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High-performance lithium battery
anodes using silicon nanowires
CANDACE K. CHAN1, HAILIN PENG2, GAO LIU3, KEVIN McILWRATH4, XIAO FENG ZHANG4,
ROBERT A. HUGGINS2 AND YI CUI2*
1Department of Chemistry, Stanford University, Stanford, California 94305, USA
2Department of Materials Science and Engineering, Stanford University, Stanford, California 94305, USA
3Environmental Energy Technologies Division, Lawrence Berkeley National Lab, 1 Cyclotron Road, Mail Stop 70R108B, Berkeley, California 94720, USA
4Electron Microscope Division, Hitachi High Technologies America, Inc., 5100 Franklin Drive, Pleasanton, California 94588, USA
*e-mail: yicui@stanford.edu
Published online: 16 December 2007; doi:10.1038/nnano.2007.411
There is great interest in developing rechargeable lithium
batteries with higher energy capacity and longer cycle life for
applications in portable electronic devices, electric vehicles
and implantable medical devices1. Silicon is an attractive
anode material for lithium batteries because it has a low
discharge potential and the highest known theoretical charge
capacity (4,200 mAh g21; ref. 2). Although this is more than
ten times higher than existing graphite anodes and much
larger than various nitride and oxide materials3,4, silicon
anodes have limited applications5 because silicon’s volume
changes by 400% upon insertion and extraction of lithium,
which results in pulverization and capacity fading2. Here, we
show that silicon nanowire battery electrodes circumvent these
issues as they can accommodate large strain without
pulverization, provide good electronic contact and conduction,
and display short lithium insertion distances. We achieved the
theoretical charge capacity for silicon anodes and maintained a
discharge capacity close to 75% of this maximum, with little
fading during cycling.
Previous studies in which Si bulk films and micrometre-sized
particles were used as electrodes in lithium batteries have shown
capacity fading and short battery lifetime due to pulverization
and loss of electrical contact between the active material and the
current collector (Fig. 1a). The use of sub-micrometre pillars6
and micro- and nanocomposite anodes5,7–9 led to only limited
improvement. Electrodes made of amorphous Si thin films have a
stable capacity over many cycles5,8, but have insufficient material
for a viable battery. The concept of using one-dimensional (1D)
nanomaterials has been demonstrated with carbon10, Co3O4
(refs 11, 12), SnO2 (ref. 13) and TiO2 (ref. 14) anodes, and has
shown improvements compared to bulk materials. A schematic of
our Si nanowire (NW) anode configuration is shown in Fig. 1b.
Nanowires are grown directly on the metallic current collector
substrate. This geometry has several advantages and has led to
improvements in rate capabilities in metal oxide cathode
materials15. First, the small NW diameter allows for better
accommodation of the large volume changes without the
initiation of fracture that can occur in bulk or micron-sized
materials (Fig. 1a). This is consistent with previous studies that
have suggested a materials-dependent terminal particle size below
which particles do not fracture further16,17. Second, each Si NW
is electrically connected to the metallic current collector so that
all the nanowires contribute to the capacity. Third, the Si NWs
have direct 1D electronic pathways allowing for efficient charge
transport. In electrode microstructures based on particles,
electronic charge carriers must move through small interparticle
contact areas. In addition, as every NW is connected to the
current-carrying electrode, the need for binders or conducting
additives, which add extra weight, is eliminated. Furthermore,
Initial substrate After cycling
XX
X
Film
Particles
Facile strain
relaxation
Good contact with current collector
Efficient 1D
electron transport
Nanowires
Figure 1 Schematic of morphological changes that occur in Si during
electrochemical cycling. a, The volume of silicon anodes changes by about
400% during cycling. As a result, Si films and particles tend to pulverize during
cycling. Much of the material loses contact with the current collector, resulting
in poor transport of electrons, as indicated by the arrow. b, NWs grown directly
on the current collector do not pulverize or break into smaller particles after
cycling. Rather, facile strain relaxation in the NWs allows them to increase in
diameter and length without breaking. This NW anode design has each NW
connecting with the current collector, allowing for efficient 1D electron transport
down the length of every NW.
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with smooth sidewalls (Fig. 3a) and had an average diameter of
89 nm (s.d., 45 nm) (Fig. 3e). Cross-sectional scanning electron
microscopy (SEM) showed that the Si NWs grew off the substrate
and had good contact with the stainless steel current collector
(Fig. 3a, inset). After charging with Li, the Si NWs had roughly
textured sidewalls (Fig. 3b), and the average diameter increased
to 141 nm (s.d., 64 nm). Despite the large volume change, the
Si NWs remained intact and did not break into smaller particles.
They also appeared to remain in contact with the current
collector, suggesting minimal capacity fade due to electrically
disconnected material during cycling.
The Si NWs may also change their length during the change in
volume. To evaluate this, 25-nm Ni was evaporated onto as-grown
Si NWs using electron beam evaporation. Because of the shadow
effect of the Si NWs, the Ni only covered part of the NW surface
(Fig. 3c), as confirmed by energy dispersed X-ray spectroscopy
(EDS) mapping (see Supplementary Information, Fig. S4). The
Ni is inert to Li and acts as a rigid backbone on the Si NWs.
After lithiation (Fig. 3d), the Si NWs changed shape and
wrapped around the Ni backbone in a three-dimensionally helical
manner. This appeared to be due to an expansion in the length
of the NW, which caused strain because the NW was attached to
the Ni and could not freely expand but rather buckled into a
helical shape. Although the NW length increased after lithiation,
the NWs remained continuous and without fractures,
maintaining a pathway for electrons all the way from the
collector to the NW tips. With both a diameter and length
increase, the Si NW volume change after Li insertion appears to
be about 400%, consistent with the literature5.
Efficient electron transport from the current collector to the Si
NWs is necessary for good battery cycling. To evaluate this, we
conducted electron transport measurements on single Si NWs
before and after lithiation (see Methods). The current versus
voltage curve on a pristine Si NW was linear, with a 25 kV
resistance (resistivity of 0.02 V-cm) (Fig. 3f ). After one cycle, the
NWs became amorphous, but still exhibited a current that was
linear with voltage with an 8 MV resistance (resistivity of
3 V-cm) (Fig. 3g). The good conductivity of pristine and cycled
NWs ensures efficient electron transport for charge and discharge.
The large volume increase in the Si NWs is driven by the
dramatic atomic structure change during lithiation. To
understand the structural evolution of NWs, we characterized the
NW electrodes at different charging potentials. The X-ray
diffraction (XRD) patterns were taken for initial pristine Si NWs,
Si NWs charged to 150 mV, 100 mV, 50 mV and 10 mV, as well as
after 5 cycles (Fig. 4a). XRD patterns of the as-grown Si NWs
Nu
m
be
r o
f n
an
ow
ire
s Initial
Voltage (V)
10 mV
Cu
rre
nt
(µ
A)
Cu
rre
nt
(n
A)
0.0–0.4 0.4
Voltage (V)
0.0–0.4 0.4
0
10
20
–10
–20
0
4
8
–4
Diameter (nm)
200
100
0
0 100 200 300
250 nm50 nm
Si Ni Si
Ni
Figure 3 Morphology and electronic changes in Si NWs from reaction with Li. a,b, SEM image of pristine Si NWs before (a) and after (b) electrochemical cycling
(same magnification). The inset in a is a cross-sectional image showing that the NWs are directly contacting the stainless steel current collector. c,d, TEM image of a
pristine Si NW with a partial Ni coating before (c) and after (d) Li cycling. e, Size distribution of NWs before and after charging to 10 mV (bin width 10 nm).
The average diameter of the NWs increased from 89 to 141 nm after lithiation. f, I –V curve for a single NW device (SEM image, inset) constructed from a pristine
Si NW. g, I –V curve for a single NW device (SEM image, inset), constructed from a NW that had been charged and discharged once at the C/20 rate.
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showed diffraction peaks associated with Si, a-FeSi2, Au (the Si NW
catalyst) and stainless steel (SS). The a-FeSi2 forms at the interface
between the SS and the Si wires during the high temperature (530
8C) NW growth process. The a-FeSi2 was not found to appreciably
react with Li during electrochemical cycling, although a small
amount of reaction has been reported24. After Si NWs were
charged to 150 mV, the higher angle Si peaks disappeared. Only
the Si(111) peak was still visible, but its intensity was greatly
decreased. This is consistent with the disappearance of the
initial crystalline Si and the start of the formation of amorphous
LixSi. The four broad peaks that appeared in the lower angles
are due to the formation of Li15Au4 (see Supplementary
Information, Fig. S5). At 100 mV, the pure Au peaks
disappeared, indicating that the Au had completely reacted with
Li. The Si(111) peak was very weak at 100 mV, and
disappeared completely at 50 mV. It appears that Si NWs
remain amorphous after the first charge, consistent with the
non-flat voltage charging/discharging curve in Fig. 2b. This
contrasts, however, with other studies on Si electrodes25,26, which
have reported the formation of crystalline, Li3.75Si at potentials
less than 30–60 mV. In situ XRD studies have determined
that this crystalline phase only forms at ,50 mV for films
thicker than 2 mm (ref. 27). We did not observe this to be
the case in our Si NWs, most likely because of their shape
and small dimensions.
The local structural features of Si NWs during the first Li
insertion were studied with transmission electron microscopy
(TEM) and selected area electron diffraction (SAED). The as-
grown Si NWs were found to be single-crystalline. Figure 4b
shows an example of a typical Si NW with a k112l growth
direction29. Figure 4c shows a Si NW with a k112l growth
direction that was charged to 100 mV. In this case there were two
phases present, as expected from the voltage profile. Both
crystalline and amorphous phases were clearly seen. The
distribution of the two phases was observed both across the
diameter (a crystalline core and an amorphous shell) and along
the length. The SAED showed the spot pattern for the crystalline
phase (Si), but weak diffuse rings from the amorphous phase
(LixSi alloy) were also observed. Li ions must diffuse radially into
the NW from the electrolyte, resulting in the core–shell phase
distribution. The reason for phase distribution along the length is
not yet understood. At 50 mV, the Si NW became mostly
amorphous with some crystalline Si regions embedded inside the
core, as seen from the dark-field image and HRTEM (Fig. 4d).
The SAED showed spotty rings representative of a polycrystalline
sample and diffuse rings for the amorphous phase. At 10 mV
(Fig. 4e), all of the Si had changed to amorphous Li4.4Si, as
indicated by the amorphous rings in the SAED. These TEM
observations were consistent with the XRD results (Fig. 4a) and
voltage charging curves (Fig. 2b).
METHODS
Si NWs were synthesized using the VLS process on stainless steel substrates using
Au catalyst. The electrochemical properties were evaluated under an argon
atmosphere by both cyclic voltammetry and galvanostatic cycling in a three-
electrode configuration, with the Si NWs on the stainless steel substrate as the
working electrode and Li foil as both reference and counter-electrodes. No
binders or conducting carbon were used. The charge capacity referred to here is
the total charge inserted into the Si NW, per mass unit, during Li insertion,
whereas the discharge capacity is the total charge removed during Li extraction.
For electrical characterization, single Si NW devices were contacted with metal
electrodes by electron-beam lithography or focused-ion beam deposition. For
more detailed descriptions of NW synthesis, TEM and XRD characterization,
electrochemical testing, and device fabrication, see the Supplementary
Information.
Received 23 July 2007; accepted 14 November 2007; published 16 December 2007.
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Acknowledgements
We thank Dr Marshall for help with TEM interpretation and Professors Brongersma and Clemens for
technical help. Y.C. acknowledges support from the Stanford New Faculty Startup Fund and Global
Climate and Energy Projects. C.K.C. acknowledges support from aNational Science FoundationGraduate
Fellowship and Stanford Graduate Fellowship.
Correspondence and requests for materials should be addressed to Y.C.
Supplementary information accompanies this paper on www.nature.com/naturenanotechnology.
Author contributions
C.K.C. conceived and carried out the experiment and data analysis. H.P., G.L., K.M. and X.F.Z. assisted in
experimental work. R.A.H. carried out data analysis. Y.C. conceived the experiment and carried out data
analysis. C.K.C., R.A.H. and Y.C. wrote the paper.
Reprints and permission information is available online at http://npg.nature.com/reprintsandpermissions/
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