Decoupling engineering of formamidinium–cesium perovskites for efficient photovoltaics

Abstract

Although pure formamidinium iodide perovskite (FAPbI₃) possesses an optimal gap for photovoltaics, their poor phase stability limits the long-term operational stability of the devices. A promising approach to enhance their phase stability is to incorporate cesium into FAPbI₃. However, state-of-the-art formamidinium–cesium (FA–Cs) iodide perovskites demonstrate much worse efficiency compared with FAPbI₃, limited by the different crystallization dynamics of formamidinium and cesium, which result in poor composition homogeneity and high trap densities. We develop a novel strategy of crystallization decoupling processes of formamidinium and cesium via a sequential cesium incorporation approach. As such, we obtain highly reproducible, highly efficient and stable solar cells based on FA_1–xCs_xPbI₃ (x = 0.05–0.16) films with uniform composition distribution in the nanoscale and low defect densities. We also revealed a new stabilization mechanism for Cs doping to stabilize FAPbI₃, i.e. the incorporation of Cs into FAPbI₃ significantly reduces the electron–phonon coupling strength to suppress ionic migration, thereby improving the stability of FA–Cs-based devices.

Sequential Cs incorporation strategy is developed to decouple crystallization of FACs perovskite with reduced electron-phonon coupling, resulting in highly stable FACs tri-iodide perovskite photovoltaics with record efficiency.

INTRODUCTION

Metal–halide perovskites with superior photophysical properties and low-cost solution technology have emerged as promising candidates for different optoelectronic devices, including solar cells, light-emitting diodes, etc. [1–6]. For perovskite solar cells (PSCs), the certified power-conversion efficiency (PCE) has reached 25.7%, which is comparable to the current commercial crystalline silicon solar cells. ABX₃ perovskites with tailoring compositions, where A is an organic or inorganic cation, B is a metal cation and X is a halide anion, have been attempted for high efficiency and stable photovoltaic devices. Among these, formamidinium lead iodide (FAPbI₃) has exhibited great potential as the absorber layer, due to its optimal band gap of ∼1.5 eV and excellent thermal stability [7–9].

However, the photoactive FAPbI₃ black phase would easily transform into a non-photoactive yellow δ-FAPbI₃ phase at room temperature, especially under humid conditions [10,11]. The poor phase stability challenges both the efficiency and long-term stability of the PSCs based on FAPbI₃ [12,13]. It is generally believed that the phase instability of FAPbI₃ perovskites originates from its unsuitable tolerant factor. To address this problem, alloying FA⁺ with MA⁺/Cs⁺ cations or partially substituting I⁻ with Br⁻ ions has been employed to tune the tolerant factor [14,15]. The resulting mixed-ion FA-based perovskites exhibit improved resistance to phase transition.

Among these different alloying approaches, formamidinium–cesium mixed-cation pure iodide (FA_1–xCs_xPbI₃) perovskites are particularly promising, because they avoid the concerns about volatile MA cations and phase segregation induced by mixed halide ions (Br–I) [16–20]. However, because of the complex crystallization kinetics of formamidinium and cesium, these pure iodide FA–Cs perovskites fabricated by one-step crystallization suffer from poor composition homogeneity and high defects/traps densities [21,22]. The PSCs based on these films are therefore facing relatively low efficiencies. Especially, strong non-radiative recombination in all reported FA–Cs-based PSCs limited the open-circuit voltage (V_oc) of the resulting devices [23–25].

Herein, we develop a novel sequential Cs incorporation (SCI) strategy to decouple the crystallization processes of formamidinium and cesium, and achieve highly efficient pure iodide FA_1–xCs_xPbI₃ (x = 0.05–0.16) perovskites (denoted as SCI-FA_1–xCs_xPbI₃). The ratio of FA and Cs in FA_1–xCs_xPbI₃ can be straightforwardly tuned by introducing different concentrations of cesium formate (HCOOCs) solution on the FA-based perovskite precursor film during the SCI process. A unique feature of our SCI-FA_1–xCs_xPbI₃ perovskites is the uniform distributions of Cs cations, in contrast to their poor uniformity in typical one-step (1S) crystallized FA–Cs perovskite films (denoted as 1S-FA_1–xCs_xPbI₃). As a result, the champion SCI-FA_0.91Cs_0.09PbI₃ PSCs yield a record PCE of 24.7% (certified 23.8%) with improved V_oc and fill factor, which is the highest value for the pure iodide FA_1–xCs_xPbI₃ perovskites. Compared with FAPbI₃, the SCI-FA_0.91Cs_0.09PbI₃ perovskite shows reduced electron–phonon coupling and lattice fluctuations, which suppress the formation of iodide-rich clusters and finally contribute to the excellent operational stability of the FA_0.91Cs_0.09PbI₃-based PSCs.

RESULT AND DISCUSSION

Figure 1a shows the schematic diagram of SCI-FA_1–xCs_xPbI₃ perovskite films prepared by decoupling the crystallization processes of formamidinium and cesium. A FAPbI₃ precursor film was first deposited by a typical anti-solvent method followed by annealing for 1 min. The Cs cation is sequentially introduced onto the FA perovskite film by spin-coating HCOOCs isopropanol (IPA) solution, followed by further annealing. For comparison, we employed different concentrations of HCOOCs solution (2.5, 5 and 10 mg mL^–1) to fabricate SCI-FA_1–xCs_xPbI₃ perovskites. The final ratios of incorporated Cs in the above SCI-FA_1–xCs_xPbI₃ perovskite films, i.e. the value of x, are 0.05, 0.09 and 0.16, as confirmed by inductively coupled plasma–mass spectrometry (ICP–MS) analysis (Supplementary Table 1). The corresponding SCI-FA_1–xCs_xPbI₃ perovskite films are noted as x = 0.05, x = 0.09 and x = 0.16 in Fig. 1.

Figure 1.

Spectroscopic, structural and morphological characterizations of SCI-FA_1–xCs_xPbI₃ films. (a) Schematic diagram of the SCI-FA_1–xCs_xPbI₃ perovskite films fabricated by decoupling the crystallization processes of formamidinium and cesium. (b) UV–vis absorption and normalized PL spectra and (c) XRD patterns of FAPbI₃ and SCI-FA_1–xCs_xPbI₃ perovskites. Inset pattern corresponding to the characteristic peaks of (001) perovskite crystal planes. (d) Top-surface SEM images of FAPbI₃ and SCI-FA_1–xCs_xPbI₃ perovskite films. The inset presents the cross-sectional morphology of the corresponding perovskite films. Scale bars: 1 μm.

Optical and structural measurements of perovskite films indicate that Cs⁺ from HCOOCs has successfully been incorporated into the lattice of FAPbI₃ perovskites. Figure 1b shows the ultraviolet–visible (UV–vis) spectra of SCI-FA_1–xCs_xPbI₃ perovskite films, in which the absorption edges of SCI-FA_1–xCs_xPbI₃ perovskites strongly depend on the amount of Cs⁺ incorporation. When x increases from 0 to 0.16, the absorption edges of SCI-FA_1–xCs_xPbI₃ perovskites gradually blue-shift from 816 to 802 nm, and the corresponding photoluminescence (PL) peaks shift from 809 to 797 nm. The X-ray diffraction (XRD) measurements are carried out to investigate the crystal structure evolution of SCI-FA_1–xCs_xPbI₃ perovskites (Fig. 1c). All the SCI-FA_1–xCs_xPbI₃ perovskites exhibit stronger peak intensity than the pure FAPbI₃ at around both 14° and 28°, corresponding to (001) and (002) perovskite crystal planes. The inset image of Fig. 1c shows that the peak between 13.8° and 14.1° shifts to a higher degree, indicating that Cs ions are incorporated into the perovskite lattice. The lattice parameter decreases with increasing the amount of Cs (Supplementary Fig. 1), further confirming the successful mixing of Cs⁺ in the perovskite lattice. The tolerance factor of SCI-FA_1–xCs_xPbI₃ perovskites is also reduced compared with pure FA perovskite, potentially contributing to an stable perovskite structure (Supplementary Fig. 2).

The Cs incorporation also significantly improves the film morphologies (Fig. 1d). All SCI-FA_1–xCs_xPbI₃ films show enlarged and pinhole-free grains compared with the FAPbI₃ film, which shows coarse grains and pinholes. As shown in the cross-sectional scanning electron microscopy (SEM) images, SCI-FA_1–xCs_xPbI₃ perovskite films (550–600 nm) with vertical growth of grains benefit efficient charge extraction.

By adopting these Cs-incorporated perovskites as light absorber layers, we fabricate PSCs with a configuration of fluorine-doped tin oxide (FTO)/electron-transport layer/perovskite/hole-transport layer/Au. All SCI-FA_1–xCs_xPbI₃ (x = 0.05, 0.09, 0.16)-based PSCs exhibit improved device efficiency compared with the FAPbI₃-based devices (Supplementary Fig. 3). Considering that x = 0.09 provides the optimal photovoltaic (PV) performance, we then chose this composition (denoted as SCI-FA_0.91Cs_0.09PbI₃) for detailed investigations on Cs incorporation and its role on film and device properties.

X-ray photoelectron spectroscopy spectra are conducted to explore the effect of SCI on the elements and their chemical states in perovskite films. All core-level peaks are assigned to Cs, Pb, N and I (Supplementary Fig. 4) elements. The characteristic Cs signals in the SCI-FA_0.91Cs_0.09PbI₃ perovskite locate at 738.5 and 724.7 eV, and show a 1.1-eV shift compared with the Cs in HCOOCs. Such a large shift is attributed to the formation of chemical bonds between Cs⁺ and [PbI₆]⁴⁻. For the Pb 4f spectra in FAPbI₃, two peaks corresponding to Pb 4f_7/2 and Pb 4f_5/2 are observed at 138.3 and 143.2 eV. However, in the SCI-FA_0.91Cs_0.09PbI₃ perovskite, both Pb 4f_7/2 and Pb 4f_5/2 shift by 0.1 eV toward higher binding energy, originating from the stronger bond energies between Cs⁺ and [PbI₆]⁴⁻ than that between FA⁺ and [PbI₆]⁴⁻. Additionally, the N and I elements representing the formamidine component have undergone a small shift. These results further confirm that Cs has been successfully incorporated into the FAPbI₃ perovskite lattice to form SCI-FA_0.91Cs_0.09PbI₃ perovskite.

For comparison, we also fabricate one-step crystallized 1S-FA_1–xCs_xPbI₃, where Cs⁺ (from HCOOCs) is directly mixed with FA⁺ in the perovskite precursor. For 1S-FA_1–xCs_xPbI₃ PSCs (Supplementary Fig. 5), the optimal efficiency is also obtained from the x = 0.09 sample, which hence will be used as the control sample for comparison with SCI-FA_0.91Cs_0.09PbI₃. The absorption edge, characteristic XRD peaks and surface morphologies of the 1S-FA_0.91Cs_0.09PbI₃ perovskite films are consistent with SCI-FA_0.91Cs_0.09PbI₃ perovskite (Supplementary Fig. 6).

The synchrotron-based grazing-incidence wide-angle X-ray scattering (GIWAXS) and time-of-flight secondary ion mass spectrometry (ToF-SIMS) measurements are further employed to explore the crystal structure and internal composition in perovskite films. As shown in Fig. 2a and b, there is a signal ring at q_xy = 8.8 nm⁻¹ corresponding to δ-FAPbI₃ in 1S-FA_0.91Cs_0.09PbI₃ perovskite, while the GIWAXS result of SCI-FA_0.91Cs_0.09PbI₃ exhibits high phase purity and obvious crystal orientation without any obvious phase impurities. Since the distribution of Cs in SCI-FA_1–xCs_xPbI₃ perovskite has a significant effect on both the phase stability and traps/defects [22,26], we proceed to investigate the distribution of Cs in the SCI-FA_0.91Cs_0.09PbI₃ and 1S-FA_0.91Cs_0.09PbI₃ perovskites. Both energy-dispersive spectroscopy mapping and ToF-SIMS have established that the incorporated Cs⁺ cations homogenously distribute in the surface and bulk of SCI-FA_0.91Cs_0.09PbI₃ perovskite, which is completely different from the inhomogeneous Cs⁺ cation distributions in 1S-FA_0.91Cs_0.09PbI₃ perovskite (Fig. 2c and Supplementary Fig. 7). The Cs aggregation at the top surface of 1S-FA_0.91Cs_0.09PbI₃ reveals the phase separation in 1S-FA_0.91Cs_0.09PbI₃, which is consistent with the GIWAXS data. Such uniform distribution of Cs ions in the SCI-FA_0.91Cs_0.09PbI₃ perovskite originates from the decoupled crystallization kinetics by the SCI strategy, which is beneficial for enhancing the phase stability and reducing defects/traps density. Different from Cs⁺ cations, other ions, including FA⁺, Pb²⁺ and I⁻, are uniformly distributed throughout both the SCI-FA_0.91Cs_0.01PbI₃ and 1S-FA_0.91Cs_0.09PbI₃ perovskite films (Supplementary Fig. 8). Although most of the HCOO⁻ anions escape from the FA_0.91Cs_0.09PbI₃ perovskite during the annealing process, residual traces of HCOO⁻ still exist at the bottom interface between the perovskite and the substrate (Supplementary Fig. 9), which can further passivate the defects on the bottom interface and improve the charge carriers’ dynamics in the devices [27,28].

Figure 2.

Characterizations of FA_0.91Cs_0.09PbI₃ perovskite films. GIWAXS data of (a) SCI-FA_0.91Cs_0.09PbI₃ and (b) 1S-FA_0.91Cs_0.09PbI₃ films. (c) 3D distribution in the SCI-FA_0.91Cs_0.09PbI₃ and 1S-FA_0.91Cs_0.09PbI₃ film by ToF-SIMS analysis, where the signal intensity in two kinds of perovskites is normalized. (d) TRPL decay curves of FAPbI₃, SCI-FA_0.91Cs_0.09PbI₃ and 1S-FA_0.91Cs_0.09PbI₃ perovskite films.

Uniform incorporation of Cs ions into FAPbI₃ has two positive effects: enhanced phase stability and decreased trap/defect densities. The enhanced phase stability is evidenced from the absence of color/structural changes under humid conditions. Under 60% relative humidity, the SCI-FA_0.91Cs_0.09PbI₃ perovskite maintains the black phase for 7 days without any changes (Supplementary Fig. 10), showing significant improvement compared with FAPbI₃. The decreased trap/defect densities are demonstrated from photophysical measurements. The photoluminescence (PL) intensities of SCI-FA_0.91Cs_0.09PbI₃ perovskites are much stronger than those of 1S-FA_0.91Cs_0.09PbI₃ and pure FA perovskite films (Supplementary Fig. 11). In addition, the time-resolved PL (TRPL) spectra in Fig. 2d show that the PL lifetime (τ) of SCI-FA_0.91Cs_0.09PbI₃ is much longer (413.84 ns) than that of 1S-FA_0.91Cs_0.09PbI₃ (240.41 ns) and FAPbI₃ (165.49 ns). Enhanced PL intensity and improved PL lifetime indicate that non-radiative recombination is suppressed in SCI-FA_0.91Cs_0.09PbI₃, attributed to decreased trap densities. Consistently, we calculated the Urbach energy (E_u) according to the equation: A = A₀ exp(E/E_u), where A is the absorbance, A₀ is a constant for data fitting and E is the excitation energy. The Urbach energy (Supplementary Fig. 12) is decreased from 22.2 meV in FAPbI₃ and 21.6 meV in 1S-FA_0.91Cs_0.09PbI₃ to 18.3 meV in SCI-FA_0.91Cs_0.09PbI₃. The smaller Urbach energy in SCI-FA_0.91Cs_0.09PbI₃ corresponds to a lower density of trap states. These results suggest that, with sequential incorporation of Cs⁺, the SCI-FA_0.91Cs_0.09PbI₃ perovskite exhibits significantly reduced non-radiative recombination via defects/traps.

Benefitting from these advantages of crystallization decoupling engineering, the resulting SCI-FA_0.91Cs_0.09PbI₃ perovskite shows much enhanced device performance. Figure 3a compares the current density–voltage (J–V) characteristics of champion PSCs based on FAPbI₃, SCI-FA_0.91Cs_0.09PbI₃ and 1S-FA_0.91Cs_0.09PbI₃ perovskites, respectively. The SCI-FA_0.91Cs_0.09PbI₃-based PSC exhibits an impressive PCE of 24.7% compared with 22.6% for FAPbI₃, representing a new record for pure iodide FA_1–xCs_xPbI₃-based PSCs (Supplementary Table 2). A certified PCE of 23.8% with negligible hysteresis has been obtained in the SCI-FA_0.91Cs_0.09PbI₃-based PSCs (Supplementary Fig. 13). The most striking difference is the V_oc, which increases from 1.09 V in FAPbI₃ to 1.18 V in SCI-FA_0.91Cs_0.09PbI₃. The counterpart PSCs based on 1S-FA_0.91Cs_0.09PbI₃ perovskite yield a lower efficiency of 23.1%. This comparison indicates that our crystallization decoupling engineering is beneficial for enhancing SCI-FA_1–xCs_xPbI₃-based PSCs performance. The incident photon to electron conversion efficiency (IPCE) (Fig. 3b) is similar for both devices, with a high value of >90% in the wavelength range of 450–650 nm. The short-circuit current density (J_sc) of the SCI-FA_0.91Cs_0.09PbI₃ device is slightly decreased compared with the FAPbI₃ device, mainly due to slight increase in the band gap upon Cs incorporation. Figure 3c compares the PV parameters of FAPbI₃-, SCI-FA_0.91Cs_0.09PbI₃- and 1S-FA_0.91Cs_0.09PbI₃-based PSCs for 18 devices, respectively, indicating that sequential Cs incorporation also improves the device reproducibility. In addition, the SCI-FA_0.91Cs_0.09PbI₃-based PSCs exhibit a smaller hysteresis (Supplementary Figs 13 and 14), resulting in a stabilized output power of 24.4% (Fig. 3d).

Figure 3.

Photovoltaic and device characterization. (a) The J–V curves of the champion devices of FAPbI₃, SCI-FA_0.91Cs_0.09PbI₃ and 1S-FA_0.91Cs_0.09PbI₃ PSCs with 0.1 cm² of effective cell area. (b) IPCE spectra of FAPbI₃, SCI-FA_0.91Cs_0.09PbI₃- and 1S-FA_0.91Cs_0.09PbI₃-based PSCs. (c) The PV performance distribution of FAPbI₃-, SCI-FA_0.91Cs_0.09PbI₃- and 1S-FA_0.91Cs_0.09PbI₃-based PSCs from 18 devices, respectively. (d) Steady-state efficiency of FAPbI₃ and SCI-FA_0.91Cs_0.09PbI₃ and 1S-FA_0.91Cs_0.09PbI₃ PSCs. (e) J–V characteristics of PSCs based on FAPbI₃, SCI-FA_0.91Cs_0.09PbI₃ and 1S-FA_0.91Cs_0.09PbI₃ with 1 cm² of effective cell area under simulated AM 1.5-G solar illumination of 100 mW cm^–2 in the reverse scan. (f) EQE_EL of FAPbI₃-, SCI-FA_0.91Cs_0.09PbI₃- and 1S-FA_0.91Cs_0.09PbI₃-based PSCs vs. the current density.

In addition to small-area PSCs, the large-area PSCs based on these sequential Cs-incorporated perovskite films also exhibit significantly improved device performance. The champion SCI-FA_0.91Cs_0.09PbI₃ device, fabricated on 2.5 × 2.5 cm² substrates with an effective cell area of 1 cm² (Fig. 3e), displays a PCE of 22.4%, which is far higher than the FAPbI₃ (∼19.9%) and 1S-FA_0.91Cs_0.09PbI₃ (∼20.7%)-based devices.

The significantly enhanced V_oc of the SCI-FA_0.91Cs_0.09PbI₃ device is mainly due to suppressed non-radiative recombination, which can be quantified by measuring the external quantum efficiency of electroluminescence (EQE_EL) values [29]. As shown in Fig. 3f, at the injection current densities corresponding to J_sc, the EQE_EL value of the SCI-FA_0.91Cs_0.09PbI₃ device is 6.38%, while that of the FAPbI₃ device is 0.16%. We calculate the voltage losses due to non-radiative recombination (ΔV_ocnon-rad,) based on the formula [30]:

where k, T and q represent the Boltzmann constant, temperature and elementary electric charge, respectively. The difference in ΔV_ocnon-rad, (0.09 V) matches well with the difference of device V_oc (0.09 V).

Suppressed non-radiative recombination in the SCI-FA_0.91Cs_0.09PbI₃ device is consistent with previous photophysical measurements on the films, which indicate that the sequential Cs incorporation can reduce the defects/traps. Further measurements on the devices also reach similar conclusions. The trap-filled limiting voltage in the space-charge limited current measurements decreases from 0.13 V in the FAPbI₃ device to 0.09 V in SCI-FA_0.91Cs_0.09PbI₃ device (Supplementary Fig. 15), indicating suppressed traps/defects upon Cs sequential incorporation [31,32]. These results are also consistent with transient photovoltage (TPV) decay and transient photocurrent (TPC) decay results (Supplementary Fig. 16), which show slower TPV decay (indicating longer recombination lifetime) and quicker TPC decay (indicating fewer trapping effects) in the SCI-FA_0.91Cs_0.09PbI₃ device [33,34].

In addition to improved PV performance, the SCI-FA_0.91Cs_0.09PbI₃ device also shows significantly enhanced stability. We first measure the shelf life by storing the unencapsulated devices in dark at 25°C and 20% relative humidity. Figure 4a shows that the PCE of the FAPbI₃ device decreases by ∼30% after 3000 h of aging, whereas the SCI-FA_0.91Cs_0.09PbI₃ device shows a degradation of only 10% over 4500 h of aging. We then investigate the long-term operational stability of the PSCs by aging the unencapsulated devices under a nitrogen atmosphere, using maximum power point (MPP) tracking under simulated 1-sun conditions. As shown in Fig. 4b, the SCI-FA_0.91Cs_0.09PbI₃ based PSCs retains >90% of the initial PCE while the FAPbI₃ device maintains only 60% PCE after 1000 h of continuous illumination (Supplementary Fig. 17). Especially, the sharp decline of the efficiency in FAPbI₃ PSCs in the initial stage should be attributed to the intrinsic instability of the FAPbI₃ perovskite layer and potential severe ion migration in the FAPbI₃ device.

Figure 4.

Stability of the SCI-FA_0.91Cs_0.09PbI₃ PSCs. (a) The shelf-life stability of unencapsulated FAPbI₃ and SCI-FA_0.91Cs_0.09PbI₃ PSCs. (b) The long-term operational stability of unencapsulated FAPbI₃ and SCI-FA_0.91Cs_0.09PbI₃ PSCs. The 2D I⁻ distribution in the middle of the perovskite layer (at ∼300 nm) for PSCs based on (c) FAPbI₃ and (d) SCI-FA_0.91Cs_0.09PbI₃ perovskites after 240 h of operational stability test at their MPP by ToF-SIMS. (e) Fit of the FWHM of the PL spectra vs. temperature.

A main reason for enhanced-stability PSCs is attributed to suppressed ionic migration. In Fig. 4c and d, we compare the I⁻ ions distribution at the ∼300-nm depth of the perovskite layer for PSCs based on FAPbI₃ and SCI-FA_0.91Cs_0.09PbI₃ after 240-h MPP test. In the FAPbI₃-based device, strong aggregation of I⁻ clusters is observed in the perovskite absorber layers; in contrast, I⁻ ions distribute uniformly in the SCI-FA_0.91Cs_0.09PbI₃-based devices. This sharp contract indicates that the ionic migration in the SCI-FA_0.91Cs_0.09PbI₃ is much suppressed upon Cs sequential incorporation.

Suppressed ionic migration in SCI-FA_0.91Cs_0.09PbI₃ is consistent with suppressed electron–phonon coupling upon Cs incorporation. Figure 4e shows the full-width half-maximum (FWHM) of the PL peak of SCI-FA_0.91Cs_0.09PbI₃ and FAPbI₃ perovskites (Supplementary Fig. 16) ranging from 110 to 296 K. The wide broadening of the PL linewidth in FAPbI₃ perovskites arises from strong electron–phonon coupling [35,36]. The electron–phonon interaction is dominated by high energy longitudinal optical (LO) phonons in the high-temperature region, where the measured FWHM data could be fitted by the Boson model (Fig. 4c, Supplementary Fig. 18 and Supplementary Table 3). Compared with FAPbI₃, both the electron–LO phonon coupling coefficient (Γ_LO) and LO phonon energy (hω) in the SCI-FA_0.91Cs_0.09PbI₃ are significantly reduced, indicating that the fluctuation of the PbI₆ octahedra cage in SCI-FA_0.91Cs_0.09PbI₃ is associated with much smaller energies upon the Cs sequential incorporation. This is consistent with the previous theoretical investigations, which indicate that mixed A-site cations could reduce the lattice fluctuations in halide perovskites [37]. As such, the suppressed lattice fluctuations and electron–phonon coupling in SCI-FA_0.91Cs_0.09PbI₃ rationalize suppressed ionic migration and hence enhanced stability in SCI-FA_0.91Cs_0.09PbI₃ PSCs, which agree with previous research results that the suppressed lattice fluctuations and reduced electron–phonon coupling could suppressed the formation of iodide-rich clusters to improve the stability of halide perovskites [38,39].

CONCLUSION

In summary, we successfully develop a novel SCI strategy to tackle the critical challenge of different crystallization dynamics of different cations in developing FA_1–xCs_xPbI₃ perovskite PSCs. The resulting pure iodide SCI-FA_1–xCs_xPbI₃ perovskites show more uniform composition distribution and reduced defects/traps density than FAPbI₃ and one-step crystallized 1S-FA_0.91Cs_0.09PbI₃. Compared with FAPbI₃, the SCI-FA_0.91Cs_0.09PbI₃ exhibits reduced electron–phonon coupling and lattice fluctuations, minimizing ion migration and hence enhancing the stability. As such, we have been able to achieve highly stable PSCs with a high efficiency of 24.7%, which is a record for SCI-FA_1–xCs_xPbI₃ PSCs. This work opens up new possibilities to develop high-quality mixed-cation perovskites, presenting a milestone towards the development of highly efficient and highly stable perovskites for various applications, including solar cells, light-emitting diodes and lasers.

METHODS

Materials

Lead iodide (PbI₂, 99.9985%), cesium formate (HCOOCs, 98%), methylamine hydrochloride (MACl, 99%) and tin (IV) oxide colloid precursor (SnO₂, 15% in H₂O colloidal dispersion) were purchased from Alfa Aesar. Formamidinium iodide (FAI) was purchased from Xi’an Polymer Light Technology Corp. Other materials were purchased from Sigma-Aldrich and used as received without any purification. N,N-dimethylformamide (DMF, anhydrous, 99.8%), dimethyl sulfoxide (DMSO, anhydrous, 99.7%), chlorobenzene (anhydrous, 99.8%) and isopropanol alcohol (IPA, 99.5%) were purchased from J&K Scientific Ltd. Ammonium solution (AR, 25%–28%) was purchased from Aladdin.

Device fabrication

A compact TiO₂ layer (20 nm) was deposited using the spray pyrolysis method using a titanium bis(ethyl acetoacetate)-diisopropoxide/1-butanol solution (1 : 9 volume ratio). The cleaned patterned fluorine-doped tin oxide (FTO, 7 Ω sq^–1) substrate was placed on a 450°C plate during the spray process followed by 1 h of annealing. The SnO₂ colloid precursor/ammonium solution (1 : 9 volume ratio) was spin-coated on a TiO₂ layer at 3000 rpm for 30 s followed by annealing for 30 min at 180°C. The SCI-FA_1–xCs_xPbI₃ perovskite precursor was prepared by mixing PbI₂, FAI and MACl (30 mol%) in DMF : DMSO = 9:1 solution to form 1.5 M FAPbI₃ precursor. The perovskite precursor films were deposited by spin-coating on FTO/TiO₂/SnO₂ substrate at 5000 rpm for 15 s. During spin-coating, 150–200 μL of chlorobenzene was dripped at the end of 10 s. For the SCI perovskite samples, the perovskite film was first annealed for 1 min at 150°C to obtain a light-brown FA perovskite film. Different concentrations of HCOOCs/IPA solution were dropped on the FAPbI₃ precursor films then spin-coated at 3000 rpm for 30 s and further annealed at 150°C for 9 min and 100°C for 10 min. The 1S-FA_1–xCs_xPbI₃ perovskite was prepared by mixing PbI₂, FAI, MACl (30 mol%) and x mol% HCOOCs in DMF : DMSO = 9 : 1 solution to form 1.5 M FA–Cs precursor. The perovskite precursor films were deposited by spin-coating on FTO/TiO₂/SnO₂ substrate at 5000 rpm for 15 s. During spin-coating, 150–200 μL of chlorobenzene was dripped at the end of 10 s. The perovskite films annealed at 150°C for 10 min and 100°C for 10 min. The spiro-OMeTAD layer was spin-coated by dissolving 90 mg of spiro-OMeTAD in 1 mL of chlorobenzene and mixing with 39.5 μL of 4-tert-butylpyridine (TBP), 23 μL of Li-bis(trifluoromethanesulfonyl) imide (Li-TFSI)/acetonitrile (520 mg/mL) and 10 μL of tris(2-(1h-pyrazol-1-yl)-4-tert-butylpyridine)-cobalt(III)tris(bis(trifluoromethylsulfonyl)imide) (FK209)/acetonitrile (375 mg mL^–1) at 4000 rpm for 25 s. Finally, a 80-nm thick gold electrode was thermal evaporated on the spiro-OMeTAD layer. The perovskite films fabrication was operated in a dry box with <15% humidity.

FUNDING

The work performed at Shanghai Jiao Tong University was supported by the National Natural Science Foundation of China (22025505) and the Program of Shanghai Academic/Technology Research Leader (20XD1422200). The work performed at the Linköping University was supported by the Swedish Government Strategic Research Area in Materials Science on Functional Materials at Linköping University (SFO-Mat-LiU #2009–00971).

AUTHOR CONTRIBUTIONS

Y.Z. and F.G. designed and directed the study. H.C. and Y.F. conceived and performed the device fabrication work. Y.W. performed photophysical measurements and analysed the results. H.C., Y.F., Y.W., Y.C., Y.M., Z.Q., X.W., X.L., K.Z., F.G. and Y.Z. participated in characterization and data analysis. All authors contributed to the discussions. H.C., Y.F., Y.W., F.G. and Y.Z. wrote the manuscript with input from all authors. All authors reviewed the paper.

Conflict of interest statement . None declared.