Understanding surface chemistry during MAPbI₃ spray deposition and its effect on photovoltaic performance

Abstract

In recent years there has been significant research into organic–inorganic halide perovskites for photovoltaic applications and has resulted in power conversion efficiencies greater than 20%. A number of deposition techniques have been utilized to produce high quality perovskite films including thermal evaporation, single and two-step spin coating and vapor assisted processes. However, these laboratory based fabrication methods are not well matched with low costing, roll to roll processing that is associated with large scale manufacture. In this work we use a low costing and rapid spray coating technique to fabricate solution processed MAPbI₃ films with controlled film thickness achieved through modifying concentration/spray volume. Investigation into the chemical composition and film morphology showed differing surface chemistries for the different film thicknesses, altering the surface morphology from large micro grained structures to small distorted grain arrangements. We have then implemented spray deposited films into solar devices (ITO/TiO₂/MAPbI₃/Au) and have studied the current density–voltage (JV) characteristics including light soaking effects. Through optimizing the absorber thickness the short-circuit current density is increased from 4.9 mA cm⁻² to 22.3 mA cm⁻², increasing overall power conversion efficiency by a factor >5. Investigation into the performance of devices under continued illumination showed an inverse relationship existed between low efficiency (thick) MAPbI₃ films and high efficiency (thin) MAPbI₃ films, where the former displayed large increases in overall performance compared to the rapid degradation in the latter. We attribute device instabilities to the increase in defect density as a consequence of the spray-induced surface chemistry, leading to trap assisted recombination at the back electrode.

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Introduction

Over recent years, organic–inorganic halide perovskites have received considerable attention due to their intense light absorption, greater charge diffusion lengths and low exciton binding energies.¹ Rapid developments in device engineering along with material processing have led to power conversion efficiencies greater than 20%.^2,3 The perovskite structure is typically represented by the general formula ABX₃, where in the case of organic–inorganic halide perovskites A is a small organic cation (CH₃NH₃⁺), B is a metal cation (Pb²⁺) and X is a monovalent halide anion (I⁻). Variations in the metal and halide components have been demonstrated to impact the perovskites electronic structure, enabling bandgap tunability between 1.5 eV and 3.1 eV.⁴ Despite the perovskites many enhanced opto-electronic properties, successful implementation into applications, mainly photovoltaic, has been hindered by a number of key operational issues. These issues arise from the complex nature of the perovskites intrinsic properties which are not yet fully understood. Under normal operating conditions, light soaking effects and hysteresis phenomena have been observed for perovskite solar cells.^5–9 Experimental studies have shown that both light soaking effect and hysteresis phenomena are related to the electrode interfaces in heterojunction perovskite devices. There is a general consensus that a number of mechanisms may be responsible and act simultaneously, such as reorientation of ferroelectric organic cations,¹⁰ trapping/de-trapping of charge carriers,¹¹ lattice distortion-induced polarization and ion migration.^12,13 In order for perovskites enhanced properties to be successfully implemented into a range of applications, high quality thin films must be produced. Several deposition techniques including thermal evaporation,¹⁴ single and two step spin coating^15,16 and vapor assisted processes¹⁷ have been optimized to achieve high quality perovskite films for use in high performing PV devices. However, these laboratory based fabrication methods for perovskite films are not well-matched with a low costing, roll to roll process that is typical for large scale manufacture.¹⁸ A solution would be to develop a one-step spray coating approach, which has already been successfully exploited for a range of electronic devices such as light emitting diodes,¹⁹ photo-detectors,²⁰ field effect transistors²¹ and also photovoltaics.^22,23 However, challenges still exist to deploy spray coating for perovskite photovoltaics as for instance thin layers as required for optimal device performance are still difficult to achieve by spray. Nonetheless, a one-step spray coating method still presents technological relevance and advantages that are highly beneficial to industrial manufacture, such as the ability to simultaneously provide high throughput, efficient use of material and uniform film coverage. Additionally, control over directional deposition means that grains grow in a vertically aligned manner, resulting in low number of grain boundaries that could benefit overall charge transport properties. A main issue concerning the implementation of methylammonium lead iodide (MAPbI₃) as a feasible and efficient absorber for solar devices is its stability. Exposure to even ambient atmospheric conditions causes severe degradation of the MAPbI₃ film and their unique opto-electronic properties diminish as a consequence. Numerous works have established the effects of moisture and oxygen ingress into MAPbI₃ films and it is well understood that as MAI is evaporated solid PbI₂ remains on the film, which turns into a yellowish shade from the original dark colour. MAI can be further broken down to CH₃NH₂ (liquid) and HI (liquid) with the latter easily reacting with O₂ to form I₂ (solid) and H₂O (liquid) or decompose into I₂ (solid) and H₂ (gas).²⁴ In this work we use a low costing and scalable one-step spray technique to fabricate solution processed MAPbI₃ films. By changing the concentration of MAPbI₃ in DMF solution (0.25 M and 0.125 M) along with the spray volume (400 μL, 300 μL and 200 μL) we were able to control the film thickness. The surface chemistry of MAPbI₃ films induced during spraying are studied in detail using X-ray photoelectron spectroscopy (XPS) and indicates increased surface reactivity for reduced film thicknesses. A direct consequence is the accelerated decomposition from MAPbI₃ to MAPbI_2+x/PbI₂ where we highlight in detail the intermediate chemical phases that are formed during this transition together with surface morphology changes. The spray deposited films are used in solar devices with TiO₂ as an electron transfer/hole blocking layer and their current density–voltage (JV) characteristics are studied to include light soaking effects over a period of 0–6 min. This work provides in-depth details of changing surface chemistry as a result of the spray method, which highlights a more rapid degradation process at the surface of thinner films. The subsequent effect on the electronic band structure for spray produced MAPbI₃ films is also analysed in details.

Experimental section

Materials

Precursor materials triethanolamine, titanium(IV) isoperoxide and titanium(IV) oxide anatase nanopowder <25 nm diameter were purchased from Sigma-Aldrich and used as received. Methylammonium lead iodide (CH₃NH₃PbI₃) perovskite powder with 99% purity was purchased from Xi’an p-OLED.

Film deposition and device fabrication

Fabrication of compact TiO₂. Solution of TiO₂ was produced using a mixture of triethanolamine (0.39 g) and Titanium(IV) isoperoxide (1.56 mL) with ethanol (18 mL) before mixing at 40 °C for 2 hours then left overnight. The solution was then spin coated at 5000 rpm for 30 s on top of the indium-tin oxide (ITO) coated glass substrate followed by annealing the film at 300 °C (Fig. 1a.1 and 2).

Fig. 1 (a) Fabrication steps for the synthesis of MAPbI₃ solar devices and (b) non-equilibrated band diagram for our cell architecture and (inset bottom right) unit cell for MAPbI₃ perovskite. (c–h) Cross sectional scanning electron microscopy (SEM) images of deposited MAPbI₃ on a mesoporous TiO₂ layer by a one-step spray method using the following MAPbI₃ in DMF concentration/spray volume: (c) 0.25 M/400 μL (d) 0.25 M/300 μL (e) 0.25 M/200 μL (f) 0.125 M/400 μL (g) 0.125 M/300 μL and (h) 0.125 M/200 μL. Scale bars 2.5 μm. (i) Measured MAPbI₃ film thickness using following deposition implementing set concentrations and spray volumes.

TiO₂ mesoporous and CH₃NH₃PbI₃ thin film fabrication using one-step spray technique. A solution of 40 mg Titanium(IV) oxide anatase nanopowder <25 nm diameter and 20 mL of ethanol/DI-water (1 : 1) is mixed and sonicated for 30 minutes followed by an additional settling period of 30 minutes. 400 μL is then sprayed in air using an airbrush (pro series BD-132) via nitrogen gas flow at 1 bar pressure positioned 15 cm above a hot plate at 0° angle incident to substrate surface and annealed for 1 h at 150 °C to form a mesoporous TiO₂ layer (Fig. 1a.3). Methylammonium lead iodide (CH₃NH₃PbI₃) perovskite powder is mixed in dimethylformamide (DMF) at 60 °C at 0.25 M and 0.125 M (1 M = 0.619.9 g mL⁻¹) and using the same spray deposition technique with nitrogen gas flow and at 1 bar pressure the solution is sprayed in air on to mesoporous TiO₂ layer and annealed for 40 min at 100 °C (Fig. 1a.4). All samples were transferred directly to XPS vacuum chamber within minutes following the annealing procedure.

Deposition of metal contacts. Finally for use in photovoltaic device gold (Au) contacts are deposited using a Moorfield minilab DC/RF magnetron box sputter system using argon plasma at constant current of 0.15 A for 60 min with a working pressure of 1.5 × 10⁻² mbar (Fig. 1a.5). Energy band diagrams show that TiO₂ film acts as an electron transport/hole blocking material with Au contact creating a Schottky barrier at the back electrode (Fig. 1b).

Characterization of sprayed MAPbI₃ films

Cross sectional scanning electron microscopy (SEM) was performed using a JEOL JSM-6010PLUS at 20 kV acceleration voltages. X-ray photoelectron spectroscopy (XPS) to determine surface chemical composition was performed using an X-ray source (Al = 1486 eV) and the Kratos AxisUltra DLD spectrometer with peak fitting performed using Kratos software. Current and voltage were 10 mA and 15 kV respectively with an operating pressure of 10⁻⁹ bar. Specific region scans were performed at a resolution of 0.05 eV and pass energy of 20 eV. X-ray spot size and hence capture area is equal to 400 μm² with a penetration depth of 10 nm. Calibration of obtained spectra was performed using the C 1s peak located at 284.8 eV. Attenuated total reflectance Fourier transform infra-red spectroscopy (FTIR) was performed and measured using a Thermo-scientific spectrometer. Spectra were obtained over a range of 600–3500 cm⁻¹ at a 2 cm⁻¹ resolution. Raman spectroscopy was carried out using an ISA Lab-ram 300 confocal spectroscope with a 632.8 nm helium-neon laser to obtain spectra from 0 cm⁻¹ to 500 cm⁻¹ with 20 s exposure time and laser power of 10 mW. The Fermi level measurements are characterized using a scanning Kelvin probe system (KPTechnology) with a total of 20 measurements are taken across the sample surface and presented as a mean value with error bars representing standard deviation. Ultraviolet-visible (UV-Vis) absorption measurements are measured using a deuterium-halogen light source and an integrating sphere. Current density–voltage (JV) curves of devices incorporating sprayed MAPbI₃ films were recorded by a Keithley 6430 sub-FA sourcemeter unit under illumination 1.5 AM in air at ambient temperature. The applied bias varied from −1 V to 1 V in steps of 0.01 V at a scan rate 700 mV s⁻¹ consistently for all devices.

Results

Varying MAPbI₃ film thickness

MAPbI₃ films of varying thickness were produced by changing concentration (0.25 M and 0.125 M) and spray volume (200 μL, 300 μL and 400 μL). To track the changes in film thickness the cross sectional scanning electron microscopy (SEM) images were taken (Fig. 1c–h). Fig. 1c image was captured using back scattered electrons and presented to differentiate between meso-TiO₂/MAPbI₃/Au layers. A number of measurements using a profilometer for compact TiO₂ and mesoporous TiO₂ show thicknesses of 32.76 nm ± 12.16 and 569.25 nm ± 145.02 nm respectively. A plot of film thickness versus spray volume for 0.25 M and 0.125 M concentrations of MAPbI₃ in DMF show that film thickness decreases in a linear fashion with decreasing spray volume (Fig. 1i). Error bars represent 8 measurements taking across full width of film in order to highlight the homogeneousness nature of deposition process.

Chemical analysis of MAPbI₃ films

The XPS spectra were recorded to investigate the elemental composition and understand the impact that decreasing film thickness may factor into the surface stability and durability of our sprayed MAPbI₃ films. The XPS survey spectrum (ESI†) for a typical MAPbI₃ perovskite shows peaks at binding energies of 532 eV, 401 eV and 284.8 eV corresponding to the photoelectron peaks of O 1s, N 1s and C 1s respectively. The doublet peaks of I and Pb are located around 618 eV and 137 eV, respectively, with their associated spin orbit splitting. To assess the chemical changes of our MAPbI₃ films produced at decreasing thicknesses, we recorded the high resolution photoelectron spectra of each sample. XPS spectra were calibrated to 284.8 eV according to the C 1s binding energy for adventitious (aliphatic) carbon.²⁵ The core level spectra obtained were fitted using Gaussian line shapes after linear background subtraction. The atomic concentrations for each specific region were determined from the integrated area under their peaks accounting for photoionization cross section and plotted as a function of MAPbI₃ film thickness (Fig. 2a); error bars in the y-axis represent average and standard deviation (SD) for 3 different spots across sample surface. There is a comparable reduction in the elemental composition for I 3d and N 1s components as the thickness of the MAPbI₃ film reduces. The elemental composition for C 1s increases slightly whereas the Pb 4f component decreases by a similar amount, also with decreasing thickness. Most noticeably is the increase in O 1s concentration for decreasing MAPbI₃ film thicknesses.

Fig. 2 (a) Atomic concentration of C 1s, O 1s, N 1s, Pb 4f and I 3d components in MAPbI₃ films as a function of film thickness, determined from XPS measurements with (b) quantified XPS results highlighting atomic mass ratio for I/Pb and oxygen atomic mass percentage for decreasing MAPbI₃ film thickness.

The atomic mass ratio of I/Pb and the total atomic mass percentage of O 1s vs. MAPbI₃ film thickness, calculated from the integrated areas under Pb 4f, I 3d and O 1s peaks, show a mirrored symmetry (Fig. 2b). As highlighted previously, for decreasing MAPbI₃ film thickness, the total atomic concentration of the I 3d peak decreased gradually relative to the total concentration of the Pb 4f peak. This reduction (we expect a 3 : 1 I/Pb ratio in the stoichiometric crystal) implies that the amount of surface oxygen species correlates directly with the decrease in iodide concentration for MAPbI₃ films.

C 1s core level spectra for our MAPbI₃ with I/Pb ratio closer to the stoichiometric 3 : 1 (i.e. thickness >2 μm, Fig. 2b) included three peaks located at 284.8 eV, 285.9 eV and 287 eV (Fig. 3a–c). The peak found at 284.8 eV is commonly used for calibration and is known as adventitious carbon or adsorbed surface hydrocarbon species from the atmosphere.²⁵ The peak located at 285.9 eV can be attributed to methyl carbons in the MAPbI₃ film or additionally adsorbed surface carbon singly bonded to a hydroxyl group.²⁶ An additional peak that corresponds to carbon doubly bonded to oxygen can be found at 287 eV, which is also attributed to surface contamination. As MAPbI₃ film thickness reduces (Fig. 3d–f, corresponding to <2 μm thickness), reducing the I/Pb surface ratio, a carbon peak located at 289 eV appears and can be assigned to PbCO₃ which confirms the formation of a carbonate as an additional by-product during the decomposition of MAPbI₃ to form solid PbI₂.²⁷

Fig. 3 XPS core level spectra for C 1s where (a–f) represents decreasing MAPbI₃ film thickness: (a) 4.99 μm, (b) 4.65 μm, (c) 2.67 μm, (d) 1.71 μm, (e) 1.37 μm and (f) 0.67 μm.

Fig. 4a–f shows the XPS analysis of the O 1s region for MAPbI₃ films of varying thickness. The O 1s spectra for our samples can be initially de-convoluted into two peaks found at 531.8 eV and 533.5 eV corresponding to O=C and O=C–O respectively.²⁸ These peaks are typical for most air exposed samples and are known as adventitious carbon, comprising of a variety of hydrocarbon species with both singly and doubly bound oxygen functionality. The appearance of a peak found at 530.4 eV (O′) is attributed mainly to weakly adsorbed OH⁻ and O²⁻ ions but has also been referenced as intercalated lattice O or PbO/Pb(OH)₂ like states.^27,28 MAPbI₃ films are known to be significantly affected by surface oxidation where chemisorbed or dissociated oxygen species can diffuse and distort the MAPbI₃ framework causing degradation. As the concentration of iodide ions reduces at the surface of our MAPbI₃ films, the intercalated oxygen (O²⁻) can readily terminate Pb dangling bonds resulting in the formation of PbO/Pb(OH)₂ like states. However, through XRD measurements it has been highlighted that Pb–O bonds in MAPbI_2+x do not lead to the formation of secondary phases of PbO.²⁹ The atomic mass percentages for our de-convoluted O 1s peaks was calculated from their integrated areas and plotted against MAPbI₃ film thickness (Fig. 4g). The decrease in the O′ component coupled with the Increase in O=C and O=C–O components, highlight the ability for adsorbed oxygen species to more easily diffuse into thin MAPbI₃ films.

Fig. 4 (a) X-ray photoelectron spectroscopy analysis of the O 1s region for decreasing MAPbI₃ film thickness: (a) 4.99 μm, (b) 4.65 μm, (c) 2.67 μm, (d) 1.71 μm, (e) 1.37 μm and (f) 0.67 μm. (g) Summarized atomic mass concentration for each fitted component in the O 1s spectra.

Binding energy of metallic Pb is typically found at lower energies in the region 136.3–136.9 eV,³⁰ however during the perovskite formation there is a substantial alteration in Pb 4f core level due to the re-structuring of surface states to form Pb ions. The non-appearance of a metallic Pb state signifies strong charge transfer between Pb and I species to form donor–acceptor complexes. The Pb 6p valence state donates its excess unpaired electrons to the more electronegative iodide atoms where oxidation state Pb²⁺ is formed with iodide reducing to anion I₂⁻.^28,31 An additional iodide ion bonded with organic MA⁺ cation form weak van der Waals interaction with Pb–I framework to make up the PbI₆ octahedral within the perovskite lattice. It follows that the Pb 4f spectra of the MAPbI₃ films generally exhibit symmetrical peaks at 137.6 eV and 142.6 eV, corresponding to the spin orbit splitting of the Pb 4f_7/2 and 4f_5/2 components respectively that represent the Pb²⁺ cation in the MAPbI₃ structure.³² The Pb 4f_7/2 component of our MAPbI₃ films with decreasing thickness is presented in Fig. 5a–f; the peak at 142.6 eV is not shown here but remains constant throughout the samples. For thicker film (Fig. 5a and b), the Pb 4f_7/2 component is present with no other peak. As the film thickness is reduced, there is significant broadening in the accompanying Pb 4f spectra (Fig. 5c and d) that allows for the fitting of an additional peak around 138.4 eV attributed to the PbI₂ component.³³ The intensity of this PbI₂ peak increases with decreasing MAPbI₃ film thickness while the peak corresponding to Pb²⁺ state (137.6 eV) decreases (Fig. 5c–f). It is worth mentioning that a peak corresponding to PbO/Pb(OH)₂ in the Pb 4f spectra has shown to also have a binding energy of around 138.4 eV³⁴ and therefore may overlap with the PbI₂ component. Reviewed literature highlights that PbO/Pb(OH)₂ and PbI₂ components can be found within binding energy ranges of 137.7–138.5 eV and 138.3–138.9 eV, respectively.^35–37 However, weakly adsorbed oxygen species (O²⁻, OH⁻) highlighted previously in O 1s spectra (BE = 530.4 eV, Fig. 4a–f) indicate that only constrained/amorphous PbO/Pb(OH)₂ like states may be formed due to the termination of Pb dangling bonds (iodide vacancies). Formation of crystalline PbO in the O 1s region is typically found at lower binding energies around 529 eV³⁸ and therefore the loss of iodine ions resulting in PbI₂ bonding arrangement is mainly responsible for the broadening observed in the Pb 4f spectra. Furthermore, the appearance of a third peak located at 139.1 eV is attributed to PbC₂O₄/PbCO₃ bonds, consistent with Fig. 3, and has recently been shown to be an additional oxidation by-product of MAPbI₃.²⁷ The overall carbon concentration was observed to increase while the overall Pb concentration decreased by a comparable amount for thin MAPbI₃ films (Fig. 2a): the C 1s/Pb 4f atomic ratio increased around 1.3 times with decreasing thickness. Since Pb cannot leave the film it can be concluded that surface polymerization produces a surface hydrocarbon layer, which may also limit the escape depth of Pb 4f photoelectrons. The fact that the carbonate (PbC₂O₄/PbCO₃) component is more pronounced in the C 1s core level spectra compared with the Pb 4f core level is due to the photo ionization cross section of C 1s electrons being around 100 times lower than the Pb 4f electrons. Fig. 5g summarizes the concentrations derived from Pb 4f spectra, which highlights drastic changes taking place at the surface of the sprayed film. Thick films (>2 μm) resembling a typical MAPbI₃ composition and thin films (<2 μm) exhibiting surface changes.

Fig. 5 (a) X-ray photoelectron spectroscopy analysis of the Pb 4f region for decreasing MAPbI₃ film thickness: (a) 4.99 μm, (b) 4.65 μm, (c) 2.67 μm, (d) 1.71 μm, (e) 1.37 μm, (f) 0.67 μm and (g) summarized atomic concentration percentages for each fitted component in the Pb 4f spectra.

XPS for core level I 3d exhibits two intense peaks for thicker films, corresponding to doublets 3d_5/2 and 3d_3/2 with the lower binding component located at 618.9 eV and is assigned to triiodide I³⁻ (Fig. 6a–c).²⁸ The shift to higher binding energy 619.1 eV (Fig. 6d–f) is associated with the loss of an iodide ion through MAI evaporation leaving the I²⁻ anion from PbI₂.³³ Additional broadening of the I 3d spectra for thinner MAPbI₃ films indicate the presence an additional peak at 620.2 eV that is credited to the oxidation of iodine at the surface which forms the I₂⁺ cation and the iodite (IO₂⁻) anion.^28,30,33

Fig. 6 XPS core level spectra for I 3d where (a–f) represents decreasing MAPbI₃ film thickness: (a) 4.99 μm, (b) 4.65 μm, (c) 2.67 μm, (d) 1.71 μm, (e) 1.37 μm and (f) 0.67 μm.

Detailed XPS analysis of MAPbI₃ films clearly indicates that spraying thinner MAPbI₃ films allows for the formation of surface intermediate MAPbI_2+x phases and finally surface PbI₂ on the very top-layer (Fig. 7). The loss of iodine ions through the breakdown of the MAI component is closely linked with the increase in oxygen concentration at the films surface. For thinner MAPbI₃ films, the concentration of adsorbed OH⁻ and O²⁻ ions increased relative to adventitious hydrocarbons (O=C and O=C–O) that are typically found on samples surface. These surface adsorbed species diffuse into the MAPbI₃ and cause distortion of the perovskites framework where the MA cation is easily dissociated due to its weak interaction with PbI₆ network. The destruction of the MAPbI₃ lattice and specifically the organic core is also evident through the decrease in N 1s atomic concentration (Fig. 2a) even before NH_x bonds are absent from the surface (see ESI†). The breakdown of the MAI component creates a number of iodine vacancies that can be easily terminated by intercalated oxygen species to create amorphous and constrained bonding arrangements such as PbO/Pb(OH)₂, IO₂⁻ and PbCO₃/PbC₂O₄. MAI can be further broken down into the volatile components CH₃NH₂ and HI and are evaporated leaving solid PbI₂ at the surface.

Fig. 7 Schematic highlighting the changing surface chemistry for MAPbI₃ films of decreasing thickness: (a) 4.99 μm, (b) 4.65 μm, (c) 2.67 μm, (d) 1.71 μm, (e) 1.37 μm and (f) 0.67.

Surface morphology and grain analysis

SEM images of MAPbI₃ films show the changing surface morphology with decreasing film thickness (Fig. 8). The grains transition from large micro sized cube like structures typical of MAPbI₃ (Fig. 8a–c, thicker films) to de-agglomerated smaller distorted arrangements in thinner films (Fig. 8d–f) is evident, as the surface undergoes decomposition to MAPbI_2+x/PbI₂. Recent work involving density functional theory calculations have shown that defects, originating within the perovskites organic core, can be easily formed at room temperature.³⁹ The generation of these defects therefore enhances interactions with atmospheric species where it is understood that the deprotonation of MA cation is the first degradation step. The creation of intrinsic volatile components (HI & CH₃NH₂) and succeeding desorption process within the MAPbI₃ structure was shown to progressively erode the MAPbI₃ lattice.⁴⁰ Structural rearrangement of PbI₆ octahedral due to the loss of I atoms to volatile components MAI and HI allows for PbI₂ crystal seed to form simultaneously and gradually transform the MAPbI₃ lattice into PbI₂.⁴¹ The deagglomeration of large micro sized MAPbI₃ grains (Fig. 8c–f) signifies this structural reorganization, occurring at the interface of MAPbI₃ crystallites or at lattice defects within the grains. The subsequent increase in surface roughness accelerates the degradation process where micro sized PbI₂ hexagonal structures being to form visibly at the surface (Fig. 8d and e). Quantitative grain analysis is performed for each film by selecting a small area with >15 grains (insets Fig. 8a–f) and using imageJ software. Plotting a histogram of grain size for each film (Fig. 8g) and a graph of grain size as a function of MAPbI₃ thickness (Fig. 8h), it is observed that grain size decreases significantly as the surface decomposes from MAPbI₃ to MAPbI_2+x/PbI₂. The size distribution of Fig. 8g has been fitted with Gaussian fits to produce mean size and standard deviation summarized in Fig. 8h. For sprayed films where the surface was experimentally determined to have I/Pb atomic mass close to 2 : 1, (Fig. 8d–f), the small distorted grains create a largely porous layer that presents a rough and non-uniform surface. This structural transformation and decreased grain sizes associated with PbI₂ formation has been observed in multiple works, regardless of deposition technique.^42–44

Fig. 8 Surface scanning electron microscopy (SEM) images of MAPbI₃ films at decreasing thickness: (a) 4.99 μm, (b) 4.65 μm, (c) 2.67 μm, (d) 1.71 μm, (e) 1.37 μm and (f) 0.67 μm. Insets show increased magnification images used for size analysis (all scale bars are set to 5 μm including insets). (g) Histogram of grain size analysis in the range 0–3750 nm and bin size 300 nm for films (a–f) with (h) plot of average grain size as a function of MAPbI₃ film thickness where error bars for y-axis represent the SD of measured grain sizes for films (a–h).

Electronic band structure

Low binding energy XPS scans (Fig. 9a) for MAPbI₃ films with decreasing film thickness give insight into valence levels and can demonstrate more sensitivity to chemical interactions compared with core electrons. The XPS signal onset begins when highly kinetic/low binding photoelectrons from the outer most valence levels of MAPbI₃ are bombarded by the X-ray source, escaping to vacuum level and to the electron detector. Valence band spectra exhibits a broad signal in the range 0–6 eV that can be de-convoluted into two peaks corresponding to I 5p non-bonding state with a right shoulder peak assigned to the anti-bonding hybrid state of Pb 6s and I 5p state.^45,46 XPS analysis previously revealed that the total oxygen atomic concentration calculated from the integrated area under the core level O 1s increased for decreasing MAPbI₃ film thickness (Fig. 2a). After peak fitting the spectra it was specified that an increase in OH⁻ and O²⁻ ions was the main cause (Fig. 4), indicating that thinner MAPbI₃ films fabricated using one step spray method exhibited increased surface reactivity. The consequence of this increased interaction resulted in the enhanced surface decomposition of MAPbI₃ to form MAPbI_2+x through the evaporation of the MAI component. Once again it can be observed that reducing MAPbI₃ film thickness increases the concentration of Pb 6s and I 5p state anti-bonding hybrid state relative to I 5p non-bonding state. Additionally, the decomposition of MAPbI₃ to PbI₂ also shifts the onset of XPS signal to higher binding energies.

Fig. 9 (a) X-ray photoelectron spectroscopy (XPS) valence region scans for decreasing MAPbI₃ film thicknesses (4.99 μm, 4.65 μm, 2.67 μm, 1.71 μm, 1.37 μm, 0.67 μm) where E_f–VBM is valence band maximum energy relative to the Fermi energy (E_f). (b) E_f–VBM plotted as a function of MAPbI₃ film thicknesses and corresponding I/Pb atomic mass ratio. (c) Absolute Fermi level measurements determined using Kelvin probe plotted as a function of MAPbI₃ film thicknesses and corresponding I/Pb atomic mass ratio.

Extrapolating the XPS signal through the x-axis enables calculation of the valence band maximum (VBM) energy relative to the Fermi level (E_f) for the MAPbI₃ films. A number of studies have suggested that XPS-based measurements to determine VBM positions for air exposed MAPbI₃ films are more consistent and reproducible when compared to measurements derived from ultraviolet photoelectron spectroscopy, since the latter provides information only to a depth of only 0.5–1 nm whereas the former provides information up to a 10 nm penetration depth.^47–49 Therefore, XPS analysis minimizes effects of surface states due to contaminants on the electronic density of states profile since a larger penetration depth ensures measurement is more representative of bulk material properties. Plotting E_f–VBM for each of the MAPbI₃ films shows decreasing energy separation from nearly 0.9 eV to 0.55 eV when going from a degraded 2 : 1 (yellow) to a stoichiometric 3 : 1 (blue) for I : Pb atomic ratio (Fig. 9b); for convenience we included here also the I : Pb atomic ratio from Fig. 2b. This implies that the Fermi level and the valence band are shifting away from each other as the film thickness is reduced and therefore moving from thicker weakly p-doped film to a more n-type thinner films. Experimental studies have highlighted the effects of Pb and I vacancies in MAPbI₃ films and concluded that MAPbI₃ films that are MAI deficient/PbI₂ rich are typically n-doped whereas PbI₂ deficient/MAI rich films are p-doped and is in good agreement with our experimental results.^50,51

The VBM values determined using XPS measurements give values relative to the Fermi level which displays the energy separation and direction of Fermi movement successfully. However in order to establish absolute VBM positions relative to vacuum level the Fermi level of MAPbI₃ films must be first determined. The Fermi level of MAPbI₃ films has been measured using a Kelvin probe (Fig. 9c) that measures the contact potential difference of material relative to a known work function reference, in this case gold (−5.1 eV). Plotting MAPbI₃ Fermi level against film thickness shows that for thicker MAPbI₃ (I : Pb = 3 : 1) the Fermi level is −5.05 eV and is agreeable with literature;⁵² also here in this case we have included the I : Pb atomic ratio from Fig. 2b. Reducing the MAPbI₃ film thickness decreases the I : Pb ratio to around 2.5 : 1 meaning that a high number of MA and I vacancies are present at the film surface with consequent reduction (in absolute values) of the Fermi level to −5 eV. This confirms the effect of MA and I deficiencies, which promote n-doping behaviour in MAPbI₃ films. Decreasing film thickness further where the I : Pb atomic mass ratio is <2.2 : 1 causes a gradual decrease (in absolute value) in the Fermi level to around −4.8 eV. The initial slow shift of the Fermi level can be attributed to the loss of MA and I concentration and associated vacancies, however it is not until after PbI₂ begins to form on our film surface (I : Pb atomic ratio <2.2) that the Fermi level noticeably shifts. It should be emphasized that depending on the thickness of the PbI₂ surface layer there may be some Fermi level pinning from the underlying MAPbI₃/PbI₂ interface which can explain the gradual increase in Fermi level as the thickness of PbI₂ layer increases. As PbI₂ crystallizes and grows the influence of the underlying MAPbI₃ electronic states become less significant until the value of the Fermi level reduces (in absolute values) to that of pure PbI₂.

Photovoltaic performance of sprayed MAPbI₃ films

In order to understand the effect of the surface chemistry for spray deposited MAPbI₃ films, typical current density–voltage (JV) characteristics were measured under AM 1.5 illumination in air. An increase in current density (J_SC), open circuit voltage (V_OC) and overall power conversion efficiency (PCE) was observed for progressively thinner MAPbI₃ absorbers. The PCE increases exponentially from 0.83% to 5.23% with decreasing film thickness due to the matching of the MAPbI₃ carrier diffusion length (Fig. 10b). Examining individual parameters of our MAPbI₃ devices (V_OC, FF, J_SC) we find that the main reason for increased PCE is the significant increase in J_SC from 4.9 mA cm⁻² to 22.3 mA cm⁻² with the V_OC also increasing from 0.52 V to 0.73 V (Fig. 10c).

Fig. 10 (a) Illuminated current density–voltage characteristics of our devices with decreasing MAPbI₃ film thickness, (A–F) represents MAPbI₃ film thickness of (a) 4.99 μm, (b) 4.65 μm, (c) 2.67 μm, (d) 1.71 μm, (e) 1.37 μm and (f) 0.67 μm, (b) power conversion efficiency vs. decreasing MAPbI₃ film thickness and (c) individual device parameters V_OC, FF and J_SCvs. decreasing MAPbI₃ film thickness. Dotted lines are the result of exponential fitting using Origin software.

Depending on processing technique, diffusion lengths in the range 200–500 nm have been observed for MAPbI₃ films which closely match typical film thicknesses seen in high efficiency MAPbI₃ devices.⁵³ However for devices that have used spray coating, optimal thin films cannot be produced yet and thicker films are generally obtained. Also in our case, reducing the film thickness further to <500 nm using spray technique proved difficult and produced extremely porous films exhibiting small rod like structures assumed to be PbI₂ (ESI†). The incorporation of these low thickness films in PV devices resulted in shunting behavior that is attributed to the non-homogeneity of the sprayed film. The increased degradation for thin sprayed films (<500 nm) is in good agreement with trend of the surface chemical changes for decreasing MAPbI₃ films (Fig. 2, 4 and 5). Our thinnest sprayed MAPbI3 film (670 nm thickness) gave the greatest PCE of 5.23% due to the optimization of absorber thickness close to carrier diffusion lengths.

The consistently low FF of around 0.33 could be in principle improved considerably by including a hole transport material (HTM), where incorporation of Spiro-OMeTAD in perovskite based devices has shown efficiencies increase beyond 20%.^2,3 However, since the spray deposition method used in this technique is fairly novel and to better assess the band alignment and energy band level changes with respect to a known stable metal (i.e. Au), it was decided to first understand the surface chemistry of as deposited MAPbI₃ films before inclusion of sequential layers. As the inclusion of an HTM such as spiro-OMeTAD has been widely reported to improve device efficiency, we expect this to happen also for these sprayed MAPbI₃ films. In any case, as our work clearly point out, spray deposition still has to overcome degradation challenges before optimal thin perovskite films can be produced to be competitive with devices produced by other methods. According to the equivalent circuit (ESI†), the JV characteristic can be described as

where J_L is the light induced constant current density, J₀ is the reverse saturated current density, R_S is the series resistance, R_sh is the shunt resistance, J is the current density flowing through the external load, n is the ideality factor, k_B is the Boltzmann constant, T is the absolute temperature, q is the elementary charge, A is the surface area of the cell and V is the direct-current (DC) bias voltage applied at the cell.⁵⁴ Based on eqn (1), some intrinsic parameters of our cells can be derived from the JV measurements. The first derivative of the voltage vs. current density can be used to extract useful parameters.

Fig. 11a shows the plots of −dV/dJ vs. (J_SC − J)⁻¹ and the linear fitting curves according to eqn (2). Our cells show a linear relationship between −dV/dJ and (J_SC − J)⁻¹ in the region approaching the V_OC. The area normalized series resistance (AR_S) of the cells can be deduced from the intercept of the linear fitting results in the region approaching (J_SC − J)⁻¹ = 0.⁵⁵ We observe AR_S of our devices reducing from 33.4 Ω cm² to 7.2 Ω cm² with decreasing MAPbI₃ film thickness and as a consequence we observe a large increase in the devices J_SC. Additionally, the JV curves can be used to determine the voltage dependant ideality factor (n) of devices for values from 0 V up to the corresponding V_OC (Fig. 11b). When plotting the natural log of the current density against the voltage, the slope q/nk_BT gives a value for n.⁵⁶ In a heterojunction solar cell, values of n can give an insight into the quality of a junction and the carrier recombination mechanisms. The ideality factor is typically in the range 1 < n < 2, for a “well-behaved” heterojunction.⁵⁶ When the junction is dominated by the carrier diffusion in the neutral region, n is at low values of 1. On the other hand, when carrier recombination in the depleted space charge region dominates the current, n is moving away from unity. In most cases, both carrier diffusion and recombination take place simultaneously.^55,56

Fig. 11 (a) Plot of −dV/dJ vs. (J_SC − J)⁻¹ for devices with decreasing MAPbI₃ film thickness and extrapolated linear fit near (J_SC − J)⁻¹ = 0 (b) voltage dependant ideality factor of the devices with abnormally low and high values emphasized. (A–F) Represents MAPbI₃ film thickness of (a) 4.99 μm, (b) 4.65 μm, (c) 2.67 μm, (d) 1.71 μm, (e) 1.37 μm and (f) 0.67 μm.

High values for ideality factor can also be caused by various other effects such as interface states, tunnelling processes and non-uniform distribution of interfacial charges.^57–59 More specifically interface dipoles due to interface doping as well as fabrication induced defects play an important role in the performance and reliability of metal/semiconductor structures.⁶⁰ Looking at Fig. 11b, we see that n increases with voltage approaching V_OC for each device. For the two devices with the greatest MAPbI₃ film thickness (black squares/red circles), the ideality factor is between values of 1 to 2 in the range 0 V to the device V_OC. As we reduce our MAPbI₃ film thickness within our solar cell, we see that the value of n beings to exceed 2 in the region approaching each devices specific V_OC and falls below 1 at voltages <0.3 V. Ideality factor moving away from unity approaching V_OC has previously been highlighted for MAPbI₃ HTM free solar cells^56,61 and suggests that the current is mainly controlled by the recombination in the space charge region of the TiO₂ and MAPbI₃ layers. However, ideality factors above 2 have been previously shown to be a clear indication of trap assisted tunnelling or field enhanced recombination through isolated point defects.^62,63

Nonetheless, the highlighted areas suggest poor Schottky contact effect in our devices for thin MAPbI₃ films. The fact that decreasing MAPbI₃ film thickness produces smaller grain boundaries as highlighted earlier, may explain why the ideality factor drops below unity. The smaller grains/increased grain boundaries clearly increase the potential for recombination through surface states before carriers can be successfully collected at the back contact. Light soaking effects over a period of 0–6 min, including effect on individual parameters were collected (ESI†). Comparing light soaking effects on PCE for 3 devices that represent the maximum, minimum and intermediate MAPbI₃ absorber thickness (Fig. 12), a distinct trend can be observed. Thick MAPbI₃ film exhibits lowest PCE (<1%) most likely due to the thickness being larger than the carrier diffusion length; however with continued light soaking over a time period of 6 min the PCE significantly increases to more than 3%. For MAPbI₃ film with intermediate thickness (i.e. 2.67 μm) the beneficial light soaking effect is diminished and the PCE drops slightly from 2.9% to around 2% over the same time period. Finally for the thinner MAPbI₃ film (0.67 μm), the PCE at 0 min illumination is more than 5%. However increasing the time under illumination induces notable negative impact with the PCE decreasing from >5% to less than 2%.

Fig. 12 Power conversion efficiency of devices with maximum, intermediate and minimum MAPbI₃ film thickness of 5 μm, 2.6 μm and 0.6 μm respectively, at 0 min, 2 min, 4 min and 6 min under illumination.

Discussion

PV performance of our thicker MAPbI₃ devices (∼5 μm) showed a relatively low J_SC measured at 0 minutes (Fig. 10c) before increasing considerably after 6 min light soaking (4.9 mA cm⁻² to 14 mA cm⁻², see ESI†), corresponding to the behaviour of the PCE (Fig. 12). This light soaking phenomena is generally observed for perovskite devices where the PCE is typically presented after stabilization at the maximum power point which takes place in a short time scale (0–3 min).^64,65 The stabilization period is understood to be caused by the filling of shallow charge trap sites by photo generated electrons or migrating ions, shifting the quasi-Fermi-level for electrons and therefore V_OC increases (Fig. 13, 0 min). The reduction in trap site density also increases J_SC accordingly due to the reduction in non-radiative pathways available.⁶⁴ Recently photo-induced brightening of MAPbI₃ photoluminescence was demonstrated and attributed to an order of magnitude reduction in trap state density due to migration of iodide ions.⁷ Due to the high number of oxygen vacancies (V_O) that are expected in TiO₂, many of the photo excited carriers (e⁻) may become trapped at the interface which can create electric fields and encourage iodide ion (I⁻) migration.⁶⁶ The mobile iodide ions migrate away from the incident illumination into the MAPbI₃ film and fill iodide vacancies (V_I) that translate as shallow states below (above) the CBM (VBM) (Fig. 13, 2–4 min).⁶⁷ The annihilation of these defects under increased light soaking would result in the shifting of quasi-Fermi level and the reduction of non-radiative recombination pathways, largely increasing open circuit voltage and photo carrier collection (Fig. 13, 6 min).

Fig. 13 Schematic showing the potential photo induced cleaning due to halide distribution that can explain light soaking stabilization period experienced in MAPbI₃ films under illumination.

Depending on processing technique, diffusion lengths in the range 200–500 nm have been observed for MAPbI₃ films which closely match typical film thicknesses seen in high efficiency MAPbI₃ devices.⁵³ Therefore, in order to optimize the photo generated current from spray deposited MAPbI₃ films presented in this work, absorber thickness was reduced to around the same magnitude of carrier diffusion lengths. Decreasing the MAPbI₃ film thickness increased overall carrier collection as a substantial increase in PV performance at 0 minutes is observed for devices (Fig. 10b and c). It is known that MAPbI₃ films are susceptible to chemical decomposition that originates from chemisorbed or dissociated oxygen species and XPS analysis has highlighted that this modification is accelerated at the surface of thin spray-coated films with the increase of OH⁻/O²⁻ surface species. Energy bands are displayed for MAPbI₃ films of decreasing thickness and changing surface chemistry (Fig. 14). Figure labels (a–f) correspond to varying film thickness of 4.99 μm, 4.65 μm, 2.67 μm, 1.71 μm, 1.37 μm and 0.67 μm, respectively. XPS, Kelvin probe and UV-Vis measurements correlate well to show that reduced PbI₂ on the surface creates a type-1 hetero-junction with the underlying MAPbI₃. Energy levels presented here are non-equilibrated however it is evident that a significant potential barrier would exist for hole extraction at the back electrode (Fig. 14c–f). This is in agreement with a previous report that claimed a type-1 alignment of PbI₂ with MAPbI₃ (VB and CB located at −5.75 and −3.45 eV) and a potential barrier of 30–50 mV.⁶⁸

Fig. 14 Energy band structures for our device architectures showing changes in surface chemistry and how subsequent electron transfer process may be affected for decreasing MAPbI₃ film thickness: (a) 4.99 μm, (b) 4.65 μm, (c) 2.67 μm, (d) 1.71 μm, (e) 1.37 μm and (f) 0.67 μm. Energy level are non-equilibrated. The absolute values of the band edges and Fermi levels for the MAPI₃ and PbI₂ layer are taken from our experimental measurements (Fig. 9b and c) and bandgap values from literature.

From theoretical simulations, the majority of defects in MAPbI₃ have low formation energies and create shallow levels within the bandgap,⁶⁷ which are unlikely to influence our Fermi level measurements. However these theoretical simulations were carried out on the assumption of pure perovskite phases. It is important to note however that the formation of different phases at the surface of MAPbI₃ films, such as those highlighted in this study (i.e. PbI2, Fig. 14c–f), can result in low energy/deep defect states and potentially influence the Fermi levels determined experimentally in this work.⁶⁹ Moreover, these defect states can provide mid gap levels at which carriers can recombine thus suppressing carrier collection at the back electrode.

Counter to the beneficial light soaking effects observed for MAPbI₃ devices in this work (for thicker films, Fig. 12 and 14a, b) and also in the literature, we notice a contradictory trend for spray-coated thin films (Fig. 12 and 14c–f) that we attribute directly to the changing surface chemistry and subsequent defect states at the back electrode. A MAPbI_2+x film (I/Pb atomic ratio = 2.5 : 1) whose surface (<10 nm) has been degraded from the initial stoichiometric crystal showed a decrease in J_SC over a period of 0–6 min under illumination (15.15 mA cm⁻² to 13.86 mA cm⁻²). This decrease in PV performance correlates directly with the changing surface chemistry where an increase in adsorbed oxygen species has resulted in the decomposition of MAPbI₃ structure. This instability under light soaking is exacerbated for a number of thinner films that show an initially high PCE before J_SC and V_OC significantly decrease over a period of 6 min. More specifically, a device incorporating a film with measured thicknesses of 0.67 μm showed the greatest PCE of 5.23% which decreased to 1.3% after 6 minutes continuous light soaking (Fig. 12). The quantified surface composition determined from XPS measurements showed an overall increase in O 1s atomic mass that was explicitly due to intercalated OH⁻ and O²⁻ species originating from adsorbed atmospheric species. The different morphologies are the determining factor for the surface chemistry where the increased surface to volume ratio that exist for thinner MAPbI₃ sprayed films imposes a large degree of surface stress and strain which limits the chemical stability under ambient conditions. In turn, the formation of different grain size distributions is a consequence of the aerosol droplet size and velocity distributions which affect how the droplets coalescence and subsequently evaporate on the dynamic growing surface. These factors are in turn influenced by input parameters (solution composition, gas flow and spraying conditions) and the evolving topography and thermal properties of the surface. Present results show that, for thinner MAPbI₃ films, a morphology with larger surface area increases the surface interactions with ambient atmospheric species, which are most likely responsible for a higher density of surface defects. These defects originate during the spray process annealing stage which allows the continuous adsorption of atmospheric species. The increased adsorption and interaction of OH⁻/O²⁻ species leads to distortion of the perovskite lattice where the MA cation is originally dissociated due to weak van der Waals bonds with PbI₆ network. Recent work has highlighted that water vapour reactions, which lead to the formation of MAPbI₃–H₂O and MAPbI₃–2H₂O hydrate phases and consequential hysteresis in solar devices, are reversible.⁷⁰ Water molecules form hydrogen bonds with surface halide, however no proton transfer happens unlike the situation with adsorbed OH radicals and OH⁻ species, where accelerated perovskite decomposition has been shown to occur.⁷⁰ The resulting hydrogen/proton extraction from the surface MA cations forming surface CH₃NH₂ is suggested as the initial step in the degradation process with further breakdown producing NH₂ and HI components. The former can be easily evaporated whereas the latter can react with O₂ to form I₂ and H₂O or decompose into I₂ and H₂.⁷¹ The same degradation path is evident here due to the experimentally observed decrease in both N 1s and I 3d atomic mass concentrations in addition to the Pb 4f core level shifting to a higher binding energies associated with the loss of iodide ions.

The interplay between oxygen and MAPbI₃ films has been well documented where films have exhibited poor morphologies and incomplete surface coverage under high oxygen concentrations. The distorted grain structure among inferior crystallinity observed in MAPbI₃ films in this study is clear evidence of oxygen detrimental effect on chemical stability. Moreover, electron–hole recombination lifetimes have been shown to dramatically shorten with increasing oxygen concentration along with substantial quenching of photoluminescence,^29,72 suggesting that intercalated oxygen species may act as deep trap centres for photo excited carriers. The high ideality factors demonstrated for thin MAPbI₃ films (Fig. 11b) indicated poor Schottky contact effect and trap assisted recombination at the back electrode. The fact that devices, where surface decomposition (<10 nm) has occurred (MAPbI₃ to MAPbI_2+x/PbI₂), exhibit good PV performance initially at 0 minutes before the photocurrent and open circuit voltage are degraded under continued light soaking (0–6 minute) suggests that trap sites become filled before consequently forming macroscopic charged regions, increasing the potential barrier between bulk MAPbI₃ and surface PbI₂.

Conclusion

A rapid one step spray technique that possesses the potential for large scale manufacture has been implemented to fabricate MAPbI₃ absorber films with varying thickness and corresponding power conversion efficiencies from 0.83% to 5.22%. Investigation into the stability of devices under constant illumination showed conflicting behaviour with low efficiency (thick) MAPbI₃ films and high efficiency (thin) MAPbI₃ films, where the former displayed large increases in overall performance compared to rapid reductions experienced in the latter. It was experimentally concluded that the surface modification was directly related to instability of device performance under light soaking conditions. Determination of the Fermi level showed a net upward shift near the surface when transitioning from pristine MAPbI₃ to MAPbI_2+x/PbI₂. The measured energy separation E_f–VBM showed a comparable increase resulting in a potential barrier at the back contact that would significantly imped hole extraction. It was predicted however that the combined effect of intercalated OH⁻/O²⁻ species, constrained amorphous states and dangling bonds due to I vacancies created a high density of shallow and deep level trap states that assisted in the non-radiative recombination of carriers. The spray deposition process demonstrated in this work can accelerate surface modification for thin films which is detrimental to PV performance under light soaking. Overall, this study suggests that deposition by spray method may enhance interactions of the solution being sprayed with atmosphere oxygen. The deployment of spray technologies will therefore need to take into account and reduce such detrimental impact. Future work will involve improvement of the one-step spray technique for MAPbI₃ thin film deposition for use in high efficiency solar cells among other applications.

Acknowledgements

This work has been partially funded by the Leverhulme Trust (IN-2012-136) and the EPSRC (EP/K022237/1 and EP/M024938/1). CR would like to acknowledge the support for the Northern Ireland Department for Employment and Learning (DEL). This work was also partially supported by a NEDO Project (Japan).

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Footnote

† Electronic supplementary information (ESI) available: XPS survey and N 1s core level scans of sprayed MAPbI₃ films, FTIR and Raman spectroscopy for atmospheric aging of pristine MAPbI₃. JV curves 0–6 minutes under illumination with plotted J_SC and V_OC. See DOI: 10.1039/c6tc04864a

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