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Excess Cesium Iodide Induces Spinodal Decomposition of CsPbIBr Perovskite Films 2

Xiangyue Meng, Zheng Wang, Wei Qian, Zonglong Zhu, Teng Zhang, Yang Bai, Chen Hu, Shuang Xiao, Yinglong Yang, and Shihe Yang J. Phys. Chem. Lett., Just Accepted Manuscript • DOI: 10.1021/acs.jpclett.8b03742 • Publication Date (Web): 31 Dec 2018 Downloaded from http://pubs.acs.org on January 1, 2019

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Excess Cesium Iodide Induces Spinodal Decomposition of CsPbI2Br Perovskite Films Xiangyue Meng,§† Zheng Wang,§‡ Wei Qian,‡ Zonglong Zhu,† Teng Zhang,† Yang Bai,† Chen Hu,† Shuang Xiao,†,‡ Yinglong Yang,† Shihe Yang*,†,‡ †Department

of Chemistry, The Hong Kong University of Science and Technology, Clear Water

Bay, Kowloon, Hong Kong, China. ‡Guangdong

Key Lab of Nano-Micro Material Research, School of Chemical Biology and

Biotechnology, Peking University Shenzhen Graduate School, Shenzhen 518055, China. §X.M.

and Z.W. contributed equally to this work.

Corresponding Author *Email: [email protected]


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ABSTRACT: We report an exploratory study on the crystal formation behavior of CsPbI2Br perovskite films by adding excess cesium iodide (CsI). Surprisingly, facile co-crystallization of CsI and CsPbI2Br in the form of spinodal decomposition is observed. Significantly, the two phases spontaneously form morphing into a remarkably uniform bi-continuous nano scale blend with a high orientational correlation through the well-matched (110) plane of CsI and the (200) plane of CsPbI2Br. The CsPbI2Br films produced by the spinodal decomposition method not only enjoy a compact surface, low defect concentration and long carrier lifetimes, and they also retain the excellent charge transport property. By employing such a CsPbI2Br film for carbon based perovskite solar cells, power conversion efficiency exceeding 10% is achieved with a remarkable thermal stability. Our results provide valuable insight into the role of CsI in the perovskite crystallization and a promising approach for designing inorganic halide perovskite based devices.

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Halide perovskite materials have become the superstar in emerging semiconductor materials, as the past decade witnessed their successful application in high performance photovoltaics.1-5 Thanks to their broad and intense light absorption,6 low exciton binding energies,7 and excellent charge transport characteristics even with a solution processed film,8 the halide perovskite materials have allowed the power conversion efficiency (PCE) of perovskite solar cells (PVSCs) to rapidly increased from 3.8% to 23.3%.9 However, the instability of organic-inorganic hybrid halide perovskite has become a bottleneck for long-term practical deployment due to the volatility of the organic cations under thermal and humidity stress.10 Meanwhile, all inorganic cesium-based perovskites (CsPbX3) have demonstrated excellent thermal stability up to their melting points at more than 400°C, which have drawn rising attention in the photovoltaic community.11-13 Among them, the mixed-halide composition CsPbI2Br, with a 1.92 eV band gap,14 appears to be a promising light absorber for the top-cell of tandem devices, which offers a better phase stability than the neat iodide composition (CsPbI3), yet a narrower band gap than the neat bromide perovskite (CsPbBr3). However, it is still challenging to prepare compact CsPbI2Br films with both sufficient thickness and high quality. Increasing efforts have been dedicated into developing high quality CsPbI2Br films with large grain size, dense surface morphology and low defect density. In the typical one-step spin-coating process, the CsPbI2Br perovskite films were often prepared by spin coating the precursor solution of stoichiometric CsBr and PbI2 in the solvent of N,N-dimethylformamide (DMF).15 Unfortunately, the solubility of CsBr in DMF limited the concentration of the CsPbI2Br precursor solution to 0.4 M, resulting in a CsPbI2Br film that was only 150 nm, and consequently a low PCE.16 In principle, using mixed solvent with dimethyl sulfoxide (DMSO) would increase the precursor concentration and film thickness. But the higher boiling point, stronger polarity and


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coordination ability of DMSO tends to slow down the crystallization of CsPbI2Br, yielding numerous pin-holes in as-formed film.17 Although many efforts have been made to fabricate high quality CsPbI2Br films, such as post annealing condition optimizing,18,19 precursor composition designing,20,21 additive incorporating,22 and anti-solvent engineering,23 it is still challenging to prepare pinhole-free and uniform CsPbI2Br films with a sufficient thickness and high quality. An additional stability problem with the previously reported CsPbI2Br PVSCs lies in the fact that they commonly contained organic hole transport materials (HTMs) and metal electrodes. The additives in organic HTMs are deliquescent and hygroscopic, detrimental to the device stability.24 Notably, carbon electrodes are inherently water-resistant and inert to ion migration, making them the most promising for addressing the stability issue of PVSCs.25-27 The simplest embodiment of this conception would be to replace the organic HTMs and the metal electrodes with a single carbon electrode, which has in fact been tested with great success. In the present work, we document a new strategy for the crystallization of CsPbI2Br films. Specifically, we introduce another bulk phase of cesium iodide (CsI) to co-crystallize with CsPbI2Br, generating a stabilized nano scale phase of CsPbI2Br that spans the bulk of the entire film. To our surprise, the excess CsI engendered a co-crystallization mode of spinodal decomposition through the well-matched (110) plane of CsI and the (200) plane of CsPbI2Br. Strikingly, the spontaneous co-crystallization of CsI and CsPbI2Br with a high orientational correlation dramatically improved the film coverage and uniformity, and reduced halide vacancies, resulting in a PCE of higher than 10% for carbon based perovskite solar cells, which is a record value for the CsPbI2Br PVSCs without HTMs. Moreover, the CsI incorporation could stabilize the CsPbI2Br phase against high thermal stress.


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Figure 1. (a) SEM image of the CsPbI2Br film without CsI (the inset shows the corresponding cross-sectional SEM image). (b) SEM image of the CsPbI2Br film with excess CsI (the inset shows the corresponding cross-sectional SEM image). (c) XRD pattern and (d) absorption spectrum of the CsPbI2Br film with excess CsI. In our experiments, the perovskite films were fabricated by a one-step spin-coating method.28 PbI2 and CsBr (1:1, 1 mmol) were dissolved in anhydrous DMSO (1 mL), and then excess 0.8 mmol CsI was added to prepare the perovskite precursor solution. Films were spin coated and then annealed at 180 °C for 60 s on a hot plate. When the CsPbI2Br thin film was prepared without CsI by using DMSO as the solvent, the coverage was poor due to the slow nucleation rate and crystallization process of perovskite (Figure 1a). In marked contrast, the CsPbI2Br thin film prepared with excess CsI displays a compact, pinhole-free and full coverage with a larger thickness of 350 nm (Figure 1b), which turns out to be crucial for optimizing the photovoltaic performance. X-ray diffraction (XRD) pattern of the as-prepared perovskite film with excess CsI


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is shown in Figure 1c. Evidently, two separate crystalline phases, namely CsPbI2Br and CsI, have formed in the as-prepared perovskite film. The Bragg peaks at 2θ = 14.6 ° and 29.5° could be indexed to a phase-pure cubic CsPbI2Br perovskite structure, and the peak located at 2θ = 27.5 ° is a characteristic feature of CsI. Interestingly, these are the only XRD peaks observed, meaning that both crystalline phases are highly oriented. Referring to the high uniformity of the film (Figure 1b), the CsI and CsPbI2Br phases appear to blend together in the nanometer size range. Moreover, the UV-Vis absorption spectrum of the as-prepared film with excess CsI exhibits an absorbance onset at about 670 nm (Figure 1d), consistent with the absorption spectrum of CsPbI2Br reported previously.29,30 The main composition of the prepared film should be CsPbI2Br rather than CsPbI3 because of more stable perovskite cubic phase of CsPbI2Br.

Figure 2. (a) AFM height image, (b) AFM phase image and (c) AFM potential image of the perovskite film with excess CsI (2 μm × 2 μm). (d) Schematic representation of the perovskite film with excess CsI.


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To map out the distribution of both CsI and CsPbI2Br phases in real space within the films, Kelvin probe force microscopy (KPFM) was applied providing information about the surface morphology. According to the height image, the as-prepared film with excess CsI is uniform (Figure 2a). However, two percolating regions are observed on the order of 100 nm from the phase image (Figure 2b) and the potential image (Figure 2c). The darker region should be CsI and the brighter region should be CsPbI2Br due to the potential difference that associated with their relative work functions. Evidently, the CsI and CsPbI2Br phases form an interpenetrating structure in the as-prepared film (Figure 2d). In other words, the as-prepared film is a bicontinuous composite of CsI and CsPbI2Br phases. It should be noted that the phase of CsPbI2Br is continuous, like the bulk heterojunction structure of donor and acceptor in organic solar cells,31 which allows effective charge transport in continuous pathways to the electrodes.

Figure 3 (a) XRD pattern of the perovskite film with structural evolution after thermal annealing at 180 °C for different time. (b) XRD intensity of CsI (110) and CsPbI2Br (100) as a function of heating time. (c) Schematics of the crystal structures of CsI phase, CsPbI2Br perovskite phase and Ruddlesden-Popper phase.


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To better understand the function of excess CsI in the crystallization behavior of the perovskite CsPbI2Br, XRD pattern and intensity of the perovskite film before and after thermal annealing at 180 °C for different time were measured and analyzed (Figure 3a and 3b). There are several interesting observations: (1) the non-annealed sample shows a relatively small peak at 27.5°, which corresponds to the (110) plane of CsI; (2) after thermal annealing at 180 °C for 10 s, the peak intensity of CsI becomes strongest, and a new diffraction peak appears at 14.6°, corresponding to the formation of cubic CsPbI2Br; (3) the diffraction peaks of the (110) plane of CsI and the (200) crystal plane of CsPbI2Br are quite close; (4) the (110) plane of CsI is gradually weakened with the further increase of the annealing time. While the peak intensity of CsPbI2Br at 14.6° is significantly increased after thermal annealing at 180 °C for 20 s, and strengthened slightly with the further increase of the annealing time. During the thermal treatment, phase segregation of CsI and CsPbI2Br without any nucleation event was observed. This type of process, known as spinodal decomposition, refers to the demixing of components in a solid solution, i.e., a single thermodynamic phase to form two coexisting phases with a specific spacial structure and length scale.32 From observations (1), (2) and (3), we can conclude that CsI and CsPbI2Br will immediately form after the solvent is evaporated at 180 °C. We deduce that the CsI first assembles, triggering spinodal decomposition. During the process, the CsI has a higher growth rate than CsPbI2Br, and promotes the formation of the cubic perovskite crystal phase through the interaction between the (110) plane of CsI and the (200) plane of CsPbI2Br owing to their structural match. As a result of the spinodal decomposition process, the CsI and CsPbI2Br phases are self-connected but mutual separated, which guarantees the high mobility of carriers in the continuous perovskite phase. Observation (4) indicates that the growth of CsPbI2Br crystal will consume CsI, suggesting that


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the (200) crystal plane of CsPbI2Br could be Pb-terminated so that CsI will be consumed as the cubic CsPbI2Br grows. From observations (3) and (4), it is reasonable for us to expect that the new perovskite CsPbI2Br phase grew on the (110) plane of CsI upon further thermal annealing. To confirm this hypothesis, we checked the lattice constants of the (110) plane of CsI (a = b = 5.55 Å) and the (200) plane of CsPbI2Br (a = 5.97 Å, b = 6.40 Å), and found that the mismatch is indeed relatively low. More importantly, all of the dangling bonds of the (110) plane of CsI and the (200) plane of CsPbI2Br would be removed after interfacing them together with a ratio of 1:1. It has been well established that the control of the lattice mismatch to within a low range and the removal of all of dangling bonds are highly favorable for crystallization. Therefore, the CsPbI2Br perovskite phase would grow on the (110) plane of CsI in the direction via spinodal decomposition. The growth of CsPbI2Br on the (110) plane of CsI in the direction is in agreement with the recent report about all inorganic halide perovskite with the Ruddlesden-Popper phase.33,34 Yu et al. reported the existence of Ruddlesden-Popper phases in CsPbBr3 nanosheets supported by atomic-level aberration-corrected scanning transmission electron microscopy observations.34 For the Ruddlesden-Popper phase, the general formula for the system is AX(ABX3)n. Specific to our system, the parent structures of the Ruddlesden-Popper phases should be the cubic rock-salt type CsI and the cubic CsPbI2Br perovskite. Therefore, the formula can be written as (CsI)0.8(CsPbI2Br) in our system, which is composed of layers of CsPbI2Br unit cells which are separated by CsI layers. To understand it more straightforwardly, we illustrate the growth process of the as-prepared film with excess CsI in Figure 3c. In the first stage of thermal annealing, CsI was formed with the (110) plane dominantly exposed. Then, cubic CsPbI2Br grow in the direction and interacts with the (110) plane of CsI in the second stage of thermal


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annealing. Finally, the RP phase consisting of composite of CsI and CsPbI2Br phases via spinodal decomposition was generated upon further thermal annealing.

Figure 4 (a) Gibbs free energy, (b) the first and (c) second derivatives of Gibbs free energy with respect to the mole fraction of CsI (α), and (d) the related phase diagram of CsI and CsPbI2Br. To gain further insight into the co-crystallization of CsI and CsPbI2Br, we plot in Figure 4 Gibbs free energy, the first and second derivatives of Gibbs free energy, as well as the phase diagram of CsI and CsPbI2Br with respect to the mole fraction of CsI (α = nCsI / (nCsI + nCsPbI2Br)). The computational details can be found in the Supporting Information. As shown in Figure 4, spinodal decomposition takes place over a wide range of mole fraction of CsI (0.13 < α < 0.87) at the experimental temperature of 453 K, corresponding to the amount of CsI in the range between 0.15 equivalents and 6.69 equivalents. This result bolsters the hypothesis that addition of excess CsI (0.8 equivalents) would promote the film quality of perovskite via spinodal decomposition.


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Figure 5 (a) Photoluminescence decay curves of perovskite films prepared by the conventional spin-coating method and spinodal decomposition method, (b) Type I energy level alignment of CsI and CsPbI2Br. To explore the novel perovskite films from spinodal decomposition for photovoltaics, we investigated the carrier transport dynamics. First, we employed the time-resolved photoluminescence decay (TRPL) to measure the carrier lifetime of the perovskite films. Conventional spin-coating processed CsPbI2Br film with good crystallinity and coverage prepared by DMF as the solvent was used as the control sample. TRPL decay curves of the perovskite films prepared by the conventional spin-coating method and our spinodal decomposition method are shown in Figure 5a. By ﬁtting the data to a bi-exponential decay model, two time constants are obtained, i.e., τ1 and τ2, as summarized in Table S1. The faster decay component τ1 might come from trap-mediated nonradiative recombination. The slower decay component τ2 should be attributed to radiative recombination. In the conventional spincoating processed CsPbI2Br film, the faster component τ1 (0.8481 ± 0.0127 ns) dominated the PL decay with negligible slower component, suggesting a severe recombination occurred. For the perovskite film prepared by our method, τ1 is slightly longer (0.8999 ± 0.0209 ns), and τ2 appears


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to be reaching 8.1104 ± 0.9958 ns indicating prolonged carrier lifetime. The enhanced PL lifetime suggests the reduced recombination in the (CsI)0.8(CsPbI2Br) film, resulting from reduced defect density by the introduction of excess CsI. Clearly, the presence of excess CsI closely correlates to the reduced carrier recombination in the (CsI)0.8(CsPbI2Br) film. First, the excess CsI introduces excess iodine into the perovskite film to compensate for any halide vacancies. The excess iodine filled these vacancies, thereby passivating the nonradiative recombination pathways and leading to the enhanced PL lifetime. Second, the excess CsI is atomically connected with the CsPbI2Br phase in terms of the RP structure, which can eliminate defects at the surface and grain boundaries of CsPbI2Br phase. Third, the reduced carrier recombination in the (CsI)0.8(CsPbI2Br) film is related to the favored energy band alignment (Figure 5b). This type I band alignment between CsI and CsPbI2Br forms energy barriers to prevent excitons from reaching the grain boundary traps states and defects, thus blocking the nonradiative channels. Together, the excess CsI effectively passivates the CsPbI2Br phase locally, resulting in a substantially reduced carrier recombination. Next, hall-effect measurements were also performed on the conventional spin-coating processed CsPbI2Br film and (CsI)0.8(CsPbI2Br) film. The samples were measured under a magnetic field of 0.48 T under a constant current of 10 nA. The carrier concentrations of the perovskite films prepared by the conventional spin-coating method and spinodal decomposition method were in the same order of 1012 cm-3. Combining resistivity and Hall effect data, the mobility was calculated to be 200 cm2 V-1 s-1 for the conventional spin-coating processed CsPbI2Br film, 36.2 cm2 V-1 s-1 for the (CsI)0.8(CsPbI2Br) film prepared by the spinodal decomposition method. These are remarkably high values given the presence of comparable amount of CsI, and are indeed comparable to the perovskite films reported previously,35


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indicating excellent charge transport property. Thus, the introduction of the CsI phase in the (CsI)0.8(CsPbI2Br) blends seem to have little impact on the charge transport, evidently because the continuous CsPbI2Br phase allows effective charge transport all the way to the electrodes as expected for the bi-continuous phases developed from the spinodal decomposition.

Figure 6. (a) Schematic view and (b) J-V plots of the carbon based PVSCs with CsPbI2Br prepared by the conventional spin-coating method and spinodal decomposition method. The as-prepared perovskite film with excess CsI was employed as the light absorber to fabricate carbon based PVSCs as the light absorber. Figure 6a shows the schematic view of carbon based PVSCs with the functional layers of ITO/SnO2/CsPbI2Br/carbon. The carbon cathode is inexpensive, is not prone to corrosion, and additionally, can act as a hydrophobic moisture barrier to enhance the stability of the PVSC.36 Typically, the carbon based PVSCs with CsPbI2Br prepared by the conventional spin-coating method exhibited a representative PCE of only 6.85% under reverse scan, mainly due to the film that was too thin with poor quality. In stark contrast, our best performing PVSC based on CsPbI2Br prepared by spinodal decomposition method showed a much better PCE of 10.13% under reverse scan. The


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enhancement in all the photovoltaic parameters, i.e., Jsc, Voc, and FF, can be attributed to the compact surface, low defect concentration, long carrier lifetime, and unobstructed charge transportation of the (CsI)0.8(CsPbI2Br) film. Furthermore, the (CsI)0.8(CsPbI2Br) based PVSCs show smaller J-V hysteresis (Figure S2 and Table S2), resulting in a stable 9.5% output (Figure S3). The external quantum efficiency (EQE) spectrum of the champion device based on (CsI)0.8(CsPbI2Br) is shown in Figure S4, which corresponds to an integrated Jsc of 12.19 mA cm-2, in good agreement with the measured Jsc from the J-V curves. Device performance was also highly reproducible (Figure S5). The improved reproducibility demonstrated by the solar cell devices based on (CsI)0.8(CsPbI2Br) should be ascribed to the better control on the film formation. We also examined photovoltaic performance of the carbon based PVSCs with different ratios of CsI to CsPbI2Br (Table S3). It was found that the (CsI)0.8(CsPbI2Br) solar cells exhibited highest PCE due to the formation of the desirable bi-continuous morphology. At a lower CsI content, the perovskite film grew with inefficient number of CsI (110) planes, and the non-ideal condition led to pin holes (Figure S6). On the other hand, at a higher CsI content, the insulating CsI phase started to dominate the perovskite phase, leading to reduced light utilization and carrier mobility. This suggests that there is an optimal CsI content of approximately 0.8 equivalents, which is a compromise between high tendency of spinodal decomposition and the retention of high light utilization and high charge-carrier mobility. Moreover, the CsI incorporation could stabilize the CsPbI2Br phase by reducing the surface Gibbs free energy.37 The long-term thermal stability of ITO/SnO2/CsPbI2Br/carbon PVSCs was monitored at 85 °C, close to the ISOS-D-2 protocol (Dark, Intermediate level (Level 2)). Clearly, the (CsI)0.8(CsPbI2Br) PVSCs just display about 10% PCE loss while being continually heated at


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85 °C for 200 h (Figure S7), which is likely associated with the humidity-induced phase change during the characterization process. No obvious change was observed in the absorption spectrum of the (CsI)0.8(CsPbI2Br) film after being heated at 85 °C for 200 h (Figure S8), indicating the good thermal stability of the (CsI)0.8(CsPbI2Br) film. We attribute the good thermal stability mainly to the interactions between the (110) plane of CsI and the (200) plane of CsPbI2Br by forming a RP phase. Therefore, the carbon based PVSCs with (CsI)0.8(CsPbI2Br) not only exhibit a decent PCE, but also enjoy a much improved thermal stability. In conclusion, we reported a new crystallization mode of the CsPbI2Br perovskite by spontaneously forming a nanoscale, oriented bi-continuous blend film with CsI through spinodal decomposition. Facile co-crystallization of CsPbI2Br with a high orientational correlation was observed due to the similar structure of the (110) plane of CsI to the (200) plane of CsPbI2Br. Strikingly, the excess CsI induces spinodal decomposition of CsPbI2Br, which not only affords excellent film coverage and reduced defects, but also allows fluent charge transportation. As a demonstration, we applied the blend film to carbon-based PVSCs and obtained a PCE of higher than 10% and a high thermal stability. This new crystallization mode of CsPbI2Br via spinodal decomposition highlights the critical role of CsI in the process and heralds a promising direction for the electronics and photonics enabled by the inorganic halide perovskite. Acknowledgments This work was financially supported by the HK-RGC General Research Funds (GRF Nos. 16312216 and 16300915), the HK Innovation and Technology Fund (ITS/219/16 and GHP/079/17SZ), the Shenzhen Peacock Plan Program (KQTD2016053015544057).


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Supporting Information. The Supporting Information is available free of charge on the ACS Publications website. Experimental section, mapping distribution of the element, EQE, steady-state photocurrent, PCE distribution, SEM, thermal stability test.

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