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Letter

Current Induced Phase Segregation in Mixed Halide Hybrid Perovskites and its Impact on Two-Terminal Tandem Solar Cell Design Ian L. Braly, Ryan J Stoddard, Adharsh Rajagopal, Alexander R. Uhl, John K Katahara, Alex K-y. Jen, and Hugh W. Hillhouse ACS Energy Lett., Just Accepted Manuscript • DOI: 10.1021/acsenergylett.7b00525 • Publication Date (Web): 24 Jul 2017 Downloaded from http://pubs.acs.org on July 25, 2017

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ACS Energy Letters

Current Induced Phase Segregation in Mixed Halide Hybrid Perovskites and its Impact on TwoTerminal Tandem Solar Cell Design Ian L. Braly1†, Ryan J. Stoddard1†, Adharsh Rajagopal2, Alexander Uhl1, John Katahara1, Alex K.-Y. Jen2 and Hugh W. Hillhouse1* 1

Department of Chemical Engineering, Clean Energy Institute, and Molecular Engineering &

Sciences Institute, University of Washington, Seattle, Washington 98195-1652, United States 2

Department of Materials Science and Engineering, University of Washington, Seattle, WA 98195-2120, United States.

†

Equal author contribution

*Corresponding author. Email: [email protected]; Ph 1-206-685-5257 (H.W.H).

1

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Abstract Mixed-halide hybrid perovskites are of significant interest since their bandgap can be tuned as a current-matched top-cell in tandem photovoltaics. However, several mixed-halide perovskites phase segregate under illumination, exhibit large voltage deficits, and produce unstable photocurrents. We investigate the origin of phase segregation and implication for tandems with mixed-halide large-bandgap (~1.75eV) perovskites. We show explicitly that MAPb(I0.6Br0.4)3 and (MA0.9,Cs0.1)Pb(I0.6,Br0.4)3, referred to as “MA” and “MACs,” rapidly phase segregate in the dark upon 1-Sun equivalent current injection. This is direct experimental evidence that conduction-band electrons or valance-band holes are the culprit behind phase segregation. In contrast, (FA0.83,Cs0.17)Pb(I0.66,Br0.34)3, or “FACs,” prepared at only 75°C resist phase segregation below 4-Suns injection. FACs prepared at 160°C yields larger grains and withstands higher injected carrier concentrations before phase-segregation. The FACs and MACs devices sustain near constant power output at 1-Sun and do not affect the current output of a CIGS bottom cell when used as an incident-light filter.

The addition of a solution processed top-cell to a silicon or Cu(In,Ga)(S,Se)2 (CIGS) bottom cell may result in increased voltage and efficiency compared to the single junction solar cell 2


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ACS Energy Letters

without significant additional processing costs.1 The caveat is the top cell must produce the same current under operating conditions as that generated by the bottom cell. It is true that current matching can be achieved by using a smaller-than-ideal bandgap for the top absorber if a thinner top absorber layer is used. However, the theoretical maximum efficiency for a two-terminal tandem is lower if a thin top absorber with smaller bandgap is employed since it results in a lower open circuit voltage. The lower voltage is a result of ineffective splitting of the solar spectrum and the presence of avoidable thermalization losses in the bottom cell. Thus, calculations show that the maximum efficiency of a two-terminal tandem with a bottom cell bandgap of ~1.1 eV will be achieved with an optically thick top cell with bandgap of ~1.75 eV,25

assuming the top and bottom cells operate in the Shockley-Queisser limit.6 However, if the top

or bottom cells deviate from ideality and have IPCE less than one over the range of photon energies they are ideally supposed to harvest, the optimum bandgaps will shift slightly. The record certified single junction lead halide hybrid perovskite solar cells have a bandgap of 1.61 eV,7 which is too low for implementation into a two-terminal tandem with a c-Si or CIGS bottom cell. Therefore, hybrid perovskite compositions with larger bandgaps are under investigation. An effective way to tune the hybrid perovskite bandgap into the ideal range is to form iodide-bromide alloys.8, 9 Large-bandgap, mixed halide perovskites have been successfully paired with c-Si,10-13 Cu2ZnSn(S,Se)4 (CZTS),14 CIGS,15,

16

and lead-tin iodide perovskites5,17 into monolithic 2-

terminal tandems; however, significant voltage losses in mixed halide top cells have prevented such devices from exceeding 80% of their respective detailed-balance limit open circuit voltage5 (for comparison, the record III-V tandem has a VOC of 2.248 V18 which is 94.5% of its respective detailed-balance limit). Indeed, the open circuit voltage of single junction mixed halide hybrid 3


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perovskites devices with bandgaps between 1.7 and 1.8 eV has not surpassed 1.24 V19-22 (the detailed balance limit VOC for a 1.75 eV bandgap absorber is 1.463 V). When the bromide composition exceeds about 20% of the halides in methylammonium lead iodobromide, MAPb(I,Br)3, the material reversibly segregates into an iodide-rich phase and a bromide rich phase under illumination.23 This phase segregation was first discovered through the observation of (a) a red-shifting photoluminescence (PL) peak position, (b) a red-shifting absorption band edge, and (c) splitting X-ray diffraction peaks with time under continuous illumination. The inferred impact light induced phase segregation can have on device performance is three-fold. First, the red-shift of the band edge can decrease the open-circuit voltage (VOC) due to increased thermalization losses in the absorber (i.e. carriers lose potential energy moving from the large-bandgap material to the small-bandgap material). Second, segregation into two phases can inhibit current collection due to carrier accumulation in the smaller bandgap (I-rich) phase, especially if the I-rich domains are spatially confined. Third, a transient band edge could inhibit effective current matching in a two-terminal tandem because it would cause the top cell transmissivity and therefore the associated bottom cell current to change with time. These implications, concurrent with observations that VOC does not increase proportionally with increasing bandgap8,

9, 24, 25

and that current decreases with time under

constant operation in many mixed halide perovskite devices,26-28 have driven significant research efforts to discover large-bandgap compositions that do not exhibit phase segregation under illumination. Several approaches have been used to suppress the phase segregation that occurs under illumination in high bromide hybrid perovskites including increasing domain size,28 increasing lattice strain,26 and mixing the A-cation site.21, 29-32 One effective approach has been to replace 4


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the single-cation methylammonium with “double-cation” formamidinium and-cesium (FACs), which exhibits a stable PL peak position with time.21, 33 A recent review summarizes the current understanding of the light induced phase segregation.27 Molecular dynamic simulations have predicted phase segregation to be caused by excess charge carriers distorting the lead halide lattice via electron-phonon coupling, while cathodoluminescence experiments studying MAPb(I,Br)3 films before and after light soaking provided evidence that the iodide-rich phase forms mostly between morphological domains during phase segregation.34 Here, we provide direct experimental evidence that excess charge carriers alone are the cause of phase segregation in mixed halide hybrid perovskites by observing spectrally resolved electroluminescence as a function of time under constant current injection conditions. MAPb(I0.6Br0.4)3 devices phase segregate when injected current densities are equivalent to the detailed

balance

limit

short-circuit

current

(JSC).

However,

we

show

that

(FA0.83,Cs0.17)Pb(I0.66,Br0.34)3 devices with a similar bandgap and annealing conditions do not phase segregate under the same current injection conditions. Furthermore, we show that elevated annealing temperatures of 160°C improve film morphology of (FA0.83,Cs0.17)Pb(I0.66,Br0.34)3, which in turn improve the phase stability under both high current injection and incident light intensity conditions. Finally, we emulate the impact of phase segregation on tandem device performance for different hybrid perovskite compositions by (1) measuring operating current under illumination with time, and (2) measuring CIGS device JSC with time while shaded by phase segregating large-bandgap perovskite thin films. Note that hereafter we use “MA”, “MACs”, and “FACs” to refer not to the isolated cations but to the entire mixed-halide hybrid

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perovskite

compositions

of

MAPb(I0.6Br0.4)3,

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(MA0.9,Cs0.1)Pb(I0.6,Br0.4)3,

and

(FA0.83,Cs0.17)Pb(I0.66,Br0.34)3, respectively. Morphology and Domain Size. We evaluated the impact of cation substitution of the monovalent cation site and film morphology on phase stability and photovoltaic device characteristics of large-bandgap hybrid perovskites. The film preparation and device fabrication methods are given in the supplemental information (SI). Improving crystallinity has been proposed as a method to kinetically limit light induced phase segregation in hybrid perovskites.33 In this study, we probe relative phase stability of large-bandgap hybrid perovskites with different compositions as well as films with different crystalline domain sizes but with the same composition. Figure 1 a-c shows top-down SEM images of MA annealed at 100 ˚C for 10 minutes, FACs annealed at 75 °C for 10 minutes (FACs-LT), and FACs annealed at 165 °C for 50 minutes (FACs-HT). Additionally, MACs thin-films were prepared (see SI for preparation details), and top-down SEM and XRD results are given in Figure S1. All perovskite film compositions in this work were assumed to be the same as the initial ink solution concentrations. The bandgaps of these materials (abstracted from UV-vis absorbance) are 1.83 eV for MA and MACs and 1.75 eV for FACs-LT and FACs-HT as shown in Figure S2. Both FACs-LT and MA films exhibit smaller morphological domains (140 ± 150 nm and 130 ± 50 nm, respectively), while the FACs-HT film annealed at a higher temperature has larger domains (320 ± 140 nm, see Figure S3 for more details of morphological domain size statistics).

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Figure 1. Top-view SEM micrographs of perovskite thin films: (a) MAPb(I0.4,Br0.6)3 annealed at 100°C for 10 minutes, (b) (FA0.83,Cs0.17)Pb(I0.66,Br0.33)3 annealed at 75 ˚C for 10 min, and (c) (FA0.83Cs0.17Pb(I0.66,Br0.33)3 annealed at 165 ˚C for 50 min. (d) corresponding XRD patterns of above shown films. (e) Normalized and smoothed high resolution data of (200) reflections of comparable samples with the same compositions as (d), with the inset showing the full-width at half-maximum (B) and the domain size d as determined by the Scherrer formula.

Electron micrographs of large morphological domains are not always indicators of crystal domain size or extent of lattice order. Therefore, we collected XRD data to observe differences in area-averaged crystallite size. Figure 1 d shows the XRD data collected on the same films as Figure 1 a-c. The MA film was observed to exhibit stronger preferential orientation with the glass substrate along the (100) plane (See Figures S4 and S5). The FWHM of the (200) peak 7


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from the MA, FACs-HT, and FACs-LT diffraction patterns are shown in Figure 1 e to be 0.09, 0.14, and 0.19 degrees 2θ, which, according to Scherrer equation, translates to crystallite domain sizes of 110, 70, and 50 nm, respectively. It is intriguing that these XRD determined domain sizes are not the same values as the SEM determined domain sizes. However, a calculated XRD domain size (from Scherrer equation) that is smaller than SEM morphological grain size for hybrid perovskite thin films is not unique to this work.21,28,33 We note that both crystallite size and strain contribute to the FWHM broadening; therefore, the Scherrer equation may underestimate the crystalline domain size for FACs films. Nevertheless, since both the FACs-HT and FACs-LT samples likely have similar strain within the crystal, we conclude that the FACsHT films have larger morphological and crystallite domains than FACs-LT.

Electroluminescence and Photoluminescence at 1 Sun Equivalent Carrier Densities. To investigate the origin of phase segregation in mixed-halide perovskites we measured spectrally resolved electroluminescence (EL) of devices and photoluminescence (PL) of neat films for each composition-anneal condition. Both the MA and FACs-HT devices used for EL studies exhibit power conversion efficiencies (PCE) above 10%, with device parameters comparable to other studies in this bandgap range17, 21, 27-29 (see Figure S6 and Table S1 for device data). Relevant 1 Sun carrier densities were achieved by setting the current injection density to the ShockleyQueisser limit short-circuit current density (JSC,SQ) for a 1.81 eV material (19.2 mA/cm2), and by setting the continuous wave (CW) photoexcitation intensity to the above-bandgap photon flux for a 1.75 eV material (1.3x1021 photons/m2/s) for EL and PL experiments, respectively. We chose to inject with JSC,SQ for EL measurements instead of device JSC to prevent from providing a

8


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ACS Energy Letters

carrier density lower than present under AM1.5 illumination for devices with transport-limited JSC. Results for MA and FACs-LT are presented in Figure 2. Figure 2 a shows absolute intensity electroluminescence evolution with time for a MA device. In five minutes, the MA electroluminescence spectrum peak position red-shifts from 1.8eV to 1.63eV while the emitted flux increases by 50-fold. This is in agreement with the formation of the smaller bandgap iodide phase and shows, for the first time, phase segregation in the absence of photoexcitation. Figure 2 b depicts the mean emission spectrum energy () for MA and FACs-LT with current injection and photo-excitation. The parameter is simply the first moment of the emission spectrum and has been employed previously to quantify peak shifts due to phase segregation33. For MA, decreases from 1.8 eV to ~1.65eV in five minutes for both EL (current injection) and PL (photoexcitation). This data strongly confirms the hypothesis that excess charge carrier density is the origin of phase segregation. Moreover, this finding aligns with mechanisms suggested previously, i.e. excited charge carriers inducing lattice strain via electron-phonon coupling which results in the nucleation of I-rich clusters near the borders of morphological domains.34 Our findings refine this mechanism by suggesting that the I-rich clusters are not light-stabilized, but rather sustained by a high carrier density. The high carrier density that results in phase segregation can be achieved by either photo-excitation or current injection.

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Figure 2. Electroluminescence and photoluminescence measurements over time. (a) Electroluminescence spectra at selected times of a MAPb(I0.6Br0.4)3 device being injected with a constant current density of 19.2 mA/cm2 for five minutes. (b) Mean electroluminescence and photoluminescence emission energy of devices utilizing MAPb(I0.6Br0.4)3 and low temperature annealed

(FA0.83,Cs0.17)Pb(I0.34Br0.66)3 absorbers. (c) External electroluminescence and

photoluminescence quantum yield from the same experiment.

Unlike MA, FACs-LT shows a stable peak position in PL and EL with time. This indicates the FACs composition is phase stable at 1 Sun relevant carrier densities, and demonstrates this 10


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ACS Energy Letters

composition has enhanced phase stability over the MA analogs even without a high temperature anneal. This highlights the non-trivial relationship between composition, crystallinity and phase stability, as has been shown previously.33 MACs and FACs-HT exhibit similar EL and PL results as MA and FACs-LT, respectively, as shown in Figure S7. Figure 2 c shows EL and PL quantum yield (ELQY and PLQY) for MA and FACs-LT devices and films. The increase in PLQY and ELQY in the phase segregating MA material has been proposed to be due to carrier funneling into the locally smaller bandgap, I-rich domains as demonstrated elsewhere35-37. This mechanism could be present, but we note that the driving force for photoemission becomes exponentially stronger for states at lower energy, thus even without any carrier funneling mechanism, one should expect significantly higher QY simply due to the appearance of lower energy states (see Figure S8 and discussion there). The several orders-of-magnitude difference between ELQY and PLQY is likely due to additional recombination pathways that are present in devices (e.g. interfacial recombination) and carrier separation at the ETM/HP and HTM/HP interfaces. Furthermore, an additional control experiment is discussed in the SI, where we show that an electric field alone is not sufficient to induce phase segregation without carrier injection (see Figures S9 and S10).

Higher Carrier Densities Required for Phase Segregation of FACs. We have shown that both the FACs-LT and FACs-HT are phase stable at 1 Sun relevant carrier densities, achieved with either photo-excitation or current injection. To probe the boundaries of the phase stability of FACs materials, we investigated phase stability under carrier densities higher than present under 1 Sun illumination. In this experiment, we observed EL and PL emission spectra for both FACsLT and FACs-HT with time and doubled the photoexcitation flux or current injection if we did 11


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not observe an emission peak low-energy shoulder. We continued this procedure until a lowenergy shoulder was finally visible. The results of this experiment are presented in Figure 3 a and b for FACs-LT and FACs-HT, respectively. Dashed vertical lines indicate times at which the photoexcitation flux or current injection was doubled. Both FACs-LT and FACs-HT show stable with time at 1 Sun, 2 Suns, and 4 Suns carrier density upon photoexcitation as well as current injection. However, both FACs-LT and FACs-HT display obvious emission spectrum red-shifts at 8 Suns and 32 Suns, respectively. This demonstrates that the FACs materials will ultimately show similar phase segregating behavior as MA materials, however, at several-fold higher excited carrier density.

Figure 3. Mean emission spectrum energy with time. The electroluminescence current density and photoluminescence illumination intensity is doubled every minute. (a) Low temperature annealed (FA0.83,Cs0.17)Pb(I0.66Br0.34)3 under 2, 4, and 8 Suns equivalent current density or illumination intensity. (b) High temperature annealed (FA0.83,Cs0.17)Pb(I0.66Br0.34)3 under 2, 4, 8, 16 and 32 Suns equivalent current density or illumination intensity. 12


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We propose that a combination of larger morphological grains and the presence of size mismatch on the perovskite A+ site in FACs materials stabilize these materials at 1 Sun excitation. However, as excited carrier induced strain increases with carrier densities at above 1 Sun, nucleation of I-rich clusters to relieve bulk strain will overcome the energy penalty of phase boundaries, and the threshold is higher in FACs than in MA. Further, Figure 3 shows that the carrier density required to induce phase segregation in FACs-HT was higher than FACs-LT, indicating that larger grains indeed provide enhanced phase stability within the FACs material class. Interestingly, the threshold carrier density required for phase segregation was equivalent for PL and EL measurements, 8 Suns in the case of FACS-LT and 32 Suns for FACs-HT. These observations confirm that phase segregation behavior is similar for photoexcitation and current injection at comparable excess carrier densities. We note that phase segregation is a function of time as well as carrier density, and one minute may not be sufficient time to observe the slow phase segregation behavior. Indeed, the decreasing slope for FACs-HT at 16 Suns indicates that phase segregation would likely be observed at longer times, suggesting our FACs-HT material is not phase stable at 16 Suns. However, the data unambiguously shows relative trends that the FACs-HT is more stable at higher excitation than FACs-LT and that has similar behavior for EL and PL in both materials. Collectively, our PL and EL measurements on four different perovskite films (FACs-LT, FACs-HT, MA, and MACs) allow us to investigate impact of carrier density profile on the rate of phase segregation (EL and PL results for MACs are shown in Figure S7). Recently, Barker et. al examined the impact of carrier density profile on rate of phase segregation and discovered that higher excitation photon energies (i.e. blue light) gives faster rate of phase segregation for a 13


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comparable absorbed photon flux in MAPb(I0.4,Br0.6)3.38 We add to this picture by comparing the rate of phase segregation under illumination with the rate of phase segregation upon current injection. Under one-sided illumination, both electron and holes have similar profiles with highest concentration close to the illuminated side, while upon current injection the electron and hole concentrations will each be highest closest their selective contacts. Our results confirm that differences in carrier density depth profiles affect the rate of phase segregation, yet we observe different behavior for different compositions. Table 1 presents the extreme slope of vs. time data for each of the four films studied, where this slope serves as a quantitative metric to examine phase segregation rate. We find that for MA and MACs, the rate of decay is much faster for photoexcitation than for current injection. Contrarily, the FACs LT and HT films (at the 8 or 32 Suns required to induce phase segregation), show slightly faster decay for current injection than for photoexcitation. Understanding the reasons behind the different relative rates of phase segregation is a complex topic since these experiments were conducted at different excitations (1 Sun, 8 Suns, and 16 Suns), with different excitation mechanisms (EL and PL), and on different materials with distinct mass and electronic transport properties. However, we speculate these observed differences between relative rates of phase segregation in FACs films and the MA/MACs films to be primarily associated with the differences in carrier diffusion lengths of the materials and differences in total absorbed/injected carrier density in the films. We acknowledge that the carrier concentrations (photoexcitation for the PL experiment and injection for the EL experiment) will be about 20-25% different due to non-unity absorptivity of these films at 532 nm. We suspect the fastest phase segregation will occur when the local carrier density of both holes and electrons is the highest (inducing the greatest lattice strain). MA films have shorter carrier diffusion lengths than FACs films,39 suggesting the highest local carrier 14


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densities of both carriers occur at the illuminated side of the substrate during PL measurements resulting in greater lattice strain than with current injection and faster phase segregation. In contrast, FACs films have longer carrier diffusion lengths which result in lower carrier density gradients within the films, causing faster phase segregation in EL due to slightly higher total carrier injection (due to non-unity absorptivity in PL experiments). Table 1. Phase Segregation Rates, d/dt [meV/s] MA

MACs FACs-LT FACs-HT

PL -6.27 -6.93

-0.28

-0.27

EL -1.22 -0.69

-0.58

-1.25

Implications of Phase Segregation for Tandem Photovoltaics. Since the development of monolithic tandem solar cells is one of the key reasons for interest in the large-bandgap hybrid perovskites, we turned our investigation to evaluate the impact of phase segregation on twoterminal monolithic tandem solar cell performance. One potential pitfall is that phase segregation can result in carrier trapping in I-rich domains, which could reduce the current output of the perovskite cell26-28. This is especially problematic in a two-terminal tandem which requires current matching. Additionally, a phase segregating top cell could also be detrimental due to changing absorption behavior, which could result in a changed transmitted photon flux and thus altered current output of the bottom cell. We evaluated both criteria (current output of top cell and transmitted photon flux to bottom cell) for MA, MACs, and FACs-HT by measuring current output with time of the perovskite top cell as well as current output with time for a CIGS bottom cell with perovskite a film filter.

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Results from these experiments are shown in Figure 4. Figure 4 a-c show the normalized PL emission spectra of FACs-HT, MA, and MACs films before (dashed line) and after (solid line) five minutes of illumination with Newport AAA Solar Simulator (the exact films used for the experiment with the CIGS device shown in Figure 4 e). As expected from previous PL results with 532nm laser excitation (Figure S7 and Figure 3b), FACs-HT exhibits no PL peak shift while both MA and MACs show red-shifts. The normalized current output at maximum power point for large-bandgap perovskite devices are presented in Figure 4 d. FACs shows stabilized current output at about 3% above starting value, while MA shows current output decreasing to below 85% the starting value before five minutes. These results are expected and agree with previous findings21,

26-28

: MA is phase segregating which results in decreased current collection upon

carrier confinement in I-rich domains contrasting the phase-stable FACs. Remarkably, the MACs device shows stable current output at ~3% below initial current (as observed in our previous work)5 despite PL and EL results showing similar phase segregation behavior as MA (see Figure 4 a-c and Figure S7). For MACs, I-rich domains clearly form as indicated by PL, but this does not have a significant impact on carrier collection.

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Figure 4. Device current with time during perovskite phase segregation. Normalized photoluminescence spectra of (a) (FA0.83,Cs0.17)Pb(I0.66,Br0.34)3, (b) MAPb(I0.6,Br0.4)3, and (c) (MA0.9,Cs0.1)Pb(I0.6,Br0.4)3 films

before (dashed) and after (solid) exposure to AAA solar

simulator in an air-free quartz assembly. (d) Normalized maximum power operating current at fixed

voltage

over

five

minutes

of

MAPb(I0.6,Br0.4)3,

high

temperature

(FA0.83,Cs0.17)Pb(I0.66,Br0.34)3, and (MA0.9,Cs0.1)Pb(I0.6,Br0.4)3 devices in nitrogen. (e) Normalized CIGS short circuit current over time while filtered by the exact films of (a-c), where the control is with the empty air-free quartz assembly as a filter.

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The JSC as a function of time from a 12.1% efficient solution-processed CIGS cell (see reference for fabrication details, and Figure S11 for JV performance data)40 with perovskite film filter is shown in Figure 4 e. This data exemplifies the impact of changing transmitted photon flux to bottom cell on bottom cell current output. In this experiment, the MA and MACs film filters are phase segregating, yet the impact of the changing transmissivity is insignificant. The inset of Figure 4 e shows a JSC decrease of 1 cm2. J. Phys. Chem. Lett. 2016, 7, 161-166. (14) Todorov, T.; Gershon, T; Gunawan, O.; Sturdevant, C.; Guha, S. Perovskite-Kesterite Monolithic Tandem Solar Cells With High Open-Circuit Voltage. Appl. Phys. Lett. 2015, 105, 173902. (15) Todorov, T.; Gershon, T.; Gunawan, O.; Lee, Y. S.; Sturdevant, C.; Chang, L.-Y.; Guha, S. Monolithic Perovskite-CIGS Tandem Solar Cells via In Situ Band Gap Engineering. Adv. Energy Mater. 2015, 5, 1500799. (16) Uhl, A. R.; Yang, Z.; Jen, A. K.-Y.; Hillhouse, H. W. Solution-Processed Chalcopyrite– Perovskite Tandem Solar Cells in Bandgap-Matched Two- and Four-Terminal Architectures. J. Mater. Chem. A 2017, 5, 3215-3220. (17) Eperon, G. E.; Leijtens, T.; Bush, K. A.; Prasanna, R.; Green, T.; Wang, J. T. W.; McMeekin, D. P.; Volonakis, G.; Milot, R. L.; May, R. et al. Perovskite-Perovskite Tandem Photovoltaics With Optimized Band Gaps. Science 2016, 354, 861-865. (18) Kayes, B. M.; Zhang, L.; Twist, R.; Ding, I. K.; Higashi, G. S. Flexible Thin-Film Tandem Solar Cells With >30% Efficiency. IEEE Journal of Photovoltaics 2014, 4, 729-733. 22


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(19) Chen, B.; Zheng, X.; Bai, Y.; Padture, N. P.; Huang, J. Progress in Tandem Solar Cells Based on Hybrid Organic–Inorganic Perovskites. Adv. Energy Mater. 2017, 7, 1602400. (20) 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.; Gratzel, M. Incorporation of Rubidium Cations into Perovskite Solar Cells Improves Photovoltaic Performance. Science 2016, 354, 206-209. (21) McMeekin, D. P.; Sadoughi, G.; Rehman, W.; Eperon, G. E.; Saliba, M.; Horantner, M. T.; Haghighirad, A.; Sakai, N.; Korte, L.; Rech, B. et al. A Mixed-Cation Lead Mixed-Halide Perovskite Absorber for Tandem Solar Cells. Science 2016, 351, 151-155. (22) Yu, Y.; Wang, C.; Grice, C. R.; Shrestha, N.; Zhao, D.; Liao, W.; Guan, L.; Awni, R. A.; Meng, W.; Cimaroli, A. J. et al. Synergistic Effects of Lead Thiocyanate Additive and Solvent Annealing on the Performance of Wide-Bandgap Perovskite Solar Cells. ACS Energy Lett. 2017, 2, 1177-1182. (23) Hoke, E. T.; Slotcavage, D. J.; Dohner, E. R.; Bowring, A. R.; Karunadasa, H. I.; McGehee, M. D. Reversible Photo-Induced Trap Formation in Mixed-Halide Hybrid Perovskites for Photovoltaics. Chem. Sci. 2015, 6, 613-617. (24) Suarez, B.; Gonzalez-Pedro, V.; Ripolles, T. S.; Sanchez, R. S.; Otero, L.; Mora-Sero, I. Recombination Study of Combined Halides (Cl, Br, I) Perovskite Solar Cells. J. Phys. Chem. Lett. 2014, 5, 1628-1635.

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(25) Unger, E. L.; Kegelmann, L.; Suchan, K.; Sorell, D.; Korte, L.; Albrecht, S. Roadmap and Roadblocks for the Band Gap Tunability of Metal Halide Perovskites. J. Mater. Chem. A 2017, 5, 11401-11409. (26) Yang, Z. B.; Rajagopal, A.; Jo, S. B.; Chueh, C. C.; Williams, S.; Huang, C. C.; Katahara, J. K.; Hillhouse, H. W.; Jen, A. K. Y. Stabilized Wide Bandgap Perovskite Solar Cells by Tin Substitution. Nano Lett. 2016, 16, 7739-7747. (27) Slotcavage, D. J.; Karunadasa, H. I.; McGehee, M. D. Light-Induced Phase Segregation in Halide-Perovskite Absorbers. ACS Energy Lett. 2016, 1, 1199-1205. (28) Hu, M.; Bi, C.; Yuan, Y.; Bai, Y.; Huang, a. J. Stabilized Wide Bandgap MAPbBrxI3–x Perovskite by Enhanced Grain Size and Improved Crystallinity. Adv. Sci. 2016, 3, 1500301. (29) Duong, T.; Wu, Y.; Shen, H.; Peng, J.; Fu, X.; Jacobs, D.; Wang, E.-C.; Kho, T. C.; Fong, K. C.; Stocks, M. et al. Rubidium Multication Perovskite With Optimized Bandgap for Perovskite-Silicon Tandem with over 26% Efficiency. Adv. Energy Mater. 2017, 7, 1700228. (30) Beal, R. E.; Slotcavage, D. J.; Leijtens, T.; Bowring, A. R.; Belisle, R. A.; Nguyen, W. H.; Burkhard, G. F.; Hoke, E. T.; McGehee, M. D. Cesium Lead Halide Perovskites With Improved Stability for Tandem Solar Cells. J. Phys. Chem. Lett. 2016, 7, 746-51. (31) Forgacs, D.; Perez-del-Rey, D.; Avila, J.; Momblona, C.; Gil-Escrig, L.; Danekamp, B.; Sessolo, M.; Bolink, a. H. J. Efficient Wide Band Gap Double Cation – Double Halide Perovskite Solar Cells. J. Mater. Chem. A 2017, 5, 3203-3207. (32) Ndione, P. F.; Li, Z.; Zhu, K. Effects of Alloying on the Optical Properties of Organic– Inorganic Lead Halide Perovskite Thin Films. J. Mater. Chem. C 2016, 4, 7775-7782. 24


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(33) Rehman, W.; McMeekin, D. P.; Patel, J. B.; Milot, R. L.; Johnston, M. B.; Snaith, H. J.; Herz, L. M. Photovoltaic Mixed-Cation Lead Mixed-Halide Perovskites: Links Between Crystallinity, Photo-Stability and Electronic Properties. Energy Environ. Sci. 2017, 10, 361-369. (34) Bischak, C. G.; Hetherington, C. L.; Wu, H.; Aloni, S.; Ogletree, D. F.; Limmer, D. T.; Ginsberg, N. S. Origin of Reversible Photoinduced Phase Separation in Hybrid Perovskites. Nano Lett. 2017, 17, 1028-1033. (35) Yoon, S. J.; Draguta, S.; Manser, J. S.; Sharia, O.; Schneider, W. F.; Kuno, M.; Kamat, P. V. Tracking Iodide and Bromide Ion Segregation in Mixed Halide Lead Perovskites During Photoirradiation. ACS Energy Lett. 2016, 1, 290-296. (36) Marongiu, D.; Chang, X.; Sarritzu, V.; Sestu, N.; Pau, R.; Lehmann, A. G.; Mattoni, A.; Quochi, F.; Saba, M.; Mura, A. et al. Self-Assembled Lead Halide Perovskite Nanocrystals in a Perovskite Matrix. ACS Energy Lett. 2017, 2, 769-775. (37) Karimata, I.; Kobori, Y.; Tachikawa, T. Direct Observation of Charge Collection at Nanometer-Scale Iodide-Rich Perovskites during Halide Exchange Reaction on CH3NH3PbBr3. J. Phys. Chem. Lett. 2017, 8, 1724-1728. (38) Barker, A. J.; Sadhanala, A.; Deschler, F.; Gandini, M.; Senanayak, S. P.; Pearce, P. M.; Mosconi, E.; Pearson, A. J.; Wu, Y.; Srimath Kandada, A. R. et al. Defect-Assisted Photoinduced Halide Segregation in Mixed-Halide Perovskite Thin Films. ACS Energy Lett. 2017, 2, 1416-1424.

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(39) Stoddard, R. J.; Eickemeyer, F. T.; Katahara, J. K.; Hillhouse, H. W. Correlation Between Photoluminescence and Carrier Transport and a Simple In-Situ Passivation Method for HighBandgap Hybrid Perovskites. J. Phys. Chem. Lett. 2017, 8, 3289-3298. (40) Uhl, A.; Katahara, J.; Hillhouse, H. Molecular-Ink Route to 13.0% Efficient Low-Bandgap CuIn(S,Se)2 and 14.7% efficient Cu(In,Ga)(S,Se)2 solar cells. Energy Environ. Sci. 2016, 9, 130134.

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Figure 1. Top-view SEM micrographs of perovskite thin films: (a) MAPb(I0.4,Br0.6)3 annealed at 100°C for 10 minutes, (b) (FA0.83,Cs0.17)Pb(I0.66,Br0.33)3 annealed at 75 ˚C for 10 min, and (c) (FA0.83Cs0.17Pb(I0.66,Br0.33)3 annealed at 165 ˚C for 50 min. (d) corresponding XRD patterns of above shown films. (e) Normalized and smoothed high resolution data of (200) reflections of comparable samples with the same compositions as (d), with the inset showing the full-width at half-maximum (B) and the domain size d as determined by the Scherrer formula. 165x102mm (220 x 220 DPI)


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Figure 2. Electroluminescence and photoluminescence measurements over time. (a) Electroluminescence spectra at selected times of a MAPb(I0.6Br0.4)3 device being injected with a constant current density of 19.2 mA/cm2 for five minutes. (b) Mean electroluminescence and photoluminescence emission energy of devices utilizing MAPb(I0.6Br0.4)3 and low temperature annealed (FA0.83,Cs0.17)Pb(I0.34Br0.66)3 absorbers. (c) External electroluminescence and photoluminescence quantum yield from the same experiment. 83x140mm (300 x 300 DPI)


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Figure 3. Mean emission spectrum energy with time. The electroluminescence current density and photoluminescence illumination intensity is doubled every minute. (a) Low temperature annealed (FA0.83,Cs0.17)Pb(I0.66Br0.34)3 under 2, 4, and 8 suns equivalent current density or illumination intensity. (b) High temperature annealed (FA0.83,Cs0.17)Pb(I0.66Br0.34)3 under 2, 4, 8, 16 and 32 suns equivalent current density or illumination intensity. 83x87mm (300 x 300 DPI)


ACS Energy Letters

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Figure 4. Device current with time during perovskite phase segregation. Normalized photoluminescence spectra of (a) (FA0.83,Cs0.17)Pb(I0.66,Br0.34)3, (b) MAPb(I0.6,Br0.4)3, and (c) (MA0.9,Cs0.1)Pb(I0.6,Br0.4)3 films before (dashed) and after (solid) exposure to AAA solar simulator in an air-free quartz assembly. (d) Normalized maximum power operating current at fixed voltage over five minutes of MAPb(I0.6,Br0.4)3, high temperature (FA0.83,Cs0.17)Pb(I0.66,Br0.34)3, and (MA0.9,Cs0.1)Pb(I0.6,Br0.4)3 devices in nitrogen. (e) Normalized CIGS short circuit current over time while filtered by the exact films of (a-c), where the control is with the empty air-free quartz assembly as a filter. 83x140mm (300 x 300 DPI)


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