Natural collagen fiber-enabled facile synthesis of carbon@Fe₃O₄ core–shell nanofiber bundles and their application as ultrahigh-rate anode materials for Li-ion batteries

Abstract

A C@Fe₃O₄ core–shell nanofiber bundle (C@Fe₃O₄NFB) was easily prepared using a natural collagen fiber (CF) as the biotemplate and carbon source. The as-prepared C@Fe₃O₄NFB features a continuous conductive pathway with dramatically accelerated electron and Li⁺ transport kinetics, thus exhibiting an ultrahigh-rate capability. Furthermore, the C@Fe₃O₄ core–shell structure is able to suppress the electrode pulverization due to the existence of the carbon nanofiber as an elastic buffering matrix, thus delivering long-term cycling stability. Based on electrochemical experiments, C@Fe₃O₄NFB delivered capacities of 839, 668, 422 and 301 mA h g⁻¹ at current densities of 0.2, 1.0, 5.0 and 10.0 A g⁻¹, respectively. At the current density of 1.0 A g⁻¹, the reversible capacity of C@Fe₃O₄NFB is as high as 632 mA h g⁻¹ in the 500th cycle, which accounts for 84.38% capacity of the 2nd cycle, suggesting a low capacity loss of 0.23 mA h g⁻¹ per cycle. Even after 2000 cycles at 5.0 A g⁻¹, the discharge capacity still reaches 354 mA h g⁻¹. This biotemplated synthesis approach provides a novel and interesting route for fabricating ultrahigh-rate anode materials of LIBs.

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1. Introduction

To replace commercial anode material of graphite in Li-ion batteries (LIBs), transition metal oxides, such as Fe₃O₄, SnO₂ and Co₃O₄, hold the promise of high energy density through conversion reactions.^1–3 Unlike graphite, dependent on Li⁺ intercalation, the storage of Li⁺ in transition metal oxides is based on the conversion reaction, which is expressed as MaOb + (bn)Li ↔ aM + bLinO, where M is transition metal and n is formal oxidation state of O.⁴ The significance of this conversion reaction is that the reduction of transition metal involves multiple electron transfers and allows the storage of multiple Li⁺ ions, thus delivering high gravimetric specific capacity. However, the high energy density of transition metal oxide is still accompanied by its intrinsic drawbacks: (i) transition metal oxides have low electrical conductivity and are unable to provide effective conductive pathways for fast electron transport. Thus, poor rate capability stemming from electrode polarization is inevitably encountered when a large current density is applied.^5,6 (ii) The particle reorganization and huge volume change of the transition metal oxide easily leads to electrode pulverization and loss of electrical contact, thereby dramatically decreasing the reversible capacity upon long-term cycling.^7–9 (iii) In addition, the high irreversible capacity loss of anode material in the first cycle renders LIBs a low coulombic efficiency.

Currently, the most possible strategies to circumvent the abovementioned issues include nanomaterials fabrication and carbon coating. These two strategies are usually combined to prepare nanocomposites, including core–shell nanoparticles (NPs), composite nanotubes and hybrid nanosheets,^10–13 so as to solve the kinetics and instability problems simultaneously. Although the utilization of these nanocomposites has successfully increased the reversible capacity to several folds of graphite, the tested cycling life is usually limited to several hundred cycles, and the current density applied during the cycles is still subjected to relatively low-rates, up to several hundred milliamperes per gram.^14–16 Nowadays, the long-term cycle-ability of thousands of cycles at real high current densities (at the ampere per gram level) is rarely achieved expect only a few examples that are based on the fabrication of ordered nanocomposite materials.^17,18 In general, the ordered nanocomposites feature shortened ionic diffusion distances and continuous electronic conductive pathways, which can significantly enhance the charge transfer kinetics of LIBs, thus minimizing the problems of electrode polarization caused by material structure-induced kinetic anisotropy.^19,20 In our recent investigations,²¹ we also found that transition metal oxide NPs encapsulated inside Ni foam-supported porous carbon can deliver ultrahigh-rate capability and super-long cycling stability at a large current density of 2.0 A g⁻¹ because of the existing of continuous pathways of charge transport.

However, current strategies used for the fabrication of ordered nanocomposites mainly rely on the use of artificial templates, which generally involve tedious procedures and strict control over experimental conditions.^17,18 Moreover, the scale-up fabrication of these ordered nanocomposites is another persisting problem due to the high cost of artificial templates. To address these issues, it is highly desirable to develop more convenient and low-cost synthetic routes. In this regard, a natural biotemplate provides us a more convenient choice than the artificial template owing to its intrinsic nature of well-defined nanostructured morphology, capability of self-assembly and abundant resource. Moreover, biotemplates usually have abundant functional groups, which can directly react with the metal precursors without the need of complex synthesis procedures or harsh reaction conditions. As a result, the synthesis is usually carried out at benign conditions such as room temperature and in aqueous solutions. From the viewpoint of green chemistry, a biotemplate synthesis approach shows a great potential in convenient and large-scale fabrication of ordered nanocomposite electrode materials.

In this study, we report a facile synthesis of 1D carbon@Fe₃O₄ core–shell nanofiber bundle (C@Fe₃O₄NFB) using a natural collagen fiber (CF) as the biotemplate. Natural CF is an abundant resource in nature, which can be conveniently prepared from the skins of domestic animals. As for the microstructure, CF has well defined hierarchical nanofibrous morphology,²² which is constituted of nanoscaled collagen fibrils. The collagen fibrils have a diameter of ∼50–200 nm, and these fibrils are further self-assembled into CF with micron scale diameters. Hence, CF is actually an ideal template for a 1D nanofiber bundle with orderly arranged hierarchical nanofibers. The use of this natural ordered template successfully eliminates the tedious procedures of template self-assembly. More importantly, CF has abundant reactive functional groups (e.g. –COOH and –NH₂), which are able to chelate with transition metal ions in aqueous solutions.^23–25 The significance of this property suggests that CF can directly react with transition metal ions in aqueous solutions at mild reaction temperatures without the need of further functionalization, thus forming a large-scale, core–shell structured nanofiber bundle with simplified procedures and low-cost. The core–shell nanofiber bundle structure of C@Fe₃O₄NFB was fully characterized by SEM, SEM-EDX, TEM, STEM, TEM-SADP, XRD and Raman techniques. The electrochemical properties of C@Fe₃O₄NFB as anode material were also thoroughly evaluated at various current densities, including the rate capabilities and long-term cycling stability. For comparison, carbonized CF and multi-wall carbon nanotubes mixed with commercial Fe₃O₄NPs (MWCNTs–Fe₃O₄NPs) were used in control experiments. Under the same experimental conditions, C@Fe₃O₄NFB exhibited considerably enhanced rate capability and cycling stability, thus manifesting the significant role played by the unique electrode configuration.

2. Material and methods

2.1 Preparation of C@Fe₃O₄NFB

Collagen fiber was prepared according to the literature.²⁶ 5.0 g of collagen fiber was soaked in 200 mL of deionized water at 25 °C, and the pH of the mixture was adjusted to pH ∼2.0–2.3. Then, 4.0 g of Fe₂(SO₄)₃ was added into the abovementioned mixture, maintained under constant stirring for 3.0 h. Furthermore, the pH of the mixture was increased to ∼3.8–4.0 using a saturated sodium bicarbonate solution, and the resultant mixture reacted at 40 °C for 12.0 h. Subsequently, the intermediates were collected by filtration, followed by washing and drying at 40 °C. The collected powder materials were treated by vacuum calcination at 600 °C (with the heating rate of 5 °C min⁻¹) for 3.0 h. The resultant materials were denoted as C@Fe₃O₄NFB. We also investigated the influences of the loading amount of Fe and thermal treatment conditions on the electrochemical properties of C@Fe₃O₄NFB by varying the amount of Fe₂(SO₄)₃ (2.0, 6.0, 8.0 and 10.0 g) and calcination temperature (500 °C, 600 °C, 700 °C and 800 °C).

2.2 Electrochemical measurement

The working electrodes comprised 70 wt% @Fe₃O₄NFB, 20% carbon nanotubes, and 10 wt% polyvinylidene fluoride (PVDF) binder and were cast onto a Cu foil and vacuum-dried at 90 °C. The electrochemical properties of the working electrodes were measured under ambient temperature using a CR2032 coin cell with a lithium foil serving as both counter and reference electrodes. The electrolyte was 1 M LiPF₆ dissolved in a 1 : 1 : 1 (volume ratio) mixture of ethylene carbonate (EC), ethyl methyl carbonate (EMC) and dimethyl carbonate (DMC). The galvanostatic charge/discharge tests were conducted using a CT-3008 Neware battery testing system with a voltage window of 0.005–3.0 V. Cyclic voltammetry (CV) measurements were carried out at a scan rate of 0.1 mV s⁻¹ using a LK9805 electrochemical work station and EIS measurements were performed using a Zahner (Zennium E) electrochemical work station. For comparison, commercial Fe₃O₄NPs (Aladdin Chemistry Co. Ltd.) were also used as the active material to prepare working electrodes and tested in CR2032 coin cells with a lithium foil serving as both counter and reference electrode.

2.3 Characterization

Surface morphologies of the samples were studied by field emission scanning electron microscopy (FESEM, Hitachi 4700, Japan) and the microstructure of the samples was analyzed using a Tecnai G2 F20 (TEM, FEI, Netherlands) operating at 200 kV. The X-ray diffraction patterns of the C@Fe₃O₄NFB composites were recorded using a Cu-K_α wide-angle X-ray diffractometer (XRD, Philips X'Pert Pro-MPD, Netherlands). X-ray photoelectron spectroscopy (XPS, Shimadzu ESCA-850, Japan) was carried out by employing Mg-K_α X-rays (hν = 1253.6 eV) and a pass energy of 31.5 eV, wherein all the binding energy peaks were calibrated by placing the principal C 1s binding energy peak at 284.7 eV. The peak fittings of all high resolution core spectra were carried out by the XPSPEAK 4.1 software, using mixed Gaussian–Lorentzian functions. Thermogravimetric analysis (TGA) was measured in air by a NETZSCH STA 409 thermogravimetric analysis (TGA) instrument at a heating rate of 10 °C min⁻¹.

3. Results and discussion

3.1 Preparation and characterization of electrode materials

Fig. 1a is the illustration showing the preparation procedures of C@Fe₃O₄NFB. Fe³⁺ ions are first chelatively adsorbed onto CF. This process does not damage the morphology of native CF according to SEM observation (Fig. S1†). In the following thermal treatment, the Fe³⁺ ions are converted to Fe₃O₄NPs and the CF is decomposed leaving the carbon nanofiber bundle, forming a typical core–shell structured C@Fe₃O₄NFB. Owing to this unique structure, the C@Fe₃O₄NFB features two distinct advantages, as shown in Fig. 1b: (i) the carbon nanofiber bundle, existing as a 1D ordered continuous conductive pathway, ensures sufficient electrode/electrolyte contact area and dramatically accelerated electron transport and Li⁺ diffusion, which substantially suppresses the electrode polarization at high current density, thus exhibiting an ultrahigh-rate capability. (ii) The C@Fe₃O₄ core–shell structure is able to suppress the electrode pulverization due to the existence of the carbon nanofiber as an elastic buffering matrix, thus delivering long-term cycling stability. (iii) The formation of solid electrolyte interface (SEI) on the fibrous bundle surface of C@Fe₃O₄NFB is enough to maintain stable lithiation/delithiation for each core–shell C@Fe₃O₄ nanofiber, which considerably reduces the amount of Li⁺ consumed by the formation of SEI, and thus, a high coulombic efficiency can be expected.

Fig. 1 The illustration showing (a) the preparation mechanism of C@Fe₃O₄NFB and (b) the advantages of the electrode designed for charge transport.

Fig. 2 shows the SEM images of native CF and C@Fe₃O₄NFB. Compared with native CF (Fig. 2a and b), C@Fe₃O₄NFB preserves well the hierarchal nanofibrous morphology of the biotemplate, of which the average diameter is in the range of 3–6 μm (Fig. 2c and d). An individual bundle of C@Fe₃O₄NFB (Fig. 2e) actually consists of a large number of nanofibers, of which the size is close to that of the collagen fibrils. Moreover, the HRSEM image in Fig. 2f confirms that each nanofiber is covered by aggregated Fe₃O₄NPs. Fig. 2g shows the SEM mapping images of C, Fe and O, wherein these elements share a similar shape as that of the nanofiber bundle, suggesting good coverage of Fe₃O₄NPs over the carbon nanofiber. The Raman spectra of C@Fe₃O₄NFB show the typical D band and G band of graphitic carbon at 1337 and 1584 cm⁻¹, respectively (Fig. S2†),²⁷ revealing the successful carbonization of the CF template to form carbon nanofibers. XRD patterns of C@Fe₃O₄NFB shows the characteristic (220), (311), (400), (422), (511) and (440) peaks of Fe₃O₄ at 30°, 36°, 43°, 54°, 57° and 63°, respectively, (Fig. S3†), demonstrating the existence of Fe₃O₄ in C@Fe₃O₄NFB.²⁸

Fig. 2 SEM images of (a, b) native CF and (c, d) C@Fe₃O₄NFB; (e, f) HRSEM images of marked area from (d) and (g) SEM-DES elemental mapping images of C@Fe₃O₄NFB.

Fig. 3 shows the TEM images of C@Fe₃O₄NFB. In Fig. 3a, an individual nanofiber bundle of C@Fe₃O₄NFB is observed, wherein the surface is covered by a rough layer of Fe₃O₄ with bright contrast.

Fig. 3 (a, b) TEM, (c) HRTEM images and (d) TEM SADP (selected area diffraction patterns) of C@Fe₃O₄NFB, (e) TEM and (f) HAADF images of C@Fe₃O₄NFB showing the clear core shell structure, wherein the area marked with a red box is analysed by TEM-DEX, (g) the XPS survey scan spectrum, (h) Fe 2p XPS spectra and (i) O 1s XPS spectra of C@Fe₃O₄NFB.

Fig. 3b shows the TEM image of a single nanofiber of C@Fe₃O₄NFB, wherein the d-period of native CF is clearly observed with a large number of Fe₃O₄NPs (the black dots) embedded onto the carbon nanofiber. The HRTEM image in Fig. 3c shows that the average particle size of Fe₃O₄NPs is ∼5–7 nm (Fig. S4†), and these NPs have high crystallinity as the (400) lattice reflections of Fe₃O₄ are clearly observed. Based on the TEM SADP (selected area diffraction patterns) in Fig. 3d, the diffraction rings are well indexed to Fe₃O₄ crystal faces of (311), (400), (422), (511) and (440).²⁹ These results confirm the formation of highly crystalline Fe₃O₄NPs on the surface of the carbon nanofiber. In Fig. 3e, the nanofiber bundle of C@Fe₃O₄NFB shows a clear carbon nanofiber core (with dark contrast) and Fe₃O₄NPs shell (with bright contrast). A further HAADF image (Fig. 3f) confirms that the nanofiber bundle indeed consists of a core–shell structured configuration. Based on TEM-EDX analysis (Fig. S5†), the shell layer with bright contrast is indeed Fe₃O₄. Hence, these TEM observations confirm the core–shell structured morphology of C@Fe₃O₄NFB. Fig. 3g shows the XPS survey scan spectrum of C@Fe₃O₄NFB, which consists of C, N, O and Fe without any other impurities. The Fe 2p XPS spectrum in Fig. 3h shows two peaks located at 710.4 and 724.3 eV (ref. 30) confirming the existence of Fe₃O₄. Further analysis of the O 1s XPS spectrum (Fig. 3i) reveals that there are still C=O (531.2 eV) and C–O (533.0 eV) functional groups existing on C@Fe₃O₄NFB,³¹ which are able to stabilize Fe₃O₄NPs during the lithiation/delithiation processes. Moreover, the core of the carbon nanofiber can serve as a buffer material to accommodate the huge volume change of Fe₃O₄ during the repeated lithiation/delithiation cycles. Hence, it is rational to expect that the C@Fe₃O₄NFB composite can exhibit a high cycling stability. TGA analysis confirms that the content of Fe₃O₄ in C@Fe₃O₄NFB is about 32.4% (Fig. S6†).

3.2 Electrochemical properties of C@Fe₃O₄NFB as an ultrahigh-rate anode material

The representative cyclic voltammetry (CV) curves of C@Fe₃O₄NFB as anode materials are shown in Fig. 4a. In the first cycle, the cathodic peak observed at 0.63 V can be attributed to the lithiation of Fe₃O₄ (Fe₃O₄ + 8Li⁺ + 8e⁻ ↔ 3Fe⁰ + 4Li₂O) and the formation of the SEI based on the irreversible reaction with the electrolyte.³² Accordingly, the peaks at 1.5 and 1.9 V are related to the anodic processing of Fe⁰ to Fe²⁺ and Fe³⁺.³³ Due to the formation of the SEI, the intensity of the cathodic peak in the first cycle is significantly higher than that in the 2nd cycle. In subsequent cycles, the CV curves overlap with only very small changes in area and shape, suggesting that the electrode materials have good reversible lithiation and delithiation abilities. This is consistent with the conclusion drawn from the following cycling stability measurements.

Fig. 4 (a) Cyclic voltammetry (CV) of C@Fe₃O₄NFB electrode for the first three cycles, (b) discharge–charge profiles of the initial three cycles of C@Fe₃O₄NFB, (c) cycling stability of C@Fe₃O₄NFB at the current density of 0.2 A g⁻¹ and (d) the electrochemical impedances before and after the cycling tests at 0.2 A g⁻¹, (e) rate capability of C@Fe₃O₄NFB at the current density of 0.2–5.0 A g⁻¹, (f) 0.2–10.0 A g⁻¹ and (g) 0.6–15.0 A g⁻¹, (h) the electrochemical impedances during the rate tests at 0.6–15.0 A g⁻¹.

Fig. 4b shows the initial three cycles of the galvanostatic discharge–charge voltage profiles of C@Fe₃O₄NFB at the current density of 0.2 A g⁻¹ within the voltage window of 0.005–3.0 V, wherein the specific capacities are calculated based on the mass of C@Fe₃O₄NFB. The 1st discharge and charge capacities of C@Fe₃O₄NFB are 1520 and 996 mA h g⁻¹, respectively, suggesting a first cycle coulombic efficiency of 65.53%. The irreversible capacity loss should be derived from the consumption of Li⁺ to form the SEI and irreversible insertion of Li⁺ into the defects of the electrode materials. Notably, this coulombic efficiency is much higher than those of electrode materials based on Fe₃O₄ nanocomposites,^34,35 although nanostructured electrode material is more likely to form a large area of SEI owing to the high specific surface area exposed to the electrolyte. We believe that the formation of SEI covering the bundle surface is enough to maintain stable lithiation/delithiation of each core–shell C@Fe₃O₄ nanofiber contained inside the nanofiber bundle (as shown in Fig. 1c). In this way, the Li⁺ consumed by the formation of SEI is considerably reduced. On the other hand, the coulombic efficiency in the 2nd cycle is quickly increased back to ∼97%, maintaining a high coulombic efficiency (∼100%) in subsequent cycles, as shown in Fig. 4c. C@Fe₃O₄NFB exhibits excellent cycling stability at 0.2 A g⁻¹, which can deliver a high reversible capacity of 847 mA h g⁻¹ in the 100th cycle, corresponding to a low capacity fading of 1.5 mA h g⁻¹ per cycle. In contrast, the multi-wall carbon nanotubes (MWCNTs) mixed with commercial Fe₃O₄NPs (Fig. S7†) (MWCNTs-Fe₃O₄NPs) exhibit much lower reversible capacity (526 mA h g⁻¹) in the 100th cycle with a current density of 0.2 A g⁻¹.

We investigated the electrochemical impedance of C@Fe₃O₄NFB after different cycles and analyzed by Z-view software using an equivalent circuit (Fig. 4d). As for the circuit, the Z′ values at the high-frequency region indicate the total ohmic resistance of the electrolyte and electrical contacts (R_e). The diameter of the semicircular region in medium-frequency is ascribed to surface film (R_sf) and the charge transfer resistance (R_ct) in the electrode reaction. Furthermore, the constant phase element (CPE₁) indicates the pure capacitance, and the inclined line in the low frequency region represents the Warburg impedance (W_o) related to lithium diffusion in the solid.^36,37 Based on the measurements of electrochemical impedances, the initial impedance resistance of C@Fe₃O₄NFB is 40 Ω, which is much lower than that of MWCNTs–Fe₃O₄NPs (67 Ω). In addition, there are two semicircles in the incipient cycle in the impedances of the MWCNTs–Fe₃O₄ electrode. The one at the high-frequency region reflects the R_e and the other is ascribed to R_sf and R_ct, which is much larger than that of the C@Fe₃O₄NFB electrode. After 100 cycles at 0.2 A g⁻¹, the impedance resistance of C@Fe₃O₄NFB is further decreased to 6 Ω. The decreased impedance resistance could be explained by the milling effect.^38,39 During the lithiation/delithiation process, electrochemical milling occurs on the C@Fe₃O₄NFB composite, and thus Fe₃O₄NPs become smaller. This pulverization process will increase the total surface area of the electrode material so that the contact area between the active material and electrode would be largely increased, and accordingly, the cell impedance decreases. Although the impedance resistance of MWCNTs–Fe₃O₄NPs is also decreased after the same cycles, the obtained value of 34 Ω is still larger than that of C@Fe₃O₄NFB. Actually, the Fe loading amount and thermal treatment conditions have significant influences on C@Fe₃O₄NFB. Excess/insufficient amount of Fe loading or inappropriate calcination temperature both lead to a low discharge capacity of C@Fe₃O₄NFB, as shown in Fig. S8 and S9.† To evaluate the rate capability of C@Fe₃O₄NFB, the rate performances were conducted in the voltage window of 0.005–3.0 V with the current density from 0.2 to 5.0 A g⁻¹. As shown in Fig. 4e, the reversible discharge capacity of C@Fe₃O₄NFB reaches 845 mA h g⁻¹ at the current density of 0.2 A g⁻¹ after 10 cycles. By increasing the current density to 0.4, 0.6, 0.8, 1.0, 2.0, 3.0 and 5.0 A g⁻¹, the reversible discharge capacities are still high up to 782, 741, 704, 658, 611, 543 and 440 mA h g⁻¹, respectively. No obvious capacity decrease is observed for each tested current density. In contrast, the discharge capacities of MWCNTs–Fe₃O₄NPs are much lower at 533, 488, 460, 420, 406, 314, 221 and 183 mA h g⁻¹ for 0.2, 0.4, 0.6, 0.8, 1.0, 2.0, 3.0 and 5.0 A g⁻¹, respectively. We also tested the discharge capacity of carbonized CF, which exhibits limited capacity, thus confirming that the high reversible capacity of C@Fe₃O₄NFB should mainly stem from the Fe₃O₄ active materials. Notably, the reversible discharge capacity of C@Fe₃O₄NFB can quickly increase back to 840 mA h g⁻¹ without any delay when the current density is reduced back to 0.2 A g⁻¹. The discharge capacity can still be maintained at 890 mA h g⁻¹ even after 20 cycles at a current density of 0.2 A g⁻¹. These results confirm that the C@Fe₃O₄NFB indeed exhibits high rate capability.

To further demonstrate the excellent rate capability of C@Fe₃O₄NFB, the rate tests were carried out at an even higher current density range from 0.2 to 10.0 A g⁻¹, wherein fast electron and charge transports are necessary to guarantee the high rate. As shown in Fig. 4f, the corresponding discharge capacities are 839, 668, 422 and 301 mA h g⁻¹ at 0.2, 1.0, 5.0 and 10.0 A g⁻¹. Moreover, the discharge capacity of C@Fe₃O₄NFB can quickly respond to the change of current density without loss of capacity when the current density is quickly switched between low and high current densities (10.0 A g⁻¹ → 1.0 A g⁻¹ → 0.2 A g⁻¹ → 1.0 A g⁻¹ → 10.0 A g⁻¹ → 0.2 A g⁻¹). Actually, the C@Fe₃O₄NFB is able to deliver the capacity of 392, 360, 350, 332, 311, 262 and 205 mA h g⁻¹ even at high current densities of 5.0, 6.0, 7.0, 8.0, 9.0, 10.0 and 15.0 A g⁻¹, respectively (Fig. 4g). At the same time, the capacity can still be recovered to ∼750 mA h g⁻¹ when the current density is decreased from 15.0 to 1.0 A g⁻¹.

As shown in Fig. 4h, the Nyquist plot of C@Fe₃O₄NFB at 5.0 A g⁻¹ shows a semicircle of 66 Ω in the medium frequency region, which is related to internal resistances in the electrode. Furthermore, the internal resistance of C@Fe₃O₄NFB is substantially decreased and reaches a stable value of 37 Ω at 10.0 A g⁻¹. The impedance resistance changes slightly at the rate test of 15.0 A g⁻¹. These results suggest that a stable SEI is quickly formed on C@Fe₃O₄NFB owing to the good accessibility of the nanofibrous bundle structure,^40,41 and the SEI is very stable even at such a high current density of 15 A g⁻¹.

In addition to rate capability, a long-term cycling stability of the C@Fe₃O₄NFB was also investigated at current densities of 1.0 and 5.0 A g⁻¹. As shown in Fig. 5a, the reversible capacity of C@Fe₃O₄NFB reaches 632 mA h g⁻¹ in the 500th cycle with the current density of 1.0 A g⁻¹, which accounts for 84.38% capacity of the 2nd cycle, suggesting a good cycling stability with a capacity loss of 0.23 mA h g⁻¹ per cycle. Similar to the rate performances, the stable SEI is quickly established at the initial cycles of cycling tests (Fig. 5b and c). Satisfactory cycling stability is also observed at the current density of 5.0 A g⁻¹ (Fig. 5d). When tested after 2000 cycles, the discharge capacity of C@Fe₃O₄NFB is as high as 354 mA h g⁻¹. Even after 6000 cycles at 5.0 A g⁻¹, the discharge capacity of C@Fe₃O₄NFB is still high up to ∼200 mA h g⁻¹. When cycled 8000 times at 5.0 A g⁻¹, the discharge capacity is around 146 mA h g⁻¹. The impedance resistance of C@Fe₃O₄NFB is very stable before the 2000 cycles, whereas this value is increased in the 8000th cycle (Fig. 5e and f), which indicates that our electrode suffered from a slow damage to electrode integration during the long-term cycling test at such a high rate. Actually, the long-term cycling stability of C@Fe₃O₄NFB presented is superior to other reported nanostructured Fe₃O₄ electrode materials.^42–47

Fig. 5 (a) Cycling stability of C@Fe₃O₄NFB at 1.0 A g⁻¹, (b and c) the electrochemical impedances of C@Fe₃O₄NFB during the cycling at 1.0 A g⁻¹, (d) a long-term cycling stability of C@Fe₃O₄NFB at 5.0 A g⁻¹, (e and f) the electrochemical impedances of C@Fe₃O₄NFB during the cycling at 5.0 A g⁻¹.

Fig. 6 shows the SEM and TEM images of the C@Fe₃O₄NFBs composite after 100 cycles at the current density of 1.0 A g⁻¹. It is clear that the fibrous structure of C@Fe₃O₄NFB composite is well retained without cracking or pulverization, which manifests the effective accommodation of the elastic carbon nanofiber of C@Fe₃O₄NFB for the Fe₃O₄ NPs.

Fig. 6 (a) SEM and (b) TEM images of C@Fe₃O₄NFB composite after 100 cycles at a current density of 1.0 A g⁻¹.

4. Conclusions

In summary, we successfully prepared one dimensional carbon@Fe₃O₄ core–shell nanofiber bundle (C@Fe₃O₄NFB) using a collagen fiber as both the biotemplate and carbon source. Owing to the unique material structure, the as-prepared C@Fe₃O₄NFB can act as a high performance anode material for Li-ion batteries. The carbon nanofiber bundle of C@Fe₃O₄NFB serving as a conductive pathway dramatically accelerates the charge transfer in LIBs. The designed core–shell structured bundle morphology of C@Fe₃O₄NFB is able to suppress electrode pulverization. The fibrous bundle morphology of C@Fe₃O₄NFB helps to decrease the required Li⁺ with improved columbic efficiency. As a consequence, the rate capability and long-term cycling stability of the electrode materials can be dramatically improved. The biotemplate of collagen fiber is an abundant resource with low-cost and thus, our strategy may be extended for convenient and scaled-up synthesis of high performance anode materials.

Acknowledgements

This study was supported by the National Natural Science Foundation of China (51507107) and the National High Technology R&D Program (2011AA06A108).

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c5ra22481k

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