Vanadium pentoxide/carbide-derived carbon core – shell hybrid particles for high performance electrochemical energy storage † 2016, 4 , 18899

A novel, two step synthesis is presented combining the formation of carbide-derived carbon (CDC) and redox-active vanadium pentoxide (V 2 O 5 ) in a core – shell manner using solely vanadium carbide (VC) as the precursor. In a ﬁ rst step, the outer part of VC particles is transformed to nanoporous CDC owing to the in situ formation of chlorine gas from NiCl 2 at 700 (cid:1) C. In a second step, the remaining VC core is calcined in synthetic air to obtain V 2 O 5 /CDC core – shell particles. Materials characterization by means of electron microscopy, Raman spectroscopy, and X-ray di ﬀ raction clearly demonstrates the partial transformation from VC to CDC, as well as the successive oxidation to V 2 O 5 /CDC core – shell particles. Electrochemical performance was tested in organic 1 M LiClO 4 in acetonitrile using half- and asymmetric full-cell con ﬁ guration. High speci ﬁ c capacities of 420 mA h g (cid:3) 1 (normalized to V 2 O 5 ) and 310 mA h g (cid:3) 1 (normalized to V 2 O 5 /CDC) were achieved. The unique nanotextured core – shell architecture enables high power retention with ultrafast charging and discharging, achieving more than 100 mA h g (cid:3) 1 at 5 A g (cid:3) 1 (rate of 12C). Asymmetric cell design with CDC on the positive polarization side leads to a high speci ﬁ c energy of up to 80 W h kg (cid:3) 1 with a superior retention of more than 80% over 10 000 cycles and an overall energy e ﬃ ciency of up to 80% at low rates.


Introduction
Supercapacitors are electrochemical energy storage devices, which present ultrafast charge-discharge rates leading to higher power compared to conventional batteries, but suffer from relatively low specic energies. [1][2][3][4] The high specic power of $10 kW kg À1 benets applications with charge-and discharge rates in the order of seconds and minutes, such as ashlights, high power machinery, or emergency doors in airplanes, to name just a few. 2,5 Depending on the energy storage mechanism, we can differentiate between two types of supercapacitors, namely electrical double-layer capacitors (EDLCs) that store energy by ion electrosorption, and pseudocapacitors that accomplish a capacitor-like charge-voltage behavior by fast redox reactions. 6 Today, most EDLCs employ activated carbon (AC) electrodes with surface areas between 1500 and 2000 m 2 g À1 , leading to typical capacitance values of 100-200 F g À1 and specic energies of up to ca. 20 W h kg À1 , depending on the electrolyte and the cell conguration. 6 By virtue of these values, the energy storage capacity of supercapacitors is signicantly smaller than conventional batteries (e.g., lithium ion batteries: >100 W h kg À1 ). 6 A facile way to enhance the electrochemical performance is to increase the specic surface area of the electrode material and to optimize the pore size distribution. 7 This is particularly effective for a material with very high pore tunability, as presented in the case of carbide-derived carbons (CDCs). 8 However, this strategy is limited by the minimum pore size required for ion accessibility and the charge screening ability of carbon. 7,9,10 Further enhancements of the energy storage capacity can be accomplished by employing redox-active materials for faradaic charge transfer. 11,12 The latter can be either in the form of fast redox processes of surface groups, ion intercalation in the electrode material, or redox reactions of the electrolyte. 13,14 Depending on the charge-voltage prole of the involved faradaic reaction(s), one differentiates between capacitor-like (i.e., pseudocapacitive) or battery-like behavior. 15 Pseudocapacitive materials, exhibiting a constant ratio of charge and voltage (i.e., capacitance, unit: farad), are, for example, ruthenium oxide, 16 manganese oxide, 17 or MXene. 18 Battery-like systems, exhibiting distinct redox peaks in the cyclic voltammogram, are, for example, quinone surface groups 19,20 or many metal oxides working as cathodic lithiumintercalation hosts (such as LiFePO 4 , LiCoO 2 , LiMn 2 O 4 , and LiFeSO 4 F; ref. 21).
A very promising electrochemical system employs the hybridization of carbon materials with vanadium pentoxide (V 2 O 5 ). [22][23][24][25][26] By this way, the benecial electrical conductivity and nanotextured surface of carbon is complemented by the high, battery-like charge storage capacity of vanadia with a theoretical capacity of 147 mA h g À1 for intercalation of 1 M of Li in the V 2 O 5 structure. 27 Several studies have investigated vanadium pentoxide as a cathode material in lithium ion batteries, exhibiting high capacities between 350 and 750 mA h g À1 . [28][29][30] Furthermore, other types of vanadia in different oxidation states can be used in batteries, such as VO 2 , 31 and also dissolved as an electrolyte in redox-ow batteries. 32 In a previous study, we used atomic layer deposition (ALD) for highly controllable decoration of carbon onions with V 2 O 5 , achieving dened thickness, morphology, and crystallinity of the metal oxide phase. 33 A high specic energy (38 W h kg À1 ) was achieved for a full-cell using V 2 O 5 /carbon onion vs. activated carbon (asymmetric cell), which is ca. 2-times higher than that of standard activated carbon in organic electrolytes. 33 Even higher values for the specic energy were found for interpenetrating networks of CNTs and V 2 O 5 nanowires (40 W h kg À1 ) 22 and vanadia-coated reduced graphene oxide sheets (76 W h kg À1 ). 34 While promising, the practical adaptation of microfabrication and ALD remains limited and more facile synthesis routes for V 2 O 5 /carbon hybrid materials remain in high demand.
The combination of redox-active metal oxides with conductive and porous carbons in a hybrid device can be achieved by mechanical mixing (thus exhibiting strongly reduced conductivity), or the decoration of the carbon with metal oxides. For the latter, especially when using highly porous carbons, issues of pore clogging may arise, leading to a strong decrease in capacitance and specic energy. These complications are avoided by not coating the carbon with the metal oxide, but rather fabricating a metal oxide core with a porous carbon shell. Inspired by carbon only core-shell architectures, which can be obtained for CDCs, 35,36 we present a novel two-step synthesis for V 2 O 5 /CDC composites in a core-shell arrangement. Our approach requires only one precursor (vanadium carbide), which is rst partially transformed to nanoporous CDC by chlorine treatment of carbide, and then partially calcined, yielding core-shell V 2 O 5 /CDC particles. The chlorination step was needed as it is an effective way to transform a metal carbide to carbon. The nanoporous carbon shell makes the vanadium pentoxide core accessible for ion intercalation and works as a conductive additive. This was accomplished by a high level of control over the synthesis steps, which prevented a completion of both the chlorination and the oxidation processes. Rigorous structural and electrochemical characterization was carried out to establish the high energy storage capacity of the material in organic electrolytes. The combination of carbon and vanadia is of mutual benet: carbon contributes electrical conductivity (and, to a lesser extent, double-layer capacitance) and vanadia provides a large energy storage capacity. However, the design of the two phases and their nanoscopic arrangement are of high importance for the resulting electrochemical performance. As our data will show, the rm of transformed VC-derived V 2 O 5 and VC-CDC enables attractive electrochemical performances.

Materials and methods
V 2 O 5 /porous carbon core-shell particles were synthesized using a two-step approach: (1) VC/CDC core-shell particles are produced by etching vanadium from the outside to the inside of VC particles with a remaining VC core. (2) The residual VC core inside the CDC matrix is calcined by oxidation in synthetic air, resulting in V 2 O 5 /CDC core-shell particles. These two steps are explained in detail in the next sections.
2.1 Synthesis of vanadium carbide/carbide-derived carbon core-shell particles Vanadium carbide (VC)/CDC core-shell particles were synthesized by reactive extraction of vanadium using in situ generation of chlorine. 1 g VC powder (purity 99.9%, <2 mm, Sigma Aldrich) was thoroughly mixed with NiCl 2 $6H 2 O (purity 99.95%, Alfa Aesar) and lled into a graphite crucible placed in an isothermal zone of the reactor (Ø ¼ 32 cm, length ¼ 1 m). The reactor was evacuated to 0.01 mbar and then heated to 700 C at a heating rate of 2.5 C min À1 . Aer a holding period of 3 h at this temperature, the reactor was cooled to room temperature under a helium purge. The resulting product was then poured into 500 mL of 1 M HCl and the solution was vigorously stirred using a magnetic stir bar. Aer ltration, the acid washing was repeated with similar procedures to ensure the removal of nickel species. The product was nally washed using distilled water until neutral pH and then dried in an oven (60 C) overnight.
To achieve incomplete carbide-to-CDC conversion, we used molar ratios of NiCl 2 $6H 2 O-to-VC of 1.4, 1.8, and 2.7. It is assumed that chlorine (Cl 2 ) is released during vacuum decomposition of NiCl 2 and reacts with VC in a stoichiometric reaction to produce VCl 4 via eqn (1) and (2).
Based on these reactions and assuming that all chlorine reacts with carbide the theoretical conversion (X theo. ) can be derived from eqn (3) assuming full reaction of VC with Cl 2 release.
where n is in molar units. Ratios of 1.4 and 1.8 correspond to a theoretical value (X theo. ) of the conversion of 70% and 90%, respectively. A ratio of 2.7 provides a surplus of evolving chlorine to facilitate complete CDC formation, theoretically 100%. The resulting samples are labelled VC-X70, VC-X90, and VC-X100.
To produce a reference material through the "classical" and not in situ chlorine generation route, VC was exposed to chlorine gas diluted in helium at 700 C for 5 h, as described elsewhere. 63 The resulting material is labelled VC-CDC.

Synthesis of V 2 O 5 /carbide-derived carbon core-shell particles
Calcination of the VC/CDC particles was performed at 450 C for 30 min in a VG Scienta furnace using synthetic air. Additionally, VC particles without prior chlorine gas treatment were oxidized at different temperatures (450, 500, and 600 C) for 4 h to produce V 2 O 5 without an additional carbon shell.

Chemical and structural characterization
The chemical composition of all samples was measured by energy-dispersive X-ray analysis (EDX) using a X-Max-150 detector from Oxford Instruments in a JSM-7500F (JEOL) scanning electron microscope. Spectra were taken at ten different positions with an acceleration voltage of 10 kV.
Transmission electron microscopy (TEM) was carried out with a JEOL 2100F microscope using 200 kV acceleration voltage. Powders were dispersed in isopropanol, tip sonicated for 10 s and dropcast on a copper grid with a lacey carbon lm (Gatan) support.
Chemical composition maps were acquired by energy ltered transmission electron microscopy (EF-TEM) using a FEI Titan Themis 60-300 at an acceleration voltage of 300 kV. Electron energy loss spectra (EELS) were recorded using a Gatan GIF Quantum ERS energy lter at an energy resolution of $1.1 eV. From the acquired EELS the edge onsets of the carbon K-edge, the vanadium L-edge and the oxygen K-edge were determined and the three-window method applied to acquire the respective elemental maps. A slit width of 10 eV was utilized for the pre-and post-edge images. Since the vanadium and oxygen edges overlap, the exact same pre-edge windows were taken for both images. More information can be found in the ESI. † Thermogravimetric analysis (TGA) was carried out with a TG 209 F1 Libra system (Netzsch). Samples were heated up to 900 C in synthetic air with a rate of 20 C min À1 as well as isothermally at 450 C for 30 min. The sample mass ranged between 10 mg and 20 mg.
The nitrogen gas sorption analysis at À196 C was carried out with a Quantachrome Autosorb iQ system and the calculations were performed with ASiQwin-soware 3.0. The samples were degassed at 100 C for 1 h and subsequently heated to 150 C and kept at this temperature for up to 20 h at a relative pressure of 0.1 Pa to remove volatile molecules from the surface. The relative pressure with nitrogen was 5 Â 10 À7 to 1.0 in 76 steps. The specic surface area (SSA) was calculated using the Brunauer-Emmett-Teller equation (BET) in the linear regime of the measured isotherms, typically 3 Â 10 À2 to 2 Â 10 À1 (relative pressure), with a Quantachrome Autosorb 6B. 64 Raman spectra were recorded with a Renishaw inVia Raman Microscope equipped with a Nd-YAG laser (532 nm). A 50Â objective was used with a power of 0.02 mW at the surface of the sample. The spectra of CDC samples were recorded with 10 accumulations and 20 s acquisition time. Vanadium pentoxide and vanadium carbide spectra were recorded with 1 accumulation and 10 s acquisition as well as in two separate measurements for the full wavenumber range (0-4000 cm À1 ). All spectra were normalized and background corrected by subtracting a linear baseline.
X-ray diffractograms were collected with a Bruker D8 Discover diffractometer using Cu-K a radiation (0.154 nm) with a step size of 0.02 and a measurement time of 1 s per step. The system was calibrated with a corundum standard. The freestanding PTFE-bound electrodes were placed on a sapphire single crystal for the measurement.

Electrochemical characterization
For electrochemical characterization, free-standing, polymerbound electrodes were used. The powder materials were dispersed in isopropanol and stirred in a mortar until most of the isopropanol was evaporated and a carbon slurry was obtained. For all samples, 10 mass% PTFE (60 mass% aqueous solution, Sigma Aldrich) was added and mixed in a mortar until a dough-like paste was formed. For the electrodes with pure V 2 O 5 (VC-500) an additional conductive additive (see also the ESI, Table S1 †), in this case 25 mass% VC-CDC was added to achieve the same amount of carbon as in the core-shell electrodes. Using a rolling machine (MTI HR01, MTI Corp.) electrodes with a thickness of 60 mm (AE10 mm) were fabricated and dried overnight at 90 C and 20 mbar.
Electrochemical characterization was carried out in 1 M LiClO 4 in acetonitrile (ACN) from BASF. For all measurements, a two-and three-electrode setup was used, corresponding to full-and half-cell conguration, respectively. The electrodes were punched out to obtain discs with a diameter of 6 mm (1-3 mg) and were separated by a glass-ber disc with a diameter of 13 mm (type GF/A, GE Healthcare). The electrode/separator/electrode arrangement was compressed between two carbon-coated aluminum discs (diameter 12 mm, type Zo 2653, Coveris Advanced Coatings) using spring-loaded titanium pistons, sealed by a polyether ether ketone (PEEK) body. The cells were dried at 90 C and 20 mbar before they were put in an argon-lled glove box (MBraun Labmaster 130; O 2 , H 2 O <1 ppm). The cells were vacuum back-lled with a syringe containing the electrolyte. Full-cell measurements were carried out in an asymmetric conguration with a charge-balanced counter VC-CDC electrode by measuring the discharge capacity at 0.05 A g À1 in a half-cell setup for both, working and counter electrodes, and adjusting the mass of the VC-CDC electrode. In the half-cell setup, a ca. ve times oversized AC counter electrode was used together with a reference electrode; the latter was PTFE-bound activated carbon YP-50F from Kuraray.
Electrochemical characterization was performed using a VSP300 and VMP300 potentiostat/galvanostat from Bio-Logic in cyclic voltammetry (CV) and galvanostatic mode with potential limitation (GCPL). In the half-cell conguration a potential window from À1.2 V to +1.2 V vs. carbon was used. For data obtained with cyclic voltammetry, a scan rate of 1 mV s À1 was chosen to clearly identify the redox peaks. In galvanostatic mode with potential limitation (GCPL), specic currents up to 20 A g À1 were applied. All specic current values were normalized to the active mass of the working electrode (i.e., without the binder). Every cycle was repeated 2 times, followed by a 10 s resting time. The specic capacity of the working electrode in the half-cell conguration was calculated according to eqn (4), using the data from GCPL.
with m the active mass of the working electrode (i.e., without the binder), I the applied current, and t 1 and t 2 the start and end of discharging (from +1.2 V to À1.2 V vs. carbon), respectively. The charge efficiency was calculated by dividing the discharge by the charge capacity and is presented in percent. The specic energy (E specic ) and specic power (P specic ) of the asymmetric full-cells (two electrodes) were calculated using eqn (5) and (6): with M the active mass of both electrodes (i.e., without the binder), I the applied current, U the applied cell voltage, t 1 and t 2 the start and end of charging or discharging. Charging is in this case dened as the potential step from 0 V to 2.5 V cell voltage. Long-time stability was measured by galvanostatic cycling at 1 A g À1 from 0 V to 2.5 V cell voltage for 10 000 times. The energy efficiencies, as well as the energy retention, were calculated dividing the respective value from discharging by the value from the charging part of the curve and are presented in percent.

Structure, chemical composition, and porosity
In this study, a novel two-step process is presented, which can be used for the synthesis of vanadium oxide combined with porous carbide-derived carbon hybrid structures in a core-shell arrangement: (1) VC/CDC core-shell particles are produced by etching vanadium from the outside to the inside of VC particles with a remaining VC core. (2) The residual VC core inside the CDC matrix is calcined by oxidation in synthetic air, resulting in V 2 O 5 /CDC core-shell particles. The resulting samples are labelled VC-X70, VC-X90, and VC-X100 for a theoretical conversion of 70%, 90%, and 100%. For additional calcination the samples were labelled VC-X70-air, VC-X90-air, and VC-X100air. Complete transformation from VC to CDC (i.e., VC-CDC) was achieved by conventional chlorine gas treatment at 700 C, resulting in highly disordered carbon (Fig. 1A). Minor amounts of VC are still visible in VC-CDC (Fig. 1A) but not quantitatively signicant according to the chemical analysis (Table 1). For the synthesis of VC/CDC core-shell particles with different ratios of VC to CDC, the amount of NiCl 2 was varied for in situ chlorine formation during thermal treatment. As expected from the stoichiometric calculation presented in eqn (3), VC-X70 had the largest amount of VC in the composite with 40.2 AE 5.3 mass%. Larger amounts of NiCl 2 led to lower VC contents of 29.0 AE 7.5 mass% and 6.9 AE 2.4 mass% for VC-X90 and VC-X100, respectively (Table 1). Although a stoichiometric surplus of chlorine in VC-X100 should result in complete transformation to CDC, TEM analysis still shows a small core region of residual VC. We explain this by a reduced vanadium etching rate because of the graphitic carbon layers wrapped around the carbide domains. 37 A calcination temperature of 450 C was chosen to completely transform residual VC to vanadium oxide aer initial benchmarking of the samples via thermogravimetric analysis (see the ESI, Fig. S1 †). Aer calcination, vanadium oxide cores inside the carbon shell are clearly visible (Fig. 1D, G and J). The smallest vanadia cores are measured for VC-X100-air (d 50 ¼ 150 AE 100 nm), similar to VC-X90-air (d 50 ¼ 160 AE 110 nm), much smaller than for VC-X70-air (d 50 ¼ 410 AE 190 nm) and the calcined VC particles (VC-500 C) (d 50 ¼ 520 AE 200 nm). As seen in Table 1, EDX measurements are in qualitative agreement with TEM and show V 2 O 5 contents of 98.1 AE 6.2 mass% for the precursor VC particles oxidized at 500 C and lower amounts from 78.2 AE 2.3 mass% for VC-X70-air to 67.4 AE 4.1 mass% for VC-X100-air.
To show the structure and morphology in more detail, the two step synthesis process is illustrated in Fig. 2 using energy ltered TEM (EF-TEM) based chemical mapping. As exemplied for VC, VC-X90, and VC-X90-air, we rst see how VC particles were successively transformed to CDC from the outside to the inside retaining a VC core. Calcination of VC inside the composite particles leads to V 2 O 5 domains engulfed in nanoporous VC-CDC. EF-TEM chemical mapping conrms the distribution of carbon around the V 2 O 5 . During calcination, a partial cracking of the carbon shell is induced due to expansion of the core, as a result of the different skeletal densities of 3.36 g cm À3 for V 2 O 5 and 5.48 g cm À3 for VC. 38,39 By use of Raman spectroscopy, we can track the structural changes aer CDC synthesis (Fig. 3A) and calcination (Fig. 3C). It is shown that the VC signal intensity at 269 cm À1 decreases with successive transformation from VC to CDC (see the inset in Fig. 3A). 40 In addition to the VC signal, the Raman spectra show incompletely graphitized carbon, with the D-mode at ca. 1340-1350 cm À1 , the G-mode at ca. 1580-1600 cm À1 , and a pronounced second order spectrum at 2200-3500 cm À1 . The degree of carbon ordering in VC/CDC core-shell particles is signicantly higher than for the fully transformed CDC by chlorine treatment (VC-CDC). This is evidenced by the broader peak shapes and transition between D-and G-mode due to amorphous carbon at $1520 cm À1 , as well as the less distinct second order spectrum for the VC-CDC sample (Fig. 3A). 41,42 However, the process stemming from NiCl 2 also leads to the local emergence of metallic Ni, which, in turn, catalyzes the formation of graphitic carbon arranged in onion-like shells ( Fig. 1I and J). 43 As seen from Fig. S1 (ESI †), precise control over the calcination process temperature and duration is required to avoid either incomplete carbide oxidation or complete CDC burn-off. For this reason, we conducted rst oxidation experiments only on the initial VC powder. The increase in mass of the initial VC particles, without prior NiCl 2 treatment, is $40 mass% (ESI, Fig. S1 †) for an annealing at 600 C, leading to crystalline V 2 O 5 (Fig. 3B). However, using this calcination temperature, more than 50 mass% of the CDC was burned-off. Therefore, the synthesis was repeated at a lower temperature of 450 C for a duration of 30 min. With this annealing process, the same mass increase ($40 mass%) was achieved, expecting the same vanadia phase, and only $20 mass% of the CDC material was removed. Synthesis at 450 C for 30 min results in a mixture of V 2 O 5 , 44 amorphous V 2 O 5 , 45 and VO 2 (ref. 46) because of the large size of the precursor VC particles ($1 mm) (Fig. 3B). The particles with the VC/CDC core-shell arrangement present much smaller cores than the initial VC particles (Fig. 1B), which leads to the formation of V 2 O 5 already at 450 C applied for 30 min (Fig. 3C). These ndings are supported by X-ray diffraction (XRD), shown in the ESI, Fig. S2. † Transformation from VC to VC/CDC is clearly evidenced by the emergence of the (002) graphite signal at 26.4 2q for VC-X70 and the continuous increase of its intensity up to the highest degree of transformation for VC-X100 (ESI, Fig. S2A †). The fully transformed VC-CDC presents a less distinct carbon signal, coming from nanocrystalline and amorphous carbon, in agreement with Raman spectroscopy. Aer calcination, only V 2 O 5 is iden-tied (ESI, Fig. S2B †).
Highly nanoporous materials with high specic surface areas typically result from chlorine treatment of VC (i.e., VC-CDC). 38 The nitrogen gas sorption isotherm of VC-CDC shows a characteristic type I(b) shape related to microporous materials with pore sizes <2.5 nm (ESI, Fig. S3A †). Both microand mesopores are found for the partial transformation of VC with type I(b) isotherms, related to the internal microporosity, combined with type IV(b) isotherms related to mesopores with a small H4 hysteresis (Fig. S3B †). 47 The combination of microand mesopores has been reported in other studies on carbon core-shell structures. 36 The BET surface area increases from 120 m 2 g À1 for VC-X70 to 577 m 2 g À1 for VC-X100, and to 1466 m 2 g À1 for VC-CDC ( Table 2). Calcination of VC/CDC core-shell particles does not change the pore type (micro-and mesopores), but leads to oxidation of the high surface area carbon, with consequent decrease in surface area. For VC-X100-air and VC-X90-air, the surface area collapses to 61 m 2 g À1 (À89%) and 124 m 2 g À1 (À38%), respectively, due to oxidation of the CDC shell. The surface area of VC-X70-air is even larger than before calcination with 310 m 2 g À1 (+158%; Table 2). The high content of VC in VC-X70 (40 mass%), which is turned into the lower density V 2 O 5 , has a stronger impact on the surface area than the CDC, compared to VC-X90 and VC-X100 with lower VC contents.

Electrochemical performance
As seen from the previous section, partial chlorination of VC to VC/CDC and successive calcination to V 2 O 5 /CDC in a facile twostep approach leads to highly crystalline V 2 O 5 embedded graphitic and porous carbon. This combination is promising for electrochemical hybrid energy storage applications, which was benchmarked in our study in 1 M LiClO 4 in ACN as the electrolyte. Cyclic voltammetry of VC-CDC using the half-cell setup shows a nearly ideal double-layer capacitor behavior with a rectangular CV shape without any redox reactions. This behavior is expected from nanoporous carbons in general and CDC materials in specic. 6 VC-X70-air is characterized by double-layer capacitance on the positive side between +0.5 V and +1.2 V vs. carbon. In the negative range, a very broad redox peak emerges (Fig. 4B and ESI, Fig. S4B †). This broad peak correlates with the larger V 2 O 5 core of VC-X70 compared to the nanotextured V 2 O 5 particles in VC-X90-air and VC-X100-air. Due to longer ion pathways in the V 2 O 5 crystal, the system presents slower redox reactions and broader redox peaks.
In stark contrast to the capacitor-like behavior of VC-CDC, the CVs of VC-X90-air and VC-X100-air are characterized by several, very sharp redox peaks from Li-ion intercalation. 30,48,49 The occurring electrochemical Li + insertion process for V 2 O 5 can be expressed by eqn (7): 50 with x as the mole fraction of Li-ions.  Typically, a set of two peaks for inter-and deintercalation is observed between +2 and +4 V vs. Li/Li + or À1.2 V and +0.8 V vs. carbon. 48 The structural changes during Li + intercalation into V 2 O 5 involve several phases like a, 3, d, and g for x # 2, with three voltage plateaus at 3.4, 3.2, and 2.3 V vs. Li/Li + (À0.7 V, 0 V, and +0.2 V vs. carbon) (Fig. 4B). 48 The plateaus seen from the galvanostatic curves (Fig. 4B) and the pairs of redox peaks in the CVs (Fig. 4A) belong to the two-phase regions a/3, 3/d, and d/g. Larger lithium contents (x > 2) for potentials below 2 V vs. Li/Li + (À1.2 V vs. carbon) would lead to irreversible reactions and the formation of the g-phase. 48,51,52 A detailed discussion of the structural changes is already given in ref. 48 and 51-53. The rst set of redox peaks emerges between À0.25 V and +0.8 V vs. carbon for inter-and deintercalation of the rst Li-ion. A second lithiation step occurs between À1.2 V and 0 V vs. carbon (ESI, Fig. S4A †). For potentials above +0.8 V vs. carbon, only doublelayer capacitance of the porous carbon is encountered. Calcined VC without the core-shell structure, but the same V 2 O 5 phase, shows very broad sets of redox peaks in the CV compared to the core-shell structure. It is expected that longer ion intercalation paths of the larger size of V 2 O 5 domains in VC-500 C result in slower redox reactions of the electrode, similar to VC-X70-air. In contrast, the core-shell structures present much smaller V 2 O 5 particles due to partial transformation into porous carbon.
Galvanostatic charge/discharge cycling (GCPL) in half-cells was further employed for assessing the power handling ability of the electrode materials. As seen from Fig. 4B, GCPL curves of VC-X100-air and VC-X90-air exhibited several voltage plateaus between À1.2 V and +1.2 V vs. carbon, in agreement with the redox peaks in the CVs. Nearly ideal double-layer behavior with triangular-shape GCPL curves was observed for VC-CDC. The Li + intercalation and deintercalation leads to phase transitions during cycling and is kinetically enhanced for nanoscopic materials because of better accessibility of intercalation pathways. 54 Therefore, smaller V 2 O 5 core sizes of V 2 O 5 /CDC coreshell samples show superior rate handling compared to VC-500 with a more than 2-times larger V 2 O 5 particle size (Fig. 4C). This is especially seen for VC-X100-air with more than 100 mA h g À1 capacity at 5 A g À1 (equivalent to a rate of ca. 28C).
The lower vanadium oxide content of VC-X100-air limits the maximum capacity at low rates. The highest value is achieved  for VC-X90-air with more than 300 mA h g À1 . Typically, in the literature capacity values are normalized to the active mass, in this case the mass of V 2 O 5 in the composite electrode. Using this way of normalization, the highest capacity for VC-X90-air is 415 mA h g À1 . Even for higher specic currents of 5 A g À1 the specic capacity reaches more than 100 mA h g À1 (comparable to a C-rate of 12C). The normalized values are given as an inset in Fig. 4D for comparison. Cycling in the potential window of À1.2 V to +1.2 V vs. carbon is expected to be completely reversible, 48,51,52 and this behavior was conrmed for all V 2 O 5 / CDC core-shell particles with a coulombic efficiency of ca. 100% (Fig. 4D). The VC particles calcined at 500 C with much larger particle sizes (Fig. 1B) and without CDC shells, present a lower charge efficiency of less than 80% at 20 A g À1 . This might be the result of V 2 O 5 degeneration due to kinetic limitations. 30 Electrochemical properties on a device level were measured using an asymmetric full-cell setup with nanoporous VC-CDC as the counter electrode. Typically, symmetric setups for electrode  materials with battery-like behavior (e.g., intercalation/deintercalation) are disadvantageous due to their polarization dependent redox-reactions. Yet, several studies used different types of V 2 O 5 or V 2 O 5 /carbon composites, such as electrospun V 2 O 5 nanobers, 55 V 2 O 5 /MWCNT core-shell aerogels, 56 and V 2 O 5 /graphene aerogels 57 with the symmetrical full-cell setup. However, in some of these studies, negative and positive polarization in the full-cell arrangement was used (e.g. À1 V to +1 V in an aqueous electrolyte), instead of 0 V to 1 V. In these cases the capacitance and specic energies are much larger reaching misleading high values of 87 W h kg À1 in aqueous 1 M Na 2 SO 4 (ref. 56) and 78 W h kg À1 in organic 1 M LiClO 4 in PC. 55 Enhanced electrochemical performance can be obtained by use of an asymmetric two-electrode setup (Fig. 5). Studies so far have shown, for aqueous systems, the highest specic energy of 29 W h kg À1 (Table 3) for hydrothermally prepared V 2 O 5 $0.6H 2 O with activated carbon as the counter electrode. 58 Due to the improved voltage stability of organic electrolytes, much higher energies can be achieved, ranging from 15 to 76 W h kg À1 with MnO 2 /rGO, AC fabric, or rGO as the counter electrode in an asymmetric arrangement. 34,59,60 In this study, an asymmetric full-cell setup with V 2 O 5 /CDC on the negative polarization side and charge-balanced VC-CDC on the positive polarization side was used. VC-X90-air as a negative electrode material presents superior performance with a specic energy in excess of 80 W h kg À1 for charging and more than 50 W h kg À1 for discharging ( Fig. 5A). Even at higher rates, with a specic power of $6.7 kW kg À1 , the specic energy is 27 W h kg À1 (12 W h kg À1 for discharging). In agreement with half-cell measurements, VC-X100air shows a lower specic energy of 66 W h kg À1 with comparable power handling because of the lower V 2 O 5 content. The energy and power performance of this core-shell system is highly competitive with current literature values and is, to the best of our knowledge, the only core-shell system reaching 50-80 W h kg À1 specic energy when using asymmetric full-cell measurements and proper application of the cell voltage. 61 The high attractiveness of this material for electrochemical energy storage is underlined by the high cycling stability with an energy retention of more than 80% and 70% over 10 000 galvanostatic cycles for VC-X90-air and VC-X100-air, respectively, measured at 1 A g À1 (Fig. 5B). The nearly ideal charge efficiency of the half-cell measurements (100%) translates to a still attractive energy efficiency of the asymmetric devices between 60 and 80% for low rates. For comparison, standard doublelayer capacitors using activated carbon electrodes only reach slightly higher values of 80-95% dependent on the electrode. 62

Conclusions
This study, for the rst time, demonstrates a two-step in situ synthesis of hybrid particles consisting of porous carbon and redox-active V 2 O 5 by solely using VC as the precursor. The Table 3 Literature values for full-cell measurements (two-electrodes) in symmetric and asymmetric fabrication. OLC: onion-like carbon; CNT: carbon nanotube; NW: nanowire; rGO: reduced graphene oxide; ACC: activated carbon fabric; NS: nanosheet; PPy: polypyrrole; particles present a V 2 O 5 /CDC core-shell structure with nanoscopic V 2 O 5 cores, completely accessible to electrolyte ions and interconnected by highly conductive carbon. High specic energies of 50 and 80 W h kg À1 for discharging and charging, respectively, were measured in a facile asymmetric full-cell setup, while simultaneously retaining ultrafast cycling. The asymmetric cell design with CDC on the positive polarization side leads to superior stability over 10 000 cycles and comparably high energy efficiency of 60-80% at low rates. The novel synthesis procedure can easily be scaled up by means of larger precursor mass and may also be very attractive for other redoxactive compounds.