Uniform phase distribution of wide bandgap perovskite for high-performance perovskite-silicon tandem solar cells

Domain: physics
Published: 2025-12-12
Editorial outcome: minor_revision

Abstract
Metal halide perovskite-silicon tandem solar cells offer a promising pathway towards surpassing the efficiency limits of single-junction devices, and the focus remains on the wide bandgap (WBG) perovskite top cell optimization. Here, by incorporating melamine additive into a rubidium-alloyed perovskite composition, we suppress film phase separation and inhomogeneity and enhance charge carrier mobility. Consequently, our modified 1.68 eV WBG perovskite solar cells exhibit superior charge transport and minimized non-radiative recombination losses, achieving state-of-the-art performance with open circuit voltage (V OC) of 1.31 V, fill factor of 86.4% and efficiency of 25%. This also derives an efficient two-terminal perovskite-silicon tandem cell with stabilized efficiency of 33.5% and high V OC of 2.02 V. Both the single-junction and tandem devices show noticeable operational stability that the WBG cells maintain 80% of initial efficiency (T 80) after over 3200 hours of 1-sun illumination under 65 °C (ISOS-L-2), while the tandem device survive T 90 lifetime exceeding 1100 hours under the same conditions.

Introduction
Crystalline silicon (c-Si) solar cells have achieved a certified photovoltaic efficiency of 27.8%, approaching the Auger-recombination-constrained Shockley-Queisser (S-Q) limit of 29.8%1,2. Metal halide perovskite-silicon tandem architectures, with two complementary absorption layers, provide a promising route to surpass the efficiency limit of single-junction (S-J) solar cells3–10. For such tandem devices, the optimal bandgap for the top perovskite cell is approximately 1.7 eV, considering the current matching with the c-Si bottom cell and maximize power output11. Achieving this normally requires a high bromide content in the mixed bromide-iodide (Br-I) halide perovskite composition, along with a balance of A-site cations and X-site halides in the ABX 3 structure for the desired WBG. However, the high bromide content often leads to rapid crystallization kinetics, posing challenges in obtaining uniform cation and halide distributed WBG perovskite films3,12–15. Consequently, spatial phase inhomogeneity could bring forth device performance losses, manifested as reduced device V OC, sacrificed efficiency, and compromised operational stability, especially under light exposure combined with thermal or electrical stress4,7,16–19.

To advance perovskite-silicon tandem technology, the primary focus remains on optimizing the WBG perovskite composition and enhancing film quality, particularly in terms of phase uniformity and purity20,21. Rubidium (Rb) alloying has been shown to improve the optoelectronic properties of perovskites, likely due to modified grain boundaries and inter-grain charge transport and reduced trap densities22–24. Additionally, Rb incorporation has demonstrated the potential for improving stability by increasing the energy barriers for ion-migration, thereby suppressing light-induced phase segregation25–30. However, it may also introduce new challenges related to the A-site inhomogeneity, which can lead to the formation of inactive non-perovskite secondary phases such as RbPb(I 1−x Br x ) 3, ultimately compromising both efficiency and stability. The underlying causes of these phenomena, along with strategies to achieve better phase uniformity control, remain under active investigation22,31–33. Therefore, it is essential to effectively harness the advantages of Rb incorporation while mitigating potential phase non-uniformity.

We develop a triple-cation RbCsFA (Cs: cesium; FA: formamidinium) WBG perovskite to replace our previously reported CsFA composition as the starting material27. The emergence of undesired non-active secondary Rb-rich phases prompted us to regulate the nucleation and crystallization through additive engineering. Amine groups and triazine core functionalized materials stand out as promising regulators, due to their ability to form strong electrostatic interactions through hydrogen bonding with the lead halides octahedra and A-site cations, thereby reinforcing lattice integrity and enabling defect passivation simultaneously8,34–37. To this end, we selected melaminium iodide (MLAI), a triazine heterocyclic compound as the regulator, which was a universal and effective additive material in some previously reported results. For example, it behaved as a good hydrogen bond acceptor and Lewis base to interact with both the organic cation and Lewis acidic lead iodide (PbI 2), which resulted in good passivation effect to the perovskite layer35,36. It was also adopted to construct porous and disordered PbI 2 morphology and then help convert the residual PbI 2 to perovskite phase in two-step sequential deposited perovskite films, which was based on the tridentate chelation effect between melamine and PbI 2 37.

In this work, MLAI is utilized as an effective phase uniformity regulator in our RbCsFA WBG perovskites. The strong coordination of melamine to the alloyed-cation perovskite phase can lower its formation energy and the nucleation barrier, suppressing the formation of δ-RbPbI 3, thus mediate homogeneous nucleation and distribution of our rubidium-containing WBG perovskite phases. The introduction of MLAI also significantly improve film and device charge transport and minimize non-radiative recombination losses. Our optimized 1.68 eV S-J WBG PSCs achieve state-of-the-art performance with a V OC of 1.31 V, a FF of 86.4% and a PCE of 25%, with both V OC and FF approaching 95% of their S-Q limit. The devices also exhibit excellent averaged operational stability, with T 90 lifetime (90% of its initial efficiency) retaining for over 1400 h and T 80 for over 3200 h under 1-sun continuous illumination at 65 °C (ISOS-L-2). Furthermore, the integrated perovskite-silicon two-terminal tandem cell reach a stabilized PCE of 33.5% (certified stabilized PCE of 32.9%), with impressive operational stability of T 90 lifetime for over 1100 h under 1-sun at 65 °C.

Results
Perovskite phase homogeneity modulation by MLAI addition
We adopted Rb 0.05 Cs 0.1 FA 0.85 Pb(I 0.75 Br 0.25 ) 3 (abbrev. as RbCsFA) to replace previously studied Cs 0.15 FA 0.85 Pb(I 0.77 Br 0.23 ) 3 (abbrev. as CsFA) for our WBG perovskites composition in this study27, and the perovskite bandgap is approximately 1.68 eV (Supplementary Fig. 1). The choice was based on the potential benefits of Rb for enhancing charge carrier mobility and improving charge transport, as evidenced by transient absorption microscopy (TAM) measurements (Supplementary Note 1, Supplementary Figs. 2, 3). The RbCsFA perovskite film exhibited obviously higher carrier diffusion coefficient (D, 0.112 cm 2 s −1) than that of CsFA (0.063 cm 2 s −1). By using Einstein relation, the carrier mobilities (μ) were translated to be 4.43 cm 2 V⁻1 s⁻1 and 2.43 cm 2 V⁻1 s⁻1 for RbCsFA and CsFA, respectively. These improvements showed the great potential of RbCsFA for better charge transport and solar cells applications. Initial solar cells fabrication also demonstrated significant advances in reducing device hysteresis and overall efficiencies, further validating the promises of RbCsFA perovskite and setting a good benchmark for further study (Supplementary Fig. 4).

Although RbCsFA composition is a potential perovskite absorber for further solar cells application, we found there were phase segregation existed in the original perovskite film. As shown in Fig. 1a, from the top view scanning electron microscopy (SEM) analyses, crystallized flake-like secondary phase precipitation was observed on the top surface. Energy dispersive X-ray spectroscopy (EDS) analyses revealed that these were aggregated Rb-rich species (Supplementary Fig. 5). Grazing incidence X-ray diffraction (GIXRD) further confirmed that these were non-perovskite δ-RbPbX 3 (X denotes the combinations of I and Br halide), indicated by a characteristic (110) diffraction peak at 10.18° and comparison with a powder diffraction file (PDF) of δ-RbPbI 3 38 (Fig. 1b). Here we used RbPbI 3 as a reference considering the undistinguishable XRD patterns of RbPbI 3 and RbPbBr 3 39,40. As for depth-dependent distribution by GIXRD, the intensity ratio of the (110) peak of δ-RbPbI 3 to the perovskite (100) peak (14.14°) decreased from 12.3 to 6.1% as the incident angle increased from 0.5° to 2.0° (Fig. 1c, Supplementary Fig. 6a). The tilt cross-sectional SEM image also visualized that the main accumulation of δ-RbPbI 3 was located at the perovskite top surface (Supplementary Fig. 6b). The preferential precipitation of δ-RbPbI 3 on the top surface can be attributed to the top-to-down crystallization during annealing, coupled with the lower formation energy of δ-RbPbI 3 phase, for which we will discuss later16,41,42. To fully exploit the advantages of Rb-containing WBG perovskites and mitigate non-uniform phase distribution, we introduced MLAI into the perovskite. Surface SEM image and GIXRD results confirmed the absence of the δ-RbPbI 3 phases on the MLAI-modified perovskite film (Fig. 1d,e, Supplementary Fig. 7). EDS analyses further validated the disappearance of Rb-rich phase and a uniform Rb distribution (Supplementary Fig. 8).

We conducted density functional theory (DFT) calculations to reveal the phase formation process of the perovskite films. The formation energy calculations revealed that the thermodynamic formation energy of δ-RbPbI 3 and alloyed cation phase of α-RbFACsPbI 3 are −3.42 eV/unit and −3.21 eV/unit, respectively, before melamine addition (Fig. 1f). The calculation details can be found in Supplementary Note 2 and Supplementary Fig. 9. The nucleation barriers calculations revealed that the nucleation barriers (ΔG*) for δ-RbPbI 3 and α-RbFACsPbI 3 are 0.28 eV and 0.42 eV, respectively (Supplementary Note 3, Fig. 1g). Since the formation energy and nucleation barrier of δ-RbPbI 3 are both lower compared with that of alloyed cation phase, δ-RbPbI 3 is more likely to nucleate first and then precipitated at perovskite top surface during the crystallization process, which was consistent with GIXRD and EDS results. After introducing melamine, the formation energy of α-RbFACsPbI 3 (−3.76 eV/unit) can be reduced to even below that of δ-RbPbI 3 (−3.62 eV/unit) (Fig. 1f). Moreover, the introduction of melamine also significantly lowers the nucleation barrier ΔG* for the alloyed cation phase (0.12 eV), compared with δ-RbPbI 3 (0.15 eV) (Fig. 1g). This finally enhanced the nucleation rate of alloyed cation phase according to the Arrhenius-type formula (Supplementary Eq. 8), which significantly promoted the preferred nucleation of α-RbFACsPbI 3, thereby inhibiting the precipitation of δ-RbPbI 3 (Supplementary Note 3). Through X-ray photoelectron spectroscopy (XPS) analysis, clear peak shifts of the Rb 3d, Pb 4f, Cs 3d and N 1s toward higher binding energies was observed, which means that the Fermi level is more distant to the core level, indicating that there was strong interaction between melamine with lead halides octahedra and A-site cations. And the strong interaction should be responsible for the formation energy changes (Supplementary Fig. 10). Note that a weak signal corresponding to the metallic Pb 0 state—identified by the Pb 4f 7/2 and Pb 4f 5/2 peaks located at around 136.06 eV and 140.89 eV, respectively, was detected in the control sample but disappeared in target sample. We attribute this signal to the reduction of PbI 2 induced by high-energy X-ray exposure during the XPS measurement, where the PbI 2 diffraction peak was also detected in the XRD patterns of control perovskite film (Fig. 1c). And we could also expect that the target perovskite phase possesses superior stability under high-energy X-ray irradiation compare to the control sample, thereby mitigating such reduction effects.

A schematic diagram of perovskite film formation comparison was depicted in Fig. 1h. The melamine addition function as a crystallization regulator to the Rb-containing composition, promote the uniform alloy-cation phase distribution in the final perovskite film and suppress the δ-RbPbI 3 non-perovskite phase precipitation.

There is optimal MLAI additive concentration window for suppressing δ-RbPbI 3, as demonstrated by the SEM surface morphologies in Supplementary Fig. 11, where the target concentration is around 4 mM with the most uniform phase distribution. At lower MLAI concentration, the suppression of δ-RbPbI 3 precipitation was incomplete, while higher concentration could hinder grain growth, reduce the perovskite grain size and potentially create transport barriers at grain boundaries. By comparing device efficiencies and photovoltaic parameters across different MLAI concentrations, we also determined 4 mM to be the target concentration for further optoelectronic investigations (Supplementary Fig. 12).

To gain deeper insights into the perovskite film growth mechanism, we performed in-situ photoluminescence (PL) spectrometry monitoring during the film formation process. As shown in Fig. 2a, b, the target perovskite film with 4 mM MLAI addition exhibited a noticeable slower PL peak shift and a more gradual intensity change during the initial crystallization stage (0–2 s), compared to the control film without MLAI. These observations suggested that MLAI effectively retards the crystallization process, which is expected to facilitate the formation of high-quality WBG perovskite films with improved phase homogeneity43,44.

We also conducted TAM to evaluate the bulk charge transport properties of perovskite film with or without MLAI. The target film exhibited an enhanced carrier diffusion coefficient (0.147 cm 2 s -1) and carrier mobility (5.69 cm 2 V⁻1 s⁻1) than that of the control film (D of 0.112 cm 2 s -1 and μ of 4.43 cm 2 V⁻1 s⁻1) (Supplementary Note 1, Supplementary Figs. 2, 3), suggesting that the perovskite films charge transport was further enhanced after MLAI regulation. In parallel, time-resolved photoluminescence (TRPL) measurements were conducted on two device configurations (as illustrated in the inserts of Fig. 2c and d). For the stack of glass/FTO/CbzNaph/perovskite, bi-exponential fitting of the decay curves yielded an average carrier lifetime of 2.47 μs for the target sample, significantly longer than the 0.8 μs observed for the control (Fig. 2c, Supplementary Table 1). This result suggests a notable reduction in defect density and suppression of non-radiative recombination in target perovskite film. Similar improvement was also observed for the stack of glass/FTO/perovskite/EDAI 2 +PEAI/C 60 (Fig. 2d, Supplementary Table 2). To further investigate interfacial charge dynamics, we performed differential lifetime (τ PL(t)) analysis based on the TRPL data by computing the inverse negative slope of the logarithmically plotted PL photon flux ϕ(t). From the differential lifetime graph (Fig. 2e, f), the first interval at shorter times of the steep rise represents the charge extraction/transfer process and the sharpness of the rise implies the speed of the charge transfer, the transition from the ascending trend to a plateau signals the cessation of charge transfer, after which the second interval at longer delay times represents the non-radiative recombination dominant process45,46. In both configurations, the target samples exhibited a sharper rise in the initial time region, indicating more efficient charge transfer and reduced trap-assisted recombination at both SAM/perovskite and perovskite/C 60 interfaces, compared to control47,48. The mitigated aggregation of δ-RbPbI 3 phase and homogeneous phase distribution greatly helped bulk and interfacial charge transfer properties.

We further evaluated the photostability of the films by conducting in-situ time-dependent PL measurements under continuous illumination. As shown in Fig. 2g, h, the PL emission peak of the control sample red-shifted by approximately 35 nm after 80 min exposure of 2-sun intensity in air, indicating severe photoinduced phase segregation. In contrast, the MLAI-modified perovskite film exhibited negligible peak shift, underscoring its improved photostability and effective suppression of phase segregation and ion migration. These results suggested that MLAI-incorporated perovskite films hold great potential for better solar cells applications.

Single junction WBG photovoltaic performance and analysis
To comprehensively evaluate the effectiveness of WBG S-J PSCs after introducing MLAI into perovskite film, we fabricated devices with the configuration of glass/FTO/CbzNaph/Rb 0.05 Cs 0.1 FA 0.85 Pb(I 0.75 Br 0.25 ) 3 /EDAI 2 +PEAI/C 60 /SnO 2 /Cu. As shown in Fig. 3a, the representative current density-voltage (J–V) curves based on MLAI-optimized devices delivered a state-of-the-art efficiency of 25.03%, featuring a V OC of 1.31 V, FF of 86.4%, and J SC of 22.17 mA/cm 2 under reverse scan direction, without hysteresis. In contrast, the control device obtained a reverse scan efficiency of 22.84% (V OC of 1.28 V, J SC of 21.68 mA/cm 2, and FF of 82.3%) (Supplementary Table 3). Note that we also fabricated devices without the surface treatment to evaluate the V OC improvement resulting from MLAI addition. As shown in Supplementary Fig. 13, the V OC also increased approximately 0.04 V (from 1.24 V without MLAI to 1.28 V with MLAI). The enhancement is consistent with that observed in devices treated with EDAI 2 and PEAI. The additional surface treatment contributed a further V OC increase of 0.03 V. What’s more, MLAI couldn’t act as an effective surface passivation material in our perovskite system (Supplementary Fig. 13). The corresponding stabilized power output (SPO) efficiency under their maximum power point was increased from 22.7% for control to 24.7% for target, respectively (Fig. 3b). The integrated J SC from the external quantum efficiency (EQE) are 21.37 mA/cm 2 and 21.53 mA/cm 2 for control and target devices, respectively, which aligned well with the J–V characteristics (Supplementary Fig. 14). The determined bandgaps both were approximately 1.68 eV from the first derivative of the EQE spectra, and the incorporation of MLAI didn’t change the perovskite bandgap. Our device efficiency was independently certified at National Photovoltaic Industry Measurement and Testing center (NPVM), where the reverse scan efficiency was 24.92%, with a V OC of 1.31 V, a J SC of 22.15 mA/cm 2, and an FF of 86.15%. The forward scan efficiency was 24.83%, with a V OC of 1.31 V, a J SC of 22.15 mA/cm 2, and an FF of 85.81%, along with the stabilized efficiency of 24.08% after 300 s maximum power point tracking (Supplementary Fig. 15). These results aligned closely with our laboratory measurements.

To assess the repeatability of the MLAI modification, we collected photovoltaic parameters from 60 control devices and 73 target devices for the statistical distributions analysis. The average device efficiency increased from 22.2 to 24%, alongside narrower parameters distributions for the target devices (Fig. 3c). The superior efficiencies with the obvious V OC and FF increasement, can be attributed to the superior bulk perovskite quality and improved interfacial charge transfer thanks to MLAI modification. The delayed crystallization helped providing a larger process window to achieve excellent device reproducibility and reliability.

We further compared our WBG PSCs results with previously reported devices (1.64–1.69 eV bandgaps), as summarized in Supplementary Table 4 and we derived the visualized results as Fig. 3d, e. For this bandgap range, three ranges were defined for V OC and FF as 95%, 90%, and 85% of the S-Q limit. Our results way exceeded literature values, achieving 95% of the S-Q limit for both the V OC and FF. Moreover, our best PCE of 25% also way beyond the reported PCEs across 1.64–1.69 eV bandgaps and surpassed 85% of its S-Q limit (Fig. 3e), demonstrating our great endeavor and progress in pushing forward the thermodynamic potential of the WBG perovskite absorbers.

To gain deeper insights into the improved device efficiencies, we investigated the carrier recombination and charge transport behaviors, by recording the device J–V responses under varied light intensities (Supplementary Fig. 16). Both the control and MLAI modified devices displayed a linear relationship of J SC versus light intensity with the slope close to 1, indicating that the bimolecular recombination is negligible and implied the unimpeded carrier transport and collecting within the solar cells49,50. For the linear relationship of V OC with the natural logarithmic of light intensity, where the slope of the plots can determine the ideality factor (n), the target device delivers an n value of 1.08, much lower than that of the control device of 1.21, suggesting a substantial attenuation of defect-assisted recombination in target device51,52. Electroluminescence (EL)-EQE further supported these findings (Supplementary Fig. 17). At an injection current density of 22.14 mA/cm 2, an EL efficiency (EQE EL) of the target device reached 12.6%, whereas only 1.1% for the control. Based on EL-EQE results, the V OC non-radiative recombination loss ΔV non-rad oc of control and target device were calculated to be 0.12 V and 0.06 V according to Supplementary Eq. 13, aligning well with the J–V results (Supplementary Note 4)53,54. The greatly enhanced EL efficiency and minimal ΔV non-rad oc confirmed the improved optoelectronic quality of perovskite and the suppressed non-radiative recombination in the entire device after MLAI addition. Dark I–V characteristics showed that the target device exhibited higher shunt resistance, reduced series resistance and lower leakage current compare to control, also suggesting more efficient overall charge transport and suppressed non-radiative recombination losses55,56 (Supplementary Fig. 18).

To assess device stability, we tracked the encapsulated devices operational stability under 1-sun light-emitting diode (LED) illumination at 65 °C. The MLAI-incorporated device retained 90% of its initial PCE (T 90) after over 1400 h and kept T 80 lifetime for over 3200 h, with 6 averaged device results (Fig. 3f). Note that one champion device survived over 3800 h of T 80 lifetime (Supplementary Fig. 19), and our stability results substantially surpassed the current literature reports for WBG (~1.68 eV) PSCs (Supplementary Table 5). In stark contrast, the control devices exhibited averaged T 90 lifetime for only 405 h. For the detail degradation comparison of each of the solar cell parameters including Voc, FF and J SC, all of them decayed faster for the control sample, indicating that both the perovskite bulk and device interfaces possibly underwent more severe degradation under light and heat combination (Supplementary Fig. 19). The spectrum of the LED light source for stability measurement was depicted in Supplementary Fig. 20. We also conducted thermal stability tests under 85 °C, five control and ten target devices were monitored that they showed similar trends without obvious efficiency drop (Supplementary Fig. 21), indicating that light soaking is a harsher condition for phase segregation in our perovskite system. The greatly improved operational stability under elevated temperature and light-soaking could be attributed to the homogenized phase distribution, improved perovskite crystallization quality, facilitated charge transport, and suppressed phase segregation.

Perovskite-silicon tandem device performance and analysis
To further check the active role of our additive strategy for WBG perovskites, we applied our optimized WBG perovskite atop of a double-textured silicon heterojunction solar cell (SHJ). A schematic stack of the silicon-perovskite two-terminal tandem structure is depicted in Fig. 4a. Before integrating with silicon cell, 1 cm 2 semitransparent single-junction WBG PSCs was fabricated to check the uniformity of perovskite film and effectiveness of MLAI regulation with the structure of ITO/CbzNaph/perovskite/EDAI 2 +PEAI/C 60 /SnO 2 /indium zinc oxide (IZO)/silver (Ag) grid. Upon illumination from the IZO side, the MLAI modified device delivered the best PCE of 22.77%, with a J SC of 21.07 mA/cm², V OC of 1.29 V, and FF of 83.48%, while the control exhibited inferior PCE of 20.51%, with a J SC of 20.59 mA/cm², V OC of 1.25 V, and FF of 79.4% (Fig. 4b and Supplementary Fig. 22a). The EQE spectra of the corresponding devices were shown in Supplementary Fig. 22b, where the integrated J SC of 20.81 mA/cm² and 20.36 mA/cm² for target and control device, respectively, also aligned well with their J–V characteristics. The obviously enhanced J SC of the target device should be attributed to the better perovskite/C 60 interface quality after involving MLAI to mitigate the aggregated Rb-rich phase on top of perovskite surface and achieve uniform phase distribution. A cross-sectional SEM image showed the high-quality deposition of the perovskite film on the textured silicon substrate (Supplementary Fig. 23). The J–V characteristics of our representative tandem device, shown in Fig. 4c, yielded a PCE of 33.91%, with a V OC of 2.02 V, J SC of 20.22 mA/cm², and FF of 83.15% under reverse scan, and a PCE of 34.02%, V OC of 2.02 V, J SC of 20.23 mA/cm², and FF of 83.38% under forward scan, along with the SPO efficiency delivering 33.5% (Fig. 4d). The insert shows a photo of the tandem device. The J SC values derived from the integration of the top WBG subcell and the bottom silicon subcell EQE tests were 20.1 mA/cm² and 19.9 mA/cm², respectively (Fig. 4e), which closely match the J SC values obtained from the J–V measurements. Compared to reported literature values, our V OC × FF product and PCE are among the same level of highest reported values, especially that our impressive V OC of 2.02 V is the highest value ever reported, affirming the success of our MLAI-modified WBG perovskites in reducing the energy loss and advancing tandem solar cells efficiency (Supplementary Fig. 24, Supplementary Table 6). The tandem device efficiency was also independently measured in NPVM certification center. It showed a certified PCE of 33.22%, with a V OC of 2.01 V, J SC of 20.2 mA/cm², and FF of 81.98% under reverse scan; PCE of 33.16%, V OC of 2.01 V, J SC of 20.2 mA/cm², and FF of 81.83% under forward scan, along with a certified stabilized efficiency of 32.85% after MPP tracking for 300 s (Fig. 4f and Supplementary Fig. 25). In addition, our tandem cell exhibited excellent operational stability, with device encapsulated devices maintaining a T 90 lifetime for 1135 h under 1-sun/65 °C (ISOS-L-2) (Fig. 4g). While the results are promising, the efficiency parameters of our as-used silicon bottom cells are far below the world’s most advanced level (Supplementary Fig. 26), which limited the final tandem device efficiencies in this study. Note that we achieved high FF of 83% for the tandem cell, indicating that the limitations observed in the standalone bottom cell—whether resistive or related to defects—are effectively compensated within the integrated, current-matched tandem stack57,58. There is still much room for the further optimization and advancement. Apart from silicon bottom cell optimization, future work could focus on optimizing the functional layers of the semitransparent conductive electrodes such as IZO, electron transport layers of C 60 and SnO 2, to improve transmission and minimize parasitic absorption, thereby increasing the current of the overall tandem device.

Methods
Materials
Dimethylformamide (DMF), dimethyl sulfoxide (DMSO), chlorobenzene (CB), isopropanol (IPA), ethanediamine dihydroiodide (EDAI 2), ethyl acetate (EA), lead iodide (PbI 2), melamine and hydriodic acid were purchased from Sigma Aldrich. cesium iodide (CsI), formamidinium iodide (FAI), phenethylammonium iodide (PEAI), lead bromide (PbBr 2), (4-(7H-dibenzo[c,g]carbazol-7-yl)butyl)phosphonic acid (CbzNaph) were purchased from Xi’an Yuri Solar Company Limited. C 60 was purchased from Luminescence Technology. All the chemicals are used directly without further purification. The synthesize of melaminium iodide (MLAI): 1 g of melamine powder (99% purity, Sigma-Aldrich) was combined with 2 mL of hydroiodic acid (57% weight in water, Sigma-Aldrich) and 25 mL of methanol in a 50 mL three-neck round-bottom flask under a nitrogen atmosphere for 12 h with stirring. The resulting suspension was transferred to a centrifuge tube, and the white precipitate was separated from the solution via centrifugation. Subsequently, the precipitate was sequentially washed with methanol, ethanol, diethyl ether, and then placed in a vacuum oven to dry at 65 °C for 12 h.

Preparation of precursor solutions
The Rb 0.05 Cs 0.1 FA 0.85 Pb(I 0.75 Br 0.25 ) 3 perovskite precursor (1.75 M), was prepared by dissolving a mixture of PbI 2 (504.2 mg), PbBr 2 (240.8 mg), FAI (255.9 mg), CsI (45.5 mg), and RbI (18.6 mg) in a mixtures solvent of DMF and DMSO (1 mL, v:v = 4:1). For the target perovskite precursor solution, melaminium iodide (MLAI) (4 mM, ~1 mg mL −1) was directly added into the precursor, and the precursor was stirred at room temperature in a N 2 glovebox for 4 h. The self-assemble hole selective molecule (SAM) solution was prepared by dissolving CbzNaph in ethanol solvent with the concentration of 1 mg mL −1.

Single-junction perovskite solar cells device fabrication
The device structure is: Glass/FTO/CbzNaph/Perovskite/EDAI 2 +PEAI/C 60 /SnO 2 /Cu. Rb 0.05 Cs 0.1 FA 0.85 Pb(I 0.75 Br 0.25 ) 3 was used as the perovskite composition. FTO glass was cleaned using detergent, deionized water, acetone and IPA before drying in a N 2 flow. Next, the FTO substrate was treated by ultraviolet ozone for 15 min and then transferred to N 2 glovebox. 1 mg mL −1 CbzNaph solution was deposited on the FTO at 3000 r.p.m. for 30 s, and annealed at 100 °C for 10 min. For the perovskite layer deposition, 1.75 M perovskite precursor solution was deposited on the FTO/SAM substrate at 2500 r.p.m. for 40 s with the accelerated speed of 700 r.p.m./s. Then the obtained perovskite precursor wet film was put on a low temperature (e.g., 20 °C) substrate and moved to a vacuum chamber, and pumped down to around 20 Pa for 60 s. After that, the films were annealed at 100 °C for 20 min. After that, 1 mg mL −1 EDAI 2 and 1 mg mL −1 PEAI mixture solution in IPA was coated on the perovskite surface at 4000 r.p.m. for 20 s and annealed at 100 °C for 5 min. 20 nm C 60, 15 nm tin oxide (SnO 2) were finished by thermally evaporation and atomic layer deposition, respectively. Finally, 150 nm Cu (2.0 Å s-1) was thermally evaporated as the electrode using a shadow mask. The device area was 0.108 cm 2, and a 0.0725 cm 2 non-reflective mask was used to define the accurate active cell area when conducting J–V measurements. For 1 cm² semi-transparent single-junction device, the device stack is ITO/CbzNaph/perovskite/EDAI 2 +PEAI/C 60 /SnO 2 /indium zinc oxide (IZO)/silver (Ag) grid. The thicknesses of C 60, SnO 2, and IZO were 15 nm, 15 nm, and 40 nm, respectively.

Perovskite-silicon tandem solar cells fabrication
Our silicon bottom cell was silicon heterojunction (SHJ) solar cells, with the layer stack shown in Fig. 4a. It was provided by Zhejiang AIKO Solar Technology Co., Ltd., and it features a front-side micro-texture with randomly distributed pyramids averaging 600–800 nm in height, and rear-side texture of 700–800 nm. The silicon bottom cell was laser-scribed with dimensions of 20 mm × 20 mm before using for perovskite-silicon tandem solar cells fabrication. 15 nm ITO recombination layer was served as the intermediate recombination layer. The wide bandgap perovskite top cell was deposited onto the silicon substrates. The perovskite processing details were the same as the single junction devices fabrication. Then 40 nm IZO was deposited by radio frequency magnetron sputtering at room temperature using a IZO target (In:Zn = 9:1), processed at 60 W and 3 mtorr in pure Argon. Silver (800 nm) at the rate of 4.0 Å s −1 was thermally evaporated through a designed mask as metal grid and the Ag electrode (800 nm) on the back of the silicon bottom cell was also deposited using this process. Finally, 100 nm MgF 2 as an antireflective layer was thermally evaporated at the rate of 2.0 Å s −1.

Characterization of perovskite layers
The XRD and GIXRD patterns of perovskite films were characterized on a D8 ADVANCE system (Bruker Nano) using Cu Kα radiation (λ = 1.5418 Å). Perovskite films were prepared under the SAM covered substrates. XPS was performed on a Thermo Scientific ESCALab 250Xi using 200 W monochromated Al Kα (1486.6 eV) radiation. A 500 μm X-ray spot was used for XPS analysis. The base pressure in the analysis chamber was about 3 × 10 −10 mbar. The SEM images and EDS mapping was acquired using a field-emission SEM (FEI NanoSEM 650). SEM uses an electron beam accelerated at 500 V to 30 kV, enabling operation at a variety of currents. An accelerating voltage of 10 kV was used throughout EDS measurement, the spectrum were acquired from a series of maps (10 × 10 μm 2) on each sample. PL spectra were measured by a FLS1000 spectrometer. The excitation wavelength was set at 475 nm. TRPL spectra were measured by F900 spectrometer with a 375 nm pulsed laser (EPL-375). Two half device structures were used: Glass/FTO/SAM/Perovskite and Glass/Perovskite/Modification layer/C 60. The decay curves were fitted by the bi-exponential equation:
ϕ(t) = t 0 + A 1 e −t/τ 1 + A 2 e −t/τ 2 (1)
τ average = (A 1 τ 1 2 + A 2 τ 2 2) / (A 1 τ 1 + A 2 τ 2) (2)
where ϕ(t) is the time-dependent PL photon flux, τ 1 represent the fast decay lifetime corresponding to carrier extraction process and τ 2 represent the slow decay lifetime relating to radiative recombination process. Based on TRPL results, we used the differential lifetime τ PL(t) to have further analysis, it defined as the inverse negative slope of the logarithmically plotted PL photon flux ϕ(t):
τ PL(t) = − (d In(ϕ(t)) / dt) −1 (3)
In-situ PL for monitoring perovskite growth in initial stage was carried out by capturing the PL signal every 10 ms using Portman portable Raman spectrometer (PR532-T) produced by Ocean Optics. In-situ time-dependent PL tracking for perovskite film stability comparison was measured under white LEDs (about 2-sun light intensity). The PL spectra were scanned and recorded every ten minutes.

Device characterizations
J–V characteristics of WBG single-junction PSCs were recorded by a Keithley 2400 Source Meter and under a solar simulator from EnliTech (SS-X50R); a KG-5 Si photodiode was used to calibrate the solar simulator light intensity to 1-sun. The photovoltaic cells were measured both in reverse scan (1.32 V → 0 V, step 0.02 V) and forward scan (0 V → 1.32 V, step 0.02 V) in a nitrogen glovebox, under room temperature. For dark I–V measurement, the voltage continuously sweeps from −1.0 to 1.5 V with a step of 0.01 V, using a Keithley 6482 Source Meter. J–V characteristics of perovskite/Si tandem solar cells were recorded by a Keithley 2450 Source Meter and under a solar simulator from EnliTech (SS-PST220R); a KG-3, and BL-7 silicon photodiode was used to synergistically calibrate the solar simulator light intensity to 1-sun. The tandem devices were measured both in reverse scan (2.04 V → 0 V, step 0.02 V) and forward scan (0 V → 2.04 V, step 0.02 V) in a nitrogen glovebox, under room temperature. Stabilized power output (SPO) results were performed under maximum power point (MPP) by fixing the cell at the MPP voltage (determined by the J–V scan in both scan directions) and tracking the changes of current over time, under room temperature. The device EQE was measured by EnliTech EQE measurement system (QE-R3018). Standard silicon and germanium solar cell were used as the references for the EQE measurement. For the tandem device measurement, bias illumination was used that the perovskite top cell was measured under an 850 nm continuous biased light and the silicon bottom cell was measured under a 470 nm continuous biased light. For the light-intensity-dependent J–V measurements, the data was analyzed with the formula:
V OC = (nk T / q) In(I) + V 0 OC (4)
where n is the ideality factor, k is the Boltzmann constant, T is the temperature, q presents the unit charge, I presents the light intensity, and V 0 OC is the constant.

Device stability measurements
The testing system was supported by Wuhan 91PVK Solar Technology Co. Ltd, China. The illumination was provided by a LED-based solar simulator, calibrated to approximately 1-sun intensity, the light spectrum was shown in Supplementary Fig. 20. For the encapsulated single-junction devices, the devices were put into a tightly closed chamber and N 2 flow was maintained during the testing. Throughout the operational tracking duration, a hot plate served as the thermal source to maintain the chamber temperature at approximately 65 °C, a temperature sensor was put into the chamber to monitor the temperature. SPO was monitored by applying a constant bias voltage of 800 mV to the single-junction device. The 85 °C thermal stability was carried out by putting unencapsulated device onto a hotplate with 85 °C, and cooled down to room temperature before J–V measurements. In the case of silicon/perovskite tandem device stability measurements, the devices were encapsulated in a nitrogen-purged glovebox using a desiccant impregnated polyisobutylene (PIB) edge seal between two glass sheets, the package method was similar as previously reported59. The packaged tandem devices were put onto a 65 °C hotplate to maintain the temperature and a bias voltage as 1500 mV was used for the tracking.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "No further comments. Good to publish as is.", "reviewer_2": "I was reviewer #2 of the original manuscript that was submitted to Nature Materials.\nHan et al. have addressed many of the points raised by me and the other reviewer. I appreciate especially the clarification\nregarding the apparent absence of surface passivation by MLAI (additional Fig. 13 in the SI), as well as additional\ndiscussion of the working mechanism of MLAI. Furthermore, I think that the scope and potential impact of the presented\nresults are well aligned with Nature Communications. However, some of the issues mentioned in the previous results\nrequire further clarification to make the manuscript ready for publication:\nI had pointed out in my previous review that S. Luo et al. recently demonstrated and discussed in detail the benefits of\nmelamine as additive in CsMAFA (DOI:10.1039/D4EE04692G). This reference has now been added (Ref. 37), but is just\nsummarily discussed together with 4 other papers on p.4: “Amine groups and triazine core functionalized materials stand out\nas promising regulators, due to their ability to form strong electrostatic interactions through hydrogen bonding with the lead\nhalides octahedra and A-site cations, thereby reinforcing lattice integrity and enabling defect passivation simultaneously\n[8,34-37].” I think that this somewhat obfuscates the importance of these previous publications. The previously published\nresults should be discussed more carefully, e.g. using the arguments now provided in the response letter: The similarities\nand differences between the results presented in the present manuscript (stronger coordination of melamine to the α-\nRbFACsPbI3 lattice lowering the formation energy and nucleation barrier of the target multi-cation phase, suppressing\nformation of δ-RbPbI3) and Luo’s results (melamine suppressing PbI2 formation by bonding with uncoordinated lead, while\nthe amino groups coordinate with the iodide anions and effectively “trichelate” PbI2) are obvious and relevant. Furthermore,\nthe precedence of the results in Ref. 37 and other works must be made clear to the reader.\nAnother point that merits more careful discussion are the DFT calculations, which are the basis for claiming the beneficial\neffects of MLAI, i.e. the suppression of δ-RbPbI3. The calculated formation energies and nucleation barriers are discussed\nas being representative for the *nucleation phase* (l.145ff), i.e. for the case of nanocrystallites of just a few unit cells in size.\nHowever, the DFT calculations represent the *bulk* case, i.e. the case of a slab of δ-RbPbI3 with laterally infinite dimensions\non an equally laterally infinite slab of perovskite. This should be clarified in the main text, and it should be discussed how the\ncalculations for the bulk case translate to the nucleation phase. Furthermore, the authors mention in their response letter the\nratio of the target concentration of melamine (4 mM) vs. that of the perovskite precursor solution (1.75 M), thus a factor of\n~500. This can indeed rationalize the use of just one MLAI molecule in a calculation of the initial nucleation process – if the\nMLAI is diluted in the film, i.e. not segregated on the crystallite surfaces. Are there any indications that this is the case? On\nthe contrary, the used model and the sketch in Fig 1h imply that the MLAI segregates on the perovskite surface (and\nprobably on inner surfaces, grain boundaries). For this case, one would have to consider the *surface* coverage by MLAI of a *large*\n(tens or hundreds of nm) perovskite crystallite. This would probably correspond to a denser coverage than the one\nshown here – with obvious implications for e.g. the orientation of the molecule, surface energy etc. Please discuss these\npoints carefully and motivate your approach, or revise and redo the calculations, if necessary.\nFurthermore, the following points should be considered:\n- Fig.1a: The caption reads now “...the marked aggregated material was confirmed to be 91 RbPbI3 by X-ray diffraction\n(XRD) results.” However, XRD does not provide local information on the µm scale. Thus, this should be rephrased (also in\nthe main text): It can be assumed that what we see here are the RbPbI3 precipitates, but you do not show direct proof.\n- l. 115: “Although RbCsFA composition is potential, we found there were phase segregation existed in the original\nperovskite film.” – please rephrase and explain what is meant by “composition is potential”\n- l. 156: “…clear peak shifts of the Rb 3d, Pb 4f, 156 Cs 3d and N 1s toward higher binding energies was observed, which\nmeans that their Fermi levels are more distant to the core levels indicating that there was strong interaction between\nmelamine with lead halides octahedra and A-site cations and it should be responsible for the formation energy changes” -\nPlease reconsider your arguments and rephrase: a) there is just one Fermi level for a homogeneous, conductive material. B)\nIn the absence of any chemical changes (changes in the core level spectral shapes), the shift of EF with respect to the\ncore levels probably indicates a different charge state (different density of electrons and\nholes) at the sample surface - not a\nchange of “formation energies”.\n- l. 227: “compared to control” -> compared to *the* control (sample).\n- The additional information that is now provided about the silicon bottom cell, i.e. that it is a SHJ; pyramid height on the front\n(what about the rear side?), the laser scribing etc., should be summarized in the corresponding “fabrication” section in the SI\n- The question raised in the previous review, how a FF >83% in the tandem can be achieved despite the low FF of the\nbottom cell and almost current matched subcells, has been discussed (somewhat unconclusively) in the response letter, but\nnot in the manuscript. Please add a brief discussion. Note, that your statement from the response letter that “there is no\nstraightforward correlation between the FF of a tandem device and that of each subcell“ is not correct: cf. for example DOI:\n10.1039/C9SE00120D."}, "author_response": "Response to Reviewers’ comments\nWe thank the reviewers for their time and efforts on evaluating our manuscript\n“Uniform phase distribution of wide bandgap perovskite for high-performance\nperovskite-silicon tandem solar cells” (NCOMMS-25-69423A). The comments help\nus further elevate the quality of the manuscript. We have addressed the reviewer’s\ncomments point-by-point and modified the manuscript accordingly. We hope our\nfurther elaboration and clarification could fully address the reviewer’s concerns.\nReviewer #1 (Remarks to the Author):\nNo further comments. Good to publish as is.\nResponse: We greatly appreciate the reviewer for the support of our work.\nReviewer #2 (Remarks to the Author):\nI was reviewer #2 of the original manuscript that was submitted to Nature Materials.\nHan et al. have addressed many of the points raised by me and the other reviewer. I\nappreciate especially the clarification regarding the apparent absence of surface\npassivation by MLAI (additional Fig. 13 in the SI), as well as additional discussion of\nthe working mechanism of MLAI. Furthermore, I think that the scope and potential\nimpact of the presented results are well aligned with Nature Communications. However,\nsome of the issues mentioned in the previous results require further clarification to make\nthe manuscript ready for publication:\nI had pointed out in my previous review that S. Luo et al. recently demonstrated and\ndiscussed in detail the benefits of melamine as additive in CsMAFA\n(DOI:10.1039/D4EE04692G). This reference has now been added (Ref. 37), but is just\nsummarily discussed together with 4 other papers on p.4: “Amine groups and triazine\ncore functionalized materials stand out as promising regulators, due to their ability to\nform strong electrostatic interactions through hydrogen bonding with the lead halides\noctahedra and A-site cations, thereby reinforcing lattice integrity and enabling defect\npassivation simultaneously [8,34-37].” I think that this somewhat obfuscates the\nimportance of these previous publications. The previously published results should be\ndiscussed more carefully, e.g. using the arguments now provided in the response letter:\nThe similarities and differences between the results presented in the present manuscript\n(stronger coordination of melamine to the α-RbFACsPbI3 lattice lowering the\nformation energy and nucleation barrier of the target multi-cation phase, suppressing\nformation of δ-RbPbI3) and Luo’s results (melamine suppressing PbI2 formation by\nbonding with uncoordinated lead, while the amino groups coordinate with the iodide\nanions and effectively “trichelate” PbI2) are obvious and relevant. Furthermore, the\nprecedence of the results in Ref. 37 and other works must be made clear to the reader.\nResponse: We thank the reviewer for the comment and we acknowledge that it would\nbe clearer to have more discussions with the previously published literatures with\nMLAI. Thus, we have added the following discussions in the revised manuscript:\n“To this end, we selected melaminium iodide (MLAI), a triazine heterocyclic\ncompound as the regulator, which was a universal and effective additive material in\nsome previously reported results. For example, it behaved as a good hydrogen bond\nacceptor and Lewis base to interact with both the organic cation and Lewis acidic lead\niodide (PbI ), which resulted in good passivation effect to the perovskite layer35,36. It\n2\nwas also adopted to construct porous and disordered PbI morphology and then help\n2\nconvert the residual of PbI to perovskite phase in two-step sequential deposited\n2\nperovskite films, which was based on the tridentate chelation effect between melamine\nand PbI 37.\nRbCsFA WBG perovskites. The strong coordination of melamine to the alloyed-cation\nperovskite phase can lower its formation energy and the nucleation barrier, suppressing\nthe formation of δ-RbPbI , thus mediate homogeneous nucleation and distribution of\n3\nour rubidium-containing WBG perovskite phases.”\nWe hope the clarification and revise could adequately satisfy the reviewer and be clear\nto the readers.\nAnother point that merits more careful discussion are the DFT calculations, which are\nthe basis for claiming the beneficial effects of MLAI, i.e. the suppression of δ-RbPbI3.\nThe calculated formation energies and nucleation barriers are discussed as being\nrepresentative for the *nucleation phase* (l.145ff), i.e. for the case of nanocrystallites\nof just a few unit cells in size. However, the DFT calculations represent the *bulk* case,\ni.e. the case of a slab of δ-RbPbI3 with laterally infinite dimensions on an equally\nlaterally infinite slab of perovskite. This should be clarified in the main text, and it\nshould be discussed how the calculations for the bulk case translate to the nucleation\nphase.\nFurthermore, the authors mention in their response letter the\nratio of the target concentration of melamine (4 mM) vs. that of the perovskite precursor solution (1.75\nM), thus a factor of ~500. This can indeed rationalize the use of just one MLAI molecule\nin a calculation of the initial nucleation process – if the MLAI is diluted in the film,\ni.e. not segregated on the crystallite surfaces. Are there any indications that this is the case?\nOn the contrary, the used model and the sketch in Fig 1h imply that the MLAI\nsegregates on the perovskite surface (and probably on inner surfaces, grain boundaries).\nFor this case, one would have to consider the *surface* coverage by MLAI of a *large*\n(tens or hundreds of nm) perovskite crystallite. This would probably correspond to a\ndenser coverage than the one shown here – with obvious implications for e.g. the\norientation of the molecule, surface energy etc. Please discuss these points carefully\nand motivate your approach, or revise and redo the calculations, if necessary.\nResponse: We thank the reviewer for the insightful comments and perspective\nregarding our DFT calculations.\n1. Constructing a cluster (nanocrystallites) model would indeed represent a more\nrigorous but also significantly more computationally demanding approach, as large-\nscale models are required. In our case, the calculations of formation energies and\nnucleation barriers involve interfacial interactions and surface energy variations. To\nsimplify the process while maintaining accuracy, we assume that both δ-RbPbI and\n3\ncorresponding to the (001) crystal plane. Accordingly, we evaluated the interfacial\nenergy of the perovskite (001) surface before and after melamine interaction using a\nslab model instead of a cluster model. This slab-based approach is well established for\nprobing specific crystal facets, as it directly represents extended surfaces. The reliability\nof the slab model has been extensively validated in surface science studies (Phys. Rev.\nB 2004, 69, 195409; Nano-Micro Lett. 2024, 16, 111; Surf. Sci. 2015, 637, 124–131; J.\nPhys.: Condens. Matter 1996, 8, 6525; Phys. Rev. B 2004, 70, 155417), confirming its\naccuracy and computational efficiency for evaluating surface energy changes induced\nby molecular interactions. In contrast, cluster models inherently contain multiple crystal\nplanes, making them more complex and unsuitable for isolating plane-specific\nphenomena such as adsorption, catalysis and molecular anchoring.\nTherefore, for purpose of interfacial energy calculations, introducing a nanocrystal\nmodel involving explicit bond formation is therefore unnecessary; the slab model alone\nprovides a sufficiently accurate and physically meaningful description.\nWe have incorporated the relevant discussion into Supplementary Note 2: “Note that\nwe adopted the slab model rather than the cluster model to perform the DFT calculation,\nbased on the purpose of the interfacial energy related study and the reliability and\ncomputational efficiency of slab model for evaluating the surface energy change\ninduced by molecular interactions.” And we hope that this clarification satisfactorily\naddresses the reviewer’s concerns\n2. Considering the small molecule size of melamine and the strong coordination of\nmelamine and lead, we think it is largely uniformly distributed in the perovskite films,\nrather than segregated on the surfaces. In Fig. 1h, we draw the sketch to show the\ndisappearance of the segregated δ-RbPbI phase on the perovskite surface, rather than\n3\ndemonstrating that the MLAI segregated on the perovskite surface, we have re-draw\nthe sketch to make it clearer.\nFurthermore, the following points should be considered:\n- Fig.1a: The caption reads now “...the marked aggregated material was confirmed to\nbe RbPbI by EDS analysis in Supplementary Fig. 5 and grazing incident\n- l. 115: “Although RbCsFA composition is potential, we found there were phase\nsegregation existed in the original perovskite film.” – please rephrase and explain what\nis meant by “composition is potential”\nResponse: In line 120-124 in the revised manuscript, we demonstrated the great\npotential of RbCsFA composition for further solar cells application and performance\noptimization, considering the better charge transport properties and device efficiencies\nwhen comparing RbCsFA and CsFA composition. It will be a good benchmark for\nfurther study, therefore, we expressed “composition is potential” in the later section.\nWe have rephrased the sentence as “Although RbCsFA composition is a potential\nperovskite absorber for further solar cells application” and hope it is clear now.\n- l. 156: “…clear peak shifts of the Rb 3d, Pb 4f, 156 Cs 3d and N 1s toward higher\nbinding energies was observed, which means that their Fermi levels are more distant to\nthe core levels indicating that there was strong interaction between melamine with lead\nhalides octahedra and A-site cations and it should be responsible for the formation\nenergy changes” - Please reconsider your arguments and rephrase: a) there is just one\nFermi level for a homogeneous, conductive material. B) In the absence of any chemical\nchanges (changes in the core level spectral shapes), the shift of Ef with respect to the\ncore levels indicates a different charge state (different density of electrons and\nholes) at the sample surface - not a change of “formation energies”.\nResponse: We thank the reviewer for the comments.\na) We have changed the expression to “the Fermi level is more distant to the core level”,\nin the revised manuscript.\nb) We agree with the reviewer that the Ef shifts with respect to the core levels indicates\na different charge state. While the XPS could also further reflect the interaction changes\nin the sample. We were aim to express the points that it is the strong interaction between\nmelamine with lead halides and A-site cations that resulted in the formation energies\nchanges. We have modified the expressions in the revised manuscript to make it clearer.\n- l. 227: “compared to control” -> compared to *the* control (sample).\nResponse: We have added “the” and “sample” in the revised manuscript.\n- The additional information that is now provided about the silicon bottom cell, i.e. that\nit is a SHJ; pyramid height on the front (what about the rear side?), the laser scribing\netc., should be summarized in the corresponding “fabrication” section in the SI\nResponse: The rear side pyramid height of the texture is about 700-800 nm. And we\nhave added the silicon bottom cell information in the “perovskite-silicon tandem solar\ncells fabrication” section in the SI.\n- The question raised in the previous review, how a FF >83% in the tandem can be\nachieved despite the low FF of the bottom cell and almost current matched subcells,\nhas been discussed (somewhat unconclusively) in the response letter, but not in the\nmanuscript. Please add a brief discussion. Note, that your statement from the response\nletter that “there is no straightforward correlation between the FF of a tandem device\nand that of each subcell“ is not correct: cf. for example DOI: 10.1039/C9SE00120D.\nResponse: We thank the reviewer for the comments. We agree with the reviewer that\nthe previous silicon bottom cell FF seems too low to guarantee an efficient tandem solar\ncell. As we mentioned in the previous response letter that we randomly selected one\npiece of silicon bottom cell to check the status, that could not be the representative\nstatus. We further checked this time by fabricating more devices. The best result we\ncould get is shown below; the FF is above 80%. Given that the high FF of our single\njunction WBG PSCs and the following possible explanations, the FF of 83% is quite\nreasonable for the tandem solar cells:\n1. Suppression of series resistance losses: Resistive losses scale with the square of the\ncurrent (I2Rs). In a tandem cell, the current at the maximum power point (I ) is\nMPP\napproximately halved compared to a single-junction device. Furthermore, the relative\nvoltage drop (ΔV/Voc) caused by a given series resistance (Rs) is significantly reduced\nbecause the tandem's Voc is roughly double that of a single junction. Consequently, a\nseries resistance that severely degrades the FF of a high-current single-junction bottom\ncell becomes largely benign in the tandem configuration.\n2. Mitigation of shunt and defect-related losses by symmetry: The tandem's current\nis limited by the lower one of the two sub-cell currents. At or near perfect current\nmatching condition, any degradation in shunt resistance or defect resistance of one sub-\ncell could be partially offset/compensated by the healthier partner, thus the tandem FF\ndoes not collapse to the lower single-junction value while yield a composite FF higher\nthan that of the weaker individual sub-cell.\n3. Increased dynamic impedance at the maximum power point: The tandem cell\noperates at a higher voltage and lower current, resulting in a significantly larger\ndynamic resistance (dV/dI) at the maximum power point. A higher impedance\nintrinsically leads to a 'squarer' J-V characteristic and thus a higher FF, independent of\nthe specific mechanisms limiting the FFs of the individual sub-cells.\nIn summary, when the tandem is almost current-matched, the tandem FF is not a simple\naverage of the sub-cell FFs but is governed by this impedance-matching interplay.\nFigure R1 (a) Representative J-V curves and (b) corresponding EQE of a single\njunction silicon heterojunction (SHJ) bottom cell (laser scribed to 20 mm×20 mm size).\nWe have updated the results in Supplementary Fig. 26, added brief discussions and cited\nrelated references in the revised manuscript as “Note that we still achieved high FF of\n83% for the tandem cell, indicating that the limitations observed in the standalone\nbottom cell—whether resistive or related to defects—are effectively compensated\nwithin the integrated, current-matched tandem stack.”\nWe sincerely hope these clarifications are adequately clear to the reviewer and readers."}, {"round": 2, "reviewer_comments": {"reviewer_2": "Thanks you very much for considering my comments and clarifying the open questions. I think that the manuscript is now\nready for publication."}, "author_response": ""}]

Decision letter