Enhanced Performance with Nano-Porous Silicon in TiFeSi₂/C Composite Anode for Lithium-Ion Batteries

Abstract

The innovative design of the microstructure of silicon-based composite anodes in Li-ion batteries holds great potential for overcoming inherent limitations, such as the significant volume change experienced by silicon particles. In this study, TiFeSi₂/C composites prepared using micro, nano, and porous silicon showed reversible capacities of 990.45 mAh.g⁻¹, 1137.69 mAh.g⁻¹, and 1045.43 mAh.g⁻¹ at C/10. The results obtained from the electrochemical characterization show that the porous structure of the composite anode material created via acid etching reduced silicon expansion during the lithiation/delithiation processes. The void spaces formed in the inner structure of the porous silicon and the presence of carbon increased the electronic conductivity between the silicon particles and, on the other hand, lowered the overall diffusion distance of Li⁺. This study confirms that TiFeSi₂/C prepared with porous silicon dispersed in a transition metal matrix delivers better electrochemical performance compared to micro and nano silicon with a retention of 80.16%.

1. Introduction

One of the growing concerns globally over a few decades has been the consumption of fossil fuels to meet increased energy demands. The emission of greenhouse gases has been the direct offshoot of fossil fuel production, harming the environment and threatening the social welfare of humanity. The direct response to these issues has been the exploration of clean and renewable energy sources [1]. Additionally, demand for electric vehicles has undergone a huge boom to gradually reduce fossil fuel vehicles. However, utilizing energy sources like wind or solar is not dependable due to their inherent intermittency nature [2]. Lithium-ion batteries have proven to be important storage devices as the demand for portable electronic devices and electric vehicles continues to rise [3]. Nonetheless, the overall performance and cost of batteries are very critical to ensure a holistic transition to cleaner energy [4]. Certain properties such as retention, coulombic efficiency, capacity, cycling performance, energy and power densities amongst others are critical in building high-energy batteries [5]. Silicon, due to its low lithiation potential (0.2 V vs. Li/Li⁺) and high theoretical capacity, has been extensively used to synthesize lithium-ion battery anodes with superior capacity, energy density (Wh/L), and specific energy (Wh/kg) [6]. Pure silicon, however, due to its innate low electrical conductivity and its large stress/strain response during lithiation/delithiation, hampers its practical use in synthesizing battery electrodes [7]. The diffusion of lithium ions in silicon is restricted to ~100 nm per hour at ambient temperature due to its low diffusion coefficient of ~10⁻¹⁷ m²/s [8]. The volume expansion of silicon leads to the fracture of particles, delamination of active materials from the current collector, loss of overall electronic and ionic conductivity [9], consumption of the electrolyte which brings about the repeated formation of a semi-conductive passivation film known as the solid electrolyte interphase (SEI) as cycling continues [10], loss of retention, and ultimately overall capacity fading [11]. To reduce the drawbacks of silicon, several approaches have been exploited to develop electrode material capable of delivering high and stable cyclability [12]. Some of these approaches include redesigning the hybrid structure, manipulation of the SEI [13], surface engineering [14], and electrolyte design.

Transition metal silicides stand out as promising anode materials capable of improving cycling stability by buffering the volumetric changes of silicon when silicon is dispersed in an inactive metal matrix [15]. Varying particle sizes of silicon with distinct morphologies and internal structures, such as micro [16], nano [17,18], and porous [19] silicon show improved electrode performance when used alongside transition metals [20]. J. Cheng et al. [21] used SiO₂ and TiO₂ powders in a molar ratio of 20:1 powder blended with 25 wt.% NH₄HCO₃ then sintered at a temperature of 500 °C for 3 h to obtain porous SiO₂/TiO₂ precursor. Further electrolysis was carried out on the product to yield porous TiSi₂/Si. The TiSi₂ alloy formed was inactive and no side reactions with lithium were detected. High initial discharge capacity, retention, and first coulombic efficiency of 2107 mAh.g⁻¹, 1703 mAh.g⁻¹, and 89.96%, respectively, were delivered by the electrode for 100 cycles at a current density rate of 0.2 A/g. An alloy matrix with nanocrystalline silicon embedded in it was prepared by S. H. Kim et al. [22] using arc melting then followed by single-roll solidification. The alloy comprised Al₄Cu₉, AlFe, and TiFeSi₂, with silicon embedded in its matrix. The alloy showed an initial discharge capacity and capacity retention of 1459.3 mAh.g⁻¹ and 85.7%, respectively, at the current rate of 300 mA/g after 200 charge/discharge cycles. The enhanced performance of the alloy was due to the unique nano-sized nature of silicon contained therein. The nanosized silicon successfully reduced the distance of Li⁺ transport as well as the fracturing of silicon during cycling. A study by W. Ren et al. [23] prepared using the ball milling and spray drying method, porous silicon/carbon microspheres as anode materials for lithium-ion batteries. Their study showed that there was an average charge capacity and current capacity of 589 mAh.g⁻¹ and 50 mA.g⁻¹, respectively, when porous silicon/carbon microspheres were used as anode material. W. Cao et al. [24] investigated the morphological structures and electrochemical performance of silicon obtained through acid-etching with aluminum-silicon (AlSi) as precursor. The study shows that the porous structure of silicon can be tuned by dealloying with Al acting as sacrificial material. The porous microstructure of silicon synthesized using C₂H₂O₄ displayed a specific capacity of 1063 mAh.g⁻¹ after 200 cycles. The formation of porous structures provides void spaces which can alleviate the volumetric expansion associated with silicon, effectively improving its micro-structural integrity and cyclability of electrodes.

Herein, we present a comparative study of TiFeSi₂ composite prepared with micro silicon as in our previous work with similar composites synthesized using nano and porous silicon sources. Using porous silicon is shown to have great potential of relieving more stress from the volumetric expansion of silicon and improving cyclability. In this study, TiFeSi₂ composite anode material for lithium-ion batteries was prepared through a two-step synthesis method. Porous silicon was obtained by acid etching AlSi₃₅ and then followed by milling with elemental titanium, iron, and carbon powders. The composite prepared with porous silicon exhibited an outstanding initial coulombic efficiency of 99.8%. This is attributed to the voids present on the silicon particles, which effectively alleviate its expansion and stabilize the electrode structure. The inactive matrix formed and the carbon conductive network also contribute to the high efficiency and overall cell performance. Additionally, ball milling of the acid-etched porous silicon not only reduced the particle sizes of silicon but also enhanced the electrochemical performance of the composite, resulting in improved cycle life. Notably, a stable capacity of 646.97 mAh.g⁻¹ was achieved from the 103rd cycle, with a retention rate of 87.01% at a 1 C rate for 200 cycles. This was achieved using a fast, easy, and highly efficient technique with low-cost and abundant starting materials, thus making pTiFeSi₂ a promising high-performance anode composite material.

2. Experimental Procedure

2.1. Sample Preparation

Synthesis of porous silicon. Al-Si alloy with a weight ratio of 65:35, purchased from Sigma Aldrich (St. Louis, MO, USA) was used as a precursor to obtain porous silicon powder by acid etching [25]. The Al-Si alloy was then dispersed into 50 mL of 1.0 M hydrochloric acid (HCl ACS 37%) aqueous solution in a stepwise manner while stirring in a beaker. Continuous slow stirring operation was performed for 24 h at 200 rpm until no bubbles were observed, signifying that aluminum was totally dissolved from the Al–Si alloy. The solution was filtered with filter paper (Whatman quantitative filter paper, Grade 540), washed 3 times with de-ionized water (1.1 µS/cm), and then rinsed with high purity ethanol to evaporate all traces of H₂O present in the porous Si powder. To confirm that the de-ionized water was pure and free from unwanted contaminants, conductivity tests were conducted using EC700 conductivity meter (Apera conductivity). The conductivity of the water was found to be 1.1 µS/cm which is suitable for preparation of the HCl solution. Finally, the porous Si powder was dried in an oven under vacuum at 80 °C overnight.

Synthesis of TFSC composites. TiFeSi₂/C (TFSC) composites were synthesized by mixing the prepared porous silicon powder with Fe (Sigma Aldrich, 99.5% purity, 5 mm), Ti (Sigma Aldrich, 99.5% purity, >300 mesh) basis with a weight ratio of 1:1:2:0.1. The mass of precursors used was Ti—0.5391 g, Fe—0.6290 g, Si—0.6325 g, and C—0.2008 g. High-energy ball milling in the presence of argon (Ar) was employed at the speed of 500 rpm for 48 h. TFSC composites made from micro-Si (Sigma Aldrich, 99.5% purity, 2 mm) and nano Si (<100 nm ≥ 98% trace metals powders) were also synthesized following the same procedure as that of porous silicon, and the same weight ratio with Ti, Fe, and SP, respectively. TFSC made with different silicon precursors were denoted as pTFSC (porous silicon), mTFSC (micro silicon), and nTFSC (nano silicon). After milling, the amount of product was about 1.91 g, showing a high product recovery of 95.5% of the starting materials. The little loss in product was from the powders sticking to the walls of the milling jars and balls. This shows that the milling process was efficient and non-destructive. Silicon reacts with Li⁺ to form an alloy with high gravimetric energy density. At 400 °C, silicon can be lithiated to form Li_4.4Si alloy with a theoretical capacity of ~4200 mAh.g⁻¹ [26]. The theoretical capacity can be evaluated using Equation (1).

$$Q={\displaystyle \frac{nF}{M_w}}$$

(1)

where Q = specific capacity, n = number of moles, F = Faraday’s constant (mAh/mol), $M_w$ = molecular weight.

The specific capacity of silicon contained in TFSC was calculated by multiplying the theoretical capacity by the percentage mass (30%) of silicon contained in the composite. A detailed description of the synthesis method used to prepare pTFSC is presented in Figure 1.

2.2. Cell Preparation

The working electrode was composed of 70 wt.% active material (pTFSC), 17 wt.% SP, and 13 wt.% carboxymethyl cellulose (CMC-average M_w~90,000) binder dissolved in water and coated on a copper foil (MSE PRO 5kg/roll Lithium Battery Grade Copper Foil (280 mm W × 9 um T) for Battery Anode Substrate). The electrolyte used was 1.0 M LiPF₆ (battery grade, ≥99.99% trace metals basis) dissolved in EC/DEC (1:1) with 2 vol.% vinylene carbonate (VC) battery grade, 99.5%, acid < 200 ppm, H₂O < 100 ppm, and 10 vol.% fluoroethylene carbonate (FEC) battery grade, ≥99%, acid < 200 ppm, anhydrous additives. Coin cells were assembled in an Ar-filled glovebox using a Celgard separator (25 µm monolayer microporous membrane (PP) and tested using a lithium foil counter electrode.

2.3. Electrochemical Tests

The galvanostatic discharge tests were all carried out on a battery test system (BTS) purchased from Neware Technology Limited between 0.05 and 1.2 V vs. Li/Li⁺ using coin cells (type CR2032). Different current rates were applied as specified in Section 3.3, Section 3.4 and Section 3.5.

2.4. Material Characterization

To determine the phase composition of TiFeSi₂/C composites, X-ray diffraction (XRD X-Pert PRO MPD) tests were carried out under Cu radiation = 0.154056 nm, while the microstructures were characterized by scanning electron microscopy (FESEM, FEI Inspect F50, accelerating voltage 200 V–30 kV).

3. Results and Discussions

3.1. Phase Characterization with XRD

XRD analyses were carried out to determine the existing crystalline phases of the samples, as shown in Figure 2a. The tests were carried out between a range of 10 to 90° with a scan step of 0.05° per step. The composites are named mTFSC, nTFSC, and pTFSC representing micro, nano, and porous silicon sources used to synthesize the TiFeSi₂/C composites. The largest Si (111) and Fe (110) peaks at 2Ɵ angles 28.04° and 44.5°, respectively, were drawn together in Figure 2b.

No aluminum peaks were observed in the XRD pattern in Figure 2a, showing that the removal of Al was successfully achieved with the acid etching. The pure Fe, Ti, and Si peaks in the three composites show that there were no side reactions with silicon, nor any contamination/formation of alloys during milling. We successfully obtained the desired Ti-Fe-Si composites using ball milling which is cheaper than using arc-milling which is comparatively expensive [26]. The silicon peaks in the composite pTFSC are the same as reported in the literature, where the porous silicon was obtained using a spray drying process and magnesiothermic reduction method [27]. Figure 2b shows that the crystallinity of silicon was higher in mTFSC followed by pTFSC and least in nTFSC after milling, as indicated by the diffraction peaks of Si (111) at 28.04° in all three composites. Notably, there is a slight shift of the peaks at 28.04° which is due to using the same milling time and energy for different sizes of silicon. Increasing milling time as reported can account for ≈ 3° peak shifts [28]. The application of mechanical energy to the particles during milling can cause microstrain in the crystal lattice resulting in peak shifting or broadening [29]. Different relative intensities of silicon are observed, though all composites contain the same stoichiometric amounts of silicon. Some smaller structures noticed between 30°–50° could be attributed to the formation of metastable intermetallic phases of silicon [30,31].

Williamson-Hall analysis was conducted to evaluate the effects of milling on microstrain in the crystal lattice of the composites as presented in Figure 3.

Microstrain values of −0.435% and −0.2345% obtained for mTFSC and nTFSC, respectively, suggest a compressive strain in the lattice, while pTFSC which has a positive value of 0.719% indicates tensile strain. The peak broadening observed in Figure 2b due to crystallite size is calculated using the Scherer’s relation as shown in Equation (2)

$$D={\displaystyle \frac{0.9\lambda }{\beta Cosϴ}}$$

(2)

where D = crystallite size, λ = wavelength, β = full width at half maximum and ϴ = diffraction angle

The crystallite sizes for the composites mTFSC, nTFSC and pTFSC were found to be 2.028 nm, 1.923 nm, and 2.856 nm, respectively, as shown in Figure 3. It can be observed that the particles in mTFSC are larger in size compared to nTFSC and pTFSC because the same milling time and energy were used to prepare all the composites. mTFSC therefore has larger particle sizes since it requires more energy to break down particles into smaller crystallites. nTFSC, which contains nano-sized silicon particles requires less milling energy to break down particles, showed the least crystallite size, while pTFSC prepared with porous silicon has voids and a comparably higher surface area. The pores contained in pTFSC act as sites for energy dissipation [32]. Instead of the silicon being broken into smaller crystallites, the energy from mechanical milling is possibly absorbed within the porous structure, leading to less particle fragmentation and higher crystallite sizes.

3.2. Microstructure Characterization with SEM

The morphology of the synthesized composites was characterized by SEM. The intensities of silicon peaks in the XRD results give some insights into the expected surface of the SEM images. High-purity double-sided conductive adhesive carbon tabs purchased from Rave Scientific having a thickness of 125 µm were used. The tests were conducted with magnification between 500 nm–10 µm to electrochemical performance of the various composites. Figure 4 shows the SEM images of pristine TFSC composites with the three different silicon sources of mTFSC (Figure 4a,b), nTFSC (Figure 4c,d), and pTFSC (Figure 4e,f).

The silicon particles in mTFSC (Figure 4a,b) are curve-edged and flaky in morphology with negligible formation of inactive silicon during milling. Figure 4c,d shows smaller particle size compared with the other three composites, and a homogenous microstructure and size distribution are well achieved in the nTFSC composite matrix. The fluffy surface structures of pTFSC in Figure 4e,f depict the presence of a porous microstructure obtained during the acid etching process. The amorphous carbon is observed to provide a conductive coating network for the inner pore spaces. A flaky inner appearance is also observed in Figure 4f. This enhanced microstructure provides more active reaction sites for the composite and promotes faster diffusion of Li-ions [33]. The voids within the silicon structure absorb silicon expansion, provide a robust conductive path for Li⁺ ions, and improve the overall performance of porous silicon [34]. More SEM images of composites with different magnifications are presented in Figure S1. While more characterizations with energy dispersive mapping (EDS) can provide a clearer picture of the chemical composition, many other works did these characterizations already in similar system with persuasive conclusions. For example, H-J. Sohn et al. [35] also showed that TiFeSi₂ matrix contained all starting elemental metals and EDS scan profiles showed that the TiFeSi₂ matrix was robust to buffer the expansion of silicon as similarly seen in outstanding cycling performance of our work. S-S. Kim et al. [26] also reported homogeneous distribution of Si, Ti, and Fe elements by EDS for silicon-silicide, Si-TiFeSi₂ nanocomposite prepared using melt-spinning. H. Kim et al. [36] showed that microstructural changes in Si/TiFeSi₂ nanocomposite were successfully achieved as observed by EDS mapping analysis. Their results show that milling for 48 h as in this work gave a more uniform dispersion of silicon in the transition metal matrix.

3.3. Performance Evaluation of Initial Cycles

The initial two discharge-charge curves of the composites mTFSC, nTFSC, and pTFSC are shown in Figure 5a–c, respectively. The first cycle results of the three composites are plotted in Figure 5d. The voltage hysteresis plots shown in Figure 5a–c are the characteristic cycling curves for silicon-based anodes. Considering the first discharge of all composites, the composite mTFSC has a very steep-sloped discharge curve. Lower steepness is observed in nTFSC and least in pTFSC. Comparing the 3 curves as shown in Figure 5d shows that the crystallinity of silicon in the composite decreases as the steepness of the curves increases [37]. This can be correlated with the XRD results and further confirms the predominant silicon phase in each composite. The first cycle capacities of the composites mTFSC, nTFSC, and pTFSC electrodes were 1288.54 mAh.g⁻¹, 1395.56 mAh.g⁻¹, and 1386.9 mAh.g⁻¹ during discharge and 1016.32 mAh.g⁻¹, 997.88 mAh.g⁻¹, and 995.66 mAh.g⁻¹ during charge, respectively, as shown in Figure 5a–c. The high first discharge capacity can be attributed to the activation cycle with a cut-off potential of 0.05 V at C/20. Successive charge/discharge tests were carried out at a voltage range of 0.05–1.2 V. The reaction plateaus observed around 0.15 V during discharge represent the lithiation process of crystalline Si to form Li_xSi alloy. A corresponding delithiation plateau is seen at ~0.5 V during charging to form amorphous Si [38].

The discharge cycles for all composites during the second cycle decreased with a capacity loss of 213.29 mAh.g⁻¹, 319.31 mAh.g⁻¹, and 309.84 mAh.g⁻¹, respectively. The charge capacity in the second cycle decreased by 29.3 mAh.g⁻¹, 38.32 mAh.g⁻¹, and 39.41 mAh.g⁻¹ for mTFSC, nTFSC, and pTFSC, respectively. The loss of capacity in the first cycle can be attributed to the trapping of lithium ions and also to the SEI passivation layer formation [39]. The discharge cycles for all composites during the second cycle all decreased because of SEI formation [40]. The charge capacity increase in the second cycle is attributed to the stability of the semiconductive passive film, which successfully reduced further decomposition of the solvent [41]. SEI films that possess such intrinsic properties show improved electrode performance [42]. The first charge and discharge, as shown in Figure 5d for all composites, confirm that nano- and porous silicon can deliver higher capacity than micro silicon.

3.4. Rate Capability Performance

The electrochemical rate capability tests for mTFSC, nTFSC, and pTFSC at current rates of C/20, C/10, C/5, C/3, C/1, and C/10 were performed. The rate capability plots as presented in Figure 6, are similar to those previously reported [43]. The curves were normalized to the second discharge capacity, considering that the first discharge is the electrode activation step. Activation discharge/charge cycles are usually performed for silicon to improve its overall electrochemical performance [44]. This process begins with the lithiation of silicon to its maximum capacity. The SEI is usually formed in this step. The delithiation and lithiation cycles help to stabilize the structure of silicon. This process minimizes the volume expansion of silicon in subsequent discharge and charge cycles [45].

To understand the initial cycle life of the composites, the coulombic efficiencies were evaluated based on Equation (3).

$$Columbic efficiency={\displaystyle \frac{Discharge capacity}{Charge capacity}} \times 100\%$$

(3)

The coulombic efficiencies of the composites after the first two cycles were 99.56%, 99.33%, and 99.8% for mTFSC, nTFSC, and pTFSC, respectively. The composites all begin with a similar capacity at a rate of C/20, and a gradual decrease is observed at C/10. The composite pTFSC demonstrates superior electrochemical performance, followed by nTFSC, with mTFSC at least at C/20 and C/10. The pTFSC exhibits the best electrochemical performance with capacities of 1150.82 mAh.g⁻¹ and 1075.31 mAh.g⁻¹ at the rates C/20 and C/10, respectively, compared to mTFSC (1059.45 mAh.g⁻¹ and 1039 mAh.g⁻¹) and nTFSC (1090.77 mAh.g⁻¹ and 1029.65 mAh.g⁻¹). This tendency is observed to be reversed at successive rates of C/5, C/3, and C/1 with notable differences observed in the charge/discharge capacities of the composites. pTFSC yields the poorest performance at rates C/5, C/3, and C/1. A similar trend is observed in the fabrication of porous silicon-carbon composites through simple mixing, calcination, and etching with NaOH [46]. Therefore, the lower electrochemical performance of pTFSC at faster current rates in this study can be attributed to the etching parameters. Another possible reason for the lower electrochemical performance in nTFSC and pTFSC could be that the same milling parameters were used for the three composites, which are optimal synthesis conditions for microsilicon, as reported earlier [15]. The milling energy and time used for micro silicon is presumed to be excessive for nano- and porous-sized silicon. However, the reversible cycle at C/10 confirms that the porous morphology of silicon can accommodate the expansion of silicon during cycling [47]. The capacities for the reversible cycles were 1137.69 mAh.g⁻¹, 1045.43 mAh.g⁻¹, and 990.45 mAh.g⁻¹ for pTFSC, nTFSC, and mTFSC, respectively. The composite pTFSC exhibits significantly improved electrochemical performance compared to nTFSC and mTFSC. The excellent cycling and rate performance of pTFSC can be attributed to the higher purity of silicon after acid etching of AlSi followed by decreased particle sizes achieved through ball milling. The void spaces created and the presence of carbon increase the electronic conductivity between silicon particles and lower the overall resistance/diffusion distance of Li⁺ ions [48]. The lattice parameter determines the microstrain of materials and leads to an increase in defects such as dislocations and vacancies. The presence of these defects enhances the electrochemical reactivity in materials. The compressive strains in mTFSC and nTFSC, as observed from the Williamson–Hall plot, depict contraction of lattices [49], which has the tendency to influence lithium-ion diffusion. This can be correlated to the comparatively higher electrochemical performance of mTFSC and pTFSC at the initial cycles and fast C-rates in Figure 6 since both composites have more active reaction sites. Conversely, nTFSC exhibits tensile strain which tends to expand the lattice and potentially lower lithium-ion diffusion. This can be attributed to the low electrochemical performance of nTFSC at fast C-rates in Figure 6 and during the initial cycles, as shown in Figure 7a. Microstrain detected in all composites influences their surface structure, reactivity, and SEI formation. Therefore, mTFSC and nTFSC, which undergo compressive microstrain and have higher surface reactivity, facilitate the formation of a stable and uniform SEI layer due to the increased availability of active sites during the initial discharge/charge cycles. The expanded lattice in pTFSC due to tensile microstrain suggests fewer reaction sites and non-homogenous passivation on the electrode surface, leading to the formation of an unstable SEI [50].

3.5. Cyclability Performance Tests

The electrochemical performance of the composites mTFSC, nTFSC, and pTFSC is presented in Figure 7a,b. The tests were conducted at C/1 for 200 cycles and at C/5 for 100 cycles, respectively.

A relatively lower initial cycling performance of ~502.03 mAh.g⁻¹ is observed in pTFSC for the first 37 cycles, as shown in Figure 7a. This could be due to the instability of the SEI during the early cycles [51]. A progressive increase to 646.97 mAh.g⁻¹ is noticed until the 103rd cycle. The capacity stabilized and maintained a final capacity of 576.98 mAh.g⁻¹ and retention of 87.01% for 200 cycles, comparably higher than mTFSC and nTFSC. The mTFSC showed the highest initial cycling performance, starting with a specific capacity of 620.81 mAh.g⁻¹ and increasing to 635.44 mAh.g⁻¹ for the first 92 cycles. Capacity fade is observed beginning from the 93rd cycle, with a decrease in capacity from 630.76 mAh.g⁻¹ to 531.98 mAh.g⁻¹ after 200 charge/discharge cycles, and 84.44% retention at C/1 is observed.

The electrochemical performance test results conducted at C/5 are presented in Figure 7b. All composites follow a similar trend as observed at the C/1 rate. A lower initial starting capacity is observed in pTFSC with a steady increase for 49 cycles to a capacity of 968.52 mAh.g⁻¹. This is followed by a stable and higher capacity of 969.24 mAh.g⁻¹ until the 100th cycle. The superior electrochemical performance of pTFSC compared to mTFSC and nTFSC is due to the porous structure of the silicon particles. The higher capacity exhibited at C/5 is due to deeper diffusion of Li⁺ into the core structure of silicon during the alloying and de-alloying process [52]. The mTFSC and nTFSC exhibited initially stable capacity followed by steady fading in capacity at C/1 and C/5. This is because there is more silicon expansion and agglomeration in micro- and nano-sized silicon compared to porous silicon [53]. The excellent electrochemical performance of pTFSC can be attributed to the hollow inner and outer surfaces of the submicron porous structure of silicon. The acid etching process successfully provided sufficient space to contain the expansion of silicon [53]. Moreover, the mesh-like inner configuration of porous silicon provides a shorter path for the diffusion of Li⁺, as seen in the optimal rate capability of composite pTFSC [54]. The carbon added to this composite also coats the surface of porous silicon particles for a more stabilized microstructural configuration [55]. Though pTFSC shows higher electrochemical performance at C/20 and C/10, the comparably lower electrochemical performance of pTFSC at faster current rates can be attributed to its larger crystallite sizes. This can directly be correlated to the use of the same milling conditions for all composites. Another important process parameter that could be responsible for lower electrochemical performance at faster C-rates is adjusting the concentration of the acid used for etching to obtain porous morphologies of silicon that will have fewer sites for energy dissipation during milling. A gradual stabilization of the electrode with a corresponding increase in capacity is observed in Figure 7 after 90 cycles and 50 cycles at C/1 and C/5, respectively. Though mTFSC, due to compressive strains in the lattice, shows higher initial electrochemical performance compared to nTFSC and pTFSC, the larger particle sizes of silicon, as observed in Figure 4a,b, led to rapid volumetric expansion during lithiation/delithiation processes. This expansion results in the degradation of the silicon particles and delamination from the current collector, leading to poor cyclability. The smaller particle sizes of nTFSC, calculated in Figure 3 and observed in Figure 4c,d, reduce the overall strain on the composite from volume expansion of silicon, thereby mitigating cracking. The diffusion path for lithium ions is also shortened due to reduced particle sizes. However, the higher surface area of nTFSC allows for more contact of the electrode with the electrolyte, which can increase SEI formation, further electrolyte decomposition after many cycles, and low electrochemical performance. The initial low electrochemical performance observed in pTFSC is attributed to its comparatively higher microstrain and crystallite sizes. This is due to the porous microstructure of silicon (Figure 4e,f), which provides void spaces that can accommodate the volume changes of silicon during cycling, thereby reducing mechanical strain on the electrode.

4. Conclusions

In summary, three samples of TiFeSi₂/C composites were synthesized using micro, nano, and porous sources of silicon by ball milling. The silicon used in the material containing porous silicon was prepared through acid etching of Al-Si in 1.0 M of HCl (ACS 37%). The physical measurement tests were carried out using XRD and SEM. No Al peak was detected in the XRD results of pTFSC showing that the etching process was effective. The SEM images reveal that the desired silicon morphologies of micro-, nano-, and porous silicon were successfully obtained.

The improved cycling performance of silicon in the three composites can be attributed to the presence of inactive matrices formed by Ti and Fe, as well as the conductive network provided by carbon. The high reversible capacities of 990.45 mAh.g⁻¹, 1045.43 mAh.g⁻¹, and 1137.69 mAh.g⁻¹ were observed at C/10 for mTFSC, nTFSC, and pTFSC, respectively, confirming that the hollow structures seen in the SEM images successfully accommodate silicon expansion during lithiation/delithiation. The easy synthesis method used successfully yielded excellent initial coulombic efficiency of pTFSC (99.8%). Additionally, the excellent performance of the composites at C/1 and C/5 can be attributed to the inactive transition metal matrix formed by titanium and iron, which also provides a buffer for the expansion of silicon. It can also be observed that the particles in mTFSC are larger in size compared to nTFSC and pTFSC because the same milling time and energy were used to prepare all the composites. mTFSC therefore has larger particle sizes since it requires more energy to break down particles into smaller crystallites. nTFSC, which contains nano-sized silicon particles, requires less milling energy to break down particles and showed the smallest crystallite size. pTFSC, prepared with porous silicon, has voids and a comparably higher surface area. The pores contained in pTFSC act as sites for energy dissipation. Instead of the silicon being broken into smaller crystallites, the energy from mechanical milling is possibly absorbed within the porous structure, leading to less particle fragmentation and higher crystallite sizes. The simple and cheap two-step synthesis route used in this work suggests that similar high-energy composites with rate performance and improved silicon cyclability can be reproduced using this method.