We deposited planar TiO₂ as the ETL and treated it with a SK solution. The SK modification increased the wettability of the TiO₂ as shown in Figure S1 (Supporting Information). The UV–vis absorption spectra and tauc‐plot curves are presented in Figure S2 (Supporting Information), corresponding to a bandgap of 3.015 eV for the control TiO₂ and TiO₂‐SK ETLs. The formamidinium lead iodide (FAPbI₃) PSCs were fabricated using a one‐step method, and further details are presented in the experimental section. Figure 1a illustrates the electrostatic potential (ESP) distributions of the SK modifier. The SK agent exhibits distinctive positive and negative surface potential on the two sides of the molecule. Figure 1b shows a schematic diagram of the PSCs based on TiO₂ ETL with the SK modifier. To comprehend the function of the SK modifier at the interface, we conducted Fourier transform infrared spectroscopy (FTIR) tests of the PbI₂, SK, and PbI₂‐SK mixtures (Figure 1c; Figure S3, Supporting Information). The results reveal that the stretching vibration of S═O is located at 1039.31 cm⁻¹ for the SK sample, while the value shifted to 1033.10 cm⁻¹ for the PbI₂‐SK sample. Additionally, the stretching vibration of ─CF₃ is located at 1165.44 cm⁻¹ for the SK sample and shifted to 1179 cm⁻¹. We further carried out the X‐ray photoelectron spectroscopy (XPS) of FAPbI₃ and SK modified FAPbI₃ perovskite films. The binding energy of F 1s increased from 687.05 to 688.70 eV while the binding energy of Pb 4f_5/2 and Pb 4f_7/2 reduced from 143.39 and 138.53 eV to 143.29 and 138.41 eV, respectively, as shown in Figure S4 (Supporting Information), indicating the interaction between the SK modifier and perovskite films. Due to the electron‐donating ability of the ─CF₃ unit, the electron cloud density of F will decrease while the electron cloud density of Pb increases after accepting electrons from the ─CF₃ units. Therefore, the binding energy of F inner shell electrons increased because of the weakened shielding effect of the outer electrons, leading to an increased binding energy of F 1s and a reduced binding energy of Pb 4f. Consequently, the ─SO₃ ⁻ and ─CF₃ units could act as Lewis‐base and tend to coordinate with the uncoordinated Pb²⁺, reducing defects in the perovskite film and charge recombination of the PSCs.^{[ 15 ]}

To investigate the interaction between the SK and TiO₂ ETL, we performed XPS tests of the control TiO₂ and SK treated TiO₂ (TiO₂‐SK) (Figure 1d–f; Figure S5, Supporting Information). Figure 1d shows the F 1s XPS spectrum, with the peak located at 687.05 eV for ─CF₃ in the SK modifier. Due to the electron‐donating ability of ─CF₃, the electron cloud density decreased after the interaction between the ─CF₃ and TiO₂, resulting in the increased binding energy of F 1s to 688.43 eV. In contrast, the electron density of Ti⁴⁺ increased after accepting electrons from ─CF₃, leading to a slight decrease in the binding energy for Ti 2p^1/2 and Ti 2p^3/2 from 464.64 and 458.82 eV to 464.43 and 458.64 eV, respectively, as shown in Figure 1e. Figure 1f shows that the binding energy of O 1s for the SK sample is located at 532.67 eV. Meanwhile, the TiO₂ ETL could be decoupled into two peaks, which are ascribed to the lattice oxygen (O_L) and oxygen vacancy (O_V), corresponding to a binding energy of O 1s at 529.95 and 531.36 eV.^{[ 28 ]} For TiO₂‐SK ETL, the two peaks were shifted to 530.05 and 531.70 eV due to the occupation of the O_V by the O atom of the ─SO₃ ⁻ unit. Overall, both the aggregative negative potential of ─CF₃ and the positive potential of ─SO₃ ⁻ could coordinate with the vacancy in the TiO₂ ETLs and perovskite layers due to its electron donating ability. As a result, the functional ─SO₃ ⁻ and ─CF₃ units of SK agents could modify the perovskite and TiO₂ layer, acting as a crosslinker at the interface to improve the interfacial contact of the PSCs.

We further investigated the surface morphology of the TiO₂ and TiO₂‐SK ETLs using atomic force microscopy (AFM) and scanning electron microscopy (SEM) shown in Figure S6 (Supporting Information). A thin and compact TiO₂ layer (≈30 nm) is deposited on the FTO substrates. The AFM results show that the root mean square roughness (RMS) of the TiO₂ and TiO₂‐SK ETLs is 37.7 and 32.9 nm, respectively, as depicted in Figure 2a,b. Additionally, we utilized Kelvin probe force microscopy (KPFM) to measure the contact potential difference (CPD = V_tip − V_sample). The CPD images are shown in Figure 2c,d, and the CPD value of the TiO₂ film decreased from 285 to 186 mV. These results indicate that TiO₂‐SK ETL displays a reduced work function and an increased Fermi level. The strong interaction of SK with TiO₂ alters the energy levels of the TiO₂ ETL. To further assess the energy levels, ultraviolet photoelectron spectroscopy (UPS) tests were conducted (Figure S7 Supporting Information). The E_cutoff edge of the UPS spectra of the TiO₂ shifted from 16.55 to 16.73 eV after modification with SK, as illustrated in Figure 2e,f, resulting in an elevated Fermi level from −4.67 to −4.49 eV, closer to that of the FAPbI₃ film (E_fermi = −4.24 eV). Consequently, the conduction band minimum (CBM) for TiO₂ and TiO₂‐SK were calculated to be −4.41 and −4.15 eV, respectively, which is beneficial for the energy level alignments between the ETL and perovskite film (CBM = −4.04 eV). Figure 2g illustrates the energy level alignments of the PSCs. The decreased energy level offset between the TiO₂‐SK ETL and FAPbI₃ perovskite film is advantageous for interfacial charge extraction and transfer, contributing to improved photovoltaic performance.

We then investigated the properties of the perovskite films. The FAPbI₃ film, when using reference TiO₂ and TiO₂‐SK ETLs, was designated as TiO₂/FAPbI₃ and TiO₂‐SK/FAPbI₃, respectively. Both the TiO₂/FAPbI₃ and TiO₂‐SK/FAPbI₃ perovskite films exhibit the same absorption onset at 811 nm, corresponding to an optical bandgap of 1.52 eV, Figures 3a and S8 (Supporting Information). The TiO₂‐SK/FAPbI₃ perovskite film shows a slightly higher absorption coefficient compared to the TiO₂/FAPbI₃ perovskite film. The X‐ray diffraction patterns (XRD) of the TiO₂/FAPbI₃ and TiO₂‐SK/FAPbI₃ perovskite films reveal peaks at 13.92° (100) and 28.01° (200), as shown in Figure 3b. The full width at half maximum (FWHM) of the TiO₂/FAPbI₃ are 0.159 and 0.177, which is larger than that of 0.110 and 0.139 for TiO₂‐SK/FAPbI₃ perovskite films shown in Figure S9 (Supporting Information). The higher intensity of the characteristic peaks and smaller FWHM for TiO₂‐SK/FAPbI₃ perovskite film indicates better crystallinity. SEM was employed to investigate the morphology of perovskite films. From SEM images in Figures 3c,d and S10 (Supporting Information), it is evident that the TiO₂‐SK/FAPbI₃ perovskite film displays larger crystal grains compared to the TiO₂/FAPbI₃ perovskite film, mainly attributed to the K⁺ diffusion into the FAPbI₃ film, promoting nucleation crystallization growth and enlarge the crystal size.^{[ 14} , ^{25 ]} The average grain size for the TiO₂‐SK/FAPbI₃ perovskite film is 1.3 µm, larger than the 1.1 µm for the TiO₂/FAPbI₃ perovskite film, as shown in Figure 3e,f. The increased grain size results in fewer defects, as the defects are generally distributed on the grain boundaries of the perovskite films.^{[ 5 ]} The improved morphology contributes to the enhancement of the photovoltaic performance of the PSCs.

To investigate the photovoltaic performance, we fabricated FAPbI₃ PSCs using a device configuration of FTO/c‐TiO₂/perovskite/Spiro‐OMeTAD/Au, and the experimental details are provided in the Experimental Section. Figures 4a and S11 (Supporting Information) illustrate the photovoltaic performance of the TiO₂‐SK PSCs compared to that of the control TiO₂ PSCs, with detailed parameters displayed in Table S1 (Supporting Information). Notably, significant improvements are observed in the V _oc and FF, which improved from 1.13 ± 0.05 V to 1.16 ± 0.05 V and 82.5 ± 1.0% to 84.0 ± 0.9%, respectively. Additionally, Figure 4b presents the current density–voltage (J–V) curves of the champion devices. The TiO₂‐SK PSC exhibits a remarkable reverse scanning PCE of 25.22% with a V _oc of 1.165 V, a J _sc of 25.51 mA cm⁻² and an FF of 84.9%. The forward scanning PCE is 24.80% with a V _oc of 1.161 V, a J _sc of 25.50, and an FF of 83.8%. Due to enhanced carrier extraction ability and suppressed nonradiative recombination, the TiO₂‐SK PSC displays less hysteresis. In contrast, the control PSC shows a reverse scanning PCE of 24.01% with a V _oc of 1.130 V, a J _sc of 25.47 mA cm⁻² and an FF of 83.5%, while the forward scanning PCE is 22.04% with a V _oc of 1.11 V and a J _sc of 25.44 mA cm⁻² and an FF of 78.1%. Furthermore, Figure 4c displays the stabilized power outputs (SPO) of the TiO₂ and TiO₂‐SK PSCs, corresponding to the PCEs of 23.51% and 24.82%. External quantum efficiency (EQE) measurements were performed to ascertain the J _sc value, resulting in integrated J _sc values of 24.85 mA cm⁻² for the control TiO₂ PSC and 24.91 mA cm⁻² for the TiO₂‐SK PSC, as shown in Figure 4d.

To investigate the carrier extraction and transport ability of the TiO₂ and TiO₂‐SK, we conducted steady‐state photoluminescence (PL) and time‐resolved photoluminescence (TRPL) tests. The PL results indicate an emission peak at ≈ 810 nm, as shown in Figure 4e,f, and the FAPbI₃‐SK film displays lower PL intensity compared to the FAPbI₃ film due to enhanced charge transfer from the perovskite layer to the TiO₂ ETL. This result was further confirmed by the TRPL test, which revealed faster PL quenching for TiO₂‐SK/FAPbI₃ film, verifying a shorter carrier lifetime of 761.7 ns compared to 1142.6 ns for the TiO₂/FAPbI₃ film (Table S2 Supporting Information). The reduced lifetime indicates efficient electron extraction at the TiO₂‐SK ETL/perovskite interface with SK modification, thereby reducing charge accumulation at the interface and suppressing device hysteresis. To quantify the trap defect density of the perovskite films, we conducted space‐charge limited current (SCLC) tests on the electron‐only devices in a structure of FTO/TiO₂/Perovskite/PCBM/BCP/Au (Figure 4g). The V_tfl values were estimated to be 0.934 and 1.812 V, respectively. The trap density was calculated using the equation N_trap = 2ɛ_rɛ₀V_TFL/eL², where e represents the elementary charge, and ɛ_r and ɛ₀ are the relative dielectric constant and vacuum permittivity, respectively.^{[ 29} , ^{30 ]} The trap densities of the TiO₂ and TiO₂‐SK PSCs were estimated to be 2.96 × 10¹⁶ and 1.51 × 10¹⁶.^{[ 31 ]}

The relationship between V _oc and light intensity was also analyzed, as depicted in Figure 4h. Linear fitting of V _oc versus log‐scaled light intensity allows the reflection of trap‐assisted recombination through the slope value and nkT/q (where k represents the Boltzmann constant, T is temperature, q is the electric charge, and n is the ideality factor). The smaller ideality factor of 1.44 for the TiO₂ PSC compared to 1.68 for the TiO₂‐SK PSC indicates a suppressed trap‐assisted recombination. Furthermore, electrical impedance spectroscopy (EIS) measurements were conducted to further investigate the interfacial charge transport behavior in PSCs (Figure 4i; Table S3, Supporting Information). The increased recombination resistance (R_rec = 1898 Ω) for TiO₂‐SK PSC indicates suppressed carrier recombination compared to the TiO₂ PSC (R_rec = 349 Ω). Additionally, the decreased transport resistance (R_tr) for the TiO₂‐SK PSC (13.78 Ω vs 15.05 Ω for TiO₂ PSC) is beneficial for the charge transport, resulting in an improved FF and reducing hysteresis.^{[ 32} , ^{33 ]}

We assessed the stability of the PSCs by storing them in an N₂ atmosphere, and the results are presented in Figures 5a and S12 (Supporting Information). The TiO₂‐SK PSC retained almost 100% of their initial PCE for 35 days, while the control device decreased to 97%. Additionally, we tested the stability of the unencapsulated PSCs in ambient conditions (RH = 20 ± 5%, RT = 25 ± 5 °C), and found that the TiO₂ and TiO₂‐SK PSCs preserved 84% and 95% of the initial efficiency, respectively, as shown in Figure 5b and Figure S13 (Supporting Information). Furthermore, we investigated the operational stability of the PSCs by aging the unencapsulated devices under a nitrogen atmosphere, using maximum power point (MPP) tracking under an LED lamp (light intensity 100 mw cm⁻²). Figure 5c shows that the normalized PCE of the TiO₂ PSC decreased to 73.7%, while the TiO₂‐SK PSC remained at 81.3% of the initial efficiency after operating for 1000 h. The improved stability of the TiO₂‐SK PSCs is attributed to the reduced interfacial defect density and improved energy level alignments after interfacial cross‐linking modification. The reduction in recombination and charge accumulation during the charge transfer process contributes to the operational stability of the PSCs under light illumination.^{[ 21} , ³⁴ , ^{35 ]} Moreover, the high crystallinity and enlarged grain size resulting from K‐doping also contribute to better stability and photovoltaic performance of the TiO₂‐SK PSCs.