Controlling Crystallization of All-Inorganic Perovskite Films for Ultralow

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Controlling Crystallization of All-Inorganic Perovskite Films for Ultralow-Threshold Amplification Spontaneous Emission Zi-Jun Yong, Yang Zhou, Ju-Ping Ma, Ya-Meng Chen, JunYi Yang, Ying-Lin Song, Jing Wang, and Hong-Tao Sun ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b10863 • Publication Date (Web): 11 Sep 2017 Downloaded from http://pubs.acs.org on September 11, 2017

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ACS Applied Materials & Interfaces

Controlling Crystallization of All-Inorganic Perovskite Films for Ultralow-Threshold Amplification Spontaneous Emission Zi-Jun Yong,†,○ Yang Zhou,†,○ Ju-Ping Ma,†,○ Ya-Meng Chen,†,○ Jun-Yi Yang,‡ Ying-Lin Song, ‡ Jing Wang,§ and Hong-Tao Sun*,† †

College of Chemistry, Chemical Engineering and Materials Science, State and Local Joint

Engineering Laboratory for Novel Functional Polymeric Materials, Soochow University, Suzhou 215123, PR China ‡

College of Physics, Optoelectronics and Energy, Soochow University, Suzhou 215123, PR

China §

Ministry of Education Key Laboratory of Bioinorganic and Synthetic Chemistry, State Key

Laboratory of Optoelectronic Materials and Technologies, School of Chemistry and Chemical Engineering, Sun Yat-sen University, Guangzhou 510275, PR China

KEYWORDS: lead halide perovskites; CsPbBr3; amplification spontaneous emission; crystallization; luminescence; polyethylene glycol

ABSTRACT: All-inorganic lead halide perovskites have gained considerable interest owing to their potential applications in an array of high-performance optoelectronic devices. However, producing highly-luminescent, nearly pinhole-free, all-inorganic perovskite films through a 1

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simple solution process remains challenging. Here, we provide a detail investigation of the crystallization control of inorganic perovskite films fabricated by a one-step spin-coating process. Our results reveal that the coating temperature in the fabrication process is of paramount importance in influencing perovskite crystallization, and that lowering the coating temperature and fine stoichiometry modification of the precursors favor the suppression of trap states in CsPbBr3 perovskite films. A broad range of experimental characterizations help us identify that non-synergistic assembly of solutes, resulting from poor diffusion capability of inorganic salts, is the dominant cause for the inhomogeneous element distribution, low luminescence yield, and poor surface coverage of the resulting films. Importantly, we find that polyethylene glycol can also be used for tailoring the crystallization process, which enables the attainment of high-quality CsPbBr3 films with a maximum luminescence yield of ~30%. Finally, we demonstrate that amplification spontaneous emission with an ultralow threshold can be readily accomplished by using the developed film as an emissive component. Our findings provide deep insights into the crystallization control of CsPbBr3 perovskite films, and establish a systematic route to high-quality all-inorganic perovskite films, paving the way for widespread optoelectronic applications.

1. Introduction Lead halide perovskites have been established as an important family of promising semiconducting materials for optoelectronic devices such as photovoltaic cells,1-5 light emitting diodes,6-8 and photodetectors.9-11 Hybrid lead halide perovskites are particularly attractive because of the low-cost fabrication process and the excellent optoelectronic properties including 2


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large absorption coefficient, long carrier diffusion length, tunable band gap, and high photoluminescence quantum yield (PLQY).12,13 However, hybrid lead halide perovskites suffer from poor stability owing to sensitivity to moisture, O2, light, electric field and thermal stress, greatly hindering their practical applications.14 While several groups have reported success toward significantly improved long-term stability of hybrid perovskite solar cells,15,16 the intrinsic instability of hybrid perovskites suggests that the ultimate avenue for improving operational stability is to replace volatile organic components entirely by inorganic cations such as cesium, since inorganic perovskites such as CsPbI3 and CsPbBr3 are compositional stable up to their melting points (~ 500 ºC).17-20 Therefore, recent efforts have focused on the construction of solar cells using inorganic perovskite nanocrystal films, achieving efficiencies over 10%.19,20 Besides photovoltaic application, the use of inorganic perovskites as light emitters has also received considerable interest since the pioneering work of Protesescu et al.,21 because of extremely high PLQYs (50-90%) of the synthesized CsPbX3 (X=halide ions) nanocrystals. High-efficiency PL,22-40 low-threshold lasers and ASE,36, 41-47 and electroluminescence have been realized by using a broad range of cesium-based nanostructures.48-50 The construction of nanostructure-derived devices involved multiple steps, typically including synthesis of nanocrystals, ligand exchange, and coating them on substrates to produce uniform films. This lengthy and time-consuming route not only increases the fabrication cost, but also introduces more uncertainty in the reproducibility of the device performance. Obviously, to circumvent this issue, directly producing highly-luminescent CsPbX3 crystalline films by means of a simple, reproducible strategy is urgently needed. Spin-coating, one of the cheapest film production methods, has been adopted to produce CsPbX3 films, but yielded films exhibiting modest PLQYs with incomplete surface coverage.51-53 Although recent work has reported that spin-coating 3



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CsPbBr3 precursor solution could produce film with a high PLQY, unfortunately, the effect of a series of fabrication factors on the crystallization and film formation was greatly underestimated,53 rendering it extremely difficult to be reproduced. Despite intensive work and abundant knowledge accumulated regarding the fabrication of high-quality hybrid perovskite films over the past several years,54 at present one-step production of highly-luminescent CsPbX3 crystalline films has indeed been hampered by the difficulty in the control of the intractable crystallization process of spin-coated inorganic precursor solutions on selected substrates. Here, we present a detailed exploration of the crystallization control of inorganic perovskite films fabricated by a one-step spin-coating process. We find that the temperature used prior the thermal treatment in the fabrication process, commonly described as ‘room temperature’ and regarded as an insignificant factor, plays a paramount role in affecting the crystal growth kinetics and film morphology. We demonstrate that suppressing the formation of trap states can be accomplished by lowering the spin-coating temperature and fine stoichiometry modification of the precursors. Additionally, we first show that, akin to hybrid perovskites, CsPbBr3 crystalline films also feature inter-grain luminescence heterogeneity. Experimental characterizations help us identify that non-synergistic assembly of solutes is the dominant reason for the low PLQY and poor surface coverage of the resulting films. We also reveal that polyethylene glycol (PEG) can be used for tailoring the crystallization process, which enables the attainment of nearly pinhole-free CsPbBr3 films with a maximum PLQY of ~30%, as a consequence of fine control of nucleation and crystal growth. Benefiting from this advance, we demonstrate that amplification spontaneous emission (ASE) with a low threshold comparable to that of nanocrystal-casted films can be accomplished.

4


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2. Experimental Section 2.1 Materials and Synthesis CsBr (99.999%, Alfa), PbBr2 (99%, Aladdin), PEG (M.W. 20000, Alfa) and dimethyl sulfoxide (DMSO) (+99.7%, Acros) were used without further purification. For the fabrication of PEG-absent CsPbBr3 perovskite films, the perovskite precursor solutions were prepared by mixing CsBr and PbBr2 with different molar ratios in DMSO, which were stirred for 4 h at a constant temperature (21 ºC, 23 ºC, 25 ºC, or 30 ºC) in a nitrogen-filled glovebox. Note that the temperature in the glovebox was strictly controlled by tailoring room temperatures and putting ice bags atop the glovebox. The temperature at different position of the glovebox was monitored by a probe thermometer, which confirmed that the temperature difference is less than 0.5 ºC lower or higher than the set temperature. The glass substrates were cleaned and treated using a UV Ozone (UVO) cleaner for 20 min. Finally, the 15 wt% precursor solutions were spin-coated at 3000 rpm for 58 seconds, and then the films were transferred to a hot plate and annealed at 70 ºC for 15 min. For the fabrication of PEG-bearing CsPbBr3 perovskite films, PEG were dissolved in DMSO with concentrations of 10, 20, 30 and 40 mg ml-1, which was stirred for 20 min at 50 ºC. Then the CsPbBr3 precursor solution (15.57 wt%, 150 µl) and PEG solution (265 µl) were premixed to form a transparent precursor solution (CsPbBr3: 6.25 wt%) with different PEG weight ratios ranging from 0.58 wt% to 2.27 wt%. The mixed solution was stirred for 20 min and spin-coated onto the glass substrate at 3000 rpm for 58 seconds in a nitrogen-filled glovebox. Additionally, the high concentration precursor solution was prepared by using PEG solution (12.5, 15, and 20 mg ml-1) as solvent, which contains CsPbBr3 with a ratio of 15 wt% and PEG ratios ranging from 5



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1.12 wt% to 1.79 wt%. Then the solution was stirred for 6 h at 21 ºC. Perovskite thin films were prepared by spin-coating the precursor solution at 4500 rpm for 58 seconds, followed by transferring the as-coated films to a hot plate and subsequent annealing at 70 ºC for 5 min to form CsPbBr3 perovskites. We emphasize that the temperature inside the glovebox is set at 21 ºC for the production of all these PEG-containing films. 2.2 Characterization An infrared thermal imaging camera (Fotric 220) with an accuracy of ± 2 ºC was used to monitor the temperature change of the film during the thermal treatment process. The X-ray diffraction (XRD) patterns were taken using a Bruker D8 ADVANCE diffractometer with Cu Kα radiation (λ=1.54056 Å). Scanning electron microscopy (SEM) images were acquired on an S-4700 microscopy (Hitachi, Japan) operating at 10 kV. The concentration of cesium and lead ions were determined by using inductively coupled plasma-mass spectrometry (ICP-MS) (Thermo ELEMENT 2). The PL spectra were taken by using a monochromator (iHR550, Horiba) equipped with a photomultiplier tube (PMT) (Hamamatsu, R928). The spectra response of the detection system was corrected using a standard sample. Emission quantum yields were acquired using an integration sphere incorporated into a spectrofluorometer (FLS980, Edinbursh Instruments). For temperature-dependent PL measurement, the samples were cooled by a closed-cycle He cryostat and the luminescence was coupled into a monochromator (iHR550, Horiba) equipped with a PMT (Hamamatsu, R928). The absorption spectra were taken by a double-beam UV-Vis-NIR spectrophotometer (Cary 5000, Agilent) equipped with an integrated sphere. Time-resolved PL measurements monitored at 526 nm were acquired on a Lifespec II setup (Edinburgh Instrument, UK) with the excitation of a picosecond-pulsed 406 nm laser. The 6


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X-ray photoelectron spectroscopy (XPS) measurements were carried out with a K-Alpha+ spectrometer (Thermofisher Scientific). The sample was analyzed using a micro-focused, monochromated Al Kα X-ray source (400 um spot size). The cathodoluminescence (CL) measurement was carried out by a Gatan Mono-CL system attached to an SEM (FEI Quanta 400F). For the ASE measurements, a femtosecond amplified laser system was employed as the pumping source. The femtosecond pulses were generated by a regenerative amplified ytterbium doped fiber femtosecond laser (Light Conversion PHAROS-SP) that produces pulses centered at 400 nm with a repetition rate of 1 kHz and an FWHM of 190 fs. The laser beam was focused into a stripe via a cylindrical lens with a focal length of 75 mm. Laser power was measured by using an energy detector (Rjp-765 energy probe) linked to an energy meter (Rj-7620, Laser Probe). The emission from the edge of the film was vertically collected by a spectrometer (SP2500, Princeton Instruments) equipped with an electrically cooled charge-coupled device (Pixis 100, Princeton Instruments). The integration time for this measurement is 200 ms.

3. Results and Discussion Controlled Growth of CsPbBr3 Perovskite Films by Tailoring the Pre-Annealing Temperature.

The

crystallization

process

consists

of

two

major

successive

events, nucleation and crystal growth. Heterogeneous nucleation, nucleation with the nucleus at a surface, is typically much faster than homogeneous nucleation because the nucleation barrier is much lower at a surface,55 which can be viewed as the first major step experienced by spin-coated, supersaturated perovskite precursors. The typical fabrication process of perovskite films involves spin-coating the precursor solution at a given temperature and treating the film 7



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containing solvent on a hot plate at a certain temperature. In most cases, the films were spin-coated at room temperature, a parameter generally not strictly defined. We note that the change of this parameter greatly affects the thermal history experienced by spin-coated films, as shown in Figure 1a, which is hypothesized to conspire towards a difference in the properties of the resulting films. We start our investigation by spin-coating stoichiometric perovskite solution and subsequent thermal treatment at 70 ºC. The temperature used before thermal treatment of films is referred to as the pre-annealing (PA) temperature. Interestingly, we find that the films demonstrate PLQYs that are closely related to this temperature (Figure S1, Supporting Information), suggesting that it plays a critically important role in influencing the crystallization process. To further improve the PLQY and to guarantee the reproducibility of the fabrication, we next tried to modify the stoichiometry of the precursor solution and strictly controlled the preparation parameters, as depicted in Figure 1b. The solutions has different molar ratios of CsBr to PbBr2 (CsBr:PbBr2=0.95, 1, 1.05, 1.1, and 1.2), and the PA temperature is set at 21 ºC. Strikingly, we find that the gradual increase in CsBr molar proportion from 0.95 to 1.05 in the precursor leads to an increase of PLQY from 0.14% to 12.22%, and further increasing the ratio cannot affect PLQYs significantly. We chose the precursor with a CsBr:PbBr2 proportion of 1.05 as a model precursor to examine the details of the fabrication process. It is observed that varying the PA temperature significantly affects the evaporation rate of DMSO. At 21 ºC, after completing spin-coating, the film is colorless and transparent (see video 1, Supporting Information); putting the as-coated film onto a hot plate at 70 ºC changes its color from colorless to yellow in less than two seconds, and the resulting films show green luminescence under 365 nm excitation (see video 2, Supporting Information). In marked contrast, at 30 ºC, the film slowly changes its color to yellow during 8


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spin-coating (see video 3, Supporting Information), implying the formation of CsPbBr3 phase upon the gradual removal of DMSO. Spin-coating at intermediate temperatures (23 and 25 ºC) leads to slower evaporation of DMSO with respect to that at 30 ºC (see videos 4 and 5, Supporting Information). The XRD patterns indicate that the thermally-treated films comprise orthorhombic CsPbBr3 phase (Figure 1c), as evidenced by sharp peaks at 15.45º, 21.87º, 30.94º, 34.69º, 38.03º and 44.29º that correspond to (101), (121), (202), (141), (321) and (242) crystallographic planes, respectively. The films produced at 21, 23, 25 and 30 ºC are denoted the 21, 23, 25 and 30 films, respectively. The crystallite size of CsPbBr3 crystallites calculated from the XRD data applying a Scherrer analysis is 51, 47, 42 and 31 nm for the 21, 23, 25, and 30 films, respectively. Apparently, the PA temperature could affect the orientation of the obtained films; compared with the 25 and 30 films, the 21 and 23 cousins are preferentially oriented with the (101) planes. This can be attributed to the variation of the thermal history experienced by these films. As is observed, for the 30 film, the CsPbBr3 phase could form before the completion of the spin-coating. However, for the 21 film, this phase occurs at the beginning of the thermal treatment. Therefore, the growth temperature of the CsPbBr3 crystallites for the 21 film is higher than that of the 30 film, thus resulting in larger crystallite size. Additionally, we note that the as-spin-coated film at 21 ºC can also convert into CsPbBr3 phase via slow evaporation of DMSO, but features the absence of preferential orientation with the (101) planes owing to the lack of energy needed for the assembly of constituents for the CsPbBr3 phase (Figure S2). Interestingly, we find that varying the PA temperature not only influences the crystallographic orientation of the resulting films, but also induces a regular change of the morphology. The films fabricated at 21 and 23 ºC comprise interconnected, irregularly-shaped grains and feature large uncovered regions (Figure 1d and Figure S3a), while those at 25 and 30 9



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ºC are covered by fusiform grains with small pinholes (Figure 1e and Figure S3b); increasing PA temperature causes better film coverage, although the films’ thicknesses are virtually identical (Figure S4). Because the grain sizes are much larger than those of the crystallite size, we conclude that all grains consist of some small crystallites. These observations, combined with the decreased intensity of diffraction peaks as the rise of PA temperature, point to lower crystallinity of the films at higher PA temperatures, thus underscoring the importance of this previously-overlooked factor.

Figure 1. (a) Temperature curves showing the thermal history experienced by the films. (b) Schematic demonstration of the experimental setup. (c) XRD patterns of the 21, 23, 25, and 30 films. (d) SEM image of the 21 film. (e) SEM image of the 30 film.

Optical Properties of CsPbBr3 Films. We investigated the optical characteristics of as-prepared 10


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CsPbBr3 films by performing absorption spectroscopy, steady-state and time-resolved PL measurements. Analysis of the absorption profiles shows that the exciton absorption edge slightly shifts to high energy from 2.335 to 2.345 eV as the increase of the PA temperature from 21 to 30 ºC (Figure S5), accompanied by notable weakened PL emissions and change of PLQYs from 12.22% to 2.03% (Figure 2a). We also measured the temperature-dependent PL for the 21 and 30 films (Figure S6), and the integrated PL emission intensity as a function of temperature for both films are plotted in Figure 2b. Assuming that the decrease in PL intensity is only due to the increase in thermal dissociation rate of exciton at higher temperatures, the exciton binding energy can be obtained by fitting the temperature-dependent PL intensity using the following equation, IT =

⁄

(1)

in which I0 is the intensity at 0 K, Eb the excition binding energy, and kB the Boltzmann constant.[56] Clearly, a relatively larger exciton binding energy of 34.02 meV for the 21 film is obtained. We note that the fitted curves for both cases display a deviation from the experimental data, which will be discussed below. The luminescence decays are fitted by a bi-exponential function, and the average lifetimes were calculated to be 3.53, 2.85, 2.59, and 2.01 ns for the 21, 23, 25, and 30 films, respectively (Figure 2c, Table S1). Collectively, these observations suggest that using a low PA temperature is capable of inhibiting the formation of structural and chemical defects, thus resulting in stronger excitonic emission. Additionally, the 21 film exhibits relatively larger exciton binding energy than the 30 film, indicating the excitonic level of the 30 film is closer to the conduction band. Therefore, the exciton absorption energy for the 30 film is higher than the 21 film, thus causing the blueshift of the absorption edge from 21 to 30 oC (Figure S5). 11



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With resolution much better than that achievable with optical microscopy, CL microscopy has been recognized as a powerful tool to investigate the nanoscale optical properties of a wide variety of functional materials. To correlate the physical structure of CsPbBr3 films with their optical properites, we took CL measurement. Secondary electron (SE) and CL images for the films at different PA temperatures are illustrated in Figure S7, and the overlaid SE/CL images are displayed in Figures 2d-g. Note that all images were taken under the same measurement conditions. Obviously, all films display surprisingly strong inter-grain CL heterogeneity. For the 21 film, the irregularly-shaped small grains show much stronger luminescence signal, forming a series of bright spots, whereas the wire-shaped regions are comparably dark (Figure 2d). For those at higher PA temperatures, most regions show comparable CL, accompanied by a limited amount of dark cousins (Figures 2e-g). This suggests that, similar to hybrid perovskites, these films probably suffer from compositional heterogeneity.[57] The CL results signify that the films consist of crystallites with different quality, thus causing that the fitted result shows a deviation from the experimental values as displayed in Figure 2b. Combined with the facts of a small PLQY and short PL lifetime for the 30 film, we could conclude that higher PA temperatures create a higher density of defect states and mediocre electronic quality. Importantly, we also note that the PA temperature can also affect the properties of other inorganic perovskite films. For instance, the PLQY of CsPb(Br0.8Cl0.2)3 film fabricated by directly spin-coating the precursor could increase from 0.77% to 1.89% when the PA temperature decreases from 30 ºC to 21 ºC, implying the importance of strict control over the fabrication factor. On the basis of these results, we recognized that producing films homogeneously composed of the highly-luminescent grains could offer the promise of an improvement of overall PLQYs. To achieve this, however, a deeper understanding of the underlying issues affecting the film quality is essential. 12


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Figure 2. (a) PL spectra and PLQYs of the CsPbBr3 films produced at different PA temperatures. (b) Temperature-dependent PL intensity for the 21 and 30 films in the temperature range of 10-290 K. Black and red lines are the fitted curves. (c) Time-resolved PL decays of the films monitored at 526 nm. (d-g) Overlaid SE/CL images for the (d) 21, (e) 23, (f) 25, and (g) 30 films.

Crystallization Process of CsPbBr3 Films. Previous work has shown that directly spin-coating precursors followed by thermal treatment can yield highly-luminescent hybrid perovskite films; the precursors used were organic halides and lead halides. By contrast, for the fabrication of CsPbBr3 perovskite films, CsBr and PbBr2 were used. The major difference between the two cases is that inorganic chemicals are stable enough at the annealing temperature used, whereas organic chemicals such as methylammonium halide and formamidinium halide can slowly decompose into volatile species during annealing, providing opportunities for full interaction between volatile, decomposed products and non-volatile inorganic species. Therefore, thermal treatment of hybrid perovskite films yields perovskite phases based on self-assembly of precursors as well as the passivation of structural defects by means of vapor-phase diffusion of 13



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decomposed species. Unfortunately, such a kind of passivation is extremely difficult to be achieved once the inorganic species are inhomogeneously distributed in the as-coated films, in view of their involatile nature at 70 ºC, although there are chances to crystallize into CsPbBr3 crystals via reaction of the most intimate Cs, Pb, and Br atoms upon removal of DMSO. Based on this consideration, we presume that the key reason for inefficient PL of CsPbBr3 films is the inhomogeneous distribution of inorganic species at the atomic scale. Examining the precipitation dynamics of inorganic species from as-spin-coated films with gradual DMSO loss is challenging, but is of vital importance for the understanding of nucleation and crystal growth of perovskite phases. It is noted that the maximum weight ratio of CsBr and PbBr2 with a proportion of 1.05 in DMSO can reach 21.4%. We thus prepared two precursor solutions by mixing CsBr and PbBr2 in DMSO with a weight percent of 24.3% and 26.1% (i.e., both are beyond the critical solubility limit), which can be viewed as 45% and 50% DMSO loss of the original 15 wt% precursor solution, respectively. After vigorous stirring and then letting the solution stand, we separated the insolute powders, and measured their compositions. The molar ratio of Cs to Pb determined by ICP-MS shows that the Cs/Pb ratio in the precipitate is much larger than 1.05 (Table S2), the value in the 15% precursor solution. Interestingly, a rise of DMSO loss results in a higher Cs/Pb ratio. Considering the complicated distribution of solutes in the DMSO solvent,58 we highlight that this analysis may not be virtually equated with the exact process experienced by solutes when DMSO evaporates. However, this observation, to a certain extent, indicates that the Cs solute tends to separate out from DMSO much faster than the Pb solute. To study the composition and possible chemical changes of CsPbBr3 perovskite films prepared at different PA temperatures, we performed high-resolution XPS measurement. The 14


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survey spectra show strong peaks of Br (~68 and 70 eV), Pb (~138 and 143 eV) and Cs (724 and 738 eV) (Figure 3a, 3b, Figure S8). As increasing the PA temperature, the Cs/Pb ratio slightly decreases from 1.49 to 1.42, all of which are larger than 1.05. Systematic deconvolution of Cs 3d, Pb 4f, and Br 3d spectra into summations of Gaussian-Lorentzian curves reveals that the nature of chemical bonds in the CsPbBr3 films with various PA temperatures are almost identical, since the peak position and FWHM do not show a significant difference (Figure S9). It is necessary to underscore that XPS provides the element information at a sample surface with a thickness in the range of 1-10 nm and with a spot diameter of 400 um, thus merely reflecting average elemental information. The strong deviation from the expected stoichiometric value of the CsPbBr3 phase implies that non-synergistic assembly of Cs, Pb, and Br occurs when DMSO gradually evaporates, and more Cs atoms accumulate at the surface region of the film. We, therefore, presume the occurrence of Cs vacancies in the CsPbBr3 phase, which can be viewed as the direct consequence of non-synergistic assembly of constituents and as the dominant reason for the low PLQYs. Recently, Nayak and coworkers provided convincing experimental evidence for the existence of lead halide colloids in the hybrid perovskite precursor solution, which could exist in the inorganic precursor solution, considering the successful growth of CsPbBr3 single crystal via the discipline gleaned from hybrid systems.58 Interestingly, we find that codissolving CsBr with PbBr2 can boost dissolvability of CsBr with respect to singly dissolving CsBr into DMSO. For instance, we find that CsBr cannot singly dissolve into DMSO at a weight ratio of 6.26 wt%, but a combination of 6.26 wt% CsBr with 9.88 wt% PbBr2 yields a clear DMSO solution after stirring for 4 h. This suggests that, beyond single-ion species, a complicated distribution of Cs, Pb, and Br exists in the DMSO solution. 15



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On the basis of the above observations, we propose the plausible mechanism for the formation of CsPbBr3 films and explain why varying PA temperatures is capable of controlling crystallization process. As shown in Figure 3c, at a low environmental temperature, e.g., at 21 ºC, only a tiny amount of DMSO evaporates during the spin-coating stage, which cannot shift the concentration of ions into the supersaturation regime. After quickly transferring it onto the hot plate, the as-spin-coated film undergoes a quick temperature rise to ca. 68 ºC, 2 ºC lower than the set temperature, 70 ºC, owing to the convective heat flow. At this stage, the DMSO solvent evaporates in less than 2 seconds, resulting in the self-assembling of dissolved species, probably via the diffusion of Cs into Pb-Br colloids, followed by subsequent assembly of Pb, Cs and Br ions. Before totally removing DMSO, the solvent environment with an average temperature over 30 ºC provides higher diffusion energies of constituents for the CsPbBr3 phase (Figure 1a). However, once DMSO is depleted, the solid phase diffusion becomes extremely difficult. The existence of CL heterogeneity in the final films may be caused by the preferential growth of CsPbBr3 phase at certain sites and/or crystal nuclei. In marked contrast, increasing the environmental temperature to 30 ºC results in quick evaporation of DMSO and induces formation of CsPbBr3 phase after spinning (Figure 3d). That is, almost all DMSO is depleted after spin-coating, thus rendering it hard to further assemble Pb, Cs and Br at the following thermal treatment stage. Obviously, in this case the average temperature corresponding to the self-assembly process is ca. 30 ºC (i.e., it occurs at the spin-coating stage), thus leading to lower diffusion energies of constituents of the CsPbBr3 phase and resulting in small crystallite sizes in comparison to that using the PA temperature of 21 ºC, as confirmed by XRD (Figure 1c). Additionally, the gradual loss of DMSO during spinning at 30 ºC causes progressive precipitation of species. This is favorable for keeping more solutes on the substrate owing to the 16


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increased surface roughness, resulting in the improved film coverage (Figure 1e). Collectively, the higher PA temperature could create a high density of defect states, thus leading to a low PLQY. The combined experimental evidences also strongly imply that bright CsPbBr3 crystallites tend to be created randomly at some preferential sites of spin-coated films, in particular for those at the PA temperature of 21 ºC.

Figure 3. (a, b) XPS survey spectra of CsPbBr3 perovskite films produced at PA temperatures of (a) 21 and (b) 30 ºC. (c-d) Schematic illustrations of the formation process for the (c) 21 and (d) 30 films.

PEG-Assisted Crystallization Control for High-Quality CsPbBr3 Films. The above 17



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investigation suggests that producing high-quality CsPbBr3 films lies in the creation of more bright crystallites interconnected with each other to yield pinhole-free films, which is synonymous with assembling of Pb, Cs and Br in a well-controlled manner. Blending polymers (e.g., poly-(ethylene oxide), polyimide, and PEG) and precursors have been used to change the crystallization process of hybrid and all-inorganic perovskite films, yielding improved film coverage and PLQYs.59,60 We thus reasoned that PEG-assisted crystallization control can potentially be unified with the PA temperature method described above, which may be favorable for the production of high-quality CsPbBr3 films with respect to that without PEG. The addition of PEG into the precursor solution aims at using its long chain to modify the crystallization dynamics, and using a low PA temperature of 21 ºC is to increase the diffusion energies of constituents to form crystallites with a low density of defects when subsequently annealed at 70 ºC. We term this method as a PEG-assisted cold-casting approach. Using a 6.25 wt% precursor solution as a starting point, we investigated the introduction of PEG on the influence of morphology, surface coverage as well as PLQYs. As illustrated in Figure 4a and Figure S10, the morphologies of the films vary with the amount of PEG introduced. The resulting films are referred to as 6.25%-x% (x=0.58, 1.15, 1.71, and 2.27), where x% represents the weight percent of PEG in the precursor solution. It is noted that all films comprise a large amount of uniform-sized grains, with improved film coverage and PLQYs with respect to PEG-absent counterparts; the maximum PLQY of the PEG-bearing film could reach 30.63% (Table S3). When the PA temperature increases to 25 ºC, the PLQY of the 6.25%-1.15% film drops to 21.44%, signifying the important role of the PA temperature. Interestingly, the XRD pattern indicates that the treated films are composed of orthorhombic CsPbBr3, as evidenced by sharp peaks at 21.73º and 43.97º that can be assigned to (121) and (242) planes, 18


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respectively (Figure 4b). Compared with PEG-absent films, it is apparent that the orientation of CsPbBr3 crystallites is significantly changed, which could be attributed to the passivation effect of PEG molecules. The size of CsPbBr3 crystallites calculated from the XRD data is 37 nm, in a good agreement with the result of SEM measurement (Figure 4a), suggesting that PEG molecules hinder the particles to merge together. Additionally, the XPS result reveals that the Cs/Pb ratio in the PEG-bearing film is 1.17 (Figure 4c), much smaller than those of PEG-absent films. All these facts suggest that PEG molecules is capable of affecting crystallization dynamics of CsPbBr3 films, and to some extend, can make the Cs and Pb atoms assemble in a synergistic manner. We next increased the weight ratio of CsPbBr3 to 15 wt% in the precursor solution to examine its effect on the properties of the films, which were denoted 15%-x% where x% is the PEG weight percent in DMSO from 1.12 to 1.79 wt%. By tailoring the PEG content to an optimal value, the film coverage and thickness is improved with respect to those of the 6.25%-x% films, and the 15%-1.79% film, with a PLQY of 24.78%, is nearly pinhole-free (Figure 4d, Figure S11). Similarly, the XRD pattern of the 15%-1.79% film also shows that the film is preferentially oriented with the (121) planes (Figure S12). The XPS result reveals that the Cs/Pb ratio in the 15%-1.79% film is 1.31 (Figure 4e), higher than that of the 6.25%-1.15% film. This evidences that the synergistic assembly of Cs, Pb and Br becomes hard with increasing precursor ratio. Interestingly, a redshift of PL by 2 nm and an increased FWHM take place when the precursor ratio increases from 6.25wt% to 15wt% (Figure 4f), which can be attributed to the change of crystal quality as evidenced by the change of composition (Figure 4c, 4e). All these observations indicate that PEG incorporation greatly improves the film coverage. Apparently, the viscous PEG polymer used here allows fine control of nucleation and crystal 19



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growth, i.e., it acts as barriers for the diffusion of the precursors, making the solutes only assemble each other in a short range and thus resulting in uniform-sized nanocrystallites capped by PEG molecules in the films, as schematically shown in Figure 4g. That is, PEG induces nanocrystal formation, which then assemble into a film through close packing of nanocrystallites. We can observe that the as-coated film quickly shows green PL under the excitation of 365 nm light after transferring onto a hot plate at 70 ºC, and the PL intensity almost keeps constant as time goes by (see video 6, Supporting Information). This implies that the assembly of Cs, Pb and Br completes in just a few seconds. Furthermore, considering that the film coverage is significantly improved once PEG is introduced, we presume that PEG molecules stabilize the Pb-Br colloidals in the precursor solution; when DMSO is gradually evaporated, the colloidals and single ions precipitate on the substrate, followed by subsequent assembly of intimate Pb, Cs and Br ions into the Pb-Br colloidals that finally form CsPbBr3 crystallites. This causes increased chances for the incorporation of Cs ions into the CsPbBr3 framework, evidenced by the decreased Cs/Pb ratio at the surface of the PEG-bearing film in comparison to PEG-absent films. We also note that the PLQYs of some PEG-bearing films do not increase so much relative to those of PEG-absent films, probably as a result of the increased amount of surface defects resulting from the decreased size of crystallites. However, it is obvious that the delicate compromise between the precursor and PEG concentrations leads to the attainment of uniform, highly luminescent, nearly pinhole-free CsPbBr3.

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Figure 4.

(a) SEM image of the 6.25%-1.15% film. (b) XRD pattern of the 6.25%-1.15% film.

(c) XPS survey spectrum of the 6.25%-1.15% film. (d) SEM image of the 15%-1.79% film. (e) XPS survey spectrum of the 15%-1.79% film. (f) Normalized PL spectra of the 6.25%-1.15% and 15%-1.79% films under the excitation of 407 nm light. (g) Schematic illustration of the formation process for the PEG-bearing CsPbBr3 perovskite films.

ASE from PEG-Bearing CsPbBr3 Perovskite Film. The achievement of uniform, nearly pinhole-free, PEG-bearing CsPbBr3 film via a simple solution process makes it attractive for optoelectronic applications. We show that ASE can be readily achieved from the PEG-bearing CsPbBr3 film under femtosecond laser excitation. The dynamics of optical gain in the film was investigated in a stripe excitation geometry. Emission from the film edge with increasing pump fluence demonstrates an obvious transition from PL to stimulated emission (Figure 5a), as 21



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evidenced by an abrupt increase in output intensity and spectral narrowing. The ASE peak has a width of 6 nm, and is redshifted by 8.5 nm from the PL maximum. Figure 5b plots the emission intensity versus the pumping fluence, from which an ASE threshold of about 3.17 uJ cm-2 can be extracted. Table S4 summarizes the recent reports of ASE employing all-inorganic CsPbBr3 perovskites as gain media. We note that the threshold achieved using our film is lower than that of individual nanowire lasers, and approaches the values reported using CsPbBr3 nanocrystal-derived films as gain media. In sharp contrast, the PEG-absent CsPbBr3 film does not show ASE. This further evidences the high quality of our CsPbBr3 films that benefit from the suppression of electronic defects, uniform crystallite sizes, as well as nearly pinhole-free coverage. We believe that continued optimization of the parameters such as the film thickness, excitation beam size and device structures can potentially decrease the ASE threshold, and that these advantages also render our films attractive for other functional applications such as light emitting diodes and photodetectors.

Figure 5. (a) Emission spectra under increasing femtosecond pulsed excitation fluence from a PEG-bearing CsPbBr3 film (15%-1.35%). (b) FWHMs of the emission spectra (solid triangles) and integrated emission intensity (hollow circles) as a function of pump energy density. 22


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4. Conclusions In summary, we have provided an insight into the crystallization control of inorganic perovskite films fabricated by a one-step spin-coating process. We discovered that the coating temperature plays a decisive role in influencing the crystallization process of CsPbBr3 films, which is of vital importance in guaranteeing reproducibility. We found that suppressing the formation of trap states can be achieved by lowering the spin-coating temperature and fine stoichiometry modification of the precursors. Furthermore, we uncovered that CsPbBr3 perovskite films show inter-grain luminescence heterogeneity. On the basis of thorough characterizations, we identified that non-synergistic assembly of solutes can be viewed as the key reason for the inhomogeneous element distribution, low PLQY, and poor surface coverage of the resulting films. This highlights the importance of fine control of nucleation and crystal growth for inorganic perovskite phases to realize synergistic assembly of solutes. Additionally, we found that PEG molecules can also be used for tailoring the crystallization process, which enables the attainment of high-quality CsPbBr3 films with a maximum PLQY of ~30%. Finally, we demonstrated that ASE with an ultralow threshold can be readily accomplished by using the developed films as emissive components. Our findings provide important information for the control of crystallization of CsPbBr3 perovskite films, and shed light on the effective strategies to improve the quality of films. All the knowledge can be potentially applied to the fabrication of other classes of inorganic perovskites. We envisage that this work could significantly improve the reproducibility of the growth of perovskites and enables the continued advancement of perovskite optoelectronics. 23



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ASSOCIATED CONTENT Supporting Information The Supporting Information is available free of charge on the ACS Publications website. Videos, steady-state and time-resolved PL, XRD, SEM, XPS, SE and CL images, absorption spectra, fitted lifetimes, ICP-MS, PLQYs, ASE results of CsPbBr3 (PDF) AUTHOR INFORMATION Corresponding Author Email: [email protected] Author Contributions ○

Z-J. Yong, Y. Zhou, J.-P. Ma, and Y.-M. Chen contributed equally to this work.

Notes The authors declare no competing financial interests. ACKNOWLEDGMENT This work is supported by the National Natural Science Foundation of China (11574225), Jiangsu Specially Appointed Professor program (SR10900214), Natural Science Foundation of Jiangsu Province for Young Scholars (BK20140336), and a project funded by the Priority Academic Program Development of Jiangsu Higher Education Institutions (PAPD).

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Controlling Crystallization of All-Inorganic Perovskite Films for Ultralow

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