u003e\n\u003cp\u003eBased on the calculations, homojunction engineering reduces both the density and variety of traps in the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e bulk phase. According to equation (2), the reduction in recombination centers consequently reduces the SRH recombination of carriers.\u003c/p\u003e\n\u003cp\u003eIn the homojunction target device, the n-type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e layer is only about 50 nm thick\u0026mdash; comparable to the thickness of CdS electron-transport layer, so p-type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e forms the main bulk of the absorber. DLTS characterization is primarily applied to probe the deep-level defects in the p-type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e region. Analysis of pure p-type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e-based devices (Supplementary Fig. S15) reveals only a single E2 trap, with energy level, density, and capture cross-section values consistent with those observed in the target device. This discovery indicates that the p-type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e intrinsically maintains a more favorable defect profile. Therefore, it is because of the superiority of p-type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e in reducing bulk SRH recombination along with p-n homojunction expanded space charge region that we achieved a momentous improvement in device performance.\u003c/p\u003e"},{"header":"Conclusion","content":"\u003cp\u003eIn conclusion, we demonstrate a cation/anion composition-driven homojunction engineering which achieves an efficiency breakthrough in Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e photovoltaics.\u0026nbsp;Our\u0026nbsp;study reveals\u0026nbsp;that the modulation of\u0026nbsp;elemental\u0026nbsp;chemical potential during thermal evaporation\u0026nbsp;enables\u0026nbsp;controls\u0026nbsp;over\u0026nbsp;charge carrier polarity, facilitating the fabrication of\u0026nbsp;a\u0026nbsp;monolithic\u0026nbsp;p-n\u0026nbsp;homojunction.\u0026nbsp;Furthermore, we elucidate the mechanism of this strategy in suppressing non-radiative recombination from two aspects: (i) the homojunction\u0026nbsp;increases\u0026nbsp;the depletion width from 131 nm to 182 nm\u0026nbsp;with introducing an additional built-in field to\u0026nbsp;enhance\u0026nbsp;charge separation; and (ii) the\u0026nbsp;Se-rich components passivate\u0026nbsp;deep-level defects, reducing trap density by over 70%. Depth-resolved KPFM and\u0026nbsp;carrier dynamic analysis\u0026nbsp;confirm\u0026nbsp;the\u0026nbsp;accelerated carrier extraction and\u0026nbsp;an\u0026nbsp;order-of-magnitude reduction in recombination\u0026nbsp;losses. The integration of homojunction into planar n-i-p devices\u0026nbsp;enables a record power conversion efficiency\u0026nbsp;of\u0026nbsp;10.15 %\u0026nbsp;for thermally evaporated\u0026nbsp;Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells, along with\u0026nbsp;a minimal \u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e deficit of 0.459 V\u0026mdash;the lowest reported\u0026nbsp;value among\u0026nbsp;Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e cells. This investigation underscores the synergistic advantages of homojunction engineering in efficient charge separation and recombination suppression, pointing out a new direction for optimizing optoelectronic performance of antimony chalcogenide thin-film solar cells.\u0026nbsp;\u003c/p\u003e"},{"header":"Declarations","content":"\u003cp\u003e\u003cstrong\u003eData availability\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eAll of the data needed to evaluate the conclusions in the paper are present in the paper and its Supplementary Information.\u003c/p\u003e\n\u003cp\u003eAcknowledgments\u003c/p\u003e\n\u003cp\u003eThis work was supported by National Natural Science Foundation of China (22275180),\u0026nbsp;the Fundamental Research Funds for the Central Universities (No. WK2490000002), Major Science and Technology Projects of Anhui Province (AHZDCYCX-LSDT2023-10), and the University Synergy Innovation Program of Anhui Province (GXXT-2023-031).\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAuthor contributions\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAuthors and Affiliations\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eDepartment of Materials Science and Engineering, School of Chemistry and Materials Science, University of Science and Technology of China, Hefei, People\u0026rsquo;s Republic of China.\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eJunjie Yang,\u003csup\u003e\u0026nbsp;\u003c/sup\u003eJianyu Li, Shuwei Sheng, Zhiyuan Cai, Zichen Ruan, Bo Che, Rongfeng Tang,\u0026nbsp;and\u003csup\u003e\u0026nbsp;\u003c/sup\u003eTao Chen\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eInstitute of Deep Space Sciences, Deep Space Exploration Laboratory, Hefei 230088,\u0026nbsp;\u003c/strong\u003e\u003cstrong\u003eP. R. China.\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eTao Chen\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eInstitute of Optoelectronic Materials and Devices, School of Materials Science and Engineering, Jiangxi University of Science and Technology, Ganzhou 341000, P. R. China.\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eKe Li\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eContributions\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eJ. Y., J. L. and S. S. contributed equally to this work.\u0026nbsp;J.Y. and T.C. conceived of the original concept and designed the experiments. J.Y., J.L. and S.S. fabricated the devices and conducted the photovoltaic and optical characterization and analysis. J. L. did the TAS measurements and performed the SEM, TEM analysis. K.L. Z. R.\u0026nbsp;and B. C. were involved in materials\u0026nbsp;characterization\u0026nbsp;and device simulation. Z. C. conducted the DFT calculations. J.Y., R.T. and T. C. co-wrote the paper. T.C., R.T., S.S. and J. Y. revised the paper with all authors commenting on the paper.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eCompeting interests\u0026nbsp;\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe authors declare no competing interests.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAdditional information\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003ePublisher\u0026rsquo;s note Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eSource data\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eSource Data\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\n\u003cli\u003eZhou, Y. \u003cem\u003eet al.\u003c/em\u003e Thin-film Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e photovoltaics with oriented one-dimensional ribbons and benign grain boundaries. \u003cem\u003eNat\u003c/em\u003e\u003cem\u003e.\u003c/em\u003e\u003cem\u003e Photon\u003c/em\u003e\u003cem\u003e.\u003c/em\u003e \u003cstrong\u003e9\u003c/strong\u003e, 409\u0026ndash;415 (2015).\u003c/li\u003e\n\u003cli\u003eZhao, Q. \u003cem\u003eet al.\u003c/em\u003e Grain‐Boundary Elimination via Liquid Medium Annealing toward High‐Efficiency Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e Solar Cells. \u003cem\u003eAdv\u003c/em\u003e\u003cem\u003e.\u003c/em\u003e\u003cem\u003e Mater\u003c/em\u003e\u003cem\u003e.\u003c/em\u003e \u003cstrong\u003e37\u003c/strong\u003e, 2414082 (2025).\u003c/li\u003e\n\u003cli\u003eWen, X. \u003cem\u003eet al.\u003c/em\u003e Vapor transport deposition of antimony selenide thin film solar cells with 7.6% efficiency. \u003cem\u003eNat\u003c/em\u003e\u003cem\u003e.\u003c/em\u003e\u003cem\u003e Commun\u003c/em\u003e\u003cem\u003e.\u003c/em\u003e \u003cstrong\u003e9\u003c/strong\u003e, 2179 (2018).\u003c/li\u003e\n\u003cli\u003eDong, J. \u003cem\u003eet al.\u003c/em\u003e Carrier management through electrode and electron-selective layer engineering for 10.70% efficiency antimony selenosulfide solar cells. \u003cem\u003eNat\u003c/em\u003e\u003cem\u003e.\u003c/em\u003e\u003cem\u003e Energy\u003c/em\u003e \u003cstrong\u003e10\u003c/strong\u003e, 857-868 (2025).\u003c/li\u003e\n\u003cli\u003eTang, R. \u003cem\u003eet al.\u003c/em\u003e Heterojunction Annealing Enabling Record Open‐Circuit Voltage in Antimony Triselenide Solar Cells. \u003cem\u003eAdv\u003c/em\u003e\u003cem\u003e.\u003c/em\u003e\u003cem\u003e Mater\u003c/em\u003e\u003cem\u003e.\u003c/em\u003e \u003cstrong\u003e34\u003c/strong\u003e, 2109078 (2022). \u003c/li\u003e\n\u003cli\u003eWang, L. \u003cem\u003eet al.\u003c/em\u003e Stable 6%-efficient Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells with a ZnO buffer layer. \u003cem\u003eNat\u003c/em\u003e\u003cem\u003e.\u003c/em\u003e\u003cem\u003e Energy\u003c/em\u003e \u003cstrong\u003e2\u003c/strong\u003e, 17046 (2017).\u003c/li\u003e\n\u003cli\u003eLi, Z. \u003cem\u003eet al.\u003c/em\u003e 9.2%-efficient core-shell structured antimony selenide nanorod array solar cells. \u003cem\u003eNat. 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All the chemicals and reagents were used without any treatment.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eSb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells Fabrication\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003ePreparation of CdS layer: All devices were deposited on cleaned FTO-coated glass. Prior to deposition, the FTO-coated glasses were ultrasonically cleaned sequentially with detergent, deionized (DI) water, isopropanol, acetone, and ethanol for 40\u0026nbsp;min, followed by cleaning with ultraviolet ozone for 20 minutes. Then, we deposited SnO\u003csub\u003e2\u003c/sub\u003e nanoparticles onto the FTO substrate by spin-coating SnO\u003csub\u003e2\u003c/sub\u003e colloidal water dispersion (15%) diluted with a water volume ratio of 1:5. The spin-coating was performed at 3000 r.p.m. for 30 seconds, then following a further annealing at 250 \u0026deg;C in air for 30 minutes to prepare the SnO\u003csub\u003e2\u003c/sub\u003e interfacial layer. Afterwards, the electron transport layer, CdS, was deposited on FTO/SnO\u003csub\u003e2\u003c/sub\u003e-coated glass via the CBD method. The preparation of CdS is as follows: First, 17.3 mL of ammonium hydroxide (NH\u003csub\u003e3\u003c/sub\u003e\u0026middot;H\u003csub\u003e2\u003c/sub\u003eO, 25\u0026ndash;28%) was added into 13.3 mL of 15 mmol L\u003csup\u003e-1\u003c/sup\u003e Cd(NO\u003csub\u003e3\u003c/sub\u003e)\u003csub\u003e2\u003c/sub\u003e solution with continuously stirring in the air for 1.5 min. Next, 8.8 mL of 1.5 mol L\u003csup\u003e-1\u003c/sup\u003e thiourea solution was added into the above solution. After stirring for 30 s, 93.3 mL of deionized water was added, with continuing stirring for 30 s. The FTO-coated glass was immersed into the mixed solution, and reacted at 65 \u0026deg;C for 16 min. After that, took out the substrates and rinsed them with deionized water and ethanol, and dried them with nitrogen gas. Then\u0026nbsp;30 mg mL\u003csup\u003e-1\u0026nbsp;\u003c/sup\u003eof\u0026nbsp;SbCl\u003csub\u003e3\u0026nbsp;\u003c/sub\u003emethanol solution was spin-coated to treat the CdS surface with 3000 r.m.p. for 30 s. The CdS-coated substrates were then heat-treated at a high temperature of 400 \u0026deg;C for 10 min to obtain crystallized CdS in an air atmosphere.\u0026nbsp;After natural cooling, the FTO/SnO\u003csub\u003e2\u003c/sub\u003e/CdS film was transferred into thermal evaporation equipment for the deposition of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e absorber layer.\u003c/p\u003e\n\u003cp\u003ePreparation of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e absorber layer:\u0026nbsp;The\u0026nbsp;Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e alloy precursor was synthesized through ball mill and solid phase reaction approach.\u0026nbsp;Due to selenium\u0026rsquo;s high vapor pressure, some Se inevitably evaporates during substrate heating and post‑annealing of Sb₂Se₃ films. To investigate this, we ball‑milled a series of Sb₂Se₃ precursor powders with systematically varied Se: Sb ratios. After film deposition and annealing, all samples exhibited Se loss, and the extent of depletion increased with higher initial Se content, yet we were still able to produce Sb₂Se₃ thin films covering a broad range of Se: Sb ratios. Firstly, Sb and Se powders were placed in a ball mill tank according to the different Se/Sb ratio from 1.45-1.65 of\u0026nbsp;Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e and ball-milled for 8 hours. The precursor was dried in vacuum environment for 30 min to obtain\u0026nbsp;Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e precursor powder, then the precursor powder was pressed into a tray and then annealed at 300 \u0026deg;C for 30 min in an N\u003csub\u003e2\u003c/sub\u003e glove box to obtain\u0026nbsp;Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e alloy.\u0026nbsp;Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e films were deposited using thermal evaporation deposition under a pressure of 1x10\u003csup\u003e-4\u003c/sup\u003e Pa.\u0026nbsp;For the Control-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e,\u0026nbsp;before the deposition, 0.2 g of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e powder\u0026nbsp;with Se/Sb ratio of 1.5\u0026nbsp;was placed in the middle of a tungsten boat, with an evaporation rate of 5-8 nm per second, and the substrate temperature was set at 315 \u0026deg;C. The thickness of the deposited film was controlled\u0026nbsp;about\u0026nbsp;400 nm. Then, the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e film was annealed in a glove box with N\u003csub\u003e2\u003c/sub\u003e atmosphere at 370 \u0026deg;C for 8 minutes.\u0026nbsp;For the Target-Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e homojunction, we deposited n-type and p-type\u0026nbsp;Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e films sequentially using thermal evaporation. For the n-type film, the preparation method is the same as that in Control-\u0026nbsp;Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e, except that the thickness is about 50 nm. Next, we deposited about 350 nm of\u0026nbsp;Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e using the precursor with a Se/Sb ratio of 1.6, and then annealed the film in a glove filled with nitrogen\u0026nbsp;at 370 \u0026deg;C for 8 minutes.\u003c/p\u003e\n\u003cp\u003eFor the fabrication of Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells: After preparing the Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e film, the spiro-OMeTAD HTL was spin-coated onto Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e film at 3000 r.p.m. for 30 seconds. The spiro-OMeTAD solution was prepared by dissolving 36.6 mg of spiro-OMeTAD in 1 mL of chlorobenzene, along with 14.5 \u0026mu;L of 4-tert-butylpyridine and 9.5 \u0026mu;L of Li-TFSI solution (520 mg mL\u003csup\u003e-1\u003c/sup\u003eLi\u003csup\u003e-\u003c/sup\u003e\u003csup\u003e1\u003c/sup\u003eTFSI in acetonitrile). After spin coating, the film was annealed at 100 \u0026deg;C in air for 10 minutes. Finally, an Au electrode was thermally evaporated onto the film.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eFilm and Device Characterizations:\u003c/strong\u003e\u003cstrong\u003e\u0026nbsp;\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe crystal structure of\u0026nbsp;Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e thin films was analyzed on a Bruker Advance D8 diffractometer using X-ray diffraction of Cu K𝛼\u0026nbsp;radiation (𝜆\u0026nbsp;= 1.5416 \u0026Aring;). The morphology and thickness of the films were characterized by scanning electron microscopy (SU8220, Hitachi High-Tech), and the element analysis was conducted by Zeiss G500 SEM equipped with an EDS (Bruker) module. The work function and valence band binding energy were measured by UPS, and the excitation source was a He light source (h\u003cem\u003ev\u003c/em\u003e = 21.22 eV) with a beam spot of 2 mm.\u0026nbsp;The Hall effect measurement was exerted by Ecopia Hall Effect Tester (HMS-7000) at 300 K.\u0026nbsp;EQE spectra were measured using a single source lighting system (halogen lamp) and a monochromator (model SPIEQ200). The surface atom chemical state was characterized using X-ray photoelectron spectroscopy (Thermo Fisher ESCALAB 250Xi) which possesses Al k𝛼\u0026nbsp;150 W with beam spot 500 \u0026mu;sm. The change in mass of the sample with temperature or time was performed by Thermogravimeter (TGA Q5000IR). The \u003cem\u003eJ-V\u003c/em\u003e and VOC-light response of devices was measured by Keithley 2400 apparatus with a standard xenon lamp-based solar simulator (Oriel Sol 3A) under solar-simulated AM 1.5 sunlight (100 mW cm\u003csup\u003e-2\u003c/sup\u003e). The calibration for the light intensity was conducted in advance by reference cell (Oriel P/N 91 150 V, with KG-5 visible color filter), and the monocrystalline silicon was also calibrated by the National Renewable Energy Laboratory (NREL). Secondary ion mass spectrometry (CAMECA IMS 7f-Auto) was adopted to analyze elemental distribution of the films.\u0026nbsp;AFM topography images and KPFM measurements were acquired on a Bruker Dimension Icon instrument. UPS was performed on a Thermo Escalab Xi+ analyzer and HeI lamp (21.22 eV).. Capacitance-voltage measurement and the point defects in the devices were investigated using the Phystech FT-1230 HERA DLTS system to record Deep-Level Transient Spectrum (DLTS). The reverse bias, pulse voltage, period width, and pulse width were set to -0.4, 0.2-0.6 V, 100 ms, and 10 ms, respectively.\u003cstrong\u003e\u0026nbsp;\u003c/strong\u003eThe CV and DLCP were characterized using a Keysight B1500A. For the CV measurements, we employed a frequency of 10 kHz, an AC bias of 30 mV, and a scanning voltage range of -1 to 1 V. For DLCP, the measurement was conducted at 10 kHz with a DC bias range from -0.2 to 0.2 V, while the sum of VDC + VAC was fixed at 30 mV. Carrier transport kinetics was studied via transient absorption spectroscopy (TAS), which was performed on a pump-probe system (Helios, Ultrafast System LLC).\u0026nbsp;\u003c/p\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":true,"hideJournal":false,"highlight":"","institution":"","isAcceptedByJournal":true,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"nature-portfolio","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"","title":"Nature Portfolio","twitterHandle":"","acdcEnabled":false,"dfaEnabled":false,"editorialSystem":"ejp","reportingPortfolio":"","inReviewEnabled":true,"inReviewRevisionsEnabled":false},"keywords":"","lastPublishedDoi":"10.21203/rs.3.rs-7257351/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-7257351/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eAntimony selenide (Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e) has emerged as a promising thin-film photovoltaic absorber due to its ideal bandgap (1.1-1.3 eV), high absorption coefficient (\u0026gt;10⁵ cm\u003csup\u003e-\u003c/sup\u003e¹), and environmentally benign composition. However, Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells (SSCs) often suffer from significant open-circuit voltage (\u003cem\u003eV\u003c/em\u003e\u003csub\u003eOC\u003c/sub\u003e) losses, attributed to weak built-in fields and severe non-radiative recombination at interfaces and the absorber layer. Here, we demonstrate a composition‑driven carrier polarity control strategy to form an n-type/p-type Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e homojunction. By precisely tuning the chemical potentials of Se and Sb, we reversibly modulate the conductivity type, achieving carrier densities exceeding 10\u003csup\u003e14\u003c/sup\u003e cm\u003csup\u003e-\u003c/sup\u003e³ for the n- and p-type states, with Hall coefficients ranging from −3.14×10\u003csup\u003e-\u003c/sup\u003e² m³ C\u003csup\u003e-\u003c/sup\u003e¹ (n-type) and +9.51×10\u003csup\u003e-\u003c/sup\u003e² m³ C\u003csup\u003e-\u003c/sup\u003e¹ (p-type). More importantly, we incorporate a homojunction structure into the planar SSC, which simultaneously enhances the built-in electric field and passivates deep-level defects. These synergistic effects promote carrier separation, reduce non-radiative recombination, and accelerate carrier extraction. As a result, the study demonstrates a record power conversion efficiency of 10.15% for thermally evaporated Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e devices, along with the lowest open-circuit voltage deficit (0.459 V) among all reported SSCs. This work not only establishes a new efficiency benchmark for Sb\u003csub\u003e2\u003c/sub\u003eSe\u003csub\u003e3\u003c/sub\u003e solar cells but also offers a universal approach for defect management and junction design in emerging chalcogenide photovoltaics.\u003c/p\u003e","manuscriptTitle":"Homojunction Sb2Se3 Solar Cell","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2025-08-22 04:27:27","doi":"10.21203/rs.3.rs-7257351/v1","editorialEvents":[],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"nature-photonics","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"nphoton","sideBox":"Learn more about [Nature Photonics](https://www.nature.com/nphoton/)","snPcode":"41566","submissionUrl":"https://mts-nphot.nature.com/cgi-bin/main.plex","title":"Nature Photonics","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"ejp","reportingPortfolio":"Nature Research","inReviewEnabled":true,"inReviewRevisionsEnabled":false}}],"origin":"","ownerIdentity":"a9e72e12-7216-4f57-a31d-5a8159894f1d","owner":[],"postedDate":"August 22nd, 2025","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"published-in-journal","subjectAreas":[{"id":53541374,"name":"Physical sciences/Materials science/Materials for energy and catalysis/Solar cells"},{"id":53541375,"name":"Physical sciences/Materials science/Materials for devices"}],"tags":[],"updatedAt":"2026-04-14T07:11:57+00:00","versionOfRecord":{"articleIdentity":"rs-7257351","link":"https://doi.org/10.1038/s41566-026-01888-1","journal":{"identity":"nature-photonics","isVorOnly":false,"title":"Nature Photonics"},"publishedOn":"2026-04-13 04:00:00","publishedOnDateReadable":"April 13th, 2026"},"versionCreatedAt":"2025-08-22 04:27:27","video":"","vorDoi":"10.1038/s41566-026-01888-1","vorDoiUrl":"https://doi.org/10.1038/s41566-026-01888-1","workflowStages":[]},"version":"v1","identity":"rs-7257351","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-7257351","identity":"rs-7257351","version":["v1"]},"buildId":"XKTyCvWXoU3ODBz1xrDgd","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}