Morphology control of perovskite in green antisolvent system for MAPbI3-based solar cells with over 20% efficiency

Morphology control of perovskite in green antisolvent system for MAPbI3-based solar cells with over 20% efficiency

Solar Energy Materials & Solar Cells 203 (2019) 110197 Contents lists available at ScienceDirect Solar Energy Materials and Solar Cells journal home...

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Solar Energy Materials & Solar Cells 203 (2019) 110197

Contents lists available at ScienceDirect

Solar Energy Materials and Solar Cells journal homepage: http://www.elsevier.com/locate/solmat

Morphology control of perovskite in green antisolvent system for MAPbI3-based solar cells with over 20% efficiency Soyoung Kim a, Inyoung Jeong b, Cheolwoo Park b, c, Gumin Kang d, Il Ki Han d, Wooyul Kim a, Minwoo Park a, * a

Department of Chemical and Biological Engineering, Sookmyung Women’s University, Seoul, 04310, South Korea Photovoltaics Laboratory, Korea Institute of Energy Research (KIER), Daejeon, 34129, South Korea Department of Energy Science, Sungkyunkwan University, Suwon, 16419, South Korea d Nanophotonics Research Center, Korea Institute of Science and Technology (KIST), Seoul, 02792, South Korea b c

A R T I C L E I N F O

A B S T R A C T

Keywords: Perovskite solar cell Green antisolvent Ethyl acetate Grain size Solvent and compositional engineering

Solvent engineering has been considered a reliable process for the fabrication of pinhole-free and highly crys­ talline perovskite thin films. Recently, green solvents have received immense attention, as the toxic antisolvents used in the conventional fabrication process cause environmental and health hazards. In this regard, ethyl ac­ etate (EA) is a promising environmentally friendly antisolvent. Here, we present the fabrication of perovskites with controlled morphologies by changing the composition of the perovskite and the volume of EA in ambient humidity. The incorporation of [HC(NH2)2]PbIBr2 into a CH3NH3PbI3 matrix results in a grain size up to ~1.5 μm. This induces a considerable reduction of trap density, leading to the suppression of charge recombi­ nation and, consequently, improvement in the photoluminescence characteristics. The resulting power conver­ sion efficiencies (PCEs) of the optimized devices are 20.93% and 19.51% for active areas of 0.12 cm2 and 0.7 cm2, respectively. The reduced diffusion of moisture along grain boundaries improves device stability. Further, by virtue of the excellent humidity resistance of EA, the film morphologies obtained at high relative humidity (50%) are similar to those obtained under dry conditions, exhibiting an impressive PCE of 20.11%. We believe that our optimized fabrication process using EA can be extended to other green antisolvent systems.

1. Introduction Studies on improving a power conversion efficiency (PCE) and environmental stability of organic-inorganic perovskite solar cells (PSCs) have been carried out intensively by a variety of approaches [1–5]. Solvent and compositional engineering have been considered key processes to produce high–quality perovskite thin films since they have contributed considerably towards realization of excellent PCEs of over 24% [6]. For deposition of perovskite thin films via spin coating, a precursor solution of the perovskite is prepared in polar aprotic solvents with high boiling points such as dimethylformamide (DMF), γ-butyr­ olactone (GBL) and dimethylsulfoxide (DMSO). Due to such nonvolatile solvents, the precursor films have a poor and non-uniform morphology after spin coating. For instance, uneven surfaces with large and thick dendrites or islands of the perovskite material are formed during

crystallization, thereby creating a large number of pinholes in the film [7–9]. This leads to the considerable increase in charge recombination and contact resistance at the interface between the perovskite and charge transport layers (CTLs). Further, beyond morphology issues, in­ homogeneity in the composition of the perovskite material also results in significant decrease of device performance and stability [9]. In one-step deposition of perovskites, the critical aspect which governs the quality of films produced is the formation of perovskite precursor-DMSO com­ plexes before crystallization. Addition of an antisolvent to the rotating substrate washes DMF, enabling the formation of pure precursor-DMSO complexes as an intermediate phase, and subsequently results in the growth of uniform pinhole-free perovskite grains during thermal annealing [10–12]. Solvent engineering has three major design parameters: (1) class of antisolvents, (2) volume of antisolvent drips, and (3) dripping time of an

* Corresponding author. E-mail address: [email protected] (M. Park). https://doi.org/10.1016/j.solmat.2019.110197 Received 5 August 2019; Received in revised form 6 September 2019; Accepted 22 September 2019 Available online 27 September 2019 0927-0248/© 2019 Elsevier B.V. All rights reserved.

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antisolvent. The influence of the nature of antisolvents on the grain size, defect density, and surface roughness of perovskite thin films has been evaluated. In particular, physical properties of the antisolvents play a key role in controlling the film morphologies. Several promising can­ didates, including diethyl ether (DE), chlorobenzene (CB), toluene, ethyl acetate (EA), and anisole have been proposed for solvent engineering [10,12,13-23]. Even though they lead to desirable perovskite mor­ phologies and excellent device performance, the use of CB and toluene are limited by their high toxicity [10,17-19]. DE, EA and anisole are relatively mild, and are hence called green solvents [12-16,20-25,26]. DE has been commonly used as an alternative to CB and toluene. DE can be an efficient antisolvent in solvent engineering as the film morphol­ ogies are similar to those obtained by CB and toluene. In addition, it also yields excellent PCEs. However, a high volume (~ 0.5 ml) of DE is required due to its low boiling point (34.6 � C) and nonpolar nature. The highly volatile DE is also sensitive to temperature and humidity, which leads to poor uniformity of the films over large areas, making it difficult to increase the grain size. EA and anisole, on the other hand, are suitable antisolvents due to their high polarity and boiling point (154 � C and 77 � C, respectively), yielding impressive PCEs over 19% for a wide range of perovskites [22-25]. A recent report has suggested that EA/n-hexane mixtures induce the slow growth of perovskite crystals, thereby inhib­ iting the generation of cracks within the films during spin coating [23]. However, the use of a single solvent would prevent effects such as increased sensitivity towards temperature or humidity due to any drastic changes in the evaporation rate of the solvent mixture. In addition to solvent engineering, control of perovskite composi­ tions is also critical to improve the device performance and stability [22, 27-32]. The incorporation of halide anions and metal cations into the main compound in the form of methylammonium (MA) and for­ mamidinium (FA) lead halides provides excellent phase stability and humidity resistance, as well as improved open-circuit voltage (Voc) and short-circuit current density (Jsc) [22,27-32]. The controlled band gaps by compositional engineering facilitate efficient charge injection and transfer from the perovskite to the CTLs. Recently, in the EA system, triple and quadruple cation mixtures have been employed in the pe­ rovskites, which exhibit enlarged grains and reduced charge trap density [22,23]. The doping of alkali metal cations (Cs, Rb, and K) also increase the size of perovskite grains, resulting in improved PCEs of over 20% and excellent device stability at ambient conditions [22,23]. However, light-induced phase segregation in multication perovskites becomes more dominant by migrating metal cations [33,34]. This would accel­ erate an increase in the charge trap density of perovskite thin films. Therefore, highly reproducible, large-grained perovskite thin films free from alkali metal cations would be more proper to solar cell applications. Here, we report the fabrication of highly efficient PSCs based on alkali metal cation-free perovskite thin films containing (MAPbI3)1-x (FAPbIBr2)x, synthesized by employing EA as a green antisolvent. We thoroughly investigate the variation in grain size, surface morphology, and charge trap density with the change in the mole fraction of FAPbIBr2 (x). Increase in the mole fraction of FAPbIBr2 facilitates the growth of large perovskite grains until when FAPbIBr2 is miscible in the MAPbI3 matrix. When the mole fraction of FAPbIBr2 exceeds the critical misci­ bility point (x ¼ 0.048), crystals of FAPbIBr2 start to separate out from the matrix degrading the device performance. Further, the volumetric effect of EA on the formation of precursor-DMSO complexes with the change in relative humidity (RH) is also investigated. The resulting perovskite morphologies determine the device performance. Under the optimum volume of EA (40 μl) at RH of 20% and x ¼ 0.048, the devices exhibit impressive PCEs of 20.93% and 19.51% with suppressed hys­ teresis for active areas of 0.12 cm2 and 0.7 cm2, respectively. The large active area device also exhibits an excellent PCE retention ratio of 0.805 (with respect to the initial value) after 30 d under ambient conditions (RH ¼ 15–20%).

2. Experimental details 2.1. Materials PbI2 (99.9985%), PbBr2 (99.999%), Li-bis(trifluoromethylsulfonyl) imide (Li-TFSI) (>98%), SnO2 colloidal solution (15% in H2O), and anhydrous dimethylformamide (DMF) (99.8%), were purchased from Alfa Aesar. Anhydrous dimethylsulfoxide (DMSO) (99.9%), 4-tert-butyl pyridine (tBP) (96%), anhydrous chlorobenzene (CB) (99.8%), and anhydrous ethylacetate (EA) (99.8%) were purchased from Sigma Aldrich. Methylammonium iodide (MAI) (99.98%) and formamidinium iodide (FAI) (99.98%) were purchased from GreatCell Solar. SpiroOMeTAD was purchased from Lumtec. The MAPbI3 precursor solution was prepared by dissolving 461 mg of PbI2, 159 mg of MAI, and 78 mg of DMSO in 600 mg of DMF. FAPbIBr2 precursor solution was prepared by dissolving 367 mg of PbBr2, 172 mg of FAI, and 78 mg of DMSO in 600 mg of DMF. MAPbI3 and FAPbIBr2 solutions were mixed with a volume ratio of 200 μl:0, 200 μl:5 μl, 200 μl:10 μl, 200 μl:15 μl, and 200 μl:20 μl for x ¼ 0, 0.024, 0.048, 0.072, and 0.096, respectively. The spiro-OMeTAD solution was prepared by dissolving 56 mg of spiroOMeTAD, 5.6 mg of Li-TFSI, and 28 mg of tBP in 1 ml of CB. 2.2. Device fabrication Etched In-doped SnO2 (ITO)/glass (2.5 cm � 2.5 cm) substrates were rinsed for 10 min with acetone and 2-propanol (1:1 vol ratio) using ultrasonication. They were then rinsed with deionized (DI) water and dried in a N2 atmosphere. Subsequently, they were treated with UVozone for 20 min. The dilute SnO2 solution with a composition of 6.5:1 (w/w) in DI water was deposited onto the ITO/glass substrates by spin coating at 3000 rpm for 30 s [35]. After thermal annealing at 150 � C for 30 min, the perovskite solution was spin-coated onto the 20 nm-thick SnO2 electron transport layers (SnO2-ETLs) in two-steps with different rotation speeds: 1000 rpm for 5 s, and 4000 rpm for 12 s. 30–50 μl of EA (best condition ¼ 40 μl) was dripped on the rotating sample using a micropipette during the remaining 5 s in the second step (RH ¼ 20%). At RH ¼ 50%, 50 μl of EA was dripped in the second step to obtain the desirable perovskite morphologies. The samples were immediately heated at 65 � C for 1 min for the slow removal of residual solvents, suppressing the generation of pinholes (which occur in case of fast evaporation of the solvents). Following this, the sample was annealed at 135 � C for 3 min to induce further crystallization. After cooling the samples, spiro-OMeTAD solution was spin coated onto the sample at 2500 rpm for 20 s. Au electrodes were deposited using a thermal evap­ orator with a constant rate of 0.5 Å/s. All samples were fabricated in a fume hood. Before the perovskites were deposited onto the substrate, the RH was found to be 30%. The RH values were thoroughly controlled from 20% to 50% using a humidifier and dehumidifier. 2.3. Characterization SEM images and elemental spectra were obtained using JSM-7600 F. UV–Vis spectra were obtained using a UV–Vis spectrometer (Scinco S–3100). XRD analyses were performed with an X-ray diffractometer (D8 Advance). FT-IR spectra were obtained using a FT-IR spectropho­ tometer (Nicolet IS50). TRPL measurements were performed at the PL maxima of the perovskite (770 nm) using a time-correlated single photon counting module (TCSPC, MPD–PDM Series DET–40 photon counting detector, and a Pendulum CNT–91 frequency counter). This module was combined with a monochromator and the second-harmonic (400 nm) derived from a Ti:Sapphire laser (Mai Tai, Spectra–Physics), and served as a detector and as an excitation source. Steady-state PL measurements were performed using the same equipment used in TRPL analyses. TPV measurements were performed using a nanosecond laser (10 Hz, NT342A, EKSPLA) as a small-perturbation light source and a Xe lamp (300 W, Newport) as a bias light source. The device was directly 2

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connected to a digital oscilloscope (350 MHz, MDO4034C, Tektronix), whose input impedance was set to 1 MΩ for TPV. The bias light intensity was controlled by neutral density filters for various open-circuit voltages (Voc). A strongly attenuated laser pulse (550 nm) was used to generate a transient voltage (ΔV) of <20 mV. J-V and SCLC curves were obtained using Keithley model 2400 source meter. A 180 W xenon arc lamp (Newport) was used as the light source in the solar simulator. The light intensity was adjusted to AM 1.5 G with the use of an NREL-calibrated Si solar cell equipped with a KG-1 filter. The active areas of 0.12 cm2, 0.7 cm2, and 1.4 cm2 were defined using metal masks. EQE spectra were obtained at wavelengths which ranged from 300 nm to 900 nm using IPCE equipment (PV Measurements Inc.).

grained morphology of the perovskite was obtained in the entire area, without any non-uniform crystallite structures such as dendrites. Fig. 1C–G shows the field emission-scanning electron microscope (FE–SEM) images of the films obtained using different mole fractions of FAPbIBr2. When pristine MAPbI3 (x ¼ 0) is deposited onto the substrate, distinct grain boundaries and small pin holes (10–50 nm) are observed (Fig. 1C). Since the growth of grains ceases when two grains come in contact with each other, control of the growth rate is very critical. Some growth inhibitors including guest organic cations and lead halide anions are required to induce the growth of large perovskite grains. Taking into account all these factors, we can estimate the nucleation and growth mechanism of perovskite grains under the variations in the mole fraction of FAPbIBr2. From the Johnson-Mehl-Avrami (JMA) model, the time-dependent nucleation rate (β) of the perovskite from the perovskite-DMSO com­ plex is simply given by the following Arrhenius equation [36]: � � Ea β ¼ k0 exp t​ ​ ​ (1) RT

3. Results and discussion Fig. 1A shows the schematic of solvent engineering process for the synthesis of monolithically grained perovskite thin films. After spreading the perovskite solution on the substrate, 40 μl of EA was dripped onto the rotating substrate to form the precursor-DMSO com­ plexes. The resulting films after thermal annealing are shown in the inset of Fig. 1A. The moderate evaporation rate of EA sustained the precursorDMSO complexes until spin coating process finished, enabling the coating of perovskite thin films to take place over a large-area with a dimension of 7.5 cm � 10 cm, by dripping 240 μl of EA (Fig. 1B). The mirror-like surface of the large substrate indicates that a smooth multi-

where k0 is a constant which is neither time nor temperature dependent, R is the gas constant, T is the temperature, Ea is the activation energy for the nucleation, and t is the time. As the mole fraction of FAPbIBr2 in­ creases, FAþ and Br substituted for MAþ and I significantly reduce the Ea. Since Br is smaller than I , the resulting variation in Ea favors the formation of perovskite unit cells [35]. A large number of nuclei during

Fig. 1. (A) Schematic of green solvent engineering for deposition of the perovskite thin films. The transparent film comprising precursor-DMSO complexes (left) obtained upon dripping EA onto the rotating substrate is immediately crystallized into the perovskite by thermal annealing (right). (B) Photograph of the large area perovskite substrate (7.5 cm � 10 cm). (C–G) SEM images of the perovskite for (C) x ¼ 0, (D) x ¼ 0.024, (E) x ¼ 0.048, (F) x ¼ 0.072, and (G) x ¼ 0.096. (H) Box chart for the grain sizes of each perovskite thin film for x ¼ 0–0.096. The scale bar is 500 nm. 3

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the initial stages of growth promote the growth of pinhole-free thin films. Once the nuclei are formed, the final grain sizes would strongly depend on the activation energy for boundary mobility. The approxi­ mate size of grains is given by the following equation [37]: � � Q d2 d20 ¼ k’ exp (2) rT

The enlarged and oriented perovskite grains decrease the charge trap density [42,43]. However, for x ¼ 0.072 and 0.096, the intensity of all XRD peaks decreases significantly, suggesting that the small grains ob­ tained at these concentrations are associated with the phase separated I-rich nanocrystallites, as shown in Fig. 1F and G. No noticeable shift of XRD peaks are observed when low concentrations of FAþ and Br are added into the MAPbI3 matrix. Fig. 2B shows the Fourier-transform infrared (FT-IR) spectra of the perovskite thin films. The two absorp­ tion peaks observed at the wavenumbers 3271 cm 1 and 3405 cm 1 correspond to the stretching vibration mode of the N–H bond in FAþ [44, 45]. The incorporation of FAþ and Br into the MAPbI3 matrix is also clearly observed by ultraviolet–visible (UV–Vis) spectroscopy (Fig. 2C). As the mole fraction of FAPbIBr2 increases, a blue-shift of the cutoff wavelengths is observed. This is mainly attributed to changes in the lattice spacing due to the incorporation of FAþ and Br into the MAPbI3 unit cells, which affects the bond energies and band gaps (Fig. S2). From (100) peaks, we simply calculated the lattice distances (a) for a-axis by Bragg’s law, yielding a ¼ 6.256, 6.269, 6.275, 6.279, and 6.283 Å for x ¼ 0, 0.024, 0.048, 0.072, and 0.096, respectively. Therefore, the size of unit cells increased by introduction of FAþ and Br . The band gap in­ creases slightly from 1.61 eV to 1.63 eV, accompanied by the blue-shift of photoluminescence (PL) peaks from 763.8 nm to 758.6 nm, as x is increased from 0 to 0.096 (Fig. 2D). The steady-state PL intensities in­ crease with increasing x up to x ¼ 0.048, and then decrease as x is increased further. The large size of grains and absence of I-rich nano­ crystallites would provide less defect sites in the thin films, suppressing charge recombination and resulting in stronger PL emission. To further investigate the charge recombination dynamics within the perovskites in detail, time-resolved PL (TRPL) spectroscopy was employed (Fig. 2E). The perovskites with different mole fractions of FAPbIBr2 were prepared on a bare glass (glass/perovskite) for TRPL measurements. The PL decay curves and the corresponding time constants, τ1 and τ2, were fitted using a bi-exponential function. The fitted parameters and average PL life­ times are summarized in Table S1 τ1 represents the bulk recombination in the perovskite, while τ2 accounts for the radiative recombination by trapped charge carriers [46,47]. In particular, the radiative recombi­ nation process was found to be more significant in the analysis, since a high τ2 to τ1 ratio was observed in all samples (>80%). The average PL lifetimes were found to be 20.67 ns, 31.14 ns, 39.89 ns, 30.22 ns, and 17.21 ns for x ¼ 0, 0.024, 0.048, 0.072, and 0.096, respectively. The delayed recombination would be of advantage to the charge extraction and transfer from the perovskite to the CTLs. Similar to the results of steady-state PL analysis, the lifetimes strongly depend on the charge trap density of the perovskites, which in turn depends on the nature of the grain boundaries and the occurrence of phase separated I-rich nano­ crystallites. Therefore, an optimal concentration of perovskite additives would help improve the device performance. Fig. 3A shows the cross-sectional SEM image of the solar cell. After a 550 nm-thick perovskite layer was deposited onto the SnO2-ETL, a 180 nm-thick 2,20 ,70 70 -tetrakis-(N,N-di-4-methoxyphenylamino)-9,90 spirobifluorene (spiro-OMeTAD) was deposited onto the perovskite as a hole transport layer (HTL). Following this, 100 nm-thick Au contacts were deposited using a thermal evaporator. Fig. 3B–F shows the photocurrent density-voltage (J-V) curves of the devices measured at a scan delay time of 200 ms. The PCEs were found to be 18.35%, 19.76%, 20.93%, 18.89%, and 17.67% for each champion cell, corresponding to x ¼ 0, 0.024, 0.048, 0.072, and 0.096, respectively. These values were obtained using the backward sweep method. For the best performing device (x ¼ 0.048), impressive Voc, Jsc, and fill factor (FF) were found to be 1.132 V, 24.07 mA cm 2, and 0.768, respectively. The steady–state photocurrent densities of the devices were obtained at the maximum power voltage (Fig. S3). The calculated PCEs (17.7%, 19.19%, 20.28%, 18.32%, and 16.92% for x ¼ 0, 0.024, 0.048, 0.072, and 0.096, respec­ tively) showed excellent agreement with those obtained from the J-V curves. Further, the suppressed J-V hysteresis was also observed in the device for x ¼ 0.048, since it strongly depends on the charge

where d and d0 are respectively the final and initial grain sizes, k’ is a constant, r is the radius of the nucleus, and Q is the activation energy for boundary mobility. It has been known that the introduction of FAþ and Br reduces the activation energy for boundary mobility [37]. We thoroughly investigated the activation energy for boundary mobility (Q) to clarify the possibility of fitting the relation between the grain size and mole fraction of FAPbIBr2 [38], Q ¼ ΔHf þ ΔHm

(3)

where ΔHf is the enthalpy of formation and ΔHm is the enthalpy of motion. For MAPbI3, ΔHf is fixed to 371 kJ mol 1 at 298 K [39]. Since the small amounts of FAPbIBr2 rarely contribute to the ΔHf , Q strongly depends on the ΔHm . The boundary mobility (M) is proportional to the ΔHm by following the relationship [40]: � � ΔHm ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ M∝exp (4) kT where k is the Boltzmann constant. The increasing crystallinity of the perovskites as the mole fraction of FAPbIBr2 increases up to x ¼ 0.048 indicates that well-oriented grains would reduce ΔHm [41]. Therefore, the boundary mobility increases by incorporating FAPbIBr2. Combining Eqs. (2)–(4) yields: ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi� ffiffiffi sffiffiffiffiffiffiffi� ΔHm ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ ​ d∝ exp kT (5) For the small variation in ΔHm , the final grain sizes would be fitted to the x values by Eq. (5). This relation is valid before I-rich nano­ crystallites are extracted, since they severely hinder the growth of the perovskite grains. As shown in Fig. 1D and E, pinhole-free large-grained morphologies are obtained after small amounts of FAPbIBr2 (x ¼ 0.024 and 0.048, respectively) were added to the MAPbI3 matrix. However, at x ¼ 0.072 and 0.096, the bright nanocrystallites shown in the SEM im­ ages were supposed to be the compound of demixed FAI and PbBr2 in the MAPbI3 matrix. The composition of the nanocrystallites was investi­ gated by energy dispersive x-ray spectroscopy (EDS) analysis (Fig. S1). For 10 particles, their average atomic ratio of I and Br was found to be 2.3:1, which clearly indicates that the I-rich nanocrystallites were formed by demixing of the precursors at above the critical concentra­ tion. These undesirable crystallites inhibited the growth of perovskite grains, thereby resulting in small-grained morphologies for the films. The white arrows in Fig. 1F indicate the extracted FAPbIBr2 crystallites with high brightness. This result clearly suggests that the complete miscibility of FAPbIBr2 in the MAPbI3 matrix is essential and the mole fraction of FAPbIBr2 should be maintained below x ¼ 0.048. At x ¼ 0.096, the grain sizes become much smaller, and numerous small FAPbIBr2 crystallites are generated (Fig. 1G). Fig. 1H shows the box chart of grain sizes for different mole fractions of FAPbIBr2. The average grain sizes were found to be 0.42 μm, 0.63 μm, 1.04 μm, 0.38 μm, and 0.23 μm for x ¼ 0, 0.024, 0.048, 0.072, and 0.096, respectively. Fig. 2A shows the x-ray diffraction (XRD) spectra of the perovskite thin films for x ¼ 0, 0.024, 0.048, 0.072, and 0.096. These spectra show the variation of crystallinity and grain sizes for various mole fractions of FAPbIBr2. It should be noted that the intensity of (100) and (220) peaks increases at x ¼ 0.024 and 0.048, while that of (312), (224), and (314) peaks decreases. It can be seen that the oriented growth of perovskite grains becomes dominant as the mole fraction of FAPbIBr2 increases. 4

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Fig. 2. Characterization of the perovskite thin films for x ¼ 0–0.096. (A) XRD spectra, (B) FT-IR spectra obtained by attenuated total reflection (ATR) mode, (C) UV–Vis spectra, (D) steady-state PL spectra, and (E) time-resolved PL (TRPL) decay curves, of the perovskite thin films. The film thickness of all samples is fixed to 550 nm, and the films were prepared on a glass substrate (EA ¼ 40 μl).

recombination in the perovskite and the interfacial layers [48-50]. In particular, fast ion migration takes place along the grain boundaries. Therefore, lesser number of grain boundaries would delay the ionic diffusion and charge recombination [49,50]. In addition to the grain boundary effect, the degradation of device performance at x ¼ 0.072 and 0.096 was also attributed to the phase separated I-rich nanocrystallites. Therefore, large grains with a pure mixed halide phase were essential to realize high-performance devices. The EA system also provides the excellent reproducibility of the device performance. By measuring fifty independent devices for x ¼ 0, 0.024, 0.048, 0.072, and 0.096, the average PCEs were found to be 17.73%, 19.15%, 20.06%, 18.22%, and 16.93%, respectively (Fig. 3G). For x ¼ 0.048, the device with the perovskite layer synthesized using DE as an antisolvent exhibited PCEs of 19.03% and 18.17% for the backward and forward sweep, respec­ tively, at a scan delay time of 200 ms (Fig. S4). Since no pinholes or other heterogeneous nanocrystallites were observed in the film, the lower PCEs could be attributed to the smaller grains compared to the sample synthesized by using EA (Fig. S5). Fig. 3H shows the external quantum efficiency (EQE) spectra of the devices. The high EQE values (>90%) were observed at wavelengths from 350 nm to 760 nm for x ¼ 0.048. The suppression of carrier recombination in the perovskite results in efficient charge extraction and transfer, which contributes to the high EQE values. The integrated Jsc values in the EQE spectra were found to be 23.17 mA cm 2, 23.49 mA cm 2, 23.7 mA cm 2, 23.38 mA cm 2, and 22.93 mA cm 2 for x ¼ 0, 0.024, 0.048, 0.072, and 0.096, respectively. The values display excellent agreement with the Jsc obtained from the J-V curves. A slight increase of the maximum wavelength, from 790 nm to 805 nm, was observed with an increase of x. This is associated with the increase in the band gap. Transient photovoltage (TPV) analysis was performed to support the charge recombination behavior in the devices. Fig. 3I shows the photovoltage decay curves of the devices. The average photovoltage decay times were found to be 1.068 μs, 1.54 μs, 1.59 μs, 1.46 μs, and 1.023 μs for x ¼ 0, 0.024, 0.048, 0.072, and 0.096, respectively. For x ¼ 0.048, the recombination rate was reduced

considerably by decreasing recombination centers. This delayed recombination, in turn, contributed to the improvement of Voc [51,52]. The fitted parameters and average photovoltage decay times are sum­ marized in Table S2. For the quantitative analysis of the charge trap density, spacecharge-limited current (SCLC) curves were measured. Fig. 4A shows schematically the device architecture used for the SCLC measurements. In the dark state, a voltage of 0 V–4 V was applied to the devices. Three distinct regions: Ohmic, trap filled limit (TFL), and Child regions were observed in the SCLC curve (Fig. 4B–F). At low voltages, the dark current increases linearly (Ohmic region). As the voltage is increased further, the current increases nonlinearly from the critical point, indicating that all trap sites are filled by the injected charge carriers (TFL region). At higher voltages, the current exhibits quadratic voltage dependence (Child region). From the voltage at the TFL point (VTFL), we can calcu­ late the charge trap density by using the following equation [53,54]: VTFL ¼

entrap L2 2ε0 ε

(6)

where e is the elementary charge, ntrap is the charge trap density, L is the thickness of the perovskite layer, ε0 is the vacuum permittivity, and ε is the relative dielectric constant of the perovskite material. From the SCLC curves, the VTFL values were found to be 0.77 V, 0.67 V, 0.48 V, 0.73 V, and 0.8 V, for x ¼ 0, 0.024, 0.048, 0.072, and 0.096, respectively. They were obtained at the extrapolated contact points of the two lines cor­ responding to the Ohmic and TFL regions. The ntrap values thus calcu­ lated were 9.16 � 1015 cm 3, 7.97 � 1015 cm 3, 5.7 � 1015 cm3, 8.68 � 1015 cm3, and 9.52 � 1015 cm 3, for x ¼ 0, 0.024, 0.048, 0.072, and 0.096, respectively. The defect level of the optimized perovskite was similar to that of the quadruple cation-employed perovskite [22]. Therefore, it can be seen that a high density of large oriented grains reduces the charge trap density significantly. The effect of film morphology on the device performance was also

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Fig. 3. (A) Cross-section SEM image of the device. The perovskites were deposited by dripping 40 μl of EA. (B–F) J-V curves of the devices for (B) x ¼ 0, (C) x ¼ 0.024, (D) x ¼ 0.048, (E) x ¼ 0.072, and (F) x ¼ 0.096. The measurements were performed under 1 sun condition. (G) Box chart for PCEs of fifty devices for each x value (total number of devices ¼ 250). (H) EQE spectra, and (I) transient photovoltage (TPV) decay curves of the devices (x ¼ 0–0.096).

respectively, exhibiting high reproducibility. For A ¼ 1.4 cm2, the PCE of the MAPbI3-device was found to be 14.53% and 12.56% for the back­ ward and forward sweeps, respectively (Fig. S6). In addition to the small size of grains present in the MAPbI3 film, the direct contact between the HTL and ETL occurs through pinholes, leading to significant degradation of device performance and a pronounced J-V hysteresis for large areas. Furthermore, compared to the MAPbI3-device, the mixed halide-device (x ¼ 0.048) exhibits excellent environmental stability for A ¼ 1.4 cm2 (Fig. 5D). The stability test of mixed halide- and MAPbI3-devices was performed for 30 d at ambient conditions (RH ¼ 15–20%). After 30 d, while the mixed halide-device retained the normalized PCE up to 0.805 with respect to its initial value, the corresponding PCE of the MAP­ bI3-device decreased drastically to 0.346. The absorption of moisture occurs preferentially at grain boundaries, which accelerates the decomposition of the perovskite grains [56,57]. Therefore, poor envi­ ronmental stability of the MAPbI3-device can be attributed to the large number of grain boundaries and small pinholes that serve as sites for the absorption of moisture. The photovoltaic parameters of the devices with A ¼ 0.7 cm2 and 1.4 cm2 are summarized in Table 1. For x ¼ 0.048, we also investigate the volumetric effect of EA on the formation of perovskite morphologies during the fabrication process. The optimized EA volume is found to be 40 μl for a substrate of size 2.5 � 2.5 cm at RH ¼ 20%, as shown in Fig. 1A. During spin coating, a smaller volume (30 μl) of EA would not completely remove DMF. Re­ sidual DMF in the precursor-DMSO complex produces coarse perovskite

observed as the area (A) of the devices increased. Grain boundaries in the perovskite act as a major recombination site at the interface of the perovskite and the CTLs (perovskite/CTLs) [49,50]. In this regard, the density of grain boundaries becomes a major factor in controlling the device performance when the active area is large. The interfacial resis­ tance at the perovskite/CTLs would increase in proportion to the active area. FF and J-V hysteresis are particularly associated with the interfa­ cial resistance [48,55]. Therefore, the decrease in the density of grain boundaries would suppress the charge recombination, yielding highly efficient devices with large active areas. Fig. 5A shows the photograph of the device with A ¼ 1.4 cm2 (active area ¼ 0.7 cm2). For A ¼ 0.7 cm2, impressive PCEs were found to be 19.51% and 18.44%, for the backward and forward sweeps, respectively, at a scan delay time of 200 ms. For this sample, a minor PCE drop of ~1% was observed with respect to that obtained at a small area of 0.12 cm2. Furthermore, excellent PCE retention was observed for the larger area sample (1.4 cm2). The PCEs were found to be 17.9% and 16.54% respectively, for the backward and forward sweeps, at a scan delay time of 200 ms. Even though the J-V hysteresis became more pronounced, no severe hysteresis was observed as the active area increased due to the high-quality of the perovskite layer. To the best of our knowledge, in the EA system, our devices exhibited the best performance reported thus far for both small and large active areas (Table S3). Fig. 5C shows the box chart of the twenty in­ dependent devices for A ¼ 0.7 cm2 and 1.4 cm2. The average PCEs were found to be 18.56% and 17.15% for A ¼ 0.7 cm2 and 1.4 cm2, 6

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Solar Energy Materials and Solar Cells 203 (2019) 110197

Fig. 4. (A) Schematic of the device structure describing the SCLC analysis. (B–F) Dark current-voltage curves for different x values: (B) x ¼ 0, (C) x ¼ 0.024, (D) x ¼ 0.048, (E) x ¼ 0.072, and (F) x ¼ 0.096.

Fig. 5. (A) Photograph of the device with an active area of 0.7 cm2 (total area ¼ 1.4 cm2). (B) J-V curves of the device (x ¼ 0.048). The measurement was performed under 1 sun condition. (C) Histogram of twenty independent devices for different areas (0.7 cm2 and 1.4 cm2). (D) Normalized PCE curves of two devices, x ¼ 0 and x ¼ 0.048. The devices were stored at RH ¼ 15–20% for 30 d.

morphologies. After thermal annealing, the resulting perovskite films become opaque and hazy (Fig. 6A). In these films, large and rough perovskite dendrites (thickness ~ 0.5 μm and length ~ 5 μm) create a number of pinholes (Fig. 6B). Meanwhile, using an excess amount of EA (50 μl) would partially remove DMSO, leading to the extraction of pre­ cursors from the intermediate phase. After excess EA is dripped onto the substrate, the perovskite nanocrystals grow immediately, and the color of the substrate changes to light brown. After thermal annealing, the mirror-like dark brown colored film obtained is shown in Fig. 6C.

However, the SEM image in Fig. 6D shows that this film has a more heterogeneous morphology compared to the optimized films. The perovskite nanocrystals having bright contrast in the SEM image are believed to be synthesized during spin coating. The major grains would then be formed during thermal annealing. The poor crystallinity of both films synthesized with non-optimal volumes of EA was observed by XRD analysis (Fig. S7). The random orientation of the perovskite crystals forming a bulky dendrite significantly reduces the peak intensity. Even though excess EA produces well-oriented perovskite grains, they still 7

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Solar Energy Materials and Solar Cells 203 (2019) 110197

devices, respectively. Meanwhile, at higher relative humidity (RH ¼ 35–50%), the optimal volume of EA increases to 50 μl. Due to the slower evaporation of DMF in the more humid atmosphere, the required volume of EA should be increased to successfully synthesize the pure precursor-DMSO complexes. At RH ¼ 50%, the resulting film morphol­ ogies are similar to the best-performing counterparts obtained under dry conditions (RH < 30%), exhibiting an excellent PCE of 20.11% and suppressed J-V hysteresis (Fig. S8). The excellent humidity resistance of EA would also be useful to fabricate high-performance devices under ambient humidity.

Table 1 Photovoltaic parameters of PSCs for an active area of 0.7 cm2 and a total area of 1.4 cm2 measured under standard AM 1.5 G illumination of 100 mW cm 2a. Area [cm2]

Scan direction

Voc [V]

Jsc [mA cm 2]

FF

PCE [%]

0.7

Backward

1.4

Forward Backward

1.11 (1.093) 1.088 1.079 (1.061) 1.052

23.75 (23.68) 23.78 23.53 (23.46) 23.48

0.74 (0.717) 0.713 0.705 (0.689) 0.67

19.51 (18.56) 18.44 17.9 (17.15) 16.54

Forward a

The values in brackets are the average values calculated from 20 devices.

4. Conclusions

exhibit lower peak intensity than the optimized films due to their very small sizes (<200 nm). The undesirable film morphologies produced under both conditions are eventually responsible for the degradation of PL characteristics. Fig. 6E shows the TRPL decay curves of the two films synthesized by dropping 30 μl and 50 μl of EA (perov-30 μl and perov-50 μl, respectively). The average PL lifetimes were found to be 5.86 ns and 12.17 ns for perov-30 μl and perov-50 μl, respectively. The fitted pa­ rameters and average PL lifetimes are summarized in Table S4. The very short PL lifetime in perov-30 μl indicates that the charge carrier recombination becomes dominant in the bulky dendrites. For perov-50 μl, a number of grain boundaries and heterogeneous perovskite nano­ crystals act as recombination sites. Therefore, the degradation of the device performance is strongly influenced by the poor quality of perovskite films synthesized under conditions that are not optimal. Fig. 6F shows the J-V curves of the devices employing perov-30 μl and perov-50 μl as solar absorbing layers. All photovoltaic parameters are extremely inferior to those of the optimized devices. The PCEs were found to be 7.62% and 13.18% for the perov-30 μl and perov-50 μl

In conclusion, we have proposed a technique involving optimization of green solvent and compositional engineering for successful deposition of monolithically grained perovskite thin films for solar cell applica­ tions. The incorporation of FAPbIBr2 into the MAPbI3 matrix contrib­ uted significantly to the formation of large perovskite grains. The increasing grain size resulted in reduced charge trap density, leading to the effective suppression of charge carrier recombination at grain boundaries and the subsequent improvement of PL characteristics. The impressive PCEs of the optimized devices (x ¼ 0.048) were found to be 20.93% and 19.51% for active areas of 0.12 cm2 and 0.7 cm2, respec­ tively. The excellent environmental stability of the devices was also demonstrated in ambient conditions (RH ¼ 15–20%) over a period of 30 d, where the normalized PCE was retained to 0.805 with respect to the initial value. Furthermore, the volume of EA dripped during spincoating deposition was controlled to obtain the desirable film mor­ phologies, even in a humid atmosphere. The excellent humidity resis­ tance of EA in the fabrication process of perovskites provides the high reproducibility of film morphologies and device performance, Fig. 6. The volumetric effect of EA drips on the film morphologies, PL characteristics, and device perfor­ mance. (A) Photographs of the perovskite substrate obtained by dropping 30 μl of EA. (B) SEM image corresponding to (A) (after thermal annealing). (C) Photographs of the perovskite substrate obtained by dropping 50 μl of EA. (D) SEM image corresponding to (C) (after thermal annealing). (E) TRPL decay curves of the perovskite thin films obtained by drop­ ping 30 μl and 50 μl of EA. (F) J-V curves of the two devices: 30 μl-perov and 50 μl-perov.

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S. Kim et al.

irrespective of the relative humidity during the growth process. The optimization engineering of perovskites based on our work would also be highly feasible in other green solvent systems.

[20]

Declaration of competing interest

[21]

The authors declare no conflict of interest.

[22]

Acknowledgement

[23]

This research was supported by the National Research Foundation of Korea grant funded by the Ministry of Science and ICT (NRF2019R1C1C1005129).

[24] [25]

Appendix A. Supplementary data

[26]

Supplementary data related to this article can be found at https ://doi.org/10.1016/j.solmat.2019.110197.

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