*2.2. Synthesis of Si@RF@SiO<sup>2</sup>*

The above prepared Si NPs were modified with 3-aminopro-pyltrimethoxysilane (APTES) (AR) to positively charge the surfaces. In total, 0.4 g Si NPs were uniformly dispersed in 300 mL ethanol (AR), containing 4 mL of APTES, and stirred for 5 h to obtain APTES-Si NPs. The above APTES-Si NPs was re-dispersed in an alkaline mixture of 150 mL deionized water and 30 mL ethanol, containing 1 mL of aqueous ammonia (AR) under magnetic stirring for 30 min. After that, 0.6 g resorcinol (AR) and 0.8 mL formaldehyde (AR) were separately added to the reaction system and continued stirring for 10 h to form a homogeneous phenolic resin (RF) coating layer under room temperature. The Si@RF powders were obtained via centrifugation treatment of the reaction solution. For the SiO2-coated Si@RF composite (named Si@RF@SiO2), 400 mg Si@RF was mixed with 1.5 mL tetraethyl orthosilicate (TEOS) (AR) hydrolysis under alkaline condition to form SiO<sup>2</sup> coating layer on the outer surface of Si@RF.

### *2.3. Synthesis of Si@C@v@CNTs*

Carbon nanotubes were grown in situ via a CVD method using Iron(III) nitrate nonahydrate (AR) as the catalyst (Fe) and acetylene (5%) as carbon precursors at 900 ◦C for 2 h under Ar/H<sup>2</sup> in a tube furnace. The catalyst was loaded on the precursor microspheres of Si@RF@SiO<sup>2</sup> prior to the deposition procedure to ensure that the CNTs could be grown in situ on the active positions during the CVD process. Finally, the SiO<sup>2</sup> coating layer was removed with dilute hydrofluoric acid (AR) solution and followed by centrifugation treatment and ethanol washing. After removing the SiO<sup>2</sup> sacrificial coating layers, the final composite was named Si@C@v@CNTs ("v" stands for "void"). For comparison purposes, the Si@C@v@C ("v" stands for "void") without in-situ grown CNTs was also synthesized via the same CVD method but in the absence of a Fe(NO3)3·9H2O catalyst, and the Si@v@CNT without an inner carbon layer was also synthesized via the same CVD method in the absence of the RF-layer.

### *2.4. Characterizations*

The XRD patterns of samples were obtained via a DX-2007 X-ray diffraction (XRD) experiment apparatus (λ = 1.5418 Å) (Dandong Haoyuan Instrument Co, Dandong, China) to confirm the phases and crystallinity. The Si content of composites was characterized by a thermogravimetric analyzer (TGA4000) (NSK LTD, Tokyo, Japan) at a heating rate of 10 ◦C min−<sup>1</sup> in air atmosphere. The nitrogen adsorption/desorption isotherm curves were obtained via a Micromeritics ASAP-2020M nitrogen adsorption/desorption apparatus (Best Instrument Technology Co, Beijing, China) to confirm the porosity character. The micro-structures and morphologies of the materials were collected via a JSM-6700F scanning electron microscope (SEM) (JEOL, Tokyo, Japan) with an IE300X energy-dispersive X-ray spectrometer and a JEM-2100F transmission electron microscope (TEM) (JEOL, Tokyo, Japan). X-ray photoelectron spectroscopy (XPS) curves were obtained by applying an electron spectrometer (ESCALab250) (Thermo Fay, Boston, MA, USA) to analyze the surface of composites. Raman spectra was collected with a Raman spectrometer (JobinYvon HR800) (Renishaw, London, UK).

### *2.5. Electrochemical Measurements*

The CR2016 coin-type cells were assembled in the glove box under inert atmosphere conditions without water and oxygen to test the electrode performance by using polypropylene films as separators (the thickness of the separator was 25 µm) to separate the working and counter electrodes (lithium wafer) in an electrolyte. The electrolyte was the 1 M LiPF<sup>6</sup>

dissolved in the solvent of ethylene carbonate and dimethyl carbonate (the amount of the electrolyte used in assembling the coin-type cell was 70 µL). The working electrode was made by coating the slurry of the above active materials containing a proportional conductive carbon black and polyvinylidene fluoride (PVDF) binder on the copper foil (the mass loading of the electrode was about 1.2 mg cm−<sup>2</sup> ), which was then dried under a 120 ◦C vacuum oven for 14 h (the thickness of the active electrode layer was about 20 µm). Finally, the copper foil was cut into wafers with a uniform size of 1 cm. The galvanostatic cycling measurements were conducted by a CT 2001A battery tester at determinate voltage windows. Cyclic voltammogram (CV) tests were performed by using an electrochemical workstation within a fixed voltage range and scan rate.

### **3. Results and Discussions**

Figure S1 shows the XRD patterns of reduced Si NPs, Si@RF, and Si@RF@SiO2. Three sharp diffraction peaks, which are located at 2θ values of 28.4, 47.2, and 56.1◦ , were attributed to the planes of (111), (220), and (311) for the crystal Si phase, respectively (JCPDS NO. 27-1402), indicating that the amorphous SiO<sup>2</sup> synthesized by the well-established Stöber method were fully reduced to crystalline Si in the MR process [43]. Another broad peak at ∼25◦ corresponded with the amorphous carbon and silica coming from the double coat with RF and silica layers. As shown in Figure 1a, for Si@C@v@CNTs and Si@v@CNTs, after the CVD process, CNT grown in situ across Si-CNPs and the diffraction peak at ∼25◦ was observed, corresponding to the (002) plane of the crystalline carbon [40]. Figure 1b displays the Raman spectra of Si@C@v@CNTs, Si@v@CNTs, and Si@C@v@C. The sharp peak at about 500 cm−<sup>1</sup> could be appointed to the Si peaks. Additionally, the weak peaks at about 1345 cm−<sup>1</sup> (D-band) and 1595 cm−<sup>1</sup> (G-band) could be associated with the vibration modes of sp<sup>3</sup> -bonded carbon atoms in amorphous carbon and sp<sup>2</sup> -bonded carbon atoms in typical graphite, respectively [44]. As calculated, the *ID/I<sup>G</sup>* was 0.98, 0.96, 0.95 for Si@C@v@C, Si@v@CNTs, and Si@C@v@CNTs, respectively. Si@C@v@CNTs had a relatively higher graphitization degree due to the microcrystalline structure of CNTs. Thermogravimetric analysis (TGA) was tested to confirm the proportion of carbon and silicon for the samples (Figure 1c). The weight losses occurred from 500 to 800 ◦C were ascribed to the carbon combustion and calculated to be 54.3%, 70.6%, and 85.9% for Si@C@v@C, Si@v@CNTs, and Si@C@v@CNTs, respectively. As the temperature continued to rise, Si NPs were further oxidized, leading to a weight increase in the TGA curves. The Si content in the three samples were 45.7%, 29.4%, and 14.1% for Si@C@v@C, Si@v@CNTs, and Si@C@v@CNTs, respectively. Compared with the other two samples, high C and CNT content in Si@C@v@CNTs provided more electric contact between particles, thus enhanced overall electronic conductivity and superior performance could be expected. The elemental compositions and valence states in the composites were determined by X-ray photoelectron spectroscopy (XPS) spectra. Figure 1d reveals the whole spectrum of Si@C@v@CNTs, confirming the coexistence of Si, C, and O. The high-resolution spectrum of Si 2p is shown in Figure 1e. Three peaks of Si-Si (98.6 eV), Si-C (101.6 eV), and Si-O (102.98 eV) originated from monatomic Si, residual silica, or slight oxidation of the Si NPs [45]. In addition, the high-resolution C 1s and O 1s in Figure S2 shows that the peak at 283.75 eV was related to the graphite-like sp<sup>2</sup> hybridized carbon and another peak at 283.3 eV was assigned to C-O. The O 1s peak at 537 eV may be attributed to adsorbed oxygen and the residual silica layer [46].

The Brunauer−Emmett−Teller (BET) specific surface area and pore volume for the samples are illustrated in Table S1. Compared with Si@C@v@C, the Si@C@v@CNTs had a relatively higher BET specific surface area of 106.98 m<sup>2</sup> g −1 , with a pore volume of 0.3212 cm<sup>3</sup> g −1 . We believe that the higher surface area was mainly due to the existence of reserved void space and in-situ grown CNT networks. The nitrogen adsorption–desorption isotherm curve of the samples is displayed in Figure 1f. All the curves had a distinct hysteresis loop, suggesting the existence of a mesoporous structure [47]. The pore size distribution curve (Figure S3) shows a diverse pore structure between 2 and 10 nm, which

incorporate a series of micro- and mesopores derived from CNT networks [48]. Furthermore, these multiple pore structures not only afford fast and shortened electronic/ionic diffusion pathways but also absorb the huge expansion in volume inside Si NPs during cycling, resulting in a significantly enhanced overall structural integrity and electrochemical performance.

**Figure 1.** XRD patterns (**a**), Raman spectra (**b**), and thermogravimetric analysis (TGA) curves (**c**) of Si@C@v@C, Si@v@CNTs, and Si@C@v@CNTs; X-ray photoelectron spectroscopy (XPS) of Si@C@v@CNTs (**d**), High-resolution XPS spectrum of Si 2p (**e**), and N<sup>2</sup> adsorption-desorption isotherms curves of Si@C@v@C, Si@v@CNTs, and Si@C@v@CNTs (**f**).

Figure 2 shows the scanning electron microscopy (SEM) images of the structural evolution of Si@C@v@CNTs at different synthesis stages. The SEM image of SiO<sup>2</sup> spheres firstly synthesized by the well-established Stöber method is shown in Figure 2a. The pristine SiO<sup>2</sup> spheres present a monodispersed spherical shape with a uniform diameter of approximately 300 nm. After the Mg-reduction process, porous Si NPs with well-preserved monodispersed spherical morphology were successfully prepared (Figure 2b). Figure 2c,d show the SEM images of Si@RF and Si@RF@SiO2. The diameters of as-prepared Si@RF and Si@RF@SiO<sup>2</sup> precursors were increased to 400 nm and 500 nm, respectively, indicating the successful coating of RF and SiO<sup>2</sup> layers on the Si cores, resulting in Si@C@SiO<sup>2</sup> particles. Figure 2e,f reveals a large amount of tangled CNTs with various diameters in the range of 50–100 nm externally grown in situ across the Si@C@SiO<sup>2</sup> particles. The multiple Si NPs were well supported by the in-situ generated flexible porous CNTs networks, contributing to better electric contact, not only between the inner Si-yolks but also among the Si-CNT microparticles. Figure 2g reveals the uniform elemental distribution of the Si, C, and Fe in Si@C@v@CNTs. It is clearly observed the Si-yolks were well coated by the outer C-shell and distributed across the flexible CNT networks. Figure S4 shows the SEM images of Si@C@v@C. The observed wrinkles over the entire surface confirm the full encapsulation of Si@C@SiO<sup>2</sup> by the carbon layer. The SEM images of Si@v@CNTs in Figure S5 reveal the similar tangled CNTs with Si@C@v@CNTs.

Figure 3a–c shows the transmission electron microscopy (TEM) images of homemade SiO<sup>2</sup> and reduced Si NPs. The high-resolution TEM image (HRTEM), taken from one of the Si NPs, reveals that the obtained Si NPs had high crystallinity [49,50]. The TEM images of Si@RF and Si@RF@SiO2, as shown in Figure 3d–e, confirm that the Si NPs were well wrapped by the RF carbon-SiO<sup>2</sup> double coating layers. The TEM and HRTEM images of Si@C@v@CNTs presented in Figure 3f–i confirm CNTs grown in situ on the outer layer of the Si@RF NPs with a void between the two. The carbon layer derived from RF was about 5 nm and could accelerate electron transfer between Si NPs and CNTs, leading to enhanced structural stability. The void generated due to the etching of SiO<sup>2</sup> could effectively alleviate the expansion of the inner Si NPs. According to the TEM images of a single CNT shown in Figure S6, the outer diameter of this single CNT was approximately 35 nm and the tube wall was 12 nm. The TEM image of Si@C@v@C in Figure S7 shows that the Si NPs were well wrapped with double carbon layers. The TEM images of Si@v@CNTs in Figure S8 confirms that existence of similar CNTs grown to Si@C@v@CNTs but without of a carbon layer on the Si NPs.

**Figure 2.** SEM images of SiO<sup>2</sup> nanoparticles (NPs), synthesized by the well-established Stöber method (**a**), reduced Si NPs in the magnesiothermic reduction (MR) process (**b**), SEM images of the in-process samples: Si@RF (**c**), Si@RF@SiO<sup>2</sup> (**d**), Si@C@v@CNTs (**e**,**f**). (**g**) the elemental mapping results of Si@C@v@CNTs.

**Figure 3.** Transmission electron microscope (TEM) images of SiO<sup>2</sup> NPs synthesized by the wellestablished Stöber method (**a**), reduced Si NPs in the MR process (**b**), high resolution TEM (HRTEM) images of Si NPs (**c**), TEM images of Si@RF (**d**), Si@RF@SiO<sup>2</sup> (**e**), and Si@C@v@CNTs (**f**–**h**), and HRTEM images of Si@C@v@CNTs (**i**).

The cyclic voltammetry (CV) curve of Si@C@v@CNTs exhibited the typical electrochemical properties of Si-based anode materials (as shown in Figure 4a). In the first cathodic branch (lithiation), a distinct broad peak between 0.5 and 0.8 V was ascribed to the generated solid-electrolyte interphase (SEI) film. However, this peak disappeared after the first lithiation, suggesting the generation of the firm and stable films during the first cycle [51]. The lithiation peak at 0.18 V can be appointed to the lithiation process of Si. In the following anodic branch (delithiation), the peaks located at 0.34 and 0.5 V were ascribed to the dealloying process from LixSi to amorphous Si [52]. Figure 4b shows the initial charge–discharge profiles of the materials at 100 mA g−<sup>1</sup> . The disappearance of the voltage plateaus between 0.5 and 0.8 V after the first cycle also confirms the generation of stable SEI films, which is in accordance with the CV results in Figure 4a [44]. The initial discharge and charge capacities were 2698.4 and 1684.2 mAh g−<sup>1</sup> for Si@C@v@C, 2546.5 and 1760.6 mAh g−<sup>1</sup> for Si@v@CNTs, and 2350.1 and 1787.0 mAh g−<sup>1</sup> for Si@C@v@CNTs at 100 mA g−<sup>1</sup> , corresponding to the initial coulombic efficiencies (ICE) of 62.41, 69.14, and 76.04%, respectively. Figure S9 shows the galvanostatic charge–discharge curves of Si@C@v@CNTs during the first five cycles at 100 mA g−<sup>1</sup> . After the first cycle, the CE increased to 92.73% in the following cycling test and reached 96.04% after the fifth cycle. Most importantly, the voltage plateaus during the cycling test was well-maintained, suggesting improved electrochemical utilization of the active electrode materials. The cycling behavior in Figure 4c shows the charge capacities at 100 mA g−<sup>1</sup> for Si@C@v@C.

**Figure 4.** Electrochemical properties: cyclic voltammetry for the first three cycles of Si@C@v@CNTs at a scan rate of 0.1 mV s−<sup>1</sup> between 0–2.0 V (**a**), the initial discharge–charge curves at 100 mA g−<sup>1</sup> (**b**), cycling performances at 100 mA g−<sup>1</sup> (**c**), the rate performances (**d**) of Si@C@v@C, Si@v@CNTs, and Si@C@v@CNTs.

Si@v@CNTs and Si@C@v@CNTs were 598.7, 736.1, and 912.8 mAh g−<sup>1</sup> after 100 cycles, respectively, indicating that the Si@C@v@CNTs were endowed with the best cycling behavior. Therefore, the outer in-situ grown CNT networks can significantly improve the cycling behavior and structural integrity of the inside-coated Si NP anodes. Figure 4d shows the rate performances of Si@C@v@C, Si@v@CNTs, and Si@C@v@CNTs performed at a series of different current densities. As expected, Si@C@v@CNTs present the best rate ability, even at high current densities. Very high reversible capacity of 907.7 mAh g−<sup>1</sup> was maintained when the current density was back to 100 mA g−<sup>1</sup> , suggesting an excellent rate ability of Si@C@v@CNTs. Furthermore, a high reversible capacity of 650 mAh g−<sup>1</sup> was retained for Si@C@v@CNTs at high 1 A g−<sup>1</sup> after 1000 cycles (Figure 5a). Therefore, it can be concluded that the introduced void space and porous CNT networks can absorb the

huge volume expansion inside Si NPs, leading to enhanced overall structural stability and integrity [52–54]. In addition, the CNT networks and inner rigid carbon coating provided more sufficient conductive contact to fast electronic/ionic diffusion pathways, resulting in significantly improved cycling stability and rate performance [55–57]. Figure S10 shows the SEM images of Si@C@v@CNTs after 1000 cycles at 1 A g−<sup>1</sup> . No cracking can be detected for Si@C@v@CNTs, and all the active materials were still well adhered to the current collector without exfoliation (Figure S10a,b). In addition, the morphology was well maintained in Si@C@v@CNTs after the long cycling test. (Figure S10c). Figure 5b shows the schematic illustration of lithiation and delithiation processes.

**Figure 5.** Cycle performance of Si@C@v@CNTs at a current density of 1000 mA g−<sup>1</sup> (**a**) and a schematic illustration of lithiation and delithiation processes (**b**).

Moreover, we fabricated a full cell with our Si@C@v@CNTs as the anode and LiNi0.6 Co0.2Mn0.2O<sup>2</sup> (NCM622) as the cathode (Si@C@v@CNTs//NCM622). The voltage profiles of NCM622 are shown in Figure S11. The NCM622 cathode exhibited a stable reversible capacity of 160 mAh g−<sup>1</sup> with a flat charging/discharging plateau at about 3.5 V. Referring to the voltage profiles of the Si@C@v@CNTs anode and the NCM622 cathode from half cells, the working potential range for the full cell was set between 2 and 4 V. The cycling performance of Si@C@v@CNTs//NCM622 is shown in Figure S12. The full cell displays a reversible capacity of 92 mAh g−<sup>1</sup> at 100 mA g−<sup>1</sup> after 100 cycles, indicating potential cycling stability for commercial viability.

### **4. Conclusions**

In summary, we synthesized a yolk-shelled structured silicon/carbon nanotube composite for high performance lithium storage application. This novel Si-based anode was made of an external grown flexible CNT network to confine the inner multiple Si NPs. The in-situ generated outer CNT networks not only accommodated the huge changes in volume space inside Si nanoparticles but also provided fast electronic/ionic diffusion pathways, resulting in markedly improved cycling stability and rate ability. Furthermore, the facial and scalable preparation method created in this work could make this new Si-based anode material promising for practical application in next generation Li-ion batteries.

**Supplementary Materials:** The following are available online at https://www.mdpi.com/2079 -4991/11/3/699/s1. Figure S1: XRD patterns of reduced Si NPs in the MR process, Si@RF, and Si@RF@SiO<sup>2</sup> . Figure S2: XPS spectrum of C1s (a) and O1s (b) of Si@C@v@CNTs. Figure S3: The pore size distribution curve of Si@C@v@C, Si@v@CNTs, and Si@C@v@CNTs. Figure S4: SEM images of Si@C@v@C. Figure S5: SEM images of Si@v@CNTs. Figure S6: TEM images of a single CNT. Figure S7: TEM images of Si@C@v@C. Figure S8: TEM images of Si@v@CNTs. Figure S9: The first five dischargecharge curves of Si@C@v@CNTs at current density of 100 mA g−<sup>1</sup> . Figure S10: Digital photograph (a), SEM image of Si@C@v@CNTs after the cycling test at 1.0 A g−<sup>1</sup> (b,c). Figure S11: Charge/discharge profiles of NCM626 between 2.0–4.3 V. Figure S12: The electrochemical performance of the full cell using Si@C@v@CNTs as anode and LiNi0.6Co0.2Mn0.2O<sup>2</sup> (NCM622) as cathode at the current density of 100 mA g−<sup>1</sup> . Table S1: The Brunauer-Emmett-Teller (BET) surface area, pore volume and average pore size of the samples.

**Author Contributions:** Conceptualization, J.L. and Z.P.; methodology, N.H., and Z.P.; software, X.W.; validation, C.Z., J.Q., and G.L.; formal analysis, X.L.; investigation, N.H.; resources, X.Z.; data curation, N.H.; writing—original draft preparation, N.H.; writing—review and editing, N.H., X.Z., and L.Z.; supervision, X.Z.; project administration, X.Z. and Z.P. All authors have read and agreed to the published version of the manuscript.

**Funding:** This work was supported by the National Natural Science Foundation of China (No. 51503109) and the China Postdoctoral Science Foundation Funded Project (No. 2016M600522).

**Data Availability Statement:** The data presented in this study are available on request from the corresponding author.

**Conflicts of Interest:** The authors declare that they have no competing interests.

### **References**


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