Mengna Xieab,
Yongheng Zhouc,
Shuai Dongb,
Fei Lia,
Fenghua Zhang*a,
Wei Wei*b and
Jinhai Cui
*b
aSchool of Petrochemical Engineering, Liaoning Petrochemical University, Fushun 113001, P. R. China. E-mail: zhangfenghua@lnpu.edu.cn
bHenan Engineering Center of New Energy Battery Materials, School of Chemistry and Chemical Engineering, Shangqiu Normal University, Shangqiu 476000, P. R. China. E-mail: cuijinhai535@gmail.com; weiweizzuli@163.com
cSchool of Material and Chemical Engineering, Kaifeng University, Kaifeng 475000, P. R. China
First published on 6th June 2025
Lower lithium-ion diffusion rates and significant volumetric expansion present serious challenges for using SnO2/SnO composites as promising anode materials in advanced lithium-ion batteries. To address this issue, we synthesized a novel Sn@C/CNT composite from a Sn-based organometallic complex with 2-methylimidazole and oxidized multi-wall carbon nanotubes. Structural analysis has confirmed that the tin-based composites consist of nano-lamellar assemblies modified by oxidized carbon nanotubes. In these composites, the tin active particles have an average size ranging from 2 to 3 nm, while the layered nano-lamellar structure has an average thickness of 6 nm. The resulting Sn@C/CNT anode material demonstrated a stable specific capacity of up to 688 mA h g−1 even after 500 cycles at a higher charging–discharging current density of 1 A g−1. The significant diffusion-controlled lithium ion diffusion coefficient of approximately 10−12 cm2 s−1 indicates vigorous dynamic activity from reversible Sn–Li alloy electrochemical reactions. Additionally, the substantial capacity-controlled lithium ion diffusion coefficient, which drops to 10−16 cm2 s−1, illustrates the predominance of the pseudo-capacitance arising from interface reaction. By coupling electrochemical impedance spectroscopy, galvanostatic intermittent titration technique, and linear sweep voltammetry, the mixed lithium-ion diffusion effect was proposed to explain the remarkable adaptability of these Sn-based anode materials for cycling performance across a wide range of specific currents. This work provides a new intention for resolving the drastic volumetric expansion and unsatisfactory dynamic activity of Sn-based anode materials.
Tin-based materials, including Tin (Sn), SnOx (where 1 ≤ x ≤ 2), and its sulfides, have gained attention as promising anode materials for advanced LIBs due to their high theoretical specific capacity and lower lithiation plateau.8 According to research on the interaction between SnOx and lithium,9 the mechanism by which SnOx stores lithium ions has been described as a two-step electrochemical reaction, as outlined by Chen and others.10
Scheme 1 shows those both SnO2 and SnO have the maximum theory-specific capacity of either 1494 mA h g−1 or 1273 mA h g−1, respectively, when alloy LiySn (y = 4.4) were obtained in reaction (2), in which electrons y can be up to 4.4 (reaction (2)) and has been confirmed by Dahn and Courtney.11 Although 783 and 398 mA h g−1 of the theory-specific capacity can be pinpointed to either SnO2 or SnO, respectively, they seem meaningless because of the irreversible interaction (1) of SnOx and Li+ in Scheme 1. Therefore, the fundamental theory-specific capacities of those SnO2, SnO, and Sn should correspondingly drop down to 783 mA h g−1, 876 mA h g−1, and 994 mA h g−1, respectively, when the alloying reaction reversibly goes on until the final alloy Li22Sn5 (y = 4.4) was obtained (eqn (2)).
However, lower dynamic activity in the diffusion of lithium ions and significant volumetric expansion of the Sn–Li alloy phase during the charging and discharging processes of Sn-based anode materials can lead to a substantially reduced specific capacity. Additionally, this can cause severe efflorescence and abnormal growth of the Solid Electrolyte Interface (SEI), ultimately decreasing the cycling life of LIBs.12 So far, exploiting elemental Sn as the anode material in LIBs has been difficult. Researchers have proposed various strategies to address the limitations of Sn-based anode materials.13 One effective approach involves minimizing the particle size of these materials. For instance, X. Ye and colleagues reported that they prepared a type of five-nanometer Sn/C composites by annealing Sn-based sludge, which was degraded by microorganisms, under an inert atmosphere. These fine Sn/C nanorod materials exhibit reduced efflorescence, stable cycling life, and excellent rate capability in LIBs.14 Another method to enhance the performance of Sn-based anode materials is to create functional composites that combine Sn-based materials with structural carbon sources, such as carbon nanotubes (CNTs), graphene (GN), and three-dimensional (3D) network carbon. Among the various Sn-based composites, Sn/CNT, Sn/GN, and Sn/3D-C composites demonstrate significant resistance to efflorescence and deformation, as well as lower impedance. This is achieved by confining the Sn-based materials within the cavities of these carbon structures,15,16 thereby localizing them within their layers17–19 and trapping them in lattices.20 Consequently, the bulk deformation of these anode composites can be effectively controlled, and their impedance can be reduced.
Unlike the mechanical mixing of carbon nanotubes or graphene sheets, 3D network-carbon composites are typically prepared through in situ growth.21 This method allows the 3D network-carbon matrix to better embed the Sn-based active materials, resulting in improved conductive capabilities.22
In this paper, we present a new network-carbon composites made from Sn-based materials. This innovative approach combines both an in situ preparation of 2D nano-lamellar assembly and the mixing preparation of carbon nanotubes, which are annealed at higher temperatures, with carbon nanotube doping, which is annealed at lower temperatures. The goal is to achieve a long-term cycling lifespan of Sn-based anode material by addressing volumetric expansion, enhancing the activity of electro-chemical reactions, and accelerating the migration rate of Li+ ions for this Sn-based anode. Additionally, this method offers a novel synthetic routine for other anode materials. To our knowledge, this proposal has not been previously reported.
With the comparison, the commercial acetylene black (35–45 nm) and nano Tin (100 nm in size, Shanghai Pantian Powder Materials Co., Ltd) were repeatedly mixed in a same dosage of the acetone–water solution, milled, annealed at 800 °C for getting the in situ Sn-based carbon nano-lamellar assemblies, and then doped with the same amount of oxidized MWCNT, dealt with as same as those processes of Sn@C/CNT for getting the in situ Sn-based carbon assemblies doped by MWCNT. The obtained composite is marked as Sn*@C/CNT.
![]() | ||
Fig. 1 XRD patterns (a) for Sn@C, Sn@C/CNT, Sn*@C/CNT, oxidized carbon nanotube (O-CNT), and nano Tin (Sn); and Raman spectra (b) for Sn@C, Sn@C/CNT and Sn*@C/CNT materials. |
Fig. 1(a) presents the XRD diffraction peaks of the Sn@C nano-lamellar assemblies, which display significant intensity at 2θ values of 30.02, 30.64, 44.90, and 55.33°. These characteristic peaks, along with several smaller peaks, align with the distinctive lines of the ICDD standard card PDF#04-00673 for tetragonal crystal Sn. This indicates that Sn2+ has been completely reduced to Sn after the in situ Sn-based carbon assemblies was annealed at 800 °C in a nitrogen atmosphere. The lattice parameters for the resulting white tin correspond to a = b = 5.83 Å and c = 3.18 Å. Whatever, the in situ Sn-based carbon annealing at either 700 °C (Sn@Ca and Sn@C/CNTa) or 900 °C (Sn@Cb and Sn@C/CNTb) exhibit either the increased Sn content or the formation of the Sn phase with larger particle sizes in the complexes (see Fig. S1(a) and (b)†). These will ultimately impair their cycling performance in LIBs. The XRD pattern of the Sn@C/CNT composites also indicates that the addition of MWCNT does not alter the crystal structure of Sn, but there is a minor presence of SnO as an impurity. This impurity is confirmed by two diffraction peaks at 2θ values of 37.15 and 50.76°, which correspond to the characteristic lines noted in the ICDD standard card PDF#06-0305 for SnO. The presence of hydroxyl or carboxyl groups in the oxidized CNT can further oxidize Sn. The same SnO impurity is also observed in the XRD patterns of both Sn@C/CNTa and Sn@C/CNTb (see Fig. S1† (b)). In comparison with the XRD patterns of the purely commercial Tin particles, Sn*@C/CNT obtained through the addition of acetylene black (annealed at 800 °C) and CNT (annealed at 300 °C) to commercial tin nanoparticles sized 100 nm also displays the presence of SnO impurity. This observation indicates that the impurity arises from the oxidation process between the commercial Tin particles and the oxidized MWCNT. Furthermore, the oxidized CNT material displays acute peaks at 2θ values of 11.65, 23.48, and 28.60, alongside a large, broad XRD diffraction peak at 2θ = 26.62°. This broad peak indicates that both the oxidized CNT and the in situ carbon composites possess a crystal structure similar to that of graphite.24 The ordered carbon structures present in both the tin-situ graphite and oxidized CNT are expected to enhance the electrochemical performance of the Sn-based anode materials.
Fig. 1(b) illustrates the relationship between the D-band peak and the G-band peak in the Raman spectra for three Sn-based carbon materials. The D-band peaks, associated with the stretching vibrations of C–C and CC bonds, are observed at 1340 cm−1. The intensity of these peaks correlates with lattice defects, structural distortions, and the amorphous state of carbon atoms.25 In contrast, the G-band peaks, which result from the in-plane stretching vibrations of sp2 hybrid carbon atoms, are located at 1585 cm−1. These peaks reflect the stratification and topological structures of carbon, such as graphene and graphite.25 The ID/IG value of 1.14 indicates that the Sn@C assembly contains a significant amount of in situ carbon materials, which exhibit a higher number of lattice defects, structural distortions, and an amorphous state of carbon atoms. This condition enhances the internal activity of both carbon matrix and Sn nano particles to combine the lithium ions.26 Reducing the ID/IG ratio by enhancing the structure regularity of the Sn@C nano-lamellar assembly is essential to balance the electronic conductivity and the diffusion efficiency of lithium ions. Lower values of ID/IG realized by incorporating the oxidized MWCNT, precisely 0.12 for Sn@C/CNT and 0.20 for Sn*@C/CNT shown in Fig. 1(a), are expected to achieve this purpose.
The BET adsorption–desorption isotherms, HK micropore analysis, and BJH pore size distributions for Sn@C, Sn@C/CNT, and Sn*@C/CNT are presented in Fig. 2. The pronounced hysteresis loop observed in each adsorption–desorption isotherms shown in Fig. 2(a) indicates that all three materials possess a mesoporous structure.27 In Fig. 2(b), it is evident that Sn@C/CNT exhibits a significantly higher value of dV/dlog
D = 0.382 cm3 g−1 nm−1 in the HK micropore analysis, while Sn@C and Sn*@C/CNT only show much lower values of 0.038 cm3 g−1 nm−1 and 0.043 cm3 g−1 nm−1, respectively. This suggests that the Sn@C/CNT composites contains nearly ten times more micropores compared to the other two composites.
![]() | ||
Fig. 2 (a) N2 adsorption–desorption isotherms, (b) HK micro-pore analysis and (c) pore size distributions and specific surface area of Sn@C, Sn@C/CNT and Sn*@C/CNT. |
The BJH pore size distributions of three composites at different scales are illustrated in Fig. 2(c) and summarized in Table 1. The results indicate that the micro-pores and meso-pores in the Sn@C/CNT composites contributed the most significantly to the total specific surface area, which measured 687.69 m2 g−1, accounting for 72.8%. In comparison, the Sn@C nano-lamellar assembly had a total specific surface area of 202.16 m2 g−1, contributing 62.4%, while the Sn*@C/CNT sample had 144.37 m2 g−1, contributing 62.3%. Additionally, the Sn@C/CNT composite features the smallest average pore size of 6.09 nm and the largest D90 hole diameter of 58.69 nm among the three composites, further supporting these observations. Overall, the combination of the largest surface area (687.69 m2 g−1), significant micro-pore contributions, and the smallest average pore diameter in the Sn@C/CNT composites may play a crucial role in effectively coating the nano Sn particles, accommodating the bulk expansion of Sn, and facilitating the diffusion rate of Li+.28
TGA and DTA curves of three composites are presented in Fig. 3(a–d). The Sn content for the three composites shown in Fig. 3(a) is 20.3 wt%, 18.3 wt%, and 26.2 wt%, respectively. The post-weight of the materials annealed in a furnace tube under an air atmosphere should be converted from that of SnO2 to the weight of pure Sn. In Fig. 3(b), the Sn@C precursor displays three exothermic peaks at 110, 208, and 294 °C. The first exothermic peak at 110 °C corresponds to the release of heat from water, while the subsequent peaks at 208 °C and 294 °C indicate heat loss associated with the breakdown of 2-methylimidazole ligands.29 Notably, the additional endothermic peak observed at 356 °C not only marks the onset of the crystallization of Sn but also signifies the reduction of Sn2+ due to the reductive gases released during ligand degradation. In Fig. 3(c) and (d), further annealing of the Sn@C and Sn*@C at 360 °C in a nitrogen atmosphere was conducted with the addition of oxidized MWCNTs. The DTA of both the Sn@C/CNT and the Sn*@C/CNT precursors show weight loss temperatures around 110 °C and 229 °C, respectively. These temperatures correspond to the release of water molecules from the surface of the oxidized MWCNTs and the decomposition of oxygen-containing functional groups such as carboxylic acids and hydroxyls present in the oxidized MWCNTs.30 This indicates that the Sn-based carbon nano-lamellar assembly interacts with the oxidized MWCNTs through decomposition to form a new three-dimensional Sn-carbon materials.
Fig. 4 illustrates the topographies of oxidized carbon nanotube, Sn@C, Sn@C/CNT, and Sn*@C/CNT materials. Fig. 4(a) highlights the distinct cracks that appear on the wall of the carbon nanotube following the oxidation process. These oxidized cracks can provide more edge carbon atoms, increasing the amount of the naked functional groups (e.g., COOH, OH), and enhancing the compatibility between carbon nanotubes and Sn@C material.31 In contrast to the spherical nano Sn embedded in a disordered carbon matrix found in the Sn*@C/CNT particles (Fig. 4(d)), the Sn@C particles resemble graphite-like nano-lamellar assemblies, as shown in Fig. 4(b). This thin graphite-like slice has an average thickness of only 6 nm (see Fig. S3†). Furthermore, Fig. 4(c) illustrates that these lamellar assemblies of Sn@C particles are bound by doped, oxidized MWCNT to form the new Sn@C/CNT particles. This design aligns with our experimental goal of enhancing the conductivity of the Sn@C composites. EDS observations and corresponding mappings of the Sn@C particles reveals the presence of Nitrogen (N) in the Sn@C material, which can be reasonably attributed to the incomplete oxidation of amino groups in the presence of Sn2+ (see Fig. S4†). The EDS analysis and mappings of Sn@C/CNT particles (see Fig. S5†), indicate the presence of oxygen (O) alongside the Tin (Sn) and nitrogen (N); this O contributes to the tiny SnO impurities observed in the composites. Both the Sn and N elements in the Sn@C particle, as well as the Sn, N, and O elements in the Sn@C/CNT particle, are uniformly distributed throughout their carbon-based composites. In contrast, the elements in the Sn*@C/CNT particle are unevenly distributed (see Fig. S6†). This suggests that the simple mechanical mixing of commercial nano-Sn with both acetylene black and oxidized MWCNT cannot ensure structural consistency in the Sn*@C/CNT composites. As a result, we can predict poor electrochemical performance for lithium-ion storage and diffusion. Additionally, the Sn content in both the Sn@C and Sn@C/CNT closely matches the calculations obtained from the TGA results (see Fig. 3(a)).
Fig. 4(e and f) displays the TEM images of the Sn@C/CNT particles. In these images, the bright area represents the carbon matrix, while the small dark spots correspond to the Sn nanoparticles. It is evident that the Sn nanoparticles, which average between 2 and 3 nm in width, are well dispersed within the carbon matrix. Additionally, the lamellar structure of the in situ carbon matrix can also be observed together with the that of HRTEM image of Fig. 4(g). Fig. 4(h) presents a HRTEM image of the Sn nanoparticles within the Sn@C/CNT composites. This image reveals lattice fringe spacings of 0.278 nm and 0.291 nm, which correspond to the (101) and (200) planes of the tetragonal Sn phase, respectively. Furthermore, the micro-area diffraction pattern diffraction patterns of the Sn phase, shown in the Fourier Transform HRTEM image (Fig. 4(i)), display perfect concentric-circle patterns. This observation indicates that various clusters of the tetragonal crystalline Sn phase are arranged in an unordered manner, resulting in a polycrystalline system. Within this system, additional crystal planes, such as (211) at the secondary ring and (312) at the outer ring, are clearly visible.
The XPS survey spectra of the Sn@C/CNT composites shown in Fig. 5 confirms the presence of tin (Sn), carbon (C), oxygen (O), and nitrogen (N). The Sn 3d spectra (Fig. 5(b)) show only two BE peaks at 487.7 eV and 495.8 eV. These peaks correspond to Sn 3d5/2 and Sn 3d3/2 of tin, as reported in the literature.32 In the O 1s spectrum of Sn@C/CNT (Fig. 5(c)), the observed BE peak splits into three simulated peaks at 531.5 eV, 532.5 eV, and 533.6 eV, which can be assigned to different categories of oxygen: carboxyl (–CO), hydroxyl (H2O), and ether groups (C–O–), respectively.33 This also indicates that the oxidized MWCNT can connect to the in situ carbon assemblies through hydroxyl groups. The two BE peaks of the N 1s spectrum for Sn@C/CNT, located at 398.5 eV and 401.1 eV, as shown in Fig. 5(d), indicate the presence of both imino and amino groups within the in situ carbon assemblies.34 This formation occurs when methylimidazole is reduced by Sn2+ during the annealing process. The C 1s spectrum of Sn@C/CNT, shown in Fig. 5(e), displays four BE peaks at 284.6, 285.7, 286.4, and 288.9 eV, as determined through deconvolution analysis. These peaks correspond to various carbon categories, including the graphite structure (sp2 C
C), nitrile groups (C–N), ether groups (C–O), and carboyl groups (C
O).35 Unlike the XPS spectra of Sn@C/CNT, the Sn@C material only exhibited peaks for Sn, N, and C (see Fig. S7†). The C 1s and N 1s spectra for Sn@C showed BE peaks consistent with those observed in Sn@C/CNT. Two deconvoluted BE peaks, depicted in Fig. 5(f), are located at 284.6 eV, corresponding to the signature of graphite (sp2 C
C) and amino groups, respectively. Additionally, the BE peak at 288.9 eV can be attributed to the unavoidable absorption of oxygen by the active carbon matrix of Sn@C material.36 The XPS spectra of Sn*@C/CNT (see Fig. S8†) did not confirm the presence of nitrogen (N) in the composites, except for the presence of tin (Sn), carbon (C), and oxygen (O). This finding further proved that the source of carbon defects only originates from the annealing process of tin(II) complex with 2-methylimidazole rather than oxidized MWCNT, which has significant implications for our understanding of the material's properties.
![]() | ||
Fig. 5 XPS spectra of Sn@C/CNT: (a) survey and deconvoluted, (b) Sn 3d, (c) O 1s, (d) N 1s, (e) C 1s; and of (f) Sn@C C 1s. |
The impressive cycling performance and rate capabilities can be attributed to the effective activation of the Sn–Li alloy reaction and the excellent pathways for lithium ion diffusion. The three-dimensional carbon network, formed from a layered in situ nano-carbon matrix woven with oxidized MWCNTs, can effectively accommodate the significant volume changes of Sn particles. This structure promotes the activation of the Sn–Li alloy reaction and maintains the structural integrity of the anode materials during the lithiation and delithiation processes.
The following summarizes the electrochemical performance of various valuable anode materials, including metallic Sn and functional matrix materials, along with Sn@C/CNT, as presented in Table 2.
Materials | Cycling stability/cycles/current density (mA h g−1/cyce/A g−1) | Rate capability/current density (mA h g−1/A g−1) | Year | Ref. |
---|---|---|---|---|
a Polyaniline.b 1,2,4,5-Benzene-tetracarboxylic acid.c Two dimensional laminar matrix of graphene composites.d Helical carbon fibers.e Cycles. | ||||
Sn-CNTs | 437/100/0.1 | 429/2 | 2013 | 21 |
Sn/MoS2/C | 625/500/1 | 630/2 | 2015 | 37 |
Mn2SnO4/Sn/C | 908/100/0.5 | 550/2 | 2016 | 38 |
Sn-PMAa | 707/400/0.8 | 226/1.6 | 2019 | 39 |
Sn/C-PANIb | 855/100/0.1 | 153/10 | 2020 | 40 |
Sn@2DLMG cc | 539/500/0.1 | 240/5 | 2021 | 19 |
C/Sn/HCNFd | 610/200/0.2 | 317/2 | 2022 | 41 |
Sn@SiOC | 547/200/1 | 538/5 | 2022 | 42 |
Sn@CNT | 616.9/100/0.1 | 558/0.5 | 2023 | 43 |
Sn@C | 542/100/0.1 | 535/0.5 | 2024 | 44 |
Sn@C/CNT | 949/150/0.1 | 341/10 | This work |
Among these materials, Sn@C/CNT demonstrates exceptional cyclic performance, achieving a remarkable specific capacity of 949 mA h g−1 at a current rate of 0.1 A g−1, even after 150 cycles. It also maintains an impressive specific capacity of 341 mA h g−1 at a highest current rate of 10 A g−1.
In comparison, some significant anode materials reported in the literature include the excellent monometallic anode material of Sn/C-PANI, which reaches a specific capacity of only 855 mA h g−1 at a current rate of 0.1 A g−1 after 100 cycles, and retains a specific capacity of 153 mA h g−1 at a 10 A g−1 rate. Additionally, the energetic bimetallic anode material Sn/MoS2/C shows a maximum specific capacity of 625 mA h g−1 at 1 A g−1 after 500 cycles, maintaining a specific capacity of 630 mA h g−1 at a current rate of 2 A g−1 for rate performance.
These results indicate that the synthesis method for the new Sn@C/CNT anode material has successfully achieved the goal of developing new Sn-based anode materials for LIBs by incorporating one-dimensional MWCNTs into the in situ carbon matrix.
EIS analysis is widely used to evaluate the resistance of electrode materials and the diffusion of lithium ions within them.45 The equivalent circuit model used to fit the EIS spectrum is shown in the inset of Fig. 7(d). The series resistance (Rs) is approximately 7.0 Ω, representing the resistance between the electrolyte and the anode materials, as illustrated in Fig. 7(c). This indicates the suitability of the selected electrolyte. In contrast, the charge transfer resistances (Rct) for Sn@C and Sn@C/CNT are 109.8 Ω and 148.1 Ω, respectively. The Rct of 85.0 Ω for the Sn@C/CNT material shows the lowest charge transfer resistance during the electrochemical reaction, indicating superior performance. The Warburg factor (σw), shown in Fig. 7(e), significantly influences the internal resistance related to lithium ion diffusion. This factor can be derived by calculating the slope of the fundamental part of resistance (Z′) versus the −1/2 exponent of angular frequency (ω−1/2) curves, as expressed in the following formula (1):47
Z = Rs + Rct + σωω−1/2 | (1) |
The diffusion-controlled lithium ion diffusion coefficient (DLi+) can subsequently be calculated using the Nyquist plot (2):48
DLi+ = R2T2/2A2F4n4C2σW2 | (2) |
The lowest Rct and the highest DLi+ confirm that Sn@C/CNT material has a substantial advantage in electrochemical performance over the other two Sn-based materials. However, assuming that lithium ion diffusion efficiency stems solely from the Li4.4Sn alloy is an over-idealization.49 Fig. 7(g) illustrates possible lithium ion diffusion pathways, incorporating the reversible charge/discharge chemical reactions from both C6Li composites50 and LixSn (where x = 1, 2, 3, or 4.4) alloys. The lithium ion diffusion coefficient for Sn@C/CNT from the C6Li composites is 1.9 × 10−10 cm2 s−1, while the coefficients from LixSn alloys (with x = 1, 2, 3, and 4) are 1.9 × 10−10, 1.07 × 10−11, 2.11 × 10−12, and 6.70 × 10−13 cm2 s−1, respectively. These diffusion coefficients decrease in order as the charge transfer number (x) increases during the reversible electrochemical reactions of the LixSn alloy, despite considering the counterbalancing effects of the lithium ion concentration and Rct on DLi+. Similarly, Sn@C and Sn@C/CNT materials exhibit DLi+ values at least one hundred times lower across all stages of the Sn–Li alloy reaction, demonstrating the same trend in lithium diffusion coefficients. Given that the order of magnitude for the highest DLi+ is at least 10−13, it is unlikely that the total reversible specific capacity for the three Sn-based anode materials is solely derived from the higher lithium-rich alloy Li4.4Sn. Instead, a series of reversible electrochemical reactions likely indicates optimal diffusion-controlled dynamics, suggesting that up to five Sn–Li alloys can emerge during these charging and discharging cycles. Therefore, the normal charge transfer number is approximately three rather than 4.4 in these SnLix phases.51 Consequently, the optimal lithium diffusion coefficient is estimated to be in the range of 10−9 to 10−10. The promising anode material Sn@C/CNT shows adequate kinetics to sustain long-term cycling and achieves a higher specific capacity of 688 mA h g−1 at a current density of 1C, confirming its suitability.
In order to verify the analysis of diffusion-controlled lithium ion diffusion based on the Nyquist plot, Lithium ion coefficients of DLi+ based on the GITT curves for three anode materials were finished. The GITT DLi+ can be calculated by using the Weppner–Huggins plot:52
![]() | (3) |
Fig. 8(a–c) demonstrate that both Sn@C/CNT and Sn@C anode materials exhibit more pronounced voltage hysteresis and higher GITT specific capacitance (1100 mA h g−1 for Sn@C/CNT and 700 mA h g−1 for Sn@C) compared to Sn*@C/CNT. This indicates that interior-controlled processes predominantly influence the lithium-ion dynamic behavior in the first two materials at a lower exchange current density of 0.1C. In contrast, surface-controlled lithium-ion dynamic behavior prevails for Sn*@C/CNT at the same current density.53 Fig. 8(d) illustrates the variations in DLi+ across different discharge and charge voltages. When comparing the evolution of DLi+ for Sn@C/CNT, these values can be categorized into four distinct ranges: (a) DLi+ values in the order of magnitude of 10−10 to 10−11 cm2 s−1, occurring within a discharge voltage of 1.8 to 1.4 V and a charge voltage of 0 to 0.5 V; (b) DLi+ values around the order of magnitude of 10−11 cm2 s−1, within a discharge voltage of 1.4 to 0.75 V and a charge voltage of 0.5 to 0.75 V; (c) DLi+ values in the order of magnitude of 10−12 to 10−13 cm2 s−1, occurring within a discharge voltage of 0.75 to 0.5 V and a charge voltage of 0.75 to 1.5 V; (d) DLi+ values in the order of magnitude of 10−14 to 10−15 cm2 s−1, occurring within a discharge voltage of 0.5 to 0 V and a charge voltage of 1.5 to 2.5 V. The four components of DLi+ closely align with the analyses of various DLi+ variations corresponding to at least four electrochemical reactions. These include the intercalation/deintercalation lithium reaction of C6Li and the alloying reaction of LixSn (where x can be 1, 2, 3, or 4.4) for the Sn@C/CNT, as illustrated in Fig. 7.
Additionally, Fig. 8(d) also clearly shows that the Sn@C anode material exhibits lower DLi+ values across all four sections of DLi+ variation, and the Sn*@C/CNT anode material displays even lower DLi+ values, making it challenging to differentiate the four parts of the DLi+ distribution within similar discharging and charging voltages.
log(i) = b![]() ![]() | (4) |
The slope b in formula (4) is a adjustable parameter which be used to analyze the different diffusion mechanism of lithium ion in electrode material. Typically, when b = 0.5, the principal contribution derives from the diffusion-controlled process. As the b-value approaches 1.0, the capacity-controlled process dominates.51 As shown in Fig. 9(b), LCV curves of the Sn@C/CNT electrode, reveals the variety of summit currents at different scan rates from 0.1 to 1.0 mV s−1. Similar to two broad oxidation current peaks at 0.53 and 1.20 V in the anodic scans, two reduction peaks at 0.08 and 1.0 V in the cathodic scans were also chose to calculate the adjustable parameter b through plotting method based on the formula (5). Fig. 9(c) shows the b-values of the four peaks are 0.925, 0.819, 0.840 and 0.760 responding to summit currents from peak 1 to 4, respectively. It means that a mixed lithium ion diffusion mechanism is dominated by diffusion-controlled and surface-controlled processes.51 Fig. 9(d)–(e) quantitatively describes the their contribution rate of the capacity-controlled diffusion to the total diffusion of lithium ion.55 It indicates that the capacity-controlled diffusion win more and more proportion in the total diffusion control for Sn@C/CNT material when the scan rate increased from 0.1 to 1.0 mV s−1, and the utmost contribution rate can be up to 79% at scan rate of 1.0 mV s−1. In contrast with those of Sn@C/CNT, the capacitive contribution of Sn@C electrode is lower at each scan rate, and only has as lower as 53% at the same scan rate of 1.0 mV s−1 (Fig. 9(f)).
The capacity-controlled diffusion of lithium ion can be known as the pseudo-capacitance control, and be studied in advanced by Randles–Sevcik eqn (5), in which the equation is always used to resolve the interface catalytic activity of the catalyst.56
Ip = 2.69 × 105n3/2A(DLi+)1/2ν1/2ΔC0 | (5) |
![]() | ||
Fig. 10 Lithium ion diffusion efficient based on the pseudo-capacitance: (a) DLi+ of Sn@C; (b) DLi+ of Sn@C/CNT anode material; (c) DLi+ of Sn*@C/CNT. Peak 1, 2, 3(a), 3(b), and 3(c) refer to a reversible Sn–Li alloy reaction between y = 2 and 3, y = 0 and 2, y = 0 and 4.4, y = 0 and 4, y = 0 and 3, respectively; and peak 4 refer to a reversible C6Li reaction between carbon matrix and Li. Tables S2–S7† provide detailed values of capacity-controlled lithium-ion diffusion coefficients for various reversible redox reactions. |
Fig. 10(a) and (c) reveal that both Sn@C and Sn*@C/CNT all possess extraordinary lower delithiation–lithiation DLi+ than those of Sn@C/CNT as shown in Fig. 10(b). The lower adjustable parameter (b = 0.705, 0.592, 0.529 and 0.402 for Sn@C, see Fig. S10,† and b = 0.486, 0.445, 0.492 and 0.488 for Sn*@C/CNT, see Fig. S11†) further confirmed this conclusion.
The significant variations in the series of DLi+ values corresponding to different peak currents, as illustrated in Fig. 10(a)–(c), indicate distinct activation dynamics for each reversible reaction, whether it involves different Sn–Li alloys or the C–Li compound (C6Li). Furthermore, this suggests that the capacity-controlled diffusion performance of Sn@C and Sn*@C/CNT is considerably inferior compared to that of Sn @C/CNT.
The capacity-controlled lithium ion diffusion coefficients (DLi+) of Sn@C/CNT, as shown in Fig. 10(b), reveal significantly higher values in the context of reversible electrochemical reactions in Sn–Li alloys. These coefficients are one or two orders of magnitude greater than those calculated through Warburg factors (refer to Fig. 7(d)). Specifically, the average capacity-controlled DLi+ of 1.72 × 10−8 cm2 s−1, which corresponds to current peak 1 around 0.53 V, is nearly one hundred times higher than the diffusion-controlled DLi+ of 1.9 × 10−10 cm2 s−1 observed between Sn and SnLi (x = 1). Similarly, the average DLi+ of 2.15 × 10−9 cm2 s−1 associated with current peak 2 around 1.20 V is approximately two hundred times higher than the diffusion-controlled DLi+ of 1.07 × 10−11 cm2 s−1 found between Sn and SnLi (x = 2).
In contrast, the average DLi+ of 1.34 × 10−11 cm2 s−1, which corresponds to current peak 3(a) around 0.73 V, is only twenty times greater than the diffusion-controlled DLi+ of 6.70 × 10−13 cm2 s−1 between Sn and SnLi4.4 (x = 4.4). This indicates higher dynamic resistance and minimal formation of SnLi4.4 species. Similar trends can be observed in the average capacity-controlled DLi+ values for both Sn@C and Sn*@C/CNT anode materials (refer to Fig. 10(a) and (c)), compared to those of the two materials shown in Fig. 7(f).
These results align well with the increasing capacity-controlled diffusion behavior in Sn@C/CNT as the charge/recharge current density rises. This implies that reversible Sn–Li alloy reactions with lower charge transfer numbers (1 ≤ y ≤ 3) can generate significantly more pseudo-capacitance due to their enhanced dynamic activity.
For the C–Li compound (C6Li), its capacity-controlled DLi+ value corresponding to current peak 4 around 0.08 V is notably higher, reaching up to 1.08 × 10−6 cm2 s−1, compared to 1.90 × 10−10 cm2 s−1 obtained from EIS testing. This indicates that the carbon matrix has sufficient active sites on the surface of the anode material, enabling the generation of a substantial amount of pseudo-capacitance under large current densities.57 The higher DLi+ value results from the low resistance of the interfacial reaction between the carbon matrix and lithium ions. Additionally, the presence of carbon defects created by nearly 9.53 wt% nitrogen (N) and 9.52 wt% oxygen (O) in the in situ carbon composites significantly enhances Li+ transportation and contributes to the formation of more active sites.58
Therefore, Fig. 11 delineates a mixed storage and diffusion mechanisms of lithium ions that are influenced by diffusion-controlled and capacity-controlled behaviors in Sn@C/CNT anode material.
The structure of the obtained Sn@C/CNT anode material, as illustrated in Fig. 11(I), consists of a graphite-like layered nano-lamellar assembly interwoven with oxidized MWCNT fibers, within which tin nanoparticles are aligned. During the charging cycle, lithium ions are sequentially introduced into the material. First, they occupy the nitrogen (N) and oxygen (O) vacancies (Fig. 11(I)). They are trapped by the external active sites of the carbon matrix (Fig. 11(II)), and finally, they combine with the internal active tin nanoparticles (Fig. 11(III)). The third step of lithium insertion is slower than the first two due to unavoidable transmission resistance.
When lithium ions are extracted from the anode material, they follow the same sequence: they separate from the N and O vacancies (Fig. 11(IV)), are released from the carbon matrix (Fig. 11(V)), and then migrate from the inner tin nanoparticles (Fig. 11(VI)). Steps I, II, IV, and V can be categorized as capacity-controlled storage and diffusion behaviors since they involve minimal resistance in the external reactions. In contrast, steps III and VI are classified as diffusion-controlled storage and diffusion behaviors due to the apparent bulk resistance during the internal response.
This mixed mechanism of lithium-ion storage and diffusion ensures that the Sn@C/CNT anode material exhibits excellent long-term charge/discharge performance, achieving a capacity of up to 946 mA h g−1 at a low current rate of 0.1C. It also demonstrates an outstanding long-term cycling performance with a capacity of up to 688 mA h g−1 at a higher current density of 1 A g−1 (1C).
The Sn@C/CNT composites exhibited impressive long-term cycling spans at both the lower current density of 0.1C and the higher current density of 1C. The dynamic properties of the Sn@C/CNT anode materials are characterized by a higher diffusion-controlled lithium ion diffusion coefficient resulting from the internal reversible charge and discharge reactions of the Sn–Li alloy, as well as a higher capacity-controlled lithium ion diffusion coefficient originating from the pseudo-capacitance induced by the reversible charge and discharge reactions of both the Sn–Li alloy and the C6Li composites.
The topographical features of Sn@C/CNT are crucial for improving both the reversible specific capacity and the long-term cycling lifespan of the material. These features include nano-sized Sn active particles that measure as small as 2 to 3 nm, which are uniformly dispersed within a nano-lamellar carbon matrix, and the matrix consisting of graphite-like layers that are nearly 6 nm thick. Additionally, a comprehensive analytical approach to determining the mixed lithium-ion diffusion mechanism provides a deeper understanding of the electrochemical behavior of electrode materials during internal or interface reactions.
Footnote |
† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d5ra02378e |
This journal is © The Royal Society of Chemistry 2025 |