Deep Defect States in Wide-Band-Gap ABX3 Halide Perovskites

http://pubs.acs.org/journal/aelccp

Deep Defect States in Wide-Band-Gap ABX3 Halide Perovskites Igal Levine,*,† Omar Garcia Vera,‡ Michael Kulbak,† Davide-Raffaele Ceratti,† Carolin Rehermann,‡ José A. Márquez,§ Sergiu Levcenko,§ Thomas Unold,§ Gary Hodes,† Isaac Balberg,∥ David Cahen,† and Thomas Dittrich‡ †

Department of Materials & Interfaces, Weizmann Institute of Science, Rehovot 76100, Israel Institut für Si-Photovoltaik, Helmholtz-Zentrum Berlin für Materialien und Energie GmbH, Kekulé-Strasse 5, 12489 Berlin, Germany § Department Structure and Dynamics of Energy Materials, Helmholtz-Zentrum Berlin, Hahn-Meitner-Platz 1, 14109 Berlin, Germany ∥ The Racah Institute of Physics, The Hebrew University, Jerusalem 91904, Israel

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‡

S Supporting Information *

ABSTRACT: Lead bromide-based halide perovskites are of interest for wide-band-gap (>1.75 eV) absorbers for low-cost solar spectrum splitting to boost solar-to-electrical energy conversion efficiency/area by adding them to c-Si or Cu(In,Ga)Se2 PV cells and for photoelectrochemical solar fuel synthesis. Deep in-gap electronic states in PV absorbers serve as recombination centers and are detrimental for the cell’s photovoltaic performance, especially for the open-circuit voltage (Voc). We find four different deep defect states in polycrystalline layers of mixed-cation lead tribromide from highsensitivity modulated surface photovoltage (SPV) spectroscopy. Measurements were performed with different contact configurations, on complete solar cells and on samples before and after aging or stressing at 85 °C under illumination. Three of the four states, with energies of ∼0.63, 0.73, and 1.35 eV below the conduction band edge, are assigned to intrinsic defects, whereas defect states in the middle of the band gap could be associated with (uncontrolled) impurities. thermally stimulated current (TSC).10 Recently, it was shown that deep defects in lead bromide-based perovskites can be detected by using a highly sensitive photocurrent technique, showing transitions at 1.4 and 1.9 eV in CH3NH3PbBr3.11 Still, to date, experimental knowledge about deep defect states in lead bromide-based perovskites is extremely scarce, probably due to the very low density of such defects in those HaPs.2−4 Surface photovoltage (SPV) techniques probe the separation of photogenerated charge carriers in space. The SPV signal is deduced from the change in the surface potential of the sample upon illumination. The following conditions should be met for measuring an SPV signal in a semiconductor: (1) photogeneration of mobile charge carriers, i.e., the incident photons need to have an energy that is equal to or larger than an allowed optical transition from or to a delocalized state; (2)

D

eep defect states in the band gap of semiconductors can crucially limit the performance of devices. For example, nonradiative recombination via deep defect states can strongly limit the open-circuit voltage (Voc) in a solar cell. Therefore, detecting and characterizing deep defect states in the band gap of semiconductors is important for identifying limiting factors in semiconductor devices, and is a necessary step for their further development. In halide perovskites (HaPs), it has been difficult to observe deep in-gap states by direct optical absorption of sub-band-gap photons. While this is consistent with the low defect densities (1013−1016 cm−3),1−4 reported in HaP films, especially for bromide HaPs, lower than expected Voc values can be explained by defect recombination or trapping centers, which, however, remained elusive. Bromide HaPs have a relatively large Eg (about 2.3 eV5) and are of great interest as parent compounds of wide-band-gap (>1.75 eV) mixed-halide HaP absorbers in the high-energy cell of tandem solar cells with bottom cells based on c-Si or CIGS.6 In general, a large number of methods is available for characterizing deep defect states in wide-gap semiconductors, for example, photoluminescence (PL),7 photoconductivity (PC),8 deep-level transient spectroscopy (DLTS),9 and © XXXX American Chemical Society

Received: April 2, 2019 Accepted: April 17, 2019

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DOI: 10.1021/acsenergylett.9b00709 ACS Energy Lett. 2019, 4, 1150−1157

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Cite This: ACS Energy Lett. 2019, 4, 1150−1157

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ACS Energy Letters

Figure 1. (a) PC of the MixAPbBr3 layer as a function of light intensity, measured in a lateral device configuration (HaP layer between two Au pads with a 400 μm gap on glass, as shown in the inset of Figure 1a). The light source was a solid-state λ = 473 nm laser. (b) TRPL curves of a MixAPbBr3 layer on glass for different excitation intensities.

relatively long carrier diffusion lengths and relatively low free carrier densities, resulting in an improved overall photovoltaic efficiency compared to that of single-cation Br-based HaPs.18 To see if upon illumination the carrier recombination mechanism in our MixAPbBr3 layers is primarily trap-assisted or band-to-band, we started our study by measuring steadystate PC and time-resolved PL (TRPL) as a function of light intensity, a combination of methods that we used successfully earlier for MAPI1. The results gave a strong indication that the dominant recombination mechanism in our films is trapassisted. To find evidence for the existence of such deep defects that can serve as efficient recombination centers and enhance trapassisted recombination, we use here modulated SPS (for details, see the experimental part and ref 19). In our experiments, contact configurations were varied and some layers were aged in N2 and stressed at increased temperature under illumination. The contact configurations included TiO2 and/or PTAA (poly(triaryl)amine (used as electron- and holeextracting contacts, respectively, in a solar cell). Because stray light resulting from photons with energies above Eg must be avoided to exclude artifacts,20 an optical cutoff filter (OG580) to allow only photons with energies below 2.14 eV, i.e., well below EG of the films, to reach the sample was used in the measurements. Furthermore, a halogen lamp and a quartz prism monochromator (SPM2) were used for illumination to avoid discontinuities in the spectra due to changes of filters and/or gratings (for more details, see the SI). In this manner, several electronic transitions including states with energies at 0.63, 0.73, 1.15, and 1.35 eV (all values have an uncertainty of ∼±50 meV) below the conduction band (CB) edge of mixedcation lead tribromide were identified, corresponding to deep defects, which probably serve as nonradiative recombination centers. Figure 1a shows the PC of Au−MixAPbBr3−Au junctions as a function of light intensity, measured over a wide excitation intensity range of ∼5 orders of magnitude. The results indicate a power law exponent of γ = 0.79 ± 0.04. Because for nonradiative recombination in a single defect-level model the expected power law exponent should be 1 for nonradiative recombination and 0.5 for radiative bimolecular recombination,1 the observed γ values suggest that one or more

the photogenerated charge carriers need to be separated in space, hence generating a photovoltage; and (3) the resulting photovoltage needs to be high enough to be detectable within the given instrumental sensitivity. Several SPV techniques exist, and the most common ones are based on the vibrating Kelvin probe (KP)12,13 setup. In our study, we have employed the less common, highly sensitive, modulated SPV spectroscopy in a fixed capacitor arrangement,14−16 which allows detection of very low levels of photovoltage signals (down to values less than 1 μV, compared to about 1 mV in standard KP SPV setups). In modulated surface photovoltage spectroscopy (SPS), only those signals are detected that can follow the charge separation and recovery of the light-induced disturbance of the space charge within the modulation period, i.e., the quasi-static charge distribution can change in a sample during the measurement of a modulated SPV spectrum. By scanning the excitation wavelength, the resulting spectroscopy can probe sub-band-gap electronic transitions due to defect states. Thus, the method can be and was used to follow passivation of defects, as was demonstrated for CIGS17 and MAPI layers.14 One of the significant advantages of SPS, over more common photocurrent-based spectroscopies such as external quantum efficiency (EQE), is that it requires only one contact, allowing one to measure partial or complete device structures with top or bottom contact, or both, to separate the contributions to charge separation of the photoactive layer and its individual interfaces. Because the contacts need not be ohmic, photoactive layers can be investigated in various configurations of contacts with insulators, semiconductors, or metals. This allows, for example, to distinguish between defect states in the bulk or at interfaces. In addition, the direction of charge separation provides information about the nature of the sub-band-gap electronic transitions, allowing one to determine which of the bands is involved in these transitions. The purpose of the present study was to search for sub-band-gap transitions in HaPs using SPV spectroscopy (SPS) and determine the location of the corresponding defect states in terms of their (average) energies within the band gap. For the present deep-level investigation of Br-based HaPs, we focus on mixed-cation lead tribromide (FA0.85MA0.1Cs0.05)PbBr3 (denoted as “MixAPbBr3“) as the photoabsorber. In earlier studies, we found that layers of MixAPbBr3 have 1151


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Figure 2. (a, left) SPV amplitude (combined in- and out-of-phase signals) spectrum of a full FTO/d-TiO2/mp-TiO2/ MixAPbBr3/ PTAA/Au n−i−p MixAPbBr3 solar cell with (blue) and without a 580 nm long-pass filter (black), using a halogen lamp. A scheme of the experimental setup is shown in the inset (b, right) for an ITO/SnO2/mixed MAPI/spiro-OMeTAD/Au n−i−p mixed MAPI solar cell with (green) and without an 810 nm long-pass filter (black), using a Xe lamp. For both samples, the illumination direction is from the FTO side.

different experiments, suggesting that if deep traps are present in the studied mixed MAPI layers their concentration is 2.2

recombination centers participate in the recombination process under 1 sun equivalent illumination (generation rate of 1021− 1022 1/cm3 s1). Steady-state PL measurements (Figure S2 in the SI) indicate a relatively broad transition close to 2.3 eV, which we attribute to radiative band−band recombination, and some inhomogeneities of the absolute PL quantum yield (PLQY) on the scale of tens of micrometers observed, with a maximum PLQY yield of 0.5%, which is about a factor of 10 lower than the best PLQY observed for actual MAPI solar cells reported recently.21 Figure 1b shows the TRPL decay curves of the MixAPbBr3 sample as a function of excitation level. The TRPL shows a multiexponential decay with an initial fast component of τ1 < 10 ns and a slow component of τ2 < 90 ns. Such a multiexponential decay is often observed in TRPL measurements and can be associated with fast trapping into defects and/or surface recombination followed by the carrier decay arising from bulk recombination. Here the initial PL amplitude is found to scale as ϕ1.4, with the excitation intensity, ϕ, indicating that the carrier decay is limited by nonradiative recombination through defects,22 consistent with the PC results. Thus, both PC and TRPL results suggest that trap-assisted e−h recombination is significant over band-to-band recombination under 1 sun equivalent conditions and one or more deep defect levels act as active nonradiative recombination centers. Next, to experimentally determine if deep defect levels are indeed present and to actually find them in the MixAPbBr3 layers, we measured the SPV spectrum of the full MixAPbBr3 solar cell, ∼4% eff., and compared it with that of Cs0.05(FA0.83MA0.17)0.95Pb(I0.83Br0.17)3 cells (mixed MAPI, which gave >15% eff.; for the J−V curves, see Figure S1 in the SI). Figure 2 shows the results obtained with and without the appropriate long-pass optical filters (580 and 810 nm for the MixAPbBr3 and mixed MAPI, respectively). In contrast to studies that used DLTS and found 2−3 deep trap states in mixed MAPI,23−25 we could not detect any evidence for deep states, within the sensitivity of the setups that we used for the 1152


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Figure 4. Spectra of the in-phase (a) and out-of-phase (b) SPV signals for Au/PTAA/MixAPbBr3, FTO/TiO2/MixAPbBr3, FTO/ TiO2/MixAPbBr3/PTAA, and FTO/TiO2/MixAPbBr3/PTAA/Au (black lines and circles, red lines and triangles, squares, and stars, respectively). The shaded vertical lines mark the characteristic transitions, associated with changes in the slope of the spectrum. For the samples on TiO2 substrates (blue, red, and green), illumination was from the FTO side, while for the Au/PTAA/ MixAPbBr3 sample, illumination was from the HaP side.

Figure 3. Spectra of the in-phase (filled triangles) and out-of-phase (open triangles) SPV signals for a FTO/TiO2/MixAPbBr3/PTAA/ Au sample. The vertical black arrow marks the filter cutoff. The blue arrows and the blue dashed lines with the letters mark characteristic transitions. Illumination was from the FTO side.

eV) were within the noise. Analysis of the sub-band-gap features is shown in Figure 3 and is based on the changes in the SPV slopes (see also refs 13 and 29). The in- and out-of-phase SPV signals started to grow between 1.15 and 1.20 eV (transition A) to −0.5 and 3.5 μV at 1.3 eV. The signals then started to change back to the zero line at photon energies above 1.30 eV, giving evidence for the onset of a transition B, which, because of the opposite SPV sign, corresponds to charge separation in the direction opposite to that indicated by transition A. At photon energies between 1.5 and 1.55 eV and between 1.55 and 1.60 eV, the slopes of the spectra of the in- and out-of-phase signals changed again, which is ascribed to the onset of additional transitions (transitions C and D, respectively). To investigate if the above four sub-band-gap transitions, which can be related to defects, arise from the MixAPbBr3 buried interface with the electron transport layer (ETL) or the hole transport layer (HTL) or from the bare HaP surface, we compared the SPV spectra of all possible layer stack combinations, i.e., FTO/TiO2/MixAPbBr3/PTAA/Au, FTO/ TiO2/MixAPbBr3/PTAA, FTO/TiO2/MixAPbBr3, and Au/ PTAA/MixAPbBr3. Figure 4a,b shows the spectra of the inand out-of-phase signals, respectively, for the different contact configurations of as-prepared HaP layers. For the in-phase signals of the top-coated HaP samples, i.e., FTO/TiO2/ MixAPbBr3/PTAA and FTO/TiO2/ MixAPbBr3/PTAA/Au configurations, transitions A−C could not be detected, and only a slight increase to about 0.15 μV was observed around transition D. In contrast, the spectrum of the out-of-phase signals of the FTO/TiO2/MixAPbBr3/PTAA configuration showed signatures for the A and B transitions, and the signal disappeared around the onset of transition C on the one hand, while on the other hand, the spectrum of the out-of-phase signals of the FTO/TiO2/MixAPbBr3/PTAA/Au configuration gave clear evidence for all four transitions. For samples with an uncovered MixAPbBr3 surface, with the FTO/TiO2/MixAPbBr3 configuration, the onsets for transitions A and B were missing in the in-phase signal spectrum. For the Au/PTAA/MixAPbBr3 configuration, evidence was found for all four transitions in the spectra of both the in- and out-of-phase signals, whereas the onset of transition A was

shifted with a tail beginning at already 0.9−1.0 eV. Considering both in- and out-of-phase signals, for samples where the MixAPbBr3 surface is uncovered (black and red curves in Figure 4), the sub-band-gap features are more clearly observed than those for samples where that surface is coated with a top layer (blue and green curves in Figure 4). This suggests that for the latter the top layer partially passivates the bare surface of the MixAPbBr3 layers. We will now try to interpret the results shown in Figures 3 and 4. As we noted earlier, each change in slope of the SPS spectrum can be viewed as indication for an electronic transition that changes the position of the quasi-Fermi level with respect to the bands in the semiconducting component of the stack, the HaP. Such a change can come about by a change in population of energy levels around the Fermi level (for a very low doped material) and by a change in band bending due to a change in space charge as a result of a change in charge distribution as a result of illumination. Thus, it is clear that we need to look at the charge distribution near and at the uncovered surface and the interfaces of the sample. The different electronic transitions should be related to their corresponding defect energies within the band gap of the MixAPbBr3. Transition A: Because, based on its energy, transition A corresponds to a defect with an electronic energy level near midgap (transition energy of ∼1.15 eV), interactions with the CB and the valence band (VB) are equally probable. The separation to in- and out-of-phase signals can serve as an indicator for charge carrier injection into the ETL/HTL layer. In general, due to the fact that the buildup/decay of the photovoltage signal depends on how fast charge separation and recombination processes are with respect to incident light modulation, a strong out-of-phase signal suggests a retarded response in time, consistent with charge injection of a free carrier to a selective contact, followed by a slow trap-assisted recombination process. Hence, by measuring samples with only ETL or HTL layers, we can in principle determine if free holes or electrons were generated due to the sub-band-gap 1153


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ACS Energy Letters photon absorption and deduce if the process involved was optical excitation of an electron from the VB to a defect state, generating a free hole, or optical excitation of an electron from a defect state to the CB, generating a free electron. Transition B: There is a strong out-of-phase signal and no inphase signal for the samples containing the ETL layer, TiO2, (red, blue, and green curves). We interpret this observation as showing that the separation of the photogenerated charges is strongly retarded in time, as was argued above when discussing transition A. The delayed SPV response associated with transition B, observed for the samples with the TiO2, suggests that a detrapping process occurred and a free electron was generated in the process, followed by injection into the TiO2 layer. Thus, it can be concluded that the defect related to transition B is about 1.35 eV below the HaP conduction band minimum (CBM). Transitions C and D: Following the same logic as above, the delayed SPV response, resulting in a strong out-of-phase signal, corresponding to charge separation caused by transitions C and D, was clear for the samples with the HTL + Au layer combination, Au/PTAA/MixAPbBr3 and FTO/TiO2/MixAPbBr3/PTAA/Au configurations (black and green curves), compared to the samples without the Au/PTAA layer combination (red and blue curves). We interpret this difference by ascribing transitions C and D to electron trapping, resulting in the generation of free holes, followed by injection of the free holes into the PTAA. Therefore, transitions C and D can be related to defect states about 1.5 and 1.65 eV above the valence band maximum (VBM) of the HaP. Hence, we conclude that, although the contact configurations used here can influence differently the relaxation of charges, separated in space, leading to a change in the magnitude and/or in the time evolution of the SPV signals, they do not eliminate the existence of the defect states related to transitions A−D. This implies that the observed defect states are not induced by the HTL or the ETL but are rather an inherent property of the HaP bulk/surface itself. (We note that the onsets of transitions could be shifted by up to about 50 meV because in the measurements we change in excitation energy in steps of 50 meV. Also, the spectral resolution for the slits used, combined with the resolution of the quartz prism monochromator, amounted to about 30−40 meV. Also, some shift of characteristic features in SPV spectra can occur due to the superposition of signals with opposite sign and different strengths). On the basis of the above analysis of the SPV results, we present in Figure 5 a band diagram model that summarizes the suggested in-gap state distribution in the MixAPbBr3 layer as well as the observed electronic transitions that are denoted by Γ. The fact that defects C and D are closer to the CBM agrees well with the results of Musiienko et al.,30 who found evidence for defects C and D also in MAPbBr3 single crystals using photo-Hall measurements and showed that the signal associated with these defects corresponds to the generation of free holes, i.e., that these defects interact with the VB, as shown in Figure 5. The PC and TRPL results shown in Figure 1 suggest that nonradiative trap-assisted recombination is the dominant recombination mechanism in the MixAPbBr3 layers. Because the SPV-deduced in-gap states illustrated in Figure 5 are rather deep, it is reasonable to assume that at least one or more of these defects serve as nonradiative recombination centers that

Figure 5. Suggested defect-level diagram based on the experimental SPV results for the MixAPbBr3 layer, showing four different defect levels: B (occupied in dark conditions) and A, C, and D (empty in dark conditions). The gray arrows indicate electronic transitions to/from the defect level and the CB or VB, which correspond to the experimentally observed electronic transitions shown in Figure 4. Electrons (holes) are noted as blue (white) circles.

is/are sufficiently efficient to be responsible for the observed PC and TRPL trends. To check if different contact layers might passivate the above-mentioned defects, we fabricated p−i−n cells in the configuration glass/ITO/PTAA/MixAPbBr3/C60/BCP/Cu and measured their SPV spectra. The results (shown in Figure S3 in the SI) indicate that not only do the same sub-band-gap features discussed above appear in the p−i−n cells, their intensity is 1−2 orders of magnitude higher than that in the n− i−p cell (with TiO2 instead of C60 as ETL), indicating that the observed defects are not induced by the TiO2. The higher defect signal in the p−i−n cell might be due to either one or a combination of a lower passivation quality of the C60/BCP ETL layer or differences in the morphology of the HaP layer when deposited on a flat substrate with PTAA compared to the deposition on a mesoporous layer. Using highly sensitive EQE measurements of single-cation MAPbBr3-based solar cells, Sutter-Fella et al. found a sub-band-gap transition at ∼1.4 eV.11 This transition could either correspond to a deep defect level lying ∼0.9 eV below the CBM, which could be attributed to the transition corresponding to defect C in our case, or, alternatively it could correspond to defect B, lying ∼0.9 eV above the VBM. Thus, considering the possible error in determining the defect-level energy within the band gap here (that was based on the change in the SPV slope, where the error is roughly 50 meV, or on the Gaussian peak fit as done in ref 11), the results obtained by these different methods on different Pb−Br HaPs are in reasonable agreement. Because we studied mixed-cation Br-based HaP layers, in which the majority of the cations present are FA (∼85%), 1154


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Figure 6. Spectra of the in-phase (filled circles and triangles in (a,c)) and out-of-phase (open circles and triangles in (a,c)) signals and of the phase angle (b,d) for Au/PTAA/MixAPbBr3 (a,b) and FTO/TiO2/MixAPbBr3/PTAA/Au (c,d) configurations. This is before (triangles) and after (circles) aging for 30 days in N2 (a,b) and stressing at 85 °C under illumination equivalent to 1 sun for 1 h (c,d).

not lead to the formation of additional transitions. Figure 6b gives the spectrum of the phase angle (defined as the arctan of the ratio of the out-of- and in-phase signals). The spectrum of the phase angle contains signatures for the A, C, and D transitions, probably due to a different influence of the PTAA/ MixAPbBr3 interface and MixAPbBr3 surface on the relaxation of trapped charge carriers. Figure 6c compares the spectra of the in- and out-of-phase signals of the full n−i−p cell configuration of FTO/TiO2/ MixAPbBr3/PTAA/Au before and after stressing at 85 °C under illumination (equivalent to 1 sun) for 1 h in air. The SPV spectra remained practically unchanged in the range of transition A but increased strongly in the range of transitions B−D. The phase spectrum of the FTO/TiO2/MixAPbBr3/ PTAA/Au configuration after stressing was constant in the range of transitions B−D but was strongly shifted for the range of transition A. We thus conclude that transitions B−D are related to the evolution of intrinsic defects in the perovskite layer, which can be further photoenhanced using steady-state illumination (similarly to the phenomena observed by Motti et al. using pulsed illumination33), whereas transition A is probably caused due to the combination of an exposed surface of the HaP and possible reaction/penetration of foreign species from the glovebox into the perovskite layer and is rather insensitive to illumination. The fact that under both inert and ambient conditions the concentrations of the B−D defects increased suggests that formation of these defects is an intrinsic feature of the solar cell device, unrelated to contributions of oxygen and/or humidity. If correct, then the intrinsic instability found here originates from the MixAPbBr3 layer itself, as suggested by others.34−36 On the basis of our results, the mechanism leading to the rapid loss in PV performance upon aging under standard operating conditions of (unencapsulated) HaPs, in comparison with other PV absorbers,

while Sutter-Fella et al. studied single-cation Br-based HaP layers in which the only cation is MA, the observed similarities in the defect-level distributions suggests that the cation composition does not strongly influence formation of electrically active defects in the Br-based HaP layer, which is reasonable in view of what we know about the electronic structure of the Pb-based HaPs.5 Thus, although the defect densities may vary significantly between Pb−Br-based HaP layers that are prepared in different ways, leading to different doping concentrations and free carrier densities,31 our conclusions regarding the defect-level distribution, shown in Figure 5, may well apply to all Br-based HaP layers (a hypothesis that requires further experimental verification). Recent theoretical work by Cohen et al.32 on CsPbBr3 predicts that Br vacancies behave as unusual, “dynamic” defects within the band gap, with time-dependent energies from shallow (near the CBM) to ∼1 eV below the CBM. Because experimentally we are measuring time-averaged electronic transitions, corresponding to defect energies of ∼1.3−0.6 below the CBM, we hypothesize, in line with Cohen et al.’s result, that the observed deep defects seen here correspond to physically a single “dynamic” Br-related defect in the MixAPbBr3 lattice. After revealing the energetic distribution of the in-gap states in the MixAPbBr3 layers, we set out to study the effect of aging on the distribution of these defects. We performed SPV measurements on a full cell and a half cell after two different aging processes, and the results are presented in Figure 6. Figure 6a compares the spectra of the in- and out-of-phase signals of a Au/PTAA/MixAPbBr3 configuration before and after aging in an N2-filled glovebox for 30 days. The signals related to all four transitions increased after aging, and the tails at photon energies below transition A became more pronounced after aging. We conclude then that aging did 1155


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ACKNOWLEDGMENTS I.L. thanks the Minerva Stiftung for a short-term research grant. C.R. thanks the HyPerCells (Hybrid Perovskite Solar Cells) joint Graduate School and the Helmholtz International Research School “Hybrid Integrated Systems for Conversion of Solar Energy” (HI-SCORE), an initiative cofunded by the “Initiative and Networking Fund of the Helmholtz Association”. In Rehovot, this work was supported by a grant from the Ullmann family foundation. We thank the reviewers for their constructive comments.

can then be related to the formation of in-gap states within the HaP absorber. The effect of defect-related recombination on the performance of solar cells can be estimated by calculating the bulk quasi-Fermi level splitting, μ, of the absorber layer from its luminescence properties, using μ = μSQ + kT ln(PLQY), where μSQ is the thermodynamically ideal quasi-Fermi level splitting for a given band gap in the Shockley−Queisser limit. The second term denotes the nonradiative loss, which can be calculated to be >140 meV for the measured maximum PLQY of 0.5%. We ascribe this loss to nonradiative recombination via the deep defects identified in the SPV study and discussed in detail above. To estimate the total Voc loss occurring in the solar cells, we compare the estimated quasi-Fermi level splitting μ ≈ 1.89 eV (μ SQ ≈ 2.03 for Eg = 2.3 eV) to the measured Voc of the best Br-based perovskite solar cells, Voc ≈ 1.65 V,37 which clearly indicates that significant additional losses are introduced in the process of building the device, e.g., by interface recombination at the perovskite/charge transport layers and/or unfavorable energy band alignment at the HaP/ charge transport layer interfaces. To summarize, by using modulated SPV, we were able to resolve the electronic defect structure of (FA0.85MA0.10Cs0.05)PbBr3 layers. Transitions related to extrinsic (A) and intrinsic (B−D) defects were observed in SPV spectra at photon energies of about 0.9−1.15, 1.30−1.35, 1.50−1.55, and 1.65− 1.70 eV, respectively. Because these four defect levels are relatively deep (>0.6 eV from the band edges), they serve as efficient recombination centers, enhancing nonradiative, trapassisted recombination under steady-state illumination conditions, as confirmed by the PC and TRPL measurements. Measurements of the PLQY indicate that these recombination losses contribute ∼140 meV to the total open-circuit voltage loss for (FA0.85MA0.10Cs0.05)PbBr3 solar cells. We further postulate that electronic passivation of these deep defect states, assuming they originate from the surface or the grain boundaries of the HaP, by means of surface treatments and/or additives to the HaP precursor solution,38,39 might be a promising route to minimize these voltage losses.

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REFERENCES

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsenergylett.9b00709. Detailed description of the sample preparation and SPV and TRPL experimental details (PDF)

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Letter

AUTHOR INFORMATION

ORCID

Igal Levine: 0000-0003-0406-2922 José A. Márquez: 0000-0002-8173-2566 Thomas Unold: 0000-0002-5750-0693 Gary Hodes: 0000-0001-7798-195X Isaac Balberg: 0000-0002-0664-8735 David Cahen: 0000-0001-8118-5446 Thomas Dittrich: 0000-0002-2698-9481 Notes

The authors declare no competing financial interest. 1156


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