Letter www.acsami.org
High-Work-Function Molybdenum Oxide Hole Extraction Contacts in Hybrid Organic−Inorganic Perovskite Solar Cells Philip Schulz,*,†,‡ Jan O. Tiepelt,† Jeffrey A. Christians,‡ Igal Levine,§ Eran Edri,§,# Erin M. Sanehira,‡,∥ Gary Hodes,§ David Cahen,§ and Antoine Kahn*,† †
Department of Electrical Engineering, Princeton University, Princeton, New Jersey 08544, United States National Renewable Energy Laboratory, Golden, Colorado 80401, United States § Department of Materials and Interfaces, Weizmann Institute of Science, Rehovot 76100, Israel ∥ Department of Electrical Engineering, University of Washington, Seattle, Washington 98195, United States ‡
S Supporting Information *
ABSTRACT: We investigate the effect of high work function contacts in halide perovskite absorber-based photovoltaic devices. Photoemission spectroscopy measurements reveal that band bending is induced in the absorber by the deposition of the high work function molybdenum trioxide (MoO3). We find that direct contact between MoO3 and the perovskite leads to a chemical reaction, which diminishes device functionality. Introducing an ultrathin spiro-MeOTAD buffer layer prevents the reaction, yet the altered evolution of the energy levels in the methylammonium lead iodide (MAPbI3) layer at the interface still negatively impacts device performance. KEYWORDS: electronic structures/processes/mechanisms, photoemission spectroscopy, hybrid materials, photovoltaic devices, band offsets, charge carrier transport
T
oxides to enable charge transfer from the perovskite layer into the HTL. In the past few years, the high work function molybdenum trioxide (MoO3) has been successfully used as hole-injection and -extraction material in organic electronics, including solar cells in which it also creates a large built-in voltage.10−13 In these devices, direct transfer of electrons from the very low lying conduction band of the n-type MoO3 (electron affinity, EA, of about 6.7 eV) into the highest occupied molecular orbital (HOMO) of the donor molecules of the adjacent active layer is responsible for the efficient hole extraction process.10 Moreover, MoO3 interlayers are of considerable interest for the realization of recombination layers in tandem cells based on perovskite subcells.14 To date, MoO3 has been considered predominantly as a hole contact on MAPbI3 solar cells on top of the organic HTL, to increase the work function of standard metals like Ag or improve device stability.15,16 Although it was found to improve carrier extraction from the HTL, the considerable spatial separation from the active perovskite layer may have diminished the presumed benefits of the high work function of the oxide, in particular with respect to the built-in voltage and field in the active layer. Liu et al. recently investigated
he outstanding performance of photovoltaic devices based on hybrid organic−inorganic perovskite absorbers has spurred a flurry of activity aimed at unravelling the fundamental electronic, chemical and physical properties of this class of materials.1−3 In particular, issues related to charge separation and collection, and to the electronic coupling to adjacent charge transport layers, are being extensively investigated. We recently demonstrated that the electronic structure of organic materials used as carrier extraction layers directly affects the device characteristics of both conventional and inverted perovskite solar cells.4,5 Specifically, a mismatch between the transport levels of the carrier extraction layer and the perovskite active layer would limit the device open circuit voltage (Voc) or introduce an unwanted carrier extraction barrier, and could as well degrade the carrier selectivity by allowing carrier crossover and recombination.4 Such instances were demonstrated with both methylammonium lead iodide (MAPbI3) and bromide (MAPbBr3). In particular, the latter exhibits a large band gap of 2.3 eV and high ionization energy (IE) of ∼5.9 eV and thus requires an organic hole-transport layer (HTL) with commensurable electronic properties.4,6,7 Even in the case of MAPbI3, which exhibits a band gap of ∼1.6 eV and an IE of ∼5.4 eV, significant improvements in cell performance were achieved using high ionization energy organic HTLs.8,9 The present study focuses on an alternative for hole extraction interlayers, based on high work function metal © 2016 American Chemical Society
Received: August 30, 2016 Accepted: November 7, 2016 Published: November 8, 2016 31491
DOI: 10.1021/acsami.6b10898 ACS Appl. Mater. Interfaces 2016, 8, 31491−31499
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ACS Applied Materials & Interfaces
Figure 1. (a−c) Ultraviolet (UPS) and (d−f) X-ray (XPS) photoemission spectra of MAPbI3 on TiO2 with MoO3 incrementally evaporated on top. (a) Secondary electron cutoff and work function (WF), (b) valence band region and (c) close-up of valence band region in semilog plot, emphasizing the molybdenum oxide interface state IS (see also Figure 3a); (d) Pb 4f, (e) I 3d, and (f) Mo 3d core level spectra. The inset in f depicts a magnification and fit of the Mo 3d core level spectrum of the 1.5 Å MoO3 film, showing the Mo4+ component (shaded, blue).
sandwiched between the perovskite and the MoO3 layer, which inhibits the chemical reaction, and we investigate the influence on the photovoltaic activity of the device. Photoemission spectra of the MAPbI3 layer, bare and with incremental thicknesses of top-deposited MoO3, are shown in Figure 1. The work function of the film, determined from the secondary electron cutoff in the UPS spectra in Figure 1a, increases monotonically with oxide thickness. With 100 Å MoO3, the vacuum level is found to be at 6.8 eV above the Fermi level (EF), in line with previous reports on the work function of evaporated MoO3 thin films.12,13 The valence band region of the film clearly shows the formation of a MoO3 film with a valence band maximum (VBM, dotted line in Figure 1b) at 2.8 eV below EF (the oxide ionization energy, IE, is 2.8 + 6.8 = 9.6 eV, again in accord with the previous reports for comparable MoO3 thin films.). The 3 eV band gap oxide is ndoped, with its conduction band minimum (CBM) only 0.2 eV above EF. The emergence of surface/interface states (IS′) in the oxide gap at energy close to EF, is clearly visible on the 20 Å spectrum (Figure 1b), and characteristic of substoichiometric evaporated MoO3.12,19 Another incidence of surface/interface states is more clearly seen in a comparison of the semilogarithmic plots of the valence band region of the MAPbI3 film in pristine condition and covered with a submonolayer of MoO3 with nominal thickness of 1.5 Å where an interface state density with onset at 0.2 eV below EF is observed (IS in Figure 1c). Yet, the exact nature of the interface state IS remains unresolved from the UPS data. Note that the determination of
chemistry and electronic structure at the MoO3/MAPbI3 interface obtained by direct evaporation of the oxide on the perovskite.17 In the present work, we show that the formation of this junction can lead to even more pronounced band bending in the MAPbI3 film and to a reduction of the MoO3 to MoO2 at the very interface. We outline, and show initial data for, an approach that can mitigate the impact of the oxide/ perovskite chemical reaction and improve the chemical integrity of the interface, while preserving the effects of the very high work function overlayer on the built-in field. This allows us to identify the impact of the chemically altered interface layer and of the band bending in the perovskite active layer, on device functionality. The MAPbI3 film is deposited from solution in a one-step process on top of a thin hole-blocking TiO2 layer, which in turn has been grown on a fluorine-doped tin oxide (FTO)-covered glass substrate by spin-casting from solution.18 We find the performance of devices comprising a direct contact between perovskite and MoO3 to be degraded, even compared to samples without hole transport layer. The result is attributed to • the chemical reaction in the MoO3 film at the very interface, and • an inadequate electronic level alignment between the two layers. To discriminate between these issues, we introduce an ultrathin buffer layer (few nm) of 2,2′,7,7′-tetrakis(N,N-di-pmethoxyphenyl-amine)-9,9′-spirobi-fluorene (spiro-MeOTAD) 31492
DOI: 10.1021/acsami.6b10898 ACS Appl. Mater. Interfaces 2016, 8, 31491−31499
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Figure 2. Ultraviolet (a−c, UPS) and X-ray (d−f, XPS) photoemission spectra of MAPbI3 on TiO2 with a 35 Å thick layer of spiro-MeOTAD with MoO3 incrementally evaporated on top. (a) Secondary electron cutoff and work function (WF), (b) valence band region, and (c) close-up of valence band region in semilog plot. (d) Pb 4f, (e) I 3d, and (f) Mo 3d core level spectra. The inset depicts a magnification and fit of the Mo 3d core level spectra of the 1.5 Å MoO3 film. Note the absence of the Mo4+ component seen in Figure 1f.
In order to suppress this interface chemical reaction, a 35 Å thick buffer layer of the hole transport material spiro-MeOTAD was evaporated on top of the perovskite film, as a means to protect the absorber from the evaporated MoO3. The thickness of the buffer layer was thick enough to prevent direct physical contact between the perovskite and the oxide, yet thin enough to maintain good electronic transport through the layer stack. The photoemission data presented in Figure 2 enable a direct assessment of the electronic energy level alignment for this system. The secondary electron cutoff (Figure 2a) shows that the work function first decreases by about 0.2 eV with the evaporation of spiro-MeOTAD, in line with our previous findings,4 then follows the same increase as in Figure 1a upon evaporation of MoO3, reaching a maximum of 6.8 eV with a MoO3 thickness of 100 Å. The valence band region (Figure 2b) shows the onset of the spiro-MeOTAD HOMO level at 1.0 eV below EF, as reported before.4 After covering the layer with 1.5 Å of MoO3, the HOMO level of the organic buffer layer is roughly 0.4 eV closer to EF than without the oxide overlayer, indicating that the spiro-MeOTAD layer is p-doped by charge transfer to the MoO3 (Figure 2c). With increasing oxide coverage, the electronic structure in the valence band region evolves with the appearance of the unperturbed Mo 4d level. Beyond 20 Å, the spectrum is indistinguishable from that of the thick MoO3 film on top of MAPbI3 without the spiroMeOTAD buffer layer. The core level analysis (Figure 2d,e) reveals the same shift of the Pb 4f and I 3d levels observed for the system without spiro-
the valence band onset of the perovskite layer is achieved here by fitting the UPS valence band spectrum to a spectrum simulated from density functional theory calculations. A detailed description of this process has been provided by Endres et al.20 Further insight in the energy level alignment at the perovskite/oxide interface is gained from XPS data. Upon deposition of increasing amounts of MoO3, both the iodine and lead core levels (I 3d5/2 and Pb 4f7/2) of the perovskite layer shift to lower binding energies (Figure 1d, e), by a total of 0.5 eV after deposition of 20 Å MoO3. This rigid shift indicates a MoO3-induced band bending in the perovskite film, consistent with the deposition of the high work function oxide. The spin− orbit split components of the Mo 3d core level do not exhibit any significant shift (Figure 1f). However, the decomposition of the core level spectrum into components corresponding to Mo6+, Mo5+, and Mo4+ oxidation states, with Mo 3d5/2 binding energies at 233.2 eV, 232.0 and 230.1 eV respectively, reveals a distinct component of substoichiometric molybdenum oxide. While the presence of Mo6+ and Mo5+ has been seen repeatedly for evaporated molybdenum oxide films,12,20 the appearance of the Mo4+ component (inset in Figure 1f), normally associated with MoO2, is surprising. This Mo4+ contribution is only apparent at very low film thickness (1.5 and 5 Å in Figure 1f), which indicates that the reduction to molybdenum dioxide is confined to the MAPbI3/oxide interface and likely involves the perovskite layer. 31493
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ACS Applied Materials & Interfaces MeOTAD. However, in contrast to the previous case, the Mo 3d core level does not exhibit any significant Mo4+ component (inset of Figure 2f), indicating that the spiro-MeOTAD layer fulfills its role of a buffer layer and prevents the chemical reaction at the perovskite/molybdenum oxide interface, and thus the formation of MoO2. The region of the valence band close to EF comprises a complex spectral superposition of various components. In the system without the spiro-MeOTAD interlayer, a DOS appears at 0.2 eV below the Fermi level with the initial deposition of one monolayer of molybdenum oxide (Figure 1c). MoO2 has been reported in the literature to be metallic with EF cutting through the Mo 4d bands.21 This picture changes with the application of the spiro-MeOTAD buffer layer on MAPbI3. The HOMO onset of the organic HTM is ∼1.0 eV below EF prior to MoO3 deposition (Figure 2c). The subsequently evaporated oxide layer p-dopes the spiro-MeOTAD, and the HTM HOMO onset shifts to 0.6 eV below EF, with a tail of states reaching almost up to the Fermi level for a 1.5 Å MoO3 overlayer thickness. Note that the same coverage of 1.5 Å of MoO3 on bare MAPbI3 led to the formation of an interface state (IS) originating from reduced MoO2 (Figure 1c). Eventually, with 20 Å of deposited molybdenum oxide, the electronic structure of the MoO3 surface, as it is also found for the system without spiro-MeOTAD interlayer, is recovered. The combined results from the PES measurements are summarized in the energy level diagrams of the interface with and without the spiro-MeOTAD buffer layer (Figure 3). The shift of the Pb 4f and I 3d core levels is the same without (Figures 1d, e) and with (Figures 2d, e) the buffer layer, and indicates a 0.5 eV band bending induced in the perovskite layer by the deposition of MoO3. This 0.5 eV shift of the perovskite VBM (and core levels) toward EF is neither enhanced nor impeded by the reduced MoO2 found in the spiro-MeOTADfree device, and seems therefore to originate solely from the electrostatic potential imposed by the high work function of the molybdenum oxide layer. Information on how far the band bending reaches into the perovskite layer is difficult to obtain, but it is assumed here that the film remains n-type, i.e., EF close to the conduction band minimum, at and near the interface with the bottom TiO2 substrate. Previous studies have demonstrated that the Fermi level position in hybrid organic lead halide perovskites is strongly influenced by the doping character and work function of the (oxide) substrate,5,22 i.e., perovskite films deposited on TiO2 have n-type character while those deposited on NiO have a slight p-type character. This adaptive change of the Fermi level position as a function of the electronic properties of the substrate also implies a very low density of deep electronic levels in the perovskite band gap and low doping concentrations.23,24 We therefore assume that the MAPbI3 film sandwiched between TiO2 and MoO3 gradually changes from n-type to intrinsic, as depicted in Figure 3. The electronic structure of these interfaces is expected to affect solar cell functionality and performance. In particular, band bending and a corresponding internal field could alter carrier transport through the perovskite film and change recombination rates. Furthermore, the role of the chemistry at the perovskite/MoO3 interface on charge carrier collection needs to be characterized in detail. To test the respective impact of MoO3 and spiro-MeOTAD/MoO3 interlayers on device characteristics, we use copper(II) phthalocyanine (CuPc) as our choice for a hole transport layer, since it can be evaporated on top of the hole extraction layer without
Figure 3. Energy level diagram of a MAPbI3 perovskite layer on top of a compact TiO2 film. (a) Subsequently, deposited MoO3 in immediate contact is reduced to MoO2, which causes the emergence of an interface state (IS), whereas (b) a spiro-MeOTAD interlayer prevents this chemical reaction. In both cases, a 0.5 eV band bending is induced in the perovskite layer making the film more intrinsic at the interface. Moreover, the formation of an interface state (IS′) in the slightly substoichiometric MoO3 layer becomes apparent with increasing oxide film thickness. For each layer, the Fermi level position (EF, dotted line), work function (in red), and electron affinity and ionization energy (in black) are indicated in eV.
redissolving constituents of that junction, such as the undoped spiro-MeOTAD (u-spiro-MeOTAD) buffer layer. Previous studies showed that CuPc is a reasonably good HTM in perovskite solar cells, even though expected efficiencies are lower than for devices with a standard doped spiro-MeOTAD (d-spiro-MeOTAD) HTL.25 Furthermore, the energy alignment of the CuPc HOMO level with the MAPbI3 VBM has been found to be similar to that of spiro-MeOTAD.26 We therefore produced the following three solar cell configurations: glass/FTO/TiO 2 /MAPbI 3 /MoO 3 /CuPc/Au, glass/FTO/ TiO2/MAPbI3/CuPc/Au, and glass/FTO/TiO2/MAPbI3/uspiro-MeOTAD/MoO3/CuPc/Au, respectively. The undoped spiro-MeOTAD interlayer is spin-coated from a dilute solution to yield a film thickness ≤5 nm. According to the photoemission experiments, any MoO3 thickness above 20 Å leads to a film with a work function larger than 5.5 eV. Thus, a layer thickness of 35 Å is used here to induce maximum band bending without significantly increasing the series resistance (cf. Figure 3 and Figure S3), and to understand the effect of such a high WF layer has on hole extraction. We note that in organic photovoltaic devices, with some 25% lower current at 31494
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Figure 4. (a) J−V curves and (b) normalized EQE measurements for MAPbI3 perovskite solar cells with and without MoO3 and u-spiro-MeOTAD/ MoO3 interlayer between active layer and CuPc HTL/Au electrode (normalization factors: 1.3, 3.6, 7.0). In the J−V curves, solid lines represent reverse and dashed lines represent forward scans.
maximum power,1,27 the resistance of a 50 Å thick layer of MoO3 alone is negligible.13 In addition, another device geometry was tested where the CuPc HTL was substituted by a spin-coated layer of doped spiro-MeOTAD to validate the performance of the MAPbI3 films (PCE > 15%, see Figure S1). HTL-free devices with Au contacts showed greatly diminished device characteristics (Figure S1). We attribute this result in part to the perovskite morphology, which might allow pathways for shorting between the Au contact and the substrate. In conventional cells, the organic HTL takes the additional role of smoothing the rough perovskite surface, thereby minimizing the amount of pinholes and shorts from the Au to the TiO2/ FTO in the device. However, using the evaporated CuPc HTL we arrive at a homogeneous pinhole free layer stack as seen by cross-sectional scanning electron microscopy images (Figure S2) The J−V curves of the three types of solar cells with CuPc HTL are shown in Figure 4a, and the device characteristics are given in Table 1. The devices with the plain CuPc HTL without
collection at the interface between perovskite active layer and the CuPc HTL and the overall hole transport properties in the CuPc.26 However, the glass/FTO/TiO2/MAPbI3/MoO3/CuPc/Au devices with bare MoO3 interlayer show significantly worse photovoltaic activity, a situation likely due to both the chemical and the electronic interaction between MoO3 and the perovskite. As the oxide is brought into direct contact with MAPbI3, it is reduced to MoO2 at the interface as directly observed in the photoemission experiment (Figure 1f). The MoO2 layer is metallic with a significant DOS at EF and has been shown to be an efficient hole-injection material at contacts with organic semiconductors.21,27,28 In the present case, however, the chemistry taking place at the MoO3/perovskite junction and leading to the reduction of Mo6+ to Mo4+ species is likely to affect the perovskite as well. Chemically one can speculate that Pb2+ → Pb4+ oxidation and/or 2I− → I2 oxidation accompany the Mo6+ → Mo4+ reduction, both of which can create defect states at the interface that could serve as potential recombination centers, resulting in a decrease of device performance. However, the open circuit voltage (Voc) shows only a small decrease compared to the device without MoO3 interlayer, which would usually indicate that recombination at the interface does not seem to be significantly increased. In contrast, we observe a drop in Jsc and FF, likely due to an increase in series resistance and a decrease in collection efficiency at low net electric field across the device, i.e., close to Voc. Also the electric field and Fermi level position in the CuPc HTL are affected by the presence of the high work function MoO3. Photoemission spectroscopy measurements reveal that the CuPc becomes heavily p-doped at the interface with the MoO3 film and less p-doped further away from the interface, i.e., toward the top electrode (see Figure S3), as expected from similar studies on MoO3/CuPc interfaces.29 This progressive shift of the HOMO level in the CuPc HTL is generally detrimental for hole transport through the CuPc film
Table 1. Device Characteristics of Perovskite Solar Cells with Various Hole Extraction Layers HTL/top electrode
Voc (V)
Jsc (mA cm−2)
PCE (%)
FF (%)
CuPc/Au MoO3/CuPc/Au u-spiro-MeOTAD/MoO3/CuPc/Au
0.92 0.86 0.81
19.5 13.8 13.7
10.4 4.2 5.0
58 35 45
additional interlayer show moderately high performance characteristics. While the current density comes close to the 21.9 mA/cm2 short circuit current density (Jsc) of the reference device (d-spiro-MeOTAD HTL, see Table S1), a ∼0.13 V loss in open circuit voltage (Voc) and a substantial loss (from 0.73 to 0.58) in fill factor (FF) are observed, which leads to a drop of over 5% in power conversion efficiency, to 10.4%. In particular, the loss in FF could be explained by nonoptimized hole 31495
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Figure 5. (a) Photoluminescence decays and (b) spectra of MAPbI3 thin films on TiO2/FTO/glass substrates for various hole transport layer configurations with and without MoO3 and u-spiro-MeOTAD/MoO3 interlayer between the active layer and the CuPc HTM.
the perovskite layer by the high work function of the adjacent MoO3 charge extraction layer. In this picture the collection at the MAPbI3/u-spiro-MeOTAD/MoO3/CuPc junction would be particularly reduced. Although devices with only the thin uspiro-MeOTAD layer but without MoO3 layer between the MAPbI3 and CuPc exhibit the same performance parameters as devices with a direct MAPbI3/CuPc junction (see Figure S5), the u-spiroMeOTAD/MoO3/CuPc system is detrimental for unimpeded hole transfer. Essentially the u-spiro-MeOTAD and CuPc become p-doped by the n-type MoO3 layer (Figure S4) in the interface region, which leads to band bending in the CuPc film. In consequence, hole transport toward the top electrode occurs against the field in the CuPc layer and is impeded. Photoluminescence spectroscopy measurements were performed in order to further investigate the collection of charges that are photogenerated in the perovskite absorber. Here, the decay time of the characteristic photoluminescence emission (Figure 5) observed at 770 nm wavelength can serve as a proxy to track the fate of photoexcited carriers in the device structures with varied HTL configuration. Only small differences in carrier quenching can be observed between the MAPbI3 with an immediate CuPc, MoO3/CuPc, or u-spiro-MeOTAD/MoO3/ CuPc HTL interface. A double exponential fit to the decay curves reveals initial fast recombination with time constants τ1 of 3 ns for all three films followed by a second order decay time τ2 which amounts to 19, 12, and 17 ns, respectively. This indicates that interface recombination is slightly increased for the samples with an immediate MAPbI3/MoO2 interface, which can in part be mitigated by introducing the thin u-spiroMeOTAD interlayer. The emission spectra depicted in Figure 5b of the three different samples look virtually identical, which leads to the conclusion that this increased interface recombination does not open new radiative decay channels. In summary, the results from these time-resolved photoluminescence measurements indicate that the observed changes in the interfacial electronic structure do not strongly contribute to recombination and carrier quenching in the MAPbI3 bulk. This is in contrast to reports of alternative charge carrier extraction interlayers (e.g., carbon nanotubes) in which changes of the interfacial electronic structure can strongly affect recombination in the perovskite active layer,32,33 In conclusion, our study demonstrates that the evaporation of MoO3 on top of a TiO2/MAPbI3 film induces band bending in the perovskite film which, as a result, becomes nearly
from the side of the MAPbI3/MoO3 junction in direction of the metal top electrode (see Figure S4) but could change with the application of the Au top electrode. This result underlines why no reasonably well-performing perovskite solar cell with direct molybdenum oxide contact has been reported so far. However, the origin of the changes in the device characteristics remains ambiguous as the effects of the interface chemistry and band bending in the perovskite active layer are conflated for this device geometry. To that point, the introduction of the undoped spiroMeOTAD interlayer in the glass/FTO/TiO2/MAPbI3/u-spiroMeOTAD/MoO3/CuPc/Au device eliminates interface chemistry and leaves the band bending in the MAPbI3 as the sole change with respect to the reference device with the plain CuPc HTL. Remarkably, the device characteristics in this geometry are comparable to the characteristics of the glass/FTO/TiO2/ MAPbI3/MoO3/CuPc/Au device. While Voc is decreased by a small amount, the overall efficiency is improved mainly due to a substantial increase in FF. One interpretation of these findings is that the series resistance of the two devices containing the interlayers are comparable, while the charge collection at low net electric fields across the device is slightly improved if the formation of the reduced MoO2 interlayer is entirely avoided by the u-spiro-MeOTAD buffer layer, leading to a better fill factor. Yet, the FF in both interlayer devices is significantly lower than in the reference devices. We note that the hysteresis effect in the J−V curves between forward and reverse scan directions is of the same order of magnitude for all devices shown here. Multiple mechanisms contribute to the observed hysteresis in these devices, including ion migration with an activation energy affected by the interfaces between perovskite active layer and charge extraction layers, as discussed in previous work.30 Tracking of the maximum power point output has been conducted in accordance to earlier reports (see Figure S1).31 The EQE curves presented in Figure 4b show that major losses occur for carriers generated with longer wavelength photons for the devices with MoO3 and u-spiro-MeOTAD/ MoO3 interlayer. This observation can be attributed to two distinct effects. First, it is another indication for increased probability of recombination close to the hole extraction contact in devices with interlayer, as the blue light is absorbed predominantly near the electron-extracting TiO2 contact. Second, in devices with the high-work-function hole extraction contact, charge collection at that side of the cell is reduced which can potentially be linked to the band bending induced in 31496
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nanoparticle paste (Dyesol, 30 NR-D) diluted in 1 mL of ethanol, and spin-coated at 700 rpm for 10 s, 1000 rpm for 10 s, and 2000 rpm for 30 s. The films were sintered at 500 °C for 1 h. The CH3NH3PbI3 photoactive layer was deposited by a procedure reported by Ahn et al.35 The precursor solution consisted of 461 mg of PbI2, 159 mg of CH3NH3I, and 78 mg of DMSO (molar ratio 1:1:1) in 600 mg of DMF. The solution was stirred at room temperature until fully dissolved and then passed through a 0.45 μm PTFE syringe filter prior to use. To deposit the CH3NH3PbI3 film, the precursor solution was spin-coated on top of the mp-TiO2 film at 4000 rpm for 25 s. While the substrate was rotating (after spinning at 4000 rpm for approximately 10 s), 0.5 mL of diethyl ether was dripped onto to the substrate in a continuous stream. The resultant transparent film was heated at 65 °C for 1 min and 100 °C for 2 min in order to obtain a highly specular CH3NH3PbI3 film. The d-spiro-OMeTAD layer for control devices was deposited by spin-coating a solution containing 72 mg spiro-OMeTAD (Lumtec), 28.8 uL 4-tert-butylpyridine,17.5 μL of LiTFSI stock solution (stock: 520 mg mL−1 LiTFSI in acetonitrile), and 1 mL chlorobenzene at 5000 rpm for 30 s. The MoO3 was deposited on MAPbI3 and MAPIbI3/u-spiroOMeTAD at a rate of 0.2 Å/s at a base pressure lower than 2 × 10−7 Torr for a total thickness of 3.5 nm. CuPc hole transport layers were evaporated in an Angstrom Evaporator setup at a base pressure of 1 × 10−8 mbar and deposition rate of 0.5 Å/s monitored on a quartz crystal microbalance. Films of incremental thicknesses (2, 5, 10, and 20 nm) on top of the MAPbI3 and MAPbI3/MoO3 films were transferred with a brief (
