Enabling Flexible All-Perovskite Tandem Solar Cells

Article

Enabling Flexible All-Perovskite Tandem Solar Cells Axel F. Palmstrom, Giles E. Eperon, Tomas Leijtens, ..., Joseph M. Luther, Joseph J. Berry, David T. Moore [email protected] (A.F.P.) [email protected] (G.E.E.) [email protected] (T.L.) [email protected] (J.J.B.) [email protected] (D.T.M.)

HIGHLIGHTS Multijunction perovskite solar cells for lightweight, flexible, high-efficiency PVs A nucleation layer-based strategy for an improved recombination layer Cation tuning for wide-band-gap perovskite solar cells with a stable and high voltage Flexible and rigid all-perovskite tandems with PCEs of 21.3% and 23.1%, respectively Two advances that address the main challenges of all-perovskite two-terminal tandem solar cell fabrication are reported. First, a nucleation layer is used to enable high-quality atomic layer deposition-based recombination layers that reduce electronic losses. Second, cation tuning is used for wide-band-gap perovskite solar cells that produce high, stable voltages. Combining these advances allows us to fabricate tandem perovskite solar cells on both rigid and flexible plastic substrates that have high efficiency and promising stability.

Palmstrom et al., Joule 3, 1–12 August 21, 2019 ª 2019 Elsevier Inc. https://doi.org/10.1016/j.joule.2019.05.009

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Article

Enabling Flexible All-Perovskite Tandem Solar Cells Axel F. Palmstrom,1,7,* Giles E. Eperon,2,7,* Tomas Leijtens,2,6,7,* Rohit Prasanna,2,4 Severin N. Habisreutinger,2 William Nemeth,2 E. Ashley Gaulding,2 Sean P. Dunfield,3 Matthew Reese,1 Sanjini Nanayakkara,1 Taylor Moot,2 Je´re´mie Werner,5 Jun Liu,1 Bobby To,1 Steven T. Christensen,1 Michael D. McGehee,5 Maikel F.A.M. van Hest,1 Joseph M. Luther,2 Joseph J. Berry,1,8,* and David T. Moore2,*

SUMMARY

Context & Scale

Multijunction all-perovskite solar cells offer a route toward efficiencies of III-V materials at low cost by combining the advantages of low thermalization loss in multijunction architectures with the beneficial properties of perovskites— namely, low processing cost, high-throughput fabrication, and compatibility with flexible substrates. However, there are two main challenges for enabling high-efficiency tandems: (1) design of a recombination layer to efficiently combine two perovskite subcells while also preventing bottom cell damage during top cell processing and (2) achieving high open-circuit voltage of the widegap subcell. Herein, we overcome both of these challenges. First, we demonstrate a nucleation layer consisting of an ultra-thin polymer with nucleophilic hydroxyl and amine functional groups for nucleating a conformal, low-conductivity aluminum zinc oxide layer by atomic layer deposition (ALD). This method enables ALD-grown recombination layers that reduce shunting as well as solvent degradation from solution processing on top of existing perovskite active layers. Next, we demonstrate a band-gap tuning strategy based on A-site cations of mismatched size (dimethylammonium and cesium) to enable a 1.7 eV perovskite with high, stable voltages. By combining these advances, we fabricate two-terminal all-perovskite tandem solar cells with 23.1% power conversion efficiency on rigid substrates and 21.3% on flexible plastic substrates.

Metal halide perovskites offer high-efficiency photovoltaics at low fabrication costs. By stacking two layers of perovskite with complimentary band gaps, even greater sunlight-to-electricity conversion efficiencies can be reached. Two-terminal tandem architectures are currently limited by challenges in the recombination layer connecting the two perovskite materials and by insufficiently high and stable voltages produced by the widerband-gap perovskite. We developed a strategy for improved recombination layers through the incorporation of a nucleation surface for atomic layer deposition and a strategy to achieve stable, wide-gap perovskites with high voltage through cation composition tuning. Combining these advances, we fabricated allperovskite tandems on glass with 23.1% efficiency and flexible tandems on plastic with 21.3% efficiency—the most-efficient flexible, thin-film solar cells reported to date.

INTRODUCTION Multijunction photovoltaics represent a route to achieve higher power conversion efficiencies (PCEs) of sunlight to electricity than those based on just one lightabsorbing semiconductor. By stacking semiconductors with different band gaps in one device, thermal losses are minimized. The best multijunction devices have proven this concept with efficiencies of over 38% attained as opposed to the best single-junction devices of around 29%.1 However, these use extremely high-quality III-V semiconductors, requiring high temperatures and slow rates of growth, making them expensive. In recent years, there has been a surge of interest in incorporating halide perovskites into multijunction devices; these materials exhibit excellent properties for a photovoltaic material yet are inexpensive and can be fabricated very rapidly via solution or vapor-based methods.2–6 Combining these with silicon in tandem architectures has recently led to efficiencies as high as 28%, exceeding the singlejunction silicon record.1,7 However, the Si-perovskite tandem is still limited by the Si subcell with regards to rigidity, weight, cost, and throughput.8 Alternatively, a

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tandem device using only perovskite absorbers in a multijunction device would fully leverage the advantages of this material system, including lower cost, higher throughput, high specific power, and use of flexible substrates.8,9 All-perovskite tandems have been demonstrated with up to 21.0% PCE in a two-terminal monolithic architecture and up to 23.1% in a four-terminal mechanically stacked architecture.10–12 The two-terminal architecture (2T) represents unique challenges as it requires more intimate integration and better optical management of all the elements of the individual subcells and for the connecting recombination layer in particular. If these hurdles can be overcome, the 2T has a higher potential efficiency limit due to having fewer parasitic absorptive layers. Furthermore, 2T devices are potentially lighter weight and cheaper to manufacture due to the removal of two transparent contacts and potentially one substrate—all of which impact balance of systems costs.2,9,13,14 The state-of-the-art all-perovskite 2T tandems have been limited by several main factors: (1) shunting due to the use of a thick (100 nm) and conductive recombination layer of indium tin oxide (ITO), enabling lateral connectivity of shunting paths in each subcell,3,7 and (2) large voltage losses in the wide-gap subcell due to iodidebromide halide segregation, resulting in lower than ideal voltages for the tandem overall.15,16 Here, we develop strategies to overcome both of these issues and attain high-efficiency all-perovskite tandem devices. One of three approaches can be taken to fabricate two perovskite films monolithically: (1) orthogonal solvent systems can be used (such that the solvent system for the second perovskite film does not dissolve the first), (2) a solvent-free approach can be used to deposit the second perovskite layer (e.g., thermal evaporation), or (3) a sufficiently good barrier layer can be deposited to prevent solvents used in the second perovskite layer from reaching the first. In previous tandem designs, deposition of a thin tin oxide layer grown by atomic layer deposition (ALD) was used to prevent sputter damage to the bottom subcell, enabling sputter processing of ITO as a recombination layer.11 This ITO layer served a dual purpose as a solvent barrier; however, a layer of 100 nm thickness was necessary. While orthogonal solvent systems or vapor-phase approaches can also be used, these preclude leveraging the most common solvents (i.e., dimethylformamide [DMF], diemethylsulfoxide [DMSO]) that have enabled the most efficient perovskite devices.17,18 Here, we develop a method to enhance the barrier properties of the ALD layer by modifying the ALD growth surface with a nucleophilic polymer. This improves nucleation of the ALD-grown buffer layer, making it dense and conformal and allowing it to act as an effective barrier against both sputter and solvent damage in its own right, thus enabling use of non-orthogonal solvents for processing. We use semicrystalline aluminum-doped zinc oxide (AZO) deposited by ALD for this buffer layer. This allows us more design flexibility, including reduction of the sputtered transparent conducting oxide (TCO) thickness, enabling control over the lateral resistivity of the recombination layer. We increase the lateral resistivity to reduce the likelihood of electrically connecting any shunt paths in either subcell and thus improve the device fill factor. In this study, FA0.75Cs0.25Sn0.5Pb0.5I3, with a band gap of 1.27 eV, is used as our lowband-gap perovskite.19 Restricting the tin content to 50% enables the low-gap cell to be almost as robust to atmospheric exposure as pure Pb devices yet still shows excellent efficiency.20 The most ideal band-gap energy for the complementary wide-band-gap subcell in a tandem is 1.7–1.85 eV, as shown by Hoerantner et al.13 Previous all-perovskite

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1Center

for Materials Science, National Renewable Energy Laboratory, 15013 Denver West Parkway, Golden, CO 80401, USA

2Center for Chemistry

and Nanoscience, National Renewable Energy Laboratory, 15013 Denver West Parkway, Golden, CO 80401, USA

3Materials

Science and Engineering Program, University of Colorado, Boulder, CO 80309, USA

4Department

of Materials Science and Engineering, Stanford University, 476 Lomita Mall, Stanford, CA 94305, USA

5Chemical

and Biological Engineering, University of Colorado, Boulder, CO 80309, USA

6Swift

Solar Inc., Golden Hills Road, Golden, CO 80401, USA

7These 8Lead

authors contributed equally

Contact

*Correspondence: [email protected] (A.F.P.), [email protected] (G.E.E.), [email protected] (T.L.), [email protected] (J.J.B.), [email protected] (D.T.M.) https://doi.org/10.1016/j.joule.2019.05.009

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tandems have typically used wide-gap subcells with high bromine content (30% in solution) to attain the desired band gaps, but these suffer from well-known halide segregation under illumination.15 At high bromine content, the segregation results in an unstable voltage as it is limited by the lower-band-gap iodide-rich regions that form.21 Here, we develop a novel method based on mixing large and small A-site cations to widen the band gap. This strategy permits us to achieve a wide band gap, but at lower bromine concentration, alleviating the halide segregation problem. Specifically, we use a dimethylammonium (DMA)-formamidinium-cesium A-site alloyed perovskite (FA0.6Cs0.3DMA0.1PbI2.4Br0.6) and find that we can fabricate 1.7 eV materials using only 20% bromine in solution. With this new concept, we observe reduced halide segregation, which in turn allows us to attain higher voltages and efficiencies that are stable under operation in excess of 1,000 h. By exploiting our improved ALD-based recombination layer, we are then able to process the low-gap material directly on the DMA-containing wide-gap cell. This approach permits us to solve the general problem of subcell integration and the specific problem of the wide-gap instability that plagued the previous implementations of all-perovskite tandems. This approach results in all-perovskite tandem solar cells with improved fill factor, voltage, and stable efficiencies of up to 23.1%. Using the flex-compatible nature of the recombination layer developed, we also build devices on flexible plastic substrates, attaining efficiencies of 21.3%. These results demonstrate the most efficient flexible thin-film solar cell of any non-III-V technology made to date, showing the potential for perovskite tandems to excel in applications where lightweight and flexible devices are required, and also providing a high-value early market for perovskite photovoltaics, such as unmanned aerial vehicle or portable photovoltaic markets where there is a premium on high specific power.9 The use of flexible substrates is crucial in particular, as this is one of the most important stepping stones toward massive scalability via roll-toroll processing.

RESULTS AND DISCUSSION Recombination Layer Carrier recombination between the subcells of a 2T device is a particular challenge resulting from the combination of electrical and practical considerations that must be realized; specifically, it must enable the appropriate device physics of the tandem with minimal loss while enabling material integration of top and bottom cells. The ideal recombination layer enables voltage addition between the perovskite subcells, results in minimal optical losses (especially in the nearinfrared, where the rear low-gap cell absorbs), and exhibits maximum out-of-plane conduction with minimal lateral conduction. These parameters can be achieved with conformal, ultra-thin oxide films. One method for thin-film oxide deposition on top of perovskite active layers is low-temperature ALD. Low-temperature (85 C) ALD enables metal oxide deposition on top of a perovskite active layer without damage to the perovskite material. Previously reported recombination layers used thin-film, ALD-grown oxides that were employed to effectively form sputter barriers on top of perovskite materials to enable subsequent sputter deposition.22 The wide-gap perovskite stack is fabricated in an inverted p-i-n architecture, terminated by an n-type C60 layer; the ALD-grown recombination layer is deposited on top of the C60 film.11 The poor barrier properties of ALD-grown metal oxides on bare C60 necessitate the use of thick (100 nm) sputtered TCO layers to provide adequate solvent resistance for processing of the second subcell. For example, we observed that perovskite films protected with just C60 (30 nm) and

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Figure 1. Nucleated AZO Atomic Layer Deposition (A) Schematic depicting AZO growth on C 60 and PEIE-treated C 60 surfaces. Transmission electron microscopy images of (B) 5 nm C 60 /4 nm AZO and (C) 5 nm C 60 /PEIE/4 nm AZO showing differences in AZO structure at the C60 interface with fast Fourier transform insets to highlight differences in crystallinity. (D) Water vapor transmission rate of 25 nm AZO as a function of time grown on bare C 60 and C 60 functionalized with a PEIE nucleation layer; the average WVTR values are 2.8 3 10 1 and 2.9 3 10 2 g/(m 2 day), respectively. (E) FA 0.6 Cs 0.3 DMA 0.1 PbI 2.4 Br 0.6 perovskite sample split in two coated with 30 nm C 60 /25 nm AZO on the left and 30 nm C 60 /PEIE/25 nm AZO on the right. (F) The same perovskite sample after 60 s of dimethylformamide (DMF) exposure.

ALD-deposited AZO (25 nm) are susceptible to solvent (both water and DMF) degradation on the timescale of seconds (Figures 1E and 1F). Similar results are observed with other ALD oxides on C60, including TiO2, Al2O3, and SnOx (Figure S1). Unlike ALD growth on inorganic surfaces, ALD precursors can diffuse into polymer or small-molecule substrates, resulting in subsurface growth. When the organic is chemically inert to the ALD metal-organic, a rough and diffuse interface often forms; nucleophilic functional groups reactive to metalorganics (e.g., hydroxyl, carbonyl, or amine) can increase the abruptness of the organic/ALD metal oxide interface.23 We suspect the weak barrier properties of the C60/AZO stack are the result of poor nucleation of ALD-grown metal oxide films on the chemically inert C60 surface, resulting in porous, non-continuous films. We therefore aim to improve the ALD growth behavior on C60 using an ultra-thin (1 nm) nucleation layer in the form of poly(ethylenimine) ethoxylated (PEIE)—a polymer chain containing both hydroxyl and amine functional groups. These nucleophilic functional groups are ideal for initiating nucleation and reaction with the metalorganic precursors used for metal oxide ALD. The 1 nm PEIE thickness is measured by spectroscopic ellipsometry, although it is likely non-continuous. The two growth modes of AZO on C60 with and without PEIE treatment are depicted in Figure 1A. Despite potential non-continuity in the PEIE, we observe significant differences in AZO nucleation under these two conditions. Figures 1B and 1C show transmission electron microscopy images of C60/4 nm AZO and C60/PEIE/4 nm AZO, illustrating the effect of PEIE on AZO growth near the C60 interface. AZO films grown on C60 nucleate as small crystallites, whereas AZO films grown on PEIE form a combination

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of larger, crystalline sheets as well as crystallites. Fast Fourier transform analysis of the TEM images (Figure S2) shows a higher degree of order with PEIE-nucleated AZO, consistent with the observation of larger crystalline AZO sheets. Further nucleation differences are observed through atomic force microscopy (AFM), as shown in Figure S3. While 4 nm AZO films grown on C60 and C60/PEIE have a similar rootmean-square surface roughness of 1.3 G 0.1 nm and 1.5 G 0.2 nm, respectively, the phase angle shows differences in surface coverage. The phase angle represents the tip response to interactions with the sample, and a uniform surface has a continuous phase-angle response. The phase angle of 4 nm AZO on C60 (Figure S3C) shows non-uniformity with islands surrounded by darker regions that we believe to be exposed C60. In contrast, the phase angle of PEIE-nucleated AZO (Figure S3D) is uniform, indicating more complete coverage at thin AZO thicknesses. We thus infer that the surface coverage depends on the mode of nucleation, with the PEIEinduced nucleation providing superior coverage. Next, we investigate the effect these nucleation properties have on device-thickness (25 nm) films. PEIE nucleation decreases the water vapor transmission rate (WVTR) nearly a full order of magnitude from 2.8 3 10 1gm 2day 1 with no nucleation layer to 2.9 3 10 2gm 2day 1 with PEIE (Figure 1D) and substantially improves barrier properties against DMF permeation for ALD-grown AZO (Figures 1E and 1F) as well as a broader range of ALD systems, including TiO2, Al2O3, and SnOx (Figure S1). The sheet resistance, measured by four-point probe, of the 30 nm C60/25 nm AZO stack decreases from 4.1 3 109 G 1.7 3 109U/, to 9.2 3 108 G 4.7 3 108U/, with PEIE treatment of the C60 surface. No significant change in the AZO film thickness was observed by spectroscopic ellipsometry when AZO was deposited in these two architectures, and PEIE itself does not offer any significant barrier properties or lateral electrical conductivity, suggesting the aforementioned differences arise from material quality rather than the film thickness. This nucleation strategy enables improved barrier layers from a range of ALD-grown oxides on C60 and should serve as a successful platform for the monolithic integration of two or more perovskite absorbers. Band-Gap Widening via Lattice Strain For wide-band-gap (1.7–1.85 eV) perovskite devices, stabilized voltages are limited by photoinduced halide segregation to around 1.15 V—far off the ideal 1.3–1.4 V.16 Often, compositions that are able to generate a high initial voltage suffer rapid voltage loss (see, for example, Figure S4), likely due to this phenomenon. With less bromine alloyed into a pure iodide film, halide segregation has been shown to be less severe.24,25 High voltages have been attained with bromine-free compositions, though these do not offer the correct band gap for tandem incorporation.26,27 It would thus be highly beneficial for a new strategy to exist for widening band gap with a reduced amount of bromine. Previously shown mechanisms for widening the band gap include the introduction of strain into the lattice, inducing tilting of the BX6 octahedra, and fabrication of lower-dimensional perovskite-like structures.28–31 The ions that can be incorporated into the ABX3 perovskite lattice can be predicted using the Goldschmidt tolerance factor. There are a small number of known A-site cations that fit within a lead halide lattice that give the tolerance factor of 0.8–1.0 necessary for a stable perovskite phase.32 The cations are most commonly Cs, methylammonium, and formamidinium. In contrast, the larger guanidinium cation has an ionic radius too large to allow the formation of a stable perovskite phase by itself. However, Hillhouse and co-workers recently demonstrated a rather surprising result. Specifically, Hillhouse and co-workers showed that some guanidinium can

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Figure 2. Stable Wide-Band-Gap Subcells via Cation Tuning (A) UV-vis absorbance spectra of perovskite films on glass made with increasing DMA percentage of the A-site, with (top) 0% bromine and (below) 20% Br. DMA addition was compensated in an equimolar manner with addition of Cs. (B) Band gap extracted from Tauc plots as a function of DMA percentage for 0% and 20% Br. The region shaded blue represents the band gaps relevant for wide-gap subcells in perovskite tandems. (C) Change in PL emission maximum as a function of time under illumination with a 1 mW 532 nm laser, spot size 1 mm 2 , for perovskite films with varying Br and DMA percentages. (D) Mobility and charge-carrier lifetime extracted from time-resolved microwave conductivity as a function of DMA percentage for films with similar band gap (attained by varying Br percentage). (E) Current-voltage characteristics of champion DMA-containing device (10% DMA, 20% Br). (F) Long-term stability of perovskite devices with varying DMA and Br percentages. Devices were held under load and constant illumination, in either air or N 2 environments as denoted in the caption. Note the change in axis scale at 260 h. Devices in N 2 had the architecture ITO/PolyTPD/PFN-Br/ perovskite/LiF/C60/BCP/Au, and the composition chosen was 10% DMA to correlate with the champion tandem (see later). The 10% and 12% DMA device data are normalized at 50 h.

be incorporated into the lattice when an equimolar amount of A-site cations with a smaller ionic radius are simultaneously incorporated. In this case, they show that Cs has a sufficiently small ionic radius that can compensate for the large ionic radius of guanidinium, allowing the formation of a black-phase perovskite with a strained lattice.31 Interestingly, it appeared that using this compensation approach to manipulate the lattice in this way exhibited a slight increase in band gap, which was noteworthy because adding large cations typically expands the lattice, having the effect of decreasing the band gap.33 This concept of compensating the addition of a large cation with a small one to maintain a perovskite phase is a promising route to increase the band gap without adding more bromine. We found that compensating for large A-site cations with cesium worked to increase the band gap for a range of larger cations, including acetamidinium, DMA, and guanidinium (see Figure S5), which all appeared to increase the band gap by a similar amount for a given fraction of the large cation. We found that films that used DMA as the large A-site cation were able to incorporate more of the large cation before forming a non-photoactive phase, so we continued with this cation as the most promising option. Our initial X-ray examination (Figure S6) provides insufficient evidence favoring any one of the three hypotheses mentioned above, while the resulting changes in optical and device metrics are quite clear. We initially tested whether we could attain wide enough band gaps for tandems using no bromine at all. Figure 2A shows the absorption of FA0.725(1 x)Cs0.2(1 x)+0.5xDMA0.5xPbI3 with increasing DMA percentage, expressed

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as a fraction of the total A site. We note that here and hereafter, the formulae used for the composition refer to the composition of the precursor solution. We started with a composition of 72.5% formamidinium, 27.5% cesium, known to be a stable and efficient composition.34 As we increased the percentage of DMA, we increased the Cs by an equimolar amount to compensate for the larger ionic radius of DMA. As the percentage of DMA increased, the absorption onset shifted to a shorter wavelength. Surprisingly, we found that we could add up to 50% DMA before films became colorless—and 40%–50% DMA enabled materials with an 1.7 eV band gap. However, X-ray powder diffraction showed that at DMA percentages this high, the majority phase was non-perovskite (see Figure S7). We found that adding a small amount of bromine facilitated the formation of black-phase perovskite (Figure S7) while also increasing the band gap. We targeted 1.7 eV as the lowest optimal band gap and found compositions where the black phase was a majority with the correct band gap existed between 0% DMA / 30% Br and 25% DMA / 10% Br. Band gap tuning by increasing the percentage of DMA and Cs with a fixed amount of Br (20%) is shown in Figure 2B. We found that tuning the band gap with DMA in this way did not have a major effect on the perovskite film morphology (see Figure S8). To test whether reduction of Br, enabled by increasing percentage of DMA and Cs, reduced the halide segregation relative to films with similar band gaps, we carried out photoluminescence measurements under constant laser illumination. As shown in Figure 2C, we found that the films with reduced Br showed a less-severe change in the photoluminescence (PL) peak maximum over time, suggesting that halide segregation has indeed been reduced. A concern when adding large amounts of bulky organic, insulating cations such as DMA into the lattice is that the charge-carrier mobility may suffer as a result. To probe whether this would limit device performance, we carried out time-resolved microwave conductivity measurements and found the lifetime of charge carriers in films of compositions with similar band gap and varying DMA/Br percentage. We show extracted mobility and lifetime as a function of DMA percentage in Figure 2C. Indeed, addition of DMA does decrease charge-carrier mobility, dropping from 27 to 8 cm2/Vs. This could be due to the inclusion of bulky organic species or a change in the band structure. However, we also observed that upon increasing DMA from 0% to 12%, the lifetime shows a large increase, before dropping as DMA percentage increases further.35 Combining the lifetime and mobility data, a composition with 12% DMA would have a higher mobility-lifetime product than that with no DMA, making it a superior material for solar cells. We postulate that this may be related to a reduction of defect density with incorporation of the large cations, as suggested by Hillhouse et al.31 We thus chose 10%–12% DMA incorporation as the optimum composition for a 1.7 eV material and fabricated PV devices with this perovskite (nominal composition of DMA0.1FA0.6Cs0.3PbI2.4Br0.6) in the p-i-n architecture ITO/PolyTPD/PFN-Br/ perovskite/LiF/C60/AZO/IZO/Ag. Ultra-thin PFN-Br and LiF act as effective recombination-suppressing interlayers, as demonstrated by Stolterfoht et al.36 We were able to attain champion devices with PCE 19% as shown in Figure 2D. Control devices with no DMA attain high efficiencies, but voltage was limited to 1.15 V (Figure S10). As expected, we found that increasing DMA content further caused a drop in PV performance (Figure S12). DMA-containing devices attained voltages of 1.2 V. Importantly, we showed that this voltage was stable after light soaking at open circuit for 10 min, demonstrating negligible impact of phase

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segregation on device performance (Figure S13). External quantum efficiency (EQE) for a DMA-containing perovskite cell is shown in the Supplemental Information (Figure S14). We carried out long-term stability tests on devices with varying DMA content. We used ALD-grown Al2O3 to act as an encapsulant and held devices under constant bias under illumination in air at a constant 25 C. The performance over time for varying DMA content is shown in Figure 2F. We observe that the 0% DMA control device drops to 80% of its initial efficiency within 50 h, having only 70% of initial efficiency after 250 h. Devices with 10%–12% DMA fared better, retaining over 90% of initial PCE over the same time. At 20% DMA, stability was worse than 0% DMA. Examination of the 20% DMA cells after aging showed that these devices had discolored and undergone a phase transition to a non-perovskite phase (Figure S15). We suspect that they are only meta-stable in the black phase at fabrication and over a long time they transitioned to a DMA-rich non-perovskite phase. The superior long-term stability of the 10%–12% DMA devices indicates that this material should be most suitable for the wide-band-gap cell in a tandem and superior to the 0% DMA material. To avoid the potential impact of atmospheric interactions, we also tested DMA-containing devices under illumination in nitrogen atmosphere. Over 1,000 h, we observe effectively no decrease in performance, as plotted in Figure 2F. Tandem Devices We built all-perovskite tandem devices by processing the low-gap subcell on top of a DMA-containing wide-gap subcell terminated with a recombination layer. As in our previous work, the low-gap device structure is PEDOT:PSS/perovskite/ C60/BCP/Ag, with a thick (850 nm) low-gap perovskite with composition FA0.75Cs0.25Sn0.5Pb0.5I3.19 Single-junction devices made using this composition were able to attain efficiencies of 16.5%, reaching short-circuit currents of over 31 mA/cm2 with 0.69 V, as shown in Figure S16. These efficiencies are not as high as those achieved with MAxFA1 xSn0.60Pb0.40I3 compositions but have more tractable stability challenges than the perovskites with higher tin content. Thus, we have chosen to use the more stable FA0.75Cs0.25Sn0.5Pb0.5I3 composition for our current work.19 As outlined, the ALD AZO layer is an effective solvent barrier to both water and DMF. We can then tailor the recombination layer with thin conductive oxides in conjunction with AZO (i.e., AZO/IZO) to enable rapid recombination and effective contact to the low-band-gap subcell, doing away with the thick TCO used in previous iterations. Maximizing sheet resistance while maintaining mobility permits this layer to provide an effective series connection between the two subcells consistent with the well-known device physics demand in tandems. Ideally, we aim to minimize lateral conductivity of the interlayer to prevent connection of shunt pathways that may be spatially dislocated in each subcell. One obvious way to do this is to reduce the thickness of the TCO layer. We found that devices with no sputtered TCO at all (the recombination layer consisted of AZO/PEDOT:PSS) suffered from large charge extraction barriers and resultantly low fill factors. This is presumably due to a reaction between the acidic PEDOT:PSS solution and the acid-sensitive AZO.37 We found that at least 10 nm of ITO or IZO was critical to prevent damage from the PEDOT and enable high fill factors when regular acidic PEDOT:PSS dispersions were used. However, using pH-neutral PEDOT:PSS dispersions enabled the fabrication of efficient tandems with just 5 nm of IZO. Such thin (5–15 nm) IZO layers, made with lower conductivity IZO (by increasing oxygen partial pressure during the sputter process), resulted in recombination layers with sheet resistances on the order

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Figure 3. Flexible and Rigid All-Perovskite Tandems (A) Cross-sectional scanning electron micrograph of all-perovskite tandem. (B) Current-voltage characteristics for an all-perovskite 2T tandem device, measured under AM1.5 illumination and ensuring minimal (<1.5%) mismatch between subcells. (C) Normalized external quantum efficiency spectra for the subcells of the same tandem. (D) Photo of flexible tandem device. (E) Current-voltage characteristics for an all-perovskite tandem fabricated on flexible PEN, including current-voltage data after performing 30 and 100 bends with 1.5 cm radius. (F) Maximum power point tracking for the rigid and flexible tandems over 30 s. Inset: photo of flexible tandem device. (G) Longer maximum power point tracking for a rigid tandem, under AM1.5 illumination in nitrogen for 13 h. (H) Long-term stability of encapsulated tandem device, held under load and constant illumination in air. Data are normalized at 75 h.

of 10–100 KU/sq. This high sheet resistance inhibited lateral conduction to shunt pathways while the 5–15 nm of resistive TCO layers had minimal effect on the series resistance (out-of-plane charge transport), resulting in a high fill factor and voltage. The full tandem device stack is shown in the Supplemental Information (Figure S17). We show a cross-sectional scanning electron micrograph of a tandem device in Figure 3A and a TEM image with energy-dispersive X-ray elemental mapping in Figure S18. Current-voltage characteristics for an all-perovskite tandem device are shown in Figure 3B. A 180-nm-thick layer of LiF was evaporated onto the glass side of the device as an antireflective coating. Optimization of the current matching in our subcells enabled us to reach 16.0 mA/cm2 in the tandem. Further, by combining our advancements in the wide-gap subcell and the recombination layer, we obtain an open circuit of 1.88 V and fill factor of 77%, which are a significant improvement over the previous generation of devices. Together, these performance parameters yield a PCE of 23.1%. In Figure 3C, we plot the normalized EQEs for the subcells of the tandem. The absolute reflectance spectrum is shown in Figure S22. The normalized EQEs were used with the solar simulator spectrum to calibrate our efficiency measurements, as done by certification laboratories and as discussed in the Supplemental Information. We also performed initial tests to validate the inherent stability of our tandem architecture—a 21.5% tandem was operated at maximum power point under 1 sun for 13 h with no observable drops in performance (Figure 3G), while another tandem device exhibited minimal degradation upon aging at 85 C in the dark in inert atmosphere for 50 h (Figure S20). We also encapsulated tandem devices using a combination of ALD-deposited alumina and an epoxy-sealed glass cover slide and held these devices in ambient conditions under constant illumination and constant load. We found that after 500 h of aging in this

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manner, our devices still retained above 90% of their initial efficiency (Figure 3H). This stability is competitive with the best perovskite-silicon tandems.7 These initial stability tests indicate that there is no rapid photoinduced halide segregation occurring in our wide-band-gap composition in the full tandem cell, nor is there rapid ion diffusion between the subcells. We also note here that while we made numerous devices in the 22%–23% range, the performance distribution is wide, and a nonnegligible fraction of devices were shunted (see Figures S23 and S24). We believe this to be due to the fact that the devices consist of many (5) spin-coated layers and are handled through many steps in heavily used multi-user facilities and not in a clean, dust-free environment. The innovations presented here (recombination layer and wide-band-gap absorber design), however, should yield highly reproducible devices with improved performances when coupled with improved deposition methods in clean environments; isolating these contributions is ongoing work. Fabrication currently takes place in multiple gloveboxes, evaporators, and fume hoods, so involves many steps requiring loading and unloading of substrates into holders, which is a process that can cause scratches and dust transfer onto the samples. Samples are left in non-dust-free glovebox atmospheres overnight in between stages of the fabrication process, during which time we expect that dust accumulates. By streamlining the process to involve fewer transfer and loading steps, and carrying out each step of the fabrication process immediately after the previous, we would expect a higher device yield. The highest-efficiency flexible thin-film devices (excluding III-V materials) to date are copper indium gallium selenide (CIGS), with an efficiency of 20.4%.38 Improving on this maximum efficiency for flexible devices with perovskites would be proof of their utility in low-cost, flexible, and lightweight applications. The highest-performing flexible perovskite single junction is 18.4%.39 Taking advantage of the low-temperature tandem processing and specifically the thin and flex-compatible recombination layer we developed, we fabricated the same all-perovskite tandem device stack on flexible polyethylene napthalate (PEN) substrates, with 200 nm high-conductivity sputtered IZO as the bottom TCO. Current-voltage characteristics for the flexible tandem are shown in Figure 3E. The flexible tandem on PEN delivers a steady-state PCE of 21.3%. This efficiency somewhat improved with 30 bends to 21.4% and only dropped to 20.4% after 100 bends at 1.5 cm bend radius. To the best of our knowledge, these results represent the most efficient 2T all-perovskite tandems reported to date and the most efficient among non-III-V (low cost) flexible photovoltaics. Conclusions We used a nucleation layer to fabricate a dense, conformal layer of aluminum zinc oxide that, in conjunction with a thin layer of ITO or IZO, acts as a very effective recombination layer, eliminating lateral shunt pathways and effectively stopping solvent and sputter damage when processing the rear subcell on top. Using a compensating A-site-based strategy of mixing large and small A-site cations to widen the band gap of perovskite materials, we were able to reduce the amount of bromine needed to attain the band gap needed for tandems, enabling a wide-band-gap composition resistant to photoinduced phase segregation with a high, stable voltage. Combining these advances, we fabricate 2T tandem all-perovskite devices with 23.1% efficiency on rigid and 21.3% on flexible substrates. These results demonstrate two advancements to overcome the challenges of monolithic all-perovskite tandems: a conformal recombination layer for facile coupling of two perovskites with non-orthogonal solvent systems in monolithic architectures and a new method for band-gap tuning of wide-gap perovskites with stable voltages. These strategies will help realize the potential of all-perovskite tandems as low-cost III-V alternatives.

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Please cite this article in press as: Palmstrom et al., Enabling Flexible All-Perovskite Tandem Solar Cells, Joule (2019), https://doi.org/10.1016/ j.joule.2019.05.009

EXPERIMENTAL PROCEDURES Full details of experimental procedures can be found in the Supplemental Information.

SUPPLEMENTAL INFORMATION Supplemental Information can be found online at https://doi.org/10.1016/j.joule. 2019.05.009.

ACKNOWLEDGMENTS We thank John Geisz and Ryan France for assistance with measuring EQE and current-voltage characteristics of tandem devices accurately, Ji Hao for 4-point probe measurements, and Al Hicks for graphics production. The work at the National Renewable Energy Laboratory is supported by the U.S. Department of Energy under Contract No. DE-AC36-08GO28308. The authors acknowledge support from the De-risking halide perovskite solar cells program of the National Center for Photovoltaics, funded by the U.S. Department of Energy, Office of Energy Efficiency and Renewable Energy, Solar Energy Technology Office. Tandem device fabrication was funded by NREL’s LDRD program. Contributions to the work from the University of Colorado were supported by the Office of Naval Research under award number N00014-17-1-2212. S.N.H. and E.A.G. acknowledge support from the NREL directors fellowship LDRD program. Efforts of T.M. and J.M.L. were provided by the Operational Energy Capability Improvement Fund. The views expressed in the article do not necessarily represent the views of the DOE or the U.S. Government. The U.S. Government retains and the publisher, by accepting the article for publication, acknowledges that the U.S. Government retains a nonexclusive, paid-up, irrevocable, worldwide license to publish or reproduce the published form of this work, or allow others to do so, for U.S. Government purposes.

AUTHOR CONTRIBUTIONS G.E.E. and D.T.M. led the efforts on developing the wide-gap material. T.L. and G.E.E. led work on tandem device construction and measurements; A.F.P., T.L., and J.J.B. led the development and characterization of the templated ALD process. R.P., W.N., T.M., J.W., S.T.C., and M.F.A.M.v.H. contributed to development of the device fabrication processes and carried out some stages of device fabrication. S.N.H., E.A.G., S.P.D., M.R., S.N., J.L., and B.T. carried out various device and material characterization techniques. J.L., M.D.M., A.F.P., G.E.E., T.L., J.J.B., and D.T.M. contributed to preparation of the manuscript.

DECLARATION OF INTERESTS G.E.E. and T.L. are founders of Swift Solar. M.D.M. is an advisor to Swift Solar. Received: February 17, 2019 Revised: April 5, 2019 Accepted: May 8, 2019 Published: May 16, 2019

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