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Nature Communications volume 17, Article number: 4193 (2026)
The performance of all-perovskite tandem solar cells is critically hindered by the defective and high-roughness surfaces of lead-tin narrow-bandgap subcells, which induce non-radiative recombination and impede carrier extraction. Herein, we report a robust and multifunctional strategy to convert the above narrow-bandgap perovskite surface into an efficient and smooth one by a picosecond ultraviolet pulsed laser polishing technology combined with surface reconstruction. The polished surface is decoded as [PbI₆]⁴⁻/[SnI₆]⁴⁻ octahedral frameworks with metastable A-site vacancies. By screening guanidinium bromide as an A-site passivator, the polished surface is reconstructed into a guanidinium-cesium-based perovskite phase, substantially enhancing carrier extraction and suppressing ion migration. The resulting single-junction lead-tin and tandem solar cells, fabricated via an antisolvent-free method, achieve efficiencies of 23.47% (certified) and 29.80%, respectively, alongside exceptional operational stability. This versatile interface engineering paradigm surmounts a pivotal barrier in the advancement of next-generation photovoltaic technologies.
Tandem solar cells (TSCs) offer the potential to reach over 43% power conversion efficiency (PCE) by minimizing thermalization losses, thereby surpassing the Shockley-Queisser limit (33%) of single-junction solar cells1,2. Currently, all-perovskite TSCs have achieved 30.1%3 certified PCE in less than two decades, ascribing to unique photoelectrical properties, demonstrating as promising candidates for the next generation of photovoltaic (PV) technology.
As the bottom subcell of all-perovskite TSCs, the narrow-bandgap (NBG) perovskite (~1.25 eV) one plays the vital role in driving the efficiency of tandem cells to new heights4. The present NBG perovskite SC performance suffered from serious open-circuit voltage (VOC) deficit and relatively low fill factor (FF), primarily induced by the NBG-perovskite/C60 interface defects related non-radiative recombination5. They were introduced from asynchronous crystallization process from SnI2 and PbI2 reaction with organic halides6,7,8,9, resulting in Sn-rich surface with Sn2+ oxidation-based self-doping defects10. Moreover, for scalable commercial production technology, vacuum/gas-assisted crystallization methods for NBG perovskite fabrication feature a solvent extraction rate that is roughly four orders of magnitude slower than antisolvent-assisted quenching11, further worsening asynchronous crystallization related defects12,13. On the other hand, high-quality polycrystalline perovskite film typically shows large grains, yet inevitably parasitizes high surface roughness14, leading to critical interfacial issues. One is non-uniform contact between electrodes and perovskites, leading to local shunting risk5. The other exacerbates carrier scattering, hindering the carrier extraction10,15,16.
To address above challenges, many studies regulated the crystallization rate of Pb-based and Sn-based perovskite through solvent engineering9 or additive strategies4 to realize a synchronized crystallization process. Instead of controlling the crystallization rate, surface chemical polishing strategies have also been developed to remove severe phase segregation surface by polishing agents, such as 1,4-butanediamine17 and 1,2-diaminopropane18. However, these polishing agents have to meet the target-selective etching and avoid damage to underlying perovskite. Thus, a universal and controllable polishing strategy is essential for addressing above challenges. Different to chemical polishing methods, physical strategies have also been developed such as tape stripping19 and mechanical nano-polishing20. While all the above removal methods are in contacting style, which leads to the application limits due to the high sensitivity of perovskite films. Inspiringly, laser processing technology is characterized with non-contact style, high robustness and high precision, widely applied in polishing, additive manufacturing, annealing21,22 etc.
Herein, we firstly develop a high-resolution and robust laser polishing strategy to remove defective Pb-Sn perovskite surface. The newly exposed surface is decoded with [PbI₆]⁴⁻/[SnI₆]⁴⁻ octahedral frameworks, rich in A-site vacancies (VA). By utilizing this distinct opening platform, we further screen guanidine hydrobromide (GABr) to reconstruct the polished surface. The resulting NBG perovskite SCs achieve a PCE of 24.07% (certified 23.47%), and its TSCs obtain a top 29.80% efficiency based on antisolvent-free method. And the optimal device can retain 80% of its initial efficiency after 650 h under operational conditions.
The Pb-Sn perovskite films prepared by vacuum-driven percrystallization (VDP) technology face a defective and high-roughness surface9. The resulting interface between perovskite layer and electron-transport layer (ETL, C60) hinders interfacial charge extraction and poses local shunting risks, thereby inducing severe non-radiative recombination14 as schematically described in Fig. 1a. In order to solve the above challenges, we devised a picosecond ultraviolet pulsed laser polishing technology (PLPT) (Fig. 1b), which was expected to remove and smooth the defective surface to enhance carrier extraction with negligible thermal effect.
a, b Schematic diagram illustrating carrier extraction behavior in control (a) and PLPT-treated (b) perovskite solar cell devices. The red region represents perovskite, and the blue region represents the electron transport layer. c Schematic diagram of PLPT. d 3D-AFM images of control (top) and PLPT-treated (bottom) NBG perovskite films. Scale bar: 500 nm. e The thickness statistics of Pb-Sn perovskite films with different PLPT recipes. f Atomic ratio values of Sn, Pb, and I of control and PLPT-treated films with different polishing depths. g ToF-SIMS of control and PLPT-treated perovskite films with 50 nm polishing depth. h Sn 3d XPS spectra of control and PLPT-treated perovskite films surface with different polishing depths.
The fabrication process of Pb-Sn perovskite films was described via VDP technology as shown in Supplementary Fig. 1. The annealed VDP film was denoted as control sample. To remove the low-quality surface layer, the control film was treated by laser polishing (Fig. 1c). A picosecond ultraviolet pulsed laser was screened for present work to mitigate laser-induced damage to the interior film. The polishing recipe was optimized for power and scanning speed as shown in Supplementary Table 1.
Firstly, we investigated the evolution of the surface morphology and chemical information for the picosecond-ultraviolet-pulsed-laser-treated Pb-Sn perovskite films (PLPT). Their atomic force microscopy (AFM) images directly illustrated that the average roughness decreased from 29.8 nm to 10.2 nm (Fig. 1d). A series of characterizations and simulations confirmed that the PLPT-treated perovskite films exhibited an overall uniform and flat surface on a large scale (Supplementary Figs. 2–6). Figure 1e statistically analyzes the thickness of Pb-Sn perovskite films with different PLPT parameters. The polishing resolution and depth could be conveniently adjusted by varying the laser power and scanning speed. It has been demonstrated that present PLPT strategy was capable of achieving ~0.90 nm nanoscale-precision in a non-contact manner. Hundreds of polishing batches further verified the strong reliability (Supplementary Fig. 7) and high compatibility (Supplementary Fig. 8) of present strategy comparing with reported results19,20.
The surface and interior chemical information of control and PLPT-treated Pb-Sn perovskite samples were further investigated using X-ray photoelectron spectroscopy (XPS) and time-of-flight secondary-ion mass spectrometry (ToF-SIMS). As shown in Fig. 1f, the ratio of I/(Pb + Sn) for the top surface of the control sample is 1.64:1, obviously deviating from the stoichiometric ratio of 3:1 (dashed line). And its Sn/Pb ratio reaches ~2.9 showing a Sn-accumulated surface state much larger than stoichiometric ratio of 1:1 (dashed line). It is noted that the above two kinds of atomic ratios gradually return to its ideal values after PLPT treatment with increasing polishing depth. When polishing to a thickness more than 50 nm, both the Sn/Pb and I/(Pb + Sn) values met ideal values. The atomic proportion profiles were further investigated by ToF-SIMS. In Fig. 1g, the element content of control film surface exhibits a gradient distribution. And it gradually stabilizes at a depth of approximately 50 nm. The oxidized Sn4+ content plays crucial role in device performance, thus it was further analyzed for different polishing depth films. XPS spectra of Sn 3d of perovskite samples were shown in Fig. 1h, and the ratio of Sn4+ in the control film surface is 20.2%. This may be attributed to the accumulation of Sn2+ readily oxidized to Sn4+ 23. The surface characteristics strongly indicated the subpar surface quality of the control film with Sn-rich and I-deficient surface states, leading to the serious non-radiative recombination losses4. When precisely polishing the top surface more than 50 nm, the surface percentage of Sn4+content was significantly decreased to 5.2%. Thus, the polishing thickness was selected as 50 nm for optimal recipe.
To assess PLPT effect to device performance, Pb-Sn NBG PSCs were fabricated based on the control and PLPT-treated films. The device structure utilized a typical inverted structure. In Fig. 2a, the PLPT-treated device without post-passivation held a VOC of 0.852 (0.852) V, a short-circuit current density (JSC) of 31.72 (31.70) mA cm−2 and a FF of 80.14% (79.01%), yielding a PCE of 21.65% (21.33%) under reverse (forward) scanning. And the control device achieved a VOC of 0.841 (0.842) V, a JSC of 31.02 (30.98) mA cm−2, and an FF of 75.32% (70.89%), yielding a PCE of 19.64% (18.49%) under reverse (forward) scanning. It demonstrated an obvious improvement of the JSC and FF values of PLPT-treated devices compared with the control ones (Fig. 2b and Supplementary Fig. 9). More interestingly, the hysteresis index (HI) decreased from 8.42% to 1.17% (Supplementary Fig. 10).
a J–V curves of control, PLPT-treated Pb-Sn NBG PSCs without post-passivation. b Statistical photovoltaic parameters of control and PLPT-treated Pb-Sn NBG PSCs. Box-plot elements: center line, median; box limits, upper and lower quartiles; whiskers, 1.5× interquartile range; data points, individual values from 9 independent devices. Error bars represent the standard deviation (SD) values. c, d The GIXRD results of the control and PLPT-treated film. e Fitted line of 2θ-sin2ψ from GIXRD results for the control and PLPT-treated film. f TRPL measurements of the PLPT-treated perovskite films with thickness varying from 250 to 850 nm. g Thickness dependence of the TRPL lifetime with analysis to extract the bulk carrier lifetime and the surface recombination velocity. h τb–Srv curves. i The dynamic interaction process of the perovskite surface structure during laser polishing (Ep: The energy was received by atoms per mole in perovskite.).
After PLPT treatment to remove the defective surface, the corresponding devices demonstrated a PCE improvement from 19.64% to 21.65%. To understand PLPT-treated working mechanism, we aimed to decode its surface information. Grazing-Incidence X-ray Diffraction (GIXRD) measurements were utilized to investigate the residual stress and thermal effects of PLPT treatment. As illustrated in Fig. 2c, d, at different tilt angles ψ, PLPT-treated films obtain much smaller systematic shift (0.04°) than control (0.13°). By fitting the relationship between 2θ and sin2ψ, the residual tensile strain of the PLPT-treated perovskite film was 21.78 MPa, one third of the control film (72.59 MPa) (Fig. 2e and Supplementary Note 1). For comparison to present UV-laser-treated film, the infrared laser treated ones showed obvious thermal effects (Supplementary Fig. 11) and large residual tensile strain (125.06 MPa). (Supplementary Fig. 12). Thermal effect of PLPT was further simulated by COMSOL software on the perovskite film, indicating a little thermal effect from PLPT process (Supplementary Figs. 13, 14 and Supplementary Note 2). Besides, XRD results showed that the main peak intensities of the PLPT-treated film increased (Supplementary Fig. 15) and the full width at half maximum (FWHM) values decreased (Supplementary Table 2) compared with the control samples. It was noted that the PbI2 impurity peak at 2θ = 12.71° of control was completely eliminated24,25. And the bandgap remained unchanged ~1.242 eV after PLPT treatment (Supplementary Fig. 16). Thus, PLPT treatment could release the surface stress, change the stoichiometric ratio to the ideal one with little thermal effects.
Next, the thickness dependence of time-resolved photoluminescence (TRPL) decay kinetics was examined to decouple the quality diagnosis for surface and bulk. Figure 2f and Supplementary Fig. 17 show the normalized TRPL decay curves with varied thickness values for control and PLPT-treated films with a step of 200 nm. The TRPL lifetimes (τTRPL) of the PLPT-treated films were longer than those of the control film at approximately the same thickness level. And the TRPL decays became faster with reduced film thickness for both samples (Supplementary Table 3). In general, the measured τTRPL are composed of the surface recombination velocity (Srv) and bulk recombination lifetime (τb) by the following equation26:
where d is the film thickness. According to the fitting relationship between 2/d and 1/τTRPL (Eq. (1)), Srv and τb values were obtained using the extracted slope and intercept parameters (Fig. 2g). The extracted τb value of PLPT-treated film reached approximately 3.0 μs, notably surpassing the value of the control film (2.2 μs) (Fig. 2h). Nonetheless, the Srv value of PLPT-treated films ~ 5.3 cm s−1 did not exhibit an obvious decrease compared to the control one (5.5 cm s−1). Thus, the enhanced τTRPL values of the PLPT-treated perovskite films were primarily attributed to the increase in bulk lifetime. Combining the detailed analysis (Supplementary Notes 3, 4), we have decoded the dynamic interaction between PLPT-treatment and perovskite. According to laser power attenuation, the laser polishing process could be roughly delineated into two distinct stages (Fig. 2i). At initial stage (stage Ⅰ), the laser energy (Ep) was much greater than the desorption energy of B-site ions (EB), where EB was the highest one among the desorption energies of A-site ions (EA), EB and X-site ions (EX). All atoms were etched away indiscriminately into ions fragments and clusters27. At ending stage (stage Ⅱ), Ep value attenuated to around EB, insufficient for perovskite etching instead of surface ion releasing19. And the surface ions with lower bond desorption energies demonstrated a greater probability of being dissociated28. Thus, PLPT-treated film showed rich of A-site vacancies (VA) according to our experimental testing and reported work results27.
According to the above decoding and experimental analysis, the PLPT-treated perovskite film surface was mainly composed of VA arrays (PLPT, Fig. 3a). By utilizing these VA arrays, it helped to support an opening platform to reconstruct polished surface with new A-site ions in order to achieve the synergistic effect of defect elimination and stability enhancement. By screening A-site library, guanidinium (GA+) can form effective hydrogen bonds two times larger than normal A-site cations (FA+, MA+)29 (Supplementary Fig. 18), possessing the ability to enhance stability and suppress ion migration29,30,31. Hence, GABr was prioritized for reconstructing the polished surface. Hereafter, the PLPT films reconstructed by GABr treatment were marked as target ones (Target, Fig. 3a). The optimization GABr concentration was 1 mg/mL and the details can be found in Supplementary Figs. 19, 21.
a Schematic diagram illustrating the evolution routes of the perovskite surface structure from control to target. b Atomic ratio data of A-site ions of PLPT-treated and target films surface after 5 nm, 10 nm, and 20 nm Ar-etching. c The A-site (FA+, Cs+, MA+ and GA+) ions distribution of PLPT-treated and target films. d Enlarged and normalized GIXRD spectra collected from the surface of the PLPT-treated and target perovskite films. The blue peak represents the VA-rich phase. e, f High-angle annular dark-field TEM images for the different regions from PLPT-treated and target samples. Scale bars are 2 nm. The second row shows the calculated interplanar spacing for each lattice.
The target sample was investigated by XPS with Ar-etching to survey the A-site and iodide atomic percentage. A-sites atomic ratio of PLPT-treated film surface (the top part of Fig. 3b) was 0.53 (Supplementary Note 5), while it stabilized around 1.00 after Ar-etching for over 10 nm. In contrast, the A-site atomic ratio of target film kept constant with uniform distribution in increasing depth via Ar-etching (the bottom part of Fig. 3b). Additionally, X-site atomic ratios of the target film surface showed slight increase compared with PLPT-treated film after GABr surface reconstruction (Supplementary Fig. 22). We further used ToF-SIMS to investigate the A-site cation composition after GABr treatment. A distinct signal of GA+ was clearly observed on the film surface comparing with PLPT-film (Fig. 3c). Through XPS and SIMS results, it can be confirmed that GA+ had entered the crystal lattice and obtained a new GA1−xCsₓ-based perovskite surface.
To study target surface information, GIXRD technology was employed to detect the crystal structure for PLPT and target samples (Fig. 3d). For the PLPT-treated surface, the characteristic peak localized at around 28.3°, which was indexed for the (200) plane of perovskite. When the incident angle was lower than 0.1° approaching surface, there was a new shoulder peak at around 28.8°. These emergent shoulder peak gradually weakened with an increase of the incident angle and ultimately vanished, attributed to substantial A-site vacancies causing lattice distortion32,33. After surface reconstruction with GABr, the shoulder peaks were completely eliminated, directly demonstrating that GA⁺ effectively incorporated into the crystal lattice and thereby resolved the structure distortion (the right part of Fig. 3d). Notably, when the incident angle was 0.05°, the diffraction peak of the (200) plane exhibited a slight shift (~0.01°) toward the lower angle. Moreover, Transmission electron microscopy (TEM) images were collected to further survey the newly recontruction surface. For the PLPT-treated film, the d values of interplanar spacing were measured to be about d = 3.48 Å (Fig. 3e), lower with the values reported in the literature attributed to partial A-site ions loss34. In contrast, the target film shows a d value of 3.57 Å (Fig. 3f), slightly higher than the literature34, attributed to larger radius of GA+ ion29. Thus, GA+ incorporation filled in A-site vacancies and reconstructed the surface composition of PLPT samples according to GIXRD and TEM results.
We further investigated the film quality and stability for target samples. Figure 4a shows enhanced (100) and (200) plane intensities of the target sample compared with the control and PLPT-treated ones, and the PbI2 peak at 2θ = 12.71° was completely eliminated. The absorption spectra of target sample remained unchanged (Supplementary Fig. 23). While new characteristic peaks were observed at 6.5° and 11.7° with a high concentration of GABr, attributed to the formation of GA2PbSnI4 phase35 (Supplementary Fig. 24). The steady-state photoluminescence (PL) and TRPL measurements were conducted to investigate the defects intensity and carrier dynamics for the perovskite films. As shown in Fig. 4b and Supplementary Fig. 25, the steady-state PL intensity for target sample is significantly enhanced comparing with the control and PLPT-treated ones from both top and bottom incidence, indicating improved film quality and reduced non-radiative recombination. The average carrier lifetime of the target film reached 2.82 µs longer than that of the PLPT film (2.44 µs) and the control sample (1.97 µs), indicating a reduced defect density (Fig. 4c and Supplementary Table 3). Furthermore, large-area PL-mapping was employed to characterize the film uniformity. The emission spectra of the control sample exhibited lower intensity and larger peak intensity fluctuation with a standard deviation (SD) of ~0.18 compared to the PLPT-treated sample (SD = 0.05) (Fig. 4d, e). In contrast, the target film displayed significantly higher emission intensity and smaller SD (0.04) relative to the PLPT-treated and control samples (Fig. 4f). This distinct difference directly confirmed that the target film possessed substantially improved lateral homogeneity, a key advantage for scalable optoelectronic application36.
a XRD patterns of perovskite films. The star mark represents the characteristic peak of PbI2. b Steady-state PL spectra of perovskite films excited from the top surfaces. c TRPL measurements of perovskite films. d–f PL intensity imaging of perovskite films deposited on glass substrates (2.5 × 2.5 cm2). The color bar shows the normalized PL intensity. g Band alignment of control, PLPT-treated and target films compared with C60. h PLQY data from control, PLPT-treated and target films, with and without C60. i The density values of mobile ions of the control, PLPT-treated and target films were obtained by BACE measurements in dark.
The obtained perovskite/C60 interfaces were further investigated by ultraviolet photoelectron spectroscopy (UPS). Figure 4g and Supplementary Fig. 26 display obvious p-type behavior for the control film surface due to the self-doping effect of Sn4+, leading to the emergence of an energy offset with C60. This led to minority carrier accumulation at the perovskite/C60 interface, and an approximately 140 meV band offset between perovskite and C60. This misalignment hindered carrier transport and aggravated defect-assisted carrier recombination losses at the interface37. The Fermi level of the target perovskite film showed a significant upshift from −4.75 eV to −4.64 eV relative to the control film. The band offset was reduced with 50 meV for perovskite/C60 interface, facilitating charge extraction. To verify the suppression of non-radiative recombination, photoluminescence quantum yield (PLQY) was carried out to assess the perovskite/ETL interface properties (Fig. 4h). The PLQY of the perovskite samples were gradually improved from 1.88% (control), to 2.72% (PLPT) and further to 4.71% (target). Upon contact with the C60 ETL, the PLQY value at the control interface dropped markedly to 0.91%, which testifies to the non-radiative recombination and energy loss at the perovskite/C60 interface. In contrast, the PLQY values of the PLPT-treated and target perovskite/C60 samples reached 1.15%, and 3.26%, respectively. Thus, the target surface reconstruction could remove the interfacial defects, favor energy level alignment, which synergistically suppressed carrier trapping and non-radiative recombination losses.
To study the defect distribution of Pb-Sn NBG PSCs, three kinds of devices from control, PLPT and target films were fabricated and analyzed by drive-level capacitance profiling (DLCP) and electroluminescence (EL)-mapping. The surface defect density for the target film was reduced by two-thirds compared to that of the control (Supplementary Fig. 27). EL-mapping of control, PLPT-treated, and target PSCs (Supplementary Fig. 28) further demonstrated that target PSCs had a higher and more uniform EL intensity at the same applied current. Therefore, the combination of DLCP and EL mapping results verified the effective suppression of trap density and carrier recombination by the present strategy.
Next, target perovskite film stability with GA1−xCsₓ-based surface was further investigated by quantifying the density of mobile ions. The density of mobile ions and their extraction time could be obtained by bias-assisted charge extraction (BACE) measurements in dark38,39 (Supplementary Fig. 29).
The amount of diffused charges (({Q}_{{{{rm{dif}}}}})) was obtained by integrating the displacement current ({I}_{{{{rm{tran}}}}}) with time t (Eq. (2)), then the density of mobile ions (ρ) was calculated by Eq. (3), where ({Q}_{{{{rm{dif}}}}}) was divided by active volume (({A}_{{{{rm{act}}}}})) and elementary charge e. For the PLPT-treated film, the (rho) value was measured to be about 1.08 × 1017 cm−3 (Fig. 4i), lower with the value of the control film (1.81 × 1017 cm−3). After GABr reconstruction, (rho) value further decreased to 0.72 × 1017 cm−3. The efficient suppression of ion diffusion contributed to the obvious improvement of device hysteresis40. We then compared the formation energy of Pb-Sn perovskite with different A-site ions. The density functional theory (DFT) calculations demonstrated that GA1−xCsₓ-based perovskite exhibited a smaller formation energy compared with Pb-Sn perovskite with other calculated A-site ions (Supplementary Fig. 30), confirming an pronouncedly improved surface stability41.
We then fabricated a series of Pb-Sn NBG PSCs for control, PLPT-treated and target devices in typical inverted structure. The champion device was obtained from target device with a VOC of 0.882 (0.881) V, a JSC of 33.25 (33.04) mA cm−2 and an FF of 82.11% (81.82%), yielding a PCE of 24.07% (23.82%) under reverse (forward) scanning (Fig. 5a and Supplementary Table 4). And the target PSCs achieved a record efficiency of 24.07% (certified 23.47%) by antisolvent-free technology (Fig. 5b, Supplementary Fig. 31 and Supplementary Table 5). Compared with control devices, the target devices showed improved values in all photovoltaic parameters (Supplementary Fig. 32). It was noted that the target devices exhibited much smaller hysteresis index (0.72%), compared with the control device (8.44%), primarily attributed to the reduction of defects acting as channels for ion migration (Supplementary Fig. 33A). The target PCEs also demonstrated a narrow distribution as shown in Supplementary Fig. 33B. The steady-state output (SPO) efficiency at the maximum power point reached 23.01%, while the control one was only 16.74% (Supplementary Fig. 34).
a J–V curves of control, PLPT-treated and target Pb-Sn NBG PSCs. b A summary of reported PCEs of Pb-Sn NBG PSCs fabricated by antisolvent and antisolvent-free methods. c EQE spectra of control, PLPT-treated and target Pb-Sn NBG PSCs. d TPV decay curves of control, PLPT-treated and target Pb-Sn NBG PSCs. e, f The analysis of VOC and FF losses of control, PLPT-treated and target Pb-Sn NBG PSCs. g Cross-sectional SEM image of a representative tandem device. Scale bar: 1 μm. h J–V curves of the target TSC and device architecture (inset). i MPPT curves of the control and target tandem devices with encapsulation under continuous 1 sun illumination.
The integrated JSC values derived from external quantum efficiency (EQE) curves of the control and target PSCs were 30.99 mA cm−2 and 32.10 mA cm−2 (Fig. 5c), respectively. The main EQE improvement stemmed from photons beyond 750 nm (dashed rectangle), which could be converted more efficiently by target device. Interestingly, the thinner target (50 nm less) absorber achieved a higher JSC value compared with the control. On one hand, it was ascribed to the improvement of perovskite film quality (Fig. 4). On the other hand, the target film surface improved the utilization yield of light. Compared with control device, when light illuminated from the ITO glass side, the smoother perovskite surface and smoother Ag electrode led to stronger specular reflection of the initially unabsorbed light (Supplementary Fig. 35), which ultimately obtained a longer optical path within the perovskite film (Supplementary Fig. 36), thereby decoupling the high requirements for both carrier lifetime and mobility of NBG film35. As shown in Fig. 5d, transient photovoltage (TPV) characterizations confirm a longer decay lifetime for the target device (761 μs) relative to the control device (342 μs), verifying the significant suppression of carrier trapping and nonradiative recombination. This could be further corroborated by light intensity-dependent photovoltage, where the ideality factor (n) of the target device (1.53) was much smaller than those of the control (1.97) and PLPT-treated (1.67) devices (Supplementary Fig. 37). Excitingly, the present strategy was not only efficient for present absorber bandgap (~1.25 eV), but also applicable to other bandgap PSCs for p-i-n and n-i-p structures, such as ~1.75 eV, ~1.55 eV, ~1.33 eV (Supplementary Fig. 38).
A quantitative loss analysis for VOC and FF was performed by referencing device performance from the ideal Shockley–Queisser limit (Supplementary Notes 6–7). Based on quasi-Fermi level splitting (QFLS) results, VOC losses were ascribed to bulk, ETL-interface, and HTL-interface loss components (Fig. 5e and Supplementary Table 6). The target film presented notably lower bulk and ETL-interface losses than control, verifying that surface polishing and reconstruction gave rise to the VOC improvement. In addition, FF losses are mainly induced by non-radiative recombination and transport loss (series resistance)42. The control device indicated high non-radiative recombination and transport loss, while PLPT-treated devices primarily reduced transport loss, explaining the FF increase. And target devices further suppressed the above two kinds of losses with negligible transport loss by GABr surface reconstruction (Fig. 5f).
Two-terminal all-perovskite TSCs were fabricated based on the target NBG perovskite device. The top subcell had an absorber composition of DMA0.1Cs0.4FA0.5Pb(I0.75Br0.25)3 with 5 mol% MAPbCl3 additives. The detailed information of the wide-bandgap (WBG) top cells can be found in Methods. The tandem devices utilized a typical inverted structure (Fig. 5g and the inset of Fig. 5h). Ultimately, our champion device obtained a PCE of 29.80% (29.52%) under reverse (forward) scanning, with a VOC of 2.16 (2.16) V, a JSC of 16.60 (16.62) mA cm−2 and an FF of 83.12% (82.24%) (Fig. 5h and Supplementary Table 7). The present PCE value represented a notable advancement for antisolvent-free methods, one percent net PCE higher than previous work9. The integrated JSC values of the WBG and NBG subcells from EQE spectra (Supplementary Fig. 39) were 16.11 and 16.04 mA cm−2, respectively, in good agreement with the JSC value from J–V measurements. The device demonstrated a SPO efficiency of 29.18% at the maximum power points (MPP, Supplementary Fig. 40). Based on the advantages of high uniformity and robustness of laser processing, we further fabricated 1 cm2-size and 20.07 cm2-size TSCs. The 1 cm2 device yielded a VOC of 2.15 (2.15) V, a JSC of 16.28 (16.15) mA cm−2, and an FF of 81.90% (81.82%), corresponding to a PCE of 28.67% (28.41%) under reverse (forward) scanning (Supplementary Fig. 41). The 20.07 cm2 mini-modules, yielded a PCE of 24.38% (24.27%) under reverse (forward) scanning, with a VOC of 17.00 (17.01) V, a JSC of 1.83 (1.82) mA cm−2, and an FF of 78.39% (78.41%) (Supplementary Figs. 42, 43 and Supplementary Table 8). The encapsulated small-area TSC retained 80% of its initial PCE for approximately 650 h of MPP tracking under AM 1.5 G in N2 ambient condition at room temperature (Fig. 5i), top value among reported data1,18, whereas the control device decayed to lower than 80% after only 172 h.
In summary, our work has demonstrated a robust strategy for converting the defective and rough surface of NBG film into an efficient and smooth one. PLPT is first developed to precisely remove the defective surface of NBG perovskite films, thereby improving surface compositional uniformity and flatness. By further decoding the treated surface, we screen GABr for reconstructing the newly exposed surface into GA1−xCsₓ-based perovskite, achieving enhanced carrier extraction efficiency and stability. The smooth perovskite surface obtains a substantial increase in optical path length, which results in a higher photocurrent. Impressively, the target Pb-Sn PSC yields a record PCE of 24.07% (certified 23.47%) and exhibits negligible hysteresis via an antisolvent-free approach. In contrast, the control device attains a PCE of only 21.58%. The target two-terminal all-perovskite TSC achieved an efficiency of 29.80%, one percent net PCE higher than previous work, demonstrating exceptional operational stability. The present surface conversion strategy effectively eliminates the key surface-effect bottleneck across various perovskite compositions, paving the way for universal performance improvement.
All raw materials were not purified and were used as received. PbI2 (99.999%), SnI2 (99.999%), PbBr2 (99.99%), PEAI (99.5%) and NiO nanocrystal (99.999%) were purchased from Advanced Election Technology Co., Ltd. FAI (99.5%), MAI (99.5%), MACl (99.5%), Pb(SCN)2 (99.5%), FABr (99.5%), CsBr (99.9%), CsI (99.999%), PbCl2 (99.9%), DMAI (99.5%), EDAI2 (99.5%), PDADI (99.5%), GAI (99.5%), GABr (99.5%), GASCN (99%), PEDOT:PSS, C60 and PC61BM were purchased from Xi’an Yuri Solar Co., Ltd. SnF2 (99%) was purchased from Aladdin Co., Ltd. 4PADCB was purchased from Vizuchem Co., Ltd (Shanghai, China). Bathocuproine (BCP) was purchased from TCI. Tin Powder (99.999%), N,N-dimethylformamide (DMF, 99.8%), dimethyl sulfoxide (DMSO, 99.9%), chlorobenzene (CB, 99.9%), toluene (TL, 99.5%) and isopropanol (IPA, 99.5%) were purchased from Sigma Aldrich.
The mixed Cs0.1FA0.6MA0.3Sn0.5Pb0.5I3 narrow-bandgap perovskite precursor with a concentration of 1.8 mol L−1 was prepared by mixing CsI (46.8 mg, 0.180 mmol), FAI (185.7 mg, 1.08 mmol), MAI (85.8 mg, 0.540 mmol), SnI2 (335.3 mg, 0.900 mmol), PbI2 (414.9 mg, 0.900 mmol), SnF2 (14.1 mg, 0.090 mmol), and Pb(SCN)2 (2.7 mg, 0.036 mmol) in mixed solvents of 0.25 mL DMSO and 0.75 mL DMF.
The wide-bandgap perovskite precursor with a concentration of 1 mol L−1 was prepared by mixing CsBr (42.56 mg), FABr (31.24 mg), DMAI (17.30 mg), MACl (3.38 mg), Pb(SCN)2 (3.25 mg), PbCl2 (13.90 mg), PEAI (2.5 mg), FAI (42.99 mg), CsI (51.99 mg), PbBr2 (55.05 mg), and PbI2 (391.85 mg) in mixed solvents of 0.2 mL DMSO and 0.8 mL DMF. The precursor solution was filtered through a 0.22 μm PTFE filter before using.
The pre-patterned indium tin oxide (ITO) substrates underwent a cleaning process involving ultrasonication in deionized water, isopropanol, and ethanol for 30 min in succession. poly(3,4-ethylenedioxythiophene) polystyrene sulfonate (PEDOT:PSS) which was diluted by n-propyl alcohol (1:5) was coated on the cleaned ITO substrate at 2,000 r.p.m. for 30 s and then heated at 125 °C for 20 min. After cooling, the substrates were transferred to an N2-filled glovebox quickly, and perovskite films were spin-coated onto the substrate at 3,500 r.p.m for 10 s. The wet films were then directly transferred into a vacuum chamber (100 mL). After the vacuum pump was turned on, the vacuum level dropped to 3 Pa within 4-5 seconds. The total vacuum quenching time was about 12 s. Subsequently, the perovskite films were annealed at 100 °C for 10 min. After cooling, the perovskite films were post-treated by spinning a solution of EDAI2 (0.5 mg mL−1) in 1:1 IPA:CB solvent at 4,000 r.p.m. for 20 s, followed by heating at 100 °C for 5 min. After cooling down to room temperature, 20 nm fullerene (C60) film and 7 nm bathocuproine (BCP) were subsequently deposited by thermal evaporation at a deposition rate of 0.15 Å s−1. Finally, 150 nm Ag electrode was deposited by thermal evaporation at a deposition rate of 0.5 Å s−1.
For the narrow-bandgap perovskite solar cell with PLPT and GABr surface passivation. After annealing NBG perovskite films at 100 °C for 10 min, the films were polished with a picosecond ultraviolet pulsed laser. GA+ surface reconstruction solution was prepared by adding GABr (GASCN and GAI) into IPA at concentrations of 0.5, 1, 2, and 5 mg mL−1. GABr was further dissolved in mixed solvents of IPA and TL with different ratios (IPA:TL = 1:0, 3:1, 1:1, 1:3). After PLPT, GABr solutions were spin-coated onto the perovskite at 4,000 rpm for 20 s, then the film was annealed at 100 °C for around 5 min. The subsequent process is the same as above.
NiO nanocrystal dissolved in deionized water (4 mg mL−1) was spin-coated onto the ITO substrates at 5,000 r.p.m. for 30 s, followed by an annealing process at 150 °C for 10 min. After cooling, oxygen plasma was used to treat the substrate for 5 min, and then the self-assembled monolayer 4PADCB dissolved in ethanol (0.5 mg mL−1) was spin-coated onto the substrates at 3,000 r.p.m for 30 s, followed by heating at 100 °C for 10 min. After cooling, 45 µL of wide-bandgap perovskite precursor was dropped on the substrate and spin-coated through a two-step process: 1,000 r.p.m. for 5 s and 4,500 r.p.m. for 40 s. At the twentieth second, hot gas flow (70 m s−1, 50 °C) started, the gas flow was maintained for 4–6 seconds, resulting in the film turning dark brown. After gas purging, the samples were then annealed at 100 °C for 10 min. After cooling to room temperature, the perovskite films were post-treated by spinning PDAI2 solution (1 mg mL−1 in 1:1 IPA: toluene) at 5,000 r.p.m. for 30 s, followed by heating at 80 °C for 5 min. Then, 20 nm C60 was deposited on top of the perovskite films by thermal evaporation at a rate of 0.15 Å s−1. The samples were then transferred to an ALD system to deposit 30 nm SnO2. After that, 1 nm Au was thermally evaporated. After that, NBG subcells were fabricated using the abovementioned methods.
The mini-modules were fabricated on the 6.0 cm × 6.0 cm sized glass/ITO substrate using a 355 nm nanosecond laser scribing with a power of 0.1 W (P1), isolating into 8 subcells with a width of 5.6 mm. All the spin-coated layers in mini-modules were spin-coated at the same spin-coating speed as that used for small-area devices. The power of P2 scribing is 0.12 W. The effective monolithically inter-connected modules were formed by laser scribing (0.5 W) to form P3 lines.
The active area of small-area devices, including narrow-bandgap, wide-bandgap, and tandem devices, was 0.0768 cm2. For the PCE measurements conducted in our laboratory, the aperture area of mask was 0.0395 cm2. For the certification tests performed at the third-party laboratory, the aperture area of the mask was 0.0401 cm2. The film surface and cross-section morphology were characterized using a Hitachi S4800 SEM. The XPS spectra for powder and film samples were conducted using the AXIS SUPRA+ instrument from Shimadzu-Kratos (Japan). The ToF-SIMS measurement (Helios 5 HX/Helios 5 UX/Helios 5 FX DualBeam) was performed with a BiMn primary ion beam (3-lens 30 keV) for the analysis. A 50 × 50 µm2 area was analyzed with a 256 × 256 primary beam raster. Sputtering depth was acquired with 1 keV Cesium ion beam (6 nA sputter current) with a raster of 150 × 150 microns. The AFM height images were obtained in the ambient atmosphere using a Bruker Dimension Icon XR AFM. UV-Vis absorption was measured by a SolidSpec-3700 spectrophotometer. XRD were recorded using Rigaku D-MAX 2200 equipment. The theta/2theta modes were conducted with a Cu Kα radiation and an anode operating at 40 kV and 250 mA. Film morphology and cross-sectional structures of devices were measured by a field-emission scanning electron microscope (ZEISS Gemini 300). Steady-state PL was measured using a laser confocal Raman spectrometer (LabRAM HR800, Horiba JobinYvon). The light was illuminated from both the top and bottom surface of the perovskite films (excited by 532 nm). Time-resolved PL was measured using a spectrofluorometer (QuantaMaster 8000 series fluorometers, Horiba), samples were excited by a 532 nm pulsed laser.
The measurements of EL-mapping were obtained by combination of lock-in amplifier (SR830), electric-meter (Keithley 2000, Keithley 2400), and photodector (Thorlabs PDA100A). The J–V characteristics were measured using a Keithley 2450 sourcemeter and a solar simulator (EnliTech, Class AAA, AM1.5 G). The AM 1.5 G was calibrated with NREL reference solar cells (KG-5 and KG-0 reference cells were used). Bias voltages for J–V measurements of single-junction NBG perovskite solar cells were scanned from −0.1 V to 1 V (forward scanning) and from 1 V to −0.1 V (reverse scanning) with a scanning step of 0.05 V. The active area was determined by the aperture shade masks (3.95 mm2) placed in front of the solar cells. Bias voltages for J–V measurements of tandem cells were scanned from −0.1 V to 2.2 V (forward scanning) and from 2.2 V to −0.1 V (reverse scanning) with a scanning step of 0.05 V. EQE measurements were performed in ambient air using a QE system (EnliTech) with monochromatic light focused on a device pixel and a chopper frequency of 20 Hz. For EQE measurements of tandem solar cells, two light-emitting diodes with emission wavelengths at 450 nm and 850 nm were used as the bias lights to measure NBG and WBG subcells, respectively. The operational stability tests were carried out under multicolor light-emitting diode illumination in N2. No UV filter was used during the stability tests.
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This work is supported by the National Natural Science Foundation of China (Grant No. 62374065, 62574096), the Wuhan Key Research and Development Program (2025010602030106), Project for Building a Science and Technology Innovation Center Facing South Asia and Southeast Asia (202403AP140015), the Innovation Project of Optics Valley Laboratory (No. OVL2021BG008, OVL2024ZD002), Shenzhen Science and Technology Program (JCYJ20250604190827037), the International Science and Technology Cooperation Projects of Hubei Province (GJHZ202500083), Natural Science Foundation of Wuhan (2025040601020188), Hubei Optical Fundamental Research Center (HBO2025TQ003). The authors thank Engineer Jun Su from the Center of Optoelectronic Micro and Nano Fabrication and Characterizing Facility, WNLO of HUST for the support in the SEM test.
These authors contributed equally: Tianjun Ma, Dingfu Luo.
Wuhan National Laboratory for Optoelectronics (WNLO), Huazhong University of Science and Technology, Wuhan, China
Tianjun Ma, Wenjiang Ye, Jun Yan, XuKe Yang, Mingyu Li, Yuheng Li, Salman Ali, Shiwu Chen, Haisheng Song & Jiang Tang
School of Optical and Electronic Information (SOEI), Huazhong University of Science and Technology, Wuhan, China
Dingfu Luo, Xinzhao Zhao, Hao Wang, Ruiheng Gao, Sifan Liu, Ying Zhou, Chao Chen, Haisheng Song & Jiang Tang
China-EU Institute for Clean and Renewable Energy (ICARE), Huazhong University of Science and Technology, Wuhan, China
Qilin Guo & Haisheng Song
Hubei Key Laboratory of Optical Information and Pattern Recognition, Wuhan Institute of Technology, Wuhan, Hubei, P. R. China
Bingxin Ding & Pingli Qin
Houston Technology Research Center, CNPC USA Corporation (CNPCUSA), Houston, TX, USA
Michael Wang & Chris Cheng
Optics Valley Laboratory, Wuhan, China
Chao Chen, Haisheng Song & Jiang Tang
Shenzhen Huazhong University of Science and Technology Research Institute, Shenzhen, China
Haisheng Song
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H.S. and J.T. conceived the idea and directed the overall project. T.M. and D.L. fabricated all the devices by PLPT and conducted the characterizations. W.Y, X.Z., Y.Y., H.W., X.Y., and M.L. helped to set up laboratory equipment and optimize the vacuum-assisted devices. Y.L., S.L., A.S, R.G., and S.C. performed maximum power points tracking measurement. B.D., S. H., M.W, C.C.(Chris), Y.Z., C.C., and P.Q. helped to discuss and analyze data. T.M. and H.S. wrote the manuscript. All authors discussed the results and commented on the paper.
Correspondence to Haisheng Song or Jiang Tang.
The authors declare no competing interests.
Nature Communications thanks Han Chen, Jiangang Liu and the other anonymous reviewer(s) for their contribution to the peer review of this work. A peer review file is available.
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Ma, T., Luo, D., Ye, W. et al. Non-contact laser polishing and reconstruction towards high-efficiency all-perovskite tandem solar cells. Nat Commun 17, 4193 (2026). https://doi.org/10.1038/s41467-026-71017-7
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