Highly Efficient Perovskite Solar Cells with Gradient Bilayer Electron

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Highly Efficient Perovskite Solar Cells with Gradient Bilayer Electron Transport Materials Xiu Gong, Qiang Sun, Shuangshuang Liu, Peizhe Liao, Yan Shen, Carole Graetzel, Shaik Mohammed Zakeeruddin, Michael Grätzel, and Mingkui Wang Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.8b01440 • Publication Date (Web): 21 May 2018 Downloaded from http://pubs.acs.org on May 21, 2018

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Highly Efficient Perovskite Solar Cells with Gradient Bilayer Electron Transport Materials Xiu Gonga, Qiang Sun a, Shuangshuang Liu a, Peizhe Liao a, Yan Shen a, Carole Grätzel b, Shaik M. Zakeeruddin b, Michael Grätzel b and Mingkui Wang *,a a

Wuhan National Laboratory for Optoelectronics, School of Optical and Electronic

Information, Huazhong University of Science and Technology, Wuhan 430074, Hubei, P. R. China b

Laboratoire de Photonique et interfaces (LPI), Ecole Polytechnique Federale de Lausanne

CH-1015 Lausanne, Switzerland

ABSTRACT: Electron transport layers (ETLs) with suitable energy level alignment for facilitating charge carrier transport as well as electron extraction are essential for planar heterojunction perovskite solar cells (PSCs) to achieve high open-circuit voltage (VOC) and short-circuit current. Herein we systematically investigate band offset between ETL and perovskite absorber by tuning F doping level in SnO2 nanocrystal. We demonstrate that gradual substitution of F- into the SnO2 ETL can effectively reduce the band offset and result in a substantial increase in device VOC. Consequently, a power conversion efficiency of 20.2% with VOC of 1.13 V can be achieved under AM 1.5 G illumination for planar heterojunction PSCs using F-doped SnO2 bilayer ETL. Our finding provides a simple pathway to tailor ETL/perovskite band offset to increase built-in electric field of planar heterojunction PSCs for maximizing VOC and charge collection simultaneously. KEYWORDS: doping; electronic material; Fermi level; perovskite; photovoltaic; charge transfer 1 ACS Paragon Plus Environment

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1. Introduction The power conversion efficiency (PCE) of emerging organic-inorganic halide perovskite solar cells (PSCs) has rapidly risen from an initial 3.8% efficiency in 2009 to a recently certified 22.7% due to their intriguing properties,1,2 including high ambipolar carrier transport,3 large light absorption coefficients,4 as well as long charge diffusion lengths.5 The photovoltaic performance of planar heterojunction (PHJ) PSCs, using recently developed techniques such as interface, composition engineering or latest incoming nanoscale perovskites, has been largely augmented to a level of being close to their mesoscopic counterparts.6-11 Indeed PSCs hold potentials for producing high efficiency at low cost. Furthermore, the characteristics of semitransparent, flexible, thermal stability and light-weight of hybrid perovskites have opened up paths to diverse applications for solar cells. 11,12 The high PCE achieved for PHJ PSCs could be attributed to a large short circuit current (JSC, 22~23 mA cm-2 under standard testing conditions by considering light scatering for instance),13,14 which is approaching the theoretical value with deducting reflective and parasitic absorption losses (23.8 mA cm-2 for CH3NH3PbI3 )15. However, we have observed that the open-circuit voltage (VOC, ~1.1 V) remains lower than the predicated value of 1.32 V from the Shockley Queisser limit for perovskite compounds with band-gap of 1.59 to 1.63 eV.16 Several approaches have been proposed to further increase the PCE of PHJ PSCs by boosting their VOC. The open-circuit voltage of perovskite devices depends on several factors like the optical band-gap of perovskite compounds, the quasi-Fermi level difference between electron transport layers (ETLs) and hole transport layers (HTLs), relative conduction/valance band energy levels, the generation rate of bound electron-hole pairs and the dissociation probability of bound electron-hole pairs into free charge carriers, interfacial recombination as well as non-radiative recombination.16-18Among these factors, the energy band difference (i.e., energy band offset) between the ETL/perovskite and HTL/perovskite interfaces, offers the 2 ACS Paragon Plus Environment

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driving force (about 150 mV for each side) for efficient charge transfer, reducing interfacial and non-radiative recombination, and thus has a significant influence on the VOC. To achieve an ideal VOC, therefore, it is crucial to build up a functional interface with well-matched energy bands and superior electron mobility and conductivity to avoid excessive band offset while accelerating carrier extraction and transport. To date the TiO2 nanocrystal (NC) thin film is one of the most often used ETLs in PSCs. However, the optoelectronic characteristics of TiO2 nanocrystal still exhibit some shortfalls, such as low electron mobility (0.1-1 cm2 V-1 s-1) and conductivity (~1.1×10-5 S cm-1). The PSCs using TiO2 as ETL are likewise sensitive to UV illumination, which has become one critical issue for practical application in terms of long-term stability.19-21 Therefore, several alternatives have been proposed as effective and promising ETLs to the conventional TiO2.22,23 For example, the semi-conducting SnO2 with a wide optical band-gap (3.6-4.0 eV) has attracted great attention because of its excellent chemical stability, high electron mobility and conductivity that can be tailored easily through low cost processes such as deposition approach, doping, and composition engineering. Fang et al. firstly reported solution-processed SnO2 ETL for efficient planar PSCs showing an overall PCE over 17%, which was attributed to good antireflection and high electron mobility.24 Nonetheless the PCE under forward voltage scans was reduced to 14.82% due to a large density of surface traps on the assynthesized SnO2 and inefficient electron transfer at the SnO2/perovskite interface. Without further surface/interface modification, these traps eventually result in photocurrent hysteresis phenomenon and poor fill factor (FF). 25-27 Lately, doping metal ions (Li+, Nb5+, Y3+, Sb5+, etc) into the SnO2 ETL has led to improvement of the device photovoltaic performance. 28-30 For instance, p-dopant such as Li+ or Y3+ induces a downward shift of the conduction band (CB) minimum of SnO2 and thus enlarges the band offset to accelerate the injection and transfer of electrons from perovskite absorber to the ETL. To avoid VOC losses by the ETL, it must be assured that at open circuit the quasi-Fermi level of electrons set up under illumination in the 3 ACS Paragon Plus Environment

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perovskite film is equal or higher than that attained in the absence of the ETL. This implies that the rate of interfacial recombination of electrons injected into the ETL with the holes in the perovskite must not exceed their intrinsic rate of carrier recombination in the perovskite. Thus, it is imperative to develop new ETLs with low interfacial carrier recombination velocity. By tailoring the Fermi energy some ETLs seem to meet this requirement by forming wellmatched energy levels with perovskite active layer while reducing interfacial carrier recombination (Table S1). Herein we reported a new ETL via fabrication of F-doped SnO2 (F:SnO2) nanocrystals using a facile low-temperature solution-process method for efficient n-i-p planar PSCs. We further demonstrated that the device VOC could be tailored by gradually moderation of the band offset at the interface of perovskite active layer and F-doped bilayer SnO2 ETL via tuning Fermi level of the latter. 31,32 Grain boundary barriers and defects within the bilayer ETL can be largely restrained due to a matched lattice constant. Consequently, efficient PHJ PSC devices using bilayer ETL can be achieved with a PCE of 20.2%, a VOC of 1.13 V, a short-circuit current density (JSC) of 22.92 mA cm-2, FF of 78%. The advantages of bilayer ETL synthesized with low-temperature processing can be easily foreseen. This represents an important proof of concept that paves the way to further realization of high-performance and large-scale production of PSCs based on bilayer transparent substrates with different carrier concentration. 2. Results and Discussion The F:SnO2 films with different doping levels were prepared by spin-coating SnCl2·2H2O and NH4F in ethanol solution onto pre-cleaned FTO substrates. Then the films were thermally annealed at 180 °C, 380 °C and 500 °C for 1 hour in air and are represented as (F:SnO2)180, (F:SnO2)380 and (F:SnO2)500, respectively. The (F:SnO2)180 films with 0.1 molar ratio of F/Sn doping are denoted as (F:SnO2)180-0.1 for clarity and brevity serving as an example. Figure S1 presents XRD patterns of the as-synthesized pristine SnO2 and F:SnO2 4 ACS Paragon Plus Environment

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films. Three major peaks located at 26.62°, 33.91°and 51.76° correspond to (110), (101), and (211) diffraction planes respectively, which are well-matched with tetragonal rutile SnO2.33 Moreover, no additional diffraction peaks from any other impurity phases were detected in the F:SnO2 thin films. This indicates the SnO2 lattice structure has not been disrupted by introducing F atoms. We found that the sample crystalline quality was improved with temperature but decreased with F-doping concentration (Figure S1). This could be caused by lower melting point of SnOF compound as well as by thermally promoted ion diffusion during crystal growth associated with incorporation of F into SnO2 NC lattices.34 Figures 1a and 1b present transmission electron microscopy (TEM) images of F:SnO2 NCs of 6-9 nm size that is evenly distributed. A clear lattice spacing of 0.332 nm, corresponding to (110) facet of rutile SnO2, can be observed in the high resolution TEM image of F:SnO2 NCs, which further demonstrates their high crystallinity. Moreover, the selected area electron diffraction pattern of F:SnO2 NCs also confirms tetragonal structure. Figure S2 shows mapping of Sn, O and F in the scanning electron microscope-energy dispersive spectrometer, further indicating the homogeneously distribution of F. Figures S3-S5 display top-view scanning electron micrographs (SEM) images of SnO2 and F:SnO2 films annealed at different temperatures. The (F:SnO2)180 films with low doping ratio exhibited conformal and pin-hole-free film coverage (Figure S3). As for the high doping ratios, at 0.7 and 1 for example, the films presented obvious small cracks (marked with red circle in Figures S3g and S3h). The samples annealed at 380 oC displayed slightly bigger particle size with tiny agglomeration (Figure S4), indicating the importance of annealing temperature for crystallization of F:SnO2 NCs. However, significant pinholes and cracks were easily observed for the samples annealed at 500 oC (Figure S5). This situation becomes even more significant for samples of high doping ratio.35 Therefore, we selected (F:SnO2)180 and (F:SnO2)380 as ETL for PSC device fabrication during the following experimental studies in avoid of high leakage current due to low coverage or high pinholes. 5 ACS Paragon Plus Environment

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To accurately obtain the band-gaps of films and to avoid the shielding effect from FTO substrates,36 we measured the transmission spectra of SnO2 and F:SnO2 films coated on quartz substrates (Figure S6). Optical band-gaps were calculated from linear line portion of the plot of (αhv)2 versus (hv) based on the equation of αhv=A(hv-Eg)2 (Table S2).37 As depicted in Figure 1c, the optical bandgap slightly broadens with not only increasing temperature but also F doping level (less than 0.2 F:Sn molar ratio). These observations can be explained with the filling of states in the conduction band due to lifting of the Fermi level as a function of increased carrier concentration according to the Bursteine-Moss effect.38 However, further augmentation in F doping could narrow the band gap due to the many-body interaction effect between ionized impurities and free carriers. 39 Figures S6a and 6b compare the transmission of the SnO2 and F:SnO2 film deposited on FTO substrates under different temperatures. The F:SnO2 film with 0.2 F doping showed a slightly higher transmittance than films without doping. This can be beneficial for device light harvesting. The thin films electrical parameters including carrier concentration (n), mobility (µ), and electrical conductivity (σ) were investigated as a function of F doping ratio (Figure S7 and Tables S3). As clearly shown in Figure S7b, the carrier concentration and electrical conductivity increases with the annealing temperature at the same F content. All these parameter values first increase and then decrease along with F doping under a constant temperature, and achieving maximum at 0.2 doping, except for electron mobility. This is consistent with the change of optical band-gap. Clearly, an appropriate F doping can improve electrical properties of SnO2 films. This once again proves that replacing O ions by F ions leads to higher free electron concentration in films. The mobility depends heavily on ionized impurities and grain boundaries according to the scattering mechanism. 40,41 The density of electrons (nc) in conduction band for semiconducting NCs is given by  E − EC nc = N C exp F  K BT

   

(1)

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where NC is the accessible density of electronic states in the conduction band, kB the Boltzmann constant, T the temperature, Ec the position of the lower edge of the CB, and EF the position of the Fermi level of the semiconductor. Therefore, assuming an identical density of electron states NC, a variation of electron concentration would induce a shift in the Fermi energy level according to Equation 2.

n E F1 = E F2 + K BT ln c1  nc2

   

(2)

where nc1 and nc2 are the carrier concentration for different SnO2 films, EF1 and EF2 are the corresponding Fermi levels, respectively. Accordingly, Figure 1d compares the Fermi levels from Hall characterization on various F:SnO2 films at different annealed temperatures, and the corresponding values are summarized in Table S3. The Fermi levels of F:SnO2 films at a constant temperature follow a similar tendency with carrier concentration by increasing with the F doping ratio, reaching the highest value at 0.2 F doping. In this case, the carrier concentration increases from 2.04×1014 and 3.49×1015 to 1.90×1015 and 9.07×1016 cm–3 for the (F:SnO2)180 and (F:SnO2)380 samples, respectively. Thus, the up-shifted of Fermi level by doping can be evaluated to be around 68 meV and 95 meV (Table S3), respectively. This shift was confirmed with Kelvin probe microscopy measurements on the work function of F:SnO2 based on FTO (Table S4). A reduced work function of FTO cathode with F:SnO2 can be of benefit to free electron extraction and transport. These findings were expected to augment VOC and FF of PSCs which have indeed been observed in this study. We fabricated PSCs with a planar heterojunction architecture of FTO/ETL (50 nm)/ (FAPbI3)0.85(MAPbBr3)0.15 (500 nm)/spiro-OMeTAD (200 nm)/Au (100 nm) by varying the ETL while maintain the rest constant. Figure 2 shows a process flow for device fabrication and a clear cross-sectional SEM image of PSC device. Figure S8 compares the photocurrent density-voltage (J-V) characteristics of PSCs measured under AM1.5G irradiation (100 mW cm–2). The relevant performance parameters are summarized in Tables S5 and S6. Under the 7 ACS Paragon Plus Environment

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same F content, the (F:SnO2)180-devices exhibited better photovoltaic performance than that the (F:SnO2)380-devices. This is mainly contributed to the enhanced JSC and FF, stemming from the slightly high electrical conduction in the ETL as verified in Figures S8c and S8d. For the low FF of (F:SnO2)380-device, however, it can be ascribed to poor contact correlated with nanoparticles agglomeration in this ETL film. Importantly, the photovoltaic parameters were increased with F doping under constant temperature and reached the maximum at 0.2 doping level, especially the VOC (Figures S8c). Specifically, as the doping level increased, the VOC increased from 1.03 and 1.04 to 1.1 V and 1.14 V for the (F:SnO2)180-devices and the (F:SnO2)380-devices, respectively. This result clearly illustrates the correlation between VOC and electronic Fermi energy level of ETL. As expected, the upward-shifting of the Fermi level by moderation F doping can effectively decrease the band offset between ETL and perovskite absorber (Figure S9). We noted that at higher F doping above 0.2 the increased value of VOC by 70 and 100 mV for both devices was obviously larger than the up-shifted Ef of 58 and 84 meV. This fact can be explained by a slight widening of bandgap (< 60 meV) for these layers, resulting in a shift of the ETL CB to vacuum. This corresponds to a decrease excessive energy loss of electron transfer in the ETL/perovskite and improvement in device VOC as payback. This argument was further verified by UPS measurements (Figures S10 and S11). Therefore, the devices with (F:SnO2)380-0.2 ETL delivered a PCE of 16.50% with VOC of 1.14 V, JSC of 20.85 mA cm-2 and FF of 0.69, which was higher than that of the (SnO2)380-0 ETL devices (the PCE being 14.04%). The boosting efficiency could be mainly ascribed to an increased VOC from 1.04 V to 1.14 V. The devices’ PCEs further increased to 18.10% when (F:SnO2)1800.2 ETL was utilized, resulting from an improvement in JSC (21.96 mA cm-2) and FF (0.75). As discussed above a high VOC could be achieved for PSC devices with (F:SnO2)380-0.2 ETL by tailoring electrical conductivity and band offset. However, the device FF must be further improved to enhance the performance. Herein, we introduced the concept of bilayer ETL with a gradient F-doped SnO2 nanocrystal for producing highly efficient PSCs. In this 8 ACS Paragon Plus Environment

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architecture, the configuration of bilayer ETL comprises an inner layer (noted as D1) employed on top of FTO substrate and outer layer (noted as D2) deposited on top of the inner layer. The D1 and D2 layers should have large differences in carrier concentrations of D1 > D2 aiming to imitating a space-charge region and generating a built-in electric field (Ebi), while both layers should possess an appropriate band alignment (i.e., CBD1CBD2 might be valid when considering the D1 layer could play a role of energy barrier to inhibit recombination of carriers,32,42 it is significant that these conditions can be closely satisfied by specially tuning F doping SnO2 ETL in this study. Therefore, a thin (F:SnO2)380-0.2 film (~20 nm, n=9.07×1016 cm–3) was applied onto the FTO substrate as D1 layer with the property of modifying surface work function and accelerating electron transport. A (F:SnO2)180-0.2 layer (~40 nm, n=1.90×1015 cm–3) was deposited as D2 layer, aimed at facilitating charge carrier extraction from the light absorber. The D1 and D2 layers have large difference in carrier concentration but small difference in CB edge energy levels. Therefore, their combination guarantees a concentration gradient within the bilayer ETL, which in turn induces a space-charge region and thus enhances Ebi as illustrated in Figure 3e. Figure 3a shows the J-V characteristics of the corresponding devices under AM 1.5 G conditions (100 mW cm-2). The device A with bilayer ETL exhibited an overall PCE of 20.2% with JSC of 22.92 mA cm−2, VOC of 1.13V, and FF of 0.78. Compared with the reference device B based on SnO2 ETL, the augmented conversion efficiency for device A could be mainly ascribed to the higher VOC and FF. There is a negligible difference in J-V curves with forward and reverse scans for device A (Figure 3a and Table 1). Moreover, the average PCE (~18.42%) of device A was about 20% higher than that of reference devices and exhibited highly efficient reproducibility (Figure 3b). After aging for 300 h, device A showed more stable PCE performance than that of the device B (Figure S12). The improved stability could be attributed to hydrophobic F--doping in bilayer ETL. Figure 3c shows the external quantum 9 ACS Paragon Plus Environment

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efficiency (EQE) spectra of the corresponding PSCs. Both devices exhibited high EQE values of above 80% in the range from 400 to 700 nm. Compared with device B, device A showed a relatively higher EQE, explaining its higher photocurrent. The perovskite films deposited onto either bilayer F-SnO2 ETLs or SnO2-ETL presents almost identical absorbance (Figure 3d). Figure 4 further compares the SEM images of two ETL films (a and b) and perovskite films (c and d) on them. Compared to the SnO2-ETL (Figure 4a), the bilayer ETL films exhibit smooth and uniform morphology (Figure 4b) with a low surface roughness (RMS=4.05 nm, Figure S13). The insets in Figure 4a and 4b are the SEM images with highmagnification. Both perovskite films exhibit similar morphology when deposited on these different ETLs. These results suggest that the high-performance of PSCs using bilayer ETL might be due to the judicious energy level alignment and electrical properties of ETL/perovskite interface. The Ef of (F:SnO2)380-0.2 and (F:SnO2)180-0.2 was calculated to be 4.41 and -4.45 eV respectively using ultraviolet photoelectron spectroscopy characterization (Figures S10 and S11). The corresponding VB (valence band) was found to be -8.03 eV and 8.07 eV, respectively. Therefore, the CB was further evaluated to be -4.09 and -4.15 eV for them, respectively. These two layers were used to construct the bilayer ETL in this study. For the SnO2 ETL, the corresponding band value of Ef, VB, and CB were found at -4.61, -8.28, and -4.41 eV, respectively. Clearly, doping SnO2 with F can up-shift its Ef and VB energy levels (versus vacuum). This indicates that device A can reduce the splitting of quasi-Fermi levels of perovskite light absorber layer under illumination due to the up-shifted Fermi level (Figure 3f), and thus decrease band offset at the bilayer ETL/perovskite interface.30,43 The decrease in band offset was reflected by the experimental observation in VOC. The built-in potential at the ETL/perovskite heterojunction interface was estimated to be 1.06 V for device A by performing Mott-Schottky plots conducted under dark condition (Figure 5a). 44-46 This value is larger than that of device B (0.95 V). This difference could be attributed to an extended depletion region due to charge diffusion motion induced by carrier 10 ACS Paragon Plus Environment

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concentration differences in bilayer F:SnO2 ETL. Figure S13 presents the schematic of builtin electric field. This built-in potential assists in the separation of electrons and holes as well as their transport. Therefore, a higher built-in potential would be expected to improve photoexcited carrier extraction, and thus decrease carrier recombination and an injected (dark) current to increase the VOC. The reverse bias saturation current J0 of bilayered ETL baseddevice was obviously suppressed in comparison to that contrast devices in the fitted dark current-voltage curves (Figure S14). A smaller J0 would result in a higher VOC as verified by Equation 3. Voc =

 KT  J sc ln + 1 q  Jo 

(3)

Moreover, the fitting of dark curves with heterojunction solar cells’ model delivered less series resistance of 0.75 Ω cm2 for device A compared with that for device B (1.23 Ω cm2), with a bigger shunt resistance of 9.5 kΩ cm2 compared to 4.2 kΩ cm2 (Figure S15). These results are consistent with a higher FF of device A. To elucidate the dynamics of charge-transfer and extraction processes, steady-state photoluminescence (PL) and time-resolved PL (TRPL) were measured with the perovskite films deposited on the bilayer doped or undoped ETL substrates (Figure 5b). The lifetimes (τpl) obtained from TRPL measurements were determined to be 3.91 and 8.42 ns for the bilayer ETL and undoped ETL, respectively (Figure 5c). The shorter PL lifetime and efficient PL quenching for the bilayer ETL suggest its electron extraction ability. This might be caused by an enhanced driving force by an enlarged Ebi to efficient collect free carriers. Nanosecond transient absorption spectroscopy (ns-TAS) characterization was further performed to understand the effect of bilayer ETL on charge transfer at the ETL/perovskite interface. As shown in Figures 5d and 5e, there are a photo-bleaching (PB) negative peak at around 760 nm and a photo-absorption (PA) positive peak of 500-600 nm. The PA positive peak at about 500600 nm can be assigned to the absorption of transient species.47-49 The feature PB negative 11 ACS Paragon Plus Environment

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peak

appearing

at

760

nm

can

be

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to

the

exciton

transition

of

(FAPbI3)0.85(MAPbBr3)0.15 film. The intensity of PB peak around 760 nm obviously decreases in glass /bilayer ETL/perovskite films, indicating that a fast exciton transition from the perovskite film (Figure 5d). The kinetics decay traces of PB peaks were further evaluated as shown in Figure 5f. The lifetime constant (τTAS from single exponential function fitting) of bilayer ETL /perovskite (~356.5 ns) is much higher than that of undoped SnO2 ETL/perovskite (~198.7 ns). Consequently, the interfacial charge recombination velocity (defined as k=1/τTAS) can be evaluated to increase from 2.81×106 s-1 to 5.03×106 s-1. This indicates that photo-excited carriers can be collected quickly and thus effectively suppressed carrier recombination by bilayer F:SnO2 as ETL. Electronic impedance spectroscopy measurements were carried out to further explore charge transfer dynamics at ETL/perovskite interfaces (Figure S16). Figures 6a and 6c compare the interfacial recombination resistance (Rct), capacitance (C), and correlated lifetime τimp (τimp=Rct×C) for two devices. Device A consistently shows larger Rct and longer τimp than those of device B under the same bias, indicating the retarded interfacial charge recombination in the former. Furthermore, less interfacial charge accumulation exists in device A reflected by smaller capacitance as shown in Figure 6b. Therefore, we plotted the density of trap state distribution in Figure 6d obtained by fitting the capacitance curves with an exponential function.50,51 In contrast, device A has an apparent smaller density of states compared with device B at a given energy level, offering less active sites for non-radiative recombination. In short, we ascribe less charge accumulation at the ETL/perovskite interface of device A to an increased built-in field and an up-shifted Fermi level, ultimately leading to less charge recombination and increased facility in charge transport. 3. Conclusion In summary, we systematically investigated highly efficient PSCs with gradient F-doped SnO2 ETL fabricated by a facile low-temperature solution-process method. We introduced 12 ACS Paragon Plus Environment

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bilayer ETLs in efficient n-i-p planar PSCs, achieving 20.2 % PCE with a VOC of 1.13 V, a JSC of 22.92 mA cm-2 and FF of 78%. The gradient F-doped SnO2 nanocrystals and the decreased band offset between ETL and perovskite are beneficial for reducing energy loss and consequently improving VOC. This novel concept of bilayer ETL paves a new path for the future in applications of PSCs. 4. Methods Synthesis SnO2 NCs materials: SnCl2·2H2O (1.128 g, 5 mmol) and 50 mL absolute ethyl alcohol were added to a flask (100 mL), and then magnetic stirring until fully dissolved under uncap. Subsequently, acetic acid (4 mL) was added drop-wise to the flask immersed in a water bath to adjust the PH to 3. Meanwhile tetramethylammonium hydroxide (3-4 mL) was also added drop-wise to the flask as a stabilizing agent to well control the hydrolysis reaction. Then the resultant clear solution was sustained stirring at 70 °C until it turned to a white cloudy dried gel mixture under uncap. Next the resultant of white cloudy dried gel was dissolved using absolute ethyl alcohol, meanwhile NH4F with different doping ratios (0-1 mol) were added to above ethanol solution. Ultimately, the mixed solutions were stirred at 50 °C until the mixture solutions induce to 50 mL, and thus achieving 0.1 M F:SnO2 solution with different F doping ratios on standby. Other experimental details including device fabrication, characterization is shown in the Supporting Information.

ASSOCIATED CONTENT Supporting Information This material is available free of charge via the Internet at The Supporting Information is available free of charge on the ACS Publications website at http://pubs.acs.org.


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Experimental details, characterization with XRD, SEM, UPS, Hall, AFM and photoelectronic properties; AUTHOR INFORMATION Corresponding Authors *E-mail: [email protected] (M.W.). Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes The authors declare no competing financial interest. ACKNOWLEDGMENTS Financial support from the NSFC Major International (Regional) Joint Research Project NSFC-SNSF (51661135023), the Natural Science Foundation of China (No. 21673091), the Natural Science Foundation of Hubei Province (No. ZRZ2015000203), and Technology Creative Project of Excellent Middle & Young Team of Hubei Province (No. T201511). MG. and AMZ thanks the King Abdulaziz City for Science and Technology (KACST) for financial support. The authors thank the Analytical and Testing Centre of Huazhong University of Science

&

Technology for the measurements of the samples. The authors thank the Prof.

Yinhua Zhou (HUST) for work function measurements. The authors thank Prof. Guojia Fang (Wuhan University) for the Hall Effect measurements.


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REFERENCES (1) Kojima, A.; Teshima, K.; Shirai, Y.; Miyasaka, T. J. Am. Chem. Soc. 2009, 131, 6050-6051; (2) NREL Efficiency Chart. http://www.nrel.gov/pv/assets/images/efficiency-chart.png. (accessed November 17, 2017). (3) Etgar, L.; Gao, P.; Xue, Z.; Peng, Q.; Chandiran, A. K.; Liu, B.; Nazeeruddin, M. K.; Grätzel, M. J. Am. Chem. Soc. 2012, 134, 17396-17399. (4) Xu, X.; Li, S.; Zhang, H.; Shen, Y.; Zakeeruddin, S. M.; Graetzel, M.; Cheng, Y.-B.; Wang, M. ACS nano 2015, 9, 1782-1787. (5) Zhang, C.; Sun, D.; Sheng, C.; Zhai, Y.; Mielczarek, K.; Zakhidov, A.; Vardeny, Z. Nat. Phys. 2015, 11, 427-434. (6) Sutherland, B. R.; Hoogland, S.; Adachi, M. M.; Kanjanaboos, P.; Wong, C. T.; McDowell, J. J.; Xu, J.; Voznyy, O.; Ning, Z.; Houtepen, A. J. Adv. Mater. 2015, 27, 53-58. (7) Jeon, N. J.; Noh, J. H.; Yang, W. S.; Kim, Y. C.; Ryu, S.; Seo, J.; Seok, S. I. Nature 2015, 517, 476. (8) Wang, Z.-K.; Gong, X.; Li, M.; Hu, Y.; Wang, J.-M.; Ma, H.; Liao, L.-S. ACS nano 2016, 10, 54795489. (9) Gong, X.; Li, M.; Shi, X. B.; Ma, H.; Wang, Z. K.; Liao, L. S. Adv. Funct. Mater. 2015, 25, 6671-6678. (10) Shai, X.; Wang, J.; Sun, P.; Huang, W.; Liao, P.; Cheng, F.; Zhu, B.; Chang, S.-Y.; Yao, E.-P.; Shen, Y. Nano Energy 2018, 48, 117-127. (11) Shi, E.; Gao, Y.; Finkenauer, B. P.; Coffey, A. H.; Dou, L. Chem. Soc. Rev. 2018, DOI:10.1039/c7cs00886d. (12) Lin, J.; Lai, M.; Dou, L.; Kley, C. S.; Chen, H.; Peng, F.; Sun, J.; Lu, D.; Hawks, S. A.; Xie, C. Nat. Mater.2018, 17, 261-267.

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(22) Tiwana, P.; Docampo, P.; Johnston, M. B.; Snaith, H. J.; Herz, L. M. ACS nano 2011, 5, 5158-5166. (23) Snaith, H. J.; Ducati, C. Nano lett. 2010, 10, 1259-1265. (24) Ke, W.; Fang, G.; Liu, Q.; Xiong, L.; Qin, P.; Tao, H.; Wang, J.; Lei, H.; Li, B.; Wan, J. J. Am. Chem. Soc. 2015, 137, 6730-6733. (25) Zhu, Z.; Bai, Y.; Liu, X.; Chueh, C. C.; Yang, S.; Jen, A. K. Y. Adv. Mater. 2016, 28, 6478-6484. (26) Ma, J.; Yang, G.; Qin, M.; Zheng, X.; Lei, H.; Chen, C.; Chen, Z.; Guo, Y.; Han, H.; Zhao, X. Adv. Sci. 2017, 4, 1700031. (27) Rao, H. S.; Chen, B. X.; Li, W. G.; Xu, Y. F.; Chen, H. Y.; Kuang, D. B.; Su, C. Y. Adv. Funct. Mater. 2015, 25, 7200-7207. (28) Park, M.; Kim, J.-Y.; Son, H. J.; Lee, C.-H.; Jang, S. S.; Ko, M. J. Nano Energy 2016, 26, 208-215. (29) Yang, G.; Lei, H.; Tao, H.; Zheng, X.; Ma, J.; Liu, Q.; Ke, W.; Chen, Z.; Xiong, L.; Qin, P. Small 2017, 13, . (30) Bai, Y.; Fang, Y.; Deng, Y.; Wang, Q.; Zhao, J.; Zheng, X.; Zhang, Y.; Huang, J. ChemSusChem 2016, 9, 2686. (31) Olson, D. C.; Shaheen, S. E.; White, M. S.; Mitchell, W. J.; van Hest, M. F.; Collins, R. T.; Ginley, D. S. Adv. Funct. Mater. 2007, 17, 264-269. (32) Schulz, P.; Edri, E.; Kirmayer, S.; Hodes, G.; Cahen, D.; Kahn, A. Energy Environ. Sci. 2014, 7, 13771381. (33) Moholkar, A.; Pawar, S.; Rajpure, K.; Bhosale, C. Mater. Lett. 2007, 61, 3030-3036 (34) Thirumoorthi, M.; Prakash, J. T. J. Superlattice Microstr. 2016, 89, 378-389. (35) Ke, W.; Zhao, D.; Cimaroli, A. J.; Grice, C. R.; Qin, P.; Liu, Q.; Xiong, L.; Yan, Y.; Fang, G. J. Mater. Chem. A 2015, 3, 24163-24168. (36) Xiong, L.; Qin, M.; Yang, G.; Guo, Y.; Lei, H.; Liu, Q.; Ke, W.; Tao, H.; Qin, P.; Li, S. J. Mater. Chem. A. 2016, 4, 8374-8383. (37) Tauc, J. Amorphous and Liquid Semiconductors, Plenum Press, New York, 1974. (38) Burstein, E. Phys. Rev. 1954, 93, 632. (39) Cao, Y.; Yang, W.; Zhang, W.; Liu, G.; Yue, P. New J. Chem. 2004, 28, 218-222. (40) Lu, J.; Ye, Z.; Zeng, Y.; Zhu, L.; Wang, L.; Yuan, J.; Zhao, B.; Liang, Q. J. Appl. Phys. 2006, 100, 073714. (41) Tran, Q.-P.; Fang, J.-S.; Chin, T.-S. Mater. Sci. Semicond. Process. 2015, 40, 664-669. (42) Lu, H.; Tian, W.; Gu, B.; Zhu, Y.; Li, L. Small 2017,13,1701535.. (43) Y. Shao, Y. Yuan, J. Huang, Nat. Energy 2016, 1, 15001. (44) Perrier, G.; de Bettignies, R.; Berson, S.; Lemaître, N.; Guillerez, S. Sol. Energy Mater. Sol. Cells. 2012, 101, 210-216. (45) Fabregat-Santiago, F.; Garcia-Belmonte, G.; Bisquert, J.; Bogdanoff, P.; Zaban, A. J. Electrochem. Soc. 2003, 150, E293-E298. (46) Zuo, L.; Guo, H.; Jariwala, S.; De Marco, N.; Dong, S.; DeBlock, R.; Ginger, D. S.; Dunn, B.; Wang, M.; Yang, Y. Sci. adv. 2017, 3, e1700106. (47) Yang, Y.; Ostrowski, D. P.; France, R. M.; Zhu, K.; Van De Lagemaat, J.; Luther, J. M.; Beard, M. C.


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Nano Letters

a)

b)

(110) (211)

0.332nm (110)

6nm 10nm

F:SnO2 0.2

c)

d)

(F:SnO2)180 (F:SnO2)380

3.95

3.90

2nm

F:SnO2 0.2

(F:SnO2)180 (F:SnO2)380

80

△ Ef (meV)

Bandgap (eV)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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40

0

3.85 -40

3.80

0

0.1 0.15 0.2

0.3

0.5

0.7

0 0.1 0.15 0.2 0.3 0.5

1

0.7

1

molar ratio of F:Sn

molar ratio of F:Sn

Figure 1. a) and b) TEM and HR-TEM images of F:SnO2 nanocrystals. The inset shows the selected-area electron diffraction. c) The bandgap of (F:SnO2)180 and (F:SnO2)380 deposited on quartz substrates versus F content. Solid squares and solid lines indicate the raw data and fitted curves with Gauss function, respectively. d) The electronic Fermi levels of (F:SnO2)180 and (F:SnO2)380 deposited on quartz substrate versus F content.


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Figure 2. a) Flowchart of fabrication process and b) Cross-sectional SEM image of a typical perovskite device based on F:SnO2.


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Figure 3. a) J-V curves of PSCs based on doped SnO2 and bilayer ETL measured under simulated AM 1.5 G sunlight of 100 mW/cm2. b) Histogram of PCEs measured from 50 bilayer ETL-based PSCs. Inset is the histogram of PCEs measured from 45 doped SnO2-based control devices. c) EQE spectra of PSCs based on undoped SnO2 and bilayer ETL and the integrated current densities of two cells from the IPCE spectra. d) Absorption spectra of perovskite films deposited on undoped SnO2 and bilayer ETL. e) The energy band diagram of the bilayer ETLs based devices in the dark. f) The quasi-Fermi level splitting of the bilayer ETLs based devices at open circuit under illumination. Ec: conduction band; Ev: valence band; Ef: fermi energy; ꭓ: electron affinity;∆Ec: conduction-band offset; ∆Ev: valence-band offset; Vbi: build-in potential; qVbi-1: band bend of interface D1/D2; qVbi-2: band bend of interface D2/perovskite; qVbi-3: band bend of interface perovskite/spiro-OMeTAD; Enf : fermi energy of ETL; Epf : fermi energy of HTL.


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a)

b) 200 nm

200 nm

500 nm

c)c)

500 nm

d) d)

(SnO2 )180

(F:SnO2 )380 /(F:SnO2 )180

500 nm

0.2/0.15

500 nm

Figure 4. SEM images of a) FTO/(SnO2)180, b) (F:SnO2)380-0.2/(F:SnO2)180-0.2. SEM images of (FAPbI3)0.85(MAPbBr3)0.15 perovskite films deposited on c) FTO/(SnO2)180, d) (F:SnO2)380-0.2/(F:SnO2)1800.2. The insets in a) and b) are the corresponding SEM images with high amplification.


Nano Letters

3

2 Vbi = 1.06V

1

Vbi =0.95 V

0

0.6

0.8

Bias (V)

c)

Single SnO2-ETL Bilayer F-SnO2 ETLs

1.0

700

d)

△O.D. 10

15

f) 1.0

Bilayer F-SnO2 ETLs

0.1

0.0 30 ns 130 ns 330 ns 430 ns

-0.1

500

600

500

20

T (ns)

Delta OD(norm.)

5

900

30 ns 130 ns 330 ns 430 ns

-0.1 0

850

0.0

10-2

10-3

800

Single SnO2-ETL

0.1

10-1

e)

750

Wavelength (nm)


100

Intensity (a.u.)

b)


PL intensity (a.u.)

C-2 (1015 F-2)

a)

△O.D.

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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600

700

Wavelength (nm)

800


τ =356.5 ns

0.5

τ =198.7 ns

700

Wavelength (nm)

0.0

800

0

100

200

300

400

Delay Times (ns)

Figure 5. a) C-2-V characteristics of solar cells based on undoped SnO2 and bilayer ETL under dark condition with revise bias. b) Steady-state PL and c) time-resolved PL spectra of perovskite films deposited on undoped SnO2 and bilayer ETL. Transient absorption spectra of d) FTO/(SnO2)180 /Perovskite films and e) FTO/ bilayer ETL/Perovskite films. Excitation at 500 nm for black (30 ns), red (130 ns), blue (330 ns), pink (430 ns), green (430 ns). f) Normalized kinetic traces for photo-bleaching probed at 760 nm for perovskite films deposited on undoped SnO2 and bilayer ETL substrates.


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a)

b)

103

C (nF cm-2)

Rct (W cm2)

103

102

102 device A device B 0.7

0.8

0.9

10 1.0

c)

Energy level (eV.vs HOMO of HTL)

10

device A device B 0.7

0.8

0.9

0.7

1.0

1.1

Voltage (V)

0.8 0.9

1.0

1.1

Voltage (V)

d)

-5

10-6

device A device B

1

1.1

Voltage (V)

τ (s)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

1.0

0.5

device A device B 0.0

0

2

4 6 16

DOS (10

8 -3

10

# cm V)

Figure 6. The interfacial charge transfer parameters for devices A based on bilayer ETL and devices A based on undoped SnO2 substrate obtained by fitting the data from electronic impedance spectroscopy measurements at around VOC under illumination with the equivalent circuit as shown in supporting information: a) Recombination resistance (Rct), b) the corresponding capacitance (C), and c) the apparent recombination lifetime (τimp) as a function of the applied voltage. d) The density of states as a function of electronic energy level calculated with the capacitance from c).


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Table 1. Photovoltaic Parameters of Perovskit Solar Cells Device A based on Bilayer ETL Substrates and Device B based on Undoped SnO2 Substrate.

Samples

Scan direction

Jsc (mA/cm2)

Voc (V)

FF (%)

PCE (%)

Rs (Ω cm2)

Rsh (KΩ cm2)

J0a (mA/cm2)

Device A

RS FS RS FS

22.92 22.46 21.74 20.95

1.13 1.12 1.03 1.01

78.05 77.58 72.28 68.72

20.20 19.65 16.25 14.54

0.75

9.5

5.38×10-7

1.23

4.2

3.25×10-6

Device B a

Rs, Rsh and J0 are obtained from the fitting of dark J-V curves. RS(FS), scan from reverse (forward) direction.


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TOC GRAPHIC


Highly Efficient Perovskite Solar Cells with Gradient Bilayer Electron

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