Field-effect passivation for minimized voltage loss in highly efficient antimony selenosulfide solar cells

Domain: physics
Published: 2026-01-06
Editorial outcome: minor_revision

Abstract
The inferior heterojunction quality and misaligned energy levels at the buffer/absorber interface cause severe interface recombination and large open-circuit voltage (V ) loss, limiting the power conversion efficiency (PCE) of antimonyselenosulfide (Sb (S,Se) ) solar cells. Here, we develop a field-effect passivation strategy by introducing a low-work-function Ta O dielectric layer between the CdS and Sb (S,Se) layers. This Ta O layer serves as an optimal substrate for growing highly crystalline Sb (S,Se) films while also enhancing interfacial charge transport. The positive fixed charges in Ta O strengthen the built-in electric field and promote electrons extraction while suppressing holes accumulation at the interface, thereby substantially suppressing non-radiative recombination probabilities. Implementing this passivation strategy yields a record PCE of 10.95% (10.65% certified) an V of 695 mV, corresponding to a remarkably low voltage deficit. This work establishes a universal physical passivation paradigm for interface quality optimization and V loss mitigation in high-performance Sb (S,Se) photovoltaics.

Introduction
In response to the depletion of traditional energy sources, there has been a growing demand for clean and renewable energy, driving the development of new thin-film solar cells. Antimony selenosulfide (Sb (S,Se) ) has recently demonstrated significant potential in emerging thin-film photovoltaics due to its high stability, low toxicity, tunable bandgap and excellent optoelectronic properties1–4. Based on these advantages, Sb (S,Se) solar cells have undergone a rapid development, achieving a reported power conversion efficiency (PCE) of 10.92%5. However, the current efficiency of these thin-film solar cells is limited by significant open-circuit voltage (V ) losses caused by severe non-radiative recombination at the absorber layer and interfaces6,7. Currently, high-efficiency Sb (S,Se) solar cells are often fabricated through direct deposition on cadmium sulfide (CdS) substrates. An unfavorable energy level arrangement at the CdS/Sb (S,Se) interface, such as a cliff-like band structure, can cause severe carrier recombination at the interface. Therefore, developing an efficient passivation strategy to improve the quality of the CdS/Sb (S,Se) heterojunction and mitigate the V losses of devices is crucial.

Researchers have conducted extensive studies on interface engineering to improve the quality of heterojunctions and suppress interface carrier recombination. To transform the band structure at the CdS/Sb (S,Se) interface from a cliff-like to spike-like structure, Liu et al. employed a heterojunction lithiation strategy and achieved a high PCE of 10.76% for the Sb (S,Se) solar cells8. As for optimizing the CdS buffer layer, a hydrazine hydrate solution was used to remove surface oxides and improve the surface morphology of CdS films, resulting in a PCE of 10.30%9. Besides, Chen group investigated the effects of potassium chloride (KCl) and ethylenediamine (EDA) post-treatments on CdS films, which yielded PCEs of 9.98% and 10.10%, respectively10,11. They found that KCl treatment effectively reduced the surface roughness of CdS films and improved their conductivity, while EDA treatment promoted the grain growth and crystallinity of CdS films and suppressed the formation of interface defects. Although these passivation approaches reduced the V losses of the device to some extent, such chemically-based passivation methods suffer from byproduct residues, poor long-term stability and potentially harmful chemical reactions at interfaces. This hinders the reproducibility of device performance and large-scale applications12.

In addition to chemical passivation technology, a physical passivation strategy, such as field-effect passivation by introducing a dielectric layer, is a promising approach. In crystalline silicon (Si) solar cells, the zinc oxide (ZnO), titanium oxide (TiO), strontium oxide (SrO ) thin films have been deposited on the Si substrate to provide moderate surface passivation and form the additional interfacial fields, which facilitate the charge-selective transport and reduce the probability of interface recombination13–15. In perovskite solar cells, a molybdenum oxide (MoO ) thin film with a high work function is inserted between the perovskite and hole transport layer. This field-effect passivation of MoOx effectively enhances the built-in field and improves interface charge transport thanks to the formation of interfacial polarization16. Furthermore, fixed charge passivation is used to modify carrier concentration distribution near the heterojunctions and selectively extract charge carriers17,18. Despite these advantages, the application of the physical passivation strategy to inorganic thin-film solar cells remains underexplored, and the impact of field-effect passivation on the charge transport and recombination dynamics in Sb (S,Se) solar cells has rarely been reported.

Herein, we fabricated an ultrathin, transparent Ta O dielectric layer between CdS and Sb (S,Se) layers to investigate the effect of field-effect passivation on the V loss in the device. Ta O is a wide-bandgap transition metal oxide that exhibits exceptional optical transparency and chemical stability19,20. Its material properties remain unaffected during subsequent high-temperature hydrothermal and annealing processes. There are positive fixed charges generated in Ta O films, which can attract electrons while concurrently repelling holes away from the interface21. Thanks to the reduced V losses resulting from this field-effect passivation, the champion device achieves a certified PCE of 10.65% with V of 692 mV, corresponding to 54.96% of the theoretical limit (V /V ), which is comparable to the highest reported value (55.14%) for Sb (S,Se) solar cells5. This type of inorganic thin-film solar cells also exhibits good operational stability.

Results
Characterization of CdS/Ta O buffer layers
In this study, a highly transparent and ultrathin Ta O dielectric layer was fabricated via the sol-gel spin-coating technique. Fig. 1a depicts the specific reaction pathway and conditions for TaO: Ta⁵⁺ ions in the solution form ethanol-tantalum complexes preferentially, followed by the generation of the desired Ta O precipitate under the influence of high temperature and hydroxyl groups. The inset shows the transparent, homogeneous Ta O precursor solution. To demonstrate the formation of Ta O, X-ray diffraction (XRD) was conducted on the thin-film samples to determine their composition and crystal structure. As shown in the XRD pattern, all diffraction peaks of the sample closely match the orthorhombic phase of Ta O (JCPDS #25-0922), with a pronounced (001) grain orientation. The Ta O dielectric layer has a high dielectric constant of ~2822, a wide bandgap of 4.25 eV (Supplementary Fig. 1), and an ultra low electrical conductivity of ~2.11×10−5 S/cm based on Hall effect measurements. Scanning electron microscopy (SEM) analysis of CdS films without and with Ta O modification revealed distinct morphological characteristics. As shown in Fig. 1b, e, the pure CdS films exhibited a rougher surface morphology with numerous pinholes. Such low-quality seed layers can degrade the quality of subsequent Sb (S,Se) films and the resultant device performance. Remarkably, the deposition of Ta O effectively filled surface pinholes and cracks in the CdS films, resulting in a smoother and more uniform surface topography. To verify the uniformity of the ultrathin Ta O layer, we performed EDS mapping analysis on the CdS/Ta O substrate. As shown in Supplementary Fig. 2, the Ta signal (representing Ta O) exhibits a continuous and homogeneous distribution on the surface of CdS films, with no observable island-like aggregation or large-scale uncovered areas. This indicates a uniform and complete coverage of the CdS film by the deposited Ta O layer. Simultaneously, the signals of Cd and S also display a homogeneous distribution, further confirming that the TaO layer is sufficiently ultra-thin, thereby allowing the underlying CdS signals to remain clearly detectable. The CdS and CdS/TaO films exhibited similar transmittance across the entire visible spectrum (Supplementary Fig. 3). We also used atomic force microscopy (AFM) to analyse the surface morphology of CdS films (Fig. 1c, f). Following the deposition of Ta O, the roughness (R ) of the CdS films was decreased from 27.8 nm to 23.1 nm, consistent with the SEM results. We also conducted the contact angle tests on CdS films (Fig. 1d). The contact angle was reduced from 61.9° to 44.0° after the Ta O layer was introduced, positively impacting the heterogeneous nucleation and grain growth of the Sb (S,Se) films. To confirm that the TaO layer facilitates carrier transport, we measured the conductivity of CdS films, and the results are shown in Fig. 1g (the inset shows the structure of device). After depositing the TaO dielectric layer, the conductivity of the CdS films was increased from 3.74×10−5 S/cm to 6.22×10−5 S/cm.

Morphological analysis of Sb (S,Se) layer
According to previous studies, the polarity of CdS surfaces has a significant influence on the growth orientation of the absorber layer1. As shown in Supplementary Fig. 4, the Sb (S,Se) films tend to form a favorable (hk1) grain orientation on the non-polar (100) plane of CdS films. In contrast, exposure of the polar (101) and (10 (cid:1) 1) planes lead to complex anisotropy in the grain orientation of Sb (S,Se) , negatively affecting the carrier transport. To demonstrate that the Ta O layer reduces the exposure of polar planes in CdS, we performed XRD measurements on CdS films. As shown in the Supplementary Fig. 5, compared to the pure CdS film, the CdS/Ta O sample exhibits significantly reduced peak intensities of polar planes, such as (101), and enhanced peak intensities of nonpolar planes, such as (100)23. To investigate the effect of Ta O dielectric layer on the quality of Sb (S,Se) films, we implemented XRD and SEM measurements. The diffraction patterns of Sb (S,Se) films matched well with the orthorhombic Sb S (JCPDS#42-1393) and Sb Se (JCPDS#15-0861) phases (Supplementary Fig. 6a). We calculated the texture coefficients (TC) of the crystal planes for Sb (S,Se) films deposited on CdS and CdS/Ta O substrates based on the XRD data, with detailed results presented in Supplementary Fig. 6b24. The Sb (S,Se) films deposited on the CdS/Ta O substrate exhibit significantly higher TC values for (hk1) planes, such as (211), (221), (301) and (421) compared to those of the control Sb (S,Se) films. This provides evidence for the optimized grain orientation of Sb (S,Se) films induced by the Ta O layer. As shown in Supplementary Fig. 7, the local XRD analysis revealed that Sb (S,Se) films deposited on Ta O substrates exhibited stronger (hk1) diffraction peaks and narrower full width at half maximum than the control Sb (S,Se) films. Regarding the slight shift in diffraction peaks, we attribute it to the changes in the chemical composition of the absorber layer (Supplementary Fig. 8 and Supplementary Table 1). The introduction of Ta O substrate modifies the crystallization kinetics, leading to an altered S/Se ratio in Sb (S,Se) 25. SEM images reflected that both the control and Ta O -modified Sb (S,Se) films had a compact morphology. However, the control films displayed a higher density of large granular precipitates, and the phenomenon was effectively mitigated in the Ta O -modified Sb (S,Se) films. Supplementary Fig. 9 showcased the AFM images of the Sb (S,Se) films, indicating that the R was reduced for films deposited on Ta O substrate (39.6 nm) compared to that of the control films (50.3 nm). We attributed this to the suppression of large granular precipitates in the absorber layer. The reduced surface roughness facilitates uniform interfacial contact and minimizes localized recombination, which is critical for enhancing charge transport and device performance. These results suggest that the insertion of Ta O layer improves the crystallinity, grain orientation and morphology of Sb (S,Se) films. Therefore, the Sb (S,Se) films deposited on Ta O substrate exhibit an enhanced light absorption due to the improved film quality (Supplementary Fig. 10). Since interfacial stability is fundamental to device reproducibility and long-term reliability, we performed HAADF-STEM imaging and corresponding EDS mapping analysis for a cross-sectional sample with structure of FTO/CdS/Ta O /Sb (S,Se) . As shown in Supplementary Fig. 11, the HAADF-STEM image reveals sharp and distinct interfaces between all functional layers. Notably, a continuous thin layer with a thickness of ~3 nm, corresponding to Ta O , is clearly observed between the CdS and Sb (S,Se) layers. EDS mapping further demonstrates a sharp and confined Ta signal precisely at the CdS/Sb (S,Se) interface, while other elements such as Cd, S, Sb, and Se exhibit negligible interdiffusion. These results confirm that the Ta O dielectric layer maintains robust chemical and structural stability during high-temperature hydrothermal and annealing processes, forming well-defined heterointerfaces with adjacent layers.

Fixed charge passivation
In crystalline silicon solar cells, TaO has been widely recognized as an efficient passivation layer, which is capable of introducing fixed charges at the silicon surface. Theoretically, when Ta O contacts with Sb (S,Se) layer, it may exhibit a similar fixed charge effect. To validate our analyses, we measured the fixed charges by preparing the metal oxide semiconductor (MOS) structured devices for C–V measurements (Fig. 2a). By differentiating the C–V data, the voltage corresponding to the extreme value represents the flat-band voltage. As shown in Fig. 2b, the flat-band voltages of the CdS and CdS/Ta O samples are −0.35 V and −0.50 V, respectively. We therefore confirm that the Ta O enables positive fixed charges at the interface, consistent with reported literatures26,27. The fixed charge (Q) density was calculated using the following formula28:

Q f= C in (V MS − V fb ) / q (1)

Here, C denotes the intrinsic capacitance of the MOS structure, V represents the metal-semiconductor potential difference determined by the work function difference between FTO (4.8 eV) and Sb (S,Se) , V is the flat-band voltage, A is the electrode area, and q is the elementary charge. The calculated fixed charge density Q of the Ta O layer is +2.24×1011 cm−2 (Supplementary Table 2). When positive fixed charges of TaO appear at the interface, more electrons of Sb (S,Se) gather therein, resulting in a downward band bending of Sb (S,Se) (Fig. 2c). These positive fixed charges enable field-effect passivation at the TaO/Sb (S,Se) interface, facilitating efficient electron extraction and suppressing hole accumulation17,18. To theoretically verify the field-effect passivation mechanism of the Ta O layer, we performed the device simulations by using the SCAPS-1D software29,30. A device model is established that accurately corresponds to the actual structure of solar cells. All the essential materials parameters are shown in Supplementary Table 3. The simulation results (Supplementary Fig. 12) demonstrate that the insertion of the Ta O layer significantly reduces the conduction band offset (ΔE ) between CdS and Sb (S,Se) from 280 to 50 meV. The TaO layer generates a favorable path for the electron transport to the electron transport layer.

In addition to the fixed positive charges, Fermi level alignment between Ta O/Sb (S,Se) also results in band bending with the same trend. We performed the ultraviolet photoelectron spectroscopy (UPS) for the Ta O and Sb (S,Se) films, respectively. Fig. 2d, e present that the secondary electron cut-off edges are 17.79 and 17.20 eV, and the distances between Fermi energy level and valence band maximum are 3.38 and 0.97 eV for the Ta O and Sb (S,Se) , respectively. The corresponding work functions are 3.43 eV and 4.02 eV, similar to the values reported in the literature11,18. As shown in Fig. 2c, the band bending of Sb (S,Se) is downward while the band bending of Ta O is upward when the Fermi-level alignment is achieved between Ta O/Sb (S,Se) . Kelvin probe force microscopy (KPFM) was conducted to probe the contact potential difference (CPD) across the Ta O/Sb (S,Se) heterojunction (Fig. 2f). A significant reduction in CPD was observed as the probe traversed from the Ta O region to the Sb (S,Se) region (Fig. 2g, h). The Ta O region exhibited an average CPD of approximately 1010 mV, while the Sb (S,Se) region exhibited a decreased CPD of approximately 593 mV. This difference in CPD is consistent with the UPS analysis. With the insertion of Ta O layer, the presence of fixed positive charges and the significant work function difference induce more pronounced downward band bending of Sb (S,Se) .

Carrier dynamics and interfacial defect study
To further verify that the Ta O dielectric layer facilitated carrier transport, we used transient absorption spectroscopy (TAS) to analyze the carrier transport dynamics. Fig. 3a, b showcased the two-dimensional pseudo-color maps of the TAS spectra for Sb (S,Se) films. Broad photo-induced absorption (PIA, ΔA>0) in the range of 525–700 nm appeared in both samples due to the formation of S•− radical and the trapping of holes within Sb (S,Se) films31,32. Compared to the control samples, the Sb (S,Se) films deposited on TaO dielectric layer exhibited lower PIA signals. Supplementary Fig. 13 displayed the time-resolved absorption spectra at different times, in which the TaO incorporated Sb (S,Se) films showed a weaker peak intensity than the control films at the same delay time. In general, the TAS intensity reflects the number of carriers in the conduction and valence bands of the absorber, so the decreased intensity can be attributed to the effective charge transport via the Ta O dielectric layer33. Fig. 3c showed the decay kinetics of the photobleaching peak intensity at 622 nm for the samples, and the carrier decay lifetimes were obtained by fitting the double-exponential equations in Supplementary Note 1. The corresponding results were listed in Supplementary Table 4. The average carrier decay lifetimes of the Sb (S,Se) films on the Ta O layer are much lower than those of the control films, suggesting that Ta O layer enables the fast electron extraction and transportation to cathode. To explore the carrier recombination behavior, we performed the transient photovoltage/photocurrent (TPV/TPC) measurements on the devices without and with TaO dielectric layers. As shown in Fig. 3d and Supplementary Table 5, the TPC results indicated a reduced photocurrent decay lifetime of 252 μs in the Ta O incorporated devices compared to that of the control devices (328 μs), consistent with the analysis of TAS measurements. Fig. 3e presented the TPV decay curves of the devices, in which the Ta O incorporated devices exhibited a longer carrier recombination lifetime of 3238 μs than that of the control devices (953 μs) (Supplementary Table 5). We attributed this to the suppressive carrier recombination rate by the field-effect passivation of TaO. Additionally, the capacitance-voltage (C-V) measurements showed a decreased capacitance in the TaO incorporated devices (Fig. 3f). This suggests that the Ta O dielectric layer can effectively suppress the aggregation of carriers at the CdS/Sb (S,Se) interface and reduce the probability of carrier recombination34, which is in agreement with the results of the TPV measurements.

We performed C–V and DLCP measurements on the devices without and with TaO layer to explore the interface defects in the CdS/Sb(S,Se) heterojunction. The defect density from C–V measurements (N ) includes the free carriers, bulk defects, and interfacial defects, and the defect density from DLCP measurements (N ) contains the free carriers and bulk defects4,35. Therefore, the interfacial defects can be determined by the difference between the N and N . The defect densities of the devices were calculated according to the equation in Supplementary Note 2, and the corresponding results were shown in Supplementary Table 6. Compared to the control devices, the TaO incorporated devices exhibited a lower interfacial defect density and a wider depletion region (Fig. 3g, h). This was due to the introduction of TaO improved the morphology and hydrophilicity of CdS films, providing a favorable substrate for the deposition of high-quality Sb(S,Se) films. The defect density and electron mobility are further investigated by the J–V curves according to the space-charge limited current model. Figure 3i displayed the dark J–V curves of the electron-only devices without and with TaO, and the defect densities were calculated according to the formulas in Supplementary Note 3. The insertion of TaO reduced the electron defect densities of the devices from 6.91×1015 to 5.05×1015 cm−3. Moreover, the TaO incorporated devices exhibit a higher electron mobility compared to that of the control devices (Supplementary Fig. 14), demonstrating that the Ta O dielectric layer effectively enhances electron extraction and collection. We conclude that the TaO dielectric layer not only improves the charge transport dynamics at the interface but also achieves a high-quality CdS/Sb(S,Se) heterojunction with fewer interface defects.

Device performance and characterization
Encouraged by the field-effect passivation of Ta O, we fabricated the solar cells with the device structure of FTO/CdS/Ta O/Sb (S,Se) /Spiro-OMeTAD/Au. The influence of Ta O layers with different thickness on device performance was analyzed, and the corresponding current density-voltage (J–V) curves and device parameters were shown in Supplementary Fig. 15 and Supplementary Table 7. As expected, the introduction of TaO dielectric layers with varying thickness significantly improved photovoltaic performance. Fig. 4a showcased the J–V curves of the control and champion TaO incorporated devices measured under AM 1.5G illumination (100 mW cm−2). The control devices achieved a PCE of 9.21% with a V of 0.647 V, a short circuit current density (J ) of 22.53 mA cm−2 and a fill factor (FF) of 63.19%. Specifically, the devices with optimal thickness of TaO reached a PCE of 10.95% with a V of 0.695 V, a J of 23.22 mA cm−2 and a FF of 67.87%, which is the highest efficiency for the Sb-based thin-film solar cells so far. Besides, the J–V curves simulated by SCAPS-1D show significant improvements in V and FF (Supplementary Fig. 16 and Supplementary Table 8), which are fully consistent with our experimental results. Moreover, the Ta O incorporated devices exhibited a good reproducibility (Supplementary Fig. 17). Fig. 4b showed the external quantum efficiency (EQE) of the control and Ta O incorporated devices, and the integrated current densities are in good agreement with J obtained from the J–V curves. The response in the spectral range of 330–500 nm is enhanced due to the optimization of the CdS buffer layer and the heterojunction interface, while the response in the spectral range of 500–750 nm is increased because of the improved quality of Sb (S,Se) absorber layer36. We further calculated the Urbach energy of the devices based on the EQE spectra. As shown in Supplementary Fig. 18, the Ta O incorporated devices displayed a reduced Urbach energy of 28.43 meV compared to 31.47 meV for the control devices. The lower Urbach energy indicates that the Sb (S,Se) films deposited on Ta O layers have a higher crystallinity and less bulk defects37,38.

To investigate the impact of the Ta O dielectric layer on the FF of devices, we decomposed the FF losses into non-radiative recombination losses and transport losses through Supplementary Note 4 (Fig. 4c and Supplementary Table 9). The theoretical maximum fill factor (FF ) values for the control and Ta O incorporated devices were 74.2% and 76.64%, respectively. Introducing the TaO dielectric layer reduced FF losses caused by non-radiative recombination from 19.90% to 13.16%. The optical bandgap calculated from EQE measurements show no variation between the control and Ta O incorporated devices (Supplementary Fig. 19). This indicates that the improvement of V is due to the field-effect passivation of Ta O and the increased built-in electric field of devices. As shown in the capacitance–voltage (C−2–V) curves, the extracted built-in voltage (V ) of the control devices is 649 mV, which is increased to 703 mV for the Ta O incorporated devices (Fig. 4d). Furthermore, the optimized devices were submitted to an authoritative, independent laboratory (South China National Center of Metrology) for performance certification (certification report provided in Supplementary Fig. 20). Consequently, the TaO incorporated devices demonstrated a certified PCE of 10.65% (10.57%) with a V of 0.692 (0.691) V, a FF of 68.70% (68.17%) and a J of 22.42 (22.45) mA cm−2 measured in reverse (forward) scan. It is worth noting that the certified PCE of 10.65% is second only to the certified record efficiency of 10.70% (Fig. 4e and Table 1). As shown in Fig. 4f (calculation method detailed in Supplementary Note 5), we calculated the product of V ×FF according to the Shockley-Queisser limit values across different bandgaps, and compared the V ×FF values of the reported certified devices. The product of V ×FF in this work reached 47.54%, which was the highest value among all the reported certified devices (Table 1). Concurrently, the V /V ratio of the certified device attained 54.96%, indicating a substantially reduced V deficit (V ) (Supplementary Fig. 21 and Table 1). These results fully demonstrated the effectiveness of the field-effect passivation strategy in enhancing device performance.

To further demonstrate that carrier recombination was suppressed, we implemented the dark J–V characterization, light-intensity-dependent J–V measurements, and electrical impedance spectroscopy (EIS) under both dark and different light intensity conditions. It can be seen in Supplementary Fig. 22a that the reverse leakage currents were significantly reduced upon the incorporation of Ta O layer. By processing and fitting the dark J–V curves based on the equations in Supplementary Note 6, we can obtain relevant device parameters, including shunt conductance (G), series resistance (R ), diode ideality factor (A) and reverse saturation current density (J ). As shown in Supplementary Fig. 22b-d, the TaO incorporated devices exhibit a G value of 0.0005 mS cm−2, an R of 2.55 Ω cm2, a diode ideality factor of 1.78 and a J of 6.28×10−6 mA cm−2. All these parameters show improvement over those of the control device, which had a G value of 0.006 mS cm−2, an R of 6.43 Ω cm2, a diode ideality factor of 1.92 and a J of 2.48×10−5 mA cm−2. We then measured the V versus light-intensity curves that followed the relationship of V /n(k T/q)ln(I), where the k , T, and q represent Boltzmann’s constant, the temperature in Kelvin, and the elementary charge, respectively. In the case of ideal trap-free recombination, the slope of the curves corresponds to k T/q. Any deviation of the measured slope from k T/q indicates the extent of trap-assisted recombination under open-circuit conditions39. As shown in Fig. 4g, the TaO incorporated devices exhibit a lower slope of 1.219 k T/q than that of the control devices (1.323 k T/q), indicating that the Ta O layer suppresses trap-assisted recombination at the CdS/Sb (S,Se) interfaces through the effective field-effect passivation. Furthermore, the light-intensity dependence of J (Supplementary Fig. 23) also demonstrates that the introduction of the Ta O dielectric layer effectively suppresses carrier recombination in the devices. We performed EIS on the devices without and with Ta O dielectric layer under dark conditions (Supplementary Fig. 24). The starting point in the high-frequency region represents the series resistance R , and the semicircle arc in the low-frequency region denotes the recombination resistance R . The extracted R of solar cells decreased from 19.79 to 3.35 Ω and the R rec increased from 1358.85 to 4676.42 Ω with the introduction of Ta O dielectric layer. To investigate the interfacial charge transfer properties of the devices, we performed EIS measurements under different light intensities. As shown in Supplementary Fig. 25a, b, all the Nyquist plots exhibit a single semicircular arc, indicating that the interfacial process is dominated by a single time constant40. The equivalent circuit model was used to fit the impedance data (Supplementary Fig. 25c), where R represents the series resistance, R corresponds to the charge transfer resistance, and the constant phase element denotes the geometric capacitive behavior of the device41. The R value of both devices decreases linearly with the increasing light intensity (Supplementary Fig. 25d), which can be attributed to that the increased photogenerated carrier concentration effectively reduces the energy barrier of charge transfer. Notably, the R ct values of the Ta 2 O 5 incorporated devices were significantly lower than those of the control device under different light intensities. This demonstrates that the fixed charges in the Ta O dielectric layer effectively reduce the charge transfer barrier through the field-effect passivation effect, thereby promoting efficient extraction and transport of charge carriers42–44. Clearly, the Ta O incorporated devices exhibit enhanced electrical properties and suppressed carrier recombination in comparison to the control devices, accounting for the improvement of V and FF45. Finally, we performed the stability measurement of an unencapsulated target device that maintained ∼96% of its initial PCE after continuous tracking at maximum power point for 200 h, demonstrating good operational stability (Fig. 4h).

Methods
Materials
All chemicals in this work were used as received without further purification unless otherwise specified. Cadmium nitrate tetrahydrate (Cd(NO ) ·4H O, AR), ammonium hydroxide (NH OH, 25–28%), thiourea (CH N S, GR), Cadmium chloride hemipentahydrate (CdCl ·2.5H O, AR), antimony potassium tartrate (C H KOSb·0.5HO, CP), and sodium thiosulfate pentahydrate (Na S O ·5HO, AR) were purchased from Sinopharm Chemical Reagent Co., Ltd. Tantalum(V) chloride (TaCl, 99.99%) was obtained from Macklin. Selenourea (CH N Se, 98%) and acetonitrile (CH N, 99.8%) were sourced from Sigma-Aldrich Trading Co., Ltd. Lithium bis(trifluoromethanesulfonyl)imide (Li-TFSI, 99%) and Spiro-OMeTAD (99.9%) were procured from Yolight Solar Technology Co., Ltd. 4-tert-Butylpyridine (tBP, 96%) and ethanol (CH O, 99.8%) were acquired from Aladdin Biochemical Technology Co., Ltd.

Devices fabrications
Fabrication of Ta 2 O 5 dielectric layer: The fabrication process initiated by dissolving 10 mg tantalum chloride (TaCl) in 1.8 mL ethanol with 30 min ultrasonication. Subsequently, 1.12 mL ethanol and 0.08 mL ultrapure water were added to the solution, followed by another 30 min ultrasonication. The homogenized mixture was filtered through a 200 nm membrane filter to remove large aggregates, after which 1 mL ethanol was introduced and magnetically stirred for 1 h to obtain the TaO precursor solution. This precursor (50 μL) was then spin-coated onto CdS substrates at varying spin speeds (4000–6000 rpm) and thermally annealed at 400–600 °C for 5 min under ambient air conditions, ultimately yielding densely compacted Ta O films. The thicknesses of the Ta O films prepared by different spin speeds were determined by spectroscopic ellipsometry to be 1±0.5, 3±0.5 and 5±0.5 nm, respectively.

Fabrication of Sb 2 (S,Se) 3 solar cells: The cleaned FTO substrates were first treated with UV-ozone for 20 min, followed by CdS film deposition via chemical bath deposition at 66 °C for 16.5 min. The CdS precursor solution was prepared by thoroughly mixing 20 mL cadmium nitrate tetrahydrate (Cd(NO ) ·4H O), 26 mL ammonium hydroxide (NH OH, 25–28%), 14 mL thiourea (CH N S), and 140 mL ultrapure water. A cadmium chloride methanol solution (20 mg/mL) was then spin-coated onto the CdS film at 3000 rpm for 30 s and annealed in air at 400 °C for 10 min to complete the CdS buffer layer. Subsequently, Ta O dielectric layers with varying thicknesses were fabricated on the CdS surface. For the deposition of Sb (S,Se) absorber layer, antimony potassium tartrate (C H KO Sb·0.5H O), sodium thiosulfate pentahydrate (Na S O ·5HO), and selenourea (CH N Se) served as Sb, S, and Se sources. Specifically, 0.260 g antimony potassium tartrate was dissolved in 39 mL ultrapure water, followed by sequential addition of 0.774 g sodium thiosulfate and 19.2 mg selenourea to form the precursor solution. The solution was transferred into a Teflon-lined autoclave containing FTO/CdS or FTO/CdS/TaO substrates and hydrothermally deposited at 135 °C for 2 h. Post annealing was conducted at 400 °C for 10 min under a nitrogen atmosphere to enhance the crystallinity of Sb (S,Se) . Spiro-OMeTAD hole transport layers were fabricated by thoroughly mixing 36.6 mg Spiro-OMeTAD, 29 μL 4-tert-butylpyridine (tBP), 9.5 μL lithium bis(trifluoromethanesulfonyl)imide (Li-TFSI, 520 mg/mL in acetonitrile), and 1 mL chlorobenzene, followed by spin-coating at 2000 rpm for 30 s under ambient conditions and subsequent annealing at 100 °C for 10 min. Device fabrication was finalized through thermal evaporation of an approximately 80 nm Au back electrode.

Characterization
Scanning electron microscopy (SEM, TESCAN MIRA LMS) was used to analyze the surface morphology of the CdS and Sb (S,Se) films. The surface roughness and CPD were measured using a Bruker Multimode 8 equipped with an AFM and a KPFM. The absorption and transmission spectra were measured by a spectrophotometer (UV-2600i). The crystal structure of the Sb (S,Se) films was measured by an X-ray diffractometer (XRD, Bruker D8 Discover) with a radiation source of Cukα. The energy level structures of Ta O and Sb (S,Se) films were measured by UV photoelectron spectroscopy (UPS, Thermo Fisher Scientific ESCALAB Xi) at −5 V bias. Core energy level variations of different elements in Ta O and Sb (S,Se) films were analyzed using an X-ray photoelectron spectrometer (XPS, Thermo Scientific Nexsa). Femtosecond TAS of Sb (S,Se) films was tested using a pumped probe system (Helios, Ultrafast System LLC). TPC, TPV, EIS, and C−2-V curves were measured using a Zahner electrochemical workstation. C–V and DLCP measurements were conducted using a Keithley 4200 semiconductor characterization system on device structures of Glass/FTO/CdS/without or with Ta O /Sb (S,Se) /Spiro-OMeTAD/Au. The C–V measurements were performed in the dark at room temperature with an alternating current (AC) amplitude of 30 mV, a frequency of 100 kHz and a direct current (DC) bias voltage range of 1 to 0.3 V. DLCP measurements were performed with an AC amplitude range of 15 to 140 mV and a DC bias voltage range of −0.2 to 0.2 V. The J–V curves of the devices were measured by combining a standard xenon sunlight simulator (LSS-55) with a test source meter (Keithley 2450) at a simulated AM 1.5G (100 mW·cm−2) light intensity. The J–V curves were measured from −0.02 V to 0.70 V with a step size of 5 mV and no wait time. The EQE of the devices was measured using a quantum efficiency test system (LST-QE).

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "In this manuscript, the authors developed a field-effect passivation strategy by inserting a low-work-function Ta2O5 dielectric layer at the heterojunction to promote charge transfer and suppress carrier recombination. This method is found to form interface dipoles between CdS and Sb2(S,Se)3 layers, which can selectively extract electron-carriers and block hole-carriers. The authors demonstrated a certified efficiency of 10.65% with low voltage deficit. The device performance and in-depth analysis make this study highly valuable for the field. Thus, it can be recommended for acceptance of publication, after several of the concerns are addressed.（1）The authors must cite more latest papers in these two years about the development of related photoelectric applications to make your introduction and discussions more comprehensive.（2）The surface SEM characterization is insufficient to demonstrate the formation of a uniform and conformal Ta2O5 layer. It is recommended to conduct EDS mapping of the CdS/Ta2O5 substrate to verify the homogeneity and coverage of the Ta2O5 films.（3）To confirm the stability of the Ta2O5 layer during the subsequent high-temperature hydrothermal deposition of Sb2(S,Se)3 and to assess potential interfacial interdiffusion, cross-sectional TEM images of the CdS/Ta2O5/Sb2(S,Se)3 interface should be provided.（4）Can a Ta2O5 dielectric layer replace the CdCl2 treatment? The Ta2O5 incorporated devices without CdCl2 treatment should be supplemented to explore whether the passivation mechanisms are complementary.（5）The authors should provide the key electrical parameters of Ta2O5 dielectric layer, such as dielectric constant and conductivity. These parameters are fundamental to evaluating the efficacy of field-effect passivation.（6）Semiconductor device simulations should be employed by introducing an interfacial dipole moment or modulating interface charge density to provide a theoretical basis for the field-effect passivation of Ta2O5.（7）The significant improvement of FF is commonly related to the interface. Therefore, the resistance characteristics are needed to analyze FF losses.（8）To gain deeper insights into the charge transport and interfacial properties within the device, it is recommended that the authors conduct impedance tests under conditions ranging from low to high light intensities. This method can effectively probe the charge extraction process and its evolution under varying light intensities.（9）Reliance solely on spin-coating rotational speed for estimating the thickness of the Ta2O5 layer introduces significant uncertainty. The authors shall employ a rigorous technique to precisely determine the thickness of the dielectric layer.", "reviewer_2": "Gong et al proposed using a field-effect passivation strategy by introduction of a low-work-function Ta2O5 interfacial layer to suppress the VOC loss and achieve a record PCE of Sb2(S,Se)3 solar cells. It is of interest and meaningful. However, there are some major issues needed to address.\n1. The authors claimed that the field-effect passivation originates from Ta2O5-induced interfacial dipole, which promotes both majority carrier extraction and suppression of minority carrier accumulation. Nevertheless, previous studies demonstrated that there exist interfacial fixed charges in Ta2O5. I strongly recommend the fixed charges should be measured, to determine whether the interfacial dipole or the fixed charges play the role of field-effect passivation.\n2. Ta2O5 with different thicknesses were used as the interfacial layer. The thickness should be measured, not in the form of spin-coating rate (Fig. S12).\n3. “we achieved a VOC of 695 mV, which is comparable to the highest reported value (700 mV) for Sb2(S,Se)3 solar cells[7], along with a state-of-the-art PCE of 10.95% and a certified PCE of 10.65%.” Since the VOC is closely related to the bandgap of photovoltaic material, the comparison of VOC is not rigorous.\n4. “the deposition of Ta2O5 effectively filled surface pinholes and cracks in the CdS films and reduced the depth of their grain boundaries”, the description of reduced the depth of their grain boundaries is not accurate.\n5. For Fig. S5b, “the local XRD analysis revealed that Sb2(S,Se)3 films deposited on Ta2O5 substrates exhibited stronger [hk1] diffraction peaks”. It is not the case for the (320) plane. Besides, I noticed that there is a shift of the diffraction peaks, for which the reasons can be analyzed.\n6. In Fig. S7, what are the large granular precipitates in the deposited Sb2(S,Se)3 films?\n7. “Electrons migrate from Ta2O5 to Sb2(S,Se)3 to achieve fermi-level alignment, inducing downward band bending in Sb2(S,Se)3 and upward band bending in Ta2O5.” I suggest the corresponding band bending should be illustrated in Fig. 2c.\n8. The shifts of the Se 3d and S 2p peaks to larger binding energies (Fig. 2d, e) and the Ta 4f peaks to lower binding energies (Supplementary Fig.10) is ascribed to the interfacial charge transfer. Normally, the shifts of binding energies in XPS is related to the changes of electron cloud densities. As the authors mentioned, XPS is limited to the depth of 5-10 nm, and the Ta2O5 films were deposited on the rear interface of Sb2(S,Se)3 films. Therefore, the results of XPS measurements could not provide the evidences of interfacial charge transfer.\n9. Fig. 3g and h show the doping density by C-V and DLCP measurements. As mentioned, defect densities measured by CV contain the bulk defects, interfacial defects, and free carriers, while the DLCP is less sensitive to the interfacial defects. How to distinguish the different defects and determine the interfacial defect density? In addition, what does the depth in Fig. 3g and h mean? The architecture or the cross-sectional image of the devices for the measurement should be given.\n10. For Fig. 4c, the improvement of EQE response should be discussed and analyzed.\n11. For Fig. 4i, the fitted RS and Rrec values of solar cells can be given.\n12. How about the long-term stability of the solar cells?", "reviewer_3": "In the manuscript titled, ‘Field-effect Passivation for Minimized Voltage Loss in Highly Efficient Antimony Selenosulfide Solar Cells,’ the authors have fabricated high-efficiency Sb2(S,Se)3 solar cells. The high efficiency of ~10.95% was possible sequentially by physical passivation strategy, such as field-effect passivation in Sb2(S,Se)3 photovoltaics. I appreciate the authors for using various characterizations to support their findings. I value the author's work since the the introduction of Ta2O5 not only optimizes the CdS buffer layer and the Sb2(S,Se)3 absorber layer, but also provides additional driving force for extracting electron-carrier to cathode and mitigating hole-carrier accumulation at the interface. In my view, this work is valuable to this field and thus I support the publication of this work in Nature Communications provided the following concerns are well addressed to make this manuscript clearer to the readers.\n1) The authors propose that the Ta2O5 layer reduces the exposure of polar facets in CdS. To make this mechanism clearly, direct evidence (e.g., TEM, XRD, or DFT analysis) is presented.\n2) The authors observe that the Ta2O5 layer suppresses the formation and deposition of Sb2(S,Se)3 cluster particles. However, the underlying mechanism remains unclear. A detailed explanation of how Ta2O5 influences the growth kinetics, nucleation behavior, or interfacial energetics of Sb2(S,Se)3 films should be provided.\n3) The authors attribute improved performance to optimized grain orientation but omit texture coefficient analysis. Quantitative comparison of preferred orientations (e.g., (hk0) vs. (hk1)) on CdS vs. CdS/Ta2O5 substrates is necessary to validate this hypothesis.\n4) To comprehensively assess the interfacial stability of the Ta2O5 layer and the structural integrity of the heterojunction, cross-sectional HRTEM should be performed.\n5) This reviewer finds the XPS analysis dubious. Specifically, the peak with a binding energy of 162-163 eV should be assigned to S 2p1/2, the peak with a binding energy of 161-162 eV should be assigned to S 2p3/2, and the peak with a binding energy near 160 eV should be assigned to Se 3p3/2. Clearly, the peak fitting in Figure 3e appears to be questionable, as the authors used selenourea as the Se source, meaning that the film should contain a certain amount of Se. I think the authors should refer to previous related works on XPS analysis of antimony selenosulfide thin films, such as the papers published by Li et al. (Adv. Mater. 2024, 36, 2410669) and Ren et al. (Adv. Mater. 2025, 37, 2416885.). Please assign each fitting peak exactly to its corresponding atomic orbital (and not simply by labeling these peaks).\n6) Further investigation is required to elucidate the influence of Ta2O5 dielectric layer on the electron mobility of the devices.\n7) For Figure S2, please provide a descriptive label for colour bars, e.g. Height (nm).\n8) Please correct the format of the references. For instance, the title of Ref [1] has capitalized initial letters, while Ref [2] has lowercase. Please recheck."}, "author_response": "Point-by-Point Response\nREVIEWER COMMENTS\nReviewer #1 (Remarks to the Author):\nIn this manuscript, the authors developed a field-effect passivation strategy by inserting\na low-work-function Ta O dielectric layer at the heterojunction to promote charge\n2 5\ntransfer and suppress carrier recombination. This method is found to form interface\ndipoles between CdS and Sb (S,Se) layers, which can selectively extract electron-\n2 3\ncarriers and block hole-carriers. The authors demonstrated a certified efficiency of\n10.65% with low voltage deficit. The device performance and in-depth analysis make\nthis study highly valuable for the field. Thus, it can be recommended for acceptance of\npublication, after several of the concerns are addressed.\nResponse: We thank the reviewer’s recognition on our paper, and the valuable\ncomments that help us improve the quality of our paper.\n(1) The authors must cite more latest papers in these two years about the development\nof related photoelectric applications to make your introduction and discussions more\ncomprehensive.\nResponse: Thank you very much for your suggestion. We have carefully studied the\nhigh-impact papers from the last two years (2024-2025) in the field of photovoltaic\napplications, particularly those related to novel photovoltaic materials, physical\npassivation, and interface engineering. Accordingly, we have added the latest papers in\nthe revised manuscript.\n(2) The surface SEM characterization is insufficient to demonstrate the formation of a\nuniform and conformalTa O layer. It is recommended to conduct EDS mapping of the\n2 5\nCdS/Ta O substrate to verify the homogeneity and coverage of the Ta O films.\n2 5 2 5\nResponse: Thank the reviewer for the suggestion. To further verify the uniformity of\nthe ultrathin Ta O layer, we have performed EDS mapping analysis on the CdS/Ta O\n2 5 2 5\nsubstrate. As shown in Supplementary Fig. 2, the Ta signal (representing Ta O )\n2 5\nexhibits a continuous and homogeneous distribution on the surface of CdS films, with\nno observable island-like aggregation or large-scale uncovered areas. This indicates a\nuniform and complete coverage of the CdS film by the deposited Ta O layer.\n2 5\nSimultaneously, the signals of Cd and S also display a homogeneous distribution,\nfurther confirming that the Ta O layer is sufficiently ultra-thin, thereby allowing the\n2 5\nunderlying CdS signals to remain clearly detectable.\nSupplementary Figure 2. Top-view SEM image and EDS mapping of the CdS/Ta O\n2 5\nfilms.\n(3) To confirm the stability of the Ta O layer during the subsequent high-temperature\n2 5\nhydrothermal deposition of Sb (S,Se) and to assess potential interfacial interdiffusion,\n2 3\ncross-sectional TEM images of the CdS/Ta O /Sb (S,Se) interface should be provided.\n2 5 2 3\nResponse: We sincerely thank the reviewer for raising this critical point. Since\ninterfacial stability is fundamental to device reproducibility and long-term reliability,\nwe performed HAADF-STEM imaging and corresponding EDS mapping analysis for\na cross-sectional sample with structure of FTO/CdS/Ta O /Sb (S,Se) . As shown in\n2 5 2 3\nSupplementary Fig. 11, the HAADF-STEM image reveals sharp and distinct interfaces\nbetween all functional layers. Notably, a continuous thin layer with a thickness of ~3\nnm, corresponding to Ta O , is clearly observed between the CdS and Sb (S,Se) layers.\n2 5 2 3\nEDS mapping further demonstrates a sharp and confined Ta signal precisely at the\nCdS/Sb (S,Se) interface, while other elements such as Cd, S, Sb, and Se exhibit\n2 3\nnegligible interdiffusion. These results confirm that the Ta O dielectric layer maintains\n2 5\nexcellent chemical and structural stability during subsequent high-temperature\nhydrothermal and annealing processes, forming well-defined heterointerfaces with\nadjacent layers. The corresponding descriptions and data have been added to the revised\nmanuscript.\nSupplementary Figure 11. Cross-sectional high-angle annular dark-field scanning\ntransmission electron microscope (HAADF-STEM) image and corresponding EDS\nelemental mappings of the Glass/FTO/CdS/Ta O /Sb (S,Se) multilayer sample.\n2 5 2 3\n(4) Can a Ta O dielectric layer replace the CdCl treatment? The Ta O incorporated\n2 5 2 2 5\ndevices without CdCl treatment should be supplemented to explore whether the\n2\npassivation mechanisms are complementary.\nResponse: Thank the reviewer for this valuable suggestion. To elucidate the effects of\nCdCl or Ta O treatments on device performance, we measured the current density–\n2 2 5\nvoltage (J–V) curves of the devices with only CdCl treatment or with only the Ta O\n2 2 5\ndielectric layer (without CdCl treatment). As shown in the below Figure and Table, the\n2\ndevice with solely Ta O layer achieved a PCE of 8.7%, slightly lower than that of the\n2 5\ntraditional CdCl treated device (9.2%). It is worth noting that the highest PCE of 10.95%\n2\nwas achieved in this work through the synergistic combination of CdCl and Ta O\n2 2 5\ntreatments. Thus, we conclude that the Ta O dielectric layer cannot replace the CdCl\n2 5 2\ntreatment.\nFigure J-V curves of devices with only CdCl treatment and only Ta O layer.\n2 2 5\nTable Photovoltaic parameters of devices with only CdCl treatment and only Ta O\n2 2 5\nlayer.\nSample J (mA cm−2) V (V) FF (%) PCE (%)\nSC OC\nOnly CdCl 22.27 0.651 63.41 9.2\n2 5\n(5) The authors should provide the key electrical parameters of Ta O dielectric layer,\n2 5\nsuch as dielectric constant and conductivity. These parameters are fundamental to\nevaluating the efficacy of field-effect passivation.\nResponse: According to the literatures, Ta O has a high dielectric constant of\n2 5\napproximately 28.[1] To determine the conductivity of Ta O dielectric layer, we\n2 5\nconducted Hall effect measurements, which revealed an ultralow electrical conductivity\n(σ ~ 2.11×10−5 S/cm). The high dielectric constant is critical for efficient field-effect\npassivation, as it enables the Ta O layer to establish and maintain a strong electric field\n2 5\nacross the interface, thereby inducing significant band bending in the semiconductor\nand providing a driving force for charge separation. These descriptions and data have\nbeen added to the revised manuscript.\n[1] Sahoo, K. K., Pradhan, D., Ghosh, S. P., Gartia, A. & Kar, J. P. Modulation of\nelectrical properties of sputtered Ta O films by variation of RF power and\n2 5\nsubstrate temperature. Phys. Scr. 99, 025934 (2024).\n(6) Semiconductor device simulations should be employed by introducing an interfacial\ndipole moment or modulating interface charge density to provide a theoretical basis for\nthe field-effect passivation of Ta O .\n2 5\nResponse: Thank the reviewer for this suggestion. To theoretically verify the field-\neffect passivation mechanism of the Ta O layer, we performed the device simulations\n2 5\nby using the SCAPS-1D software.[1-4] A device model is established that accurately\ncorresponds to the actual structure of the solar cells. All the essential materials\nparameters are shown in Supplementary Table 3. The simulation results\n(Supplementary Fig. 12) demonstrate that the insertion of the Ta O layer significantly\n2 5\nreduces the conduction band offset (ΔE ) between CdS and Sb (S,Se) from 280\nC,ETL 2 3\nmeV to 50 meV. The Ta O layer provides a more favorable path for efficient transport\n2 5\nof photogenerated electrons from the absorber layer to the electron transport layer.\nFurthermore, the simulated J-V curves show significant improvements in V and FF\nOC\n(Supplementary Fig. 16), which is fully consistent with our experimental results.\nDevice simulations demonstrate that field-effect passivation induced by the Ta O layer\n2 5\nserves as a key factor for the enhanced device performance. The relevant simulation\ndetails and results have been added to the revised manuscript.\nSupplementary Figure 12. Simulated band diagrams of (a) the control device and (b)\nthe Ta O incorporated device using SCAPS-1D software.\n2 5\nSupplementary Figure 16. Simulated J-V curves of the control device and the Ta O\n2 5\nincorporated device obtained from SCAPS-1D software.\nSupplementary Table 8. Simulated photovoltaic parameters of the control device and\nthe Ta O incorporated device obtained from SCAPS software.\n2 5\nSample J (mA/cm2) V (V) FF (%) PCE (%)\nSC OC\nControl 22.71 0.655 62.14 9.24\nWith Ta O 23.25 0.693 69.60 11.21\n2 5\n[1] Wang, X., Tang, R., Wu, C., Zhu, C. & Chen, T. Development of antimony sulfide–\nselenide Sb (S,Se) -based solar cells. J. Energy Chem. 27, 713-721 (2018).\n2 3\n[2] Ali, H., Warsame, A., Penrose, C., John, A. & John, F. Simulation and analysis of\nmethylammonium lead iodide (CH NH PbI ) perovskite solar cell with Au contact\n3 3 3\nusing SCAPS 1D simulator. American Journal of Optics and Photonics 7, 33-40 (2019).\n[3] Chen, S. et al. Simultaneous band alignment modulation and carrier dynamics\noptimization enable highest efficiency in Cd-free Sb Se solar cells. Adv. Funct. Mater.\n2 3\n34, 2403934 (2024).\n[4] Burgelman, M., Nollet, P. & Degrave, S. Modelling polycrystalline semiconductor\nsolar cells. Thin Solid Films 361-362, 527-532 (2000).\n(7) The significant improvement of FF is commonly related to the interface. Therefore,\nthe resistance characteristics are needed to analyze FF losses.\nResponse: Thank you very much for your suggestion. We have added the series\nresistance (R ) and shunt resistance (R ) of the devices in the revised manuscript. The\ns sh\nlower series resistance and higher shunt resistance contribute to improving the FF of\ndevices.\nTable Series resistance (R ) and shunt resistance (R ) of the control and Ta O\ns sh 2 5\nincorporated devices.\nSample Rs (ohm) Rsh (ohm)\nControl 110.09 4837.53\nWith Ta O 68.39 29177.22\n2 5\n(8) To gain deeper insights into the charge transport and interfacial properties within\nthe device, it is recommended that the authors conduct impedance tests under conditions\nranging from low to high light intensities. This method can effectively probe the charge\nextraction process and its evolution under varying light intensities.\nResponse: Thanks for the good suggestion. To investigate the interfacial charge transfer\nproperties of the devices, we performed electrochemical impedance spectroscopy (EIS)\nmeasurements under different light intensities. As shown in Supplementary Fig. 25a,b,\nall the Nyquist plots exhibit a single semicircular arc, indicating that the interfacial\nprocess is dominated by a single time constant.[1,2] The equivalent circuit model was\nused to fit the impedance data (Supplementary Fig. 25c), where R represents the series\nS\nresistance, R corresponds to the charge transfer resistance, and the constant phase\nct\nelement (CPE) denotes the geometric capacitive behavior of the device.[3-5] The R\nct\nvalue of both devices decreases linearly with the increasing light intensity\n(Supplementary Fig. 25d), which can be attributed to that the increased photogenerated\ncarrier concentration effectively reduces the energy barrier of charge transfer. Notably,\nthe R values of the Ta O incorporated device were significantly lower than that of the\nct 2 5\ncontrol device under different light intensities. This demonstrates that the fixed charges\nin the Ta O dielectric layer effectively reduce the charge transfer barrier through field-\n2 5\neffect passivation effect, thereby promoting efficient extraction and transport of charge\ncarriers.[6-7]\nSupplementary Figure 25. Nyquist plots of a) control and b) Ta O incorporated devices\n2 5\nmeasured under different light intensities. c) Equivalent circuit model. d) Variation of\ncharge transfer resistance (R ) with light intensity.\nct\n[1] Hermoso-Diaz, I. A. et al. Experimental and theoretical studies of α-linolenic acid\nas green corrosion inhibitor for carbon steel in 0.5 M sulfuric acid. Molecules 26\n(2021).\n[2] Lazanas, A. C. & Prodromidis, M. I. Electrochemical impedance spectroscopy–a\ntutorial. ACS Meas. Sci. Au 3, 162-193 (2023).\n[3] Asada, T. et al. Influence of the electron transport layer on the performance of\nperovskite solar cells under low illuminance conditions. ACS Omega 9, 32893-\n32900 (2024).\n[4] Chen, J., Seo, J.-Y. & Park, N.-G. Simultaneous improvement of photovoltaic\nperformance and stability by in situ formation of 2D perovskite at\n(FAPbI ) (CsPbBr ) /CuSCN interface. Adv. Energy Mater. 8, 1702714 (2018).\n3 0.88 3 0.12\n[5] Zarazúa, I. et al. Operating mechanisms of mesoscopic perovskite solar cells\nthrough impedance spectroscopy and J–V modeling. J. Phys. Chem. Lett. 8, 6073-\n6079 (2017).\n[6] Hu, Y. et al. Construction of charge transport channels at the NiO /perovskite\nx\ninterface through moderate dipoles toward highly efficient inverted solar cells.\nACS Appl. Mater. Interfaces 14, 13431-13439 (2022).\n[7] Chen, J., Zhao, X., Kim, S.-G. & Park, N.-G. Multifunctional chemical linker\nimidazoleacetic acid hydrochloride for 21% efficient and stable planar perovskite\nsolar cells. Adv. Mater. 31, 1902902 (2019).\n(9) Reliance solely on spin-coating rotational speed for estimating the thickness of the\nTa O layer introduces significant uncertainty. The authors shall employ a rigorous\n2 5\ntechnique to precisely determine the thickness of the dielectric layer.\nResponse: Thank the reviewer for this suggestion. To precisely determine the thickness\nof Ta O layers, we used spectroscopic ellipsometry to measure the thickness of the\n2 5\nTa O dielectric layer deposited on a silicon substrate. The thickness of Ta O layers\n2 5 2 5\nprepared by different rotational speeds (4000, 5000 and 6000 rpm) is determined to be\n5 ± 0.5, 3 ± 0.5 and 1 ± 0.5 nm, respectively. Moreover, the cross-sectional TEM image\nshows a thickness of ~3 nm for the Ta O layer prepared under rotational speed of 5000\n2 5\nrpm. The consistency between the two independent techniques demonstrates that the\noptimal thickness of Ta O layer is ~3 nm. The corresponding contents have been\n2 5\nshown in the revised manuscript.\nReviewer #2 (Remarks to the Author):\nGong et al proposed using a field-effect passivation strategy by introduction of a low-\nwork-function Ta O interfacial layer to suppress the V loss and achieve a record PCE\n2 5 OC\nof Sb (S,Se) solar cells. It is of interest and meaningful. However, there are some major\n2 3\nissues needed to address.\nResponse: We thank the reviewer for the positive comments on our work. We also\nappreciate the reviewer’s valuable and constructive comments to help us significantly\nimprove the manuscript quality.\n1. The authors claimed that the field-effect passivation originates from Ta O -induced\n2 5\ninterfacial dipole, which promotes both majority carrier extraction and suppression of\nminority carrier accumulation. Nevertheless, previous studies demonstrated that there\nexist interfacial fixed charges in Ta O . I strongly recommend the fixed charges should\n2 5\nbe measured, to determine whether the interfacial dipole or the fixed charges play the\nrole of field-effect passivation.\nResponse: We sincerely thank the reviewer for their insightful comments. To\ninvestigate the fixed charges in the Ta O layer, we fabricated metal oxide\n2 5\nsemiconductor (MOS) structured devices for C-V measurements. The corresponding\ndevice structure is FTO/CdS/Sb (S,Se) /Au or FTO/CdS/Ta O /Sb (S,Se) /Au. By\n2 3 2 5 2 3\ndifferentiating the C-V data, the voltage corresponding to the extreme value represents\nthe flat-band voltage. As shown in Fig. 2a,b, the flat-band voltages of the CdS and\nCdS/Ta\n5\nsamples are -0.35 V and -0.50 V, respectively. We therefore confirm that the\nTa O enables positive fixed charges at the interface, consistent with reported\n2 5\nliteratures.[1,2] The fixed charge (Q) density was calculated using the following\nf\nformula:[3]\n \nC V V\nQ  in MS fb (1)\nf Aq\nHere, C denotes the intrinsic capacitance of the MOS structure, V represents the\nin MS\nmetal-semiconductor potential difference determined by the work function difference\nbetween FTO (4.8 eV) and Sb (S,Se) , V is the flat-band voltage, A is the electrode\n2 3 fb\narea, and q is the elementary charge. The calculated fixed charge density Q of the Ta O\nf 2 5\nlayer is +2.24 × 1011 cm−2 (Supplementary table 2).\nThese positive fixed charges facilitate efficient electron extraction while suppressing\nhole accumulation at the interface.[4,5] Based on these results, we conclude that the\npositive fixed charges in Ta O are the primary factor to induce the field-effect\n2 5\npassivation. The corresponding description and data have been revised in the\nmanuscript. We once again thank the reviewer for their valuable guidance.\nFigure 2. (a) C-V curves, (b) flat band voltage of the CdS and CdS/Ta O devices.\n2 5\nSupplementary Table 2. Parameters used to calculate the Q.\nf\nSample C (nF) V (eV)a V (V) Q (cm-2)b\nin MS fb f\nCdS 11.19 0.78 -0.35 9.86×1011\nCdS/Ta O 12.15 0.78 -0.50 1.21×1012\n2 5\naThe W of FTO was taken as 4.80 eV,[6] while the work function of the Sb (S,Se)\nf 2 3\nabsorber layer, determined by UPS measurements, is 4.02 eV in this work.\nbThe Q of the Ta O dielectric layer was calculated as the difference between the Q of\nf 2 5 f\nthe CdS/Ta O sample and that of the bare CdS sample.\n2 5\n[1] Noh, Y. S. et al. Characteristics of MIS capacitors using Ta O films deposited on\n2 5\nZnO/p-Si. Microelectron. Eng. 66, 637-642 (2003).\n[2] Novkovski, N. & Atanassova, E. Charge trapping during constant current stress in\nHf-doped Ta O films sputtered on nitrided Si. Thin Solid Films 519, 2262-2267\n2 5\n(2011).\n[3] Cunha, J. M. V. et al. Perovskite metal–oxide–semiconductor structures for\ninterface characterization. Adv. Mater. Interfaces 8, 2101004 (2021).\n[4] Yang, Y. et al. Inverted perovskite solar cells with over 2,000 h operational stability\nat 85 °C using fixed charge passivation. Nat. Energy 9, 37-46 (2024).\n[5] Zhang, D. et al. Mitigated front contact energy barrier for efficient and stable\nperovskite solar cells. Energy Environ. Sci. 17, 3848-3854 (2024).\n[6] Gu, Y. et al. Solvent annealing enabling reconstruction of cadmium sulfide film\nfor improved heterojunction quality and photovoltaic performance of antimony\nselenosulfide solar cells. Adv. Funct. Mater. 34, 2311577, (2024).\n2. Ta O with different thicknesses were used as the interfacial layer. The thickness\n2 5\nshould be measured, not in the form of spin-coating rate (Fig. S12).\nResponse: To precisely determine the thickness of Ta O layers, we used spectroscopic\n2 5\nellipsometry to measure the thickness of the Ta O dielectric layer deposited on a silicon\n2 5\nsubstrate. The thickness of Ta O layers prepared by different rotational speeds (4000,\n2 5\n5000 and 6000 rpm) is determined to be 5 ± 0.5, 3 ± 0.5 and 1 ± 0.5 nm, respectively.\nMoreover, the cross-sectional TEM image shows a thickness of ~3 nm for the Ta O\n2 5\nlayer prepared under rotational speed of 5000 rpm. The consistency between the two\nindependent techniques demonstrates that the optimal thickness of Ta O layer is ~3\n2 5\nnm. The corresponding contents have been shown in the revised manuscript.\n3. “we achieved a V of 695 mV, which is comparable to the highest reported value\nOC\n(700 mV) for Sb (S,Se) solar cells, along with a state-of-the-art PCE of 10.95% and a\n2 3\ncertified PCE of 10.65%.” Since the V is closely related to the bandgap of\nOC\nphotovoltaic material, the comparison of V is not rigorous.\nOC\nResponse: Thank the reviewer for raising this important point regarding analytical rigor.\nWe have deleted the direct comparison of absolute V values, and corresponding\nOC\ncontents have been shown in the revised manuscript.\n4. “the deposition of Ta O effectively filled surface pinholes and cracks in the CdS\n2 5\nfilms and reduced the depth of their grain boundaries”, the description of reduced the\ndepth of their grain boundaries is not accurate.\nResponse: Thank the reviewer for pointing out this issue. We have deleted the\ndescription of reduced the depth of their grain boundaries.\n5. For Fig. S5b, “the local XRD analysis revealed that Sb (S,Se) films deposited on\n2 3\nTa O substrates exhibited stronger [hk1] diffraction peaks”. It is not the case for the\n2 5\n(320) plane. Besides, I noticed that there is a shift of the diffraction peaks, for which\nthe reasons can be analyzed.\nResponse: Thanks for this comprehensive comment. We have carefully analyzed the\nXRD data and calculated the texture coefficient (TC) of Sb (S,Se) films.As shown in\n2 3\nSupplementary Fig. 6, the TC value of (230) plane decreases and the (hk1) orientations,\nsuch as (211), (221), (301) and (421) orientations increase with the insertion of Ta O\n2 5\nlayer. This trend aligns fully with our expectations, as (hk1) orientations are beneficial\nto the carrier transport across the absorber compared to (hk0) orientations.[1,2]\nRegarding the slight shift in diffraction peaks, we attribute it to the changes in the\nchemical composition of the absorber layer (Supplementary Fig. 8 and Supplementary\nTable 1). The introduction of Ta O substrate modifies the crystallization kinetics,\n2 5\nleading to an altered S/Se ratio in Sb (S,Se) and the resultant shift of diffraction peak.[3]\n2 3\nThe corresponding description and data have been added to the revised manuscript.\nSupplementary Figure 6. (a) XRD patterns and (b) corresponding texture coefficients\nof Sb (S,Se) films deposited on CdS and CdS/Ta O substrates.\n2 3 2 5\nSupplementary Figure 7. (a) Selected-region XRD patterns, (b) corresponding full\nwidth at half maximum (FWHM) values of Sb (S,Se) films deposited on CdS and\n2 3\nCdS/Ta O substrates.\n2 5\nSupplementary Figure 8. SEM images and corresponding EDS mapping of Sb (S,Se)\n2 3\nfilms deposited on CdS and CdS/Ta O substrates.\n2 5\nSupplementary Table 1. Elemental content and S/Se ratios derived from EDS mapping.\nSample Element Wt (%) Atom (%) S/Se\nS 21.23 46.62\nCdS/Sb (S,Se) Se 24.92 22.23 2.10\n2 3\nSb 53.85 31.15\nS 17.84 41.78\nCdS/Ta O /Sb (S,Se) Se 22.57 21.46 1.95\n2 5 2 3\nSb 59.60 36.76\n[1] Ren, D. et al. PO 3− tetrahedron assisted chelate engineering for 10.67%-efficient\n4\nantimony selenosulfide solar cells. Adv. Mater. 37, 2416885 (2025).\n[2] Huang, L. et al. Temperature-gradient solution deposition amends unfavorable\nband structure of Sb (S,Se) film for highly efficient solar cells. Angew. Chem. Int.\n2 3\nEd. 63, e202406512 (2024).\n[3] Dai, G. et al. Selenium-rich Sb (S,Se) thin films deposited via sequential chemical\n2 3\nbath deposition for high-efficiency solar cells. Adv. Funct. Mater. 35, 2415215\n(2025).\n6. In Fig. S7, what are the large granular precipitates in the deposited Sb (S,Se) films?\n2 3\nResponse: We thank the reviewer for the valuable comments. We have performed the\nSEM image and EDS mapping of the Sb (S,Se) film. As shown in Supplementary Fig.\n2 3\n8, the large particulates observed on the surface of Sb (S,Se) film were identified as\n2 3\naggregates of Sb (S,Se) particles. Their formation is attributed to non-uniform\n2 3\nheterogeneous nucleation on the CdS surface during the hydrothermal process. With\nthe introduction of the Ta O layer, the Sb (S,Se) film exhibit more uniform\n2 5 2 3\nmorphology with fewer large particles. This improvement indicates that the Ta O layer\n2 5\nprovides a superior substrate that effectively modulates the crystallization kinetics,\nwhich is consistent with the water contact angle measurements on CdS and CdS/Ta O\n2 5\nsubstrates (Fig. 1d).\nSupplementary Figure 8. SEM images and corresponding EDS mapping of Sb (S,Se)\n2 3\nfilms deposited on CdS and CdS/Ta O substrates.\n2 5\n7. “Electrons migrate from Ta O to Sb (S,Se) to achieve fermi-level alignment,\n2 5 2 3\ninducing downward band bending in Sb (S,Se) and upward band bending in Ta O .” I\n2 3 2 5\nsuggest the corresponding band bending should be illustrated in Fig. 2c.\nResponse: We thank the reviewer for this valuable suggestion. As suggested, we have\nillustrated the band bending of Ta O /Sb (S,Se) in Fig. 2c. When positive fixed charges\n2 5 2 3\nof Ta O appear at the interface, more electrons of Sb (S,Se) gather therein, resulting\n2 5 2 3\nin a downward band bending of Sb (S,Se) .[1] In addition to the fixed positive charges,\n2 3\nFermi level alignment between Ta O /Sb (S,Se) also results in band bending with the\n2 5 2 3\nsame trend. As shown in Fig. 2c, the band bending of Sb (S,Se) is downward while the\n2 3\nband bending of Ta O is upward when the fermi-level alignment is achieved between\n2 5\nTa O /Sb (S,Se) . The control device exhibits a large conduction band (CB) offset\n2 5 2 3\n(ΔE = 280 meV) between CdS and Sb (S,Se) . With the insertion of Ta O layer,\nC,ETL 2 3 2 5\nthe presence of fixed charges and the significant work function difference induce more\npronounced downward band bending in Sb (S,Se) and reduce the ΔE to 50 meV.\n2 3 C,ETL\nThe Ta O layer provides a more favorable path for efficient transport of\n2 5\nphotogenerated electrons from the absorber layer to the electron transport layer. The\ncorresponding description and data have been revised in the manuscript.\nFigure 2c. Schematic band diagrams of (a) the reference device and (b) the device with\nthe Ta O dielectric layer.\n2 5\n8. The shifts of the Se 3d and S 2p peaks to larger binding energies (Fig. 2d, e) and the\nTa 4f peaks to lower binding energies (Supplementary Fig.10) is ascribed to the\ninterfacial charge transfer. Normally, the shifts of binding energies in XPS is related to\nthe changes of electron cloud densities. As the authors mentioned, XPS is limited to the\ndepth of 5-10 nm, and the Ta O films were deposited on the rear interface of Sb (S,Se)\n2 5 2 3\nfilms. Therefore, the results of XPS measurements could not provide the evidences of\ninterfacial charge transfer.\nResponse: Thank the reviewer for this important comment. We apologize for our\ninaccurate analysis of the XPS measurements. Since the structure of XPS sample does\nnot correspond to the actual device structure, which cannot accurately reflect the true\nstate. We have deleted the corresponding description and data in the revised manuscript.\n9. Fig. 3g and h show the doping density by C-V and DLCP measurements. As\nmentioned, defect densities measured by C-V contain the bulk defects, interfacial\ndefects, and free carriers, while the DLCP is less sensitive to the interfacial defects.\nHow to distinguish the different defects and determine the interfacial defect density? In\naddition, what does the depth in Fig. 3g and h mean? The architecture or the cross-\nsectional image of the devices for the measurement should be given.\nResponse: Thanks for this comprehensive comment. Generally, low-temperature DLCP\nmeasurements can reflect the concentration of free charge carriers, while the defect\ndensity from high-temperature DLCP measurements (N ) contains the free carriers\nDLCP\nand bulk defects.[1] The defect density from C-V measurements (N ) includes the free\nCV\ncarriers, bulk defects, and interfacial defects.[2,3] Therefore, the interfacial defects can\nbe determined by the difference between the N and N . The depth in Fig. 3g and\nCV DLCP\nh represents the depletion width (W ) of space-charge region (SCR). The architecture\nd\nof devices for the measurement is Glass/FTO/CdS/without or with\nTa O /Sb (S,Se) /Spiro-OMeTAD/Au. The corresponding contents have been shown in\n2 5 2 3\nexperimental section of the revised manuscript.\n[1] Duan, H.-S. et al. The role of sulfur in solution-processed Cu ZnSn(S,Se) and its\n2 4\neffect on defect properties. Adv. Funct. Mater. 23, 1466-1471 (2013).\n[2] Fu, B. et al. Reaction kinetics regulation suppressed carrier recombination loss for\nhigh-efficient solution-based antimony selenosulfide photovoltaic devices. Adv.\nEnergy Mater. 15, 2500586 (2025).\n[3] Cao, L. et al. Modifying surface termination by bidentate chelating strategy enables\n13.77% efficient kesterite solar cells. Adv. Mater. 36, 2311918 (2024).\n10. For Fig. 4c, the improvement of EQE response should be discussed and analyzed.\nResponse: In the EQE of solar cells, the short-wavelength region primarily shows the\nphotoelectric response within buffer layer and at the front interface. The mid-\nwavelength region mainly indicates the photoelectric response within the absorption\nlayer and near the p-n junction interface. The long-wavelength region primarily reflects\nthe photoelectric response at the back contact of device.[1,2] As shown in Fig. 4b, the\nresponse in the spectral range of 330-500 nm is enhanced due to the optimization of the\nCdS buffer layer and the heterojunction interface, while the response in the spectral\nrange of 500-750 nm is increased because of the improved quality of Sb (S,Se)\n2 3\nabsorber layer. The corresponding discussion has been added to the revised manuscript.\nFigure 4b. EQE spectra of the reference device and the device with the Ta O dielectric\n2 5\nlayer.\n[1] Thiesbrummel, J. et al. Universal current losses in perovskite solar cells due to\nmobile ions. Adv. Energy Mater. 11, 2101447 (2021).\n[2] Liu, X. et al. Enhanced Sb Se solar cell performance through theory-guided defect\n2 3\ncontrol. Prog. Photovoltaics 25, 861-870 (2017).\n11. For Fig. 4i, the fitted R and R values of solar cells can be given.\nS rec\nResponse: As suggested, we have fitted the impedance spectra using the equivalent\ncircuit model to acquire the relevant parameters. As shown in Supplementary Fig. 24,\nthe extracted R of devices decreased from 19.79 to 3.35 Ω, and the R of devices\nS rec\nincreased from 1358.85 to 4676.42 Ω with the insertion of Ta O dielectric layer.\nThe\n2 5\ncorresponding data have been added to the revised manuscript.\nSupplementary Figure 24. Electrochemical impedance spectroscopy (EIS) for control\nand Ta O incorporated devices.\n2 5\n12. How about the long-term stability of the solar cells?\nResponse: Thank the reviewer for their attention to device stability. In fact, the stability\nof this type of inorganic thin-film solar cell is exceptionally good, and numerous studies\nhave confirmed this (Nat. Energy 2025, 10, 857–868). As shown in Fig. 4h, we also\nperformed the operational stability measurement of an unencapsulated target devices\nthat maintaining ∼96% of its initial PCE after continuous tracking at maximum power\npoint (MPP) for 200 h. The corresponding description and data have been added to the\nrevised manuscript.\nFigure 4h. Maximum power point (MPP) tracking of the unencapsulated device with\nthe Ta O dielectric layer under ambient conditions.\n2 5\nReviewer #3 (Remarks to the Author):\nIn the manuscript titled ‘Field-effect Passivation for Minimized Voltage Loss in Highly\nEfficient Antimony Selenosulfide Solar Cells’, the authors have fabricated high-\nefficiency Sb (S,Se) solar cells. The high efficiency of ~10.95% was possible\n2 3\nsequentially by physical passivation strategy, such as field-effect passivation in\nSb (S,Se) photovoltaics. I appreciate the authors for using various characterizations to\n2 3\nsupport their findings. I value the author's work since the introduction of Ta O not only\n2 5\noptimizes the CdS buffer layer and the Sb (S,Se) absorber layer, but also provides\n2 3\nadditional driving force for extracting electron-carrier to cathode and mitigating hole-\ncarrier accumulation at the interface. In my view, this work is valuable to this field and\nthus I support the publication of this work in Nature Communications provided the\nfollowing concerns are well addressed to make this manuscript clearer to the readers.\nResponse: We appreciate the positive comments and strong support for the impact and\nnovelty of our work.\n1) The authors propose that the Ta O layer reduces the exposure of polar facets in CdS.\n2 5\nTo make this mechanism clearly, direct evidence (e.g., TEM, XRD, or DFT analysis)\nis presented.\nResponse: Thank the reviewer for this suggestion. To demonstrate that the Ta O layer\n2 5\nreduces the exposure of polar facets in CdS, we performed XRD measurements on CdS\nfilms without and with the Ta O layer. As shown in the Supplementary Fig. 5,\n2 5\ncompared to the bare CdS film, the CdS/Ta O sample exhibits significantly reduced\n2 5\npeak intensities of polar plane, such as (101), and enhanced peak intensities of nonpolar\nplane, such as (100).[1] These descriptions and data have been added to the revised\nmanuscript.\nSupplementary Figure 5. XRD patterns of CdS and CdS/Ta O samples showing (a) the\n2 5\nfull scan range and (b) a magnified view of selected diffraction peaks.\n[1] Li, K. et al. Molecular beam epitaxy deposition of in situ O-doped CdS films for\nhighly efficient Sb (S,Se) solar cells. Adv. Funct. Mater. 33, 2304141 (2023).\n2 3\n2) The authors observe that the Ta O layer suppresses the formation and deposition of\n2 5\nSb (S,Se) cluster particles. However, the underlying mechanism remains unclear. A\n2 3\ndetailed explanation of how Ta O influences the growth kinetics, nucleation behavior,\n2 5\nor interfacial energetics of Sb (S,Se) films should be provided.\n2 3\nResponse: Thank the reviewer for raising this insightful question. Regarding the\nsuppression of Sb (S,Se) cluster formation by the Ta O interlayer, we speculate that\n2 3 2 5\nthe mechanism is primarily associated with the interfacial energy modification and the\nnucleation behavior regulation. According to the classical nucleation theory, the\nformation of Sb (S,Se) film consists of a four-step process: (i) the formation of stable\n2 3\ncrystal nuclei; (ii) evolution of nuclei into an island structure; (iii) the generation of\nnetwork structure; and (iv) the formation of the continuous and compact film.[1] In the\ninitial nucleation step, the nucleation rate R exponentially decreases with the\nN\nnucleation barrier Δ:[2]\n \nR exp  (2)\nN   T  \nwhere T is temperaturein Kelvin.\nIn the hydrothermal process, the homogeneous and heterogeneous nucleation rates of\nSb (S,Se) exponentially decrease with the homogeneous Gibbs free energy (ΔG )\n2 3 Hom\nand the heterogeneous Gibbs free energy (ΔG ), respectively.The homogeneous and\nHet\nthe heterogeneous Gibbs free energy can be expressed by the equations below:[3,4]\n4\nG  r3nG 4r2n (3)\nHom 3 k V k SL\nG G f  (4)\nHet Hom\nwhere r is the critical radius of the nuclei, n is the number of critical nuclei per unit\nk\nvolume, ΔG is the change in free energy between the nuclei and solution per unit\nv\nvolume, γ is the particle surface energy, θ is the contact angle, and 0 < f(θ) < 1. Based\nSL\non the classical nucleation theory, a smaller contact angle can reduce Gibbs free energy\nfor heterogeneous nucleation, thereby increasing nucleating density to promote the film\ndensification process and preventing the formation of coarse grains.[3,4] We measured\nthe water contact angle of pure CdS and CdS/Ta O substrates. As shown in Fig. 1d, the\n2 5\nCdS/Ta O substrate exhibits a smaller water contact angle of 44.0º than that of pure\n2 5\nCdS substrate (61.9º), suggesting the reduced energy barrier for heterogeneous\nnucleation of Sb (S,Se) . Consequently, this suppresses abnormal grain coarsening and\n2 3\ncluster formation induced by local concentration inhomogeneity or defect-related\naggregation. These theoretical and experimental findings consistently support our\nhypothesis. The corresponding discussion and data have been added to the revised\nmanuscript.\nFigure 1d. Contact angle measurements of CdS versus CdS/Ta O samples.\n2 5\n[1] Li, J. et al. Manipulating the morphology of CdS/Sb S heterojunction using a Mg-\n2 3\ndoped tin oxide buffer layer for highly efficient solar cells. J. Energy Chem. 66,\n374-381 (2022).\n[2] Zhumekenov, A. A. et al. The role of surface tension in the crystallization of metal\nhalide perovskites. ACS Energy Lett. 2, 1782-1788 (2017).\n[3] Zhao, H. et al. Enhanced stability and optoelectronic properties of MAPbI films\n3\nby a cationic surface-active agent for perovskite solar cells. J. Mater. Chem. A 6,\n10825-10834 (2018).\n[4] Yang, D. et al. High efficiency planar-type perovskite solar cells with negligible\nhysteresis using EDTA-complexed SnO . Nat. Commun. 9, 3239 (2018).\n2\n3) The authors attribute improved performance to optimized grain orientation but omit\ntexture coefficient analysis. Quantitative comparison of preferred orientations (e.g.,\n(hk0) vs. (hk1)) on CdS vs. CdS/Ta O substrates is necessary to validate this\n2 5\nhypothesis.\nResponse: Thank the reviewer for this suggestion. We have calculated the texture\ncoefficients (TC) of the crystal planes for Sb (S,Se) films deposited on CdS and\n2 3\nCdS/Ta O substrates based on the XRD data (Supplementary Fig. 6a), with detailed\n2 5\nresults presented in Supplementary Fig. 6b.[1] The Sb (S,Se) films deposited on the\n2 3\nCdS/Ta O substrate exhibit significantly higher TC values for (hk1) planes, such as\n2 5\n(211), (221), (301) and (421) compared to those of Sb (S,Se) films grown on the bare\n2 3\nCdS substrate. This provides evidence for the optimized grain orientation of Sb (S,Se)\n2 3\nfilms induced by the Ta O layer. The corresponding descriptions and data have been\n2 5\nadded to the revised manuscript.\nSupplementary Figure 6b. Texture coefficients calculated from XRD patterns of\nSb (S,Se) films deposited on CdS and CdS/Ta O substrates.\n2 3 2 5\n[1] Yao, L.-Q. et al. A liquid medium annealing strategy for highly [041]/[141]-\noriented planar antimony sulfide solar cells with 7.23% efficiency. Nano Energy\n106, 108064 (2023).\n4) To comprehensively assess the interfacial stability of the Ta O layer and the\n2 5\nstructural integrity of the heterojunction, cross-sectional HRTEM should be performed.\nResponse: Thank the reviewer for this important suggestion. Since interfacial stability\nis fundamental to device reproducibility and long-term reliability, we performed\nHAADF-STEM imaging and corresponding EDS mapping analysis for a cross-\nsectional sample with structure of FTO/CdS/Ta O /Sb (S,Se) . As shown in\n2 5 2 3\nSupplementary Fig. 11, the HAADF-STEM image reveals sharp and distinct interfaces\nbetween all functional layers. Notably, a continuous thin layer with a thickness of ~3\nnm, corresponding to Ta O , is clearly observed between the CdS and Sb (S,Se) layers.\n2 5 2 3\nEDS mapping further demonstrates a sharp and confined Ta signal precisely at the\nCdS/Sb (S,Se) interface, while other elements such as Cd, S, Sb, and Se exhibit\n2 3\nnegligible interdiffusion. These results confirm that the Ta O dielectric layer maintains\n2 5\nexcellent chemical and structural stability during subsequent high-temperature\nhydrothermal and annealing processes, forming well-defined heterointerfaces with\nadjacent layers. The corresponding descriptions and data have been added to the revised\nmanuscript.\nSupplementary Figure 11. Cross-sectional high-angle annular dark-field scanning\ntransmission electron microscope (HAADF-STEM) image and corresponding EDS\nelemental maps of the Glass/FTO/CdS/Ta O /Sb (S,Se) multilayer sample.\n2 5 2 3\n5) This reviewer finds the XPS analysis dubious. Specifically, the peak with a binding\nenergy of 162-163 eV should be assigned to S 2p1/2, the peak with a binding energy of\n161-162 eV should be assigned to S 2p3/2, and the peak with a binding energy near 160\neV should be assigned to Se 3p3/2. Clearly, the peak fitting in Figure 3e appears to be\nquestionable, as the authors used selenourea as the Se source, meaning that the film\nshould contain a certain amount of Se. I think the authors should refer to previous\nrelated works on XPS analysis of antimony selenosulfide thin films, such as the papers\npublished by Li et al. (Adv. Mater. 2024, 36, 2410669) and Ren et al. (Adv. Mater.\n2025, 37, 2416885.). Please assign each fitting peak exactly to its corresponding atomic\norbital (and not simply by labeling these peaks).\nResponse: Thank the reviewer for their expert corrections regarding the XPS analysis.\nWe have carefully fitted the XPS data, and the assignment of characteristic peaks has\nbeen clarified. The peaks located at 161.5 eV and 162.6 eV are belonged to S 2p\nand\n3/2\nS 2p , respectively. The peaks located at 160.1 eV and 166.0 eV are ascribed to Se\n1/2\n3p and Se 3p , respectively. The peaks located at 53.9 eV and 54.8 eV are assigned\n3/2 1/2\nto Se 3d and Se 3d , respectively.[1,2]\n5/2 3/2\nFigure XPS spectra of (a) Se 3d and (b) S 2p core levels for Sb (S,Se) films deposited\n2 3\non CdS and CdS/Ta O substrates.\n2 5\n[1] Zhao, Y. et al. Innovative in situ passivation strategy for high-efficiency Sb (S,Se)\n2 3\nsolar cells. Adv. Mater. 36, 2410669 (2024).\n[2] Ren, D. et al. PO 3− tetrahedron assisted chelate engineering for 10.67%-efficient\n4\nantimony selenosulfide solar cells. Adv. Mater. 37, 2416885 (2025).\n6) Further investigation is required to elucidate the influence of Ta O dielectric layer\n2 5\non the electron mobility of the devices.\nResponse: Thank the reviewer for this valuable suggestion. To elucidate the influence\nof Ta O dielectric layer on the electron mobility of devices, we measured the electron-\n2 5\nonly devices by using space-charge-limited current (SCLC) model. The Ta O\n2 5\nincorporated devices exhibit a higher electron mobility of 8.78 cm2 V−1 s−1 than that of\nthe control devices (7.84 cm2 V−1 s−1), indicating that the introduction of Ta O\n2 5\neffectively facilitates the electron transport and collection. The corresponding\ndescriptions and data have been added to the revised manuscript.\nSupplementary Figure 14. Electron mobility extracted from the space-charge-limited\ncurrent (SCLC) model for the control and Ta O -incorporated electron-only devices.\n2 5\n7) For Figure S2, please provide a descriptive label for colour bars, e.g. Height (nm).\nResponse: Thank you very much for your reminding. We have added the descriptive\nlabel \"Height (nm)\" for the color bars of AFM images, and the corresponding contents\nhave been shown in the revised manuscript.\nFigure 1c,f. AFM images for (a) CdS films and (b) CdS/Ta O films.\n2 5\nSupplementary Figure 9. AFM images of Sb (S,Se) films deposited on (a) CdS and (b)\n2 3\nCdS/Ta O substrates.\n2 5\n8) Please correct the format of the references. For instance, the title of Ref [1] has\ncapitalized initial letters, while Ref [2] has lowercase. Please recheck.\nResponse: Thank you very much for your reminding. We have thoroughly checked and\nstandardized the formatting of all references in the revised manuscript."}, {"round": 2, "reviewer_comments": {"reviewer_1": "The revised version shall be accepted.", "reviewer_2": "The authors have well addressed the comments. I suggest it can be accepted in Nature Communications.", "reviewer_3": "In this manuscript, the authors have fabricated high efficiency Sb2(S,Se)3 solar cells. The high efficiency of ~10.95% was possible sequentially by physical passivation strategy, such as field-effect passivation in Sb2(S,Se)3 photovoltaics. The introduction of Ta2O5 not only optimizes the CdS buffer layer and the Sb2(S,Se)3 absorber layer, but also provides additional driving force for extracting electron-carrier to cathode and mitigating hole carrier accumulation at the interface. The authors have addressed all my concerned quesitons properly. And now the quality of the manuscirpt has been improved greatly. I think the work reported in the manuscript has reached the requirement of Nat. comm. Thus, I recommend to accept it with present quality."}, "author_response": ""}]

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