Probing lithium germanide phase evolution and structural change in a germanium-in-carbon nanotube energy storage system

Article abstract of DOI:10.1021/ja5116259

Lithium alloys of group IV elements such as silicon and germanium are attractive candidates for use as anodes in high-energy-density lithium-ion batteries. However, the poor capacity retention arising from volume swing during lithium cycling restricts their widespread application. Herein, we report high reversible capacity and superior rate capability from core-shell structure consisting of germanium nanorods embedded in multiwall carbon nanotubes. To understand how the core-shell structure helps to mitigate volume swings and buffer against mechanical instability, transmission electron microscopy, X-ray diffraction, and in situ ⁷Li nuclear magnetic resonance were used to probe the structural rearrangements and phase evolution of various Li-Ge alloy phases during (de)alloying reactions with lithium. The results provide insights into amorphous-to-crystalline transition and lithium germanide alloy phase transformation, which are important reactions controlling performance in this system.

Full text of DOI:10.1021/ja5116259

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Article
Probing Lithium Germanide Phase Evolution and Structural Change
in a Germanium-in-Carbon Nanotube Energy Storage System
Wei Tang, Yanpeng Liu, Chengxin Peng, Mary Y. Hu, Xuchu Deng, Ming Lin, Jian Zhi Hu, and Kian Ping Loh
J. Am. Chem. Soc., Just Accepted Manuscript • Publication Date (Web): 03 Feb 2015
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Probing Lithium Germanide Phase Evolution and Structural Change
in a Germanium-in-Carbon Nanotube Energy Storage System
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a,d
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Wei Tang , Yanpeng Liu , Chengxin Peng , Mary Y. Hu , Xuchu Deng , Ming Lin* , Jian Zhi Hu*
a
and Kian Ping Loh*
a
Department of Chemistry and Graphene Research Centre, National University of Singapore, 3 Science Drive 3, Singapore
1
b
17543.
Pacific Northwest National Laboratory, Richland, WA 99352, USA.
c
Institute of Materials Research and Engineering, 3 Research Link, Singapore 117602.
d
NUS Graduate School for Integrative Sciences and Engineering, 28 Medical Drive #05ꢀ01, Singapore 117597.
KEYWORDS Germanium; CNTs; Energy storage; inꢀsitu NMR; phase transition
ABSTRACT: Lithium alloys of group IV elements such as silicon and germanium are attractive candidates for use as anodes in
highꢀenergyꢀdensity lithiumꢀion batteries. However, the poor capacity retention arising from volume swing during lithium cycling
restricts their widespread application. Herein, we report high reversible capacity and superior rate capability from core–shell strucꢀ
ture consisting of germanium nanorods embedded in multiwall carbon nanotubes. To understand how the core–shell structure helps
to mitigate volume swings and buffer against mechanical instability, transmission electron microscopy, Xꢀray diffraction and inꢀsitu
7
Li nuclear magnetic resonance were used to probe the structural rearrangements and phase evolution of various Li–Ge alloy phases
during (de)alloying reactions with lithium. The results provide insights into amorphousꢀtoꢀcrystalline transition and lithium germaꢀ
nide alloy phase transformation, which are important reactions controlling performance in this system.
INTRODUCTION
Group IV elements such as tin (Sn), germanium (Ge) and
require the GeH /GeCl precursors to be operated at very diꢀ
4 4
1
ꢀ3
4ꢀ10
lute concentrations. Using solution phased SFLS and SLS
methods, the production can be scaled up to gram level. Howꢀ
ever, the high operation pressure in SFLS and requirement of
high boilingꢀpoint reaction solvents in SLS hinder upward
1
1ꢀ17
silicon (Si),
with theoretical capacities of 992, 1623 and
ꢀ
1
4200 mAh g respectively, are attractive alternatives to carꢀ
bonꢀbased anodes in lithiumꢀion batteries (LIBs). Despite havꢀ
ing a lower capacity than Si, Ge has 400 times higher diffusivꢀ
36
scaling. Using liquid organometallic compounds as precurꢀ
sors, MOCVD approach is capable of scaling up the producꢀ
4
ity of lithium and 1×10 times higher electrical conductivity
36
tion of Ge NWs more safely and economically. Asꢀprepared
and thus has advantages for use in highꢀpower LIBs for elecꢀ
1
8ꢀ24
Ge NWs/NRs (nanorods) were modified with carbonꢀbased
protection layer to improve their electrochemical performance.
trical vehicles and large energyꢀstorage stations.
However,
similar to Si, drastic volume changes (370%) during lithiation
and delithiation, together with the loss of electrical contact or
unstable formation of surface SEI film arising from these,
Despite the rapid progress on the material engineering of
Ge coreꢀshell structure, the mechanism of reaction of Ge in
lithium cycling remains unclear, which makes attempts to
optimize performance based on electrochemical–structural
correlations difficult. The problem is compounded by the
18
result in rapid capacity decay. Strategies to mitigate these
problems generally involve nanostructuring the Ge to miniꢀ
mise volume swing, and embedding it on an electrically conꢀ
ducting scaffold such as amorphous carbon or reduced graꢀ
phene oxide to buffer against mechanical instability and enꢀ
7
,35,41
amorphous nature of the lithium germanide phases,
which
precludes clear elucidation of the structural changes that occur
during4 lithiation/delithiation using Xꢀray diffraction
6
,8,9,17,25ꢀ33
hance the rate capability.
Despite the enhanced perꢀ
2,43
44ꢀ48
formance, the inevitable inhomogeneity of these carbon coatꢀ
ings in physically mixed systems negatively affects the capaciꢀ
(XRD)
or electron microscopy (EM).
By providing
detailed local structural information and detecting both crysꢀ
talline and the amorphous phase simultaneously, nuclear magꢀ
netic resonance (NMR) spectroscopy can be combined with
5
ty retention and rate capability of the system. Recently “coreꢀ
shell structure” such as Ge NWs (core)/carbon (shell) as a
generic concept has emerged as a favored material design in
traditional methods to provide a powerful tool for elucidating
5
,34,35
49ꢀ57
energy storage system.
The common route to synthesize
the electrochemical–structural correlations.
Herein, we
such core/shell structure is based on a postꢀcarbonꢀcoating
successfully develop a strategy to fabricate a capsuleꢀlike
structure consisting of Ge nanorods encapsulated by bambooꢀ
type multiwall carbon nanotubes (Ge@CNT), using Ge nanoꢀ
particles as seeds. The unique capsuleꢀlike compartment of
Ge@CNT protects against volume swing and allows dynamic
structural change and phase evolution (reversible amorphousꢀ
toꢀcrystalline transition) of the Ge during lithiation/
method on the asꢀprepared Ge NWs. Metalꢀseeded vapor–
5,36,37
liquid–solid
(VLS),
supercriticalꢀfluid–liquid–solid
34,38
(SFLS),
solution–liquid–solid (SLS) and metal–organic
chemical vapor deposition(MOCVD) have been applied to
9
,35,36,39,40
synthesize Ge NWs.
Conventional VLS type CVD
suffers from low synthetic yield due to safety concerns which
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Figure 1. Structural change in the Ge@CNT composite during the first lithiation/ delithiation cycle. (a) Illustration of the TEM coinꢀcell
configuration; (b),(e) TEM and (c) STEM images of the Ge@CNT composite, and (d) elemental mapping and STEMꢀHAADF plot across
line M1 in the Ge@CNT composite at full lithiation where an interfacial layer with reduced Ge density is indicated by the red arrows; (f),(i)
TEM and (g) STEM images of the Ge@CNT composite at full delithiation; (h) SAED pattern for the Ge@CNT composite, showing that it
is amorphous with Ge crystalline ‘islands’ dispersed within; (j) elemental mapping and STEMꢀHAADF plot across line M2 in the
Ge@CNT composite at full delithiation, where a buffer zone is indicated by the blue arrows. At full lithiation there is a 10 nm interfacial
layer between the Ge NRs and the carbon layers, with a lower Ge density based on STEM and elemental mapping. At full delithiation,
parts of the Ge NRs contract from the carbon layers and a buffer zone is created.
7
delithiation to be revealed by inꢀsitu Li NMR and TEM. Unꢀ
locking phase evolution during Li alloying and deꢀalloying in
the lithium germanide systems shed light on the origin of the
highꢀcapacity Li storage and its reversibility.
was introduced by bubbling part of the reaction gases through
58
distilled water during the growth process. Carbon, hydrogen,
nitrogen and sulphur (CHNS) analysis showed that the
Ge@CNT composite contained only ~5% by weight of carbon.
As a control, bare Ge particles were obtained by using HCl to
dissolve the MgO from the Ge NP/MgO composite directly
after ballꢀmilling. The detailed synthetic procedure, morpholꢀ
ogy and characterization of Ge@CNT and bare Ge have been
added in Fig.S1ꢀS5.
RESULTS AND DISCUSSION
1. Synthesis and electrochemical performance of Ge@CNT
Ge @CNT was prepared through a oneꢀstep CVD growth apꢀ
proach from Ge nanoparticles directly: A Ge nanopartiꢀ
cle/magnesium oxide catalyst (Ge NP/MgO) was first preꢀ
pared by ballꢀmilling magnesium (Mg) and germanium dioxꢀ
ide (GeO ) together, during which a solidꢀstate reduction of
GeO resulted in the formation of Ge NPs coated with MgO.
Supported on an alumina substrate, the composite was used to
catalyse carbon nanotube (CNT) growth at 880 C in a CVD
When used as anode materials for LIBs, the Ge@CNT system
can deliver promising performance. The detailed electrochemꢀ
ical characterizations and discussions are included in Fig.S6ꢀ
S12. The performance of this Ge@CNT system as anode for
LIBs can be clearly seen in Table S1, in which the rate capaꢀ
bility, capacity retention and high specific capacity of the
Ge@CNT composite are compared with previously reported
Geꢀbased systems, including Ge coated with amorphous carꢀ
2
2
o
furnace using a mixture of methane (CH ) and hydrogen (H ).
4
2
After CVD, A high yield of the core (Ge) plus shell (CNT)
structure was obtained; the Ge NPs were reshaped into nanoꢀ
rods around 1.5ꢁm long inside CNT walls with thickness of 5ꢀ
10nm. To suppress growth of amorphous carbon, water vapour
4ꢀ8,25,32ꢀ35,59ꢀ61
bon or reduced graphene oxide,
the best performance nano Ge/rGO/C system reported to
and even one of
39
date. In contrast, bare Ge exhibits poor electrochemistry
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Figure 2. Reversible amorphousꢀcrystalline transitions during the lithiation/delithiation of Ge@CNT. (a) First cycle voltage plot of
Ge@CNT at 0.2C to indicate stages where XRD was measured; (b) Stacked Xꢀray diffraction patterns collected from Ge@CNT composite
electrode during the first discharge/charge cycle (XRD1ꢀXRD8), reꢀdischarge to 0.005V (XRD9) and charge to 1 V after 200 cycles
(
XRD10). Reversible amorphousꢀcrystalline transition is revealed during the lithiation and delithiation of Ge@CNT. There is no Li Ge
22 5
crystalline phase shown in all the XRD spectra. The reference patterns for Li Ge (red; JCPDS card NO.89ꢀ3034), Ge (blue; JCPDS card
15
4
NO.89ꢀ5011) and Li Ge (dark yellow; JCPDS card NO.17ꢀ0402) are indicated as bars in the Figure. The diffraction peaks of Mylar seal
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5
film and the Copper current collector are indicated in the figure.
beyond the first discharge, in part due to problems associated
with maintaining electrical contacts between the bare Ge parꢀ
The TEM observations revealed a dynamic structural rearꢀ
rangement of the Ge NRs during lithium cycling, with several
salient features. The morphology of the Ge@CNT composite
was well preserved throughout the lithiation (Fig. 1b–d) and
3
5
ticles and the current collector (Fig.S10). This limitation
precludes detailed and operando study of phase transition on
Geꢀbased electrode during lithium cycling. In this report, the
Ge@CNT structure provides an effective model for multiꢀ
probe studies (TEM, XRD and inꢀsitu NMR) as electrical conꢀ
7,9,35
delithiation (Fig. 1e–i) cycles.
The carbon sheath over the
Ge ensured that electrical contact was maintained throughout
the lithiation/delithiation processes. Second, rather than isoꢀ
tropic volume expansion, the confinement of the Ge within the
carbon walls meant that some volume expansion occurred
along the axis of the carbon wall, with the latter acting as a
1
2
tact was maintained throughout the lithium cycling.
2. Structural change in the Ge@CNT composite during
lithiation/delithiation using a TEM coin-cell configuration
Stressꢀinduced pulverisation or unstable formation of surface
SEI film due to volume swing has been reported to be the
main cause of degradation in Geꢀbased batteries, and thus a
direct visualisation of the rearrangement process may provide
63
supporting scaffold. In the initial discharging to 0.17 V,
about half of the capacity was delivered and the bulk of the Ge
+
NRs were alloyed with Li to become an amorphous Li2.26Ge
phase (as later verified by NMR) (Fig.S13a,b). As shown in
the TEM image in Fig. 1e, at full lithiation the amorphous
phase was transformed into a crystalline Li Ge phase (veriꢀ
4
ꢀ10
insights into stress points in the structure. Herein, the strucꢀ
tural changes in Ge@CNT during lithiation/delithiation were
15
4
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4,62
fied by exꢀsitu XRD, as discussed below).As verified by
STEM (Fig. 1c) and elemental mapping (Fig. 1d), a 10 nmꢀ
thick amorphous Li–Ge phase (indicated by red arrows to
show Ge with a relative low distribution) was formed between
studied using a TEM coinꢀcell configuration.
As illustrated
in Fig. 1a, the Ge@CNT composite was first loaded onto a
TEM grid, and the entire structure was then incorporated into
a conventional coin cell, which was subjected to galvanostatic
charging/discharging (0.2C). The charge/ discharge process
was interrupted at different stages and the TEM grids were
disassembled from the cell for TEM characterisation. This
method minimises damage to the sample during preparation,
as direct viewing is possible without mechanical agitation or
dispersion of solutions.
the CNT walls and the crystalline Li Ge phase (also obꢀ
15
4
45
served in previous inꢀsitu TEM study ). The coꢀexistence of
amorphous and crystalline phases is consistent with the
STEMꢀHAADF (highꢀangle annular dark field) plot. Around
the shoulder of HAADF plot (5ꢀ15nm and 45ꢀ55nm), the
counts have reduced to the background level of CNT, indicatꢀ
ing a relatively low Ge concentration in the amorphous phase.
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Upon delithiation (charging to 0.5V), the crystalline phase was correlate with capacity retention. Combining data from TEM
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converted back to a mostly amorphous lithium germanide
phase (Fig.S13c,d). Importantly, at the end of delithiation
(charging to 1 V), a contraction of the deꢀalloyed Ge from the
and inꢀsitu Li NMR, we discovered that the phase interꢀ
conversion during cycling was mediated by coꢀexisting amorꢀ
phous and crystalline phases, and that the high capacity obꢀ
served was correlated with an overꢀlithiated lithium germanide
phase.
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carbon wall was apparent. Fig. 1gꢀh shows that the deꢀ
alloyed Ge became largely amorphous, with several crystalline
Ge ‘islands’ dispersed within it. Similar to the mechanism
proposed for the nucleation of amorphous Si clusters from
existing seeds or defects, the presence of these Ge nanocrysꢀ
tals may act as seeds for the growth of amorphous Ge domains
7
Inꢀsitu Li NMR experiments on the Ge@CNT composite
were performed at 0.2C between 0.005 V and 1 V. In Fig.3, 4,
7
a series of inꢀsitu Li NMR spectra is plotted. Generally, four
7
Li resonances, centred at 24 ppm, 13 ppm, 10 ppm and ꢀ24
+
64
after Li have diffused away. The amorphous Ge domains
are believed to allow a more facile alloying reaction kinetics
ppm (marked R1, R2, R3 and R4, respectively) were identified
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in addition to the Li resonance of the electrolyte and the SEI
+
5
0,65
in subsequent cycles as it will be easier to achieve Li inserꢀ
peak centred around 0 ppm(ꢀ3 ppm, ꢀ1ppm and +4ppm).
52
tion in a loose Ge tetrahedral network.
According to the
As shown in Fig.3, the linewidth and lineshape of the signals
from electrolyte and SEI film are quite narrow and resolved,
TEM image (Fig. 1i) and the elemental mapping profile (Fig.
j), a 10 nm buffer zone (indicated by blue arrows to show
1
thus they can be separated from the Li Ge signals during the
x
empty carbon sheath without Ge; the plateau in HAADF plot,
Fig.1j, around 40ꢀ50 nm with only background counts of
CNTs is consistent with the emerging of buffer zone) was
present throughout the lithiation and delithiation cycles and
remained structurally similar when it was analysed at the end
of 200 cycles (Fig.S14). This buffer zone is believed to acꢀ
commodate the stress associated with volume change during
fitting. Mostly importantly, the signals from electrolytes and
SEI are almost constant during the dischargeꢀcharge process
(the experimental design, data fitting method and quantitative
analysis of the phase compositions following the methods of
Eckert and Pöttgen, et al are included in Fig.S15 and Table S2.
6
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).
(de)lithiation cycles.
3. X-ray diffraction study of the reversible amorphous-to-
crystalline transition in Ge@CNT composite during lithi-
4
3
um cycling According to previous XRD studies, the lithiaꢀ
tion mechanism in Ge occurs via four distinct crystal phases,
Li Ge , Li Ge and Li Ge /Li Ge , although XRD is not caꢀ
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pable of detecting the presence of the amorphous alloy phase.
To investigate the amorphousꢀtoꢀcrystalline transition of the
Ge@CNT composite during lithium cycling, exꢀsitu XRD
analyses were performed at selected potentials as indicated in
the first discharge/charge voltage plot (Fig. 2a). At the first
discharge to 0.17 V, the crystalline diffraction peaks of Ge
(XRD1ꢀXRD4, Fig. 2b) vanished. At 0.11 V, near the end of
discharge, peaks attributable to crystalline Li Ge phase apꢀ
1
5
4
peared (XRD5). There were no new crystalline phases at the
end of the discharge (XRD6). Upon delithiation, the crystalꢀ
line phase was quickly converted to the amorphous phase at
the charging stages of 0.5 V (XRD7) and 1 V (XRD8). At the
7
Figure 3. Colourꢀmapped in situ Li NMR spectra of Ge@CNT
end of the second discharge, the Li Ge crystalline phase was
15
4
during lithiation and delithiation cycles.
observed to reꢀappear (XRD9). These observations suggest a
reversible amorphousꢀtoꢀcrystalline transition during the lithiꢀ
um cycling of the Ge@CNT composite: upon first lithiation,
crystalline Ge becomes amorphous and at the end of the disꢀ
charge it partly transforms into the crystalline Li Ge phase;
According to the evolution of resonance peaks R1–R4 (see
Table 1) during the discharge/charge process, lithiation was
1
5
4
cꢀGe → aꢀLiꢀ.ꢀꢁGe → aꢀLiꢂ.ꢃGe → cꢀLi Ge +aꢀLiꢂ.ꢃGe →
ꢅ
ꢄꢃ
upon delithiation, the crystalline Li Ge phase changes into
1
5
4
cꢀLiꢄꢃꢆꢇGe (metastable)+aꢀLiꢂ.ꢃGe(metastable)
(1)
ꢅ
amorphous Li Ge or Ge. As seen in the TEM images in Fig.
x
During the first discharge (Fig.4a), there was no obvious
change in the spectra until 0.26 V due to carbon lithiation and
1
e, amorphous and crystalline phases coꢀexisted during the
interꢀconversion. Note that the reaction kinetics of the
50,52
18
the initial formation of passive SEI layers.
The R1 peak at
Ge@CNT composite were much faster than those of Si; thus,
after one cycle, almost all of the active mass involved in the
electrochemical reaction became amorphous upon delithiation.
In subsequent cycles, the reversible amorphousꢀtoꢀcrystalline
transition between amorphous Ge and crystalline Li Ge conꢀ
0.17 V was attributed to the amorphous aꢀLi2.26Ge phase or
another early amorphous lithiated phase of Ge (as no XRD
peaks were observed in the exꢀsitu study at this point and after
the capacity of the carbon process and the initial formation of
15
4
ꢀ
1
the SEI were subtracted: 1184ꢀ345 = 839 mAh g , equivalent
tinued during alloying and dealloying (XRD9, XRD10).
+
7
to 2.26 Li inserted per Ge, this phase agreed well with cꢀ
4. In situ Li NMR study of the Lithium germanide alloy
43
Li Ge . ).
9
4
phase transformation in Ge@CNT To understand the deꢀ
7
At around 0.14 V, an R2 peak emerged in the spectra, reflectꢀ
tailed local structural evolution, an inꢀsitu Li NMR study was
ing the transformation of aꢀLi2.26Ge to aꢀLi Ge (based on the
performed. An operando identification of the lithium germaꢀ
nide phases under various cycling regimes permit understandꢀ
ing of the kinetics of phase transition between different strucꢀ
tural phases, including the amorphous phases, and how these
3.5
calculation of the peak area ratio of R1/R2 = 0.45 and the corꢀ
responding capacity). The phase was similar to the previous
43
exꢀsitu XRD observation of cꢀLi Ge , but our XRD study
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Figure 4. Stacked inꢀsitu Li NMR spectra of the Ge@CNT composite obtained during the first two galvanostatic charge/discharge cycles:
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(
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b) First cycle; (c) second cycle. Along with the lithiation/delithiation cycles, four Li resonances centred at 24 ppm, 13 ppm, 10 ppm and ꢀ
4 ppm (marked R1, R2, R3 and R4, respectively) were identified and monitored together with the Li resonance of the electrolyte and that
of the SEI centred around 0 ppm. The high Liꢀshield phase (R4) was attributed to an overꢀlithiated Li15+ƟGe phase.
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Table 1 Summary of in-situ Li NMR observations
Resonance
index
Assigned
phase
Chemiꢀ
cal
Potential
capacity*
/
Crystal property
Observed by
shift/ppm
R1
R2
R3
Li2.26Ge
Li Ge
24
0.17/839
0.14/1086
0.11/1325
Amorphous
NMR
(inꢀsitu
7
7
7
4
2
Li)
NMR
Li)
NMR
Li),
inꢀsitu XRD
NMR
(inꢀsitu XRD and
13
10
TEM)
(inꢀsitu
3
.5
Li Ge
4
Crystalline phase
(inꢀsitu
1
5
42
(
inꢀsitu XRD and
68,69
42
TEM
)
R4
Li15+ƟGe
ꢀ24
0.005/1653
Overꢀlithiated
(inꢀsitu
7
Li)
4
*
The capacities of carbon lithiation and the initial formation of SEI passivation layers have been subtracted.
(
Fig.2 XRD4) did not observe any crystalline phases and loss of R3 intensity, although exꢀsitu XRD confirmed that an
hence we denote this as an amorphous phase. Upon dischargꢀ
ing to 0.11 V, a new resonance R3 peak appeared at the disꢀ
charging capacity of 1325 mAh g (3.6 Li inserted per Ge).
According to XRD (Fig. 2) and TEM observations, this phase
should be assigned to Li Ge (3.75 Li inserted per Ge) .
After careful investigation of the NMR spectra, we did not
find a pronounced peaking point for R2, indicating that transꢀ
formation from Li2.26Ge to Li Ge through Li Ge may be a
nonꢀlinear process. At the end of the discharge, a new resoꢀ
nance appeared at ꢀ24 ppm (R4) at 0.07 V, accompanied by a
4
2
Li Ge crystalline structure was still present (Fig. 2). Peak
15 4
R4 was strongest in intensity at the end of the discharge, at
~1653 mAh g . We assigned R4 to the overꢀlithiated phase
ꢀ1
+
ꢀ1
Li15+ƟGe4 (0 < Ɵ < 1, according to the second discharging
+
42
ꢀ1 18
capacity of 1468 mAh g . ). This assignment is analogous to
1
5
4
52
overꢀlithiated lithium silicide phases found previously. At
the end of the first discharge, we found that R2 reꢀappeared in
the spectra accompanied by the vanishing of R3. We believe
that the conversion of the aꢀLi Ge phase to the cꢀLi Ge
15
4
3.5
3.5
15
4
phase is incomplete during lithiation. The stoichiometric phase
5
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Journal of the American Chemical Society
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of Li Ge at 0.11 V consists of both Li Ge and Li Ge phasꢀ ner, suggesting the robustness of the phase transformation; this
3.6
3.5
15
4
1
2
3
4
5
6
7
8
9
1
1
1
1
1
1
1
1
1
1
2
2
2
2
2
2
2
2
2
2
3
3
3
3
3
3
3
3
3
3
4
4
4
4
4
4
4
4
4
4
5
5
5
5
5
5
5
5
5
5
6
es, suggesting that there may be a kinetic barrier to Li Ge
4
was in good agreement with the highly reversible electroꢀ
chemical performance observed previously. The formation of
15
(crystalline phase) formation from Li Ge (amorphous
3
.5
52
phase). We believe that formation of crystalline cluster phase
would require not only GeꢀGe bond breakage but also the rearꢀ
rangement of Ge (Ge migration), which involves a large active
the overꢀlithiated phase Li15+ƟGe and the amorphous phase
4
Li Ge seemed to be a terminal process once the cell reached
3.5
the end of the third discharge, as no significant change was
observed in the spectra even when the cell potential was mainꢀ
tained at 0.005 V for 1 hour (Fig. 5(ii)). It is important to note
that both phases vanished quickly once the electrical contact
was removed (Fig. 5 (iii)–(v)), indicating that these two phases
are electrically driven, metastable phases. The good reversibilꢀ
64
barrier. At the end of lithiation, the electrode material should
be a mixture of the aꢀLi Ge phase and the overꢀlithiated
3.5
phase cꢀLi15+ƟGe (see Fig.1e, a TEM image of the grid disꢀ
4
charged to 0.005 V) rather than single crystalline Li Ge or
1
5
4
4
2,51,52
overꢀlithiated Li15+ƟGe4 phases.
The existence of aꢀ
0
1
2
3
4
5
6
7
8
9
0
1
2
3
4
5
6
7
8
9
0
1
2
3
4
5
6
7
8
9
0
1
2
3
4
5
6
7
8
9
0
1
2
3
4
5
6
7
8
9
0
7
Li Ge may further reduce the kinetic barrier for the transforꢀ
ity in the charge/discharge cycles observed using Li NMR
3
.5
64
mation of cꢀLi15+ƟGe to aꢀLi Ge upon delithiation, which
was consistent with the observed battery performance, and we
can infer that the ability of the Ge@CNT composite to attain a
high capacity may be associated with the presence of the overꢀ
lithiated phase.
4
x
helps to reduce the ‘overꢀpotential’ between lithiation and
delithiation.
In the charging process, the R4 peak disappeared completely
by 0.42 V while the R2 resonance became pronounced. Before
the end of delithiation, the resonance shifted to high frequenꢀ
cies at R1, indicating a rapid delithiation process in a short
potential window (see the sharp oxidation peak in CVs in
Fig.S 6a). In the second cycle (Fig.4b), the phase transforꢀ
mation appeared to be similar to the first cycle except that R1
(Li2.26Ge) occurred at a somewhat more positive potential durꢀ
The electrically driven, reversible interꢀconversion and coꢀ
existence of amorphous Li Ge and crystalline Li Ge /overꢀ
3.5
15
4
lithiated Li15+ƟGe phase is a key electrochemical process unꢀ
4
derpinning the superior battery performance of Ge@CNT. The
existence of the amorphous phase reduces the activation enerꢀ
gy for lithium insertion and facilitates the reꢀcrystallization of
+
Li Ge and extra insertion of Li to form overꢀlithiated phase.
15
4
ing lithiation (0.25 V) and R4 (Li15+ƟGe ) was converted to R2
The formation of the overꢀlithiated phase allows the capacity
of Ge@CNT to exceed the theoretical limitation of 1385 mAh
4
(
Li Ge) more rapidly upon delithiation (0.005–0.37V). These
3.5
ꢀ
1 42
g , which is based on Li Ge as the terminal phase. Li Ge
differences may be attributed to the amorphous property of Ge
upon delithiation (see Fig. 1g, h, i), which speeds up the reacꢀ
15
4
15
4
is an unusual electronꢀdeficient phase which lacks one elecꢀ
+
3.75ꢀ 51
5
2,70
tron per formula unit (Li 15Ge 4), thus it can accommodate
tion kinetics of (de)lithiation.
+
extra charges. It has been reported that extra Li can be stored
based on a “job sharing” mechanism in the crystalline defects
71
or amorphousꢀcrystalline interface. Upon delithiation, there
exists a kinetic barrier for the formation of aꢀGe from crystalꢀ
line Li Ge , since GeꢀGe bonds have been disrupted during
15
4
the alloying. The existence of aꢀLi Ge (with some residual
3.5
GeꢀGe clusters) as well as sparsely distributed crystalline Ge
nanoꢀdomains (observed in TEM) help to nucleate larger Ge
domains and reduce the overpotential needed in the phase
6
4
conversion. Along with enhanced electrical conductivity
provided by CNT, the improved ionꢀassociated reaction kinetꢀ
ics facilitates reversibility of Ge@CNT especially under high
current operation.
CONCLUSIONS
In conclusion, we have synthesized Ge NRs encapsulated by
bamboo type multiwall CNTs. Ge@CNT delivers high capaciꢀ
ty, superior rate capability (discharge in seconds) and good
cycle stability. The unique structure of the CNT scaffold meꢀ
chanically protects the Ge NRs from drastic volume swing
during (de)lithiation processes and provides an effective modꢀ
el for multiꢀprobe studies (TEM, XRD and inꢀsitu NMR).
Structural studies using TEM using coin cell configuration, Xꢀ
Figure 5. Stacked inꢀsitu 7Li NMR spectra of the Ge@CNT
composite for the ‘electricalꢀoff’ experiment at the end of the
third galvanostatic charge/discharge cycle. Once discharging had
been completed at the end of the third cycle, the cell was mainꢀ
tained at 0.005 V for 1 h. After the external electrical contact was
7
ray diffraction and inꢀsitu Li NMR studies reveal that the
reversibility of Li (de)lithiation in Ge@CNT during cycling is
mediated by coꢀexisting amorphous and crystalline phases.
The high capacity observed may be related to electrically drivꢀ
en, metastable, overꢀlithiated LiꢀGe alloy, whose existence
and reversibility depends on robust electrical interfaces afꢀ
forded by the carbon walls encapsulating the Ge. The design
and synthesis of such coreꢀshell structure affords a generic
strategy for protecting structurally unstable alloy phases in
high energy and power density LIBs.
then removed, the overꢀlithiated Li15+ƟGe phase vanished immeꢀ
4
diately. The broadening and positive shifting of the overꢀlithiated
phase peak may be attributed to the amorphisation of Ge (see Fig.
1
g,h,i,), which may speed the reaction kinetics of lithiaꢀ
tion/delithiation.
7
Interestingly, peaks R2 and R4 in the Li NMR spectra at
0.005 V at the end of the second and third lithiations (0.005 V,
Fig. 5(i)) appeared to interconvert in a highly reversible manꢀ
EXPERIMENTAL SECTION
6
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Journal of the American Chemical Society
Details on the growth, electrochemical characterisation, XRD,
TEM and inꢀsitu NMR studies are provided in the supporting
information.
(
3) Xu, Y.; Liu, Q.; Zhu, Y.; Liu, Y.; Langrock, A.; Zachariah, M.
1
2
3
4
5
6
7
8
9
1
1
1
1
1
1
1
1
1
1
2
2
2
2
2
2
2
2
2
2
3
3
3
3
3
3
3
3
3
3
4
4
4
4
4
4
4
4
4
4
5
5
5
5
5
5
5
5
5
5
6
R.; Wang, C. Nano Lett. 2013, 13, 470.
7
(4) Seng, K. H.; Park, M.ꢀH.; Guo, Z. P.; Liu, H. K.; Cho, J. Angew.
Chem. Int. Ed. 2012, 51, 5.
NMR experiments The inꢀsitu Li NMR experiments were
performed in a VarianꢀAgilent 300 MHz NMR spectrometer,
7
corresponding to Li Larmor frequencies of 116.57 MHz. A
7
(5) Kim, H.; Son, Y.; Park, C.; Cho, J.; Choi, H. C. Angew. Chem.
Int. Ed. 2013, 52, 5997.
home made inꢀsitu Li NMR probe with a fiveꢀturn solenoid
radiofrequency coil (internal diameter 12 mm), where the exꢀ
citation magnetic field (B ) was perpendicular to the main
1
(
6) Cui, G.; Gu, L.; Zhi, L.; Kaskhedikar, N.; van Aken, P. A.;
magnetic field (B ), was used for the measurement. 1M LiCl
0
in distilled water was used as the reference (0 ppm). Each
spectrum was acquired using a single pulse with a tip angle of
Müllen, K.; Maier, J. Adv. Mater. 2008, 20, 3079.
0
1
2
3
4
5
6
7
8
9
0
1
2
3
4
5
6
7
8
9
0
1
2
3
4
5
6
7
8
9
0
1
2
3
4
5
6
7
8
9
0
1
2
3
4
5
6
7
8
9
0
(7) Chan, C. K.; Zhang, X. F.; Cui, Y. Nano Lett. 2007, 8, 307.
4
5° and pulse width of 4 ꢁs, an acquisition time of 30 ms, a
recycle delay time of 1 s and a total accumulation number of
56, corresponding to a total measurement time of about 4.4
(
8) Xue, D. J.; Xin, S.; Yan, Y.; Jiang, K.ꢀC.; Yin, Y.ꢀX.; Guo, Y.ꢀG.;
2
Wan, L.ꢀJ. J. Am. Chem. Soc. 2012, 134, 2512.
minutes for each spectrum. For the specified charge–discharge
current (0.18 mA) discussed below, a series of spectra were
acquired consecutively.
(9) Kennedy, T.; Mullane, E.; Geaney, H.; Osiak, M.; O’Dwyer, C.;
Ryan, K. M. Nano Lett. 2014, 14, 716.
The inꢀsitu NMR battery cell was a planar Li halfꢀcell battery
with the Ge@CNT composite as the working electrode and Li
metal foil as the counterꢀelectrode. The composition and the
preparation of the working electrode were exactly as described
above. More details are provided in the supporting inforꢀ
mation.
(
10) Song, T.; Cheng, H.; Town, K.; Park, H.; Black, R. W.; Lee, S.;
Park, W. I.; Huang, Y.; Rogers, J. A.; Nazar, L. F.; Paik, U. Adv.
Funct. Mater. 2014, 24, 1458.
(
11) Magasinski, A.; Dixon, P.; Hertzberg, B.; Kvit, A.; Ayala, J.;
Yushin, G. Nat. Mater. 2010, 9, 353.
ASSOCIATED CONTENT
Supporting Information. Detailed synthetic process, electroꢀ
chemical performance and quantitative analysis of the phase
compositions based on inꢀsitu NMR has been provided in supportꢀ
ing information. This material is available free of charge via the
Internet at http://pubs.acs.org.”
(12) Liu, N.; Lu, Z.; Zhao, J.; McDowell, M. T.; Lee, H.ꢀW.; Zhao,
W.; Cui, Y. Nat. Nanotechnol. 2014, 9, 187.
(13) Guo, J.; Chen, X.; Wang, C. J. Mater. Chem. 2010, 20, 5035.
(
14) Luo, J.; Zhao, X.; Wu, J.; Jang, H. D.; Kung, H. H.; Huang, J. J.
AUTHOR INFORMATION
Phys. Chem. Lett. 2012, 3, 1824.
Corresponding Author
(15) Chen, X.; Li, X.; Ding, F.; Xu, W.; Xiao, J.; Cao, Y.; Meduri,
P.; Liu, J.; Graff, G. L.; Zhang, J.ꢀG. Nano Lett. 2012, 12, 4124.
*
Kian Ping Loh, Email: chmlohkp@nus.edu.sg; Jian Zhi Hu,
Email: Jianzhi.Hu@pnnl.gov; Ming Lin, Email: mꢀlin@imre.aꢀ
star.edu.sg.
(
16) Yi, R.; Dai, F.; Gordin, M. L.; Sohn, H.; Wang, D. Adv. Energy
Mater. 2013, 3, 1507.
17) Ge, M.; Lu, Y.; Ercius, P.; Rong, J.; Fang, X.; Mecklenburg, M.;
Zhou, C. Nano Lett. 2013, 14, 261.
18) Park, C. M.; Kim, J. H.; Kim, H.; Sohn, H. J. Chem. Soc. Rev.
010, 39, 3115.
Author Contributions
The manuscript was written through contributions of all authors. /
All authors have given approval to the final version of the manuꢀ
script. /
(
(
ACKNOWLEDGMENT
2
MOE Tier II grant “Interface Engineering of Graphene Hybrids
for Energy Conversion.” Grant Number: Rꢀ143ꢀ000ꢀ488ꢀ112.
(19) Tarascon, J. M.; Armand, M. Nature 2001, 414, 359.
(20) Goodenough, J. B. Energy Environ. Sci. 2014, 7, 14.
7
Inꢀsitu Li NMR experiments were performed in the Environmenꢀ
tal Molecular Sciences Laboratory, a national scientific user facilꢀ
ity sponsored by the DOE’s Office of Biological and Environꢀ
mental Research, and located at Pacific Northwest National Laꢀ
boratory (PNNL). PNNL is a multiꢀprogram national laboratory
operated for the DOE by Battelle Memorial Institute under Conꢀ
tract DEꢀAC06ꢀ76RLO 1830. The authors receive assistance from
Dr. Xilin Chen and Dr. Xiaolin Li in PNNL, Ms Lili Liu and Prof
Yuping Wu in Fudan University for conducting ballꢀmilling and
and Dr Tan Teck Leong in Institute of High Performance Compuꢀ
ting (Singapore) for theoretical calculations.
(
21) Yang, Z.; Zhang, J.; KintnerꢀMeyer, M. C. W.; Lu, X.; Choi, D.;
Lemmon, J. P.; Liu, J. Chem. Rev. 2011, 111, 3577.
22) Yang, S.; Bachman, R. E.; Feng, X.; Müllen, K. Acc. Chem. Res.
012, 46, 116.
23) Thackeray, M. M.; Wolverton, C.; Isaacs, E. D. Energy Environ.
Sci. 2012, 5, 7854.
24) Wang, Y.; Wang, Y.; Hosono, E.; Wang, K.; Zhou, H. Angew.
Chem. Int. Ed. 2008, 47, 7461.
(
2
(
(
REFERENCES
(
(
25) Park, M.; Kim, K.; Kim, J.; Cho, J. Adv. Mater. 2010, 22, 415.
26) Seng, K. H.; Park, M.ꢀh.; Guo, Z. P.; Liu, H. K.; Cho, J. Nano
(
1) Wang, Y.; Lee, J. Y. Angew. Chem. Int. Ed. 2006, 45, 7039.
(2) Zhang, W. M.; Hu, J. S.; Guo, Y.ꢀG.; Zheng, S.ꢀF.; Zhong, L.ꢀS.;
Lett. 2013, 13, 1230.
Song, W.ꢀG.; Wan, L.ꢀJ. Adv. Mater. 2008, 20, 1160.
7
ACS Paragon Plus Environment

Journal of the American Chemical Society
Page 8 of 10
(
27) Sun, X.; Chen, T.; Yang, Z.; Peng, H. Acc. Chem. Res. 2012, 46,
39.
(49) Chevallier, F.; Letellier, M.; Morcrette, M.; Tarascon, J.ꢀM.;
1
2
3
4
5
6
7
8
9
1
1
1
1
1
1
1
1
1
1
2
2
2
2
2
2
2
2
2
2
3
3
3
3
3
3
3
3
3
3
4
4
4
4
4
4
4
4
4
4
5
5
5
5
5
5
5
5
5
5
6
5
Frackowiak, E.; Rouzaud, J.ꢀN.; Béguin, F. Electrochem. SolidꢀState
Lett. 2003, 6, A225.
(28) Lv, D.; Gordin, M., L.; Yi, R.; Xu, T.; Song, J.; Jiang, Y.; Choi,
D.; Wang, D. Adv. Funct. Mater. 2014, 24, 1059.
(50) Letellier, M.; Chevallier, F.; Béguin, F. J. Phys. Chem. Solids
2006, 67, 1228.
(
29) Ren, J. G.; Wu, Q. H.; Tang, H.; Hong, G.; Zhang, W.; Lee, S.
T. J. Mater. Chem. A 2013, 1, 1821.
30) Liu, J.; Song, K.; Zhu, C.; Chen, C.ꢀC.; van Aken, P. A.; Maier,
(51) Key, B.; Bhattacharyya, R.; Morcrette, M.; Seznéc, V.;
Tarascon, J.ꢀM.; Grey, C. P. J. Am. Chem. Soc. 2009, 131, 9239.
(52) Ogata, K.; Salager, E.; Kerr, C. J.; Fraser, A. E.; Ducati, C.;
Morris, A. J.; Hofmann, S.; Grey, C. P. Nat. Commun. 2014, 5, 3217.
(53) Trill, J.ꢀH.; Tao, C.; Winter, M.; Passerini, S.; Eckert, H. J. Solid
State Electrochem. 2011, 15, 349.
(
J.; Yu, Y. ACS Nano 2014, 8, 7051.
0
1
2
3
4
5
6
7
8
9
0
1
2
3
4
5
6
7
8
9
0
1
2
3
4
5
6
7
8
9
0
1
2
3
4
5
6
7
8
9
0
1
2
3
4
5
6
7
8
9
0
(31) Cho, Y. J.; Im, H. S.; Kim, H. S.; Myung, Y.; Back, S. H.; Lim,
Y. R.; Jung, C. S.; Jang, D. M.; Park, J.; Cha, E. H.; Cho, W. I.;
Shojaei, F.; Kang, H. S. ACS Nano 2013, 7, 9075.
(32) Cui, G.; Gu, L.; Kaskhedikar, N.; van Aken, P. A.; Maier, J.
Electrochim. Acta 2010, 55, 985.
(54) Chen, H.; Hao, Q.; Zivkovic, O.; Hautier, G.; Du, L.ꢀS.; Tang,
Y.; Hu, Y.ꢀY.; Ma, X.; Grey, C. P.; Ceder, G. Chem. Mater. 2013, 25,
2777.
(
33) Li, W.; Yang, Z.; Cheng, J.; Zhong, X.; Gu, L.; Yu, Y.
Nanoscale 2014, 6, 4532.
(55) Wang, W. D.; Lucier, B. E. G.; Terskikh, V. V.; Wang, W.;
Huang, Y. J. Phys. Chem. Lett. 2014, 5, 3360.
(
(
34) Yuan, F.ꢀW.; Yang, H.ꢀJ.; Tuan, H.ꢀY. ACS Nano 2012, 6, 9932.
35) Seo, M. H.; Park, M.; Lee, K. T.; Kim, K.; Kim, J.; Cho, J.
(56) Feng, N.; Wang, Q.; Zheng, A.; Zhang, Z.; Fan, J.; Liu, S.ꢀB.;
Amoureux, J.ꢀP.; Deng, F. J. Am. Chem. Soc. 2013, 135, 1607.
(57) Huo, H.; Chamas, M.; Lippens, P.ꢀE.; Ménétrier, M. J. Phys.
Chem. C 2012, 116, 2390.
Energy Environ. Sci. 2011, 4, 425.
(
(
36) Yang, H.ꢀJ.; Tuan, H.ꢀY. J. Mater. Chem. 2012, 22, 2215.
37) Wang, D.; Chang, Y.ꢀL.; Liu, Z.; Dai, H. J. Am. Chem. Soc.
2005, 127, 11871.
38) Tuan, H.ꢀY.; Lee, D. C.; Korgel, B. A. Angew. Chem. 2006, 118,
308.
(39) Yuan, F.ꢀW.; Tuan, H.ꢀY. Chem. Mater. 2014, 26, 2172.
40) Pandurangan, A.; Morin, C.; Qian, D.; Andrews, R.; Crocker,
M. Carbon 2009, 47, 1708.
41) Weker, J. N.; Liu, N.; Misra, S.; Andrews, J. C.; Cui, Y.; Toney,
M. F. Energy Environ. Sci. 2014, 7, 2771.
42) Baggetto, L.; Notten, P. H. L. J. Electrochem. Soc. 2009, 156,
A169.
43) Yoon, S.; Park, C.ꢀM.; Sohn, H.ꢀJ. Electrochem. SolidꢀState
Lett. 2008, 11, A42.
44) Lin, F.; Nordlund, D.; Weng, T.ꢀC.; Zhu, Y.; Ban, C.; Richards,
R., M.; L.Xin, H. Nat. Commun. 2014, 5, 3358.
45) Liu, X. H.; Huang, S.; Picraux, S. T.; Li, J.; Zhu, T.; Huang, J.
Y. Nano Lett. 2011, 11, 3991.
46) Huang, J. Y.; Zhong, L.; Wang, C. M.; Sullivan, J. P.; Xu, W.;
(58) Hata, K.; Futaba, D. N.; Mizuno, K.; Namai, T.; Yumura, M.;
Iijima, S. Science 2004, 306, 1362.
(
5
(59) Yan, C.; Xi, W.; Si, W.; Deng, J.; Schmidt, O. G. Adv. Mater.
2013, 25, 539.
(
(60) Hwang, I. S.; Kim, J.ꢀC.; Seo, S.ꢀD.; Lee, S.; Lee, J.ꢀH.; Kim,
D.ꢀW. Chem. Commun. 2012, 48, 7061.
(
(61) Lee, H.; Kim, H.; Doo, S.ꢀG.; Cho, J. J. Electrochem. Soc. 2007,
154, A343.
(
(62) Lin, F.; Markus, I. M.; Nordlund, D.; Weng, T.ꢀC.; Asta, M. D.;
Xin, H. L.; Doeff, M. M. Nat. Commun. 2014, 5, 3529.
(63) Zhang, L. Q.; Liu, X. H.; Liu, Y.; Huang, S.; Zhu, T.; Gui, L.;
Mao, S. X.; Ye, Z. Z.; Wang, C. M.; Sullivan, J. P.; Huang, J. Y. ACS
Nano 2011, 5, 4800.
(
(
(64) Key, B.; Morcrette, M.; Tarascon, J.ꢀM.; Grey, C. P. J. Am.
Chem. Soc. 2010, 133, 503.
(
(65) Rangeet, B.; Baris, K.; Hailong, C.; Adam, S. B.; Anthony, F.
H.; Clare, P. G. Nat. Mater. 2010, 9, 504.
(
Zhang, L. Q.; Mao, S. X.; Hudak, N. S.; Liu, X. H.; Subramanian, A.;
(66) Cattaneo, A. S.; Dupke, S.; Schmitz, A.; Badillo, J. P.; Winter,
M.; Wiggers, H.; Eckert, H. Solid State Ionics 2013, 249–250, 41.
(67) Dinges, T.; Hoffmann, R.ꢀD.; van Wüllen, L.; Henry, P.; Eckert,
H.; Pöttgen, R. J. Solid State Electrochem. 2011, 15, 237.
(68) Liu, X. H.; Zhang, L. Q.; Zhong, L.; Liu, Y.; Zheng, H.; Wang,
J. W.; Cho, J.ꢀH.; Dayeh, S. A.; Picraux, S. T.; Sullivan, J. P.; Mao, S.
X.; Ye, Z. Z.; Huang, J. Y. Nano Lett. 2011, 11, 2251.
Fan, H.; Qi, L.; Kushima, A.; Li, J. Science 2010, 330, 1515.
(
47) Liu, X. H.; Wang, J. W.; Huang, S.; Fan, F.; Huang, X.; Liu, Y.;
Krylyuk, S.; Yoo, J.; Dayeh, S. A.; Davydov, A. V.; Mao, S. X.;
Picraux, S. T.; Zhang, S.; Li, J.; Zhu, T.; Huang, J. Y. Nat.
Nanotechnol. 2012, 7, 749.
(
48) Gu, M.; Yang, H.; Perea, D. E.; Zhang, J.ꢀG.; Zhang, S.; Wang,
C. Nano Lett. 2014, 14, 4622.
8
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Journal of the American Chemical Society
(
69) Liang, W.; Yang, H.; Fan, F.; Liu, Y.; Liu, X. H.; Huang, J. Y.;
(71) Maier, J. Angew. Chem. Int. Ed. 2013, 52, 4998.
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(70) McDowell, M. T.; Lee, S. W.; Nix, W. D.; Cui, Y. Adv. Mater.
013, 25, 4966.
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