Tailoring tin-based perovskite crystallization via large cations and pseudo-halide anions for high mobility and high stable transistors

Abstract

Tin-based perovskites, renowned for their eco-friendliness, intrinsic high hole mobility, and low effective mass, hold great potential for p-type thin-film transistors (TFTs). However, their propensity for rapid crystallization and oxidation severely limits stability and carrier mobility. Here, we strategically enhance perovskite TFT performance by incorporating 2-thiopheneethylamine thiocyanate (TEASCN) into 3D tin-based perovskites. The induction of the pseudo-halide SCN⁻ into a bilayer quasi-2D perovskite intermediate phase, combined with the strong interaction between sulfur-bearing thiophene rings (TEA⁺) and Sn-I octahedra, effectively reorients perovskite crystallization while inhibiting Sn²⁺ oxidation and reducing trap density. Consequently, TEASCN-based TFTs achieve an average hole mobility of more than 60 square centimeters per volt per second and an on/off current ratio surpassing 10⁸, standing out among state-of-the-art p-type perovskite TFTs. Furthermore, unencapsulated devices preserve 84% of their initial mobility after 30 days in an N₂ atmosphere, underscoring their remarkable stability. This work opens a straightforward path toward high-mobility and highly stable tin-based perovskite transistors.

TEASCN-induced intermediate phase enhances crystallization and stability, enabling high-performance tin-based perovskite TFTs.

INTRODUCTION

N-type oxide thin-film transistors (TFTs) have superior mobility and stability, which are widely applied in electronic industrialization, with low-temperature operation, good economy, and mature technology (1–3). However, the mismatch between n-type and p-type TFTs poses a substantial challenge in balancing electron and hole carrier transport, which critically impedes data processing speed and hinders the realization of high-performance, low-power complementary metal oxide semiconductors devices, and integrated circuits (4–6). Consequently, there is an urgent need to explore and develop nascent p-type materials that meet the performance standards (mobility and on/off ratio) of n-type oxide materials, thereby advancing their application in cutting-edge electronic devices and systems (7, 8). Among various p-type semiconductor materials, tin-based perovskites are a rising star material because they have weaker Frohlich interactions and higher 5s orbital energy levels (9). This unique combination results in highly dispersed valence bands and enhanced hole mobility, making them suitable for high-performance p-type TFTs.

Among the two-dimensional (2D) layered perovskites, (PEA)₂SnI₄ was system researched as the channel for perovskite transistors, owing to their structural stability and versatility (10, 11). Subsequent advancements in film morphology, device structure, and crystal formation led to improved transistor performance (12–14). However, further improvements are constrained by quantum and dielectric confinement, where charge tunneling perpendicular to the inorganic layers is severely inhibited, thereby limiting their charge transport performance (15–17). In contrast, the advantage of 3D perovskites lies in their continuous 3D network structure, which holds the potential for faster charge transport (16, 18). Recently, this was demonstrated by a high carrier mobility of over 50 cm² V⁻¹ s⁻¹ in a 3D Sn-Pb hybrid perovskite system (19, 20). However, pure 3D tin-perovskite transistors still exhibit relatively poor carrier mobility and long-term stability, owing to the ease of Sn2⁺ oxidation and uncontrollable crystal growth kinetics (21–24). To improve device performance and stability, large cations are incorporated into 3D structures to form low-dimensional phases (25–27), However, because of the coexistence of different dimensional structures, it will result in disordered film structures and an increase in film defects (28, 29), leading to poor stability and notable device hysteresis. Furthermore, the device’s performance will still degrade, even when stored in nitrogen-filled glove boxes with only trace amounts of oxygen. To obtain high-mobility and high-stability transistors, it is necessary to control the crystallization dynamics to achieve the desired crystal orientation and improve the stability of the tin-based perovskite film.

In this study, we develop high mobility and stable perovskite TFTs by introducing 2-thiopheneethylamine thiocyanate (TEASCN) into 3D Sn-based perovskites (CsFASnI₃). The amino and sulfur groups within the TEASCN additive form robust bonds with the Sn-I octahedrons, while the thiocyanate ions induce the formation of a bilayer quasi-2D intermediate phase, which facilitates the vertical orientation of the 3D structure perpendicular to the substrate. These interactions enhance the morphology and crystallinity of tin-based perovskites, promote charge transfer characteristics at the interface and within perovskite layers, and markedly reduce the concentration of Sn⁴⁺ in the tin-based perovskite film, leading to a decrease in defect density. The 3D perovskite TFTs incorporating TEASCN exhibit an average hole mobility (μ_h) exceeding 60 cm² V⁻¹ s⁻¹. After storage in the nitrogen glove box for 30 days, the unencapsulated TFTs still exhibit excellent stability.

RESULTS

To explore the influence of TEASCN on the electrical properties, Sn-based perovskite TFTs were fabricated with a bottom-gate, top-contact structure, as illustrated in Fig. 1A. The mixed-cation Cs_0.15FA_0.85SnI₃ perovskite was used as the reference material, primarily due to the larger size of the formamidinium (FA) ion in comparison to other 3D perovskites, which weakens the antibonding interactions between the 5s orbitals of tin (Sn) and the 5p orbitals of iodine (I) in FASnI₃. This results in a higher formation energy for tin vacancy defects (V_Sn) in FASnI₃ (30). Moreover, the incorporation of Cs into the FASnI₃ lattice further bolsters the structural stability of the 3D perovskite (31–33).

Fig. 1. Electrical performance characterization of perovskite TFTs.

(A) Schematic device structure of the perovskite TFT. Typical transfer curves of TFTs based on (B) CsFASnI₃, (C) TEAI-CsFASnI₃, and (D) TEASCN-CsFASnI₃ films. Output curves of (E) TEAI-CsFASnI₃ and (F) TEASCN-CsFASnI₃ TFTs. (G) Statistical histograms of mobility of TEAI-CsFASnI₃ and TEASCN-CsFASnI₃ TFTs. (H) Comparison of the average mobility of the TEASCN-CsFASnI₃ TFTs with previously reported perovskite TFTs.

Figure 1B illustrates the transfer curves of freshly prepared CsFASnI₃ perovskite TFTs, measured in both forward and reverse scanning directions after the perovskite semiconductor layer was patterned as shown in fig. S1. Consistent with previous reports, the CsFASnI₃ TFTs exhibit weak gate modulation, which is attributed to the inherent instability of the divalent tin (Sn²⁺) state, making it prone to oxidation. This oxidation results in unwanted P-type self-doping and an increased number of charge carriers, making the transistor difficult to turn off (25, 34, 35). Hall effect measurements confirm that the carrier concentration in the CsFASnI₃ films reaches as high as 10¹⁹ cm⁻³, indicating an excess of holes (fig. S2). In contrast, the TEAI-CsFASnI₃ perovskite TFTs exhibit notably improved performance, with a mobility of 12 cm² V⁻¹ s⁻¹, a subthreshold swing of 0.58 V/dec, a threshold voltage of 12 V, and a current on/off ratio of 1.6 × 10⁷ (Fig. 1C). The TEASCN-CsFASnI₃ perovskite TFTs show even further enhancement, with a mobility of 74 cm² V⁻¹ s⁻¹, a subthreshold swing of 0.48 V/dec, a threshold voltage of 7 V, and a current on/off ratio of 2.1 × 10⁸ (Fig. 1D). The observed reduction in threshold voltage is likely due to a decrease in hole sources, particularly a reduction in tin vacancies (V_Sn). Hall effect measurements indicate that, compared to the CsFASnI₃ film, the carrier concentration in the TEAI- and TEASCN-treated films has decreased by one to two orders of magnitude (fig. S2). This suggests that the additives effectively lower the carrier concentration, resulting in improved TFT performance. A more detailed discussion of the relationship between hole concentration and threshold voltage can be found in the Supplementary Materials. In addition, the details of mobility extraction are shown in fig. S3, and the specific optimization details of the additive content are provided in figs. S4 and S5. To further investigate the role of the TEA⁺ cation in modulating transistor performance, we explored the effects of other thiocyanate compounds with different cations, such as formamidine thiocyanate (FASCN), guanidine thiocyanate (GuaSCN), and phenylethyl ammonium thiocyanate (PEASCN), on the performance of tin-based perovskite films, as shown in fig. S6. Incorporating FASCN and GuaSCN into CsFASnI₃ films did not markedly improve the gate modulation, as evidenced by the weak gate response in the corresponding transfer curves. In contrast, the inclusion of PEASCN led to a noticeable enhancement in gate modulation, with the transistor exhibiting a mobility of 18 cm² V⁻¹ s⁻¹, although still notably lower than that of the TEASCN-based perovskite TFTs. These findings suggest that both the TEA⁺ cation and SCN⁻ anion in TEASCN play crucial roles in enhancing the electrical performance of tin-based perovskite transistors. Figure 1 (E and F) presents the output characteristic curves of the TEAI-CsFASnI₃ and TEASCN-CsFASnI₃ perovskite TFTs. The TEASCN-CsFASnI₃ TFT demonstrates a higher drain current, which implies an increased transconductance and enhanced device gain, thereby indicating more efficient modulation of the drain current in response to variations in the gate voltage. In addition, both devices operate in the accumulation-depletion mode, featuring distinct regions of low V_DS (linear) and high V_DS (saturation).

The contact resistance of TEASCN-CsFASnI₃ TFTs was extracted from the transfer curves of devices with different channel lengths by the transmission line method method (fig. S7) (36), exhibiting a small resistance of 35 ohm·cm. Such a low resistance suggests superb contact between the gold and perovskite semiconductor materials, enabling efficient carrier transport. Furthermore, we fabricated 100 separate TEAI-CsFASnI₃ and TEASCN-CsFASnI₃ perovskite TFTs to ensure repeatability, as illustrated in Fig. 1G and fig. S8. The corresponding distribution diagram shows a narrow range of mobility values for both TEAI-CsFASnI₃ and TEASCN-CsFASnI₃ perovskite TFTs. The TEASCN-CsFASnI₃ TFTs demonstrate a notably high average mobility of 63 cm² V⁻¹ s⁻¹, indicating the consistent attainment of high-quality tin-based perovskite TFTs. Overall, the TEASCN-CsFASnI₃ TFTs demonstrate substantial enhancements in all electrical parameters compared to the CsFASnI₃ TFTs. Compared to all types of perovskite transistors, the mobility and on/off ratio values are an excellent demonstration of superior electrical performance (Fig. 1H and table S1).

To investigate the contribution of TEASCN to the promoted hole mobility, we further assessed the chemical interaction mechanism between TEASCN and tin-based perovskites. Comprising an aminocation functional group (─NH₃⁺), a thiophene ring, and a thiocyanogen root (─S─C≡N), the TEASCN theoretically facilitates the formation of coordination bonds in tin-based perovskites. Specifically, the ─S─ bonds in TEA⁺ and the S and N atoms in SCN⁻ can form coordination bonds with Sn²⁺, stabilizing the tin atoms in the [SnX₆]⁴⁻ octahedron. In addition, the ─NH₃⁺ group in TEA⁺ can interact with negatively charged defects by electrostatic forces, inhibiting their migration and disabling their charge-trapping capabilities (37, 38). Fourier transform infrared (FTIR) spectroscopy was used to examine the functional groups in the additives and their coordination with SnI₂, as depicted in Fig. 2 (A and B). For the TEASCN powder, the characteristic stretching vibration peak of the C≡N bond of SCN⁻ is located at 2065.6 cm⁻¹. Because of the strong conjugation between the lone pair electrons of S and the thiophene ring, the stretching vibration of the C─S bond occurs at 706.9 cm⁻¹, representing a notable shift compared to that of TEAI (at 698.6 cm⁻¹). After the introduction of SnI₂, the N─H bending vibration peaks of TEASCN and TEAI experienced varying degrees of attenuation. Notably, the attenuation was more pronounced in TEASCN-SnI₂, suggesting a stronger molecular interaction between TEASCN and SnI₂. This interaction in turn exerted an influence on the molecular structure, consequently constraining the N─H bending (39, 40). Simultaneously, a notable shift in the thiophene signal originating from TEA can be observed, indicating the involvement of the thiophene ring in the interaction with Sn²⁺. When TEASCN and TEASCN + SnI₂ are dissolved in a mixed solution of N,N′-dimethylformamide (DMF) and dimethyl sulfoxide (DMSO), infrared spectra show a red shift in the SCN⁻ vibrational lines after SnI₂ is introduced. This red shift can be attributed to the partial donation of lone-pair electron density from the sulfur atom of SCN⁻ to a vacant orbital of the tin element (fig. S9). These interactions are a result of the synergistic action of coordination and hydrogen bonding, enhancing stability and promoting better interaction between TEASCN and the tin-based perovskite structure.

Fig. 2. Strong interactions between the TEASCN and SnI₂ and inhibition of Sn²⁺ oxidation.

(A and B) FTIR spectra of TEASCN, TEASCN + SnI₂, TEAI, and TEAI + SnI₂. (C) XRD diffractions of CsFASnI₃, TEAI-CsFASnI₃, and TEASCN-CsFASnI₃ perovskite films before annealing. XPS spectra in the Sn 3d regions of (D) TEASCN-CsFASnI₃, (E) TEAI-CsFASnI₃, and (F) CsFASnI₃ perovskite films. a.u., arbitrary units.

Furthermore, the impact of TEASCN on the crystallization process of tin-based perovskites was investigated. Before film annealing, the strong bonding between TEASCN and Sn²⁺ facilitated the nucleation and crystallization of a high-quality perovskite intermediate phase, controlling the direction of crystal growth to achieve the preferred (100) crystal orientation (Fig. 2C). The introduction of TEA⁺ induced the high-oriented crystallization of the 3D perovskite phase, while SCN⁻ promoted the generation of a bilayer (2L) quasi-2D perovskite intermediate phase [TEA₂(FACs)Sn₂SCN₂I₅]. This was further corroborated by PL and absorption spectra (fig. S10), which confirmed the substantial formation of the quasi-2D intermediate phase before annealing. After annealing, the 2L peak was weakened, and the 3D peak was enhanced (fig. S11). To further confirm the presence of any residual quasi-2D phase, we performed grazing-incidence x-ray diffraction (XRD) at different angles, which did not reveal any small peaks corresponding to the 2L phase (fig. S12A). Moreover, we conducted photoluminescence (PL) measurements (fig. S12B) and converted the data to a logarithmic scale (fig. S12C). A peak observed around 700 nm suggests the existence of a trace amount of the 2L structure within the bulk of the perovskite film. Subsequently, time-of-flight–secondary ion mass spectrometry (TOF-SIMS) characterization of the annealed perovskite films revealed that TEA⁺ is enriched at the surface (fig. S13). Given the “top-down” crystallization mode of the one-step spin-coated perovskite films and the substantial reduction of the 2L phase after annealing (41), we hypothesize that the initially formed 2L phase, aligned parallel to the substrate, acts as a template for the highly oriented vertical growth of the 3D structure. This 2L phase provides an ordered foundation for the vertical alignment of the 3D perovskite while simultaneously transitioning into a higher n value 3D phase. As TEA⁺ is incorporated during the formation of the 2L phase and enriched at the surface, it remains on the surface as the perovskite transitions to a higher n value phase. This phenomenon was observed exclusively when the cations of SCN-based compounds were capable of forming low-dimensional perovskite phases. As shown in fig. S6, incorporating FASCN and GuaSCN, which form 3D perovskite structures, did not result in notable improvements in crystallinity before and after annealing. In contrast, both PEASCN and TEASCN promoted the formation of low-dimensional intermediate phases, which, after annealing, led to a notable enhancement in the crystallinity of the perovskite films. In particular, the TEA⁺ cation, with its electron-withdrawing sulfur effect and its smaller, more compact structure compared to PEA⁺, caused a more marked reduction in the interlayer spacing of the perovskite, thereby improving carrier transport efficiency (42). Therefore, unlike other pseudo-halides, which regulate crystallization primarily through interactions with Sn²⁺ to slow down nucleation and reduce defects (43), thiocyanates with larger cations induce quasi-2D bilayer templates, guiding vertical crystal growth and phase evolution. This enhances carrier transport, while surface hydrophobic organic cations provide additional protection, improving both performance and stability.

The antioxidant capacity of the film was evaluated using x-ray photoelectron spectroscopy (XPS), as shown in Fig. 2 (D to F). The Sn⁴⁺ content in CsFASnI₃ films reaches as high as 28.1%. Upon introducing TEAI, the Sn⁴⁺ content in the film decreases to 20.8%, indicating the effective antioxidative capability of TEA⁺. In addition, the presence of SCN⁻ enhances the proportion of TEA⁺ cations at the surface, as evidenced by the TEA⁺ surface spectrum from TOF-SIMS of the perovskite films (figs. S13 and S14). This increased concentration of TEA⁺ on the surface, combined with its ability to block water and oxygen, contributes markedly to the improved antioxidative properties. As a result, the TEASCN-CsFASnI₃ film contains only 13% Sn⁴⁺, demonstrating superior antioxidative performance. The S 2p XPS spectrum of the TEASCN-CsFASnI₃ film reveals a doublet after 162.5 eV, corresponding to sulfur in the thiophene ring (44, 45), and another doublet before 162.5 eV, assigned to sulfur in SCN⁻, confirming the presence of SCN⁻ on the perovskite film (46), as shown in fig. S15. The SCN⁻/I molar ratio is ~4.1%, indicating a trace residual amount of SCN⁻ in the perovskite film. The TOF-SIMS depth profiles (fig. S16) of SCN⁻ in perovskite films annealed for different durations support this observation, confirming that SCN⁻ may evaporate during the annealing process. It also plays a crucial role in improving the surface concentration of TEA⁺, thereby enhancing the film’s stability and antioxidative properties (47–49).

With the increasing of films’ antioxidant capacity, we further ascertain the difference in crystallinity and crystal orientation between perovskite films with and without additives, using XRD and grazing-incidence wide-angle x-ray scattering (GIWAXS). The CsFASnI₃ film shows the weak (100) and (200) peaks from the cubic phase (Fig. 3E). Its GIWAXS pattern exhibits relatively weak and dispersed Debye-Scherrer rings corresponding to the (100) crystal facet, indicating poor crystallinity and random crystal orientation (Fig. 3A). The introduction of TEAI and TEASCN in the perovskite films results in markedly enhanced (100) diffraction peaks in the XRD spectrum and bright Bragg spots or short arcs corresponding to the respective diffraction planes in the GIWAXS pattern (Fig. 3, B and C), indicating an improvement in both crystallinity and orientation of the films.

Fig. 3. Enhanced orientation, crystallinity, and morphology of the perovskite films.

GIWAXS patterns of the (A) CsFASnI₃, (B) TEAI-CsFASnI₃, and (C) TEASCN-CsFASnI₃ perovskite films. (D) Azimuth curves of the integrated ring at q_r = 1.0 Å⁻¹. (E) XRD patterns of CsFASnI₃, TEAI-CsFASnI₃, and TEASCN-CsFASnI₃ perovskite films. (F) XRD patterns of CsFASnI₃, TEAI-CsFASnI₃, and TEASCN-CsFASnI₃ perovskite films with a partially enlarged view of the (100) plane. Typical SEM images of (G) CsFASnI₃, (H) TEAI-CsFASnI₃, and (I) TEASCN-CsFASnI₃ films.

Numerical calculations were performed to integrate the diffraction intensity of the (100) plane concerning the azimuthal angle (Fig. 3D). The results indicate that the diffraction peak signals of TEAI-CsFASnI₃ and TEASCN-CsFASnI₃ films are concentrated around 90°, while CsFASnI₃ exhibits a comparatively broader distribution. Notably, the TEASCN addition yields a superior enhancement in diffraction peak intensity compared to TEAI, underscoring the vital role of the SCN-induced 2L intermediate phase formation in promoting crystallinity. In addition, the comparison of XRD patterns between perovskite films incorporating TEASCN and TEAI reveals a small-angle shift in the 3D main peak (Fig. 3F). This shift suggests that SCN⁻ is successfully incorporated into the perovskite lattice, influencing the crystal structure.

From scanning electron microscopy (SEM) results, the CsFASnI₃ films display poor coverage, numerous pores, and a rough surface, which can be attributed to rapid crystal growth preceding nucleation completion (Fig. 3G). The incorporation of TEAI enhanced the film morphology, leading to notably broader coverage (Fig. 3H). Further, the film incorporating TEASCN shows a higher coverage and fewer pinholes, indicating improved uniformity and stability during transistor fabrication (Fig. 3I). The grain size statistics, presented in fig. S17, provides additional insight into the crystallization behavior of the films. The enhancement was corroborated by atomic force microscopy (AFM) images. As shown in fig. S18, the smoother surface facilitates enhanced interface quality between the dielectric layer, semiconductor layer, and electrode, reducing interface trap density and mitigating leakage current, enhancing the transistor’s switching speed and voltage driving capability (20, 50).

To further elucidate the suppressed defect density in tin-based perovskites introduced with TEASCN, we conducted steady-state PL measurements. The PL spectra in Fig. 4A indicate a twofold and fivefold increase in peak intensity with the introduction of TEAI and TEASCN, respectively, suggesting reduced nonradiative charge recombination and a more pronounced inhibitory effect of TEASCN on recombination (51, 52). To gain deeper insights into the carrier dynamics of perovskite films, we further analyzed their photoluminescence quantum efficiency (PLQE) and time-resolved photoluminescence (TRPL) spectra. The results demonstrate that the PLQE of the TEASCN-CsFASnI₃ film reaches 6.76%, which is higher than the 5.40% of the TEAI-CsFASnI₃ film and notably surpasses the 3.32% of the CsFASnI₃ film (fig. S19A). The TRPL decay curves, fitted using a biexponential function (Fig. 4B), show that the PL lifetime of the TEAI-treated film extends to 1.47 ns, surpassing the 0.94 ns of the pristine CsFASnI₃ film. The introduction of TEASCN further improves the PL lifetime to 3.85 ns, more than four times longer than that of the CsFASnI₃ film. These results suggest that the TEASCN treatment effectively suppresses nonradiative recombination, facilitating better interfacial carrier transport and contributing to enhanced transistor performance (53, 54). The electroluminescent device fabricated with TEASCN-treated films exhibited markedly improved EL intensity and a higher EQE_EL (2.09%), compared to the pristine CsFASnI₃ device (1.32%), further confirming the reduced nonradiative recombination (55). In addition, the current density-voltage (J-V) characteristics reveal lower turn-on voltages and suppressed leakage currents, which contribute to better efficiency and stability (fig. S19 and table S2) (56, 57). To characterize the intrinsic defects of the devices, we conducted low-frequency noise testing and temperature-dependent transfer curve measurements. The noise of all devices exhibits frequency dependence, manifesting as low-frequency 1/f noise, associated with voltage fluctuations arising from defect capture and release processes (58). Moreover, the noise current of the TEAI-CsFASnI₃ and TEASCN-CsFASnI₃ devices is not only lower than that of the CsFASnI₃ device at V_GS = 0 V and V_DS = −1 V but also exhibits lower normalized postsynaptic density values and more uniform voltage dependence (Fig. 4C and fig. S20), indicating that the additives effectively suppress device defects. Furthermore, we calculated the interface defect density of the perovskite TFTs (fig. S20). The results demonstrate that the defect density in the optimized perovskite transistors is on the order of 10¹⁴ to 10¹⁵ cm⁻² eV⁻¹, representing a reduction of three to four orders of magnitude compared to the pristine CsFASnI₃ devices. These findings confirm the effective defect passivation induced by TEASCN.

Fig. 4. Suppressed defect densities.

(A) PL and (B) TRPL spectra of CsFASnI₃, TEAI-CsFASnI₃, and TEASCN-CsFASnI₃ films. (C) Noise current curves of CsFASnI₃, TEAI-CsFASnI₃, and TEASCN-CsFASnI₃ devices at 0 V. (D) Temperature-dependent electrical transfer curves of typical CsFASnI₃, TEAI-CsFASnI₃, and TEASCN-CsFASnI₃ TFTs. (E) Temperature-dependent threshold voltage diagrams of TEAI-CsFASnI₃ and TEASCN-CsFASnI₃ TFTs.

From Fig. 4 (D and E), we observe that the CsFASnI₃ TFTs lack gate control ability even at 150 K, indicating a high defect density. In contrast, the TEAI-CsFASnI₃ TFTs exhibit a steeper slope in the V_TH versus T curve compared to the TEASCN-CsFASnI₃ devices within the temperature range of 250 to 300 K. This suggests that the introduction of TEAI affects the temperature dependence of the threshold voltage more notably than TEASCN. Since the FASnI₃ undergoes a cubic-to-tetragonal phase transition at around 225 K, we calculated the defect density at the temperatures between 250 and 300 K by the following equation: $D_{\mathrm{t}}=\left(C_{\mathrm{i}}\text{∂}{V}_{\text{TH}}\right)/\left(\mathit{ek}\text{∂}T\right)$ (59, 60), where D_t is the trap density, C_i represents the areal capacitance of the 100-nm SiO₂ dielectric layer, k is the Boltzmann’s constant, e is the basic charge, and T represents the temperature. The calculated defect density of TEASCN-CsFASnI₃ TFTs is 1.086 × 10¹⁴ cm⁻² eV⁻¹, which is much lower than that of TEAI-CsFASnI₃ (3.741 × 10¹⁴ cm⁻² eV⁻¹), indicating the influential role of SCN⁻ in reducing the device defect density.

According to the above discussion, the introduction of TEASCN in tin-based perovskite films can enhance oxidation resistance, reduce defect density, improve morphology, and increase crystallinity. These benefits contribute to the stability and reliability enhancement of transistors, prolonging the device’s life span. To further investigate the device stability, various gate voltages (V_GS = ±30 V) were applied to observe the switching behavior of the transistors. From Fig. 5 (A and B), the TFTs incorporating TEASCN and TEAI exhibited clear on/off states with consistent current responses throughout 1500 consecutive cycles. The current response with TEASCN showed greater uniformity, indicating excellent stability.

Fig. 5. Assessment of operational stability and environmental stability.

Continuous on/off switching test diagram (A) and local magnification diagram (B) of TEAI-CsFASnI₃ and TEASCN-CsFASnI₃ TFTs for 1500 consecutive cycles. (C) Changes of threshold voltage and on-state current of TEAI-CsFASnI₃ and TEASCN-CsFASnI₃ TFTs during cyclic transfer curve measurement. (D) Negative bias stress stability of TEAI-CsFASnI₃ and TEASCN-CsFASnI₃ TFTs (V_GS = V_DS = −30 V). (E) The transition curve of the TEASCN-CsFASnI₃ TFTs changes with the storage time in the nitrogen atmosphere. (F) The on-state current and mobility of the TEASCN-CsFASnI₃ TFTs change with storage time in the nitrogen atmosphere.

Figure 5C and fig. S21 further demonstrates the reliability and consistency of the TEASCN-CsFASnI₃ TFTs under various operating conditions, showing negligible variation even after 100 scans. Under negative bias stress of V_GS = V_DS = −30 V for 1000s (Fig. 5D), the output current of the TEAI-CsFASnI₃ TFTs dropped to 63% of its initial value, while the output current of the TEASCN-CsFASnI₃ TFT is remained at 97% of its initial value. It indicates the reduced charge capture, which is further confirmed by the nearly unchanged transfer curve after bias stress testing (fig. S22). Last, the TFT electrical performance changes after long-term storage was evaluated. As depicted in fig. S23, encapsulated TFTs by paraffin and stored in a vacuum for 15 days showed nearly identical transfer curves, while unencapsulated TFTs with TEAI stored in a nitrogen glove box exhibited notable degradation after 15 days, due to trace oxygen exposure (fig. S24). In contrast, unencapsulated TFTs incorporating TEASCN showed only a slight increase in on-state current after 30 days, with mobility retention at 84% of its original value (Fig. 5, E and F).

DISCUSSION

High mobility and stable tin-based perovskite TFTs are fabricated by incorporating TEASCN into CsFASnI₃ perovskites. The strong bonding interaction between TEASCN and CsFASnI₃ effectively suppressed the oxidation of Sn²⁺ to Sn⁴⁺, reducing defect density and prolonging carrier lifetime. Furthermore, this strong interaction induced the formation of a bilayer quasi-2D perovskite intermediate phase, controlling the crystallization kinetics of perovskite. It substantially promotes highly oriented growth and improves the film morphology and crystallinity, thus facilitating more efficient carrier transport. Benefiting from these improvements, the optimized devices exhibited an average mobility of 63 cm² V⁻¹ s⁻¹ and high operational stability, maintaining stable performance in a nitrogen environment for up to 30 days. This study paves the way for developing high-performance perovskite TFTs with excellent stability.

MATERIALS AND METHODS

Preparation of the 0.8 M SnI₂ solution

First, 2.538 g of iodine (I₂) was dissolved in 2.5 ml of DMSO and stirred until completely dissolved. Then, 10 ml of DMF was added to the solution, followed by further stirring to ensure thorough mixing, resulting in a dark purple solution. Subsequently, an excess amount of 1.5 g of tin powder was added to the mixture, and stirring was continued. The solution gradually turned yellow-green and was stirred further until fully reacted. Afterward, it was left undisturbed until the excess tin powder completely settled. The supernatant was collected as needed for further use. The comparison between SnI₂ prepared by solution and commercial SnI₂ is shown in fig. S25 and discussed in detail in the Supplementary Materials.

Preparation of perovskite films

The synthesis of TEASCN is detailed in the Supplementary Materials, with its nuclear magnetic resonance data presented in fig. S26. TEAI (1.6 M was dissolved in DMF), TEASCN (1.6 M was dissolved in DMF), formamidinium iodide (FAI) (0.8 M was dissolved in DMF), CsI (0.8 M was dissolved in DMSO), and SnI₂ [SnF₂ (0.08 M) was added] were used as the mother solution and heated at 60°C overnight. For CsFASnI₃, CsI, FAI, and SnI₂ were mixed in a precise volume ratio of 0.15:0.85:1 and stirred thoroughly to ensure uniformity, forming a 0.2 M solution in a DMF:DMSO (4:1) solvent mixture. For TEAI-CsFASnI₃, TEAI, CsI, FAI, and SnI₂ were mixed in a precise volume ratio of x:0.15(1 − x):0.85(1 − x):1, where x represents the volume fraction of TEAI, and stirred thoroughly to ensure uniformity. For TEASCN-CsFASnI₃, TEASCN, CsI, FAI, and SnI₂ were mixed in a precise volume ratio of x:0.15(1 − x):0.85(1 − x):1, where x represents the volume fraction of TEASCN. For example, the optimal ratio of TEASCN in this study is 12.5%, corresponding to x = 0.125. The respective volume ratios of TEASCN, CsI, FAI, and SnI₂ in this case are 0.125:0.131:0.744:1.

TFT device fabrication

Highly doped P-type silicon with 100 nm of SiO₂ was used as the substrate. The Si/SiO₂ substrate was treated under argon plasma for 20 s to improve the surface wettability. The perovskite precursor was spin-coated onto the substrate at 5000 rpm for 60 s, respectively. Then, 100 μl of antisolvent chlorobenzene was dropped onto the surface, followed by thermal annealing at 100°C for 10 min. The source/drain electrodes of Au (50 nm) were deposited by thermal evaporation through shadow masks. The channel length/width of the TFTs was 150/1200 μm. All preparation processes were completed in a nitrogen-filled glove box.

Perovskite film characterization

The GIWAXS measurement was performed by using a beam energy of 10 keV and a PILATUS detector at the BL02U1 beamline of Shanghai Synchrotron Radiation Facility, Shanghai, China. The SEM images were taken using JSM-7800, JEOL. XRD pattern data for 2θ values were collected with a Bruker AX D8 Advance diffractometer with nickel-filtered Cu Kα radiation (λ = 1.5406 Å). AFM measurements were taken using Bruker’s Dimension Edge03040155. The TRPL spectra in the manuscript were obtained using a Horiba Jobin Yvon TRPL system. The excitation source was a Ti:sapphire laser with a pulse width of 35 fs and a frequency of 2 kHz. The excitation light was frequency-doubled to 400 nm to excite the thin films, and the emitted signal was collected by a streak camera (Hamamatsu, C6860). The PL and PLQE measurements were recorded on an Edinburgh Instruments FLS1000 spectrofluorometer equipped with a 450-W xenon lamp and an integrating sphere. The excitation wavelength was 475 nm, with the laser intensity set to 200,000 counts per second. The spectrofluorometer’s detector, a PMT-1010, had a response width of 600 ps, and the spectral range was from 300 to 1000 nm. The thickness of the film measured with the stepper (Zeptools JS100A) is between 35 and 40 nm. The electrical and low-frequency noise signals of the perovskite transistors are measured in a probe station equipped with a semiconductor parameter analyzer (FS-Pro) in a nitrogen-filled glove box.

Supplementary Materials

This PDF file includes:

Figs. S1 to S26

Tables S1 and S2

References