li alloying nanomaterials - greenlionproject.eu · where m represents a group iv alloying...
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EUROPEAN LI-ION BATTERY ADVANCED
MANUFACTURING FOR ELECTRIC VEHICLES
Li Alloying Nanomaterials Driving anode performance …
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Li Alloying Nanomaterials
Driving anode performance …
Introduction
There is a need to develop more powerful, smaller and lighter lithium ion batteries (LIBs) with longer
discharge times, in order to satisfy the requirement for longer range and better driving performance in
electric vehicles. Advances in LIB performance through improved design and manufacturing have now
reached a limit imposed by the fundamental chemistry of the cell components.[1, 2] Therefore the further
required developments in LIB cell technology will have to be achieved largely through innovation in
anode, cathode and electrolyte materials.
Despite extensive research, the charge storage capacity of practical cathode materials remains low (< 200
mAh.g-1). In view of this, increasing the anode capacity is considered to be the most promising medium
term approach to achieving overall LIB performance improvements. Carbon (especially graphite) is
presently the ubiquitous choice of anode material, but it also has a rather limited theoretical gravimetric
capacity of 372 mAh.g-1. Amongst the most attractive graphite alternatives are the group IV elements
Silicon (Si), Germanium (Ge) and Tin (Sn), which offer room-temperature Li charge storage capacities of
3579 mAh.g-1, 1348 mAh.g-1 and 994 mAh.g-1, respectively. [2-4] In contrast to graphite, which stores
charge by intercalating Li atoms between its constituent graphene layers in a 1:6 lithium to carbon atomic
ratio, these group IV elements electrochemically alloy Li according to the general reaction,
M + xLi+ + xe- ↔ LixM (0 ≤ x ≤ 3.75)
where M represents a group IV alloying element.[5] This equation implies a 3.75:1 lithium to alloying
element atomic ratio at full charge, thereby accounting for the superior capacity achievable by alloying
compared to intercalation.
Despite their impressive theoretical capacity values, lithium alloying elements have, as yet, failed to find
application in practical LIB anodes. This arises principally from the large degree of volume
expansion/contraction (> 300% in the case of Si) that accompanies their lithiation/delithiation,
respectively. Such volume changes lead to a variety of problems for conventional anode formulations,[6]
which consist of micron-scale active particles mixed with sub-micron conductive carbon particles held
together by a polymeric binder. These negative effects may be classified into three main groupings: (i)
whole electrode level, (ii) individual particle level, and (iii) solid-electrolyte interface issues.
At the whole-electrode level, the repeated expansion and contraction of the active mass leads to the loss of
electrical contact between the particles and additionally can cause cracking of the electrode film and/or
its delamination from the separator. The polymeric binders used in graphitic anodes (where the volume
change during charge/discharge is only approximately 10%) have proven to be incapable of
accommodating the volume changes inherent to lithium alloying electrodes, while a viable alternative
binder with the desired properties has yet to be identified. At single particle level, expansion upon
lithiation causes the build-up of internal stresses, which may be relieved through crack formation and
propagation. The continuous occurrence of this process over a number of cycles leads to particle
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pulverisation, again resulting in the loss of electrical contact between significant proportions of the active
matter. This is in turn manifested as a rapid decrease in anode capacity with cycling.
Particle expansion and shrinkage also leads to the undesirable phenomenon of continuous solid-
electrolyte interface (SEI) layer formation. SEI layers form on anode surfaces owing to the decomposition
of electrolyte at potentials below ca. 1 V vs. Li/Li+. Formation of a stable SEI layer during the first few
charge cycles can passivate against further electrolyte breakdown. However, where large particle volume
changes occur it is difficult to form a stable layer. The SEI that forms as the particle expands upon
lithiation is no-longer supported as the particle contracts with delithiation. It therefore breaks and
crumbles , leaving the particle surface exposed for further electrolyte decomposition on the next charge.
This leads to low Coulombic efficiencies and electrolyte depletion.
It has been proposed that utilising nanoscale forms of lithium alloying elements could alleviate the
fracture and pulverisation issues, since the total elastic energy stored in a nanostructure during
deformation may not be sufficient to cause crack development.[7] Anodes prepared with nanoparticles of
lithium alloying elements and conventional binders have indeed shown superior capacity retention
relative to those utilising microscale particles of the same material, however their performances remain
far short of commercial thresholds. An alternative approach involves combining nanoscale lithium
alloying materials with innovative anode designs. In this regard, a significant advance came in 2007 when
it was demonstrated that gold seeded silicon nanowires (NWs) grown directly onto a stainless steel
current collector by chemical vapour deposition (CVD) could maintain a capacity of over 3000 mAh.g-1 for
10 charge/discharge cycles.[8] In addition to the resistance of the NWs to elastic strain, the retention of
high capacity over a number of cycles was attributed to the space available between adjacent wires to
facilitate expansion, and to the robust electrical connection between each wire and the current collector. A
further advantage of this architecture is that it dispenses with inactive components such as binders and
conductive additives. The same workers subsequently reported a similarly prepared, gold seeded
germanium NW anode that held a capacity of approximately 1000 mAh.g-1 over 20 cycles, albeit at a
modest C/20 rate.[9] While these early NW anodes did not offer acceptable performance over extended
cycles, they did inspire further research into nanostructured, group IV element, based electrode
architectures. Accordingly, improved capacity retention has been reported for a range of anodes based on
more sophisticated nanoscopic Si or Ge based materials.[6, 10-12] These include carbon sheathed Ge[13]
and Si[14] NWS, Ge[15] and Si[16] nanotubes, carbon-fibre core/ Si shell NWs,[17] graphene supported
Ge and Si NWs[18] and Si/carbon yolk-shell structures.[19] However, despite their scientific interest,
most of these materials are unsuitable for practical application, because their production is complex or
limited to a small scale (e.g. templated growth). Indeed, cost and scalability concerns also mitigate against
the adoption of techniques such as conventional CVD in the battery industry, even for the production of
relatively simple solid NW anodes.
In view of this, our goal in the Greenlion Project, was to develop lower cost and more scalable routes to
the fabrication of high performance Ge and Si NW anodes. A solvent vapour growth method is outlined
herein, which dispenses with the use of costly, low-throughput CVD equipment, and also replaces gold
with less expensive catalyst metals (tin or copper). It is also demonstrated that the performance of
electrodes produced by these methods far outstrips that of the most comparable NW anodes reported
prior to the start of Greenlion.
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Greenlion breakthrough
1/ Production and characterisation of nanowire anodes
1.1/ Solvent vapour growth of Si and Ge nanowires
A variety of nanowire anode systems were synthesised and electrochemically evaluated during the course
of the Greenlion project. Tin (Sn) seeded Si or Ge NWs were grown by a vapour-liquid-solid (VLS)
mechanism, the concept of which is depicted schematically in Figure 1. The Si or Ge containing precursor
is injected into a reactor, at a controlled temperature above 400°C. Here it thermally decomposes to yield
SiH4 or GeH4 in the vapour (V) phase, which are the monomers for NW growth. During the ramp up to
this temperature, a Sn (melting point: 232°C) layer, previously evaporated onto the stainless steel current
collector, melts to form discrete islands. The monomer infuses into these liquid (L) phase Sn islands
causing their supersaturation with Si or Ge atoms and the extrusion of a solid (S) NW. Copper germanide
(Cu3Ge) seeded Ge NWs are similarly formed by a vapour-solid-solid (VSS) mechanism, the only
difference being that the Cu3Ge seeding material exists in the form of solid (S) nanoparticles at the NW
growth temperature.
Fig. 1: Schematic representation of the VLS growth of Sn seeded Ge NWs on a stainless steel current
collector by the solvent vapour method. Adapted from Kennedy et al.[20]
Most usually VLS or VSS growth of NWs is achieved through conventional chemical vapour deposition
(CVD) which requires expensive equipment, but only offers low yields.[21] However the solvent vapour
growth (SVG) method, developed and applied during Greenlion, uses a simple glassware reactor (as in
Figure 1), or a stainless steel confiner, and is thus adaptable to upscaling. In contrast to traditional CVD,
where the reaction medium is a heated inert gas stream, the vapour phase of a high boiling point solvent
facilitates NW growth in the SVG protocol. This simple implementation of VLS or VSS NW growth offers
a direct route to the fabrication of nanostructured lithium alloying element anodes, where electrical and
mechanical connection to the current collector is accomplished, in-situ, during the production of the
active material, without the need for binders or conductive additives.
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1.2/ Sn seeded Ge or Si nanowire anodes
These types of NW anodes were produced by solvent vapour growth in a glassware reactor as presented
schematically in Figure 2.
Fig. 2: Schematic illustrating the glassware based apparatus for solvent vapour growth of Si or Ge
NW anodes. Reproduced from Mullane et al.[22]
Fig. 3: a) SEM image showing Sn seeded Ge NWs growing from the stainless steel substrate. b) High
magnification SEM image of a Ge NW with a Sn seed visible at its end. c) TEM image of the Sn/Ge
interface of a NW with inset SAED patterns indexed for cubic Ge (i) and tetragonal Sn (ii)
respectively. d) XRD of Sn seeded Ge NWs. Reproduced from Kennedy et al.[20]
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Prior to NW growth, catalytic Sn layers with an optimal thickness of 20 nm were deposited onto the
stainless steel current collector by thermal evaporation. The current collector substrate was transferred to
a long-necked, round-bottomed flask, where it was placed in a vertical position. A quantity of the high
boiling point solvent squalane (b.p: 176°C) was added and the flask was attached to a schlenk line via a
reflux condenser. This reactor was placed in an upright three-zone furnace, the temperature was ramped
to 125°C and a vacuum was applied to remove any moisture. Subsequent to this, the reactor was purged
with Ar gas and the temperature was raised to the correct reaction temperature (430°C for Ge NWs or
460°C for Si NWs) under a constant Ar flow. Upon attainment of the desired temperature, the liquid
precursor (phenylsilane for Si or diphenylgermane for Ge) was injected into the reactor. The density of
NW growth on the substrate was found to depend on the reaction time. The anodes discussed in the
present document had active material loading densities of 0.22 mg.cm-2 (Ge) and 0.20 mg.cm-2 (Si),
corresponding to reaction times of 10 and 60 minutes respectively, for Ge and Si anodes. The higher
growth temperature and reaction duration required for Si NW growth arises because phenylsilane is a less
reactive precursor than diphenylgermane. The reactions were terminated by opening the furnace and
allowing the setup to cool to room temperature.
Fig. 4: a) SEM image showing Sn seeded Si NWs growing from the stainless steel substrate. b) High
magnification TEM image of a Si/Sn interface with the corresponding low magnification image inset
(i) and an SAED pattern of a region of the Si stem inset (ii). c) Dark Field Scanning TEM image of the
Sn/Si interface from (b) with overlaid EDX line profile. Adapted from Mullane et al.[22]
A scanning electron microscope (SEM) image of Ge NWs produced by the SVG method is presented in
Figure 3 (a). At higher resolution (Figure 3 (b)) a spherical Sn seed can clearly be seen attached to the end
of a Ge NW. A statistical analysis of several SEM images indicated an average NW diameter of 73 nm.
Additionally, the average ratio of seed to wire diameter was approximately 1.75:1, implying a 5:1 mass
ratio of Ge to Sn. The high resolution transmission electron microscopy (TEM) image of Figure 3 (c)
reveals good contact between wire and seed. The selected area electron diffraction (SAED) pattern
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recorded from an area on the Ge wire (Figure 3 (c) - inset i) is indexed with spots corresponding to those
expected for a cubic Ge lattice and are consistent with a <111> growth direction. The observation of
discrete spots confirms the single-crystal nature of the NWs. The SAED pattern for the Sn catalyst (Figure
3 (c) - inset ii) is indexed with spots corresponding to tetragonal Sn. These conclusions were supported by
XRD analysis of a NW array (Figure 3.3 (d)) which yielded reflections consistent with cubic Ge (space
group mFd3 ) and tetragonal Sn (space group I41/amd), with the remaining peaks arising due to the
underlying stainless steel current collector.
The SEM image of Figure 4 (a) shows that, as for Ge case, there is a dense growth of Sn seeded Si NWs on
the stainless steel substrate when phenylsilane precursor is used in the solvent vapour system. These Si
NWs were on average slightly thicker than their Ge counterparts with an observed diameter distribution
of 105 ± 30 nm. The seed to wire diameter ratio was ≈ 2.25:1, corresponding to a 2:1 mass ratio of Si to Sn
in these NWs. The high and low magnification TEM images of Figure 4 (b) and Figure 4 (b) inset (i)
respectively show good attachment across a crystalline interface between a Si wire stem and the Sn seed.
The elemental composition in the junction region is illustrated in Figure 4 (c) by overlaying an energy-
dispersive X-ray (EDX) line profile over a dark-field scanning TEM image of the NW from Figure 4 (b). An
SAED pattern from a region of the Si NW is presented in Figure 4 (b) inset (ii) and is consistent with
single crystalline, diamond cubic Si with a <111> growth direction.
1.3/ Rapid pyrolysis method for nanowire anode production
While the glassware based implementation of the solvent vapour growth method can produce high
quality NW anodes as outlined in section 1.2, its batch nature may ultimately limit its utility in a high
throughput manufacturing environment. With this in mind, an alternative SVG approach, denoted as the
rapid pyrolysis method was developed, targeting greater amenability to upscaling for high volume
production in a semi-continuous, roll-to-roll type process. The operating principal is illustrated
schematically in Figure 5 using the example of VSS growth of Cu3Ge seeded Ge NWs, however it should be
noted that a range of Si or Ge NWS with different seeding metals can be grown by the rapid-pyrolysis
protocol.[23]
Fig. 5: Schematic of the process flow involved in the rapid pyrolysis implementation of solvent
vapour growth of nanowire anodes. Adapted from Mullane et al.[24]
A thin layer of seeding metal is evaporated onto stainless steel current collector foil which is then placed
on a hot-plate surface in an inert gas chamber. A stainless steel 'confiner' chamber is then placed over the
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area where NW growth is targeted and the precursor is injected. The precursor decomposes into the
monomer for NW formation with the hot current collector and VSS or VLS growth occurs as described in
section 1.1. In addition to its obvious compatibility with production line manufacturing, the hot-plate
pyrolysis technique offers fast NW electrode fabrication times (< 2 min) in comparison to the glassware
approach where reactions may take up to 2 hours in total (encompassing solvent degassing, temperature
ramping, reaction time and cool-down).
1.4/ Cu3Ge seeded Ge nanowire anodes by the rapid pyrolysis method
For the production of Cu3Ge seeded Ge NWs, a Cu layer was thermally evaporated onto stainless steel foil,
with a thickness of 1 nm found to be optimum for the growth of wires with good adhesion to the substrate.
The NW growth was conducted in an Ar filled dry glovebox. The Cu coated stainless steel foil was fed onto
a hotplate at 425°C, and left for 1 minute to aquire thermal equilibrium. A stainless steel 'heatsink
confiner' was then pressed around the area to be coated with NWs - see Figure 6. Next, the appropriate
volume of diphenylgermane precursor was dropped onto the substrate. This droplet was allowed to
evaporate, thus terminating the reaction. The finned design of the confiner was found to be effective at
cooling its side walls and thereby reducing loss of vaporised precursor through the injection point before
reaction had occurred.
Fig. 6: Design of the stainless steel 'Heatsink Confiner' used in a laboratory scale implementation of
the rapid pyrolysis method of nanowire anode fabrication.
An SEM image is presented in Figure 7 of a typical Ge NW test anode produced by the rapid pyrolysis
method. The diameter of these wires is approximately half of that exhibited by the Sn seeded Ge NWs
grown in the glassware reactor(section 1.2). The XRD diffractogram of Figure 8 (a) shows reflections that
are consistent with cubic Ge. The underlying SS substrate reflections are also indexed along with the
(200) and (002) reflections from copper germanide (Cu3Ge), suggesting that the latter is the seeding
material. The results of EDX analysis performed in the vicinity of the end of one of these NWs is shown in
Figure 8 (b). The region which shows a high Cu signal and non-zero Ge signal corresponds to the Cu3Ge
seed with the NW composed solely of Ge. High resolution TEM analysis of a typical straight NW in Figure
8 (c) illustrates the <110> growth direction typically seen for Cu3Ge seeded NWs. The inset SAED patterns
(i and ii) are indexed for cubic Ge and orthorhombic Cu3Ge respectively. Considering the evidence
presented in Figure 8, we postulate that the growth process is initiated when the diphenylgermane
precursor decomposes to GeH4 upon contact with the heated growth substrate. This reacts with the 1 nm
thick copper layer to form discrete Cu3Ge catalytic particles on the stainless steel surface.[24] Further
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saturation of these particles with Ge atoms leads to extrusion of Ge NWs in a typical VSS process. The
Cu3Ge seeded Ge NW anodes had an active material loading density of 0.19 mg.cm-2. Owing to the small
size of the seeds, it was calculated that they contribute only approximately 1% of the NW mass.
Fig. 7: SEM image of Ge NWs grown from a Cu coated stainless steel current collector. Adapted from
Mullane et al.[24]
Fig. 8: a) XRD analysis of the NW covered substrate showing reflections consistent with the presence
of Ge NWs, SS current collector and Cu3Ge seeds. b) EDX line profile collected from a Cu3Ge seeded
Ge NW. c) High resolution TEM image of a Ge NW with inset FFTs, i and ii, corresponding to the
crystalline Ge NW and Cu3Ge seed respectively. Adapted from Mullane et al.[24]
Towards the end of Greenlion, a reaction rig was designed and commissioned to implement the rapid
pyrolysis production of NW anodes on a significantly increased scale. This semi-automated apparatus
features an entirely sealed confiner chamber with a diameter of 73 mm, which implies a 65-fold increase
in reaction surface area compared to the initial confiner depicted in Figure 6. This development indicates
that the rapid pyrolysis technique can be successfully up-scaled and marks another milestone in our
continuing commitment to achieve commercialisation of SVG nanowire production.
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1.5/ Ge/Si heterostructured nanowire anodes
As revealed in sections 1.6, 1.8 and 1.9, Si and Ge NWs have distinct strengths when applied as LIB
anodes. While Si is inexpensive and can deliver very high charge storage capacities, Ge offers better rate
capability and superior levels of capacity retention with cycling. In an attempt to combine the positive
attributes of both elements, we developed a novel NW anode architecture consisting of Ge NW 'stems'
with attached Si NW 'branches'. A schematic overview of the production process for such anodes is
presented in Figure 9.
Fig. 9: Schematic of the steps involved in the synthesis of Ge/Si heterostructure NW anodes.
Reproduced from Kennedy et al.[25]
In the first step, Cu3Ge seeded Ge NW 'truncks' were prepared via the rapid pyrolysis method - Figure 10
(a). The Ge NW covered substrate was then soaked in ethanedithiol (EDT) for before immersion in a
colloidal tin suspension. The EDT acts as a linker molecule, affecting the decoration of the Ge NWs with
tin nanoparticles - Figure 10 (b).
Fig. 10: SEM images of a) Cu3Ge seeded Ge NWs. b) Ge NWs decorated with Sn nanoparticles. c) and
d) Si NWs seeded from the Sn nanoparticles growing from the original Ge NWs.
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Growth of the Si NW 'branches' was then achieved using the solvent vapour growth with a phenylsilane
precursor. Here the Sn nanoparticles provided the seeds for the Si NW growth as opposed to the
evaporated Sn layer used in the direct growth of Si wires from the current collector (see section 1.2). The
SEM images in Figures 10 (c) and (d) show that in general the Si 'branches' tend to wrap around the Ge
'stems' rather than radiating outwards. The active mass loading of the heterostructured anodes was
approximately 0.30 mg.cm-2.
1.6/ Half-cell electrochemical performance of Sn seeded Ge nanowire anodes
The long term cycling behaviour of a typical Sn seeded Ge NW anode in a half cell test against a Li counter
electrode is summarised in Figure 11 (a), while selected voltage profiles recorded during this experiment
are included in Figure 11 (b). The mass of both the Sn seed and the Ge NW were taken into account when
calculating the gravimetric capacities, giving a maximum theoretical specific capacity for the composite
anode of 1320 mAh.g-1. Based on this, the anode was charged and discharged at a C/2 rate. The NWs
exhibited an initial discharge capacity of 1103 mAh.g-1 with an average Coulombic efficiency (C.E.) of
97.0%. The electrode retained a reversible capacity of 888 mAh.g-1 after 1100 cycles, with most of the
fading occurring during the first 100 cycles. Beyond this point the capacity dropped by only 0.01% per
cycle. This level of consistent performance for a binder free, solid Ge nanowire based anode is completely
unprecedented, with the previous best comparable report only detailing stable performance up to 50
cycles.[26]
Fig. 11: a) Discharge capacities of Sn seeded Ge NW electrode. The active material was charged and
discharged at a C/2 rate. The electrolyte was 1M LiPF6 in EC/DMC (1:1 v/v) + 3wt% VC. b) Voltage
profiles of 1st, 10th, 50th, 100th, 300th and 1000th cycle of the electrode cycled in a). The profiles show
characteristic plateaus corresponding to the lithiation and delithiation of both Ge and Sn.
Reproduced from Kennedy et al.[20]
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Characteristic plateaus for the lithiation of Sn appear at potentials of 620 mV and 400 mV in the first
charge cycle in Figure 11 (b), before the onset of Ge lithiation at 350 mV. Similarly, plateaus for Sn
delithiation are present at 600 mV, 725 mV and 800 mV on the first discharge. This observation implies
that in addition to Ge, the Sn seeds also reversibly cycle lithium in these anodes. In a more detailed
analysis published elsewhere, [20, 22] we have constructed differential capacity plots from the charge and
discharge curves of Sn seeded Ge NW anodes, which clearly show that Sn contributes to the charge
storage capacity of such electrodes, in both the first and subsequent cycles. The fact that the Sn seed is
electrochemically active contrasts with previously reported on-current collector grown Ge NWs which
were seeded by Au,[9, 26] which is more expensive and doesn't alloy with Li, thereby adding weight but
no charge storage performance to the anode.
The electrolyte additive vinylene carbonate (VC) plays an important role in maintaining the high capacity
performance of the Ge NW anode. Prior to commencing the extended experiment detailed in Figure 11, a
number of additives to the standard 1 M LiPF6 in ethylene carbonate (EC) and dimethylcarbonate (DMC)
electrolyte were screened over a more limited number of cycles. VC was identified as the most promising
of these - the performances of Ge NW anodes cycled in 1M LiPF6 in EC/DMC (1:1 v/v) with, and without, 3
wt% VC are compared in Figure 12. The results show that 80.2% of the initial discharge capacity is
retained after 200 cycles using the VC additive, compared to only 33.1% for the additive free electrolyte.
Fig. 12: Comparison of the capacity data and C.E. values of two Sn seeded Ge NW electrodes using
two electrolytes, one with VC (1M LiPF6 in EC/DMC + 3wt% VC) and one without VC (1M LiPF6 in
EC/DMC). The electrodes were cycled at a 1C rate. Reproduced from Kennedy et al. [20]
The C.E. also benefits, with the VC containing electrolyte exhibiting 99.5% after 200 cycles, compared to
96.3% for its VC-free equivalent. While nanostructuring of lithium alloying elements may alleviate
expansion/contraction driven material and electrode degradation, the issue of continuous SEI layer
formation, identified in the Introduction, persists. It is this problem, essentially chemical in origin, that is
addressed by the VC additive. VC is known to produce a more durable and cohesive SEI layer, preventing
cracking and continuous re- exposure of the active material to the electrolyte with each cycle.[27, 28] The
positive impact of VC on the cycling performance in Figure 12 suggests, that for these Ge NWs, the
minimisation of repeated surface exposure and SEI layer growth has a beneficial effect on the
morphological evolution of the active matter.
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The rate capability of the Sn seeded Ge NWs was tested by charging and discharging an anode for 5 cycles
each at rates of C/10, C/5, C/2, 1C and 2C in a sequential experiment (Figure 13 (a)). Average discharge
capacities of 1250, 1174, 1050, 821 and 722 mAh.g-1 were observed at these rates, respectively.
Fig. 13: Rate capability data for Sn seeded Ge NW anodes cycled in 1M LiPF6 in EC/DMC + 3wt% VC
electrolyte. a) Capacities and C.E. values, where charge and discharge were conducted at the same
rate for 5 successive cycles. b) As for a) but with higher charge/discharge rates. c) Discharge
capacities measured for 5 cycles at each of 6 different discharge rates. The charge rate was C/2 for all
cycles. d) 20C and 100C discharge capacities for anodes charged at a 2C rate. Reproduced from
Kennedy et al. [20]
The rate capability analysis was extended to higher C rates (Figure 13 (b)). The delivered capacity
continued to outstrip the theoretical graphite value until the rate exceeded 10 C. By comparison with
previous literature on Ge anodes,[29, 30] it was suspected that performance was limited by the lithiation
kinetics at higher charge rates. This was verified by maintaining the charging rate at C/2 while
discharging at very high current rates up to 100C (Figure 13 (c)). To examine whether this high discharge
rate capability could be sustained over many cycles, two NW anodes were charged at a 2C rate, with one
discharged at 20C and the other at 100C - Figure 13 (d). After 80 cycles, the former offered a stable
reversible capacity of 610mAh.g-1, with the latter delivering 435 mAh.g-1. This implies that even at 100C
discharge rates, Sn seeded Ge NW anodes can outperform the maximum achievable capacity at 1C rates
for graphite based electrodes (372 mAh.g-1). This data demonstrates the potential for the utilisation of Ge
NW anodes in LIBs for high power applications.
1.7/ Morphological evolution of Sn seeded Ge NWs anodes with cycling
In order to understand the exceptional cycling stability of the Sn seeded Ge NWs compared to other
lithium alloying materials, ex-situ HRSEM of anodes was conducted after various numbers of
charge/discharge cycles - Figure 14. The electrodes were subjected to a washing procedure to remove the
SEI layer prior to imaging.[20]
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Fig. 14: (a-d) SEM images of Sn seeded Ge NWs after 1, 10, 20 and 100 cycles, respectively. Adapted
from Kennedy et al. [20]
After 1 cycle (Figure 14 (a)) the general form of the NWs remains intact, although some fusion between
adjacent wires is apparent. Following 10 cycles (Figure 14 (b)) the outline shape of the wires is still evident
but significant opening of the surface has occurred through channel formation. By the 20th cycle (Figure
14 (c)), the original NW form has effectively disappeared, with the agglomeration and surface texturing
phenomena combining to lead to the emergence of a porous layer of Sn and Ge material. At the 100th
cycle (Figure 14 (d)) a network architecture of interconnected Ge strands is evident. SEM images recorded
after further cycling showed no further overall restructuring, suggesting that it is this network
morphology that gives rise to the extremely consistent cycling performance noted after ca. 100 cycles in
Figure 11 (a). The initial drop in capacity over the first ca. 80 cycles is therefore envisaged to arise during
the morphological transition from NWs to network - once the latter is formed, it is evidently structurally
robust and permits stable Li cycling for over 1000 cycles. TEM analysis of this porous network, reported
elsewhere,[20] reveals that the constituent Ge ligaments have a diameter of 5.6 ± 1.0 nm. The electron-
microscopy study summarised in Figure 3.11 is significant, in that it reveals for the first time the
morphological origin of stable cycling performance from lithium alloying NW anode structures. The
transformation of the original discrete wires into an extremely stable network of nanoscale Ge ligaments
forms the basis for the consistent high capacity performance of Sn seeded Ge NW anodes with cycling.
This restructuring process is depicted schematically in Figure 15.
Fig. 15: Schematic representation of the structural reorganisation of Sn-seeded Ge NWs into a stable
porous network of Ge ligaments. The process is driven by repeated lithiation/delithiation over < 100
cycles. Adapted from Kennedy et al. [20]
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Interestingly, as we have outlined elsewhere,[20] the stable porous network structure did not emerge
when Ge NW anodes were cycled without vinylene carbonate additive. Instead, the NWs morphed into a
rather dense agglomeration of amorphous Ge and Sn, with delamination from the current collector also
evident. It would therefore seem, that in forming a more stable SEI layer and reducing the fresh exposure
of the Ge surface on each cycle, the VC indirectly assists in the development of the desirable porous,
interwoven Ge ligament structure of Figure 14.
1.8/ Half-cell electrochemical performance of Sn seeded Si nanowire anodes
Despite its higher theoretical capacity (3579 vs. 1348 mAh.g-1), achieving stable anode performance is
more challenging with Si than Ge. The diffusivity of Li is lower in Si (× 400), it has inferior electrical
conductivity (× 105),[30] while its lithiated alloy is also more brittle than that of Ge. Nonetheless it was
envisaged that previously reported performance for solid Si NW electrodes could be bettered through a
combination of the mechanical properties of our Sn seeded Si NWs and a judicious choice of electrolyte.
Fig. 16: (a) Discharge capacities of Sn seeded Si NW anodes over 100 cycles in 1 M LiPF6 electrolyte
solutions with different additives. Charging and discharging were conducted at a C/5 rate. (b)
Summary of the charge and discharge capacities and Coulombic efficiencies from (a) for the 1st and
100th cycles.
With this in mind, Si NW anodes were cycled in 1 M LiPF6 solutions of the following solvents: (a) ethylene
carbonate (EC) /diethyl carbonate (DEC) (1:1 v/v), (b) EC/DEC (1:1 v/v) + 3 wt% vinylene carbonate (VC)
+ 1wt% Lithium bis(oxalato)borate (LiBOB), (c) EC/DEC (1:1 v/v) + 2 wt% VC + 2 wt% vinylethylene
carbonate (VEC) + 2 wt% fluoroethylene carbonate (FEC), (d) EC/DEC (1:1 v/v) + 3 wt% VC, and (e)
FEC/DEC (1:1 v/v). The masses of both the Sn and Ge components were considered, leading to a
maximum theoretical capacity for the anodes of 2784 mAh.g-1, on which the C rate currents are based. The
results of these long-term cycling experiments are summarised in Figure 16. It is immediately apparent
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that the EC/DEC (1:1 v/v) + 3 wt% VC solvent significantly outperforms the other electrolyte solutions
over the first 100 lithiation/delithiation cycles. This parallels the situation for the Sn seeded Ge NWs,
where 3% VC additive was also optimum. As in that case, it is likely that the VC improves cycling capacity
retention and C.E. by helping to stabilise the SEI layer. The 1st and 100th cycle charge and discharge
capacities and C.E. values are plotted in the bar charts of Figure 16 (b) and suggest a capacity retention of
90.5% for the EC/DEC (1:1 v/v) + 3 wt% VC electrolyte after 100 cycles. The performance of our Sn seeded
Si NWs is very similar to that of the best comparable literature report (ca. 1600 mAh.g-1 at 100 cycles),
which however, employs a convoluted conventional CVD preparation of Au seeded Si NWs, requiring the
use of a sacrificial anodic aluminium oxide (AAO) hard template.[31]
Fig. 17: Rate capability data for Sn seeded Si NW anodes cycled in 1M LiPF6 solutions with various
solvent formulations.
Rate capability data for Sn seeded Si NW electrodes in the various electrolytes are presented in Figure 17.
Here again EC/DEC (1:1 v/v) + 3 wt% VC proves to be the best performing electrolyte, retaining a
discharge capacity of 1387 mAh.g-1 (cycle 25) at a 1C rate, which is 71% of the value at C/20 (1941 mAh.g-1
at cycle 5). A differential capacity plot analysis that we have published elsewhere demonstrates that, as for
the Ge NWs discussed in section 1.6, the Sn seeds are active for reversible lithium cycling in the case of
these Si NW anodes.[22]
1.9/ Half-cell electrochemical performance of Cu3Ge seeded Ge nanowire
anodes
The long-term cycling performance of Cu3Ge seeded Ge NWs is summarised in Figure 18. Capacity
retention in this case exceeds even that of our Sn seeded Ge NWs, which as discussed in section 1.6,
exhibited much superior performance to any previous report on a binder free, solid Ge NW anode. The
Cu3Ge seeded Ge NW anode delivers a discharge capacity of 958 mAh.g-1 after 1100 cycles at 1C, compared
to the value of 888 mAh.g-1 achieved by the Sn seeded Ge NW electrode at the same stage at C/2. The
experiment was continued to 1900 cycles, by which stage the Cu3Ge seeded Ge NWs were still offering a
reversible capacity of 866 mAh.g-1. The general form of the capacity vs. cycle plot in Figure 18 is similar to
that for the Sn seeded Ge NW material in Figure 11 (a), in that there is a gradual decline in capacity from
1333 to 1128 mAh.g-1 over the first 50 cycles followed by more stable performance over the next 1850
cycles, during which the anode loses less than 1/4 of its capacity and exhibits an average C.E. of 99.7%.
This behaviour suggests that, as for the Sn seeded wires, restructuring of the active mass to a more stable
morphology may occur during the first tens of lithiation/delithiation cycles. The SEM image of Figure 19
16
supports this viewpoint and shows, that after 50 cycles, the NWs have evolved into a porous, sponge-like
network of Ge nano-fibres, reminiscent of that observed for the Sn seeded Ge NWs in Figure 14 (d). Given
the comparative stability in reversible capacity beyond ca. 50 cycles, it again seems that this is a stable
morphology that permits extensive cycling with only very gradual degradation, or delamination from the
current collector.
Fig. 18: Charge and discharge capacity and Coulombic efficiencies for a Cu3Ge seeded Ge NW anode
cycled at 1C rate. The electrolyte was 1 M LiPF6 in EC/DMC (1:1 v/v) + 3% VC. Reproduced from
Mullane et al.[24]
Fig. 19: SEM image of the Ge material after 50 lithiation/delithiation cycles. The transition from
nanowires to a porous nano-fibre network is obvious. Reproduced from Mullane et al.[24]
The rate capability of the Cu3Ge seeded Ge NW anode was tested by charging and discharging for 5 cycles
at rates of C/10, C/5, C/2, 1C, 2C and then back to C/10 (Figure 20 (a)). For the 5th cycle at each rate the
observed discharge capacities were 1318, 1277, 1210, 1177, 1081 and 1285 mAh.g-1 respectively. Notably
there was no significant abrupt decrease in capacity when increasing the rate from C/10 to 1C. In view of
the previous experience (section 1.6) that charging (lithiation) kinetics limits the high rate capability of Ge
NWs, it was again elected to probe the anode performance at very high discharge rates by charging at a
constant C/2 rate - Figure 20 (b).
17
Fig. 20: Rate capability data for Cu3Ge seeded Ge NW anodes cycled in 1M LiPF6 in EC/DMC + 3wt%
VC electrolyte. a) Capacities and C.E. values, where charge and discharge were conducted at the
same rate for 5 successive cycles . b) High rate discharge capacities measured at each of 6 different
discharge rates. The charge rate was C/2 for all cycles. Reproduced from Mullane et al.[24]
The electrode delivered capacities of 1148, 908, 876, 857, 828, 745 and 439 mAh.g-1 in the 5th cycle at
discharge rates of C/2, 10C, 20C, 40C, 60C, 100C, and 250C, respectively. The attainment by the NWs of
twice the 1C theoretical capacity of graphite at the ultra fast discharge rate of 100C again points to the
potential application of this type of anode for high power applications. Cu3Ge seeded Ge NWs outperform
those seeded with Sn in high rate capability tests - recall (Figure 13 (c)) that the 1000 C discharge capacity
for Sn seeded Ge was 354 mAh.g-1, less than ½ the value exhibited at this rate in Figure 3.20 (b). The
superior rate capability of the Cu3Ge seeded wires is probably largely attributable to their smaller
diameter (approximately ½ that of the Sn seeded wires) which suggests shorter Li diffusion distances.
1.10/ Half-cell electrochemical performance of Ge/Si heterostructure NW
anodes
The principal motivation for pursuing a study of a nano-structured anode material containing both Ge
and Si, was to ascertain if performance synergies could be achieved relative to the single element
materials. Specially it was hoped that the mechanical stability of germanium could be coupled with the
higher charge storage capacity of Si in the composite. Additionally it was possible to vary the mass ratio of
the elements by varying the duration of the Si 'branch' growth step.
Galvanostatic cycling data for heterostructure NWs with three different Ge/Si mass proportions are
summarised in Figure 21 (a). Discharge capacities of 1612, 1459 and 1255 mAh.g-1 were noted after 100
cycles for the 2:1, 3:1 and 4:1 Ge:Si ratios, respectively. These values for the two most Si rich
18
heterostructures exceed the maximum theoretical capacity of a Ge-only anode (1348 mAh.g-1), while
Figure 21 (b) shows that the characteristic fading associated with Si-only anodes is noticeably reduced. As
might be expected, the greatest degree of cycling stability is exhibited by the most Ge rich composite,
while that with the highest Si content offers the highest capacity.
Fig. 21: a) Discharge capacity data, and b) relative capacity retention for Ge/Si branched
heterostructure anodes. The electrolyte was 1M LiPF6 in EC/DEC + 3wt% VC and the
charge/discharge rate was C/10. Reproduced from Kennedy et al.[25]
The rate capability data for these anodes presented in Figure 22 is particularly interesting. At the slower
rates of C/10 and C/5 the intrinsic high charge storage capability of Si comes to the fore, and it is the pure
(Sn seeded) Si NW anode that delivers the highest capacities. However as the cycling rate is increased the
heterostructured anodes outperform the pure Si NWs due to the greater electrical conductivity of Ge and
its higher rate of Li+ diffusivity. Indeed at the highest tested rate of 10C, the most Si poor anode (4:1
Ge:Si) offers the highest capacity of 802 mAh.g-1, compared to a mere 130 mAh.g-1 delivered by the pure Si
NW electrode. The Ge/Si heterostructured NW architecture therefore offers the intriguing possibility of
being able to tailor a lithium alloying anode to either energy (Si rich) or power (Ge rich) applications
simply by altering the relative amounts of Ge and Si during electrode fabrication.
Fig. 22: Rate capability data for several Ge/Si heterostructured anodes and a pure Sn seeded Si
anode. The electrodes were charged and discharged for 5 cycles at the indicated rate in the potential
range of 0.01 - 1.0 V vs. Li/Li+. The electrolyte was 1M LiPF6 in EC/DEC + 3wt% VC. Adapted from
Kennedy et al.[25]
19
Conclusions and Perspectives for the future
During the Greenlion project, the team at The University of Limerick developed methodology to achieve
the direct growth of lithium alloying NWs (Si, Ge or Ge/Si branched heterostructures) onto stainless steel
current collectors without the need for polymeric binders or conductive additives. These techniques
facilitate NW production through either vapour-solid-solid (VSS) or vapour-liquid-solid (VLS)
mechanisms. In this chapter it was illustrated how Sn seeded Ge or Si NWs could be grown by Solvent
Vapour Growth in a glassware based batch reactor approach. Additionally the production of Cu3Ge seeded
Ge NWs by the Rapid Pyrolysis method was outlined. In fact both Sn or Cu3Ge seeded materials can be
produced by either technique, with the rapid pyrolysis method having the advantage of been suitable for
up-scaling to a roll-to-roll type, semi-continuous manufacturing process. A combination of the two
procedures was also used to produce a novel nanoscale Ge 'stem' - Si 'branch' heterostructed anode
material.
The electrochemical performances of these various NW anodes were evaluated in half cell tests. For both
types of Ge material, high capacities with low fading rates were maintained for many more cycles than had
previously been reported for binder free, Ge NW electrodes. The potential for the application of Ge NW
anodes in LIBs with high power demands was highlighted by the fact that these materials can offer greater
discharge performance than the 1C capacity of graphite (372 mAh.g-1) even at 100C. Despite possessing a
higher theoretical capacity than Ge, it is more challenging to achieve stable long term cycle of Si, since the
latter is not as mechanically robust. Nevertheless, we have shown that Sn seeded Si NW electrodes can
maintain a discharge capacity of 1500 mAh.g-1 after 100 cycles at a C/5 rate. Testing of such anodes is
ongoing and we plan to publish a study on their performance over a much greater number of cycles in due
course. Investigations on Ge/Si heterostructured NW electrodes have indicated that they combine the
high capacity of Si with the long term cycling stability and excellent rate capability of Ge. These anodes
can therefore be directed towards high energy or high power applications by tuning the mass ratio of Ge to
Si. A summary of the various anodes described in this chapter is provided in Table 1.
Table 1: Summary of discharge capacities obtained for NW anodes in both long-term cycling and rate
capability experiments.
Material Electrolyte
1 M LiPF6 in: Initial
capacity/ mAh.g-1, rate
Extended capacity/ mAh.g-1, C.E. / %, rate
Rate capability capacity/ mAh.g-1
Sn seeded Ge NWs EC/DMC (1:1 v/v)
+ 3wt% VC 1103, C/2
888 (1100 cycles), 99.5%, C/2
930 at 60C (C/2 charge)
Cu3Ge seeded Ge NWs
EC/DMC (1:1 v/v) + 3wt% VC
1333, 1C 866 (1900 cycles), 99.7%, 1C 745 at 100C (C/2
charge)
Sn seeded Si NWs EC/DEC (1:1 v/v)
+ 3wt% VC 1658, C/5
1500 (100 cycles), 98.9%, C/5
1387 at 1C (1C charge)
2:1 Ge/Si NWs EC/DEC (1:1 v/v)
+ 3wt% VC 1484, C/10
1612 (100 cycles), 97.5%, C/10
582 at 10C (10C charge)
3:1 Ge/Si NWs EC/DEC (1:1 v/v)
+ 3wt% VC 1255, C/10
1459 (100 cycles), 99.3%, C/10
704 at 10C (10C charge)
4:1 Ge/Si NWs EC/DEC (1:1 v/v)
+ 3wt% VC 1249, C/10
1255 (100 cycles), 98.9%, C/10
802 at 10C (10C charge)
The impressive capacity maintenance reported here for the various nanostructured anodes can be
attributed to two main factors. Firstly, the NWs restructure over the first 50 - 100 lithiation/delithiation
cycles to form a porous network of interconnected nano-ligaments. This material is resistant to further
20
morphological change, while remaining mechanically and electrically connected to the current collector.
Secondly, the issue of continuous SEI layer formation has been addressed through a systematic study of
electrolyte additives. It was found that vinylene carbonate additive was particularly helpful in promoting
consistent long term performance. This agent is believed to stabilise the SEI layer and thereby reduce the
problem of ongoing electrolyte decomposition known to accompany the large volume changes inherent to
lithium alloying anodes.
The work described herein is significant on both applied and fundamental levels. From a technological
viewpoint it has been demonstrated that, through nanostructuring, the high charge capacity potential of
lithium alloying elements can be harnessed over extended numbers of cycles. At a more fundamental
level, valuable insights have been acquired into the morphological evolution of these nanomaterials with
lithium cycling. This research brings us closer to the practical application of nanostructured Si and/or Ge,
as full or partial replacements for graphite in next generation LIB anodes. We have recently produced an
overview article[32] which considers the remaining challenges to the commercialisation of lithium
alloying NW anodes (e.g. mass loading, first cycle Coulombic efficiency and cost) and the ongoing
progress in alleviating these difficulties. Research and development work on Si and Ge NW based anodes
continues at The University of Limerick, both in terms of fundamentals and the up-scaling of the
production of these materials towards commercial quantities. Finally we note that the NW materials
described in this chapter can also be readily harvested from their growth substrate and utilised in a
conventional slurry coating process for anode fabrication.
21
Contacts and References
Dr Kevin M. Ryan,
Materials and Surface Science Institute and Department of Chemical and
Environmental Sciences,
University of Limerick,
Limerick,
Ireland.
Authors
Dr Michael Brandon: [email protected]
Dr Tadhg Kennedy: [email protected]
Dr Emma Mullane
Dr Kevin M. Ryan: [email protected]
[1] N. Dimov, in: M. Yoshio, R.J. Brodd, A. Kozawa, (Eds.)Lithium Ion Batteries, Science and Technologies; Springer, New York, 2009, p 241-243. [2] B. Scrosati, J. Garche, Journal of Power Sources 195 (2010) 2419-2430. [3] L. Ji, Z. Lin, M. Alcoutlabi, X. Zhang, Energy & Environmental Science 4 (2011) 2682-2699. [4] N. Nitta, G. Yushin, Particle & Particle Systems Characterization 31 (2014) 317-336. [5] C.-M. Park, Jae-Hun Kim, H. Kim, H.-J. Sohn, Chemical Society Reviews 39 (2010) 3115-3141. [6] H. Wu, Y. Cui, Nano Today 7 (2012) 414-429. [7] U. Kasavajjula, C. Wang, A.J. Appleby, Journal of Power Sources 163 (2007) 1003-1039. [8] C.K. Chan, H.L. Peng, G. Liu, K. McIlwrath, X.F. Zhang, R.A. Huggins, Y. Cui, Nature Nanotechnology 3 (2008) 31-35. [9] C.K. Chan, X.F. Zhang, Y. Cui, Nano Letters 8 (2008) 307-309. [10] M. Osiak, H. Geaney, E. Armstrong, C. O'Dwyer, Journal of Materials Chemistry A 2 (2014) 9433–9460. [11] T. Song, L. Hu, U. Paik, Journal of Physical Chemistry Letters 5 (2014) 720−731. [12] M.J. Armstrong, C. O'Dwyer, W.J. Macklin, J.D. Holmes, Nano Research 7 (2014) 1-62. [13] M.H. Seo, M. Park, K.T. Lee, K. Kim, J. Kim, J. Cho, Energy & Environmental Science 4 (2011) 425-428. [14] R. Huang, X. Fan, W. Shen, J. Zhu, Applied Physics Letters 95 (2009) 133119. [15] M.-H. Park, Y. Cho, K. Kim, J. Kim, M. Liu, J. Cho, Angewandte Chemie 123 (2011) 9821–9824. [16] M.-H. Park, M.G. Kim, J. Joo, K. Kim, J. Kim, S. Ahn, Y. Cui, J. Cho, Nano Letters 9 (2009) 3844–3847. [17] C.K. Chan, R. Ruffo, S.S. Hong, Y. Cui, Journal of Power Sources 189 (2009) 1132–1140. [18] A.M. Chockla, M.G. Panthani, V.C. Holmberg, C.M. Hessel, D.K. Reid, T.D. Bogart, J.T. Harris, C.B. Mullins, B.A. Korgel, Journal of Physical Chemistry C 116 (2012) 11917–11923. [19] N. Liu, H. Wu, M.T. McDowell, Y. Yao, C. Wang, Y. Cui, Nano Letters 12 (2012) 3315-3321. [20] T. Kennedy, E. Mullane, H. Geaney, M. Osiak, C. O’Dwyer, K.M. Ryan, Nano Letters 14 (2014) 716–723. [21] H. Geaney, E. Mullane, K.M. Ryan, Journal of Materials Chemistry C 1 (2013) 4996-5007.
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[22] E. Mullane, T. Kennedy, H. Geaney, C. Dickinson, K.M. Ryan, Chemistry of Materials 25 (2013) 1816–1822. [23] M. Bezuidenhout, T. Kennedy, S. Belochapkine, Y. Guo, E. Mullane, P.A. Kiely, K.M. Ryan, Journal of Materials Chemistry C 3 (2015) 7455-7462. [24] E. Mullane, T. Kennedy, H. Geaney, K.M. Ryan, ACS Applied Materials & Interfaces 6 (2014) 18800-18807. [25] T. Kennedy, M. Bezuidenhout, K. Palaniappan, K. Stokes, M. Brandon, K.M. Ryan, ACS Nano 9 (2015) 7456–7465. [26] Y.-D. Ko, J.-G. Kang, G.-H. Lee, J.-G. Park, K.-S. Park, Y.-H. Jin, D.-W. Kim, Nanoscale 3 (2011) 3371-3375. [27] H. Wu, G. Chan, J.W. Choi, I. Ryu, Y. Yao, M.T. McDowell, S.W. Lee, A. Jackson, Y. Yang, L. Hu, Y. Cui, Nature Nanotechnology 7 (2012) 310-315. [28] M. Ulldemolins, F. Le Cras, B. Pecquenard, V.P. Phan, L. Martin, H. Martinez, Journal of Power Sources 206 (2012) 245-252. [29] A.M. Chockla, K.C. Klavetter, C.B. Mullins, B.A. Korgel, ACS Applied Materials & Interfaces 4 (2012) 4658−4664. [30] J. Graetz, C. Ahn, R. Yazami, B. Fultz, Journal of The Electrochemical Society 151 (2004) A698-A702. [31] J.-H. Cho, S.T. Picraux, Nano Letters 13 (2013) 5740-5747. [32] T. Kennedy, M. Brandon, K.M. Ryan, Advanced Materials (2016) Accepted.