Photoinduced Stark Effects and Mechanism of Ion Displacement in

Photoinduced Stark Effects and Mechanism of Ion Displacement in Perovskite Solar Cell Materials Meysam Pazoki,†,⊥ T. Jesper Jacobsson,‡ Jolla Kullgren,⊥ Erik M. J. Johansson,† Anders Hagfeldt,‡ Gerrit Boschloo,*,† and Tomas Edvinsson*,∥ †

Department of Chemistry − Ångström Laboratory, Physical Chemistry, Uppsala University, Box 523, SE 75120 Uppsala, Sweden Laboratory of Photomolecular Science, Department of Chemistry and Chemical Engineering, Swiss Federal Institute of Technology, Station 6, CH-1015 Lausanne, Switzerland ⊥ Department of Chemistry − Ångström Laboratory, Structural Chemistry, Uppsala University, Box 538, SE 75120 Uppsala, Sweden ∥ Department of Engineering Sciences, Solid State Physics, Uppsala University, Box 534, SE 751 21 Uppsala, Sweden ‡

S Supporting Information *

ABSTRACT: Organometallic halide perovskites (OMHPs) have recently emerged as a promising class of materials in photovoltaic technology. Here, we present an in-depth investigation of the physics in these systems by measuring the photoinduced absorption (PIA) in OMHPs as a function of materials composition, excitation wavelength, and modulation frequency. We report a photoinduced Stark effect that depends on the excitation wavelength and on the dipole strength of the monovalent cations in the A position of the ABX3 perovskite. The results presented are corroborated by density functional theory calculations and provide fundamental information about the photoinduced local electric field change under blue and red excitation as well as insights into the mechanism of lightinduced ion displacement in OMHPs. For optimized perovskite solar cell devices beyond 19% efficiency, we show that excess thermalization energy of blue photons plays a role in overcoming the activation energy for ion diffusion. KEYWORDS: Stark effect, photoinduced ion migration, perovskite solar cells, CH3NH3PbI3, mixed halide perovskites, cation-dependent ion movement conversion efficiencies (PCE) of more than 20% feasible.2 Perovskite solar cells have rapidly developed from a reported efficiency of 9.7% in 20127 to a certified efficiency of 22.1% in 2016,8 soon after a report for implementation of a lead halide perovskite material as a sensitizer in a dye solar cell configuration was published in 2009 by Miyasaka et al.6 Hybrid lead halide perovskites exhibit a plethora of interesting fundamental phenomena such as ferroelectric effects,9,10 anomalous hysteresis,11,12Stark effects,13 giant dielectric constants,14 giant switchable photovoltaic effects,15 and extremely slow photoconductivity response,16 which may play key roles in the final device performance. The precise origin and importance of these effects are still largely unclear, and an increased fundamental understanding of the excited-

S

olar cell technologies are anticipated to play a key role in clean and sustainable energy production in the future. Recently, organometallic halide perovskite materials have appeared as promising light-absorbant and charge-transport materials, leading to the emergence of hybrid perovskite solar cells.1,2 High-efficiency and low-cost solution-processed fabrication methods together with spectacular material properties make hybrid perovskite solar cells a serious candidate for largescale solar energy conversion,3 which recently was distinguished in Nature Materials research highlights.4,5 Organometallic halide materials with general formula ABX3, where A represents an organic molecular cation, X is a halide, and B is lead or tin, crystallize in the perovskite structure. These hybrid perovskites have a combination of favorable characteristics for thin-film photovoltaics, such as long charge carrier lifetime, strong optical absorption, low trap densities and high dielectric constant. In solar cells, the photoactive perovskite layer is sandwiched between two selective contacts3 where a few hundred nanometer thick active layers make power © 2017 American Chemical Society

Received: November 24, 2016 Accepted: February 27, 2017 Published: February 27, 2017 2823

DOI: 10.1021/acsnano.6b07916 ACS Nano 2017, 11, 2823−2834

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Figure 1. Schematic illustration of the experimental (a) EA and (b) PIA setup, (c) first-order derivative of the absorption spectra of MAPbI3, FAPbI3, and CsPbI3, (d) PIA spectra of MAPbI3, FAPbI3, and CsPbI3, and (e) linear dependency of observed stark effect versus light intensity for MAPbI3. The excitation wavelength in (d) is 470 nm. The presented data in (d) are representative of spectral features of the stark effect and are not normalized versus the calibration factors.

state properties and charge migration processes will be important for further device optimization.17,18 Investigations of the interaction of the different perovskite cations with the photogenerated charge carriers in lead iodide perovskites are important for understanding the device behavior9,16,19 and are, together with the photoinduced ion displacement, the main topic of this paper. Stark effects, spectral changes in the presence of an external electric field, have previously been observed in both quantumconfined semiconductor structures20 and in dye-sensitized solar cells (DSC)21 and have been used for charge compensation studies in DSCs.22,23 Stark effects in hybrid perovskite materials were reported by Listorti et al.13 in 2013, who attributed the observed spectral changes to interfacial dipole moment at the TiO2/MAPbI3 interface. More recently, Wu et al. investigated second harmonic electro-reflectance spectra of MAPbI3 (CH3NH3PbI3, methylammonium lead iodide) and FAPbI3 (CH(NH2)2PbI3, formamidinium lead iodide). They instead attributed the observed spectral changes to changes in the light induced dipole moment in the bulk of the perovskite material.24

Here, we perform a detailed investigation of photoinduced Stark effects in perovskite solar cell materials, which hold noteworthy information about the local electric fields within the material. The origin of the effect is explored by varying the monovalent cation dipole and measuring the excitation energy and modulation frequency dependence of the Stark effect using photoinduced absorption (PIA) spectroscopy combined with electro-absorption (EA) measurements. The molecular cations utilized in the hybrid lead iodide perovskites are methylammonium (MA) and formamidinium (FA), which have different dipole moments. To investigate the role of the dipole moment of the positive cation in the perovskite, the results were compared with an analogous system with a cation carrying no net dipole moment (Cs+). We present an experimental inspection of the photoinduced Stark shifts and thus the local electric fields in the polycrystalline perovskite films where the charge response density and the corresponding dynamics dependency on the surrounding cation dipoles under red and blue light illumination are investigated in a frequency range from 10 to 105 Hz. Time-dependent DFT calculations with 2824


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Figure 2. Frequency dependence of the normalized photoinduced bleach (ΔT/T) for OMHP films under (a) blue light excitation (470 nm) and (b) red light excitation (630 nm).

Figure 3. Frequency dependence of the observed photoinduced bleach and the effect of excitation wavelength and A cation in OMHP thin films: (a) MAPbI3, (b) FAPbI3, and (c) CsPbI3.The excitation wavelengths for blue and red are 470 and 630 nm, respectively.

by spectral shifts as a function of excitation energy attributed to a Burstein−Moss (BM) shift,17 i.e. a shift of the optical absorption by filling of the valence/conduction band states by charge carriers. In this study, however, no such shifts are found in the PIA spectra (Figure S2). The PIA signals found here also persists up to millisecond−second time scale, much longer than the lifetime of free charge carriers in the perovskite, effectively showing another origin of the observed bleach. An external applied electric field can also cause an exponential decay of absorption coefficient below the band edge and oscillations of absorption above the band edge in a so-called Franz-Keldysh (FK) effect.25 A FK effect can be detected by oscillations in EA spectra for energies higher than the band gap of the semiconductor while no such effects are found in the PIA spectra recorded here. Hence, all data show a photoinduced Stark effect arising from photoinduced internal electric fields (see also section S3 in the Supporting Information). To determine whether the photoinduced change in absorption is an intrinsic property of the perovskite material, or caused by interfacial effects as suggested previously,13 PIA spectra were recorded for MAPbI3 films deposited on a variety of mesoporous scaffold layers, TiO2, ZrO2 and Al2O3 with different dimensions, and on bare glass. In all cases, significant photoinduced bleach was found (Figure S3 and section S1 of Supporting Information). These observations suggest that the observed photoinduced effect is intrinsic to the perovskite material and not sensitive to the proximity or type of another surface. This in agreement with the recent work of Wu et al.,24 who observed a Stark effect in an electroreflectance study of MAPbI3 and FAPbI3 thin films.

photoinduced charge response as well as thermally induced ion displacements collaborate the experiments and provide insights into the fundamental thermalization mechanisms of these materials during illumination.

RESULTS Thin films of hybrid lead iodide perovskites, MAPbI3 and FAPbI3, as well as an inorganic perovskite, CsPbI3, were deposited by spin coating on glass substrates with a thin Al2O3 nanoparticle scaffold layer (see the Supporting Information for synthetic details and optical and XRD characterization). PIA measurements displayed significant signals at wavelengths near the absorption onset of the perovskite films (Figure 1). A negative absorption peak is found that corresponds to a slight blue-shift of the absorption spectra upon the excitation. The PIA spectra closely parallel the first-order derivative of the absorption spectra (Figure 1c), showing minima at 765 nm for MAPbI3, 795 nm for FAPbI3, and 700 nm for CsMAPbI3. Additional EA measurements, where an external electric field is applied across the perovskite layer, yielded very similar spectra (Figure S1). These observations provide evidence that the photoinduced optical changes in the perovskite films are due to the Stark effect where an electric field perturbs the system (Supporting Information, section S4), resulting in a change in the optical absorption within the material. The amplitude of the PIA signal increases linearly with excitation power (Figure 1e), consistent with first-order kinetics of the disappearance of the photoinduced state. Recently, on the picosecond time scale, a bleach in the photoinduced absorption spectra of MAPbI3, was corroborated 2825


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Figure 4. DFT-optimized ground-state structures and experimental determination of the frequency-dependent Stark effect. (a) Ground-state geometry of MAPbI3 in super cell calculations showing the local position of MA; the dipole of the MA cation calculated in the gas phase. (b) Ground-state geometry of FAPbI3 in super cell calculations showing the local position of FA; the dipole of the FA cation calculated in gas phase. (c) PIA spectra of MAPbI3 in intermediate (IF, 1 kHz) and high frequencies (HF, 10 kHz) measured in first (1H), second (2H), and third harmonics (3H). (d) PIA spectra of FAPbI3 in intermediate (1 kHz) and high frequencies (10 kHz) measured in first, second, and third harmonics. The excitation wavelength in (c) and (d) is 470 nm.

The effect of the monovalent “A” cation on the observed PIA of the APbI3 perovskite thin film was investigated in more detail using frequency-dependent PIA measurements (Figure 2). In this series, the dipole moment of the A cation decreases from MA (2.22 D), via FA (0.19 D) to Cs (no dipole). Cs with no dipole delivers the strongest PIA signal where the dipoles in the material seem to contribute in shielding of the changes in the local electric fields that cause the photoinduced Stark effect. Interestingly, the wavelength of excitation has significant effect on the PIA response of perovskite thin films. The blue light excitation response is quite flat in the modulation frequency range of 10−105 Hz, while the response under red light excitation was found to be larger when the same photon flux was used for excitation, and it displayed an increase in amplitude with increasing frequency. Comparing the frequency response under red and blue excitation for the different cations (Figure 3), a shift to higher frequencies of the crossing between the bleaching response under blue and red excitation is seen for the perovskites with a low/absent dipole in the A cation. Different features can appear in the delta absorption spectrum where the background level may shift and interfere with the Stark features. Here, experiments quantifying the signals that can be accounted for by the background in comparison to the change in absorption have to be performed. In our case, all the Stark features are clearly distinguishable from the background shifts for all frequencies except the highest frequencies where the relative intensities of the background and the Stark shifts become comparable (See Supporting Information S7 and Figure S16). The different orders of the Stark effects can be quantified by analyzing the first, second, and third harmonics in the PIA measurements for high and low dipolar A cations would give

some extra information on an overall mechanism (see section S3 and Figure S4−5 of Supporting Information). The second harmonic PIA spectrum, which should correspond to the second order Stark effect, interestingly changes sign for MAPbI3 and not for FAPbI3 (Figure 4) and thus reveal a reversal of the local polarizability for the material with the strongest dipolar cation. The Stark effect response is also analogous to the shape and wavelength of the first derivative of the absorption spectra (Figure 1c) giving further support that it is consistent with a photoinduced effect. This effect also correlates with the different transition dipole moment of FAPbI3 and MAPbI3, and could thus be related to both the higher binding energy of FA in the FAPbI3 structure which has a higher number of possible hydrogen bonds compared to MA,26 as well as secondary effects investigated in more detail below. To investigate the Stark effect in more complete devices and its relevance for understanding the device physics, state-of-theart perovskite solar cells with efficiencies beyond 19% were fabricated (Figure 5a, here mixed perovskite27 with compositional formula MA0.33FA0.77PbI2.5Br0.5 that shows the highest reported efficiencies so far). The Stark effect was clearly seen for the full device with mixed organic cations as well, and the response closely resembles the Stark spectrum of MAPbI3 (Figure 5b), which is as expected since the charge-transporting layers have no interfering features near the wavelengths of band edge absorption.28 We can therefore conclude that the observations made for the perovskite films above also are valid for perovskite films in high-efficiency devices with mixed cations. A hysteresis effect11 in the current voltage curves is observed here when the device is scanned in forward or backward scans and give an overall power conversion efficiency (PCE) of η = 19.2% in the backward direction and η = 16.1% in 2826


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Figure 5. Perovskite solar cell device: performance, photoinduced absorption, and photovoltage decay. (a) Current−voltage characteristics under AM1.5G light (100 mW cm−2); (b) PIA spectrum; (c) voltage decay; (d) frequency response of the Stark effect with real (Re) and imaginary (Im) parts. Voltage decay is measured after different illumination times by red and blue LEDs. The measured device is from the mixed perovskite family with compositional formulas MA0.33FA0.77PbI2.5Br0.5 that shows the highest reported efficiencies so far.

structure. Calculations were also utilized to extract the energy barriers for A cation related ion diffusion from first-principles (Figures 6 and 7 and Table 1) and quantification of the iodide diffusivity in the different systems. The bottom of the conduction band (CB) of the hybrid perovskites consists of molecular orbitals dominated by Pb states, while the valence band (VB) upper states have a strong contribution from iodide (Figure 6), in agreement with recent studies.29,30 The organic cations MA/FA yield energy levels more than 1 eV above the CB band edge (Figure 6a). The energetic position of the unoccupied organic states is here significantly lowered in the perovskites crystals compared to the virtual states of MA and FA in the gas phase, which can be ascribed to a stabilization of the organic cations in the negatively charged inorganic framework. The virtual states of MA are higher up than FA and Cs states, consistent with the relative strength of the dipole of the cations (Figure 5a,c and Figure S16). The exact determination of the absolute energetic positions of the states has to be taken with care though, due to the well-known band gap uncertainty in the generalized gradient approximation (GGA). The charge density response to an incoming light was calculated by TD-DFT to quantify the charge response in red and blue light (Figure 6c,d). Here, red (730 nm) and UV/blue (395 nm) excitation were chosen to match the two distinct characters from the partial density of states (PDOS). Red light excitation mainly leads to an electronic charge transfer from Ibased to Pb-based orbitals. In contrast, the charge response shows that higher energy excitation (blue/UV light) has a possible charge transfer to MA/FA localized states (I → MA/

the forward direction with the largest hysteresis difference in the fill factor (FF) (Figure 5a). Encouraged by the large differences observed for blue and red light excitation in PIA measurements, we investigated the Voc decay in dark after illumination with blue or red light. Intriguing differences were found: After red light (630 nm) illumination the Voc decayed relatively rapidly, but after blue illumination (470 nm) with the same photon flux, a much slower decay was found, even though the initial voltage was approximately the same (Figure 5c). When the illumination time with the blue photons was varied from 3 to 30 s, an extended decay time was observed (Figure 5c) and complied well with an increased number of migrated ions. To further verify the interdependence of the activation energy in blue light and the drift in a photovoltaic field, a blue light illuminated solar cell was switched into dark and then into red light illumination (Figure 5c). The red light applied directly after the blue light revealed a markedly extended tail in Voc, in comparison to only red light illumination; its relevance to Stark spectrum and ion movement would be further discussed. Varying the illumination wavelength and monitoring the Voc decay would thus provide a promising method to pinpoint the ion movement in different perovskite compositions assisted by thermalization of charge carriers. Quantum chemical calculations, ground-state DFT, and timedependent (TD) DFT were performed in order to relate the experimental observations to optical transitions, excitation wavelength dependency of charge response, and implications for bond length relaxations and ion displacements in the crystal 2827


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Figure 6. PDOS of DFT-optimized ground-state structures and charge response from TD-DFT. (a) Calculated MAPbI3 (P)DOS with energy states available upon excitation indicated with red and blue arrows. (b) TD-DFT calculations of the charge response for z-axis polarization upon excitation using UV/blue light and red light for MAPbI3. (c) Calculated FAPbI3 (P)DOS with energy states available upon excitation indicated. (d) TD-DFT calculations of the charge response for z-axis polarization upon excitation using UV/blue light and red light for FAPbI3.

Figure 7. Schematic of suggested mechanism for the excited-state charge response inferred from the DFT calculations. (a) (i) Ligand-to-metal charge transfer with bond length change and tilting in the inorganic octahedron sublattice; (ii) change of local hydrogen−iodide bonding situation for the organic ligand; (iii) organic cation displacement or rotation in response to the new local environment. (b) Band diagram of MaPbI3 and illustration of excess thermalization energy during red and blue excitations and thermally activated ion movement in the perovskite lattice.

Table 1. Values of Diffusivity, Energy Barriers, Movement of Counterion (A-Cation) and the Vicinity Octahedron Distortion Factor for Iodide Vacancy Diffusion in Perovskite Solar Cell Materials material

constraint path barrier for I− movement (meV)

NEB barrier for iodide movement (meV)

iodide diffusivity using NEB (cm2 s−1)

coupled movement of counterion (Å)

rel octahedron distortion factor

MAPbI3 FAPbI3 CsPbI3

315 191 212

372 257 213

0.35 × 10−9 26.05 × 10−9 173.75 × 10−9

0.14 0.24 0.26

0.535 0.736 1.963

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energy, one can calculate the effective diffusivity (D) via the Einstein−Smoluchowski relation

FA) as well as to the Pb states. Here, partial charge transfer from iodide to FA+ is clearly seen in FAPbI3 under blue light, but only minor charge transfer contribution to MA+ in MAPbI3 is observed. These results then indicate that the FA+ cation perovskite have a partial charge transfer to the organic part when illuminated with light sources including UV/blue light and thus a possibility to change of the effective dipole of the A cation in this materials excited state in comparison to the MAPbI3 and CsPbI3 systems. As the local field also is dependent on possible ion displacement and its dependence on the type of A cation, the energy barriers of iodide displacement in the perovskites were calculated for the MAPbI3, FAPbI3, and CsPbI3 and, thus, as a function of monovalent cation dipolar strength. The energy barrier for iodide displacement was determined (Table 1) by means of two methods, constrained path and nudged elastic band (NEB). To evaluate the ionic movement, a transitionstate model was considered in which the excess energy of the excited electrons is thermalized and transferred to vibrations and thus could assist the ion movement in overcoming the energy barriers of the iodide vacancy diffusion (see Table 1 and section S9 of Supporting Information for details). Previously reported activation energies for ionic movement in MAPbI3 have ranged from 80 to 600 meV31,32 and thus in a rather wide range, where our data comply well with recent calculations that are in the middle of that range.33 Here, we find that constrained path calculations do not capture the expected decrease in the activation energy of the dipolar series in MA+, FA+, and Cs+. The different hydrogen-bonding situations during the iodide displacements thus have to be taken into account with lattice relaxations during the ion movement. Therefore, nudge elastic band (NEB) calculations with lattice relaxation during the iodide movement were performed, showing 372, 257, and 213 meV activation energies for iodide displacements in MAPbI3, FAPbI3, and CsPbI3, respectively. Recalling the red light excitations using 630 nm and the blue light excitation using 470 nm in the experiments in Figure 3, this translates into 1.97 and 2.64 eV, respectively. With a band gap of 1.55 eV for MaPbI3, this implies a limiting thermalization energy of a maximum of 420 and 1.09 meV if relaxed to the band edge with no deconstructive lattice vibrations. The available thermalization energy would be lower for other, higher band gap perovskite materials. The available thermalization energy under red light excitation is seen to barely be enough to thermally activate iodide movement, in contrast to the blue excitation light that carries enough excess thermal energy after relaxation to the band edge states to allow thermally activated ion movement as discussed below. The dipolar MA+ and FA+ cations that also form I−H bonds showed a coupled motion and rotation with the movement of the iodide presented as a displacement of the center of mass with 0.14 and 0.24 Å for MA+ and FA+. Interestingly, the monovalent cations Cs+ with no inherent dipole and less possibilities for directive hydrogen bonding show the largest mean displacement of 0.26 Å. Here, the magnitude of the dipole and the possibility to charge compensate with a rotation seem to effectively counter balance the center-of-mass movement of the cation. As a result, a large spherical cation with no inherent dipole shows a larger displacement in response to the iodide displacement (Table 1). Using a transition-state model for the thermal activation of the iodide movement via the jump rate Γ = ν0 exp(−ΔE/kT),34 where ν0 is the attempt frequency and the ΔE activation

D=

1 2 1 d Γ = d 2ν0 exp(−ΔE / kT ) 6 6

(1)

where d is the jump distance, k is Boltzmann’s constant, and T is the temperature. The attempt frequency is related to the I− Pb−I thermal motion, where one of the strongest intensity low wavenumber phonons can be found at 96 cm−1.35 The energy of this mode translates to 2.88 THz and is here considered as the major contributor to the thermally activated attempts. In addition, modes at lower wavenumbers can contribute but are lower in intensity and as the attempt frequencies constitutes the prefactor, the lower wavenumber modes as well as possible shifts of the dominating I−Pb−I phonon band with different cations in the cuboctahedra voids in between the octahedral inorganic framework are here neglected compared to the different contributions in the exponential term. The total jump distance over the NEB path is 4.614, 3.977 , and 4.248 Å for MAPbI3, FAPbI3, and CsPbI3, respectively. The effective iodide diffusivities for the different cations (Table 1) show a factor of 500 higher diffusivity for iodide in CsPbI3 in comparison to MAPbI3, with FAPbI3 in the middle, consistent with the relative strengths of the dipoles of the A cations. In addition, the local structural changes including the neighboring PbI6 octahedron distortions due to the ion movement seems to be correlated to the observed absorption changes and magnitude of Stark effect as well as the dipole of the A cation that would be further discussed (Table 1). The reported distortion factor (DF) in the Table 1 is here defined as the averaged and squared difference of Pb−I bond lengths (Pb···I) within the octahedron minus the average Pb−I bond length (Pb···Iavg) with DF =

1 6

i=1

∑ (Pb ··· Ii − Pb ··· Iavg)2 × 100 6

(2)

The reported DF numbers in Table 1 correspond to the vacancy neighboring octahedron showing a rather local effect. DF for far-away octahedrons were significantly lower.

DISCUSSION The spectral change caused by the Stark effect can be analyzed in terms of a frequency shift Δν of an optical transition due to an electric field E, which is related to the change in dipole moment between ground-state and excited-state Δμ and the change in polarizability Δα (Supporting Information, section S4) h × Δυ = −E ⃗μ ⃗ −

⎯ 1 ⃗ ⎯→ E ΔαE ⃗ 2

(3)

where h is Planck’s constant and the products are the scalar products if not explicitly denoted as cross-products. The resulting experimentally measured absorption change ΔA is a function of electric field E (Supporting Information, section S4): ⎯ 1 dA ⃗ ⎯→ 1 d 2A ⃗ 2 dA ⃗ (E μ ⃗ ) E μ⃗ − E ΔαE ⃗ + 2 dv 2 dυ 2 dv ⎯→ ⎯ 1 d 2A ⃗ 1 d3A ⃗ 3 (E μ ⃗ )(E ⃗ΔαE ⃗) − (E μ ⃗ ) + 2 2 dυ 6 dυ 3

ΔA × h = −

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(4)


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ACS Nano Analysis of the PIA spectra for MAPbI3 and FAPbI3 films shows that the dominating part of the observed phenomena can be attributed to the first-order Stark effect described by the product of the electric field and the dipole moment change (Supporting Information, section S3, Figures S4−5, and Table S1). Different mechanisms seem to take part in the low- and high-frequency parts of the PIA spectra and were seen to depend on the incoming photon energy. The red excitation gives features in the PIA response that get weaker with lower frequency, whereas the stable imaginary part of blue observed in the whole frequency range is indicative of a second phenomenon occurring with a delay relative to the incoming blue light. The frequency dependence of the Stark bleach observed in the PIA spectrum corresponds to a relaxation of the photoinduced charge response and the electric fields inside the material (eq 3). Although the lifetime of the charge carriers in the MAPbI3 film is below the millisecond time regime, the relaxation of photoinduced structural changes and photoinduced halide redistribution has been observed on considerably longer times in these systems and approaches the second time scale.16,36,37 Neglecting the initial charge response contributions, the frequency-dependent first-order Stark effect would contain information on the electric field dynamics in the film. The behavior of the measured frequency dependent PIA for the blue excitation follows the electric field relaxation (dielectric constant) previously reported inside a semiconductor with a dipolar response38 and further validating the ability of Stark spectroscopy as a probe for the local electric field inside a photoactive dielectric material. Excitation of hybrid lead halide perovskite films leads to the formation of free charge carriers on the picosecond time scale,18and the emission lifetime of the charge carriers is on the order of 100 ns for high-quality perovskite films.39 Reported trap-assisted recombination lifetimes are in the sub-microsecond and nanosecond regimes.40 The photoinduced Stark effect observed herein is on the microsecond to second time scale and must therefore be caused by long-lived states rather than free charge carriers or vacancy trap states. Photoinduced trap states, PbI6-octahedra distortions, and charge trapping at interfaces together with ion displacements are also possible candidates for the origin of the Stark effect and are discussed below in relation to the experimental results presented in this study. Trapping of charge carriers can result in long-lived states in the film. If such states are randomly distributed in the film, a second-order effect, which results in broadening of the absorption peak, could occur, but this effect was not observed experimentally here. A net electric field can can occur if electrons and holes are trapped at different interfaces of the crystals. For instance, electrons can trap at Pb-terminated surfaces and holes at iodide terminated interfaces. Under such conditions, an electric field will appear in the crystal in a specific orientation, and a first-order Stark effect can appear. However, this picture cannot explain the observed excitation light dependency of the PIA spectrum and is not in agreement with expected terminated surfaces of MAPbI3 crystals.41 Rotation of the MA cation in MAPbI3 has been identified as a process occurring in devices under working conditions and suggested to affect local metal−halogen bond length,26,42 the charge collection,9 the recombination,19 the VB and CB positions,19 and the hysteresis observed in current voltage measurements.43 The time constant of this rotation has been estimated theoretically19 for single MA in the lattice, and there

are experimental measurements evaluating the time constant in these materials,44,45 but until now there have been no direct measurements of how this depends on the available total energy or spectral distribution of the incoming light. The photoinduced charges and ion displacements play important roles in and are coupled to local orientation of the MA in the lattice.42 Considering the ensemble average of the dipole orientations, any change in the orientation of the MA dipoles would change the local electric fields and relaxation dynamics of Pb−I bond length in the material and result in a local Stark effect. In contrast, the observation of a dominating first-order Stark effect for the perovskite suggests a fixed angle between the photoinduced electric field and the dipole moment change.39 Consequently, the main contribution to the Stark effect does not belong to randomly oriented organic dipoles within the film, but those could instead contribute to the second-order Stark effect, which depends on the change in the local polarizability. Observation of the Stark effect in CsPbI3 perovskite with no A-cation dipole further justifies this view, and although the reorientation of the A-cation dipole indirectly contributes to the local field change, it is not the main origin of the corresponding local electric fields. The observed cation dipole dependency of the Stark effect (Figures 2 and 3) instead points toward a local movement that will shield local charge response processes in the inorganic framework of the iodide− lead octahedral (see also Table 1). Recently suggested photoinduced ion movement37 and consequent structural changes such as shortening of the Pb−I bond length26,36 present long-lived relaxation of local electric fields in the lattice that certainly can provoke the Stark effect within the films. Slightly different Pb−I bond lengths together with different orientation ordering of organic cation has previously been reported to give meaningful changes in CB and VB band edges19 and can subsequently change the dipole transition moment of the red optical transition. The latter can be another possible explanation for the observed absorption changes quantified by eq 4. Ionic movement has previously been suggested as one possible main effect responsible for the current voltage hysteresis observed in the lead perovskite solar cells31,32,46 and it has also been suggested to be responsible for variations in the low frequency part of perovskite dielectric constant (ε) .47 In a recent publication by De Angelis et al., it has been suggested that the ionic movement in the device is coupled with the rotation of MA dipoles that can provide local electric fields (E) through Pb−I bond length change.42 The shielding of the electric field can be formulated via the dielectric constant (ε), which is related to the polarizability (α) through the vacuum dielectric constant (ε0) by ε = ε0 + α

(5)

Therefore, ionic movement can not only create an E and affect the Stark effect through the field but also influence the secondorder Stark effect through the changes in the dielectric constant via Δα (eq 3). That is expected to be perturbed with a delay with respect to the incoming light and can thus participate in the imaginary part of the Stark effect absorption changes and dielectric constant at low frequencies (Figure 5d). The participation of ionic movement in the low frequency part of the Stark spectrum would be different for red and blue light excitations at the same photon number or the same intensity conditions (Figure 5d). As a result of the thermalization, heating the samples with more energetic phonons result in an 2830


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(Table 1), the charge compensation is also shifting to higher frequencies as shown in Figure 3b,c also showing that an increase in the thermally activated ion movement under blue light excitation help in the shielding over all of the measured frequencies. The processes at longer time scales also displayed a spectral dependence where blue light excitation carry significant energy for the thermalization process to overcome the energy barrier of the iodide migration in contrast to just a small effect for red light. The spectral dependence and the frequency dependency of these effects with the dependence on the inherent dipole moments of the cuboctahedrally coordinated cations together with first-principle calculations, thus given fundamental information on the mechanism of the local thermalization, as well as how this is coupled to the ion movement dynamics in OMHPs.

increased ionic transport within the sample that have consequences on the hysteresis behavior of the device. A full behavior of the hysteresis effect and its dependency to the wavelength of incoming photons is under investigation and will be published in a separate study. It should also be noted that the observed PIA signal represents the ensemble average of the absorption changes of the individual octahedrons within the film but where displaced ions and also locally under stoichiometric structures will experience a different local field. It is worth noting that apart from translational displacements, the photoinduced local electric fields in the octahedron can also be affected by the orientation of the organic molecule, which depends on the different preferential orientations of the perovskite grains and in turn on the illumination history and film preparation conditions.48 Based on the experimental results and the theoretical assessments, a plausible mechanism for the excitation energy dependent Stark effect can be formulated via thermalization effects from octahedral distortion and ion (A cation and iodide) displacements (Figure 7). Upon excitation or photoinduced ionic movement,37 (i) the local electric field change as the bond lengths change in the inorganic octahedron and also affects the neighboring octahedrons with an induced tilting during a ligand-to-metal charge transfer similar to the mechanism for the photoinduced piezoelectric effects reported in inorganic perovskites49 and here amplified with a thermally activated ion movement. This in turn leads to (ii), a change in the hydrogen bonding between the organic cation’s hydrogen atoms and the halide. This gives important information concerning the magnitude of the coupling strength between the organic cation in the cuboctahedral void and the lightinduced charge reorganization in the Pb−I octahedral. In the final step, (iii) the organic cation rotation is initiated by a change of the charge of a neighboring iodide in an I− → Pb2+ charge-transfer process and also the subsequent change in Pb−I bond length.42 In UV/blue light there is also a possibility that a charge transfer accompanies this effect to the A cation in the deep absorption (blue light) for FAPbI3 and CsPbI3 but only to a minor extent for the higher-laying bands for the A cation in MAPbI3. As the Stark effect is strongest for the Cs+ ion, without strong hydrogen bonding or inherent dipole, this reveals that the main part of the Stark effect is initiated from process i with displacement of the iodide due to a changed bond length and possible tilting of the octahedrals and at a longer time scales also iodide migration, changing the local electric field. The experimental data show that the cations with the stronger dipoles shield the change in local electric field more effectively. The dependence of the distortion factor and the local movement of the A cation during the ion movement (Table 1) from the density functional theory calculations are in line with the trends of the Stark effect and the dipole moment of A cation, which supports this idea. The weaker second-order stark effect is instead dependent on processes ii and iii where the strongest dipolar cation (MA) displays a flip and thus a reversal of the local polarizability indicative of MA rotation. On longer time scales, thermalization energy can induce ion movement and is interconnected to both steps i and iii and can be considered as mainly responsible for decreasing the observed Stark effect on longer time scales. This is also consistent with the experimental data for the red light in Figure 3 where there is no charge compensation response at high frequencies. With lower magnitude of the A-cation dipole, and higher diffusivity

CONCLUSIONS In summary, from the data presented one can attribute the observed Stark shift in the perovskite materials to a local electric field change upon excitation. An A cation with higher dipole in the perovskite material shows a more effective shielding of the photoinduced Stark effect where red light excitation (630 nm) shows a frequency dependence of the bleaching in contrast to the photoinduced Stark effect under blue light excitation (470 nm). Following the experimental trends and collaborating DFT calculations, an optically and thermally activated mechanism for changes in the local structure leading to less molecular interaction from the surrounding dipolar monovalent cation and, last, on a longer time scale, a local change in polarizability related to the ion displacements is suggested. In agreement with our proposed mechanism, the activation energies for iodide displacements were different for the different monovalent cations following the magnitude of the dipole moments where DFT calculations showed that the iodide movements were coupled to the movements of the monovalent cations (MA+, FA+, and Cs+) and the rotation of the MA and FA dipoles. The change in polarizability in domains can affect the observed Stark effect on the second order, while our results for the first- and secondorder Stark effects are in line with the picture related to the ion migration and the impacts on the I−V hysteresis behavior of the device. The experimental quantifications of the photoinduced Stark effect for different cations reveal how the local electric field depends on the molecular components where mechanisms of charge compensation as well as ion movement can be analyzed. Adding spectroscopically resolved responses and extended frequency ranges for the future bears promise of a Stark spectroscopy that can be used to analyze many aspects of dynamics of the local electric fields and their dependence on local order as well as thermally induced ion migration in hybrid perovskites and possibly also in photoactive semiconductor materials, in general. METHODS Film Preparation. Substrates of microscope glass or fluorinedoped tin oxide conducting glass were ultrasonically cleaned, and if needed, FTO was etched with zinc powder (and 2 M HCl aqueous solution) for 10 min with subsequent polishing and washing with DI water. Thin dense TiO2 films (blocking layers) were deposited (if needed) with either spray pyrolysis or spin coating. Different mesoscopic scaffold layers were deposited on the substrates, dried on a hot plate (80 °C, 10 min), sintered for 30 min (150 °C for Al2O3) and 450 °C for the others, and transferred to the drybox. 2831


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ACS Nano MAPbI3 was deposited on the substrates following either the twostep deposition method to obtain iodide MAPbI3 or a one-step deposition method for “mixed chloride” MAPbI3. Briefly, for the twostep deposition method, a 1 M solution of PbI2 in anhydrous N,Ndimethylformamide (DMF) was spin coated on the substrate. Substrates were heated to 70 °C on a hot plate placed in a drybox; after being cooling to room temperature, they were dipped in 10 mg/ mL MAI solution in anhydrous 2-propanol for 30 s. Finally, the substrates were washed with 2-propanol and dichloromethane and heated to 100 °C in a drybox. In summary, for mixed chloride MAPbI3, a 3:1 molar ratio of MAI/PbCl2 in DMF was prepared in glovebox with overnight stirring. The solution was spin coated on the substrates in a drybox at 2000 rpm, and the substrates were heated to 120 °C on a hot plate in a drybox. The same procedure was used for FAPbI3 material using the FAI instead of MAI in a mixed halide procedure. The device fabrication procedure for mixed perovskite films is described in the Supporting Information as well. Film Characterizations. The prepared films were characterized by XRD, UV−vis, PIA, and EA (details in the Supporting Information). The PIA spectroscopy setup utilized a modulated light produced by a blue LED (470 nm) or a red LED (630 nm) connected to a wave generator and coupled to a lock-in amplifier. A white lamp produces the probe light. A Si detector together with a monochromator records the wavelength dependent absorption of the sample through the computer-connected current amplifier. The PIA spectra of the bulk perovskite materials were recorded using a glass-covered perovskite with an Al2O3 scaffold using halide perovskites. To certify the validity of the observed frequency-dependent signal, the accuracy of measured data was extensively double-checked and corrected for any instrument response delay (Supporting Information, text and Figure S10). For EA, a semitransparent thin layer of silver (10 nm) was applied as top electrode. For the PIA measurements, the samples were excited using a blue LED (with a maximum wavelength of 470 nm) using a square wave on/off modulation with a frequency of 93 Hz where the change in light absorption was monitored using a lock-in detection system. For the EA measurement, the change in light absorption was measured for square-wave modulation of an electric field. DFT Calculations. The computations were performed on resources provided by SNIC through the Uppsala Multidisciplinary Center for Advanced Computational Science (UPPMAX) under Project Nos. sinc2014-3-71 and snic2015-6-65. The Quantum Espresso package was used for DFT calculations of the MAPbI3 and FA PbI3 structure using the General Gradient Approximation (GGA) with the Perdew−Burke−Ernzerhof (PBE) pseudopotential. Lead 5d10/6s2/6p2, nitrogen 2s2/2p3, iodide 5s2/5p5, and carbon 2s2/ 2p2 electrons were considered as valence electrons. The super cells consisted of 48 atoms corresponding to four unit cells with different orientations of MA and FA dipoles in order to obtain the most stable phases. Lattice parameters of the super cells were a = 8.71 Å and c = 12.46 Å in the tetragonal Bravai lattice, while the basic unit cell of MALIP inside the super cell is near cubic with a = 8.86 Å lattice constant. Full details of the DFT calculations and experimental methods can be found in the Supporting Information.

structural geometry of the MAPbI3 and FAPbI3, and normalized UV−vis absorption of the perovskite films (PDF)

AUTHOR INFORMATION Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. ORCID

Meysam Pazoki: 0000-0001-6776-5460 T. Jesper Jacobsson: 0000-0002-4317-2879 Jolla Kullgren: 0000-0003-3570-0050 Erik M. J. Johansson: 0000-0001-9358-8277 Tomas Edvinsson: 0000-0003-2759-7356 Author Contributions

M.P. initiated the work, carried out the experimental part except the device fabrication, main part of the DFT calculations, data analysis and wrote the draft. M.P., E.M.J.J, G.B, T.J.J., J.K., A.H., and T.E. discussed and analyzed the results in the context of possible mechanisms. M.P. and J.K. performed the energy-barrier calculations, T.J.J. provided the mixed halide perovskite devices. T.E. performed part of the DFT calculations, and G.B. and T.E. guided the work. Notes

The authors declare no competing financial interest.

ACKNOWLEDGMENTS M.P. thanks Michael Wang for his valuable help in constructing the amplifier circuit for highly resolved EA measurements, dedicates this paper to Dr Roghayeh Imani for her support and help, and appreciates Filippo De Angelis for his kind help regarding the DFT calculations. The Uppsala Multidisciplinary Center for Advanced Computational Science (UPPMAX) is acknowledged for providing computational resources under projects snic2014-3-71 and snic2015-6-65. We thank the Swedish Energy Agency, the Swedish Research Council, and the STandUp for Energy program for financial support. REFERENCES (1) Perovskite Fever. Nat. Mater. 2014, 13, 837−837 10.1038/ nmat4079. (2) Green, M. A.; Ho-Baillie, A.; Snaith, H. J. The Emergence of Perovskite Solar Cells. Nat. Photonics 2014, 8, 506−514. (3) Grätzel, M. The Light and Shade of Perovskite Solar Cells. Nat. Mater. 2014, 13, 838−842. (4) Martin, C. Perovskite Photovoltaics. Nat. Mater. 2012, 11, 1002− 1002. (5) McGehee, M. D. Perovskite Solar Cells: Continuing to Soar. Nat. Mater. 2014, 13, 845−846. (6) Kojima, A.; Teshima, K.; Shirai, Y.; Miyasaka, T. Organometal Halide Perovskites as Visible-Light Sensitizers for Photovoltaic Cells. J. Am. Chem. Soc. 2009, 131, 6050−6051. (7) Kim, H. S.; Lee, C.-R.; Im, J.-H.; Lee, K.-B.; Moehl, T.; Marchioro, A.; Moon, S. J.; Humphry-Baker, R.; Yum, J. H.; Moser, J. E.; Grätzel, M.; Park, N. Lead Iodide Perovskite Sensitized All-SolidState Submicron Thin Film Mesoscopic Solar Cell with Efficiency Exceeding 9%. Sci. Rep. 2012, 2, 591−597. (8) Nationl Cent. Photovoltaics, www.nrel.gov/pv/assets/images/ efficiency-chart.png (accessed Nov 24, 2016). (9) Frost, J. M.; Butler, K. T.; Brivio, F.; Hendon, C. H.; van Schilfgaarde, M.; Walsh, A. Atomistic Origins of High-Performance in Hybrid Halide Perovskite Solar Cells. Nano Lett. 2014, 14, 2584− 2590.

ASSOCIATED CONTENT S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.6b07916. Full description of theoretical and experimental methods, experiments with different scaffolds/interfaces, XRD graphs of the perovskite films, details of curve fitting for different harmonics of PIA spectra, derivation of Stark effect, AC and DC bias dependency of EA spectra for FAPbI3 and MAPbI3, light intensity dependence of photo induced Stark effect, comment about timed PIA and calibration of frequency-dependent PIA, hopping model for iodide movement in the perovskite lattice, additional EA spectra of the perovskite films, estimated PDOS and 2832


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