skite CH3NH3PbBr3 Thin Films for High-Performance Green Light

Subscriber access provided by READING UNIV

Article

Vapor-Assisted Solution Approach for High-Quality Perovskite CH3NH3PbBr3 Thin Films for High-Performance Green Light-Emitting Diodes Applications Huifang Ji, Zhifeng Shi, Xuguang Sun, Ying Li, Sen Li, Lingzhi Lei, Di Wu, Tingting Xu, Xinjian Li, and Guotong Du ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b13260 • Publication Date (Web): 15 Nov 2017 Downloaded from http://pubs.acs.org on November 17, 2017

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a free service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are accessible to all readers and citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

ACS Applied Materials & Interfaces is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 14

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

ACS Applied Materials & Interfaces

Vapor-Assisted Solution Approach for High-Quality Perovskite CH3NH3PbBr3 Thin Films for High-Performance Green Light-Emitting Diodes Applications Huifang Ji,1 Zhifeng Shi,*1 Xuguang Sun,1 Ying Li,1 Sen Li,1 Lingzhi Lei,1 Di Wu,1 Tingting Xu,1 Xinjian Li*1 and Guotong Du2 1

Key Laboratory of Materials Physics of Ministry of Education, Department of Physics and Engineering, Zhengzhou University, Daxue Road 75, Zhengzhou 450052, China. 2 State Key Laboratory on Integrated Optoelectronics, College of Electronic Science and Engineering, Jilin University, Qianjin Street 2699, Changchun 130012, China. ABSTRACT: Vapor-assisted solution method was developed to prepare the high-quality organic-inorganic halide perovskite CH3NH3PbBr3 (MAPbBr3) thin films. We detailedly investigated the effect of evaporation time and temperature of MABr powder on the microstructure, crystallinity and optical characterizations of MAPbBr3 thin films, and a controllable morphology evolution with varying surface coverage was observed. Temperature-dependent and time-resolved photoluminescence (PL) measurements were carried out to investigate the optical transition mechanisms and carrier recombination dynamics of MAPbBr3 thin films. Our results revealed that no structural phase transition occurred within the heating process (10 K–300 K). In addition to the exciton related emission, a trapped chargecarrier emission appeared at a critical temperature of 140 K. The corresponding temperature sensitivity coefficient of bandgap, exciton binding energy, and optical phonon energy of the MAPbBr3 thin films were extracted from the experimental data. Further, planar perovskite light-emitting diodes (PeLEDs) based on Al/LiF/TPBi/MAPbBr3/NiO/ITO structure were fabricated, and a high-purity green emission at~532 nm with a low linewidth (25 nm) was achieved. The devices demonstrated a remarkable performances with high luminance (6530 cd/m2), current efficiency (8.16 cd/A), external quantum efficiency (EQE, 4.36%), and power efficiency (4.49 lm/W). This research will provide valuable information for the preparation of high-quality perovskite thin films, facilitating their future applications in novel high-performance PeLEDs. KEYWORDS: perovskite, CH3NH3PbBr3, vapor-assisted solution approach, surface coverage, light-emitting diodes

INTRODUCTION Organic-inorganic hybrid perovskite materials (CH3NH3PbX3, MAPbX3, X=I, Br, Cl) have increasingly attracted considerable attention because of their exceptional optoelectronic properties.1−4 This upsurge has not faded away because solar cells employing a perovskite thin layer have achieved an excellent power conversion efficiency above 20% in short time.5 More recently, MAPbX3 perovskites have been shown to have great application potentials in luminescent fields, such as LEDs and lasing devices.6−8 The appealing properties of such MAPbX3 perovskites that enable advances in LEDs and lasing devices are high emission efficiency, wavelength-tunable emissions, ambipolar charge transport, and high emission color purity. Compared to the intensive efforts on photovoltaic applications, the investigation of perovskites as light emitters in solid-state light source applications is relatively less, not to mentioning significant breakthrough that could propel PeLEDs into the commercialization stage although a high EQE of 8.53% has been achieved based on MAPbBr3 perovskites.9 This issue is primarily due to the difficulty in achieving a high-quality perovskite emissive layer with a complete surface coverage and a high optical property, and also the absence of a reasonable device structure design. Therefore, improvement in the device performance of the emerging PeLEDs should start with the controllable preparation of high-quality perovskite emissive layer because its un-

desired discontinuity would inevitably induce shunt paths for the current transport, further degrading the device performance. It is known to all that the morphology of perovskite is very sensitive to its synthesis method, composition and structure details,6,9,10 therefore, optimizing preparation parameters to promote the quality of the perovskite thin films still requires systematic investigation in the progression toward highefficient PeLEDs.11,12 Moreover, compared to the intensive efforts in optoelectronic device applications, the inherent optical property of halide perovskite materials has been less studied. Given the rapid development in perovskite-based optoelectronic devices, exploring and deeply understanding the inherent optical properties of halide perovskites is imperative and highly desired. Current research concentrating on the optical properties of perovskites was always performed on MAPbI3, while there is little research on the optical behaviors of MAPbBr3 thin films, not to mentioning their optical transition mechanisms and carrier recombination dynamics performed by temperature-dependent and time-resolved PL measurements. From the device point of view, the spectroscopic studies at low temperatures and transition state are highly desired because they can provide us valuable insights into the fundamental photophysics of perovskites, which is favorable for the subsequent applications of MAPbBr3 thin films in LEDs. In this work, vapor-assisted solution method was developed for the preparation of high-quality MAPbBr3 thin films. We have detailedly investigated the effect of the evaporation time

ACS Paragon Plus Environment


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

and temperature of MABr powder on the microstructure, crystallinity and optical characterizations of MAPbBr3 thin films, and a controllable morphology evolution was observed. The optical transition mechanisms and carrier recombination dynamics of MAPbBr3 thin films were also detailedly studied by using the temperature-dependent and time-resolved PL measurements. Moreover, we estimated the temperature sensitivity

Page 2 of 14

coefficient of bandgap, exciton binding energy (EB), and optical phonon energy (ħωop) of MAPbBr3 thin films based on the obtained data. Further, a planar PeLED based on Al/LiF/TPBi/MAPbBr3/NiO/ITO structures were fabricated and characterized, and the device performance was remarkable in terms of its high luminance (6530 cd/m2), current efficiency (8.16 cd/A), EQE (4.36%), and power efficiency (4.49 lm/W).

Table 1. Preparation parameters of MAPbBr3 thin films. Sample No.

Evaporation time (min@MABr)

S1

5

S2

10

S3

15

S4

20

S5

25

S6

30

M1 M2 M3 M4

Evaporation Temp. (°C@MABr)

Pressure (Torr)

PbBr2 source (mol/L)

MABr source (mol)

300

1.0

0.02

135

120 25

150 165 180

EXPERIMENTAL SECTION Growth of MAPbBr3 thin films. In the experiment, MAPbBr3 thin films were prepared by using a vapor-assisted solution method, and commercially available ITO (200 nm, square resistance: 8 Ω/□) glasses were used as the substrates, which were ultrasonic cleaned sequentially in the acetone, ethyl alcohol, and deionized water three times. The detailed processing procedures of MAPbBr3 thin films were presented in Figure 1. Firstly, a PbBr2 solution dissolved in N,Ndimethylformamide with the concentration of 1 mol/L was spin-coated onto the substrates (3000 rpm, 30 s), and then the samples were annealed at a temperature of 100 °C (10 min) to evaporate the solvent in a glove box. Followed that, the resulting PbBr2/ITO template was transferred into a tubular furnace to evaporate the MABr layer. Note that the MABr powder is evenly laid on the bottom of the ceramic boat, with the PbBr2 layer on the front of the substrate facing the MABr powder, and the distance between the substrate and the MABr powder is about 2 mm. Before the evaporation process, the pressure of the furnace was evacuated to a low value (< 10−3 Pa) using a vacuum pump. During the evaporation process, the growth parameters (evaporation time and temperature of MABr) were systematically varied to evaluate their effects on the microstructure, crystallinity and optical characterizations of MAPbBr3 products. Consequently, two groups of MAPbBr3 products were prepared, and the corresponding preparation parameters were summarized in Table 1. Group I: MAPbBr3 products were prepared at different evaporation time (5, 10, 15, 20, 25, and 30 min) of MABr powder, and the resulting samples were named as S1−S6, respectively. Note that the evaporation temperature of MABr was set at 135 °C during the evaporation process with the reaction pressure of about 300 Torr for the tubular furnace was maintained at 300 Torr. Group II: based on the experimental results of S5 (group I), further experiments were performed to determine the effects

of evaporation temperature of MABr on the physical properties of MAPbBr3 samples. The evaporation experiment was carried out over an evaporation temperature range of 120−180 °C at a typical evaporation time of 25 min, and other growth parameters remained unchanged compared with those in group I. Consequently, four MAPbBr3 samples denoted as M1, M2, M3, and M4, were prepared. Measurement and Characterization. The morphologies of the MAPbBr3 samples were analyzed by using scanning electron microscope (SEM; JSM-7500F). The crystallinity of the produced MAPbBr3 samples were characterized by using a Panalytical X’ Pert Pro X-ray diffraction (XRD). The absorption spectra of the MAPbBr3 thin films were obtained by using an UV–visible absorption spectrophotometer (Shimadzu; UV3150). A fluorescence spectrofluorometer (Horiba; Fluorolog3) was employed to record the PL spectra of the MAPbBr3 thin films at the temperature range of 10–300 K. Transient PL measurement was carried out by using a single photon counting spectrometer with a pulsed nano-LED as the excitation source (371 nm). The chemical composition and bond states of the produced MAPbBr3 samples were estimated by using the X-ray photoelectron spectrometer (XPS, SPECS XR50 system) measurements.

2


Page 3 of 14

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


Figure 1. Preparation procedures for the MAPbBr3 thin films by using vapor-assisted solution method.

RESULTS AND DISCUSSION The variation of the surface morphology of MAPbBr3 thin films with different evaporation time of MABr powder were examined by SEM measurements, and one can observe some distinct changing trends from Figure 2a. For S1 with the evaporation time of 5 min (Figure 2a), the sample displayed a smooth surface without obvious grain boundary. This may be due to that the evaporation time of MABr is too short and a stoichiometric reaction of MABr and PbBr2 could not be guaranteed. As the evaporation time was extended to 10 min (S2, Figure 2b), some large grains appeared on the surface of the sample, but the surface coverage of MAPbBr3 thin films was poor and some micro-pores existed. The formation of these micro-pores was mainly due to the restricted diffusivity ability of MABr vapor onto the surface of PbBr2/ITO template. Besides, the adjacent nano-grains were linked, and the grain size and grain boundary were difficult to identify from the SEM image with a high-magnification. As the evaporation time of MABr powder was further increased to 15 min shown in Figure 2c (S3), some cubic grains with the average size of ~150 nm were formed, but the undesired micro-pores still existed in the MAPbBr3 thin films. Further increasing the evaporation time of MABr powder to 20 min (S4) and 25 min (S5) increased the grain size of MAPbBr3 to ~400 nm, but the grains could not uniformly cover on the entire substrate surface. That was to say, the surface coverage had almost no change compared with S3. In the case of an increased evaporation time of 30 min (S6, Figure 2f), a nonuniform aggregation of the nanograins emerged. The grain size turned to be uneven and the typical grain size ranged from 200 to 800 nm. Also, the undesired micro-pores still stayed on the surface. It should be mentioned that the undesired micro-pores did not penetrate throughout the entire thin films, and they only extended a certain depth into the layer because the diffraction peaks from the underlying PbBr2 layer could not be identified, as seen in the corresponding XRD pattern later. Apparently, the obvious but regular changes of the microstructure for the produced MAPbBr3 thin films should be related with the increased evaporation time of MABr powder as the growth of PbBr2 layer was the same for all samples.

Figure 2. Surface SEM images of the MAPbBr3 samples grown with different evaporation time of MABr powder: (a) S1, 5 min, (b) S2, 10 min, (c) S3, 15 min, (d) S4, 20 min, (e) S5, 25 min, and (f) S6, 30 min.

To investigate the variation of the crystallinity properties of the MAPbBr3 thin films with different evaporation time of MABr, XRD measurements were carried out. Note that the XRD patterns of PbBr2 and MABr thin films were also put together for a better comparison (Black and red curves in Figure 3a). When the evaporation time of MABr powder was 5 min, only the diffraction peaks from PbBr2 can be identified, indicating the inadequate reaction process that only a small amount of MAPbBr3 components formed, matching with the observation in corresponding SEM image. With increasing the evaporation time to 10 min and above, the characteristic diffraction peak from MAPbBr3 crystal appeared at 14.92°, corresponding to the (100) plane of crystalline cubic MAPbBr3, and the diffraction intensity increased gradually with the evaporation time. Inversely, the diffraction intensities from PbBr2 component were weakening with the increasing evaporation time of MABr. These results suggested that the MAPbBr3 grains had formed and their crystallinity could be improved with more MABr vapor on the surface.13 However, owing to the limited evaporation time of MABr, the underlying PbBr2 layer could not be fully converted into MAPbBr3. Except for an insufficient evaporation time, another possible factor of solid-vapor interfacial conversion reaction should also be considered. At a low concentration of MABr vapor, the conversion into MAPbBr3 occurred though a solid-vapor interfacial reaction process, by which the MABr diffusion into the underlying PbBr2 layer resulted in the formation of crystalline MAPbBr3, and the reaction process can be expressed by equation (1) below. However, the proposed solid-vapor interfacial reaction process would be suppressed to some extent by the initially formed MAPbBr3, which would obstruct the subsequent diffusion of MABr vapor into the underlying PbBr2.14 PbBr2 ( s ) + MA+ ( v ) + Br − ( v ) → MAPbBr3 ( s )

(1)

3



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

Page 4 of 14

a complete conversion of the underlying PbBr2 into the crystalline cubic MAPbBr3. To investigate the effect of evaporation time of MABr on the optical characterizations of the MAPbBr3 thin films, PL measurements at room-temperature (RT) were performed. Note that the emission peak position in the PL spectra has been calibrated by the excitation line. From Figure 3b, one can observe that six MAPbBr3 samples presented dominant green emission at around 530 nm, and the emission intensity of the MAPbBr3 samples increased with increasing evaporation time, consistent with their morphology and crystalline evolution trends discussed above. Since the layers of PbBr2 and MABr did not exhibit luminescence characteristics, the measured PL emission deservedly came from the resulting MAPbBr3 thin films. The inset of Figure 3b plotted the full-width at halfmaximum (FWHM) of the green emission of MAPbBr3 products, and an ever-reduced FWHM value also indicated an improved emission performance with longer evaporation time of MABr. An optimized FWHM of 23.8 nm was achieved for S6 with an evaporation time of 30 min. A good optical quality with a relatively high emission color purity suggests the formation of high-quality MAPbBr3 thin films, which are therefore considered as an effective emission layer used in novel PeLEDs.17

Figure 3. (a) XRD patterns and (b) PL spectra measured at RT of the MAPbBr3 samples grown with different evaporation time of MABr powder: S1 (5 min); S2 (10 min); S3 (15 min); S4 (20 min); S5 (25 min); S6 (30 min). The inset in Fig. 3b shows the integrated intensity (blue square) and FWHM (red circle) of the emission peak versus sample number.

Such undesired process could be eliminated at a high concentration of MABr vapor, in which a dissolutionrecrystallization reaction process might take effect. A sufficient MA+ and Br– ions would favor for the thermodynamic formation of PbBr42– complex on the surface. Under this circumstance, the PbBr2 became more soluble in concentrated bromine vapor.15,16 In detail, the excessive amounts of Br– could provide an additional driving force that supported the dissolution of the initially formed MAPbBr3 products and the unconverted PbBr2 by the reactions (2) and (3) until the local concentration of PbBr42– reached saturation: MAPbBr3 ( s ) + Br − ( v ) → MA+ ( v ) + PbBr4 2− ( v ) −

(2)

2−

(3) PbBr2 ( s ) + 2 Br ( v ) → PbBr4 (v ) 2– After the MABr vapor was oversaturated with PbBr4 complexes, PbBr42– ions will react with MA+ ions and slowly recrystallize to growth MAPbBr3 consequently, following the reaction (4) below. (4) PbBr4 2 − ( v ) + MA+ ( v ) → MAPbBr3 ( s ) + Br − ( v ) The proposed dissolution-recrystallization growth mechanism was further verified by the experimental results. As the evaporation time of MABr powder was extended to 25 min and above, the characteristic diffraction peak of MAPbBr3 obviously dominated the XRD spectra, and no other peaks from PbBr2 and MABr components were detected, suggesting

Figure 4. Surface SEM images of the MAPbBr3 thin films prepared with different evaporation temperature of MABr powder: (a) M1, 120 °C, (b) S5, 135 °C, (c) M2, 150 °C, (d) M3, 165 °C, and (e) M4, 180 °C. (f) Absorption spectrum of the MAPbBr3 thin films (M4) measured at RT.

The adjustment of other processing parameters in the MAPbBr3 preparation process may strengthen an additional understanding of the growth kinetics for MAPbBr3 thin films. Take S5 as a standard, a series of samples at different evaporation temperature of MABr powder ranging from 120 to 180 °C were prepared, and the evaporation time was kept at 25 min. Four samples were denoted as M1 (120 °C), M2 (150 °C), M3 (165 °C), and M4 (180 °C), respectively. As shown in Figure 4a–e, the MAPbBr3 thin films presented a regular morphological evolution. With increasing evaporation temperature of MABr powder, the grain size and surface coverage of MAPbBr3 increased gradually. For M1 with a low evaporation temperature (120 °C), a similar morphology feature can be observed as that of S5 except for a relatively smaller grain size. As the evaporation temperature was elevated to 150 °C and above (M2, M3), the amount and size of micro-pores were substantially reduced, and the grain consolidation played a primary role. One can see that densely packed MAPbBr3 thin films had been achieved with less grain boundaries although a

4


Page 5 of 14

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


low density of micro-pores can be still observed on the film surface. At a higher evaporation temperature of 180 °C (M4), an uniform and compact surface without any micro-pores was observed, and a two-dimensional MAPbBr3 thin films was achieved. The inset of Figure 4f displayed a representative photograph of the MAPbBr3 sample (M4) on ITO/glass substrate, and its corresponding optical absorption spectrum was also measured. One can see that the absorption onset of MAPbBr3 occurred at 530 nm, well consistent with the previous studies on solution-processed MAPbBr3 thin films.9,18

Figure 5. (a) XRD patterns and (b) PL results of the MAPbBr3 thin films prepared with different evaporation time of MABr powder. (c) Transient PL results and the fitting curves of the MAPbBr3 thin films. The corresponding τave. of five samples were presented in the inset. (d) Total XPS spectrum of the produced MAPbBr3 thin films (M2), and the corresponding (e) Br 3d and (f) Pb 4f core levels.

Further, the structural characterizations of the produced MAPbBr3 thin films were investigated by XRD. As shown in Figure 5a, five samples presented almost the same dominant diffraction peaks at 14.92°, 30.11°, 33.78°, and 37.91°, which can be assigned to the (100), (200), (021), and (211) planes of cubic crystalline MAPbBr3, respectively. The absence of diffractions from PbBr2 or MABr components indicated a high level of phase purity of the products, which suggested that MABr vapor had fully reacted with the bottom PbBr2 layer at such an evaporation temperature window. In addition, as the evaporation temperature was increased gradually, the FWHM values of (100) and (200) planes were observed to diminish, and the corresponding diffraction intensities were observed to increase, which suggests the crystallinity of the MAPbBr3 products was improved as the elevated evaporation temperature of MABr promoted the surface coverage and also the coalescence of the adjacent nano-grains. To examine the optical characterizations of the MAPbBr3 thin films under different evaporation temperature, PL measurement were performed at RT. As shown in Figure 5b, five samples presented excellent optical quality with single emission peak at ~530 nm, and the corresponding emission intensity increased gradually with increasing evaporation temperature, well consistent with the morphology and crystalline evolution trends observed above. To investigate the carrier recombination dynamics of the MAPbBr3 thin films, a time-resolved PL measurement was performed. As shown in Figure 5c, M2 sample was taken as

the representative considering that five curves displayed the similar spectra features and they all can be fitted with a biexponential decay, and an average PL decay lifetime (τave.) of 45 ns was derived for M2. An additional observation was that the τave. of five samples differed significantly and depended strongly on their processing conditions. As summarized in the inset (Figure 5c), the τave. increased gradually with the increase of evaporation temperature of MABr. As deduced from foregone research, the lengthening of the τave. can be explained by the decrease of charge carrier trapping states as the relaxation pathways.19,20 The previous studies had recognized that the τave. was inversely proportional to the trap state density, where longer τave. correlated with less state density. The observed changing trend of τave. was well consistent with the results of SEM and PL performances. We further performed the XPS measurements to investigate the chemical composition and bond states of the produced MAPbBr3 samples. As shown in Figure 5d, a total XPS spectrum of the MAPbBr3 thin films (M2) was displayed, where the signal peaks of Br, Pb, C, and N elements were obtained in the range of specific detection. Note that the molar ratio of Pb and Br elements incorporated was 1:3.05. As shown in Figure 5e, the Br 3d spectrum was investigated to examine the chemical bond state of Br element. The spectrum can be fitted with two components, with their binding energies at 68.12 eV (Br-1) and 69.23 eV (Br-2). The Br-1 peak corresponded to the level of Br 3d5/2, splitting to the level of Br 3d3/2 at 1.11 eV. Obviously, both two components originate from the bromide ions in MAPbBr3.9 The Pb 4f spectrum was presented in Figure 5f, which can be fitted well with two separated peaks, Pb-1 (137.94 eV) and Pb-2 (143.83 eV), corresponding to the levels of Pb 4f7/2 and Pb 4f5/2. Because there exists obvious differences of the Pb 4f7/2 and Pb 4f5/2 levels with those in metallic Pb and PbBr2,21,22 we therefore considered that the contributions for the two bands were associated with the Pb2+ in PbBr3–. For a better understanding on the optical recombination properties of produced MAPbBr3 thin films, PL spectra at different measurement temperatures (10 to 300 K) were acquired for M4, a representative sample. As shown in Figure 6a, a clear changing trend of the luminescence characteristics was revealed, mainly involving the emergence/disappearance of the emission peak, the peak shift, and the peak intensity variation. At the high-energy side of the PL spectra, there are two important aspects that deserved to note: first, as the temperature increased, the corresponding PL intensity decreased, and the peak position monotonically blue-shifted. As for the unusual blue-shift behavior with increasing temperature, two important factors should be considered, including the thermal expansion of crystal lattice and the electron-phonon interaction that induced the renormalization of bandgap energy of semiconductors, and the similar phenomena had also been observed in lead/copper chalcogenide semiconductors such as CuCl(Br/I) and PbS(Se/Te).23 Secondly, at a critical temperature of 140 K, an asymmetric structure of the PL emission band can be observed and an additional emission peak at around 549.5 nm emerged. At the same time, the ratio of the relative intensity of two emission peaks (Peak I, Peak II) was changing with temperature. Herein, the Peak I was attributed to the free exciton related emission in MAPbBr3 thin films.24 The Peak II was located at the low-energy side of Peak I, and

5



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

Page 6 of 14

its energy difference between Peak I was obviously different with the temperature changing, which suggested that peak II was not the phonon replica of Peak I. Based on the observations in previous studies, the low-energy Peak II appearing at 140 K should be attributed to the trapped charge-carrier pairs.23 In our case, the MAPbBr3 polycrystalline thin films were characterized by a high density of nano-sized crystallites with a poorly defined crystal facet, and the undesired trap states probably located at the interfacial areas among them. Therefore, trap-mediated recombinaiton or trapped chargecarrier recombination may exist in the MAPbBr3 thin films. In addition, it was noteworthy that no structural phase transition occurred within the heating process from 10 to 300 K because no unnormal PL shift resulting from structural phase transition appeared, unlike other reports on MAPbBr3 nanowires with obvious reversal of the PL emission.26 At the low-energy side, another new emission peak emerged at around 620 nm at 130 K, and the corresponding intensity of this emission component (Peak III) increased with decreasing temperature. The intensity enhancement behavior of Peak III can be explained by the effective suppression of nonradiative recombination at low temperature. We speculated that the origin of this peak may be related to the defect states of Br vacancy in MAPbBr3 thin films.27 The position of Br vacancy energy level was close to the conduction band of MAPbBr3 thin films, and electrons transition from the Br vacancy level to valence band resulted in the recombination luminescence. The corresponding physical model could be described as follows: the incident light was absorbed by the MAPbBr3 thin films, and thus the electrons in the valence band could be promoted to the excitation states in the conduction band of MAPbBr3. Upon relaxation, the free electrons in the conduction band will be trapped or captured by a localized Br vacancy defect. Consequently, the combination of trapped electrons with free holes in the valence band of MAPbBr3 thin films occurs. The above evolution behaviors of PL spectra suggested that three recombination channels existed in the MAPbBr3 thin films, and they took effect at different temperature windows. At the temperature window of 140−300 K, Peak I and II dominated the overall emission spectra, which can be more clearly observed by a typical Gaussian deconvolution of the PL spectra obtained at 170 K (Figure 6c, bottom pane). While, at another temperature window of 10−130 K, Peak I and III contributed the emission spectra, as seen in a typical Gaussian deconvolution of the PL spectra obtained at 60 K (Figure 6c, upper pane). Figure 6d summarized the shift of three emission components (Peak I, II, and III) of the MAPbBr3 thin films versus measured temperature. From the wavelength shift of Peak I, the sensitivity of the measured photon energy of MAPbBr3 with temperature was estimated based on the empirical equation of ET = E0 + αT, where ET is the measured photon energy, α presents the temperature sensitivity coefficient, and E0 (T = 0) is the fitting constant. Finally, the fitting results generate a value of 0.319 ± 0.005 meV/K for α.

Figure 6. (a) PL spectra of the MAPbBr3 thin films (M4) at different measured temperatures. (b) Magnified view of the PL spectra at the low-energy side taken from 130 K to 10 K. (c) Gaussian deconvolution of the PL spectra measured at 60 and 170 K, respectively. (d) Variation of the peak position of Peak I, II, and III versus measured temperature.

Considering the fact that the cubic phase of MAPbBr3 thin films was stable in the investigated temperature region (10– 300 K), important physical parameters such as EB and excitonphonon interaction were therefore estimated according to the obtained results above. We firstly plotted the integrated emission intensity of Peak I versus temperature to determine the EB of MAPbBr3 thin films. As presented in Figure 7a, the integrated PL emission intensity of Peak I increased exponentially with decreasing temperature, and this phenomenon can be assigned to the thermally activated nonradiative recombination process, which was increasing with the temperature. The overall temperature quenching behavior of the MAPbBr3 could be well described by the following expression:28 (5) EB 1 + A exp( − ) KT where I0 was the emission intensity at a temperature of 0 K, K was the Boltzmann constant, A was the fitted constant, and EB was the calculated exciton binding energy. As displayed in Figure 7a, The fitting curve yielded a value for EB as 50.2±1.8 meV. Theoretically, the relatively higher EB of MAPbBr3 thin films than that of thermal ionization energy at RT (~26 meV) ensured excitons survival well above RT and their high-rate recombination, highlighting the strong potentiality of such materials in exciton-related optoelectronic devices, such as LEDs and laser diodes. I (T ) =

I0

6


Page 7 of 14

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

ACS Applied Materials & Interfaces carefully measured the PL spectra at 10 K under different excitation lines (325, 365, 405, and 467 nm). As displayed in Figure 7c, the obtained emission spectra at 10 K were almost unchanged in their spectral shape even with obviously different excitation energy, and the variation in their emission intensity may be due to the different power density of four excitation lines. Such feature was obviously distinct from the wellknown performance from carbon and ZnO NCs,31,32 in which the emission peak position depended strongly on the excitation energy.

Figure 7. (a) Integrated PL intensity and (b) FWHM of Peak I of the MAPbBr3 thin films (M4) versus reciprocal temperature. (c) PL spectra (measured at 10 K) of the MAPbBr3 thin films (M4) with different excitation wavelength. (d) Excitation power dependent PL spectra of the MAPbBr3 thin films (M4) from 0.25 mW to 1.0 mW. (e) Logarithm plot of the integrated emission intensity of the MAPbBr3 thin films (M4) at 300 K and 10 K as a function of excitation power. (f) Excitation power dependent spectrum bandwidth of the PL emission peak (M4).

Figure 7b shows the summarized date of the measured FWHM of the PL emission peak in the temperature range of 160−300 K. From the broadening behavior of the FWHM value, the exciton-phonon interaction can be examined and the discrete data can be fitted well according to the Boson model:29 Γ(T ) = Γ0 + σT +

Γop

(6) exp(hωop / KT ) − 1

in which Γ0 was the inhomogeneous broadening contribution, σ was the interaction of exciton-acoustic phonon to the linewidth broadening, Γop described the exciton-optical phonon contribution to the linewidth broadening. Obviously, at the low temperature level, Γ0 dominated the right side of the formula; while, at relatively high temperatures, the contributions from σ and Γop were larger than that of Γ0. Therefore, the FWHM value of the emission peak of MAPbBr3 thin films would nonlinearly increased with temperature. The fitting results indicated that the optical phonon involved in the interaction of exciton-phonon mainly, and the constants of Γop and optical phonon energy (ħωop) values were calculated to be 87.2±6.6 meV and 35.0±2.9 meV, respectively, which implied a strong exciton-phonon interaction and agreed well with the thermal antiquenching effects of MAPbBr3 thin films at high temperature.30 To better understand the characteristics of the observed excitonic emission from MAPbBr3 thin films, we

Further, power-dependent PL measurements were performed at 300 and 10 K to verify the exciton emission characteristics of the MAPbBr3 thin films, in which the excitation line of 405 nm was employed and the excitation power was adjusted from 0.25 mW to 1.0 mW. As shown in Figure 7d, six PL spectra were put together, from which one can observe that six spectra shared almost the same emission peak position and spectral shape. Further, the relationship of integrated emission intensity (IEm) and the excitation light intensity (IEx) was established in Figure 7e, and the relationship can be described by an empirical equation of IEm = IExα, in which α denoted the non-linear component. Considering the fact that the MAPbBr3 thin films were optically excited at non-resonant condition, the above equation can be fitted by the power-law function, and a value of 1.57 for β was obtained at 300 K, confirming the exciton behavior of the recombination emission from MAPbBr3 thin films (It should be mentioned herein that 2 > β > 1 represented the free/bound excitons related recombination emission).34 At a low temperature of 10 K, a reduced thermal quenching can be expected, so a smaller value of β = 1.22 held the curve. Besides, with the increase of light excitation power, the bandwidth of the PL emission peak was observed to increase gradually, as displayed in Figure 7f. The observed band broadening phenomenon maybe a signature of the change of carrier recombination path in MAPbBr3. Herein, we considered that the generation of free carrier related recombination under high excitation power was responsible for the above phenomenon. It was because that the free carrier related recombination generally occurred at high excitation conditions, which acted as a competitor to the dominant excitonic recombination. Thus the newly generated free carriers would inevitably induce a screening effect, and the direct consequence was that the bandgap for MAPbBr3 perovskites was reduced and the excited states were more likely to dissociate into free carriers. As a result, the probability of radiatively excitonic recombination was decreased. In spatial spectra, the component or contribution of excitonic emission decreased gradually with increasing the excitation power.

7



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

Page 8 of 14

Figure 8. (a) Schematic device structure and (b) energy band diagram of the proposed PeLEDs. (c) Current density (red) and luminance (blue) of the PeLEDs at different bias voltage. (d) EL spectra of the device under different applied voltage. A typical emission photograph of the LED with an active area of 2 × 2 mm2 (at 8.0 V) was presented in the inset. (e) EQE, power efficiency, and current efficiency of the PeLEDs at different bias voltage. (f) CIE coordinates of the PeLED (8.0 V). (g) Integrated EL intensity of the PeLED at different viewing angles. (h) Response of the EL intensity for six switch-on/switch-off pulses at 8.0 V (100 s between each switch). (i) Statistical diagram of the maximum EQE measured from thirty PeLEDs with the same device structure.

The excitonic emission characteristics of the as-synthesized MAPbBr3 products with a good crystalline and a high surface coverage made themselves a suitable green emitter in LEDs applications. For device preparations, a multi-layered device configuration based on Al/LiF/TPBi/MAPbBr3/NiO/ITO structure was designed, in which NiO (TPBi) thin films were employed as the hole (electron) injector. The detailed processing procedures can be found in the Experimental Section. Figure 8a displayed a schematic device structure of the proposed PeLED, and the its energy band diagram was presented in Figure 8b. The NiO layer with an p-type conductivity played the roles of electron-blocking and hole-providing layer simultaneously in virtue of its relatively high ionization potential energy and also matched electron affinity (χ) with MAPbBr3,35 and analogously, the TPBi layer served as the hole-blocking and electron-providing layer because of its relatively small χ and low energy level of valence band.18 Therefore, the carrier injection processes from bottom p-NiO and top TPBi layers to the middle MAPbBr3 active layer can be expected, allowing the electrons and holes to recombine effectively, and further producing the dominant green emission with the photon energy consistent with the bandgap of MAPbBr3.

For electroluminescence (EL) measurements, the fabricated device was electrically driven in a continuous current mode, and the corresponding EL spectra were detected from the ITO/glass substrate side. Figure 8c presented the luminance (blue) and current density (red) curves of the PeLED at different bias voltage. From the current density–voltage curve, a turn-on voltage of ~2.5 V can be observed (Note that the turnon voltage was commonly defined in literature as the voltage necessary to detect luminance of 1 cd/m2), much smaller than many reported PeLEDs,3,17 which presumably benefited from a reasonable device architecture with well-designed barrier-free charge injectors. Above the onset voltage, the luminance and current density of the device increased quickly, which suggested that the undesired series resistance and/or nonradiative recombination in the device would not increase with the driving voltage proportionally.36 At a bias voltage of 10.0 V, a peak luminance (6530 cd/m2) was achieved, corresponding to a current density of 236 mA/cm2. Typical EL spectra of the PeLED was shown in Figure 8d, displaying a bright green emission (532 nm) with the FWHM as small as 25 nm, and no obvious broadening behavior can be observed with the increasing applied bias. Such a low FWHM for color-pure devices opened up enormous opportunities for their ultrahigh-

8


Page 9 of 14

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


definition display applications. The inset of Figure 8d displayed a typical emission photograph of the PeLED captured by a digital camera (8.0 V), and the emission was uniformly over the entire active electrode unit. The Commission International de I’Eclairage (CIE) color of the green emission is measured with the coordinates of (0.10, 0.71), as shown in Figure 8f. In addition, we also measured other key device parameters, such as EQE, power efficiency, and current efficiency, of the PeLED, which were plotted in Figure 8e as a function of bias voltage. These parameters rose rapidly with the increasing voltage, indicating that the radiative bimolecular recombination dominated at high excitation densities.37 The best-performing device reached up an EQE of ~4.36%, a power efficiency of 4.49 lm/W, and a current efficiency of 8.16 cd/A, much better than most PeLEDs from MAPbBr3 thin films or quantum dots (Table 2).[7,9,10,13,17,38–47] It should be mentioned that the device suffered a degradation above 10.0 V, and we considered that the heating effect induced by high excitation intensity was probably the leading cause. Further, we estimated the internal quantum efficiency (IQE) of the PeLED according to the empirical formulae of IQE = 2n2EQE,48 which gave an IQE of 19.6% (a refractive index of ~1.5 for glass was used for calculation). As shown in Figure

8g, angle-dependent EL measurements were carried out to further verify the emission profile of the PeLED. For such measurements, the detector was placed at different angles (– 90º–90º) with respect to the surface of the device. From Figure 8g, we can observe that the maximum EL intensity occurred at the normal direction of the electrode unit, and the EL intensity decreased gradually with decreasing inclination angels, similar with the observations in other studies. From the application point of view, the solid-state light source with a relatively large emission profile and a high brightness may have a potential application in daily life. Moreover, we further performed the kinetic experiments to characterize the stability and time response of the studied PeLED. Figure 8h shows the response of the EL intensity for six switch-on/switch-off pulses at 8.0 V, and the time decay is 100 s between each switch. No any degradation on the device performance indicates the good stability of the studied PeLEDs. To assess the reproducibility of the PeLEDs, thirty devices were randomly selected for identical measurements in a continuous current mode. As shown in Figure 8i, an average peak EQE of 4.32% with a relative standard deviation of ~12.6% was derived, manifesting a good reproducibility of our devices.

Table 2. Summary of the device performance of the prepared PeLEDs.

MAPbBr3 particles

EL λmaxa (nm) 520

Linewidth (nm) –

Max. Lub (cd/m2) 3515

Max. CEc (cd/A) 11.49

Max. EQE (%) 3.8

Max. PEd (lm/W) 7.84

[38]

MAPbBr3 QDse)

524

24

2503

4.5

1.1

3.50

[39]

MAPbBr3 QDs

520

20

2398

3.72

1.06

–

[40]

MAPbBr3 QDs

534

19

~3000

–

1.20

–

[41]

MAPbBr3 films

535

30

186

0.12

0.065

–

[17]

MAPbBr3 films

517

≥30

364

0.3

0.24

–

[7]

MAPbBr3 films

530

23

4064

0.74

0.165

–

[42]

MAPbBr3 films

513

–

2900

17.1

9.3

13.0

[10]

MAPbBr3 films

528

~25

8794

5.1

–

–

[13]

MAPbBr3 films

543

~20

417

0.577

0.125

–

[43]

MAPbBr3 films

530

~41

~545

0.22

0.051

0.11

[44]

MAPbBr3 films

420

20

6942

0.72

0.17

–

[45]

MAPbBr3 films

540

20

>10000

42.9

8.53

–

[9]

MAPbBr3 films

550

≥100

1.8

0.013

–

–

[46]

MAPbBr3 films

540

~25

3490

0.43

0.10

0.31

[47]

MAPbBr3 films

532

25

6530

8.16

4.36

4.49

This work

Emission materials

a

Ref.

λmax: Peak position; bLu: Luminance; cCE: Current efficiency; dPE: Power efficiency; eQDs: Quantum dots.

CONCLUSION In conclusion, organic-inorganic halide perovskite MAPbBr3 thin films were synthesized by vapor-assisted solution method. We have detailedly investigated the effect of the evaporation time and temperature of MABr powder on the microstructure, crystallinity and optical characterizations of MAPbBr3 thin films, and a controllable morphology evolution was observed. Further, we performed the temperaturedependent and time-resolved PL measurements to investigate the optical transition mechanisms and carrier recombination dynamics of MAPbBr3 thin films. Moreover, we estimated the temperature sensitivity coefficient of bandgap, EB, and optical

phonon energy of MAPbBr3 thin films based on the obtained data. More importantly, a planar green PeLED based on Al/LiF/TPBi/MAPbBr3/NiO/ITO structure was designed and fabricated. The device performance was remarkable in terms of its high luminance (6530 cd/m2), current efficiency (8.16 cd/A), EQE (4.36%), and power efficiency (4.49 lm/W). We believe that the experimental results obtained may have a guiding significance for the synthesis of high-quality perovskite thin films than can be employed in reliable and high-performance PeLEDs.

9



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

EXPERIMENTAL SECTION Device Fabrication. The fabrication procedure of the Al/LiF/TPBi/MAPbBr3/NiO/ITO structure started with the pre-cleaned ITO (200 nm, square resistance: 8 Ω/□) glass substrates, which were followed by the oxygen plasma treatment for 5 min. Firstly, a densely packed NiO thin films were grown on the ITO/glass substrates by using the spraying technique, in which the solutions of C10H14NiO4 in acetonitrile were employed as the precursors and the ITO glass substrate were mounted on a hot template (450 °C). Then, the substrates were transferred into a glove box for subsequent perovskite processing step. Followed that, vapor-assisted solution method was employed for MAPbBr3 active layer preparation according to the processing conditions of M4 sample. Finally, the substrates were taken into a evaporator to thermally deposit TPBi (40 nm), LiF (1 nm), and metal Al (100 nm). The active area of singe emission unit was 4 mm2 as defined by the overlapping area of the Al electrode and patterned ITO. Device Characterization. The current-voltage properties of the PeLEDs were recorded by a semiconductor characterization analyzer (Keithley 2400). The EL spectra were acquired by using an acquisition equipment including a lock-in amplifier systems (Stanford SR830-DSP) and a photomultiplier tube (PMTH-S1-R1527). The luminance-voltage characterizations were measured by using a PR650 SpcetraScan spectrophotometer (Photo Research) in air. The EQE of the PeLEDs was measured by a photodetector (THORLABS, S120VC) and a digital optical power meter (THORLABS, PM100D).

AUTHOR INFORMATION Corresponding Authors [email protected]; [email protected]

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENTS This work was financially supported by the National Natural Science Foundation of China (Nos. 11774318, 11604302, 61176044 and 11504331), the China Postdoctoral Science Foundation (2017T100535 and 2015M582193), the Postdoctoral Research Sponsorship in Henan Province (2015008), the Science and Technology Research Project of Henan Province (162300410229), the Outstanding Young Talent Research Fund of Zhengzhou University (1521317001), and the Startup Research Fund of Zhengzhou University (1512317003).

REFERENCES (1) Luo, J.; Im, J. H.; Mayer, M. T.; Schreier, M.; Nazeeruddin, M. K.; Park, N. G.; Tilley, S. D.; Fan; H. J.; Gratzel, M. Water Photolysis at 12.3% Efficiency via Perovskite Photovoltaics and EarthAbundant Catalysts. Science 2014, 345, 1593. (2) Green, M. A.; Ho-Baillie, A.; Snaith, H. J. The Emergence of Perovskite Solar Cells. Nat. Photon. 2014, 8, 506–514. (3) Song, J. Z.; Li, J. H.; Li, X. M.; Xu, L. M.; Dong, Y. H.; Zeng, H. B. Quantum Dot Light-Emitting Diodes Based on Inorganic Perovskite Cesium Lead Halides (CsPbX3). Adv. Mater. 2015, 27, 7162– 7167. (4) Deschler, F.; Price, M.; Pathak, S.; Klintberg, L. E.; Jarausch, D. D.; Higler, R.; Hüttner, S.; Leijtens, T.; Stranks, S. D.; Snaith, H.

Page 10 of 14

J.; Atatüre, M.; Phillips, R. T.; Friend, R. H. High Photoluminescence Efficiency and Optically Pumped Lasing in Solution-Processed Mixed Halide Perovskite Semiconductors. J. Phys. Chem. Lett. 2014, 5, 1421–1426. (5) Yang, W. S.; Noh, J. H.; Jeon, N. J.; Kim, Y. C.; Ryu, S.; Seo, J. High-Performance Photovoltaic Perovskite Layers Fabricated Through Intramolecular Exchange. Science 2015, 348, 1234. (6) Zhu, H. M.; Fu, Y. P.; Meng, F.; Wu, X. X.; Gong, Z. Z.; Ding, Q.; Gustafsson, M. V.; Trinh, M. T.; Jin, S.; Zhu, X. Y. Lead Halide Perovskite Nanowire Lasers with Low Lasing Thresholds and High Quality Factors. Nat. Mater. 2015, 14, 636–642. (7) Tan, Z. K.; Moghaddam, R. S.; Lai, M. L.; Docampo, P.; Higler, R.; Deschler, F.; Price, M.; Sadhanala, A.; Pazos, L. M.; Credgington, D.; Hanusch, F.; Bein, T.; Snaith, H. J.; Friend, R. H. Bright LightEmitting Diodes Based on Organometal Halide Perovskite. Nat. Nanotechnol. 2014, 29, 687–692. (8) Quintero, O. A. J.; Sanchez, R. S.; Rincon, M.; Sero, I. M. Bright Visible-Infrared Light Emitting Diodes Based on Hybrid Halide Perovskite with Spiro-OMeTAD as a Hole-Injecting Layer. J. Phys. Chem. Lett. 2015, 6, 1883–1890. (9) Cho, H.; Jeong, S. H.; Park, M. H.; Kim, Y. H.; Wolf, C.; Lee, C. L.; Heo, J. H.; Sadhanala, A.; Myoung, N.; Yoo, S.; Im, S. H.; Friend, R. H.; Lee, T. W. Overcoming the Electroluminescence Efficiency Limitations of Perovskite Light-Emitting Diodes. Science 2015, 350, 1222–1225. (10) Xiao, Z. G.; Kerner, R. A.; Zhao, L. F.; Tran, N. L.; Lee, K. M.; Koh, T. W.; Scholes, G. D.; Rand, B. P. Efficient Perovskite Light-Emitting Diodes Featuring Nanometer-Sized Crystallites. Nat. Photon. 2017, 11, 108–115. (11) Liu, M. Z.; Johnston, M. B.; Snaith, H. J. Efficient Planar Heterojunction Perovskite Solar Cells by Vapor Deposition. Nature 2013, 501, 395–398. (12) Shi, Z.; Sun, X.; Wu, D.; Xu, T. T.; Tian, Y.; Zhang, Y.; Li, X.; Du, G. T. Near-Infrared Random Lasing Realized in a Perovskite CH3NH3PbI3 Thin Film. J. Mater. Chem. C 2016, 4, 8373–8379. (13) Wang, Z.; Cheng, T.; Wang, F.; Dai, S.; Tan, Z. Morphology Engineering for High-Performance and Multicolored Perovskite Light-Emitting Diodes with Simple Device Structure. Small 2016, 12, 4412–4420. (14) Fu, Y.; Meng, F.; Rowley, M. B.; Thompson, B. J.; Shearer, M. J.; Ma, D.; Hamers, R. J.; Wright, J. C.; Jin, S. Solution Growth of Single Crystal Methylammonium Lead Halide Perovskite Nanostructures for Optoelectronic and Photovoltaic Applications. J. Am. Chem. Soc. 2015, 137, 5810–5818. (15) Chen, C. W.; Kang, H. W.; Hsiao, S. Y.; Yang, P. F.; Chiang, K. M.; Lin, H. W. Efficient and Uniform Planar-Type Perovskite Solar Cells by Simple Sequential Vacuum Deposition. Adv. Mater. 2014, 26, 6647–6652. (16) Tong, X.; Lin, F.; Wu, J.; Wang, Z. M. High Performance Perovskite Solar Cells. Adv. Sci. 2016, 3, 1500201. (17) Shi, Z.; Sun, X.; Wu, D.; Xu, T.; Zhuang, S.; Tian, Y.; Li, X.; Du, G. T. High-Performance Planar Green Light-Emitting Diodes Based on a PEDOT:PSS/CH3NH3PbBr3/ZnO Sandwich Structure. Nanoscale 2016, 8, 10035–10042. (18) Chih, Y.; Wang, J.; Yang, R. T.; Liu, C. C.; Chang, Y. C.; Fu, Y. S.; Lai, W.; Chen, P.; Wen, T. C.; Huang, Y.; Tsao, C.; Guo, T. F. NiOx Electrode Interlayer and CH3NH2/CH3NH3PbBr3 Interface Treatment to Markedly Advance Hybrid Perovskite-Based LightEmitting Diodes. Adv. Mater. 2016, 28, 8687–8694. (19) Jones, M.; Lo, S. S.; Scholes, G. D. Signatures of Exciton Dynamics and Carrier Trapping in the Time-Resolved Photoluminescence of Colloidal CdSe Nanocrystals. J. Phys. Chem. C 2009, 113, 18632–18642. (20) Wu, X.; Trinh, M. T.; Niesner, D.; Zhu, H. M.; Norman, Z.; Owen, J. S.; Yaffe, O.; Kudisch, B. J.; Zhu, X. Y. Trap States in Lead Iodide Perovskites. J. Am. Chem. Soc. 2015, 137, 2089–2096.

10


Page 11 of 14

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


(21) Pederson, L. R. Two-Dimensional Chemical-State Plot for Lead Using XPS. J. Electron Spectrosc. Relat. Phenom. 1982, 28, 203–239. (22) Nefedov, V. I. X-Ray Photoelectron Spectra of Halogens in Coordination Compounds. J. Electron Spectrosc. Relat. Phenom. 1977, 12, 459–476. (23) Dey, P.; Paul, J.; Bylsma, J.; Karaiskaj, D.; Luther, J. M.; Beard, M. C.; Romero, A. H. Origin of the Temperature Dependence of the Band Gap of PbS and PbSe Quantum Dots. Solid State Commun. 2013, 165, 49–54. (24) Shi, Z. F.; Li, Y.; Zhang, Y. T.; Chen, Y. S.; Li, X. J.; Wu, D.; Xu, T. T.; Shan, C. X.; Du, G. T. High-Efficiency and Air-Stable Perovskite Quantum Dots Light-Emitting Diodes with an AllInorganic Heterostructure. Nano Lett. 2017, 17, 313–321. (25) Dai, J.; Zheng, H. G.; Zhu, C.; Lu, J. F.; Xu, C. X. Comparative Investigation on Temperature-Dependent Photoluminescence of CH3NH3PbBr3 and CH(NH2)2PbBr3 Microstructures. J. Mater. Chem. C 2016, 4, 4408–4413. (26) Xing, J.; Liu, X. F.; Zhang, Q.; Ha, S. T.; Yuan, Y. W.; Shen, C.; Sum, T. C.; Xiong, Q. Vapor Phase Synthesis of Organometal Halide Perovskite Nanowires for Tunable Room-Temperature Nanolasers. Nano Lett. 2015, 15, 4571–4577. (27) Sebastian, M.; Peters, J. A.; Stoumpos, C. C.; Im, J.; Kostina, S. S.; Liu, Z.; Kanatzidis, M. G.; Freeman, A. J.; Wessels, B. W. Excitonic Emission and Above-Band-Gap Luminescence in the Single-Crystal Perovskite Semiconductors CsPbBr3 and CsPbCl3. Phys. Rev. B 2015, 92, 235210. (28) Jiang, D. S.; Jung, H.; Ploog, K. Temperature Dependence of Photoluminescence from GaAs Single and Multiple Quantum-Well Heterostructures Grown by Molecular-Beam Epitaxy. J. Appl. Phys. 1988, 64, 1371–1377. (29) Tanaka, K.; Takahashi, T.; Ban, T.; Kondo, T.; Uchida, K.; Miura, N. Comparative Study on the Excitons in Lead-Halide-Based Perovskite-Type Crystals CH3NH3PbBr3 CH3NH3PbI3. Solid State Commun. 2003, 127, 619–623. (30) Diroll, T. B.; Murray, C. B. High-Temperature Photoluminescence of CdSe/CdS Core/Shell Nanostructures. ACS Nano 2014, 8, 6466–6474. (31) Sun, Y. P.; Zhou, B.; Lin, Y.; Wang, W.; Fernando, K. A. S.; Pathak, P.; Meziani, M. J.; Harruff, B. A.; Wang, H.; Luo, P. G.; Yang, H.; Kose, M. E.; Chen, B.; Veca, L. M.; Xie, S. Y. QuantumSized Carbon Dots for Bright and Colorful Photoluminescence. J. Am. Chem. Soc. 2006, 128, 7756–7757. (32) Zeng, H. B.; Duan, G. T.; Li, Y.; Yang, S. K.; Xu, X. X.; Cai. W. P. Blue Luminescence of ZnO Nanoparticles Based on NonEquilibrium Processes: Defect Origins and Emission Controls. Adv. Funct. Mater. 2010, 20, 561–572. (33) Li, X. M.; Wu, Y.; Zhang, S. L.; Cai, B.; Gu, Y.; Song, J. Z.; Zeng, H. B. CsPbX3 Quantum Dots for Lighting and Displays: RoomTemperature Synthesis, Photoluminescence Superiorities, Underlying Origins and White Light-Emitting Diodes. Adv. Funct. Mater. 2016, 26, 2435–2445. (34) Shibata, H.; Sakai, M.; Yamada, A.; Matsubara, K.; Sakurai, K.; Tampo, H.; Ishizuka, S.; Kim, K. K.; Niki, S. Excitation-Power Dependence of Free Exciton Photoluminescence of Semiconductors. Jpn. J. Appl. Phys. 2005, 44, 6113–6114. (35) Kwon, U.; Kim, B. G.; Nguyen, D. C.; Park, J. H.; Ha, N. Y.; Kim, S. J.; Ko, S. H.; Lee, S.; Lee, D.; Park, H. J. SolutionProcessible Crystalline NiO Nanoparticles for High-Performance Planar Perovskite Photovoltaic Cells. Sci. Rep. 2016, 6, 30759. (36) Xu, S.; Xu, C.; Liu, Y.; Hu, Y.; Yang, R.; Yang, Q.; Ryou, J. H.; Kim, H. J.; Lochner, Z.; Choi, S.; Dupuis, R.; Wang, Z. L. Ordered Nanowire Array Blue/Near-UV Light Emitting Diodes. Adv. Mater. 2010, 22, 4749–4753. (37) Wang, J. P.; Wang, N.; Jin, Y.; Si, J.; Tan, Z. K.; Du, H.; Cheng, L.; Dai, X.; Bai, S.; He, H.; Ye, Z.; Lai, M. L.; Friend, R. H.; Huang, W. Interfacial Control Toward Efficient and Low-Voltage Perovskite Light-Emitting Diodes. Adv. Mater. 2015, 27, 2311–2316.

(38) Xing, J.; Yan, F.; Zhao, Y.; Chen, S.; Yu, H. K.; Zhang, Q.; Zeng, R.; Demir, H. V.; Sun, X. W.; Huan, A.; Xiong, Q. H. HighEfficiency Light-Emitting Diodes of Organometal Halide Perovskite Amorphous Nanoparticles. ACS Nano 2016, 10, 6623–6630. (39) Huang, H. L.; Zhao, F. C.; Liu, L.; Zhang, F.; Wu, X.; Shi, L. J.; Zou, B. S.; Pei, Q. B.; Zhong, H. Z. Emulsion Synthesis of SizeTunable CH3NH3PbBr3 Quantum Dots: An Alternative Route toward Efficient Light-Emitting Diodes. ACS Appl. Mater. Interfaces 2015, 7, 28128–28133. (40) Deng, W.; Xu, X.; Zhang, X. J; Zhang, Y. D.; Jin, X. C.; Wang, L.; Lee, S. T.; Jie, J. S. Organometal Halide Perovskite Quantum Dot Light-Emitting Diodes. Adv. Funct. Mater. 2016, 26, 4797– 4802. (41) Li, G. R.; Tan, Z. K.; Di, D.; Lai, M. L.; Jiang, L.; Lim, J. H. W.; Friend, R. H.; Greenham, N. C. Efficient Light-Emitting Diodes Based on Nanocrystalline Perovskite in a Dielectric Polymer Matrix. Nano Lett. 2015, 15, 2640–2644. (42) Li, J. Q.; Bade, S. G. R.; Shan, X.; Yu, Z. B. Single-Layer Light-Emitting Diodes Using Organometal Halide Perovskite/Poly(ethylene oxide) Composite Thin Films. Adv. Mater. 2015, 27, 5196–5202. (43) Kim, Y. H.; Cho, H.; Heo, J. H.; Kim, T. S.; Myoung, N.; Lee, C. L.; Im, S. H.; Lee, T. W. Multicolored Organic/Inorganic Hybrid Perovskite Light-Emitting Diodes. Adv. Mater. 2015, 27, 1248–1254. (44) Yu, J. C.; Kim, D. B.; Baek, G.; Lee, B. R.; Jung, E. D.; Lee, S.; Chu, J. H.; Lee, D. K.; Choi, K. J.; Cho, S.; Song, M. H. HighPerformance Planar Perovskite Optoelectronic Devices: A Morphological and Interfacial Control by Polar Solvent Treatment. Adv. Mater. 2015, 27, 3492–3500. (45) Li, J. H.; Cai, F.; Yang, L. Y.; Ye, F. H.; Zhang, J. H.; Gurney, R. S.; Liu, D.; Wang, T. Sodium Bromide Additive Improved Film Morphology and Performance in Perovskite Light-Emitting Diodes. Appl. Phys. Lett. 2017, 111, 053301. (46) Aygüler, M. F.; Weber, M. D.; Puscher, B. M. D.; Medina, D. D.; Docampo, P.; Costa, R. D. Light-Emitting Electrochemical Cells Based on Hybrid Lead Halide Perovskite Nanoparticles. J. Phys. Chem. C 2015, 119, 12047–12054. (47) Yu, J.; Kim, D. B.; Jung, E. D.; Lee, B. R.; Song, M. H. HighPerformance Perovskite Light-Emitting Diodes via Morphological Control of Perovskite Films. Nanoscale 2015, 8, 7036–7042. (48) Yuan, M.; Quan, L. N.; Comin, R.; Walters, G.; Sabatini, R.; Voznyy, O.; Hoogland, S.; Zhao, Y.; Beauregard, E. M.; Kanjanaboos, P.; Lu, Z.; Kim, D. H.; Sargent, E. H. Perovskite Energy Funnels for Efficient Light-Emitting Diodes. Nat. Nanotechnol. 2016, 11, 872.

11



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

Page 12 of 14

TOC figure

12


Page 13 of 14

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


49x28mm (300 x 300 DPI)



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

119x97mm (300 x 300 DPI)


Page 14 of 14

skite CH3NH3PbBr3 Thin Films for High-Performance Green Light

Recommend Documents