Electrochemical performance of carbide-derived carbon anodes for lithium-ion batteries

Author's Accepted Manuscript

Electrochemical Performance of Carbide-DerivedCarbon Anodes for Lithium-Ion Batteries

Sun-Hwa Yeon, Kyu-Nam Jung, Sukeun Yoon,Kyoung-Hee Shin, Chang-Soo Jin

PII: S0022-3697(13)00093-0DOI: http://dx.doi.org/10.1016/j.jpcs.2013.02.028Reference: PCS7020

To appear in: Journal of Physics and Chemistry of Solids

Received date: 10 December 2012Revised date: 5 February 2013Accepted date: 21 February 2013

Cite this article as: Sun-Hwa Yeon, Kyu-Nam Jung, Sukeun Yoon, Kyoung-Hee Shin,Chang-Soo Jin, Electrochemical Performance of Carbide-Derived Carbon Anodes forLithium-Ion Batteries, Journal of Physics and Chemistry of Solids, http://dx.doi.org/10.1016/j.jpcs.2013.02.028

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Electrochemical Performance of Carbide-Derived

Carbon Anodes for Lithium-Ion Batteries

Sun-Hwa Yeon*, Kyu-Nam Jung, Sukeun Yoon, Kyoung-Hee Shin, and Chang-Soo Jin

Korea Institute of Energy Research, 102, Gajeong-ro, Yuseong, Daejeon, 305-343, Republic

of Korea.

Abstract

Carbide-derived carbons (CDCs), part of a large family of carbon materials derived from carbide, are attractive

for energy-related applications, such as batteries, supercapacitors, and fuel cells. Pore textures (micro-, meso-,

and macro-pores) and structures (from amorphous to highly ordered graphite) of CDCs can be controlled by

changing the synthesis conditions and carbide precursor. Adequate control of the carbon structure, and the

porosity in terms of application as an anode can be exploited to maximize the electrochemical capacity in a

lithium ion batteries. In this study, the use of CDC as anodes by chlorine treatment of B4C and TiC7N3 in a

synthesis temperature range from 600 to 1200 oC has been explored. The discharge capacity of TiC7N3-CDC

reaches the highest value, 462 mA h g-1, at 100 cycles, which is 25 % higher than the theoretical capacity of

graphite (375 mA h g-1). B4C-CDC meanwhile affords a value of 453 mA h g-1 at 100 cycles. These results show

that B4C-CDC and TiC7N3-CDC have excellent potential as the negative electrode in Li battery applications, and

* Corresponding Author: Tel: +82-042-860-3271. Fax: +82-042-860-3134. E-mail address: ys93@kier.re.kr (S.-H. Yeon).

can be exposed to a practical low synthesis temperature range of 600 oC ~ 1200 oC. B4C-CDC and TiC7N3-CDC

can also provide 2-3 times better performance than existing graphite or hard carbon for lithium battery systems.

Keywords

Carbide-derived carbon; Li-ion battery; Anode; Porous carbon; Crystallinity

1. Introduction

Carbon is the basis of all life on earth in that entire branches of chemistry are associated with its

compounds. One of the most recent applications of carbonaceous material is as anode electrodes in rechargeable

lithium (Li) -ion batteries, a key technology that play an important role in transportation and renewable energy

storage [1-3]. They possess desirable features such as higher specific charges and more negative redox potentials

than other candidate materials (e.g., metal oxide, chalcogenides, and polymers), in addition to better cycling

behavior than Li alloys due to their dimensional stability [4]. Also, carbonaceous anodes undergo little

formation of Li dendrites on the surface of the anode, thereby providing high reliability and safety [5].

The structures of carbon, which are ideally suitable for use in the anode of Li-ion batteries, are

classified as graphitic or non-graphitic (disordered). For highly crystalline graphitic carbons of the Li-graphite

intercalation compounds, LixCn, the maximum Li content is one Li guest atom per six carbon host atoms (n� 6

in LiCn or x�1 in LixC6) at ambient pressure. While the intercalation reaction occurs at arm-chair and zigzag

faces of prismatic surfaces, intercalation through the basal planes is possible at defect sites only [6]. However,

the electrochemical performance of carbon electrodes varies largely with their morphologies and characteristics

which depend on the nature of the carbon precursors and heat-treatment conditions. The disordered (non-

graphitic) carbons possibly possessing hydrogen content, in the form of surface functional groups, or high

porosity are characterized by an important irreversible capacity and polarization between charge and discharge,

otherwise known as hysteresis. These carbons show striking correlations between functional groups along with

their porosities and electrochemical capacities [7-9]. One of the main reasons for these features is the formation

of a Solid Electrolyte Interface (SEI) during the first reduction cycle at ca. 0.8 V vs Li. The electrolyte is

decomposed with the formation of LiF, Li2O, LiOH, Li2CO3, ROCO2Li, and so on [9]. There is another

hypothesis that the capacities of carbons with suitable pores for the accommodation of lithium species would

surpass the theoretical limit (372 mA h g-1 for LiC6 composition), because the pores in the carbons would

contribute to the charge-discharge of lithium species for lithium-ion batteries with high porosities [10]. In the

application of porous carbon as an anode of Li-ion batteries, a new family of exfoliated transition metal carbides

and carbonitrides called “MXene” (where M is an early transition metal and X is C and/or N – that includes

Ti3C2, Ta4C3, TiNbC, and (V0.5,Cr0.5)3C2) has been introduced, which was synthesized by wet HF treatment of

Al-containing MAX phases (Ti3AlC) at room temperature. The produced Ti2C-based material (MXene) has

resulted in 225 mAhg-1 at C/25, suggesting that MXenes are promising as anode materials for Li-ion batteries

[11, 12]. Meanwhile, 6H-SiC (0001) has received strong interest as a substrate for epitaxial growth of graphene

following graphitization at 1350 °C in an ultrahigh vacuum, because the graphene material brought substantial

enhancement of the electrochemical lithiation capacity of surface graphitization via Li diffusion through

graphene into the surface (6�3 × 6�3)R30°, region leading to Li�Si bonding [13, 14]. Li is present in a

stoichiometry as high as (1 ± 0.2): 1 ratio of Li to Si, and this apparent stoichiometry corresponds to a Li-ion

capacity of 670 ± 130 mA h g-1, which is approximately double that of the graphite anodes currently used in Li-

ion batteries. This work has shown that a relatively straightforward surface treatment is sufficient to convert

doped SiC from an inactive, inert substrate into an electrochemically active host for Li bonding [13].

Carbide-derived carbons (CDCs), which can be produced by chlorine treatment of crystalline carbides

such as TiC, SiC, and Mo2C, possess high porosities, tunable pore size, and narrow pore size distribution. The

chemical etching method by chlorine gas generates micro- and meso-porosities that can afford large adsorption

capacities of target molecules matching the pore sizes due to the controlled pore size and high specific surface

areas [15-17]. Also, the graphitic crystallinity of the CDCs can be controlled by varying the synthesis conditions

and carbide precursors. Suitable matching between the ion size of the used electrolyte and the controlled pore

size of CDC has been shown to significantly increase the electrode capacitance of a supercapacitor [18],

especially when the pore diameter approaches the size of the desolvated ions [19]. However, there has been little

research on CDC anode applications for Li-ion batteries, despite that the controlled crystalline structure and

pore texture of CDC greatly influence the performance of the battery system.

In this work, we demonstrate significantly improved electrochemical capacity of mesoporous CDCs

produced by B4C and TiC7N3 carbide precursors as lithium negative electrodes. The binary and ternary carbides

lead to mesopore CDC with a highly-ordered graphite structure at the Cl2 temperature range of 600 oC ~ 1200

oC. For electrochemical capacity as negative electrodes, the discharge capacity of TiC7N3-CDC reaches the

highest value, 462 mA h g-1 at 100 cycles, which is 25 % higher than the theoretical capacity of graphite, 375

mA h g-1. Also, B4C-CDC offers a value of 453 mA h g-1 at 100 cycles.

2. Experimental

Carbide-derived carbon (CDC) powders were produced by the chlorine of B4C (particle size 1.7 �m),

TiC7N3 (MKNano Co., 97+%, particle size 150 nm), and TiC powder (particle size 1-4 �m). This precursor was

placed in a horizontal tube furnace, purged in argon flowand heated to temperatures between 600 C and 1200

C under flowing chlorine (10–15 cm3 min-1) for 3 h. CDC powders treated at 600 C ~ 1200 C were then

annealed at 600 C for 2h under flowing hydrogen to remove residual chlorine and chloride trapped in pores

Previous studies have showed that this procedure was not necessary at higher treatment temperatures [19, 20]. In

CDC preparation, the most applicable temperature ranges are 200 oC- 1200 oC. While the CDCs synthesized at

low temperature, 200 oC ~ 400 oC, usually show highly disorder amorphous structure, graphite fringe starts to

occur at the high temperature range from 600 oC ~ 1200 oC. Therefore, in this study, the temperature range from

600 oC ~ 1200 oC has been chosen for the anode applications of the graphitizable CDCs. Commercial MSP-20

carbon (Kansai & Chemicals Co.), MCMB (Osaka Chemical Co), hard carbon were obtained. Air activation

(500 oC and 5h) for MSP carbon was performed in order to remove some surface functional groups, which can

cause SEI layer formation during charge-discharge operation, and to control the porosities. As-received MSP

carbon is already high purity carbon with small oxygen functional groups of the MSP surface (EDS data showed

only C and O and no other impurity). After activation, the change of porosity such as reduction of micropores

occurred by carbon loss. The cathode electrode material was prepared by adding Super-P (10 wt%) as a

conductive additive and polyvinyl fluoride (PVDF, 10 wt%) as a binder in all cases.

To evaluate the electrochemical performance of carbon composites, CDCs, MCMB, and hard carbon,

2032-type coin cells (MTI) were fabricated. The prepared sample was pressed onto copper substrate as a

working electrode, including binder and conductive additives. Li was used as the counter electrode. The

separator used CelgardR 3501. The electrolyte, M LiPF6/ethylene carbonate (EC) + diethyl carbonate (DEC)

(50:50 vol%) was used.

Gas adsorption analysis was performed using BELSORP-Max MP (BEL Japan Inc.) with N2 adsorbate

at -196 oC for the porous carbon. Approximately 70 mg of the carbon was evacuated at 5 mTorr at 300 oC for 16

h. The measured pressure (P/P0) was 0.05-0.999. The Brunauer-Emmett-Teller specific surface area (BET SSA)

and PSD (pore size distribution) were calculated by using the BET theory and the Barrett-Joyner-Halenda (BJH)

method based on the adsorption branches of the isotherms. Micropore analysis by the MP method was based on

the t-curve. The specific pore volumes were measured at a relative pressure of 0.95. Also another PSD method,

the non-local density functional theory (NLDFT), was tried [21]. The NLDFT model assumes slit-shaped pores

with uniformly dense carbon walls; the adsorbate is considered as a fluid of hard spheres [21]. TEM samples

were prepared by dispersing each sample in ethanol and placing the solution over a copper mesy grid with a

carbon film. A TEM study was performed using the Tecnai F20 microscope at 200 kV. SEM was performed

using a Zeiss Supra 50VP scanning electron microscope equipped with an energy dispersive spectroscopy (EDS).

XRD analysis was done using a Rigaku diffractometer with CuK� radiation (� = 0.154 nm) operating at 30 mA

and 40 kV. XRD patterns were collected using step scans with a step size of 0.01° (2) and a count time of 2s

per step between 5 (2) and 80 (2). Samples were analyzed by micro-Raman spectroscopy (Renishaw 1000)

using an Ar ion laser (514.5 nm) at 20 X magnification (NA: 0.75: ~2 �m spot size), and <2 mW power. TGA

was performed by the TA Q-500 TGA system (TA instruments). The samples were loaded onto alumina (Al2O3)

sample pans of around 10 mg and heated to 800 oC ~ 900 oC at 10 oC min-1 under an air stream of 50 cm3 min-1.

3. Results and Discussion

Figs. 1-3 and Table 1 provide information on the porosity of various CDCs and two commercial

carbons of MCMB and hard carbon. In the N2 adsorption isotherms shown in Fig. 1, marked hysteresis was

observed in (a) B4C-CDC, (b) TiC7N3-CDC, and (d) TiC-CDC@1200 oC Cl2, which confirms the presence of

mesoprosity corresponding with typical type IV isotherm characteristics. Also, air-activated MSP carbon

(including MSP carbon (Fig. 1c)) and TiC-CDC @ 1000 oC (Fig. 1d) showed type I isotherms for the

microporosity. On the other hand, hard carbon (Fig. 1e) and MCMB (Fig. 1f) exhibited nonporous solids of type

III isotherm. Fig. 2 shows the estimated pore size distributions (PSDs) of the materials by using the MD method

for a micropore range (0.4 nm < pore width < 2.0 nm) and the BJH method for mesopore and macropore ranges

(2.4 nm < pore width < 100 nm). Average pore sizes increased with increasing Cl2 temperature, consistent with

the general pattern for various carbide precursors processed at high temperatures [15, 22, 23]. In B4C-CDCs, the

PSD is quite narrow with a peak centered at around 0.8-0.9 nm (micropore) at 800 oC Cl2, and shifts to larger

peaks centered at around 1.5 nm with an increasing Cl2 temperature of 1000 oC (Figs. 2(e and f)). These PSD

patterns are similar to those of typical TiC-CDC with increasing Cl2 temperature (Figs. 1(c and d)). However, in

B4C-CDCs, a larger pore volume is observed in a pore range between 10 nm and 100 nm (meso- and macro-

pores, respectively) relative to that of TiC-CDCs. The isotherm pattern of B4C-CDC (mesoporus type IV,

precursor particle size ~1.7 �m) shown in this study is slightly different from that of the B4C-CDC shown in a

previous study (microporous type I), which may necessitate the use of a large particle size precursor (~ 6 �m)

[24]. In the case of TiC7N3-CDC, chlorine treatment at 800 oC shows a bimodal PSD with peaks centered around

0.5 nm (primary micropore) and 1.4 nm (secondary micropore), subsequently showing a broad mesopore and

macropore range from 3 nm to 100 nm (Figs. 2(g and h)). Chlorine treatment at 1000 oC Cl2 led to a smaller

micropore and mesopore volume than at 800 oC with a broad pore width in a range from 3 nm to 100 nm. The

PSDs of MSP carbon are quite narrow with a peak at 0.8 nm, following the typical microporous PSD pattern.

The activated MSP by air at 500 oC showed a slight loss of micropore volume without any changes in the peak

center compared to the MSP sample. Table 1 shows small pore volumes (< 0.01 cm3 g-1) by the MP and BJH

methods with BET SSA, 2~4 m2 g-1 in hard carbon and MCMB.

Another PSD analysis method, the NLDFT model (N2 - carbon at 77 K based on a slit-pore model),

was used, where the applicable pore diameter range is 0.3 nm ~ 40 nm, as shown in Fig 3. MSP and MSP 500 oC

activated by air showed a similar PSD pattern with a peak centered at 1.2 nm. In the case of CDCs, TiC@1200

oC and B4C@1000 oC led to a bimodal pattern with micropores (0.5 nm ~2 nm) and mesopores (2 nm ~ 4 nm).

In TiC@1000 oC and B4C@800 oC, micropores in a range of 0.5 nm ~ 2 nm are dominant. However, all PSDs of

TiC7N3-CDC prepared at 800 oC and 1200 oC showed mesopores (3 nm ~ 5 nm) with a peak centered at 4 nm.

XRD patterns of the various carbons are shown in Fig. 4. Fig. 4a shows the XRD results of B4C-CDCs,

which were almost fully converted into CDC after chlorine treatment with no B4C crystalline peaks. The B4C

carbide precursor has a rhombohedral lattice (space group R-3m) with crystal parameters a = 0.56003 nm and c

= 1.2086 nm. Its structure can be described as an �-rombohedral boron, in which almost regular B12 icosahedra

are located in the corners of a rhombohedron and with C atoms in the three-atom chain and as part of the

icosahedra in a solid solution range [25]. After chlorine treatment at 800 oC and 1000 oC, the Bragg reflection at

around 2 ~ 26º corresponds to the diffraction from the (002) planes of graphite. Titanium carbonitride (Ti,C, N)

can be prepared by TiC and TiN. Both TiC and TiN have a cubic space group of Fm3m, where the corner of the

face-centered-cubic (fcc) lattice formed by C atoms (or N atoms in the case of TiN) is located at the point

(1/2,0,0) of the fcc superlattice formed by Ti atoms. Since TiC and TiN are isomorphous, the carbon atoms on

the TiC superlattice can be replaced by nitrogen atoms in any proportion (and vice versa). Therefore, a

continuous series of solid solutions can be formed. The lattice parameter of TiN (4.240 Å) is slightly smaller

than that of TiC (4.320 Å); therefore, with addition of TiN in TiC, the lattice constant will decrease [26]. While

no produced TiC-CDCs show a residual TiC precursor at 1000 oC or 1200 oC (Fig. 4c), some residual TiC7N3

was identified with chlorine treatment of TiC7N3 at both 600 oC and 1200 oC, as shown in Fig. 4b. The Bragg

reflection at around 2 ~ 26º, which corresponds to diffuse scattering from disordered amorphous carbon, shows

a large full width of half maximum (FWHM), indicating the amorphous nature of TiC7N3-CDC. In XRD profiles

of hard carbon and MCMB, two broad diffraction peaks were observed near 25 and 43o, which correspond to

(002) and (100) planes, respectively, indicating a low degree of crystallinity.

Raman spectra (Fig. 5) were normalized to the intensity of the G-band at 1582 cm-1. Graphitic

materials are characterized by bands labelled D (~1356 cm-1, disorder mode), G (~1579 cm-1, in-plane

vibrational mode), and 2D (~2706 cm-1, second order mode). B4C-CDCs (Fig. 5a) exhibited a broad D band

typical of amorphous disordered carbon at 800 oC Cl2. The B4C-CDCs become more disordered with an

increasing Cl2 temperature up to 1000 oC, which exhibits a larger D than G peak intensity. In the case of TiC7N3-

CDCs (Fig. 5b), while the 600 oC spectrum reveals typical amorphous disordered carbon, the 1200 oC spectrum

exhibits narrow G and D peaks along with a very intense G band, which are attributable to disordered graphite.

In the second-order region, three combination modes 2689 (2D), 2938 (D + G), and 3219 cm-1 (D`) were

observed. These Raman trends of TiC7N3-CDCs with an increasing temperature are similar to typical TiC-CDCs

as shown in Fig. 5c. Also, hard carbon and MCMB show typical D and G peaks in disordered carbon.

TGA was performed in an air environment to determine more precise purity of the prepared CDCs.

Fig. 6 show the TGA curves of the TiC-CDCs prepared at 1000 oC and 1200 oC, B4C-CDCs prepared at 1000 oC

and 800 oC, and TiC7N3-CDCs prepared at 1000 oC and 1200 oC. As shown in Fig. 6, the TiC-CDCs started to

oxidize at 450 oC and were burned completely at 500 oC ~ 600 oC, as indicated by the ~ 100 % weight loss at

that temperature. B4C-CDCs started to oxidize at 500 oC ~ 600 oC, which is the same temperature range as TiC-

CDCs, and were burned at 600 oC ~ 700 oC, indicating unreacted precursors or impurities of 2~3 %. TiC7N3-

CDCs started to oxidize at more than 600 oC and were burned at ~ 700 oC. However, these samples showed

unreacted carbides of 27 ~ 32 %, as shown in the XRD results of Fig. 2b.

Fig. 7 shows high-resolution TEM images of the B4C and TiC7N3 carbide precursor synthesized at 800

oC ~ 1200 oC, respectively. As-received TiC7N3 nanoparticles are uniform idiomorphic cubes that are often

surrounded by a few layers of graphitic carbon. Chlorine treatment at 600 oC and subsequent H2 annealing at

600 oC results in a completely amorphous carbon (Figs. 7 (a and b)) with graphite fringes covering the edge part

of the particle, which already existed before chlorine treatment as the graphitic layer surrounds the TiC7N3

nanoparticles. The TiC7N3 precursor contains amorphous carbon, highly ordered sheets of graphite with an

interplanar distance of 0.3–0.4 nm, and crystalline structures based on TiC and TiN. After chlorine treatment, the

TiC7N3-CDC formed micropores of 0.4-2.0 nm, mesopores of 2-50 nm, and macropores 100 nm in diameter,

caused by inter void and inter particle spacing due to the conformal transformation of the carbide-to-CDC. The

removal of the atomic Ti and N from the TiC7N3 molecule leads to larger pores than those derived from the TiC

precursor. The mesopores and macropores originated from etching TiN, while etched TiC leaves behind a micro-

porous carbon phase. A high temperature of 1200 oC Cl2 shows numerous lots graphite fringes surrounding the

nanoparticles. In the inset of Fig. 7, the graphene exhibits a typical wrinkled structure with corrugation and

scrolling intrinsic to graphene [27]. The B4C-CDC prepared at 800 oC and 1000 oC consists of a major

proportion of highly disordered and interlaced graphene layers. This type of carbon nanostructure with

disordered, curved, and interlaced graphene layers can be described as being composed of imperfect fragments

(Figs. 7(c and d)). The images of TiC-CDCs at 600 oC and 1200 oC (Figs. 8a and 8b, respectively) show typical

disordered carbon with some porosity, in which the TiC-CDC treated at 1200 oC contains amorphous carbon and

highly ordered curved sheets of graphite with an interplanar distance of 0.3-0.4 nm. On the other hand, the pore

texture of commercial MCMB and hard carbon (Figs. 8c and 8d.) shows that the basic structural units, made of

stacks of 2-3 nanometric graphene layers, are strongly misoriented and not completely uniform.

The cycling behaviors of various porous carbons as negative electrodes, such as TiC-CDCs, TiC7N3-

CDCs, B4C-CDC, and MSP carbon activated at 500 oC by air, are presented in Fig. 9. TiC-CDCs can tailor

various pore textures and structures according to the reaction temperatures, and show microporosity with a

disordered or amorphous structure at 600-1000 oC Cl2 and microporosity + mesoporosity with highly

crystallinity at 1200 oC Cl2, respectively [15, 23].

In TiC-CDCs treated at 1000 oC, the discharge capacities were 231 mA h g-1 at the 100th cycle,

indicating large irreversibility of the first and second cycles of 1352 mA h g-1 and 617 mAhg-1, respectively. On

the other hand, TiC-CDC treated at a high temperature of 1200 oC showed the largest discharge capacity, 317

mA h g-1, among TiC-CDCs, where the first and second discharge capacities reached 1426 mA h g-1 and 504 mA

h g-1, respectively. The 100th discharge capacity at TiC-CDC@1200 oC was 317 mA h g-1, which was about 84 %

of the theoretical capacity of 376 mA h g-1. This largest 100th discharge capacity of TiC-CDC@1200 oC among

various TiC-CDCs may have been caused by large amounts of highly ordered sheets of graphite created during a

high chlorine treatment reaction of 1200 oC. These results closely agree with those of previous studies on TiC-

CDC anodes for Li-ion batteries [28]. The specific charge and discharge capacities in a TiC-CDC 1000 °C/Li

metal half-cell were found to reach 275 mA h g-1 after the fourth cycle at a 0.125 C rate. In the case of TiC-CDC

1200 oC, the specific discharge capacity reached ~170 mA h g-1 in the fifth cycle at 1C. In both samples, a large

irreversible capacity still occurred in the first cycle [28].

Microporous air-activated (500 oC) MSP carbon with a large BET SSA area (~1800 m2 g-1) also

showed initial large first and second discharge capacities of 1512 mA h g-1 and 954 mA h g-1, respectively.

However, the discharge capacity decreased to 300 mA h g-1 in the 40th cycle, and finally remained at 253 mA h

g-1 up to the 100th cycle. These micro-pore textures usually possess a disordered amorphous structure with low

crystallinity, which can cause deterioration of the electrochemical capacity.

Unlike the TiC carbide, CDCs derived from B4C and TiC7N3 carbides possess a mesoporous type pore

texture after the low chlorine treatment process. In mesoporous B4C-CDCs and TiC7N3-CDCs, enhanced charge-

discharge capacities were shown relative to microporous TiC-CDCs. B4C-CDC @ 800 oC Cl2 indicated a high

discharge capacity of 1176 mA h g-1 at the first cycle, and decreased to 586 mA h g-1 at the second cycle, which

is similar to discharge patterns shown in microporous CDCs. However, the specific capacity of the B4C-CDC @

800 oC Cl2 stabilized after 10 cycles was about 380 mA h g-1, which is close to the theoretical capacity of

graphite. The reversible capacity of B4C-CDC @ 800 oC Cl2 became constant from the 10th cycle, which

actually increased slightly with cycling. The 100th discharge capacity reached 437 mA h g-1 which is about 1.2

times the theoretical capacity (372 mA h g-1) of graphite. Also, B4C-CDC @ 1000 Cl2 showed capacity of about

353 mA h g-1. In the other mesoporous type of TiC7N3-CDCs, the discharge capacity of TiC7N3-CDC @ 600 oC

Cl2 was 1557 mA h g-1 at the first cycle and 671 mA h g-1 at the second cycle. The 100th cycle was 462 mA h g-1.

In the TiC7N3-CDC@ 1200 oC Cl2, a noticeable graphitic carbon layer and amorphous carbon were present, but

a large amount of unreacted TiC7N3 was found at high temperature. The first discharge capacity of TiC7N3-

CDC@ 1200 oC Cl2 was 741 mA h g-1. The second discharge capacity decreased to 280 mA h g-1 with increasing

305 mA h g-1 at 100th cycles, which is about 34 % lower than that of TiC7N3-CDC@ 600 oC Cl2. The low

capacity of TiC7N3-CDC@ 1200 oC Cl2 can be caused by the existence of TiC or TiCN nanoparticle residue by

incomplete conversion of CDC at a high Cl2 temperature. The specific values in mesoporous carbon (B4C-CDCs

and TiC7N3-CDCs at 10 cycles) stabilize faster than those of microporous carbon (for example, MSP @ 500 oC

air activation at 40 cycles and TiC-CDCs at 20 cycles, respectively). In order to compare the discharge capacity

of CDCs with a commercial negative electrode, the electrochemical performances of MCMB and hard carbon

were conducted. The discharge capacity of the 100th cycle corresponds to 240 mA h g-1 for MCMB and 265 mA

h g-1 for hard carbon, respectively, showing larger values than the first cycle (217 mA h g-1 and for MCMB and

175 mA h g-1 for hard carbon, respectively). The lower capacity in the initial cycles can be attributed to an

initially higher polarization due to the carbon layers obtained from the carbonization. After further cycling, the

carbon layer becomes more facile for Li-ion transport, resulting in increased capacity [29, 30].

The voltage profiles of the first two cycles (C/10) of 0.005-2.8 V in various carbon materials are

presented in Fig. 10. In all kinds of prepared CDCs, the hysteresis between the charge and discharge curve was

observed at the first cycle, as shown in Figs. 10(a-d). The capacity losses in the first cycle can be attributed to

SEI layer formation at potential between 0.9 V and 0.3 V vs. Li+/Li [31], as well as to the irreversible reduction

of electrochemically active surface groups such as fluorine or possibly hydroxyls. TiC-CDC@1200 °C Cl2 (Fig.

10a) shows a stable discharge capacity of ~ 320 mA h g-1 in the 20th cycle, corresponding to Li0.85 C6. A high

irreversible capacity loss of ~900 mA h g-1 in the first cycle demonstrates the main drawback of the 0.005-2.8 V

potential window. In TiC-CDC, the first discharge curve has two plateaus at 1.0 V and 0.4 V, and afterward a

slopelike decay down to 0.005 V, whereas the charge curve only has a slopelike behavior, suggesting a structural

change in the first discharge. All subsequent cycles show the same slopelike behavior in both discharge and

charge, which is typical for solid solution lithium insertion/extraction. Also, in micro- and mesoporous carbons,

the high Li storage capacities are dependent on both the disorder structure [32, 33] and hydrogen content [34,

35]. The short-range order of disordered carbon makes it difficult to clarify the relationship between the

structure and reversible lithium insertion properties of disordered carbons. Despite the large insertion capacities,

of disordered carbons, their applicability is hampered by large irreversible capacities and polarization between

the charge and discharge (hysteresis) in the practical cell [9]. The most influential factor on the irreversible

capacity has already been identified as the surface functionality [36]. Although it has been reported that the

addition of H2 to a chlorine treatment reaction stabilizes dangling bonds and increases sp3 bonding, evidence of

surface functionality has not been found in experiments reported thus far [16, 37].

In the case of TiC7N3-CDCs (Figs. 10b and 10c), both TiC7N3-CDC @1200 and @600 still show

irreversible capacity loss of ~ 400 mAhg-1 and ~ 900 mAhg-1 in the first cycle, respectively. The two plateaus at

1.0 V and 0.4 V were observed in the 600 oC Cl2 sample, and were absent in subsequent cycles. It is reasonable

to preliminarily assign this irreversible peak to the formation of a solid electrolyte interphase (SEI) and to an

irreversible reaction with the electrode material. On the other hand, TiC7N3-CDC @1200 has four plateaus at 2.2

V, 1.5 V, 0.9 V, and 0.4 V in the first discharge curve with two plateaus at 1.6 V and 0.9 V in the second

discharge curve. In subsequent cycles, two plateaus were observed at 1.6 V and 0.9 V during lithiation and

delithiation, respectively. The various plateaus of the TiC7N3-CDC @1200 may result from an abundance of

unreacted residues of TiC or TiCN nanoparticles after chlorine treatment, where the graphite covering the edge

of the precursor TiC7N3 nanoparticles causes unreacted TiC7N3 nanoparticles. Also, the potentials of 2.2 V and

1.5 V are similar to those reported for TiO2 and lithiated titania [38, 39].

In B4C-CDC, the 800 oC Cl2 sample (Fig. 10d) shows irreversible capacity loss of ~ 500 mA h g-1 in

the first cycle, where the discharge curves had no potential plateaus at approximately 1.5 V, which is the largest

value among prepared CDCs. The commercial MCMB and hard carbon have first discharge capacities of ~217

mA h g-1 and ~154 mA h g-1, respectively. While they do not suffer from large first-cycle irreversible losses, the

discharge capacity is limited due to non-porous pore textures.

4. Conclusions

We produced CDCs by chlorine treatment of B4C (micron size, 1.7 �m) and TiC7N3 (nano size, 150 nm) in

a temperature range from 600 to 1200 oC. It was confirmed through N2 gas sorption measurement that B4C-

CDCs and TiC7N3-CDCs have marked hysteresis in their N2 sorption isotherms, which confirms a broader pore

size distribution of meso- and macro-porosity. While the conversion of B4C to CDC was almost 100%, the

residue of the TiC7N3-CDCs remained at 15 % and 25 % at 600 oC and 1200 oC, respectively, due to the graphite

shells covering the nanoparticle TiC7N3. For electrochemical capacity as a negative electrode, the discharge

capacity of TiC7N3-CDC@ 600 Cl2 reached the highest value, 462 mA h g-1 at 100 cycles, which was about 25 %

higher than the theoretical capacity of graphite, 375 mA h g-1. Also, B4C-CDC@800 oC Cl2 showed a value of

453 mA h g-1 at 100 cycles. Although the CDCs, including those with mesopore texture, indicated irreversible

capacity and polarization between charging and discharging at the first cycle, which is a typical pattern in a

porous material with a high surface area, the specific capacity of TiC7N3-CDC and B4C-CDC stabilized faster

(about 10 cycles) than the microporous material (about 20 cycle) and was 2-3 times higher than hard carbon and

MCMB in terms of electrochemical capacity. These results show that B4C-CDC and TiC7N3-CDC, where a high

crystalline structure and meso- and macro-porosity coexist, have strong as a potential for negative electrode of

Li battery applications through control over their porosity and crystallinity. Furthermore, they can be exposed

to a practical synthesis temperature range of 600 oC ~ 1200 oC, and provide 2-3 times better performance than

existing graphite or hard carbon anodes for lithium battery systems. However, different surface modifications

and different preparation processes, such as graphite or other carbon coating on CDC, also need to be taken into

consideration, in order to overcome the irreversible capacity loss in the initial cycle for practical application as

negative electrode materials.

Acknowledgments This work was supported by the Next Generation Military Battery Research Center program

of The Defense Acquisition Program Administration and Agency for Defense Development. We acknowledge

KAIST Central Research Instrument Facility for the use of Raman and TEM, and KIER R&D Activity Center

for SEM and XRD.

References

�

Figure Captions

�Carbide-derived carbon (CDC) was synthesized by Cl2 treatment of B4C and TiC7N3. �The CDC was applied to anode of Li-ion battery. �The capacity of TiC7N3-CDC reaches the highest value, 462 mAh/g at 100 cycles. � B4C-CDC@800 oC Cl2 possesses 453 mAh/g at 100 cycles.

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