CH3NH3PbBr3 Quantum Dots Induced Nucleation for High

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CH3NH3PbBr3 Quantum Dots Induced Nucleation for High Performance Perovskite Light-emitting Solar Cells Peng Wang, Jiangsheng Xie, Ke Xiao, Haihua Hu, Can Cui, Yaping Qiang, Ping Lin, Arivazhagan V, Lingbo Xu, Zhengrui Yang, Yuxin Yao, Tao Lu, Zihan Wang, Xuegong Yu, and Deren Yang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b06595 • Publication Date (Web): 11 Jun 2018 Downloaded from http://pubs.acs.org on June 12, 2018

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ACS Applied Materials & Interfaces

CH3NH3PbBr3 Quantum Dots Induced Nucleation for High Performance Perovskite Light-emitting Solar Cells Peng Wang,1,2‡ Jiangsheng Xie,2‡ Ke Xiao,1 Haihua Hu, 3 Can Cui,1* Yaping Qiang,1 Ping Lin,1 Arivazhagan V.,2 Lingbo Xu,1 Zhengrui Yang,2 Yuxin Yao, 1 Tao Lu,

1

Zihan Wang,

1

Xuegong

Yu,2* and Deren Yang2 1

Center for Optoelectronics Materials and Devices, Department of Physics, Zhejiang Sci-Tech

University, Hangzhou 310018, China 2

State Key Laboratory of Silicon Materials and School of Material Science and Engineering,

Zhejiang University, Hangzhou 310027, China 3

Zhejiang University City College, Hangzhou 310015, China

*

Corresponding author: [email protected] (Can Cui), [email protected] (Xuegong Yu)

KEYWORDS: perovskite light emitting solar cells, CH3NH3PbBr3 quantum dots, anti-solvent, heterogeneous nucleation, crystallinity

ABSTRACT

Solution-processed organometallic halide perovskites have developed rapidly in both light emitting diodes (LEDs) and solar cells (SCs). These devices are fabricated with similar materials and architectures, leading to the emergence of perovskite-based light emitting solar cells

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(LESCs). The high-quality perovskite layer with reduced non-radiative recombination is crucial for achieving high performance device, even though the carrier behaviors are fundamentally different in those both functions. Here, CH3NH3PbBr3 quantum dots (QDs) are first introduced into the anti-solvent in the solution phase, serving as nucleation centers and inducing the growth of CH3NH3PbI3 films. The heterogeneous nucleation based on high lattice matching and low free-energy barrier significantly improves the crystallinity of CH3NH3PbI3 films with decreased grain sizes, resulting in longer carrier lifetime and lower trap-state density in the films. Therefore, the LESCs based on the CH3NH3PbI3 films with reduced recombination exhibit improved electroluminescence and external quantum efficiency. The current efficiency is enhanced by an order of magnitude as LEDs, and meanwhile, the power conversion efficiency increases from 14.49% to 17.10% as SCs, compared to the reference device without QDs. Our study provides a feasible method to grow high quality perovskite films for high performance optoelectronic devices.


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INTRODUCTION

Organic-inorganic hybrid perovskite (OIHP) has attracted extensive attention in recent years owing to its superior optoelectronic properties.1 In particular, solution-processed OIHPs of low cost have obtained a rapid progress in solar cells (SCs), with the power conversion efficiency (PCE) exceeding 22% recently.2 In addition to photovoltaic applications, OIHPs exhibit high photoluminescence (PL) quantum efficiency, low exciton binding energy and high color purity, showing great potential for high performance light-emitting diodes (LEDs).3-5 Interestingly, given the similar architectures of perovskite LED and SCs, the emerging device called perovskite light emitting solar cells (LESCs) can perform both functions by acting as a reversible transducer between light and electricity.6 LESC is promising in practical applications, such as solar lamp, which harvests solar energy during the day as SCs and emits light at night as LEDs when it is integrated with a rechargeable battery. Though perovskite SCs and LEDs share the same structure, they perform fundamentally different functions. The SCs require light-generated electron-hole pairs being effectively separated and simultaneously extracted out of the perovskite layer, whereas LEDs require enhanced formation of excitons through the recombination of electrons and holes. Kim et al6 have investigated the energy band alignment of electron transport layers of LESCs to benefit not only the extraction of electrons under illumination to generate power but injection of electrons into the same active layer to produce radiative-recombination induced luminescence. However, despite of the different carrier behaviors in SCs and LEDs, the non-radiative recombination in the bulk or at the surface of perovskite active layer must be suppressed for satisfactory device performance.7-10 A flat and dense perovskite film is essential to attain high efficiency for SCs and excellent electroluminescence (EL) for LEDs.10-13 To date, various


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methods have been developed to fabricate high quality perovskite films.14-21 Dripping antisolvent during the spinning of the perovskite precursor has turned out to be successful and repeatable to fabricate high performance SCs.18 To further improve the quality of perovskite films, a variety of additives have been added in the anti-solvent to control nucleation and crystal growth. For example, poly (methyl methacrylate) (PMMA) was used as a template to enable both faster nucleation and slower crystal growth, modifying the morphology of mixed cation perovskite films with larger oriented grains and fewer defects for high efficiency SCs.22 The αbis-PCBM was added in chlorobenzene anti-solvent to enlarge the grain sizes, passivate the trap states at the grain boundaries, and improve the carrier separation within the mixed cation perovskite films.23 Besides, an organic small molecule, 2,2′,2′′-(1,3,5-benzinetriyl)-tris (1phenyl-1-H-benzimidazole) (TPBI) was successfully used as an additive in chloroform antisolvent to hinder the crystal growth and uniform CH3NH3PbBr3 nanograins were formed, resulting in greatly enhanced performance of LEDs.24 Typically, to improve the current efficiency (CE) of LEDs, the perovskite grain size must be decreased to confine excitons and reduce their dissociation, whereas in the fabrication of high efficiency SCs, large perovskite grains are favorable to achieve good exciton dissociation and carrier diffusion. This arising the critical problem about how to control the grain morphology and size to develop high performance LESCs. Recently, intensive attention has been devoted to quantum dots (QDs) due to the tunable size, the increased exciton binding energy as well as the high PL quantum yields (PLQYs).25-27 Especially, inorganic or organometallic halide perovskite QDs have been reported to serve as the light absorption layer or an interlayer between the perovskite and hole-transporting layer for SCs because of the tunable band edge.28-30 Moreover, inorganic perovskite QDs demonstrate higher


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PLQYs and stability which are highly desired for LEDs.31-32 However, the solution-processed growth of OIHP film based on organometallic halide perovskite QDs has not been explored to date. With the highly similar crystal lattice and composition (CH3NH3PbX3, X = Br, I), the organometallic halide QDs are more compatible with OIHP film and should be beneficial for its growth. Here, we first introduced CH3NH3PbBr3 QDs prepared with short alkyl amine ligands into anti-solvent to control the nucleation and growth of CH3NH3PbI3 films. With similar lattice constants, the CH3NH3PbBr3 QDs act as heterogeneous nucleation centers for CH3NH3PbI3 films, which reduce nucleation free-energy barrier in comparison with the homogeneous nucleation.33 As a result, the obtained CH3NH3PbI3 films exhibit an enhanced crystallinity with reduced grain sizes, leading to significant improvements in the performances of perovskite LESCs.

RESULTS AND DISCUSSION

CH3NH3PbBr3 QDs were firstly synthesized by a convenient and efficient process at room temperature. The short alkyl amine of n-butyl amine (BA) was used as both ligand and stabilizer which was different with the methods reported previously

25, 34

(The method is described in the

supporting information). When BA was added into the CH3NH3PbBr3 precursor solution, the growth of CH3NH3PbBr3 crystals was suppressed and thus nanometer-size crystals (QDs) were produced, similar to the case of addition of n-butylammonium halides35 or other alkylamines with longer chains25 in the former work. Moreover, BA can stabilize the formed colloidal CH3NH3PbBr3 QDs and the QD solution keeps stable after storage for three months at room temperature (Fig. S1).


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Figure 1a and 1b show the typical transmission electron microscopy (TEM) images of CH3NH3PbBr3 QDs. These QDs have excellent dispersion in diethyl ether (DE) (Fig. 1a), with an average diameter of 4.9 nm from the statistic histogram (Fig. S2). As seen in the highresolution TEM (HRTEM) image (Fig. 1b), the measured interplanar distances are 2.12 Å and 1.99 Å, corresponding to the (220) and (300) crystalline plane of cubic CH3NH3PbBr3, respectively. The inset in Fig. 1b is the corresponding selected area diffraction pattern of the QD crystal. The X-ray diffraction (XRD) patterns of the CH3NH3PbBr3 QDs and bulk CH3NH3PbBr3 (Fig. 1c) show that the QDs have boarder peaks with lower intensity corresponding to the small size of QDs, in comparison to the bulk CH3NH3PbBr3. UV-Vis absorption and normalized PL emission spectra of CH3NH3PbBr3 QDs and bulk CH3NH3PbBr3 are shown in Fig. 1d. The absorption spectrum with a band edge at 531 nm and a sharp PL emission peak at 521 nm are detected for CH3NH3PbBr3 QDs. The sample has a relatively small Stokes shift of ~ 45 meV, implying that the PL emission of QDs originates from the exciton recombination. Stronger absorption and blue shift (~ 14 nm) in the PL spectrum are observed for CH3NH3PbBr3 QDs, compared with bulk CH3NH3PbBr3. The full width at the half-maximum (FWHM) of the PL peak is 21 nm for CH3NH3PbBr3 QDs, smaller than that for the bulk CH3NH3PbBr3 (28 nm), indicating the superior color saturation which is comparable to the reported CH3NH3PbBr3 QDs.25, 37 The higher luminous purity of QDs is attributed to the generation and recombination of more stable excitons with increased exciton binding energy.25 Figure 1e shows the schematic diagram of solution-processed preparation of CH3NH3PbI3 films using DE as the anti-solvent, with or without CH3NH3PbBr3 QDs. During the normal preparation, the anti-solvent is dropped onto the spinning precursor films at an appropriate time to extract solvent, followed by the annealing to promote the growth of CH3NH3PbI3 films.18 In


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this work, CH3NH3PbBr3 QDs added into DE provide a large number of nanocrystals with the similar crystal lattices as CH3NH3PbI3, would affect the nucleation and crystallization process of the CH3NH3PbI3 films upon heat treatment. Fig. 1f shows the speculative mechanism of nucleation and growth of CH3NH3PbI3 films based on CH3NH3PbBr3 QDs. When DE with QDs is dropped onto the spinning precursor, the organic ligands (BA) from QD solution are dissolved by residual solvent (dimethylformamide, DMF) in precursor on the substrate, whereas the QDs are retained in the precursor under the protection of DE even if the ligands are dissolved. Thereafter, the QDs act as heterogeneous nucleation centers which reduce the free-energy barrier36 and the CH3NH3PbI3 crystals are grown around the QDs more easily when the substrate with the spinning-coated precursor film is subjected to an annealing. Figure 2 shows the scanning electron microscopy (SEM) images of CH3NH3PbI3 films induced with various concentration of CH3NH3PbBr3 QDs (CQDs). The referenced CH3NH3PbI3 film (CQDs = 0.0) exhibits non-uniform grains, and the average grain size is 323 nm (Fig. 2a). By introducing QDs with different concentration (CQDs = 0.005, 0.01, 0.02 and 0.03 mmol/L) into the CH3NH3PbI3 precursor films, the grain sizes in the grown films decrease because of the QD nuclei with nano-meter sizes. For CQDs = 0.005, the film becomes dense accompanied with the appearance of smaller grains with an average size of 231 nm (Fig. 2b). By increasing CQDs to 0.01, a flat film is obtained and the grain sizes decrease further to an average size of 138 nm (Fig. 2c), since more foreign seeds lead to smaller grains. As CQDs increases to 0.02, the grain sizes decrease slightly with an average size of 119 nm (Fig. 2d). Whereas, some pinholes appear between the grains. For CQDs = 0.03, more pinholes appear and the grain shape even becomes ambiguous (Fig. 2e), indicating excessive nucleation centers accompanied with the dissolved BA could lead to poor morphology of CH3NH3PbI3 film. Figure 2f summarizes the average grain


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size of CH3NH3PbI3 films induced by various CQDs. The crystal lattice constants of CH3NH3PbI3 (a = b = 8.67 Å, c = 12.86 Å) and CH3NH3PbBr3 (a = b = 8.19 Å, c = 12.15 Å) are nearly the same, with the lattice mismatches of 5.5% at (110) crystalline plane.37 It is known that heterogeneous nucleation has lower nucleation free-energy barrier than the homogeneous nucleation.38 Therefore, the growth of CH3NH3PbI3 on CH3NH3PbBr3 heterogeneous nuclei is more favorable than the growth of CH3NH3PbI3 crystals in CH3NH3PbI3 precursor with homogeneous nucleation. The XRD patterns of CH3NH3PbI3 films induced with different CQDs are shown in Fig. 3a. Compared to the reference film, no shift is observed in XRD peak positions in the CH3NH3PbI3 films with CH3NH3PbBr3 QDs. Moreover, no Br element can be detected by X-ray photoelectron spectroscopy (XPS) in Fig. S3, indicating that the small amount of Br (theoretically I : Br = 6400 : 1 for CQDs = 0.01) added into the precursor generates negligible CH3NH3PbBrxI3-x. An enlarged view of the dominant peaks corresponding to (110) and (220) crystalline planes, and the FWHM of the corresponding peaks are shown in Fig. 3b and 3c, respectively. It is clearly seen that for an optimal concentration of CQDs = 0.01, the intensities of the dominant peaks increases more than two times with reduced FWHM, compared with those of the reference film. The results strongly confirm that an enhanced crystallinity of CH3NH3PbI3 films is obtained by nucleating from CH3NH3PbBr3 QDs. The steady-state PL and time resolved PL (TRPL) measurements were performed on the perovskite films treated with different CQDs in order to investigate the carrier behaviors, as shown in Fig. 4a and 4b, respectively. The PL intensity increases dramatically as the CQDs increases, and it reaches the maximum value when CQDs = 0.01, afterwards it decreases with further increasing CQDs (Fig. 4a). The PLQY of the CH3NH3PbI3 films with different CQDs is shown in Fig. S4. The


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change tendency of PL intensities or PLQY is in good agreement with that of XRD patterns of the perovskite films, inferring that an enhanced PL emission is ascribed to the improved crystallinity. When CQDs increases to 0.03, the PL intensity becomes weaker compared to the reference film, due to the increased non-radiative recombination of the film with many pinholes (Fig. 2e). Thus, it is concluded that the enhanced PL or PLQY is attributed to the excellent crystallinity by heterogeneous nucleation and the improved radiative recombination in small grains.24 Fig. 4b shows the TRPL spectra fitted by a mono-exponential decay model for perovskite films with different CQDs on glass substrate, except the curve of the film with CQDs = 0.03, which shows a faster decay at the initial stage and should be fitted by the bi-exponential decay fitting. The characteristic parameters of time constants can be expressed by the following formula,

I (t ) = A1 exp(−t / τ 1 ) + y0 (1) or I (t ) = A1 exp(−t / τ 1 ) + A2 exp(−t / τ 2 ) + y0 (2) where A1 and A2 are the relative amplitudes, ߬1 represents the slow decay from radiative recombination of trapped charges in the bulk of perovskite film, and ߬2 represents the fast decay from quenching of free carriers due to the surface/interface recombination of perovskite film, respectively. The calculated parameters of carrier lifetime are listed in Table S1. All the films treated with QDs have longer bulk carrier lifetime than the reference film (࣎1 = 97.3 ns). More strikingly, the film with the best crystallinity (CQDs = 0.01) has the longest carrier lifetime of 157.1 ns, confirming that an enhanced crystallinity effectively reduces non-radiative recombination within the grains. However, the curve for the film with CQDs = 0.03 exhibits a much faster decay at the initial stage (࣎2 = 3.6 ns), indicating that the appeared pinholes act as the


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severe non-radiative recombination centers. In addition to the enhanced crystallinity, the passivation effect of BA could contribute to the increased carrier lifetime in the films.39 Therefore, we designed an experiment of using BA (0.01 mmol/L) instead of QDs in DE antisolvent to prepare CH3NH3PbI3 films. The film treated with BA exhibits reduced grain sizes and vague morphology (Fig. S5), and it shows slightly enhanced PL intensity and a longer bulk lifetime (࣎1 = 117.8 ns) compared to the reference film, but with a similar fast decay (࣎2 = 3.1 ns) as the film with CQDs = 0.03 (Fig. S6 and Table S1). It confirms that BA has passivation effect on CH3NH3PbI3 film, as reported in our former work,40 while the poor morphologies of the film are primarily due to the lack of effective nucleation centers which the QDs could act as. Furthermore, electron-only devices with the structure of ITO/PCBM/CH3NH3PbI3QDs/PCBM/Ag were fabricated to investigate the electron trap-state density in the CH3NH3PbI3 films (Fig. 4c). The trap-state density is calculated by using the equation of VTFL = entL2/(2εε0),4142

where VTFL is the trap-filled limited voltage, nt is the trap-state density, L is the thickness of

CH3NH3PbI3 layer, ε is the relative dielectric constant of CH3NH3PbI3 (ε = 28.8)43 and ε0 is the vacuum permittivity. The calculated trap-state density for the films with different CQDs is shown in Fig. 4d. All the perovskite films with QDs have lower trap-state densities compared to the reference film, in accordance with the improved crystallinity and increased bulk carrier lifetime, while the lowest trap-state density of 5.77×1015 cm-3 is obtained for CQDs = 0.01 (Table S2). The CH3NH3PbI3 film with optimized CQDs is further applied in perovskite LESC with the configuration of glass/FTO/NiO/CH3NH3PbI3-QDs/PCBM/Ag, and the cross-sectional SEM image of the multi-functional device is shown in Fig. 4e. The CH3NH3PbI3 film is sandwiched between an electron transport/injection layer (PCBM) and hole transport/injection layer (NiO), with Ag and FTO as the electrodes, respectively. The cross-sectional SEM view of perovskite


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devices shows that the reference CH3NH3PbI3 film possesses uniform and large grains, however, in the CH3NH3PbI3 film with the introduction of QDs, smaller grains appear in the upper layer and the grain sizes increase beneath the upper layer (Fig. S7). The evolution of grain morphology confirms that the QDs act as heterogeneous nucleation centers for CH3NH3PbI3 film growth. The corresponding energy band diagram of the device is illustrated in Figure 4f. The incorporation of QDs does not change the bandgap of CH3NH3PbI3 films, because the amount of Br added into the films is so small and even lower than the detection limit of XPS measurement. However, the work function (Ws) of perovskite films shifts from -3.93 to -4.12 eV and the conduction band (EC) shifts from -3.90 to -4.09 eV with the increased CQDs, as shown in Fig. S8. The current density-voltage (J-V) characteristics of the devices in dark condition were tested to analyze the carrier recombination loss, and the results are shown in Fig. 5a. For CQDs ≤ 0.01, the devices show rectifying characteristics with low reverse leakage currents. For CQDs > 0.01, the devices exhibit large leakage currents and even non-ideal diode behaviors (CQDs = 0.03), which originate from the increased recombination in the CH3NH3PbI3 films with pinholes.44 The EL spectra of the LESC devices with different CQDs at the bias voltage of 2 V are shown in Fig. 5b. The devices give narrow EL spectra with the peak at the wavelength of 778 nm, which are solely attributed to the band-edge emission of CH3NH3PbI3 without any detectable emissions originating from the carrier injecting layer (i.e., NiO or PCBM). Meanwhile, the EL intensities are enhanced in the CH3NH3PbI3 films with appropriate CQDs and the maximum value is obtained when CQDs = 0.01. Besides, the EL spectra of the devices with different CQDs show similar trends at higher bias voltages (3, 4 and 5 V), as shown in Fig. S9. Particularly, Fig. 5c shows the EL spectra of the device with CQDs = 0.01 at varying bias voltages. With the increased bias voltage, a continuously enhanced luminescence is observed until 5 V, then it begins to decrease at 6 V.


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Uniform red light is emitted from the device at the bias voltage of 5 V, as shown in the inset in Fig. 5c. The color coordinate of this sample is labeled as a black point in the commission international de L'Eclarage (CIE) chromaticity diagram (Fig. S10), showing high color saturation which is consistent with its relatively narrow emission. Figure 5d-f show the optoelectronic performances of the perovskite LEDs, i.e., CE, luminance and average external quantum efficiency (EQE), respectively. The corresponding characteristic parameters are summarized in Table 1. As CQDs increases (CQDs ≤ 0.02), the CE of the devices substantially increases at all the applied voltages. The CE has been improved by nearly an order of magnitude with a maximum value of 2.3 ×10-5 cd A-1 obtained when CQDs = 0.01 (Fig. 5d). The maximum luminance of 0.54 cd m-2 is achieved in the device with CQDs = 0.01 at the applied voltage of 5.5 V (Fig. 5e). Likewise, the average EQE with different CQDs shows that the maximum EQE is obtained in the device with CQDs = 0.01 (Fig. 5f). The enhanced luminance and EQE are attributed to the increased carrier radiative recombination caused by high crystallinity within small crystal grains of perovskite films. Though CE and EQE of the LESCs are low at present, the performances of LEDs should be further improved if the energetic barriers are minimized with optimized interfacial layers.6 On the other hand, the photovoltaic performances of the LESC devices with different CQDs are investigated. The illuminated J-V curves of the champion devices with different CQDs under AM 1.5G illumination (100 mW cm-2) are shown in Fig. 6a. The characteristic parameters derived from the J-V measurements are summarized in Table 1. The typical reference device gives the short-circuit current density (JSC) of 20.1 mA cm-2, the open-circuit voltage (VOC) of 1053 mV and the fill factor (FF) of 68%, yielding a PCE of 14.49%. All the JSC, VOC and PCE of the reference device are comparable to the reported ones using the similar structure45-47, except a


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smaller FF which will be further improved with doping45 or UV-O3 treatment47 on the pristine NiO. The photovoltaic performances are enhanced by introducing appropriate concentration of QDs (CQDs = 0.005, 0.01) into the perovskite absorption layer. Specially, the highest performance is achieved in the device with CQDs = 0.01, with the JSC of 21.1 mA cm-2, the VOC of 1080 mV, the FF of 75%, resulting in the PCE up to 17.10%. Note that the device we fabricated using p-i-n inverted structure for planar perovskite solar cell has negligible hysteresis (Fig. S11). The dense and highly crystalline CH3NH3PbI3 film with reduced carrier recombination effectively promotes the VOC and FF of the device. It is thus inferred that the crystallinity rather than the grain size of CH3NH3PbI3 film is a key factor for high SC performance. Besides, the shift of conduction band (EC) of CH3NH3PbI3 films with QDs benefits the extraction and transportation of electrons from the CH3NH3PbI3 layer to the PCBM layer with CQDs = 0.01 (Fig. S8d), leading to the increase of the JSC for the device. The EQE spectra and integrated JSC of the devices with different CQDs are shown in Fig. 6b. The EQE is improved in the range of 400-700 nm as the CQDs increases, and it achieves the maximum value when CQDs = 0.01, whereas it decreases upon further increasing the CQDs. The mismatch factors of integrated JSC with respect to the J-V characteristics (Fig. 6a) are 0%, 1.5%, 0.9%, 1.5% and 14.2%, for the devices with CQDs of 0.0, 0.005, 0.01, 0.02 and 0.03, respectively. The change of PCE or EQE on different CQDs is in agreement with that of the XRD patterns or TRPL spectra for the films, confirming that the improved SC performance is attributed to the enhanced crystallinity and reduced non-radiative recombination of grown films. Furthermore, the statistical distribution of the PCEs from 20 SCs for each group with different CQDs is shown in Fig. 6c. Upon increasing the CQDs, the PCE first increases with a peak at CQDs = 0.01 and then decreases, exhibiting good repeatability.


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To further investigate the stability of PSCs, the unpackaged devices were intentionally exposed to ambient air with relative humidity of 60 ± 5 % for 6 days, and the evolutions of normalized PCEs of the devices without and with QDs (CQDs = 0.01) are shown in Fig. 6d. The PCEs of the reference devices have reduced ~ 45% in average within 6 days, while those of the devices with CQDs = 0.01 have only lost ~ 25%. To explore the mechanism, the element contents of CH3NH3PbI3 films without and with QDs after exposure to ambient air for 2 days were measured by XPS (Fig. S12). The results show that the intensities of I, Pb and N in the reference film are lower than those in the film with QDs, however, the content of oxygen in the former is much larger than that of the latter. It is thus suggested that a highly crystalline film can effectively prevent the degradation of perovskite film originating from the oxygen or water in air.

CONCLUSIONS

The highly dispersed CH3NH3PbBr3 QDs with average diameter of 4.9 nm were successfully prepared using short alkyl ligand of BA and applied in anti-solvent as the nucleation centers for growth of CH3NH3PbI3 films. With optimal concentration of QDs (0.01 mmol/L), the crystallinity of CH3NH3PbI3 films is obviously improved, leading to higher carrier lifetime and lower trap-state density. Meanwhile, the corresponding perovskite LESC with p-i-n structure using the QDs induced CH3NH3PbI3 film exhibits improved EL and CE as LEDs, and enhanced photovoltaic characteristics and stability as SCs. To conclude, the results strongly suggest an effective strategy for the growth of high quality CH3NH3PbI3 film, which benefits the optoelectronic device applications in the near future. ASSOCIATED CONTENT Supporting Information


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Materials, equipment and methods are available in the supporting information. Figure S1-S11 and Table S1-S2 give more details on characterization of the CH3NH3PbBr3 QDs, CH3NH3PbI3 films, and the device. Photographs of the QDs, SEM images of the films, PLQY of the films, absorption spectra of the films, XPS analysis, PL and PL decay, EL intensity and CIE color, TRPL fitting results and calculated trap-state density are included. AUTHOR INFORMATION Corresponding Author *[email protected] (Can Cui), [email protected] (Xuegong Yu) Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. ‡ P. Wang and J. Xie contributed equally to this work. Notes The authors declare no competing financial interest. ACKNOWLEDGMENTS This work was supported by the National Natural Science Foundation of China (No. 61604131, 61422404 and 61704154), Natural Science Foundation of Zhejiang Province (No. LY17F040005), Central basic scientific research in colleges and universities operating expenses, Program for Innovative Research Team in University of Ministry of Education of China (IRT13R54), Visiting Scholar Foundation of State Key lab of Silicon Materials (No. SKL20161), Science Foundation of Zhejiang Sci-Tech University (No. 16062067-Y), National Undergraduate Training Program for Innovation (201710338011) and Xinmiao Undergraduate


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Student Talents Program of Zhejiang Province (2017R406037), 521 Talents Project of Zhejiang Sci-Tech University. REFERENCES (1)

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Figure 1 (a) TEM image of CH3NH3PbBr3 QDs dispersed in DE. (b) HRTEM image of a typical CH3NH3PbBr3 QD crystal. The inset is the corresponding selected area diffraction pattern. (c) XRD patterns of CH3NH3PbBr3 QDs and the bulk CH3NH3PbBr3 film. (d) UV-Vis absorption spectra and normalized PL spectra for the CH3NH3PbBr3 QDs and bulk CH3NH3PbBr3 film, respectively. (e) Schematic illustration of the solution-processed preparing CH3NH3PbI3 films using DE without (top) and with CH3NH3PbBr3 QDs (bottom) as the antisolvent. (f) Speculative mechanism of nucleation and growth of CH3NH3PbI3 films based on CH3NH3PbBr3 QDs.

Figure 2 (a)-(e) Top-view SEM images of CH3NH3PbI3 films deposited on FTO/NiO layers with different CQDs. Inset figures are the corresponding histograms of grain sizes. (f) The average grain sizes with varying CQDs obtained from Fig. 2a-e.


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Figure 3 (a) XRD patterns and (b) enlarged part of XRD patterns from 14.0° to 14.8° (left) and 28.4° to 29.2° (right) of CH3NH3PbI3 films with different CQDs grown on FTO/NiO layers. (c) The FWHM of XRD peaks for (110) and (220) crystalline planes of CH3NH3PbI3 films with different CQDs.


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Figure 4 (a) Steady PL and (b) normalized TRPL spectra of CH3NH3PbI3 films deposited on glass with different CQDs. (c) Dark I-V measurement of the electron-only devices. The inset shows the device with the structure of ITO/PCBM/CH3NH3PbI3-QDs/PCBM/Ag. (d) The measured trap-state densities for the films with different CQDs. (e) Cross-sectional SEM image of a typical LESC device with the structure of FTO/NiO/CH3NH3PbI3-QDs/PCBM/Ag. The thickness of CH3NH3PbI3 layer is ~ 420 nm. (f) The energy band structure and carrier transport behaviors of the perovskite LESC.


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Figure 5 (a) Dark J-V characteristics of the devices with different CQDs. (b) EL spectra of the devices with different CQDs at the applied bias voltage of 2 V. (c) EL spectra of the device (CQDs = 0.01) under varying bias voltages. The inset shows emission of uniform red light from the device at the bias voltage of 5 V. (d) CE, (e) luminance and (f) average EQE of 10 devices for each group with different CQDs applied with the bias voltage from 2 to 6 V.


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Figure 6 (a) Illuminated J-V curves of champion LESCs with different CQDs. (b) EQE and integrated JSC of corresponding devices with different CQDs. (c) The statistical distribution of PCEs for the devices (20 devices for each group) with different CQDs. (d) The evolution of PCE for unpackaged devices (10 devices for each group) without and with QDs (CQDs = 0.01) exposure to the air with 60 ± 5% humidity at 25 °C for 6 days.


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Table 1 Summary of perovskite LESC parameters derived from the champion devices with different CQDs.

Reference 0.005 0.01 0.02 0.03

LED Characteristics Max. CE (cd A-1) Max. EQE 3.6×10-6 0.020 1.8×10-5 0.054 2.3×10-5 0.084 1.7×10-5 0.052 2.3×10-6 0.007

VOC (mV) 1053 1067 1080 1045 598

Solar Cell Characteristics JSC (mA cm-2) FF (%) 20.1 68 20.6 74 21.1 75 19.9 65 10.6 39

PCE (%) 14.49 16.29 17.10 13.58 2.45


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Table of Contents CH3NH3PbBr3 QDs are synthesized using short alkyl ligand of BA, which act as nucleation centers and induce the growth of CH3NH3PbI3 films, leading to significantly improved crystallinity and reduced non-radiative recombination. The corresponding perovskite light emitting solar cells exhibit improved current efficiency and power conversion efficiency.

TOC


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CH3NH3PbBr3 Quantum Dots Induced Nucleation for High

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