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Letter

Amine-Based Passivating Materials for Enhanced Optical Properties and Performance of Organic-Inorganic Perovskite in Light-Emitting Diodes Seungjin Lee, Jong Hyun Park, Bo Ram Lee, Eui Dae Jung, Jae Choul Yu, Daniele Di Nuzzo, Richard H. Friend, and Myoung Hoon Song J. Phys. Chem. Lett., Just Accepted Manuscript • DOI: 10.1021/acs.jpclett.7b00372 • Publication Date (Web): 05 Apr 2017 Downloaded from http://pubs.acs.org on April 6, 2017

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The Journal of Physical Chemistry Letters

Amine-Based Passivating Materials for Enhanced Optical Properties and Performance of OrganicInorganic Perovskite in Light-Emitting Diodes Seungjin Lee1, Jong Hyun Park1, Bo Ram Lee2, Eui Dae Jung1, Jae Choul Yu1, Daniele Di Nuzzo2, Richard H. Friend2, Myoung Hoon Song1* 1

School of Materials Science Engineering and Low Dimensional Carbon Center and KISTUNIST Ulsan Center for Convergent Materials, Ulsan National Institute of Science and Technology (UNIST), UNIST-gil 50, Ulsan, 44919, Republic of Korea 2

Cavendish Laboratory, JJ Thomson Avenue, Cambridge, CB3 0HE, United Kingdom *To whom correspondence should be addressed. E-mail: [email protected]

AUTHOR INFORMATION Corresponding Author Tel: +82-52-217-2316 Fax: +82-52-217-2309 *E-mail: [email protected]

ACS Paragon Plus Environment

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ABSTRACT

The use of hybrid organic-inorganic perovskites in optoelectronic applications are attracting an interest because of their outstanding characteristics, which enable a remarkable enhancement of device efficiency. However, solution-processed perovskite crystals unavoidably contain defect sites that cause hysteresis in perovskite solar cells (PeSCs) and blinking in perovskite lightemitting diodes (PeLEDs). Here, we report significant beneficial effects using a new treatment based on amine-based passivating materials (APMs) to passivate the defect sites of methyl ammonium lead tribromide (MAPbBr3) through coordinate bonding between the nitrogen atoms and under-coordinated lead ions. This treatment greatly enhanced PeLEDs efficiency, with an external quantum efficiency (EQE) of 6.2 %, enhanced photoluminescence (PL), a lower threshold for amplified spontaneous emission (ASE), a longer PL lifetime and enhanced device stability. Using confocal microscopy, we observed the cessation of PL blinking in perovskite films treated with ethylenediamine (EDA) due to the passivation of the defect sites in the MAPbBr3. TOC GRAPHICS


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Perovskite optoelectronic devices (PeOEDs) such as solar cells (SCs), light-emitting diodes (LEDs) and photodetectors have been extensively investigated because of the excellent optical and electrical properties of hybrid organic-inorganic perovskite materials.1-8 Recently, remarkable enhancements have been realized in the power conversion efficiency (PCE) of perovskite solar cells (PeSCs), reaching values of up to 22.1 %,6 and in the external quantum efficiency (EQE) of perovskite LEDs (PeLEDs), reaching values of up to 11.7 % for nearinfrared emission7 and 9.3 % for green emission.8 High-quality perovskite films are required to realize highly efficient and stable perovskitebased optoelectronic devices.9-12 However, solution-processed perovskite films inevitably contain defect sites such as voids, pinholes, grain boundaries and under-coordinated ions, creating a large number of undesired electronic trap sites.13-16 Many researchers have demonstrated that defect-induced electronic trap sites are a major cause of undesired phenomena in PeOEDs, such as hysteresis behaviour in PeSCs16-19 and blinking behaviour of PeLEDs,20-26 which limit their applicability and commercialization. In this regard, considerable efforts have been made to reduce the presence of defects in perovskite films through the control of their crystallization or through post-treatment to suppress these undesired phenomena.10,16,19,27-29 Jeon et al. have demonstrated the preparation of uniform and dense perovskite layers using a solventengineering technique, which significantly reduces hysteresis behaviour and improves the PCEs of PeSCs.10 Noel et al. have reported that organic Lewis bases can be used to passivate the defect sites on a perovskite surface via coordinate bonding, resulting in reduced hysteresis behaviour and improved device efficiency in PeSCs.16 Tachikawa et al. have demonstrated that the PL blinking behaviour of perovskite nanocrystals that is induced by trap sites can be reduced by passivating the traps via treatment with Lewis bases.25


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Moreover, the poor stability of current PeOEDs remains a serious issue that must be resolved before they can be broadly applied and commercialized. In essence, the instability of hybrid organic-inorganic perovskite originates from its soft nature, which is related to the low interaction energy between metal cations and halide anions, allowing ionic defects to be easily generated.30-33 Ionic defects in perovskite facilitate the decomposition of the material due to extrinsic factors such as moisture, light and heat as well as ion migration to adjacent layers or metal electrodes, leading to intrinsic instabilities such as interface interplay and the corrosion of metal electrodes.34-38 Consequently, many researchers have struggled to improve the stability of PeOEDs for commercialization.29,38-43 F. Wang et al. have demonstrated that the molecular passivation of perovskite films improves device stability with regard to the moisture resistance of the perovskites, and it simultaneously reduces hysteresis behaviour and improves device efficiency in PeSCs.29 Smith et al. have reported that device stability can be improved by exchanging methylammonium (CH3NH3) groups with long and hydrophobic C6H5(CH2)2NH3 to enhance the resistance to humidity.43 Back et al. have suppressed the corrosion of metal electrodes by inhibiting ion migration from the perovskite layer using an amine-mediated titanium oxide as a chemical inhibitor layer to improve device stability.38 Here, we present an effective method of passivating the defect sites in methyl ammonium lead tribromide (MAPbBr3) by introducing amine-based passivating materials (APMs) to enable the fabrication of highly efficient PeLEDs with excellent stability. To investigate the effect of the molecular size of the APMs on defect passivation in the perovskite crystals, we used two different APMs, namely, branched polyethyleneimine (PEI) and ethylenediamine (EDA). We observed that perovskite films with APMs showed enhanced photoluminescence (PL) intensities, long PL lifetimes and reduced PL blinking because of the significant suppression of non-


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radiative recombination. Moreover, the APMs treatment significantly suppressed electrode corrosion by inhibiting ion migration from the perovskite layer to the metal electrode, thereby improving the device stability. Finally, by means of EDA treatment, we obtained highly efficient and stable PeLEDs.

Figure 1. Device schematic, cross-sectional image and energy level diagram of a PeLEDs as well as the chemical structures of the APMs. (a) Schematic illustration of the device structure for the PeLEDs with APMs. (b) Cross-sectional scanning electron micrograph of a PeLEDs. (c) Energy levels of the components of a PeLEDs. (d) Chemical structures of the APMs: (i) EDA and (ii) PEI. Figure 1a,b presents a schematic illustration of the device design and a cross-sectional scanning electron microscopy (SEM) image of a PeLEDs. An energy level diagram of a PeLEDs


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and the chemical structures of the two APMs showing their different molecular sizes are depicted in Figure 1c,d. The APMs contain amine groups, including lone pairs, which can passivate under-coordinated lead ions by donating electrons. Schematic illustrations of MAPbBr3 with and without APMs are helpful for understanding the passivation mechanism (Figure S1). There are many under-coordinated lead ions in MAPbBr3 because MAPbBr3 is crystallized through a solution process, and bromide and methylammonium ions can be lost from the MAPbBr3 during post-deposition thermal annealing process. Bromide ion vacancies result in positively charged under-coordinated lead ions; these unintentional defects can be passivated through coordinate bonding with the nitrogen atoms of amine groups, resulting in charge neutralization and a consequent reduction in the number of electronic trap sites.


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Figure 2. Negative ion TOF-SIMS spectra, depth profiles, and schematic illustrations of the depth penetration of APMs deposited on MAPbBr3. (a) TOF-SIMS spectra of CN- ions in MAPbBr3 with and without APMs. (b) Depth profiles of CN- ions in MAPbBr3 with and without APMs. (c) Schematic illustration of the depthwise penetration of PEI into MAPbBr3. (d) Schematic illustration of the depthwise penetration of EDA into MAPbBr3. To investigate the number of existing amine groups on the surface of the MAPbBr3 passivated with APMs and the penetration depth of the APMs into the MAPbBr3, time-of-flight secondary ion mass spectrometry (TOF-SIMS) was performed (Figure 2a,b). The number of existing amine groups on the surface of PEI-passivated MAPbBr3 was larger than that on EDA-passivated MAPbBr3. However, the effective passivation of the amine groups of EDA on the MAPbBr3 surface and the passivation depth of EDA in the MAPbBr3 film were superior to those of PEI because the relatively small molecules could more effectively passivate the top surface of the MAPbBr3 and penetrate more deeply; schematic illustrations of the penetration of the two APMs into MAPbBr3 are shown in Figure 2c,d.


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Figure 3. Time-resolved, steady-state PL spectra and ASE measurement of MAPbBr3 with and without APMs. (a) Time-resolved PL spectra for MAPbBr3 with and without APMs. (b) Steady-


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state PL spectra of MAPbBr3 with and without APMs. (c) PL intensities and FWHM of MAPbBr3 without APMs at various excitation intensities. (d) PL spectra of MAPbBr3 without APMs at various excitation intensities illustrating the transition from spontaneous emission to ASE. (e) PL intensities and FWHM of MAPbBr3 with EDA at various excitation intensities. (f) PL spectra of MAPbBr3 with APMs under at various excitation intensities illustrating the transition from spontaneous emission to ASE. Table 1. Summarized PL lifetime of MAPbBr3 with and without APMs.

τavr [ns]

χ2

Glass / PEDOT:PSS / MAPbBr3

13.8

1.112

Glass / PEDOT:PSS / MAPbBr3 / PEI

26.7

1.078

Glass / PEDOT:PSS / MAPbBr3 / EDA

62.5

1.001

Film configuration

Figure 3 show the improvement of optical properties of MAPbBr3 with APMs, which are strong evidence for the defect passivation of MAPbBr3. MAPbBr3 materials passivated with APMs showed long PL lifetimes and high steady-state PL intensities compared with MAPbBr3 without APMs (Figure 3a,b and Table 1), indicating the suppression of non-radiative recombination through a reduction in the number of electronic trap sites. In particular, EDApassivated MAPbBr3 showed a longer PL lifetime and a higher PL intensity compared with PEIpassivated MAPbBr3, which suggests that EDA passivates the defect sites in MAPbBr3 more effectively than PEI does. We also measured the amplified spontaneous emission (ASE) of the MAPbBr3 materials with and without EDA to confirm the improvement of the optical properties of the MAPbBr3 upon treatment with EDA. The MAPbBr3 with EDA showed a lower threshold for ASE (10.6 µJ cm-2) and a smaller full width at half maximum (FWHM) of its ASE (3.60 nm)


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than those (28.6 µJ cm-2, 4.63 nm) of the MAPbBr3 without EDA; these findings indicate a reduction in optical loss caused by electronic trap sites with EDA treatment. The normalized PL spectra of the MAPbBr3 with APMs are slightly blue-shifted compared with that of MAPbBr3 without APMs (Figure S2), and they show sharp band-edge emission because of the passivation of shallow trapping levels.44 However, these slight peak shifts cannot be attributed to any structural change in the MAPbBr3 with APMs because the X-ray diffraction (XRD) patterns of MAPbBr3 with and without APMs are nearly identical (Figure S3).

Figure 4. Confocal PL images and variations in PL intensities observed over time for MAPbBr3 materials with and without APMs. (a) Confocal PL image of MAPbBr3 without APMs. (b) Confocal PL image of MAPbBr3 with PEI. (c) Confocal PL image of MAPbBr3 with EDA. (d) Variations in the PL intensities of MAPbBr3 without APMs. (e) Variations in the PL intensities of MAPbBr3 with PEI. f) Variations in the PL intensities of MAPbBr3 with EDA.


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Many scientists have reported a blinking phenomenon in hybrid organic-inorganic perovskites, which limits the applicability of PeOEDs. Although several mechanisms to explain blinking phenomenon have been proposed, but the details are not yet fully understood. Several researchers have suggested that the blinking phenomenon can be attributed to non-radiative Auger recombination between photogenerated electron-hole pairs and charges captured by electronic trap sites (Figure S4a).20,21,45 Other researchers have reported a mechanism such that the blinking phenomenon is attributed to the photo-induced activation and deactivation of electronic trap sites associated with defect sites acting as PL quenchers (Figure S4b).23,25,46 Both mechanisms suggest that the blinking phenomenon is closely related to electronic trap sites; therefore, the reduction of electronic trap sites through defect passivation is crucial for blinking suppression. To investigate the passivation effect of the studied APMs on the PL blinking of MAPbBr3, we acquired confocal PL images and PL intensity trajectories of the MAPbBr3 materials with and without APMs over time (Figure 4 and Movie S1). The confocal PL images of the MAPbBr3 without APMs revealed large dark regions dominated by non-radiative recombination, whereas the confocal PL images of the MAPbBr3 materials passivated with APMs showed a reduction of the dark regions and an increase in brightness, with almost full coverage in the case of EDA. (Figure 4a-c). Moreover, the MAPbBr3 without APMs showed a noticeable deviation in its PL intensity as a function of time, whereas the MAPbBr3 materials passivated with APMs showed reduced (PEI) or no (EDA) deviation of their PL intensities (Figure 4d-f). In Movie S1, vigorous PL blinking can be seen in the MAPbBr3 without APMs as a result of non-radiative recombination in the presence of electronic trap sites, whereas no PL blinking can be clearly observed in the EDA-passivated MAPbBr3, which unequivocally confirms the effective passivation of electronic trap sites due to EDA treatment.


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Recently, deQuilettes et al. and Mosconi et al. have reported that perovskites with higher trap densities exhibit more photo-induced PL enhancement because of the reduction in the trap density caused by halide migration.47,48 To observe the electronic trap densities of the MAPbBr3 materials with and without APMs, we measured the variations in their PL intensities over time both under illumination and in the dark (Figure S5a,b). The MAPbBr3 without APMs showed a significant increase in PL intensity compared with those of the MAPbBr3 materials with APMs, indicating that the MAPbBr3 materials passivated with APMs had a lower electronic trap density than did the MAPbBr3 without APMs. Moreover, the higher initial PL intensity of the MAPbBr3 passivated with APMs compared with that of the photo-stabilized MAPbBr3 without APMs demonstrates the effective passivation capability of the APMs. When the illumination was removed, the photo-enhanced PL intensity of the MAPbBr3 without APMs rapidly relaxed within hours (Figure S5b), which indicates that the photo-induced enhancement is reversible. By contrast, the MAPbBr3 materials passivated with APMs showed almost no change in their PL intensities over time, as confirmed by the confocal microscopy images (Figure S5c and Figure S6); these findings indicate that the defect passivation induced by the APMs was maintained over a long duration. We believe that the sustainability of the APMs effects can be attributed to a strong interaction between the nitrogen atoms and the under-coordinated lead ions, such as coordinate bonding. Therefore, this method of passivating the defect sites in MAPbBr3 using EDA is very simple, effective and long lasting.


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Figure 5. Device performance of PeLEDs and SEM top-surface images of MAPbBr3 with and without APMs. (a) Current density versus voltage (J-V) characteristics. (b) Luminance versus voltage (L-V) characteristics. (c) Luminous efficiency versus voltage (LE-V) characteristics. (d) EQE versus luminance (EQE-V) characteristics. (e) SEM top-surface image of MAPbBr3 without APMs. (f) SEM top-surface image of MAPbBr3 with PEI. (g) SEM top-surface image of MAPbBr3 with EDA. Table 2. Summarized device performance of PeLEDs with and without APMs. Turn-on L max [cd/m2]

LE max [cd/A]

EQE max [%]

@ bias

@ bias

@ bias

Device configuration (PeLEDs)

voltage [V] @ 0.1 cd/m2

ITO / PEDOT:PSS / MAPbBr3 / SPW-111 / LiF / Ag

7,080 (4.8 V)

0.559 (4.8 V)

0.12 (4.8 V)

2.2 V

ITO / PEDOT:PSS / MAPbBr3 / PEI / SPW-111 / LiF / Ag

10,700 (4.8 V)

2.72 (4.8 V)

0.58 (4.8 V)

2.4 V

ITO / PEDOT:PSS / MAPbBr3 / EDA / SPW-111 / LiF / Ag

22,800 (3.4 V)

28.9 (3.0 V)

6.19 (3.0 V)

2.4 V


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Figure 5 and Table 2 show the current density and luminance versus voltage (J-V-L) characteristics and the device efficiencies of PeLEDs with and without optimized concentrations of APMs. The PeLEDs were optimized by testing different APMs concentrations. The optimized weight concentration for PEI was found to be 0.3 wt. %. A PEI layer with such a high concentration effectively passivates the defect sites in MAPbBr3, whereas a thicker insulating layer reduces electron injection from the electron transport layer into the emissive layer, leading to unbalanced charge transport. Moreover, the J-V-L characteristics and device efficiencies of PeLEDs with various concentrations of EDA are shown in Figure S7 and Table S1. A highconcentration EDA layer effectively passivates the defect sites in MAPbBr3, but a high concentration of EDA also results in partial melting of the MAPbBr3 crystal (Figure S8), which causes the crystal quality to deteriorate. Thus, we optimized the weight concentration of EDA to be 0.10 wt. %. At low voltages, the current densities of the PeLEDs with EDA were substantially reduced compared with that of the control device, indicating that EDA caused the leakage current to decrease as a result of the defect passivation of the MAPbBr3, as shown in Figure 5a. The luminance and device efficiencies of the PeLEDs with EDA were much higher than those of the control device, indicating effective defect passivation. The morphology of the MAPbBr3 passivated with APMs with the optimized APMs concentrations did not change compared with that of the MAPbBr3 without APMs, as shown in Figure 5d-f. The optimized EDA-passivated device achieved a luminance of 22,800 cd m-2 and an EQE of 6.2 %. The histogram of EQE for 15 devices with EDA are shown in Figure S9.


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Figure 6. Corrosion of Ag on MAPbBr3 and operational stability of PeLEDs with and without APMs. (a) Photographs of Ag on MAPbBr3 materials with and without APMs under ambient conditions over time. (b) XPS spectra of Ag on MAPbBr3 materials with and without APMs after 30 days. (c) XRD patterns of Ag on MAPbBr3 materials with and without APMs after 30 days. (d) Normalized luminances of encapsulated PeLEDs with and without APMs under ambient conditions as functions of operation time. The ionic defects that form in perovskite because of the low interaction energy lead to decomposition induced by moisture, light and heat as well as ion migration to adjacent layers or metal electrodes, resulting in poor stability of the material. In particular, the intrinsic instability


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of PeOEDs that is associated with ion migration from the perovskite layer cannot be eliminated by means of device encapsulation, and thus additional strategies are required for improving the intrinsic device stability, such as passivation of the perovskite to inhibit ion migration and modification of the components to enhance the interaction energy. We investigated the effect of the APMs on the corrosion resistance of Ag on MAPbBr3 and the device stability of PeLEDs under continuous operation. Figure 6a presents photographs of Ag on MAPbBr3 materials with and without APMs under ambient conditions over time without any encapsulation. Changes in the electrode colour indicate the corrosion of Ag. The Ag on MAPbBr3 without APMs showed a colour change within 1 day (24 h) and additional colour changes after 15 and 30 days, indicating that Ag on MAPbBr3 without APMs can be easily corroded. By contrast, the Ag on MAPbBr3 with APMs exhibited excellent stability, showing almost no colour change even after 30 days (720 h); these findings demonstrate that the APMs treatments effectively suppressed the corrosion of the electrodes. To investigate the specific origin of the corrosion of Ag on MAPbBr3, we performed X-ray photoelectron spectroscopy (XPS) and XRD analyses of the Ag on the MAPbBr3 materials with and without APMs after 30 days (Figure 6b,c). The Ag 3d spectra of pristine Ag and of the Ag on the MAPbBr3 materials with and without APMs show two peaks at 374.1 eV (3d3/2) and 368.1 eV (3d5/2) (Figure 6b). These two peaks can each be further divided into two additional peaks at 374.1 and 373.3 eV (3d3/2) and at 368.1 and 367.6 eV (3d5/2), respectively, where the 374.1 and 368.1 eV peaks are assigned to metal Ag0 and the 373.3 and 367.6 eV peaks are assigned to the Ag+ in AgBr.49 The spectra of the Ag on the MAPbBr3 materials with APMs are nearly the same as that of the pristine Ag, whereas the spectrum of the Ag on the MAPbBr3 without APMs shows clear contributions from the peaks at 373.3 and 367.6 eV corresponding to the Ag+ in AgBr. These results are consistent with the


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XRD analysis (Figure 6c). The peaks at 38.1° and 31.0° correspond to the (111) crystal plane of Ag (JCPDS no. 04-0783) and the (200) crystal plane of AgBr (JCPDS no. 06-0438), respectively.50 The XRD pattern of the Ag on the MAPbBr3 without APMs shows a clear signal from the (200) peak of AgBr in addition to the (111) peak of Ag, whereas the XRD patterns of the Ag on the MAPbBr3 materials with APMs show only the (111) peak of Ag, with no evidence of the (200) peak of AgBr. The XPS and XRD analyses indicate that the APMs treatments completely prevent electrode corrosion by inhibiting ion migration from the perovskite layer to the metal electrode because the APMs treatments effectively passivate the surface of the MAPbBr3 and block the ion migration paths, thereby improving the device stability. The luminances of encapsulated PeLEDs with and without APMs were measured at 20 mA cm-2 under ambient air conditions as functions of operation time (Figure 6d). The operational stability of the PeLEDs with APMs was found to be significantly better than that of the control device. The PeLEDs without APMs exhibited a sharp decrease to less than 10 % of the initial luminance after only 2,500 sec. By contrast, the PeLEDs with APMs retained over 70 % of their initial luminance after more than 14,000 sec. In summary, we demonstrated the fabrication of highly efficient and stable PeLEDs using a simple and effective method employing APMs. APMs can passivate defect sites in MAPbBr3 through coordinate bonding between the nitrogen atoms and the under-coordinated lead ions, leading to a reduction in the electronic trap density. Therefore, APMs treatment results in enhanced PL intensity with a lower threshold for ASE, a long PL lifetime and reduced PL blinking, ultimately leading to enhanced device performance. In particular, EDA can effectively passivate defect sites in MAPbBr3 on both the surface and the inside of the MAPbBr3 crystal. In contrast to the reversible photo-induced enhancement of PL, the PL enhancement achieved


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through the introduction of the APMs was found to be maintained for a long time because of a strong interaction between the nitrogen atoms and the under-coordinated lead ions, such as coordinate bonding. This finding indicates that APMs treatment is an irreversible and effective method of passivating defect sites in MAPbBr3. Moreover, APMs treatment completely prevents electrode corrosion by inhibiting ion migration from the perovskite layer to the metal electrode, leading to improved device stability. Consequently, this method allowed us to fabricate optimized PeLEDs with EDA which exhibited a remarkable maximum luminance of 22,800 cd m-2 and EQE of 6.2 %, with excellent device stability. We believe that this enhancement in the device efficiency and long-term stability of PeLEDs with the simultaneous suppression of PL blinking can contribute to the commercialization of PeLEDs.

Materials. The poly(3,4-ethylenedioxythiophene):poly(styrenesulfonate) (PEDOT:PSS, AI 4083, Clevios) and the SPW-111 polymer (Merck Co.) were used without any purification. The lead bromide (PbBr2, 99.999 %, Alfa Aesar), methylammonium bromide (MABr, Dyesol), ethylenediamine (EDA, 99.5 %, Aldrich) and branched polyethyleneimine (PEI, average Mw ~25,000, Aldrich) were used without further purification. Device fabrication. The PEDOT:PSS was spin coated at 5,000 rpm for 40 s onto a patterned ITO glass slide and then annealed at 130 °C for 10 min. To prepare the perovskite precursor solution, PbBr2 and MABr at a molar ratio of 1:1.05 were dissolved in N,N-dimethylformamide (37.8 wt. %), and hydrobromic acid (48 wt. % in water, 99.99%) was added at 6 vol. % as a stabilizer. The precursor solution was spin coated onto the PEDOT:PSS via a one-step process at 3,000 rpm for 45 s and then annealed at 80 °C for 30 min. The amine-based interfacial materials EDA and PEI were dissolved in chloroform at 0.10 vol. % and 3.0 wt. %, respectively, and then spin


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coated onto the perovskite layer at 2,000 and 5,000 rpm, respectively, for 45 s. The SPW-111 was dissolved in chlorobenzene (5.0 mg ml-1) and then spin coated at 2,000 rpm for 45 s. Finally, for electrode formation, lithium fluoride (1 nm) and silver (80 nm) were successively deposited using a thermal evaporation system. Device characterization. The J-V-L characteristics and efficiencies of encapsulated PeLEDs were measured using a computer-controlled Keithley 2,400 Source Meter and a Konica Minolta spectroradiometer (CS-2000, Minolta) under ambient air conditions.

ACKNOWLEDGMENT This work is financially supported by the KIST-UNIST partnership program (1.160097.01/2.160482.01)

and

Mid-Career

Researcher

Program

(2015R1A2A2A01003263). This work was supported by the Human Resource Training Program for Regional Innovation and Creativity through the Ministry of Education and National Research Foundation of Korea (NRF-2014H1C1A1073051).

ASSOCIATED CONTENT Supporting Information. Supporting experimental methods, figures and table (PDF) Movie S1 (AVI) AUTHOR INFORMATION Notes


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The authors declare no competing financial interest. REFERENCES (1) Lee, M. M.; Teuscher, J.; Miyasaka, T.; Murakami, T. N.; Snaith, H. J. Efficient Hybrid Solar Cells Based on Meso-Superstructured Organometal Halide Perovskites. Science 2012, 338, 643-647. (2) Burschka, J.; Pellet, N.; Moon, S. J.; Humphry-Baker, R.; Gao, P.; Nazeeruddin, M. K.; Gratzel, M. Sequential Deposition as a Route to High-Performance Perovskite-Sensitized Solar Cells. Nature 2013, 499, 316-319. (3) Liu, M. Z.; Johnston, M. B.; Snaith, H. J. Efficient Planar Heterojunction Perovskite Solar Cells by Vapor Deposition. Nature 2013, 501, 395-398. (4) Stranks, S. D.; Eperon, G. E.; Grancini, G.; Menelaou, C.; Alcocer, M. J. P.; Leijtens, T.; Herz, L. M.; Petrozza, A.; Snaith, H. J. Electron-Hole Diffusion Lengths Exceeding 1 Micrometer in an Organometal Trihalide Perovskite Absorber. Science 2013, 342, 341-344. (5) Yuan, M. J.; Quan, L. N.; Comin, R.; Walters, G.; Sabatini, R.; Voznyy, O.; Hoogland, S.; Zhao, Y. B.; Beauregard, E. M.; Kanjanaboos, P. et al. Perovskite Energy Funnels for Efficient Light-Emitting Diodes. Nat. Nanotechnol. 2016, 11, 872-877. (6) Zheng, H. M.; Rivest, J. B.; Miller, T. A.; Sadtler, B.; Lindenberg, A.; Toney, M. F.; Wang, L. W.; Kisielowski, C.; Alivisatos, A. P. Observation of Transient Structural-Transformation Dynamics in a Cu2S Nanorod. Science 2011, 333, 206-209.


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