Self-assembled manganese acetate@tin dioxide colloidal quantum dots as an electron-transport layer for efficient and stable perovskite solar cells

Abstract

An SnO₂ electron-transport layer (ETL) is a critical and mainstream component utilized in perovskite solar cells (PSCs) on account of its exceptional photoelectric performances and low-temperature fabrication process. Nevertheless, a significant number of defects on the surface of SnO₂ quantum dots (QDs) can lead to QD agglomeration in dispersion and poor film quality, ultimately compromising the power conversion efficiency (PCE) and stability of devices. In this study, we developed multifunctional manganese acetate (Mn(Ac)₂)-stabilized SnO₂ quantum dots (Mac-SnO₂ QDs) and utilized them to prepare ETLs for PSCs. The C=O groups of Mn(Ac)₂ can terminate the Sn⁴⁺ dangling bonds on Mac-SnO₂ QDs, thereby preventing QD agglomeration, facilitating the preparation of a uniform film with good electrical properties. In addition, Mn²⁺ can permeate upward into the perovskite film, passivating I⁻ at the grain boundary and reducing perovskite bulk defects. Moreover, the perovskite grown on Mac-SnO₂ releases residual stress, which can improve device performance and stability. As a consequence, Mac-SnO₂ PSCs achieved a high PCE of 23.36% with improved stability. This successful approach of utilizing an organic acid salt to modify SnO₂ QDs provides a facile and efficient strategy for enhancing the efficiency and stability of PSCs.

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1. Introduction

Interest in organic–inorganic hybrid perovskite solar cells (PSCs) has recently surged in the photovoltaic field owing to their remarkable optoelectronic performances, such as high carrier mobility, low exciton binding energy, long carrier diffusion length, and excellent absorption coefficient.^1–3 Since the advent of the first PSCs, their certified power conversion efficiency (PCE) has boomed to 26.1% over the past dozen years, which is comparable to mature commercial silicon cells.^4–6 One of the critical factors for achieving advanced devices is the application of inorganic metal oxides, such as ZnO, SnO₂, and TiO₂, as electron-transport layers (ETLs).^7–10 These ETLs play a key role in extracting and separating carriers, have high transmittance to allow light into the absorbed layer, and display well-matched energy levels with the perovskite material. Currently, SnO₂ stands out among the many ETLs because of its excellent photoelectric properties, high optical transmittance, exceptional ultraviolet (UV) stability, and suitability for low-temperature preparation.^11–15 However, similar to other inorganic metal oxides, the surface of SnO₂ quantum dots (QDs) possesses a significant number of metallic dangling bonds, leading to disordered surface charges. Consequently, QDs will attract each other because of van der Waals interaction, resulting in instability and agglomeration of the QDs.^16,17 Upon film formation, the material suffers from pinholes and non-uniformity, leading to poor film quality. Moreover, the presence of Sn dangling bonds on the surface can capture hydroxyl groups during film formation and induce perovskite degradation.^18–20 These surface defects inevitably trap electrons, resulting in carrier non-radiative recombination, thus affecting interface electron transport.

To mitigate film surface defects, some researchers have adopted the method of surface modification for SnO₂ films. For instance, a multi-fluorine-containing C₇₀-porphyrin dyad molecule (F₇₀PD), bathocuproine (BCP)-core-based nonconjugated polyelectrolytes, histamine diiodate (HADI), and ionic potassium L-aspartate (PL-A) are considered suitable modification layers for the SnO₂ surface.^21–24 Their multifunctional groups can interact with surface-uncoordinated Sn⁴⁺ of the SnO₂ layer, enhance interfacial charge transport, and ultimately restrain carrier non-radiative recombination. Moreover, fullerene dimers prepared by combining phenanthroline with two fullerenes (2C₆₀-Bphen), oteracil potassium (OP), and a potassium sulfonate salt (3Cl-BSAK) have been used to match the energy level between the SnO₂ ETL and perovskite layer, thus effectively decreasing the interface energy barrier and enhancing interfacial charge extraction.^25–27 However, the surface modification of SnO₂ films may not sufficiently passivate the bulk defects of the SnO₂ film. As a result, some researchers have focused on modifying the surface of individual SnO₂ QDs for the sake of enhancing the overall quality of SnO₂ films. Organic large molecules, such as streptomycin sulfate, heparin potassium, and poly(ethylene glycol) diacrylate, have been incorporated to significantly passivate the defects of SnO₂ QDs and inhibit their agglomeration through their multiple functional groups.^28–30 This method can result in a uniform SnO₂ film. However, the large organic molecules may induce additional steric hindrance, affecting the ETL conductivity.³¹ Consequently, utilizing small ions to stabilize QDs seems to be a potential way to relax the trade-off between the surface passivation and reducing the steric hindrance.

In this study, we developed multifunctional manganese acetate (Mn(Ac)₂)-stabilized SnO₂ QDs (Mac-SnO₂ QDs) and utilized them to prepare ETLs for PSCs. The molecular formula of Mn(Ac)₂ is (CH₃COO)₂Mn. The C=O groups of the Mac-SnO₂ QDs could terminate Sn dangling bonds, which reduces particle interactions, suppressing QDs agglomeration, and allowing forming a uniform film. Moreover, the Mac-SnO₂ QDs could reduce the defects that exist in the ETL/perovskite interface due to the C=O group. Interestingly, during the experiment, we observed that the perovskite grown on Mac-SnO₂ released residual stress during annealing and the Mn²⁺ could spread upward into the perovskite layer, passivating the I⁻ that appear in the grain boundary. In consequence, the PCE could be elevated from 21.11% to 23.36%. More importantly, the Mac-SnO₂ devices demonstrated good long-term stability, with the PCE maintaining 90% of its original value under environmental conditions for nearly 600 h.

2. Results and discussion

In Fig. S1,† X-ray diffraction (XRD) was utilized to confirm the ETLs formation and crystallographic characteristics. The results demonstrate that the addition of Mn(Ac)₂ did not alter the SnO₂ crystal structure. Subsequently, the size variation between Mac-SnO₂ and SnO₂ QDs was investigated using transmission electron microscopy (TEM). The distinct electron diffraction circles observed for both samples indicated a polycrystalline nature.³² It was found that both the SnO₂ and Mac-SnO₂ QDs had sizes of ∼5 nm (Fig. 1a and b). One step further, the particle size analysis in the SnO₂ and Mac-SnO₂ colloidal dispersion was performed by dynamic light scattering (DLS) (Fig. 1c and d). The average particle sizes for the fresh SnO₂ and Mac-SnO₂ dispersions were similar, respectively. After 72 h of aging, the particle size of the fresh SnO₂ solution significantly increased, while that of the Mac-SnO₂ solution remained almost unchanged under identical conditions. Additionally, we measured the zeta potential to exhibit the electrostatic potential of the particles in solution. The fresh SnO₂ dispersion showed a zeta potential value of −23.60 mV, which declined to −9.94 mV after aging for 72 hours in storage (Fig. S2†). This indicates the electrostatic potential could not provide enough repulsion to disperse the particles, leading to unfavorable agglomeration (a common occurrence for inorganic metal oxide particles in colloidal dispersions).³³ On the contrary, the stabilized Mac-SnO₂ solution exhibited a large zeta potential (−40.5 mV), and was little changed after aging, suggesting that the addition of Mn(Ac)₂ stabilized the SnO₂ QDs by terminating the surface defects, thereby inhibiting QDs agglomeration.³⁴ Furthermore, the more evenly dispersed Mac-SnO₂ particles benefit the formation of a denser and more uniform film. Top-view scanning electron microscopy (SEM) and atomic force microscopy (AFM) were carried out to illustrate the ETLs surface morphology. The Mac-SnO₂ film displayed a compact and dense morphology, superior to that of the SnO₂ film (Fig. S3†). The surface root-mean-square (R_q) roughness decreased from the SnO₂ film (1.26 nm) to the Mac-SnO₂ film (0.95 nm), manifesting the smoother film (Fig. 1e and f). Moreover, UV–visible (UV–vis) transmittance spectroscopy was performed to probe the optical properties of the SnO₂ and Mac-SnO₂ films (Fig. 1g). The optical transmittance of the ITO/Mac-SnO₂ film increased in the range from 470 to 800 nm, verifying that the Mac-SnO₂ film was uniform and compact due to the more dispersed QDs.

Fig. 1 HRTEM images and their SAED images of (a) SnO₂ and (b) Mac-SnO₂. (c) Particle-size distribution of fresh SnO₂ and Mac-SnO₂ solutions. (d) Particle-size distribution of 72 h-aged SnO₂ and Mac-SnO₂ solutions. AFM topographical images of (e) SnO₂ and (f) Mac-SnO₂ films. (g) Optical transmittance spectra of ITO/SnO₂ and ITO/Mac-SnO₂ films.

In order to further elucidate the mechanism behind the differences between SnO₂ and Mac-SnO₂, we conducted a comprehensive study from various perspectives. Our findings revealed that the dispersion of SnO₂ was transparent and colorless, whereas the Mac-SnO₂ dispersion exhibited an orange hue, indicative of an interaction between Mn(Ac)₂ and SnO₂ (Fig. S4†). The interaction was analyzed through X-ray photoelectron spectroscopy (XPS) measurements (Fig. S5a†). The Mn peak of 2p could be observed in the Mac-SnO₂ film (Fig. S5b†). Notably, after the incorporation of Mn(Ac)₂, the peaks for Sn 3d_3/2 and Sn 3d_5/2 in SnO₂ were shifted toward lower binding energies, from 494.88 and 486.48 eV to 494.45 and 486.04 eV, respectively (Fig. 2a). This shift indicated the increased electron density around the Sn atoms, attributable to the interaction between Sn and Mn(Ac)₂.³⁵ Additionally, XPS for O 1s was also performed (Fig. 2b and c). The pristine SnO₂ film displayed a distinct peak at 530.38 eV, confirming the presence of hydroxyl groups (OH⁻), while the peak of saturated oxygen (O²⁻) was at 531.47 eV. However, the OH⁻ peak was diminished in the Mac-SnO₂ film, suggesting that C=O effectively reduced the density of absorbed OH⁻ sites. To further confirm the bonding between SnO₂ and Mn(Ac)₂, Fourier-transform infrared (FTIR) spectroscopy was conducted (Fig. 2d). Mn(Ac)₂ exhibited two characteristic peaks for C=O bonds.³⁶ After adding Mn(Ac)₂ into the SnO₂ QDs dispersion solution, the C=O peaks moved from 1558 and 1444 cm⁻¹ to 1543 and 1412 cm⁻¹, consistent with the XPS results mentioned earlier, and further supporting the formation of Mac-SnO₂ resulting from the interaction between C=O and the SnO₂ QDs surface. This effect can reduce the following presence of OH⁻ groups absorbed on the oxygen vacancies and inhibit the non-radiative recombination of the SnO₂ surface, offering potential benefits.³⁷ We further verified our experimental results via density functional theory (DFT) calculations. As displayed in Fig. 2e, there were oxygen vacancies on the SnO₂ surface. After adding acetic acid groups (Ac⁻), the C=O in acetic acid could combine with Sn on the SnO₂ surface (Fig. S6†). The interaction was then explored by calculating the charge density difference of Ac⁻ adsorbed on the SnO₂ surface (Fig. 2f). In the figure, the blue regions denote electron depletion after charge redistribution, and the yellow regions represent electron accumulation. It could be concluded that the Sn atoms obtained electrons from the C=O group through the Sn–O bond from the electron clouds at the interface, which was in accord with the XPS and FTIR results. A schematic drawing shows the whole process in Fig. 2g; whereby, the pristine SnO₂ QDs agglomerated due to a good deal of Sn⁴⁺ dangling bonds existing on the surface, while the C=O groups of Mn(Ac)₂ could passivate the Sn⁴⁺ dangling bonds, so that Ac⁻ inhibited the QDs agglomeration, like the ligands around SnO₂, thus allowing obtaining a dense Mac-SnO₂ film without pinholes.

Fig. 2 XPS spectra of (a) Sn 3d and (b and c) O 1s for SnO₂ and Mac-SnO₂. (d) FTIR spectra of SnO₂ and Mac-SnO₂. Side view of (e) SnO₂–O_v and the charge density difference of (f) SnO₂–O_v/Ac⁻. (g) Schematic of the distribution of SnO₂ and Mac-SnO₂ QDs in the solutions and films.

Electrical properties are a key factor to evaluate the performance of ETLs. We determined the electrical conductivities of the two films by using the current–voltage (I–V) characteristic curves. The device structure of ITO/ETL/Ag was used in the test, and the specific values were as 2.81 × 10⁻⁵ S cm⁻¹ (SnO₂) and 5.68 × 10⁻⁵ S cm⁻¹ (Mac-SnO₂), respectively (Fig. S7a†).³⁸ Additionally, we obtained the electron mobility by fabricating the space charge-limited current (SCLC) model (structure: ITO/ETL/PCBM/Ag) and by using the Mott–Gurney equation (Fig. S7b†):³⁹

where the physical quantities ε₀, ε_r, L, J, and V in the formula represent the vacuum permittivity, relative dielectric constant, ETL thickness, dark current density, and applied voltage, respectively. It was verified that the electron mobility of the Mac-SnO₂ ETL was 5.23 × 10⁻³ cm² V⁻¹ S⁻¹, which was significantly larger than that of the SnO₂ ETL (9.03 × 10⁻³ cm² V⁻¹ S⁻¹). The increased conductivity and electron mobility of the Mac-SnO₂ film were ascribed to inhibited non-radiative recombination and the denser ETL film without pinholes, demonstrating the electron extraction and transfer efficiency were higher.⁴⁰

We next investigated the surface morphology and structure of the perovskite layers based on the two types of ETLs by SEM (Fig. 3a and b). Their cross-sectional images are displayed in Fig. 3c and d. The Mac-SnO₂/perovskite grains were larger than those of the SnO₂, and the Mac-SnO₂/perovskite layer appeared denser and more uniform, while the control film exhibited undesirable pinholes. Consistent with a more homogeneous morphology, the AFM images show the R_q of the Mac-SnO₂/perovskite films (8.56 nm) was slightly reduced compared to that of the control film (9.69 nm), indicating a smoother and superior surface quality (Fig. S8†). The effect of different ETLs on the crystallization of perovskite was further explored by XRD (Fig. 3e). Two characteristic peaks at 14.0° and 28.2° were observed, indicating the (110) and (220) cubic phase of the perovskite film. Compared to the diffractions of the (110) and (220) lattice planes on the SnO₂/perovskite film, those on the Mac-SnO₂/perovskite film were stronger, which means the Mac-SnO₂/perovskite possessed enhanced crystallinity. The different crystallizations of the perovskite films may stem from the difference in the surface groups on the ETLs. We mirrored the surface Gibbs free energies via performing contact angle analysis of the two types of ETLs (Fig. 3f). The contact angle on Mac-SnO₂ (13.64°) was found to be lower than that on SnO₂ (22.89°), indicating the perovskite deposited on Mac-SnO₂/ETL had a lower surface Gibbs free energy and higher surface wettability.⁴¹ Lower Gibbs free energies are beneficial for the spreading of the perovskite precursor solution, uniform nucleation, and the formation of a high-crystallinity perovskite layer.⁴²

Fig. 3 Top-view and cross-section SEM images of perovskite layers grown on (a and c) SnO₂ and (b and d) Mac-SnO₂ ETL. (e) XRD patterns of the SnO₂/perovskite and Mac-SnO₂/perovskite. (f) The water contact angles on SnO₂ and Mac-SnO₂ films.

Furthermore, we conducted cross-sectional elemental mapping of the Mac-SnO₂/perovskite films (Fig. S9†), and found that Pb and I were evenly distributed within the perovskite, indicating uniform film growth. Surprisingly, we also detected the presence of Mn²⁺ in the perovskite film, suggesting that Mn²⁺ had diffused upward into the perovskite films. Previous reports indicate that Mn²⁺ can passivate I⁻ dangling bonds at perovskite grain boundaries, leading us to speculate that the diffused Mn²⁺ may passivate the I⁻ at the grain boundary.^43,44 XPS measurements were next applied to investigate the interaction between Mac-SnO₂ and the perovskite (Fig. S10†), and so Mn(Ac)₂ was added to the perovskite for XPS characterization. As exhibited in the I 3d spectrum (Fig. 4a), the peaks for I 3d_5/2 and I 3d_3/2 in the perovskite film (618.65 and 630.13 eV) moved toward higher binding energies (618.91 and 630.40 eV) after adding Mn(Ac)₂, which verified the Mn²⁺ and the uncoordinated I⁻ interact with each other, as we had estimated. Meanwhile, it could be noted that the Pb 4f_5/2 and Pb 4f_7/2 peaks (142.86 and 137.98 eV) were shifted to lower binding energies of 142.69 and 137.81 eV in the Mn(Ac)₂-added perovskite, implying the interactions between Pb²⁺ of perovskite and the C=O of Mn(Ac)₂ (Fig. 4b).⁴⁵ In addition, we found a peak for Pb⁰ in the perovskite sample, which disappeared in the Mn(Ac)₂-added perovskite sample, indicating that the interaction of Pb²⁺ dangling bonds and the C=O groups suppressed the Pb⁰ production. Afterwards, we evaluated the charge trap density (N_trap) of the perovskite films by adopting the SCLC model (electron-only device: ITO/ETL/perovskite/PCBM/Ag) (Fig. 4c). The formula used to calculate this is as follows:⁴⁶

where L, e, ε, and V_TFL represent the thickness of the perovskite films, elementary charge, relative dielectric constant of the perovskite, and the onset voltage of the trap-filled limit region, respectively. The calculated N_trap of the perovskite was appreciably reduced from 3.37 × 10¹⁵ (SnO₂/perovskite) to 2.27 × 10¹⁵ cm⁻³ (Mac-SnO₂/perovskite), which corresponded to the higher quality perovskite layer without pinholes and the reduced non-radiative recombination.

Fig. 4 XPS spectra of (a) Pb 4f and (b) I 3d for the perovskite and Mn(Ac)₂ + perovskite. (c) Dark I–V curves based on SnO₂ and Mac-SnO₂. Depth-dependent GIXRD patterns of the (d) SnO₂/perovskite and (e) Mac-SnO₂/perovskite, and their (f) linear fits of 2θ − sin² Ψ. (g) Residual stress diagram of the SnO₂/perovskite and Mac-SnO₂/perovskite.

In addition, it is known that film stress can accelerate perovskite film degradation and device failure.⁴⁷ Consequently, we explored the residual stress of the SnO₂/perovskite and Mac-SnO₂/perovskite films by depth-dependent grazing incident X-ray diffraction (GIXRD).^48–50 Variable grazing tilt angles (ψ) from 10° to 50° were selected to detect the different depths of the perovskite films. As the incidence angle increased, the characteristic peak (110) in the SnO₂/perovskite film gradually shifted to lower angles due to the compressive stress caused by lattice contraction (Fig. 4d). On the contrary, there was a slight peak shift for the Mac-SnO₂/perovskite film, implying a release of interfacial stress owing to the effective passivation of the uncoordinated Pb²⁺ at the perovskite buried interface (Fig. 4e).⁵¹ Subsequently, we fitted the 2θ as a function of sin² Ψ to evaluate the lattice strain variation more accurately, where the slopes of the fitting line represent the residual strain and were calculated by the 2θ − sin² Ψ curves (Fig. 4f). Also, the elastic stress ε can be calculated as follows:⁵²

where C is the constant and E is the Young's modulus. The fitted curve had a negative slope in the SnO₂/perovskite film, demonstrating severe strain; while the slope was closer to 0 for the Mac-SnO₂/perovskite film, implying the residual strain was effectively eliminated. Fig. 4 g shows that the residual strain in the perovskite was significant, and there was little residual strain in the Mac-SnO₂/perovskite film.

The energy level alignment between the ETL and perovskite determines the effectiveness of charge extraction and the recombination rate. Consequently, we examined the energy level of SnO₂ and Mac-SnO₂ by ultraviolet photoelectron spectroscopy (UPS). A decreased work function (W_F) of the Mac-SnO₂ was observed, calculated to be −3.12 eV, while the value for the pristine SnO₂ was −3.16 eV (Fig. 5a and b).⁵³ The valence band maximum (VBM) for both were −7.90 eV (SnO₂) and −7.81 eV (Mac-SnO₂), respectively. The optical band gap (E_g) values were determined using a Tauc plot derived from the UV–vis absorption spectroscopy analysis (Fig. 5c), showing values of 3.78 eV (SnO₂) and 3.77 eV (Mac-SnO₂). The conduction band minimum (CBM) was derived from the VBM and the E_g, and were calculated to be −4.12 eV for SnO₂ and −4.04 eV for Mac-SnO₂. The CBM of Mac-SnO₂ was closer to the CBM of the perovskite, resulting in a smaller energy loss during electron transfer. As displayed in Fig. 5d, the energy band diagram demonstrated that the Mac-SnO₂ was more beneficial for enhancing electron extraction and preventing hole transport. We further probed the surface potential of the SnO₂ and Mac-SnO₂ ETL via applying Kelvin probe force microscopy (KPFM) (Fig. S11†). The average surface potential of the Mac-SnO₂ surface (575 mV) was higher than that of the SnO₂ film (460 mV), which was in line with the consequences of UPS. Additionally, we examined the interfacial electron transmission processes of the control and Mac-SnO₂/perovskite films by measuring the steady-state photoluminescence (PL) (Fig. 5e). Significant PL quenching was noticed in the Mac-SnO₂/perovskite film compared to the control samples, indicating a shortened carrier recombination lifetime and enhanced electron extraction.⁵⁴ Besides the PL, in order to further explore the carrier dynamics of the perovskite, time-resolved photoluminescence (TRPL) analysis was performed. As depicted in Fig. 5f, the Mac-SnO₂/perovskite film displayed a shorter lifetime in the first stage (τ₁) (10.86 ns) than that of SnO₂ (21.79 ns). The shortened lifetime indicated efficient charge transport at the Mac-SnO₂/perovskite interface due to the improved interface contact and band alignment.⁵⁵

Fig. 5 UPS spectra of (a) SnO₂ and (b) Mac-SnO₂. (c) Tauc plots of SnO₂ (left) and Mac-SnO₂ (right). (d) Energy band structure diagram of the ETL and perovskite. (e) PL and (f) TRPL spectra of perovskite layers deposited on glass/SnO₂ and glass/Mac-SnO₂ substrates.

We fabricated complete PSCs with a structure of ITO/ETL/perovskite/HTL (Spiro-OMeTAD)/Ag to explore the influence on the device photovoltaic performance (Fig. 6a). The corresponding SEM images of devices based on the two ETLs are presented in Fig. 6b. Fig. 6c shows the reverse scanning current density–voltage (J–V) curve of the two champion devices. The introduction of Mac-SnO₂ led to a champion PCE of 23.36% [short-circuit current density (J_sc) = 24.96 mA cm⁻², open-circuit voltage (V_oc) = 1.16 V, fill factor (FF) = 80.69%], obviously higher than the reference PSC based on SnO₂ (PCE = 21.11%, J_sc = 24.25 mA cm⁻², V_oc = 1.12 V, FF = 77.66%). The photovoltaic parameters of the forward scans are shown in Table S1.† The V_oc and FF displayed significant improvements, which could be mainly ascribed to the reduced bulk and interfacial non-radiative recombination, and the suitable energy band alignment of the Mac-SnO₂/perovskite film, leading to rapid charge carrier extraction. Compared with the pristine sample, the PSC devices based on Mac-SnO₂ demonstrated lower hysteresis, which was likely due to the efficient carrier transport and reduced non-radiative recombination (Fig. S12†). Furthermore, we varied the concentrations of the Mac-SnO₂ solution to determine the optimal PCE (Fig. S13†). The relevant parameters are exhibited in Table S2,† indicating that the best PCE was obtained when the Mn(Ac)₂ concentration was 7.5 mg mL⁻¹. Subsequently, we fabricated 20 individual devices based on the SnO₂ and the optimized Mac-SnO₂ layers. The statistical histogram of the PCEs and other parameters are displayed in Fig. 6d, and Tables S3 and S4.† Based on the parameters, it could be concluded that the devices based on different ETLs all exhibited excellent reproducibility, with the Mac-SnO₂ ETL devices showing higher PCEs. A comparative analysis of the maximum power point (MPP) was conducted to understand the steady PCE outputs and PCEs for SnO₂ and Mac-SnO₂ devices, respectively (Fig. 6e). The Mac-SnO₂ device exhibited a stable PCE of 21.61% over 100 s, which the J–V results above accorded with. Meanwhile, the SnO₂ device demonstrated a stable PCE of 19.51% over time, indicating that devices with Mac-SnO₂ ETLs offered greater operational stability. The external quantum efficiency (EQE) spectra of the PSCs and the integrated current density curve are exhibited in Fig. 6f, while the EQE spectra of cells with different Mac-SnO₂ concentrations are shown in Fig. S14.†

Fig. 6 (a) Schematic structure of a Mac-SnO₂-based PSC. (b) SEM images of the cross-sections of the SnO₂- and Mac-SnO₂-based devices. (c) J–V curves of the optimal SnO₂ and Mac-SnO₂ cells. (d) PCE distributions of the SnO₂- and Mac-SnO₂-based devices from 20-device statistics, respectively. (e) Steady-state output of the SnO₂ and Mac-SnO₂ PSCs measured at MPP. (f) EQE spectra of the SnO₂- and Mac-SnO₂-based devices.

The performance of cells is closely associated with the trap-assisted recombination behavior. We performed light intensity-dependent V_oc analysis to gain insights into the mechanism of the carrier recombination process (Fig. S15a†). The relationship where V_oc is dependent on the light intensity is as follows:⁵⁶

where n, k, and I represent the ideal coefficient, Boltzmann constant, and light intensity, respectively. Compared to the n value of 1.36 for PSCs with pristine SnO₂, the lower n value for the Mac-SnO₂ device (1.12) indicated that trap-assisted Shockley–Read–Hall (SRH) monomolecular recombination was inhibited. Meanwhile, the light intensity-dependent J_sc analysis was performed (Fig. S15b†). In general, the expression J_sc ∝ I^α describes the relationship of the J_sc with the light intensity.⁵⁷ The α of the Mac-SnO₂ PSC (0.99) was closer to 1 compared to the pristine SnO₂ device (0.98), which indicated the negligible bimolecular recombination of the Mac-SnO₂ device. The recombination rates of PSCs were also explored by conducting transient photovoltage (TPV) measurements (Fig. S16†). The Mac-SnO₂ device demonstrated a longer lifetime (263 μs), while that of the SnO₂ device was 126 μs. The slower carrier recombination was attributed to the reduced trap-assisted recombination, leading to the higher V_oc in the Mac-SnO₂ devices.⁵⁸ Furthermore, carrier migration owing to defects in PSCs will result in the formation of dark current. The Mac-SnO₂ devices displayed a lower leakage current, indicating inhibited non-radiative recombination at the interface (Fig. S17†).⁵⁹ Subsequently, electrochemical impedance spectroscopy (EIS) measurements were utilized to explore the carrier recombination and carrier-transport behaviors in the two types of devices (Fig. S18†). It is well known that the high-frequency ranges in a Nyquist diagram is associated with transfer resistance (R_tr) at the ETL/perovskite interface, and the high frequency represents the recombination resistance (R_rec).⁶⁰ The Mac-SnO₂ devices exhibited a relatively lower R_tr value but higher R_rec value than that of the control devices, demonstrating a low electron-transfer barrier and restrained recombination at the interfaces. The built-in potential (V_bi) and the carrier-transport behavior of PSCs can be estimated via capacitance–voltage (C–V) measurements (Fig. S19†). The extracted V_bi values of the SnO₂ and Mac-SnO₂ devices were approximately 0.95 and 1.07 V, respectively. A higher V_bi is conducive to extending the depletion region to increase the driving force for photogenerated carriers and accelerate carrier transport, which were in accordance with the V_oc values for the PSCs.⁶¹ The above discussions of the PSCs shows that the Mac-SnO₂ ETL could effectively improve the interface electron transfer and suppress non-radiative recombination.

Besides the photovoltaic properties, the stability of cells is also crucial for practical applications. Consequently, the long-term stability of the two cells under the ambient environment (25–30 °C, 30%–40% humidity) was monitored. As depicted in Fig. 7a, the Mac-SnO₂ PSC possessed an increased life. After storage for 600 h, the Mac-SnO₂-based PSCs maintained 90% of the initial PCE, whereas the SnO₂ PSC degraded faster, dropping to 55% of its initial efficiency. Moreover, the XRD patterns of the SnO₂/perovskite film exhibited a clear peak of PbI₂ after 20 days under ambient conditions (Fig. 7b). It was also observed that part of the SnO₂/perovskite film changed color from black to yellow, while the Mac-SnO₂/perovskite film maintained a well-preserved structure and black color (Fig. S20†). We also tested the light stability and thermal stability of devices. After storing the device in a N₂ atmosphere at 85 °C for 150 h, the efficiency of the Mac-SnO₂ device retained 84% of its original PCE, while the SnO₂ device quickly dropped to 59% of its original PCE (Fig. S21a†). In addition, after lighting the devices continuously for 60 h under 1 sun illumination, the Mac-SnO₂-based device retained 78% of the original PCE, while the SnO₂-based device retained only 61% of the original efficiency (Fig. S21b†). Consequently, the introduction of Mac-SnO₂ proved to be an efficient way to improve the stability of the device.

Fig. 7 (a) Normalized PCEs of SnO₂ and Mac-SnO₂ PSCs placed in the ambient environment (25–30 °C, 30%–40% humidity). (b) XRD spectra of the SnO₂/perovskite and Mac-SnO₂/perovskite before and after exposure to an air environment with 30%–40% humidity for 20 days.

3. Conclusion

In summary, we developed multifunctional Mac-SnO₂ QDs and employed them to fabricate ETLs for PSCs. The C=O groups in Mn(Ac)₂ could passivate the Sn⁴⁺ on the surface of the SnO₂ QDs, thereby reducing the van der Waals forces between the QDs, inhibiting QDs agglomeration, and contributing to the increased conductivity of the Mac-SnO₂ ETLs. Meanwhile, the Mn²⁺ spread upward in the perovskite layer, passivating the I⁻ at the grain boundary and reducing perovskite bulk defects. The perovskite grown on the Mac-SnO₂ ETL could release residual stress and suppress non-radiative recombination at the interface. Moreover, the well-matched energy level between Mac-SnO₂ and the perovskite is beneficial for extracting and transporting electrons. Ultimately, the champion device showed an increased PCE from 21.11% to 23.36%. More importantly, the Mac-SnO₂ devices also displayed enhanced long-term stability, with the PCE maintaining 90% of the original value under environmental conditions for nearly 600 h. Therefore, using multifunctional organic acid salts to stabilize SnO₂ QDs offers an effective method to further improve ETL quality and PSC performance. In the future, we should further explore small ions with better passivation effects and smaller resistance to stabilize SnO₂ QDs, which will further enhance the conductivity of the ETL. In addition, the use of ionic liquids to stabilize SnO₂ QDs is also a method worth exploring in depth, as it can not only passivate the surface defects of QDs, but also effectively inhibit ion migration at the ETL/perovskite interface through organic ligands, thereby improving device performance and stability.

4. Experimental section

Preparation of the ETL

The SnO₂ colloidal dispersion was diluted with deionized water at a ratio of 1 : 3. In the case of the Mac-SnO₂ solution, different concentrations of Mn(Ac)₂ (2.5, 5.0, 7.5, 10.0, and 12.5 mg mL⁻¹ in deionized water) were added to the SnO2 colloid dispersion, also at a ratio of 1 : 3.

Fabrication of the PSCs

The ITO substrate needed to undergo washing four times with deionized water and alcohol, respectively, with each cycle requiring 20 min of ultrasonic treatment. Subsequently, the ITO substrates were treated with a 20 min exposure to ultraviolet ozone (UV/O₃). The SnO₂ solution and Mac-SnO₂ solution were spin-coated on the ITO substrate at 3000 rpm for 30 s, then annealed at 150 °C for half an hour and then cooled to room temperature. The ITO/ETL samples were further subjected to a 20 min UV/O₃ treatment. Next, in order to grow the perovskite layer, the ITO/ETL material was transferred into a nitrogen-filled glove box (with H₂O and O₂ concentrations of <1 ppm). In the case of the perovskite solution, it was formed by dissolving 217 mg of FAI (1.26 mmol), 643 mg of PbI₂ (1.39 mmol), 27 mg of PbBr, 19 mg of CsI, 7.4 mg of MABr, and 20 mg of MACl in 1 mL mixed solvent (DMSO to DMF volume ratio of 2 : 8). A total of 85 μL of the Cs_0.05(FA_0.85MA_0.15)_0.95Pb(I_0.85Br_0.15)₃ precursor solution was then deposited onto the ITO/SnO₂ layer using a spin-coating process of 10 s at 1000 rpm and 30 s at 5000 rpm. At 27 s after the start of the procedure, 120 μL of CB was injected onto the rotating substrate for 2 s. The resulting films were heated for 1 h at 100 °C. Subsequently, to produce the HTL mixed solution, 72.3 mg of Spiro-OMeTAD was combined with 35 μL of Li-TFSI solution (250 mg of Li-TFSI in 1 mL of acetonitrile), 10 μL of FK209 solution (300 mg of FK209 in 1 mL of acetonitrile), and 30 μL of TBP in 1 mL of CB. Next, 85 μL of the HTL mixture was spin-coated onto the perovskite films for 30 s at 4000 rpm. The final step involved plating a 100 nm Ag layer under vacuum conditions.

For more comprehensive information regarding the materials used and the characterization processes, refer to ESI notes 1 and 2.†

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

The authors gratefully acknowledge the support from the National Natural Science Foundation of China (Grant No. 62275101, 22075101), Program for the Development of Science and Technology of Jilin Province (Item No. YDZJ202201ZYTS300, 20210509050RQ).

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4qi00166d

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