Ultrathin oxygen deficient SnO_x films as electron extraction layers for perovskite solar modules

Abstract

The design of high-quality junctions capable of efficiently extracting carriers from perovskite-based absorbers is key in the transition from lab-scale devices to modules. In the so-called n–i–p configuration, SnO₂ nanoparticle (np-SnO₂) films have been thoroughly investigated as electron transporting layers (ETLs) in view of their good optimal band alignment, chemical stability and appropriate surface chemistry for nucleating high-quality perovskite films. In this report, we show for the first time that np-SnO₂ films are characterized by a heterogeneous surface electronic landscape and introducing quasi-monoenergetic conformal layers between the transparent conducting oxide (TCO) and the np-SnO₂ film can lead to significant improvement in perovskite solar modules. These SnO_x extraction layers are developed using a highly innovative plasma-modified atomic layer deposition (PMALD) tool, which enables tuning the Sn : O ratio, conductivity, and effective work function. Energy-filtered photoemission electron microscopy (EF-PEEM) shows a remarkably homogeneous surface electronic landscape of PMALD SnO_x. We examine the impact of PMALD-SnO_x in an n–i–p device configuration, with poly(triarylamine) (PTAA) as the hole transporting layer, which leads to the improvement in perovskite module power conversion efficiency from 17.9% to 20.1%, with an active area of 23.2 cm². Furthermore, the devices retained 92% of their initial efficiency for 2700 h at 85 °C and 85% relative humidity and 96% for 1000 h under continuous illumination with maximum power point tracking.

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Introduction

The transition from high performance laboratory scale devices to interconnected modules not only requires scalable methods and deposition of organo-halide perovskite precursors but also requires homogeneous hole transporting layers (HTLs) and electron transporting layers (ETLs).^1–7 The so-called n–i–p perovskite solar cell (PSC) architecture is the most common configuration investigated at module levels given the fact that ETL deposition often requires high processing temperatures, leading to devices with higher chemical stability, good band alignment and high overall performance.^8–11

SnO₂ has emerged as one of the best performing ETLs based on its high transmittance in the visible range, electron mobility, and chemical stability.^12–17 A variety of methods have been proposed for the deposition of SnO_x films including pulsed laser deposition, chemical bath deposition, atomic layer deposition (ALD), e-beam evaporation, and electrochemical deposition.^12–14 Atomic layer deposition (ALD), including both thermal ALD and plasma-enhanced ALD (PEALD), is widely used due to its ability to produce highly conformal, ultrathin films with excellent thickness control, which is beneficial for ensuring uniform coverage on complex device architectures.^18–25 Thermal ALD and plasma-enhanced ALD (PEALD) deposited SnO_x films commonly generate highly conformal ultrathin films with a high oxygen content, which surprisingly show poorer device performance in PSCs in comparison to nanoparticle SnO₂ (np-SnO₂) films deposited by solution-based methods.²⁶ However, several reports have shown that introducing ALD SnO₂ films prior to deposition of np-SnO₂ films promotes improvements in device performance.^27–32 These interesting observations raise important questions about the role of morphology and electronic structure of the ETL construct and how parameters such as oxygen content can affect the optoelectronic properties and device performance.

In this report, we demonstrate that tuning oxygen content of highly conformal SnO_x electron extraction layers, deposited prior to the nanoparticle SnO₂ (np-SnO₂) layer, leads to significant improvement in module performance and stability. These extraction layers are deposited by a novel ALD approach incorporating an argon (Ar) plasma treatment step in between two different tin (Sn) precursor pulse steps, followed by a deionized water pulse step. This approach, referred to as plasma-modified ALD (PMALD),³³ operates in an entirely different fashion to conventional PEALD, and is capable of generating SnO_x films, with x values as low as 1.2. For the first time, the surface electronic landscape of these ultrathin SnO_x films is mapped by energy-filtered photoemission electron microscopy (EF-PEEM), showing a remarkable narrow distribution of work function, which contrasts with a significantly broader distribution observed for the np-SnO₂ ETL layer. We demonstrate that this quasi-monoenergetic electron extraction layer improves the efficiency and fill factor of n–i–p devices featuring (FAPbI₃)_0.95(MAPbBr₃)_0.05 as the solar absorber and poly(triarylamine) (PTAA) as the hole transporting layer, with 25 cm² modules achieving PCE values over 20%.

Results and discussion

Fig. 1a schematically shows how the pulse sequence for conventional ALD and PEALD differs from the process designed for PMALD. All processes start with a pulse of the Sn precursor, tetrakis(dimethylamino)tin(IV) (TDMASn), followed by an Ar purge step. Full details of the various ALD protocols and thin-film deposition can be found in Experimental Section S1 (ESI†). In thermal ALD and PEALD, oxygen is introduced by water or plasma activated oxygen pulses, respectively, followed by another Ar purge before repeating the sequence. By contrast, PMALD undergoes an Ar-plasma step with varying power, followed by another Ar purge – TDMASn pulse – Ar purge sequence before the water pulse is introduced. As illustrated in Fig. 1b, the PMALD strategy leads to significant changes in the O/Sn ratio as determined by Rutherford backscattering spectrometry (RBS). Thermal ALD (displayed as 0 W) shows an O/Sn ratio of 2.17, consistent with the Sn(IV) oxidation state. Moreover, introducing a 200 W Ar plasma step, (200 W PMALD) leads to oxygen deficient films with an O/Sn ratio as low as 1.38. The O/Sn ratio in PEALD films exhibits high oxygen content, varying from 2.27 at 100 W O₂ plasma power to 2.40 at 300 W.

Fig. 1 Composition and carrier transport properties of SnO_x films obtained by various ALD approaches. (a) Comparison of growth sequence for SnO_x by thermal ALD, plasma enhanced ALD (PEALD), and plasma-modified ALD (PMALD). (b) O/Sn ratio in SnO_x thin-films obtained using various ALD routes as a function of the plasma (O or Ar) power. (c) Carrier concentration, mobility, and resistivity, as probed by Hall measurements, in terms of water pulse time for thermal ALD, oxygen plasma in PEALD and argon plasma power in PMALD.

AFM images of the PMALD SnO_x films (Fig. S1†) deposited with various plasma power values onto Si wafers revealed that the films are composed of nanoscale grains in highly compact films with apparent roughness below 3 nm. Interestingly, we can see that increasing the plasma power shows a slight increase in grain size. The apparent roughness of the np-SnO₂ film is comparable to that of the PMALD SnO_x films.

Carrier concentration, mobility, and resistivity of the SnO_x films estimated from Hall measurements are displayed in Fig. 1c. These parameters are plotted as a function of the plasma power (Ar or O₂) and the water pulse time in the case of thermal ALD. Carrier concentrations in thermal ALD films are several orders of magnitude lower than those in PEALD and PMALD films, with the latter achieving values in the range of 10¹⁹ cm⁻³ for 100 W and 200 W PMALD SnO_x. Interestingly, the carrier mobility increases from 28.6 ± 9 to 58.2 ± 7 cm² V⁻¹ s⁻¹ upon increasing the Ar plasma power from 100 to 200 W, decreasing to 19.1 ± 6 cm² V⁻¹ s⁻¹ at 300 W for PMALD films. PEALD films show a monotonic decrease in mobility from 50.1 ± 6 to 16.5 ± 5 cm² V⁻¹ s⁻¹ with increasing O₂ plasma power from 100 to 300 W. The resistivity values for thermal ALD and PEALD films, under all conditions, are in the range of 10³ to 10⁴ Ω cm, while PMALD films show values that are 4 to 5 orders of magnitude lower.

These trends in composition and electronic properties can be rationalized in terms of ligand removal at the surface of SnO_x during the Ar-plasma step. The generated dangling bonds on the growth surface act as adsorption sites for further TDMASn molecules in the second pulse.³³ This mechanism will also promote chemisorption of Sn precursors with different terminations and hence different reactivity towards H₂O. The Ar plasma can result in the reduction of Sn⁴⁺ to Sn²⁺, owing to the partial removal/decomposition of ligands in the first TDMASn step, which will also increase oxygen vacancies and enhance n-type conductivity. It is interesting to note that Koida et al. recently concluded that amorphous SnO_x with low oxygen content, used as TCO in Si heterojunction devices, can be generated by reactive plasma deposition, while ALD techniques are unsuited for this approach as they operate at temperatures below 200 °C.³⁴ However, our results also show that employing plasma power above 200 W can produce electronic defects that compromise carrier mobility. Previous studies employing PEALD have also shown deterioration in carrier mobility at high plasma power.^18,35

Fig. 2a–c contrast the local effective work function (WF) maps of np-SnO₂ (used as the ETL) and the PMALD SnO_x obtained at 0 and 200 W plasma step, as probed by EF-PEEM over a field of view of 22 μm × 22 μm. As described in the ESI (Section S1†), we have conducted systematic surface pre-treatment to eliminate the carbon and oxygen contamination prior to recording the EF-PEEM maps. The first key observation is the 0.4 eV difference in the mean WF between the np-SnO₂ and the PMALD SnO_x films, which reflects significant differences in surface dipoles associated with Sn electronic configuration. Introducing the 200 W plasma step in PMALD leads to a slight decrease in the work function which can be partly rationalized in terms of the observed increase in carrier concentration (Fig. 1c). Fig. 2d–f also show that np-SnO₂ has a broader distribution of WF values observed in comparison to PMALD SnO_x, with the former displaying a tail towards values as low as 3.5 eV. These areas of low WF, which may act as shunting paths, may arise from complex electronic and geometric factors associated with the surface chemical environment of Sn in np-SnO₂. Given that EF-PEEM probes valence band electrons, it is rather complex to extract precise information on various chemical environments. Examples in the literature of Sn containing semiconductors have shown areas of low work functions associated with surface confined Sn phases^36–38 which contributes to shunting paths as well as disorder in opto-electronic properties (Urbach tails).^39,40 These measurements also highlight the excellent electronic homogeneity of the SnO_x films obtained by PMALD, which reflects their compositional homogeneity.

Fig. 2 Energy-filtered photoemission electron microscopy (EF-PEEM) maps of the np-SnO₂ and PMALD SnO_x films obtained using a monochromatic He I (21.2 eV) excitation source. 22 μm × 22 μm work function (WF) maps extracted from EF-PEEM data obtained for the (a) np-SnO₂ ETL layer, (b) ALD SnO_x film (PMALD – 0 W), and (c) 200 W PMALD. (d)–(f) Histograms of the WF values of the corresponding EF-PEEM maps.

Prototype cells with n–i–p configuration, as illustrated in Fig. 3a, were constructed using np-SnO₂ as the ETL, (FAPbI₃)_0.95(MAPbBr₃)_0.05 as the perovskite solar absorber and PTAA as the HTL. Fig. 3b compares the performance of the best cells with thermal ALD (0 W) and 200 W PMALD SnO_x layers as electron extraction layers, as well as devices without extraction layers. The best device without ALD extraction layers showed an open-circuit voltage (V_OC) of 1.08 V, a short-circuit current density (J_SC) of 23.6 mA cm⁻², a fill factor (FF) of 77.8%, and a PCE of 19.8%. The introduction of a 200 W PMALD extraction layer resulted in the champion device featuring a V_OC of 1.13 V, a J_SC of 23.7 mA cm⁻², an FF of 81.5%, and a PCE of 21.8%. As shown in Fig. 3c, the external quantum efficiency (EQE) spectrum of the champion device shows an integrated J_SC of 24.1 mA cm⁻², which is within 1.7% of the measured J_SC value, while a bandgap energy of 1.52 eV was obtained from the inflection point of the absorption threshold of the EQE spectrum. The champion device exhibited very small hysteresis as illustrated in Fig. S2,† with PCE values estimated from forward and reverse J–V scans within 3.7% (Table S1†).

Fig. 3 Impact of ALD SnO_x extraction layers on cell device parameters. (a) n–i–p device stack, (b) J–V curves of champion devices with and without ALD extraction layers, and (c) external quantum efficiency spectrum and integrated short-circuit current for the champion device featuring a 200 W PMALD layer. Box plots of (d) open-circuit voltage (V_OC), (e) short-circuit current density (J_SC), (f) power conversion efficiency (η), (g) fill factor (FF), (h) series resistance (R_S), and (i) shunt resistance (R_SH).

The effect of the PMALD layer on device performance can be better visualized in terms of the dispersion of PV parameters as shown by the box plots in Fig. 3d–i (Table S2†). We can see a significant decrease in dispersion and an increase in the mean FF, V_OC and PCE upon introducing ALD SnO_x as the electron extraction layer. This observation can be linked to the narrow spatial distribution of WF values shown in Fig. 2, which is a consequence of the highly homogeneous composition film of the PMALD films. The device metrics improve with increasing plasma power up to 200 W, while a substantial drop in performance is observed upon increasing the power to 300 W. Interestingly, this drop in performance is consistent with the change in resistivity, mobility, and carrier concentration of the ALD layer extracted from the Hall measurements (Fig. 1c). Another device metric affected by the extraction layer is the series resistance (R_S) as shown in Fig. 3h. Fig. S3† displays measurements of current vs. voltage across ALD and np-SnO₂ films, confirming that the oxygen deficient 200 W PMALD films has the lowest series resistance value in agreement with the data in Fig. 3h. These observations strongly suggest that introducing a homogeneous ultrathin layer with a very smooth electronic landscape significantly improves the quality of the perovskite device, while the plasma power in PMALD tunes the efficiency of carrier extraction leading to improvement in FF. As shown in Fig. S4,† the improvement in carrier extraction brought about by the PMALD films also manifests itself by a decrease in room temperature PL efficiency as well as a decrease in the time constant of time-resolved photoluminescence spectroscopy. Fig. S5† compares the key device metrics as a function of the extraction layer thickness (200 W), which can be tuned between 5 and 20 nm by the number of PMALD cycles. The comparison clearly shows that the extraction layer thickness mainly affects the device FF. There is a systematic increase in FF with increasing SnO_x film thickness between 5 and 10 nm, while the 20 nm film shows a significant decrease which we link to an increase in series resistance (Table S3†). Furthermore, Fig. S6† compares representative dark J–V curves of modules with and without the 200 W PMALD SnO_x electron extraction layer. The significantly lower current observed in the presence of the 200 W PMALD SnO_x layer further demonstrates its role as an electrical barrier preventing physical contact between ITO and the perovskite layer. This barrier reduces undesirable current leakage, which is often a source of efficiency losses in perovskite devices. The suppression of leakage current not only improves the device stability but also contributes to a higher overall power conversion efficiency, as it minimizes recombination losses associated with unintended current pathways between the ITO and perovskite.^41,42

Finally, Fig. 4a compares the performance of 5 cm × 5 cm modules, with an aperture area of 24.5 cm² and an active area of 23.2 cm², with and without the (reference) oxygen deficient 200 W PMALD extraction layer, and the performance metrics are summarized in Table S4.† The data show a small increase in V_OC and J_SC in the presence of the 200 W PMALD SnO_x extraction layer; however, the main difference arises in the FF increase from 71.0 to 74.8%, which leads to an increase in PCE from 17.9% to 20.1%. As illustrated in Fig. 4b, the module performance is among the highest reported in the literature for PTAA-based modules in terms of PCE. Table S5† provides a more detailed comparison of module architecture and performance metrics, showing that parameters such as FF and V_OC are among the highest for PTAA based devices.⁴³Fig. 4c shows the stability of encapsulated devices exposed to extreme conditions of 85 °C and 85% relative humidity for 2700 hours. Devices with the oxygen deficient 200 W PMALD SnO_x layer retained 92% of their initial efficiency, whereas the device with only the np-SnO₂ layer retained 83% of its initial efficiency. Enhanced working stability is also observed for encapsulated devices with the 200 W PMALD SnO_x layer retaining 96% of their initial efficiency compared to the reference retaining 72% under continuous LED illumination with maximum power point tracking (MPPT) for 1000 h (Fig. 4d).

Fig. 4 Module performance and stability. (a) Illuminated J–V curves of 5 cm × 5 cm modules with an np-SnO₂ layer (reference) and upon including a 200 W PMALD SnO_x electron extraction layer under 1 sun irradiation. (b) PCE values of the highest performance perovskite modules as a function of the device area, including this work. (c) Long-term damp-heat stability tests at 85 °C and 85% relative humidity. (d) Maximum power point tracking measurement of the reference device and the device with PMALD SnO_x (200 W) under continuous LED illumination.

Conclusion

In conclusion, we have shown for the first time that oxygen deficient SnO_x obtained by plasma-modified ALD acts as a highly efficient quasi-monoenergetic electron extraction layer in high performance n–i–p PSC devices and modules. By inserting an argon plasma pulse with varying power between two TDMASn pulses during the ALD growth sequence, thin SnO_x films, used as an additional electron transporting layer on top of np-SnO₂, featured a narrower distribution of local effective work function values than np-SnO₂, at a value of approximately 0.4 eV higher. These films efficiently extract electrons from the ETL, improving the overall series resistance and fill factor of the PSC device. The PMALD SnO_x-based PSC also shows high stability at 85 °C and 85% relative humidity retaining 92% of its initial efficiency for 2700 h, and high working stability retaining 96% for 1000 h under continuous illumination with MPPT. Employing PMALD SnO_x between np-SnO₂ and the TCO layer can also greatly improve efficiencies over larger areas of approximately 5 cm × 5 cm, resulting in efficiencies over 20% at an active area of 23.2 cm² using PTAA as the HTL. Our study demonstrates the importance of efficient extraction layers with the appropriate band alignment in the performance and stability of PSCs, as well as the capabilities of PMALD to control composition and transport properties of the oxide layer.

Data availability

The data supporting this article have been included as part of the ESI.† No primary research results, software or code have been included.

Conflicts of interest

The authors declare no competing financial interest.

Acknowledgements

This research was supported by the National R&D Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Science and ICT (No. 2022K1A4A8A02079724) and the Korea Research Institute of Chemical Technology (KRICT), Republic of Korea (No. KS2422-10). R. E. A., A. S., J. L. and D. J. F. gratefully acknowledge the support from the Engineering and Physical Sciences Research Council (EPSRC) through the SolPV programme (EP/V008676/1). A. S. is also grateful for the EPSRC support via the grant EP/T517872/1.

References

1. H. S. Jung and N.-G. Park, Small, 2015, 11, 10–25.
2. Z. Li, T. R. Klein, D. H. Kim, M. Yang, J. J. Berry, M. F. A. M. van Hest and K. Zhu, Nat. Rev. Mater., 2018, 3, 18017.
3. Z. Liu, L. Qiu, L. K. Ono, S. He, Z. Hu, M. Jiang, G. Tong, Z. Wu, Y. Jiang, D.-Y. Son, Y. Dang, S. Kazaoui and Y. Qi, Nat. Energy, 2020, 5, 596–604.
4. S. Razza, F. D. Giacomo, F. Matteocci, L. Cina, A. L. Palma, S. Casaluci, P. Cameron, A. D’epifanio, S. Licoccia, A. Reale, T. M. Brown and A. D. Carlo, J. Power Sources, 2015, 277, 286–291.
5. S. Ulicna, B. Dou, D. H. Kim, K. Zhu, J. M. Walls, J. W. Bowers and M. F. A. M. van Hest, ACS Appl. Energy Mater., 2018, 1, 1853–1857.
6. J.-W. Lee, S.-I. Na and S.-S. Kim, J. Power Sources, 2017, 339, 33–40.
7. H. H. Park and D. J. Fermin, Nanomaterials, 2023, 13, 3112.
8. D. Burkitt, J. Searle and T. Watson, R. Soc. Open Sci., 2018, 5, 172158.
9. S. Y. Park and K. Zhu, Adv. Mater., 2022, 34, 2110438.
10. D. Kim, H. J. Jung, I. J. Park, B. W. Larson, S. P. Dunfield, C. Xiao, J. Kim, J. Tong, P. Boonmongkolras, S. G. Ji, F. Zhang, S. R. Pae, M. Kim, S. B. Kang, V. Dravid, J. J. Berry, J. Y. Kim, K. Zhu, D. H. Kim and B. Shin, Science, 2020, 368, 155–160.
11. J. Kim, T. Hwang, B. Lee, S. Lee, K. Park, H. H. Park and B. Park, Small Methods, 2019, 3, 1800361.
12. G. Yang, X. Li, B. Zhao, C. Liu, T. Zhang, Z. Li, Z. Liu and X. Li, Langmuir, 2022, 38, 6752–6760.
13. A. Uddin and H. Yi, Sol. RRL, 2022, 6, 2100983.
14. Y. Lee, S. Lee, G. Seo, S. Paek, K. T. Cho, A. J. Huckaba, M. Calizzi, D. Choi, J.-S. Park, D. Lee, H. J. Lee, A. M. Asiri and M. K. Nazeeruddin, Adv. Sci., 2018, 5, 1800130.
15. Q. Jiang, X. Zhang and J. You, Small, 2018, 14, 1801154.
16. O. Gunawan, S. R. Pae, D. M. Bishop, Y. Virgus, J. H. Noh, N. J. Jeon, Y. S. Lee, X. Shao, T. Todorov, D. B. Mitzi and B. Shin, Nature, 2019, 575, 151–155.
17. A. J. Yun, J. Kim, B. Gil, H. Woo, K. Park, J. Cho and B. Park, ACS Appl. Mater. Interfaces, 2020, 12, 50418–50425.
18. C. Wang, D. Zhao, C. R. Grice, W. Liao, Y. Yu, A. Cimaroli, N. Shrestha, P. J. Roland, J. Chen, Z. Yu, P. Liu, N. Cheng, R. J. Ellingson, X. Zhao and Y. Yan, J. Mater. Chem. A, 2016, 4, 12080–12087.
19. C. Wang, L. Guan, D. Zhao, Y. Yu, C. R. Grice, Z. Song, R. A. Awni, J. Chen, J. Wang, X. Zhao and Y. Yan, ACS Energy Lett., 2017, 2, 2118–2124.
20. P.-H. Huang, Z.-X. Zhang, C.-H. Hsu, W.-Y. Wu, C.-J. Huang and S.-Y. Lien, Materials, 2021, 14, 690.
21. S. Seo, S. Jeong, H. Park, H. Shin and N.-G. Park, Chem. Commun., 2019, 55, 2403–2416.
22. Y. Kuang, V. Zardetto, R. Gils, S. Karwal, D. Koushik, M. A. Verheijen, L. E. Black, C. Weijtens, S. Veenstra, R. Andriessen, W. M. M. Kessels and M. Creatore, ACS Appl. Mater. Interfaces, 2018, 10, 30367–30378.
23. T. Eom, S. Kim, R. E. Agbenyeke, H. Jung, S. M. Shin, Y. K. Lee, C. G. Kim, T.-M. Chung, N. J. Jeon, H. H. Park and J. Seo, Adv. Mater. Interfaces, 2021, 8, 2001482.
24. H. H. Park, A. Jayaraman, R. Heasley, C. Yang, L. Hartle, R. Mankad, R. Haight, D. B. Mitzi, O. Gunawan and R. G. Gordon, Appl. Phys. Lett., 2014, 105, 202101.
25. H. H. Park, R. Heasley and R. G. Gordon, Appl. Phys. Lett., 2013, 102, 132110.
26. F. Gayot, E. Bruhat, M. Manceau, E. D. Vito, D. Mariolle, N. Nguyen and S. Cros, AIP Conf. Proc., 2023, 2826, 100002.
27. X. Zhang, Y. Zhou, M. Chen, D. Wang, L. Chao, Y. Lv, H. Zhang, Y. Xia, M. Li, Z. Hu and Y. Chen, Small, 2023, 19, 2303254.
28. S.-U. Lee, H. Park, H. Shin and N.-G. Park, Nanoscale, 2023, 15, 5044–5052.
29. N. Ren, C. Zhu, R. Li, S. Mazumdar, C. Sun, B. Chen, Q. Xu, P. Wang, B. Shi, Q. Huang, S. Xu, T. Li, Y. Zhao and X. Zhang, Appl. Phys. Lett., 2022, 121, 033502.
30. E. Choi, J.-W. Lee, M. Anaya, A. Mirabelli, H. Shim, J. Strzalka, J. Lim, S. Yun, M. Dubajic, J. Lim, J. Seidel, R. E. Agbenyeke, C. G. Kim, N. J. Jeon, A. M. Soufiani, H. H. Park and J. S. Yun, Adv. Energy Mater., 2023, 13, 2301717.
31. H. Li, B. Yu and H. Yu, Adv. Funct. Mater., 2024, 34, 2402128.
32. Y. Zhang, B. Yu, Y. Sun, J. Zhang, Z. Su and H. Yu, Angew. Chem., Int. Ed., 2024, 63, e202404385.
33. H. H. Park, T. J. Larrabee, L. B. Ruppalt, J. C. Culbertson and S. M. Prokes, ACS Omega, 2017, 2, 1259–1264.
34. T. Koida, T. Matsui and H. Sai, Sol. RRL, 2023, 7, 2300381.
35. M. Mattinen, J. J. P. M. Schulpen, R. Dawley, F. Gity, M. A. Verheijen, W. M. M. Kessels and A. A. Bol, ACS Appl. Mater. Interfaces, 2023, 15, 35565–35579.
36. D. Tiwari, M. Cattelan, R. L. Harniman, A. Sarua, A. Abbas, J. W. Bowers, N. A. Fox and D. J. Fermin, iScience, 2018, 9, P36–P46.
37. M. C. Naylor, D. Tiwari, A. Sheppard, J. Laverock, S. Campbell, B. Ford, X. Xu, M. D. K. Jones, Y. Qu, P. Maiello, V. Barrioz, N. S. Beattie, N. A. Fox, D. J. Fermin and G. Zoppi, Faraday Discuss., 2022, 239, 70–84.
38. R. Agbenyeke, A. Sheppard, J. Keynon, N. Benhaddou, N. Fleck, V. Corsetti, M. A. Alkhalifah, D. Tiwari, J. W. Bowers and D. J. Fermin, ACS Appl. Mater. Interfaces, 2024, 16, 35315–35322.
39. D. Tiwari, M. V. Yakushev, T. Koehler, M. Cattelan, N. Fox, R. W. Martin, R. Klenk and D. J. Fermin, ACS Appl. Energy Mater., 2022, 5, 3933–3940.
40. D. Tiwari, M. Cattelan, R. L. Harniman, A. Sarua, N. Fox, T. Koehler, R. Klenk and D. J. Fermin, ACS Energy Lett., 2018, 3, 2977–2982.
41. Y. Xiao, X. Cui, B. Xiang, Y. Chen, C. Zhao, L. Wang, C. Yang, G. Zhang, C. Xie, Y. Han, M. Qiu, S. Li and P. You, Molecules, 2023, 28, 2668.
42. K. Tvingstedt, L. Gil-Escrig, C. Momblona, P. Rieder, D. Kiermasch, M. Sessolo, A. Baumann, H. J. Bolink and V. Dyakonov, ACS Energy Lett., 2017, 2, 424–430.
43. J. Xia, P. Luizys, M. Daskeviciene, C. Xiao, K. Kantminiene, V. Jankauskas, K. Rakstys, G. Kreiza, X.-X. Gao, H. Kanda, K. G. Brooks, I. R. Alwani, Q. U. Ain, J. Zou, G. Shao, R. Hu, Z. Qiu, A. Slonopas, A. M. Asiri, Y. Zhang, P. J. Dyson, V. Getautis and M. K. Nazeeruddin, Adv. Mater., 2023, 35, 2300720.

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Footnotes

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4ta06871h

‡ These authors contributed equally to this work.

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