Tailored Silicon Nanostructures in Conductive-based Hydrogel Binders: Impact of Size, Structure, and Surface Chemistry on Li-Ion Battery Performance

Tailored silicon nanostructures in hydrogel-derived conductive binders: Role of size, structure, and surface chemistry in enhancing Li-ion battery performance Gabriela Soukupov´ a a , Filip Matˇ ejka a,b , Zuzana Vlˇ ckov´ a ˇ Zivcov´ a c , Abdelghani Laachachi d , Pavel Gal´ aˇ r b , Miloslav Lhotka e , Otakar Frank c , Jiˇ rí ˇ Cervenka b , Fatima Hassouna a,* a Faculty of Chemical Engineering, University of Chemistry and Technology, Prague, Technick´ a 5, 166 28, Prague 6, Czech Republic b FZU - Institute of Physics of the Czech Academy of Sciences, Cukrovarnick´ a 10/112, 162 00, Prague 6, Czech Republic c J. Heyrovsky Institute of Physical Chemistry, Czech Academy of Sciences, Dolejskova 2155/3, 182 00, Prague 8, Czech Republic d Luxembourg Institute of Science and Technology, 5, rue Bommel, L-4940, Hautcharage, Luxembourg e Faculty of Chemical Technology, University of Chemistry and Technology, Prague, Technick´ a 5, 166 28, Prague 6, Czech Republic HIGHLIGHTS GRAPHICAL ABSTRACT •Larger Si nanocrystals (100 nm) achieve higher initial capacity in LIB anodes. •Smaller Si nanocrystals (6 nm) improve cycling stability. •Amorphous Si nanoparticles outperform crystalline Si in cycling stability. •3D crosslinked PPy accommodates Si volume expansion, enhancing anode performance. •Optimal Si size (6–20 nm) in PPy networks improves electrochemical performance. ARTICLE INFO Keywords: Li-ion battery Si nanoparticles Electrically conductive polymer 3D network Electrochemical properties ABSTRACT Silicon (Si) is a promising anode material for Li-ion batteries (LIBs), but its practical application is limited by volume expansion during lithiation/delithiation, leading to poor cycling stability. While Si nanostructuring mitigates this issue, it remains only a partial solution. This study systematically investigates the effects of Si particle size (6, 20, 55, or 100 nm), surface chemistry (type and degree of oxidation), and solid-state properties (amorphous vs. crystalline) on the electrochemical performance of Si-based anodes using a three-dimensional (3D) crosslinked polypyrrole (PPy) binder. In situ PPy polymerization around Si nanoparticles forms a 3D interconnected conductive network within the PPy/Si anodes, effectively accommodating volume changes and maintaining electrical contact during the galvanostatic cycling. The particle size dependence shows that larger Si nanoparticles provide higher initial charge capacity (2975 mAh/g), whereas smaller ones improve cycling stability (85 % capacity retention after 100 cycles). * Corresponding author. E-mail address: [email protected] (F. Hassouna). Contents lists available at ScienceDirect Journal of Power Sources journal homepage: www.elsevier.com/locate/jpowsour https://doi.org/10.1016/j.jpowsour.2025.238620 Received 4 March 2025; Received in revised form 3 October 2025; Accepted 13 October 2025 Journal of Power Sources 661 (2026) 238620 Available online 23 October 2025 0378-7753/© 2025 The Authors. Published by Elsevier B.V. This is an open access article under the CC BY license ( http://creativecommons.org/licenses/by/4.0/ ). Amorphous Si exhibits significantly lower specific capacity but superior capacity retention (~100 % after 100 cycles) compared to crystalline Si. Cyclic voltammetry and electrochemical impedance spectroscopy demonstrate that integrating 6 or 20 nm Si nanocrystals into a 3D crosslinked PPy enhances anode performance. These findings highlight the importance of optimizing Si properties in designing conductive hydrogel-derived anodes for high-performance LIBs. 1. Introduction In recent years, the demand for cost-effective, high-performance rechargeable batteries with extended cycle life has grown significantly, driven by their applications in electronic devices, electric vehicles, and large-scale energy storage systems. Among all, Li-ion batteries (LIB) have attracted significant attention due to their exceptional energy density, high power density, and long cycle life [1–4]. The commercially used graphitic anode offers only a limited theoretical specific capacity (~370 mAh/g) [2,5–7], which has prompted extensive research into alternative anode materials. Si is considered one of the most promising anode materials for LIB due to its high theoretical specific lithiation capacity (~3579 mAh/g for Li 15 Si 4 ), low discharge potential (~0.4 V vs Li/Li + ), environmental sustainability, and high elemental abundance [1, 8–10]. However, the practical application of Si as an anode material in LIB is significantly limited due to its poor electrical conductivity and substantial volume expansion (~300 % for Li 15 Si 4 [11]) during lithiation. This expansion leads to severe issues such as material cracking, anode pulverization, overgrowth of a solid electrolyte interphase (SEI), and capacity fading [7–9,12,13]. To address these challenges, several strategies have been explored. One promising approach is the nanostructuring of Si (e.g., nanowires or nanoparticles), which helps accommodate the volume changes, thereby enhancing the mechanical stability of the material and improving its cycling performance [14–18]. It has been reported that when the particle size is reduced below a critical threshold of 150 nm, the cracking of the material is significantly reduced [19–22]. Although nanostructuring improves the mechanical properties of Si anodes, several significant challenges remain. These include the higher surface activity of Si, which can lead to particle agglomeration, the formation of a thicker SEI layer, and the high cost associated with the preparation process [17,19, 22–24]. To overcome the poor electrical conductivity of nanostructured Si and facilitate more efficient Li-ion diffusion, conductive carbon (C) additives such as carbon black (Super P), carbon nanotubes, and graphene are commonly incorporated. The addition of these C materials not only helps to buffer the volume expansion during lithiation but also improves the overall conductivity, thereby providing an effective electronic pathway for electron transfer [8,17,23]. Generally, polymer binders such as carboxymethyl cellulose, polyamide imide, or poly (acrylic acid) (PAA) are incorporated with Si/C to enhance mechanical stability and hence, the cycling performance of the anode material [7,8, 10,25]. Effective binders for Si-based anodes must meet several critical requirements. They must be capable of forming uniform mixtures with Si and C, establishing strong binding interactions with both Si and C to ensure stable electrical pathways to the current collector, and demonstrating electrochemical stability within the operational potential window. Moreover, they should exhibit high elastic modulus values to accommodate the significant volume expansion of Si, resist excessive swelling in the electrolyte, facilitate the formation of a stable and thin SEI layer, and ensure a robust connection between the anode material and the current collector throughout cycling. In addition to these technical characteristics, an ideal binder should also be economically feasible, straightforward to manufacture in large quantities, affordable, and compatible with existing production techniques [25,26]. The use of the aforementioned insulating polymer binders may result in a weak interface between Si and C, leading to contact loss during cycling [25, 27]. In this context, replacing the insulating conventional binders with electrically conductive polymers presents a promising alternative, as it would eliminate the Si/C interface issues. The conductive polymer would not only serve as a binder but also enhance the conductivity of the anode material, thereby improving overall electrochemical performance. This approach was clearly demonstrated by Li et al. [1], who compared the electrochemical performance of Si-based anodes employing either an insulating binder (PAA) or a conductive polymer binder. Their study revealed that conductive polymer binders significantly enhance cycling stability and rate capability compared to their insulating counterparts. Conductive polymers, such as polyaniline (PANI), polypyrrole (PPy), and poly(3,4-ethylenedioxythiophene) (PEDOT), offer several advantages, including mechanical flexibility, cost-effectiveness, ease of synthesis and processing, environmental friendliness, high electrical conductivity, and intrinsic electrochemical activity [28–30]. Each polymer, however, presents unique strengths and limitations. For instance, PPy binders exhibit enhanced Li-ion diffusion but suffer from continuous degradation; PANI binders provide good adhesion but are prone to volume expansion during cycling; PEDOT binders offer chemical stability and self-healing properties but require optimization to balance conductivity and mechanical flexibility [31–33]. Although conductive polymers have received significant interest for their application in LIB, relatively limited research has been conducted on integrating them with Si nanostructures to develop commercially viable anodes for LIB [1,2,5]. A viable and promising strategy for integrating Si nanostructures into conductive polymer binders involves a hydrogel-based preparation technique. In this approach, in situ polymerization and crosslinking of the conductive polymer take place in the presence of Si, often accompanied by an additional electrically conductive additive. The process results in the formation of a three-dimensional (3D) network, wherein Si is uniformly coated and interconnected by the conductive polymer matrix [1,2,5,9,14,25,34, 35]. This 3D interconnected network improves electrochemical performance through several beneficial features. The 3D framework functions as both a conductivity enhancer and a conductive binder, improving the particle-to-particle contact. Its porous structure helps buffer the volume changes of Si during lithiation [35], while the 3D conductive network facilitates improved electron and ion diffusion [14]. Furthermore, embedding Si within the framework promotes the formation of a more stable SEI during lithiation by better isolating Si from the electrolyte [34]. Despite their importance, the impact of key factors such as the size, surface chemistry, and solid-state properties of Si on the electrochemical performance of resulting anodes remains poorly understood and has not been systematically investigated in Si/conductive polymer-based anodes, including those prepared using hydrogel-based approaches. Previous studies investigating the size effect of Si nanoparticles in anode materials have primarily focused on physical blending methods, where Si particles or carbonized core-shell Si particles (ranging in size from 30 nm to 5 μ m) were mechanically mixed with a binder and carbon additive [20,24,36–38], or prepared via gas deposition techniques [39]. The studies on Si-based anodes using 3D crosslinked conductive polymeric binders have used different types and sizes of Si nanoparticles, including nanoparticles with average diameters ranging from 40 to 300 nm [1,5,9,14,25,35,40,41], micro-sized porous Si particles [42], or Si dendrites [34]. However, the selection of Si type and size has often been arbitrary, lacking a systematic methodology. Furthermore, these studies have predominantly used commercial Si materials with poorly controlled surface chemistry. This lack of systematic investigation also applies to the emerging class of Si quantum dots (SiQD), which, despite G. Soukupov´ a et al. Journal of Power Sources 661 (2026) 238620 2 offering advantages, such as high surface area, shorter diffusion pathways, and reduced volume expansion, have received only limited attention in the context of LIB anodes [43–50]. Practical implementation of SiQD remains challenging due to difficulties in synthesizing monodisperse particles and mitigating undesirable side reactions arising from their high surface reactivity due to their significantly curved surface and presence of hydrogen surface termination. To the best of our knowledge, no comprehensive study has systematically examined the combined influence of Si particle size, surface chemistry, and intrinsic solid-state properties on the electrochemical performance of Si-based anodes incorporating 3D crosslinked conductive polymeric binders in LIB. The present study aims to address this important knowledge gap by providing a fundamental understanding of structure-property relationships that govern the behavior of such advanced and complex anode systems. This study systematically investigates the effects of size (6, 20, 55, and 100 nm), surface chemistry (oxide type and degree) and solid-state properties (amorphous versus crystalline) using both commercial and laboratory-synthesized Si nanoparticles, including highly monodisperse 6 and 20 nm particles, on the structural, morphological, and electrochemical performance of the resulting anodes. The anode materials were prepared via the in situ polymerization of a conductive PPy hydrogel, serving as a model crosslinked conductive polymer binder, using an environmentally friendly, waterborne method performed at near-zero temperatures. Phytic acid (PhA), a naturally occurring molecule, was employed as a crosslinking agent for the PPy chains, while PAA served as a stabilizing agent for the Si nanoparticles. The resulting 3D crosslinked PPy network functioned as a high-performance conductive binder, conformally coating the Si nanoparticles. This integration of the 3D crosslinked PPy network with the Si nanoparticles was achieved at low temperatures, eliminating the need for high-temperature carbonization, which is commonly employed to create stable conductive matrices. A clear relationship was established between the Si nanoparticle size, surface chemistry, solid-state properties, and their collective impact on the electrochemical performance of the resulting anodes. These insights provide a foundation for the rational design and optimization of Si-based anodes for enhanced electrochemical performance. 2. Experimental part 2.1. Materials Two types of Si nanocrystals (SiNC 100 with 100 nm average diameter, stock keeping unit NG04CO28095; SiNC 55 with 55 nm average diameter, stock keeping unit NG04EO1804) were purchased from Nanografi. Two other types of SiNC (SiNC 6 with 6 nm average diameter; SiNC 20 with 20 nm average diameter) and one type of amorphous Si nanoparticles (SiNA 20 with 20 nm average diameter) were synthesized in the frame of this study. Diluted silane (1 % in argon, Linde, Ar 5.0, SiH 4 5.0, UN1954), hydrogen (H 2 7.0, Linde, UN1049), argon (Ar, Ar 6.0, UN1006), and pure silane (SiH 4 , UN2203) were used for the Si synthesis. Conductive carbon filler Super P (40 nm average diameter) was kindly donated by Imerys S.A. Pyrrole monomer (reagent grade, 98 %, M w =67.09), phytic acid solution (PhA, 50 % (w/w) in H 2 O, M w = 660.04), poly(acrylic acid) (PAA, M w =450,000), ammonium persulfate (APS, ≥98.0 %, M w =228.20), lithium hexafluorophosphate solution in ethylene carbonate and dimethyl carbonate (1.0 M LiPF 6 in EC/DMC = 50/50 (v/v), battery grade), and Li-metal foil (thickness: 0.6 mm, 99.9 %) were purchased from Merck. Current collector copper foil (thickness: 25 μ m, 99.8 %) was purchased from Thermo Fisher. Deionized (DI) water was used as an aqueous medium in all the experiments. 2.2. Materials preparation 2.2.1. Synthesis of SiNC SiNC 20 , SiNC 6 , and SiNA 20 were synthesized using a non-commercial non-thermal plasma flow-through reactor operated under low pressure. The system was created by adapting the apparatus of Korsthagen et al. [51]. The operational pressures were within 10 Pa in the standby mode and lower than 500 Pa in the synthesis mode. The plasma was generated with a radiofrequency (RF) power source (RFG 600W, Coaxial Power Systems) operating at 13.56 MHz equipped with automatic match box (Coaxial Power Systems). For the synthesis of 6 mm crystalline particles (SiNC 6 ), a narrow glass tube reactor was used, with an internal diameter of 0.8 cm. The plasma discharge was generated using planar electrodes (dimensions 5 cm–12 cm). The output RF power was 150 W, and the SiNCs were synthesized using a 1 % diluted silane in argon mixed with hydrogen. For the synthesis of the 20 nm crystalline particles (SiNC 20 ), a glass reactor with an internal diameter of 2.1 cm, equipped with ring electrodes (height 2.7 cm), was used. The reaction mixture was composed of argon and pure silane, and the output RF power was set to 107 W. The 20 nm amorphous particles (SiNA 20 ) were synthesized using a two-stage reactor. The first stage consisted of a narrow glass tube reactor with an internal diameter of 0.8 cm and a length of 50 cm. Planar electrodes (measuring 5 cm by 12 cm) were attached to the reactor. A flow of diluted silane was introduced into this stage, and the RF output power was set to 150 W. The second stage was implemented using a broad glass tube with an internal diameter of 2.1 cm and a length of 50 cm. Double-helix electrodes were attached to the tube, with one electrode grounded, and the wires spaced 1 cm apart. During this stage, flows of pure silane and argon were introduced, and the second RF power output was set to 250 W. 2.2.2. Synthesis of PPy hydrogel-based Si composites and anode preparation PPy hydrogel-based composites were prepared via in situ oxidative polymerization of pyrrole monomer using APS in the presence of Si nanoparticles (SiNC or SiNA). The composites were prepared according to the following procedure. Firstly, 11 μ l of pyrrole monomer was transferred to a small glass vial with 35.4 μ l of PhA and DI water. The solution was mixed using a magnetic stirrer. Next, 38.4 mg of Si nanoparticles were ground in a mortar for 10 min. Subsequently, 15.9 mg of Super P was added and mixed with the Si nanoparticles for an additional 10 min. Then, 15.4 mg of PAA was incorporated into the Si and Super P mixture and blended for another 10 min. The resulting powder mixture was transferred into a small glass vial, and DI water was added. The powder mixture was dispersed in DI water using a sonication bath (PS 3000, PowerSonic) and further mixed with a magnetic stirrer to achieve a homogeneous dispersion. Once the powders were fully dispersed and a uniform ink was obtained, the solution of pyrrole with PhA and DI water was added to the ink. This mixture was stirred for 10 min with a magnetic stirrer, and the vial was labeled as solution A. In parallel, an APS solution was prepared by dissolving 11 mg of APS in DI water. Both solutions, solution A and APS solution, were cooled separately in an ice bath for 10 min. Finally, the APS solution was poured into solution An under continuous stirring in the ice bath, initiating the in situ polymerization of the pyrrole monomer. After 4 h, the resulting ink was cast onto copper foil and allowed to dry at room temperature. The obtained thin films were then pressed for 1 min at 60 kPa and dried in a vacuum oven at 60 ◦C overnight. The prepared anodes were labeled as PPy/Si, specifying the type of Si used, i.e., SiNC or SiNA. 2.3. Characterization methods The chemical structure of all Si nanoparticles was analyzed by Raman spectroscopy (633 nm, 1.96 eV, objective 100x, He-Ne laser; Horiba LabRAM HR spectrometer integrated with an Olympus microscope) and Fourier-transform infrared spectroscopy (FTIR, Nicolet iS50 ABX). The surface chemistry was examined using X-ray photoelectron spectroscopy (XPS, X-ray beam Al K α with E =1486.6 eV and power of 60 W, photoelectron emission take-off angle of 0◦; Thermofisher Nexsa G2). The obtained XPS spectra were deconvoluted using the CasaXPS G. Soukupov´ a et al. Journal of Power Sources 661 (2026) 238620 3 software. The micro/nanostructure was visualized using scanning electron microscopy (SEM, Mira3 LMH, Tescan, secondary electrons at 3 kV) and high-resolution transmission electron microscopy (TEM, EFTEM Jeol 2200 FS, copper mesh). The image analysis was performed using Image J software. Energy-dispersive X-ray spectroscopy (EDS) analysis was conducted on the composites using SEM Mira 3 LMH (Tescan) and Bruker XFlash 6|10 detector with software Esprit 2.1. The crystallinity of the Si nanoparticles was characterized using X-ray diffraction (XRD, Cu lamp; PANalytical X’Pert PRO with PIXcel1D_1D detector). The specific surface area was analyzed using the nitrogen physisorption technique on a 3Flex analyzer (Micromeritics, Norcross). To assemble Li-ion half-coin cell batteries, the anode materials prepared in this work were dried at 80 ◦C in a vacuum oven for 12 h and cut into 1.5 cm diameter circles. Each cut circle was weighed and transferred to an Ar-filled glovebox for the battery assembly. Li metal foil was used as the reference and counter electrode, and 1.0 M LiPF 6 in EC/DMC served as the electrolyte. The electrochemical performance of the anodes was evaluated through galvanostatic charge/discharge (GCD) cycling within the potential range of 0.05–1 V versus Li/Li + . This was performed using a battery tester (Neware BTS-4008-5V50mA) at a charging rate of 0.1C (0.4 A/g), calculated for each sample individually based on the mass of Si. Further characterization involved cyclic voltammetry (CV) using a potentiostat ( μ Autolab, Metrohm), and electrochemical impedance spectroscopy (EIS) in the discharged state (0.05 V), with frequencies ranging from 80 kHz to 0.03 Hz (Ivium CompactStat, Ivium Technologird B. V.). 3. Results and discussion The 3D crosslinked PPy-based anodes were synthesized through a straightforward in situ oxidative polymerization process in an aqueous medium (Fig. 1). This method facilitates the formation of a 3D interconnected network, wherein the conductive polymer embeds, connects, and stabilizes Si nanoparticles and Super P. PPy was selected due to its excellent electrical conductivity, mechanical flexibility, and its ability to form a conductive, 3D crosslinked network when combined with a crosslinking agent such as PhA. Different types of Si nanoparticles with varying sizes and physicochemical properties were integrated within the anodes to examine the effect of Si particle size, surface chemistry, and intrinsic solid-state properties on the electrochemical performance of Sibased anodes with conductive polymeric binder in LIB. 3.1. Chemical and physical properties of studied Si nanoparticles and micro-structure of anodes The chemical and physical properties of all Si nanoparticles were analyzed and compared. The XRD diffractograms, shown in Fig. 2a, revealed the crystalline nature of all SiNC (SiNC 6 , SiNC 20 , SiNC 55 , and SiNC 100 ) and the amorphous state of SiNA 20 . Notably, a slight shift in all peaks toward higher 2θ values is observed in the diffractogram of SiNC 55 compared to other SiNC, indicating a structural difference [52]. In addition, the SiNC 55 diffractogram exhibits a small peak at 22.34◦, which is attributed to the presence of SiO 2 [53–55]. Variations in peak shape and sharpness across the SiNC are associated with the size of the nanocrystals, with peak broadening indicative of smaller nanocrystals, specifically in the case of SiNC 6 [56]. In contrast to SiNC, the diffractogram of SiNA 20 displays two broad bands, confirming its amorphous nature. A similar pattern has been previously reported in studies of amorphous Si and SiO 2 [57]. Fig. 2b depicts the Raman spectra of the SiNC and SiNA. The Raman spectra of SiNC exhibit the characteristic sharp crystalline Si peak in the range of 508–516 cm −1 , i.e., 516 cm −1 for SiNC 6 and SiNC 20 , 508 cm −1 for SiNC 55 , and 511 cm −1 for SiNC 100 . These peaks are clearly distinguishable from the broad band of SiNA 20 centered at 498 cm −1 , which is characteristic of amorphous Si [58,59]. The Raman spectra further corroborate the findings of the XRD analysis. The FTIR spectra of the SiNC and SiNA are presented in Fig. 2c. All spectra exhibit absorption bands characteristic of Si. The bands observed at 3355 cm −1 and 1625 cm −1 are ascribed to Si-OH stretching and bending vibrations, respectively [60–62], indicating the presence of surface oxidation in all the Si nanoparticles. A shoulder at 2246 and a band at 975 cm −1 are associated with the Si-H vibrations in O y -Si-H x [63]. The bands in the ranges of 2000–2200, and 600–700 cm −1 are assigned to Si-H vibrations, specifically, Si-H stretching at 2076 cm −1 , Si-H 2 stretching at 2103 cm −1 , and Si-H 3 stretching at 2134 cm −1 , and Si-H bending and/or Si-H 2 wagging at 655 cm −1 [64]. Bands at 902 and 851 cm −1 can be ascribed to Si-H 3 deformation modes or to Si-H 2 vibrations [64]. Notably, the band at 1042 cm −1 , corresponding to Si-O-Si stretching, further confirms the surface oxidation of Si nanoparticles [61,63]. Si-O-Si stretching is also represented by another shoulder at 794 cm −1 [60,65], supporting the observations. The spectrum of SiNC 55 displays an additional band at 1230 cm −1 , attributed to Si-O-Si stretching, confirming the presence of SiO 2 impurities [66,67]. This finding aligns with the XRD results, where only SiNC 55 exhibited a SiO 2 peak, suggesting a higher degree of oxidation. Interestingly, the spectrum of SiNC 100 also reveals a shifted band at 1206 cm −1 , which may indicate either the presence of C [68] or SiO 2 impurities. The XPS spectra further confirm surface oxidation across all Si Fig. 1. Schematic illustration of the preparation of 3D crosslinked PPy-based anodes. G. Soukupov´ a et al. Journal of Power Sources 661 (2026) 238620 4 nanoparticles, which is consistent with the findings from the FTIR analysis. Table S1 summarizes the elemental atomic composition of the surface of each sample, revealing peaks at binding energies of approximately 532, 285, 154, and 100 eV, corresponding to the O(1s), C(1s), Si (2s), and Si(2p) core levels. The survey spectra, combined with the deconvolution of the Si(2p) and C(1s) core-level peaks (Table S1, Fig. S1), reveal distinct differences between the commercial and labsynthesized Si nanoparticles. The commercial Si nanoparticles exhibit a higher surface O content and lower C content compared to the labsynthesized counterparts. The deconvolution of the Si core-level spectra reveals several distinct peaks: Si 2p 3/2 (99.4 eV) and Si 2p 1/2 (100.0 eV), both characteristic of elemental Si, along with peaks corresponding to oxidized species, i.e., Si 4+ (103.5 eV), Si 3+ (102.5 eV), Si 2+ (101.6 eV),and Si + (100.6 eV). The presence of these oxidized species is attributed to oxygen termination resulting from atmospheric exposure [69]. The spectra of the commercial Si nanoparticles exhibit a higher proportion of Si 4+ , originating in SiO 2 , compared to their lab-synthesized counterparts. This observation confirms differences in the surface chemistry and indicates a greater degree of oxidation, and thus a higher SiO 2 amount, in the commercial samples. This finding is further corroborated by FTIR spectra, which lack the peak at 2246 cm −1 , corresponding to Si-H stretching vibrations, in both SiNC 100 and SiNC 55 . The absence of this peak provides additional evidence for the oxidation of these Si nanoparticles. The nature of C on the surface of Si nanoparticles was analyzed through the deconvolution of the C core-level spectra, revealing five distinct peaks: C-(C, H) (285.0 eV), C-(O,N) (286.5 eV), C=O/O-C-O (287.9 eV), O=C-O (289.2 eV), and C-COO (285.4 eV). These findings closely align with previously reported XPS spectra of adventitious C [70], which rises from contamination by C compounds present in the air. The higher C content in the lab-synthesized Si nanoparticles can be ascribed to their highly curved surfaces, which are more susceptible to surface contamination. Upon exposure to air during analysis, the highly reactive Si nanoparticle surface readily adsorbs atmospheric C, leading to the formation of surface-bound C species. The exact conditions of the preparation process of the commercial Si nanoparticles is not known. However, the presence of SiO 2 suggests differences in synthesis and handling procedures, likely involving more extensive surface oxidation. The TEM and SEM images of the Si nanoparticles are presented in Fig. 3a – e and Fig. S2, respectively. The TEM images reveal a spherical morphology of all Si nanoparticles. Image analyses of the TEM data indicate that the lab-synthesized SiNC 20 , SiNA 20 , and SiNC 6 exhibit a monodisperse size distribution, with minimal standard deviations (6 ± 1 nm for SiNC 6 , and 20 ±5 nm for SiNC 20 and SiNA 20 ), confirming the uniformity of the nanoparticle formation. In contrast, the commercial SiNC 100 and SiNC 55 display a significantly broader size distribution with large standard deviations (100 ±45 nm for SiNC 100 and 55 ±30 nm for SiNC 55 ). The corresponding particle size distribution curves are presented in Fig. S3. The SEM images of the anodes are shown in Fig. 3f–i and Fig. S4. The images reveal good dispersion and embedding of Si nanoparticles and Super P within the 3D porous, foam-like network composed of interconnected dendritic nanofibers. No morphological differences were observed between PPy/SiNC 20 and PPy/SiNA 20 , as both SiNC 20 and SiNA 20 have similar sizes and surface chemistries (side-by-side comparison of SiNC 20 and SiNA 20 , along with Si-free anode is shown in Fig. S4). Notably, owing to the QD size of SiNC 6 , the SEM image of the PPy/SiNC 6 anode shows clusters of SiNC 6 and Super P embedded within PPy network (Fig. 3i), while the other anodes exhibit uniformly dispersed, well-defined, and interconnected nanofillers. Fig. 2. Physico-chemical properties of the Si nanoparticles: a) XRD diffractograms, b) Raman spectra (* the Raman spectrum of SiNC 6 was multiplied by factor 2), and c) FTIR spectra. G. Soukupov´ a et al. Journal of Power Sources 661 (2026) 238620 5 The EDX elemental maps shown in Fig. S5 reveal a homogeneous distribution of PPy throughout the anode cross-sections, as indicated by the uniform presence of nitrogen, which is specific to the PPy structure. Phosphorus is also detected due to the use of PhA as a cross-linking agent during binder synthesis, while copper originates from the Copper foil used as the current collector. These results confirm the effective and uniform incorporation of the PPy binder within the electrode structure. The EDX maps further demonstrate the homogeneous distribution of Si across all anode materials, regardless of nanoparticle size. 3.2. Performance-related characteristics of the anodes To evaluate the electrochemical performance of PPy/SiNC and PPy/ SiNA anodes in half coin-cell LIB, GCD cycling tests were conducted in the potential window of 0.05–1 V vs. Li/Li + (PPy decomposes above 1V) at a current density of 0.4 A/g. Fig. 4a shows the GCD cycling performance up to 500 cycles, while Fig. 4b provides a zoomed-in view of the first 100 cycles. A correlation between SiNC size and initial specific capacity is observed, with the capacity increasing as the average particle size increases. According to Fig. 4c and Table S2, the highest initial charge capacity of 2975 mAh/g (discharge capacity 4732 mAh/g) is displayed by the PPy/SiNC 100 , while the lowest, 1013 mAh/g (discharge capacity 3291 mAh/g), is observed for PPy/SiNC 6 . Fig. 4c also shows that SiNC with sizes below 100 nm exhibit lower initial reversible Li storage capacity. The reason can be found in their higher surface area (Table S3), which leads to a relatively more prominent irreversible SEI formation despite partial protection by PPy [19,20]. In addition, their reduced active volume can limit Li alloying capacity, and the strong Si–PPy interface can trap Li or hinder diffusion, lowering reversibility [71]. To evaluate the influence of Si solid-state properties (crystalline vs. amorphous) on electrochemical performance, the specific capacities of PPy/SiNC 20 and PPy/SiNA 20 composites, both containing Si nanoparticles of identical average size and comparable specific surface area, were compared. PPy/SiNA 20 exhibits remarkable capacity retention of 78 % after 500 cycles. This enhanced cycling stability is consistent with previous reports and is primarily attributed to the isotropic, disordered structure of amorphous Si, which accommodates volume changes during lithiation/delithiation more uniformly than crystalline Si [72]. In Fig. 3. TEM images of Si nanoparticles: a) SiNC 100 , b) SiNC 55 , c) SiNC 20 , d) SiNA 20 , and e) SiNC 6 . SEM images of the anodes: f) PPy/SiNC 100 , g) PPy/SiNC 55 , h) PPy/ SiNC 20 , and i) PPy/SiNC 6 . G. Soukupov´ a et al. Journal of Power Sources 661 (2026) 238620 6 contrast, crystalline Si undergoes sharp phase transitions (e.g., to Li 15 Si 4 ), leading to localized mechanical stress and eventual fracturing. The gradual Li + insertion in amorphous Si, without distinct phase boundaries, reduces internal stress and helps preserve structural integrity over long-term cycling. Despite its stability advantage, PPy/SiNA 20 delivers a significantly lower initial charge specific capacity (247 mAh/g) than PPy/SiNC 20 , highlighting the role of Si crystallinity in achieving higher capacities. Interestingly, with extended cycling, the capacity of PPy/SiNC 20 gradually converges toward that of PPy/SiNA 20 , likely due to the progressive amorphization of crystalline Si under repeated lithiation. In general, all investigated anodes exhibit a capacity drop during the first cycle, ascribed to the formation of the SEI layer, which results from the decomposition of the electrolyte [73,74]. Although the formation of the SEI layer leads to irreversible capacity loss, its presence is beneficial, as a stable SEI layer subdues ongoing side reactions with the electrolyte that would otherwise contribute to Li consumption [74]. The capacity loss during the initial cycles can be correlated with the size and the specific surface area of the SiNC, as summarized in Table S3. Nitrogen physisorption measurements confirm that the specific surface area Fig. 4. Electrochemical performance in potential window 0.05–1 V at current density 0.4 A/g of anodes containing SiNC or SiNA: a) GCD cycling for 500 cycles, b) zoom of the GCD cycling up to 100 cycles, c) galvanostatic voltage profiles from the 1st cycle, d) coulombic efficiency; e) capacity retention after 100 cycles and initial specific capacity (measured at 2nd cycle) of anodes containing SiNC, and f) evolution of capacity retention during the cycling for each PPy/SiNC anode material. G. Soukupov´ a et al. Journal of Power Sources 661 (2026) 238620 7 increases as the particle size decreases. For instance, PPy/SiNC 6 , which contains SiNC 6 with the smallest particle size and the highest specific surface area (509.4 m 2 /g), displays the most substantial initial capacity loss. As reported by Kong et al. [74], while Si nanostructuring reduces volume expansion, it also increases the formation of the SEI layer due to the increased specific surface area, leading to enhanced Li consumption. Similarly, the coulombic efficiency (CE) (Fig. 4d) in the first cycle is significantly below 100 % for all anodes due to the SEI layer formation. Among them, PPy/SiNC 100 exhibits the highest initial CE (63 %), while PPy/SiNC 6 shows the lowest (31 %). Nevertheless, in subsequent cycles, the CE increases to nearly 100 % for all anodes. This trend of increasing capacity loss in the first cycle with decreasing SiNC particle size aligns with the findings of Nadimpalli et al. [75], who reported that a higher specific surface area leads to greater initial capacity drop, ultimately reducing energy density in later cycles. The evolution of capacity retention during GCD cycling is shown in Fig. 4e–f and Table S2. PPy/SiNC 100 exhibits the lowest capacity retention (24 % after 100 cycles), while PPy/SiNC 6 shows the highest retention (85 % after 100 cycles). All anode materials experience capacity fade during repeated charge/discharge cycling, primarily due to the degradation of SiNC, which is linked to the volume expansion of SiNC and the formation of an unstable SEI [74]. After 500 cycles, PPy/SiNC 100 retains only 3 % of its capacity, whereas PPy/SiNC 6 maintains a capacity retention of 44 %. In order to determine whether the improvement in cycling stability is solely attributed to the decreasing size of SiNC, a PPy hydrogel-derived anode free of SiNC was prepared using the same experimental procedure. Although the measured capacity (Fig. S6) was inherently low, given that SiNC is the primary active material in this system, no capacity degradation was observed over 500 cycles. Instead, the capacity gradually increased. This emphasizes the significant role of SiNC on the overall capacity decay. To further confirm the beneficial role of the 3D interconnected PPy structure in combination with SiNC, an anode material composed of non-crosslinked PPy and SiNC 20 (a representative type of SiNC) was prepared. The GCD cycling of the resulting anode, labeled as PPy non-crosslinked /SiNC 20 (Fig. S7), reveals poor cycling stability, with the specific capacity dropping from approximately 800 mAh/g to nearly 0 mAh/g after 300 cycles. These findings highlight the importance of combining small SiNC with a 3D crosslinked PPy matrix to achieve good capacity retention. The SiNC serves as the primary active material in the anode, while the crosslinked PPy matrix buffers mechanical stress and enhances electrical conductivity. It is worth noting that the pronounced capacity fade observed in the PPy/SiNC 100 could also be partially attributed to the size heterogeneity of SiNC 100 [76,77]. The GCD profiles were measured at various current densities (Fig. S8) of 0.4, 0.7, 1.8, and 3.6 A/g and compared for two selected representative anodes, PPy/SiNC 20 and PPy/SiNC 100 . The charge capacity of PPy/SiNC 20 ranges from ~1280 to 165 mAh/g, demonstrating good stability and reversibility, especially at lower current rates. After returning to 0.4 A/g, it recovers its previous capacity values. In contrast, the charge capacity of PPy/SiNC 100 varies from 2698 to 69 mAh/g. Interestingly, this anode does not show the same level of reversibility and stability as the one with a smaller SiNC 20 . While the reversibility of PPy/SiNC 100 remains fairly high, it does not fully recover its previous capacity upon returning to 0.4 A/g. This behavior could be partially related to the broader size distribution of SiNC 100 , which may lead to uneven lithiation and non-uniform SEI formation. Overall, these observations emphasize the beneficial role of small, more uniform SiNC (20 nm) in enhancing cycling stability and capacity retention. To investigate the structural and morphological evolution of the anode materials after the GCD cycling, representative anodes were analyzed by Raman spectroscopy and SEM before and after cycling. SEM images (Fig. S9) of the anode materials on their cross-sections after the cycling show a loss of the structure and/or the buildup of the SEI; however, this might also be caused by the electrolyte remnants. The Raman spectra of PPy/SiNC 100 and PPy/SiNC 20 before and after the cycling are compared to those of PPy/SiNA 20 (Fig. S10). All anodes display the characteristic peaks of PPy (e.g., the peak at 970 cm −1 , corresponding to the in-plane ring deformation of the polaron state, and the peak at 938 cm −1 , assigned to bipolaron state) [78,79] along with the distinct Si peak. Prior to cycling, the Si peak appears at 515 cm −1 for both PPy/SiNC 100 and PPy/SiNC 20 , corresponding to the crystalline Si. However, after cycling, a noticeable downshift of the SiNC peak (~472 cm −1 ), along with an asymmetry towards lower wavenumbers, is observed, aligning with the peak shape and position of PPy/SiNA 20 . These changes confirm the amorphization of SiNC after cycling. Table S4 presents the electrochemical performance of our developed anodes containing SiNC 6 and SiNC 20 (i.e., PPy/SiNC 6 and PPy/SiNC 20 ) in comparison with the most relevant hydrogel-derived conductive polymer binder/Si anodes reported in the literature. It is worth mentioning that two key factors must be considered before making any direct comparison. First of all, the Si used in those studies varies significantly in terms of particle size and surface chemistry. In many cases, properties such as crystallinity and surface chemistry are neither analyzed nor discussed. Often, commercially available Si is selected without thorough examination and used as is, which can introduce significant variability in the results. As demonstrated in our study, any surface modification, pronounced oxidation, or alteration in the crystalline structure can have a tremendous impact on the resulting electrochemical properties. Secondly, the conditions under which the electrochemical measurements are performed, such as the potential window, current rate, electrolyte used, and the number of cycles, also vary across studies. In addition, important details, such as whether the reported initial specific capacity refers to discharge or charge capacity or whether the stated capacity retention is related to the 1st, 2nd, or later cycles, are not always clearly specified. Therefore, direct comparisons of these results should be made with caution, as they may conceal several potential pitfalls. Both anodes (PPy/SiNC 6 and PPy/SiNC 20 ) presented in this study exhibit promising characteristics and properties across various parameters. The preparation procedure was conducted in an aqueous medium using a conducting polymer as a binder, both of which are environmentally friendly. As shown in Table S4, these anodes achieved initial discharge capacities of 3291 mAh/g for PPy/SiNC 6 and 2978 mAh/g for PPy/SiNC 20 , which are among the highest values reported for 3D crosslinked conductive polymeric binder/Si anode materials. Although their initial charge capacities are relatively modest (1014 mAh/g for PPy/SiNC 6 and 1273 mAh/g for PPy/SiNC 20 ), they remain significant and are comparable to those reported in the literature [5,9, 35]. Furthermore, the obtained capacity retentions of this work are satisfactory and comparable with previously reported results (Table S4). For deeper understanding of the electrochemical performance, CV was conducted on coin half-cell batteries (Fig. 5 and Fig. S11). Measurements were performed in the potential window of 0.05–1 V at a scan rate of 0.1 mV/s. Fig. 5a–c compares 1st, 3rd, and 10th cycles of all the anodes and Fig. S11 presents CV of each anode separately. In the 1st cycle, a reduction peak appears at 0.4 V for PPy/SiNC 20 , 0.5 V for PPy/ SiNC 6 , and approximately 0.6 V for PPy/SiNC 100 and PPy/SiNC 55 . This peak disappears in subsequent cycles and corresponds to the irreversible reaction of electrolyte decomposition on the anode surface, and to the formation of the SEI layer [80–82]. In the PPy/SiNA 20 , anode, the characteristic reduction peak associated with SEI formation is not distinctly observed during the first cycle of CV. When amorphous Si is used as the anode material, the absence of a pronounced reduction peak after the first cycle, despite a significant capacity loss, could be attributed to the nature of Li insertion into amorphous Si [83–85]. Unlike crystalline Si, which undergoes a clear phase transformation during initial lithiation [86], amorphous Si accommodates Li via a gradual, solid-solution mechanism without sharp phase transitions [83–85]. This results in broad, indistinct electrochemical features rather than well-defined reduction peaks. The capacity fade observed in GCD measurements (Fig. 4b) is primarily due to irreversible SEI formation, mechanical stress, and Li becoming trapped in electrochemically inactive phases, rather than alterations in the fundamental Li mechanism. G. Soukupov´ a et al. Journal of Power Sources 661 (2026) 238620 8 Nevertheless, SEI formation is evident, as indicated by the capacity decline in GCD cycling (Fig. 4b) and the impedance increase observed in EIS (Fig. 6d). Interestingly, PPy/SiNC 55 exhibits a broad reduction peak at around 0.9 V, which is significantly more intense than the peak at 0.6 V, suggesting the occurrence of another irreversible reaction. This peak (around 0.9 V and higher) has been previously connected with the irreversible reaction between SiO 2 and Li + , resulting in the conversion of SiO 2 to Si and lithium silicates (e.g., Li 4 SiO 4 ) and lithium oxide (e.g., Li 2 O) [87,88]. This observation aligns well with the earlier analysis of SiNC 55 , which showed a higher SiO 2 content. During continued CV cycling, four peaks emerged [50]. Reduction cathodic peaks at 0.17 V and 0.05 V are linked to the conversion of crystalline Si to amorphous Li-Si phases (a-Li x Si) during the lithiation and the subsequent crystallization and formation of c-Li 3.75 Si [8,81]. Oxidation anodic peaks at 0.35 V (a-Li 3.5 Si to a-Li 2 Si) and 0.5 V (a-Li 2 Si to a-Si) correspond to the transition of crystalline to an amorphous phase, followed by a dealloying reaction [8,89,90]. However, PPy/SiNC 6 and PPy/SiNA 20 (Fig. S11d–e) initially showed only one reduction and one oxidation peak. This behavior can be expected for PPy/SiNA 20 as it contains amorphous Si, meaning there is no transition from crystalline to amorphous Si, and only alloying and dealloying processes occur. As CV cycling continues, a reduction peak around 0.1 V starts to appear, indicating a gradual activation of lithiation sites. In the case of PPy/SiNC 6 , only one reduction and one oxidation peak are observed until the 7th cycle, when the reduction peak at 0.17 V first appears. The anodic peak at 0.35 V remains barely noticeable throughout the 10 cycles. These observations suggest that the lithiation connected with phase transformation of both PPy/SiNC 6 and PPy/SiNA 20 is slower compared to those of PPy/SiNC 20 . The positions of both reduction and oxidation peaks remain stable throughout cycling for all the studied anodes, indicating good reversibility of the Si-Li reactions [89]. CV cycling tests of PPy/SiNC 20 , PPy/SiNC 6 , and PPy/SiNC 55 reveal increasing peak intensities over time, a phenomenon previously reported in the literature and ascribed to the activation process of the anode material [89]. This activation process is facilitated by the amorphization of Si, which enhances lithiation capacity with each cycle [91]. Notably, this effect is less pronounced for PPy/SiNC 100 , where peak intensities initially increase but begin to decline after 5 cycles. This behavior can be attributed to the larger size of SiNC 100 , as well as its size heterogeneity, and a limited lithiation time, which restricts the depth of lithiation and stress buildup to the sub-surface regions in the larger nanoparticles. This localized stress can cause nanoparticle fracture, deactivation, and a subsequent decrease in peak intensity [91]. In contrast, the peak intensities of PPy/SiNC 55 (Fig. S11b) increase slightly over cycling, as a result of the smaller size of SiNC 55 . Among all the anodes, PPy/SiNC 20 demonstrates the best CV performance, which can be linked to the optimal size of the SiNC 20 , which balances surface area, Li + diffusion, and mechanical stability. In contrast, PPy/SiNC 6 exhibits lower peak intensities, reflecting reduced electrochemical activity and correlating with its initially lower specific capacity (Fig. 4a). This diminished performance is likely a result of nanoparticle agglomeration, which negates the typical advantages of ultrasmall particle sizes (<20 nm) [92]. Indeed, as previously observed, SEM images of PPy/SiNC 6 (Fig. 3i) reveal significant agglomeration of SiNC 6 , a feature unique to this anode. Meanwhile, PPy/SiNC 55 shows lower electrochemical activity consistent with the higher degree of oxidation exhibited by SiNC 55 (i.e., greater SiO 2 content). Since SiO 2 is largely electrochemically inactive, it contributes minimally to reversible Li storage and can hinder performance by forming insulating byproducts during lithiation. As noted above, during lithiation, SiO 2 can undergo an irreversible reaction with Li + to form lithium silicates and lithium oxide [93], which consume Li + without contributing to reversible capacity. These reaction products are typically electrically insulating and can block Li diffusion pathways, further reducing the active use of surrounding Si [94]. In addition, SiO 2 on the surface of SiNC 55 can hinder the formation of Li–Si alloys by acting as a barrier layer, reducing both Li accessibility and overall Fig. 5. CV measurements at a scan rate 0.1 mV/s of all the anodes in a) 1st, b) 3rd cycle, and c) 10th cycles. G. Soukupov´ a et al. Journal of Power Sources 661 (2026) 238620 9