Achieving high open-circuit voltage on planar perovskite solar cells via

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Achieving high open-circuit voltage on planar perovskite solar cells via chlorine doped tin oxide electron transport layers Jiwei Liang, Zhiliang Chen, Guang Yang, Haibing Wang, Feihong Ye, Chen Tao, and Guojia Fang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b03873 • Publication Date (Web): 11 Jun 2019 Downloaded from http://pubs.acs.org on June 11, 2019

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ACS Applied Materials & Interfaces

Achieving high open-circuit voltage on planar perovskite solar cells via chlorine doped tin oxide electron transport layers Jiwei Liang,†,‡ Zhiliang Chen,† Guang Yang,† Haibing Wang,† Feihong Ye,† Chen Tao*,† and Guojia Fang*,†,‡ School of Physics and Technology, Wuhan University, Wuhan 430072, People's Republic of China. †

‡Shenzhen

Institute, Wuhan University, Shenzhen 518055, People's Republic of

China.

ABSTRACT: The open-circuit voltage deficit is one of the main limiting factors for the further performance improvement in planar structured perovskite solar cells. In this work, we elaborately develop chlorine binding on the surface of tin oxide electron transport layer for high open-circuit voltage device (1.195 V). The chlorine passivation on SnO2 not only effectively mitigates the interfacial charge recombination between SnO2 and perovskite, but also enhances the binding of chlorine with lead at the SnO2/perovskite interface. The chlorine passivated SnO2 electron transport layer exhibits a better energy alignment with perovskite layer and an improved electron mobility, which will promote efficient electron transfer at the interface. In addition, the elevated Fermi level of SnO2 electron transport layer increases carrier extraction and suppresses interfacial recombination, which is responsible for the open circuit voltage enhancement. Planar perovskite solar cells with chlorine passivated SnO2 exhibit a higher open-circuit voltage of 1.195 V than that of reference ones (1.135 V) for a lower bandgap of 1.58 eV perovskite absorber, which achieve a PCE of 20%

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with negligible hysteresis.

KEYWORDS: chloride doping, tin oxide, interfacial passivation, open circuit voltage, perovskite solar cells

1. INTRODUCTION: Organic-inorganic lead halide perovskites have been widely studying as a promising photovoltaic absorber thanks to their strong absorption coefficient, tunable band gap, high carrier mobility and long carrier lifetime. 1-4 Tremendous efforts, such as compositional engineering 5, interface engineering 6, defect passivation 7, have been making to push the record power conversion efficiencies (PCEs) of perovskite solar cells (PSCs) over 23%, approaching crystalline silicon solar cells. Among these high-efficiency cells, one typical structure is the so-called n-i-p structure, in which the electron transport layer (ETL) inserts between perovskite and the front transparent electrode. Stemmed from dye-sensitized solar cells, titanium oxide (TiO2) frequently serves as ETL in PSCs with well-matched energy levels with perovskites. 8 However, the necessity of high-temperature (500 ℃ ) sintering to form TiO2 crystalline layers precludes its application on temperature sensitive substrates. Moreover, its ultraviolet light sensitivity has been demonstrated to degrade the perovskite absorber and hence to reduce the lifetime of PSCs. Particularly the abnormal hysteresis in planar PSCs with TiO2 prevents the precise evaluation of PSC performance. 9-11


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Tin oxide has emerged as the state-of-the-art alternative ETL in planar n-i-p PSCs due to its higher carrier mobility, optical transmission, and wider bandgap. 12-15 Planar structured PSCs with SnO2 have been reported to yield considerable efficiencies with more than 22%. The higher carrier mobility and therefore better balanced extraction of carriers than TiO2 are well accepted to cause less hysteresis in planar PSCs.

16.17

There are various methods to grow SnO2 ETL, for example, spin

casting, chemical bath deposition, and atomic layer deposition. 18-30 However, it is not widespread to obtain hysteresis-free PSCs from all these methods. Bu et al. uncovered that the addition of potassium in perovskite precursor efficiently eliminated the hysteresis in planar PSCs. 31 It has been proved that passivated SnO2 ETLs could exhibit better performance compared to the pristine ETLs, such as higher power conversion efficiency (PCE), preferable long-term stability and extensive environmental suitability.

32

Yang et al.

discovered that Ethylene Diamine Tetraacetic Acid (EDTA) complexed SnO2 ETLs not only eliminated the notoriety hysteresis but also improved the stability of the devices. 33 Recently, Liu et al. found that inorganic alkaline halide could simultaneously passivate different kinds ions defect at the perovskite/SnO2 interfaces, in which the passivated device shown negligible hysteresis and enhanced long-term stability.

34

The PSCs equipped by bilayer electron transport shown obviously advantage, which reduce the energy barrier and achieve a higher Voc. 35,36 Our group has achieved a high PCE using SnO2 quantum dots (QDs) as ETLs. 37



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However, the n-i-p structure PSCs still exist J-V hysteresis and relatively lower Voc, which is partially ascribed to defect and carrier recombination at the SnO2/Perovskite interfaces. Interface engineering, such as charge transportation, defect passivation,

41, 42

band alignment,

43

38-40

is a very significant strategy to decrease

the defect and suppress the recombination. Our group has reported that the chlorine passivation on the interface of perovskite/FTO can enhance the device performance.44 Other groups found that the additive of NH4Cl was beneficial for high-quality perovskite crystals, leading to better device performance. 45, 46 Tan et al. proposed that the Cl-capped TiO2 nanocrystal film could effectively mitigates interfacial recombination and improves the stability of device.47 In this work, we proposed an elaborate and simple strategy to retard interface defects and carrier recombination. Here, we introduced NH4Cl as solution dopant, and we

devised

chlorine

passivated

SnO2

QDs

via

an

environment-friendly,

room-temperature synthetic method and employed it as ETL for planar PSCs. The devices using chlorine passivated SnO2 exhibits a stunning Voc of 1.195 V, which is among the highest reported Voc for the 1.58 eV bandgap PSCs in the n-i-p structure. The best-performing device shows a PCE of 20%. We demonstrated that chlorine passivated SnO2 can effectively reduce the interface defects, which contribute to reduce the probability of interface recombination. Therefore, the elevated Fermi level of SnO2 ETL, reduced interface defects and suppressed interfacial recombination between SnO2 and the perovskite absorber layer are all responsible for the increased open circuit voltage. Furthermore, the ammonium cations (NH4+) balance the charge


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of SnO2 QDs surface, which is beneficial for the stability of QDs solution and the formation of high-quality ETL film.

2. EXPERIMENTAL SECTION: 2.1 Materials: In this work, we devised chlorine passivated SnO2 QDs via environment-friendly, room-temperature synthesis and employed it as ETL for efficient planar PSCs. SnO2 QDs colloidal solution was prepared by using SnCl2.2H2O (purity99.99% Sigma-Aldrich) as precursor material, and deionized water as solvent. First, 1 mmol (225.65 mg) SnCl2.2H2O was added to 10 mL ultrapure water with continuous stirring to form pristine QD solution. The passivated QD solutions prepared through mixing 5.35 mg, 10.7 mg, 16.05 mg and 21.4 mg NH4Cl (purity99% Aladdin) with 225.65 mg SnCl2.2H2O in 10 mL ultrapure water to form 10%, 20%, 30%, and 40% doped QD solutions. The suspension was keep stirring for 48 h to form SnO2 colloidal precursor solution. All the above processes are performed at room temperature without any assisted procedure. A mixture of 1.3 mmol formamidinium iodide (FAI purity99.99%

Sigma-Aldrich),

lead

iodide

(PbI2

purity99.9985%

TCI),

methylammonium bromine (MABr purity99.9% Sigma-Aldrich), lead bromine (PbBr2

purity99.9%

Sigma-Aldrich),

cesium

iodide

(CsI

purity99.99%

Sigma-Aldrich) and potassium iodide (KI purity99.99% Sigma-Aldrich) were dissolved in 800 L N, N-Dimethylformamide (DMF purity99.8% Alfa) and 200 L Dimethyl

sulfoxide

(DMSO

purity99.8%

Alfa)

solutions

to

form



K0.035Cs0.05(FA0.85MA0.15)0.95Pb(I0.85Br0.15)3 perovskite. The Spiro solution was prepared by addition of 72.3 mg Spiro-OMeTAD to 1 mL chlorobenzene (99.9%, Sigma-Aldrich) with additional 17.5 L lithium bis (trifluoromethanesulfonyl) imide (Li-TFSI, 99%, Xi'an Polymer Light Technology Corp.) solution (520 mg in 1 mL acetonitrile) and 28.8 L 4-tert-butylpyridine (tBP, 99.9%, Sigma-Aldrich)

2.2 Fabrication of perovskite solar cells: The fluorine-doped tin oxide (FTO) glasses were ultra-sonicated in deionized water, acetone, and alcohol for 15 min, respectively. The samples were then blown dry with nitrogen and treated with 15-min UV plasma. The SnO2 ETLs were prepared by spin coating the precursors at 3000 rpm for 30 s and calcined at different temperature from 100-250 ℃ for 1 hour. After cooling down, the samples were transferred into an Ar-filled glove box. The perovskite precursor solution was then deposited on ETL at the first step is 2000 rpm for 10 s and the second step is 6000 rpm for 30 s. During the second step, a 150 L chlorobenzene anti-solution was dropped within the last 15 s. Then the perovskite films were annealed at 100 ℃ for 1 h followed by spinning Spiro-OMeTAD solution at 3000 rpm for 30 s. Finally, 50 nm Au was thermally evaporated by thermal evaporation under 2.010-4 Pa vacuum pressure to complete the cell fabrication.

2.3 Characterization: The J–V characteristics, space charge limit current (SCLC) and EIS of the PSCs were measured using a CHI 660D electrochemical work station with a 450 W xenon lamp and an AM 1.5G filter (Shanghai Chenhua Instruments, China). The


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external quantum efficiency (EQE) spectrum was recorded with a QE/IPCE system (Enli Technology Co. Ltd). The time resolved photoluminescence (TRPL) were obtained with Delta Flex fluorescence spectrum spectroscopy (HORIBA). The decay curves were simulated according to the formula τavg = (A1τ12 + A2τ 2 2 ) / (A1τ1 + A 2 τ 2 ) (1). The absorption and transmission spectra were measured using a SHIMADZU mini

1280

UV-visible

spectrophotometer.

X-ray

photoelectron

spectrum

measurements were performed with an XPS/UPS system (Thermo Scientific, Escalab 250Xi). The surface roughness of thin film was determined with atomic force microscopy (AFM) (SPM-9500J3, Shimadzu, Japan).

3. RESULTS AND DISCUSSION: The morphologies of SnO2 QDs collected from the solution and dried at room temperature for 6 h are presented in Figure 1a and b through transmission electron microscopy (TEM) analysis, which represent the sizes of pristine SnO2 and the NH4Cl passivated SnO2 QDs reaching 3-4 nm, respectively. Figure S1 shows the digital optical images of SnO2 quantum dot solutions irradiated with a red laser. The visible Tyndall effect illustrates that the obtained solutions are homogeneous quantum dots. It is well known that the crystal quality of SnO2 QDs significantly determines the carrier concentration, charge conductivity and Hall mobility, thereby affecting charge transport, transfer and collection. The HR-TEM images show that the interplanar spacing of (110) plane is 0.33 nm in both samples, indicating that the chlorine passivation hardly influences the crystalline phase of rutile SnO2 QDs. Soon



afterwards, the SnO2 QD solutions without/with NH4Cl are used for preparation of the electron transport films by spin coating method as shown in Figure 1d and e, where the passivated SnO2 film shows smoother and more uniform surface compared with that of pristine SnO2 (the bare FTO as shown in Figure 1c). Meanwhile, the atomic force microscope (AFM) measurement results confirm that the addition of NH4Cl can distinctly decrease the root mean square (RMS) of surface roughness from 7.50 to 6.81 nm as presented in Figure S2. At the same time, the N and Cl elements are uniformly distributed in the SnO2 film as SEM-EDS mapping depicted in Figure S3, where we speculate that the following chemical reaction mechanism could happen at the surface of SnO2 QDs in the annealing process.

 Sn - OH + NH 4 +   Sn + NH 3  +H 2 O (2)  Sn + Cl -   SnCl (3) The surface of SnO2 QDs contain a certain amount of –OH, which are sensitive to NH4+ in the solution. When annealing the film containing NH4+, the surface –OH will split away off easily to release NH3 and form exposed Sn single bond as presented in equation (2). Meanwhile, the chlorine ion existed in solution will rapidly react with Sn as presented in equation (3). Therefore, the chlorine cross bindings on the surface of SnO2 will be formed. Apart from the degree of crystallinity and the surface morphology, the surface elements and chemical states are analyzed through the X-ray photoelectron spectroscopy (XPS) for the SnO2 ETLs with or without the chlorine passivation. The binding energy of pristine and passivated SnO2 are surveyed and the full XPS spectra


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are shown in Figure 2a, suggesting that the introduction of chlorine didn’t change the SnO2 QDs primary chemical component. At the same time, the high resolution Sn 3d, O 1s and Cl 2p XPS spectra are shown in Figure 2b, 2c and 2d. From Figure 2b, the pristine peaks of Sn 3d5/2 and Sn 3d3/2 appear at 486.98 eV and 495.43 eV, but the peak position shift by 0.2 eV with the chlorine passivation. Then, the components of O 1s have been analyzed with XPS Gaussian distribution, which contributed to reveal the valence state, vacancy and compositions. According to Figure 2c, two oxygen states are shown in SnO2 film, where the 530.88 eV and 532.23 eV represent the oxygen tin binding energy (OSn) 48, 49 and the hydroxyl oxygen binding energy (OH) 6 in pristine SnO2 film. Nevertheless, the binding energy of OSn and OH increased to 531.08 eV and 532.48 eV for chlorine passivated SnO2 film, owing to the stronger binding energy of OSn compared with OH. It is shown that the OH relative peak intensity decrease obviously, while the OSn peak is expected to enhance with the chlorine passivation. Considering that the reaction discussed above, we conclude that the O-Sn bonding energy increase and the O-H peak decrease compare with the result of Sn 3d in Figure 2b. In order to reveal the relative content of oxygen, we have estimated the two type peak areas as presented in Table 1, where the single OH content (OH /OH+OSn) decreases from 29.5% to 22.5% with the chlorine passivation. To the contrary, the OSn content increases significantly from 70.5% to 77.5% on account of the chlorine binding function. The weakening surface defect is beneficial for the enhancement of carrier mobility, the decrease of interface recombination and the acceleration of charge transport. Figure 2d shows the XPS of Cl 2p, we could see



that the signal of chlorine is clearly enhanced with the addition of NH4Cl, proving the more chlorine is bound on surface of SnO2 QDs. The formation of SnCl2-termination could promote the chlorinated interface binding with lead in the perovskite layer, which can effectively reduce the interface defect at ETLs/perovskite. 47 Perovskite films were obtained through spin coating of perovskite precursor solution on pristine and passivated SnO2 ETL substrates. The corresponding top surface and the cross-sectional SEM images with or without chlorine passivation are shown in Figure 3a and 3b. Figure 3b shows that the morphology of perovskite films deposited on passivated SnO2 ETL, where the continuous pinhole-free and the denser grains of perovskite film are obtained compared with pristine SnO2 (Figure 3a). Meanwhile, the cross-sectional SEM image also proves this character as shown in Figure 3c and 3d. Through incorporating chlorine into the ETL film, we obtained higher quality perovskite film and better interface between SnO2 ETL and perovskite absorber. Figure 4a shows the J-V characteristics measured under AM1.5G simulated sunlight illumination with both forward and reverse scans, in which a champion PCE reaches 20% (forward scan direction) compared to 19.4% (reverse scan direction) with negligible hysteresis. The corresponding of PSCs parameters for all presented J-V curves are shown in Table 2, where we can clear see that the Voc, FF and PCE increase obviously with the chlorine passivation. We optimize the performance of the device using the different ETL annealing temperatures and the results are shown in Figure S4, where the Voc increases initially and decreases afterwards with the


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temperature increases and the maximum PCE obtained at 200℃. The performance of PSCs with different NH4Cl concentration passivated ETL are shown in Figure S5, in which we can clearly see that the optimal NH4Cl concentration is 20%. We consider that the moderate chlorine passivation could promote the performance of PSCs, while the excessive Cl would influence the electronic property as shown in Table S1. Meanwhile, we find that interface chlorine passivation increase the Voc from 1.135 to 1.195 V and the fill factor (FF) from 72% to 75%, while the current density increases inconspicuously compared to pristine SnO2. Figure 4b shows the EQE curve of the devices without or with chlorine passivation, and we can clearly see that the integrated current density slightly increases from 21.05 to 21.10 mA/cm2. The steady state current output densities and the corresponding maximum power point are shown in Figure 4c, the device shows maximal stable photoelectric output increase by 0.3 mA/cm2 at different bias voltage, in which we calculate that the stable state PCE increase from 18.2% to 19.6%. Figure 4d shows that the PCE distribution of PSCs without or with chlorine passivation and the average PCE value increases from 17% to 19% and a maximum PCE reach at 20% under air atmosphere at room temperature. At the same time, we investigate the Voc distribution without or with interfacial passivation and the results are shown on the upper part of Figure 4d, indicating the increased the average open-circuit voltage overtly from 1.15 to 1.18 V with chlorine passivation. The schematic view of the device structure employed the pristine and passivated SnO2 as ETL are shown in Figure 5a and 5b, in which the interfacial chlorine



passivation enhances the contact of ETLs and perovskite at the interface. The energy band level diagrams of the PSCs are shown in Figure 5c and 5d. The chlorine passivated SnO2 ETL possesses better alignment with perovskite layer, which can efficiently reduce barrier of the energy level and enhance extraction of the electron. In order to describe the energy structure of the device, the UPS spectra and the absorption spectra of the pristine and passivated SnO2 are measured and the results are shown in Figure S6. When the chlorine passivates the surface of SnO2 ETL, the Fermi energy elevates from -4.34 eV to -4.26 eV as shown in Figure S6b. Meanwhile, Figure S6c shows that the elevated Fermi energy level reduces the difference value between the Fermi energy level and the conduction band (from 0.21 eV to 0.19 eV in Figure S8), which means the increased conductivity with chlorine passivation, conforming to the Hall measurement results in Table S1. Figure S6d shows that the absorption spectra of SnO2 films with or without chlorine passivation and the homologous bandgap of pristine SnO2 is 3.88 eV and the passivated SnO2 is 3.89 eV. Figure S7a shows the transmission spectra of SnO2 films spin-coated on quartz glass substrates with or without chlorine passivation (absorption of quartz glass has been subtracted). Figure S7b shows that there is little difference for the Tauc plot spectra of perovskite layer based on pristine and passivated SnO2 ETL. And we calculate the bandgap of the perovskite is 1.58 eV deposited on both ETLs. According to the results of UPS and Tauc plot, we draw the energy level diagram as shown in Figure S8. Here, the potential barrier between perovskite and ETL decreases from 0.26 eV to 0.20 eV. The reduced potential barrier would quicken the transport of electron and enhance the


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extraction of charge. In order to survey the interfacial carrier transport property between perovskite and ETL, the steady-state photoluminescence spectra are measured based on the structure of Glass/FTO/ETLs/perovskite and the results are shown in Figure 6a. We think that the higher electron mobility (Table S1) could lead to the weaker luminescence peak intensity compared with pristine SnO2 at 780 nm.33 Meanwhile, the time-resolved PL spectra are presented in Figure 6b. The decay curve is fitted as following function (4)

y(t) = yo + A1e

(- t ) τ1

+ A2 e

(- t ) τ2

(4)

The main fitted data are shown in Table 2. The two decay processes could happen: ① the free carrier quench at the interface of perovskite and ETLs could lead to the fast decay process. ② the radiative recombination of carrier could cause the slow decay. The fast decay decreases from 13.2 ns to 11.5 ns, indicating that the charge transport from perovskite to passivated SnO2 is faster than pristine SnO2. Meanwhile, the fraction of the fast decay process increase from 31% to 37%, which attribute to the passivated SnO2 illustrates better electron extraction capacity. On the contrary, the fraction of the slow decay process decreases from 69% to 63%, which could originate from the reduction of the traps in the bulk with minimal carrier loss 50, 51.

To further investigate the interface electronic property of PSCs, the

Electrochemical Impedance Spectroscopy (EIS) measurements of devices were performed and the results with/without chlorine passivation are presented in Figure 6c. (The inset shows the fitted equivalent circuit in Figure 6c) It is obvious that the



Rrec and Rs resistive elements displayed an essentially dependence on the potential of open-voltage because the increasing bias accelerates the electron accumulation at the interface of perovskite and ETL, thus increases the recombination of the electrons and the holes. In addition, the series resistance (Rs) is originated from the structure of device that represents the resistance of electron transport. However, the recombination resistance (Rrec) means the resistance of carrier recombination.52 As listed at Table 3, the Rs decreases from 26 to 13  in comparison to the Rrec increases from 270 to 517  with the surface of chlorine passivation, suggesting that the interface passivation reduces the carrier recombination and enhances the electron transportation. The Hall measurements shows that the chlorine passivated SnO2 ETLs possess suitable carrier density and higher electron mobility, which are beneficial for extraction of carriers and transport of electron (As shown in Table S1). To further explore the defect state density at interface of ETLs and perovskite layer, we employ the space charge-limited current (SCLC) measurements of single carrier device as presented in Figure 6d. The Ohmic region, trap-filling limited region and SCLC region are fitted in the SCLC curve, where the linear relation is broken with the bias voltage increase. The trap-filled limit voltages (VTFLs) decreases from 0.18 to 0.11 V by the chlorine passivation, where the trap state density (ntrap) can be estimated as equation (5)

VTFL =

entrapL2 (5) 2εοε

The e is elementary charge (e=1.610-19 C), the L represents the thickness of perovskite layer, o is the permittivity of vacuum (o=8.810-12 F/m),  is the relative


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dielectric constant (here with reference to 62.23). According to the result of VTFLs, we calculate the ntrap that decreases from 4.931015 to 3.011015 cm-3. We think that the surface chlorine of SnO2 ETLs could connect with lead atom of perovskite, which is beneficial for the formation of PbCl2-terminated interface. Owing to the higher formation energy of Pb-Cl antisite defect state, the deep energy level defect and nonradiative recombination are suppressed in the presence of Cl atom.47 Generally speaking, the Cl atoms at the interface reduce the interfacial carrier recombination and enhance the binding energy between perovskite and SnO2.

4. CONCLUSIONS: In summary, the chlorine has successfully been bound on the surface of SnO2 QDs using a convenient method. And then, we found that the introduction of chlorine not only regulate the energy level structure but also enhance the contact between ETLs and perovskite layer. Through this passivation strategy, the elevated the Fermi-level of ETL make better energy alignment between the ETL and perovskite layer, which will reduce the energy level barrier, increase carrier extraction and enhance open-circuit voltage. Moreover, profiting from the stronger Cl-Pb bonding effect, the chlorine passivation effectively decreases the perovskite defect state density at the interface. As a result, the Voc of the devices increased from 1.135 to 1.195 V and the champion PCE increased from 18.2% to 20% attribute to chlorine passivation.



ASSOCIATED CONTENT Supporting Information Digital optical images of SnO2 quantum dot solutions; AFM image; SEM-EDS mapping image; Tauc plots of perovskite layer; cross-sectional SEM image and top surface SEM image; photovoltaic J-V curves of fabricated PSCs using passivated SnO2 with different annealing temperatures; photovoltaic J-V curves of PSCs using different concentration of NH4Cl passivated ETLs; fitted data of EIS; Hall measurement.

AUTHOR INFORMATION Corresponding Authors *E-mail: [email protected] (G.F.). *E-mail:[email protected] (C.T.). ORCID Chen Tao: 0000-0001-6410-6208 Guojia Fang: 0000-0002-3880-9943 Notes The authors declare no competing financial interest.

ACKNOWLEDGMENTS The work was supported by the National Natural Science Foundation of China (11674252) and Special Funds for the Development of Strategic Emerging Industries


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in

Shenzhen

(JCYJ20170818113036217).



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Figure 1 TEM images of pristine and chlorine passivated SnO2 QDs as shown in (a) and (b), SEM image of bare FTO as presented in (c), pristine and chlorine passivated SnO2 ETL films on FTO annealed at 200℃ for 1 hour as shown in (d) and (e).



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Figure 2 XPS of pristine SnO2 and passivated SnO2, fully spectra (a), Sn3d, O1s and Cl2p XPS spectra as presented in (b), (c) and (d).

Table 1 XPS peak position and area ratio of OH and OSn. O 1s

OSn

OH

Position (eV)

530.88

532.23

Area Ratio

70.5%

29.5%

Position (eV)

531.08

532.58

Area Ratio

77.5%

22.5%

SnO2

SnO2:NH4Cl


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Figure 3 (a) and (b) are the top surface SEM images of perovskite deposited on pristine and passivated SnO2 ETLs and (c) and (d) are the cross-sectional SEM images of the device with pristine and passivated SnO2 ETLs.



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Figure 4 The forward and reverse J-V curve of PSCs based on pristine SnO2 and chlorine passivated SnO2 ETL (a), the EQE spectra and integrated current density (b), the steady state PCE at different bias voltage (c), the statistics of 20 devices in Voc and PCE both for pristine and chlorine passivated SnO2 (d).

Table 2 Photovoltaic parameters obtained from the J–V curves. Samples

Scan Direction Voc (V) Jsc (mA/cm2) FF (%) PCE (%) Forward

1.135

21.8

74

18.3

Reverse

1.135

22.1

71

17.8

Forward

1.195

22.1

75.6

20

Reverse

1.195

22.2

73

19.4

SnO2

SnO2:NH4Cl


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Figure 5 (a) and (b) Schematic view of the device structures using pristine and passivated SnO2 as ETLs. (c) and (d) energy band diagram of the devices using pristine and passivated SnO2 as ETLs.



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Figure 6 The photoluminescence spectra of perovskite deposited on pristine and passivated SnO2 ETLs (a), the corresponding time-resolved PL (b), the Nyquist plots of SnO2 and NH4Cl passivated SnO2 (c), the SCLC curves in the dark situation (d).

Table 2 The fitted parameters of TRPL spectra. Samples

A1

1

A2

2

avg

SnO2

31%

13.2 ns

69%

338.8 ns

333 ns

SnO2:NH4Cl

37%

11.5 ns

63%

282.2 ns

276 ns


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Table 3 the EIS parameters for the PSCs based on the pristine and passivated SnO2 ETLs. Samples

Rs ()

Rrec ()

SnO2

26.231.07

270.56.8

SnO2:NH4Cl

12.892.37

517.138.1



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Page 34 of 34

Achieving high open-circuit voltage on planar perovskite solar cells via

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