What Limits the Open-Circuit Voltage of Bromide ... - ACS Publications

Nov 19, 2018 - Si, InP, and CIGS cells (only GaAs has a significantly smaller. VOC loss, of ..... electrode). Note that this is a relative small forwa...
0 downloads 0 Views 1MB Size
Download PDF
http://pubs.acs.org/journal/aelccp

What Limits the Open-Circuit Voltage of Bromide Perovskite-Based Solar Cells? Arava Zohar,†,# Michael Kulbak,†,# Igal Levine,† Gary Hodes,*,† Antoine Kahn,*,‡ and David Cahen*,† †

Department of Materials & Interfaces, Weizmann Institute of Science, Rehovot 76100, Israel Department of Electrical Engineering, Princeton University, Princeton, New Jersey 08544, United States

Downloaded via IOWA STATE UNIV on January 11, 2019 at 13:04:26 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.

‡

S Supporting Information *

ABSTRACT: High band gap Pb bromide perovskite (APbBr3)-based solar cells, where A is a mixture of formamidinium, methylammonium, and Cs, show significantly higher, relative, VOC losses than their iodide analogs. Using photoluminescence-, quantum efficiency-, and surface photovoltage-spectroscopy measurements, we show the absence of any significant electronically active tail states within the bulk of the (FA0.85MA0.1Cs0.05)PbBr3 absorber. All methods confirm that EG = 2.28 eV for this halide perovskite, HaP. Contact potential difference measurements for this HaP, on different substrates, reveal a Z-shape dependence between the substrate work functions and that of the HaP, deposited on it, indicating that the HaP is relatively low doped and that its Fermi level is affected by the substrate onto which it is deposited. We confirm results from electron beam-induced current (EBIC) and other measurements that most voltage loss of cells, made with these HaP films, is at the HaP/selective-contact interface, specifically the TiO2/HaP one, and provide a complete account of these cells’ VOC losses. Capacitance measurements indicate that 350 mV VOC could be gained by eliminating (fast) interfacial states, emphasizing the importance of interface passivation. Still, even passivating the TiO2/HaP interface cannot eliminate the band misalignment with Br-based HaPs.

A

can be extracted from a given device. Vbimax is based on the differences between the electron electrochemical potentials of the materials that make up the junction. Vbimax can be influenced by states at the interfaces between the absorber (HaP in this case) and the selective charge transport layers (SCTLs), and especially by the energy distribution of these states. The third possible cause, also related to these interfaces, is the alignment of the energy levels between the different layers, i.e., the valence band maximum, VBM/HOMO (conduction band minimum, CBM/LUMO) energy of the SCTL used in the device relative to the absorber’s VBM (CBM) maximum (minimum). A few groups have tried to assess the defect density in the bulk and at the interfaces of the Br-HaP. Shi et al.6 characterized defects in a MAPbBr3 single crystal and found a remarkably low trap density, ntraps = 5.8 × 109 cm−3, as deduced from current− voltage measurements which, together with absorption and PL profiles, points to a nearly defect-free material (ascribed to high formation energy of deep traps). Rosenberg et al.7 studied MAPbBr3 single crystals in the dark, using Laplace current-deep level transient spectroscopy (DLTS) and found defect states in the band gap at energies ∼0.17 and ∼0.2 eV from the band edges with density of 108 and 109 cm−3, respectively. Using high-

fter several years of extensive efforts, laboratory-scale halide perovskite (HaP)-based photovoltaic (PV) devices have reached >23% photoconversion efficiency (PCE). Their tunable optical band gap, EG ≈ 1.45−1.6 eV for lead iodide-based perovskites to ∼2.3 V for lead bromide-based ones, allows tailoring this parameter for cells that are optically optimal for Si/HaP tandem PV devices. For the best I-based perovskite solar cells (PSCs), the opencircuit voltage loss (EG/q − VOC) is a remarkably low ∼0.35 V,1 with the highest reported VOC ≈ 1.24 V.2 This is on a par with cSi, InP, and CIGS cells (only GaAs has a significantly smaller VOC loss, of 10% power conversion efficiency, which is to our knowledge the highest voltage reported for a single-junction Br-HaP device.3 Wu et al.11 reported inverted MAPbBr3 devices that reach 1.61 V (with 7.5% power conversion efficiency) by employing PCBM and ICBA as the ETLs (and PEDOT:PSS as the HTL). While a better energy alignment between the CBM of Br-based HaPs (3.3−3.6 eV below vacuum level) and the PCBM/ICBA ETL LUMO could improve the VOC by 200 mV compared to a cell with TiO2 as ETL (CBM ∼3.9 eV below vacuum level), the VOC improvement was only half of that (100 mV). Zhang et al.12 compared the SnO2/ and TiO2/MAPbBr3 interfaces in PV devices and showed an increase of VOC by 200 mV when SnO2 was used as ETL. These results were surprising because SnO2 has a deeper CBM (4.5 eV) than TiO2 and should yield a lower VOC. Those experimental findings, especially the latter, suggest that an energy offset between the CBM of the HaP and the LUMO of the ETL was, at best, only partially responsible for the VOC losses. In general, in HaP-based solar cells, the vacuum level alignment approach (electron affinity rule or Anderson’s rule) to predict voltage losses at ETL− and HTL−HaP interfaces has not been very successful (unpublished results), a finding that, while not very surprising, warrants serious further consideration but is outside the scope of the present study. In our study we used a mixed cation Br-HaP, namely, (FA0.85MA0.1Cs0.05)PbBr3 (MixAPbBr3), as the photoabsorber for studying the origins of the voltage loss in PSCs. MixAPbBr3 has relatively long carrier diffusion lengths and low carrier density, compared to single-cation Br-based HaP, as previously shown by us.13 These two properties have proven to be highly beneficial for the lead iodide-based HaP PSCs. We used MixAPbBr3 in a cell configuration as well as on different substrates for characterization by optoelectronic techniques such as surface photovoltage (SPV), temperature dependence of VOC, and external quantum efficiency (EQE). In addition, we used pure electronic characterization methods, namely, contact potential difference (CPD) and capacitance voltage (C−V) measurements, to pinpoint causes for the voltage loss.

Voc = Eg − T ΔS − kT ln[iQY]

(1)

where T is temperature, S the entropy, k Boltzmann constant, and iQY the internal quantum yield of the absorber. The optical band gap of a MixAPbBr3 film (with a thickness of ∼250 nm), spin-coated on glass, was measured by photoluminescence (PL) and reflection-corrected transmission, as presented in Figure 1.

Figure 1. Reflection-corrected optical transmission (black) and photoluminescence (red) as a function of wavelength of MixAPbBr3 on glass.

The normalized PL spectrum shows a peak at 543 nm (2.28 eV), in good agreement with the maximum dT/dλ from the transmission measurement, 2.29 eV. Taking into account thermodynamic losses (∼330 meV at room temperature) and assuming a 100% iQY, the resulting VOC should be, according to eq 1, ∼1.95 V under 1 sun, as also calculated by Sutter-Fella et al.8 Using their experimentally measured internal quantum yield for MAPbBr3 (under ∼1 sun equivalent illumination) of 1− 6%8,14 reduces the VOC by about another 70−120 mV. We note that the external QY, which can be measured by electroluminescence, would be even lower because of other devicerelated losses. Therefore, the maximum VOC that can realistically be obtained from a Br-based HaP PV device is ∼1.83 V. However, the VOC at 300 K under 1 sun, of the FTO/TiO2/mpTiO2/MixAPbBr3/PTAA/Au cell that we prepared, reached only ∼1.4 V (for a 5% PCE; cf. panels A and B of Figure S1 for open-circuit voltage and steady-state voltage measurement, respectively), which is ∼430 mV below the above-estimated realistic maximum value. The best VOC ever reported for BrHaPs is, as noted earlier, 1.65 V, which is still ∼250 mV below the ∼1.9 V value estimated here. Compared to the situation encountered with HaPs with mainly iodide as halide, this is a large loss. In the optical measurements presented above, the HaP film had only one interface, glass. However, because of a lack of interface passivation leading to interfacial tail states and/or to a misalignment between the energy levels of the different layers, the effective energy gap (PV band gap) could potentially decrease when the HaP contacts different SCTLs. To check if such changes occur if the HaP is in a solar cell configuration, we measured EQE and surface photovoltage spectroscopy (SPS) upon illumination with photons from sub- to supra-band gap wavelengths (λ = 1200−400 nm), both on complete cells. 2

DOI: 10.1021/acsenergylett.8b01920 ACS Energy Lett. 2019, 4, 1−7

Letter

ACS Energy Letters

Figure 2. Semilogarithmic, surface-photovoltage spectroscopy (SPS) data and external quantum efficiency (EQE) spectrum of MixAPbBr3 as a function of excitation energy, both normalized to the maximum values, i.e., with values between 0 and 1 (A) and the first derivative dA/d(eV) of the EQE spectrum (B).

⌀b = ⌀m − χs

While both methods directly measure the processes due to photogeneration of carriers, one (EQE) does so at short circuit and the other (SPS) at open circuit. SPS measures the change in surface potential (or surface work function, WF) as a function of wavelength of illumination at open circuit. SPS is considered more sensitive than EQE because the photovoltage depends exponentially on the excitation power, in contrast to the linear dependence of the photocurrent from which the EQE is obtained. Normalized SPS and EQE as a function of the photon energy in eV are presented in Figure 2A and the first derivative of the EQE, as an intended approximation of Eg, is plotted in Figure 2B. The PV band gap, observed from the results of the first derivative of both measurements (for SPS, see Figure S2), agrees well with the optical measurement results presented in Figure 1, 2.28 ± 0.03 eV. Those results hint that there is little, if any, contribution from bulk tail states within the limits of sensitivity of these measurements in our experimental setups. Next, we look at the contribution of the HaP/ETL and HaP/ HTL interfaces to assess the presence of deep in-gap states. Trap states at one or both of the interfaces can facilitate nonradiative recombination and cause voltage losses. To the best of our knowledge, an experimental study of interfacial defect states in Br-based HaP PV devices has not been conducted. The only direct electronic measurements that can probe defect levels in a semiconductor are DLTS and thermally stimulated current (TSC). Both methods require temperature sweeping, which may pose a problem, as the phase transition of this HaP occurs close to room temperature (250−275 K).15 This fact limits what we could extract from DLTS measurements7 and complicates the interpretation of the DLTS and TSC results regarding trap states. Therefore, we turn to other techniques. The metal−semiconductor junction, known also as “Schottky” or “Schottky−Mott junction”, can highlight interfacial in-gap states by their effect on charge transport between the two materials upon intimate contact, as electrochemical potential equilibrium must be achieved (electrons transfer from the low WF material to the high WF one). The barrier height formed at a gap state-free metal−semiconductor junction, in the absence of any interaction between the two phases, should follow the Schottky−Mott theory, as shown in eq 2

(2)

where ⌀b is the barrier height for electron injection, ⌀m the metal WF, and χs the semiconductor electron affinity (equivalent to an alignment of vacuum levels across the interface). As an outcome, electrons are transferred from the semiconductor to the metal when the metal WF is high and in the opposite direction when it is low. As the semiconductor film becomes thinner and/or its doping density is low, its Fermi level, and thus its WF, responds to the charge transfer. Following deposition on a conducting substrate, and given the requirements of gap state-free interface and low doping mentioned above, the semiconductor WF will ultimately be dictated by the substrate WF (semiconductor space charge region width ≫ semiconductor film width). Such a phenomenon was already discussed for organic semiconductors (OSCs), which are known to be low-doped13 and to have very low intrinsic carrier concentrations, which allows tuning of their WF by varying the substrate WF. The change in WF can then be monitored via Kelvin probe CPD (KP-CPD) or UPS measurements.16,17 However, if defect states form at the metal− semiconductor interface, the semiconductor WF does not follow the Schottky−Mott model, as the charges exchanged between the two materials mainly fill traps. Such experiments have yielded valuable information regarding in-gap states close to a band edge of an OSC. Thus, metal−semiconductor junctions can be used in the case of HaPs as well as to find evidence for in-gap states. This argument holds especially true for the mixed cation Br, in which the free carrier concentration was found to be very low ( 2000 nm, i.e., much larger than the sample thickness (up to 500 nm). Figure 3 presents the WF measured at the exposed surface of MixAPbBr3 films deposited on a number of different substrates as a function of the WF of these substrates, measured using KPCPD. The substrates used here are not all metals, but in all cases their free carrier density is significantly higher than that of the HaP, which allows us to probe a wide range of WFs that cover the HaP band gap. The VBM (shown in blue in Figure 3) was measured by photo-emission yield spectroscopy (PEYS) under controlled humidity ambient (air) conditions (RH < 10%), and the ionization energy (IE), i.e. the position of the VBM below the vacuum level, is found to be 5.42 ± 0.03 eV. In all cases, we found the HaP films to have uniform perovskite structure 3

DOI: 10.1021/acsenergylett.8b01920 ACS Energy Lett. 2019, 4, 1−7

Letter

ACS Energy Letters

CsPbBr3), do not show such Z-shape like behavior (Figure S6), which we ascribe to their higher doping density and higher free carrier densities, which we reported on earlier.13 The Z-shape in Figure 3 is not symmetrical around the midgap of MixAPbBr3 (4.25 eV). The saturation plateau is close to the VBM when the film is deposited on high WF substrates but remains pinned at 900 meV below the CBM when deposited on low WF substrates. Our results not only are consistent with the deeper defect level reported by Sutter-Fella et al. for their thin polycrystalline films of MAPb(I1−xBrx) with relatively large values of x,9 which was not observed in single-crystal MAPbBr3 by Rosenberg et al.,7 but also reveal for the f irst time that the defect energy level is closest to the CBM. We therefore speculate that the reason for the Fermi level and WF pinning on low WF substrates is the existence of a sufficient density of in-gap states. The WF span obtained from the film correlates well with the measured VOC of our full device, 1.4 V, which points at interface states as the main cause for the VOC limitation. Thus, in-gap states at this energy are likely responsible for Fermi level pinning. Given their depth in the band gap, gap states presumably serve as recombination centers and can enhance nonradiative recombination, leading to losses that limit the VOC of the device. Because these in-gap states appear to be closer to the CBM than to the VBM of the gap, it is likely that they originate at the HaP/ETL interface. To test this hypothesis, we can estimate the defect density at the interfaces and the built-in voltage, Vbi, at the ETL/HaP junction. Capacitance−voltage (C−V) measurements at different frequencies can provide information about changes in band bending at the junction. A caveat of the analysis and interpretation of C−V measurements is that it requires a model to describe the capacitance and extract the junction parameters. As we reported before for MAPbBr3(Cl) layers,18 the high WF metal (Au) inverts the HaP film to p-type, yielding a p−n like junction with the TiO2, in agreement with the crosssectional EBIC results that show only one junction, at the TiO2/ HaP interface (Figure S7), in the devices studied here. We previously reported that the EBIC profile of the MixAPbBr3 junction resembles a p−i−n junction.13 However, for the present study we use MixAPbBr3 films of only half the thickness as that in the previous study (250 nm instead of ∼500 nm). With the thinner HaP film, the junction is closer to the TiO2/HaP interface, resulting in a relatively sharp, distinct junction. Thus, for our C−V analysis, we use the p−n junction model, i.e., C−2 ∝

Figure 3. Work function measurements of a MixAPbBr3 film on different substrates as a function of the WF of the substrate. The VB value is from photoemission yield spectroscopy (PEYS) measurements. The CBM is calculated from the VBM result by subtracting the measured band gap. The measurement errors are smaller than the widths of the data points that are shown.

(measured by XRD, Figure S3), to be pinhole-free (measured by plan-view scanning electron microscopy, Figure S4) and the IE to be independent of the substrate, as measured by PEYS (see Figure S5). In spite of an optical band gap of 2.28 eV (Figure 1), the maximum shift of the MixAPbBr3 surface potential is 1.4 V, with the WF varying between 5.3 and 3.9 eV. The resulting curve (Figure 3) has a so-called “Z-Shape” (or Zoro curve), which illustrates that the semiconductor WF is modified across a specific range of substrate WF values and reaches a lower and a higher plateau (ranges of substrate WFs, over which the semiconductor WF does not change anymore, irrespective of substrate WF). Over the linearly changing part of the curve, the slope of the film WF is approximately equal to 1, emphasizing that the electrochemical potential (or EFermi) of the HaP film seems to be def ined by the substrate. This behavior is consistent with the assumption that the defect density in MixAPbBr3 is sufficiently low in the energy range of the linear regime (5.3−3.9 eV) to allow such dependence on the substrate WF. Other Brbased perovskite films (MAPbBr 3 , MAPbBr 3 (Cl), and

Figure 4. Capacitance vs bias voltage, applied to the Au contact of a FTO/TiO2/mp-TiO2/MixAPbBr3/PTAA/Au cell at different frequencies (A); (B) the corresponding interface defect density as a function of bias from −0.2 to 1.2 V. The top x-axis gives the energy position where the defects are probed, as a result of the applied bias voltage. 4

DOI: 10.1021/acsenergylett.8b01920 ACS Energy Lett. 2019, 4, 1−7

Letter

ACS Energy Letters V (eq 5). The chosen model takes the two sides of the junction as two plates of a capacitor, and the width between those plates corresponds to the space charge region of the p−n junction. The depletion width, W, is defined by eq 3 2εs ijj kT zyz jjVbi − V − z j qNA k q zz{

W=

trapped electrons have enough time to leave the in-gap states and contribute to the cell capacitance. Strictly speaking, Vbi is always defined as the total bending obtained at equilibrium in the cell at zero frequency, as a result of the semiconductor parameters and interface properties. We assume here that the “Vbi” values, obtained at higher frequencies (Figure 4), represent the band bending that in principle would be achieved without the presence of the trap states. Under this assumption, the C−V results can predict, from Vbi, the VOC that can be obtained from the cell in the absence of the fast interface states (1.76 V). In practice, as the solar cell functions essentially in DC steady-state, Vbi will be only 1.41 V, because we need to take into account the carriers, trapped in interface states. This value, extracted from the LF C−V measurements, is in excellent agreement with the observed WF span shown in Figure 3 as well as with the VOC and SPV of the device under 1 sun illumination. The Vbi extracted from the HF C−V measurements indicates that the energy level alignment between the various materials at the p−n junction would limit VOC to 1.76 V. This value is 520 mV below the 2.28 eV band gap of the material, of which 330 mV is a pure thermodynamic loss at room temperature for nonconcentrated sunlight; ∼120 mV is lost because of the less than 100% PL quantum yield; and the remaining 70 mV must be other losses, which we ascribe to energy level misalignment at the interfaces. The doping density of the HaP was calculated from the slope of the 1/C2 curve and found to be Na= 2 × 1015 cm−3 for all frequencies (but is orders of magnitude higher than the free carrier density derived from conductivity measurements13). The capacitance versus voltage data can be used to estimate the defect energy within the HaP band gap by calculating the excess capacitance (LF-HF) at different applied voltages, as the applied bias modifies the HaP Fermi level position with respect to the grounded electron-collecting electrode (FTO). The Fermi level sweeps through the HaP gap at the interface and enables mapping of the energy gap states at the junction. By subtracting the LF capacitance from the HF capacitance, we can estimate the trapped electron or defect density per cubic cm. Figure 4b shows the interface defect concentration (Dit, cm−2) at the TiO2/ MixAPbBr3 junction (assuming that the Vbi in the entire cell originates from the ETL/HaP junction) from sweeping the bias from −0.2 to 1.2 V. The bottom x-axis shows the applied bias, which is limited to a narrow range around 0 V to prevent current from passing though the diode. The upper x-axis scale presents the energies probed within the HaP band gap, with respect to the HaP CBM at the junction’s interface; V = 0 represents EFermi in the dark, which we estimate to be 4.9 eV from the vacuum level, based on the assumption that the HF Vbi originates from the HaP/TiO2 junction (hence, EFermi = CBM + Vbi = 3.14 + 1.76 = 4.9 eV). As we apply forward bias, we probe energies closer to the CBM, resulting in a Dit analysis that shows a low defect concentration below mid gap, which increases from mid gap toward the CBM, in agreement with the Z-shape plot shown in Figure 3. Next, because both C−V and the Z-shape plot were measured under dark conditions, we check if the observed interfacial defects affect the VOC losses of the cell under illumination. A simple way to assess the VOC losses under illumination, which does not require any assumptions or a model, is to extrapolate the VOC to T = 0 K. VOC at 0 K should be equal to the activation energy for e−h recombination in the cell if all the thermalrelated losses are inactive. In the ideal case, where the dominant e−h recombination mechanism is within the HaP space charge region, the activation energy is equivalent to the semiconductor

(3)

where εs is the semiconductor permittivity; Vbi and V are the built-in voltage and the applied bias [V], respectively, and NA is the acceptor impurity density [cm−3]. As we deal with a p-type semiconductor, holes, due to ionized acceptors, are the majority carriers. Using the calculated W, we can derive the junction capacitance with eq 4 C=

∂Q SC ∂V

=

∂(qNAW ) ∂V

qεsNA ij y jjV − V − 2kT zzz bi j 2 jk q zz{

−1/2

→C=

(4) 2

where Qsc is the semiconductor charge density in [C/cm ] in the space charge region (Qsc = qNAW). To extract the electronic parameters of the junction, such as NA and Vbi, the capacitance for an abrupt p−n junction is often presented as shown in eq 5 1 = C2

(

2 Vbi − V − qϵsNA

kT q

) →N

A

=

2 1 · qϵs d( C12 ) dV

(5)

where Vbi is calculated from the intercept of 1/C2 with the x-axis and NA is calculated from the slope of 1/C2 vs the bias V. Figure 4A shows the capacitance, presented in F−2 units, of the full cell, measured under dark conditions at different frequencies and for an applied bias between 0.5 and 1.5 V (the bias applied to the Au electrode). Note that this is a relative small forward bias, which must be applied to probe the junction because the device capacitance is fixed up to this point as reported and partly explained by others.19−21 A recent explanation for this deviation from the normal Mott−Schottky characteristic behavior invokes mobile ions.22 Fischer et al. found that when prebiasing solar cells (2 V up to 60 s, i.e., somewhat higher than the bias applied here), hysteresis is observed in the forward direction, which disappeared during the second forward sweep. This can happen when the ions are pushed from the interface toward the bulk to form a p−n junction in the perovskite absorber, which can be correlated to a linear C−2(V) plot. The measurements were done in the dark, because MixAPbBr3 is a good photoconductor and we need to minimize the current between the contacts during capacitance measurements (there is almost no current below 1.4 V in the dark, as shown in Figure S1). Figure 4A presents the measured capacitance versus bias of the FTO/TiO2/mp-TiO2/MixAPbBr3/PTAA/Au cell at different frequencies, as well as the values of Vbi, calculated using eq 3, from the linear fit marked by the black dashed line, superimposed on the data at each frequency. The value changes from Vbi = 1.76 V at high frequency (HF, 100 kHz) to Vbi = 1.41 V for low frequency (LF, 3 eV, and EFermi < 100 meV below the CBM. These requirements are needed to keep the selectivity of the n-type ETL without losing voltage via electron transfer. Unfortunately, we are not familiar with such material, but other inorganic compounds such as ZnS or (Zn, Mg)O may be better in terms of alignment than TiO2. We note that there are not many selective materials with the combination of properties possessed by TiO2.

Figure 5. Open-circuit voltage versus temperature for a FTO/TiO2/ mp-TiO2/MixAPbBr3/PTAA/Au cell at high (0.5 sun, red) and low (0.16 sun, black) light intensities.

different intensities, reaching almost 300 meV, suggests that under lower light conditions, interfacial recombination becomes dominant, reducing the theoretical maximum VOC (we can, at T = 0 K, neglect thermodynamic losses) by about 300 meV to ∼1.85 V. This result indicates a more than 400 mV loss (2.28 V − 1.85 V) due to interfacial recombination via trap states, in good agreement with the losses in the built-in potential deduced from the C−V measurements (∼350 meV). Under higher light intensity (∼0.5 sun), a large part of the interfacial traps is filled; therefore, band-to-band bimolecular recombination in the HaP becomes more dominant, enabling the device to reach a higher VOC value of ∼2.13 V. Still, though, we identify ∼150 mV losses related to interfacial recombination because VOC at 0 K did not reach the Eg value of the HaP, as opposed to MAPbI3 devices.24 Moreover, considering the actual working conditions of the solar cell, it is quite likely that during most of the cell operation under real-life operating conditions, the light intensity will be lower than 1 sun, suggesting that the losses in the Voc due to interfacial recombination would be up to 150 mV.5 To conclude, we identified two reasons for the voltage losses in PV devices based on MixAPbBr3 and, generally, for other Brbased HaPs and illustrate these losses in Figure 6; the main voltage loss is the poor passivation of the HaP at the TiO2 interface, resulting in interfacial recombination via trap states. This loss amounts to roughly ∼350−400 mV with respect to the maximal built-in potential, measured by C−V and VOC vs T. To overcome this problem, one should add a passivation layer, e.g., a thin high band gap layer that electrically and chemically decouples the TiO2 from the HaP or a chemical passivation treatment of the TiO2 surface, such as C60-SAM.25 Wojciechowski et al. showed that such a SAM can passivate the trap states at the HaP/TiO2 interface and therefore reduce the nonradiative recombination channels. In practice, when subtracting all the experimentally proven losses from the theoretical maximum VOC (which takes into account the

■

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsenergylett.8b01920. Experimental details, J−V plots; steady-state voltage measurement; X-ray diffractograms; plan-view scanning electron microscopy images; PEYS measurement; work function of MAPbBr(Cl)3, CsPbBr3, and MAPbBr3 on different substrates; EBIC of a full solar cell (PDF)

■

AUTHOR INFORMATION

Corresponding Authors

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

Arava Zohar: 0000-0002-8292-5968 Gary Hodes: 0000-0001-7798-195X David Cahen: 0000-0001-8118-5446 6

DOI: 10.1021/acsenergylett.8b01920 ACS Energy Lett. 2019, 4, 1−7

Letter

ACS Energy Letters Author Contributions

Perovskite Films for Efficient Photovoltaic Devices. ACS Energy Lett. 2018, 3 (6), 1233−1240. (15) Schueller, E. C.; Laurita, G.; Fabini, D. H.; Stoumpos, C. C.; Kanatzidis, M. G.; Seshadri, R. Crystal Structure Evolution and Notable Thermal Expansion in Hybrid Perovskites Formamidinium Tin Iodide and Formamidinium Lead Bromide. Inorg. Chem. 2018, 57 (2), 695− 701. (16) Steim, R.; Kogler, F. R.; Brabec, C. J. Interface Materials for Organic Solar Cells. J. Mater. Chem. 2010, 20 (13), 2499−2512. (17) Yang, J.-P.; Shang, L.-T.; Bussolotti, F.; Cheng, L.-W.; Wang, W.Q.; Zeng, X.-H.; Kera, S.; Li, Y.-Q.; Tang, J.-X.; Ueno, N. Fermi-Level Pinning Appears upon Weak Electrode-Organic Contact without Gap States: A Universal Phenomenon. Org. Electron. 2017, 48, 172−178. (18) Kedem, N.; Kulbak, M.; Brenner, T. M.; Hodes, G.; Cahen, D. Type-Inversion as a Working Mechanism of High Voltage MAPbBr3(Cl)-Based Halide Perovskite Solar Cells. Phys. Chem. Chem. Phys. 2017, 19 (8), 5753−5762. (19) Almora, O.; Aranda, C.; Mas-Marzá, E.; Garcia-Belmonte, G. On Mott-Schottky Analysis Interpretation of Capacitance Measurements in Organometal Perovskite Solar Cells. Appl. Phys. Lett. 2016, 109 (17), 173903. (20) Mora-Seró, I.; Garcia-Belmonte, G.; Boix, P. P.; Vázquez, M. A.; Bisquert, J. Impedance Spectroscopy Characterisation of Highly Efficient Silicon Solar Cells under Different Light Illumination Intensities. Energy Environ. Sci. 2009, 2 (6), 678. (21) O’Malley, K. M.; Li, C.-Z.; Yip, H.-L.; Jen, A. K.-Y. Enhanced Open-Circuit Voltage in High Performance Polymer/Fullerene BulkHeterojunction Solar Cells by Cathode Modification with a C60 Surfactant. Adv. Energy Mater. 2012, 2 (1), 82−86. (22) Fischer, M.; Tvingstedt, K.; Baumann, A.; Dyakonov, V. Doping Profile in Planar Hybrid Perovskite Solar Cells Identifying Mobile Ions. ACS Appl. Energy Mater. 2018, 1 (10), 5129−5134. (23) Sze, S. m.; Ng, K. K. Physics and Properties of Semiconductors A Review. In Physics of Semiconductor Devices; John Wiley & Sons, Inc., 2006; pp 5−75. (24) Levine, I.; Nayak, P. K.; Wang, J. T.-W.; Sakai, N.; Van Reenen, S.; Brenner, T. M.; Mukhopadhyay, S.; Snaith, H. J.; Hodes, G.; Cahen, D. Interface-Dependent Ion Migration/Accumulation Controls Hysteresis in MAPbI3 Solar Cells. J. Phys. Chem. C 2016, 120 (30), 16399− 16411. (25) Wojciechowski, K.; Stranks, S. D.; Abate, A.; Sadoughi, G.; Sadhanala, A.; Kopidakis, N.; Rumbles, G.; Li, C.-Z.; Friend, R. H.; Jen, A. K.-Y.; et al. Heterojunction Modification for Highly Efficient Organic−Inorganic Perovskite Solar Cells. ACS Nano 2014, 8 (12), 12701−12709.

#

A.Z. and M.K. contributed equally to this work.

Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS We thank Dr. T. Dittrich (HZB) for collaboration on the SPS measurements and Dr. Pabitra K. Nayak for fruitful discussions. We acknowledge support by US-Israel Binational Science Foundation (BSF-Grant No. 2014357) and the Ullmann Family Foundation (@Weizmann Institute), for partial support. D.C. held the Schaefer Chair in Energy Research.

■

REFERENCES

(1) Nayak, P. K.; Cahen, D. Updated Assessment of Possibilities and Limits for Solar Cells. Adv. Mater. 2014, 26 (10), 1622−1628. (2) Saliba, M.; Matsui, T.; Domanski, K.; Seo, J.-Y.; Ummadisingu, A.; Zakeeruddin, S. M.; Correa-Baena, J.-P.; Tress, W. R.; Abate, A.; Hagfeldt, A.; et al. Incorporation of Rubidium Cations into Perovskite Solar Cells Improves Photovoltaic Performance. Science 2016, 354 (6309), 206−209. (3) Hu, X.; Jiang, X.-F.; Xing, X.; Nian, L.; Liu, X.; Huang, R.; Wang, K.; Yip, H.-L.; Zhou, G. Wide-Bandgap Perovskite Solar Cells With Large Open-Circuit Voltage of 1653 mV Through Interfacial Engineering. Sol. RRL 2018, 2 (8), 1800083. (4) Nayak, P. K.; Bisquert, J.; Cahen, D. Assessing Possibilities and Limits for Solar Cells. Adv. Mater. 2011, 23 (25), 2870−2876. (5) Zheng, X.; Chen, B.; Yang, M.; Wu, C.; Orler, B.; Moore, R. B.; Zhu, K.; Priya, S. The Controlling Mechanism for Potential Loss in CH3NH3PbBr3 Hybrid Solar Cells. ACS Energy Lett. 2016, 1 (2), 424− 430. (6) Shi, D.; Adinolfi, V.; Comin, R.; Yuan, M.; Alarousu, E.; Buin, A.; Chen, Y.; Hoogland, S.; Rothenberger, A.; Katsiev, K.; et al. Low TrapState Density and Long Carrier Diffusion in Organolead Trihalide Perovskite Single Crystals. Science 2015, 347 (6221), 519−522. (7) Rosenberg, J. W.; Legodi, M. J.; Rakita, Y.; Cahen, D.; Diale, M. Laplace Current Deep Level Transient Spectroscopy Measurements of Defect States in Methylammonium Lead Bromide Single Crystals. J. Appl. Phys. 2017, 122 (14), 145701. (8) Sutter-Fella, C. M.; Li, Y.; Amani, M.; Ager, J. W.; Toma, F. M.; Yablonovitch, E.; Sharp, I. D.; Javey, A. High Photoluminescence Quantum Yield in Band Gap Tunable Bromide Containing Mixed Halide Perovskites. Nano Lett. 2016, 16 (1), 800−806. (9) Sutter-Fella, C. M.; Miller, D. W.; Ngo, Q. P.; Roe, E. T.; Toma, F. M.; Sharp, I. D.; Lonergan, M. C.; Javey, A. Band Tailing and Deep Defect States in CH3NH3Pb(I1−XBrx)3 Perovskites As Revealed by SubBandgap Photocurrent. ACS Energy Lett. 2017, 2 (3), 709−715. (10) Arora, N.; Orlandi, S.; Dar, M. I.; Aghazada, S.; Jacopin, G.; Cavazzini, M.; Mosconi, E.; Gratia, P.; De Angelis, F.; Pozzi, G.; et al. High Open-Circuit Voltage: Fabrication of Formamidinium Lead Bromide Perovskite Solar Cells Using Fluorene−Dithiophene Derivatives as Hole-Transporting Materials. ACS Energy Lett. 2016, 1 (1), 107−112. (11) Wu, C.-G.; Chiang, C.-H.; Chang, S. H. A Perovskite Cell with a Record-High-V O c of 1.61 V Based on Solvent Annealed CH3NH3PbBr3/ICBA Active Layer. Nanoscale 2016, 8 (7), 4077− 4085. (12) Zhang, Y.; Liang, Y.; Wang, Y.; Guo, F.; Sun, L.; Xu, D. Planar FAPbBr3 Solar Cells with Power Conversion Efficiency above 10%. ACS Energy Lett. 2018, 3 (8), 1808−1814. (13) Kulbak, M.; Levine, I.; Barak-Kulbak, E.; Gupta, S.; Zohar, A.; Balberg, I.; Hodes, G.; Cahen, D. Control over Self-Doping in High Band Gap Perovskite Films. Adv. Energy Mater. 2018, 8 (23), 1800398. (14) Noel, N. K.; Wenger, B.; Habisreutinger, S. N.; Patel, J. B.; Crothers, T.; Wang, Z.; Nicholas, R. J.; Johnston, M. B.; Herz, L. M.; Snaith, H. J. Highly Crystalline Methylammonium Lead Tribromide 7

DOI: 10.1021/acsenergylett.8b01920 ACS Energy Lett. 2019, 4, 1−7