2.1 Synthesis of Free-Standing Tin Phosphide/Phosphate Carbon Composite Nanofiber Mats

Free-standing tin phosphide/phosphate carbon composite nanofiber mats were synthesized by electrospinning with heat treatments as illustrated in Figure 1.

As can be seen from Figure 2a, the X-Ray diffraction (XRD) patterns of the sample, prepared from the solution without a P source, corresponds to a tetragonal Sn with a space group I41/amd. Adding P leads to a shift of Sn peaks to the higher angles, possibly implying the formation of P-doped Sn and Sn_xP.^[11] At an Sn:P molar ratio of 1:1, additional small-intensity peaks of tetragonal SnP with a space group of I4mm appear at 27.65° and 33.04°. However, after further increment of the P content, all crystalline peaks disappear, confirming the formation of an amorphous structure. Increments of the H₃PO₄ content in the precursor solution to an Sn:P molar ratio of more than 1:1.3 lead to difficulties in preparing uniform fibers of good flexibility (Figure S1, Supporting Information). Since the 1:1 sample annealed at 600 °C also show an amorphous structure and SnP peaks disappear after annealing at 800 °C (Figure S2a, Supporting Information), 700 °C is admitted as an optimal annealing temperature.

Figure 2b shows the Raman shifts of different samples. All samples possess typical D and G bands at 1347 and 1588 cm⁻¹, respectively, implying the formation of graphitic carbon.^[22] The highest graphitization degree of carbon is observed for pure Sn and 1:1 samples with the lowest I_D/I_G ratio (Table 1). Additional small-intensity peaks of phosphate appear at around 950 cm⁻¹ on the spectra of P-containing samples when Sn:P molar ratio is equal to or higher than 1:1, implying the presence of residual Sn_xPO₄. The absence of this peak in the 1:0.5 sample may indicate its complete reduction. Additional peaks between D and G bands at 1444 cm⁻¹ in the spectra of P-containing samples represent -CH₂- stretching common for P- and/or Sn-containing compounds.^{[23, 24]} Such compounds could be formed during the high-temperature evaporation of P and Sn, as the corresponding peak disappears in the sample annealed at 600 °C and intensifies at 800 °C (Figure S2b, Supporting Information). According to the CHNS analysis results, the carbon content increases with the addition and increment of the P content but decreases back when the ratio reaches 1:1.3, varying from 48.56 to 56.83 wt% (Table 1). It could be attributed to the decomposition features of the H₃PO₄-reach precursor fibers during annealing, as H₃PO₄ has a catalytic effect on the decomposition and carbonization of PVP.

Figure S3a, Supporting Information presents the Sn 3d X-Ray photoelectron spectroscopy (XPS) spectra of various samples derived from electrospinning solutions containing different Sn:P molar ratios. In the pure Sn sample, a distinct peak corresponding to metallic tin appears at ≈485 eV. However, upon introducing a P source, this metallic tin peak becomes indiscernible. The absence of metallic tin signals in the XPS spectra of P-containing samples provides further evidence for the formation of P-doped Sn (denoted as Sn-P) and Sn_xP compounds. Moreover, a shift towards higher binding energies in the prominent peak originating from the surface oxidation layer (≈496 eV) confirms alterations in the surface chemistry. Additionally, the intensity of the P–O band arising from the surface oxidation layer, along with residual Sn_xPO₄, increases as the P content rises (Figure S3b, Supporting Information).

To reconfirm the formation of the Sn_xP and the presence of Sn_xPO₄ in the 1:1 sample, an XPS depth study was performed, and the results are shown in Figure 2c,d. On the Sn 3d XPS spectra, a shift from 495.6 and 487.3 to 495.1 and 486.8 eV is observed after surface etching due to the appearance of the SnP peak and its intensification.^[11] The formation of the SnP is more obvious from the phosphide peak that appeared at 130.1 eV on the P 2p XPS spectra in Figure 2d. The presence of a P–O bond at 133.8 eV even after etching confirms the residue of some Sn_xPO₄ phase. An additional peak of the P–C bond is observed before etching at 132.5 eV, reconfirming the formation of P-doped carbon during the evaporation of phosphorus upon high-temperature annealing.^[19] The fact of phosphorus evaporation is clear from the EDS mapping and elemental composition of the sample surface after annealing at different temperatures presented in Figure S4 and Table S1, Supporting Information, respectively. From these results, the formation of tin phosphide/phosphate carbon composites with a gradient concentration of phosphorus could be confirmed with a P-rich inner Sn_xP layer and a Sn-rich outer layer (Figure 1). Such a structure is expected to possess improved electrochemical properties, as a combination of a small amount of Sn with Sn_xP has been proven as an effective approach to stabilize the cycle performance of Sn_xP.^[25] Notably, no peak of P-doped carbon is observed in the 1:1.3 sample in Figure S3b, Supporting Information, implying restricted evaporation of phosphorus upon increasing H₃PO₄ content in the precursor solution. Based on these results, the effect of the Sn:P molar ratio on the possible compositions of the samples can be reconfirmed and summarized in Table 1.

The morphology of the materials was observed by scanning electron microscopy (SEM). All samples possess a uniform fibrous structure (Figure 3a). Pure Sn nanoparticles of about 50 nm are mainly decorated on the surface of carbon nanofibers in the P-free sample. However, such nanoparticles disappear after introducing H₃PO₄ to the precursor solution, suggesting their encapsulation within carbon nanofibers. This fact is reconfirmed by the transmission electron microscopy (TEM) images in Figure 3b. Aggregates of nanoparticles of about 20 nm are encapsulated in the carbon nanofibers when Sn:P molar ratio is equal to 1:0.5. Increasing the P content leads to the formation of finer nanoparticles of about 5–10 nm with a more uniform distribution within the fiber. Furthermore, nanoparticles in the prepared composite have enough room within the carbon fiber matrix for mitigating the volume expansion upon lithiation, keeping the structural integrity, and expanding the cyclic stability. Too fine nanoparticles in the case of the 1:1.3 sample could be a possible reason for the absence of crystalline peaks in its XRD patterns in Figure 2.

Figure 3c shows energy-dispersive spectroscopy (TEM-EDS) elemental mapping in a single fiber of the 1:1 sample. Sn mapping overlaps well with P and O mappings, reconfirming the formation of tin phosphide/phosphate within the carbon nanofiber matrix. From high-resolution TEM images in Figure 3d, the encapsulation of nanoparticles of different crystallinity in the amorphous carbon matrix can be observed. The dark-colored amorphous phase could be attributed to Sn_xPO_4, while different interplanar spacings of 0.225 and 0.29 nm imply the (110) plane of Sn_xP and (200) plane of P-doped Sn, respectively.^{[10, 26]}

Figure 4 illustrates the possible effect mechanism of the Sn:P molar ratio on the composition and particle size of samples. The formation of the finer nanoparticles with an increase of the H₃PO₄ content could be attributed to the formation of a strong phosphate carcass that interacted with functional groups of PVP and formed isolated nucleation sites for the restricted growth of particles.^{[27, 28]} Lack of phosphate anions in the precursor structure, on the other hand, allows free diffusion of ions, growth of bigger particles, and phosphorus evaporation.

2.2 Electrochemical Properties of Free-Standing Tin Phosphide/Phosphate Carbon Composite Nanofiber Mats at Room Temperature

Figure 5a demonstrates cyclic voltammetry (CV) curves of the samples prepared from electrospinning solutions with different Sn:P molar ratio at a scan rate of 0.1 mV s⁻¹. The CV curves of the 1:1 and 1:1.3 P-containing samples are typical for the Sn_xP and Sn_xP/C composites (Figure S5b,c, Supporting Information).^[29] Additional broad peaks at higher potential regions could be responsible for the reversible lithiation of Sn_xPO₄ with the further formation of Li₃PO₄ and Sn (1).^[30] Wide peaks that appear between 1.1 and 0.35 V during the cathodic scan indicate a multi-step reaction that involves Li⁺ intercalation into SnP and conversion of the intercalation products to Sn and Li₃P phases (2).^[7] As observed in the cathodic scan, for the pure Sn sample three small peaks located below 0.5 V can be attributed to the alloying reaction between lithium and tin, forming Li_xSn alloys (3), while a peak located at 0.8 V might be a sign of a solid electrolyte interphase (SEI) formation and which disappears in the consequent cycles (Figure S5a, Supporting Information).^[31] On the other hand, the shape of the CV curves of the 1:0.5 sample looks more like a mixture of that of Sn_xP and free Sn (Figure S5d, Supporting Information) with a lack of Sn_xPO₄, which could cause a smaller shift of Sn peaks on its XRD pattern and smaller intensity of phosphate peak on the Raman spectrum in Figure 2. During the first anodic scan of the pure Sn sample, four oxidation peaks between 0.4 and 1.0 V are assigned to the delithiation process of Li_xSn alloys that are reversible (4).^[32] The anodic peaks at 0.5, 0.6, and 1.21 V on the CV curves of the P-containing samples suggest a stepwise delithiation process of Li_xSn, Li₃P, and Li₃PO₄ phases with conversion to tin phosphide/phosphate (5) and (6).

The capacity of the tin phosphide/phosphate carbon composite anodes strongly depends on their composition, morphology, and microstructure. The difference in charge-discharge curves in Figure 5b also confirms the changes in the composition of the electrodes compared to pure Sn, although changes in microstructure in Figure 3b could be another reason. The highest initial discharge and charge capacities of 2313 and 1453 mAh g⁻¹, respectively, have been achieved at 50 mA g⁻¹ for the 1:1 sample. In contrast, a pure Sn sample has only 992 mAh g⁻¹ reversible capacity, which is very close to its theoretical value. Furthermore, it has poor cyclability with a rapid capacity degradation in the initial 20 cycles, retaining only 524 mAh g⁻¹ capacity after 100 cycles (Figure 5c). It could be because of a significant volume expansion and detachment of Sn particles from the surface of carbon fibers. Introducing P leads to higher reversible capacities and more stable cycle performance than for pure Sn samples. All three P-containing samples have retained almost similar charge capacities of about 940 mAh g⁻¹ after 100 cycles at 50 mA g⁻¹ with 100% Coulombic efficiency.

Figure 5d shows the Nyquist plots of samples after 100 cycles that consist of two semicircles at low-to-medium frequency regions implying resistances coming from the SEI (R_SEI or R_surface) and charge transfer (R_ct), and an oblique straight line at a high-frequency region corresponding to the Warburg impedance (W) (Figure S6, Supporting Information). As it can be seen from Table 2, the pure Sn sample possesses the highest R_ct, probably because of the loss of good contact between the Sn nanoparticles and carbon fibers upon the volume expansion during lithiation. Introducing P and increasing its content reduce the R_ct significantly. Despite the similarities in the cycle performance and retained capacity of P-containing samples, the 1:1 sample exhibited the lowest overall resistance.

Figure 5e shows the cycle performance of the samples cycled at a current density of 250 mA g⁻¹ at room temperature. Despite a small capacity decay within the initial 20 cycles, the sample of 1:1 ratio shows the best electrochemical performance with further cycle-by-cycle increasing capacities and stabilization at 1141 mAh g⁻¹ at the 300^th cycle. Such a trend is typical for the Sn_xP anode and can be explained by the activation of the electrode.^[7] On the other hand, the samples of 1:0.5 and 1:1.3 ratios have much lower stability, probably originating from higher resistances that compromise the fast transfer of electrons and lithium ions.

The stability of the prepared samples was checked by post-mortem ex situ characterizations after 100 cycles as an anode for lithium-ion batteries at a current density of 50 mA g⁻¹ at room temperature. Figure S7a, Supporting Information shows the XRD patterns of the samples prepared from electrospinning solutions with different Sn:P molar ratios after the 100^th charge. The crystallinity of samples has almost not changed (except for the 1:1.05 sample), indicating their good structural stability. The absence of 50-nm nanoparticles (observed in Figure 3a) on the surface of deformed nanofibers of a pure Sn sample in Figure S7b, Supporting Information confirms their detachment during cycling which could be responsible for its rapid capacity decay. The appearance of additional nanoparticles of smaller sizes may indicate the expansion of partially encapsulated particles. Even though, well-encapsulated nanoparticles can still be observed within the nanofiber matrices of all samples in Figure S7c, Supporting Information. On the other hand, unlike other P-containing samples, preservation of the original fibrous morphology and formation of a uniform SEI layer on the surface of the 1:1 sample can be observed. These results confirm the structural stability of the designed materials and the formation of a uniform stable SEI layer facilitating fast charge transfer even after 100 cycles. The improved performance of the 1:1 sample compared to other samples can be attributed to the synergistic effect of P-doped Sn, Sn_xP, and Sn_xPO₄, and the uniform distribution of fine nanoparticles of sufficient crystallinity within the carbon fiber matrix.

2.3 Low-Temperature Performance of Free-Standing Tin Phosphide/Phosphate Carbon Composite Nanofiber Mats

The electrochemical properties of the free-standing tin phosphide/phosphate carbon composite nanofiber mats, prepared from electrospinning solutions with different Sn:P ratios, were further investigated at 0 and −20 °C. Figure 6a,b show the initial potential profiles at 25 mA g⁻¹ at 0 and −20 °C, respectively. All potential plateaus of the samples at these temperatures resemble the profiles obtained at room temperature with a small shift in potential, which might be ascribed to a high overpotential due to the slower diffusion of ions at lower temperatures. Despite lowering the operating temperature to 0 °C, all samples have high initial discharge and charge capacities of above 1500 and 950 mAh g⁻¹, respectively. The highest reversible capacity of 1200 mAh g⁻¹ is achieved for the sample with a ratio of 1:1.

The temperature-capability of the 1:1 sample at 100 mA g⁻¹ is shown in Figure 6c in comparison to the commercial graphite anode. The developed electrode maintains about 200 mAh g⁻¹ capacity even at −30 °C, while commercial graphite completely loses its lithium storage capability at −20 °C. The decreased capacities result from the sluggish kinetics of the electrodes and the contribution from the electrolyte at low temperatures, which hinders the fast lithium-ion diffusion. However, the fact that the capacity values of the developed electrode largely exceed those for graphite anode at low temperatures, allows us to assume that the approach described in this work opens the perspectives to target the LT performance issues of the electrode materials for LIBs. Furthermore, an intriguing phenomenon of an inverse increase in capacity at subzero temperatures is observed for the developed material. This phenomenon is not uncommon for tin phosphides, even at room temperature, and is attributed to the gradual activation of the electrode with each subsequent cycle.^[7] The observation of such a trend at subzero temperatures provides further evidence for the promotion of electrode activation and the reversibility of electrochemical reactions facilitated by low temperatures. It is speculated that the underlying mechanism behind this behavior could be related to allotropic changes of tin, resulting in the formation of α-Sn that enhances the ionic conductivity within the material.^[3]

Figure 6d shows the cycle performance of the developed electrodes at −20 °C and 50 mA g⁻¹. After the increase of the capacity in the initial five cycles, the sample shows stable cycle performance up to 100 cycles retaining 580 mAh g⁻¹ capacity with about 100% Coulombic efficiency. The inverse increase of the capacity in the initial five cycles could probably be attributed to the activation of the electrode with gradual stabilization of the SEI layer at low temperatures.

Ex situ surface analysis was conducted after charge–discharge states of different cycles at 50 mA g⁻¹ at −20 °C to find out the nature of the activation processes in the initial five cycles. Cycled cells were disassembled in the glove box, the electrodes were washed with DMC solvent and dried at 60 °C in a vacuum. Figure 7 shows ex situ XRD patterns of developed electrodes after full discharge and charge at different cycles at −20 °C. The reversible formation of Li₃PO₄ peaks is observed, suggesting that electrochemical reactions on the electrode surface are mainly dominated by lithiation of Sn₃(PO₄)₂ (1).^[30] Although the irreversible formation of Li₃P is observed in the initial five cycles,^[33] the surface of the electrode seems to be Li₃P-free after the 100^th charge. It may indicate the importance of Li₃P formation on the further activation of the electrode after five cycles. Sn peaks are present on all patterns regardless of charge-discharge state and cycle number, indicating hindered kinetics of the alloying reaction at low temperatures, although the intensity of Li_xSn increases by the 100^th discharge.

To clarify the processes that occurred during charging, dQ/dV curves and ex situ XPS spectra after the first and fifth charges were compared. Figure 8a shows charge dQ/dV curves of developed electrodes cycled at 50 mA g⁻¹ at −20 °C. The number and positions of oxidation-reduction peaks change drastically with the change in the temperature, reconfirming incomplete electrochemical reactions. Starting point of the charge peaks gradually shift to a lower potential region, suggesting a decrease in the overpotential and activation of additional electrochemical reactions and/or their increased reversibility.

Ex situ depth XPS analysis was further conducted to better understand the activation mechanisms of the developed electrode at low temperatures. The surfaces of the electrodes were etched for 60 s to eliminate the oxidation layer. The Sn_xP peak in Sn 3d XPS spectra in Figure 8b intensifies after five cycles, confirming the increased reversibility of the electrochemical reactions. A shift of Sn_xP and Sn_xPO₄ peaks to higher binding energies suggests an increased oxidation state of tin.^[26] The appearance of LiF of high ionic conductivity in the F 1s XPS spectrum (Figure 8c) after five cycles can imply stabilization of the SEI layer, which also facilitates electrode activation.

From these results, increased reversibility of the electrochemical reactions with the formation of LiF-rich SEI layer after five cycles at −20 °C without prior room-temperature or low-current-density activation and using commercial electrolyte solution without any additives can be confirmed. These findings contribute to our understanding of the intricate interplay between temperature, material properties, and electrochemical behavior, potentially paving the way for improved energy storage systems in extreme environments.