Solvothermal synthesis of carbon-coated tin nanorods for superior reversible lithium ion storage

Article abstract of DOI:10.1016/j.materresbull.2011.08.060

This paper reports the large-scale synthesis of carbon-coated tin (Sn-C) nanorods through a facile solvothermal carbonization approach by using Sn nanorods as templates and ethanol as a solvent. The as-synthesized Sn-C nanorods have been applied as anode materials for lithium-ion batteries. The Sn-C anode, which exhibits relatively higher initial Coulombic efficiency and lithium storage capacities, is more attractive than the previously reported tin oxide or sulfide anodes. The superior performance of the Sn-C nanorods makes it a promising anode material in high-energy density batteries.

Full text of DOI:10.1016/j.materresbull.2011.08.060

Materials Research Bulletin 46 (2011) 2278–2282
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Materials Research Bulletin
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Solvothermal synthesis of carbon-coated tin nanorods for superior reversible
lithium ion storage
*
Ping Wu, Ning Du, Jie Liu, Hui Zhang, Jingxue Yu, Deren Yang
State Key Lab of Silicon Materials and Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, People’s Republic of China
A R T I C L E I N F O
A B S T R A C T
Article history:
This paper reports the large-scale synthesis of carbon-coated tin (Sn–C) nanorods through a facile
solvothermal carbonization approach by using Sn nanorods as templates and ethanol as a solvent. The
as-synthesized Sn–C nanorods have been applied as anode materials for lithium-ion batteries. The Sn–C
anode, which exhibits relatively higher initial Coulombic efficiency and lithium storage capacities, is
more attractive than the previously reported tin oxide or sulfide anodes. The superior performance of the
Sn–C nanorods makes it a promising anode material in high-energy density batteries.
ß 2011 Elsevier Ltd. All rights reserved.
Received 18 June 2011
Received in revised form 16 August 2011
Accepted 30 August 2011
Available online 6 September 2011
Keywords:
A. Composites
A. Nanostructures
B. Chemical synthesis
D. Energy storage
1. Introduction
during the first cycle, seems to be more attractive compared to
other tin-based anodes.
Rechargeable lithium-ion batteries have been generally accept-
ed as the most promising electrochemical storage devices for
potential applications in cellular telephones, laptop computers,
and digital cameras. However, lithium-ion batteries still face huge
challenges, such as energy density, safety, rate performance, and so
forth [1,2]. As the central part of the batteries, electrode materials
play a critical role in achieving the improved performances [3].
Therefore, much effort has been focused on searching for new
anode materials to replace commercial graphite anodes with
limited capacities (372 mA h gꢀ1) [4,5]. Tin-based materials have
been considered as suitable graphite replacements to attain high
energy density batteries primarily owing to their high theoretical
specific capacities [6].
Recently, the structures of tin-based materials (SnO2 [7,8], SnO
[9], SnS2 [10], SnS [11], and so forth [12]) have been carefully
designed for superior lithium storage performances. For example,
we have prepared carbon nanotubes@SnO2@carbon coaxial
nanocables [7] and ultrathin SnS2 nanosheets [10] which exhibit
stable capacities and excellent cycling performances. However, the
initial Coulombic efficiencies of these tin-based materials are
commonly rather low since the irreversible formation of Li2O, Li2S
and so forth in the first cycle, which restricts the practical
applications of the tin oxide or sulfide anodes [5]. Therefore,
metallic Sn, which does not suffer huge irreversible capacity loss
Up to now, more and more research interest has been paid to
design novel micro/nanostructures of Sn anodes [13–27]. Among
them, Sn–C hybrid structures have been generally accepted as the
most promising choices since the introduction of carbon matrixes
which have high electronic conductivity and good buffering effect
[16–27]. Various types of Sn–C hybrid anodes, including hollow
spheres [16–18], bamboo-like hollow nanofibers [21], rambutan-
like nanocomposites [23,24], and mesoporous nanowires [27],
have been demonstrated to exhibit excellent lithium ion storage
performances. However, toxic organotin precursors [18–22],
special devices [23,24], harsh reaction conditions [19–24], or
multisynthetic processes [16,17] are often needed in the formation
of such the Sn–C hybrid structures, which hinders their mass
production.
Herein, we report the large-scale synthesis of Sn–C nanorods
through a facile glucose solvothermal and subsequent carboniza-
tion approach using pristine Sn nanorods as templates. The as-
synthesized Sn–C core–shell nanorods have been applied as an
anode material for lithium-ion batteries, which exhibit relatively
high initial Coulombic efficiency and lithium storage capacities.
2. Experimental
2.1. Synthesis of Sn–C nanorods
All the chemicals were analytical grade without further
purification. Sn nanorods were synthesized by a simple reaction
using SnCl4 as a precursor and poly(diallyldimethylammonium
* Corresponding author. Fax: +86 571 87952322.
E-mail address: mseyang@zju.edu.cn (D. Yang).
0025-5408/$ – see front matter ß 2011 Elsevier Ltd. All rights reserved.
doi:10.1016/j.materresbull.2011.08.060

P. Wu et al. / Materials Research Bulletin 46 (2011) 2278–2282
2279
chloride) (PDDA) as soft templates at room temperature [28]. Sn–C
nanorods were obtained by a facile solvothermal carbonization
approach. Briefly, the pristine Sn nanorods and glucose were fully
dispersed in 40 ml ethanol under mild sonication. Then, the
solution was transferred into a 50 ml Teflon-lined stainless steel
autoclave, sealed, and maintained at 180 8C for 3 h. After the
reaction was finished, the resulting black solid products were
centrifugalized, washed with distilled water and ethanol to remove
the ions possibly remaining in the final products, and then dried at
80 8C under vacuum. Finally, the brown products were kept in a
tube furnace at 500 8C for 3 h under N2 at a ramping rate of
nanorods. The shell thickness, together with the oxidation degree
of Sn nanorods, increases with the solvothermal time, whereas the
core–shell structures can be retained. Therefore, the reaction
conditionsshouldbecarefullydeterminedforhigh-qualityproducts.
Fig. 1 shows the morphological and structural characterizations
of the Sn and Sn–CP nanorods. As observed from their XRD patterns
(Fig. 1b), the main crystalline phase can all be assigned to
tetragonal structured Sn (JCPDS: 04-0673), indicating the phase of
the nanorods is preserved during the solvothermal treatment. Two
additional broad peaks, which correspond to the (1 1 0) and (2 1 1)
planes of SnO2, respectively, appear in the product by the 6 h
solvothermal treatment (curve c). Therefore, solvothermal time
plays an important role in the formation of high-purity Sn–CP
nanorods (curve b). Fig. 1c–f gives the TEM images of Sn–CP
nanorods. As observed, the one-dimensional (1D) morphology of
the templates (Fig. 1a) is retained in the products. In addition, the
Sn nanorods in the core remain intact by the 3 h solvothermal
treatment (Fig. 1c and d), whereas some of the hybrid nanorods
have converted to tube-like composites by 6 h treatment (Fig. 1e
and f), which is in good agreement with the XRD patterns. In our
previous report, we have prepared SnO2–C hybrid nanotubes
through a hydrothermal carbonization approach using Sn nanor-
ods as sacrificing templates [29]. Compared to the hydrothermal
process, the solvothermal approach could effectively restrain the
oxidation of Sn templates, yielding Sn–C core–shell nanorods
rather than SnO2–C hybrid nanotubes. Therefore, this solvothermal
carbonization procedure can be applied to the large-scale
formation of various types of Sn–C and Si–C hybrid structures,
facilitating their practical application in lithium-ion batteries.
The carbonaceous shell in Sn–CP nanorods (3 h) can be fully
carbonized into carbon shell under inert atmosphere [29,32].
Fig. 2 shows the morphological, structural and compositional
characterizations of the as-synthesized Sn–C nanorods. As can be
seen from the FESEM image (Fig. 2a), the 1D morphology still
remains after carbonization. The TEM image (Fig. 2b) clearly
reveals the rod-in-tube structure of Sn–C nanorods. As observed,
most of the hybrid nanorods exist in the form of partially filled Sn
rods in carbon tubes since the inner core has low melting points.
Similar morphologies have been observed in Sn–C rambutan-like
composites [23], nanorods [33], and aligned nanoarches [34]
prepared at high temperatures. The observed crystalline phase
from XRD pattern (Fig. 2c) is in accordance with the pristine Sn
templates (tetragonal Sn), while the outer carbon shell is not well
crystallized. No peaks related to tin oxides (SnO, SnO2) can be
detected, suggesting high-purity of the Sn–C nanorods. It can be
seen from the magnified TEM image (Fig. 2d) that the carbon shell
is smooth, continuous and the shell thickness is about 5–10 nm,
whereas the Sn core is single crystalline in nature, which can be
verified by its selected-area electron diffraction (SAED) pattern
(Fig. 2d, inset). The HRTEM images of the middle and end parts of a
Sn–C nanorod clearly demonstrate the interface regions of the
core–shell structure. As observed from Fig. 2e and f, the lattice
fringes in the core region can be ascribed to the (1 0 1) and (2 1 ꢀ1)
planes of metallic tin, respectively. Their Fast Fourier Transform
(FFT) patterns also show strong reflections of these planes (Fig. 2e
and f, insets). In addition, there are some SnO2 nanocrystallites
with small size between the tin core and carbon shell due to the
slight oxidation of Sn nanorods prepared through a solution-
based method [28].
5 8C minꢀ1
.
2.2. Characterization
The obtained samples were characterized by X-ray powder
diffraction (XRD) using a Rigaku D/max-ga X-ray diffractometer
˚
with graphite monochromatized Cu Ka radiation (g = 1.54178 A).
The morphology and structure of the samples were examined by
transmission electron microscopy (TEM, JEM-200 CX, 160 kV),
high-resolution transmission electron microscopy (HRTEM, JEOL
JEM-2010) and field emission scanning electron microscopy
(FESEM, Hitachi S-4800). The differential scanning calorimetry
and thermogravimetric analysis (DSC–TGA) were tested on an SDT
Q600 V8.2 Bulid 100.
2.3. Electrochemical measurements of Sn–C nanorods
Electrochemical measurements were carried out using two-
electrode 2025 type coin cells with lithium metal as the counter
electrodes. The working electrodes were composed of the active
material (Sn–C nanorods), conductive materials (acetylene black,
AB), and binder (polyvinyldifluoride, PVDF) in a weight ratio of Sn–
C nanorods/AB/PVDF = 80:10:10, and pasted on a copper foil. The
amount of active material loading on each copper foil (13 mm in
diameter) was about 2 mg. The weight of carbon layer was
inclusive when calculating the specific capacity of Sn–C nanorod
anode. The electrolyte solution was 1 M LiPF6 dissolved in a
mixture of ethylene carbonate (EC), propylene carbonate (PC), and
diethyl carbonate (DEC) with the volume ratio of EC/PC/
DEC = 3:1:1. The cell assembly was performed in a glovebox filled
with pure argon (99.999%) in the presence of an oxygen scavenger
and a sodium-drying agent. The electrode capacity was measured
by a galvanostatic discharge–charge method at a current density of
100 mA gꢀ1 in the potential range of 0.01–2 V at 20 8C. Cyclic
voltammetry (CV) were recorded on a MSTAT4 (Arbin Instruments)
system in the potential range of 0.0–2.0 V at a scan rate of
0.1 mV sꢀ1
.
3. Results and discussion
Hydrothermal and solvothermal carbonization of saccharides
(glucose, sucrose, and starch) have been widely accepted as one of
the most effective and economic routes for fabricating carbon
spheres and carbon-coated hybrid nanostructures [7,8,29–31].
Among them, solvothermal approaches have their own superiority
for the formation of some special structures which are hard to
preparethroughhydrothermalprocesses[30,31]. Forexample, CoO–
C [30] and SnS2–C [31] composites have been prepared through
solvothermal approaches using ethanol and ethylene glycol as
solvents, respectively. Herein, Sn–C core–shell nanorods can be
obtained through a similar solvothermal approach by using Sn
nanorods as templates and ethanol as a solvent. During this
solvothermal process, the carbon precursor (CP) materials, which
derived from glucose in the solution, are deposited heterogeneously
onto the surface of Sn nanorods, resulting Sn–CP core–shell
Sn–C hybrid materials have been regarded as potential anode
materials in lithium-ion batteries since the combination of tin with
high specific capacity and carbon with good cycle life [5].
Therefore, carbon contents are very important for improved
lithium storage performances. Herein, DSC–TGA was performed to
determine the carbon content presented in the Sn–C nanorods.
Fig. 3a shows the DSC and TGA curves of the Sn–C nanorods under

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P. Wu et al. / Materials Research Bulletin 46 (2011) 2278–2282
Fig. 1. (a) TEM image of pristine Sn nanorods; (b) XRD patterns of pristine Sn nanorods (curve a) and Sn–CP nanorods prepared with different solvothermal times: 3 h (curve b)
and 6 h (curve c); and TEM images of Sn–CP nanorods prepared with different solvothermal times: 3 h (c and d) and 6 h (e and f).
synthetic air at a ramping rate of 10 8C minꢀ1. As can be seen from
the TGA curve, the initial weight loss (up to 150 8C) of the sample is
due to the evaporation of physically adsorbed water, whereas the
subsequent weight variation could be mainly attributed to the
oxidation of Sn core and carbon shell. The corresponding DSC curve
shows three exothermic peaks at 234.2, 297.6, and 402.3 8C, which
can be assigned to the oxidation of tin to tin oxides (SnO and SnO2)
and removal of carbon shell, respectively. The oxidation of Sn to
SnO2 (200–800 8C) leads to a weight increase and the removal of
carbon shell (300–500 8C) leads to a weight decrease of the
product, resulting in two creeping up processes between 200 and
800 8C in the TGA curve [17]. Thus, the carbon content of Sn–C
nanorods can be calculated to be about 30.3% by weight.
the current peaks from CV curves are consistent with the voltage
plateaus from discharge/charge curves, and the profiles of these
curves are characteristic of Sn-based anodes [17–21]. In the first
discharge process, the broad peak at about 0.8 V could be ascribed
mainly to the formation of solid electrolyte interface (SEI) layer on
carbon surface [23], while the peak at 0.0–0.4 V is associated with
the formation of LixSn alloys (Sn + xLi+ + xeꢀ ! LixSn). Whereas, the
peaks at 0.4–1.0 V in the charge processes are related to lithium
extraction from LixSn alloys [17]. In addition, the small anodic peaks
at about 1.3 V probably related to the partial reversible electro-
chemical processes, such as the formation of SEI layer, the
intercalation of lithium ions in defect sites, and so forth [35,36].
Fig. 3d shows the discharge and charge capacities as a function
of cycle number for the Sn–C nanorod anode at a current density of
100 mA gꢀ1 in the potential range of 0.01–2 V. For comparison,
Fig. 3b and c displays the first three cyclic voltammogram (CV),
discharge, and charge curves of a Sn–C nanorod anode. As observed,

P. Wu et al. / Materials Research Bulletin 46 (2011) 2278–2282
2281
Fig. 2. Morphological, structural and compositional characterizations of Sn–C nanorods: (a) FESEM image; (b) TEM image; (c) XRD pattern; (d) magnified TEM image and its
SAED pattern (inset); (e and f) HRTEM images and their FFT patterns (insets) of a Sn–C nanorod recorded along the [1 ꢀ3 ꢀ1] zone axis.
cycling performance of pristine Sn nanorods was also investigated
under the same conditions. As observed, the discharge and charge
capacities of Sn–C nanorod anode in the first cycle are 1477.4 and
957.1 mA h gꢀ1, respectively. The initial capacity loss (35.2%) is
mainly due to the formation of SEI layer on carbon surface in the
first discharge process. It is worthy of mention that the reversible
capacities in subsequent several cycles are still higher than the
reversible capacity of 806.6 mA h gꢀ1, which is comparable to the
previously reported Sn–C composite anodes [16–27]. In sharp
contrast, the pristine Sn nanorod anode exhibits very low initial
Coulombic efficiency (32.9%) and discharge capacities (307 and
71.3 mA h gꢀ1 in the first and 20th cycle, respectively). The reason
is as follows. The huge volume change of the Sn nanorod induces
both mechanical degradation and loss of electric contact at the
electrode, resulting in its poor electrochemical performance.
Therefore, the structural design of tin-based anodes plays a critical
role for superior lithium storage performance. Herein, the Sn–C
nanorod anode can be a good example due to the enhanced electric
conductivity and structural stability during cycling. Moreover, the
lithium storage performance can be further improved by adjusting
the carbonization temperature [26] or fabricating highly pure
products [19,20].
theoretical values of both tin (994 mA h gꢀ1
) and carbon
(372 mA h gꢀ1) anodes. Similar phenomena have been observed
in nanostructured tin-based anodes by different groups
[17,29,37,38]. The extra reversible capacities can be assigned to
the partial reversible electrochemical processes during cycling
[36–38]. For example, the formation/decomposition of SEI layer on
carbon surface has proven to be partially reversible in NiO–C anode
[36]. After 20 cycles, the Sn–C anode is able to exhibit a high

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P. Wu et al. / Materials Research Bulletin 46 (2011) 2278–2282
Fig. 3. Electrochemical performance of the Sn–C nanorod anode for lithium-ion battery: (a) DSC–TGA curves; (b) the first three CV curves; (c) the first three discharge and
charge curves; (d) cycling performances of the Sn–C nanorod and pristine Sn nanorod anodes.
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In summary, we have synthesized Sn–C nanorods through a
facile glucose solvothermal and subsequent carbonization ap-
proach by using Sn nanorods as templates and ethanol as a solvent.
The Sn–C nanorod anode exhibits relatively high initial Coulombic
efficiency (64.8%) and lithium storage capacities (806.6 mA h gꢀ1
in the 20th cycle), which makes it a promising anode material in
high energy density batteries.
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