Hole-Injection Layer-Free Perovskite Light-Emitting Diodes - ACS

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Functional Inorganic Materials and Devices

Hole-Injection Layer-Free Perovskite Light-Emitting Diodes Zhifeng Shi, Lingzhi Lei, Ying Li, Fei Zhang, Zhuangzhuang Ma, Xinjian Li, Di Wu, Tingting Xu, Yongtao Tian, Baolin Zhang, Zhiqiang Yao, and Guotong Du ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b07048 • Publication Date (Web): 05 Sep 2018 Downloaded from http://pubs.acs.org on September 5, 2018

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ACS Applied Materials & Interfaces

Hole-Injection Layer-Free Perovskite Light-Emitting Diodes Zhifeng Shi,1 Lingzhi Lei,1 Ying Li,1 Fei Zhang,1 Zhuangzhuang Ma,1 Xinjian Li,*1 Di Wu,1 Tingting Xu,*1 Yongtao Tian,1 Baolin Zhang,2 Zhiqiang Yao,3 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 3

College of Material Science and Engineering, Zhengzhou University, Kexue Road 100,

Zhengzhou 450001, China AUTHOR INFORMATION Corresponding Author *E-mail: (X.L.) [email protected]. *E-mail: (T.X.) [email protected]. KEYWORDS: perovskite, light-emitting diodes, CsPbBr3, vapor evaporation, impact ionization

ABSTRACT: In this study, dual-source vapor evaporation method was employed to fabricate the high-quality CsPbBr3 thin films with a good crystalline and a high surface coverage. Temperature-dependent and excitation power-dependent photoluminescence measurement were

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performed to study the optical properties of the CsPbBr3 material. Further, based on the experimental data, the temperature sensitivity coefficient of bandgap and exciton binding energy were estimated. More importantly, for the first time, we designed and prepared a hole-injection layer-free perovskite light-emitting diode (LED) based on the Au/MgO/CsPbBr3/n-MgZnO/n+GaN structure, producing an intense green emission (~538 nm) with a high-purity. Besides, the device demonstrated a high luminance of 5025 cd/m2, an external quantum efficiency of 1.46%, a current efficiency of 1.92 cd/A, and a power efficiency of 1.76 lm/W. We studied in detail the current-voltage and electroluminescence properties of the prepared device and proposed the hole generation models and the carrier transport/recombination mechanisms to make certain these interesting characteristics. The results obtained would provide a new and effective strategy for the design and preparation of perovskite LEDs.


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INTRODUCTION Recently, the organometal halide perovskites have received considerable attention because of their great application potentials in optoelectronic fields.1–6 For example, by employing a thin CH3NH3PbI3 layer as the absorber, perovskite solar cell has achieved a high conversion efficiency (>22%), approaching the commercial performance of polycrystalline thin films solar cells.7 Besides, the CH3NH3PbX3 materials possess good optical characteristics, thus they are potentially valuable in low-cost and light-weight light-emitting diodes (LEDs) for information displays and lighting. However, the stability of such organic-inorganic perovskite materials is always a matter of criticism, which are the major barriers for future practical applications of such materials.8–11 In this circumstance, all-inorganic perovskite system without any organic cations emerged and was considered as a possible route to address the drawbacks of aforementioned issues. Typically, red, green, and blue perovskite LEDs based on inorganic CsPbX3 nanocrystals have been reported by Zhang et al.12 However, there are still many challenges or potential breakthroughs for perovskite LEDs although a relatively high external quantum efficiency (EQE, 6.27%) has been realized by Zeng’s group.13 Generally, for the spin-coating of CsPbX3 nanocrystals, the photoluminescence (PL) quantum yield of them will suffer from a drastic decrease because of the aggregation of the nanocrysals.14 So, further promotion of the device performance should rest upon the controllable synthesis or preparation of high-quality CsPbX3 materials. It is known to all that the physical properties (optical, morphology, and crystalline characteristics) of perovskites depend strongly on their synthesis approach, and various processing methods have been developed to fabricate high-efficient perovskite LEDs, such as the one-step or two-step spin-coating,15 vapor-assisted solution process,1 and vapor deposition,16,17 By contrast, the vacuum thermal evaporation method is advantageous in obtaining uniform


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perovskite thin films with a large area, and also the study of vacuum-evaporated CsPbX3 as emissive layer for LED fabrications is rare presently. Up to now, all reported perovskite LEDs contain a multilayered heterostructure, involving a perovskite emitter, a p-type hole-providing layer, and a n-type electron-providing layer.18 In such heterostructures, the perovskite LEDs operate following the carrier injection mechanism from n/p-type carrier-injectors. Moreover, in some cases additional blocking or modified layers have to be inserted to optimize the carrier injection process,19,20 involving multiple processing steps in forming a multilayer structure. Thus the structure of perovskite LEDs becomes more complicated, and simultaneously higher requirements for the interfacial control are needed. Consequently, the complexity of fabrication process would increase inevitably. Therefore, development of simplified device structure for perovskite LEDs applications is highly desired in terms of the processing flexibility at the manufacturing stage. Especially for p-type hole injector, increasing efforts are being devoted to search simple alternatives to replace or eliminate unstable organic carrier-providing layer, such as 2,2´,7,7´-tetrakis (N,N´-di-p-methoxyphenylamine)-9,9´spirobifluorene,21

polytriarylamine,22

and

poly(3,4-ethylenedioxythiophen)

polystyrene

sulphonate.4,12,23 If these issues could not be adequately addressed, the production and future applications of such perovskite LEDs will be hindered. Aiming at a simple device architecture and a vacuum-evaporated inorganic perovskite as light emitters, we fabricated for the first time a hole-injection layer-free planar perovskite LED. Dualsource vapor evaporation approach was employed to fabricate the high-quality CsPbBr3 thin films with a good crystalline and a high surface coverage. In spite of the simple device structure without conventional hole injectors, the proposed perovskite LEDs demonstrate a remarkable performance (luminance, 5025 cd/m2; EQE, 1.46%; current efficiency, 1.92 cd/A; power


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efficiency, 1.76 lm/W). Finally, the generation models of holes and the carrier transport/recombination mechanisms were proposed to make certain the interesting characteristics of the devices. RESULTS AND DISCUSSION

Figure 1. (a) Schematic diagram of the preparation methods for the CsPbBr3 thin films (A, B, and C represent the different placement regions of substrates). (b) Photographs of three CsPbBr3 products, including the optical microscope images (under ambient conditions and the UV lamp irradiation) and SEM images. The scale bar in SEM images is 1.0 μm. In this study, inorganic CsPbBr3 thin films were fabricated by a custom-designed chemical vapor deposition system, and the corresponding experimental setup diagram was illustrated in Figure 1a. The detailed procedures for materials preparation were described in the Supporting Information. Figure 1b presents the photographs of three produced CsPbBr3 samples at different placement regions. Maybe due to the non-uniform thickness, three samples display a relatively


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rough surface. Under UV lamp (365 nm) irradiation, three CsPbBr3 samples show bright green emission color across the entire surface. By contrast, sample A possesses the best emission uniformity than that of sample B and C because some emission dark spots existed in two samples. Further, we performed the scanning electron microscope (SEM) measurements to evaluate the effect of sample placement region on the morphology of CsPbBr3 thin films. As shown in the right panes of Figure 1b, some distinct trends were revealed. One can observe that sample A displays an uniform and compact surface with the grain size of about 300–500 nm. Note that some undesired pores exist in the CsPbBr3 thin films, which may result from the limited diffusivity of CsBr and PbBr2 vapor on the substrate surface.1 Besides, the sample A was characterized by a low root-mean-square surface roughness of ~8.76 nm, indicating a relatively smooth surface (Figure S1, Supporting Information). For sample B with a large distance between evaporation sources, the grain consolidation takes effect. The adjacent nano-grains are linked, and the grain boundary is hard to be identified, so does the grain size. However, the amount and size of pores in the CsPbBr3 thin films are substantially increased. Typically, the size of pores distributes in a wide range of 50–200 nm. In the case of a larger distance between evaporation sources, the uniformity and surface coverage of CsPbBr3 thin films deteriorate. The grain size is more uneven with a larger size distribution range of 150–500 nm. More seriously, above 10% area of the sample surface was exposed, impeding their applications as an effective light emitter in LEDs. Obviously, the regular microstructure changes of the CsPbBr3 products result from their different placement regions as other processing parameters were the same for all samples. Although samples A, B and C share the same temperature, their distances away from the hightemperature zone (zone I) is different. The different surface morphology obtained may be related with the different evaporation source onto the substrate surface. On the premise of constant


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growth temperature and reaction pressure, the larger distance between the evaporation source, the more opportunities the CsBr and PbBr2 vapor has to find a site to stay on surface of the substrates for the purpose of a high surface coverage. In reality, added experiments have been performed to prove the above viewpoints. Other two samples (sample D and E) were prepared at larger and smaller distances in low-temperature zone (zone II). As shown in Figure S2 (Supporting Information), there is a regular morphology evolution for five samples with different placement regions. By contrast, sample E possesses the smallest surface coverage with over 50% area of the sample surface exposed. Although the sample D is characterized by a large surface coverage, the CsPbBr3 product presents parasitic diffraction from PbBr2, as displayed in Figure S3 (Supporting Information). This observation indicates that the sample D has an unsatisfying phase purity. Further, the optical properties of five samples were investigated by PL measurements. As shown in Figure S4 (Supporting Information), samples A, B, C, and E present good optical quality with dominant emission at ~535 nm. However, for sample D a rather weak emission was observed, which may be caused by the incomplete chemical reaction process. The above observations indicate that the placement region of substrates plays an important role in the film quality; if the distance between sample and evaporation source is shortened further, the quality of the CsPbBr3 products deteriorates instead. Besides, our experiments also confirm that the reaction pressure and temperature of low-temperature zone (zone II) have great influence on the morphology characteristics of the CsPbBr3 products, and the detailed results can be found in Figure S5 (Supporting Information). Further, we performed the X-ray diffraction (XRD) measurements to investigate the crystalline characteristics of the CsPbBr3 samples. Here, sample A was taken as an example because three samples share almost the same spectra characteristics. As shown in Figure 2a, two characteristic


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diffractions can be observed at 14.89° and 30.35°, which are indexed to the (100) and (200) diffractions of cubic CsPbBr3 (space group pm/ 3 m, a=2.95 Å).24 By using the Debye-Scherrer equation, the grain size was calculated to be ~352 nm, well consistent with the observations in top-view SEM image of CsPbBr3 thin films. The measured XPS results of the produced CsPbBr3 were displayed in Figure 2b, in which Cs, Br, Pb, N, and C elements were detected with a stoichiometric ratio of 1.00:0.93:2.79 for Cs, Pb, and Br elements. To analyze the optical properties of sample A, absorption and PL measurements at room-temperature (RT) were conducted. As presented in Figure 2c, the absorption edge occurs at ~533 nm (dotted blue line), matching well with other studies on solution-processed CsPbBr3 products,15,25 and the corresponding PL spectrum (dotted green line) peaks at ~535 nm with a narrow linewidth (20.6 nm). Moreover, we measured the absolute PL quantum efficiency of sample A using an integrating sphere, and a relatively high value (~19.6%) suggests the remarkable optical features of the resulting CsPbBr3 thin films. Figure S6 (Supporting Information) presents the timeresolved PL spectrum of sample A, and an average lifetime of 54.37 ns was obtained, much higher than those of sample B (46.9 ns) and C (17.8 ns), which may be due to the increased charge-carrier trapping states in the CsPbBr3 thin films (sample B and C). Further, temperature-dependent PL spectra were acquired for sample A, a representative sample, to better understand the optical recombination properties of CsPbBr3 products. Figure 2d shows the corresponding PL spectra on a semilogarithmic scale, and a regular PL evolution was revealed. With the decrease of temperature, the corresponding PL intensity increases accordingly, and the overall emission band monotonically red-shifts. Such red-shift phenomenon was also observed in other CuCl(Br/I) and PbS(Se/Te) semiconductors,26,27 and the possible influencing factors include the thermal expansion of lattice and electron–phonon renormalization. Unlike


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other reports, the PL spectra measured at RT and low temperatures (100–300 K) are characterized by an asymmetric configuration, demonstrating two emission peaks with a competitive relation. Owing to the reason that Peak II, as marked in Figure 2d, can be sustained at a low temperature of 10 K, and always dominates the PL spectra in the investigated temperature region, we therefore attributed it to the free exciton emission. While the emission above the conduction-band edge (Peak I) has rarely reported in previous studies. In theory, Sebastian et al. calculated an energy level state of ~0.23 eV above the conduction-band edge, which was assigned to the bromine vacancy related defect states.28 This prediction coincides well with early density functional theory calculations by Shi et al.29 Therefore, the emission contribution of Peak I is presumably related with the carrier recombination involving the bromine vacancy defect. Besides, it should be pointed out that no structural phase transition appeared during the cooling process because of the absence of phase transition-induced reversal PL shift.30 As the temperature was decreased to 80 K and below, another emission component (Peak III) emerges, and its intensity increases with decreasing temperature. From the changing tendency of the PL spectra observed above, a conclusion can be drawn that there exists three different channels for carrier recombination in the CsPbBr3 thin films. Herein, the emission component of Peak III is likely associated with the trap-mediated exciton radiative recombination. Because of the premature disappearance of Peak III in the temperature-rising process, a relatively smaller binding energy of trapped excitons for Peak III could be deduced. Note that the blue-shift phenomenon for Peak III is in contrast to those in Peak I and II. As we stated above, the bandgap energy of perovskites with changing temperature rests with two key factors, the lattice constantinduced variation of bandgap, and the renormalization of bandgap through the electron-phonon


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interaction. The interplay of both two factors determines the evolution trend of the emission peak, and thus the blue-shift of Peak III with decreasing temperature is likely caused by the different electron-phonon coupling.31 For a better comparison on three carrier recombination channels, three representative PL spectra at 220, 140, and 50 K were put together. As shown in Figure 2e, both two PL spectra at 220 and 140 K can be resolved into two components, Peak I at ~534 nm and Peak II at ~548 nm, respectively. Obviously, with the decreasing temperature, the contribution of carrier recombination channel from Peak II increases gradually. While, at 50 K, the emission component of Peak I disappears and another component of Peak III emerges. The changing trends of three emissions of the CsPbBr3 thin films versus temperature were summarized in Figure 2f.


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Figure 2. (a) XRD data of the CsPbBr3 thin films (sample A) and standard XRD pattern of cubic CsPbBr3. (b) XPS spectrum of sample A. (c) Absorption spectrum and PL spectrum of sample A measured at RT. (d) Temperature-dependent PL spectra of sample A. (e) Gaussian deconvolution of three PL spectra measured at 50, 140, and 220 K. (f) Shift of the emission position of Peak I, II, and III versus temperature. (g) Temperature-dependent PL intensity of the CsPbBr3 thin films. (h) The relationship between the PL intensity and the excitation power of sample A measured at 10 and 300 K, respectively. (i) Three proposed recombination channels for photo-generated carriers in CsPbBr3 thin films.


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Because the CsPbBr3 thin films are always characterized by the stable cubic phase in the temperature-rising process, physical parameters of bandgap and exciton binding energy (EB) are therefore estimated according to the obtained data above. Figure 2f presents the changing trend of the photon energy (Peak II) versus temperature, and an approximately linear relation can be observed with the temperature sensitivity of the photon energy of 0.112 ± 0.003 meV/K. In addition, the integrated PL intensity versus temperature was also plotted to estimate the EB of CsPbBr3, as displayed in Figure 2g, and such temperature quenching phenomenon of the CsPbBr3 thin films can be described by the following expression:32

I (T ) 

I0 E 1  A exp(  B ) KT

(1)

in which I0 was the emission intensity at 0 K, A was a coefficient, and K was the Boltzmann constant. Finally, a value of 40.5 ± 1.2 meV for EB was produced. Note that the relatively higher value for CsPbBr3 thin films than the RT thermal ionization energy supports their potential applications in exciton-related luminescent devices.18,33 Excitation power-dependent PL measurement was further performed to confirm the excitonic emission feature of the CsPbBr3 thin films, and a power law dependence with β = 1.36 (IPL = IEXβ) has been observed at RT (Figure 2h), implying the excitonic emission feature.34 At 10 K, a relatively smaller value (β = 1.15) was obtained, which was because of the weakened thermal quenching effect. Based on the above discussions, we summarized the different recombination channels for photo-generated carriers in CsPbBr3 thin films, as shown in Figure 2i, corresponding to different emission components in PL spectra.


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Figure 3. (a) Schematic of the studied perovskite LED configuration. (b, c) Cross-sectional SEM images of the heterostructure, and (d) the corresponding elemental mapping images. The scale bars in (b) and (c) are 1.0 μm and 50 nm, respectively. For LED fabrications, we designed a multi-layered Au/MgO/CsPbBr3/n-MgZnO/n+-GaN heterostructure (the detailed fabrication procedures were presented in the Supporting Information), and its schematic diagram was shown in Figure 3a. In this structure, the n+-GaN template can be deemed as the electron source layer. The n-type MgZnO thin layer serves as the electron-injection material, without introducing any conventional hole-injection materials, such as NiO, PEDOT:PSS, and spiro-OMeTAD. The cross-sectional SEM image of the heterostructure was presented in Figure 3b, in which a compact architecture can be observed. Figure 3c is a high-magnification SEM image taken from the upper region of Figure 3b (marked by a red square), from which the thicknesses of functional layers were estimated to be MgO (40 nm, red), CsPbBr3 (110 nm, green), MgZnO (140 nm, purple), and n+-GaN (3.0 μm, yellow), respectively. In addition, energy-dispersive X-ray spectroscopy element mapping measurements were conducted on the cross-sectional Au/MgO/CsPbBr3/n-MgZnO/n+-GaN structure by employing Mg, Pb, Zn, and Ga as the detection signals, and the orderly distribution of four elements shown in Figure 3d indicates the well-defined hetero-interfaces of the multi-layered structure.


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Figure 4. (a) Dependences of current density and luminance of the perovskite LEDs on bias voltage. (b) EL spectra measured at different biases. The inset displays the shift of the emission peak versus bias voltage. (c) EQE, current efficiency, and power efficiency versus the bias voltage. The inset presents an emission photo of the device at 10.0 V. (d) Variation in the EQE and current efficiency with different MgO thickness. (e) CIE coordinates of the device at 10.0 V. (f) Statistics of the peak EQE (30 devices) with the same structure showing the reproducibility of the proposed devices. (g) Emission intensity of the perovskite LEDs versus operating time (10.0 V). (h) Device performance of the perovskite LEDs as a function of storage time in air ambient.


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For device measurements, the negative voltage was connected with the In electrode. The current density-voltage curve of the device was presented Figure 4a, and a nonlinear rectifying behavior was observed with a relatively large turn-on voltage (~8.0 V), which may results from the voltage drop on the insulating MgO layer.35 The corresponding current density-voltage curve in a semilogarithmic scale was displayed in Figure S7 (Supporting Information). For conventional perovskite LEDs, a p-i-n heterostructure was always employed, and such perovskite LEDs can be operated with the carrier injection mechanism from n-/p-type carrier-injectors because the injection barriers for electrons and holes are very small. Therefore, the conventional perovskite LEDs are often characterized by a relatively small turn-on voltage. In our case, the device design concept (Au/MgO/CsPbBr3/n-MgZnO/n+-GaN structure) differs from the conventional p-i-n model. The dielectric MgO layer induces large energy barriers at the MgO/CsPbBr3 (ΔEC=3.0 eV) and Au/MgO (ΔEV=3.38 eV) interfaces. Therefore, the proposed device structured with Au/MgO/CsPbBr3/n-MgZnO/n+-GaN can be deemed as a Schottky diode. Thus, the effective carrier injection and transport could not be achieved at a low bias voltage. For the Schottky diode in this case, the transport current is contributed by electrons, including the thermionic current and the tunneling current. Only at a relatively large bias voltage was it possible to realize the above two electron transport processes. Therefore, the proposed perovskite LEDs are characterized by a large turn-on voltage. Herein, the relation between the transport current and applied voltage could be described by the following equation:36 I  k1V exp(

q  V  B )  k2V 2 exp( k3 / V ) k BT

(2)

in which q is the elementary charge, and T is the absolute temperature; k1, k2, and k3 are the variable constants; ψB is the Schottky barrier. η can be derived from the equation


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  q / 4 i d (d represents the thickness of insulating layer, and εi is the dielectric constant). In equation 2, the first (second) part of the formula represents the contribution of the thermionic (tunneling) current. Interestingly and excitingly, the perovskite LEDs can generate an intense green emission under forward bias. The measured luminance of the devices versus applied voltage was displayed in Figure 4a (blue line), and a maximum luminance value of ~5025 cd/m2 can be obtained at 15.0 V. EL spectra of the device at different applied voltages were presented, and an intense green emission at ~538 nm was achieved (10.0 V) with a small linewidth of ~23.9 nm. The corresponding emission photo of the perovskite LEDs (10.0 V) was displayed in Figure 4c, and the light distributes uniformly on the electrode area. With the increase of the bias voltage, the emission intensity gradually increased with a red-shift trend of the peak position, as shown in the inset of Figure 4b. The red-shift phenomenon may be caused by the following two reasons. On the one hand, because of the high-resistance natures of the MgO and CsPbBr3 materials, their energy bands would bend downwards under bias polarity. This behavior will become more and more serious for a higher bias, lowering the effective energy of the emissive photon. Therefore, the corresponding emission wavelength of the device will red-shift slightly with the increase of bias voltage. On the other hand, the proposed device was established on the dielectric Al2O3 substrates with a small thermal conductivity,37 thus the undesirable heating effect should be considered in the proposed device structures, especially for a device operated at relatively high bias. The heating effect would also induce a red-shift of the emission peak, as reported in Niu’s study.38 Even so, the linewidth of the measured EL spectra is not changed obviously with the increase of bias voltage (from 7.5 to 13.0 V), which only broadens from 23.6 to 25.5 nm (Figure S8, Supporting Information). In addition, an important observation is that no parasitic emission


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was detected from the MgZnO electron-injection layer (Figure S9, Supporting Information), indicating that the recombination zone of injected carriers exists primarily in the CsPbBr3 thin layer.39 Besides, other parameters of the perovskite LEDs (EQE and current efficiency) were also measured and the obtained data were plotted in Figure 4c. One can observe that the EQE and current efficiency increase parallelly with the voltage initially and then drop, matching well with the observations in other studies.4,12,15,19,40,41 The best-performing device reaches up an EQE of 1.46%, a current efficiency of 1.92 cd/A, and a power efficiency of 1.76 lm/W. It is noteworthy that the above results are achieved at a current injection of ~1/50 lower than that in other studies, implying a small energy losses from charge injection and transport. Considering the fact that the proposed device structure lacks of a fairly standard hole injector, the remarkable performances of the perovskite LEDs were deservedly associated to the inserted MgO insulating layer. Expectedly, the thin MgO layer plays a very important role in the carrier generation, transport, confinement and recombination processes in the multi-layered heterostructure, and the detailed mechanisms would be discussed later. In our experiments, six devices with different thicknesses (20, 30, 40, 50, 60, and 70 nm) were fabricated for comparative measurement. As shown in Figure 4d, one can see that the device performs best when the thickness of MgO layer is ~50 nm. We consider that the MgO layer with thinner thickness has the advantages of low series resistance, but the weakened carrier-blocking effect and compromised coverage uniformity would induce an increased probability of nonradiative recombination and also an increased leakage current. Similarly, the thicker MgO layer would lead to a higher series resistance, and meanwhile, the probability of hole generation through the impact ionization process would be reduced because this process depends strongly on the electric field on the MgO layer. Figure 4e displays the Commission International de I’Eclairage (CIE)


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coordinates (0.11, 0.71) of the device at 10.0 V. Moreover, the reproducibility of the studied perovskite LEDs was also assessed by fabricating 30 devices for identical measurement. As shown in Figure 4f, the statistics data demonstrate an average EQE of 1.45%, indicating a good reproducibility of the proposed perovskite LEDs. For stability study of the unencapsulated devices, the emission decay of the device was measured in air ambient (25 ºC, 30‒45% humidity). During the measurements, the applied voltage was fixed at 10 V and the emission intensity was recorded intermittently. As displayed in Figure 4g, the unencapsulated devices can continuously operate for 18 h with a 43.9% emission decay. Besides, a 10-day storage of the device in air ambient produces an emission decay of about 25.2% (Figure 4h). The above results confirm the resistance of the proposed all-inorganic structured perovskite LEDs against oxygen and water degradation. The above EL spectra characteristics are similar to the observations in most previous studies, but the device design concept in present case is obviously different from the conventional p-i-n structure. We therefore draw the energy band diagrams of the Au/MgO/CsPbBr3/n-MgZnO/n+GaN structure to interpret the possible emission mechanisms. Figure 5a shows the energy band alignment at zero bias. Under forward bias, the band alignment is significantly changed (Figure 5b). Because the CsPbBr3 and MgO materials in the device are high-resistance, so most of the bias voltage will drop across both two layers, and thus the energy bands of CsPbBr3 and MgO layers would bend downwards. With a low bias, electrons in n-MgZnO and n+-GaN will flow into the CsPbBr3 emitter, and they will accumulate at the MgO/CsPbBr3 interface because of a large offset for electron injection (ΔΕC ≈ 0.3 eV). Observation of green emission from the proposed devices implies that hole carriers have been injected into the CsPbBr3 active layer effectively. To interpret the hole generation behaviors, three possible hole generation and


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transportation pathways were proposed for discussion. Firstly, at a sufficiently high bias, the Fermi level of Au would move downwards, which results in the accumulation of holes at the Au/MgO interface. With a certain probability, the accumulated holes can tunnel into the corresponding levels in CsPbBr3. Through the recombination with injected electrons, a prominent green emission corresponding to the bandgap of CsPbBr3 active layer generates, as illustrated in Figure 5c. Besides, thermal emission of carriers from the defect levels in MgO to CsPbBr3 layer is another important hole transportation process, and such a hole conduction behavior can be ascribed to the Poole-Frenkel effect.42,43 Figure 5d illustrates the Poole-Frenkel conduction process in the MgO layer. Assisted by the external electric field, holes accumulated at the Au/MgO interface have a chance to tunnel into the shallow defect levels, independent of the energy barrier at the Au/MgO interface. Last but not least, through the impact ionization process at high electric field to generate hole carriers in MgO layer must be taken into account.35,44,45 In our case, the electric field intensity in MgO layer is as high as ~106 V·cm–1 for the device operation at >10.0 V. Thus, electrons in the MgO layer would gain much kinetic energy and be greatly accelerated. As a result, the MgO lattice will be impacted, exciting the electrons in the valence-band of CsPbBr3. As displayed in Figure 5e, the generated hole carriers can be swept into the valence-band of CsPbBr3 active layer, generating the green emission consequently by recombining with the confined electrons at MgO/CsPbBr3 interface. From the above discussions, one can recognize that the probabilities of three proposed hole generation and transportation processes increase with the bias voltage. In reality, a lot of repetitive experiments have been conducted, and the I–V and EL characteristics of these tested devices under forward bias were in consonance with the above observations and interpretations.


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Figure 5. (a) Schematic band alignments of the heterostructure (a) in the thermal equilibrium and (b) under forward bias. Generation mechanisms of holes with the positive voltage connected to the Au electrode: (c) Hole tunneling through the MgO insulating layer; (d) Thermal emission of carriers from the defect levels in MgO into the valence band of CsPbBr 3 layer. (e) Generation of electron-hole pairs through the impact ionization process.


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Here, one point worth emphasizing is that we demonstrate a conceptually novel device design for perovskite LEDs without the introduction of conventional hole-injection layer, and the resulting device can be efficiently operated at an ultralow injection current density. However, some other issues are still worth mentioning here. Firstly, we believe that the device performance could be further promoted because the thickness of CsPbBr3 active layer, the bandgap of MgZnO layer determined by Mg content, and the conductive ability of n-MgZnO injector were not optimized at present. On the one hand, a thin CsPbBr3 active layer is appreciated for a weakened reabsorption of the generated photons. On the other hand, n-MgZnO layer with a large electron concentration and a reasonable band alignment with CsPbBr3 would favor for an improved electron injection efficiency. Secondly, although the CsPbBr3 layer is characterized by a uniform surface and a large coverage, the undesirable pinholes, though they are few, still exist with a typical size of ~10 nm, which would inevitably induce other current pathways and even electrical short-circuits. Besides, the previous study has reported that a current channel through such undesirable pinholes may form, and local Joule heating would generate,46 thus the emission action would be quenched substantially, so the premature saturation of EQE and current efficiency is understandable. Thirdly, the PL quantum efficiency (19.6%) of the CsPbBr3 active layer is still lower than many quantum dot films, so more work should be done to optimize the optical quality and crystallinity integrity of the CsPbBr3 thin films. Fourthly, an obvious disadvantage of such hole-injection layer-free device structure is a relatively large turn-on voltage compared with conventional p-i-n structure, which is because of a large potential barrier induced by the insulating MgO layer. Additional experiments have revealed that the turn-on voltage of the device can be lowered by decreasing the thickness of MgO layer, as displayed in Figure S10 (Supporting Information). In terms of the turn-on voltage only, an decreased


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thickness for MgO or other dielectric materials with a relatively smaller bandgap than that of MgO may be a better choice. Fifthly, the generation of hole carriers via the impact ionization process in the Au/MgO/CsPbBr3/n-MgZnO/n+-GaN structure suffers from many natural drawbacks. The dielectric MgO layer could not sustain as the voltage applied on MgO reaches and exceeds its critical breakdown electrical field. Therefore, the selection of dielectric material is essential and deserves more attention not only for a lowered turn-on voltage, but also for device operation under a high excitation level without premature breakdown.47 CONCLUSIONS In conclusion, high-quality CsPbBr3 thin films with a good crystalline and a high surface coverage were successfully synthesized by using the vapor evaporation approach. Further the optical properties of the CsPbBr3 thin films were studied detailedly. Note that no structural phase transition appeared during the cooling process. Except for the exciton related emission, a trapped charge-carrier emission and a bromine vacancy defect related emission were observed. Based on the experimental data, the EB of the CsPbBr3 thin films and temperature sensitivity coefficient of bandgap were derived. More importantly, a hole-injection layer-free perovskite LED based on Au/MgO/CsPbBr3/n-MgZnO/n+-GaN structure was fabricated, and an intense green emission was achieved with the luminance of 5025 cd/m2, EQE of 1.46%, current efficiency of 1.92 cd/A, and a power efficiency of 1.76 lm/W. We studied in detail the I–V and EL characteristics of the device and proposed the hole generation models and the carrier transport/recombination mechanisms to make certain these interesting observations. The above results obtained would provide a new and effective strategy for the design and preparation of perovskite LEDs.


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ASSOCIATED CONTENT Supporting Information The following files are available free of charge. Experimental section, 2D and 3D AFM images of the CsPbBr3 thin films (sample A), effect of substrate placement region on the morphology of the CsPbBr3 thin films, crystallinity properties of the sample D, comparison of the optical quality of five samples with different placement regions, influence of other parameters on the morphology characteristics of CsPbBr3, timeresolved PL decay of three CsPbBr3 thin films samples, current density-voltage behavior of the device, broadening behavior of the EL spectra at different applied voltages, indication of no parasitic emissions from the carrier-providing layers, and current density-voltage behavior of the devices with different MgO thickness. (PDF) AUTHOR INFORMATION Notes The authors declare no competing financial interests. ACKNOWLEDGMENT This work was supported by the National Natural Science Foundation of China (Nos. 11774318, 11604302, 61176044, 11504331 and 51771276), the China Postdoctoral Science Foundation (2015M582193 and 2017T100535), the Science and Technology Research Project of Henan Province (162300410229), the Postdoctoral Research Sponsorship in Henan Province (2015008), the Outstanding Young Talent Research Fund of Zhengzhou University (1521317001), and the Startup Research Fund of Zhengzhou University (1512317003).


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