Sequentially vacuum evaporated high-quality CsPbBr3 films for efficient carbon-based planar heterojunction perovskite solar cells

Sequentially vacuum evaporated high-quality CsPbBr3 films for efficient carbon-based planar heterojunction perovskite solar cells

Journal of Power Sources 443 (2019) 227269 Contents lists available at ScienceDirect Journal of Power Sources journal homepage: www.elsevier.com/loc...

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Journal of Power Sources 443 (2019) 227269

Contents lists available at ScienceDirect

Journal of Power Sources journal homepage: www.elsevier.com/locate/jpowsour

Sequentially vacuum evaporated high-quality CsPbBr3 films for efficient carbon-based planar heterojunction perovskite solar cells Xingyue Liu a, Xianhua Tan a, Zhiyong Liu a, Bo Sun a, Junjie Li a, Shuang Xi b, Tielin Shi a, Guanglan Liao a, c, * a b c

State Key Laboratory of Digital Manufacturing Equipment and Technology, Huazhong University of Science and Technology, Wuhan, 430074, China School of Mechanical and Electronic Engineering, Nanjing Forestry University, Nanjing, 210037, China Shenzhen Huazhong University of Science and Technology Research Institute, China

H I G H L I G H T S

G R A P H I C A L A B S T R A C T

� A sequential evaporation method is devised to prepare high-quality CsPbBr3 films. � An excellent PCE of 7.58% is achieved for carbon-based CsPbBr3 PSCs. � A PCE of 6.21% is obtained for largearea devices with an active area of 1 cm2. � Superior moisture and thermal stabil­ ities are obtained for the devices. � The whole production costs of the de­ vices are very low.

A R T I C L E I N F O

A B S T R A C T

Keywords: Non-solution Sequential evaporation CsPbBr3 Perovskite solar cells Carbon-based Highly efficient and stable

All-inorganic CsPbBr3 perovskite has triggered great interests in photovoltaic field owing to its superior stability. However, the uncontrollable CsPbBr3 film growth in solution always leads to a poor film quality with low phasepurity as well as many surface and bulk defects. Herein, we demonstrate an environmentally friendly nonsolution route to fabricate high-quality CsPbBr3 films for carbon-based planar perovskite solar cells. By pre­ cisely tuning the thickness ratio of the evaporated CsBr to PbBr2 precursors (r), the dominant phase conversion of the cesium lead bromide perovskites from PbBr2-rich CsPb2Br5 (r � 12:7) to CsPbBr3 (r ¼ 12:8), and further to CsBr-rich Cs4PbBr6 (r � 12:9) are achieved. The optimized CsPbBr3 perovskites are highly phase-pure and crystallized with ultra-high light absorption ability. The as-prepared CsPbBr3 films also exhibit a dense and uniform morphology with large grain sizes and monolayer-vertical aligned grains. The corresponding devices deliver a champion PCE of 7.58%, which is an excellent efficiency among carbon-based CsPbBr3 cells with evaporated CsPbBr3 light absorbers. The large-area (1 cm2) devices also achieve an efficiency of 6.21%. More­ over, the unencapsulated CsPbBr3 devices present superior moisture and thermal stabilities. Our work provides a facile approach to fabricate high-quality and large-area CsPbBr3 films for highly efficient solar cells, lightemitting diodes and photodetectors.

* Corresponding author. State Key Laboratory of Digital Manufacturing Equipment and Technology, Huazhong University of Science and Technology, Wuhan, 430074, China. E-mail address: [email protected] (G. Liao). https://doi.org/10.1016/j.jpowsour.2019.227269 Received 23 April 2019; Received in revised form 6 September 2019; Accepted 3 October 2019 Available online 9 October 2019 0378-7753/© 2019 Elsevier B.V. All rights reserved.

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1. Introduction

dual-source coevaporation is an effective way to fabricate dense and highly crystallized CsPbBr3 films, it has a special requirement for the preparation equipment. In addition, it is hard to precisely control the molar ratio of the PbBr2 to CsBr at the same time since each sensor can be easily disturbed by the other precursor vapor in the coevaporation process. Ajjouri et al. proposed a single-source vacuum deposition of CsPbBr3 film, in which the PbBr2 and CsBr precursor were mixed at a molar ratio of 1:1 in a crucible and then were evaporated to form CsPbBr3 [21]. However, the actual molar ratio of PbBr2 to CsBr reaching to the substrate was hardly to detect for there was a distinction between their melting points. Jiang’s group developed a sequential evaporation method to prepare mixed CsPbBr3–CsPb2Br5 phase as light absorber, in which 100 nm CsBr layer was firstly evaporated and the PbBr2 film was subsequently evaporated on it with varied thickness [22]. Their devices obtained an efficiency of 8.34%, which was the highest efficiency of CsPbBr3 PSCs based on vacuum-evaporated CsPbBr3 light absorbers to date. However, the CsPbBr3–CsPb2Br5 layer was too thin to harvest the incident light efficiently and the existence of much CsPb2Br5 phase in the CsPbBr3 film increased its bandgap to 2.4 eV. Both of these issues lead to a low light absorption coefficient of the perovskite layer as evi­ denced by the absorption spectra. Besides, the high-temperature (500 � C) sintered TiO2 as well as expensive Spiro-OMeTAD hole trans­ port layers (HTLs) and noble metal Ag electrodes used in their devices greatly increased the production costs, energy-consuming and decreased the stability of the devices (only maintained ~86% of its initial PCE after 1000 h). As far as we know, the carbon-based CsPbBr3 PSCs with evaporated CsPbBr3 light absorbers have been rarely reported yet. In this work, we present a non-solution strategy to fabricate highquality CsPbBr3 films. The PbBr2 was firstly evaporated on the lowtemperature processed c-TiO2 film and CsBr precursor are then evapo­ rated on the PbBr2 film with different thickness. The abandonment of N, N-Dimethylformamide (DMF) and methanol solvents is beneficial to both human health and environment protection. The impact of the component ratio on perovskite crystallinity, light absorption coefficient and morphology are meticulously and systematically investigated. By precisely tuning the thickness ratio of the evaporated CsBr to PbBr2 precursor films (r), the dominant phase of the cesium lead bromide can be transformed from CsPb2Br5 to CsPbBr3 and further to Cs4PbBr6. When it comes to the optimized ratio of 12:8, a highly phase-pure and crys­ tallized CsPbBr3 film is obtained. The as-prepared CsPbBr3 film exhibit a dense and uniform morphology with large average grain sizes of over 1 μm, extremely smooth surface and ultra-high light absorption ability. The thickness of the CsPbBr3 perovskite layer is also optimized. The corresponding PSC achieves a champion power conversion efficiency (PCE) of 7.58%, which is an excellent performance for carbon-based CsPbBr3 PSCs with evaporated CsPbBr3 light absorbers. The devices also achieve a high efficiency of 6.21% with a large active area of 1 cm2. Many characterizations have been employed to get insight into the in­ fluence of the compositional structure of the light absorber on the device performance. Besides, the unencapsulated CsPbBr3 PSCs deliver supe­ rior moisture and thermal stabilities. Our work provides a feasible approach to precisely fabricate CsPb2Br5, CsPbBr3 and Cs4PbBr6 perovskite, which can also be applied in the light emitting and photo­ detector fields.

In the past few years, organic-inorganic hybrid perovskites have attracted tremendous attentions as photosensitive materials due to their unique photoelectric properties, such as high light absorption coefficient and carrier mobility, tunable bandgap, long carrier diffusion length as well as remarkable tolerance to defects [1–4]. The power conversion efficiency (PCE) of perovskite solar cells (PSCs) has been boosted to over 25% [5], comparable to commercial silicon solar cells. However, hybrid perovskites always suffer from poor stabilities which is relevant to the size of organic cation, crystal lattice and grain boundaries, etc, making PSCs unstable when working in a high humidity and thermal condition [6]. One considerable solution to this problem is to substitute the organic parts with inorganic cations (Csþ, Rbþ, etc.) to form all-inorganic perovskites [7]. CsPbI3 and CsPbI2Br are the most widely studied inorganic perovskites these days due to their suitable bandgap and superior thermal stability together with comparable light absorp­ tion ability [8]. The fast conversions to non-perovskite phase when exposed to air at room temperature set an obstacle for their practical applications [9]. Among all the inorganic perovskites, CsPbBr3 pos­ sessing both the best moisture and thermal stabilities is a promising alternative to hybrid perovskite light harvesters. The film quality of CsPbBr3 has a significant impact on the device performance and many efforts have been paid on ameliorating its crystallization process. Kulbak et al. first reported CsPbBr3 PSC in 2015 with the CsPbBr3 film prepared by a two-step solution method [10]. The device only obtained a PCE of 5.95%. Yu’s group found that CsPbBr3 could decompose quickly in CsBr methanol solution in the traditional two-step deposition process, resulting in a poor film morphology and low phase-purity [11]. They then introduced a face-down liquid-­ space-restricted deposition method to suppress the decomposition pro­ cess of CsPbBr3 and obtained a high PCE of 5.86%, which was the highest efficiency for planar CsPbBr3 PSC at that time. Tang’s group created a multistep solution-processing technique to optimize the ma­ terial components, crystal structure and morphologies of CsPbBr3 pe­ rovskites [12]. Their PSCs achieved a high PCE of 9.72%. The same group further boosted the PCEs of CsPbBr3 PSCs to over 10% by quan­ tum dots modification, ion doping as well as spectra and interface en­ gineering, etc [13–15]. Zeng et al. reported a space-confined growth method to gain CsPbBr3 films with high carrier mobility and low trap-state densities, which was further used in photodetectors [16]. Our group also proposed a multistep deposition strategy, involving an im­ mersion and spin-coating of CsBr process, to improve the crystallinity and phase-purity of CsPbBr3 films [17]. In combination with interface engineering, our planar-heterojunction device demonstrated an excel­ lent PCE of 8.79%. Notably, there were still some impurity phase (CsPb2Br5 and Cs4PbBr6) existing in abovementioned solution-processed CsPbBr3 films, let alone that spin-coated method is not so suitable for fabricating large-area films. Vapor-assisted deposition technique have also been developed to fabricate continuous CsPbBr3 films in large area. Luo et al. created a Br2 vapor-assisted deposition method to get a fast anion-exchange from CsPbI3 to CsPbBr3 and get a champion PCE of 5.38% [9]. Duan et al. reported a spray-assisted deposition method for CsPbBr3 preparation, in which the PbBr2 was spin-coated and CsBr was sequentially evaporated for four times [18]. Their devices got a PCE of 6.8% in small area and only 4.12% in an active area of 1 cm2. Chen et al. first deposited high-quality CsPbBr3 films by a dual-source coevapora­ tion process and optimized the evaporation rates of the PbBr2 and CsBr precursors [19]. The planar PSC with a structure of FTO/Z­ nO/CsPbBr3/Spiro-OMeTAD/Au achieved an excellent PCE of 7.78%. Liu’s group also used the same coevaporation method to fabricate CsPbBr3 films and investigated the impact of substrate and post-annealing temperature on the films’ crystallinities [20]. The cor­ responding PSC with a structure of FTO/compact TiO2 (c-TiO2)/CsPbBr3/Spiro-OMeTAD/Au yielded a PCE of 6.95% and 5.37% with an active area of 0.09 and 1 cm2, respectively. Although

2. Experimental section 2.1. Materials Cesium bromide (CsBr, �99.99%) and lead bromide (PbBr2, �99.99%) powder are purchased from Xi’an p-OLED. Titanium tetra­ chloride (TiCl4, �99.5%) and Copper (II) phthalocyanine (CuPc) are from Aladdin. NiCl2⋅6H2O (�98.0%) is bought from Sinopharm Chem­ ical Reagent Co., Ltd. The commercial carbon paste is from Shenzhen Dongdalai Chemical Co., Ltd. All the chemicals and reagents are directly used without any further purification. 2

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2.2. Device fabrication

3. Results and discussion

The conductive fluorine-doped tin oxide (FTO) glass substrates (NSG-10) were etched by laser. The FTO substrates were then sequen­ tially cleaned with detergent, acetone, anhydrous alcohol and deionized water in an ultrasonic bath each for 15 min, followed by an O3/ultra­ violet treatment for 30 min to remove the organic residues. The Nidoped c-TiO2 electron transport layers (ETLs) were deposited via the hydrolysis of the TiCl4 aqueous solution as we previously reported [23]. In specific, the TiCl4 solution was diluted to 200 mM by deionized water in ice bath with 0.01 M NiCl2⋅6H2O being added. The cleaned FTO substrates were then placed vertically in a glass container filled with TiCl4 precursor solution and kept at 70 � C for 3 h in a thermostat water bath. Subsequently, the substrates were washed with deionized water in a sonication bath for 5 min to remove any loosely bound materials. After dried by N2 flow, the substrates were annealed at 200 � C on a hotplate for 1 h. The CsPbBr3 perovskites were fabricated by a sequential evap­ oration method. 270 nm PbBr2 films were firstly evaporated on the TiO2 ETL under a pressure of ~9 � 10 4 Pa at a speed of ~5 Å/s determined by a quartz crystal monitor. Subsequently, the CsBr films with different thickness ratios to PbBr2 were evaporated on the pre-deposited PbBr2 films under the same Conditions. The substrates were then annealed at 250 � C in ambient air for 5 min to improve the crystallinity of the CsPbBr3 films. For the fabrication of the hole transport layer (HTL), 35 nm CuPc was further thermally evaporated on the CsPbBr3 films under a base pressure of ~9 � 10 4 Pa at a low speed of ~0.5 Å/s. Ul­ timately, the commercial carbon paste was doctor-bladed on the CuPc HTL and then dried at 100 � C for 15 min to form the counter electrode. All the processes were performed in ambient air.

Fig. 1a depicts the schematic illustration of the sequential evapora­ tion process to prepare high-quality CsPbBr3 films. Since the evaporated film is extremely homogeneous, the thickness of the film is approxi­ mately in direct proportion to its molar quantity. Thus the constituent of the final cesium lead bromide film can be precisely manipulated by tuning the r value. PbBr2 layers are firstly evaporated on the substrates and CsBr films with varied thickness ratios to PbBr2 films are then evaporated onto the PbBr2 films, followed by annealed at 250 � C in air to get the films crystallized. As can be seen from the digital camera image of the as-prepared cesium lead bromide films with different r values (Fig. 1b), all the films are dense and uniform (certified by the homo­ geneity of their colors) on macroscale. When the amount of the evapo­ rated PbBr2 is much more than CsBr (r ¼ 12:4), the cesium lead bromide film exhibits a white-grey color, which is assigned to the main genera­ tion of the PbBr2-rich CsPb2Br5 phase. Gradually increasing r to 12:8, the color of the as-prepared film become orange-yellow with a continuously reduced transparency, suggesting the formation of large amount of CsPbBr3 phase. Further increasing the amount of CsBr (r ¼ 12:10), the film appears a more yellow and rougher surface as witnessed by eyes. The phase-transition dynamics in this process can be well described by following three reactions [12]: 2PbBr2 þ CsBr → CsPb2Br5 (with excessive PbBr2)

(1)

CsPb2Br5 þ CsBr → 2CsPbBr3

(2)

CsPbBr3 þ 3CsBr → Cs4PbBr6 (with excessive CsBr)

(3)

The schematic illustration of the corresponding crystal phase tran­ sition from CsPb2Br5 to CsPbBr3 and further to Cs4PbBr6 with increased CsBr amount are highlighted in Fig. 1c. The PbBr2-rich CsPb2Br5 exhibits a tetragonal structure, a 3D framework of corner-connected PbBr46 octahedra with Csþ ions located between the octahedra, and the Csþ in the CsPb2Br5 crystal are sandwiched between the two layers of Pb–Brcoordinated polyhedral [22]. The CsPbBr3 possess a symmetry cubic crystal structure with connected corner-sharing [PbBr6]4 octahedra, while the CsBr-rich Cs4PbBr6 phase possess a rhombohedral structure with the [PbBr6]4 octahedra separated by Csþ cations [24,25]. The XRD patterns of the as-evaporated cesium lead bromide films with varied thickness ratios of the precursors are given in Fig. 2a. There exists prominent peaks at 11.65� , 18.82� , 23.39� , 24.03� , 27.77� , 29.38� , 33.40� and 35.47� for the 12:4 film, corresponding to the (002), (112), (210), (202), (114), (213), (310) and (312) lattice faces of CsPb2Br5 phase (PDF#25-0211), respectively. This indicates a serious insufficiency of CsBr and that CsPb2Br5 acts as the dominant phase in the film (eqn. (1)). By contrast, negligible CsPbBr3 peaks are observed. CsPb2Br5 phase is reported to possess an inactive photoluminescence behavior and a large indirect bandgap of about 3.1 eV, which is unfa­ vorable for the device performance [26]. Augmenting the r value but no more than 12:8, the intensity of the peaks at 15.18� , 21.58� , 30.68� , 34.49� , 44.14� and 49.56� increases continuously accompanied by the gradual disappearance of CsPb2Br5 peaks. These peaks can be assigned to (100), (110), (200), (210), (220) and (310) faces of the CsPbBr3 phase (PDF#54-0752), signifying the enhanced formation of CsPbBr3 perov­ skite. When r comes to 12:8, the CsPb2Br5 impurity almost vanished and the intensity of all the CsPbBr3 peaks reaches to the highest level. A highly phase-pure and crystallized CsPbBr3 film is obtained, of which the crystals mainly grow along the (100), (110) and (200) directions. In this case, the evaporated molar ratio of the CsBr to PbBr2 precursors is about 1:1 (eqn. (2)). Further increasing the CsBr amount, the peaks at 12.68� , 12.89� , 22.46� , 25.44� , 27.75� , 28.68� , 29.03� , 30.35� , 39.07� and 45.88� occur and then get stronger, revealing the formation of the CsBr-rich Cs4PbBr6 phase (PDF#54-0750). When the r value rises to 12:11, Cs4PbBr6 becomes the main phase with the CsPbBr3 phase almost disappearing in the final cesium lead bromide film, as reflected by eqn.

2.3. Device characterization The X-ray diffraction (XRD) measurements were performed on a x’pert3 powder X-ray diffractometer (PANalytical, Netherland) with Cu Kα radiation (λ ¼ 1.5418 Å) at 25� and the data were collected with a 0.013� step size (2θ). The absorption spectra of the cesium lead bromide films were obtained by a UV–visible spectrophotometer (UV 2600, Shimadzu). The X-ray photoelectron spectroscopy (XPS) measurements of the as-prepared CsPbBr3 films were performed on a photoelectron spectrometer (AXIS-ULTRA DLD-600W, Kratos, Shimadzu, Japan). The EDS elemental mapping of the CsPbBr3 films were examined by the scanning electron microscopy (SEM, Helios NanoLab G3 CX, FEI). The surface morphologies of the cesium lead bromide films, and the crosssectional images of the cesium lead bromide films and whole devices were examined by the scanning electron microscopy (SEM, Sirion 200, FEI, Heland and GeminiSEM 300, Carl Zeiss, German). The atomic force microscopy (AFM) images and root-mean-square roughness of the ce­ sium lead bromide films and perovskite films were measured by an Innova SPM 9700 (Shimadzu, Japan) in tapping mode. The current density-voltage (J-V) characteristics of the devices were recorded at a scan rate of 20 mV/s by using a computer-controlled electrochemical station (Autolab PGSTA302 N, Netherlands) under simulated AM 1.5G one sunlight illumination (100 mW/cm2) generated by a solar simulator (Oriel 94043A, Newport Corporation, Irvine, CA, USA), which was calibrated by a NREL-traceable KG5 filtered silicon reference cell. The active area of the regular cells in the measurements were defined by a metal mask with an aperture area of 0.071 cm2, while the testing area of the large-area devices is set as 1 cm2. Moreover, the electrochemical impedance spectroscopy (EIS), open-circuit photovoltage decay (OCVD), as well as capacitance-voltage measurements were also carried out on this equipment. The EIS spectra were fitted via the Nova software. The steady-state photoluminescence (PL) and time-resolved photo­ luminescence (TR-PL) decay spectra were obtained via FluoTime300 (PicoQuant, German).

3

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Fig. 1. (a) Schematic illustration of the sequential evaporation technique to prepare CsPbBr3 perovskites. (b) Digital camera image of the as-prepared cesium lead bromide films with different r values. (c) Schematic illustration of the crystal structure transition from CsPb2Br5 to CsPbBr3 and further to Cs4PbBr6 phase with varied r values. The arrow at the bottom indicates the phase transition process.

Fig. 2. (a) XRD patterns and (b) UV–vis absorption spectra of the cesium lead bromide films with varied r values deposited on the FTO substrates. (c) The plots of (Ahv)2 versus the photon energy (hv) of the representative PbBr2-rich phase, pure CsPbBr3 phase and CsBr-rich phase with the r of 12:6, 12:8 and 12:10, respectively. 4

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(3). The UV–vis absorption spectra of the as-prepared cesium lead bro­ mide perovskites with varied r values (Fig. 2b) show a high absorbance in the ultraviolet and near-ultraviolet regions for these films. The light absorption of the films increases with rising the r value but no more than 12:8. This can be ascribed to the reduction of the Cs2PbBr5 phase with larger bandgap [26] and increment of the CsPbBr3 phase. The 12:8 film possesses the best absorption ability, even much higher than the solution-processed and co-evaporated CsPbBr3 as well as sequentially evaporated CsPbBr3–CsPb2Br5 perovskite films reported before [11,19, 22], due to its highest phase-purity and crystallinity [17]. The higher absorption coefficient of the film is beneficial to achieve a higher short-circuit current density (Jsc) and open-circuit voltage (Voc) for the PSC. With the generation of Cs4PbBr6 and decline of CsPbBr3 phase (r � 12:9), the light absorption of the films falls back quickly. The (Ahv)2 versus the photon energy (hv) plots of the representative PbBr2-rich (r ¼ 12:6), highly-pure CsPbBr3 (r ¼ 12:8) and CsBr-rich Cs4PbBr6 (r ¼ 12:10) films are also drawn in Fig. 2c. The mixed CsPbBr3–CsPb2Br5 film exhibits the largest bandgap of 2.4 eV, coinciding well with the previous report [22]. The high-purity CsPbBr3 and mixed CsPbBr3–Cs4PbBr6 films presents a similar absorption band edge at about 531.5 nm, corresponding to a bandgap of 2.33 eV. X-ray photoelectron spectroscopy (XPS) analysis is performed to identify the actual composition of the 12:8 CsPbBr3 film (Fig. 3a). The peaks at the binding energies of 724.5 and 738.5 eV in the highresolution XPS spectrum (Fig. 3b) correspond to the Cs 3d5/2 and Cs 3d3/2 whilst the peaks at around 137.0 and 141.8 eV (Fig. 3c) are assigned to the Pb 4f7/2 and Pb 4f5/2, respectively. The peaks at lower binding energies of 68.4 and 69.5 eV (Fig. 3d) belongs to Br 3d5/2 and Br 3d3/2 [27]. In addition, the relative atomic ratio of the elements Cs, Pb and Br is 20.38%: 20.47%: 59.15%, approximating to 1:1:3. This verifies the accurate manipulation of the film components for our deposition technique and highly phase-pure CsPbBr3 are successfully fabricated. The compositional distribution of the 12:8 film is also studied by EDS elemental mapping (Fig. 3e). The atomic ratio of Cs, Pb and Br is also about 1:1:3 and the elements are homogeneously distributed in the

scanning area, suggesting the high purity and uniformity of the as-evaporated CsPbBr3 film. The morphology evolutions of the as-deposited cesium lead bromide films composed of different thickness ratios of the precursors are revealed by scanning electron microscopy (SEM) and atomic force mi­ croscopy (AFM), as shown in Fig. 4. It appears that all these evaporated films exhibit a high coverage with densely packed crystals over the substrates. There are many small crystals with great number of grain boundaries existing in the 12:4 film (Fig. 4a), which can be explained by the main formation of the CsPb2Br5 phase with low crystallization in the surface [22]. The cross-sectional SEM image of the 12:4 film (Fig. 4e) also shows the stacking of many small crystals in the internal crystal, also yielding massive grain boundaries. Since grain boundaries are widely considered as the major recombination sites in the perovskites due to the existence of many surface impurities and trap states [28,29], a severe charge carrier recombination loss may occur in the 12:4 films. Besides, the grain boundaries-induced shallow states near the valence band edge will hamper the hole diffusion [30], which is also unfavorable for charge transport. The grain sizes gradually increase while the grain boundaries decrease with the increase of r within 12:8 (Fig. 4b and c), leading to reduced surface and bulk defect densities and elongated charge-carrier lifetimes. When r ¼ 12:8, the average grain sizes of the highly pure and crystallized CsPbBr3 film reaches 1.05 μm with verticaland monolayer-aligned grains (Fig. 4g). The photogenerated carriers can transfer in the out-of-plane directions across the ETL/HTL towards the electrodes without passing through grain boundaries, conducive to obtaining a higher Jsc and Voc. Further adding CsBr amount (r ¼ 12:10), the CsBr-rich film become uneven and many tiny crystals appears with augmented grain boundaries (Fig. 4d), which will significantly increase the current shunting pathways and non-radiative recombination loss. This can be mainly attributed to that the phase separation of Cs4PbBr6 from CsPbBr3 with excessive CsBr may split large grains into smaller grains [12]. AFM images displayed in Fig. 4i-l further unveil that the 12:8 film exhibits a more homogeneous morphology with larger grain sizes compared to the mixed CsPbBr3–CsPb2Br5 and CsPbBr3–Cs4PbBr6 counterparts. It also possesses the slickest surface with a low

Fig. 3. (a) XPS survey of the 12:8 CsPbBr3 film. High-resolution XPS spectra of (b) Cs 3d, (c) Pb 4f and (c) Br 3d peaks of the 12:8 film. (d) EDS elemental mapping of the components (Cs, Pb, and Br) of the 12:8 film. 5

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Fig. 4. (a–d) Top-view SEM images, (e–h) cross-sectional SEM images and (i–l) the corresponding AFM images of the cesium lead bromide films with the r values to be 12:4, 12:6, 12:8 and 12:10, respectively.

root-mean-square roughness (RMS) of only 37.6 nm, which is much lower than solution-processed and co-evaporated CsPbBr3 films [17,20]. To the best of our knowledge, this is the lowest surface roughness for

CsPbBr3 light absorbers applied in PSCs reported so far. A lower surface roughness of the perovskite films is always beneficial to form a better contact with the post-deposited HTL, yielding a lower series resistance

Fig. 5. (a) Steady-state PL and (b) TR-PL spectra of the cesium lead bromide films deposited on FTO substrates with the r values as 12:4, 12:6, 12:8 and 12:10, respectively. (c) Cross-sectional SEM image and (d) the corresponding energy diagram of the as-prepared device with a structure of FTO/c-TiO2/CsPbBr3/ CuPc/Carbon. 6

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Journal of Power Sources 443 (2019) 227269

(Rs) and a higher device performance [23,31]. The charge transport and extraction mechanisms of the representa­ tive PbBr2-rich, pure CsPbBr3 and CsBr-rich films are examined by steady-state photoluminescence (PL) and time-resolved photo­ luminescence (TR-PL) spectroscopy. Since no charge transport layer is applied, it is hard for the photoinduced carriers in the excited state to be extracted out quickly, provoking a radiative recombination behavior. A faster PL quenching is related to more trap states in the perovskite film [32–34]. The 12:4 sample exhibits an emission peak at around 533.5 nm and the strongest PL quenching, indicating the highest defect density in the film (Fig. 5a). This may be aroused by the main existence of CsPb2Br5 impurity phase with many grain boundaries. Increasing the thickness ratio of CsBr to PbBr2 to 12:8, the film gets a greatly enhanced PL in­ tensity, reflecting a suppressed carrier recombination rates. This can be mainly ascribed to the ameliorated crystallinity and phase-purity as well as reduced grain boundaries as observed from the SEM image (Fig. 4c). Moreover, since the CsPb2Br5 phase is PL-inactive [35], the blue-shift of the PL peak for the 12:8 CsPbBr3 film further confirms the suppressed trap states of the film [36], in favor of achieving a higher carrier lifetime and charge extraction efficiency. For the mixed CsPbBr3–Cs4PbBr6 counterpart (r ¼ 12:10), there is a drop in the PL intensity compared to the 12:8 film, due to the poor morphology and decreased grain sizes of the film. The corresponding TR-PL decay curves of the same samples are presented in Fig. 5b, from which the carrier lifetimes can be evaluated. The impurity level at the grain boundaries have a salient impact on the intralayer charge transfer process in the perovskites [37]. A longer carrier lifetime always corresponds to a slower intralayer carrier recombination rate caused by a lower impurity level in the perovskite film. According to previous report, the PL lifetime can be preferably fitted by a bi-exponential decay function [38,39]:

by a chemical bath deposition method at low temperature as described in the experimental section. The relatively rougher surface of the rutile TiO2 film conduces to increase the contact area between the ETL and perovskite, beneficial for the electron transport [41]. The thickness of the CsPbBr3 absorber layer is about 500 nm, which is the optimized thickness in our experiment as illustrated in Fig. S1. This thickness value is also comparable to that of the solution-processed and dual-source co-evaporated CsPbBr3 films applied in the PSCs [17,20]. Unfortu­ nately, the CuPc HTL is too thin to be observed. The energy diagram of the device is plotted in Fig. 5d. The cascade-like conduction band (CB) alignment between the FTO, c-TiO2 and CsPbBr3 is in favor of the electron extraction and collection, while the higher CB and valence band (VB) edges of CuPc than that of the CsPbBr3 perovskite are conducive to blocking electrons and extracting holes. The J-V characteristics recorded under one sunlight illumination (100 mW/cm2) of the best-performing devices composed of varied thickness ratios of CsBr to PbBr2 precursors are depicted in Fig. 6a. The statistical J-V parameters extracted from 25 devices for each kind of PSCs with varied r values are summarized in Table 2, from which the relatively high reproducibility of the as-prepared devices can be observed. The 12:8 film-based device demonstrates a champion PCE of 7.58%, with a Jsc of 7.59 mA cm 2, a Voc of 1.328 V and a fill factor (FF) of 0.752. This stems from the higher quality of the 12:8 film with enhanced crystallinity, phase-purity, light absorption ability and elon­ gated carrier lifetime caused by reduced trap state densities as discussed above. By contrast, the mixed CsPb2Br5–CsPbBr3 (r � 12:7) and Cs4PbBr6–CsPbBr3 film-based counterparts exhibit much poor perfor­ mance due to the existence of many impurities and defects. Therefore, the optimized thickness ratio of CsBr to PbBr2 precursor is 12:8. In addition, the PSC based on traditionally solution-processed CsPbBr3 film merely gained a PCE of 6.81%, with a Jsc of 7.14 mA cm 2, a Voc of 1.261 V and a FF of 0.756 (Fig. S2). It is clear that our optimized evaporated CsPbBr3-based devices exhibit a much better performance than the solution-processed counterparts, especially in Jsc and Voc, which is mainly due to the improved film quality and light absorption ability [17], etc. The incident photon-to-current conversion efficiency (IPCE) spectrum of the best-performing 12:8 PSC are drawn in Fig. 6b. It presents a high utilization of photos in the short-wavelength (300-530 nm) region and a cut-off wavelength at approximately 535 nm, in good agreement with its absorption edge in the UV–vis absorbance measurement (Fig. 2b). The integrated Jsc extracted from the IPCE curve is about 7.01 mA/cm2, which is close to the value obtain from the J-V curves. To further validate the reproducibility of the de­ vices, the PCE histogram of 25 individual PSCs based on the 12:8 films is also drawn in Fig. 6c. The PCE of the devices show a narrow distribution with an average value of 6.93%. The high reproducibility of the PSCs are ascribed to the highly homogeneous and uniform CsPbBr3 films, which further signifies the superiority of the sequential evaporation technique in preparing high-quality CsPbBr3 films. Given that evaporation tech­ nique is suitable for depositing large-area films (Fig. 1b), devices with large active areas (1 cm2) are also fabricated. A champion PCE of 6.21% is achieved, with a Jsc of 6.65 mA cm 2, a Voc of 1.375 V and a FF of 0.679 (Fig. 6d), which is higher than that of the devices (1 cm2) with dual-source evaporated CsPbBr3 films [20]. The histogram of PCEs for 25 devices with active area of 1 cm2 is added in Fig. S3. The PCEs of the large-area devices are mainly distributed in 5.1%–6.2% with an average value of 5.72%, also revealing a high reproducibility of the large-area CsPbBr3 devices. This mainly originates from the excellent uniformity of the as-evaporated CsPbBr3 film, as verified by SEM images in a large scope (Fig. 6e). Since device stability is a widely concerned issue before practical applications, the moisture and thermal stability tests of the 12:8 devices without encapsulation are conducted. When stored in ambient air with a relative humidity of ~40%, the device exhibits no any performance degradation (Fig. 6f). This can be mainly attributed to the protection from the highly hydrophobic and chemically stable CuPc HTL and

(4)

f(t) ¼ A1exp(-t/τ1) þ A2exp(-t/τ1)þK

where τ1 and τ2 are the slow and fast decay time constants with a cor­ responding fractional amplitude of A1, A2, and K is a constant for the base-line offset. In general, the slow decay closely relates to the trapassisted radiative recombination in the bulk perovskite phase, whilst the fast decay component is considered as a reflection of the carrier quenching process at the interface, respectively [40]. The resultant lifetime parameters extracted from TR-PL spectroscopy of the sequen­ tially evaporated films with different r values are listed in Table 1. The τ1 and τ2 of the CsPb2Br5-dominated film (r ¼ 12:4) are 15.80 and 1.51 ns, yielding a short average carrier lifetime of only 3.81 ns caused by severe non-radiative recombinations. With the augmentation of r but no more than 12:8, the carrier lifetimes get a pronounced improvement, in consistent with the slower carrier quenching observed from the steady-state PL spectra (Fig. 5b). The 12:8 film delivers the longest carrier lifetime of 16.79 ns, revealing the highest film quality of the CsPbBr3 perovskite with the least surface and bulk defects existing in the grain boundaries. This is favorable for the charge extraction and collection, contributing to a higher Jsc and Voc. Further increasing the ratio to 12:10, the average carrier lifetime declines to 5.96 ns due to the high impurity level in the film. Based on the as-evaporated CsPbBr3 films, the carbon-based planar heterojunction PSCs with a structure of FTO/c-TiO2/CsPbBr3/CuPc/ carbon are constructed in ambient air. The cross-sectional SEM image of the whole device is given in Fig. 5c. The Ni-doped c-TiO2 ETL is prepared Table 1 Lifetime parameters extracted from the TR-PL spectroscopy for the CsPbBr3 perovskites with different r values. Sample

τave [ns]

τ1 [ns]

A1

τ2 [ns]

A2

12:4 12:6 12:8 12:10

3.81 7.19 16.79 5.96

15.80 16.63 25.79 17.36

16.1% 13.6% 47.6% 12.5%

1.51 5.70 8.61 4.33

83.9% 86.4% 52.4% 87.5%

7

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Journal of Power Sources 443 (2019) 227269

Fig. 6. (a) J-V characteristics recorded under AM 1.5G (100 mW/cm2) illumination of the best-performing PSCs with varied r values. (b) IPCE spectra of the bestperforming 12:8 device. (c) Histograms of the PCEs for 25 devices based on the 12:8 CsPbBr3 films. (d) J-V characteristics of the best-performing 12:8 film-based device with a large active area of 1 cm2. (e) SEM image of 12:8 CsPbBr3 film in a large scope. (f) Durability of the 12:8 film-based device when stored in ambient air with a relative humidity of ~40% at room temperature (25 � C) for over 1000 h and under persistent thermal attack at 60 � C for one month. Table 2 Photovoltaic parameters derived from 25 devices for each kind of PSCs with varied r values. Device 12:4 12:5 12:6 12:7 12:8 12:9 12:10 12:11

average champion average champion average champion average champion average champion average champion average champion average champion

Jsc (mA/cm2)

Voc (V)

FF

PCE (%)

0.754 � 0.315 1.026 2.61 � 0.77 3.38 5.74 � 0.76 6.33 6.96 � 0.48 7.29 7.30 � 0.59 7.59 6.43 � 0.65 6.90 4.15 � 0.82 4.83 0.427 � 0.221 0.644

0.728 � 0.041 0.769 0.95 � 0.052 1.003 1.205 � 0.053 1.264 1.268 � 0.042 1.307 1.296 � 0.037 1.328 1.257 � 0.046 1.281 1.115 � 0.052 1.163 0.567 � 0.088 0.650

0.372 � 0.043 0.415 0.421 � 0.052 0.460 0.624 � 0.049 0.675 0.686 � 0.0.031 0.714 0.732 � 0.293 0.752 0.661 � 0.041 0.702 0.524 � 0.048 0.571 0.332 � 0.069 0.401

0.21 � 0.11 0.33 1.04 � 0.66 1.56 4.32 � 1.06 5.40 6.06 � 0.73 6.81 6.93 � 0.62 7.58 5.34 � 0.87 6.21 2.42 � 0.080 3.21 0.081 � 0.072 0.168

carbon counter electrode (CE) against the oxygen and moisture [31,42]. Upon persistent thermal attack at 60 � C (commonly used for solar cell stability testing [43]) for one month, the unencapsulated devices shows an excellent thermal durability, retaining 98.6% of its initial PCE. Apart from the protection of the thermally inert carbon electrode, the superior thermal stability can also be ascribed to the intrinsically thermal-stable CsPbBr3 itself with a high decomposition temperature over 467 � C [17]. Besides, the employment of cost-effective CuPc and commercial carbon to substitute expensive Spiro-OMeTAD or PTAA hole transport materials and noble metal electrode materials is beneficial for reducing the pro­ duction costs. Our devices show great potentials in practical applications with the merits of cost-effective, environmentally friendly and highly efficient and stable. Unfortunately, the light soaking stability of the as-fabricated devices is poorer compared to the moisture and thermal stability, only retaining about 78% after being exposed to one sunlight illumination for one week (Fig. S4). The relatively poor light soaking stability can be attributed to the employment of the photoactive TiO2

ETL [44] and it may be ameliorated by introducing other photo-stable ETLs (such as SnO2) in our future work. Some characterizations are further performed for the devices based on representative PbBr2-rich (r ¼ 12:7), pure CsPbBr3 (r ¼ 12:8) and CsBr-rich (r ¼ 12:9) films to assess the influence of the compositional structure on device performance. Since the stability of a PSC under the maximum power (MMP) point is more relevant to its operational state, the steady-state photocurrent density and PCE output of the devices are measured under MMP points. As displayed in Fig. 7a, all the devices show a fast response to the incident light once the light is turned on. The best-performing 12:7 device gets a steady-state current output of 5.59 mA/cm2 under a bias voltage of 1.08 V, yielding a PCE of 6.04%. By contrast, the 12:8 device achieves a much increased steady-state PCE output of 6.85%, with a current density of 6.11 mA/cm2 under a bias voltage of 1.12 V. For the CsBr-rich device, the steady-state PCE output drops to 5.29% with a low current density of 4.90 mA/cm2. It is note­ worthy that the PCE and current outputs of the 12:8 device stay stable 8

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Journal of Power Sources 443 (2019) 227269

Fig. 7. (a) Steady-state photocurrent and PCE output, (b) dark J-V characteristics and (c) OCVD curves of the 12:7, 12:8 and 12:9 PSCs. (d) Nyquist plots of the PSCs with varied r values measured under light illumination, with the equivalent circuit drawn in the inset. (e) Voc variations of the PSCs based on the 12:7, 12:8 and 12:9 films recorded under changed light intensities. (f) Mott-Schottky plots under different applied bias voltages extracted from the impedance analysis of the devices with varied r values.

during the whole test while that of the 12:7 and 12:8 devices show a slight decrease. This confirms the pure CsPbBr3 perovskite possesses a higher operational stability than the CsPb2Br5–CsPbBr3 and Cs4PbBr6–CsPbBr3 counterparts. The dark J-V characteristics of the same PSCs are given in Fig. 7b. It appears that the 12:8 device shows the smallest leakage current, suggesting the highest shunt resistance of the device. This may result from the higher film quality and more closely packed crystals of the pure CsPbBr3 film than that of the PbBr2-rich and CsBr-rich films. The decreased dark current contributes to an improve­ ment of the Jsc and FF [32], in good accordance with the J-V measure­ ments. The Cs4PbBr6–CsPbBr3 film-based device exhibits the highest leakage current due to the existence of much grain boundaries serving as current shunting pathways, corresponding to the worst device perfor­ mance. The open-circuit photovoltage decay (OCVD) measurements of these devices (Fig. 7c) reveal that the 12:8 cell obtains a higher intrinsic Voc and a longer Voc decay time than the other two counterparts. This indicates that the 12:8 device possesses a much longer carrier lifetime and lower interface recombination rate, attributed to less trap states in the pure CsPbBr3 film as analyzed above. Thus, the 12:8 devices are expected to obtain an enhanced FF and Voc [45], coinciding well with the J-V measurements. Electrochemical impedance spectroscopy (EIS) measurements are further performed to get insight into the interfacial charge transfer and recombination dynamics of the PSCs [46]. Nyquist plots measured from 2 MHz to 0.01 Hz at an applied bias voltage of 0.8 V under light illu­ mination are drawn in Fig. 7d. As can be observed, there exits two main semicircles in each plot. The first small semicircle in high frequency is widely considered as a reflection of the hole transport at the per­ ovskite/HTL or HTL/electrode interfaces, corresponding to a charge transfer resistance (Rct) in parallel with a HTM capacitance (CPE1). The large semicircle at low frequency is closely related to the carrier recombination process at the ETL/perovskite interface, reflecting a recombination resistance (Rrec) in parallel with a chemical capacitance (CPE2) [47]. Thus, the equivalent circuit can be simplified as the inset model in Fig. 7d. The 12:8 device obtains a lower fitted Rct of 61.2 Ω compared to that of the 12:7 (76.4 Ω) and 12:9 (122 Ω) counterparts.

This contributes to a faster charge transport and extraction efficiency for the 12:8 devices, resulting in a higher Jsc. The Rrec of the 12:8 cell is 614 Ω, much higher than that of the 12:7 (444 Ω) and 12:9 (375 Ω) devices, suggesting a more suppressed charge recombination rate inside the 12:8 device. In addition, the Rs derived from the starting point in the real part of the Nyquist plot for the PSCs based on CsPb2Br5–CsPbBr3, pure CsPbBr3 and Cs4PbBr6–CsPbBr3 films is about 36.8, 33.4 and 41.8 Ω [48], respectively. The lower Rs of the 12:8 device originates from a better interfacial contact and is conducive to the realization of a higher FF. The smallest CPE2 value of the optimized 12:8 device (52.8 nF/cm2) than that of the 12:7 (59.8 nF/cm2) and 12:8 (75.6 nF/cm2) PSCs also reveals the lowest interfacial charge accumulation existing in the 12:8 device. Furthermore, Nyquist plots of the PSCs with varied r values measured under dark condition are drawn in Fig. S5 with the resistance of them summarized in the inset. It is obvious that the resistance of the devices obtained under dark and illumination condition show a similar tendency, namely, the 12:8 devices possess the lowest Rs and Rct as well as highest Rrec among all the PSCs, favorable for the carrier extraction and collection process. Therefore, the pure CsPbBr3-based PSCs exhibit a much better performance than the mixed phase-based devices. The interfacial charge recombination process of a PSC can also be evaluated by ideality factor (n), which can be reflected by the Voc var­ iations upon the change of light intensities [49]. As previously reported, the Voc of a PSC usually decreases rapidly at a low light intensity (ψ ) of less than 10 mW/cm2, owing to the severe current shunting and trap-assisted recombinations in the devices [50]. The reduction in Voc of a PSC under an illumination higher than 10 mW/cm2 closely relates to the interfacial recombinations. The Voc vs. light intensity of the 12:7, 12:8 and 12:9 devices are plotted in Fig. 7e. The n can be deduced by the equation as follow [51]. n¼

q dVoc kB T dInðψ Þ

(5)

where q is the positive elementary charge, while kB and T are the Bolzman constant and absolute temperature, respectively. The calcu­ lated n of the 12:8 device is 1.43, much lower than that of the 12:7 (1.82) 9

Journal of Power Sources 443 (2019) 227269

X. Liu et al.

and 12:8 (2.36) counterparts. This certifies the lowest interfacial re­ combinations of the 12:8 devices, which can be mainly ascribed to the fewest defects in the pure CsPbBr3 films. Furthermore, capacitancevoltage measurements are carried out for the same devices under dark conditions. The built-in potential extracted from the intercept of the linear portion of the Mott-Schottky plots (Fig. 7f) for the 12:8 device is about 1.37 V, close to the Voc obtained from the J-V curves. By contrast, the 12:7 and 12:8 devices possess lower built-in potentials of 1.30 and 1.33 V, respectively. This enhancement also unveil the fewer trap states and charge carrier accumulations in the 12:8 device. The enlarged builtin potential is favorable for facilitating the separations of photo­ generated charge carriers and providing an extended depletion region to hinder the back transfer of electrons from the ETL to the CsPbBr3 light absorbers [52]. Thus, there exists lower interfacial recombination loss in the higher phase-purity CsPbBr3-based PSCs, yielding a higher voltage output.

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4. Conclusions We demonstrate a facile and environmentally friendly sequential evaporation technique to fabricate high-quality CsPbBr3 films. By pre­ cisely tuning the thickness ratio of CsBr to PbBr2 precursors to 12:8, highly phase-pure and crystallized CsPbBr3 films with low defect den­ sities are obtained. The as-prepared CsPbBr3 films exhibit dense and uniform morphologies with large average grain sizes of over 1 μm, extremely smooth surface and ultra-high light absorption ability. The PSC with an optimized thickness (500 nm) and compositional structure achieves a champion PCE of 7.58%, which is a superior efficiency for carbon-based CsPbBr3 PSCs with evaporated CsPbBr3 light absorbers. The as-prepared device also gains a high efficiency of 6.21% in a large area of 1 cm2. According to the applied characterizations, such as PL, EIS and OCVD measurements, the pure CsPbBr3-based devices exhibit much higher charge transfer rates and more suppressed interfacial carrier re­ combinations and accumulations than the mixed phase (CsPb2Br5–CsPbBr3 and Cs4PbBr6–CsPbBr3)-based counterparts. More­ over, the unencapsulated CsPbBr3 PSCs deliver a superb stability with almost no performance degradation when stored in air for 1000 h and upon persistent thermal attack at 60 � C for one month. This can be mainly attributed to the protection from the hydrophobic and chemi­ cally stable CuPc HTL and carbon CE. Our work provides a feasible approach to precisely fabricate CsPb2Br5, CsPbBr3 and Cs4PbBr6 perovskite, which also show great potential in the light emitting and photodetector fields. Declaration of competing interest There are no conflicts of interest to declare. Acknowledgements The authors acknowledge the financial support from the National Natural Science Foundation of China (Grant nos. 51675210, 51805195 and 51905203), the China Postdoctoral Science Foundation (Grant no. 2018M640691) and Fund from Science, Technology and Innovation Commission of Shenzhen Municipality (Grant no. JCYJ20170818165724025). We also appreciate the Analytical and Testing Center and Flexible Electronics Research Center of Huazhong University of Science and Technology for the SEM, XRD and XPS measurements. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.jpowsour.2019.227269.

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[29] [30]

[31]

[32] [33]

[34] [35] [36] [37] [38] [39]

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