Tunable Crystallization and Nucleation of Planar CH3NH3PbI3

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Tunable Crystallization and Nucleation of Planar CH3NH3PbI3 through Solvent-Modified Interdiffusion Zhibo Yao, Timothy William Jones, Mihaela Grigore, Noel W. Duffy, Kenrick F Anderson, Ricky B. Dunbar, Krishna Feron, Feng Hao, Hong Lin, and Gregory J Wilson ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b00887 • Publication Date (Web): 10 Apr 2018 Downloaded from http://pubs.acs.org on April 10, 2018

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ACS Applied Materials & Interfaces

Tunable Crystallization and Nucleation of Planar CH3NH3PbI3 through Solvent-Modified Interdiffusion

Zhibo Yao1,2, Timothy W. Jones2, Mihaela Grigore3, Noel W. Duffy4, Kenrick F. Anderson2, Ricky B. Dunbar2, Krishna Feron2, Feng Hao5, Hong Lin1* and Gregory J. Wilson2* 1

State Key Laboratory of New Ceramics & Fine Processing, School of Material Science and Engineering, Tsinghua University, Beijing 100084, PR China 2

3

4

5

CSIRO Energy Centre, Mayfield West, NSW 2304, Australia

CSIRO Energy, 11 Julius Ave, North Ryde NSW 2113, Australia

CSIRO Energy, Clayton Laboratories, Clayton South VIC 3168, Australia

State Key Laboratory of Electronic Thin Films and Integrated Devices, University of Electronic Science and Technology of China, Chengdu 610054, PR China

Corresponding Author

*E-mail:

[email protected]

[email protected] 1 ACS Paragon Plus Environment

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KEYWORDS: Perovskite solar cell, solvent additives, two-step method, MAI diffusion, CH3NH3PbI3 crystallization, XRD

Abstract

A smooth and compact light absorption perovskite layer is a highly-desirable prerequisite for efficient planar perovskite solar cells. However, the rapid reaction between CH3NH3I (MAI) and PbI2 often leads to an inconsistent CH3NH3PbI3 crystal nucleation and growth rate along the film depth during the two-step sequential deposition process. Herein, a facile solvent additive strategy is reported to retard the crystallization kinetics of perovskite formation and accelerate the MAI diffusion across the PbI2 layer. It was found that the ultra-smooth perovskite thin film with narrow crystallite size variation can be achieved by introducing favourable solvent additives into the MAI solution. The effects of DMF, DMSO, γ-Butyrolactone, chlorobenzene and diethyl ether additives on the morphological properties and cross-sectional crystallite size distribution were investigated using Atomic Force Microscopy, X-Ray Diffraction, and Scanning Electron Microscopy. Furthermore, the light absorption and band structure of the as-prepared CH3NH3PbI3 films were investigated and correlated with the photovoltaic performance of the equivalent solar cell devices. Details of perovskite nucleation and crystal growth processes are presented, which opens new avenues for the fabrication of more efficient planar solar cell devices with these ultra-smooth perovskite layers. Introduction 2 ACS Paragon Plus Environment

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Organometal lead trihalide perovskites APbX3 (A= formamidinium (FA) or methylammonium (MA), X = Cl, Br, or I) are emerging as a new generation of solution processable, low cost photovoltaic materials.1-4 Significant progress on perovskite solar cells (PSCs) has been achieved in the last five years, and a certified power conversion efficiency (PCE) exceeding 22% has been recently reported (adopting FAPbI3+MAPbBr3 as light harvester).5-6 The application of perovskite semiconductors in photovoltaics would not have been realized without the development of deposition strategies that gives rise to materials with superior opto-electronic properties. However, our understanding on the organometal trihalide perovskites (in our work we mainly focus on the widely used CH3NH3PbI3) nucleation and crystallization mechanism is limited. This knowledge is crucial for rationally-designing and optimizing high-quality perovskite layers. The perovskite layers with full surface coverage, well-defined grain morphology, and small surface roughness are crucial in achieving efficient solar cells.7-8 Versatile film formation technologies such as solution-based techniques,9, 10 vacuum evaporation,11 and vapour-assisted deposition12, 13 have been employed to prepare the perovskite thin film. In the solution-based deposition, CH3NH3PbI3 film could be deposited by one-step method using the solution containing PbI2 and CH3NH3I.14-15 However, the unbalance between nucleation and growth rate in one-step method would result in a wide spread of filming morphology, which hampers the prospects for practical applications. It’s worth noting that the one-step deposition method could be optimized and widely used to fabricate high-quality 3 ACS Paragon Plus Environment


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perovskite film, together with some advanced technologies for example, anti-solvents method16, gas assisted method17 and solvents additives like HI18, NH4Cl19 et al.20 In the solution deposition method, a significant development in perovskite fabrication process is the two-step sequential deposition method, developed by Mitzi and Huang.21-22 This technique is demonstrated to be particularly useful for preparing dense films of organic−inorganic systems.23 However, the two-step solution process often results in films with significantly enhanced surface roughness due to the quick reaction of lead iodide (PbI2) and methylamonium iodide (MAI) and the rapid crystallization of these perovskite materials.10, 13 This may lead to increased deleterious perovskite|electrode contact, and decreased performance. To solve the issue of rapid crystallization of perovskite, a solvent vapour annealing process was employed into the fabrication where DMF vapour was introduced during the crystallization of the thin perovskite film.24 Since both PbI2 and MAI have high solubility in DMF, the DMF vapour provided a favourable environment and interface so that the CH3NH3PbI3 crystals could be re-dissolved and precursor ions could diffuse further under more ambient conditions than in conventional all solid-state thermal annealing. In other words, CH3NH3PbI3 crystals went through a crystallization – dissolution – reconstruction process during the solvent annealing process. Herein, we present a solvent additive strategy to retard the first crystallization kinetics, so that CH3NH3PbI3 crystals would be directly formed from the first stage of controlled uniform nucleation along with the MAI fully diffused throughout the PbI2

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layer to the last stage of simultaneous grain growth and yielding a small variation of crystallite sizes presenting the low surface roughness.

In the present study, we adopted DMF, DMSO, γ-Butyrolactone (GBL) (three representative solvents which can dissolve both PbI2 and CH3NH3PbI3) and chlorobenzene, diethyl ether (abbreviated as CB and DE, two representative anti-solvents which are commonly used to promote fast nucleation in one-step deposition method 16, 25) as additives in the MAI solution. The solvent/anti-solvent modifiers had substantial effects on the CH3NH3PbI3 film nucleation process, crystalline quality, morphology and photovoltaic performance.

Experimental Section

An optimized two-step deposition method (as illustrated by Fig. 1a) was employed to prepare the perovskite thin film on FTO/bl-TiO2 substrate. First of all, a mixture solvent of DMF/DMSO was adopted to produce PbI2 solution, since the PbI2(DMSO)x complex was proved to be a beneficial intermediate in the formation of ultraflat and superdense perovskite film.23 In the second step, we added various solvent additives into MAI solution (1% volume ratio). Once the MAI/IPA/solvent additive contacting with PbI2, the MAI would begin to diffuse from PbI2|N2 interface throughout the PbI2 layer, along with the chemical reaction between PbI2 and MAI. At the annealing stage, the activation barrier for MAI diffusion was overcome and CH3NH3PbI3 film formation was achieved via grain growth. Device preparation 5 ACS Paragon Plus Environment


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Laser-patterned conductive fluorine doped tin oxide glasses (FTO, sheet resistance: 15 Ω/ □, Dyesol) were cleaned sequentially by ultrasonication in 1% Hellmanex III detergent, deionized water and isopropanol (99.7%, Aldrich), dried under nitrogen flow, and further activated by an air plasma cleaning treatment (Harrick, USA) before use. A compact TiO2 blocking layer (bl- TiO2) was deposited onto the surface of FTO substrate by spin-coating a mildly acidic solution of titanium isopropoxide in ethanol at 2000 rpm for 60 s. The film was then sintered at 500 °C for 30 min. 26 The PbI2 (99%, Sigma-Aldrich) solution (1M, dissolved in mixed solvents with dimethylformamide (DMF, Aldrich)/ dimethyl sulfoxide (DMSO, Sigma) is 85/15 v/v) was deposited onto the bl-TiO2 film by spin-coating at 5000 rpm for 30 s in a glovebox. The films were then dried at 100 ºC for 30 min. After cooling to room temperature, 300 µL of a solution of methylammonium iodide (MAI, Dyesol) in isopropanol (50 mg/ml) with or without solvent modifier additive was dispensed directly to the film and spin-coated at 5000 rpm for 30 s. Afterward, the as-prepared films were heated at 100 °C for 2 h. In our experiment, solvent additives such as DMF, DMSO, γ-Butyrolactone (GBL, Sigma), chlorobenzene (CB, Sigma-Aldrich), diethyl ether (DE, Sigma-Aldrich) were added into the MAI solution.

MAI solutions

containing different concentrations of DMF 0.2 and 5%) were also prepared. The hole transport material (HTM) solution was prepared by dissolving 72.3 mg 2,2′,7,7′-tetrakis(N,N-di-p-methoxyphenylamine)-9,9′-spirobifluorene (spiro-OMeTAD, Luminescence Technology Corp.), 17.5 µL of a solution of 520 mg/ml

lithium

bis(trifluoromethylsulfon)imide

(LiTFSI,

Sigma-Aldrich)

in 6

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acetonitrile (Sigma-Aldrich) and 28.8 µL 4-tert-butylpyridine (TBP, 96%, Sigma-Aldrich)

in

1

ml

anhydrous

chlorobenzene.

The

as-prepared

FTO/bl-TiO2/CH3NH3PbI3 films were coated with HTM solution using spin-coating method at 4000 rpm for 30 s. Devices were kept in dry box (air humidity < 5%) for 10 h (to ensure the Spiro to be well oxidized by the oxygen in air as to improve the photovoltaic performance27) prior to thermal evaporation of 80 nm Au electrodes (under vacuum of ~10-7 Torr, at a rate of ~ 0.1 nm/s) to complete the solar cells. Film and Device characterization The morphological information of the samples was investigated with a field emission scanning electron microscope (Zeiss Auriga FIB-FESEM) operated at an accelerating voltage of 5 kV. Phase analysis of the films were performed on a PANalytical Empyrean X-Ray Diffractometer using CuKα radiation at 40 kV and 40 mA. The analyses were carried out using two configurations of the XRD such as Bragg-Brentano geometry (symmetric scan) and Grazing Incidence X-Ray Diffraction (asymmetric scans). To prevent the degradation of the samples during the analysis, an XRK-900 reactor chamber (Anton-Paar) was used to keep the samples under N2 flow. Step scans were undertaken from 7 to 70° 2θ, with a step interval of 0.02° 2θ. The asymmetric scans were run at different incident beam angles ranging from 0.2° to 1.0° 2θ and an increment of 0.1°. The crystallite size was estimated for the (110) planes using Scherrer equation:

B (2θ ) =

Kλ L cos θ

(1)



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where B = B (observed) – B (standard, LaB6) represents the Full Width at Half Maximum (FWHM) of certain plane, L is the crystallite size, K is the Scherrer constant (0.9 is applied assuming the spherical crystallites), λ is the wavelength of radiation (1.5419 Å of Cu Kα) and θ is the diffraction angle. Lanthanum hexaboride was used as standard sample (no strain or size broadening) to correct FWHM for instrumental broadening. UV-vis reflectance and transmittance spectra were obtained by an UV/vis/NIR spectrophotometer (Lambda 950, Perkin-Elmer, USA) together with an integrating sphere module. The valence band maximum was obtained by the Riken Keiki AC-2 photo-electron spectrometer in air (PESA system). The photoelectron emitted from sample surface is detected when the energy of UV photon 28 becomes larger than the work function of sample. The surface roughness of perovskite films was characterized using an Asylum Research Cypher atomic force microscope (AFM). AFM measurements were conducted in tapping mode. Cell

efficiency

and

performance

parameters

were

determined

via

photocurrent−voltage (J−V) curves recorded under a simulated AM1.5 spectrum. A Newport Class A solar simulator (150 W Xe arc lamp) was calibrated to an irradiance of 1.00 ± 0.02 suns using a silicon reference diode fitted with a KG-5 filter (Fraunhoffer ISE, WPVS). Electrical measurements were performed with a Keithley 2400 Sourcemeter, with 4-point terminal connections. Active area was accurately defined with mask apertures (laser cut 0.55 mm stainless steel, 0.25 cm2). The cells were kept under light soaking for 3 min prior to the J-V test. The scanning voltage 8 ACS Paragon Plus Environment

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step during the measurement is 50 mV, and the dwell time at each voltage is 200 ms. Forward and backward scanning directions were taken between 1.2 V and -0.2 V. Steady-state photocurrent density and PCE output at the maximum power point were also measured, where the Vmax was time-dependent varied as the device working, so each J and Pmax data collected at specific time was matched with the time-dependent Vmax, which was acquired from a cyclic scanning around the latest Vmax. Results and Discussion

Top-view SEM images (Fig. 1b) of the CH3NH3PbI3 films show that the size and the size distribution of the CH3NH3PbI3 grains are strongly influenced by the type of additives. To enable comparison between samples, the grain size distributions were determined from the SEM images and summarized by boxplot in Fig. S1. The perovskite morphology is highly sensitive to the presence of the solution additives, where only 1% volume ratio (100-150 mM) is enough to influence grain sizes. Despite the obvious morphological differences imparted by the solvent modifiers, a dense perovskite layer was successfully formed in all circumstances, with no pinholes exposing the underlying bl-TiO2 for a dense perovskite layer enhances blocking of the parasitic contact between the bl-TiO2 and the HTM which is favourable for device performance. The roughness of the CH3NH3PbI3 films measured by Atomic Force Microscopy is shown in the topographical images in Fig. 2. The root mean square (RMS) roughness of samples (Fig. S2) modified with solvating DMF, DMSO and GBL (9.4, 15.2 and 13.3 nm) were much lower than that of the control (32.9 nm). This level of roughness is commensurate with the lowest reported solution-processed 9 ACS Paragon Plus Environment


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(12.3 nm)29 and thermally-evaporated (11.3 nm)30 films in the literature. In contrast, by adding CB and DE as anti-solvent additives the RMS increased slightly to 33.6 and 38.6 nm. Such phenomenon implies that the introduction of solvents favourable for PbI2 play a crucial role in the process of the nucleation and crystallization of CH3NH3PbI3 thin film. Since PbI2 has negligible solubility in isopropanol, it does not influence the PbI2 morphology. Perovskite nucleation occurs prior to the annealing stage, once the MAI solution contacts with the PbI2 film due to the partial diffusion of MAI into the PbI2 layer, confirmed by the change in colour of the bilayer films observed whilst spin coating (Fig. S3a). That is, a heterogeneous nucleation process. During the annealing stage, the interdiffusion of unreacted PbI2 and MAI occurred and PbI2 appeared to be fully converted to CH3NH3PbI3 as the colour of the films changed to dark black (Fig. S3b), characteristic of compact CH3NH3PbI3. However, the perovskite layer formed prior to the annealing stage would in some degree hinder the MAI diffusing throughout the PbI2 layer. Since grain growth favours initial reacted large-size nuclei (the free energy of volume expansion eclipses that of interface formation)31, this inconsistent reaction time of the “top” and “underneath” layer of PbI2 with MAI would result in asymmetric nucleation across the cross-sectional film and further grain size variation. In addition, the introduction of favourable solvent additives into MAI solution, that has high solubility both of PbI2 and CH3NH3PbI3 (for example DMF), would retard the early crystallization kinetics of perovskite heterogeneous nucleation at the MAI solution|PbI2 interface and result in a uniform nucleation rate of CH3NH3PbI3 crystal through a homogeneous 10 ACS Paragon Plus Environment

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solid-state reaction. This is readily observed by the sample with DMF (before annealing, see Fig. S3a) showing more specular reflective surface than the control sample (no solvent additive). Furthermore, the residual solvating effect of DMF (BP: 152.8 °C) would provide an environment whereby the precursor ions could diffuse further into the film through a solvent-assisted process than would otherwise occur in a heterogeneous nucleation/diffusion process. This promotes a structure whereby nucleation and grain growth occur in unison, yielding a low roughness film surface, similar to that observed for the solvent vapour annealing process as demonstrated by Xiao et al.24 As evidence, we find that the average grain size of sample DMF and GBL was increased to 255 nm and 249 nm respectively from 220 nm from the control sample without solvent modification (See Fig. S1).

We characterized the film crystal properties by XRD, to determine the effect of solvent additives on the formation and evolution of the perovskite. The XRD analyses of the films were assessed using both Bragg-Brentano geometry (symmetric θ-2θ scans) and Grazing Incidence X-Ray Diffraction (GIXRD – asymmetric ω-2θ scans). The diffraction patterns acquired using these two instrumental geometries are given by crystals of different orientation to the substrate surface. In the symmetric scans, only the crystals that have lattice planes (hkl) orientated parallel to the sample surface contribute to the peak intensity (Fig. 3a, b). These lattice planes do not contribute to the intensity of the diffraction peaks in the GIXRD patterns. The lattice planes have to be tilted with respect to the substrate surface in order to satisfy Bragg condition and



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diffraction to occur in GIXRD (Fig. 3c, d). Thus, complimentary crystallographic information is obtained by performing scans in both geometries.

The diffractograms of the perovskite films acquired using Bragg-Brentano geometry are shown in Fig. 4a. The diffraction peaks observed in the measured patterns are consistent with those of tetragonal perovskite reported by other studies 32. The highest intensity peaks of perovskite are from the (110), (220) and (310) lattice planes. The perovskite crystals appear to be strongly orientated on the (110) plane. The presence of SnO2 and the broad peak between 25º and 35º 2θ (glass substrate) indicate that the beam penetrated the CH3NH3PbI3 layer, including the underlying FTO substrate. Therefore, the penetration depth of the film by the X-ray beam is generally greater than the film thickness in the Bragg-Brentano geometry. As the incidence angle increases during the scan, the beam penetration depth increases whereas its footprint on the surface of the film decreases (Fig. 3). The relative intensity of the (220) peak in all scans was the second highest, varying between 31 and 34% (Table S1). The diffraction peaks from (110) and (220) planes are given by the same population of crystallites in Bragg-Brentano geometry. Although the relative intensity of the (220) peak in a powder perovskite XRD pattern was reported to be about 60-70 %33, one reason for the reduced relative intensity of the (220) peak in this study is the smaller volume of the perovskite film analysed at 28.5º 2θ than that at 14.5º 2θ (Fig. 3). In addition, preferred orientation of crystallites on (110) planes will contribute to increase intensity of (110) reflection. The influence of these two factors on the intensity of the peaks in the XRD patterns of films acquired using 12 ACS Paragon Plus Environment

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Bragg-Brentano geometry cannot be discriminated. PbI2 was present in all films. All films but the one prepared using DMSO contained very small amounts of PbI2. The film prepared using DMSO as additive contained greater amount of PbI2 and the orientation of the perovskite crystals was more random than in the other films. The presence of PbI2 indicates incomplete phase conversion,10 which could be attributed by the decomposition of residual PbI2(DMSO)x complex during the annealing stage or partial over-annealing through volatilization of MAI.

GIXRD analyses provide additional information regarding the degree of randomness or preferred orientation of the perovskite crystals. In GIXRD, the incident angle can be very low so the beam effectively penetrates only the perovskite film (Fig. 3c and d). The incident angle is kept stationary and only the detector moves over the 2θ range during the scan. A depth profile analysis of the film can be obtained by running consecutive scans of the film at different incident angles (Fig. 4b). The footprint of the beam on the film surface changes very little for consecutive scans, but the volume of material analysed increases as the incident angle increases. The intensity of the peaks in the GIXRD patterns increased with the angle of incidence due to the increased volume of material analysed. The peaks in the GIXRD patterns are broader than those in the patterns acquired in Bragg-Brentano geometry due to greater contribution from instrumental broadening.

The films were phase pure, with perovskite the only phase identified in most films analysed up to 0.7º 2θ incident angle. Small amounts of PbI2 were found in most the 13 ACS Paragon Plus Environment


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GIXRD patterns of the perovskite layer prepared using 1% DMSO. A weak PbI2 peak is observed in the scan acquired at 0.4º 2θ incident angle, and its intensity increased with increasing incident angle (Fig S4). This indicates that residual PbI2 occurred within the perovskite layer prepared using 1% DMSO. Crystals of PbI2 were not observed in the top view SEM images of any films. Based on the XRD (GIXRD and BB geometry) and SEM data, PbI2 is likely to occur at the interface between perovskite layer and the conductive layer (FTO) in all films, excepting the film prepared using DMSO where PbI2 also occurred within the film. These results reveal the advantage of GIXRD in identifying the PbI2 within film by depth profiling. Such PbI2 impurity has been widely observed even in high efficiency devices reported by other groups, which is believed to have a positive effect on passivating the perovskite grain boundaries. 34

To compare preferred orientation in the perovskite films, two criteria need to be met. The bulk of the depth of the film should be analysed, and the volume of the film analyzed is required to be constant across the entire 2θ range. These conditions are met with GIXRD under increased incidence angles. We choose the incidence angle 0.5 ° since the difference in penetration depth (and hence sampling volume) varies minimally across the 10–70° 2θ range (see Table S2), whilst still analysing the bulk of the depth of the films. Under such conditions, the perovskite films measured (Fig. S5) show preferential orientation of the crystals on the (110) planes. The relative intensities of the highest intensity peaks ((110), (220) and (310) planes) in the GIXRD scans from our study are shown in Table S3. The data obtained from GIXRD scans 14 ACS Paragon Plus Environment

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indicate that the perovskite crystals in the films prepared using additives that dissolve/coordinate perovskite (DMF, DMSO and GBL) show less preferred orientation than in those prepared with additives that are anti-solvent for perovskite. This may be attributed to the partial dissolution of perovskite formed before annealing stage by the additives such as DMF, so that the newly nucleated crystals could orient in more directions during the annealing process as shown by the less preferred orientation on (110) planes.

We estimated the thickness of the perovskite layer using the Beer-Lambert law to be 230-260 nm. The Beer-Lambert law depends on the mass attenuation coefficient, specific mass of the material, wavelength and the incident angle.35 The angles of incidence used in the calculation were those where the SnO2 peaks became noticeable in the GIXRD patterns. The FTO peaks were noticed at angles of incidence starting from 0.70-0.75º in the GIXRD patterns of all samples. This is in good agreement with the values (approximately 240 nm) obtained from stylus profilometry for the film thickness.

Crystallite size and lattice distortion (microstrain) can be determined from peak profile.36 The decrease of crystallite size and microstrain leads to peak broadening. Line Profile Analysis (LPA) can be used to discriminate between size and microstrain effect on the peak profile, which assumes Gaussian and Lorentzian profiles to fit the data. In most symmetric scans, perovskite had only three peaks ((110), (220) and (310)) of relative intensities above 5%. Due to the small number or peaks of high 15 ACS Paragon Plus Environment


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relative intensity the calculated values for crystallite size and strain had high error. Although we could not confirm crystal strain we estimated the crystallite size by applying the Scherrer equation to the main peak ((110) reflection) to enable comparison between samples. The estimated average crystallite size determined using the symmetric scans ranged between 203 and 247 nm (Fig. 5a). The crystallite sizes for the (110) reflection was also measured using the GIXRD patterns collected at different angles of incidence (0.2-0.7º 2θ, Fig. 5b). The increase of crystallite size with increasing incident angle is due to penetration to a greater depth of the crystal and not an actual increase of the crystallite size. However, these values should be used only as a guide since they have larger error than those measured using the XRD patterns acquired in BB geometry, as there is greater instrumental error contribution to the peak broadening. The average crystallite sizes for the (110) reflections at 0.7º 2θ, where the beam started to penetrate the top part of the conductive layer (SnO2), ranged between 195 and 275 nm. The estimated average crystallite sizes for the (110) reflections from both symmetric and asymmetric scans have relative close values to the thickness value of the perovskite layers. This suggests that the perovskite grains that contributed to the XRD patterns are likely to consist of single crystals. Large perovskite crystallite size with minimal crystal boundaries are preferred in fabricating PSCs, due to defects at the interfaces such as grain boundaries that are detrimental to the charge transporting process.37

The optical property and band structure of these CH3NH3PbI3 films were carefully characterized as shown in Fig. 6. The as-prepared CH3NH3PbI3 films exhibited strong 16 ACS Paragon Plus Environment

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light absorption from 400 to 780 nm, indicating adequate light harvesting (Fig. 6a). The absorbance of sample with no doping, CB and DE were slightly higher than the sample of DMF, DMSO and GBL around the bandgap, which might be caused by the light scattering effect of larger surface roughness in the sample of no doping, CB and DE. Optical band gap and valence band (VB) maximum were determined based on reflectance and photo-electron spectrometer in air (PESA). Depicted in Fig. S6a are the transformed Kubelka-Munk spectrum based on the reflectance spectra of these CH3NH3PbI3 films. The bandgap energy Eg is obtained by the Kubelka-Munk 2 equation,14, 38 F ( R) = α = (1 − R) 2 / 2 R , [ F (R)hν ] = A(hν − Eg ) . Shown in Fig. S6b

is the square root of the photoelectron emission intensity (yield1/2) as a function of irradiated photon energy (eV). The photoemission threshold energy represents the ionization potential, which is the energy difference between the vacuum level and VBM in semiconductors.39-40 The obtained Eg and VB of CH3NH3PbI3 films were around 1.5 eV and −5.4 eV, which were in agreement with the reported values14. The close value of Eg and VB of different samples indicates that the superficial morphology and crystallinity differences along the film depth of our CH3NH3PbI3 samples did not necessarily affect the bandgap structure.

Consequently, the correct choice of solvent additives (like DMF, GBL and DMSO) help decouple the nucleation and grain growth process, coordinate the precursor distribution and produce an ultraflat film. Alternatively, the CH3NH3PbI3 film morphology could not be optimized if we chose CB and DE as additives into MAI solution. Among these solvent additives, DMF solvent additive showed the smoothest 17 ACS Paragon Plus Environment


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film surface, the smallest crystallite size distribution and the largest average CH3NH3PbI3 crystallite size during the CH3NH3PbI3 film production.

To further clarify the effect of DMF additive on the CH3NH3PbI3 crystal nucleation, growth and film properties, we adjusted the concentration of DMF in the MAI solution with volume ratio as 0.2, 1 and 5%. Topographical SEM images of the as-prepared CH3NH3PbI3 films are shown in Fig. 7a, and the average CH3NH3PbI3 grain size was calculated to have increased as a function of increased DMF additive concentration, (Fig. S7). In addition, the topographical images of these CH3NH3PbI3 film were measured by AFM (Fig. S8). The RMS values of the samples containing 0, 0.2 and 1% DMF were 32.9, 13.8 and 9.4 nm respectively. This shows that the addition of DMF would significantly decrease the film roughness and help produce a uniform CH3NH3PbI3 film. It could also be observed in Fig. S9 that both CH3NH3PbI3 films with 0.2 and 1% DMF additives have less reflective surface than the films with no solvent additives. This infers the preferred orientation of perovskite crystals in the film and their crystallite sizes are influenced by the concentration of DMF. The crystallite size of perovskite for the (110) reflection measured in Bragg-Brentano geometry increased as the concentration of DMF increased, namely 203 nm (no DMF), 220nm (0.2% DMF) and 247 nm (1% DMF) respectively (see Fig. 7d).

However, increasing the DMF ratio to 5% (beyond the optimum level) resulted in a rough perovskite film with poor coverage of the FTO/bl-TiO2 substrate. By comparing the XRD results of these films (Fig. 7c), we found the films prepared with 18 ACS Paragon Plus Environment

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DMF showed less preferred orientation on the 110 plane than that prepared with no DMF. Also, a new phase was identified in the sample prepared with 5% DMF that was not present in any of the other samples. The peaks in the XRD pattern at 11.41, 16.87, 18.58, and 28.1°, matched well to those of (CH3NH3)4PbI6·2H2O (PDF# 01-076-5812

–

PDF-4+

database).

We

believe

that

the

presence

of

(CH3NH3)4PbI6·2H2O is detrimental to the PSCs performance, since it shows a weak light absorbance (as shown in the UV-Vis absorbance spectra in Fig. S10) with pale yellow colour.41 It has been suggested that the conversion from CH3NH3PbI3 to (CH3NH3)4PbI6·2H2O can occur in the presence of water.

42

We believe the trace

amounts of H2O might be caused by the impurity of DMF reagent (99%, Aldrich), since the moisture level in the glove box used for film preparation was well controlled to be lower than 0.1 ppm. When the MAI solution with 5% DMF gets in contact with the PbI2 film, the increased concentration of DMF would enable significant amounts of the PbI2 film to re-dissolve, effectively restructuring the film. The resultant CH3NH3PbI3 and hence the crystallisation process is likely to be more closely related to the unoptimized one-step method, which does not provide ready access to a condense film with full coverage.22 As a result, the PbI2 film was re-dissolved into the MAI solution where the molar ratio of MAI/ PbI2 would be very high, and the presence of residual water in DMF is likely to react with PbI2 and MAI, producing the observed (CH3NH3)4PbI6·2H2O.

As shown above, applying appropriate solvent modifiers with optimized concentration of additives into MAI solution enables the promotion of MAI diffusion 19 ACS Paragon Plus Environment


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throughout the film, decouples the nucleation and grain growth process, and produces an ultraflat film. Furthermore, the as-prepared CH3NH3PbI3 films with different morphology and crystal properties influence the resultant solar cell performance. Photovoltaic performance for PSCs fabricated by these films is shown in Fig. 8, Fig. S11 and the detailed J−V characteristics are tabulated in Table S4.

The active area of our devices was accurately defined with a mask aperture of 0.25 cm2. By applying this relatively large active area, we could intentionally avoid the random error caused by the film quality variation in the measurements, however, average efficiencies are consequently lower due to the higher charge losses associated with lateral transport. When no solvent additives were applied, the photovoltaic performance was rather poor and the photon to electricity conversion efficiency (PCE) was only 5.0%. By increasing the DMF adding concentration in MAI solution from 0 to 0.2% and to 1%, the short-circuit photocurrent density (Jsc) was increased from 9.4 to 11.4 and to 16.5 mA/cm2, open-circuit voltage (Voc) was increased from 935 to 990 mV, fill factor (FF) was increased from 0.57 to 0.67 and the PCE was increased from 5.0% to 10.9% (acquired by the average of at least 10 samples from the sweeping direction Voc to Jsc). These enhancements were attributed to the DMF additives, promoting MAI rapidly diffusing throughout PbI2 layer resulting in a uniform CH3NH3PbI3 layer formed with a narrow crystallite size distribution along the film depth. The CH3NH3PbI3 layer with low roughness is crucial to achieve high performance PSCs in this planar configuration as a rough CH3NH3PbI3 film will be more difficult for HTM layer to provide complete coverage, which may lead 20 ACS Paragon Plus Environment

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micro-domain shorting of the cell due to direct contact of CH3NH3PbI3 and Au. Further, the increase of CH3NH3PbI3 crystallite size as a result of the DMF additive (as illustrated in Fig. 5a) decreases crystal boundaries, which is beneficial for charge transport in PSCs.37 Further increasing the DMF concentration, the coverage of CH3NH3PbI3 on the bl-TiO2 was largely decreased resulting in a lower Voc of ~700mV and a reduced PCE of 1% of the device made by 5% DMF. The huge hysteresis observed can be attributed to the perovskite crystal properties (defects were hypothesized for the origin of the J-V hysteresis and planar structure with TiO2 as electron transporting layer.26 It was found that J-V hysteresis was significantly dependent on the p- and n-type contact materials and cell structures like mesoporous versus planar structures.43-45 Moreover, the observed increase in the severity of hysteresis with increasing grain size was attributed to an increase in the accumulation of mobile ions at the perovskite interfaces with charge-selective layers, which has been previously studied for this device architecture46, 47.

Depicted in Fig. S11 are the effect of different solvent additives on the J-V performance. The photovoltaic parameters were improved when 1% DMF, DMSO and GBL were applied in the fabrication compared with that of no solvent modifier additives, which benefited from the retarded crystallization kinetics of perovskite by the favourable solvent additives. In addition, we found that the FF of devices made by CB and DE additives were 0.84 and 0.87 and the J-V curves sweeping from Voc to Jsc featured a ‘hump’, as has been observed elsewhere48. The discrepancy between J-V curves collected in forward and reverse directions is a direct indication that the 21 ACS Paragon Plus Environment


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measured J-V curves do not represent the real stabilised photovoltaic performance. To investigate the steady-state behaviour, maximum power-point tracking was used to measure the photocurrent density and PCE for the CB additive PSCs (Fig. 8b). The PSCs made by 1% DMSO additives was also tested as a comparison. The steady-state PCE of sample CB stabilised to 2.9% after light illumination of 200 s, which is inferior to the PCE (7.4% and 10.9% for forward and backward sweeping direction, respectively) acquired from J-V curves and even lower than the PCE of sample no additive.

In contrast, the steady-state PCE of sample DMSO stabilised to 7.1% after light illumination of 200 s, which lies between the PCEs evaluated from J-V measurement in forward and backward sweeping direction (6.1% and 9.0%, respectively). The steady-state current density of 12.8 mA/cm2, was in agreement with the Jsc of J-V curves as shown in Table S4. These results indicate that the crystal properties and the uniformity of perovskite layer were crucial factors affecting J-V hysteresis phenomenon. It is worth noting the lower efficiency of the cells used in this study, which we attributed to the larger cell area chosen to reduce errors associated with under representation of average film morphology in smaller active area cells. More recent studies for p-i-n architectures have demonstrated reduced hysteresis through the use of PCBM or other electron transport materials49-53 however the intention of this study was to decouple this specific interface enhancement from that of the film morphology and grain size of the active perovskite film. Our findings in a method for fabricating highly homogenous perovskite films indicate this can be applied more 22 ACS Paragon Plus Environment

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broadly to use of this material in optoelectronic devices and enhance performance of related solar cell architectures. Conclusion

The fast reaction of MAI with PbI2 when contacted during a two-step process hinders MAI diffusion throughout the PbI2 layer, which can lead to inconsistent nucleation and growth of the CH3NH3PbI3 crystals. We have demonstrated with the introduction of favourable solvents (which dissolve both PbI2 and CH3NH3PbI3) into MAI solution, the crystallization kinetics of perovskite formation was successfully retarded, promoting MAI diffusion homogenously throughout the PbI2 layer. The use of solvent modifiers additives had an observable effect on grain size distribution of the resultant perovskite crystals, film roughness, preferred orientation of the perovskite crystals and the overall crystallite size. From the additives assessed, a 1% DMF addition to the MAI solution produced the best quality CH3NH3PbI3 film, which resulted in enhanced relative photovoltaic performance for the as-prepared planar PSCs. We have demonstrated that rational control of the nucleation and crystal growth is crucial to achieve compact and smooth perovskite layer for photovoltaic devices and the methodologies outlined can be applied to perovskite semiconductor film formation more broadly to applications of these materials to other optoelectronic fields. Supporting Information



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Brief statement in nonsentence format listing the contents of XRD details, surface roughness, UV-vis spectra and J-V parameters, supplied as Supporting Information.

Acknowledgment

The authors would like to express their gratitude to the financial support provided by the National Natural Science Foundation of China (NSFC, 51772166), the projects of International Cooperation and exchanges NSFC (51561145007) and funding support from the CSIRO Research Office. GJW would like to acknowledge the support of the CSIRO Julius Award toward this work and the support of ZY during his placement. This work was performed in part at the Materials node of the Australian National Fabrication Facility, a company established under the National Collaborative Research Infrastructure Strategy to provide nano and micro-fabrication facilities for Australia's researchers.

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Figure 1. (a) Schematic illustration of the optimized two-step method in fabricating CH3NH3PbI3 thin film, (b) FESEM top views of CH3NH3PbI3 films with no solvents doping, 1% DMF, DMSO, GBL, CB and DE as additives respectively.



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Figure 2. The topographical images of CH3NH3PbI3 films with no solvents doping, 1% DMF, DMSO, GBL, CB and DE as additives respectively. The measurements were conducted under ambient air with an atomic probe force microscope.


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Figure 3. Schematics of X-ray diffraction of thin films in Bragg-Brentano geometry at low (a) and high (b) incident angles, and GIXRD at low (c) and high (d) incident angles, respectively. The footprint of the x-ray beam on the film surface and the volume of material analysed by x-rays are inserted as well. The x-rays are diffracted only by the lattice planes parallel to the film surface in BB geometry (a, b). Tilted lattice planes with respect to the film surface contribute to diffraction peaks in GIXRD (c, d). 31 ACS Paragon Plus Environment


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Figure 4. (a) Diffraction patterns of CH3NH3PbI3 films (Bragg-Brentano geometry) with no solvents doping, 1% DMF, 1% DMSO, 1% GBL, 1% CB and 1% DE as additives respectively. (b) XRD patterns of sample with no solvents doping acquired in Bragg-Brentano geometry and Grazing Incidence at different incident angles (omega 0.2, 0.5 and 1.0). The intensity of the peaks in GIXRD patterns increased with increasing angle of incidence.


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Figure 5. (a) Estimated average crystallite sizes of CH3NH3PbI3 in films with no solvents doping, 1% DMF, 1% DMSO, 1% GBL, 1% CB and 1% DE as additives, calculated from the FWHM of (110) peak of symmetric scans. (b) Estimated average crystallite size for the (110) reflection of perovskite measured at different angles of incidence (GIXRD).

Figure 6. (a) UV-Vis absorbance spectra and (b) band structure of CH3NH3PbI3 films with no solvents doping, 1% DMF, DMSO, GBL, CB and DE as additives respectively.



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Figure 7. SEM and XRD analyses of CH3NH3PbI3 films with no solvents doping, 0.2% DMF, 1% DMF and 5% DMF (a) FESEM top views of CH3NH3PbI3 films (b) Crystallite size distributions as a function of angle of incidence (omega) (c) Diffraction patterns acquired in symmetric scans (d) Estimated average crystallite sizes of CH3NH3PbI3.


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Figure 8. (a) J–V curves (under AM 1.5G illumination of 1000 W/m2 intensity, active area of 0.25 cm2) of PSCs prepared by the CH3NH3PbI3 films with 0, 0.2, 1 and 5% DMF as additives respectively. (b) Steady-state photocurrent density and PCE output at the maximum power point of PSCs prepared by the CH3NH3PbI3 films with 1% DMSO and 1% CB as additives. The Vmax was time-dependent varied as the device working, so each J and Pmax data collected at specific time was matched with the Vmax, which was acquired from a cyclic scanning around the latest Vmax.



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Tunable Crystallization and Nucleation of Planar CH3NH3PbI3

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