Local halide heterogeneity drives surface wrinkling in mixed-halide wide-bandgap perovskites

doi:10.21203/rs.3.rs-4814295/v1

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Local halide heterogeneity drives surface wrinkling in mixed-halide wide-bandgap perovskites

https://doi.org/10.21203/rs.3.rs-4814295/v1

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Compositional heterogeneity in wide-bandgap (1.8 − 2.1 eV) mixed-halide perovskites is a key bottleneck in the processing of high-quality solution-processed thin films, and prevents their application in efficient multijunction solar cells. Notably, mixed-cation (formamidinium-methylammonium) wide-bandgap perovskite films are prone to form micrometer-scale wrinkles which can interfere with the smooth surfaces ideal for multijunction devices. Here, we study the formation dynamics of wrinkled mixed-halide perovskite films and its impact on the local composition and optoelectronic properties. We use in situ X-ray scattering during perovskite film formation to show that crystallization of bromide-rich perovskites precedes that of mixed-halide phases in wrinkled films cast using an antisolvent-based process. Using nanoscopic Xray fluorescence and hyperspectral photoluminescence imaging, we also demonstrate the formation of iodide- and bromide-rich phases in the wrinkled domains. This intrinsic spatial halide segregation results in an increased local bandgap disorder and Urbach energy. Morphological and compositional heterogeneity also aggravate the formation of sub-bandgap electronic defects reducing photostability and accelerating light-induced segregation of iodide and bromide ions in thin films and solar cells.

Physical sciences/Energy science and technology/Renewable energy/Solar energy/Photovoltaics/Solar cells

Physical sciences/Materials science/Materials for energy and catalysis/Solar cells

Mixed-halide wide-bandgap APbX₃ (A is a monovalent cation and X is a halide ion (iodide or bromide)) perovskite semiconductors are promising candidates for use in monolithic multijunction photovoltaic devices where the use of complementary absorber layers enables an increase in photovoltaic performance¹. Using compositions with bromide contents of 30% – 40% in the top-cell, monolithic tandem solar cells have been developed using perovskite and c-Si bottom sub-cells, which have exceeded 28% and 33% power conversion efficiency respectively^2,3. Wide-bandgap perovskites with even higher bromide contents (~ 60%) can potentially surpass the 40% efficiency mark in monolithic triple-junction devices^4–7.

However, solution-processed mixed-halide perovskites undergo complex crystallization routes that influence their compositional and morphological homogeneity^8–10. A key challenge is the presence of distinct wrinkled morphological domains in mixed-halide perovskite thin films^11–19. The formation of these wrinkles, with peak- and valley-like features, has been attributed to compressive stress in the film developed during crystallization^16,20. The morphological disorder and the residual stress created during film formation then reduce device performance and stability^21,22. The extent of wrinkling can be controlled by steering the crystallization rate and interactions with the substrate and solvent environments^14,16,18,23. Wrinkling is especially challenging in developing efficient perovskite-based multijunction devices where several ancillary thin (often less than 10 nm) films are needed for charge extraction, surface passivation, and as a recombination layer. To ensure conformal coverage and form continuous layers, surface planarity is important²⁴.

Solution-processed perovskites also suffer from stochastic compositional disorder resulting in local richness or deficiency in ionic species^25–31. This compositional disorder is particularly crucial in mixed-halide compositions where local variations in halide distribution affect the bandgap³². In turn, this variation in local halide concentration may influence the local chemical potential, defect distribution, and charge-carrier dynamics^32,33. Morphological disorder has previously been associated with compositional heterogeneity in cesium-containing perovskite thin films with bandgaps < 1.60 eV¹³. However, the correspondence between morphological and compositional heterogeneity in wide-bandgap compositions has not been investigated extensively. This correlation is especially critical because most wide-bandgap, mixed-halide compositions currently suffer from higher defect densities, resulting in lower radiative recombination efficiencies and a stronger tendency for ion migration^34,35. As a result, the role of such disorder in determining defect behavior and subsequent effect on photostability is also lacking.

Herein, we study perovskite thin films and devices using mixed-cation (formamidinium (FA) - methylammonium (MA)) mixed-halide (iodide - bromide) compositions prepared using an antisolvent-based deposition method. We first identify the role of the [FA]/[MA] and [I]/[Br] ratios in determining morphological disorder and their influence on crystallographic structure and orientation. Thereafter, we study the crystallization dynamics using synchrotron-based in situ structural characterization and identify key stages during processing where such heterogeneity begins to develop. Furthermore, using synchrotron-based nanoscopic X-ray fluorescence mapping, we correlate morphological heterogeneity to the compositional inhomogeneity in local halide distribution. This compositional heterogeneity causes local changes in the bandgap and photoluminescence energy. Finally, using sensitive photocurrent spectroscopy, we associate disordered compositions with a higher defect density and a corresponding instability under illumination due to light-induced halide segregation.

Mixed-cation lead mixed-halide perovskite (nominal composition (FA_1 − xMA_x)Pb(I_1 − yBr_y)₃) thin films (compositions hereby denoted as {x/y}) were prepared using an ethyl acetate antisolvent-based spin-coating route as detailed in the Methods section. Independently changing the methylammonium (x) or bromide (y) contents in the precursor solution alters the bandgap between 1.55 eV (25% MA, 0% Br) and 2.38 eV (75% MA, 100% Br), as determined by UV-vis-NIR spectroscopy (Fig. 1a and Supplementary Fig. 1)¹⁹. Calculations indicate that the ideal perovskite bandgap is 1.80–1.90 eV or 1.90–2.00 eV for the top-cell in all-perovskite tandem or perovskite-based triple-junction solar cells, respectively¹. This bandgap corresponds to a bromide content of 40% – 60% for the two applications (identified by shaded regions in Fig. 1a).

We used scanning electron microscopy (SEM) to characterize the surface morphology of the resulting perovskite thin films. Figures 1b – 1g show low magnification surface SEM images of films with compositions {x/y} = 0.25/0.40, 0.50/0.40, 0.75/0.40, 0.25/0.60, 0.50/0.60, and 0.75/0.60. The images show a smooth surface in films with lower MA and Br contents ({x/y} = 0.25/0.40, 0.50/0.40) whereas other compositions that are richer in MA and/or Br ({x/y} = 0.75/0.40, 0.25/0.60, 0.50/0.60, and 0.75/0.60) show additional morphological features on the surface. Specifically, we observe a heterogeneous distribution of large peaks and valleys on the surface of the films with {x/y} = 0.75/0.40, 0.50/0.60, and 0.75/0.60 whereas the film with {x/y} = 0.25/0.60 exhibits a morphology intermediate to the smooth and rough surfaces observed. We note that the appearance of morphological heterogeneity does not significantly affect SEM features at 1–2 µm or reveal voids in the film surface (Supplementary Fig. 2).

Figures 1h and 1i show three-dimensional atomic force microscopy (AFM) profiles of two films with compositions {x/y} = 0.25/0.40 (smooth), and with higher MA and Br content {x/y} = 0.50/0.60 (rough) to better visualize the distinct peak- and valley-like features observed in rough films. Figures 1j – 1o show line profiles of surface atomic force micrographs (Supplementary Figs. 3 and 4). The feature sizes in the smooth films ({x/y} = 0.25/0.40, 0.50/40, and 0.25/0.60) are on the order of 100 nm whereas compositions that yield heterogeneous films ({x/y} = 0.75/0.40, 0.50/0.60, and 0.75/0.60) exhibit feature sizes as large as 1.50–2.0 µm. The line cuts also show that the widths of the features in the more heterogeneous films are on the order of 2.0–5.0 µm^14,16.

We used synchrotron-based grazing-incidence wide-angle X-ray scattering (GIWAXS) measurements to understand the crystallographic properties of smooth and rough perovskite thin films (Fig. 2). A shift to a higher scattering vector (q) for the peak corresponding to the (100) plane (from q ≈ 1.02 to 1.06 Å⁻¹) with increasing MA or Br content suggests the lattice contraction upon the substitution of small cation/anion species (circular averages are shown in Supplementary Fig. 5)¹⁹. The peak corresponding to unreacted PbI₂ (q ≈ 0.9 Å⁻¹) remains unchanged across compositions.

Notably, we find that compositions that show roughness > 1 µm (Fig. 1) also show a preferential crystallographic orientation. For example, compositions yielding smooth films ({x/y} = 0.25/0.40, 0.50/0.40, and 0.25/0.50) show GIWAXS patterns (Fig. 2a, 2b, and 2d) with broad and complete ring-like features, consistent with randomly oriented crystallites (Fig. 2j). In contrast, films with roughness > 1 µm with composition ({x/y} = 0.75/0.50, 0.50/0.60, and 0.75/60) show GIWAXS patterns indicating preferential orientation of the (100) plane due to the appearance of intense spots on the ring at q ≈ 1.0 Å⁻¹ (Fig. 2f, 2h, 2i, 2j and 2l), especially in the out-of-plane direction¹⁸. Here, the inclusion of Br has a stronger influence in driving preferential orientation than MA (Fig. 2j, 2k and 2l). The change in preferential orientation does not arise from compositional changes only. For example, perovskite films of nominally identical compositions deposited using a two-step interdiffusion method show the random orientation of crystallites for compositions with high MA and Br content, as opposed to the preferential orientation observed in films coated using the one-step antisolvent method (Supplementary Fig. 6)^16,18,36.

We then performed synchrotron-based in situ GIWAXS measurements during film formation to characterize the crystallization dynamics of smooth and rough perovskite thin films. For these measurements, we processed the perovskite layer on an indium tin oxide (ITO)-coated glass substrate at a spin-coater placed in the X-ray beam path. We chose the compositions (MA-poor {x/y} = 0.25/0.60 and MA-rich {x/y} = 0.75/0.60) such that they yield different morphological outcomes ({x/y} = 0.25/0.60 yields a smooth film while {x/y} = 0.75/0.60 yields a rough film) but are similar in their nominal Br content. Figure 3 shows in situ GIWAXS patterns as a function of spin-coating time, focusing on the time range near the antisolvent casting at t = 25 s.

Both compositions show no scattering features in the initial stages (up to 21 s, Fig. 3a and 3b) of the spin-coating process when the wet film is largely disordered. Upon adding the antisolvent (Fig. 3c and 3d), the film with composition {x/y} = 0.75/60 immediately shows a strong scattering feature in the out-of-plane direction, corresponding to the (100) crystallographic plane of the perovskite phase. In contrast, the (100) feature is absent in the film with composition {x/y} = 0.25/60, appearing later at 29 s (Fig. 3e and 3f) and increasing in intensity thereafter (Fig. 3g and 3h). Here, as seen in Fig. 2g, the feature resembles a ring which is indicative of a random orientation of crystalline perovskite phases. We also observe similar crystallization dynamics in the characterization of other smooth ({x/y} = 0.25/0.50) and rough ({x/y} = 0.75/0.50) perovskite films (Supplementary Figs. 7 and 8).

These results indicate a faster crystallization and an earlier onset of preferential orientation of the MA-rich {x/y} = 0.75/60 thin film compared to the MA-poor {x/y} = 0.25/60 perovskite^31,37. This result agrees with prior observations of faster crystallization of Br-rich phases in a MA-based environment than in a FA-based environment³¹. The faster crystallization of Br-rich phases can also be attributed to weak interactions of bromide precursors and DMF, their corresponding poor solubility, and the poor stability of lead bromide complexes in solution, which collectively favor faster nucleation^9,38,39. Furthermore, it has been demonstrated that Br-based phases form the perovskite phase directly whereas I-based phases undergo crystallization through intermediate complexes³¹. MA-containing Br-based perovskite phases have also been shown to grow as highly oriented phases³⁷. As a result, based on the GIWAXS data showing faster crystallization and more preferential (100) orientation in films with roughness > 1 µm ({x/y} = 0.75/0.50 and 0.75/0.60), we propose that in an MA-rich environment, heterogeneous crystal nucleation leads to the formation of oriented bromide-rich perovskites immediately after antisolvent casting (Supplementary Fig. 9) followed by the incorporation of iodide-containing phases¹⁸, and that the heterogeneous crystallization causes the film wrinkling. The larger change in the scattering vector (Δq) for the perovskite composition {x/y} = 0.75/0.60 during the film-coating and thermal-annealing stages of the crystallization compared to that for {x/y} = 0.25/0.60, supports the hypothesis that Br- and I-rich phases crystallize at different rates in rough films (Supplementary Fig. 10).

Subsequently, we used synchrotron-based nanoscopic X-ray fluorescence (nano-XRF) microscopy to study the bulk composition of perovskite layers and observe the influence of morphological heterogeneity on compositional disorder (Fig. 4). The technique relies on elemental X-ray emission signatures to map the microscopic distribution of constituents^25,32. Here, we used the elemental map of Pb (Supplementary Fig. 11) as an indicator for film thickness since peak-like features increase the local film volume that corresponds to a higher X-ray fluorescence intensity.

We used nano-XRF mapping to analyze the spatial elemental distribution in smooth and rough perovskite thin films. For example, a film with composition {x/y} = 0.25/0.40, which yields a uniform surface (Fig. 1b), shows a homogeneous distribution of Pb, I, and Br (Fig. 4a). As a result, the iodide-to-bromide ratio in the film also shows a homogeneous distribution (Fig. 4b). We examined four highlighted regions (dashed boxes in Fig. 4b) to confirm that the iodide-to-bromide ratio is largely independent of the Pb content in the region, i.e., the local thickness of the film. Line cuts of iodide-to-bromide ratio overlapped with the Pb content (Fig. 4c) further confirm the lack of correspondence between the two quantities for a smooth perovskite film. Similar homogeneous distribution of ions is observed for other smooth films with compositions {x/y} = 0.50/0.40, 0.50/0.50, and 0.25/0.60, as shown in Supplementary Fig. 12.

In contrast, compositions that yield rough surfaces {x/y} = 0.75/0.40, 0.75/0.50, 0.50/0.60, 0.75/0.60, show a heterogeneous distribution of Pb, Br, and I (Fig. 4d and Supplementary Fig. 13). For example, the elemental maps (Pb, I, and Br) for the perovskite film with composition {x/y} = 0.50/0.60 (Fig. 4d) show the presence of peak- and valley-like features, as observed in the surface SEM characterization (Fig. 1). Furthermore, the iodide-to-bromide ratio map (Fig. 4e) shows regions of iodide richness and deficiencies which follow the peaks and valleys observed through the Pb elemental map (highlighted regions in Fig. 4d). Line cuts at four distinct locations of the film and two-dimensional maps (Fig. 4f, 4g and Supplementary Fig. 14) further confirm the positive correlation between increasing layer thickness and higher iodide concentration. Based on these observations, we hypothesize that compositional differences (MA- and Br-richness) influencing the rates of crystallization of iodide- and bromide-rich phases^31,32,38 drive the development of spatial halide heterogeneity during film crystallization⁴⁰, with iodide-rich phases forming in the peak-like regions of the films whereas bromide-containing phases predominantly crystallize at the valley-like regions. Compositional heterogeneity has previously been discussed in lead halide perovskites^25,32,41; however, this work is the first to report its association with surface morphology in wide-bandgap compositions relevant for multijunction devices.

Because of the link between perovskite bandgap and stoichiometry, we expect that compositional heterogeneity will also cause spatially heterogeneous bandgap disorder with regions of high iodide content exhibiting lower bandgaps, and high relative bromide content exhibiting wider bandgaps respectively^17,30,42. To test this hypothesis, we used hyperspectral photoluminescence (PL) imaging to characterize perovskite thin films’ emissive properties. We performed hyperspectral PL mapping of as-prepared films (referred to as “pristine”), excited with a mercury halide lamp in the 350–450 nm range at 130 mW cm^− 2 excitation intensity, with an acquisition time of 1 min. We then continuously illuminated the film for another 5 min., and measured again with a 1 min. acquisition time (referred to as “illuminated”). Figure 5 displays the PL emission peak wavelength maps from those hyperspectral PL imaging of perovskite films deposited on glass with compositions {x/y} = 0.25/0.40, 0.75/0.40, and 0.75/0.50 (Supplementary Fig. S15), illuminating the top surface (all films were encapsulated in the glovebox with UV-curable glue).

As we demonstrated with AFM and SEM earlier, films with low Br content {y = 0.40} and increasing MA content (x) show more pronounced wrinkles (Fig. 1). Using hyperspectral PL imaging, we also observe that wrinkling increases with MA content in low Br films {y = 0.40}. In a homogeneous film with low MA and Br content ({x/y} = 0.25/0.40), the PL emission wavelength is centered at ~ 690 nm (Fig. 5a, 5c, and 5d). The spatially averaged PL spectrum also shows a single peak (Fig. 5c), indicating that the emission is the same across the scanned area. On the other hand, for a rougher film with composition {x/y} = 0.75/0.40, the PL map and spatially averaged spectrum show two distinct emission peaks, indicative of wide- and narrow-bandgap domains within the scanned region (Fig. 5e, 5g, and 5h). The emission is primarily dominated by wide-bandgap domains, with a peak maximum at ~ 685 nm (Fig. 5g). Sparsely distributed, narrow-bandgap domains contribute to the second PL peak maximum at 730–750 nm. Additionally, we note that narrow-bandgap sites (λ_long) are comparatively brighter than wide-bandgap sites (λ_short) (Supplementary Fig. 16). We attribute the increased brightness of the lower bandgap regions to charge-carrier funneling from wider bandgap sites, as proposed previously⁴³. This emission heterogeneity is in good agreement with the compositional heterogeneity we observed through nano-XRF mapping (Supplementary Fig. 13), where we found the raised peak-like regions to be comparatively iodide-rich and correspondingly emit at lower energies/longer wavelengths.

Over the range of compositions studied, we observed that PL heterogeneity in the pristine film increases further with increasing the bromide content (Fig. 5i, 5k and 5l), in {x/y} = 0.75/0.50, with the maximum emission wavelength distribution broadening from 650 to 670 nm (Fig. 5l). The spatially averaged emission spectrum shows a short- (655 nm) and long-wavelength (745 nm) emission contribution, consistent with halide heterogeneity as a result of morphological disorder (Supplementary Fig. 13). However, the redshifted emission at the peak-like regions could also result from self-absorption or optical interference effects that occur at increased layer thickness^44,45. To rule out optical effects, we used time-of-flight secondary ion mass spectrometry (TOF-SIMS) mapping in the hyperspectral imaging area (Supplementary Fig. 17) and found that stronger long-wavelength emission regions correspond to local increase in the iodide-to-bromide ratio. This correlation is strong evidence that the local redshifts of emission wavelength and increase in PL intensity are due to the high iodide content in peak-like regions. Together, TOF-SIMS and hyperspectral mapping correlation verified that PL wavelength mapping can provide important qualitative information about local compositional heterogeneity in these films.

Subsequently, we continuously illuminated the same samples for 5 min. with blue (450 nm) light with an intensity of 130 mW cm^− 2 (Supplementary Fig. 18). This illumination was done in order to cause ion migration that is known to lead to light-induced halide segregation³⁴. In Figs. 5b, 5f and 5j, we show the spatially resolved PL redshift after 5 min. of illumination, calculated from Figs. 5a, 5e, 5i and Supplementary Fig. 18. In a compositionally homogeneous, smooth film with composition {x/y} = 0.25/0.40, continuous illumination does not affect the PL emission maximum (Fig. 5b and 5d), but only causes a slight broadening of the PL full-width at half-maximum (FWHM) from 35 nm to 45 nm (Fig. 5c). In contrast, films with heterogeneous PL emission prior to illumination ({x/y} = 0.75/0.40 and 0.75/0.50) show a redshift in the maximum emission wavelength by approximately 50–100 nm (Fig. 5f, 5h, 5j, and 5l). We note that the change in the emission wavelength after illumination is lower in the peak-like regions (Δλ ~ 10 nm) compared to the valley-like regions (Δλ ~ 50 nm) (Fig. 5f and 5j). This is because in peak-like areas the iodide-to-bromide ratio is already high before illumination, as we show in Supplementary Fig. 17 via correlated TOF-SIMS. For the perovskite film with composition {x/y} = 0.75/0.40, the initial emission with two peak features evolves to one broad redshifted emission peak indicating the formation of a broad distribution of emissive iodide-rich species (Fig. 5g). The illuminated {x/y} = 0.75/0.50 film shows a higher low-energy emission of 745 nm (Fig. 5k), and a weak emission from short-wavelength range of 655 nm.

We performed TOF-SIMS mapping of the perovskite film ({x/y} = 0.75/0.50) before and after continuous illumination to further understand the effect of ion migration as a function of morphological heterogeneity. By tracking the iodide content before and after illumination, we found that peak-like regions in the film undergo an increase in iodide content on the film surface (Supplementary Fig. 19a). In contrast, a homogeneous film ({x/y} = 0.25/0.50) shows a small and homogeneous increase in iodide content across the film surface upon continuous illumination (Supplementary Fig. 19b), indicating comparatively reduced halide migration as expected from hyperspectral PL imaging (Fig. 5b).

Finally, we fabricated perovskite solar cells (Fig. 6a) in an inverted (p-i-n) device architecture using [2-(9H-carbazol-9-yl)ethyl]phosphonic acid (2PACz) and C₆₀ as hole- and electron-transport layers, respectively (Supplementary Fig. 20). The transport layers were especially chosen due to their compatibility with multijunction device architectures and the conformal deposition of C₆₀ with thermal evaporation⁴⁶. Sensitive photocurrent spectroscopy was used to characterize defect dynamics in solar cells. The increase in defect density as a function of Br content has previously been studied^35,38,47, but the role of morphological and compositional disorder on electronic defects of wide-bandgap perovskites is relatively poorly understood. For a solar cell with composition {x/y} = 0.25/0.40 (Fig. 6b), the external quantum efficiency (EQE), calculated from the photocurrent spectrum, shows a flat above-bandgap (> 1.8 eV) EQE profile followed by an exponential drop (Urbach tail) at the bandgap-edge of the active layer. Following that, in the sub-bandgap 0.80–1.60 eV region, a clear EQE contribution of approx. 10^− 7 – 10^− 6 can be observed (at approx. 1.35 eV). A second, less prominent, feature can also be observed at lower energies (at approx. 1.0 eV). Previous work has shown that such sub-bandgap features originate from electronic defects near the perovskite/C₆₀ interface^48–50, and that changes in the sub-bandgap photocurrent intensity and photocurrent contribution at the Urbach tail upon prolonged illumination can be associated with light-induced halide segregation^47,51,52.

With increasing MA content, the optical bandgap increases, causing a blueshift in the EQE onset. At x = 0.75, additionally, the sub-bandgap contribution (< 1.5 eV) appears less pronounced, manifesting as a broad feature instead of a peak as observed in compositions {x/y} = 0.25/0.40 and 0.50/0.40. This broadening is likely related to the increased roughness resulting from morphological heterogeneity in such compositions (Fig. 1), which reduces optical interference by increasing light scattering.⁴⁸ The morphological heterogeneity as a function of increasing MA content also manifests as an increase in the Urbach energy, indicating an increase in band-edge energetic disorder.^32,53,54

These effects are much more pronounced upon increasing the bromide content ({x/y} = 0.75/0.50 and 0.75/0.60). Both compositions, for example, show that an additional shoulder appears on the Urbach tail (1.6–2.0 eV), likely indicating the presence of low-energy iodide-rich domains in the pristine device (Fig. 6c)^47,51,52. These observations agree with the non-uniform halide distribution, as observed using nano-XRF microscopy (Fig. 4) and hyperspectral PL imaging (Fig. 5). Such an increase in band-edge disorder has been shown to limit the open-circuit voltage in solar cells⁵⁵. Furthermore, the sub-bandgap photocurrent contribution in the 0.8–1.6 eV increases by approx. an order of magnitude for {x/y} = 0.75/0.60 composition, indicating a higher defect density as compared to the composition {x/y} = 0.25/0.40. The presence of sub-bandgap defects has been consistently associated with high non-radiative recombination in perovskite solar cells^50,56,57. Taken together, these observations indicate a higher degree of band-edge disorder and sub-bandgap defect density with increasing compositional/morphological disorder in wide-bandgap perovskites.

We used continuous illumination (532 nm, 1-Sun equivalent intensity, 5 min.) to induce defect migration and drive halide segregation in the solar cells⁵⁸. Halide migration in mixed-halide wide-bandgap perovskites can result in redshift of the band-edge, related to the formation of iodide-rich phases, and also cause an increase in the sub-bandgap defect density^34,47,51,52. Figures 6d – 6h show EQE spectra of solar cells following the photo-stress. In solar cells with low MA and low Br content ({x/y} = 0.25/0.40 and 0.50/0.40), that show smooth morphology and minimal compositional heterogeneity (Figs. 6d and 6e), light-induced changes to the EQE spectrum are minimal, consistent with our observations from hyperspectral PL and TOF-SIMS imaging. In contrast, high MA and high Br containing solar cells with rough perovskite layers ({x/y} = 0.75/0.40) show a change in the band-edge signal (Fig. 6f), indicating an increase in iodide-rich phase concentration, in agreement with changes observed through PL imaging. Furthermore, in compositions such as {x/y} = 0.75/0.50 and 0.75/0.60 (Figs. 6g and 6h), a large increase in the band-edge EQE contribution indicates the formation of a large density of iodide-rich phases because of light-induced halide segregation. Moreover, in composition {x/y} = 0.75/0.60, the change in the Urbach tail is accompanied by an increase in the sub-bandgap defect contribution (0.8–1.6 eV) to the EQE spectrum, indicating an increased defect density upon continuous illumination⁴⁷. Similar trends in photocurrent contribution from band-edge states and sub-bandgap electronic defects are observed in other solar cells using smooth ({x/y} = 0.25/0.50) or rough ({x/y} = 0.50/0.60) perovskite layers (Supplementary Fig. 21).

Controlling morphological and compositional heterogeneity of wide-bandgap perovskite thin films is key to the development of efficient perovskite-based multijunction solar cells. Using the perovskite composition as a variable, we find that MA- and Br-rich compositions are vulnerable to the formation of wrinkled morphology. We use GIWAXS characterization to associate the nominal composition of the film, the morphology of the layer and its structural properties such as preferential orientation. We identify that fast nucleation of Br-rich phases upon antisolvent casting is associated to the development of heterogeneity due to heterogeneous crystallization of bromide- and iodide-containing perovskite phases. The resulting features can be 1–2 µm in height and can span 2–5 µm laterally, making them unsuitable for solution-processed multijunction device stacks.

Using nano-XRF microscopy, we find a correspondence between surface wrinkling and halide heterogeneity wherein iodide-rich phases preferentially form at peak-like regions of rough films, influencing the local bandgap distribution across the film. The resulting local narrow-bandgap sites that are rich in iodide ions cause a broadening of band-edges, a higher Urbach energy, and a larger distribution in photoluminescence emission energies, which are potentially detrimental to device performance. Crucially, we also find that compositional/morphological heterogeneity aggravates sub-bandgap defect density. Finally, the combination of high defect density and local halide heterogeneity further reduces material stability to defect-driven light-induced halide segregation under illumination.

The results show how fast nucleation of bromide-rich regions, mediated by the presence of methylammonium, followed by the slower crystallization of iodide-rich regions leads to compositional and morphological heterogeneity that degrades both the performance and stability of the perovskite semiconductor and the resulting solar cells. This insight emphasizes that controlling crystallization kinetics through processing variables such as precursor solubility, film drying conditions and the use of agents that mediate crystallite nucleation and growth can be used to suppress film wrinkling and improve the compositional homogeneity of bromide-rich perovskite compositions^16,18. These methods provide potential routes to enhance the performance and stability of wide-bandgap perovskites that currently limit the development of multijunction photovoltaics.

Materials

All materials were used as received and were stored in an inert environment prior to use. FAI, MAI, FABr and MABr were purchased from Greatcell Solar Materials. 2PACz (> 98%), PbI₂ (99.99%, trace metal basis) and PbBr₂ (> 98%) were purchased from TCI Chemicals. DMF (99.8%), DMSO (99.9%, anhydrous) and ethyl acetate (99.8%) were purchased from Sigma Aldrich. C₆₀ was purchased from SES Research and bathocuproine (BCP) was purchased from Lumtec. Ethanol (< 0.1% H₂O) was purchased from Merck Millipore.

Sample Preparation

2PACz was dissolved in anhydrous ethanol at a concentration of 0.33 mg mL^− 1 by sonication prior to use. To prepare the perovskite solution, PbI₂ (691.5 mg mL^− 1) and PbBr₂ (550.5 mg mL^− 1) were each dissolved overnight at 60°C in solvent mixtures containing DMF and DMSO in a volumetric ratio of 4:1. The solutions were cooled to room temperature, following which 458 µL PbI₂ was added to each 98.4 mg FAI and 91.0 mg MAI respectively. 458 µL PbBr₂ was added to each 75.1 mg FABr and 64.1 mg MABr respectively. The resulting solutions (FAPbI₃, MAPbI₃, FAPbBr₃, and MAPbBr₃) were stirred at 60°C for approx. 1 h. Following that, the solutions were mixed based on the desired FA:MA and I:Br ratios and stirred at 60°C for approx. 1 h. The solutions were then cooled to room temperature prior to use.

Glass and glass/ITO (Naranjo Substrates, 15–17 Ω sq. ⁻¹) substrates were cleaned by sequential cleaning in acetone, followed by scrubbing with sodium dodecyl sulphate (Acros, 99%) soap solution in deionized water, sonication in soap solution, rinsing in deionized water followed by sonication in 2-propanol. Prior to use, the substrates were exposed to UV-ozone treatment for 30 min after which they were transferred to a N₂-filled glove-box.

Perovskite thin films were deposited by spin-coating 150 µL of the precursor at 4000 rpm (5 s to ramp to 4000 rpm) for 35 s. At approx. 25 s from the beginning of the spin-coating, 300 µL of ethyl acetate was cast onto the spinning substrate. The substrates were immediately annealed at 100°C for 30 min.

For solar cells, 2PACz was spin-coated at 3000 rpm for 30 s followed by thermal annealing at 100°C for 10 min. 20 nm C₆₀ and 8 nm BCP were sequentially evaporated at a rate of 0.5 Å s^− 1. 100 nm Ag electrode was thermally evaporated to complete the device. The nominal area of the solar cells determined by the overlap of the ITO and Ag areas is 9 mm².

Thin film and solar cell characterization

UV-vis-NIR spectra of perovskite thin films were measured using PerkinElmer Lambda 1050 UV-vis-NIR spectrophotometer. SEM images were acquired using FEI Quanta 3D FEG microscope, operated with a 5 kV electron beam and a secondary electron detector. AFM measurements were conducted using a Dimension 3100 AFM in tapping mode. For determining the J − V characteristics a Keithley 2400 SMU was used. A tungsten halogen lamp, filtered by a Schott GG385 UV filter and a Hoya LB120 daylight filter, was used to simulate 100 mW cm^− 2 of visible light. A shadow mask with 0.0676 cm² aperture was used to define the illuminated cell area. J − V scans involved sweeping the applied voltage (with no pre-biasing) from + 1.5 to − 0.5 V at a rate of 0.25 V s^− 1. EQE measurement was performed in a nitrogen atmosphere. The probe light source was generated by a 50 W tungsten-halogen lamp (Philips focusline), which was modulated at 160 Hz with a mechanical chopper (Stanford Research, SR 540) before passing into a monochromator (Oriel, Cornerstone 130). The spectral response of the device was recorded as a voltage from a pre-amplifier (Stanford Research, SR 570) using a lock-in amplifier (Stanford Research, SR 830), and was calibrated by a reference silicon cell. To accurately determine the current density, a green LED (530 nm, Thorlabs M530L3, driven by a DC4104 driver) was used as a light bias to provide the solar cell with approximately one sun illumination intensity.

GIWAXS characterization

GIWAXS measurements were done at beamline 11-BM at Brookhaven National Laboratory on perovskite films deposited on glass substrates cut to a size of approx. 0.5 cm × 0.5 cm. The samples were measured at an incident angle (α_i) 0.5° with a 10 s exposure time. The X-ray beam had an energy of 13.5 keV, 0.2 mm (height) × 0.05 mm (width) size, 1 mrad divergence and an energy resolution of 0.7%. Diffraction images were recorded using a Pilatus 800k detector. Data was analyzed using the SciAnalysis package provided by the beamline. In situ GIWAXS was performed during spin coating and thermal annealing in a custom-made spin-coater attached to beamline 12.3.2 at the Advanced Light Source (ALS), Lawrence Berkeley National Laboratory. The incoming X-ray beam was at an angle of 0.5° with a beam energy of 10 keV. A DECTRIS Pilatus1 M X-ray detector at an angle of 35° to the sample plane and a sample − detector distance of ∼ 186 mm was used. Measurements were carried out on an area of 0.1 mm² (10 mm × 0.01 mm) under an overpressure of N₂ in the spin-coater chamber. Samples were heated using a ramp rate of 4°C s^− 1. The diffraction data was collected with a framerate of approximately 1.875 s^− 1.

X-ray fluorescence microscopy

X-ray fluorescence (XRF) microscopy measurements were conducted using the Advanced Photon Source (APS) at beamline 2-ID-D at Argonne National Laboratory. A synchrotron X-ray energy of 14 keV was used and a step size of 0.15 µm and 50 ms dwell time. Data were analyzed using the MAPS software and spectrum fitting was used to deconvolute overlapping peaks and background from fluorescence data. In addition, after a standard calibration, it was possible to quantify the mass concentration in the sample to calculate molar ratios. The NIST thin-film standards SRM 1832 and 1833 were used for calibration.

Hyperspectral PL

Hyperspectral measurements were performed using a Photon etc. IMA upright microscope fitted with a transmitted darkfield condenser and a 60X objective (Nikon Plan RT, NA 0.7, CC 0-1.2). The excitation was done using a mercury halide lamp (Nikon ultrahigh pressure 130 W mercury lamp) passing through a 450 nm short-pass filter and emission was collected through a 500 dichroic filter and 550 nm long-pass filter. The lamp has six levels of light intensity, and all the measurements were taken using the lowest intensity (ND32) with total incident power on the sample of 130 mW cm^− 2. The scan duration was 1 min 15 s per frame. Post-processing was done in the proprietary Photon etc. PHySpec software.

Time-of-Flight Secondary Ion Mass Spectrometry

TOF-SIMS negative ion data were acquired on an IONTOF TOF.SIMS5 spectrometer using a 25 keV Bi³⁺ cluster ion source in the delayed extraction mode. The ion source was operated at a current of 0.04 pA that was rastered over a 100 µm × 100 µm area at 256 pixels × 256 pixels for a primary ion dose of 2.5 × 10¹¹ ions cm^− 2. A low energy flood gun was used for charge neutralization. Spectra were acquired over a mass range of m/z = 0 to 800 amu. The data was calibrated using the C₄H-, Br-, I- and PbI₂- peaks. Mass resolution of the m/z = 49.005 (C₄H-) peak was around 3500.

Sensitive EQE spectroscopy

Sensitive EQE measurements to characterize sub-bandgap states were conducted using an Osram 64655 HLX 250 W halogen lamp as the illumination source. The light was mechanically chopped at 333 Hz using an Oriel 3502 chopper and was subsequently passed through a monochromator (Oriel, Cornerstone 260) and appropriate sorting filters. The solar cell was mounted into a nitrogen-filled sample holder and its response was recorded as a voltage from a pre-amplifier (Stanford Research, SR 570) using a lock-in amplifier (Stanford Research, SR 830). The measurements were calibrated using Si and InGaAs reference cells. To illuminate the solar cells, a 532 nm CW laser (B&W Tek Inc., BWN-532-20E/56486) was used and its intensity and illumination area was adjusted to match 1-Sun equivalent intensity and the solar cell area using a set of neutral density filters and fish-eye lenses. The data was normalized to the drop in photocurrent that marks the band edge in the spectrum of the pristine solar cell. In photostability studies, spectra were scaled for the signal produced from pristine solar cells in order to estimate the change in the above-bandgap EQE.

Data Availability

All relevant data in this study are available from the corresponding author upon request. Source data are provided with this paper.

Acknowledgements

The authors acknowledge the Ministry of Education, Culture, and Science (Gravity program 024.001.035) and the Netherlands Organization for Scientific Research (Joint Solar Programme III Project 680.91.011 and Spinoza grant) for funding. M.T. acknowledges the Office of Naval Research (Award # N00014-20-1-2587) for funding. TOF-SIMS was carried out at the Molecular Analysis Facility, a National Nanotechnology Coordinated Infrastructure site at the University of Washington which is supported in part by the National Science Foundation (awards NNCI-2025489, NNCI-1542101), the Molecular Engineering and Sciences Institute, and the Clean Energy Institute. The research used the CMS 11-BM beamline of the National Synchrotron Light Source II, a U.S. Department of Energy (DOE) Office Science User Facility operated for DOE Office of Science by Brookhaven National Laboratory under contract no. DE-SC0012704. The research used beamline 12.3.2, a resource from the Advanced Light Source, a DOE Office of Science User Facility under Contract No. DE-AC02-05CH11231. Work at the Molecular Foundry was supported by the Office of Science, Office of Basic Energy Sciences, of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231. This research used resources of the Advanced Photon Source, a U.S. DOE Office of Science User Facility operated for the DOE Office of Science by Argonne National Laboratory under Contract No. DE-AC02-06CH11357.

Author Contributions

K.D., S.C.W.v.L., M.M.W. and R.A.J.J. planned the research. K.D. and R.L. measured and analyzed GIWAXS. T.K. and N.T. measured and analyzed in situ GIWAXS. J.H., B.L. and K.D. measured and analyzed nano-XRF microscopy. M.T. conceived of, measured, and analyzed the hyperspectral PL, with input from D. S. G., and M.T. and D.J.G. performed and analyzed the TOF-SIMS imaging data. S.C.W.v.L, R.W. and K.D. helped analyze hyperspectral PL microscopy and TOF-SIMS data. G.J.W.A. measured and analyzed sensitive photocurrent spectroscopy. K.D. wrote the manuscript with input from all co-authors. C.M.S.-F. contributed to the revision of the manuscript.

Competing Interest

The authors declare no competing interests.

Additional Information

Supplementary information The online version contains supplementary material available at

Correspondence and requests for materials should be addressed to René Janssen.

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Local halide heterogeneity drives surface wrinkling in mixed-halide wide-bandgap perovskites

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Abstract

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Introduction

Results

Summary

Methods

Materials

Sample Preparation

Thin film and solar cell characterization

GIWAXS characterization

X-ray fluorescence microscopy

Hyperspectral PL

Time-of-Flight Secondary Ion Mass Spectrometry

Sensitive EQE spectroscopy

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References

Additional Declarations

Supplementary Files

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