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Extremely low cost and green cellulose passivating perovskites for stable and high-performance solar cells Jianming Yang, Shaobing Xiong, Tianyi Qu, Yuexing Zhang, Xiaoxiao He, Xuewen Guo, Qiuhua Zhao, Slawomir Braun, Jinquan Chen, Jianhua Xu, Yanqing Li, Xianjie Liu, Chun-Gang Duan, Jian-Xin Tang, Mats Fahlman, and Qinye Bao ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b01740 • Publication Date (Web): 18 Mar 2019 Downloaded from http://pubs.acs.org on March 18, 2019

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ACS Applied Materials & Interfaces

Extremely Low Cost and Green Cellulose Passivating Perovskites for Stable and High-performance Solar Cells

Jianming Yang, Shaobing Xiong, Tianyi Qu, Yuexing Zhang, Xiaoxiao He, Xuewen Guo, Qiuhua Zhao, Slawomir Braun, Jinquan Chen, Jianhua Xu, Yanqing Li, Xianjie Liu, Chungang Duan, Jianxin Tang, Mats Fahlman, Qinye Bao* J. Yang, S. Xiong, X. Guo, Prof. C. Duan, Prof. Q. Bao Key Laboratory of Polar Materials and Devices, Department of Optoelectronics, East China Normal University, 200241, Shanghai, P.R. China E-mail: [email protected] Prof. C. Duan, Prof. Q. Bao Collaborative Innovation Center of Extreme Optics, Shanxi University, Taiyuan, Shanxi, 030006, P.R. China T. Ou, Y. Zhang, Prof. Y. Li, Prof. J. Tang Institute of Functional Nano & Soft Materials, Soochow University, Suzhou 215123, P. R. China X. He, Prof. J. Chen, Prof. J. Xu State Key Laboratory of Precision Spectroscopy, Zhongshan Campus, East China Normal University, 200062, Shanghai, P. R. China Dr. S. Braun, Dr. X. Liu, Prof. Bao, Prof. M. Fahlman Department of Physics, Chemistry and Biology, IFM, Linköping University SE58183 Linköping, Sweden

Keywords: cellulose, passivation, perovskite solar cells, efficiency, stability

Abstract The fast evolution of metal halide perovskite solar cells has opened a new chapter in the field of renewable energy. High-quality perovskite films as the active layers are essential for both high efficiency and long-term stability. Here, the perovskite films with enlarged crystal grain size and decreased defect density are fabricated by

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introducing the extremely low cost and green polymer, Ethyl cellulose (EC), into the perovskite layer. The addition of EC triggers hydrogen bonding interactions between EC and the perovskite, passivating the charge defect traps at the grain boundaries. The long chain of the EC furthermore acts as scaffold for the perovskite structure, eliminating the annealing-induced lattice strain during the film fabrication process. The resulting devices with EC additive exhibit a remarkably enhanced average power conversion efficiency from 17.11% to 19.27% and an improvement of all device parameters. The hysteresis index is found to decrease by 3 times from 0.081 to 0.027, attributed to suppressed ion migration and surface charge trapping. In addition, the defect passivation by EC significantly improves the environmental stability of perovskite films, yielding devices that retain 80% of their initial efficiencies after 30 days in ambient air at 45% relative humidity, whereas the pristine devices without EC fully degrade. This work provides a low-cost and green avenue for passivating defects that improves both the efficiency and operational stability of perovskite solar cells.

Introduction Organic-inorganic hybrid perovskites have been demonstrated as promising materials for optoelectronics, because they have large absorption coefficients, low exciton binding energy and long carrier diffusion lengths.1-3 Tremendous efforts have been focus on perovskite material design and device architecture optimization,4-6 and the certified power conversion efficiency (PCE) of the perovskite solar cells has reached as high as 23.3%,7 which is comparable to that of the established crystalline


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silicon cells of 26%.7 However, it is noted that solution-processed polycrystalline perovskite films inevitably will have under-coordinated ions as defects at grain boundaries and on surfaces. These defects work as nonradiative recombination centers and thus cause energy loss.8, 9 Besides being detrimental to device efficiency, defects are also known to be significant contributors to the intrinsic instability of perovskite films10 and hence lead to the unsatisfactory lifetime of perovskite devices under working conditions.11, 12 The poor stability is currently one of the most critical obstacles for the commercializing perovskite solar cells. The polycrystalline perovskite films has several orders of magnitude higher defect density than the single-crystal film.1 Therefore, it is essential to reduce the defects in the perovskite films to further increase both device efficiency and long-term durability.

Additive engineering is an effective technology to improve the quality of perovskite films, yielding larger crystal grain sizes and lower defect densities.13-15 For example, metal halide salts provide selected ions, e.g., Cs+ 、 Rb+ and Co2+ with comparative radius to partially substitute MA+ or Pb2+, thus modifying the GoldSchmidt tolerance factor to create a more perfect perovskite crystal structure with enhanced morphology and charge transport properties.16-19 Inorganic acids, typically HI and HCl, control the formation of Pb-I-Pb bonds to modulate the crystallization kinetics.20, 21 These reports demonstrate the reduction of defect states through regulation of the crystal growth process. On the other hand, defects can be directly passivated through various bonding interactions by incorporating additional molecules/polymers



into the perovskite film.13 Recently, the molecules/polymer PCBM, ITIC and PBDB-T have proven effective at eliminating Pb-I anti-site defects at grain boundaries, attributed to hydrogen bonding and Lewis base group interactions.22-25 Compared to small molecules, it should be stressed that long-chain polymers with multiple functional groups can offer more stable and reliable interactions and conceivably serve as perovskite grain boundary cross-linkers to improve film stability, as have been proposed for polyethylene glycol (PEG) by Zhao et al.26 and Lu et al.,27 as well as for polyacrylic acid (PAA), poly(4-vinylpyridine) (PVP) and polyethyleneimine (b-PEI) investigated by Yang and co-workers.28

Ethyl Cellulose (EC) is a derivative of celluloses which are non-toxic, natural, earth-abundant and cheap polymers,29 its chemical structure is depicted in Figure 1a. EC has a larger number of hydrogen bonds and it is thermally stable and water insoluble, ideal properties for a perovskite-passivating additive. In this work, we introduce EC to the perovskite precursor solution to form high quality polycrystalline films, and multiple beneficial effects of the EC additive, including trap passivation, hydrogen bonding interaction, balanced charge mobility, scaffold behavior, etc., are demonstrated. These findings further improve both in PCE and stability of the perovskite solar cells. The EC passivated MAPbI3 device in the planar n-i-p architecture reaches an average of PCE of 19.41%. The open-circuit voltage (Voc), the short-circuit current density (Jsc) and the fill factor (FF) are 1.1 V, 22.89 mA cm-2, 0.77, respectively. In addition, the hysteresis index (HI) significantly decreases from 0.081 of the control device without


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EC to 0.027, attributed to suppressed ion migration and surface charge trapping. The defect passivation by EC significantly improves the environmental stability of perovskite films, enabling the devices to retain 80% of the initial efficiencies in ambient air at 45% relative humidity after 30 days, whereas the pristine devices fully degrade at the same time. Our results suggest that using the green and cheap cellulose as a filmforming additive is a promising method to improve the efficiency and stability of perovskite solar cells, especially for their large-scale fabrication in the future.

Results and Discussion We prepared the perovskite films following the one-step spin-coating process with a mixture of PbI2, MAI and EC as precursor composition, as shown in Figure 1a. EC is a long-chain polymer formed by a six membered heterocyclic unit containing hydroxyl and ethyl ether groups, and the electron lone pairs of the O atoms are forecasted to be able to passivate the under-coordinated Pb2+ vacancies.30 To confirm the existence of EC in the perovskite film, the survey and the core level spectra of bare EC films, perovskite films with and without EC are detected via X-ray photoelectron spectroscopy (XPS), respectively, see Figure 1b, 1c and Figure S1. The dominant feature of O1s peak from the bare EC film is the main peak at the binding energy of 534 eV, whereas the weak feature at 531.5 eV is from the underlying Indium Tin Oxide substrate in bare EC film. The pristine perovskite film shows no oxygen signal. The appearance of O1s core level peak at 534 eV in the perovskite film with 1mg mL-1 EC additive originates from the EC additive. There is no visible change in the Pb 4f, N 1s



and I 3d spectra after EC additive. Furthermore, the bonding interaction between perovskite and EC is evaluated by Fourier transform infrared (FTIR) spectra, see Figure 1d. In the bare EC film, the C-O stretching vibration shows a broad peak at around 1099 cm-1 contributed by oxygen in the ether, hydroxyl and heterocyclic units. Perovskite with EC exhibits the characteristic of a C-O bond located at 1035 cm-1. Such a shift indicates that oxygen in EC interacts with some component in perovskite, possibly the Pb2+ in the I-loss region of Pb-I framework through Lewis acid-base reaction.31

The crystallization and high quality of perovskite film is the key issue for the performance in device.32, 33 The color evolution of the perovskite film under 100 ℃ annealing was studied by optical microscopy with the time between 3s and 9s, see Figure 1e. Compared to the pristine films, the introduction of EC clearly slows down the growth of the polycrystalline perovskite films to finish the whole black phase transformation. The prolonged time suggests the high quality perovskite films are formed.34, 35 The top-view scanning electron microscopy (SEM) was further to study the film morphology. In Figure 1f-1i, SEM images illustrated highlight the morphological differences induced by increasing EC concentrations from 0, 0.01, 0.1 and 1 mg L-1. All surfaces of the films are dense and smooth without cracks and pinholes. The average crystal grain size clearly increases after EC modification, to 300 nm at 0.1 and 1 mg mL-1 in contrast to the 200 nm of pristine film. The largest grains exhibit a size over 400 nm, see the statistical plots in the insets. Larger grain sizes with reduced grain boundaries yield less defects.36


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The crystalline properties of the corresponding perovskite films are further explored by X-ray diffraction (XRD). In Figure 2a the XRD patterns show clean and sharp diffraction peaks for the pristine and different EC added perovskite films. There is no additional peak in the EC added perovskite films, indicating that the amorphous nature of EC additive does not change the cubic perovskite phase crystal structure. The intensities of the main peaks at 14.1°, 28.4° and 31.8°, corresponding respectively to the (110), (220) and (310) crystal planes, all increase, further confirming an improved perovskite crystallization upon EC addition. Noteworthy, the (110) diffraction peak shifts to lower angle in the amplified picture of Figure 2b, suggesting that the perovskite lattice space becomes larger. Perovskite films generally are treated with ～ 100 ℃ thermal annealing to enhance crystallization after spin-coating. During this process, the crystal lattice first expands with temperature and then shrinks upon cooling down to room temperature, see the sketch of Figure 2c. Because the huge mismatch of the thermal expansion coefficients between perovskite MAPbI3 (αv = 1.57 x 10-4 K-1) and substrate TiO2 (αv = 17.3. x 10-6 K-1), the unsynchronized expansion/shrinkage process causes tensile strain along the in-plane direction and compressive strain along the outof-plane direction. Such strains can accelerate the degradation of the perovskite films.37 However, the presence of long-chain EC serves as a scaffold for the perovskite lattice as illustrated in Figure 2c, which is speculated to reduce expansion/shrinkage during the annealing process and thus decreases the strain.

The optical properties also are enhanced in the EC passivated perovskites films.



In Figure 2d, ultraviolet-visible absorption spectra display a pronounced increased absorption over the wavelength range of 400 to 550 nm after EC modification, but a slight drop over 550 to 750 nm. The former originates from the enlarged grain size of the perovskite crystals, and the latter possibly results from the bonding interaction between perovskite and EC. Figure 2e displays the steady-state photoluminescence (PL) quenching of the perovskite films with and without EC. The change in PL provides strong evidence that EC helps enhance charge extraction and suppress charge recombination by reducing defect state density. Moreover, the PL peaks show a blue shift

from 777 to 772 nm, which can be attributed to EC-induced passivation of traps

on the film top surface through filling of deep traps.38 Time-resolved PL is employed to further study the charge carrier dynamics. In Figure 2f, perovskite films containing the EC additive remarkably boost a fast decay component (τ1) from 1.63 (pristine) to 5.10, 5.30 and 4.12 ns with EC concentrations at 0.01, 0.1 and 1 mg L-1, respectively, demonstrating highly effective trap passivation on the surface. The slow decay component (τ2) component also shows an increase from 123.32 (pristine) to 125.47, 138.00 and 132.2 ns, respectively. The prolonged carrier lifetimes confirm the high quality of the EC perovskite films with lower trap densities, which should result in improved charge transport properties and increased PCE of the corresponding solar cells. From the results in Figure 2, the best crystalline perovskite film is obtained when the EC concentration is 0.1 mg mL-1, in agreement with the largest average grain size observed in Figure1h.


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Having demonstrated remarkable improvements of EC perovskite films in terms of morphology, structure and dynamics, we now test the films in perovskite solar cells. The perovskite devices of FTO/c-TiO2/MAPbI3/Spiro-OMeTAD/MoO3/Ag are fabricated. The cross-sectional SEM images of the devices are illustrated in Figure S2. Figure 3a depicts the J-V curves of the devices with pristine and EC-passivated perovskite films at different EC concentrations. The pristine devices feature an average PCE of 17.11% with a Voc of 1.05 V, a Jsc of 22.77 mA cm-2, and a FF of 71.23%. When introducing the EC additive into the perovskite films, the PCE dramatically increases, as expected. Particularly, the devices with 0.1 mg L-1 EC exhibits the highest average PCE of 19.41% with a Voc of 1.10 V, a Jsc of 22.89 mA cm-2, and a FF of 77.1%, again in agreement with the increased grain sizes and reduced trap densities in these films. The performance parameters of the devices are summarized in table 1. Over 20 devices are fabricated to show excellent reproducibility (Figure S3). A clear enhancement in the external quantum efficiency (EQE) spectrum for the EC passivated devices are also observed (Figures S4) in line with the increase in Jsc in Figure 3a. The photocurrent density calculated with EQE is close to the measured value of Jsc. Meanwhile, the J-V hysteresis of the optimized device with EC modification is dramatically suppressed as shown in Figure 3b. The pristine devices present a large PCE difference between 15.7l% (forward scan) and 17.11% (reverse scan), while in contrast, the EC passivated devices exhibit a small hysteresis with PCE of 18.87% (forward scan) and 19.41% (reverse scan), i.e., the hysteresis index decreases by three times from 0.081 to 0.027. We attribute the effective suppression of device hysteresis to suppressed ion migration and



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surface charge trapping, induced by the interaction between EC and perovskite.39 We additionally study the time-dependent power output of the optimized device in Figure 3c. The PCE and the steady-state photocurrent of the EC passivated device are stabilized at 18.81% and 21.64 mA cm-2, respectively, ruling out deterioration cell stability due to EC modification.

To gain deep mechanistic understanding on the improved device performance, the dark J-V characteristics for both hole-only and electron-only devices are collected to assess the defect density and charge mobility, as shown in Figure 3d and 3e, respectively. The trap density is determined according to the following equation 𝑁𝑡𝑟𝑎𝑝𝑠 =

40:

2𝜀0𝜀𝑟𝑉𝑇𝐹𝐿 𝑒𝐿2

, where ε0 represents the vacuum permittivity, εr is the relative

permittivity of MAPbI3 (~ 60)41, L is the thickness of MAPbI3 layer (~ 500 nm), and VTFL is the onset voltage of trap-filling limited (TFL) current regime. With EC passivation, the hole trap density in the perovskite layer decreases from 0.57 × 1016 to 0.30 × 1016 cm-3. Synchronously, the electron trap density diminishes from 0.28 × 1016 to 0.14 × 1016 cm-3. The reduced charge trap density (50%) proves effective passivation at perovskite grain boundaries, which is expected to suppress charge recombination and thus result in the increased device efficiency. The mobility (µ) of the perovskite in the 9

a space-charge limited current (SCLC) regime are calculated by the equation 42: 𝐽 = 8 𝑉2

𝜀𝑟𝜀0𝜇 𝐿3 , where J is the current density, and V is the base voltage. The hole mobilities in the pristine and EC passivated perovskite are estimated to be 1.31 × 10-2 and 4.11 × 10-2 cm2 V−1 S−1, respectively. The electron mobilities are estimated to be 3.07 × 10-2 and


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5.26× 10-2 cm2 V−1 S−1, respectively. The increased hole and electron mobility become more balanced, accounting for the enhanced FF in the device. Figure 3f compares the Nyquist phots of the impedance spectroscopy in dark condition. The inset gives the equivalent circuit where RS is the sheet resistance of the conductive electrode and RCT is the charge transfer resistance.43 Because the devices with and without EC possess the same structure, the difference of RS can be negligible. It is found that the major difference is RCT, decreasing from 996 to 528 Ω after EC modification. The lower RCT for the passivated device indicates improved charge carrier transport in the bulk film and likely also improved charge injection at the contacts between the perovskite and the adjacent layers, which we attribute to EC-suppressed trap densities in the bulk perovskite films and on their film surface.19

The recombination mechanisms are further explored by the correlation of Voc and Jsc dependent on light intensity (Plight). As shown in Figure S5a, the linear fitting of Voc versus Plight gives an ideality factor n of 1.38 and 1.71 for devices with and without EC, respectively. If n is close to 2, Shockley-Reed-Hall (SRH) recombination induced by trap states is the dominant process in the active layer, while n is close to 1 where bimolecular recombination is the primary process.44 Clearly, the reduced ideal factor confirms the suppressed SRH recombination due to effective defect passivation, in agreement with the observed PL results. The Jsc dependence on Plight follows 𝐽𝑆𝐶 ∝ 𝑃𝛼𝑙𝑖𝑔ℎ𝑡, where α is close to 1, it means that bimolecular recombination is negligible. The fitting values of α are estimated to be 0.974 and 0.962 for the devices with and without



EC, respectively, see Figure S5b. The increased α suggests that bimolecular recombination is reduced after EC modification due to higher and more balanced charge mobilities, benefiting charge transport and collection.

Figure 4 shows conductive atomic force microscopy (c-AFM) images of the perovskite films with and without EC under bias voltage. The overall electronic conductivity of the pristine film is slightly larger than that of the EC passivated perovskite due to the insulating EC. Many dark spots were observed on the surface of the pristine film, suggesting an uneven current distribution and a current loss in those regions. In contrast, the EC passivated film shows more homogeneous current distribution with less charge recombination. Comparing with the corresponding topographical images (Figure S6), the roughness of the perovskite film decreases from 11.32 to 5.86 nm after EC modification, which further proves that the scaffold function of EC during the growth of perovskite crystal results in a smooth surface enabling enhanced surface current and top-down charge transport, i.e., charge injection at contacts with reduced charge recombination as discussed in the previous session.

Environmental stability is another key metric for perovskite solar cells in terms of realistic application. We first discuss the effect of EC on the chemical stability of the perovskite film. Liquid-state 1H nuclear magnetic resonance (NMR) spectra of the pure EC, MAPbI3 solution with and without EC are collected in Figure 5a. The peak at δ = 7.88 ppm that belongs to the proton resonance signal of –NH3+ in the MA cation moves


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towards the lower value of 7.86 ppm after incorporation of EC. The observed downfield chemical shift can be attributed to hydrogen bonding between O of -OH in EC and H of -NH3+ in MA via van der Waals interaction. It is also expected that I- in MAI forms hydrogen bond with H of -OH in EC where the proton signal shifts downfield from δ = 1.21 to 1.12 ppm. Such interactions between EC and perovskite could suppress the degradation of perovskite and thus stabilize its crystal structure,45 see the schematic diagram in Figure 5b. After EC passivation, the film degradation is retarded, which is affirmed by the contact angle measurement shown in Figure 5c. The water contact angle of the pristine perovskite film is 33.0°. For EC passivated perovskite films at different concentrations of 0.01, 0.1 and 1 mg mL-1, the contact angles significantly increase to 42.6, 67.8 and 83.5°, respectively, suppressing the moisture attacking to the perovskite films. The improved hydrophobicity is supported by the stability of the perovskite films that was exposed to 95% relative humidity in dark for 4 h, where the pristine film undergoes strong degradation with the color fading from black to yellow, while the passivated film is nearly unaffected, see Figure 5d. Finally, the environmental stability of the bare solar cells was evaluated under ambient air at 45% relative humidity in dark. Figure 5e shows the normalized PCE decay against storage time. The pristine device completely degrades after 30 days. The passivated device has the better stability in which 80% of the initial PCE is maintained for the same time period. We attribute the enhanced stability to decrease in lattice strain, lower trap densities and a more hydrophobic surface and grain boundaries after EC passivation.



Conclusion In conclusion, we have demonstrated a strategy for passivating defects in perovskite films by introducing an extremely low-cost and green Ethyl Cellulose. The synergetic effects of EC on Lewis acid defect passivation and crystal growth promotion result in high quality perovskite films with enlarged crystal grain sizes. In addition, the EC additive increases charge mobility in the perovskite layers and makes it more balanced. The average PCE of the optimized devices increases from 17.11% to 19.27% accompanied by an overall increase in Jsc, Voc and FF. The hysteresis index is found to decrease significantly from 0.081 to as low as 0.027, attributed to suppressed ion migration and surface charge trapping. Meanwhile, the long-chain EC scaffold-assisted perovskite crystal releases the lattice strain and stabilize the perovskite phase upon cooling during the post-annealing process. Hydrogen bonding between EC and perovskite via van der Waals interaction improves the stability of the perovskites. A pronounced retardation on degradation is realized where the resulting devices without encapsulation retain 80% of their initial efficiencies after 30 days in ambient air under 45% relative humidity while the control devices are fully degraded. This work shows the versatile effects of EC additive on the perovskites. The finding provides a promising way using low cost and green biomaterials to improve both efficiency and stability of the perovskite photovoltaics.

Experimental section Materials


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Lead iodide (PbI2) and n-butyl alcohol were obtained from TCI. Methylammonium iodide (MAI), lithium bis(trifluoromethanesulfonyl)imide (Li-TFSI) and 4-tertbutylpyridine (tbp) were received from Xi’an Polymer Light Technology Corp, and Ethyl cellulose (EC) powder from Aladdin. N, N-dimethylformamide (DMF), chlorobenzene (CB), dimethylsulphoxide (DMSO) and Titanium diisopropoxide bis(acetylacetonate) solution were purchased from Sigma Aldrich. The perovskite precursor solution was prepared by mixing MAI and PbI2 (1:1) with concentration of 1 M in the mixed solvent of DMF and DMSO (4:1 v/v). EC with different concentrations were added in the precursor solution. It was spin coated at 1000 rpm for 10 s, and 4000 rpm for 30 s to form perovskite films. During the second step, chlorobenzene was poured as the usual anti-solvent method. Finally, the films were annealed at 100 ℃ for 10 min. The processes of film fabrication and annealing are done in N2-protected glove box.

Film Characterization XPS spectra were collected with an ultrahigh vacuum (UHV) surface analysis system with a monochromatic Al Kα (1,486.6 eV) as the excitation source. The FTIR measurements were carried out with a Bruker Fourier Transform Infrared Spectrometer in transmittance mode. SEM images were captured by Merlin field emission-SEM instrument. C-AFM images were collected using Veeco MultiMode V with a bias voltage of 2 V to acquire the surface current distribution. The quality and crystal structure of perovskite films were obtained by XRD (PANalytical). Steady-state PL



analysis were carried out with Horiba Jobin-Yvon FL-3. 470 nm light source was used to excite perovskite films. The absorbance spectra were performed using PerkinElmer Lambda 750. TRPL measurements were performed via a Time-Correlated Single Photon Counting system.

1

H Nuclear Magnetic Resonance (NMR)

The NMR spectra were obtained in solution of N, N-dimethylformamide d7 by highresolution liquid nuclear magnetic resonance spectrometer (Bruker Avance 500 MHz).

Device Fabrication and Characterization Pre-cleaned FTO glasses were exposed to ultra violet-ozone for 15 min before using. For compact TiO2 layer, a dilute titanium diisopropoxide (1:16 in n-butyl alcohol) was spin coated on FTO at 2000 rpm for 30 s, then the films were placed in a 500 ℃ for 30 min. After that, they were transferred to N2-filled glove box immediately. The perovskite layers were spin coated on FTO/TiO2. For the hole transport layer, 72.3 mg of spiro-OMeTAD, 17.5 mL of Li-TFSI (520 mg mL-1 in acetonitrile) and 28.8 mL of tbp were dissolved in 1 mL of CB solution. Then the mixture solution was spin coated on perovskite film at 4000 rpm for 30 s. Finally, MoO3 (7 nm) and Ag (100 nm) were thermal evaporated as electrodes. The current density-voltage (J-V) curves of PSCs were acquired via a Keithley 2612 under Air Mass 1.5G light source. EQE spectra were collected using a modified spectrometer with a calibrated silicon photodetector. Electrochemical impedance spectroscopy was performed using an electrochemical


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station based on Keithley 2400 Source Meters.

Acknowledgements This work was financially supported by the National Science Foundation of China grant (No. 11604099, No. 21875067, No. 51873138, No. 51811530011), Shanghai Science and Technology Innovation Action Plan (No. 19QA1403100, No. 17JC1402500), and National Key Project for Basic Research of China (2017YFA0303403). We also acknowledges the support from the Swedish Research Council (project grant No. 201605498) and the STINT grant (CH2017-7163).

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Figure 1. (a) Schematic illustration of the perovskite film preparation and EC chemical structure. XPS spectra of (b) survey and (c) O1s core level of the bare EC film, perovskite films with and without EC at concentration of 1 mg mL-1. (d) FTIR spectra. (e) Evolution of optical photographs of perovskite films with and without EC under 100 ℃ annealing. (f-i) Top view SEM images of perovskite films with different EC concentrations, where the insets represent statistical crystal grain sizes.


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Figure 2. (a) XRD patterns of perovskite films with different EC concentration. (b) Higher magnification XRD patterns around the (110) crystal plane. (c) Schematic illustration of the EC scaffold acting on strain prevention. (d) UV-Vis absorption spectra, (e) Steady-state PL spectra. (f) Time-resolved PL spectra



Figure 3. (a) J-V curves of the perovskite solar cells with and without EC. (b) J-V curves of the optimized devices with and without EC at the concentration of 0.1 mg mL-1 by forward and reverse scan. (c) Maximal steady-state photocurrent and power output of the optimized perovskite device with EC. Dark J-V curves of (d) hole-only and (e) electron-only device. (f) Nyquist plots of impedance spectroscopy of the perovskite device with and without EC with the inserted equivalent circuit.


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Figure 4. C-AFM images of the perovskite films: (a) pristine film without EC, (b) film with EC at the concentration of 0.1 mg mL-1.



Figure 5. (a) Liquid-state 1H NMR spectra of EC and perovskite with and without EC at the concentration of 0.1 mg ML-1. (b) Schematic diagram of hydrogen bonding formation between EC and perovskite. (c) Water contact angle on perovskite films at different concentrations of EC. (d) The change of optical pictures of perovskite films with and without EC aging to 95% relative humidity. (e) Stability comparison of the bare perovskite solar cells with and without EC in ambient air with 45% relative humidity.


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Table 1. Performance parameters of perovskite solar cells employing the pristine and the EC modified perovskite layers with different concentrations. RS is reverse scan, and FS is forward scan.

0 mg

Jsc [mA cm-2]

Voc [V]

FF [%]

PCE [%]

RS

22.77

1.05

71.23

17.11

FS

22.59

1.05

66.23

15.71

RS

22.87

1.07

74.98

18.35

RS

22.89

1.10

77.10

19.41

FS

23.00

1.09

75.26

18.87

RS

22.44

1.06

73.80

17.55

mL-1

0.01 mg mL-1

0.1 mg

0.081

mL-1

1 mg mL-1

HI

-

0.027


-


TOC Figure


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