Solution epitaxy of single-crystal ferroelectric p-i-n structure with ultrahigh photovoltaic response

Domain: chemistry
Published: 2026-04-06
Editorial outcome: minor_revision

Abstract
Ferroelectric materials with a switchable polarization are appealing for optoelectronic applications because of their above-bandgap photovoltage with a potential to exceed the Shockley-Queisser limit efficiency within classic p-n junctions. However, the intrinsic photovoltaic current in ferroelectrics remains far below the junction ones by orders, hindering their practical applications. Here we report an extremely large photovoltaic current density of 130 mA cm-2 with a photovoltage of 1.31 V under 375 nm light illumination in single-domain lead titanate films. Its photoresponsivity surpasses all available ferroelectric materials by two orders of magnitude. Systematic investigations reveal that an interfacial Ta dopant layer during solution epitaxy drives a polarization converting of the film via the flexoelectric effect, where surface charged Pb vacancy layer was generated for polarization screening. These two layers produce a p-i-n structure throughout the film, accounting for the ultrahigh photovoltaic response. Furthermore, a self-powered X-ray detector made of such film was explored to achieve a sensitivity of 4.73 C Gy –1 air cm–3, three orders of magnitude higher than that of commercial amorphous Se detectors. Our findings may pave the way for exploring novel detection and imaging devices based on ferroelectric materials.

Introduction
The photovoltaic (PV) effect in semiconductors is the foundation of many technologies, such as solar energy conversion1, photocatalysis2 and optoelectronic sensing3. The key step to these applications is efficiently driving charge separation of electron-hole pair carriers or excitons by a built-in electric field and subsequent carrier transport. Towards this purpose, classic p-n junction was fabricated in the crystalline silicon solar cells. Nevertheless, the efficiency of solar-cell is curtailed by inherent energy losses, subjected to the Shockley-Queisser (S-Q) limit in single p-n junction4,5. Alternatively, perovskite oxide ferroelectric semiconductors, have triggered great interests for solar energy conversion6–8, where the above-bandgap voltage from several to several thousand volts can even be produced in single-phase ferroelectrics such as LiNbO 9, holding a potential to exceed such efficiency limit5. Interestingly, an electric polarization-direction-dependent PV effect has been demonstrated in BiFeO 3 (BFO)10. However, the resulting PV current density turns out to be quite low mainly due to poor bulk conductivity. Different strategies ranging from the enhanced absorption by doping11, high-efficient carrier collection by atomic force microscope (AFM) tips12,13 and polar order engineering14 have been developed to improve the PV current. Despite great efforts towards polarization-based PV effect in the past decades, the intrinsic PV current in ferroelectrics is still at the level of μA cm-2. The PV current mechanisms in ferroelectrics are still complex and controversial, involving shift current, ballistic current, and incompletely screened depolarization field or domain walls15. Despite of great efforts, an effective strategy to significantly improve PV current in ferroelectrics towards practical devices remain a big challenge16.

For typical crystalline silicon or halide perovskite solar cells, the p-n junction or p-i-n structure plays a decisive role in achieving efficient energy conversion17,18, where electron and hole transport layers as well as built-in electric field design are achieved simultaneously. For example, in halide perovskite solar cells, meticulous selection of transport layers and precise band structure matching are essential for optimizing p-i-n functionality. Nevertheless, analogous transport layers and their derived built-in electric field throughout single-phase ferroelectrics have not been explored yet. But it is very encouraging that intriguing electrical conductivity phenomena near domain walls have been discovered19 and illustrated systematically20. N-type or p-type conducting layer near domain walls in ferroelectrics could be generated through spatial confinement of charged defects, such as Bi vacancies and oxygen vacancies, compensating bound charges due to a spontaneous polarization discontinuity21. Interestingly, p-n junction based on single domain wall have been designed experimentally22. These advances suggest a possibility to enhance the PV current through domain wall confining charged defects. However, non-uniform direction of built-in electric field near every domain wall is not beneficial for unidirectional photogenerated carrier separation and transport. Moreover, domain walls have been identified to contribute to the high photovoltage rather than the PV current23.

In addition to domain walls, the polarization surface of single-domain ferroelectrics have been demonstrated to energetically confine various charged point defects for screening bound charges24. In principle, these defect layers near the opposite polarization surface can serve as electron and hole transport layers, as well as constructing a p-i-n structure to enhance the PV effect. Especially, it has been proved that oxygen vacancy layer is prevailed as a screening manner near the negative polarization surface of ferroelectrics25–27. However, oxygen vacancies hardly contribute to n-type conduction in perovskite ferroelectric oxides because of forming polarons by trapping electrons28,29.

In this work, we developed a strategy to control the polarization direction of single-domain ferroelectric lead titanate (PbTiO , PTO) films by an interfacial Ta dopant during a solution epitaxy. This enables constructing a p-i-n structure throughout the films, where the interfacial Ta doping layer and the polarization screening induced surface Pb vacancy layer can be employed as efficient n-type electron and p-type hole transport layers, respectively. As a result, an ultrahigh PV current density of 130 mA cm-2 with a photovoltage of 1.31 V under 375 nm light illumination was obtained, two orders of magnitude higher than that of emerging ferroelectric materials. Using these films, a self-powered and direct-type X-ray detector was fabricated to demonstrate high sensitivity and stability which is highly desired for next generation X-ray imaging devices.

Results
Single-crystal single-domain pure PTO and Ta doped PTO (Ta-PTO) films have been grown with a solution epitaxy method30 on single-crystal 0.7% Nb doped SrTiO (NSTO) substrates at a relatively low temperature (~200°C) (see methods for more details). X-ray diffraction (XRD) analysis confirmed that both films exhibit a tetragonal perovskite structure (JCPDS#70-0746) (Supplementary Fig. 1a). As shown in Supplementary Fig. 1b, a single Ta-PTO (103) diffraction peak is observed together with the NSTO (103) substrate peak, indicating epitaxial growth of the Ta-PTO film on the NSTO substrate. Additionally, 𝜙 scan revealed four peaks at 90° intervals, indicating the single-crystal nature of the Ta-PTO film (Supplementary Fig. 1c), which also has a smooth and continuous surface (Supplementary Fig. 1d). For tetragonal perovskite PTO, the spontaneous polarization (P ) arises from the off-center displacements of Ti ions31, was analyzed using atomic-resolution high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) images (Supplementary Fig. 2 and Fig. 3). As shown in Fig. 1a, Ti ions in the pure PTO film display upward displacements relative to the Pb ion columns, indicating a downward polarization along the film's thickness direction. The thickness of the pure PTO film is ~1030 nm. In contrast, Fig. 1b reveals opposite displacements of Ti ions in the Ta-PTO film, with a thickness of ~1100 nm. The out-of-plane displacement of Ti ions in Ta-PTO increases to ~0.20 Å at the surface, slightly smaller than the ~0.26 Å observed for pure PTO. Piezoresponse force microscopy (PFM) confirms that the Ta-PTO film exhibits typical ferroelectric characteristics with a single-domain configuration (Supplementary Fig. 4). Furthermore, the rocking curves scanned around the (002) diffraction peak further confirms the single-crystal nature of the films (Supplementary Fig. 5), with a full width at half maximum (FWHM) of 0.385° for Ta-PTO and 0.237° for pure PTO, further supporting the single-domain configuration.

To uncover the mechanism behind the opposite polarization direction in pure PTO and Ta-PTO films, we analyzed their interface structures. Fig. 1c,d show the lattice parameters a and c. The interfacial c-value in Ta-PTO (~3.99 Å) is smaller than in pure PTO (~4.11 Å). Notably, it increases to ~4.15 Å within the first three unit cells above the interface, indicating an inhomogeneous compressive strain (~3.86%) between the interface and the inner regions of the Ta-PTO film. The interfacial strain was further analyzed using geometric phase analysis (GPA) based on atomic-level HAADF-STEM images (Supplementary Fig. 6). As shown in Fig. 1e,f, unlike pure PTO film, Ta-PTO film exhibits a distinct interfacial compressive strain (ε ) layer, likely caused by the incorporation of Ta dopants, considering the slightly larger ionic radius of Ta5+ (0.640 Å) compared to Ti4+ (0.605 Å)32. At the nanoscale, a large stress gradient can induce a local electric field via the flexoelectric effect. Such gradients can be generated artificially, for example by localized mechanical loading with an AFM tip, as widely reported to induce flexoelectric fields and control ferroelectric polarization33,34. This effect arises from the coupling between a strain gradient and ferroelectric polarization35. Moreover, phase-field modeling was used to calculate the flexoelectric field within the PTO layer caused by strain gradient from local concentration inhomogeneity36 (Supplementary Note 1). As shown in Fig. 1g, the calculated field strength at six unit cells can be up to 130 kV cm-1, much higher than the coercive field (E ) of PTO37, enabling polarization reversal. This flexoelectric effect energetically favors the reversal of the polarization direction of the Ta-PTO film layer near the interface from downward to upward, creating a built-in field that could switch the polarization throughout the film (Fig. 1h).

The opposite polarization direction of the films shown in Fig. 1a,b necessitate different screening mechanisms. Below the Curie temperature (T ), P in PTO generates a depolarization field due to opposite surface-bound charges, which destabilizes the polarization38. Effective screening or compensation of these charges is crucial for stabilizing a single-domain structure24,39. For pure PTO film, positively charged oxygen vacancies typically screen negatively charged surfaces, while electrons from the electrode compensate for positively charged surfaces, as revealed in PbZr Ti O and BFO films25,27. As a comparison, the cross-section microstructure of Ta-PTO film was systemically characterized by using advanced electron microscopy techniques to understand its screening mechanism.

Fig. 2a,d indicate an atomically sharp interface and a smooth surface of the Ta-PTO film. Energy-dispersive X-ray spectroscopy (EDS) line scanning (Fig. 2b) shows a decrease in the Ti signal and an increase in the Ta signal within three unit cells near the interface, while the Pb signal remains nearly unchanged (Supplementary Fig. 7a–c). In addition, a significant enhancement of Ta signals near the interface (Fig. 2c) can be confirmed by time-of-flight secondary ion mass spectrometry (TOF-SIMS). These results support that Ta dopants probably occupy Ti sites in the perovskite tetragonal PTO lattice. The electron energy loss spectroscopy (EELS) characterizations indicate that Ta doping induces oxygen vacancies near the interface (Supplementary Fig. 8). The relatively higher peak 2 compared to peak 1 in the upper O-K spectra confirms that oxygen vacancies are largely confined within ~4 nm near the Ta-PTO film interface40. Meanwhile, the Pb intensity decreases towards the surface within ~5 nm of the Ta-PTO thin film (Fig. 2e and Supplementary Fig. 7d–e). Additionally, X-ray photoelectron spectroscopy (XPS) confirms the presence of Pb vacancies in the surface layer (Supplementary Fig. 9) and reveals that the Pb/Ti ratio in Ta-PTO film is ~14.1% lower than in pure PTO film. Therefore, in single-domain Ta-PTO thin films, the polarization screening by negatively charged Pb vacancies and positively charged oxygen vacancies stabilizes the upward polarization, as shown in Fig. 2f.

In addition, the Ti ion displacement related to Pb position has been used to determine the polarization intensity. For example, the P intensity of PTO is ~67 μC cm-2 when the displacement is ~0.38 Å24. In contrast, Ta-PTO shows a reduced polarization intensity with a ~0.2 Å displacement, corresponding to a Pb vacancy concentration of ~16% that can be inferred by complete polarization screening (34 μC cm-2) near the surface.

The screening mechanism in Fig. 2f could mediate a band structure of Ta-PTO films. To understand this issue, we performed density functional theory (DFT) calculations (Supplementary Note 2). The calculations were conducted at a 12.5% concentration of both Pb vacancies and Ta substitutions (Supplementary Fig. 10). The results revealed that Ta substitution for Ti (Ta· ) is more easily formed acting as an effective donor, while Pb vacancies behave as an acceptor. The local density of states (LDOS) across Ta-doped, pristine, and Pb-vacancy-rich regions in Fig. 2g,h indicate a type-II band alignment. In the Pb-vacancy-rich region, the Fermi level is closer to the valence band maximum (VBM), raising both the VBM and conduction band minimum (CBM) by ~0.2 eV. Conversely, in the Ta-doped region, the Fermi level is closer to the CBM, lowering both the CBM and VBM by a similar energy offset. To this end, we have illustrated the energy band structure of Ta-PTO (Fig. 2i). The synergistic effect of surface Pb vacancies, the pristine phase, and the Ta-doped interfacial layer collectively forms a band bending structure and corresponding built-in electric field that is analogy to a p-i-n structure. Moreover, the presence of a built-in electric field was further verified by kelvin probe force microscopy (KPFM) (Supplementary Fig. 11 and Supplementary Fig. 12)41, where the surface potential was found and decreased gradually from the Ta-PTO film surface towards the NSTO substrate.

Since a p-i-n structure is available in Ta-PTO film, it is interesting to explore PV performance of Ta-PTO film (~2.8 eV in bandgap) using a 375 nm linearly polarized laser. Fig. 3a,b show the current density-voltage (J-V) characteristics of the films under light excitation and dark field, respectively. Under zero bias, Ta-PTO film achieves an ultrahigh short-circuit current (J ) of 1018.5 μA (~130 mA cm-2), remarkably surpassing the 31.6 μA (~4.03 mA cm- 2) observed in pure PTO film under 500 mW cm-2 irradiation. The corresponding open-circuit voltage (V ) is ~1.31 V, exceeding the ~1.22 V of pure PTO film. Notably, the pure PTO film exhibits a linear J-V curve in the -1.5 to 1.5 V voltage range, indicating typical bulk PV behavior42. In contrast, Ta-PTO film shows a nonlinear J-V response, which is probably attributed to a p-i-n structure, as shown in Fig. 2i.

To further evaluate the PV performance of the Ta-PTO film, the response under varying light intensities (I ) were measured. J increases almost linearly with I due to the generation of more photogenerated carriers. In contrast, V gradually saturates as the quasi-Fermi level separation reaches its limit (Fig. 3c)4. As shown in Fig. 3d, the J response to the on-off switching of light at zero bias demonstrated excellent control and stability. To accurately assess the response speed of the Ta-PTO film, a time-resolved pulse response system was utilized with a 266 nm pulsed laser. The zero bias pulse response, shown in Fig. 3e, reveals a rise time (t) of ~0.49 μs (10% to 90%) and a decay time (t ) of ~15.66 μs (90% to 10%). Fig. 3f shows the photoresponsivity (R = J /I ) under zero bias and different wavelengths, with a maximum value reaching 0.5078 A W-1 at 375 nm (I = 50 mW cm-2) and 0.1477 A W-1 at 254 nm (I = 4.09 mW cm-2) (Supplementary Fig. 13). Moreover, R of Ta-PTO film was compared with different UV detectors based on ferroelectric and classic semiconductors from literatures, as presented in Fig. 3g. Remarkably, its self-powered response surpasses emerging ferroelectric materials by over two orders of magnitude. In the deep ultraviolet region, the responsivity of the Ta-PTO detector surpasses that of p-i-n structured third-generation semiconductors, such as GaN and SiC (Supplementary Table 1).

Ta-PTO film device exhibits outstanding PV performance not only in the UV range but also in the hard X-ray band (10–100 keV). For long-term device applications, the structural stability of perovskite ferroelectric oxides is highly desirable compared to halide perovskite43. More attractively, its high density (~7.52 g cm-3) and the large atomic number for Pb (82), makes the film attractive for direct-type X-ray detector. To evaluate the photon absorption properties, the photon cross-section database44 was used in conjunction with the material densities to calculate the photon absorption spectra of Si, amorphous Se, CdZnTe (CZT), organic-inorganic hybrid halide perovskite methylammonium lead iodide (MAPbBr ; MA = CH NH +), purely inorganic halide perovskite CsPbBr , and PTO in the X-ray energy range of 0.001 to 1 MeV (Fig. 4a). Due to its high theoretical density, and disregarding K, L, and M edge resonant absorptions, PTO demonstrates an obvious advantage in absorption coefficient across the entire photon energy range.

The J-V characteristics of the Ta-PTO film detector under X-ray irradiation and in dark conditions are shown in Fig. 4b. A key performance for X-ray detectors is sensitivity (C Gy – 1 air cm–3), which is calculated by determining the slope of the linear region in the current density-dose rate relationship and multiplying it by the active layer thickness. The Ta-PTO film detector achieved an impressive sensitivity of 4.73 C Gy –1 cm–3 (Fig. 4c). As compared in Supplementary Table 2, the sensitivity of Ta-PTO detector under zero bias is much higher than those of typical X-ray detection materials. It is approximately three orders of magnitude higher than that of commercial amorphous Se detectors (0.00125 C Gy –1 cm–3) operating under a 2 kV bias, and about 50 times higher than that of CZT detectors (0.087 C Gy –1 cm–3). Furthermore, it surpasses the performance of CsPbBr₃ thin films (0.26 C Gy –1 cm–3) and two-dimensional layered perovskite (BA) (MA) Pb I (0.276 C Gy –1 cm–3) under zero bias, indicating its superior sensitivity for advanced X-ray detection.

An ultra-low dark current of 1.8 nA cm-2 has been obtained in the detector at zero bias, resulting in an on/off ratio exceeding 500 and stable switching behavior (Supplementary Fig. 14). The radiation tolerance of the detector was evaluated by subjecting it to X-ray radiation at a dose rate of 1.46 mGy s-1 for a total duration of 2.5 h, where a cumulative dose is up to 14 Gyₐᵢᵣ, approximately equivalent to the dose received from 10,000 chest digital radiography (DR) scans45. These performances confirm the excellent radiation stability of the Ta-PTO detector, highlighting its reliability under extended exposure to X-ray irradiation (Supplementary Fig. 15). As shown in Fig. 4d, we further constructed an automated scanning system to evaluate the self-powered X-ray imaging capability of the detector. A seahorse specimen and coil spring were selected as the imaging target (Fig. 4e and Supplementary Fig. 16a,b), successfully visualizing both the external outline and internal structure of the samples under zero bias (Fig. 4f and Supplementary Fig. 16c). In contrast, a commercial amorphous Se detector failed to obtain similar imaging resolution by using the same seahorse specimen (Supplementary Fig. 17).

In summary, we report a polarization-controlled PV response in single-domain PTO films by interfacial Ta dopant. An ultrahigh PV current of ~130 mA cm-2 at 375 nm was obtained and the corresponding photoresponsivity under zero bias was 100 times higher than those of the emerged ferroelectrics. Interestingly, a self-powered X-ray detector with excellent sensitivity has been fabricated using this film, enabling high-quality specimen imaging under X-ray exposure. Theoretical and experimental studies reveal that interfacial Ta doping during solution epitaxy controls the polarization direction via flexoelectric effect, while simultaneously creating a Pb vacancy layer near the film surface for polarization screening. Both n-type interfacial doping and p-type surface defect layers generate a p-i-n structure throughout the film, leading to the observed photovoltaic properties and enhanced device performance. Our findings could pave a new way for energy conversion and optoelectronic devices ranging from UV to X-ray band by using ferroelectric materials.

Methods
Materials.
Lead nitrate (Pb(NO ) , ≥99.999%), tetrabutyl titanate (TBOT, ≥98%) were purchased from Aladdin Chemistry Co. Tantalum pentoxide (Ta O , ≥99.999%) was purchased from Alfa Aesar. Potassium hydroxide (KOH, ≥85.0%), acetone (≥99.5%) and ethanol (≥99.7%) were purchased from Sinopharm Chemical Reagent Co. Ltd. All the reagents were used as received. single-crystal 0.7 wt.% Nb doped SrTiO (NSTO) substrates with dimensions 10 × 10 × 0.5 mm3 and (100) orientation were obtained from Shanghai Institute of Optics and Mechanics.

Film preparation.
As shown in the Supplementary Fig. 18, PTO and Ta-PTO thin films were synthesized by a hydrothermal solution-epitaxy method. Pb(NO ) , TBOT and Ta O were used as metal precursors, and KOH served as the mineralizer. NSTO single-crystal substrates were ultrasonically cleaned in acetone, ethanol, deionized water, and ethanol sequentially, and then placed in a Teflon holder for drying.

In a Teflon liner, 15 mL of deionized water was first added. KOH (9.54 g) was dissolved to prepare a solution with a total precursor concentration of 6.8 M. Subsequently, 1.708 mL of TBOT was added under vigorous stirring. For Ta-PTO growth, Ta O was introduced at 0.11 g, corresponding to a Ta substitution level of 10 mol% (Ta/(Ta+Ti) = 10%). The Pb/Ti molar ratio was maintained at 1.25:1 to compensate for Pb volatility. Accordingly, Pb(NO ) (2.04 g) was dissolved separately in 10 mL of deionized water and then added dropwise into the precursor solution under continuous stirring.

After stirring for 2 h at room temperature, the NSTO substrate was placed horizontally in the precursor solution, positioned approximately 15 mm above the bottom of the Teflon liner. The Teflon liner was then sealed inside a stainless-steel autoclave, and the hydrothermal epitaxial growth was carried out at 200 °C for 12 h. After the reaction, the samples were thoroughly cleaned by repeated ultrasonication in deionized water and ethanol, and then dried in air at 60 °C.

Characterization.
The structures of the sample were characterized by Scanning electron microscope (SEM, Hitachi SU-70 Analytical FESEM), X-ray Diffraction (XRD, Thermo ARL X’TRA powder diffractometer, Cu Kα radiation, λ=1.54056 Å), Piezoresponse force microscopy (PFM, Asylum Research, Cypher S and MFP-3D), X-Ray photoelectron spectroscopy (XPS, Thermo Scientific ESCALAB 250Xi). An integrated Focused Ion Beam (FIB, Helios nanolab 600) and SEM (Hatachi S4700) was used to extract thin-film lamellae from the bulk sample, Pt and Au layers were deposited on film surface to protect from damage and increase the conductivity, respectively. Subsequent cross-sectional observations of the film were carried out using aberration-corrected transmission electron microscope (FEI Titan G2 80-200 Chemi STEM, ThermoFisher Spectra 300). The variation in composition across the thickness of PTO films were characterized by Time-of-flight secondary-ion mass (TOTF-SIMS, IONTOF GmbH TOF.SIMS 5-100).

Ti-ion displacement analysis.
Ti off-center displacements were quantified from atomic-level HAADF-STEM images by measuring the offset between each Ti atomic column and the geometric center of its four nearest Pb neighbors. Atomic column positions were determined using Gaussian peak fitting, and the out-of-plane displacement component was extracted to evaluate local polarization. Displacement statistics were collected over multiple unit cells to analyze polarization variation across the film thickness.

Photovoltaic measurement.
Photovoltaic properties of PTO and Ta-PTO films were collected using optical units from Thorlabs and source measure units (Keithley 2450). Top Ag electrodes (diameter: ~1 mm, thickness: ~20 nm, determined using an ellipsometer) were deposited by magnetron sputtering (VCT300) on the PTO surface as the top electrode and NSTO substrate acted as the bottom electrode. The films were illuminated with light sources including a 254 nm UV mercury lamp (XingChuangUV) and lasers at 266 nm (MPL-F-266), 375 nm (MDL-III-375), 405 nm (MDL-III-405), and 532 nm (MGL-III-532). The response time was measured using a preamplifier (Stanford Research Systems SR570) in conjunction with an oscilloscope (Tektronix TBS2204B).

X-ray detection performance measurement.
The X-ray detection performance evaluation was carried out by using an Amptek Mini-X2 tube with the energy of ≤50 keV and peak intensity at 22 keV. The output cone angle of X-ray was 120°, and the tube voltage was kept at 50 kV. A Radcal Accu-Gold+10X6-180 ion chamber dosimeter was used to calibrate the dose rate, which was adjusted by varying the tube current. The X-ray detector was placed at the same position as the dosimeter. The photocurrent and dark current from the device were recorded using a source measure units (Keithley 2635B).

For the X-ray imaging test, a single-pixel imaging system was constructed. A lead mask with a 0.5×0.5 mm² hole was used to collimate the incident X-rays. The object to be imaged was moved between the X-ray source and the detector using an X-Y linear stage (PI M-406.2DG) with a step size of 300 µm. The photocurrent at each point was recorded using a source measure units (Keithley 2635B). For comparison, the remote-vision mobile DR uDR 380i Pro amorphous Se imaging device from Shanghai United Imaging Healthcare was used.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The authors investigate single-crystal Ta-doped PbTiO3 films grown by solution epitaxy for the purpose of energy\nconversion and optoelectronic applications. They first characterize the structural and ferroelectric properties of the Ta-PTO\nfilms using a variety of complementary techniques, such as transmission electron microscopy, scanning probe microscopy\nand X-ray based techniques and compare them with pristine PTO films of the same thickness. This allows them to\nunderstand the screening mechanisms which lead to the formation of a single-domain state in Ta-PTO.\nBeyond this characterization, the authors continue by investigating the photo voltaic performance of their Ta-PTO film, and\ncompare their results with other ferroelectric systems, as well as commercially available products. It is revealed, that Ta-PTO\nperforms significantly better (by approx. 2 orders of magnitude) than other ferroelectric systems that have been investigated\nin literature so far. Finally, the authors developed a proof-of-concept self-powered X-ray detector, using the Ta-PTO films,\nand demonstrate its performance under various conditions and compare them with the current state-of-the-art detectors.\nThe manuscript is well-written, rigorous and their conclusions are well-supported by their findings. I recommend the\nmanuscript for publication in Nature Communications after some minor revisions and answers to some remaining questions:\n1. Can the authors provide an estimate on how high the Ta-doping of the PTO films is and include it when introducing them?\n2. In figures 1 a,b the authors show the displacement of the Ti ion with respect to the Pb ions. Could the authors provide a\ncouple of sentences in the methods section on how these displacements were obtained from the HAADF-STEM images.\n3. In figure 1 e,f the authors show GPA-based strain maps. Could the authors include where the 0-strain reference for these\nmaps was chosen (I assume in the substrate).\n4. Also in figure 1 e, from the position of the arrow indicating the interface with the substrate in this map, it appears that the\ncompressive strain is primarily located in the substrate, while the EDS lineprofiles provided in Figures 2b and S6 show the\ntantalum is incorporated in the first layers of the PTO film. Is the interface marked incorrectly in the strain map, or does\nthe incorporation of Ta really compressively strain the STO substrate?\n5. The authors mention, that the substitution of Ti4+ with Ta5+ induces the formation of oxygen vacancies. Do they have an\nexplanation for the mechanism that leads to this or provide a reference that explains it?\n6. The presence of oxygen vacancies at the substrate interface is supported by EELS-data of the O-K edge in Fig. S7. Could\nthe authors provide the ratio of intnsity between peaks 1/2 to highlight the changes in the O-K edge with distance to the\ninterface, as they are a bit difficult to make out from the line-plot alone.\n7. In lines 194-195 the authors state \"polarization screening from downward to upward switching can be realized...\". I believe\nwhat they meant to say is \"charge screening of the upward facing polarization can be realized/is realized by\", as they do not\nperform polarization switching experiments to my understanding.\n8. In order to PV performance of the PTO films, the authors have to cap their films with a metal (here Ag) electrode. Have the\nauthors considered that the introduction of this additional metal-interface could influence the polarization state of the\nferroelectric below? How would this interface influence the presence/formation of lead vacancies at the surface?", "reviewer_2": "The present manuscript titled “Solution epitaxy of single crystal ferroelectric p-i-n structure with ultra high photovoltaic\nresponse” is about demonstrating the high short-circuit current, which is two orders of magnitude larger than the value\nreported in the literature pertaining to ferroelectric photovoltaic effect. Although the films fabricated by solution epitaxy show\nbetter photovoltaic results, the manuscript has serious flaws in terms of the film geometry proposed for the PV studies, the\nclaimed ferroelectric photovoltaic effect and the mechanism proposed. The following are the detailed comments.\n1. The fabrication method given in the supplementary section is not very clear. What do they mean by single-crystal\nferroelectric film? Is it epitaxy? If so, further measurements like reciprocal space mapping have to be given.\n2. If I understood correctly, they have doped Ta into the PbTiO3. Is it doped throughout the film or only at the interface? As\nper the manuscript, I presume that the Ta is doped throughout the film. In that case, what do they mean by Ta interfacial\ndoping? The entire mechanism explained in this manuscript is based on the Ta at the interface. However, it is very difficult to\nfollow the formation of Ta doping only at the interface.\n3. Again, the Pb vacancies at the top surface are not clearly brought out in this manuscript, which is also crucial to\nunderstand the mechanism proposed in this manuscript.\n4. Overall, the formation of p-i-n structure is difficult to understand based on the fabrication details given in this manuscript.\n5. What is the role of ferroelectricity in the observed PV response? A lot has been discussed in the introduction related to the\nferroelectric photovoltaic effect. But the proof for ferroelectric PV is still not clear in this manuscript.\n6. Proof for the ferroelectric nature of the grown film, like P-E measurement, domain switching studied by PFM, etc, which\nare essential to ascertain the ferroelectric nature of the film, are missing in this manuscript.\n7. In case the observed PV effect is purely based on the p-i-n effect, then the value of the short-circuit current should not be\ncompared to the literature value on the ferroelectric photovoltaic system.\n8. If it is due to the ferroelectric PV effect, proof for the switchable PV effect should be given, which is missing in this\nmanuscript.\n9. In addition, some of the sentences are not clear. Examples are given below\n“The photovoltaic (PV) effect for semiconductors affords versatile applications from solar\nCells…..”\n“Piezoelectric force microscopy (PFM) confirms a single-domain structure of Ta-PTO film with typical ferroelectric and\npiezoelectric characteristics (Supplementary Fig. 1d).” Contrary to this statement, the figure does not give any details about\nthe piezoelectric characteristics.\n“ At the nanoscale, a large stress gradient (e.g., with the force applied through an AFM tip) can induce a local electric field\nthrough a flexoelectric effect, which is capable to switch the local polarization”. This is confusing. Where from the AFM tip\ninduced flexoelectric field is coming in this study? The flexoelectric field could be the compressive strain gradient at the\ninterface.\n“The results revealed that Ta substitution for Ti (TaTi · ) is more easily formed acting as an effective donor, while Pb\nvacancies behave as an acceptor.” Where is the proof?\nThe discussion about the Pb-vacancy-rich region and the Ta-doped region is very vague from the perspective of the grown\nTa-doped PbTiO3.\nOverall, the clarity about the samples subjected to this study itself not very clear in this manuscript. Moreover, the\nmechanism proposed for the observed ultra-high PV response based on the p-i-n structure is not clearly brought out in this\nmanuscript. Without the proof for the ferroelectric PV effect, the value of short-circuit current reported in this manuscript is\nindeed comparable only to the semiconductor-based PV effect, and it should not be compared to the ferroelectric PV effect.\nDue top these flaws, I am not recommending the manuscript for publication."}, "author_response": "Response to Reviewers’ Comments\nWe would like to thank both reviewers for all the constructive comments and suggestions,\nwhich significantly help us improve the quality of the present work. We made a detailed\nresponse to the reviewers’ comments point by point. Accordingly, the manuscript has been\ncarefully revised by providing additional experiments and discussions. The revised parts of the\nmanuscript are highlighted in red.\nReviewer #1\nQuestion 1: Can the authors provide an estimate on how high the Ta-doping of the PTO films\nis and include it when introducing them?\nResponse: Thank you for your constructive comments. In this work, the Ta-PTO films were\nprepared by using a precursor solution with a nominal Ta content of 10 mol% (Ta/(Ta+Ti) =\n10%). Regarding the absolute Ta concentration in the interfacial layer, we argue that\nquantitative determination by conventional EDS is rather challenging because the Ta-rich\nregion is a few unit cells in thickness, and the EDS analysis inevitably involves signals from\nthe adjacent NSTO substrate. This made an estimation of Ta/Ti ratio being difficult to quantify.\nAccordingly, we employed the nominal Ta content in the precursor.\nRevision: The following sentence has been modified in the Film preparation section of the\nmethod.\n……For Ta-PTO growth, Ta O was introduced at 0.11 g, corresponding to a Ta substitution\n2 5\nlevel of 10 mol% (Ta/(Ta+Ti) = 10%).\nQuestion 2: In figures 1 a,b the authors show the displacement of the Ti ion with respect to the\nPb ions. Could the authors provide a couple of sentences in the methods section on how these\ndisplacements were obtained from the HAADF-STEM images.\nResponse: Thank you for your constructive comments. In PTO, both the oxygen octahedra and\nthe titanium ions undergo cooperative displacements relative to the center of the lead-sublattice\ntetragonal unit cell, giving rise to the spontaneous polarization. Therefore, the off-center\ndisplacement of the Ti ions can be directly used to quantitatively determine the polarization\nstate of a PTO unit cell1. Fig. R1 shows a schematic of the measurement for determining the\naverage displacement of a Ti ion relative to the centers of its four nearest neighboring Pb ion\ncolumns. Based on this, we experimentally extracted the Ti-ion displacements in both pure\nPTO (Fig. 1a) and Ta-PTO (Fig. 1b) thin films using cross-sectional high-angle annular dark-\nfield scanning transmission electron microscopy (HAADF-STEM) images, and quantitatively\nanalyzed the distribution of these displacement values as a function of distance from the\ninterface.\nWe have now added a concise description of this displacement-extraction procedure in\nthe Methods section to improve clarity and reproducibility.\nig. R1 | a,b, Atomic-level HAADF-STEM image (a) and Schematic illustration (b) showing the measurement\nof the average displacement of Ti ions relative to the center of the four nearest Pb atomic columns in each\nregion. δ denotes the off-center displacement of Ti ions along the out-of-plane direction.\nc\nRevision: The following paragraphs have been modified in the methods.\nTi-ion displacement analysis.\nTi off-center displacements were quantified from atomic-level HAADF-STEM images by\nmeasuring the offset between each Ti atomic column and the geometric center of its four nearest\nPb neighbors. Atomic column positions were determined using Gaussian peak fitting, and the\nout-of-plane displacement component (δ ) was extracted to evaluate local polarization.\nc\nDisplacement statistics were collected over multiple unit cells to analyze polarization variation\nacross the film thickness.\nQuestion 3: In figure 1 e,f the authors show GPA-based strain maps. Could the authors include\nwhere the 0-strain reference for these maps was chosen (I assume in the substrate).\nResponse: Thank you for your constructive comments. In the GPA results shown in the original\nmanuscript, due to the default calculation mode of the data-processing software\n(DigitalMicrograph), the overall average lattice of the selected area was used as the strain\nreference, which led to an offset in the strain maps. To more rigorously reflect the physical\nstrain state, we provided the GPA data by taking the NSTO substrate as the zero-strain reference.\nThe corresponding results are shown in Fig. R2.\nFig. R2 | a,b, GPA-based c-axis strain mapping at the interface for (a) pure PTO and (b) Ta-PTO films.\nQuestion 4: Also in figure 1 e, from the position of the arrow indicating the interface with the\nsubstrate in this map, it appears that the compressive strain is primarily located in the substrate,\nwhile the EDS lineprofiles provided in Figures 2b and S6 show the tantalum is incorporated in\nthe first layers of the PTO film. Is the interface marked incorrectly in the strain map, or does\nthe incorporation of Ta really compressively strain the STO substrate?\nResponse: Thank you for your constructive comments. The interface in the strain map is\nmarked correctly. Consistent with the EDS line profiles in Figs. 2b and S6, Ta is incorporated\nwithin a few unit cells on the PTO side adjacent to the interface, and there is no evidence that\nthe STO substrate is compressively strained.\nQuantitative analysis of unit cell parameters from atomic-level HAADF-STEM images\nconfirms that the NSTO substrate remains unstrained near the interface, and that the\ncompressive strain is confined to the interfacial Ta-PTO layers (Fig. 1d). The strain field\nextracted by GPA is the result of Fourier filtering and averaging on a limited spatial scale.\nWhen the substrate is used as a reference, a mask of appropriate size can better track the phase\ninformation of the NSTO region. However, near the abrupt interface, the reduction of the c-\naxis spacing in the Ta-rich PTO layer makes its lattice frequency closer to that of the substrate,\ncausing some PTO signals to be included in the same phase field. This amplifies the contrast\nof compressive and tensile strains at the interface and shifts the apparent strain extreme\npositions towards the substrate side by approximately 1–2 unit cells.\nQuestion 5: The authors mention, that the substitution of Ti4+ with Ta5+ induces the formation\nof oxygen vacancies. Do they have an explanation for the mechanism that leads to this or\nprovide a reference that explains it?\nResponse: Thank you for your constructive comments. Ta substitution does not intrinsically\ninduce oxygen vacancies. From a defect-thermodynamic perspective, Ta preferentially\nsubstitutes on the Ti site as a donor, contributes extra electrons and shifts the Fermi level\nupward toward the conduction band, making PTO more n-type2. Under donor-induced n-type\nconditions, the formation energy of positively charged oxygen vacancies increases with the\nFermi level. Hence, this leads to a thermodynamic suppression of oxygen vacancy formation,\nwhich has been verified experimentally in Ta-doped PTO with a reduced oxygen-vacancy\nconcentration3.\nHowever, First-principles studies have investigated that oxygen-vacancy formation energies\nare sensitive to an applied strain, with both the sign and magnitude of the effect governed by\nthe vacancy formation volume and elastic coupling4. In our case, the Ta-induced interfacial\nlattice strain gradient is proposed to drive local oxygen-vacancy formation near the interface.\nAdditionally, the substitution of Ti4+ by Ta5+ in PTO must satisfy charge neutrality. Atomic-\nlevel HAADF-STEM analysis near the interface (Fig. R3) shows a clear reduction of the Ti-\ncolumn intensity within several unit cells. This suggests that the positive charges introduced\nby Ta substitution and oxygen vacancies are mainly compensated by negatively charged Ti\nvacancies in the interfacial region. This observation is consistent with first-principles\ncalculations for PTO, which show that under n-type (high-E ) conditions Ti vacancies are much\nF\nmore stable than Pb vacancies5.\nFig. R3 | a, Atomic-level cross-sectional HAADF-STEM image of Ta-PTO/NSTO interface with a yellow\nrectangular box indicating the scan region. b, Intensity profiles of Pb/Sr columns across the interface. c,\nIntensity profiles of Ti column across the interface showing the transition between Ta-PTO and NSTO layers.\nQuestion 6: The presence of oxygen vacancies at the substrate interface is supported by EELS-\ndata of the O-K edge in Fig. S7. Could the authors provide the ratio of intensity between peaks\n1/2 to highlight the changes in the O-K edge with distance to the interface, as they are a bit\ndifficult to make out from the line-plot alone.\nResponse: Thank you for your constructive comments. We further performed a quantitative\nanalysis of the integrated intensity ratio of the Peak 1 to the Peak 2 in the O-K edge EELS\nspectra (Fig. R4b) profile within the 10 nm region on the Ta-PTO side of the interface (Fig.\nR4a), aiming to more directly demonstrate the local oxygen vacancy concentration variation\nnear the interface. As shown in Fig. R4c, the integrated intensity ratio is measured to be 0.877\n± 0.005 at the interface, and then progressively increases over a distance of approximately 4\nnm, stabilizing at a value of 0.934 ± 0.006. These results are consistent with the trend observed\nin Fig. R4b and confirm that Ta-doping-induced oxygen vacancies are predominantly confined\nwithin ~3 nm of the Ta-PTO/NSTO interface.\nFig. R4 | a, Atomic-level HAADF-STEM image of Ta-PTO/NSTO interface with a yellow dashed scan\nregion. b, The EELS result showing the O-K edge spectra acquired along the yellow dashed line in (a). c,\nThe integrated intensity ratio of the EELS O-K edge spectral peak 1 to peak 2 within a 10 nm range on the\nTa-PTO side of the interface.\nRevision: Supplementary Fig. 8 has been updated to include this quantitative intensity-\nratio analysis to improve clarity and readability.\nQuestion 7: In lines 194-195 the authors state \"polarization screening from downward to\nupward switching can be realized...\". I believe what they meant to say is \"charge screening of\nthe upward facing polarization can be realized/is realized by\", as they do not perform\npolarization switching experiments to my understanding.\nResponse: We thank the reviewer for the careful reading and agree that the sentence in the\noriginal manuscript could be misleading. We attempted to explain that compared with pure\nPTO the Ta-doped structure creates a different charge-screening environment that stabilizes the\nupward polarization state (the reversed polarization orientation) due to interfacial oxygen\nvacancy and surface Pb vacancy. Therefore, we have revised the text by replacing\n“polarization screening from downward to upward switching” with “charge screening\nthat stabilizes the upward polarization”, and we have removed the word “switching” to\navoid confusion.\nRevision: The following sentences have been modified in the manuscript at line 201–204.\nTherefore, in single-domain Ta-PTO thin films, polarization screening from downward to\nupward switching can be realized by negatively charged Pb vacancies and positively charged\noxygen vacancies the polarization screening by negatively charged Pb vacancies and positively\ncharged oxygen vacancies stabilizes the upward polarization, as shown in Fig. 2f.\nQuestion 8: In order to PV performance of the PTO films, the authors have to cap their films\nwith a metal (here Ag) electrode. Have the authors considered that the introduction of this\nadditional metal-interface could influence the polarization state of the ferroelectric below?\nHow would this interface influence the presence/formation of lead vacancies at the surface?\nResponse: Thank you for your constructive comments. We clarify that the polarization state in\nTa-PTO films is not affected by the metal electrode.\nFirst, the polarization orientation is confirmed before metal deposition by PFM on the free\nsurface. As shown in Fig. R5a–c, the written regions show distinct phase contrast consistent\nwith polarization reversal (Fig. R5c) after applying programmed bias patterns (Fig. R5b).\nNotably, the region poled with a negative bias at the center showed the same phase state as the\nunpoled background, indicating that the as-prepared Ta-PTO film is initially in a single-domain\nstate with an upward polarization. The focused ion beam (FIB) sample was protected by an\nAg/W multilayer before processing, HAADF-STEM reveals downward Ti off-centering\ndirection (Fig. R5d), demonstrating that the polarization direction remains upward below the\nmetal cap region. This provides direct evidence that the polarization state is not affected by the\nadditional metal interface.\nSecond, surface sensitive XPS shows that a Pb vacancy rich surface already exists in the as-\ngrown Ta-PTO film (Supplementary Fig. 9). In addition, we performed atomic-level HAADF-\nSTEM analysis on samples prepared with a protective metal capping layer and compared Ta-\nPTO with pure PTO (Fig. R6). A clear reduction of the Pb column intensity is observed near\nthe Ta-PTO surface, whereas the pure PTO sample remains nearly unchanged. This indicates\nthat the Pb vacancies at surface is an intrinsic feature of Ta-PTO formed during growth, and\nwill not be affected by additional metal-interface.\nThird, it is unlikely that metal deposition introduces additional Pb vacancies. Pb vacancies\nexhibit a much higher formation energy (~6.1 eV) than that of oxygen vacancies (~2.9 eV), and\nare therefore not expected to form during conventional Ag sputtering or evaporation6.\nConsistently, sputtering processes in oxides are well known to favor oxygen vacancy formation\nover A-site cation loss7.\nFig. R5 | Determining the polarization direction of Ta-PTO thin film. a-c, PFM phase maps before and after\nlocal electrical poling of Ta-PTO film. (a) Initial out-of-plane PFM phase image. (b) Applied bias writing\npattern (outer frame: 0 V; center square: -60 V; intermediate region: +60 V). (c) PFM phase image acquired\nimmediately after bias writing. d, Atomic-level HAADF-STEM image of Ta-PTO showing upward\npolarization.\nFig. R6 | Surface HAADF-STEM and Pb/Ti column-intensity profiles of pure PTO and Ta-PTO after FIB\nsample preparation with a protective metal capping layer. a,d, HAADF-STEM image of (a) pure PTO and\n(d) Ta-PTO surface region with a yellow rectangular box indicating the scan region. b,e, Intensity profiles\nof Pb column near the (b) pure PTO (e)Ta-PTO surface. c,f, Intensity profiles of Ti column near the surface\nof (c) pure PTO (f) Ta-PTO.\nReviewer #2:\nQuestion 1: The fabrication method given in the supplementary section is not very clear. What\ndo they mean by single-crystal ferroelectric film? Is it epitaxy? If so, further measurements like\nreciprocal space mapping have to be given.\nResponse: We thank the reviewer for pointing out that the fabrication description in the\nSupplementary Methods section was not sufficiently clear. We have rewritten the “Film\npreparation” sections to provide a complete description of the hydrothermal solution-epitaxial\ngrowth procedure (precursor formulation, mineralizer, substrate cleaning, growth temperature\nand time, and post-treatment), replacing the previous abbreviated text. In addition, we have\nadded a schematic workflow diagram (Fig. R7) in the method to illustrate the full\nhydrothermal growth process.\nFig. R7 | Schematic workflow of the hydrothermal solution-epitaxial growth of Ta-PTO thin films on NSTO\nsubstrates.\nIn the revised manuscript, we explain that the term “single-crystal ferroelectric film” refers to\na single-crystalline epitaxial film grown on a single-crystal substrate. Specifically, the film\ndisplays a clearly defined in-plane epitaxial registry with the substrate. As suggested, we have\nadded a reciprocal space mapping (RSM) analysis. The RSM around the asymmetric (103)\nreflection (Fig. R8a) shows a single Ta-PTO (103) diffraction peak together with the NSTO\n(103) substrate peak, and the film peak aligns in the in-plane reciprocal coordinate with the\nsubstrate, supporting in-plane epitaxy8. The -scan (Fig. R8b) exhibits the expected four-fold\nsymmetry, consistent with the tetragonal structure of PTO (space group P4mm) and its in-plane\n𝜙\nepitaxial alignment with the perovskite substrate9. The rocking curve (Fig. R78) shows a\nnarrow full width at half maximum (FWHM), indicating the high crystalline quality of the Ta-\nPTO film. These results support the single-crystal Ta-PTO film.\nIn the revised manuscript, the main text has been updated to include the RSM analysis,\nand the corresponding supplementary figures have been renumbered accordingly. This\nstructural evidence provides solid evidence to support the single-crystalline epitaxial\nnature of the Ta-PTO film.\nFig. R8 | Structural characterization of epitaxial Ta-PTO film on NSTO substrate using off-axis X-ray\ndiffraction. a, X-ray RSM study about the pseudocubic 103-diffraction condition for Ta-PTO/NSTO. b, phi\n( ) scan spectrum of Ta-PTO film along the pseudocubic <101> direction. c, X-ray rocking curves of the\n(𝜙002) plane for Ta-PTO film.\nRevision:\nThe following paragraphs have been modified in the Film preparation section of the\nmethod.\nAs shown in the Supplementary Fig. 18, PTO and Ta-PTO thin films were synthesized by a\nhydrothermal solution-epitaxy method. Pb(NO ) , TBOT and Ta O were used as metal\n3 2 2 5\nprecursors, and KOH served as the mineralizer. NSTO single-crystal substrates were\nultrasonically cleaned in acetone, ethanol, deionized water, and ethanol sequentially, and then\nplaced in a Teflon holder for drying.\nIn a Teflon liner, 15 mL of deionized water was first added. KOH (9.54 g) was dissolved to\nprepare a solution with a total precursor concentration of 6.8 M. Subsequently, 1.708 mL of\nTBOT was added under vigorous stirring. For Ta-PTO growth, Ta O was introduced at 0.11 g,\n2 5\ncorresponding to a Ta substitution level of 10 mol% (Ta/(Ta+Ti) = 10%). The Pb/Ti molar ratio\nwas maintained at 1.25:1 to compensate for Pb volatility. Accordingly, Pb(NO ) (2.04 g) was\n3 2\ndissolved separately in 10 mL of deionized water and then added dropwise into the precursor\nsolution under continuous stirring.\nAfter stirring for 2 h at room temperature, the NSTO substrate was placed horizontally in the\nprecursor solution, positioned approximately 15 mm above the bottom of the Teflon liner. The\nTeflon liner was then sealed inside a stainless-steel autoclave, and the hydrothermal epitaxial\ngrowth was carried out at 200 °C for 12 h. After the reaction, the samples were thoroughly\ncleaned by repeated ultrasonication in deionized water and ethanol, and then dried in air at\n60 °C.\nThe following sentences have been modified in the manuscript at line 126–130.\nX-ray diffraction (XRD) analysis confirmed that both films exhibit a tetragonal perovskite\nstructure (JCPDS#70-0746) (Supplementary Fig. 1a). As shown in Supplementary Fig. 1b, a\nsingle Ta-PTO (103) diffraction peak is observed together with the NSTO (103) substrate peak,\nindicating an epitaxial growth of the Ta-PTO film on the NSTO substrate. Additionally, scan\nrevealed four peaks at 90° intervals, indicating the single-crystal nature of the Ta-PTO film\n𝜙\n(Supplementary Fig. 1c), which also has a smooth and continuous surface (Supplementary Fig.\n1d).\nQuestion 2: If I understood correctly, they have doped Ta into the PbTiO .\ndoped throughout the film. In that case, what do they mean by Ta interfacial doping? The entire\nmechanism explained in this manuscript is based on the Ta at the interface. However, it is very\ndifficult to follow the formation of Ta doping only at the interface.\nResponse: Thank you for your constructive comments. We clarify that the doping occurred at\nthe interface rather than homogeneous bulk doping of the PTO film based on our experimental\nresults. Firstly, the cross-sectional STEM-EDS line scan (Fig. R9b) reveals a clear Ta signal\nconfined to the interfacial region, with Ta intensity close to the background level away from\nthe interface. Second, the TOF-SIMS depth profile (Fig. R9c) shows a sharp Ta peak precisely\nat the interface, while the Ta signal remains close to the detection limit throughout the bulk of\nthe PTO film. Therefore, by “Ta interfacial doping,” we specifically refer to the formation of a\nTa-rich interfacial layer with a thickness of only a few unit cells, instead of Ta doping\nthroughout PTO film.\nFig. R9 | Atomic-scale evidence for interfacial Ta doping in Ta-PTO film. a, Atomic-level cross-sectional\nHAADF-STEM image of the Ta-PTO/NSTO interface. b, Elemental linear scanning profiles of Ti and Ta\nelement throughout Ta-PTO/NSTO interface. c, TOF-SIMS for the depth profiling of Ta-PTO film.\nBased on Fig. R10, we attribute the interfacial Ta incorporation to a kinetically driven\ncompetitive growth process during solution epitaxy. In the early stage (≈6 h), the solution is\nstill Ta-rich and the film is in its initial nucleation stage, so Ta can be easily incorporated into\nthe bottom layers that form first, creating a Ta-rich interfacial region. As the reaction progresses,\nlarge numbers of Ta-rich particulate phases appear. These particles effectively consume the\navailable Ta species in the solution. EDS statistics show that the Ta fraction in the Ta-rich\nparticles increases sharply with reaction time. When a reaction time is above 6h, the Ta species\nare consumed by the particulate phase, which strongly limits their availability for further\nincorporation into the growing film10,11. As a result, the film naturally forms a Ta-rich\ninterfacial region, rather than Ta doping throughout PTO film.\nFig. R10 | Competitive growth kinetically drives the emergence of interfacial Ta doping. a-h, SEM images\nof Ta-PTO films and particles obtained in the same autoclaves by hydrothermal synthesis after (a, e) 4 h, (b,\nf) 6 h, (c, d) 8 h, (d, h) 12 h, respectively. i-k, Schematic illustrations of the growth process. l, Ta/(Ta + Pb +\nTi) molar ratio of Ta-PTO particles spontaneously synthesized in the same autoclaves, measured by EDS\nafter different hydrothermal reaction times.\nQuestion 3: Again, the Pb vacancies at the top surface are not clearly brought out in this\nmanuscript, which is also crucial to understand the mechanism proposed in this manuscript.\nResponse: Thank you for your constructive comments. In our low-temperature solution epitaxy,\nthe Ta-PTO film has a single-domain state with upward polarization, which creates a positively\npolarized top surface. Such a surface should be compensated by negative charges to stabilize\nthe single-domain configuration. These negative charges can be effectively provided by surface\nPb vacancies, which serve as an efficient screening mechanism for polarization-bound charges.\nThis screening reduces the electrostatic energy of the polarized surface and is therefore\nthermodynamically favorable for stabilizing the single-domain configuration, as previously\nreported for ferroelectric films12.\nThe existing Pb vacancies have been well established by the following solid evidences. Our\nHAADF-STEM analysis shows a clear reduction of the Pb signal near the top surface (Fig.\nR11b). Consistently, the Pb atomic columns intensity decreases markedly toward the surface,\nwhereas the Ti-column intensity remains comparatively stable (Fig. R11d,e). In addition,\nsurface-sensitive XPS further supports Pb deficiency at the film surface, the measured Pb/Ti\nratio is lower for Ta-PTO than for pure PTO (Supplementary Fig. 9), consistent with the\npresence of a Pb vacancies.\nFig. R11 | Atomic-scale evidence for a Pb vacancy rich surface region in Ta-PTO. a, Atomic-level cross-\nsectional HAADF-STEM image of the Ta-PTO surface region. b, Elemental linear scanning profiles of Ti\nand Pb element near the Ta-PTO surface. c, HAADF-STEM image of Ta-PTO surface region with a yellow\nrectangular box indicating the scan region. d, Intensity profiles of Pb column near the surface, showing a\nreduction in signal. e, Intensity profiles of Ti column near the surface of Ta-PTO.\nQuestion 4: Overall, the formation of p-i-n structure is difficult to understand based on the\nfabrication details given in this manuscript.\nResponse: Thank you for your constructive comments. In our work, the p-i-n structure is\nconstructed by the combined effects of ferroelectric polarization screening and aliovalent\ndoping, which has been realized by Pb vacancies near the surface and Ta substitution near the\ninterface, respectively (Fig. R12).\nFig. R12 | Schematic illustration of the defect induced p-i-n band structure in Ta-PTO.\nFirst, the n-type interfacial region originates from a Ta-rich donor layer localized at the Ta-\nPTO/NSTO interface. This is directly evidenced by cross-sectional STEM-EDS line scans and\nTOF-SIMS depth profiles, both of which show pronounced Ta enrichment confined to the\ninterfacial region (Fig. R9). Since Ta substitutes on the Ti site as a donor, this interfacial layer\nnaturally exhibits n-type electronic character.\nSecond, the p-type near-surface region is associated with a Pb vacancy rich top layer formed\nduring film growth. Atomic-level HAADF-STEM images, Pb-column intensity statistics and\nsurface-sensitive XPS measurements reveal a clear reduction of Pb occupancy near the film\nsurface (Fig. R11). The presence of this Pb vacancy rich layer modifies the surface electronic\nstructure and work function of PTO, which is independently confirmed by Kelvin probe force\nmicroscopy (KPFM) measurements (Fig. R13). Under identical KPFM conditions using the\nsame probe (Fig. R10b,d), Ta-PTO shows a higher surface potential than pure PTO (average\n~0.951 V vs ~0.868 V). This indicates a higher surface work function. Since ,\nvac\nthe higher Φ implies a lower Fermi level closer to the valence band, consistent\nΦ\nwi\n=\nth\n𝐸\na mo\n−\nre\n𝐸\np𝜹-\ntype surface state (Fig.R13e)13.\nBetween the donor-rich interface and the acceptor-rich surface, the film interior behaves as an\nintrinsic region. The coexistence of these layers results in band bending across the film and\nforms a p-i-n structure. This internal field efficiently separates carriers and attributed to the\nenhanced photovoltaic response in Ta-PTO devices.\nFig. R13 | KPFM characterization and work-function schematic for Ta-PTO and pure PTO. a,c, AFM\ntopography and b,d, KPFM surface potential maps of Ta-PTO (top row) and pure PTO (bottom row). e,\nSchematic band diagram illustrating the work-function difference.\nRevision: We have revised the manuscript by adding a KPFM work-function analysis of\nthe surface (Supplementary Fig. 11) to the main text.\nQuestion 5: What is the role of ferroelectricity in the observed PV response? A lot has been\ndiscussed in the introduction related to the ferroelectric photovoltaic effect. But the proof for\nferroelectric PV is still not clear in this manuscript.\nResponse: Thank you for your constructive comments. As discussed in Question 3 and 4, we\nargue that the ultrahigh photocurrent reported here in Ta-PTO film originates from a p-i-n\nstructure constructed by ferroelectric polarization screening, rather than from the conventional\nferroelectric photovoltaic effect. Notably, this unique architecture is realized in a single growth\nstep in solution. Such p-i-n configuration induced a strong built-in electric field (~12 kV·cm-1,\nestimated from /d) that dominates carrier separation and collection, yielding a short-circuit\nOC\ncurrent density far exceeding that of conventional ferroelectric photovoltaic systems ( 130\n𝑉\nmA·cm-2 for Ta-PTO vs 4.03 mA·cm-2 for pure PTO). In contrast, pure PTO film exhibits a\n∼\nphotovoltaic response mainly governed by the bulk ferroelectric photovoltaic effect, where the\n∼\nphotocurrent direction can be controlled by switching polarization direction10.\nQuestion 6: Proof for the ferroelectric nature of the grown film, like P-E measurement, domain\nswitching studied by PFM, etc, which are essential to ascertain the ferroelectric nature of the\nfilm, are missing in this manuscript.\nResponse: Thank you for your constructive comments. We agree that a direct experimental\nverification of the ferroelectric nature of the grown Ta-PTO film is essential. We attempted to\nmeasure the polarization-electric field (P-E) loop using conventional macroscopic methods\nwith regular top electrodes. However, the film tends to be broken down below the coercive\nelectric field (E ≈110 kV cm-1)14. This could be attributed to a built-in electric field within the\nc\nfilms due to the p-i-n structure with negative and positive screening defects, making it very\ndifficult to obtain the desired P-E loop in Ta-PTO film.\nTo address this issue, we performed local Switching Spectroscopy Piezoresponse Force\nMicroscopy (SS-PFM) measurements (maximum applied voltage ~150 V). As shown in Fig.\nR14, the PFM phase exhibits a clear ~180° reversal with bias, accompanied by the\ncharacteristic butterfly-like evolution in the amplitude. The positive and negative switching\nvoltages are ~ +45.4 V (+413 kV cm-1) and ~ −38.5 V (−350 kV cm-1), respectively. These\nfields are higher than the typical E of PTO.\nc\nThe asymmetric switching voltages further support\nthe built-in electric field. Furthermore, we conducted PFM domain-writing experiments (Fig.\nR15). After applying programmed bias patterns, the written regions show distinct phase\ncontrast consistent with polarization reversal. Notably, the region poled with a negative bias at\nthe center showed the same phase state as the unpoled background, indicating that the Ta-PTO\nfilm is initially in a single-domain state with an upward polarization. These results are fully\nconsistent with the polarization direction independently inferred from atomic-scale Ti off-\ncentering analysis.\nFig. R14 | a,b, phase-voltage(a) and amplitude-voltage(b) curves of Ta-PTO film.\nFig. R15 | PFM phase maps before and after local electrical poling of Ta-PTO film. a, Initial out-of-plane\nPFM phase image. b, Applied bias writing pattern (outer frame: 0 V; center square: -60 V; intermediate\nregion: +60 V). c, PFM phase image acquired after bias writing.\nRevision: In the revised manuscript, we replaced the original single PFM phase image\n(previously Supplementary Fig. 1d) with a complete set of PFM phase maps and domain-\nwriting results (now as Supplementary Fig. 4). These results more convincingly\ndemonstrate the single-domain nature and switchable ferroelectric polarization of the Ta-\nPTO film. The following sentence has been modified in the manuscript at line 139–142.\nPiezoelectric force microscopy (PFM) confirms a single-domain structure of Ta-PTO film with\ntypical ferroelectric and piezoelectric characteristics (Supplementary Fig. 1d) that the Ta-PTO\nfilm exhibits typical ferroelectric characteristics with a single-domain configuration\n(Supplementary Fig. 4).\nQuestion 7: In case the observed PV effect is purely based on the p-i-n effect, then the\nvalue of the short-circuit current should not be compared to the literature value on the\nferroelectric photovoltaic system.\nResponse: Thank you for your constructive comments. We agree with the reviewer that a direct\ncomparison between our results and conventional ferroelectric photovoltaic current could be\nnot appropriate. Accordingly, we have added more representative p-n and p-i-n structured UV\nphotodetectors as a comparison, including both ferroelectric-based junction devices and\nconventional wide-bandgap semiconductors (e.g., GaN, Ga O , SiC) at zero bias responsivity.\n2 3\nAs summarized in table R1, most self-powered p-n or p-i-n UV photodetectors exhibit\nresponsivities below ~0.3 A W-1, and the corresponding value of ferroelectric junction is much\nlower due to a low conductivity. In contrast, the Ta-PTO device in this work achieves a\nresponsivity of ~0.51 A W-1 at 375 nm at zero bias, exceeding both conventional semiconductor\npn detectors and previously reported ferroelectric-based devices under similar conditions.\nTable. R1 | Comparison of UV Photodetectors Based on p-n / p-i-n Structures\nMaterial / Device Structure Bias Wavelength Responsivity Reference\n(nm) (A W-1)\nTa-PTO (p-i-n) self-powered 375 0.508 This work\nTa-PTO (p-i-n) self-powered 254 0.148 This work\nBaTiO /p-Si self-powered 365 0.027 Han et al., Appl. Surf. Sci.,\n3\n(ferroelectric/semiconductor p-n) 202315\nBaTiO @MXene/4H-SiC self-powered 310 0.188 Sha et al., ACS Photonics,\n3\n(ferroelectric/2D p-n) 202416\nBiOCl nanosheets/TiO film self-powered 340 0.17 Zhang et al., Appl. Surf. Sci.,\n2\n(thin film p-n) 202217\np-CuO/n-β-Ga O oxide self-powered 254 0.0303 Park et al., Nanomaterials,\n2 3\n(oxide p-n) 202318\nPANI/Mg Zn O bilayer self-powered 250 1.6×10-4 Chen et al., Small, 201619\n0.31 0.69\n(organic/inorganic p-n)\nAu/ZnO/PLZT/FTO self-powered 360 3.96×10-3 Chen et al., Adv. Electron.\n(oxide/ferroelectric p-n) Mater., 202120\nNiO/β-Ga O self-powered 254 1.25×10-3 Tang et al., Physica B:\n2 3\n(oxide p-n) Condens. Matter, 202521\np-GaN/n-β-Ga O self-powered 365 5.44×10-2 Li et al., J. Mater. Chem. C.,\n2 3\n(oxide p-n) 201722\nSc Al N self-powered 193 0.015 Mondal et al., APL Mater.,\nx 1−x\n(thin film) 202223\nQuestion 8: If it is due to the ferroelectric PV effect, proof for the switchable PV effect should\nbe given, which is missing in this manuscript.\nResponse: Thank you for your constructive comments. We agree that a switchable photovoltaic\nresponse is a key signature when the measured photocurrent is dominated by the intrinsic\nferroelectric photovoltaic effect, such as bulk photovoltaic effect contributions. In our device,\nhowever, the ultrahigh zero-bias photocurrent is primarily governed by an internally formed p-\ni-n junction field extending across the film thickness, rather than by a polarization controlled\nferroelectric photovoltaic effect.\nWe also note that polarization reversal in Ta-PTO is very challenging. The presence of detective\nlayers and the p-i-n structure creates a strong built-in electric field where a stable and single-\ndomain electrostatic boundary condition can be achieved through the polarization screening\nduring the growth. Once this structure was formed, an external poling suffers from\nreconfiguring the defects distributed along the film thickness direction. This makes the\nswitching of the polarization very difficult and thus reversing the photovoltaic response.\nQuestion 9: In addition, some of the sentences are not clear. Examples are given below\n“The photovoltaic (PV) effect for semiconductors affords versatile applications from solar\nCells…..”\nResponse: We agree that this sentence is not clear. We have rewritten it to make the general\nrelevance of semiconductor PV effects clearer and improve readability.\nRevision: The following sentence has been modified in the manuscript at line 57–60.\nThe photovoltaic (PV) effect for semiconductors affords versatile applications from solar cells,\nphoto-electro-catalysis to optoelectronic detection\nThe photovoltaic (PV) effect in semiconductors is the foundation of many technologies, such\nas solar energy conversion, photocatalysis and optoelectronic sensing.\n“Piezoelectric force microscopy (PFM) confirms a single-domain structure of Ta-PTO film\nwith typical ferroelectric and piezoelectric characteristics (Supplementary Fig. 1d).” Contrary\nto this statement, the figure does not give any details about the piezoelectric characteristics.\nResponse: We agree that such sentence is not suitable. The PFM data was provided in\nSupplementary Fig. 4, supporting a single-domain polarization configuration of Ta-PTO film\nby out-of-plane phase contrast. We have therefore revised the text to accurately describe the\nPFM measurements and removed the potentially misleading phrasing.\nRevision: The following sentence has been modified in the manuscript at line 139–142.\nPiezoelectric force microscopy (PFM) confirms a single-domain structure of Ta-PTO film with\ntypical ferroelectric and piezoelectric characteristics (Supplementary Fig. 1d) that the Ta-PTO\nfilm exhibits typical ferroelectric characteristics with a single-domain configuration\n(Supplementary Fig. 4).\n“ At the nanoscale, a large stress gradient (e.g., with the force applied through an AFM tip)\ncan induce a local electric field through a flexoelectric effect, which is capable to switch the\nlocal polarization”. This is confusing. Where from the AFM tip induced flexoelectric field is\ncoming in this study? The flexoelectric field could be the compressive strain gradient at the\ninterface.\nResponse: We agree that the original sentence could be confusing. AFM-tip-induced strain\ngradients are well known to generate flexoelectric fields and can even directly control\nferroelectric polarization at the nanoscale, as reported in the literature. This example was cited\nonly for illustration and does not represent the mechanism operative in the present work24,25.\nHere, the flexoelectric field arises from the intrinsic strain gradient at the interface in Ta-\nPTO(Fig. 1h). The text has been revised accordingly.\nRevision: The following sentence has been modified in the manuscript at line 155–160.\nAt the nanoscale, a large stress gradient (e.g., with the force applied through an AFM tip) can\ninduce a local electric field through a flexoelectric effect, which is capable to switch the local\npolarization can induce a local electric field via the flexoelectric effect. Such gradients can be\ngenerated artificially, for example by localized mechanical loading with an AFM tip, as widely\nreported to induce flexoelectric fields and control ferroelectric polarization.\n“The results revealed that Ta substitution for Ti (Ta ·) is more easily formed acting as an\nTi\neffective donor, while Pb vacancies behave as an acceptor.” Where is the proof?\nResponse: Fig. R14 directly supports the donor and acceptor assignment from defect-\nthermodynamic signatures. In a defect-formation-energy plot, the slope with respect to E\nF\nequals the defect charge state q. The Ta branch shows a positive slope and is stable in a\nTi\npositively charged state over a wide E range (donor-like behavior), whereas the lowest-energy\nF\nV branch shows a negative slope, indicating stability in negatively charged states (acceptor-\nPb\nlike behavior). Therefore, Ta is classified as a donor and Pb vacancies as an acceptor.\nTi\nFig. R14 | Calculated via DFT, formation energies of lead vacancies (V ) and tantalum substituents at Pb\nPb\n(Ta ) and Ti (Ta ) sites under oxygen-rich and metal-poor conditions are shown as a function of the Fermi\nPb Ti\nlevel, with slope changes marking transitions between charge states, denoted as \"+\" (electron-deficient), \"–\n\" (electron-rich), or \"0\" (neutral).\nReferences\n1. Jia, C.-L. et al. Unit-cell scale mapping of ferroelectricity and tetragonality in epitaxial ultrathin\nferroelectric films. Nat. Mater. 6, 64–69 (2007).\n2. Nazir, S. & Tariq, A. Thermodynamics and multiferroicity in PbTiO due to 4/5d electrons doping.\n3. Qi, H. et al. Reducing the vacancies associated with ferroelectric polarization to promote\nphotocatalytic overall water splitting. Sci. China: Chem. 67, 3258–3264 (2024).\n4. Mayeshiba, T. & Morgan, D. Strain effects on oxygen vacancy formation energy in perovskites. Solid\nState Ionics 311, 105–117 (2017).\n5. Yao, Y. & Fu, H. Charged vacancies in ferroelectric PbTiO : formation energies, optimal fermi region,\n3\nand influence on local polarization. Phys. Rev. B 84, 64112 (2011).\n6. Shimada, T., Ueda, T., Wang, J. & Kitamura, T. Hybrid hartree-fock density functional study of\ncharged point defects in ferroelectric PbTiO . Phys. Rev. B 87, 174111 (2013).\n3\n7. Kim, J.-Y. et al. Reducing the oxygen vacancy concentration in SrTiO thin films via an optimized\n3-δ\nO plasma treatment for enhancing device properties. Appl. Surf. Sci. 639, 158271 (2023).\n2\n8. Rispens, G. & Noheda, B. Ultra-thin lead titanate films grown by molecular beam epitaxy. Integrated\nFerroelectrics 92, 30–39 (2007).\n9. Sun, L. et al. X-ray analysis of epitaxial thin film grown on an (001) substrate by metal-organic\nchemical vapour deposition. J. Phys.: Condens. Matter 8, 10185 (1996).\n10. Lin, C. et al. Solution epitaxy of polarization-gradient ferroelectric oxide films with colossal\nphotovoltaic current. Nat. Commun. 14, 2341 (2023).\n11. Zhang, R. et al. Compositionally-graded ferroelectric thin films by solution epitaxy produce\nexcellent dielectric stability. Nat. Commun. 16, 98 (2025).\n12. Gao, P. et al. Atomic mechanism of polarization-controlled surface reconstruction in ferroelectric\nthin films. Nat. Commun. 7, 11318 (2016).\n13. Kahn, A. Fermi level, work function and vacuum level. Mater. Horiz. 3, 7–10 (2016).\n14. Nishino, R., Fujita, T. C., Kagawa, F. & Kawasaki, M. Evolution of ferroelectricity in ultrathin\nPbTiO films as revealed by electric double layer gating. Sci. Rep. 10, 10864 (2020).\n3\n15. Han, W. et al. High-performance self-powered amorphous-BaTiO /p-Si heterojunction\n3\nphotodetector controlled by ferroelectric effect. Appl. Surf. Sci. 615, 156371 (2023).\n16. Sha, S. et al. Ferro-pyro-phototronic effect enhanced self-powered UV photodetectors based on\nBaTiO @MXene/4H-SiC heterojunction. ACS Photonics 11, 4472–4485 (2024).\n3\n17. Zhang, Q. et al. High-performance self-powered ultraviolet photodetector based on BiOCl/TiO\n2\nheterojunctions: carrier engineering of TiO . Appl. Surf. Sci. 592, 153350 (2022).\n2\n18. Park, S. et al. A self-powered high-responsivity, fast-response-speed solar-blind ultraviolet\nphotodetector based on CuO/β-Ga O heterojunction with built-In potential control. Nanomaterials\n2 3\n13, (2023).\n19. Chen, H. et al. Ultrasensitive self-powered solar-blind deep-ultraviolet photodetector based on all-\nsolid-state polyaniline/MgZnO bilayer. Small 12, 5809–5816 (2016).\n20. Chen, J. et al. High-performance self-powered ultraviolet photodetector based on coupled\nferroelectric depolarization field and heterojunction built-In potential. Adv. Electron. Mater. 7,\n2100717 (2021).\n21. Tang, K. et al. NiO/Ga O heterojunction for transparent and flexible solar-blind UV photodetector\n2 3\napplication. Physica B 713, 417366 (2025).\n22. Li, P. et al. Construction of GaN/Ga2O p–n junction for an extremely high responsivity self-\n3\npowered UV photodetector. J. Mater. Chem. C 5, 10562–10570 (2017).\n23. Mondal, S. et al. Reconfigurable self-powered deep UV photodetectors based on ultrawide\nbandgap ferroelectric ScAlN. APL Mater. 10, 121101 (2022).\n24. Lu, H. et al. Mechanical writing of ferroelectric polarization. Science 336, 59–61 (2012).\n25. Ahn, Y. & Son, J. Y. Flexoelectric effect via piezoresponse force microscopy of domain switching\nin epitaxial PbTiO thin films. J. Korean Ceram. Soc. 61, 55–62 (2024).\n3"}, {"round": 2, "reviewer_comments": {"reviewer_1": "The Authors have addressed my comments to my satisfaction and revised their manuscript accordingly. Therefore, I recommend the manuscript for publication in Nature Communication.", "reviewer_2": "The authors have addressed all the queries raised by the referees. I recommend the manuscript for publication."}, "author_response": "Response to Reviewers’ Comments\nWe sincerely appreciate reviewers for their constructive and positive comments, which\nundoubtedly guide us to largely improve the quality of this work.\nReviewer #1:\nThe Authors have addressed my comments to my satisfaction and revised their\nmanuscript accordingly. Therefore, I recommend the manuscript for publication in\nNature Communication.\nResponse: We are delighted that our revisions have satisfactorily addressed your\nconcerns. We sincerely thank the reviewer for the positive evaluation and kind\nrecommendation for publication.\nReviewer #2:\nThe authors have addressed all the queries raised by the referees. I recommend the\nmanuscript for publication.\nResponse: We are pleased that all the concerns raised by the referees have been\nadequately addressed. We sincerely thank the reviewer for the positive assessment and\nrecommendation for publication."}]

Decision letter