2D and 3D organic/inorganic hybrid perovskites for electrochemical energy storage applications

Abstract

Organic/inorganic hybrid perovskites are attracting research attention due to their great potential in energy storage applications. This study reported three-dimensional (3D) and two-dimensional (2D) lead halide inorganic/organic perovskites used as anode-active materials for Li-ion batteries. A doping strategy was successfully applied to enhance the cycling performance of the 3D hybrid perovskite anode. A moderate amount of chlorine (Cl) was doped into the 3D hybrid perovskite crystals to increase the polarization inside the network and support Li-ion transfer. The Cl-doped 3D hybrid perovskite anodes showed an outstanding improvement in cycling performance up to 1000 cycles at a high current density of 500 mA g⁻¹. Furthermore, three 2D hybrid perovskites (BAPbCl₄, BZAPbCl_4, and PEAPbCl₄) were applied as anode-active materials for the first time. Results revealed that aromatic organic ammonium made a great contribution to battery capacity and stability. The 2D hybrid perovskite anode delivered an excellent capacity of over 1000 mA h g⁻¹ under 100 mA g⁻¹, which is, by far, the highest reported Li-ion storage capacity for perovskite-based anode materials.

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Introduction

The development of smart electronics devices, electric vehicles, and smart grids has drawn tremendous research attention, and these devices have set to become an essential part of our modern society. To date, these advanced technologies strongly depend on rechargeable batteries as their main energy storage devices. In this regard, Li-ion batteries (LiBs) have shown promising capability due to their light weight, high power and energy densities.^1–4 To improve the storage capacity and efficiency and minimize the cost of LiBs, selecting electrode materials with high electrochemical stability and appropriate energy levels has become essential to meeting the demand for energy storage devices.^5–7 Among the emerging materials, organic/inorganic perovskite materials play a prominent role in the high-performance energy applications due to the synergic advantages of organic materials (e.g. chemical diversity) and inorganic components (e.g. high electronic mobility and a wide range of bandgap).^8–15

In paerticluar interest of LiB anode materials, organic/inorganic lead halide perovskites have emerged as promising candidates due to their high theoretical capacity, resource abundance, low-cost production and fabrication, as well as their rapid Li⁺ diffusion.^16–19 In contrast to the intercalation-limited capacities of graphite and spinel Li₄Ti₅O₁₂ (LTO), hybrid organic/inorganic lead halide perovskites can utilize multi-mechanism storage such as insertion, conversion, and alloying that can boost the battery performance.^20–23 Together with their compositional/dimensional tunability (A-site cation, 2D/3D frameworks, and halide composition), these features allow perovskites to attain higher specific capacities and robust rate performance²¹ while avoiding the dendrite formation and safety concerns of Li metal,^24,25 as well as the volume expansion of silicon-based anodes.²²

Generally, organic/inorganic lead halide perovskites form an octahedron structure with the coordination of six halogen anions around a lead cation ([PbX₆]²⁻).^26,27 The arrangement of organic ammonium and [PbX₆]²⁻ can be varied depending on the organic chain, thus leading to differences in the crystal structure. In 3D crystals, denoted as APbX₃ (A: organic ammonium cation, X: halogenic anion), the A cations are located between an octahedron to form a cubic framework around [PbX₆]²⁻.¹⁷ Meanwhile, in the 2D perovskite structure, the crystal consists of a layer structure with the interleaved arrangement of [PbX₆]²⁻ and the organic ammonium cations layers, which is usually a longer side chain, that forms cross-links via hydrogen bonds or van der Waals force.^16,28,29 Furthermore, besides the dimensionality of lead halide perovskites, it has been reported that other structural properties such as size, shape and composition of such micro/nanostructures could have a strong influence on the electrochemical performance as a LiB electrode.³⁰

The first use of an organic/inorganic lead halide perovskite as an active anode material for Li-ion batteries was reported by Xia et al. via a facile hydrothermal process.³¹ Therein, the prepared CH₃NH₃PbBr₃ anode delivered an initial capacity of 331.8 mA h g⁻¹ before it decreased rapidly in the first 30 cycles and reached 121 mA h g⁻¹ after 200 cycles.³¹ After that, another nanostructured form of CH₃NH₃PbBr₃ was prepared by a slow precipitation method combined with thermal treatment.³² The resulting CH₃NH₃PbBr₃ crystal showed electrochemical behavior comparable to that of the one prepared by the hydrothermal process. However, an improve Li⁺ storage capacity of up to 200 mA h g⁻¹ was achieved,³² implying that different structural properties impact the lithium-ion storage capability. However, the low specific capacity and poor cycling stability beyond 500 cycles is still major challenges for its practical applications.³⁰ To improve both storage capability and cycling stability, several strategies such as doping,^33–35 tuning dimensionality,^11,36 and changing the type of organic cation and anion^15,37 have been proposed.

Herein, we report the electrochemical performance of 3D/2D hybrid organic/inorganic lead halide perovskites as a LiB anode material. To improve the electrochemical performance of 3D organic/inorganic lead halide perovskites, chlorine (Cl) doping was introduced, along with varying material preparation via slow evaporation (slow evap.) and inverse temperature crystallization (ITC) methods. The effect of changing the organic ammonium cation was also examined by introducing phenylethylammonium (PEA) and formamidinium (FA) to replace the conventional methyl ammonium (MA) organic cation. Thus, series of 3D organic/inorganic lead halide perovskites, namely MAPbBr₃-5%Cl slow evap., MAPbBr₃-5%Cl ITC, MAPbBr₃-5%PEACl, and FAPbBr₃-10%Cl, were successfully prepared and applied as LiB anodes. The results show that after Cl doping, a major capacity fading of the 3D perovskite anode could be greatly avoided. Interestingly, the Cl-doped anodes could retain nearly their first capacities even up to 1000 cycles at a current density of 500 mA g⁻¹. In addition, 2D organic/inorganic lead halide perovskites, namely BAPbBr₄, BZAPbBr₄, and PEAPbBr₄, were prepared by varying the organic cation (BA: butyl ammonium; BZA: benzyl ammonium; PEA: phenyl ethyl ammonium) to evaluate the effect of organic cation on electrochemical performance. Surprisingly, the maximum capacity could reach over 1000 mA h g⁻¹ at a current density of 100 mA g⁻¹ for the 2D perovskite anode and stayed stable up to 100 cycles. Systematic spectroscopic analyses were performed to understand the charge storage phenomena in the 3D/2D hybrid perovskites during the electrochemical process. This study paves the way for the development of perovskite-based materials for energy storage applications.

Results and discussion

Material characterization

In this study, we intentionally selected two practical Cl contents (5% for MAPbBr₃ and 10% for FAPbBr₃) to probe how modest halide incorporation tunes the Li-ion transport and electrochemical performance without introducing secondary phases. These levels serve as representative points for performance benchmarking rather than a full compositional sweep. X-ray diffraction (XRD) spectroscopy was employed to probe the structural properties of the 2D/3D hybrid perovskites. As shown in Fig. 1a and b, there is no significant difference in the diffraction spectra of MAPbBr₃-5%Cl prepared by either the ITC or the slow evaporation method. The diffraction peaks corresponding to the (100), (110), (111), (200), (210), (211), (220), (300), (310), (320), and (410) crystal planes were present in both MAPbBr₃-5%Cl prepared by the ITC and the slow evaporation method, suggesting a cubic structure (space group of Pm3̄m), consistent with previous reports.^38,39 This further indicates that 5% doping of Cl does not change the crystal structure of 3D MAPbBr₃. Interestingly, the crystal structure of MAPbBr₃ did not significantly change and still matched the reported literature,^38,39 even after additional organic ammonium doping of PEA was introduced into the system (Fig. 1c). Additionally, the crystal structure of FAPbBr₃ with the increased doping content of Cl up to 10% (Fig. 1d) was maintained and was consistent with previous reports,^39,40 indicating high structural stability.

Fig. 1 Material characterization of 3D hybrid perovskites. (a–d) XRD spectra. (e–h) FESEM images. The inset images show the zoom-in images of the red-dashed area.

To investigate the morphological properties, the field-emission scanning electron microscope (FESEM) was employed. As shown in Fig. 1e–h, the 3D hybrid perovskites showed similar features, comprising non-uniform grains agglomerated together. Some small grains were stuck on the surface of large grains, resulting in a rough surface for 3D perovskites. Notably, the FAPbBr₃ crystals displayed a smoother surface compared to the MAPbBr₃ crystals (Fig. 1h). Energy-dispersive spectroscopy (FESEM/EDS) further confirmed Cl doping into the 3D hybrid perovskites (Fig. S1). As shown in Fig. S1, the major elements detected were carbon (C), nitrogen (N), lead (Pb), and bromide (Br) along with the presence of the Cl element.

In addition, XRD spectroscopy was used to determine the structural properties of the 2D hybrid perovskites. As shown in Fig. 2a–c, the diffraction spectra of the 2D hybrid perovskites showed a good fit with the simulated peaks reported in previous studies,^41,42 suggesting the successful formation and excellent crystallization of the 2D perovskites. The XRD peaks confirmed the layered structure of the 2D inorganic/organic lead perovskites, in which all the peaks were indexed to the monoclinic structure (space group C2/m). The assigned peaks at (200), (400), (600), (800), (1000), (1200), and (1400) in all three XRD patterns could be ascribed to the [PbCl₆]²⁻ crystal structure, while the other peaks might be related to organic ammonium. Furthermore, FESEM images showed that the 2D hybrid perovskites exhibited typical Ruddlesden–Popper 2D perovskites with irregular shapes and distinct features depending on the organic cation (Fig. 2d–f). As shown in Fig. 2d–f, increasing the size of the aromatic ammonium spacer (e.g., phenethylammonium, PEA) led to smoother terraces and more defined edge steps, reflecting cation-dependent nucleation and growth kinetics. The elemental compositions are confirmed by FESEM/EDS (Fig. S2), showing elements C, N, Pb, and Cl in the 2D hybrid perovskite crystal.

Fig. 2 Material characterization of 2D hybrid perovskites. (a–c) XRD spectra. (d–f) FESEM images. The inset images show the zoom-in images of the red-dashed area.

Electrochemical performance of 3D hybrid perovskite anode

Cyclic voltammetry (CV) measurements were performed in the potential window of 0.02 to 2.5 V (against Li/Li⁺) at a scan rate of 0.1 mV s⁻¹ to examine the electrochemical response of the 3D hybrid perovskite anode. As shown in Fig. 3a, four oxidation/reduction reaction peaks were observed for the Cl-doped MAPbBr₃ anodes, indicating both insertion and de-insertion as well as alloying and dealloying mechanisms of Pb with the alkali metal.^43–46 Interestingly, differences in synthesis method and Cl doping did not significantly alter the redox properties of the MAPbBr₃ anodes. The redox pair at 1.27/2.15 V could be attributed to the reduction of Pb(II) to Pb(0), whereas the three redox couples at 0.50/0.64, 0.26/0.53, and 0.02/0.1 V correspond to the Li⁺ insertion and the formation of Li–Pb alloys. The redox reaction can be described as follows:^11,43

Pb(II) + 2e → Pb(0)

Pb(0) + Li⁺ → PbLi_x (x ≤ 4.5)

Fig. 3 Electrochemical performance of 3D perovskites. (a) CV curves of MAPbBr₃-based perovskites. (b) CV curves of FAPbBr₃-based perovskites. (c) Capacity profile at a current density of 100 mA g⁻¹. (d) Rate performance. (e) Long-term cycling performance. The error bars represent a 95% confidence interval (CI) based on at least three replicates.

Furthermore, the FAPbBr₃ anode with 10% Cl doping exhibited distinct electrochemical behavior from that of the MAPbBr₃ anodes (Fig. 3b). As shown in Fig. 3b, the reduction of Pb(II) to Pb(0) shifted to a higher voltage of 1.89 V relative to the MAPbBr₃ anodes. Additionally, a new redox peak appeared at 0.9/1.04 V, which can be attributed to Li⁺ adsorption on the FAPbBr₃ anode. We believe these distinct redox properties are associated with the different ammonium cations and the higher Cl doping content. Furthermore, the oxidation peaks associated with the Li⁺ de-insertion and de-alloying mechanism were slightly shifted to higher voltages compared to MAPbBr₃ anodes (Fig. 3b), implying strong Li⁺ binding during the lithiation process. These distinctions in the CV profile of the FAPbBr₃-10%Cl anode could indicate an exceptional Li⁺ storage mechanism during the electrochemical process. Consistent with the CV behavior, the galvanostatic charge/discharge profile shows four plateaus for the Cl-doped MAPbBr₃ anodes (Fig. S3a–c), indicating a successful lithiation/delithiation process. Unlike the Cl-doped MAPbBr₃ anodes, the FAPbBr₃-10%Cl shows a different profile with a slope at a lower potential of 1.0 V due to the different ammonium cation and higher Cl doping content (Fig. S3d).

To further investigate the effects of Cl doping on cycling performance, LiB half-cells were then fabricated using 3D hybrid perovskite anode and tested over a voltage range of 0.02–2.5 V at a current rate of 100 mA g⁻¹. As shown in Fig. 3c and a rapid capacity drop was observed for the undoped MAPbBr₃ anode from 556 to 295.1 mA h g⁻¹ after the first 10 cycles.³² Interestingly, all Cl-doped MAPbBr₃ anodes demonstrated significant improvement on cycling stability. As shown in Fig. 3c, the capacities of MAPbBr₃-5%Cl prepared via slow evaporation and ITC methods and MAPbBr₃-5%PEACl anodes were maintained at 352, 349.6, and 322 mA h g⁻¹, respectively, after 100 cycles at a current density of 100 mA g⁻¹. This phenomenon could be associated with the higher electronegativity of Cl (Cl: 3.5 vs. Br: 3.45) and the smaller radius of Cl than those of Br, leading to a charge polarization in the Cl-doped active material and strengthening the ionic bonding between MA and [PbBr₆]²⁻.⁴⁷ Additionally, the combined effects of a different organic ammonium (FA) and higher Cl doping content were revealed by testing the electrochemical performance of the FAPbBr₃ with a 10% Cl anode. As shown in Fig. 3c, the FAPbBr₃-10%Cl anode achieved the largest capacity among the 3D electrodes studied, reaching up to 523.8 mA h g⁻¹ and showing only 32.2 mA h g⁻¹ capacity loss after 100 cycles at a current density of 100 mA g⁻¹. This remarkable enhancement could be attributed to the increased electronic conductivity of FAPbBr₃ (E_g: MAPbBr₃: 2.332 eV, FAPbBr₃: 2.292 eV) and the higher amount of Cl, which increases charge polarization, therefore supporting Li-ion diffusion kinetics.⁴⁸

Furthermore, as shown in Fig. 3d, the Cl-doped 3D hybrid perovskite anode demonstrated excellent rate capability up to a current density of 5000 mA g⁻¹. In line with the cycling performance, the FAPbBr₃-10%Cl anode showed the highest storage capacities of 570.6, 436.7, 372.2, 314.1, and 235.1 mA h g⁻¹ at current densities of 100, 500, 1000, 2000, and 5000 mA g⁻¹, respectively. Meanwhile, the MAPbBr₃-5%Cl anode prepared by slow evaporation exhibited specific capacities of 309, 225, 195, 162, and 118 mA h g⁻¹ at current densities of 100, 500, 1000, 2000, and 5000 mA g⁻¹, respectively. The MAPbBr₃-5%Cl anode prepared by the ITC method showed specific capacities of 336, 252, 214, 175, and 124 mA h g⁻¹ at current densities of 100, 500, 1000, 2000, and 5000 mA g⁻¹, respectively. By contrast, the MAPbBr₃-5%PEACl anode exhibited capacities of 262, 194, 167, 139, and 100 mA h g⁻¹ at current densities of 100, 500, 1000, 2000, and 5000 mA g⁻¹, respectively. The excellent rate capability was demonstrated as their highest capacity was recovered after the current density was reduced back to 100 mA g⁻¹.

The long-term cycling performance of the Cl-doped 3D hybrid perovskite anodes again proved the considerable improvement in material and cycling stability. As depicted in Fig. 3e and a slight growth in capacities at around 360 mA h g⁻¹ at a current density of 500 mA g⁻¹ during the first 700 cycles could be observed for MAPbBr₃-5%Cl anodes prepared by slow evaporation and the ITC method. Interestingly, these performances were stable up to 1000 cycles. By contrast, the capacity of MAPbBr₃-5%PEACl anode slowly increased to over 340 mA h g⁻¹ at a current density of 500 mA g⁻¹ during the first 600 cycles and then dropped quickly to 217.7 mA h g⁻¹ at the 1000th cycle, suggesting the effect of PEA cation on cycling stability under long-term operation. Notably, although the FAPbBr₃-10%Cl anode achieved the highest capacity among the four anodes, it showed a slightly reduced stability under long-term cycling compared with the MAPbBr₃-5%Cl anodes after ∼700 cycles. Specifically, the FAPbBr₃-10%Cl anode revealed a remarkable increase in its capacity from 260.5 to 474 mA h g⁻¹ at the first 400 cycles; its capacity dropped to 324 mA h g⁻¹, slightly lower than the capacities of the MAPbBr₃-5%Cl anodes after 1000 cycles (Fig. 3e). Additionally, the performance of our 3D hybrid perovskite anodes is comparable to the best-reported hybrid perovskite anodes (Table S3).

To probe the charge transport phenomena, electrochemical impedance spectroscopy (EIS) was first performed to examine Li-ion transfer at different cycling intervals (Fig. S4). The fitted parameters are summarized in Table S1. Two types of charge transfer resistance (R_ct) were observed in the mid-frequency region, associated with Li⁺ transport across different interfaces during electrochemical operation: (i) electrode–electrolyte interfaces/solid–electrolyte interphase (SEI) (R_cta), and (ii) material interfaces within the electrode (R_ctb). As presented in Table S1, a slight change in the R_ctb was observed for the MAPbBr₃-5%Cl anodes, which reduced from 147.1 (before cycling) to 141.6 Ω (after 100 cycles) for MAPbBr₃-5%Cl prepared by the slow evaporation method and increased from 164.8 (before cycling) to 175.2 Ω (after 100 cycles) for MAPbBr₃-5%Cl prepared by the ITC method. Interestingly, the Li-ion diffusion co-efficient (D_Li) was significantly enhanced from 8.82 × 10⁻¹⁴ (before cycling) to 1.54 × 10⁻¹³ cm² s⁻¹ (after 100 cycles) for MAPbBr₃-5%Cl (slow evaporation) and from 1.36 × 10⁻¹³ (before cycling) to 3.06 × 10⁻¹³ cm² s⁻¹ (after 100 cycles) for MAPbBr₃-5%Cl (ITC method). After 1000 cycles, these two 3D perovskites show significant reduction in R_ctb and an increase in D_Li (Table S1). This indicates that lower interfacial resistances on the materials inside the electrode boosted Li⁺ mobility and helped in maintaining the excellent performance of the MAPbBr₃-5%Cl anodes prepared by slow evaporation and ITC methods after 1000 cycles.

Furthermore, the MAPbBr₃-5%PEACl and FAPbBr₃-10%Cl anodes showed distinct features compared to the MAPbBr₃-5%Cl anodes prepared by slow evaporation and ITC methods. As shown in Fig. S4 and Table S1, the MAPbBr₃-5%PEACl anode showed a considerable reduction in R_ctb values before and after 100 cycles from 410.5 to 155.4 Ω, respectively, along with an increase in D_Li, implying a significant increase in Li⁺ mobility. However, the R_ctb in the MAPbBr₃-5%PEACl anode increased to 248 Ω after 1000 cycles, suggesting an increase internal resistance possibly due to structural instability, thus reduced its performance after 1000 cycles (Fig. 3e). Similarly, the FAPbBr₃-10%Cl anode demonstrated a significant reduction in R_ctb from 215.4 (before cycling) to 140.9 Ω (after 100 cycles), along with an increase in D_Li from 6.22 × 10⁻¹³ to 9.40 × 10⁻¹³. Furthermore, the R_cta was lower than that of the MAPbBr₃-5%Cl anodes prepared by the ITC method and the MAPbBr₃-5%PEACl anode (Table S1). These phenomena confirm that during the first 100 cycles, progressive wetting, SEI stabilization and reduced interfacial/transport resistances enhanced Li⁺ mobility and storage capacity for FAPbBr₃-10%Cl in the first 100 cycles (Fig. 3e). Unlike the MAPbBr₃-5%PEACl anode, FAPbBr₃-10%Cl showed an increase in the R_cta value and D_Li after 1000 cycles (Fig. S4 and Table S1). This indicates that cumulative SEI thickening during repeated cycling hindered Li⁺ mobility and reduced storage capacity (Fig. 3e). These results further suggest that doping large organic cation A, such as PEACl, in the 3D hybrid perovskite anodes impact structural integrity under the long-term cycling test. Meanwhile, increasing halogen doping in 3D hybrid perovskite anodes can contribute to significantly increased resistive SEI formation at the electrolyte–electrode interface during long-term operation.

To further examine the effect of the doping element on charge storage behavior, sweep-rate CV measurements were performed at different scan rates ranging from 0.1 to 1.0 mV s⁻¹ (Fig. 4a–d). The total accumulated charge in the active materials during electrochemical process generally consist of different faradaic contributions, from the diffusion controlled (ion insertion) or surface-controlled process, which is referred to pseudocapacitance or the non-faradaic contribution (double layer effect).^49–52 These faradaic processes, which are related to reaction kinetics, can be expressed as follows:^53,54

i = av^b

and

log(i) = log(a) + b log(v),

where i is the peak current at a scan rate of v, while a and b are adjustment parameters. If the b-value is close to 0.5, it indicates a semi-infinite linear diffusion-controlled process. When a b value is close to 1, it suggests that the current generation is a surface-controlled process.^{49,50,55–57}

Fig. 4 Charge storage behavior of 3D hybrid perovskites. (a–d) CV curves at different scan rates and (e–h) charge storage contributions at different scan rates.

Moreover, the total contribution at a specific scan rate can be calculated using^57,58

i = k₁v + k₂v^1/2,

where k₁v and k₂v^1/2 are capacitive and diffusion contributions, respectively.

As shown in Fig. S5, the b-values for MAPbBr₃-5%Cl prepared by slow evaporation, MAPbBr₃-5%Cl prepared by the ITC method, MAPbBr₃-5%PEACl, and FAPbBr₃-10%Cl were 0.557, 0.693, 0.827, and 0.857, respectively, suggesting that the MAPbBr₃-5%Cl anode prepared by the slow evaporation method has a stronger contribution from diffusion-controlled processes rather than surface-controlled (capacitive) contributions. Meanwhile the higher b-values observed for MAPbBr₃-5%Cl prepared by the ITC method, MAPbBr₃-5%PEACl, and FAPbBr₃-10%Cl indicate that these 3D hybrid perovskites have a synergic contribution from both capacitive and diffusion-controlled processes during electrochemical operation. Fig. 4e–h show the distribution of diffusion and capacitive contributions in the 3D hybrid perovskites at various scan rates. In good agreement with the b-values, the charge storage behavior of the MAPbBr₃-5%Cl anode prepared by the slow evaporation method was dominated by diffusion contribution (more than 80%) rather than the capacitive contribution (Fig. 4e). By contrast, although having the same chemical composition, MAPbBr₃-5%Cl prepared by the ITC method exhibited a nearly 40% capacitive contribution at high scan rates of 0.5–1 mV s⁻¹, before gradually decreased at slow scan rate 0.01 mV s⁻¹ (Fig. 4f). This suggests that different synthesis methods could significantly impact the charge storage behavior of 3D hybrid perovskites. Furthermore, doping a larger ammonium cation (PEA) into the MAPbBr₃-5%PEACl system further increased the capacitive behavior of the anode (Fig. 4g). Meanwhile, doping 10% Cl into the FAPbBr₃ anode also enhanced the capacitive behavior. These results imply that higher-electronegativity doping materials significantly tune the charge storage characteristic by increasing the capacitive contribution.

Electrochemical performance of 2D hybrid perovskite anode

In order to verify the redox properties of the 2D hybrid perovskite anodes, CV was again performed within a potential window of 0.02–3.0 V (against Li/Li⁺) at a scan rate of 0.1 mV s⁻¹. As shown in Fig. 5a, all 2D hybrid perovskites exhibited comparable CV behavior. Four pair redox peaks were observed at 1.60/2.43, 0.50/0.65, 0.26/0.53, and 0.02/0.15 V, suggesting the reduction/oxidation of Pb(II) to Pb(0), the formation of Li–Pb alloys and Li⁺ insertion. Moreover, Fig. S6 shows the galvanostatic profile of the 2D hybrid perovskites within the potential range of 0.02–3.0 V at a current density of 100 mA g⁻¹. As shown in Fig. S6, four plateaus were observed during lithiation, consistent with the CV results (Fig. 5a), indicating successful Li storage. Notably, gradual increase in capacity was observed for BZAPbCl₄ and PEAPbCl₄, while the BAPbCl₄ showed relatively stable performance up to 100 cycles. These results suggest that different organic ammonium cations have a significant impact on the Li storage capability of 2D hybrid perovskites.

Fig. 5 Electrochemical performance of 2D perovskites. (a) CV curves. (b) Capacity profile at a current density of 100 mA g⁻¹. (c) Rate performance. (d) Long-term cycling performance. The error bars represent a 95% confidence interval (CI), based on at least three replicates.

Fig. 5b shows the cycling performance of the 2D hybrid perovskites at a current density of 100 mA g⁻¹. As shown in Fig. 5b, the initial cycles of the 2D hybrid perovskite anodes exhibited comparable capacities of 582.6 (BAPbCl₄), 584.8 (BZAPbCl₄), and 616.9 (PEAPbCl₄) mA h g⁻¹ at the 10th cycle. However, after 25 cycles, the cycling performance showed gradual upward capacity trends for BZAPbCl₄ and PEAPbCl₄, consistent with the galvanostatic profiles (Fig. S6). BZAPbCl₄ and PEAPbCl₄ anodes showed optimal capacities of 1008.3 and 874 mA h g⁻¹, respectively, after 100 cycles under a current density of 100 mA g⁻¹, respectively. This specific capacity by far is the highest Li⁺ storage capacity reported for perovskites-based anodes (Table S3). Meanwhile, BAPbCl₄ exhibited steady performance and delivered a specific capacity of 679.4 mA h g⁻¹ after 100 cycles (Fig. 5b). The significant increase in capacity of BZAPbCl₄ and PEAPbCl₄ can be associated with the additional Li⁺ binding sites on the benzene rings of the organic ammonium cation (BZA and PEA) compared with BAPbCl₄.⁵⁹ Additionally, the rate capability of the 2D hybrid perovskite anode was examined at various current densities ranging from 100 to 5000 mA g⁻¹. As shown in Fig. 5c, BZAPbCl₄ revealed the highest rate capability, delivering storage capacities of 667.8, 680.0, 621.4, 551.8, and 425.0 mA h g⁻¹ at current densities of 100, 500, 1000, 2000, and 5000 mA g⁻¹, respectively. BAPbCl₄ and PEAPbCl₄ anodes delivered capacities of 562 498.5, 425.9, 343.1, and 225.2 mA h g⁻¹ and 443.7, 409.4, 375, 337.5, and 471.9 mA h g⁻¹, respectively, under current densities of 100, 500, 1000, 2000, and 5000 mA g⁻¹. Furthermore, excellent rate capability was confirmed as the capacities recovered to 957, 672, and 537 mA h g⁻¹ when the current density was returned to 100 mA g⁻¹. Interestingly, an upward capacity trend was also observed during the rate capability test.

To probe the cycling stability, long-term cycling performance was tested at a current rate of 500 mA g⁻¹ up to 1000 cycles (Fig. 5d). As shown in Fig. 5d, the 2D hybrid perovskite anodes exhibited increasing trend up to an optimum storage capability and then decreased slowly up to 1000 cycles. The BAPbCl₄ anode showed an increase in capacity from 505.6 to 622.2 mA h g⁻¹ during the first 50 cycles and then gradually decreased to 250 mA h g⁻¹ after 1000 cycles. Meanwhile, BZAPbCl₄ and PEAPbCl₄ anodes demonstrated better long-term cycling stability compared with BAPbCl₄. Specifically, the PEAPbCl₄ anode achieved its optimum capacity of over 760 mA h g⁻¹ after 250 cycles, before gradually decreasing to 146.9 mA h g⁻¹ at 1000 cycles. Meanwhile, the BZAPbCl₄ anode reached its optimum capacity of 740 mA h g⁻¹ after 180 cycles before it stepwise dropped to 181.3 mA h g⁻¹ over 1000 cycles.

To understand the Li⁺ transport properties during long-term operation of 2D hybrid perovskites, EIS measurement were carried out at different cycling intervals (Fig. S7 and Table S2). As shown in Table S2, a gradual decrease in R_ctb was observed for the BAPbCl₄ anode during cycling, from 214.4 (before cycling) to 203.4 Ω (after 100 cycles) and then reached 188.1 Ω (after 1000 cycles). However, a significant increase in R_cta was observed for BAPbCl₄ from 80.9 (after 100 cycles) to 159.5 Ω (after 1000 cycles), suggesting that a huge resistive electrolyte–electrode interface significantly hindered Li⁺ mobility. Furthermore, consistent with the increasing R_cta, the D_Li was also significantly dropped by one order of magnitude, from 1.27 × 10⁻¹² (after 100 cycles) to 8.38 × 10⁻¹³ cm² s⁻¹ after 1000 cycles, implying significantly hindered Li⁺ transport after 1000 cycles. A similar phenomenon was observed for BZAPbCl₄ and PEAPbCl₄, where the R_cta value significantly increased, and the D_Li was dramatically restricted after 100 and 1000 cycles (Table S2). Notably, BAPbCl₄ demonstrated a significant increase in R_cta than that of BZAPbCl₄ and PEAPbCl₄ (Fig. S7 and Table S2). The BZAPbCl₄ anode showed an increase in R_cta from 9.146 Ω (100 cycles) to 47.77 Ω (after 1000 cycles), while PEAPbCl₄ showed an increase from 18.22 Ω (100 cycles) to 39.9 Ω (after 1000 cycles). This further suggests that the shorter organic ammonium, such as butylammonium in BAPbCl₄, results in higher excessive growth of the resistive SEI layer compared with larger organic ammonium benzylammonium (BZA) and phenyl ethyl ammonium (PEA) due to the presence of delocalized electron on the benzene ring that helps stabilize interfaces and prevent excessive SEI growth.^60–62 These observations further indicate that the significant capacity drop in 2D hybrid perovskites is associated with excessive growth of the resistive SEI layer, which restricts Li⁺ transport. This phenomenon, which hinders Li⁺ transport between the electrolyte and the anode during long-term operation, is known as one of the main aging mechanisms in LiBs.^19,63,64

X-ray photoelectron spectroscopy (XPS) analysis was performed on BZAPbCl₄ at different charge/discharge states to understand the Li⁺ storage mechanism in the 2D hybrid perovskite. As shown in Fig. 6a, three peaks were observed in the C 1s spectra at binding energy (BE) of 284.5, 285.4, and 290.4 eV, which correspond to C–C, C–N, and π–π bonding, respectively. A new BE associated with C–Li can be observed at 283.9 eV in the fully charged state, suggesting successful insertion of Li-ions into the benzene rings. Meanwhile, the BE associated with Pb(0) was observed at BEs of 136.4 (Pb 4f_5/2) and 141.2 (Pb 4f_7/2) eV in the full charge state (Fig. 6b), indicating Li⁺ storage contribution from Pb via the conversion reaction of Pb(II) to Pb(0), followed by the alloying mechanism (Li_yPb). However, the binding energy of the Li_yPb alloy could not be identified due to the overlap BE with the Pb(0) signals. Furthermore, when the BZAPbCl₄ anode was fully discharged, the BEs of Pb(0) were still present, indicating unrecovered Pb(0) after discharge. Additionally, BEs of Li–Cl were observed in the Cl 2p spectra during both full charge and full discharge states (Fig. 6c).

Fig. 6 Ex situ XPS spectra of BZAPbCl₄ at different stages, and charge storage behavior of 2D perovskites. (a) C 1s; (b) Pb 4f; (c) Cl 2p and (d) Li 1s spectra. (e–g) CV curves at different scan rates and (h–j) charge storage contributions at different scan rates for 2D perovskites.

Consistent with the C 1s core spectra, the Li 1s spectrum exhibited additional specific BE of C–Li at 53.8 eV during the full charge state, which disappeared in the full discharge state (Fig. 6d), suggesting successful Li-ion insertion/de-insertion into the organic cation. Meanwhile, the BE of Li–Cl was observed at 55.8 eV in the full charge state, indicating Li-ion interaction with Cl atoms within the BZAPbCl₄ structure, which is consistent with the Pb 4f and Cl 2p core spectra. Notably, the BE of Li–Cl persisted alongside specific BEs of Li₂CO₃ (55.2 eV) and LiF (56.5 eV) in the full discharge state (Fig. 6d). This indicates the detachment of Cl atoms from the octahedron coordination within the BZAPbCl₄ structure. The additional BEs of Li₂CO₃ and LiF are ascribed to the formation of the SEI layer. These XPS results further suggested that during lithiation, Li⁺ can be stored in BZAPbCl₄via insertion into the organic moiety (benzene ring), followed by conversion and alloying reactions. Unfortunately, the octahedron coordination between Pb and Cl seems to be disturbed and unrecovered during electrochemical cycling.

To gain a better understanding of the charge storage characteristics, sweet-rate CV measurements were again performed at scan rates ranging from 0.1 to 1.0 mV s⁻¹ (Fig. 6e–g). As shown in Fig. S8a, the b-values of the 2D hybrid perovskites were estimated to be around 0.743, 0.797, and 0.794 for BAPbCl₄, BZAPbCl₄, and PEAPbCl₄ anodes, respectively. This indicates that the total charge stored in the 2D hybrid perovskites are contributed from both diffusion- and surface (capacitive)-controlled reactions. Fig. 6h–j show the charge storage distribution between diffusion and capacitive contributions at various scan rates. As depicted in Fig. 6h, the BAPbCl₄ anode exhibited nearly 50% capacitive contribution at high scan rates of 0.6–1.0 mV s⁻¹ and then dropped to 18% at 0.1 mV s⁻¹ along with an increasing diffusion contribution. Additionally, the BZAPbCl₄ (Fig. 6i) and PEAPbCl₄ (Fig. 6j) anodes displayed relatively higher capacitive contributions of over 60 and 70%, respectively, at high scan rates of 0.6–1 mV s⁻¹ and then decreased gradually to 33 and 36% at 0.1 mV s⁻¹, suggesting that additional aromatic systems in the 2D hybrid perovskites could promote higher capacitive contributions. In brief, this study indicates that different organic ammonium cations have a significant impact on the charge storage behavior of 2D hybrid perovskites.

Furthermore, EIS analysis was again employed to probe the activation energy (E_a) of the 2D hybrid perovskite anodes. The relationship between the current response (i₀) and the charge transfer resistance at the electrode interface (R_ct) over different temperature ranges can be used to estimate the E_a using the following Arrhenius equation:^65–69

i₀ = RT/(nFR_ct)

and

i₀ = A exp(−E_a/RT),

where A is the temperature-independent coefficient, F is the Faraday constant, R is the gas constant, and T is the temperature (K). Furthermore, the E_a of electrode materials can be estimated from E_a = −Rk ln 10, where k is the slope of the Arrhenius plots (log₁₀ i₀ as a function of 1000/T).^65–68 As shown in Fig. 7a–c, the semicircle on the Nyquist plot of the 2D perovskite decreases as the test temperature increases, suggesting temperature dependent of the R_ct. Based on the Arrhenius plot (Fig. 7d), the E_a values of the 2D hybrid perovskite anodes were estimated to be around 28.07, 9.82, and 33.83 kJ mol⁻¹ for the BAPbCl₄, BZAPbCl₄, and PEAPbCl₄ anodes, respectively. The lowest E_a for the BZAPbCl₄ anode suggests that the minimum energy is required to activate the structure for charge transfer reactions inside the battery, resulting in the largest rate performance. In contrast, the highest energy is required for PEAPbCl₄ to promote charge transfer reactions. This may be ascribed to the more stable structure of PEAPbCl₄ compared with BZAPbCl₄ due to stronger van der Waals forces inside the PEAPbCl₄ network. In summary, we found that the organic ammonium cation plays an important role in tuning the charge storage characteristics and E_a of 2D hybrid perovskite anode, which further enhances their electrochemical performance as LiB electrodes.

Fig. 7 Activation energy calculation of 2D hybrid perovskites. (a), (b) and (c) Nyquist plots of BAPbCl₄, BZAPbCl₄, and PEAPbCl₄ at various temperatures, respectively. (d) Log i₀vs. 1000/T plot.

Conclusions

In summary, we introduced an effective strategy to prevent rapid capacity loss during cycling in 3D organic/inorganic hybrid perovskites. Doping with 5% Cl or 5% PEACl significantly stabilized the battery performance of MAPbBr₃ anodes, which maintained capacities of 338.8, 347.1, and 324.7 mA h g⁻¹ for MAPbBr₃-5%Cl prepared by slow evaporation, MAPbBr₃-5%Cl prepared by the ITC method, and MAPbBr₃-5%PEACl, respectively, up to 1000 cycles at a current density of 500 mA g⁻¹. Moreover, the performance of 3D hybrid perovskite FAPbBr₃ with 10% Cl doping proved the effects of organic ammonium and halogen doping levels, which changed the conductivity and polarization of the materials, thus significantly boosting Li⁺ storage capability. Interestingly, 2D hybrid perovskite anodes demonstrated better performance than 3D anodes, which is ascribable to the effect of crystal structure and the organic ammonium cation. BAPbCl₄ anodes achieved a maximum capacity of 670.4 up to 100 cycles (current density = 100 mA g⁻¹), larger than the capacity of the 3D anodes. Especially, the BZAPbCl₄ anode exhibited an excellent capacity of over 1000 mA h g⁻¹ up to 100 cycles at a current density of 100 mA g⁻¹, which is by far the highest capacity reported for a perovskite-based battery anode.

Experimental section

Materials

Lead(II) bromide (PbBr₂, 99.999% trance metal basis), lead(II) oxide (PbO, powder, ≥99.9% trace metals basis), methylammonium bromide (MABr, ≥98%), methylammonium chloride (MACl, ≥98%), formamidinium bromide (FABr, ≥98%), butylamine (BA, ≥99%), benzylamine (BZA, ≥99.5%), phenylethylamine (PEA, ≥99%), phenylethylammonium chloride (PEACl, ≥98%), hydrochloric acid (HCl, 37 wt% in H₂O, trace metal basis), and hypophosphorous acid solution (H₃PO₂, 50 wt% in H₂O) were purchased from Millipore Sigma. The conductive carbon (Super P®; >99% (metal basis)), polyvinylidene fluoride (PVDF), and 1 M lithium hexafluorophosphate (LiPF₆) in ethylene carbonate (EC) and diethyl carbonate (DEC) (1 : 1 v/v) were purchased from UBIQ Technology Co., Ltd. N-Methylpyrrolidone (NMP; >99%) was purchased from Alfa Aesar (Thermo Fisher Scientific). All materials were used as received without further purification.

Synthetic procedure of 3D hybrid perovskites

The 3D hybrid perovskite materials were synthesized through slow evaporation and ITC methods. For the slow evaporation of MAPbBr₃-5%Cl, we dissolved PbBr₂, MABr, and MACl powders in a DMF solution at a molar ratio of 1 : 1.2 : 0.5. The precursor solution was filtered through a 0.2 µm PTFE syringe filter to remove undissolved particles and then placed in a 20 ml scintillation vial for slow evaporation at room temperature. MAPbBr₃-5%PEACl and FAPbBr₃-5%Cl were prepared following the same procedure, with MACl replaced by PEACl for MAPbBr₃-5%PEACl, and MABr replaced by FABr for FAPbBr₃-5%Cl.

For the ITC growth of MAPbBr₃-5%Cl, we dissolved PbBr₂, MABr, and MACl powders in DMF solution at a molar ratio of 1 : 1.2 : 0.5. The precursor solution was filtered through a 0.2 µm PTFE syringe filter to remove undissolved particles. The solution temperature was ramped slowly from room temperature to 80 °C at a rate of 4 °C per hour, yielding large crystals with high transparency. The raw materials were characterized with an X-ray diffraction to assess their purity and crystallinity. The absorption spectra were measured in reflection mode by a Jasco V-730 UV-Visible Spectrophotometer for characterization of doped samples.

Synthetic procedure of 2D hybrid perovskites

The synthesis of 2D layered perovskite materials followed previous reports.^70,71 Generally, the raw 2D hybrid perovskite flakes were prepared by first dissolving PbO in a concentrated HCl/H₃PO₂ solvent mixture and heating to 120 °C. After the PbO₂ was completely dissolved in solution, organic amines (RA, where R = BA, BZA, and PEA) were added into the solution at an appropriate ratio of PbO : RA = 1 : 2. After constant stirring and evaporating the HCl solution, the raw 2D perovskites started to crystalize, and the beaker was removed from the hot plate to cool to room temperature overnight. The raw materials were characterized with X-ray diffraction to assess their purity and crystallinity.

Material characterization

X-ray diffraction (XRD) measurements were performed on a Bruker D8 Advance X-ray Diffractometer at 40 kV and 40 mA using Cu Kα radiation (λ = 1.5406 Å). Scanning electron microscopy (SEM) was performed using a field-emission scanning electron microscope (Ultra Plus – Carl Zeiss). The X-ray photoelectron spectroscopy (XPS) was performed using a high-resolution XPS system (PHI-Quantera II, ULVAC-PHI, Inc.).

Coin cell preparation

The 2D/3D hybrid perovskite anodes were prepared by a slurry method. In brief, 40% of active materials was mixed with 40% conductive carbon (Super P®; >99% (metal basis); UBIQ Technology Co., Ltd) and 20% poly(vinylidene fluoride) (PVDF; UBIQ Technology Co., Ltd) as a binder, using N-methyl-2-pyrrolidone (NMP; Thermo Fisher Scientific) as the solvent. After forming a homogeneous slurry, it was cast onto Cu foil and dried on a hot plate at 60 °C for 12 hours, followed by another 8 hours at 80 °C under vacuum. The resulting electrodes were then cut into 12 mm-diameter disks with an average loading of 0.32 mg cm⁻². Finally, CR2032-type coin cells were assembled in an Ar-filled glove box (Vigor, Vigor Tech USA) with water (H₂O) and oxygen (O₂) concentrations kept below 0.5 ppm. The resulting half-cell LiB consisted of the 2D/3D hybrid perovskite anode as the working electrode, Li metal foil as the counter/reference electrode, Celgard 2325 as the separator, and 40 µL of 1 M LiPF₆ in ethylene carbonate (EC)/diethyl carbonate (DEC) (1 : 1 v/v) (UBIQ Technology Co., Ltd) as the electrolyte.

Electrochemical measurement

Cyclic voltammetry (CV) was performed using MultiPalmSens4 electrochemical analyzer, PalmSens BV. Electrochemical impedance spectroscopy (EIS) analyses were conducted before and after battery cycling using a CHI electrochemical workstation model 760e (CH Instruments, Inc.) with an alternating current (AC) voltage signal of 10 mV over a frequency range between 10 mHz and 1 MHz. The LiB half-cells were galvanostatically charged and discharged using an AcuTech battery station system (AcuTech Systems Co. Ltd). All measurements were conducted in a room-temperature environment.

Author contributions

H.-J. Y., W. N., and H. T. designed the experiments. H. T. performed the perovskite synthesis. F. B. performed the electrochemical and characterization experiments. H. Q. W performed characterization experiments. P.-C. C. performed the electrochemical experiment. The manuscript was written and edited by F. B., T. T. B. T., A. S., H. J. Y., and H. T. All authors discussed the results and reviewed the manuscript.

Conflicts of interest

There are no conflicts to declare.

Data availability

The data supporting this article have been included as part of the supplementary information (SI). Supplementary information: SEM-EDS images, galvanostatic profiles, EIS spectra, log peak current vs. log scan rate plots, SI tables. See DOI: https://doi.org/10.1039/d5ta06470h.

Acknowledgements

H.-J. Yen acknowledges the financial support by Innovative Materials and Analysis Technology Exploration in Academia Sinica (AS-iMATE-114-24) and the National Science and Technology Council in Taiwan (NSTC 114-2113-M-001-018; NSTC 114-2113-M-001-015; NSTC 113-2113-M-001-036). This research is funded by the Indonesian Endowment Fund for Education (LPDP) on behalf of the Indonesian Ministry of Higher Education, Science and Technology and managed under the EQUITY Program (Contract No. 4298/B3/DT.03.08/2025).

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