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: michael.brandon@ul.ie

Dr Tadhg Kennedy: tadhg.kennedy@ul.ie

Dr Emma Mullane

Dr Kevin M. Ryan: kevin.m.ryan@ul.ie

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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.