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Enhanced thermal stability of electron transport layer-free perovskite solar cells via interface strain releasing
Journal of Power Sources 439 (2019) 227091
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Enhanced thermal stability of electron transport layer-free perovskite solar cells via interface strain releasing Peng Zhang a, Ting Zhang a, Yafei Wang a, Detao Liu a, Hao Xu b, Li Chen a, Yanbo Li c, Jiang Wu b, **, Zhi David Chen a, d, Shibin Li a, * a
School of Optoelectronic Science and Engineering, University of Electronic Science and Technology of China, Chengdu, 610054, China Department of Electronic and Electrical Engineering, University College London, Torrington Place, London, WC1E 7JE, UK Institute of Fundamental and Frontier Sciences, University of Electronic Science and Technology of China, Chengdu, 610054, China d Department of Electrical & Computer Engineering and Center for Nanoscale Science & Engineering, University of Kentucky, Lexington, KY, 40506, USA b c
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
� Lattice mismatch will promote the thermal degradation of perovskite materials. � Ba(OH)2 modification could release the lattice mismatch. � Ba(OH)2 modification could passivate the surface defects of Al-doped ZnO.
A R T I C L E I N F O
A B S T R A C T
Keywords: Perovskite solar cell Electron transport layer-free Lattice mismatch Thermal stability
The thermal decomposition of perovskite films on ZnO surfaces is generally believed to originate from specific surface states of ZnO and the impact from the lattice mismatch between ZnO and perovskite films on this process has long been ignored. In this research, the role of lattice mismatch in the thermal degradation process of cesiumcontaining perovskite films on Al doped ZnO (AZO) is studied. A Ba(OH)2 buffer layer on the surface of AZO is employed to release the lattice mismatch and suppress the thermal degradation of perovskite films resulted from ZnO. Consequently, perovskite films with enhanced thermal stability and crystalline properties are obtained. Meanwhile, the Ba(OH)2 films efficiently passivate the surface trap states and reduce the vacuum level of the AZO surfaces. On this basis, electron transport layer-free perovskite solar cells yield the best efficiency of 18.18% and the thermal stability is obviously improved.
1. Introduction Possessing the advantages of high electron mobility and various
fabrication methods, ZnO has been widely used as the electron transport layer (ETL) in perovskite solar cells (PSCs) [1–6]. However, the thermal instability of perovskite/ZnO interfaces has always been one of the
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (J. Wu), [email protected] (S. Li). https://doi.org/10.1016/j.jpowsour.2019.227091 Received 12 February 2019; Received in revised form 22 August 2019; Accepted 31 August 2019 Available online 1 October 2019 0378-7753/© 2019 Elsevier B.V. All rights reserved.
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problems that limit the performance of PSCs based on ZnO [7–9]. This apparent thermal instability primarily originates from the high surface reactivity of ZnO. At the same time, considering the long annealing time for perovskite materials containing dual/triple organic cations (100–120 � C, 30–60 min) and the normal operating temperature of solar cells (50–85 � C, depending on the regions), the thermal instability induced by ZnO inevitably affects the long term stability of PSCs [8,10, 11]. Moreover, the power conversion efficiency (PCE) of the PSCs based on ZnO is still inferior to that of the PSCs based on TiO2, which can be mainly attributed to the serious surface recombination between ZnO ETLs and perovskite films [6,12–14]. Although the use of Al doped ZnO (AZO) helps to alleviate the problems mentioned above, interfacial en gineering is still the most efficient method to suppress the thermal instability and recombination [1,14,15]. Compared with the widely researched stability and surface recom bination, the lattice mismatch between perovskite materials and metal oxides and its impact on the properties of perovskite films have long been ignored. The mismatch ratio between a substrate and a film can be roughly described by formula (1): δ ¼ (as-af)/af
degradation of perovskite films. An ultrathin Ba(OH)2 film, which had efficient hole blocking property, outstanding charge injection ability and relatively poor diffusive properties, was used to release the large lattice mismatch and obviously enhanced the thermal stability of the CsFA film on AZO [29,30]. ETL-free PSCs were fabricated with the structure of Ba(OH)2 modified AZO/CsFA/spiro-OMeTAD/Au. With the help of surface passivation and thickness optimization of the Ba(OH)2 film, the best PCE increased from 9.64% of the pristine AZO sample to a remarkable 18.18% of the modified AZO sample and the thermal sta bility of the PSCs was also enhanced. 2. Methods 2.1. Materials PbI2, PbBr2, MABr, FAI, CsI and all other chemicals were purchased from Sigma-Aldrich Co. and used as received. AZO (2 wt% Al2O3 in ZnO) covered glass substrates with a sheet resistance of 8 Ω/sq were pur chased from Huanan Xiangcheng Co.
(1)
2.2. Device preparation
, where δ is the lattice mismatch ratio between the substrate and the film, as is the lattice constant of the substrate and af is the lattice con stant of the film [16]. According to this formula, the absolute value of the lattice mismatch between ZnO and CH3NH3PbI3 (MAPbI3) is 48.52% and the value increases to 63.82% when MAPbI3 is replaced by – NH2PbI3 (FAPbI3). As for the widely used TiO2 (anatase), the NH2CH– mismatch between TiO2 and MAPbI3 (FAPbI3) is 40.17% (57.95%) [17–20]. In comparison, the lattice mismatch between a ScMgAlO4 substrate and an epitaxial GaN film is only 1.8% [21–23]. Obviously, the lattice mismatch between perovskite materials and the widely used metal oxide carrier transport materials is quite serious. A larger mismatch will cause larger stress and more interface defects, which leaves unavoidably negative impact on the properties of films [24]. Nevertheless, the effects of lattice mismatch on stability, crystallinity, structure and the resulting interface quality of perovskite films grown on the top of metal oxides are still not clear. Few researches on this issue were reported. Furthermore, as an efficient method to reduce the fabrication cost and optimize the manufacturing process of PSCs, eliminating the ETL has received increasingly more attention in recent years. The recorded PCE of ETL-free PSCs has exceeded 19% [25–28]. Such high perfor mance indicated that ETL-free PSCs can have comparable performance with the PSCs with traditional structures (electro de/ETL/perovskite/HTL, HTL ¼ hole transport layer). Meanwhile, in ETL-free PSCs, perovskite films are directly grown on transparent con ducting oxides (TCO). The widely used TCO in experimental preparation are still Indium tin oxide (ITO) and Fluorine doped tin oxide (FTO) at present. ITO has excellent properties, but the high cost and toxicity of indium have severely restricted its wide application. As to FTO, although the toxicity and cost are no longer problems, its optical and electrical properties are worse than that of ITO. Due to similar high transmissivity and low resistivity properties to ITO, nontoxic and low cost Aluminium-doped zinc oxide (AZO) has received growing attention in the recent years. However, due to the characteristics of ZnO mentioned above, the thermal instability of perovskite films and the large lattice mismatch between AZO and perovskite films still cannot be avoided. Here, we studied the impact of the lattice mismatch on the thermal stability of Cs0.05 (MA0.17FA0.83)0.95Pb(I0.83Br0.17)3 (marked as CsFA) films on the AZO surfaces. The thermal degradation processes were researched by means of microscopy and spectrum analysis. The results obtained from X-ray diffraction (XRD) and energy dispersive spec trometer (EDS) element mapping indicated that the large lattice mismatch of the CsFA/AZO interfaces accelerated the thermal
Patterned AZO glasses were cleaned sequentially by detergent, acetone, ethanol and deionized water in an ultrasonic bath and dried in N2 flow before use. All the substrates were transferred to a N2 filled glove box to deposit CsFA films using the single step method reported by Saliba et al. [8] FAI (1 M), PbI2 (1.1 M), MABr (0.2 M) and PbBr (0.2 M) were dissolved in the anhydrous N,N-Dimethylformamide (DMF)/Dimethyl sulfoxide (DMSO) mixed solution (volume ratio, 4:1) to form a mixed perovskite precursor solution. CsI solution (1.5 M in DMSO) was added into the mixed perovskite solution to achieve desired compositions. The perov skite precursor solution was then spin-coated on the AZO surface by a two-step coating progress at 1000 rpm and 5000 rpm for 10 s and 45 s, respectively. During the second step, 300 μL chlorobenzene was dropped on the rotating substrate 10 s before the end of the spin coating progress. The samples were then annealed at 100 � C for 1 h. After cooling down to room temperature (25 � C), the hole-transporting material was spin-coated from a solution consisted of 72.3 mg spiro-MeOTAD, 28.8 μl 4-tert-butylpyridine and 17.5 μl Li-TFSI solution (520 mg Li-TFSI in 1 ml acetonitrile) in 1 ml of chlorobenzene [31–35]. Finally, a 60 nm Au electrode was deposited by thermal evaporation. The active area for all solar cells was 0.04 cm2. For the Ba(OH)2 modified PSCs, Ba(OH)2 solutions with different concentrations (2 mg ml 1, 10 mg ml 1 and 15 mg ml 1) were spin coated on AZO surfaces at 3000 rpm for 40 s before the deposition of CsFA films. The Ba(OH)2 solutions were prepared by dissolving Ba(OH)2 in 2-Methoxyethanol. After spin coating, the substrates were annealed at 120 � C for 15 min. The ETL-free PSCs based on Ba(OH)2 modification were marked as B2, B10 and B15, corresponding to the concentrations of Ba(OH)2 solutions. 2.3. Measurement and characterization Current-voltage data was recorded by a Keithley 2400 source meter under AM 1.5G illumination (100 mW cm 2) from a Newport Oriel Solar 3A solar simulator. The AM 1.5G sun light was calibrated by a standard Si-solar cell each time before use. The external quantum efficiency (EQE) measurements were measured by a QEX10 solar cell quantum efficiency measurement system. X-Ray Diffraction (XRD) tests were measured using a Panalytical Empyrean system with Cu Kα radiation (Time per step ¼ 0.44 s per step, Scan step size ¼ 0.04� per step). The morphology and element mapping of perovskite films were obtained by a FEI-Inspect F50 field emission scanning electron microscope (FESEM) equipped with an energy dispersive spectrometer (EDS). Ultraviolet–visible (UV–Vis) absorption 2
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spectra were recorded on a Schimadzu UV-3101 PC UV–Vis spectro photometer (Scan range: 550 nm–850 nm). Fourier Transform Infrared Spectroscopy (FTIR) spectra were recorded on a Nicolet 6700 spectro photometer (Scan range: 1000 wavenumbers - 4000 wavenumbers). Photoluminescence (PL) was measured by a F-4600 FL fluorescence spectrophotometer (Excitation wavelength ¼ 510 nm). A ThermoFisher Ultraviolet Photoelectron Spectrometer (UPS) was used to determine the energy band position of CsFA. Surface roughness was measured by an atomic force microscope (AFM, Seiko SP13800 N). The distribution of surface groups was analyzed by a Kratos XSAM800 X-ray Photoelectron Spectrometer (XPS) system.
single crystal α-FAPbI3 (a ¼ 8.98 Å, c ¼ 10.1 Å), the lattice constant of the CsFA film decreases due to the introduction of Cs and MA cations, which was favorable to stabilize the perovskite structure [8,22,36,37]. At the same time, the characteristic peak corresponding to PbI2 around 12.6� was found in the spectra of the AZO and the B2 samples, which were resulted from the thermal decomposition of the CsFA films at the presence of the AZO substrates. The five dominant diffraction peaks of the CsFA film samples are exhibited in Fig. 1(b). In addition, to determine the presence of lattice mismatch, the diffraction patterns of standalone CsFA powder was used as the reference. The CsFA powder was prepared by scraping the asprepared perovskite films from the FTO substrates. Without the sub strate, the powder sample was free of lattice mismatch and stress [38]. As shown in the figure, the XRD peaks of the CsFA powder corre sponding to the (001), (002), (012), (022) and (003) planes are located at 14.053� , 28.366� , 31.804� , 40.574� and 43.148� , respectively. Comparatively, the diffraction peaks of all the CsFA films shift to lower angles and the shift of the high index planes is larger. For example, the diffraction peaks of the AZO sample shift to 13.99� , 28.32� , 31.761� , 40.559� and 43.12� , and for the B10 sample, the peaks shift to 14� , 28.321� , 31.797� , 40.56� and 43.138� . In addition, the full width half maximum (FWHM) of the (001) peak decreases from 0.472 of the AZO sample to 0.452 of the B15 and the same trend can be observed from the other peaks. The evolution of FWHM relates to the variation of crys tallinity and smaller FWHM demonstrates better crystalline quality. Meanwhile, based on the previous reports and Bragg formula, the change of grain size and grain boundaries could not affect the position of diffraction peaks. As a result, the shift of the diffraction angles corre sponds to the change of the crystal interplanar spacing, indicating the presence of stress in the films [38–40]. Although the peak shift and broadening were observed in all the films, it is worth noting that the changes are smaller after the Ba(OH)2 modification, revealing the release of stress and the improved crystal quality. Furthermore, in Fig. 1(a), no diffraction peak corresponding to
3. Results and discussion In order to investigate the thermal stability of CsFA on different substrates, CsFA films were deposited on the FTO, AZO and Ba(OH)2 modified AZO substrates, respectively. After film deposition, all the samples were subsequently annealed at 100 � C in dark under N2 atmo sphere for 1 h. Fig. S1 displays the photos of the CsFA films on different substrates at different annealing stages. Obviously, the CsFA film on the FTO and the modified AZO had no significant change throughout the entire annealing process. Whereas, for the CsFA film on the bare AZO, yellow spots began to appear after 30 min annealing. After 1 h anneal ing, the color of the CsFA film on the bare AZO changed from dark brown to yellowish, suggesting the serious decomposition of the perovskite materials. To study the impact of the lattice mismatch on the thermal stability of CsFA films, XRD experiments were carried out. Fig. 1(a) shows the diffractograms obtained from the CsFA films grown on different sub strates. Based on the different substrates used, the films are also marked as AZO, B2, B10, B15, respectively. The XRD patterns confirm the presence of the tetragonal perovskite phase in all the CsFA films and the average lattice constant of the AZO sample was calculated to be a ¼ 6.2228 Å, c ¼ 6.3096 Å. Compared with the lattice constant of the
Fig. 1. (a) The XRD patterns of the CsFA films on the AZO and Ba(OH)2 modified substrates, respectively. (b) The individual peaks of the CsFA films on different substrates compared with those of the scraped CsFA powder. Halder-Wagner plots of the CsFA film on (c) AZO and (d) B10, respectively. (e) The relationship between incident angle and detection depth of XRD. 3
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the Ba(OH)2 was found. Whereas, according to the XPS spectra of the AZO and the Ba(OH)2 surfaces shown in Fig. S2, two peaks located at 795.98 eV and 779.58 eV corresponding to Ba3d were clearly observed in the Ba(OH)2 modified AZO, which indicated the existence of the Ba (OH)2. The results of XRD and XPS revealed that the ultrathin Ba(OH)2 films on the AZO surfaces were amorphous. As a result, we speculated that the amorphous Ba(OH)2 buffer layers played an important role in alleviating lattice mismatch and hence releasing stress and improving the crystal quality of CsFA films. To get more details of the stress and crystallite size distribution in the films, Halder-Wagner method was performed, expressed as follows: � * �2 � � � � β 1 β* ε 2 ¼ (2) þ L δ2 δ 2
corresponding to the different incident angles are shown in Fig. S3). The details of the fitting data are shown in Table S1. For the AZO sample, with the increased detection depth, the stress ε increased from 0.3846 around the top surface to 0.5159 at the depth close to the substrate and the correlation length L decreases from 3.274 to 2.575. The same trend was observed in the data of the B10 sample, the stress ε increased from 0.2089 to 0.4926 while the correlation length L decreased from 3.943 to 2.685. Obviously, both of the two samples showed the largest stress and the smallest grain sizes in the parts of the film close to the substrates. The only difference between the AZO sample and the B10 sample was that the latter had smaller stress and larger correlation. Moreover, the surface morphology and grain size distribution of the AZO sample and the B10 sample are shown in Fig. 2. Both the CsFA films show good coverage with dense and pin hole free surfaces, but the morphology of each film is quite different. The grain size of the AZO sample is smaller and has a larger size distribution. In contrast, the B10 sample shows a larger grain size and has a relatively uniform grain size distribution ranging from 200 nm to 350 nm. Additionally, as shown in Fig. S4, the surface roughness of the AZO and the B10 sample is roughly the same. Therefore, the difference in the CsFA morphology can be attributed to the change of lattice mismatch after modification. The UV–Vis absorption spectroscopy was used to investigate the evolution of the CsFA films during the annealing process. Fig. 3(a) shows the spectrum of the CsFA film on the FTO substrate. The spectra of the CsFA films were nearly unchanged and so was the optical band gap at about 750 nm during the annealing process. For the AZO sample, as shown in Fig. 3(b), the absorption in the range from 550 nm to 700 nm decreased rapidly with the annealing time. The optical band gap decreased at the same time. These results clearly demonstrated that the CsFA itself appeared to be thermal stable under normal circumstances, but in the CsFA/AZO film, the CsFA gradually decomposed during the annealing process, being consistent with the previous reports [8,9,41].
2sinθ , where β* ¼ βcosθ λ , δ ¼ λ , β is the FWHM of the XRD peak (in order to get accurate numbers, the FWHM was obtained by fitting the diffraction peaks with Pseudo-Voigt functions and the widening of the diffraction peaks originated from the XRD equipment was excluded), θ is the Bragg angle, λ ¼ 1.5406 Å is the wavelength of the CuKα radiation, L is the correlation length (corresponding to the crystalline of the film) and ε is the stress parameter. The plots of (β*/δ) [2] as a function of (β*/δ [2]) corresponding to the 4 dominant peaks of the AZO and the B10 are shown in Fig. 1(c) and (d), respectively. The correlation length L was extracted from the slope of the linear fitting of the Halder-Wagner plot and stress ε was obtained from the Y-axis intercept. Generally, when the incident angle of X-Ray is fixed, the detection depth of XRD is fixed and the detection depth of XRD will increase with the increase of the X-Ray incident angle. As shown in Fig. 1(e), in order to study the distribution of the structural parameter L and stress parameter ε along the vertical di rection of the substrate, XRD test was used to characterize the perovskite film with different incident angles of 1� , 2� and 3.5� . (The XRD patterns
Fig. 2. The surface SEM images (left) and histogram of grain size (right) for the CsFA grown on (a), (b) the AZO and (c), (d) the B10, respectively. (e) The decomposition process of the CsFA on the AZO surface with lattice mismatch. 4
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Journal of Power Sources 439 (2019) 227091
Fig. 3. The UV–Vis absorbance spectra of the perovskite films on (a) FTO, (b) AZO, (c) B2, (d) B10, (e) B15, respectively. (f) FTIR spectra of the CsFA films on FTO, B10 and AZO after annealing, respectively.
For the B2 sample, as shown in Fig. 3(c), although the appearance of the film had no obvious change during the annealing process, the charac teristic absorption peak of perovskite material at around 750 nm degraded at the end of the annealing process. The evolution of the characteristic absorption peak indicated that the decomposition process in the B2 sample was inhibited. However, due to the low concentration of the precursor solution, the Ba(OH)2 film cannot fully cover the AZO surface, and the CsFA film still cannot withstand an hour of heat treat ment. Thus, with the increase of the concentration of the Ba(OH)2 pre cursor solution, as shown in Fig. 3(d) and (e), absorption spectra of the B10 and the B15 remained nearly unchanged and the optical band gap kept stable for the whole annealing process. These results demonstrated that the introduction of Ba(OH)2 interlayer with suitable thickness could effectively improve the stability of CsFA films on AZO surfaces. FTIR was used to further confirm this result. Fig. 3(f) shows the FTIR spectra of the FTO, B10 and AZO samples after the annealing process. As can be seen from the spectrum of the FTO sample, a strong stretching – N from FAþ at 1710 cm 1 can be observed. Meanwhile, vibration of C– two peaks at 3267 cm 1 and 3396 cm 1 corresponding to the stretching vibrations of N–H for FAþ can be identified. This FTIR spectrum revealed the existence of CsFA [42]. For the B10 sample, the spectrum was almost identical with that of the FTO sample, suggesting the B10 sample remained stable after annealing. However, the peak corresponding to hydroxyl group around 3500 cm 1 was observed in the spectrum of the AZO sample. The existence of hydroxyl group was an important indi cator of the decomposition of the CsFA film [15,43]. The FTIR experi ment further confirmed the Ba(OH)2 modification can efficiently suppress the thermal decomposition of the CsFA. To further test the result, the AZO sample and B10 sample were annealed at 130 � C for 30 min. The optical images in Fig. S5 visually exhibit the change of thermal stability after the Ba(OH)2 modification. Although both films began to decompose under annealing, the color evolution was much
more rapid in the CsFA film on the pristine AZO substrate than that in the B10 sample, further confirming the enhanced thermal stability of the CsFA/AZO interface after the Ba(OH)2 modification. The improved thermal stability after modification mainly originates from two aspects. On one hand, the Ba(OH)2 can suppress the reaction between the AZO and the CsFA films. Considering the annealing process in this study was carried out in a N2 filled glovebox, the effect of the air on the thermal stability of perovskites can be excluded. Therefore, the interaction between the ZnO and the organic part of perovskite materials became the main reason for the thermal decomposition of the CsFA [44]. After the modification, Ba2þ reacted with the AZO and a thin layer of O–Ba replaced the absorbed hydroxyl group at the surface of the AZO by the reaction: OH þ Baþ→-O-Ba þ Hþ [45]. The reduction of the surface hydroxyl group concentration can effectively improve the thermal sta bility of perovskite films on the ZnO surfaces [15]. Meanwhile, the re action will form a interface dipoles with negative charge toward AZO surface and positive charge toward the CsFA film, which will enhance the carrier transport efficiency. On the other hand, the release of lattice mismatch also improves the thermal stability of CsFA films. Without the modification, lattice mismatch leads to the accumulation of mismatch stress at the CsFA/AZO interface during the fabrication process of the CsFA film [46]. Defects and dislocations are formed in the CsFA film to release the stress. In addition, when the lattice mismatch is large enough and no modification is carried out, the threading dislocations are formed at the substrate surface and extend into the film or evolve to misfit dislocations at the interface. Fortunately, due to the soft nature of the hybrid perovskites, the dislocations and defects do not penetrate through the CsFA film and most of the stress and defects are concentrated near the substrate [47, 48]. This is in well consistent with the result of the Halder-Wagner plot and grain size distribution. A large amount of grain boundaries in perovskite films make thermal decomposition product diffuse more 5
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easily and more rapidly. The quick loss of decomposition product pro motes degradation of perovskites [49]. In addition, because of the weak chemical bonds at the grain boundaries and low ion migration activation energy, the halogen ions near the gain boundaries are moved apart by annealing and can migrate along the grain boundaries easily [50–52]. At the same time, according to the previous report, large stress can reduce the activation energy of ion migration, further promoting the ion migration in CsFA films [38]. Furthermore, MA cations in CsFA films have weaker interaction with the surrounding PbI6 octahedra than that of FA cation due to their small cation size. Owing to the poor thermal stability, MA cations are gradually released from the lattice structures through grain boundaries under high temperature annealing, leading to the collapse of perovskite lattice structures and the further degradation of CsFA films [7,9,53]. The schematic of this process is shown in Fig. 2 (e). As displayed in Fig. S6(a), showing the cross section of the CsFA film before annealing, the monolithic grains were observed and trended to extend from the bottom to the top. Comparatively, after annealing, (see Fig. S6(b), the uniform and monolithic morphology was changed to the amorphous one and the film thickness decreased at the same time. On the basis of previous description, combined with Fig. 2(e), the morphology evolution of CsFA film during annealing process could be attributed to the degradation of the CsFA film caused by the AZO sub strate and lattice mismatch. During the annealing process, the reaction between CsFA and AZO substrate will lead to the decomposition of CsFA material. At the same time, the lattice mismatch between CsFA film and AZO substrate will introduce more defects into the CsFA film and enhance the ion migration in perovskite film, both of which will pro mote the decomposition of CsFA materials and lead to the collapse of the CsFA lattice. At the same time, due to the main solid decomposition product was PbI2 and the smaller lattice size of PbI2, the morphology of CsFA film will change from monolithic to the amorphous after annealing. To further testify our assumption, EDS element mapping was carried out to determine the distribution of halogen at the surfaces of CsFA films. Fig. 4(a) shows the halogen distribution of the AZO sample, and
the inset shows where the measurement was taken. As can be seen, the Br and I accumulated at the grain boundaries after annealing, which was caused by the diffusion characteristics of the decomposition product and the ion migration paths. The high magnification SEM photo of the AZO sample after annealing is shown in Fig. 4(b). Compared with the SEM picture of the as prepared CsFA shown in Fig. 2(a), the grain size of the perovskite film increased obviously and the grain boundaries became larger, indicating the increased amount of PbI2 [7,54]. As previously stated, the mechanism of this phenomenon can be mainly ascribed to the diffusion of the decomposition product. In comparison, the Br and I distribution in the B10 sample are shown in Fig. 4(c). The Br and I were evenly distributed across the surface of the CsFA film and the concen tration of halogen in the detected area was much lower than that of the AZO sample, suggesting the improved stability of the CsFA film. The further zoom-in SEM photo of the B10 sample is shown in Fig. 4(d), in which, the morphology of the CsFA film remained roughly unchanged compared with Fig. 2(c). After understanding how the lattice mismatch affects the perovskite films, ETL free PSCs based on CsFA perovskite materials and AZO TCO were fabricated. The configuration and energy level diagrams of the ETL free PSC are shown in Fig. 5(a) and (b), respectively. The energy level of CsFA was determined by the UPS experiment, and the raw data is shown in Fig. S7. The J-V curves of the best performing PSCs based on the AZO, B2, B10 and B15 are exhibited in Fig. 5(c). Without modification, the best PCE based on the AZO sample was 9.64% with VOC ¼ 1.006 V, JSC ¼ 14.81 mA cm 2 and FF ¼ 64.68%. As the control, the B2 sample obtained the best PCE of 14.08%, with VOC ¼ 1.046 V, JSC ¼ 19.78 mA cm 2 and FF ¼ 67.58%. With the increase of the Ba (OH)2 film thickness, the best PCE of the B10 dramatically increased to 18.18%, with the VOC ¼ 1.082 V, JSC ¼ 22.88 mA cm 2 and FF ¼ 73.46%. However, with the film thickness continuously increasing, the best PCE of the B15 dropped to 15.78%, with VOC ¼ 1.043 V, JSC ¼ 21.60 mA cm 2 and FF ¼ 69.84%. The EQE spectra of PSCs based on the B2, B10 and B15 are shown in Fig. 5(d). All the PSCs show strong photo-response in the visible light range. The PSCs based on the AZO
Fig. 4. (a) Low magnification SEM pictures of the CsFA film on AZO after annealing and the distribution of Br and I at the surface of the CsFA film on AZO after annealing, respectively. (b) High magnification SEM picture of the CsFA film on AZO after annealing. (c) Low magnification SEM pictures of the CsFA film on B10 after annealing and the distribution of Br and I at the surface of the CsFA film on B10 after annealing, respectively. (d) High magnification SEM picture of the CsFA film on B10 after annealing. 6
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Fig. 5. (a) The structure of the ETL free PSCs. (b) The energy level diagram of the ETL free PSCs. (c) The J-V curves of the best performing PSCs. (d) The EQE spectra and the integrated JSC of the best performing PSCs. (e) J-V curves of the B10 recorded in reverse (from VOC to JSC) and forward (form JSC to VOC) scanning directions. (f) The power output under maximum power point tracking for 25 s. The O1S spectra of pristine (g) AZO and (h) AZO with modification. (i) PCE of the FTO, B10 and AZO sample heated at 100 � C in N2 atmosphere.
show the weakest spectrum intensity in the whole visible light range compared with the other sets of the PSCs. The JSC values extracted from the EQE data were 14.68 mA cm 2 (AZO), 19.64 mA cm 2 (B2), 21.59 mA cm 2 (B10) and 20.61 mA cm 2 (B15), respectively, which are in good agreement with the results of the J-V tests. The J-V curves of the B10 sample obtained from different scan directions were shown in Fig. 5 (e). Due to the improved interface quality and crystal property of the CsFA film, there was only a small difference in the PCE obtained from the forward scan (17.62%) and the revised scan (18.18%). Fig. 5(f) depicted the stabilized power output under the maximum power point tracking reaching 18.1% (at 0.857 V), indicating a good agreement with J-V tests. As previously stated, the poor performance of the AZO and the B2 samples can be mainly attributed to the thermal decomposition of the CsFA. When the Ba(OH)2 layer was thick enough to cover the surface of the AZO, the thermal decomposition of the CsFA was suppressed. As a result, the performance of B10 and B15 increased. On the other hand, Ba (OH)2 can efficiently passivate the surface traps and oxygen vacancies of AZO surfaces at the same time. The XPS spectra of the AZO film and the
Ba(OH)2 modified AZO film are shown in Fig. 5(g) and (h). In Fig. 5(g), the O1S peak of the pristine AZO sample was deconvoluted into two peaks located at 531.50 eV and 530.28 eV, respectively. The low binding energy (530.28 eV) corresponded to the oxygen in ZnO lattice and the higher binding energy (531.50 eV) was originated from the oxygen va cancies and chemisorbed oxygen. The area ratio of these two peaks (S530.28/S531.50) is 0.8. On the other hand, as shown in Fig. 5(h), the O1S peak of the modified AZO film can also be deconvoluted into two peaks located at 530.92 eV to and 531.33 eV, corresponding to the oxygen in ZnO lattice and the oxygen vacancy/chemisorbed oxygen, respectively. The area ratio of these two peaks (S530.92/S531.33) is 0.89. The larger area ratio in the AZO with modification indicated that the concentration of oxygen vacancies and chemisorbed oxygen at the AZO surface was efficiently suppressed [55–57]. Notably, the reaction between Ba(OH)2 and AZO leads to the for mation of interface dipoles, lowering the vacuum level of AZO surfaces and improving the electron extraction efficiency [45,58]. The PL spectra shown in Fig. S8(a) further confirm this feature. As the figure exhibited, the PL peaks located at 772 nm corresponding to CsFA were observed. 7
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Due to the different timescale of electron extraction and radiative recombination, the PL intensity was largely determined by the carrier extraction [15]. Therefore, the weakest PL intensity of the B10 sample indicated that the electron extraction efficiency was higher than the other two samples [59,60]. To demonstrate the improved crystal properties of the CsFA film is another reason for the improvement of PCE, the I–V curves of the uni polar devices (AZO/CsFA/PCBM/Ag) only with the electron charge carriers were recorded [61]. As can be seen from Fig. S8(b), the current increased linearly at the low bias voltage region. While at the high bias voltage region, the current increased with a larger slope. The transition point is the ohmic to trap filled limit transition point (VTFL), which has a linear relationship with trap density (nt). The trap-state is given by formula (3): 2
VTFL ¼
ent L 2εε0
Table 1 The performance of ETL free PSCs from other literatures. Perovskite materials
Structure
PCE
Reference
CH3NH3PbI3 CH3NH3PbI3-xClx CH3NH3PbI3 CH3NH3PbI3 CH3NH3PbI3 CH3NH3PbI3 CH3NH3PbI3
ITO/Perovskite/HTL/Ag FTO/Perovskite/HTL/Au FTO/Perovskite:C70/HTL/Au FTO/Perovskite:C60/HTL/Au ITO/Perovskite/HTL/MoO3/Ag FTO/Perovskite/HTL/Au FTO/Perovskite/HTL/Au
13.5% 14.14% 13.60% 14.30% 10.80% 18.20% 19.22%
[64] 27 [65] [66] [67] 25 26
4. Conclusion The large lattice mismatch induced stress led to the accumulation of defects at the CsFA/AZO interfaces and deterioration of crystal quality of perovskite films, further accelerating the thermal decomposition process caused by the reactive AZO surfaces. With the Ba(OH)2 modi fication, the lattice mismatch was partially relaxed and the thermal instability originated from the AZO was suppressed as well, which improved the crystal quality and thermal stability of CsFA films. On this basis, the thermal stability of ETL-free PSCs based on the AZO was enhanced as well. Additionally, due to its efficient hole blocking prop erty and surface reaction with the AZO, the Ba(OH)2 films efficiently passivated the surface trap states and reduced the vacuum level of the AZO surfaces. Finally, the best PCE of 18.18% was achieved by opti mizing the thickness of the Ba(OH)2 film.
(3)
, where L is the thickness of the CsFA, ε and ε0 are the relative dielectric constant of the CsFA and the vacuum permittivity, respectively. Since the thickness of all the CsFA films is almost the same, the trap density nt is proportional to the VTFL. As calculated from the figure, the VTFL value of the B2, B10 and B15 is 0.55 V, 0.51 V and 0.35 V, respectively, revealing the lower trap densities in the B10 and B15 samples. Mean while, it was noticed that the B15 had the thickest Ba(OH)2 film and lowest trap density but not the best PCE, which can be attributed to the larger series resistance introduced by the thicker Ba(OH)2 interlayer. Shunt Resistance (Rsh) and Series Resistance (Rs) of B10 were 1982869 Ω and 131.9 Ω, respectively. As control, the Rsh and Rs of B15 were 18629 Ω and 136.8 Ω, respectively, which is consistent with our inference. As a result, the Ba(OH)2 modification can improve the ther mal stability of the CsFA on the AZO together with enhancing the per formance of the ETL-free PSCs. Furthermore, in Fig. 5(i), we tested the thermal stability of ETL-free PSCs (the AZO, the B10 and the FTO used as the control) in a N2 at mosphere held at 100 � C. J-V, which was carried out periodically to extract the photovoltaic parameters of the PSCs. As the figure exhibited, the PCE of all the samples began to decrease under thermal stress within a few minutes. Notably, as described above, the CsFA film can remain stable on the FTO surface under annealing. As a result, the performance degradation in the FTO sample can be mainly ascribed to the deterio ration of spiro-OMeTAD and the diffusion of metal electrodes under high temperature [51,62]. As to the AZO sample, the thermal instability of CsFA/AZO interface further exacerbated the thermal stability of PSC and after about 50 min the device was no longer functioning. Clearly, with the help of the Ba(OH)2 modification, the degradation induced by AZO was suppressed and the thermal stability of PSC was also improved. Meanwhile, considering the research on the thermal stability of PSCs based on ZnO or AZO ETL is not a new topic, there have been many reports on this area. Zhao et al. fabricated ETL-free PSCs based AZO and MAPbI3. According their research, the use of AZO efficiently improved the thermal stability of PSCs and a best PCE of 12.6% was obtained [13]. Bush et al. used ITO top electrode to encapsulate the PSCs based on ZnO nanoparticles. The PSCs they obtained could maintain roughly stable after long time annealing under 100 � C and a best PCE of 13.5% [63]. Song et al. researched the thermal stability of perovskite/ZnO film and developed Cs containing triple cation PSCs based on ZnO ETL. In their research, they observed similar CsFAMA/ZnO degradation process. By reducing the annealing temperature time, they obtained a best PCE of 18.9% [9]. Obviously, in comparison with previous researches, the ETL-free PSCs we made obtained enhanced thermal stability with simplified structure and fabrication process. As exhibited in Table 1, the performance of our ETL free PSCs is competitive to the previous reported ones.
Conflicts of interest The authors declare no conflict of interest. Acknowledgements This work was supported by National Natural Science Foundation of China under Grant Nos. 61421002, 61574029, 61471085, and 61874150. This work was also partially supported by University of Kentucky. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.jpowsour.2019.227091. References [1] K. Mahmood, B.S. Swain, A. Amassian, Adv. Energy Mater. 5 (2015), 1500568. [2] J. You, L. Meng, T.B. Song, T.F. Guo, Y.M. Yang, W.H. Chang, Z. Hong, H. Chen, H. Zhou, Q. Chen, Y. Liu, N. De Marco, Y. Yang, Nat. Nanotechnol. 11 (2016) 75–81. [3] D. Liu, T.L. Kelly, Nat. Photonics 8 (2013) 133–138. [4] D.-Y. Son, J.-H. Im, H.-S. Kim, N.-G. Park, J. Phys. Chem. C 118 (2014) 16567–16573. [5] S. Li, P. Zhang, Y. Wang, H. Sarvari, D. Liu, J. Wu, Y. Yang, Z. Wang, Z.D. Chen, Nano Res. 10 (2016) 1092–1103. [6] S. Li, P. Zhang, H. Chen, Y. Wang, D. Liu, J. Wu, H. Sarvari, Z.D. Chen, J. Power Sources 342 (2017) 990–997. [7] B. Conings, J. Drijkoningen, N. Gauquelin, A. Babayigit, J. D’Haen, L. D’Olieslaeger, A. Ethirajan, J. Verbeeck, J. Manca, E. Mosconi, F.D. Angelis, H.G. Boyen, Adv. Energy Mater. 5 (2015), 1500477. [8] M. Saliba, T. Matsui, J.Y. Seo, K. Domanski, J.P. Correa-Baena, M.K. Nazeeruddin, S.M. Zakeeruddin, W. Tress, A. Abate, A. Hagfeldt, M. Gratzel, Energy Environ. Sci. 9 (2016) 1989–1997. [9] J. Song, L. Liu, X.-F. Wang, G. Chen, W. Tian, T. Miyasaka, J. Mater. Chem. 5 (2017) 13439–13447. [10] N.J. Jeon, J.H. Noh, W.S. Yang, Y.C. Kim, S. Ryu, J. Seo, S.I. Seok, Nature 517 (2015) 476–480. [11] W.S. Yang, J.H. Noh, N.J. Jeon, Y.C. Kim, S. Ryu, J. Seo, S.I. Seok, Science 348 (2015) 1234–1237. [12] Z.-L. Tseng, C.-H. Chiang, S.-H. Chang, C.-G. Wu, Nano Energy 28 (2016) 311–318. [13] X. Zhao, H. Shen, Y. Zhang, X. Li, X. Zhao, M. Tai, J. Li, J. Li, X. Li, H. Lin, ACS Appl. Mater. Interfaces 8 (2016) 7826–7833.
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Journal of Power Sources journal homepage: www.elsevier.com/locate/jpowsour
Enhanced thermal stability of electron transport layer-free perovskite solar cells via interface strain releasing Peng Zhang a, Ting Zhang a, Yafei Wang a, Detao Liu a, Hao Xu b, Li Chen a, Yanbo Li c, Jiang Wu b, **, Zhi David Chen a, d, Shibin Li a, * a
School of Optoelectronic Science and Engineering, University of Electronic Science and Technology of China, Chengdu, 610054, China Department of Electronic and Electrical Engineering, University College London, Torrington Place, London, WC1E 7JE, UK Institute of Fundamental and Frontier Sciences, University of Electronic Science and Technology of China, Chengdu, 610054, China d Department of Electrical & Computer Engineering and Center for Nanoscale Science & Engineering, University of Kentucky, Lexington, KY, 40506, USA b c
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
� Lattice mismatch will promote the thermal degradation of perovskite materials. � Ba(OH)2 modification could release the lattice mismatch. � Ba(OH)2 modification could passivate the surface defects of Al-doped ZnO.
A R T I C L E I N F O
A B S T R A C T
Keywords: Perovskite solar cell Electron transport layer-free Lattice mismatch Thermal stability
The thermal decomposition of perovskite films on ZnO surfaces is generally believed to originate from specific surface states of ZnO and the impact from the lattice mismatch between ZnO and perovskite films on this process has long been ignored. In this research, the role of lattice mismatch in the thermal degradation process of cesiumcontaining perovskite films on Al doped ZnO (AZO) is studied. A Ba(OH)2 buffer layer on the surface of AZO is employed to release the lattice mismatch and suppress the thermal degradation of perovskite films resulted from ZnO. Consequently, perovskite films with enhanced thermal stability and crystalline properties are obtained. Meanwhile, the Ba(OH)2 films efficiently passivate the surface trap states and reduce the vacuum level of the AZO surfaces. On this basis, electron transport layer-free perovskite solar cells yield the best efficiency of 18.18% and the thermal stability is obviously improved.
1. Introduction Possessing the advantages of high electron mobility and various
fabrication methods, ZnO has been widely used as the electron transport layer (ETL) in perovskite solar cells (PSCs) [1–6]. However, the thermal instability of perovskite/ZnO interfaces has always been one of the
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (J. Wu), [email protected] (S. Li). https://doi.org/10.1016/j.jpowsour.2019.227091 Received 12 February 2019; Received in revised form 22 August 2019; Accepted 31 August 2019 Available online 1 October 2019 0378-7753/© 2019 Elsevier B.V. All rights reserved.
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problems that limit the performance of PSCs based on ZnO [7–9]. This apparent thermal instability primarily originates from the high surface reactivity of ZnO. At the same time, considering the long annealing time for perovskite materials containing dual/triple organic cations (100–120 � C, 30–60 min) and the normal operating temperature of solar cells (50–85 � C, depending on the regions), the thermal instability induced by ZnO inevitably affects the long term stability of PSCs [8,10, 11]. Moreover, the power conversion efficiency (PCE) of the PSCs based on ZnO is still inferior to that of the PSCs based on TiO2, which can be mainly attributed to the serious surface recombination between ZnO ETLs and perovskite films [6,12–14]. Although the use of Al doped ZnO (AZO) helps to alleviate the problems mentioned above, interfacial en gineering is still the most efficient method to suppress the thermal instability and recombination [1,14,15]. Compared with the widely researched stability and surface recom bination, the lattice mismatch between perovskite materials and metal oxides and its impact on the properties of perovskite films have long been ignored. The mismatch ratio between a substrate and a film can be roughly described by formula (1): δ ¼ (as-af)/af
degradation of perovskite films. An ultrathin Ba(OH)2 film, which had efficient hole blocking property, outstanding charge injection ability and relatively poor diffusive properties, was used to release the large lattice mismatch and obviously enhanced the thermal stability of the CsFA film on AZO [29,30]. ETL-free PSCs were fabricated with the structure of Ba(OH)2 modified AZO/CsFA/spiro-OMeTAD/Au. With the help of surface passivation and thickness optimization of the Ba(OH)2 film, the best PCE increased from 9.64% of the pristine AZO sample to a remarkable 18.18% of the modified AZO sample and the thermal sta bility of the PSCs was also enhanced. 2. Methods 2.1. Materials PbI2, PbBr2, MABr, FAI, CsI and all other chemicals were purchased from Sigma-Aldrich Co. and used as received. AZO (2 wt% Al2O3 in ZnO) covered glass substrates with a sheet resistance of 8 Ω/sq were pur chased from Huanan Xiangcheng Co.
(1)
2.2. Device preparation
, where δ is the lattice mismatch ratio between the substrate and the film, as is the lattice constant of the substrate and af is the lattice con stant of the film [16]. According to this formula, the absolute value of the lattice mismatch between ZnO and CH3NH3PbI3 (MAPbI3) is 48.52% and the value increases to 63.82% when MAPbI3 is replaced by – NH2PbI3 (FAPbI3). As for the widely used TiO2 (anatase), the NH2CH– mismatch between TiO2 and MAPbI3 (FAPbI3) is 40.17% (57.95%) [17–20]. In comparison, the lattice mismatch between a ScMgAlO4 substrate and an epitaxial GaN film is only 1.8% [21–23]. Obviously, the lattice mismatch between perovskite materials and the widely used metal oxide carrier transport materials is quite serious. A larger mismatch will cause larger stress and more interface defects, which leaves unavoidably negative impact on the properties of films [24]. Nevertheless, the effects of lattice mismatch on stability, crystallinity, structure and the resulting interface quality of perovskite films grown on the top of metal oxides are still not clear. Few researches on this issue were reported. Furthermore, as an efficient method to reduce the fabrication cost and optimize the manufacturing process of PSCs, eliminating the ETL has received increasingly more attention in recent years. The recorded PCE of ETL-free PSCs has exceeded 19% [25–28]. Such high perfor mance indicated that ETL-free PSCs can have comparable performance with the PSCs with traditional structures (electro de/ETL/perovskite/HTL, HTL ¼ hole transport layer). Meanwhile, in ETL-free PSCs, perovskite films are directly grown on transparent con ducting oxides (TCO). The widely used TCO in experimental preparation are still Indium tin oxide (ITO) and Fluorine doped tin oxide (FTO) at present. ITO has excellent properties, but the high cost and toxicity of indium have severely restricted its wide application. As to FTO, although the toxicity and cost are no longer problems, its optical and electrical properties are worse than that of ITO. Due to similar high transmissivity and low resistivity properties to ITO, nontoxic and low cost Aluminium-doped zinc oxide (AZO) has received growing attention in the recent years. However, due to the characteristics of ZnO mentioned above, the thermal instability of perovskite films and the large lattice mismatch between AZO and perovskite films still cannot be avoided. Here, we studied the impact of the lattice mismatch on the thermal stability of Cs0.05 (MA0.17FA0.83)0.95Pb(I0.83Br0.17)3 (marked as CsFA) films on the AZO surfaces. The thermal degradation processes were researched by means of microscopy and spectrum analysis. The results obtained from X-ray diffraction (XRD) and energy dispersive spec trometer (EDS) element mapping indicated that the large lattice mismatch of the CsFA/AZO interfaces accelerated the thermal
Patterned AZO glasses were cleaned sequentially by detergent, acetone, ethanol and deionized water in an ultrasonic bath and dried in N2 flow before use. All the substrates were transferred to a N2 filled glove box to deposit CsFA films using the single step method reported by Saliba et al. [8] FAI (1 M), PbI2 (1.1 M), MABr (0.2 M) and PbBr (0.2 M) were dissolved in the anhydrous N,N-Dimethylformamide (DMF)/Dimethyl sulfoxide (DMSO) mixed solution (volume ratio, 4:1) to form a mixed perovskite precursor solution. CsI solution (1.5 M in DMSO) was added into the mixed perovskite solution to achieve desired compositions. The perov skite precursor solution was then spin-coated on the AZO surface by a two-step coating progress at 1000 rpm and 5000 rpm for 10 s and 45 s, respectively. During the second step, 300 μL chlorobenzene was dropped on the rotating substrate 10 s before the end of the spin coating progress. The samples were then annealed at 100 � C for 1 h. After cooling down to room temperature (25 � C), the hole-transporting material was spin-coated from a solution consisted of 72.3 mg spiro-MeOTAD, 28.8 μl 4-tert-butylpyridine and 17.5 μl Li-TFSI solution (520 mg Li-TFSI in 1 ml acetonitrile) in 1 ml of chlorobenzene [31–35]. Finally, a 60 nm Au electrode was deposited by thermal evaporation. The active area for all solar cells was 0.04 cm2. For the Ba(OH)2 modified PSCs, Ba(OH)2 solutions with different concentrations (2 mg ml 1, 10 mg ml 1 and 15 mg ml 1) were spin coated on AZO surfaces at 3000 rpm for 40 s before the deposition of CsFA films. The Ba(OH)2 solutions were prepared by dissolving Ba(OH)2 in 2-Methoxyethanol. After spin coating, the substrates were annealed at 120 � C for 15 min. The ETL-free PSCs based on Ba(OH)2 modification were marked as B2, B10 and B15, corresponding to the concentrations of Ba(OH)2 solutions. 2.3. Measurement and characterization Current-voltage data was recorded by a Keithley 2400 source meter under AM 1.5G illumination (100 mW cm 2) from a Newport Oriel Solar 3A solar simulator. The AM 1.5G sun light was calibrated by a standard Si-solar cell each time before use. The external quantum efficiency (EQE) measurements were measured by a QEX10 solar cell quantum efficiency measurement system. X-Ray Diffraction (XRD) tests were measured using a Panalytical Empyrean system with Cu Kα radiation (Time per step ¼ 0.44 s per step, Scan step size ¼ 0.04� per step). The morphology and element mapping of perovskite films were obtained by a FEI-Inspect F50 field emission scanning electron microscope (FESEM) equipped with an energy dispersive spectrometer (EDS). Ultraviolet–visible (UV–Vis) absorption 2
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spectra were recorded on a Schimadzu UV-3101 PC UV–Vis spectro photometer (Scan range: 550 nm–850 nm). Fourier Transform Infrared Spectroscopy (FTIR) spectra were recorded on a Nicolet 6700 spectro photometer (Scan range: 1000 wavenumbers - 4000 wavenumbers). Photoluminescence (PL) was measured by a F-4600 FL fluorescence spectrophotometer (Excitation wavelength ¼ 510 nm). A ThermoFisher Ultraviolet Photoelectron Spectrometer (UPS) was used to determine the energy band position of CsFA. Surface roughness was measured by an atomic force microscope (AFM, Seiko SP13800 N). The distribution of surface groups was analyzed by a Kratos XSAM800 X-ray Photoelectron Spectrometer (XPS) system.
single crystal α-FAPbI3 (a ¼ 8.98 Å, c ¼ 10.1 Å), the lattice constant of the CsFA film decreases due to the introduction of Cs and MA cations, which was favorable to stabilize the perovskite structure [8,22,36,37]. At the same time, the characteristic peak corresponding to PbI2 around 12.6� was found in the spectra of the AZO and the B2 samples, which were resulted from the thermal decomposition of the CsFA films at the presence of the AZO substrates. The five dominant diffraction peaks of the CsFA film samples are exhibited in Fig. 1(b). In addition, to determine the presence of lattice mismatch, the diffraction patterns of standalone CsFA powder was used as the reference. The CsFA powder was prepared by scraping the asprepared perovskite films from the FTO substrates. Without the sub strate, the powder sample was free of lattice mismatch and stress [38]. As shown in the figure, the XRD peaks of the CsFA powder corre sponding to the (001), (002), (012), (022) and (003) planes are located at 14.053� , 28.366� , 31.804� , 40.574� and 43.148� , respectively. Comparatively, the diffraction peaks of all the CsFA films shift to lower angles and the shift of the high index planes is larger. For example, the diffraction peaks of the AZO sample shift to 13.99� , 28.32� , 31.761� , 40.559� and 43.12� , and for the B10 sample, the peaks shift to 14� , 28.321� , 31.797� , 40.56� and 43.138� . In addition, the full width half maximum (FWHM) of the (001) peak decreases from 0.472 of the AZO sample to 0.452 of the B15 and the same trend can be observed from the other peaks. The evolution of FWHM relates to the variation of crys tallinity and smaller FWHM demonstrates better crystalline quality. Meanwhile, based on the previous reports and Bragg formula, the change of grain size and grain boundaries could not affect the position of diffraction peaks. As a result, the shift of the diffraction angles corre sponds to the change of the crystal interplanar spacing, indicating the presence of stress in the films [38–40]. Although the peak shift and broadening were observed in all the films, it is worth noting that the changes are smaller after the Ba(OH)2 modification, revealing the release of stress and the improved crystal quality. Furthermore, in Fig. 1(a), no diffraction peak corresponding to
3. Results and discussion In order to investigate the thermal stability of CsFA on different substrates, CsFA films were deposited on the FTO, AZO and Ba(OH)2 modified AZO substrates, respectively. After film deposition, all the samples were subsequently annealed at 100 � C in dark under N2 atmo sphere for 1 h. Fig. S1 displays the photos of the CsFA films on different substrates at different annealing stages. Obviously, the CsFA film on the FTO and the modified AZO had no significant change throughout the entire annealing process. Whereas, for the CsFA film on the bare AZO, yellow spots began to appear after 30 min annealing. After 1 h anneal ing, the color of the CsFA film on the bare AZO changed from dark brown to yellowish, suggesting the serious decomposition of the perovskite materials. To study the impact of the lattice mismatch on the thermal stability of CsFA films, XRD experiments were carried out. Fig. 1(a) shows the diffractograms obtained from the CsFA films grown on different sub strates. Based on the different substrates used, the films are also marked as AZO, B2, B10, B15, respectively. The XRD patterns confirm the presence of the tetragonal perovskite phase in all the CsFA films and the average lattice constant of the AZO sample was calculated to be a ¼ 6.2228 Å, c ¼ 6.3096 Å. Compared with the lattice constant of the
Fig. 1. (a) The XRD patterns of the CsFA films on the AZO and Ba(OH)2 modified substrates, respectively. (b) The individual peaks of the CsFA films on different substrates compared with those of the scraped CsFA powder. Halder-Wagner plots of the CsFA film on (c) AZO and (d) B10, respectively. (e) The relationship between incident angle and detection depth of XRD. 3
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the Ba(OH)2 was found. Whereas, according to the XPS spectra of the AZO and the Ba(OH)2 surfaces shown in Fig. S2, two peaks located at 795.98 eV and 779.58 eV corresponding to Ba3d were clearly observed in the Ba(OH)2 modified AZO, which indicated the existence of the Ba (OH)2. The results of XRD and XPS revealed that the ultrathin Ba(OH)2 films on the AZO surfaces were amorphous. As a result, we speculated that the amorphous Ba(OH)2 buffer layers played an important role in alleviating lattice mismatch and hence releasing stress and improving the crystal quality of CsFA films. To get more details of the stress and crystallite size distribution in the films, Halder-Wagner method was performed, expressed as follows: � * �2 � � � � β 1 β* ε 2 ¼ (2) þ L δ2 δ 2
corresponding to the different incident angles are shown in Fig. S3). The details of the fitting data are shown in Table S1. For the AZO sample, with the increased detection depth, the stress ε increased from 0.3846 around the top surface to 0.5159 at the depth close to the substrate and the correlation length L decreases from 3.274 to 2.575. The same trend was observed in the data of the B10 sample, the stress ε increased from 0.2089 to 0.4926 while the correlation length L decreased from 3.943 to 2.685. Obviously, both of the two samples showed the largest stress and the smallest grain sizes in the parts of the film close to the substrates. The only difference between the AZO sample and the B10 sample was that the latter had smaller stress and larger correlation. Moreover, the surface morphology and grain size distribution of the AZO sample and the B10 sample are shown in Fig. 2. Both the CsFA films show good coverage with dense and pin hole free surfaces, but the morphology of each film is quite different. The grain size of the AZO sample is smaller and has a larger size distribution. In contrast, the B10 sample shows a larger grain size and has a relatively uniform grain size distribution ranging from 200 nm to 350 nm. Additionally, as shown in Fig. S4, the surface roughness of the AZO and the B10 sample is roughly the same. Therefore, the difference in the CsFA morphology can be attributed to the change of lattice mismatch after modification. The UV–Vis absorption spectroscopy was used to investigate the evolution of the CsFA films during the annealing process. Fig. 3(a) shows the spectrum of the CsFA film on the FTO substrate. The spectra of the CsFA films were nearly unchanged and so was the optical band gap at about 750 nm during the annealing process. For the AZO sample, as shown in Fig. 3(b), the absorption in the range from 550 nm to 700 nm decreased rapidly with the annealing time. The optical band gap decreased at the same time. These results clearly demonstrated that the CsFA itself appeared to be thermal stable under normal circumstances, but in the CsFA/AZO film, the CsFA gradually decomposed during the annealing process, being consistent with the previous reports [8,9,41].
2sinθ , where β* ¼ βcosθ λ , δ ¼ λ , β is the FWHM of the XRD peak (in order to get accurate numbers, the FWHM was obtained by fitting the diffraction peaks with Pseudo-Voigt functions and the widening of the diffraction peaks originated from the XRD equipment was excluded), θ is the Bragg angle, λ ¼ 1.5406 Å is the wavelength of the CuKα radiation, L is the correlation length (corresponding to the crystalline of the film) and ε is the stress parameter. The plots of (β*/δ) [2] as a function of (β*/δ [2]) corresponding to the 4 dominant peaks of the AZO and the B10 are shown in Fig. 1(c) and (d), respectively. The correlation length L was extracted from the slope of the linear fitting of the Halder-Wagner plot and stress ε was obtained from the Y-axis intercept. Generally, when the incident angle of X-Ray is fixed, the detection depth of XRD is fixed and the detection depth of XRD will increase with the increase of the X-Ray incident angle. As shown in Fig. 1(e), in order to study the distribution of the structural parameter L and stress parameter ε along the vertical di rection of the substrate, XRD test was used to characterize the perovskite film with different incident angles of 1� , 2� and 3.5� . (The XRD patterns
Fig. 2. The surface SEM images (left) and histogram of grain size (right) for the CsFA grown on (a), (b) the AZO and (c), (d) the B10, respectively. (e) The decomposition process of the CsFA on the AZO surface with lattice mismatch. 4
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Fig. 3. The UV–Vis absorbance spectra of the perovskite films on (a) FTO, (b) AZO, (c) B2, (d) B10, (e) B15, respectively. (f) FTIR spectra of the CsFA films on FTO, B10 and AZO after annealing, respectively.
For the B2 sample, as shown in Fig. 3(c), although the appearance of the film had no obvious change during the annealing process, the charac teristic absorption peak of perovskite material at around 750 nm degraded at the end of the annealing process. The evolution of the characteristic absorption peak indicated that the decomposition process in the B2 sample was inhibited. However, due to the low concentration of the precursor solution, the Ba(OH)2 film cannot fully cover the AZO surface, and the CsFA film still cannot withstand an hour of heat treat ment. Thus, with the increase of the concentration of the Ba(OH)2 pre cursor solution, as shown in Fig. 3(d) and (e), absorption spectra of the B10 and the B15 remained nearly unchanged and the optical band gap kept stable for the whole annealing process. These results demonstrated that the introduction of Ba(OH)2 interlayer with suitable thickness could effectively improve the stability of CsFA films on AZO surfaces. FTIR was used to further confirm this result. Fig. 3(f) shows the FTIR spectra of the FTO, B10 and AZO samples after the annealing process. As can be seen from the spectrum of the FTO sample, a strong stretching – N from FAþ at 1710 cm 1 can be observed. Meanwhile, vibration of C– two peaks at 3267 cm 1 and 3396 cm 1 corresponding to the stretching vibrations of N–H for FAþ can be identified. This FTIR spectrum revealed the existence of CsFA [42]. For the B10 sample, the spectrum was almost identical with that of the FTO sample, suggesting the B10 sample remained stable after annealing. However, the peak corresponding to hydroxyl group around 3500 cm 1 was observed in the spectrum of the AZO sample. The existence of hydroxyl group was an important indi cator of the decomposition of the CsFA film [15,43]. The FTIR experi ment further confirmed the Ba(OH)2 modification can efficiently suppress the thermal decomposition of the CsFA. To further test the result, the AZO sample and B10 sample were annealed at 130 � C for 30 min. The optical images in Fig. S5 visually exhibit the change of thermal stability after the Ba(OH)2 modification. Although both films began to decompose under annealing, the color evolution was much
more rapid in the CsFA film on the pristine AZO substrate than that in the B10 sample, further confirming the enhanced thermal stability of the CsFA/AZO interface after the Ba(OH)2 modification. The improved thermal stability after modification mainly originates from two aspects. On one hand, the Ba(OH)2 can suppress the reaction between the AZO and the CsFA films. Considering the annealing process in this study was carried out in a N2 filled glovebox, the effect of the air on the thermal stability of perovskites can be excluded. Therefore, the interaction between the ZnO and the organic part of perovskite materials became the main reason for the thermal decomposition of the CsFA [44]. After the modification, Ba2þ reacted with the AZO and a thin layer of O–Ba replaced the absorbed hydroxyl group at the surface of the AZO by the reaction: OH þ Baþ→-O-Ba þ Hþ [45]. The reduction of the surface hydroxyl group concentration can effectively improve the thermal sta bility of perovskite films on the ZnO surfaces [15]. Meanwhile, the re action will form a interface dipoles with negative charge toward AZO surface and positive charge toward the CsFA film, which will enhance the carrier transport efficiency. On the other hand, the release of lattice mismatch also improves the thermal stability of CsFA films. Without the modification, lattice mismatch leads to the accumulation of mismatch stress at the CsFA/AZO interface during the fabrication process of the CsFA film [46]. Defects and dislocations are formed in the CsFA film to release the stress. In addition, when the lattice mismatch is large enough and no modification is carried out, the threading dislocations are formed at the substrate surface and extend into the film or evolve to misfit dislocations at the interface. Fortunately, due to the soft nature of the hybrid perovskites, the dislocations and defects do not penetrate through the CsFA film and most of the stress and defects are concentrated near the substrate [47, 48]. This is in well consistent with the result of the Halder-Wagner plot and grain size distribution. A large amount of grain boundaries in perovskite films make thermal decomposition product diffuse more 5
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easily and more rapidly. The quick loss of decomposition product pro motes degradation of perovskites [49]. In addition, because of the weak chemical bonds at the grain boundaries and low ion migration activation energy, the halogen ions near the gain boundaries are moved apart by annealing and can migrate along the grain boundaries easily [50–52]. At the same time, according to the previous report, large stress can reduce the activation energy of ion migration, further promoting the ion migration in CsFA films [38]. Furthermore, MA cations in CsFA films have weaker interaction with the surrounding PbI6 octahedra than that of FA cation due to their small cation size. Owing to the poor thermal stability, MA cations are gradually released from the lattice structures through grain boundaries under high temperature annealing, leading to the collapse of perovskite lattice structures and the further degradation of CsFA films [7,9,53]. The schematic of this process is shown in Fig. 2 (e). As displayed in Fig. S6(a), showing the cross section of the CsFA film before annealing, the monolithic grains were observed and trended to extend from the bottom to the top. Comparatively, after annealing, (see Fig. S6(b), the uniform and monolithic morphology was changed to the amorphous one and the film thickness decreased at the same time. On the basis of previous description, combined with Fig. 2(e), the morphology evolution of CsFA film during annealing process could be attributed to the degradation of the CsFA film caused by the AZO sub strate and lattice mismatch. During the annealing process, the reaction between CsFA and AZO substrate will lead to the decomposition of CsFA material. At the same time, the lattice mismatch between CsFA film and AZO substrate will introduce more defects into the CsFA film and enhance the ion migration in perovskite film, both of which will pro mote the decomposition of CsFA materials and lead to the collapse of the CsFA lattice. At the same time, due to the main solid decomposition product was PbI2 and the smaller lattice size of PbI2, the morphology of CsFA film will change from monolithic to the amorphous after annealing. To further testify our assumption, EDS element mapping was carried out to determine the distribution of halogen at the surfaces of CsFA films. Fig. 4(a) shows the halogen distribution of the AZO sample, and
the inset shows where the measurement was taken. As can be seen, the Br and I accumulated at the grain boundaries after annealing, which was caused by the diffusion characteristics of the decomposition product and the ion migration paths. The high magnification SEM photo of the AZO sample after annealing is shown in Fig. 4(b). Compared with the SEM picture of the as prepared CsFA shown in Fig. 2(a), the grain size of the perovskite film increased obviously and the grain boundaries became larger, indicating the increased amount of PbI2 [7,54]. As previously stated, the mechanism of this phenomenon can be mainly ascribed to the diffusion of the decomposition product. In comparison, the Br and I distribution in the B10 sample are shown in Fig. 4(c). The Br and I were evenly distributed across the surface of the CsFA film and the concen tration of halogen in the detected area was much lower than that of the AZO sample, suggesting the improved stability of the CsFA film. The further zoom-in SEM photo of the B10 sample is shown in Fig. 4(d), in which, the morphology of the CsFA film remained roughly unchanged compared with Fig. 2(c). After understanding how the lattice mismatch affects the perovskite films, ETL free PSCs based on CsFA perovskite materials and AZO TCO were fabricated. The configuration and energy level diagrams of the ETL free PSC are shown in Fig. 5(a) and (b), respectively. The energy level of CsFA was determined by the UPS experiment, and the raw data is shown in Fig. S7. The J-V curves of the best performing PSCs based on the AZO, B2, B10 and B15 are exhibited in Fig. 5(c). Without modification, the best PCE based on the AZO sample was 9.64% with VOC ¼ 1.006 V, JSC ¼ 14.81 mA cm 2 and FF ¼ 64.68%. As the control, the B2 sample obtained the best PCE of 14.08%, with VOC ¼ 1.046 V, JSC ¼ 19.78 mA cm 2 and FF ¼ 67.58%. With the increase of the Ba (OH)2 film thickness, the best PCE of the B10 dramatically increased to 18.18%, with the VOC ¼ 1.082 V, JSC ¼ 22.88 mA cm 2 and FF ¼ 73.46%. However, with the film thickness continuously increasing, the best PCE of the B15 dropped to 15.78%, with VOC ¼ 1.043 V, JSC ¼ 21.60 mA cm 2 and FF ¼ 69.84%. The EQE spectra of PSCs based on the B2, B10 and B15 are shown in Fig. 5(d). All the PSCs show strong photo-response in the visible light range. The PSCs based on the AZO
Fig. 4. (a) Low magnification SEM pictures of the CsFA film on AZO after annealing and the distribution of Br and I at the surface of the CsFA film on AZO after annealing, respectively. (b) High magnification SEM picture of the CsFA film on AZO after annealing. (c) Low magnification SEM pictures of the CsFA film on B10 after annealing and the distribution of Br and I at the surface of the CsFA film on B10 after annealing, respectively. (d) High magnification SEM picture of the CsFA film on B10 after annealing. 6
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Fig. 5. (a) The structure of the ETL free PSCs. (b) The energy level diagram of the ETL free PSCs. (c) The J-V curves of the best performing PSCs. (d) The EQE spectra and the integrated JSC of the best performing PSCs. (e) J-V curves of the B10 recorded in reverse (from VOC to JSC) and forward (form JSC to VOC) scanning directions. (f) The power output under maximum power point tracking for 25 s. The O1S spectra of pristine (g) AZO and (h) AZO with modification. (i) PCE of the FTO, B10 and AZO sample heated at 100 � C in N2 atmosphere.
show the weakest spectrum intensity in the whole visible light range compared with the other sets of the PSCs. The JSC values extracted from the EQE data were 14.68 mA cm 2 (AZO), 19.64 mA cm 2 (B2), 21.59 mA cm 2 (B10) and 20.61 mA cm 2 (B15), respectively, which are in good agreement with the results of the J-V tests. The J-V curves of the B10 sample obtained from different scan directions were shown in Fig. 5 (e). Due to the improved interface quality and crystal property of the CsFA film, there was only a small difference in the PCE obtained from the forward scan (17.62%) and the revised scan (18.18%). Fig. 5(f) depicted the stabilized power output under the maximum power point tracking reaching 18.1% (at 0.857 V), indicating a good agreement with J-V tests. As previously stated, the poor performance of the AZO and the B2 samples can be mainly attributed to the thermal decomposition of the CsFA. When the Ba(OH)2 layer was thick enough to cover the surface of the AZO, the thermal decomposition of the CsFA was suppressed. As a result, the performance of B10 and B15 increased. On the other hand, Ba (OH)2 can efficiently passivate the surface traps and oxygen vacancies of AZO surfaces at the same time. The XPS spectra of the AZO film and the
Ba(OH)2 modified AZO film are shown in Fig. 5(g) and (h). In Fig. 5(g), the O1S peak of the pristine AZO sample was deconvoluted into two peaks located at 531.50 eV and 530.28 eV, respectively. The low binding energy (530.28 eV) corresponded to the oxygen in ZnO lattice and the higher binding energy (531.50 eV) was originated from the oxygen va cancies and chemisorbed oxygen. The area ratio of these two peaks (S530.28/S531.50) is 0.8. On the other hand, as shown in Fig. 5(h), the O1S peak of the modified AZO film can also be deconvoluted into two peaks located at 530.92 eV to and 531.33 eV, corresponding to the oxygen in ZnO lattice and the oxygen vacancy/chemisorbed oxygen, respectively. The area ratio of these two peaks (S530.92/S531.33) is 0.89. The larger area ratio in the AZO with modification indicated that the concentration of oxygen vacancies and chemisorbed oxygen at the AZO surface was efficiently suppressed [55–57]. Notably, the reaction between Ba(OH)2 and AZO leads to the for mation of interface dipoles, lowering the vacuum level of AZO surfaces and improving the electron extraction efficiency [45,58]. The PL spectra shown in Fig. S8(a) further confirm this feature. As the figure exhibited, the PL peaks located at 772 nm corresponding to CsFA were observed. 7
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Due to the different timescale of electron extraction and radiative recombination, the PL intensity was largely determined by the carrier extraction [15]. Therefore, the weakest PL intensity of the B10 sample indicated that the electron extraction efficiency was higher than the other two samples [59,60]. To demonstrate the improved crystal properties of the CsFA film is another reason for the improvement of PCE, the I–V curves of the uni polar devices (AZO/CsFA/PCBM/Ag) only with the electron charge carriers were recorded [61]. As can be seen from Fig. S8(b), the current increased linearly at the low bias voltage region. While at the high bias voltage region, the current increased with a larger slope. The transition point is the ohmic to trap filled limit transition point (VTFL), which has a linear relationship with trap density (nt). The trap-state is given by formula (3): 2
VTFL ¼
ent L 2εε0
Table 1 The performance of ETL free PSCs from other literatures. Perovskite materials
Structure
PCE
Reference
CH3NH3PbI3 CH3NH3PbI3-xClx CH3NH3PbI3 CH3NH3PbI3 CH3NH3PbI3 CH3NH3PbI3 CH3NH3PbI3
ITO/Perovskite/HTL/Ag FTO/Perovskite/HTL/Au FTO/Perovskite:C70/HTL/Au FTO/Perovskite:C60/HTL/Au ITO/Perovskite/HTL/MoO3/Ag FTO/Perovskite/HTL/Au FTO/Perovskite/HTL/Au
13.5% 14.14% 13.60% 14.30% 10.80% 18.20% 19.22%
[64] 27 [65] [66] [67] 25 26
4. Conclusion The large lattice mismatch induced stress led to the accumulation of defects at the CsFA/AZO interfaces and deterioration of crystal quality of perovskite films, further accelerating the thermal decomposition process caused by the reactive AZO surfaces. With the Ba(OH)2 modi fication, the lattice mismatch was partially relaxed and the thermal instability originated from the AZO was suppressed as well, which improved the crystal quality and thermal stability of CsFA films. On this basis, the thermal stability of ETL-free PSCs based on the AZO was enhanced as well. Additionally, due to its efficient hole blocking prop erty and surface reaction with the AZO, the Ba(OH)2 films efficiently passivated the surface trap states and reduced the vacuum level of the AZO surfaces. Finally, the best PCE of 18.18% was achieved by opti mizing the thickness of the Ba(OH)2 film.
(3)
, where L is the thickness of the CsFA, ε and ε0 are the relative dielectric constant of the CsFA and the vacuum permittivity, respectively. Since the thickness of all the CsFA films is almost the same, the trap density nt is proportional to the VTFL. As calculated from the figure, the VTFL value of the B2, B10 and B15 is 0.55 V, 0.51 V and 0.35 V, respectively, revealing the lower trap densities in the B10 and B15 samples. Mean while, it was noticed that the B15 had the thickest Ba(OH)2 film and lowest trap density but not the best PCE, which can be attributed to the larger series resistance introduced by the thicker Ba(OH)2 interlayer. Shunt Resistance (Rsh) and Series Resistance (Rs) of B10 were 1982869 Ω and 131.9 Ω, respectively. As control, the Rsh and Rs of B15 were 18629 Ω and 136.8 Ω, respectively, which is consistent with our inference. As a result, the Ba(OH)2 modification can improve the ther mal stability of the CsFA on the AZO together with enhancing the per formance of the ETL-free PSCs. Furthermore, in Fig. 5(i), we tested the thermal stability of ETL-free PSCs (the AZO, the B10 and the FTO used as the control) in a N2 at mosphere held at 100 � C. J-V, which was carried out periodically to extract the photovoltaic parameters of the PSCs. As the figure exhibited, the PCE of all the samples began to decrease under thermal stress within a few minutes. Notably, as described above, the CsFA film can remain stable on the FTO surface under annealing. As a result, the performance degradation in the FTO sample can be mainly ascribed to the deterio ration of spiro-OMeTAD and the diffusion of metal electrodes under high temperature [51,62]. As to the AZO sample, the thermal instability of CsFA/AZO interface further exacerbated the thermal stability of PSC and after about 50 min the device was no longer functioning. Clearly, with the help of the Ba(OH)2 modification, the degradation induced by AZO was suppressed and the thermal stability of PSC was also improved. Meanwhile, considering the research on the thermal stability of PSCs based on ZnO or AZO ETL is not a new topic, there have been many reports on this area. Zhao et al. fabricated ETL-free PSCs based AZO and MAPbI3. According their research, the use of AZO efficiently improved the thermal stability of PSCs and a best PCE of 12.6% was obtained [13]. Bush et al. used ITO top electrode to encapsulate the PSCs based on ZnO nanoparticles. The PSCs they obtained could maintain roughly stable after long time annealing under 100 � C and a best PCE of 13.5% [63]. Song et al. researched the thermal stability of perovskite/ZnO film and developed Cs containing triple cation PSCs based on ZnO ETL. In their research, they observed similar CsFAMA/ZnO degradation process. By reducing the annealing temperature time, they obtained a best PCE of 18.9% [9]. Obviously, in comparison with previous researches, the ETL-free PSCs we made obtained enhanced thermal stability with simplified structure and fabrication process. As exhibited in Table 1, the performance of our ETL free PSCs is competitive to the previous reported ones.
Conflicts of interest The authors declare no conflict of interest. Acknowledgements This work was supported by National Natural Science Foundation of China under Grant Nos. 61421002, 61574029, 61471085, and 61874150. This work was also partially supported by University of Kentucky. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.jpowsour.2019.227091. References [1] K. Mahmood, B.S. Swain, A. Amassian, Adv. Energy Mater. 5 (2015), 1500568. [2] J. You, L. Meng, T.B. Song, T.F. Guo, Y.M. Yang, W.H. Chang, Z. Hong, H. Chen, H. Zhou, Q. Chen, Y. Liu, N. De Marco, Y. Yang, Nat. Nanotechnol. 11 (2016) 75–81. [3] D. Liu, T.L. Kelly, Nat. Photonics 8 (2013) 133–138. [4] D.-Y. Son, J.-H. Im, H.-S. Kim, N.-G. Park, J. Phys. Chem. C 118 (2014) 16567–16573. [5] S. Li, P. Zhang, Y. Wang, H. Sarvari, D. Liu, J. Wu, Y. Yang, Z. Wang, Z.D. Chen, Nano Res. 10 (2016) 1092–1103. [6] S. Li, P. Zhang, H. Chen, Y. Wang, D. Liu, J. Wu, H. Sarvari, Z.D. Chen, J. Power Sources 342 (2017) 990–997. [7] B. Conings, J. Drijkoningen, N. Gauquelin, A. Babayigit, J. D’Haen, L. D’Olieslaeger, A. Ethirajan, J. Verbeeck, J. Manca, E. Mosconi, F.D. Angelis, H.G. Boyen, Adv. Energy Mater. 5 (2015), 1500477. [8] M. Saliba, T. Matsui, J.Y. Seo, K. Domanski, J.P. Correa-Baena, M.K. Nazeeruddin, S.M. Zakeeruddin, W. Tress, A. Abate, A. Hagfeldt, M. Gratzel, Energy Environ. Sci. 9 (2016) 1989–1997. [9] J. Song, L. Liu, X.-F. Wang, G. Chen, W. Tian, T. Miyasaka, J. Mater. Chem. 5 (2017) 13439–13447. [10] N.J. Jeon, J.H. Noh, W.S. Yang, Y.C. Kim, S. Ryu, J. Seo, S.I. Seok, Nature 517 (2015) 476–480. [11] W.S. Yang, J.H. Noh, N.J. Jeon, Y.C. Kim, S. Ryu, J. Seo, S.I. Seok, Science 348 (2015) 1234–1237. [12] Z.-L. Tseng, C.-H. Chiang, S.-H. Chang, C.-G. Wu, Nano Energy 28 (2016) 311–318. [13] X. Zhao, H. Shen, Y. Zhang, X. Li, X. Zhao, M. Tai, J. Li, J. Li, X. Li, H. Lin, ACS Appl. Mater. Interfaces 8 (2016) 7826–7833.
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