X-ray radiog raphic and to mograph ic investiga tions o f cycled lithium
ion batteries

vorgelegt von
Master of Science
Fu Sun
ge b. in Shandong

von der Fa kultät III – Prozessw issenscha ften der T echnische n U niversitä t Berlin
zur Erlang ung des a kadem ischen Gra des

Doktor der Na turwiss enschaften

-Dr. rer. nat.-

Genehmigte Disserta tion

Promotionsausschuss:

Vorsitzender: Prof. Dr. Walter Reimers
Gutachter: Prof . Dr. J ohn Banhart
Gutachter: Dr. Yan L u

Tag der wissenschaftlichen Aussprache: 3. Februar y 2017

Berlin 2017

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Abstract

Lithium ion batteries (LIBs) ha ve become the predomi nant power supplies in portable electronic
devices such as c ell phones, laptops and tablets. They are regarded as promisi ng candidates to
power future electric vehicles and to store the intermitt ent and fluctuating energy harvested f rom
solar and wind sources . Unfortunately, the current LI B technolog y fails to meet the gro wing
energy d emand required by th e automotive indust ry and ope rators of po wer stations. In addition,
safety concerns pr event a simple scale-up from t he currentl y small sc ale LIBs to future lar ge-
scale LIBs. Future LI B t echnolog y will undoubtedl y rel y on a sound understanding of und erl y in g
work principles of current LI Bs. In this dissertation, X-ray im aging both utilizes X-ray tub e and
sync hrotron sources is e mployed to investigate the underlying r eactions of operating LIBs.

Firstly, li thium dendrites which are in -homogeneously d eposited/stripped on/from the surface o f
lithium ( L i) electrodes d uring c y cling are ch aracterized. The morphological evolution of lithium
dendrites is investigated . In addition, it is observed that both lithium dissolution during lithium
stripping and lithi um deposition during li thium plating c ontribute to the formation of the porous
lithium interfac e (PS I) . Secondly, the commercialized trila y er Celgard ® 2325 separator is
investigated during b attery c y cling . It is observed that the t rila ye r separato r can del aminate under
mechanical forces arising from grow ing lithi um microstructures. I n addit ion, partial melti ng of
the separator resulting from an internal short circuit (ISC) is demo nstrated. Thirdl y, gas
development inside an operating LIB is inv estigated. It is found that ga s is pr eferentiall y
ge nerated in regions close to the separator. Moreover, it is observed that g as generation is a
continuous process extending over man y cyc les . Fourthly, degradation of LI Bs based on Si
particles has been studi ed. The volume expansion during lithiation and volume contraction during
delithiation of Si particles have been observed. It is foun d that the distribution of frac tured Si
particles is het erogeneous. Furthermore, electrochemical deactivation of ori ginally
electroc hemicall y active Si particles is observed. Moreover, elec trochemically ina ctive Si
particles are believed to decrease the energy d ensi t y of an LIB. Fifthly, different (de)lithiation
behaviors of Sn particles are pr esented. In addi tion to the conventional “ core-shell ” reaction
mode, an unusual fractur e of a Sn particl e is obser ved during c y cling. It is observed that some S n
particles need some “ incubation ” time to become electroc hemicall y active.

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Kurzfassun g

Lithium-Ionen-Batterie n (aus dem englischen: Lithium ion batteries LI Bs) h aben sich als
vorherrschende Energiequelle in portablen elektronischen G eräten, wie zum Beispiel
Smartphones, Laptops und Tablets, e tabliert. S ie werde n als vielversprechende Möglichke it
betrachtet, zukünftige el ektrische Autos anzutrei ben als auch überschüssige Energi e aus Solar -
und Windkraftanlage n z wischenzuspeichern. Di e heutige LIB- Technologie schaff t es derzeitig
nicht den technischen Anforderungen der Automobilindustrie und der En ergiewirtschaft gerec ht
zu werden. Hinzu kommen Sicherheitsbedenken großvolumige LIBs zu b etreiben. In der Zukunft
wird diese Technologie zweifelsohne auf dem fundierten Verständnis der zugrunde liegenden
Prozesse derzeitiger LIBs basieren. I n dieser Dissertation werde n bild gebende Verfahren mittels
einer kommerziellen R öntgen-CT Anlage als auch mi ttels S ynchrotron Röntgenstrahlung
angewe ndet, um die gr undlegenden Reaktionen in LI Bs wä hrend des Be -und Entladens zu
untersuchen.

Im ersten Teil wurden Lithium-Dendrite charakterisiert, die währ end des Zyklierens inhomogen
auf der Lithiumelektrodenoberfläche ab geschieden und abgetragen w erden. Die morphologische
Entwicklung d er Lithium-Dendrite wurde dab ei untersucht. Es wurde dabei beobachtet, dass
sowohl das Auflösen während des Abtragens als auch die Besc hichtung wä hrend des
Abscheidens zur Bildung einer porösen Lithiumoberfläche b eitragen. Im zweiten Teil wu rde ein
kommerziell erhältlicher dreilagiger Celgard® 2325 Separator während de s Zyklierens untersucht.
Dabei wurde beobachtet, dass der dreilagige Separator durch die mec hanischen Kräfte der
anwac hsenden Lithiumstrukturen delaminieren kann. Zusätzli ch wird ein teilweises Schmelzen
des Separators durch einen internen Kurzschluss beobachtet. Der dritte Teil widmet sich der
Gasentwicklung in LIBs während des Z y kli eren. Es konnte gezeigt werde n, dass Gas
überwiege nd in Separator-nahen Regionen erzeugt wird. Die Gasentstehung wird als
kontinuierlicher P rozess charakterisiert, der sich über viele Zyklen erstreckt. Im vierten Teil
wurde die De gradation von Silizium-basierten LI Bs studi ert. Die Expansion während der
Lithiierung und die Kontraktion während der Delithiierung wurde n im Detail quantitativ
analy siert. Die R eaktion der Siliz iumpartikel ist dabei räumlich heterogen, das heißt, die Partikel
reagieren individuell auf verschiedene Weise und zu verschiedenen Ze itpunkten. Darüber hinaus
wurde eine elektrochemi sche Deaktivierung von ursprünglich aktiven P artikeln beobachtet, was
darauf hindeutet, dass elektrochemisch inaktive P artikel die Ener giedi chte der LIBs entscheidend
verringern. Schließlich, im fünften Teil, wird das (De -)Lithiierungsverhalten von Sn -Partikeln
untersucht. Z usätzlich zur konventionellen „core - shell“ Reaktion, wurdd ein ungewöhnliches
Verhalte n b ei der D egradation der Sn -Partikel während des Z y klierens beobachtet: Unter
andere m hat sich ge zeigt, dass eini ge Sn - Partikel eine gewisse „ Inkubationszeit“ benötigen um
elektroche misch aktiv zu werden.

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Contents
1 I ntroduction .............................................................................................................................. - 1 -
1.1 I ntroduction to lithium ion batteries .................................................................................. - 1 -
1.1.1 Cathode material ........................................................................................................ - 4 -
1.1.2 Electroly te .................................................................................................................. - 5 -
1.1.3 Anode materia l ........................................................................................................... - 6 -
1.1.4 Separator ................................................................................................................... - 11 -
1.2 X-ray imag ing setup ........................................................................................................ - 14 -
1.2.1 laboratory X-ray imag in g setup ................................................................................ - 14 -
1.2.2 Synchr otron X-ra y imaging setup ............................................................................ - 15 -
1.3 The customized cells ....................................................................................................... - 16 -
2 Published parts of work .......................................................................................................... - 17 -
2.1 Morphological evolution of lithium microstructures ...................................................... - 18 -
2.1.1 Supporting information ............................................................................................ - 31 -
2.2 Break-down of the separa tor ........................................................................................... - 39 -
2.3 Gas developme nt ............................................................................................................. - 57 -
2.3.1 Supporting Inf ormati on ............................................................................................ - 75 -
2.4 Degra dation of lithiu m ion batteries based on ~100 µm -sized Si particles .................... - 81 -
2.4.1 Supporting Inf ormati on ............................................................................................ - 99 -
2.5 Fracture behavior of ~20 µm S i particles ...................................................................... - 108 -
2.5.1 Supporting Inf ormati on .......................................................................................... - 124 -
2.6 Different (de)lithiation behaviors of Sn particles .......................................................... - 126 -
2.6.1 Supplementary Inf ormation .................................................................................... - 138 -
3 Summa ry .............................................................................................................................. - 144 -
4 Outlook ................................................................................................................................. - 147 -
5 Acknowledge ments .............................................................................................................. - 148 -
6 Refere nces ............................................................................................................................ - 149 -

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1 Introduc tion

1.1 I nt rodu ction to lit hium ion batt eries

Since the successful commercialization by SO NY in 1991 and a joint venture of Asahi Kasei and
Toshiba in 1992, the lithi um ion battery ( LI B) h as penetrated ubiquitousl y into our dail y lif e [1].
The predominant reason that LIB can revolutioni ze portable energy stor age te chnology is due to
its higher en ergy density c ompared with oth er battery s y stems such as l ead -acid and Ni-MH
(nickel-meta l h ybride), as shown in Fig. 1 a [2] . In contrast to other battery s y stems, LIBs d o not
suffer from the memor y effect problem. In addition, LI Bs have voltages nearl y three times these
of t ypical Ni-based batteries. The hi gh sin gle-cell voltag e obtained from a LI B reduces the
number of cells required in a battery s y stem. Finall y , the s elf-discharge rate in LIBs is very low
(< 5% per month) compared with Ni-based batteries (~ 20-30% per month).
Moreover, in 2005, SONY launched the Nex elion h y brid LIB, which further increased the batter y
energy densit y b y 30% compared to conventional LIBs [3] . With such high energ y densit y ,
coupled with their long cyc le life and rate capabil it y , LIB s are now consi dered promising
candidates to power the upcoming electric vehicles (EVs) or plug-in hy brid electric vehicles
(PHEVs) [ 4]. Another important potential market for LIBs is to stor e int ermittent and fluctuating
green ener gy suppl ies from renewa ble sour ces su ch as solar and wind sources [ 5]. From Fig. 1b ,
one can observe that the market for various kinds of LIBs has ex panded si gnificantl y from 2000
to 2015 [6] and this trend is going to continue.

Fig. 1 a) Com par ison of di fferent battery types in terms of energ y densi ty. Reprinted with pe rmission fr om
Ref. 2. b) Ev olution of t he lithium ion battery sale in various m arkets. Reprinted with permission from
Ref.6.

A more detailed hist orical evolution and present review of LIBs can be f ound elsewhere [ 2, 6] .
Typically , a b asic LIB consists of a cathode ( pos itive electrode), an anode (negative electrode),
electrolyte containing lithium ions and a separator isol ating the anod e from the cathode. The
components of the first commercialized LI B are schematicall y illust rated in Fig. 2 [ 7]. During
1 Introduction

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discharge and charge, lithi um ions shuttle between the anode and the cathode while electrons are
transported throu gh the ex ternal circ uit. The driving force fo r discharge of LIBs is the different
electroc hemical potential μ A and μ C of the anode and the cathode. The open-circuit voltage of a
cell can be expressed by [7]:
V = 𝜇 𝐴 − 𝜇 𝐶
𝑒 (1.1)

More details of the op en circuit volta ge with re lati on to the Fermi ene rg y leve l and Gibbs free
energy of electrde materials can be found in previous reports [ 2, 8]. The effect of lithium ion
tranportation behavior within electrode materials on open circ uit voltage can also be found in
previous investiga tions [9, 10] .
During th e char ging pro cess, the two electrodes are connected ex ternally to an electrical suppl y .
The electrons are forced to be relea sed from the cathode and transported externall y t o the anode
side. Simultaneously, lit hium ions move internally from the cathode to the anode through the
electrolyte. Hence, the external energ y is electrochemica ll y stored in the batter y in form of
chemical ener gy in both the anode and cathode [ 5]. During the dischar ge process, electrons move
from the anode back to the cathode throu gh external load to do work and lithium ions move
through the electrol y te. The stored chemic al energ y inside the batter y is released vi a
electroc hemical reactions at two electrodes.

Fig. 2 Schem at ic illustrat ion of the first l ithium ion battery . Reprin ted with perm ission from Ref. 7.

The amount of electrical energ y that can be con verted b y the electrochemical rea ction of LI B
electrode s during discharge c an be expressed either pe r unit of weight (Wh/kg) or pe r unit of
volume (Wh/l). This electrical energ y that a batter y can deliever is a function of the open-cell
potential (V) and specifi c capacit y (Ah/k g or mAh/g) [ 2]. The specific energ y (Wh/kg) c an be
obtained b y multipl ying the specific capacit y with the opera ting batter y voltage. The specific
capacity measures the a mount of charge that can be r eversibl y stored/r eleased pe r unit mass. It is
1 Introduction

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closely related to the nu mber of electrons transferred during electrochemical reactions and th e
atomic weig ht of the electrode material. The theoretical capacit y of electrode materia ls can b e
estimated based on electrochemical reactions. For example, the reversible electroche mical
reac tion of a gra phite anode LIB is [6] :

Li + + e - + 6C ↔ L iC 6 (1.2)

The theore tical specific capacit y ( mAh/g) of a graphit e anode can be estimated as:

C specific = xF/nM = 1 × (96485 C/mol)/6 × (12g/mol) = 372 mAh/g (1.3)

where x is the number of electrons transferred in formula (1.2), F = 96485 C /mol is Faraday ’ s
constant, n is the number of moles of the chosen graphite anode taking place in the reaction and
M is the mol ecular w eight of graphite. I n practice, to ev aluate the spe cific c apacity o f a battery,
not onl y integration of c athode and anode mate rials have to be tak en into consideration, but also
other essential componet s such as binders, conductive enhance rs, sepa rators, electrol ytes, current
collectors, cases, tabs as well as battery management s y st ems. As a result, the practica l energy
density is alway s less than that estimated one based onl y on battery chemistry.
Cyclability of a battery measures the reversi bilit y o f lithium ion ins ertion and extraction
processes in terms of th e number of charge and dischar ge c ycles before a battery fails or can no
longer sust ain devices [5]. S pecifically , t he c y cle life of LIBs is closel y affec ted b y the st ate of
charge (SoC) and depth of discharge (DoD) as w ell as operatin g temperature and current rate. I t
has been demonstrated that the cycle life can be enhanced b y a shallow DoD and a lower SoC
[11] .
Finally , abuse tol erance of a Li-ion battery is a critical re quirement for prac tical appli cations
expecially for hi gh-energy and high-densit y EVs and PHEVs. Before commercialization,
mechanical, thermal and electrica l abus e evalu ations have to be conducted on protot y pes of Li-
ion batteries to evaluate abuse tolerance. T ypically, mechanical abuse evaluation includes roll -
over, nail penetration, mechanical shock and drop and immersion in water tests. The thermal
abuse evaluation include s thermal stabil it y test, radiant heat test, overheat test and extreme cold
test. The electrical abuse evaluation in cludes short circuit, ov ercharge and over -dis charge tests.
Those abuse tolerance tests are extremel y import before final commercializaion. Actuall y , the fire
catastrophe of the B oeing 787 Dreamliner is reporte d to have result ed from short circuits of
employe d LIBs [12] .
1 Introduction

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Fig. 3 Vol tage versus capacity for negat ive and positive elec trode m aterials pr esently used o r under
condiderat ion for the next- generation of LI Bs. Reproduced with p ermission f rom Ref. [2 ]
The volt age range and capaci ty of differen t electrode materials are shown i n F ig. 3. In the f ollow ing, the
focus w ill be on a brief introduction to cathode, electrolyte, anod e and se parator, followed by an
introduction o f prim ar y studies of the present di ssertat ion.

1.1.1 Cathode material

Generally speakin g, cathode materials are t y pically ox ides of transition metals, which can
experience ox idation to higher v alences when li thium is removed duri ng char ge [13]. W hile
oxidation of the transition metal can cons erve charge neut ralit y in the compound, large
compositional changes can lead to phase chan ges. As a result, cathode materials have to be
crystal-structurally st able over wide ranges of composit ion. During discharge, lithium is inserted
into the cathode along with electrons fro m the anode to reduce the transit ion metal ions in the
cathode to a lowe r valen ce [14] . The rat e of this ox idation and reduction of cathode materials as
well as the rate of transporation of lithium ions in the electroly te control the max ium discharge
curre nt. It has to be note d that ex chang e of lithium ions between cathode and elect rol yte occurs at
electrode-electroly te interfaces, so cathode p erformance depends d irectly on electr ode
architec ture and morpho log y as well as on electrochemical properties of cathode m aterials [14] .
From the perspective of materials, there are a nu mber of candidates that h ave been investi gated as
cathode materials for LIBs. The ca thode mat erials can be categ orized based on their voltage
versus L i/Li + , namely [5]: TiS 2 and MoS 2 with 2- D la y ered structur e ( ~2 Volt); MnO 2 and V 2 O 5
(~3 Volt); L iCoO 2 , LiNiO 2 with 2-D layered structure and 3-D spinel LiMn 2 O 4 a nd olivine
LiFePO 4 (~4 Volt); olivine LiMnPO 4 , LiCoPO 4 and Li 2 MxMn 4-x O 8 (M=Fe, Co) spinel 3- D
structure (~5 Volt). T ypica lly, hi gh cathode voltage is d esirable for electrode materials because
the ener gy densit y is pr oportional to cell voltage. Currentl y , L iCoO 2 and L i FePO 4 are most
widely used in commercial LI Bs be cause of their excellent cyc le life and capacity r etention.
However, investigation of cathode ma terials is not the focus in the current dissertation.

1 Introduction

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1.1.2 Electrolyte

The electrol y te is crucial for LIBs. First of all , the electroly te h as to be chemicall y and
electroc hemicall y stable to withstand a redox environment at both the cathode and anode sides
besides possessing high ionic conductivity for lith ium ions. I n addition, the electroly te has to be
inert and stable in an acceptable temperature range. A detailed d escription of the potential
window of th e electrol yte with respect to the electrochemical potentials of elec trodes can b e
found in previous reports [15] .
In commercial Li-ion batteries, t ypically a liqu id electroly te is a solution of li thium salts in
organic solvents. However, it should be noted that there are othe r t y pes of electrol y t e such as
ionic electrolytes and inorganic solid electrol ytes [16] . Most of the ex isting or ganic li quid
electrolytes can potential l y catch fir es under condi tions of thermal runawa y or short circuit due to
the volatile and flammable nature of the solvents. On the one hand, polar aprotic organic solvents
such as carbonate solven ts with a high dielectic c onstant are s elected to solv e lithium salts at a
high concentration (normall y 1 M ) [ 5]. On the o ther hand, solvents with a low viscosit y and low
melting point are necessar y to meet the requirement for high ioni c mobility during c ycling. Until
now, various or ganic sol vents have bee n explored for batter y usage such as dimeth y l carbonate,
diethyl carbonate, eth y l met hyl carbonate, ethylene carbonate, etc [17] . I n addition, man y li thium
salts have been investi gated such as L iPF 6 , L iBF 4 , L iAsF 6 , L iClO 4 , etc [18]. Currently , the
wi dely used electrol yte f or LIBs is a lithium hex afluor ophosphate solut ion (LiPF 6 ): 1 M L iPF 6 in
EC/DMC (ethy lene carbonate and dim eth y l carbonate, v/v=50/50) or 1 M LiPF 6 in EC/EMC
(ethy lene carbonate and eth y l methy l carbonate, v/v=50/50), commercialized b y dif ferent
companies [5] . These electrolytes offer reasonable stabilit y over a wide ran ge of potential.
Neverthe less, during c ycling of LIBs contain g these electrolytes, a S EI (soli d electroly t e interface)
forms from the side reaction s of the electrol yte with the electrodes [19] . Meanwhile, a si gnificant
amount of gas is generated [20]. The generated gas can cause an inter nal pressure build-up and
even re sult in severe gas leakage [21].
Various characterization tool s have be en emplo ye d to investigate the decomposition of
electrolyte such as gas chromatog raphy, a flame ioni zation detector and a thermal conductive
detector [18, 22]. Until now, the dec omposition of solvent and lithium salts during battery
operation has been thoroughl y studied. In the following, the d ecomposition of EMC, DMC, EC
and L iP F 6 salt will be prese nted.
Reduction of ethy l meth yl ca rbonate (EMC) following one-electron reduct ion:

CH 3 CH 2 O( C =O )OCH 3 + e - → CH 3 CH 2 O(C • – O - )OCH 3 (1.4 )

Then CH 3 CH 2 O( C • – O - )OCH 3 rea cts with Li + to produce CH 3 CH 2 OL i acc ording to:

CH 3 CH 2 O( C • – O - )OCH 3 + e - +2L i + → L iO(C= O)CH 3 + CH 3 CH 2 OL i (1.5)

Alternatively, a two-elec tron reduction process can occur as for dimeth y l c arbonate (DMC):
1 Introduction

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CH 3 OCO 2 CH 3 + 2e - + 2 Li + → Li 2 CO 3 ↓ + C 2 H 6(g) ↑ (1.6)

Reduction of ethy lene carbonate (EC) involves:

(CH 2 O) 2 CO + 2e - + 2Li + → Li 2 CO 3 ↓ +C 2 H 4(g) ↑ (1.7)

Or alternatively :

(CH2O) 2 CO + 2e - +2Li + → (CH 2 OCO 2 L i) 2 ↓ + CH 2 =CH 2(g) ↑ (1.8)

The L iPF 6 salt dissociates into L iF and th e lewis acid PF 5 :

LiPF 6 ↔ L i F + PF 5 (1.9)

or a Li + and a nion PF 6 - as below:

LiPF 6 ↔ Li + + P F 6 - (1.10)

The dissociated lew is aci d PF 5 can react with H 2 O:

PF 5 +H 2 O → 2HF + POF 3 (1.11)

PF 5 +H 2 O → PF 4 OH + HF (1.12)

The Lewis acid PF 5 can also react with dialk y l carbonat e to form a variety of decomposition
products [23].
Apart from these extensive characterizations of decomposed prod ucts by compositional
character izations, neutron radiograph y imaging was also employed to in vestigate gas generation
and movement inside LIBs [ 20]. Unfortunately, radiographic imaging always y ields two-
dimensional (2D) information onl y , which prevents us from further comprehending gas evolution
kinetics in its inherent three-dimension (3D) sta te and from further qua ntitativel y anal y zin g its
complex evolution . In the curre nt dissertation, s ection 2.3 presents a detailed stud y of three-
dimensional gas evolut ion in an operating LIB using non-destructive s y nchrotron X- ra y
tomogra ph y.

1.1.3 Anode material

Anode materials have been ex tensively investi gated and there are man y d ifferent materials and
candidate materials. Different anode materials exhibi t different electrochemical performance such
1 Introduction

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as cyclability , current rate and energ y densit y . From the mate rial point of view, there are
basically three t y pes of materials based on their reaction t y pes with Li [4, 24, 25] :
1, L i intercalation/de-insertion materia ls su ch as carbon-based materials (porous c arbon, carb o
nanotubes and graphene) and TiO 2 ;
2, alloying/de-alloying sy s tems such as Ge, Sn, Si Al, Bi and SnO 2 ;
3, conversion materials such as transition met al oxide materials ( Mn x O y , NiO, Fe x O y , CuO, Cu 2 O
and MoO 2 ), metal sulphides, metal phosphides and metal nitrides (M x X y , h ere X= S, P, N).
Since the first commercialization of carbonaceous anodes, carbon is sti ll dominant in commercial
LI Bs due to its excellent cy cle perfor mance. Graphitic carbon with a layered struc ture can
effectively store and release l ithium ions from its lattice space, resultin g in excellent cy clabilit y .
However, carbon anod es have been approaching their theoretical capacit y of 372 mAh/g over the
past two decades of development. During the se arch for hi gh ener gy densit y anode materials ,
conversion materials and alloy m aterials have bee n discovered [ 26, 27] . I n the following section,
we will mainl y focus on the allo ying/de-allo ying anode mate rials as the y n ormally possess hi gher
energy density th an con version materials, specificall y , tin and silicon. But before that, lithi um
metal as an anode material for LI Bs is briefl y revi ewed and discussed.

1.1.3.1 Lithiu m anode

The motivation for usin g lithium metal as anode is because lit hium is the most electronegative (-
3.04 V ve rsus standard h ydrogen electrode) and the lightest (equivalent weight M = 6.94 g /mol,
and specific gravity ρ = 0.53 g/cm -3 ) m etal, thus facilitating the design o f storage s y stems with
high en ergy densit y (theoretical capacit y is 3860 mAh/g) [ 28]. The usage of lithi um metal was
first demonstrated in the 1970s in primar y (non -rechargeable) lithi um cells [29] . In 19 72, Exx on
embarked on a large p roject using TiS 2 as the positive electrode, lithium metal as the negative
electrode in a rechargeable lithium cell [30]. However, the y soon en countered the disadvantages
of a lithium metal/liquid electrol yte combination: the uncontrolled growth of dendritic lithium
microstructures during each discharge/ch arge c ycle as shown in Fig. 4a. The electrochemicall y
ge nerated lithium microstruc tures such as dend rites, fibers and moss during inhomo gene ous
lithium stripping/plating results in an internal short circuit leading to b atter y f ailure [31]. A
review of lithium dendrite formation can b e found elsew here [32]. To circumvent the s afet y
issues associated with the usage of lithium metal, an alternative approach of su bstituting metallic
lithium b y a ins ertion ma terial was proposed, which led to the discov ery o f the hi ghl y reversible,
low-voltage L i intercalation -deintercalation carbonaceous materi al as schematically shown in Fig.
4b [33] .

1 Introduction

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Fig. 4 Schem at ic illustrat ion of li thium bat teries. a) lithium dendr ite growt h durin g cy cles. b) propo sed
approach of us ing an inte rcalatio n-deinterca lation hos t m at erial as a negativ e electr ode. From ref.[2 ]

Recently , metallic lithium has again attracted res earch attention from all over the world because
of it s potential use in next -generation technologies such as Li-S and Li -O 2 batterie s. Various
character ization tools su ch as optical mi croscopy [34] , atomic force microscop y (AFM) [35],
scanning electron micros cope (SEM) [36] , transmission electron microscopy (TEM) [37] , nuclear
magnetic resonance (NMR ) [38] and mag netic r esonance im aging (MR I ) [39] have been
employe d to und erstand the growth of the lithium microstructure ( LmS) in lithium -ion batteries.
Meanwhile, different strate gies such as employ in g L i surf ace coating [40], additives [ 41] ,
polymer- [42] and/or inorganic-based [ 28] electroly tes and emplo y in g differe nt electrode/cell
config urations [43] have be en adopted to reduce and/or control t he growth o f li thium
microstructures durin g cycling. D es pite th ese int ensive efforts, a fundamental understanding of
the underl ying evolution mechanisms remains elusive. Herein, in s ection 2 .1, synchrotron in-line
phase contrast X- ra y tomograph y was employed to investigate the mor phological evolution of
electroc hemicall y d eposit ed/dissolved li thium microstructures non-destructively. For the first
time , we present a 3D characterization of electrochemically stripped Li electrodes with regard s to
electroc hemicall y plated lithium mi crostructures. W e also clarif y fundamentall y the origin of the
porous lithium interface g rowing towards Li electrodes.

1.1.3.2 Silicon anode

Since Dey et al . demonstrated in 1971 that L i met al can electrochemicall y alloy with other metals
such as Sn, P b, Al, Zn, Si, etc. at room tempera t ure, Li allo y in g/de -allo y ing materials have b een
intensively inv estigated during the past few deca des [44]. Among various allo y in g pa rtners for Li,
Si has been considered as one of the most promising candidates for nex t -generation LIBs due to
its high specific c apacity (3500 mAh/ g) [ 45], low cost, abundance and environm entally beni gn
properties. Mor eover, th e discharging potential i s about 0.1 V with respect to Li/Li + , which is
lower than for most of other alloy-type and conversion metal oxide anodes [46].
The mechanism of elec trochemical lithi ation of S i is critical to improve the performance o f a S i
anode. The Li -Si allo y ing/de-allo y ing process during discharge/charge has been extensivel y
1 Introduction

- 9 -

investigated b y several groups [46, 47]. It has been found that the reactions follow the
equilibrium L i-Si binar y ph ase dia gram at hi gh temperature, forming different intermetallic
compounds and showing distinct voltage plateaus [48] . The detailed reac tion investiga ted b y X-
ray diffrac tion (XRD) anal y sis is ex plained as follows [49]:
During discharge:

Si (cry stalline) + x L i + + x e - → yLi + +y e - + Li 15 Si 4 (crystalline) (1.13 )

During charge:

Li 15 Si 4 (cry stalli ne) → Si (amorphous) + yLi + + ye - + Li 15 Si 4 (r esidual) (1.14)

Fig. 5 Voltage- capacity curve of Si - based LI B. Revised from ref [50]

The t ypica l discharge cu rve of a Si -based LI B is shown in Fig. 5. The crystalline Si becomes an
amorphous Li-Si allo y during the first lithi ation a nd the highl y lithiated amorphous Li x Si phase is
suddenly found to cr ystallize int o L i 15 Si 4 phase. A t y pical long voltage plateau at 0.1 V is also
clearly shown (1 st discharge curve, red li ne). During delithiation, the formed cr y stalline L i 15 Si 4
turns into amorphous Si and onl y a small fraction of Li 15 Si 4 phase remains . After the first c ycle,
Li ions re act with the amorphous S i formed at the end of the first delithiation. During the
abovementioned L i stor age and r elease process, the Si electrode undergoes volume expansion
~400% [ 51]. Th is significant volume expansion contributes to the capacit y degradation of Si -
based LIBs. In order to alleviate the significant volume expansion leading to irreversible capacit y
decay, the underlying morphological ch ange of the Si electrode with regard to the
electroc hemical b ehavior should be inv estigated. In the current dissertation, section 2.4 pres ents
1 Introduction

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the detailed investigation of the failure m echanisms of ~ 100- µm -sized Si particle LI Bs b y usin g a
laboratory X-ray source. Section 2.5 presents a det ailed study of the fracture behavior of ~20 - µm -
sized Si particle LIBs by using synchr otron X-r a y source.

1.1.3.3 Tin anode

Compared with Si, the theoretical specific capacit y of Sn is calculated to be 990 mAh/g [52]
when li thiated to Li 4 .4 Sn . The volume expansion o f a Sn electrod e during lithium insertion and
de-insertion is reported to be around 300% [53] , which is smaller than for Si. Actuall y , the
Nexelion LI B launched by S ONY in 2005 wa s b ased on Sn due to its high specific capacit y and
relatively low volume expansion . Even though, cu rrent LIB technolog y fall s short of meetin g the
demands b y energ y s y st ems because of its poor c y cl e perfor mance. Significant capacity decay
has been observed in Sn -based LIBs during prolonged cycle tests [45, 54] .
The electrochemical r eduction of a Sn el ectrode le ads to the subsequent for mation of a numbe r of
intermetallic phases Li x Sn y at room temper ature. The d etailed Li-Sn allo y reaction durin g
discharge has been suggested as follows:

5Sn + 2Li → Li 2 Sn 5 (1.15)

Li 2 Sn 5 + 3 L i → 5LiSn (1.16)

5LiSn + 17Li → Li 22 Sn 5 (1.17)

Some studi es also r eveal that the formation of Li 22 Sn 5 from LiSn also invol ves other int ermediate
phases [55]. The corresponding discharge c urve of a Sn electrode is shown in Fig. 6.

1 Introduction

- 11 -

Fig. 6 Disch arge cu rve of S n based LI B. Revised from ref. [56]

Similar to the S i anode, t he detailed kno wledge of the morpholo gical changes of Sn electrodes
with regard to the electrochemical behavior can guide future b atter y desi gn fo r next -generation
Sn anode LIBs. In the cu rrent dissert ation, section 2.6 presents detailed (de)lithi ation behavior of
Sn particles during (dis)charge b y in situ sy nchrotron X-ra y radiograph y .

1.1.4 Separator

The separator pla y s an important role in LIBs . The primar y function of a separator is to p revent
physica l contact b etween anode and cathode to prevent an int ernal short circuit ( ISC). Meanwhile,
the separator allows the lithi um ions to pass through to finish (dis)charge process. G enerally,
separa tors with sub-micro pore siz es have proven adequate to block an y p enetration of electrode
particles while providing efficient transport of l ithium ions [57]. Although th e sep arator itself
does not take p art in the cell reaction , it s structure and properties directl y affect b atter y capacity,
cy cle lif e and saf ety. In addition, an ideal separator has to be electronical ly isolative ,
mechanically , dim ensionally, chemically and ele ctrochemica ll y stable. Furthermore, the separator
should not produce impurities, which could cause interference with the function of batteries. In
some applications, separators have to withstand the corrosive n ature of th e electrol y te at elevated
temperatures [58-60].
The major manufacturers of LI B s eparators alon g with their t y pical product s are listed in Table 1
[59] .

Table 1 Comm erci al separator propertie s from di fferent companies. Reprinte d with permission from Ref.
58.

The detailed require ments for a suitable separator in LI Bs are listed below:
1, the separator material must be chemically sta ble against the electrol y te especiall y , under the
strongly r eductive enviro nment of the anode and the ox idative environment of the c athode wh en
the battery is full y charged. At the s ame time, the separ ator shoul d not degrade and lose
mechanical strength dur ing batter y oper ation. An easy wa y to verif y chemical stabilit y is by
calendar life testing.
1 Introduction

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2, the sepa rator should possess a high wettability . Separators shoul d have the ability to absorb
significa nt amounts of liquid electroly te and also to retain the electrol yte d uring cell operation in
order to facilitate ion transportation.
3, a low thi ckness of a separa tor is beneficial for high energy and hi gh rate capability. How ever,
this adversely affects mechanical stren gth and safet y . In current technologies, the standard
thickness of separators for consumer r echargeable LIB s is ~ 25 µm . In addition, thickness
homogeneity is also desirable as variations in thickness may lead to cell failure.
4, an appropriate porosit y is necessar y to hold suf ficient liquid electrol yte and ensure ionic
conductivity betw een the cathode and anode. A t oo low porosit y will retain less electrol y te,
leading to internal resistanc e inc rease. However, a too high porosit y will adversel y impact the
mechanical performance of the separator. T ypically , the s eparators for LIBs have a po rosit y o f 40%
[59] .
5, a uniform pore size distribution is desirable within the total area of a separator film to avoid
performance d eca y. In a ddition, the pore size must be small er than the p article siz e of electrode
materials so no electrode particles can be transported. In practical cases, sub- µm pore size have
proven adequate in reducing the possibilit y of an internal short circuit b y bl ocking the penet ration
of larger elec trode particles meanwhile providing efficient conductivity f or li thium ions [57].
6, a separator should possess a mediate permeability for lithi um ions, i.e. t he separator should not
limit the transporta tion of lithium ions during battery operation. T ypically , the pr esence of a
separa tor increases th e effective resistance of electrolyte b y a fa ctor of 4 -5 [ 57]. The separator
with uniform permeability is essential for lon g c yc le li fe of a batter y . V ariations in permeability
will lead to an un even current densit y distribution, which has been v erified as one o f the reasons
for the for mation of lithi um dendrites on the anode [61] .
7, a sepa rator should p ossess a high mech anical stren gth and a high dimensional stabilit y .
Mechanical strength is characterized in terms of t ensile stren gth along the machine dir ection and
the transverse direction, the tear r esistance and the puncture str ength. The puncture strength is
defined as the max imum load required for a given needle to pun cture the separator. The st andard
minimum requirement for the me chanical and p uncture stren gth is 100 k g/cm -2 and 300 g for a
25 µm separa tor [61].
La stl y, the separator sho uld possess a low thermal shrinkage. In practice, the separator tends to
shrink when the temperature rises to the softening temperature. For example, polyethylene (PE)
can shrink as much as 10% when exposed to a temperature of 120 ℃ for 10 min. Thermal
shrinkage should be minimized. For LIB s, shrinkage of a separator is required to be not more
than 5% after 60 min at 90 ℃ [61 ] .
From the point view of c omposition and structure, batter y separators can b e broadly categorized
into:
1, microporous poly mer separators,
2, non-woven mats,
3, composite membrane s and
4, ge l-po y mer electrol yte separa tors [59] .
1 Introduction

- 13 -

Among them, the microporous pol y m er separators have been wide l y used in liquid electrolyte
batteries due to their comprehensive advanta ges of performance , s afet y and cost [ 58]. I n the
following section, we will focus on the trilayer separator developed by Celgar d ® .

1.1.4.1 Celgard ® 2325 separator

The purpose of the proposed trila y e r structure is to enhance inherent bat tery safet y b y shutti ng
down the pores o f the P E la yer (135 ℃ melting temperature) while maintaining the mechanical
stabilit y of th e sep arator through two PP (pol y pro pylene) la y e rs (165 ℃ melting temperature) in
case of internal chort circuit [58]. Fig.7 shows the currentl y investi gated three-l a y ered PP /PE/PP
Celgard® 2325 separator.

Fig. 7 Top view and cros s-sectional view of the Celgard® 2325 separa tor characte rized by scanning
electron microscopy (SEM). Sc ale ba rs i n left and rig ht figure are 1 µm and 5 µm, r espectiv ely. T he
imag e a re revised from r ef. [58]

In the case of over-heatin g, the middle PE la y er wi ll shut down the cell auto matically b y blocking
ionic pathwa y s upon melting, while the two PP lay ers can still provide mecha nical stren gth to
prevent ph y sical contact of the electrodes. From our investi gation, it is concluded that the
proposed thre e-la y er structure of the sep arator cannot withstand the force arising from the
growing lithium microstructures. Actually, delamination of the trilay er Celgard® 2325 separator
into three native lay ers can develop under the high pressure exerted by the grow ing li thium
microstructures. I n addition, melting of the PE separa tor can occur due to localized excessive
Joule heating re sulting from ISC. Section 2.2 presents a detailed investigation of this topic.

1 Introduction

- 14 -

1.2 X-ray imagin g set up

Recently , X-ra y ima ging based on either a laborator y X-ra y source or a s ynchrotron X-ra y sou rce
has been wid ely used as a powerful characterization to ol in material science [62, 63] . S pecifically ,
X-ray ima ging enables researchers to obtain fundamental insights into the evolution of electrodes
non-destructively. In the current disse rtation, all the investigations are conducted b y X -ray
radiography or tomography.
The underlying working principle of X-r a y imaging is based on different X-r a y absorption
coeff icients of dif ferent materials. During X -ra y imaging, an X-ra y beam produced b y the X- ray
tube or the s y nchrotron X-ra y facility p asses t hrough the LI B to b e c haracterized. Different
materials such as the cathode, anode and s eparator inside the LI B have different X -ra y absorption
coeff icients. The attenuated X -ray is finally captured b y a detector. Th en the obtained picture is
further anal y zed. In conventional X-ra y radio graphy , one 2D ima ge is obtained [64] . Through X-
ray tomography, three dimensional structures o f internal components of LI B s can be obtained
[65] .

1.2.1 laboratory X-ray imaging setup

The labora tor y X-ra y imag in g setup used in t his dissertation is shown in Fig. 8 and is located at
the Helmholtz-Zentrum Berlin, Germany.

Fig. 8 Photog raphy of the em ployed laboratory X-ray im a ging setup

The employed labor ator y X -ray im aging setup co nsists of an X-ra y tube (Hamamatsu, L8121-03)
and a fl at panel d etector (Hamamatsu, C7942SK-05). The micro focus X -ray source produc es a
cone beam which allow s for diffe rent ratios in de pendence of the relation betwee n source -
detector and source-sample distances. In this way, the field of view and the spatial resolution are
1 Introduction

- 15 -

tunable. The detailed parameters for th e labo ratory X -ra y tomography and radiograph y (section
2.4) are : 60 kV and 166 µA for the X-ra y tube v oltage and cu rrent, 58 mm and 500 mm for the
source- to -sample distance and sou rce - to -detector distance. Ea ch pixel represents 5.76 µm of the
sample. The obtained datasets are then normalized, filtered and reconstructed. Mor e detailed
information can be found in section 2.4.

1.2.2 Synchrotron X-ray imaging setup

The schematic il lustration of the BAMline used f or s y nchrotron X -ra y imaging is shown in Fi g. 9.

Fig. 9 Schem at ic illustrat ion of the em ployed s ynchro tron X- r ay imag ing setup at the BAM line.

The synchrotron bea m was monochromatized to 20 keV using a double multilayer
monochromator with an energy resolution of about 1.5 %. The detector s y stem comprised a 60 -
µm thick C dWO 4 scintillator, a microscopic optic and a pco4000 camera equipped with a
4008×2672 pixels CCD chip that is kept out of the direct beam b y using a mirror. In the p resent
dissertation, the synchrotron X-ra y imaging setup is employed to character iz e:
1, the morphological evolution of lithium microstructure s in LIBs(section 2.1);
2, the break-down of the commerc ial Celgard ® 2325 separator during c y cl es (section 2.2);
3, the three dimensiona l development of gas inside an operating battery (section 2.3);
4, the fracture b ehavior of ~ 20 µm siz ed S i particles investigated by s y nchrotron X -ra y
tomogra ph y (section 2.5);
5, different (de)lithiation behaviors of Sn p artic les investi gated b y in situ s ync hrotron X -ray
radiography.
The detailed parameters for the above measure me nts can be found in each section.

1 Introduction

- 16 -

1.3 The c ustomized cells

In order to ch aracterize LIBs b y X-ra y imaging, special LIBs h ad to be d esigned and fab ricated
because commerciall y a vailable LIBs are not compatible with this technique. For this purpose,
two protot y pes of a tomograph y cell (tomo -cell) a nd a radiograph y cell (radio -cell) are fabricated
as shown in Fig. 10, along with the schematic illustrations.

Fig. 10 Photog raphy and schem at ic illustration of the t wo cell proto t ypes employed. a), b) pho t ography of
the fabricate d tomo- cell and radio- cell, respec tively. c), sch em atic illustrat ion of the tom o -cell. d) and g )
schematic i llustration of the radio- cel l. e) enlarg ed elec trode part o f both c ells.

Both cell protot ype s are made of pol y amide-imide due to it s low X -ray absorption. S pecifically ,
the tomo-cell consists of a pol y amide-imide housing (brown), two screw electrodes and a
retaining s crews each on top (light grey), two sea ling rin gs ( y ellow), a negative electrode (blue),
a porous separator (white) and the counter positi ve elec trode ( green), as shown in Fig . 10 c. For
the radio-cell, from top to bottom, it consists of an upper housing (orange) , sealing ring ( yellow),
a ne gative electrode (blue) with copp er wi re, sep arator (light gray), a posit ive counter el ectrode
(gree n), a titanium foil current collector ( gray), a n annular copper cur rent collector (copper), a
lower housing ( orange), as shown in Fig. 10g. The dire ction of the penetrating X-ra y is
illustrated b y yellow arro ws in Fig. 10c and Fi g. 10d . More det ailed inform ation can b e found in
each section of chapter 2.

1 Introduction

- 17 -

2 Published parts o f work

2.1 Morphological evolution of lithium m icrostru ctures (paper titled “ Morpholog ical
evolution of electrochemically plated/stripped lithium mi crostructure s investigated by
sync hrotron X-ra y pha se contrast tomography ” , ACS NANO , 2016 , 10: 7990)

2.1.1 Supporting information

2.2 Break-down of the separator (paper titled “ Stud y of the Mechanisms of Inter nal Short
Circuit in a L i/Li cell b y S y nchrotron X-ra y Phase Contrast Tomogra ph y ” , ACS Energy Letters ,
2016 , 2: 94)

2.3 Gas developme nt (paper titled “ Three dimensional visualization of gas evolution and
channel formation inside a lithium-ion battery ” , ACS Applied Materials & Interfaces , 2016 , 8:
7156)

2.3.1 Supporting information

2.4 Degradation of lithium ion batteries base d on ~100 -µm-sized Si particles (paper titl ed
“ Inve stigation of failure mechanism s in silicon based half cells durin g the first c yc le b y micro X -
ray tomography a nd radiograph y ” , Journal of Power Sources , 2016 , 321: 174)

2.4.1 Supporting information

2.5 Fracture behavior of ~20- µm Si particles (paper ti tled “ S ync hrotron X-r a y tomographic
study of a sil icon ele ctrode befo re and after discharge and the effect of cavities on particle
frac turing ” , Chem EltroChem , 2016 , 3 : 1170)

2.5.1 Supporting information

2.6 Dif ferent (de)lithiation behaviors of Sn particles (p aper titled “ In Situ Radiographic
Investigation of (De)Lithiation Mechanisms in a Ti n-Electrode L ithium -Ion Battery ” ,
ChemSusChem , 2016 , 9: 946)

2.6.1 Supporting information

2 Published parts of work

- 18 -

2.1 Morp hological evolu tion of l ithium microstr uctures

Reprinted wit h perm i ssion from DOI : 10.1021 / acsnano.6b03939 . Copyrig ht (2016) Am erica n Chem ical
Society.
Morphological evolution of electr ochemically plated/stripp ed lithium microstructures
investigated by synchrotron X-ray phase contrast tomogr aphy

Fu Sun, *† ,§ L ukas Zielke, ǂ Henning Markötter, † ,§ Andre Hilger, † ,§ Dong Zhou, † ,§ Riko Moroni, ǂ
Roland Ze ngerle, ǂ Simon Thiele, ǂ ,# John Banha rt † ,§ and Ingo Manke §

† I nstitute of Mat erial Scien ce and Techn ologies
Technical Univ er sity Berlin
Strasse des 1 7. Juni 135, 10 623 Berlin, Germany
ǂ Laboratory f or MEMS Ap pl ications, I MT EK Depa rtm ent of Microsystem s Engineering
University of Freiburg
Georges- Koehler- Allee 103, 79110 F reiburg, G ermany
E-mail: Lukas.Z i elk e@im tek.de
§ Helm holtz Centre B erlin f or Material s and Energ y
Hahn- Meitner- Platz 1,
14109 Berl in, Germany
# FI T, University of F reiburg
Georges- Köhler-Allee 105, 79110 Freiburg, G ermany ,
*Em ail: fu.sun@helm hol tz-berlin.d e

Due to its low redox potential and high theo retical specific capacit y Li metal has d rawn
worldwide research atten tions because of its potential use in next-gene ration battery technolo gies
such as Li-S and Li-O 2 . Unfortunately, un controllable growth o f Li micro structures (LmSs, e.g.
dendrites, fibers ) during el ectrochemical Li s tripping/plating h as prevented their practical
commercialization. Desp ite various strategies pro posed to mi tigate LmS nucleation and/or blo ck
its growth, a fundam ental understanding of the un derly ing evolution mech anism s remains elusive.
Herein, s ynchrotron in -line phase contrast X-ra y tomogra ph y w as employed to investi gate th e
morphologica l evolution of electroc hemicall y de posited/dissolved L mSs non -destructivel y . For
the first time w e present a 3D characterization o f ele ctrochemically strip ped Li el ectrodes with
regard to electrochemicall y plated LmSs. We clarif y fundamentall y, th e ori gin of the porous
lithium interface growing towards Li electrodes. Moreover, cleavage o f the sepa rator caused b y
growing L mS wa s experimentall y observed and visualized in 3D. Our systematic investigation
provides fundamental insights into LmS evolution and enable s us to understa nd the evolution
mechanisms in L i electrodes more profoundl y.
2.1 Morphological evolution of lithium microstructures

- 19 -

Abstract Graph ic

Keywords : lithium microsctruc tures; morphological evolution; lithium ion battery ; silicon;
separa tor; lithium strip/plate

The further development of characterization methods has often promoted a better und erstanding
of the structure and function of mate rials. 1 For example, optical microscopy, 2 atomic force
microscopy (AFM), 3 scanning elec tron mi croscope (SEM), 4 transmission electron micros cop y
(TEM), 5 nuclear ma gnetic resonance (NMR) 6 and magnetic r esonance ima ging (MRI) 7 h ave been
employe d to und erstand the growth of the lithium microstructure ( LmS) in lithium -ion batteries.
Our apprehension of the growth mechanisms of LmSs has greatly impro ved 8 since usage of L i
metal as nega tive electro de began in the 1970s. 9 Nevertheless, prior anal ytical tools are inherently
limited. For instance, most techniques usually req uire prior sample removal from their as grown
environment for anal y sis, which is problematic since accidental ex posure of LmSs to air during
sample tra nsfer 10 and the L i -rinsing procedures applied fo r sample preparation 11 may
fundamentally change their morpholog y . Furthermore, n uclear resonance techniques suff er from
limited resolution of around 100 µm, the limited penetration depth of radiofrequenc y irradiation
and additional concerns arising from artefa cts introduced b y metals. 12 Recently , Eastwood et al.
have characterized the 3D microstructure of electrodeposited L mS b y s ynchrotron in -line phase
2.1 Morphological evolution of lithium microstructures

- 20 -

contrast X-ra y tomograph y at sub -µm resolution, 12 demonstrating that this m ethod is a suitable
and powerf ul tool for t he study of LmSs. I n another stud y , Harry et al. found that most Li
dendrites reside within t he ele ctrode instead o f the pol y mer s eparator during th e earl y stage of
dendrite development. They conclude that these d endrites preferentiall y g row from cr ystalline
impurities existing within an elec trode. 13
In this work, w e implement a fu rther s ystematic study by usin g in -line phase contrast X -ra y
tomogra ph y . We propose that the Li d endrites residing on th e surface of th e electrode arise from
the nascent electrochemically deposited LmS occupying the previously dissolved initial Li bulk
electrode . Furthermore, the present stud y clarifies, from a fund amental point of view, the origin
of the Porous L ithium Interfa ce (P LI) growin g towards the Li electrode, which has been
discovered previousl y. 14 In addition, an unexpec ted cleavage of the sepa rator caused by the
growing L mS is experimentally shown for the first ti me. These unprecedented findings, which
were not accessible fr om conventional mi croscopic characterizations and electrochemical
measurements, fundamentall y d eepen our understanding o f the evolution mechanism of the LmS
during el ectrochemical strippi ng/plating . The insights obtained could ope n up the wa y towards
new design principles and opportunities to further ameliora te or eliminate LmSs.
A proof-of-concept electrochemical cell, which is fully compatible with X -ra y ima ging and at the
same time representative for commercial lith ium -ion batteries ( LIBs), was designed and
manufactured as shown in Figure 1 along with a schematic illustration of the s y nchrotron X -ra y
imaging setup at BAMl ine of BESSY II, Berlin, Germany. 15 Two kinds of electrochemical c ells
were investi gated: a Li electrode coupled with a silicon -based composit e (Si/C) elec trode ( Li/Si
half-cell) and a L i electrode coupled with a Li electrode (Li/L i s y mmetrical cell). All c ells
employe d commer cial Celgard ® 2325 separators and were filled with 0.2 ml standard 1 M L iP F 6
in eth y lene carbonate (EC) and eth y l m eth y l carbonate (EMC) (EC/EMC= 50/50 (v/v)) electrolyte.
Details of parameters a nd cy cling routines of the investiga ted cells are given in the Methods
section and the S upporting Infor mation (S I). The validation of the electrochemical performance
of this proof-of- concept ce ll can be found in a previous report. 16 The cycling cu rves of all the
cells are shown in the SI. After c y cling, the cells were tr ansferred to the beamline to conduct
tomogra ph y without pri or disassembl y . Eve r y t omograph y was recorded with 2200 projections
covering a 180° rotation angle. A d etector s ystem with 0.438 µm pix el size was us ed . 17 The field
of view (FoV) is 1.7 × 1.2 mm 2 large (len gth × he ight), which enables the simultaneous imaging
of morphological changes occurring in all c omponent s of the electrodes. The detailed
normalization, tom ogra ph y reconstruction and 3 D presentation procedure is given in the Methods
section and in SI . 18

2.1 Morphological evolution of lithium microstructures

- 21 -

Figure 1 Schematic illustration of t he proo f -of- concept cell and the beam line setup. a, photograph and b ,
corresponding i llustra tion of th e e lectrochem ical cel l: po lyam i de -imide housin g (brown), two sc rews
serving as current collecto rs (light gray), two sealing r ings (yellow), cathode ( blu e), separato r (white) and
anode (green). c, schematic representati on o f the experim ental setup at BAMline BESSY II , Berlin,
Germ any. More inform at ion abou t the cell and t he setu p parameters can be fou nd in Me thods and SI .

Cross-sectional views of unc y cled Li/Li (Li/Li -N: N denotes sample cell nu mber. In this case, the
sample is Li/L i-1) and Li/Si (Li/Si-N: N d enotes sample cell number. I n t his case, the sample is
Li/Si-1) cells ar e shown in Figure 2a-c and Figure 3a-c, respectivel y , where we c an clearl y
discern the Li electrode, the separator and the co unter Li electrode (Si /C electrode) b y th e li ght
and dark interfaces between them arisin g from the in -line phase contrast. 19 I n both cases, the
interfaces between electrode and electrolyte of uncycled cells exhibi t a few minor noticeable
features (unev en surfaces shown in Fig ur e 2c an d Figure 3 c) caused b y t he high electrochemical
activity of Li. 13 The anode and cathode are l ab eled separately in Fig ur e 2 and Figure 3 (the
arra ngement is taken only for eas y demonstration: in sample Li/L i -2 (Figure 2d -l), the Li anode is
the bottom electrode , while the Li cathode is the upper. In sample L i/ Li-3 (Fig ure 2m-o), the Li
anode is the top, the Li cathode the bottom electrode. For all Li/Si half ce lls, the Li anode is the
top, the Si /C cathode th e bottom electrode ( Figure 3) ). Fi gure 2d,e show images of the Li/Li- 2
sample subjected to a constant current of 0 .3 mA · cm -2 for 10 h. T he corresponding 3D
re presentation is shown in Figure 2f. Co mpared to Figure 2a-c, the initiall y flat
electrode /electrol y te interfaces have turned in to rug ged interfaces with numerous
electroc hemicall y d eposited LmSs on the cathode side. In addition, different ly siz ed c avities can
be observed in the parti ally electrochemicall y di ssolved Li anode. I t has been reported that the
appearance of such cavities stems from t he uneven ly formed solid electroly te interface (SEI) and
that re gions containin g caviti es are elec tro chemicall y more active than flat surfaces. 20
Further more, the s eparator deforms int o an arched shape a nd adhere s clos ely to th e
electroc hemicall y dissolved porous stru cture, indi cating that significant stress is generated durin g
LmS formation.
In a nex t step, the Li/L i- 2 sample wa s continuousl y discharged galvanostat ically for another 19 h.
Fig ure 2 g-I illust rates that a large amount of Li is further dissol ved f rom the L i anode and
deposited onto the L i c athode, as indi cated by the increased volume of cavities within the anode
and by the inc reasing LmS s on top of the cathode . Consequently, some d elamination between the
separa tor and the two Li electrodes can be observed, which is comparable to previous reports. 11
After charging the L i/Li-2 sample for 17 h, the structure shown in Figure 2j-l is obtained. One
observes that : i) durin g Li strippi ng from the cathode, it is the remainin g Li bulk instead o f the
electroc hemicall y plated LmS that undergoe s electroche mical Li dissolut ion, as indicated b y the
newly formed void in the L i cathode (indicated b y white triangles in Figure 2j,k) , ii) durin g
plating of the anode wit h Li, Li deposition o ccurs p referentially in the s y stem of c avities
structures as indicated by th e newl y formed L mSs on the surfac e of the existing cavities (note
white tetragons in Figure 2j,k). I t has be en reported that the nascent electrochemically ina ctive
LmSs are still chemica lly active and can easily reac t with the el ectrol y t e to form a new SE I
2.1 Morphological evolution of lithium microstructures

- 22 -

coverage. 14 Consequently , t he SE I fo rmed electricall y insulates most of the nascent L mSs ,
thereby deactivating th em electrochemically du ring Li stripping. 21 The re ason for the preferential
deposition of L mSs in cavity-containing regions is that such areas are e lectrochemically more
active 22 or possess a hig h local ionic conductivity. 20
Fig ure 2m -o shows a c ross section of th e Li/L i-3 sample which has be en d ischarged at a constant
curre nt o f 0.3 mA · cm -2 for 30 h. Here, the unex pected phenomenon of cl eavage of the separator
due to the puncture by th e LmS is observed. The corresponding 3D re presentation is displa ye d in
Fig ure 2o. The pr esence of Li filaments runnin g across the separator is observ ed , as indicated b y
white arrows b etween the punc tu red separator in Figure 2m. The cleavage can be attributed to the
anisotropic mechanical properties of the Celgard ® separator in the transverse direction due to the
manufacturing process emplo y ed to create the micro -porous membr anes. 23 This is the first time
that the puncture of a sep arator is reported experimentally and displ ay ed in 3D. 3D
repre sentations of the morphological evolution of these cells are shown in Figure 2c, f, i, l, o . The
locations of the cross sec ti ons shown in Figure 2, along with the corresponding videos displa ying
the 3D nature of L mSs are available in S I.
2.1 Morphological evolution of lithium microstructures

- 23 -

Figure 2 Morpho logical evolution of Lm S in L i/ Li sy mmetrical ce l ls. a -c, uncycled pr istine Li/Li-1
sample. The Celgard ® sep arator i s sandw i ched betw een tw o Li electrodes. d- f, Li/L i-2 sample after
galv anostatic discharge fo r 10 h at 0.3 mA · cm -2 . Li metal has been disso lved from the Li anode (evidenc ed
by the visible cavities) and deposited onto the Li cathode (evidenced by the LmS). g -i, Li/Li-2 sa mple
after galvanosta tic dischar ge for another 19 h at 0.0 03 mA · cm -2 . An increasing amount of Li has bee n
dissolved from the Li anode and deposited onto the Li ca thode. j - l, Li/ Li-2 cell after galvanos tatic charge
for 19 h at 0.3 mA · cm -2 showing t hat t he residual Li bulk (indicated by white triangles ) instead of Lm Ss
has been dissolved and preferentially de posited onto the surface of existing cav ities (marked by white
tetragons). m- o, Li/ Li-3 sam ple galvanosta tically discharged for 30 h at 0.3 mA · cm -2 showing filamentous
Lm Ss running across the separator (indicated by arrows) and the torn Celg ard ® separa tor. The first colum n
(a, d, g, j , m) contains r ec onstructed X - ray tomographies. The second column (b, e, h, k, n) shows the
same slice afte r a com bination of manual and autom ated phase filtering and color labeling . The third
column (c, f, i, l and o) show s the corresponding 3D representations. A ll the scale bars are 50 µm long.
The cycling procedures of all cells, the detailed phase filtering procedure an d 3D presentation methods are
described in Me t hods and SI .
2.1 Morphological evolution of lithium microstructures

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More information about the evolution of L mSs c an be obtained from tom ographies showin g the
Li/Si half-cells. Figure 3d-f displa ys c ross-sectional views of the L i /S i-2 sample that wa s
ga lvanostaticall y dischar ged for 15 h at 0.3 mA · cm -2 . Compared to the unc y c led Li electrode
sample (Figure 3a - c) , a good deal o f Li has be en electrochemicall y stripped from the bulk anod e
Li ( evidenced b y th e cavities formed as marked b y white trian gles in Figure 3d,e) and
subsequently allo y ed with Si particles (evidenced by the cracks formed as marked by white
diamonds in Figure 3d,e) . Moreov er, after the Li/Si -1 sample has been di scharged for 15 h and
then charged for 5 h at 0.7 mA · cm -2 , some but not all voids generated during electrochemical
dissolution are domi nate d by the nascent electrochemicall y deposited L mSs, as shown in Figure
3g- i. This non-uniform filling of voids may be the result of the uneven distribution of L i ion
flux. 8b Finally, the L i/Si-3 sample, which has experie nced 5 cy cl es, is shown in Figure 3j-l. We
observe th at a nearl y 1 00-µm thick portion of the initial Li anode has been diss olved and is
thereafter occupied by electrochemicall y generated L mSs. On the Si cathode, almost all o f the Si
particles have been fully lithiated to SiL i x (1 <x<4) and are no longer obs ervable due to the low
X-ray attenuation coefficient . 24 3D re prese ntations of the morphological evolution of the L mS
formed in Li/Si cells are shown in Fi gure 3c, f, i, l. The locations o f the cross sections shown in
Fig ure 3, along with the corresponding videos sh owing the 3D natur e of t he L mS are available in
SI.

Figure 3 Morphologica l ev olut ion of LmS in Li/Si half cells. a -c, uncycled pristine Li /Si - 1 cell. The
Celgard ® separator is sandw iched be tween the Li anode an d the Si/C composi te cath ode. d - f, Li/Si- 2
sample that has been galv anos tatically discharged f or 15 h at 0.3 mA · cm -2 . Here, Li has been dissolved
from t he Li anode ( evid enced by cavity/void, white triangles) and alloyed wit h t he Si/C cathode
(evidenced by Si crack form ati on, white tetragons). g -i, Li/Si-1 cel l after 15 h of discharge and 5 h of
charge at 0.7 mA · cm -2 . The cavities formed are partially occupied by newly formed LmSs. j -l, Li/Si - 3 cell
that has experienced 5 cycles ( see SI). Here, nearly half of the orig inal Li anode has been dissolved and is
now occupied by deposited LmSs. T he first column (a, d, g, j) contains r econstru cted tomog raphy slices.
The second colum n (b, e, h, k) shows the sam e slice after a combination of m anual and automated phase
2.1 Morphological evolution of lithium microstructures

- 25 -

filtering and color labeling. The third colum n (c, f, i , l ) contains the correspond ing 3D representations. All
the scale bars are 50 µm l ong . The cycling procedu res of all cel ls can b e found in S I.

Base d on the se s y st ema tic inv estigations, we suggest a model that describes the phenomena
observed in Fi gures 2 a nd 3 . A s chematic illust ration of the proposed evolution is shown in
Fig ure 4. Durin g the 1 st electrochemical Li stripping, a si gnifica nt amount of Li dissolve s from
the Li anode ( yellow), and numerous cavities (white) are generated. Mea nwhile, these L i ions, i )
are electrochemicall y de posited int o the L mS on t he hostless L i cathode (as observed in Fi gure 2)
or, ii ) form an alloy with the Si/C composite electrode (as obse rved in Figure 3). The purple-
y e llow color gradient in Figure 4 indicates the two cases observed in Fig ure 2 and Figure 3.
During subsequent electrochemical Li plating, th e Li ions are transported back to the Li anode
and are preferentially de posited in the previously formed cavities due to the high electrochemical
local activity there. However, during the followin g (the 2 nd ) L i stripping from the anode, it is the
remaining bulk Li ( y ell ow) that undergoes electroche mical dissolution instead of the nascent
electroc hemicall y deposited L mS (ocre, formed during previous L i plating ) which is covered by
electron- blocking SEI. Summarizing this process one can sa y that the in itial bulk L i anode is
electroc hemicall y dissol ved continuousl y in each L i stripping step, which r esults in numerous
cavities. During subsequent L i plating, lar ge amounts of the LmS are generated and form the P LI ,
which occupies the voids formed during previous Li dissolution. It has been r eported that a
continuously growing PLI c an substantia lly in crease the intern al re sis tance of a cell, which
eventuall y r esults in poor cell p erformance and ultimately in cell failure. 14 Meanwhile, thin LmS
protruding fr om the PLI towards the separator conducts high currents and consequently
experiences large lo calized Joule heating, which results in the melting of t he separator and fin ally
puncture of the separator. 25

Figure 4 Schem atic illustration of t he morphological evolution of LmSs in a Li anode: fr om left to right, a
pristine state ce ll; Li anode then under goes the 1 st stripping during discharge; Li anode undergoing the 1 st
plating dur ing charge; Li anode undergoing the 2 nd stripping during another discharge; Li anode after N th
stripping/pl ating (N ≥ 2).

The observations of the evolution of electrochemically stripped/plated LmSs presented here shed
new li ght on a ran ge of proc esses th at pr eviousl y h ave not been acknowledged or b arely
2.1 Morphological evolution of lithium microstructures

- 26 -

considered. First ly , the proposed mechanism of LmS evolution can explain the ori gin of the
electroc hemicall y generated P LI th at grows towards the Li bulk electrode. Based on ou r
observations, both the electrochemical dissolut ion of the initial Li bulk electrodes and th e
following electrochemical deposition of nascent LmSs contribute to the inward growth of the P LI.
Secondly, the ev aluation of using Coulombic e fficienc y to investigate the c y cling efficienc y of L i
in L i s y mmet rical cells is compromised by the el ectrochemically inert LmSs . 20 According to our
observations, a significant amount of L i is stripped from the ini tial L i bulk instead from the
ge nerated nascent (electrochemicall y inac tive) LmSs to compensate for the depletion of Li ions in
the electrol y te used to be plated. Thirdly, concerns about using Li m etal as a standard counter
electrode for the evaluation of the performance of new battery materials using half-cells are
raised. In fact, it has been confirmed that the PLI, which continuously g rows in each c y cle, can
markedly increase the in ner cell resistance and is the actual origin of the o nset of cell degradation
and failure. 14 Finally , considering that the dissol ution of the initi al L i b ulk electrodes and the
deposition of nascent LmSs (nece ssitated b y th e char ge tr ansfer within cell s ) remain intrinsic and
unavoidable, the pr esent results sugge st that no vel fundamental st rategies that involve direct
eng ineering of Li electrodes are desirable and necessary. For ex ample, strategies such as
introducing surface patterns to the L i electrode to direct and control Li deposition , 26 replacing
conventional Li foil electrodes by organic-coated Li powder 27 or emplo y i ng different substrates
to selectively deposit and then enca psulate the LmS, 28 appea r more suitable than conventional
mechanical suppression or formation of protective passivation la y ers. 29
Our work represe nts a major step fo rward in understanding the evolution of electrochemicall y
stripped/plated L mSs and may speed up the process of introducing Li in to commerc ial lithium
metal batteries (LMBs) with enhanced performance. On the one h and, ground breakin g
electrolytes a re hi ghl y needed to facilit ate the nas cent electrochemically deposited LmSs to plate
smoothl y onto the Li bulk electrode inst ead of fo rming a hi gh ly resistive PLI. On the other h and,
it is also worth investigating electrochemical Li stripping to enable a more uniform ac cessibility
of L i ions from not onl y the initi al bulk L i electrode but also from the electrochemicall y plated
LmS. In addition, engineerin g a sophi sticated architecture o f the Li electrode will be
technolog icall y crucial. Finall y, ne w separator s with high ion conductivit ies and improved
rigidities are hig hl y desirable.
Mater ials and Methods
Materials
Lithium, carbon bla ck, Pol y vin ylidene difluoride (PVDF) binder and C elgard separator w ere
purchase d from MTI Cor. USA. Silicon (Si) was rece ived from El kem AS , Norwa y . Electrol yte
1M LiPF 6 in a volume-ratio mix ture (1:1) of ethy l ene carbonate (EC) and e thyl meth y l carbonate
(EMC) was pur chased from Sigma Aldrich as well as N -meth y l p y rrolidone solvent (NMP). The
housing of the proof -of-concept beamline batter y is made of pol y amide-imide (Torlon) provid ed
by Drake Plastics Europe.
Battery Preparation
2.1 Morphological evolution of lithium microstructures

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The li thium electrode s in both L i/ L i s y mmetrical ce lls and L i/S i half cells were punch ed out
(2.5 mm diameter) from a lithi um plate (1 mm th ick). The Si composite electrode was made o f
electrode slurries contai ning w eight ratios of Si:carbon bla ck:binder o f 70:20:10 in NMP. Two
kinds of Si/C electrodes were prepared: ele ctrodes containing i) Si particles with a diameter of
~20 µm (for Li/Si-1 and L i/Si-3); ii) S i pa rticles with diameters o f ~20, ~80 and ~150 µ m in a
weig ht ratio of 1:1:1 (for Li/Si-2). The slurr y was cast onto copper foil. To remove the NMP , the
cast copper foils were dried in an oven at 60 ℃ for 12 h. After dry ing, th e composite electrode
was diced into small blocks of around 1.7 mm × 1.7 mm × 0.2 mm size with a razor blade. For
the L i/Si half cells, L i was used as the anode and a S i/C elec trode as the cathode. For the Li/Li
symmetrical cells, L i w as used both as anode and cathode. All cells (6 in total, 3 Li/L i c ells from
Li/Li-1 to L i/ Li-3, 3 Li/Si cells with L i/Si -1 to L i/ Si-3) were assembled in an a rg on- filled
glovebox with humi dit y and ox yge n levels belo w 0.1 ppm. The poly mer separator ( 3.5 mm in
diameter and ~25 µm thick) was placed between the Li/Li electrode (in Li/Li c ells) and the Li/Si
electrode (in Li/Si c ells). All cells were ass embled manuall y without exert ing force. Finally, both
cells were filled with the liquid electroly te and were sealed off before taking out of glovebox.
Electrochemical Measurement
Galvanostatic charge and discharge of all the batteries was carried out using an IviumStat from
Ivium T echnologies, N etherlerlands. The cells were subj ected to different current d e nsities for
discharge and charge . The c y cling procedure for each cell is detailed in SI.
Setting of Tomography Measurement
Synchr otron X-ra y tomo graph y was carried out at the BAMline at BESSY II of the H elmholtz -
Centre Berlin, G ermany. The s ync hrotron be am was monochromatized to 20 keV usin g a double
multilayer monochromat or with an energ y resolut ion of about 1.5 %. The detector s ystem
comprised a 60-µm thi ck CdW O 4 scintillator, a mic roscopic optic and a p co4000 camera
equipped with a 4008×2672 pixel s CCD chip that is kept out of the dir ect beam b y using a mi rror.
For tomography measurements of all cells, 2200 projections within a 180° batter y rotation we re
recor ded. For the Li/ L i s ymmetrical cells, the exposure time was 2.5 s, for Li/Si half cells 3 s.
Data Processing
The raw tomography dat a was filtered, normaliz ed and reconstructed using code pro grammed in
IDL 8.2. Three-dimensio nal segmentations of the separators were made using a grid of manuall y
marked points that were fitted with a biharmonic equation using MA T LAB. For the se gmentation
of L mSs and voids, the s tatistical region merger t ool implemented in Fiji 30 was used followed b y
manual removal of the bulk lithium backg round. F or the three-dimen sional pre sentations in
Fig ure 2 and 3, th e se gmentation of all material s except the L i bulk (Li bulk set to 0, all other
materials to 1) was used as a mask for th e a ccording X -ra y dataset in order to b e able to show
LmSs and voids without the Li bulk in the background. Two -dimension al segmentations w ere
made b y first manuall y segmenting the separator, followed b y a statis tical reg ion me rging and
2.1 Morphological evolution of lithium microstructures

- 28 -

manual removal of the Li bulk back ground. The tomograms shown in the first column of Fig ure 2
are absorption contrast tom ogra ms while the second column displ ays the aver a ges of the
accor ding absorption and phase contrast tomo grams. All remaining tomograms show absorption
contrast only .
Remark concerning the expe rimental program
The effect of the pressure subjected to both cells and the electrode morpholog y on LmS
formation was currentl y not the scope of thi s pap er. Moreover, we have used a relatively small
curre nt and lon g cyc le t ime to induce as much as possible Li diss olution and deposition. The
effect of different C rates on the morpholo gy of L i diss olution and deposit ion was not
investigated. Nevertheless, these important par ameters are worth to be studied in the future.
A quantitative analysis of the amount of lithium that has been dissolved from the Li anode and
the amount of L mS that has been deposited onto the sur face of the Li c athode (allo y ed with the
Si/C cathode) has presently not b een condu cted mainl y because the currentl y used Li electrode
(2.5 mm diameter) is bi gger than the field of view (FoV), which is 1.7 mm wide. The current th at
passes through the chemical workstation may give rise to dissolut ion or deposition of L i outside
the FoV. Considering t his, the present research focuses on the qualitative analysis o f the
morphologica l evolution of L mSs in Li electrodes. However, cells assembled with L i electrode s
smaller than the F oV could be measured in the future and anal y sed quantitatively.
Acknowledgements
We thank Dr. Heinrich Riesemeier, the beamline scientist at BESSY II, for his valuable
assistance and engineer Norbert Beck for f abricating the be amline ba tter y. We also thank Elkem
AS for providing us with the Si particles. This work is sponsored b y the Helmholtz Association
and the China Scholarship Council.
Author Contributions
Fu S un, Henning Markö tter and Ingo Manke designed the cell. Fu Sun, Henning M a rk ötter and
Dong Zhou assembled and tested the cells. Andre Hilger and Henning Markött er aided in
sync hrotron characterizations and the discussion of results. Lukas Zielke, Andre Hilger, Roland
Ze ngerle, Riko Moroni and Sim on Thiele rec onstructed the data and pr epared the figure s. Fu Sun,
Henning Markötter, Ingo Manke and J ohn Banhart contributed to the data interpretation and the
discussion of the results. Fu Sun, Henning Mark ötter, I n go Manke and J ohn Banhart composed
the manuscript. Ingo Manke a nd John Banhart directed the work.
Additional inform ation
Supplementary information is available.
Competing F inancial Interests statement
The authors declare no competing f inancial interests.
2.1 Morphological evolution of lithium microstructures

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Ion Battery Separ ators. J. Power Sources 20 11, 196 , 2905-2912.
24. Gonzalez, J.; Sun, K.; Huang , M .; Lambro s, J.; Dill o n, S.; Chasiotis, I., Three Dimensi o nal Studies of
Particle Failure in Silic on Based Co mposite Ele ctrodes for Lithium I on Batteri es. J. Power Sources 2014,
269 , 334-3 43.
25. Jana, A.; Ely, D. R.; García, R. E., Dendrite-Separa to r Interacti o ns in Lithium-B ased Batte ries. J.
Power Sources 2015, 2 75 , 912- 921.
26. Ryou, M.-H.; Le e, Y. M. ; Le e, Y.; Wint er, M.; Bi eker, P., Mechanical Surface Modification of
Lithium Me tal: Towards I mproved Li Metal An ode Performance b y D irected Li P lat ing. Adv. Funct. Mater.
2015, 25 , 834-841.
27. Heine, J.; Krüg er, S.; Har tni g, C.; Wietel mann, U.; Wint er, M.; Bi eker, P., Coated Lithiu m Powder
(Clip) Electrod es for Li thium-Metal B atteries. Adv. Energy Mater. 2014, 4 , 1 3008 15-130082 2.
28. Yan, K.; Lu, Z. ; Lee, H.-W.; Xiong , F.; Hsu, P.-C.; Li, Y.; Zh ao, J.; Chu, S.; Cui, Y ., Selec tive Depositi o n
and Stable Encapsul ation of Lithium thr ough Heteroge neo us Seeded Gr o wth. Nat. Energy. 2016 , 1 ,
16010-16018.
29. Zheng, G.; Lee, S. W. ; Liang, Z.; Lee, H.-W.; Yan, K. ; Yao, H.; Wang , H. ; Li, W.; Chu , S.; Cui, Y.,
Interconnected H o llow Car bon Nanospheres f or Stable Lithiu m Metal Anodes. Nat N ano 2014, 9 , 6 18-
623.
30. Abramoff, M. D.; Magalhã es, P. J.; Ram, S. J., I mage Processing with Imagej. Biophotonic s
international 2 004, 11 , 36- 42.

2.1 Morphological evolution of lithium microstructures

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2.1.1 Suppor ting informa ti on

Morphological evolution of electrochem ically plated/stripped li thium microstructures
investigated by synchrotron X-ray phase contrast tomogr aphy

Fu Sun, *† ,§ L ukas Zielke, ǂ Henning Markötter, † ,§ Andre Hilger, † ,§ Dong Zhou, † ,§ Riko Moroni, ǂ
Roland Ze ngerle, ǂ Simon Thiele, ǂ ,# John Banha rt † ,§ and Ingo Manke §

† I nstitute of Mat erial Scien ce and Techn ologies
Technical Univ er sity Berlin
Strasse des 1 7. Juni 135 , 10623 B erlin, G ermany
ǂ Laboratory f or MEMS Ap pl ications, I MT EK Depa rtm ent of Microsystem s Engineering
University of Freiburg
Georges- Koehler- Allee 103, 79110 F reiburg, G ermany
E-mail: Lukas.Z i elk e@im tek.de
§ Helm holtz Centre B erlin f or Material s and Energ y
Hahn- Meitner- Platz 1,
14109 Berl in, Germany
# FI T, University of F reiburg
Georges- Köhler-Allee 105, 79110 Freiburg, G ermany ,
*Em ail: fu.sun@helm hol tz-berlin.d e

This section includes:
Detailed cycling procedure s for al l cells
Figure s S1- S8
Captions for the S upplem ent ary Movies

Fig ure S1 and S2 specify the locations of the cros s-sections of the Li/L i s ymmetrica l cells shown
in Figure 2 and Figure 3, respectivel y.
All the sy mmetrical L i/ Li cells and Li/Si half c ells are g alvanostaticly discharge d and char ged
using an I viumStat electrochemical wo rkstation. The L i/Li-2 cell w as discharged at 0.3 mA cm -2
for 10 h and a tomograp hy was conducted as shown in Figure 2 d-f. Th e corresponding discha rge
curve is displayed in Figure S 3a. After thi s, the Li/Li-2 cell was continued to be galvanostatically
discharged for 19 h at 0.3 mA cm -2 and another t omograph y was conducted, as shown in Figure
2g-i. Th e corresponding discharge curve is shown in Figure S3b. Finall y, the Li/Li-2 cell was
charge d for 19 h at 0.3 mA cm -2 and the third tomograph y of this cell w as conducted as shown in
Fig ure 2j -l. The corres ponding charge curve is shown in Figure S3c. Note that during the
2.1 Morphological evolution of lithium microstructures

- 32 -

prolonged discha rge and cha rge p rocess, there is a rise in th e dischar ge or ch arge voltage. The
observed rise in the overpotential is simil ar to tha t reported previous ly 1 an d the r eason ma y b e a
continuous electrolyte loss during prolong ed c yc ling 2 .
The L i/Li-3 cell was discharge d at a constant current of 0.3 mA cm -2 for 30 h and after that a
tomogra ph y w as conducted as shown in Figure 2m -o. The corresponding discharge curve is
shown in Figure S4 . I t has to be noted that the voltag e flu ctuation may a rise from the growing
LmS in agreement with previous reports 3 .
For the Li/Si-1cell, a pristine -state (without an y di scharge or charge) tomograph y was conducted
as shown in Figure3a-c. After that, the L i/Si-1 cell was discharged fo r 15 h and then charge d for
5 h at 0.7 mA cm -2 . After this c yc le, ano ther tom ograph y was conducted as shown in Figure 3 g-i.
The corresponding discharge a nd charge curve is shown in Figure S5 .
The L i/Si-2 cell was galvanostatically disch arged for 15 h at 0.3 mA cm -2 and then a tomogra ph y
was conducted as shown in Fi gure 3 d-f. Th e co rresponding discharge curve is shown in Figure
S6 .
For L i/Si-3 cell, it experienc ed 5 c y cles with different current densities: for the 1 st cy cle the
curre nt was of 0.714 mA cm -2 for 1.3 hours; for the 2 nd c y cle the current was of 0.706 mA cm -2
for 1.7 hours; for th e 3 rd cycle the current was of 0.3 mA cm -2 for 2.4 hours; for the 4 th cycle the
curre nt was of 0.159 mA cm -2 for 13 hours; for the 5 th cy cle the current was of 0.083 mA cm -2 for
44 hours. The dis charge/charge curves are sho wn in Fi gure S7 . After these procedures, one
tomogra ph y was conducted as shown in Figure 3j-l .
For a 3D representation of L i/Li symmetrical cells it is useful to enhance the phase contrast
between LmS, lthium and separator 4 . W e us ed a phase backpropagation filter prior t o
tomogra phic reconstructi on of all the Li/Li cells as displa y ed in Figure S8 .

Figure S1 Loc ation of the cross-sections shown in Fig ur e 2 in the main text. a, a 3D presentation of the
assembled Li/Li cells. b, t he das hed l ines a), d), g), j) and m) ar e indicativ e of t he l ocations of the cross -
section planes shown in Figure 2a, d, g , j and m .
2.1 Morphological evolution of lithium microstructures

- 33 -

Figure S2 Loc ation of the cross-sections shown in Fig ur e 3 in the main text. a, a 3D presentation of the
assembled Li/Si cells. b, the dashed a), d), and g)/j) l ines ar e indicative of the locations of t he cross -
section planes i n Fig ure 3a, d, g and j.
2.1 Morphological evolution of lithium microstructures

- 34 -

Figure S3 Cycling profiles of Li/L i - 2 cell. a , discharged at 0.3 mA cm -2 for 10 h. b, discharged for anoth er
19 h at 0.3 mA cm -2 . c, finally charg ed for 19 h at 0.3 mA cm -2 .
2.1 Morphological evolution of lithium microstructures

- 35 -

Figure S4 Cy cling profiles of Li/Li-3 cell di scharged a t 0.3 mA cm -2 for 30 h.

Figure S5 Cy cling profiles of Li/Si-1 cell f irst discharg ed for 15 h and then cha rged for 5h at 0.7 mA cm -2 .

Figure S6 Cy cling profile o f Li/Si-2 cell discha rged fo r 15 h a t 0.26 mA cm -2 .
2.1 Morphological evolution of lithium microstructures

- 36 -

Figure S7 Cycling profiles of Li/Si - 3 ce ll. T he inset shows t he 1 st , 2 nd and 3 rd c ycle; the 4 th cycle and 5 th
discharge cu rve are shown i n the m ain figure.

2.1 Morphological evolution of lithium microstructures

- 37 -

2.1 Morphological evolution of lithium microstructures

- 38 -

Figure S8 C ross- sections of r ec onstruc ted Li/L i sym m etrical cells after using phase backpropagati on
filtering . a, tomog r aphy of pri stine Li/Li-1 cell . b, tomog raphy sli ce of Li/Li-2 cell after 10 h of discharg e
at 0.3 mA cm -2 . c , tom ography slice of Li/Li-2 cell after another 19 h of discharg e at the sam e current
density. d, tomography sli ce of Li/Li-2 cell after 19 h of charge at the sam e curr ent density. e, tomography
slice of Li/Li- 3 sa m ple after 30 h of discharg e at 0.3 mA cm -2 .

Caption s for the Su pplementar y Movies
http://pubs.acs.org / doi/ab s/ 10.1021/a csnano. 6b0393 9

SM -1- Li - Li -1_Pristine
SM -1- Li - Li -1_Pristine-Colored-Figure2a-c
SM -2- Li - Li -2_Discharge for 10 Hours
SM -2- Li - Li -2_Discharge for 10 Hours-Colored-Figure 2d-f
SM -3- Li - Li -2_Discharge for another 19 Hours
SM -3- Li - Li -2_Discharge for another 19 Hours-Colored
SM -4- Li - Li -2_Charge for 17 Hours
SM -4- Li - Li -2_Charge for 17 Hours-Colored-Figure 2j-l
SM -5- Li - Li -3_Discharge for 30 Hours
SM -5- Li - Li -3_Discharge for 30 Hours-Colored-Figure 2m-o
SM -6- Li - Si -1_Pristine
SM -6- Li - Si -1_Pristine-C olored-Figure 3a-c
SM -7- Li - Si -2_Discharged for 15 Hours
SM -7- Li - Si -2_Discharged for 15 Hours-Colored-Figre 3d-f
SM -8- Li - Si -1_Discharged for 15 Hours and then Charged for 5 Hours
SM -8- Li - Si -1_Discharged for 15 Hours and then C harged for 5 Hours-Col ored-Figure 3g-i
SM -9- Li - Si -3_C yc led for 5 c ycles
SM -9- Li - Si -3_C yc led for 5 c ycles-Colored

Reference s
1. Harry, K. J. ; Hallinan, D. T.; Parkins on, D. Y.; Ma cDowel l, A. A.; Balsara, N. P., Detection of
Subsurface Structure s Underneath Dendrites For med on Cycled Lithium Me tal Electr o des. Nat Mate r
2014, 13 , 69- 73.
2. (a) Sarasketa-Zab ala, E.; Aguesse, F.; Villarreal, I.; Rodriguez-Martinez, L. M.; Ló pez , C. M.; Kubia k,
P., Understand ing Lithium I nventory L oss and Sudden P erformance Fad e in Cylind rical Cells Durin g
Cycling with Deep-Discha rge Steps. J. Phys. Chem. C 2015, 119 , 8 96-906; (b) Cai, L. ; An, K. ; Feng, Z.; Liang ,
C.; Harris, S. J., In-Si tu Obse rvation of Inhomog eneous Degradati on in Large F o rmat Li-Ion Cells by
Neutron Diffracti on. J. Power Sources 2013, 2 36 , 163- 168.
3. Wu, H.; Zhu o, D.; Kong, D.; Cui, Y., I m proving Batt ery S afety by E arly Detecti o n of In ternal
Shorting with a Bifuncti onal Sep arator. Nat Co mmun 2 0 14, 5 , 51 93-5199.
4. Eastwood, D. S.; Bayle y, P. M.; Chang, H. J.; Tai wo, O. O.; Vila-Comamala, J.; Brett, D. J. L.; Rau, C.;
Withers, P. J.; Shearing, P . R.; Gre y, C. P. ; Lee, P. D., Three-Dimensi onal Charac terization o f
Electrodep osited Lithiu m Micr ostructures Using Sync hrotron X -Ray Phase Contrast Im aging. Chem.
Commun. 2015, 51 , 266-26 8.
2.1 Morphological evolution of lithium microstructures

- 39 -

2.2 Break-do wn of t he separator

Reprinted with permission fr om DOI : 10.1021/ac senerg ylett.6b00589. Copy ri ght (2017) Am erican
Chem ical Society .
A Study of th e Mec hanisms of Internal Short C ircu it in a Li /Li cell by
Synchrotron X-r ay Phase C o ntrast Tom og raphy

Fu Sun,* ;†;§ Riko Moroni, ǂ Kang Dong , †;§ Henning Markötter, § Dong Zhou, †;§ André Hilger, §
Luka s Zielke, ǂ Roland Zengerle, ǂ Simon Thiele, ǂ ,# John Banhart, †,§ and I ngo Manke §

† Inst itute of Mater ial Sci ence and Techn ologies
Technical Univ ersity Berlin
Strasse des 17. Juni 135, 106 23 Berlin, Germ any
ǂ L aboratory for MEMS Appli cations, I MTEK Departm en t of Micr osy stem s Engineeri ng
Univer sity of Frei burg
Georges- K oehler- Allee 103, 79110 F reiburg, Germ any
§ Helm holtz Cent re Berli n for Materi als and En ergy
Hahn-Mei tner- P latz 1,
14109 Berl in, Germ any
# FI T, U niver sity of Frei burg
Georges- K öhler- Allee 105, 7 9110 F reiburg, Germ any

ABSTRACT
Synchr otron in -line phas e contrast X -ra y tomograph y w as emplo yed to inv estigate the underl y ing
internal cell deformation and degradation caused by an inter nal short circuit (I SC). B y comparing
the reconstructed tom ography of one unc y cl ed L i/ Li cell and one short -circ uited L i/Li cell for the
first time, we experime ntally demonstrate: 1) most of the electrochemicall y deposited lithium
microstructures ( LmSs) are electroc hemically inert during the following Li stripping, 2) the
electroc hemical stripping and plating pro cess during discharge and charg e are highly
inhomogene ous, 3) delamination into three native la yers and the partial melting of th e
commercial trila y er Celgard ® 2325 separator can develop during real I SC event, 4)
decomposition of the solid electrolyte interface (SE I ) coverin g L mS s and partial melting and re -
solidify ing o f porous LmS can also occur du e to the localiz ed ex cessive J oule heating resulted
from the ISC ev ent. These unexpected insi ghts int o the int ernal cell de gradation and deformation
mechanisms caused b y ISC shed new lights on enhancing the prop erties of separators and could
open up new design principles and opportunities to fundamentally impr ove the reliabilit y and
safety of c urrent- and/or nex t-generation LIBs.
2.2 Break-down of the separator

- 40 -

Abstract Graph ic

Keywords: lithium mi crostructure ; internal sho rt circuit; dendrites; separator; lithium ion battery;
sync hrotron X-ra y pha se contrast tomography
Introduction

Safety and reliability of lithium ion batteries (LIBs) that are targeted for high power and hi gh
energy applications are of major concern. [1 - 8] For example, the most promising L i metal-based
battery technologies such as lithium -sulphur and lithium-air batteries hav e not been successfull y
commercialized mainl y because o f the uncontrolled grow th of lithi um microstructures (LmSs)
such as dendrites, fibe rs, whiskers and moss etc., which can cause internal short circuit ( ISC) of a
cell and result in catastrophic fires or even explosions. [9 -11] Recent field inc idents such as fires on
a Boeing 787 D reamliner flight and in a Tesl a electric v ehicle (EV) or even the S amsung note 7
phone’ s exploding are be lieved to be closel y li nked to ISC in LIBs. [ 12-14] Furthermore, the soaring
number of EVs in use additionally adds tremendous weight to safety and reliabilit y in present -
and next-generation LIB technolog y. In order to gain a detailed understanding of int ernal cell
deformation and degradation leading to ISCs, vari ous t esting techniques su ch as usin g a blunt rod
to indent a cell as d eveloped b y the Unde rwrites L aboratory and the U.S. Advanced Batter y
Consortium, surface ind entation or pinch test developed b y Motorola, and the forced ISC
approac h developed b y t he Batter y Associ ation of J apan have been employed. [1 5] The purpose of
these tests is to create a s mall break in the separator by exerting an ex ternal force, mimicking the
ISC events that ma y lead to failed accidents in LIBs. Unfortunatel y, recent investigations have
demonstrated that such mecha nical intrusion methods are not e ntirel y representative of true
spontaneous I SCs due to the destructive and post -mortem nature of characterization, which
2.2 Break-down of the separator

- 41 -

eliminates much of the in sit u deformation information. [ 16, 17] Althoug h a non-m echanical
approac h to initiate an ISC has been reporte d by Ch ristopher et al. by using different melting
point metals and metal alloys, [15] the c hallenge of a detailed characterization of an ISC by
sophisticated charac terization tool in a non-destructive way remains.
To this end, various investigations emplo y ing li ght microscopy , [ 18] nuclear magnetic resonance [1 9]
(NMR) and 7 Li magnet ic resonance imaging [ 20] (MRI), scanning electron microscopy [21 , 22 ]
(SEM) and transmission electron microscopy [23 , 24] (TEM) have been carried out to chara cterize
the complicated LmS growth, which is regarded to lead to the ultim ate I S C event. However, these
reports mainl y focus o n the growth of LmSs and ignore the accompanying morphological
evolution of the separators as well as the correlation between growing LmS and cell failure .
Hitherto, no ex periments have been reported in the battery communit y that aims at investi gating
the mechanisms of I SC s leading to (internal) cell degradation. W hilst diffe rent strategies have
been proposed to avoid ISC-induced batter y fail ures, [25-32] it remains of fundamental interest to
directly visualize the ISC caused b y growin g LmS s in order to fundamentall y understand the
failure mechanism of LIBs and thereafter be able to improve their p rop erties fo r current and
future usage .
Recently , s ync hrotron X-ra y ima ging h as evo lved into a powerful characterization tool in
materials science [33] and has enabled batter y researchers to obtain unprecedented insights into the
underly ing degradation mechanisms of LIBs non-destructivel y. [34, 35] Among th ese, Eastwood et
al . have charac terized the 3D microstructure of elec trodeposited L mS s by sy nchrotron in -line
phase contrast X-ra y t omograph y , demonstrating that this technique is a suit able tool for
investigating LmSs without removing the s ample from its as - grown envir onment. [36] Moreover,
Harry et al . have charact erized the LmS and the LmS -punctured separator simult aneously , further
demonstrating that this technique is also a suitable and pow erful tool fo r studying the LmS and
the separator sim ultaneousl y in a short circuited cell. [3 7] Nevertheless, Harry et al . have employed
a self-made pol ystyrene-block-poly (eth ylene ox ide) copol y mer electrolyte as sepa rator and the y
conducted their characterization ex sit u . From a practical point of view though, non -destructive
investigations usin g widel y adopted commerc ial separators would allow for more general
conclusions on how and why separators fail and ISC develops. Obviousl y , fundamental research
on actual kinematic processes of I SC associated with commercial se parators is highl y desired.
Currently , the most widely us ed commercial separators in LI B s are made of pol yolef in materials
such as poly ethylene (PE) and pol ypr op y lene (PP) due to their proper pore structu re, good
mechanical strength and acceptable costs. [26, 2 8] I n thi s work, th e commerciall y availa ble trila ye r
Celgard ® 2325 s eparator [38, 3 9] is emplo y ed and its failure mechanism leading to an internal ISC ,
caused b y th e growin g LmS, is investigated via in -line phase contrast X-ra y tomo graph y . Th e
separa tor investi gated h ere is made of a PE layer of a low m elting point ( m.p. 135 ℃ ),
sandwiched between two PP lay ers of higher melting point ( m.p. 165 ℃ ). I n the case of ove r-
heating , the mi ddle P E layer will shut down the cell automatically b y bl ocking ionic pathwa y s
upon melting, while the two PP laye rs can still provide mechanical s trength to prevent phy sical
contact of the el ectrodes. This proposed shutdown function is meant to dispel safe t y concerns
related to ISCs in LIB s.
2.2 Break-down of the separator

- 42 -

Here, we present a non -d estructive ch aracterization of an I SC in a Li/Li s ymmetrical cell f o r the
first ti me. We demonstrate that: 1) most of the electrochemicall y deposited LmSs are
electroc hemicall y inert during e nsuing L i stripping, 2) electrochemical stripping and platin g
during discharge and ch arge are highl y inhomogeneous, 3) delamination int o three native la y ers
and partial meltin g of th e trila yer C elgard ® 2325 separator c an develop i n a real ISC event, 4 )
decomposition of the solid electrolyte interface (SE I ) coverin g L mS s and partial melting and re -
solidify ing of porous LmSs can also occur due to localized ex cessive Joule heating resulting from
the I SC. These unexpected findings, which are not accessible by conventional electrochemical
and morpholo gical characterizations, shed new li ght on the kinematic process of ISC and could
open up new d esign pr incipl es and opportunities to fundamentall y improve the safet y and
reliability of current- and next-generation LIBs.

Cell preparation, im aging setup and data acquisition
A proof-of-concept cell that is compatible with s y nchrotron X -ra y tomo graph y is designed a nd
fabricated, as shown in Figure 1 alon g with a schematic il lustration of the s y nchrotron setup at
the BAMline, BESSY II in Berlin, German y . [ 40-42] The electrochemical validation of the presentl y
designe d cell can be found in our previous report s. [43, 4 4] SEM im age s of the surface and c ross-
section of the emplo ye d t rila y er C elgard ® 2325 separator are shown in Figure 2. From Figure 2a,
the distinct sli t-pore structure that results f rom extrusion and unidirectional stretching during th e
dry p rocess manufactu re method is clearl y observed. [ 38] The three-layer structure of the PP/P E/PP
sandwich with a total thi ckness of ~25 µ m is shown in Figure 2b. [38, 39] Sy mmetrical cells made
of L i metal electrodes, the trilayered separator and the standard electrol yte (1 M LiPF 6 in
ethylene carbonate (EC) and ethyl methyl carbo nate (EMC) (EC/EMC=50/50 (v ol /v ol )) wer e
assembled. Two Li s y mmetrical c ells we re investigated here, on e bein g the pristine state (without
any cycling, h ereafter na med L i/ Li-1) and th e oth er one being short circuited after 13.4 h charge
(evidenced b y a sudden voltag e d rop shown in Fig ure 2 d, h ereafter named L i/ Li-2). Both cells
were mounted on the set -up for characterization without prior disassembl y. Ever y tom ography
was recorded b y a detect or system w ith 0.438 µm pix el siz e with 2200 projections covering 180°
rotation angle. The field of view (FoV) was (1.7 × 1.2) mm 2 (length × height). Detailed c y cling
parameters and the p rocedure of no rmalization, tomograph y reconstruction and 3D presentation
are given in Experimental Methods. [45, 46]
2.2 Break-down of the separator

- 43 -

Figure 1 . a) Photograph of the fabr icated proof- of-concept cell. b) Correspond ing schematic
representation of the cell cons isting of a polyam ide -im ide housing (brown ), tw o screw electrode s and
retaining screws each on top ( light grey), two sealing ri ngs (yellow), a porous separator (white)
sandwiched between two electrodes ( blue a nd green ). c) S chem atic repre sentat ion o f the experim ental
setup of the tom ography st ation at the BA M line at BESSY II, Helmholtz-Zentrum Berli n, Germany .
Figure 1 is ad apted from ref [44] , with permission from the American Ch em ical Society.
Morphological characterizat ion of Li/Li-1 cell and electr ochemical characterization of
Li/Li-2 cell
A cross-s ectional X- ra y t omographic slice of the uncycled Li/Li-1 cell is shown in Figure 2c, in
which the two Li electrodes and the sep arator ar e clearl y dis cernable du e to the li ght and dark
boundaries between the m arising from in -line p hase contrast. [36, 4 7] The interface b etween the Li
electrode and the s eparator of Li/Li-1 cell is flat and gapless. The Li/Li-2 cell was
ga lvanostatica ll y “c y cled” until an internal short circuit occurred. The volt age curve of the Li/Li -
2 cell measured at a current density of 0.3 mA c m -2 for 18 h of disch arge and then for 13.4 h of
charge is shown in Figure 2d (dis charge corresponds to Li strippi ng fro m the L i anode, cha rge
corre sponds to Li plating onto the L i anode. Th e anode and cathode are defined durin g the 1 st
discharge and are also used during the 1 st charge). As indicated b y two b lack arrows, there is a
steep voltage dip at the beginning of L i metal de position during both discha rge and charge. The
observed overpotential h ere is typica l of a nucleation mechanism, which sig nifies the initiation of
the growth of L i nuclei. [4 8, 49] During the charg e process, the sudden voltage drop pointed b y the
red arrow in Figure 2d indicates that ISC has occurred. T he c ross section of the short-circuited
Li/Li-2 c ell displa y s a si gnificantl y different mor pholog y compared to that of the Li/ Li -1 cell. A
2.2 Break-down of the separator

- 44 -

panoramic view of this situation is shown in Fig ure 3 (Fi gure 3b is a reconstructed r aw data
tomogra m allowing for a comparison with Figures 3c-h showing da ta after filterin g and
segmentation). Enlarged details are shown in Fi gure 4 showing bo th re constructed raw d ata as
well as segmented data. The complete three-dime nsional visuali zation of the internal structure of
the L i/Li-1 and L i/Li-2 c ells is shown in a supporting movie.

Figure 2 . a) Plane view and b) Cross-sectiona l view of the investigated trilayer Celgard ® 2325 se parato r
as obtained by scanning electron m icroscopy (SEM). c) Reconstructed i n- line X - ray phase contra st
tomog r aphy sl ice of the pristine Li/Li- 1 cell. d) Electrochem ical charact erization of the investigated Li/L i-
2 cell ( the black arrows point an abrupt voltage dip; the red arrow points a sudden internal short circuit ).
Scale bars in a), b), and c) are 1 µm, 5 µm, and 125 µm long, respectively. Fig ure 2b is adopte d from
Ref. [38] , with perm iss ion fro m the American C hemical Society.

Preliminary inspection of mor phological structures inside Li/Li -2 cell

Base d on ou r pr evious investigation of m orphological evolution of electrochemically
plated/stripped LmSs in Li/Li s y mmetrical cells, [4 3] kinetic I SC formation induced by growin g
LmSs in Li/Li-2 cell is elaborately studied. Fi gures 3 and 4 sho w various horizontal and cross-
sectional slices of the sh ort-circuited Li/Li-2 cell. The ini tially flat electrode/separator interfaces
have turned into dist inctively ru gge d int erfaces with a large cavit y wi thin the Li anode and
2.2 Break-down of the separator

- 45 -

numerous electrochemic ally d eposited LmSs on the surface of the Li cathode as shown in Figure
3b-e and Figure 4 a,f (middl e area of da rk and li ght y ellow LmS and deformed sepa rator o f gray) .
During the 1 st discharge, the total electron transfer (calculated from the external circuit) is
0.986 C (w it h the assumption that the total charge transfer of Li + inside the L i/ Li-2 sample durin g
the 1 st discharge occurs within the FoV, the same assumption for the 1 st charge), this results in
M cathode = 10 -5 mol of Li that is stripped f rom the anode (generating a large cavit y ) and
subsequently pl ated on to the surface of the L i cathode (developing numerous LmSs).
Simultaneousl y , the separator is stretched and pushed in to the cavity forme d withi n the anode by
the g rowing LmSs. These phenomena are clearl y shown in Figure 3 and Figure 4. Furthermore,
as can b e clearl y obs erved from Fi gure 3c,d (pink arrows) and Fi gure 4e (li ght blue lines),
cleavage of the separator into its three native la y e rs takes place, which con tributes to a decreased
chord length of the sep arator, see Fig u re 5b ( orange line). The significant de formation a nd
unexpected cleavage of the separator implies that enormous mechanical stress es are generated b y
the g rowin g L m Ss. According to the manuf acturer, [38, 3 9] the C elga rd ® separa tor ha s a hi gh
mechanical strength (in terms of tensile streng th alon g both the ma chine direction (MD ,
1900 kg/cm 2 ) and the transverse dir ection (TD , 135 kg/cm 2 )) and puncture strength
(300 g/cm 2 ) . [39] This observa ti on suggests that the current me chanical property of the separa tor
cannot withstand the enormous mechanical stress caused b y th e growing LmS s and thus has to be
further improved. During the 1 st charge, th e total charge tr ansfer after t he i mmediate short circuit
is 0.788 C, i.e. about M anode = 0.82×10 -5 mol of Li is stripped from the ca thode and then plated
onto the surface of the Li anode.
Notably, two observations regarding the previousl y formed L mSs (on the surface of the L i
cathode, b eneath th e separator) and n ewl y forme d LmSs (on th e surface of the Li anod e, above
the separa tor) are noted: 1) most of the previously formed LmS undergoes no elec trochemical
dissolution (evidence d by th e remaining L mSs, beneath the separator in the midd le of the FoV,
dark ye llow and light yellow shown in Fi gure 3b -e and Figure 4a,f). It is the original L i bulk
cathode to compensate for the depletion of L i ions in the electroly te used to conduc t charge
transfer (as evidenced by the nascentl y formed ca viti es within the L i cathode as shown by th e
white diamonds in Fi gure 3c -e and enlarged Fi gure 4d); 2) It is the p eripheral region of the Li
anode onto which newl y LmS is p referentially el ectrochemically deposite d (above the s eparator,
shown b y the white trian gles in Fi gure 3c -e and enlarged Figure 4d ), in correspondence to the
regions of nascently fo rmed cavities within the L i cathode (as marked by white diamonds in
Fig ure 3c- e and Figure 4d). The inhomogeneous Li stripping/plating during th e 1 st c ycle is clearl y
shown in Fi gure 3h wher e LmSs in the middle surrounded b y th e sep arator are ge nerated durin g
the 1 st discharg e by dep osition on to the surface of the L i cathode; in t he regions outside the
separa tor, the LmSs are ge nerated durin g the 1 st charge by d eposition on to the surface of the Li
anode. By further c alculating the volum e fraction of LmSs (deposited on the cathode durin g the
1 st discharg e and deposited on the anode during the 1 st charge) as a fun ction of radius r with
respec t to the centra l poi nt loca te d along z directi on as shown in Figure 5a, one obtains volum e
frac tions (blue straight li ne and dashed li ne) dist ribution as shown in Figure 5b. Obviousl y, the
2.2 Break-down of the separator

- 46 -

cathode LmSs are concentrated in the middle while the anode LmSs are distributed in the
peripheral region. From Figure 5b (or ange line ), one can also obse rve th e calculated chord length
of the separator (orange li ne): the separator located in the middle (small r) has delaminated into
layers with small average thickness as deduced from the low avera ge chord length. Th e reason
lies that the separator i n the middle underg oes cleavage (see above) and partial melting (see
below for more detailed explanation), while the separa tor in the peripheral re gion (large r) is
nearly int act. The ob served fluctuations between ~25 µm to ~30 µm result from its
morphologica l distortion.
2.2 Break-down of the separator

- 47 -

Figure 3 . a) 3D visualizat ion of the short circuited Li/Li- 2 cell . Letters b -e denote selected cross- sectional
planes; lett ers f- h selected horiz ontal plane s. These planes of Fig ure b)- h) are all defin ed in a ). b)
Reconstructe d raw data to which neither phase filter nor segm ent ation is applied, serves as com parison for
the seg m ented fi gures c)- h). In the enlarged green r ectang l e, black arrows point at a f iber that may be the
remnant of a Li dendrite or a m ol ten separator, white arrows point at an irr egular b oundary r esulting
probably from a fused lithium dendrite. In c) -f), l ight blue arrows point at the partially molten separato r;
purple arrows point at dang l ing se parator fibers ; pink a rrows point at the cleavag e of the separator. The
purple dotted li ne separates t he r esolidifi ed Li and the por ous LmSs. T he interesting region i n f igure f) i s
2.2 Break-down of the separator

- 48 -

enlarged as i ndicat ed by white dotted rectangle in inset figure. In g), the pink, black and green dot ted lines
denote t he conto ur of the anodic Li, t he separator or the Lm Ss, r espectively. In h), the electroch em ically
deposited Lm Ss cl early app ea rs heterog eneous. A l l scale bars are 250 µm long.

Comparison of the amount of LmS between morphological an d electrochem ical
characterizat ions

It has b een report ed that the electrochemicall y formed L mSs are still ch emicall y active and th at
they can easil y react with the elec trol y te to fo rm a solid electrol yte interf ace (SEI) cov erage. [50]
As a result, the SEI formed on the surfac e of L m Ss during the 1 st discharg e el ectrically insul ates
most of the L mSs, thereb y deactivating them electrochemically du rin g the 1 st charge. [51] Further
confirmation of the electrochemical inertness of newl y formed LmSs during ensuing
electroc hemical reaction is conducted b y comp aring the volume r atio of LmSs (plated on to the
cathode and LmSs plated onto the anode, both calculated from X-ra y tomography), with that of
the amount of quantit y ra tio of L mSs (calculated from the external el ectron transfer du ring the 1 st
discharge and c harge). The calculation of the volume ratio of the a mount of L mSs via the
reconstruc ted tom ography d ataset was done b y counti ng the number of v ox els belonging to the
LmSs deposited on to the cathode (V cathode ) and the L mS s deposited on to the anode (V anode ). The
reconstruc ted volum e rat io V cath ode : V anode = 10: 5.5 is obtained. The ratio calculated from the
electroc hemical characteriz ation with M cathode : M anode being 10:8. The ten denc y of the apparent
deviation between these two ratios is expected and the reason stems from the limited FoV (see
above). Nevertheless, t he comparison between these two ratios confirms that most of the
electroc hemicall y f ormed L mSs during 1 st discharge become electrochemically inactive during 1 st
charge , forming the we ll-know “ dead LmSs ” . [50] This is the first time that the amount of
electroc hemicall y deposi ted L mS s calculated from morphological characterizations is directly
corre lated with the amount of external electron tr ansfer m easur ed electrochemical ly . Concerning
the locations of preferential L i diss olution and LmS deposition during cy c ling, it has been
suggested that such locations are electrochemically mo re active [ 9] or poss ess a high local ionic
conductivity. [21] L i/ Li cells emplo y in g di fferently shap ed [52, 53] Li el ectrodes are planned to be
further investigated in the future.
2.2 Break-down of the separator

- 49 -

Figure 4 . a) Exemplary tom o gram of the short circui ted Li/Li- 2 cel l. b)- e) Zoom in into regions defined in
a) . b) T he origina l Li anode i s electrochem ic ally dissolved and partially filled by the separator and LmSs.
The partial ly molten separator form ing dangling fibers can also be seen (light blue and purple arrows) . c)
Exem plary sli ce shows electrochem ically plated L m Ss for med during t he first disch arge. Deposited
porous LmS and com pac t solid Lm S can be observed (separated by a purple dotted line). d) N ew ly formed
cavities (diamonds) and Lm Ss ( triangles) during the fi rst charge. e) Cleavag e and par tial melting of the
separator. Figure s 4a- e are reconstructed r aw data without phase filtering and segmentation. f) T hree-
dimensional visualiz ation of the segm ented data set. The scal e bar in a) is 150 µm and scale bars i n b)-e)
are 75 µm long.
Close-up inspection of m orphologi cal structures inside Li/Li -2 cell

Apart from these findings, a close -up inspection of Figure 3 and Figure 4 f urther demonstrates: 3),
there a re two di fferent phases of electrodeposited LmSs, one being poro us LmS and the other
being solid and compact LmS (both are shown in Figure 3b -e and Fi gure 4 c, f b y li ght ye llow and
dark y ellow, respectively, s eparated b y pu rple dotted line); 4 ) S ome of the originall y inta ct
separa tor has completel y disappeared (light blue arrows in Figure 3 c - e and enlarg ed Fi gure 4b).
2.2 Break-down of the separator

- 50 -

At other positi ons the se parator has molten int o filamentar y fibers that dangle into the solid and
the compact LmS phase (purple arrows in Figure 3c -e and enlarged Figure 4b ). This also
contributes to the decre ased average chord length of the separator in the middle region, see
Fig ure 5b (oran ge line). The currentl y obse rved porous LmS phase agrees well with previous
character izations [ 36] and the high lighted pore stru ctures are attributed b y Eastwood et al. to the
high-surface area lithium (HSA L ) compound that is composed of decomposed li thium salt
precipitates. [36] We assu me that these HSAL compound are porou s LmSs composed of SEI
formed during electrochemical plating of L i and numerous voids within. [50] If one compares the
porous LmS (light y ellow) with the solid and compact L mS (dark y e llow ) one can conclude that
the S EI and the voids in the porous L mS have b een completel y dissolve d or depleted. Richard
and Dahn have indeed identified an exothermic peak due to SEI decomposition at ~100 ℃ in
accelera ting rate calorimetry (ARC) studies of the thermal st abilit y of LI Bs. [54] Furthermore,
Maleki et al. have also confirmed the ex othermic SE I decomposition peak at ~100 ℃ by
differe ntial scanning calorimetr y (DSC). [55] The currentl y observed disappeara nce of S EI and
voids in porous LmS may b e attrib uted to hi gh temperature melting. Regarding the disappearance
and the meltdown of the separator, Cai et al. [56] and Malekei et al. [5 7] have independentl y
observed separator melti ng around the ISC location inside LIB s b y post-m ortem li ght microscop y.
They concluded that the thermal energ y induc ed b y an ISC is sufficient to locally increase the
temperature of a cell by 200 K , which can easil y melt the PE (m.p. 135 ℃ ), PP (m.p. 1 65 ℃ ) and
Li (m.p. 180 ℃ ). [58] Considering the decomposition of SEI and the melt ing of the separator
together with the electrochemica l characterization (a sudden volta ge drop in Figure 2d), it can be
safely concluded that an I SC occurs durin g charge. This is the first experimental demonstration of
the melting of the separator and the porous LmS phase in a non -destructive and three-
dimensional way .

Figure 5. a) 3D demonstrat ion of the calculated LmS volum e frac tion (VF) deposited on both the cathode
(here i nvis ible, beneath the gray se parator) and the anode (light yellow, above the separator). The r and z
directions are also shown. b) Calculated Lm S volum e fraction (VF) deposited on both t he Li cathode (blue
solid line) and Li anode ( blue dashed line) as a functi on of r ; t he calcula ted chord l ength of the separator
in t he z direct ion as a func tion of r is also shown by the orange line (the chord length of the separato r in
2.2 Break-down of the separator

- 51 -

the central r egion is small due to partial separator cleavag e and melt ing, is large i n the periphe ral region
because the s eparator is intact there ).

Discussion based on inspection of short-circuited Li/Li -2 cell

By checkin g intensivel y the separator br eak-do wn re gion, it is notable to find that one fibe r
remnant is located inside the delaminated separator, pointing to th e cathode LmS with
surrounding separator m olten away. This phenomenon is marked b y black arrows in the green
rectang le in Figure 3b. In the green rectangle, the irregular boundaries, point ed b y white arrows ,
is also shown. The diam eter o f this fiber is about 2 µm . We suspect th at t he irregular boundaries
probably result from the fus ing of some L i dendrites and that the remnant fiber is closel y r elated
to the I SC (due to the simil ar X-ra y absorption this fiber can not b e assigned to the separator or to
Li with certaint y ). Du ring charging most of the LmSs are electrochemicall y deposited
preferentially in the oute r region, but one dendrit e is el ectrochemically generated in this spe cific
area (lithium is probably supplied from the electrolyte). Du ring its growth it can easil y penetrate
through the pores of this part of the separator since the separator has alread y been significantly
stretched and its mechanical robustness has alread y be en weakened b y th e L mS grow ing during
discharge. [38] Once this dendrite eventuall y perforates the sep arator and bridges th e Li anode and
cathode the entire cell cu rrent (0.015 mA) concentrates onl y on ~2 µm dia meter piercin g-through
dendrite, which results in a significantly hi gh current densit y o f a bout 470 A cm -2 and
pronounced localized J oule heating around the ISC location. According to previous calculation, [5 9]
within less than ~0.1 s the induced Joule heat will lead to the decomposition of the SEI and the
melting of both the sep arator and porous L mS phase. Most of th e energ y is dissipated b y the
endothermic porous LmS and separa tor melti ng events and within ~0.4 s, the temperature drops
sharply . [59] The above-observe d results agree well with previous simulations. [9, 60, 61] Addition a lly ,
th e direct internal view of the electroche micall y short circuited Li/L i -2 cell suggests that the
shutdown mechanism provided by th e sandwich s tructure can hardl y provide an y prot ection from
a real ISC, i.e. the considerable force accompanied b y the grow ing LmSs and the I SC -induced
high Joule heating easily destro y the trila y er separator which is made b y l aminating PP and PE
layers together b y adhesion or welding. [3 9],[62] The decomposition of the SE I and the melti ng of
the porous LmS to a more compa ct resolidified LmS will release lots of the mechanical stress that
acted on the separator (th e conformal gaps betwee n the contours of Lm Ss , separator and Li anode,
as marked b y green dot line, black dot line and pink dot line in Figure 3g, act as tr acers of the
movement of the separator under stress r elief). In the investigated Li/Li -2 cell, most of the Joule
heat is diss ipated by the endothermic separator melting [ 63] after the penetrating L i dendrite has
been fused. [64] However, in real comme rcial LIBs co ntainin g normal cathode materials such as
LiCoO 2 , LiMn 2 O 4 and L iFePO 4 etc. and anode ma terials such as lithiated c arbon, an I SC -induced
Joule heating will initi ate the exothermic decomposition of some of the active materials includin g
the relea se of ox yge n and flammable gas. [62] Such reac tions can further increase the inner
temperature to above 200 ℃ and lead to the ex othermic decomposition or ignition of the
flammable electroly t e. [ 14, 65] Ultimatel y , without an effective heat dissipation or propag ation
2.2 Break-down of the separator

- 52 -

within LI Bs, the lithium -dendrite induced ISC can eventuall y r esult in ser ious incidents. Table 1
summary some particular parameters of commercial LIB components . [38, 66, 67]

Table 1 . Spec ific prope rty paramete rs of selec ted com mercial LIB com ponents.

Tensile strength
(kg cm -2 )

Puncture strength
(g cm -2 )

Melting
temperature ( ℃ )

Decomposition
temperature ( ℃ )

Celgard ® 2325
PP/PE/PP
separa tor

1900/135
(MD/TD)

300

135/165 (PE/PP)

-

SEI ( solid
electrolyte
interface)

-

-

-

90-120 b

Electroly te

-

-

-

130-230 a

Lithium

-

-

180

-

Carbon (fully
charge d)

-

-

-

100-130

LiCoO 2 (f ully
charge d)

-

-

-

178-250

a, elect rolyte decompositi on t empera ture varies among dif ferent electrolyte, this tem per ature range is
based on PC/EC / DMC (1/1 /3)+LiPF 6 (1M ) [66] [PC: C 4 H 6 O 3 ; EC: C 3 H 4 O 3 ; DMC: C 3 H 6 O 3 ]
b, SEI decom position temperatur e depends on the el ectroly te used, more inform at ion can be found
elsewhere. [14]

Conclusion

In summa r y, for the first time the underl y in g inter nal cell deformation and degradation caused b y
an I SC is experimentall y visualized by using s y nchrotron X -ra y pha se contrast tomogr aph y . From
the batter y engineer ’ s point of view, t h e present stud y su ggests t hre e po tential wa ys to further
enhance the safet y and reliability of c urrent- a nd/or nex t-generation LIBs: i) L mS s that are
electroc hemicall y plated on the surface o f Li electrode s shoul d be pe rfectly manipulated and/or
even completel y avoided . There have b een reports about the dire ct controll ed plating of LmSs on
the surface of L i electrodes [ 52, 68] and the direct engineering of Li electrodes. [69] ii) New kinds of
separa tors with considerably improved m echanical and pun cture strength as well as higher
thermal stabil ity (m.p. above 165 ℃ ) are highly desired. Promising candidates are ceramic
separa tors which comb ine the characteristics of flexibl e pol y m ers and excellent thermal
stabilit y . [7 0, 71] iii) Novel electrol y tes with enhanced thermal stability and non -flammabilit y are
desired. Although adding fire retardant (FR) addit ives to lower the fla mm ability of the liquid
electrolytes can, to som e ex tent, reduce the liquid organic electrol y te flammability, [14] y et th e
rece ntl y emerging solid state electrol y te s, which function as both electrolyte and separator, ma y
be the most promising o ption for current- and/or next -ge neration LI Bs wit h enhanced safet y and
reliability. [6, 72- 74] From the battery tester ’ s point of view, the present inve stigation ma y op en up
new opportunities for future standard battery testing procedures. Con ventional mechanical
2.2 Break-down of the separator

- 53 -

evaluation techniques h ave been p redominantl y criticiz ed by th e fact that the y onl y show how th e
cells behave under an abuse condition instea d of trul y replicating the conditions of a filed
failure. [16] The firstl y successful dir ect visuali zation of an I SC conducted by s y n chrotron X -ra y
imaging ma y help updati ng existing testing standards and creati n g ne w ones for next -generation
LI B technolog y .

EXPERIM ENTAL MET HODS
Lithium and Cel gard ® 2325 separator w ere purchased from MT I Corp. U SA. The ele ctrol y te is
1M LiPF 6 in a volume-ratio mix ture (1:1) of ethy l ene carbonate (EC) and e thyl meth y l carbonate
(EMC) and was purchased from S igma Aldrich. The housing of the proo f -of-concept be amline
battery wa s made of polyamide-imide (Torlon) provided b y D rake Plastics Europe.
The li thium electrodes i n both L i/ Li cells (2.5 mm diameter) were pun ched out from a 1-m m
thick lithium plate. Both cells were assembled i n an argon -filled glovebox with humidity and
ox y ge n l evels below 0. 1 ppm. The trila y er Cel gard ® 2325 s eparator (3. 5 mm in diameter and
~25 µm thick) was placed between the Li/ Li electrodes. All c ells were as sembled manuall y
without exerting force. Finally , both cells were filled with the standard liquid electroly t e and
were sealed off be fore taking them out of the glovebox.
Galvanostatic charge and discharge of th e Li/L i-2 battery was carried out using an I viumS tat
from Ivium Technologies, Netherlands. The cells were discharged for 18 h and charged for 13.4 h
at a current density of 0.3 mA·cm -2 .
Synchr otron X-ra y tomo graph y was carried out at the BAMline at BESSY II of the H elmholtz -
Centre Berlin, G ermany. Th e s y nchrotron beam was monochromatized to 20 keV usin g a double
multilayer monochromat or with an energ y resolut ion of about 1.5 %. The detector s ystem
comprised a 60-µm thi ck CdW O 4 scintillator, a mic roscopic optic and a p co4000 camera
equipped with a 4008×2672 pixels CCD chip that is kept out of the direct bea m b y using a
mirror. For tomograph y measurements of both c ells, 2200 projections wi th each 2.5 s exposure
time within a 180° battery rotation were recorded.
The raw tomography dat a was filtered, norm alized and reconstruc ted usin g code programmed in
IDL 8.2. Three-dimensio nal segmentations of the separators were made using a grid of manuall y
marked points that were fitted with a biharmonic equation using MA TLAB. For the se gmentation
of L mSs, the statisti cal region mer ger tool implemented in Fiji [75] was used followed by manual
removal of the bulk lithium background .
ACKNOWLEDGEMENTS
We thank Dr. Heinrich Riesemeier, the beamline scientist at BESSY II, for his valuable
assistance and en gineer Norbert B eck for fabr icating th e be amline batter y . We also thank
Christiane Förster for p reparing the SEM sample and Manz oni Anna for conductin g the SEM
measurement. This work is sponsored b y th e Helmholt z Association and the C hina Scholarship
Council.
2.2 Break-down of the separator

- 54 -

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2.3 Gas developme nt

Reprinted w ith per m ission from DOI : 10.1021/ac sam i.6b00708 . Copy right (2016) American Chemical
Society.

Three Dime nsional Visuali zation of Gas Evolution and Channel F or mation inside a
Lithium-ion Battery

Fu Sun* ,†,ǂ , Henning Markötter †,ǂ , Ingo Manke ǂ , Andre Hilger ǂ , Nikolay Kardjilov ǂ and
John Banhart †,ǂ

† I nsti tute of Materials Science and Technology
Technische Universität Berlin
10623 Berlin, Germany
ǂ Helmholtz Cent re Berlin for Materials and Energy
Hahn-Meitner-Platz 1
14109 Berlin, Germany
*Corresponding Author: fu.sun@helmholtz -berlin.de

Abstract
Gas ge neration within lith ium ion batteries (LIBs) gives ris e to safet y concerns that question their
applicability. By emplo ying s ynchrotron X - ray imaging, the gas and channel evolution occurring
in an operating LIB ha ve been directly visualiz ed in their inherent 3D state as a function of
discharge and charge. U sing the spatial 3D distribution of gas bubbles and channels, the active
particles that dictate the performance o f a functional LIB were identified and visualized in 3D.
Delithiation and lithi ation are int erpreted as the proc ess of activating particles continuousl y in a
step- by -step wa y. Th e p resent wo rk not only d emonstrates the generation and evolution of g as
within LIB in 3D, but also reveals the dist ribution of active particles for the first time. These
fundamentally findin gs presented here shed light on a range of processes that could not
previously be charac teriz ed in 3D and ca n provide practical guidance for the desi gn of next -
ge neration LIBs with improved safety.
2.3 Gas development

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Abstract Graph ic
K EY WORDS
ga s evolution; channel formation; lithium ion battery ; s ync hrotron X -ra y imaging; tomography ;
radiography; silicon anode
INTRODUCTION
Rechargeable lithium -ion batterie s (LI Bs) have penetra ted p rofoundl y in to products such as
portable electronic devices, electric vehicles or other lar ge-sized power sources. 1-5 However,
safety concerns still limit the full practica l utilization of these batteries . 6 Especiall y gas evolution
is a formidable technological and fundamental challenge 7 since gas ge nerated in herme ticall y
sealed batteries can le ad to detrimental ef fects: on the one h and, gas generation durin g stor age
results in a dimini shed shelf li fe and eventuall y a markedl y reduced c ycle lifetime . 8 On the other
hand, gas evolution duri ng c ycling leads to electrol yte displacement , 9 which causes a decrease o f
Li -ion diffusion and/or Li-ion conduction in the electrolyte and ultimatel y contributes to a
tremendous increase of b atter y resistance. 10 I n both cases, the int ernal pres sure build -up may also
induce batter y bul ging, mechanical stress inside the electrodes and ev en s evere gas leakage , 11 all
of which are detriment al to longevit y and reliability o f en erg y sto rage batter y s ystems. Therefore,
2.3 Gas development

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understanding the gas e volution mechanisms are of grea t importance from a practical point o f
view.
To date, man y experimental investi gations have been carried out in ex tensive detail and, as a
result, our knowled ge of the gas evolution mechanism has substantiall y dee pened. For instance, it
is already known that the gas evolution, which is intim ately related with soli d electrol yte
interphase (SEI) formation in the electroc hemical batter y, 12 originate s from reductive and
oxidative electroly t e decomposition reactions due to the fact that the electroche mical potential of
both electrode materials is far beyond the thermodynamic stabilit y window of the commonly used
organic electrol y tes. 13 Actually, si gnificant evolution of gaseous p roducts such as CO 2 , CO, O 2 ,
H 2 , C H 4 , C 2 H 4 , C 2 H 6 , C 3 H 6 and C 3 H 8 from the decomposition of carbonate solvents and li thium
during b atter y op eration has been detected by v arious characterizati on techniques . 14 - 21 I n
addition, utilizing isotope anal ysis, Onuk et al. ha ve unambiguousl y identified the ori gin of gase s
evolving in LIBs. 22 Furt hermore, b y adopting in sit u transmission electron micr oscop y (TEM) 23
and neutron radiograph y (NR), 9, 11 the generation of gaseous bubbles channels formed by the gas
have recentl y been visua lized. These studi es have enhanced our und erstanding of gas evolution
but, unfortunately , the applied radiographic im agin g techniqu es yield two-dimensional (2D)
information and do not help us in further comprehending gas evolution kinetics in its inherently
three-dimensiona l (3D) s tate and in quantitativel y anal y zing its complex e volution. Moreover, no
study of the spatial gas distribution in relation to the el ectrochemically active particle
population 24 in electrodes has been conducted so far, even though it is th e a ctive pa rticles that
directly determine battery performance and cycle life . 25
Due to the high fluxes generated b y s ynchrotron sources, s y nch rotron X-ra y imaging has
evolved into a powerful chara ct erization tool in the field of materials sc ience. 26 - 36 Especially
sync hrotron X -ray tomograph y has enabled rese archers to obtain unprecedented insights into
LI Bs from the level of i ndividual particles to the scale of entire el ectrodes . 37 - 40 Most rece ntly,
Ebner et al. , have directly observed and quantifie d electrochemical and mechanical degradation
in a SnO anode. 41
Using a non -destructive 3D X -ra y imaginin g t echnique, herein w e investigate in 3Ds the gas
evolution scenario in a LI B b ased on a silicon ( Si) electrode sin ce Si is considered on e of th e
most promising anode materials for next -generation power sources. 42 - 43 This is the first study that
also visualizes the distribution of electroc hemicall y active electrode particles with respect to the
spatial ga s development and interprets delithi ation and lithiation as a proce ss of activatin g
particles in a step - by -step wa y. This work sets a refined ex ample for the 3D g as evolu tion
investigation and the spatially evolved active particle distribution, providi ng in -depth knowledge
for the electrode en gineers and numer ical simulation experts.

EXPERIM ENTAL SECTION
Mater ials. A ctive sil icon particles were received from Elkem AS, Norway . Conductive carbon
black, Poly vin y lidene fluoride (PVD F) binder, Celg ard separator and li thium were purchased
from MTI Cor. USA. N -meth yl p yrrolidone solvent (NMP) and 1M LiPF 6 in a volume-ratio (1:1)
2.3 Gas development

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mixture of eth y len e carbonate (EC) and dimeth yl car bonate (DMC) were purchased from Sigma
Aldrich. The housing o f the proof -of-concept bea mline ba ttery is made of polya mide -imide
(Torlon), from McMaster-Carr.
Mater ials characterization . Scanning electron microscop y (SEM) images were taken using a
Ze iss ultraplus microscope. The particle siz e distribution of Si particles was measured b y laser
diffrac tion using a Malvern Mastersizer 2000 analy z er.
Battery pr eparation. The electrode is made of electrode slurries with weight ratios of
Si:carbon black:binder of 80:10:10 in NMP. The slurry was cast directl y onto the head of the
bottom screw shown in Figure 1. To remove the NMP, the cast slurry was dried in an oven at 60
℃ overni ght. Before and afte r casting, the scre w was weighed to determine the mass of the
electrode materials. The proof-of-concept batter y (shown in Fi gure 1a) des igned for beamline use
was assembled in an argon filled g lovebox wit h humidity and ox ygen levels below 0.1 ppm.
Metallic lithium was pla ced onto th e head o f the top screw in Fig ure 1 a ac ting as both a counter
and reference electrode. A polymer sep arator soaked with the electrol y te was plac ed b etween the
lithium electrode and the Si electrode. Curr ent leads are connected to a potentiostat for
electroc hemical tests.
Electroc hemical measurem ents . C y clic volt ametry (CV) and galvanostatic charge/discharge
of the batter y were c arried out using a n IviumStat voltameter. A fresh batter y after assembly was
measured for CV curves from 0 V to 2.5 V at a scan rate of 1 mV/s. During ga lvanostatic c y clin g
at the beamline, the assembled batter y was measured at 1.75 A g -1 based on the mass o f a ctive
silicon. Different discharge and charge current rates will be used to further investigate the
corre lation between the current rate and the gas releasing and channel formation behavior.
Settings of tom ography and radiography . S y n chrotron X -ra y tomo graphy and radiograph y
were recorded at the BAMline at the storage rin g BESSY I I of the Hel mholtz -Zentrum Berlin,
Germany. The s ynchrotron beam w as monochromatiz ed to 20 KeV using a double multilayer
monochromator with an energy resolution of about 1.5 %. The detector s y stem comprised a 60 -
µm thick CdWO 4 scintillator, a mi croscopic opti c and a pco4000 ca mera with a 4008×2672
pixel 2 CCD chip that is kept out of the direct b eam b y usin g a mirror. For tomograph y , 2200
projections during a 180° batter y rotation each w ith 4 s exposure time were recorded b efore/after
discharge/charg e. The b atter y was r adiographically ima ged during dischar ge and charge with 4 s
exposure time. 4 s exposur e time was required because of the relative low photon flux provided
by the BAMline comp ared to other ima ging stations. It has to be n oted that in order to
character ize the morphological change of the whole Si based electrode (3.2 mm × 0.4 mm, len gth
× width), we have used an optic and camer a combination with a corre sponding field of view of
3.3 mm × 2.2 mm. And the resultant vox el resolution is ~0.876 µm. However, -10 er camera
system with hig her resolution of ~0.438 µm will be used in future investigations.
Data processing. Filtering, binariz ation and segmentation was performed in I ma geJ and Avizo
Fire. Data visualiz ation was done using VGStudio MAX and Avizo Fire. Analy sis of the particles
is conducted by ImageJ. The procedure is presented in Supp orting Information.

RESULTS AND DISCUSSION
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An electrochemical cell that ca n full y represent a workin g batter y and at the same time is
compatible with s y n chrotron X-ra y im aging is crucial for the characteriz ation of gas evolving
during operation. For this purpose, we developed an electrochemical cell that is X -ra y transparent
so that imaging can be c onducted in operando to directl y stud y gas evolu tion during dis/charge
processes, a llowing us to c orrelate the gas evolution with elec trochemical activity. F ig u re 1
displays a photograph a nd the corresponding schematic illustration of the batter y as well as a
schematic repre sentation of the synchrotron X-ra y tomography setup. 44

Figure 1. T he custom ized el ectrochem ical batter y and the employ ed X - ray to mography se tup: a)
Photograph of the battery investigated. b) Corr esponding sc hematic representation of the battery cell
consisting of a poly am ide-im i de housing (brown), two sc rew electrodes and a retai ning screws each on top
(light g rey), two sealing ri ng s (yellow), a lithium metal electr ode (blu e), a porous s eparator (wh i te) and th e
Si -based anode (green). c) Schem atic representa tion of the experim ental setup o f the tom ography station at
the BAM line a t BESSY I I, Helmholtz-Z ent rum B erlin, Germ any

After assembl y of the battery, c y clic voltametr y (CV) wa s performed to verif y the reduction
and ox idation characteristics of silicon. Figure 2a shows C V curves of the batter y s canned at a
rate of 1 mVs -1 in the potential window of 0-2.5 V. The c athodic current increase fr om 1.4 V to
0.7 V in the CV curve s is re lated to small -scale reduction of electrol y te and/or surfa ce
contaminations such as trace w ater. 45 The small hump around 0.5 V is attributed to a strong
reductive decomposition of the electrolyte associated to the re lease of large amounts of gases, 46
while the peak above 0 V is related to Li alloying with Si. During delithiation, the anodic peak
observed slightly abov e 0.5 V is for de -allo y in g of Li-Si phase and th e broad anodic peak at
1.0 V is related to the oxidati on of electrolyte and/or compounds previousl y reduced durin g the
2.3 Gas development

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cathodic process. The results are in good agreement with previously r eported Si/carbon
composites, 43, 47 thus confirming the electrochemical reactions inside the batter y .

Figure 2. Electrochem ical characterization of the investigated battery: a) CV curves of the first thre e
cycles of the battery at a scan rate of 1 mVs -1 . b) The first discharge, the first charge and t he secon d
discharge v oltage- time profiles of the ba ttery at 1.7 5 Ag -1 rat e.

After this, we conducted X -ray radiograph y while simultaneously galvanostatically dischar ging
or charging the c ell at a r ate of 1.75Ag -1 as shown in Figure 2b. The discharge and ch arge curves
are also in good agreement with previous results for Si elec trodes. 43 I n addit ion, another batter y
was galvanostatically discharged/charged at a small curre nt of 0.01Ag -1 for more than 20 h as
shown in Supporting I nformation ( SI ) Figure S 1, showing t y pical characteristic features of Si.
Thus, the battery design use d for tom ographic measurements exhibits an authentic
electroc hemical behavior.
High-resolution synchrotron X-r a y tomographic imaging grants us the ability to probe the
internal structure and th e distribution of elements in the electrode nonde structively. 40 First, the
pristine battery without being exposed to an y discharge or cha rge is characterized. The t y pical
procedure of tom ographic data acquisition and pr ocessing is demonstrat ed in Figur e 3. Figure 3a
shows a pr ojection imag e of the Li electrode/s eparator/Si elec trode assembly. While the L i
electrode and the porous plastic separator appear nearl y invisible due to their low X -ra y
absorption coefficients, the Si electrode is clearly visible. As the batter y rotates during
tomogra ph y , a series of projec tions are recorded, which are late r reconstructed into a three
dimensional volume. Figure 3b shows a horizontal slice image throu gh the reconstruc ted volume,
which displays that the Si particles are well mix ed with the carbon black. I n ord er to further
analy ze the Si p articles, the gra y scale slices were segmented into bi nary images, y ielding
Fig ure 3c, in which the bright regions correspon d to higher -absorbing material, the Si particles,
whereas the black regions are assigned to less - absorbing mat erial such as c arbon black and
polymer binder. Segmentation is performed via a combined app roach of filtering and
thresholding 48 and produces Figure 3d after labeling. In order to validate the X -ray tomograph y
2.3 Gas development

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results qualitatively and quantitatively we compare scannin g electron mic roscopy im ages (SEM)
(Figure 3e) with gray scale sli ce image s as well a s the particle size dist ribution (PSD) measured
using las er diffraction w ith the PS D obtained from the tomographic d ata (Fig ure 3 f). Th e good
agree ment between the experimentally and numericall y obtained results indicates that the
segmentation applied accurately captures the electrode material propertie s. Finall y , a full -view
3D representation of the segmented particles is shown in Figure 3 g and in the Supporting Movie
1 (SM 1) (More de tails c an be found in the SI.)

Figure 3. a) X -ray projection imag e of t he assembled Li electrode/separator/Si electrode stack wit hin th e
battery. T he scale bar is 200 µm long. b) Example of a reconstruc ted gray scale slice i m age. c) Sam e slice
after binariz ation. d) Slice after se paratio n and labeling . e) Scanning elect ron m icrographs (SEM) of the
m ixed Si and carbon. Scale bars in b) – e) are 50 µm long. f) Particle size distribution (PSD) obtaine d
from the labeled particles (red) compared to value s m easured by laser dif fraction (black ). g) 3D
visualizatio n of all Si partic les with labeled r egions R0, R1 and R2. The scale bar is 100 µm long .

Gas evolution is directl y observed in 3D after each discharge or charge step and the
reconstruc ted gas evolution is straightforwardl y correlated to the ele ctrochemical process. Du ring
data processing, we can easil y separate the gas e volved of the channel from the soli d materials
due to the near zero abso rption of gas. Aft er the first discharge process, we c aptured the int ernal
morphologica l change by tom ographic ima ging. I n Figure S2, different 3D views of the
composite of Si electrode and gas or channel are display ed. Figure 4a shows a cropped region
after the first discharge, labeled as R0 in Figu re 3g, in which Si particles are shown in gra y and
2.3 Gas development

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ga s or channel in green. We then charge the battery and obtain the first charged state tom ogram.
The corresponding 3D rendering s are displ a y ed in Figure S2b and SM 2b and the corresponding
cropped region is shown i n Figure 4b. In the f inal step, we discharge the batter y again and
acquire the s econd dis c harged state tomo gram shown in Fig ure S2c, SM 2c and Figure 4c. In
addition, during the discharge and charge process, we simultaneously capture the gas movement
inside the batter y in 2D by in sit u X -ray radiograph y. The results are s hown in Figure 5 and
SM 3.

Figure 4. 3D visualization of gas evolution in a c utout at position R0 in the bat tery after each di scharg e or
charge process (see Figure 3g ). a) After t he firs t discharge. b) After the f irst charg e. c) Afte r the secon d
discharge. In all pictures , from lef t to right is fr om the periphery to the central part of the electrode. T he
left column shows t he Si partic les (gray) and the gas or channel (green) in a composit e image. In the right
2.3 Gas development

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column, one hal f of the particles i s rendered transp aren t to show the g as part. The scale bar is 100 µm long
in all the pic tures.

Figure 5. A series of projecti ons showing the mov ement of gas as imag ed by X -ray radiography. a) Fir st
discharge step. b) Firs t charg e step. c) Second discharge st ep. The scale ba r is 100 µm long in all the
2.3 Gas development

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pictures. Ellipses m ar k ar eas of gas generation, arrows point at areas where gas bubbles h ave moved t o.
Colum ns a) and c) are in the m iddl e, colum n b) near th e edge of the electrode. See m ore in SI and SM 3.

During cyc ling, lithi um ions will pass through the separator, reduce or oxidiz e the electroly te
on the surface of the electrode, form a solid elec trol y te interface (SEI) la y er containin g a lar ge
amount of generated gas, and finall y lithiate or delithiate the electrode 49 . In Figu re S2, Figure 4,
and Fi gure 5 as well as in the in op erando 2D movie SM 3, it is clearl y shown that a pronounced
ga s ge neration and accumulation takes place wit hin the elect rode du ring the first discharge step,
which is consistent with previous experimental results . 15, 50 However, some phenomena were
observed that are in contradiction to what h ad be en assumed previousl y . One is the non -uni form
distribution of gas generation s it es. During electrolyte reduction/ox idization b y active li thium, on
the one hand, the reducti ve/oxidative decompos ed products form the SE I la yer with organic
carbona te outer surface and an inne r inor ganic s alt 45 full y coverin g the e lectrode. On the other
hand, notable amounts of gas ar e released. Th e po lymer S EI l a y er remains on the surfaces of th e
electrode s and to some extent shapes into the wall of the channels thr ough which the newly
ge nerated ga s readil y migrates out and finall y me rges with the accumulated gas, which eventually
displaces the electrol y te and then takes th e form of a skeleton structure within the electrode as
clearly shown in Figure S2a and Figure 4a. Actually, because of the higher flow of
electroc hemicall y reactive L i ions it is plausible to find larger ga s agglomerations close to the
separa tor and wide ch annels grow into the electr ode from the separator, which is consistent with
the literature. 18 However, there are also a large number of gas bubbles or channels that are
ge nerated in th e periph ery of th e electrode. The reason for this ma y stem from the hi gh Li ion
concentration in the electrolyte, which su rrounds the electrode. 51 Another reason m a y stem from
the S i electrode morphol og y that can modi f y th e distribution of the electric fie ld in the electrode.
As shown in Figure 3a and g, the su rface o f the S i electrode is not flat, with the circumferential
edge slightl y hi gher than the center. Actuall y , it has been simulated th at the electric filed has its
highest value in the electrol y te a t th e tip of lit hium dendrites. 52 In our case, the non -flat Si
electrode morpholo gy ma y result in an inhomogeneous electric field, with higher intensit y in the
circumferential edge than the center area. As a r esult, the lithium ions, driven b y the differen t
electric field, will preferentiall y react with the el ectrode materials in the peripher y region. In the
region near the current collector, the development of gas bubbles or channels is also observed,
most li kely b ecause the Li ions are transported v ia the contacted particles 53 or by interstit ials or
vacancies in the SEI. 54 I n a series o f radiographi c images durin g the first discharge p rocess, s ee
Fig ure 5a and SM 3a, we can clearl y observe the gas movement. Fig ure 5b and SM 3b along with
Fig ure 5c and SM 3c show the first charge and the second disch arge p rocess, respectivel y . Th e y
also clea rl y reveal unexpected d y namics of ga s evolution and channe l development during the
electroc hemical p rocess. Note that r adiographic i mages might give misleading info rmatio n on the
direction of gas movement because of the ignored third dim ension. I n contrast, the full view of
ga s evolution in 3D provided b y tomograph y allows us access the three dimensional cha nnel
structures.
2.3 Gas development

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Another unexpected phenomenon is the continuous p roduction of gas in successive cycles.
Contrary to the widespread belief that li thium i ons diffuse uniforml y from the separator to the
curre nt collector and for m a lithiation front extending throu gh the entire electrode , 55 the bubble s
evolve successivel y from the circumferential surfa ce into the bulk of the inner electrode and form
a branched s y stem of gas bubbles and channels. It follows from the observed gas evolution after
delithiation or li thiation – as clearly shown in Figure S2b -c and Figure 4b - c – that lithium ions
preferentially enter the el ectrode in an inwardl y radial passage. The direct full -view visuali zation
of the three-dimensional gas evolution pres ented here suggests diffusion pathwa ys of li thium ions
inside the anode material that ar e non-uniforml y d istributed .
In fact, r ecently mor e and more attentions are being devoted to the investigation of the
electroc hemicall y active particles in r elation to the overall electrode curre nt and/or the d egree of
curre nt homogeneit y. 24 - 25, 56 For example, a discrepa nc y is observed between electrochemica l
measurements that r epresent the overall state of the cell and sp ectroscopic data that r epresent a
local state. 57-58 A rece nt experiment b y Delmas et al. showed the coexistence of fully lithiated a nd
fully delithiated individual LiFePO 4 cathode particles b y X-ra y di ffraction after full lithiation . 59
Sugar et al. and Brunetti et al. further con cluded that at an y given time during c y cling, onl y a
small fraction of the total ensemble of p articles is activel y charge d or di scharge d. 57 , 60 L i et al.
conducted an in-d epth characterization and observed that onl y 5% to 8% of particles are actively
intercalating du ring lithiation, while during delithiation the a ctive population ranges from 8 to
32%. 24 The y further confirm that the current is heterogeneousl y dist ributed in the electrode, that
is to sa y , onl y a sm all number of active p articles carries most of the current regardless of the tot al
electrode current. Taking into account that lithi ation of Si will not occur until the native sil icon
oxide is at least partially reduced b y Li ions 49 and the gas evolution behavior presented h ere, w e
assume that diffusion of Li ions takes place as follows: d uring lithiation, Li ions that pass though
the separator and/or the native Li io ns in the nati ve lithi um salt electrol yte will first diffus e to the
potentially active electrode pa rticles, reduce the electrol y te and/o r lithium salts with the obtain of
electrons from the ele ctrode and successively reduce the native silicon o xide, thus f orming the
SEI la y er with inorganic inner species and or ganics outer species, releasing ga s, and finall y
lithiate the active particles of the electrode, leavi ng non - active local re gions intact. 10 Vice versa
during d elithiation. Using this picture, we can c o nvincingly co rrelate gas evolution and channel
formation with the active or potentially active p articles during electroch emical delithiation and
lithiation. More specifically, where gas is generated, where the active particle s are , where the
electroc hemical current is concentra ted. This finding is fundamental. F or the first time, the
spatially distributed elect rochemica ll y active particles in an op erated LIB are dir ectl y visuali zed
via the gas evolution.
For a more conveni ent visuali zation of the dist ribut ed population of active particles and/or
potentially active particles , we separated them from the remaining particles. The results a re
shown in Fig ure 6b and Table 1. After each del ithiation and lithiation step, the active p article
population incre ases concurrently with the electrochemical process. From above anal y sis, we
interpret delithiation and lithiation as a process of pro gressivel y activating potentiall y active
particles.
2.3 Gas development

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Table 1. Relative gas or ch annel volume fraction and number of activ e particles (both in % ) as a function
of the cycle state in region 1 and r egion 2 (see Fig ure 3g ).
Fraction State
Name

First discharge

First charge

Second discharge

Region 1
Gas/Channel

20

31

39

Region 2
Gas/Channel

14

20

39

Region 1 Active
Particles

29

46

64

Region 2 Active
Particles

18

29

57

2.3 Gas development

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Figure 6. 3D visualization of r egion1 as defined in Fig. 3g. a) gas or channel. b) active particles. From top
to bottom : after first discha rge, after first charge, after second discharge. The scale bar is 10 0 µm l ong i n
all the pictu res.

To further quantif y the relationship between both gas or channel volu me chan ges and the
population of a ctive p articles with electrochemical reactions, w e calculated the volume f raction
of gas or channel volume and the fra ction of active particle s after eac h charge and discharge
process. Two independ ent region s – labeled R1 and R2 in Figur 3g – are extracted for
quantitative anal y sis. Th e evolution of the g as or channel volume in region R 1 is presented in
Tab le 1 and in Figure 6a. After th e first discharge step, the fraction of gas or channel volume is
only 20% compared to 3 1% after the first charge step. After the s econd discharge step, the gas o r
channel volume fraction is  39%. Along with the se changes of gas or channel volu me fraction,
the increase of the fraction of active particles is evident in Figur e 6b and Table 1, from 29% in
the first discharged state to 46% in the first charged state and 6 4% in the second discharged state.
Both trends impl y th e aforementioned delithi ation/lit hiation activation mechanism that progresses
in each c yc le. The anal y sis of R2 y ields a similar result, increasing confidence in the applicability
of the mecha nism to the entire electrode.
The ex periment conducted here connects the ga s bubble or channel evolution with the
evolution of effectively or potentially active particles in a functional LIB. This ope ns new
perspec tives in optimizing the overall performance of a batte r y . From the viewpoint of electrode
eng ineering, it is important to find opt imal battery operating conditions or to develop new
materials with reduced g as formation. There a re a lread y some reports of adding electrol yte
additives 8, 10 and using new electrol ytes 13 to contr ol gas generation and in crease Li ion transport
rates to enha nce the LIBs’s reversible performance. It is also crucial to optimize the electrode
architec ture to significantl y increase the popula tion of active particles to sustain the overall
ga lvanostatic current and at the same time to improve the homogeneity of th e distribution of
active particles to prevent the extent of shocks and fractures induced by hi gh local cu rrents . 24
From the perspective of numerical simulation, i t is important to identif y th e im pact of SEI
formation with the associated gas release withi n the electrode on the tran sport properties of Li
ions and c yc ling performance of LIBs during discharge and charge 12 to further sp eed up the
optimization of electrode structures. Altogether, the three dimensional microscale investi gations
of gas in the present stud y contribute to the und erstanding of the complex discharge and ch arge
processes.

CONCLUSIONS
Novel insi ghts int o the s patial distribution and kinetics of gas evolution are unraveled b y 3D
sync hrotron X -ray imagin g for the first time. The kno wledge obtained here enriches our
understanding of the elec trochemical activit y in a rea l functional LI B. The ga s evolution and
channel formation occurring in a functional LIB have been dire ctly visualiz ed in a full-3D v iew
after discharge and charge. Gas channel s successivel y evolve from the surface of the anode
2.3 Gas development

- 70 -

material into the bulk of the electrode and form a branched s ystem of ga s bubbles and channels.
The observed gas evoluti on point s at an activation of individual particles progressing from c y cle
to c y cle. These novel f indings presented here also highlight the unique s y n chrotron X - ray
character ization tool for revealin g underl ying mecha nisms of LIBs and shed light on a whole
range o f processes that could not previously b e characterized in 3D. Indeed, our r es ults and the
value of information we obtained should speed up the optimization of the architecture and/or
material of the current LIBs to design more stable electrodes to further enhance their performance
for next-ge neration demand.

ACKNOWLEDGEMENTS
The assistance of the bea mline scientist of the BAMline, Dr. Heinrich Riesemeier, is gratefull y
acknowledged. We thank Norbert Be ck for fabri cating the beamline batter y and Elkem AS for
providing us with Si particles. This work is sponsored b y the Helmholtz -Zentrum Berlin and
China Scholarship Council.
ASSOCIATED CONTENT
Supporting Inf orm ation Available: including Methods, Data Processi ng, Figure S1 -S3 and
captions for the Supporti ng Movies. This material is available free of charge via the Inter net at
http://pubs.acs.org .

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Bartelt, N. C. High- Resolution C hemical Analy sis on C ycled Lifepo4 Battery Electrode s Using
Energy-Filtered Transmission Electron Micr oscopy . J. Power Sources 2014, 246 , 512-521.

2.3 Gas development

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2.3.1 Suppor ting Inf ormation

Three Dimensional Visualization of Gas Evolu tion and Channel Formation ins ide a
Lithium-ion Battery

Fu Sun* ,†,ǂ , Henning Markötter †,ǂ , Ingo Manke ǂ , Andre Hilger ǂ , Nikolay Kardjilov ǂ and
John Banhart †,ǂ
† I nsti tute of Materials Science a nd Technolog y
Technische Universität Berlin
10623 Berlin, Germany
ǂ Helmholtz Cent re Berlin for Materials and Energy
Hahn-Meitner-Platz 1
14109 Berlin, Germany
*Corresponding Author: fu.sun@helmholtz -berlin.de

This section includes:
Methods
Data Processing
Figu re S 1, Figure S2 and Figure S3
Captions for t he Supportin g Movies
Methods
Battery Preparat ion:
A normally fun ctioning battery is crucial for the sync hrotron X-ra y ima ging technique. The
widely used coin cell is not convenient fo r the X -ray because it blocks beam from
character ization. During the rotation in the characterization, the batter y has to be fully X -ray
transpare nt to allow the imaging p rocedure. In addit ion, properly sealin g is critical to ensure that
the cell can work normall y and meanwhile be stable during the characterization process. To
satisfy th e technological needs, w e d evelop a pol ymeric tube proof -of- concept batter y, as shown
in Figure 1a. On th e one hand, th e pol y amide -imide tube housing ensures hi gh X -ra y
transmission. On the oth er h and, the proof -of-concept batter y also has good sealing ability so th at
the electroc hemical battery remains functional for months. To further improve X-ra y
2.3 Gas development

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transmission though the batter y, we machine a groove around the entire circumference of the tube
housing.
A conventional ele ctrode is made from the electrode slurries with weight ratios of S i :carbon
black:binder of 80:10:10 in NMP. The slurr y was cast directly onto the top of one of the screws.
To remove the NMP, the casted screws were dried in an oven at 60 ℃ overnight. Before and after
casting process, the screw was weighe d to det ermine the wei ght of the electrode materials.
Metallic lithium was placed on the top of another screw, as the counter and reference electrode. A
polymer sep arator soaked with the el ectrolyte was placed between the li thium electrode and Si
electrode . The beamline proof-of -concept batter y was assembled in an argon filled glovebox with
humidit y and ox y gen levels below 0.1 ppm.
Our developed proof- of -c oncept batter y also allows performi ng various electrochemica l
character izations and anal y sis such as c y clic voltagram and dis/charge profile. Furthermore, the
beamline battery can undergo considerable stability durin g the characterization process.
Electroc hemical Measurements:
Cyclic voltammetr y (CV) and galvanostatic charge/discharge of the batter y were carried out
with IviumStat. A fresh batter y after assembly was measured for CV curves from 2 .5 V to 0 V at
a scan rate of 1 mV/s. I n the beamline, the battery is mounted on the centering and rot ation table
through the small screw on top of the electrode screw, thus confirmin g t he full stabilit y of th e
battery durin g the characterization process. Thin unclad copper wires are wrapped around the
small screws connecting each electrode and are s ec ured between th e two s mall screws. The wires
are then connected to the I viumS tat potentiostat, as shown in Figure 1c. The wire lengths are also
long enou gh to enable the cell to rotate in the beamline without any obst ructions. During X -ra y
imaging process, electro chemical testin g is perf ormed b y remotel y controlling the potentiostat
within the beamline hutch. The schematic illust ration of the battery in X BAM line in BESSY is
briefly described in F i gure 1.
Tomogr aphy Settings an d Data Acquisition:
The X-ra y tomography and radiography ima ging are conducted in the X BAM li ne in BESSY,
Berlin, Germany. XBAM li ne is a dedicated tom ogra ph y beamline, of which the usable X -r a y
spectrum is 5-100 KeV. I n this experiment, monochromatic s ync hrotron radiation at 20 KeV
(selected usin g a [W/Si] 100 mul ti-lay er monochromator), a 20 µm CdWO 4 scintillator and an
optical 10X objective in conjuction with a pco4000 camera (with a 4008 x 2672 pix el 2 CCD chip)
are sele cted 1-2 . The ultimate resolution is 0.876 µm. For each tomograph y , 2200 projections over
an angular range of 180° batte r y rotation with 4 s exposure time are recorded. As for the
radiographies, the battery w as directl y im aged during th e ch arge/discharge process with 4 s
exposure time.

Data Proce ssing
Data proce ss & 3D visu alization method:
2.3 Gas development

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The obtained raw data are firstly flat - and darkfield corrected and rearranged into sinograms. A
median filter with a kernel siz e of 3 ×3×3 voxels was applied to the ori ginal image for noise
reduction. After that, the Si/Si x L i and the other regions (carbon bla ck and poly mer binder) are
labeled b y simple threshold seg mentation (Ima ge J 1.49i). The histogram of the reconstruction
image s consist of two peaks for these Si/ Si x Li and the other regions and therefore the threshold
value can be chos en as the minimum value between two peaks. The reconstructed volume is
registere d using commercial software Avizo Fire, version 8.0 and for 3D presentation. We also
adopted commercial software VGStudio Max, version 2.2 for 3D visualization.
Analysis of (pote ntially) active partic les:
After labeling process via Aviz o Fire, a lot of properties of particles can be obtained through
the buli lt-in measure ments. Herein, the diameter of each particle is obtained by the built -in
equivalent diameter: EqDiamete r:
EqD= √ 4×𝐴𝑟𝑒𝑎
𝜋
To successfull y s eparate the particles connected with the gas/channel from these particles that
don’t connect with void/channel, we use the co mmercial software Image J . The idea is this:
suppose particle A is connected with void a, i f we Max imize the diameter of void a, then the
particle A will overlap with void a and we can use the Minimum operation between particle A
and void a to ge t the overlap zone A '. So far, we have onl y small part A ' of the particles that are
connected with the voids. After that, we increase the gra y value of A ' a nd add it with all th e
particles. In the l ast step, we us e the 3D H y steresis Thresholding to disc ern the particles that h as
A' from those that do no t have, namely the partic les that are not connec ted with voi ds. Throug h
this procedure, we can successfull y s eparate the particles that a re connected with voids from
those that are not.

SI Figure s

2.3 Gas development

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Figure S1. The first disc harge, the first c har ge and the second discharge voltag e -tim e profiles of a
beamline b attery at 0.01A g -1 .

Figure S2. 3D visualiz atio n of the evolu tion of developed gas i n the battery : a) After fir st discharge. b )
After first char ge. c) Afte r se cond discharge. In all sections, from the top to the bottom, the row shows the
3D visualization, t he plane view, the cross section view of the Si and gas/void com posite and the solitary
gas/void 3D rendering . The scale bar is 100µm in all pictures.
2.3 Gas development

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Figure S3. a) 3D visualiz ation of the evolu tion of gas/channel on t he l eft. b) and the active particle
population, on t he right in Region2: from top to bottom i s the res ult of, after first discharg e; after f irst
charge; af ter second d ischa rge. The sc ale bar is 100µm in all pic tures.

Caption s for the Su pportin g Movie s

http://pubs.acs.org/doi/abs/10.1021/acsami.6b00708

SM1:
Three Dimension Visualization of Reconstructed Silicon Electrode

SM2:
SM2a-Three Dimen sion Visualization of Reconstructed Firstly Discharge d Si Electrode
SM2b-Three Dimension Visualization of Reconstructed F irstl y Char ged Si Electrode
2.3 Gas development

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SM2c-Three Dimen sion Visualization of Reconstructed Secondly Discharge d Si Electrode
SM3a:
SM3a-1-First-Discharge-Movie_Region_Left
SM3a-2-First-Discharge-Movie_Region_Right
SM3b:
SM3b-1-First-Charge-Movie_Region_Le ft
SM3b-2-First-Charge-Movie_Region_Rig ht
SM3c:
SM3c-1-Second-Discharg e-Movie_Region_Left
SM3c-2-Second-Discharg e-Movie_Region_Right

References

(1) Arlt, T.; Mai er, W.; Tötz ke , C.; Wanne k, C.; Ma rkötter, H.; Wieder, F.; Banhart, J.; Lehnert, W.;
Manke, I. Synchr o tron X-R ay Radi oscopic in Situ Stud y o f Hig h-Temperatur e Polymer El ect rolyte Fuel
Cells - Effect of Operati on Conditions on Structur e of Me mbrane. J. Power Sourc es 2014, 246 , 2 90-298.
(2) Manke, I.; Banhart , J.; H aib el, A.; Rack, A.; Zable r, S.; K ardjilov, N. ; Hilger, A.; Melz er, A.;
Riesemeier, H . In Situ Investigation o f the Discharg e o f Alkaline Zn – Mno2 Ba tteries wi th Synchr otron X-
Ray and Neut ron Tom ogra phies. Appl. Phys. Lett. 200 7, 90 , 2 14102-214 105.

2.3 Gas development

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2.4 Degrada tion of l ithium ion batter ies based on ~1 00 µm-sized Si par t icles

Reprinted w ith perm iss ion from DOI : 1 0.1016/ j.jpowsour.201 6.04.12 6. © 2016 E lsevier B.V . All rig hts
reserved.

Investigation of failure me chanisms in silicon b ased half cells during the first cyc le by micro
X-ray tomography and radiography

Fu Sun a,b, *, Henning Markötter b , Kang Dong a,b , Ingo M anke b , Andre Hilger b , Nikolay
Kardjilov b and John Banhart a,b
a I nstitute of Mater ial Scien ce and Techn ol ogies
Technical Univ er sity Berlin
10623 Berl in, Germany
b Helmholtz Cent re Berlin for M aterial s and Energ y
Hahn- Meitner- Platz 1
14109 Berl in, Germany
* Correspond ing Author : fu.sun@helmholtz- berli n.de

Abstract
Two proof-of-concept batteries were designed and prepare d for X- ra y mi crotomography and
radiography charac terizations to investi ga te the degradation mechanisms of silicon (Si) based half
cells during the first c yc le. I t is hi ghlighted here for the first time that, apart from the significant
volume expansion-induced pulverization, the electrochemica l “ deactivation ” mechanism
contributes significantl y to the capacity loss during t he first charge process. In addition, the
unexpected electrochemicall y in active Si particles are also believed to substantially decrease the
energy densit y due to the inefficient utiliz ation of loaded active material. These unexpected
findings, which cannot be deduced from macroscopic electroc hemical ch aracterizations, expand
the inherent explanations for pe rformance d eterioration of Si -anode material based lithium ion
batteries ( LI Bs) and emphasize the vital value of microscopic techniques in revealing the
corre lation between macroscopic electrode structure and the overall electrochemical performance.
Key words
lithium ion b attery; si licon part icles; degrada tion mec hanisms; X-ray m i cro tomo graphy; X-ray m i cro
radiography

Introduction
Following their commercial int roduction in the earl y 1990s, lithium ion batterie s (LIBs) have
penetrated ubiquitousl y into the mar ket for energy storage s y st ems, e .g. laptops and mobi le
phones [1, 2]. More recentl y , advanced LIBs with a larger specific ener gy , hi gher power d ensit y
and longer c ycle life h ave been considered for powering cle an electric v ehicles (EVs) and plu g -in
hybr id vehicles (PHVs), as well as for the stora ge and dist ribution of energy from sustainable
sources, such as solar and wind energy [3 -6]. The key to fabricate such n ex t-generation LIB is to
exploit high-pe rformance electrode materials. From the material ’ s point of view, sil icon (Si) is a
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promising candidate for the anode o f LIBs becaus e it possesses the hi ghe st theoretical capacit y o f
3579 mAh/g when lithiated to L i 21 Si 5 , that is about 10 ti mes larger than the currently
commercialized carbona ceous anode, whi ch features 372 mAh/g fo r LiC 6 [7 -9]. This promising
potential of the high specific capacit y has spurred considerable investigations of Si thin fil ms [10] ,
carbon mixed [11, 12] and metal coated [ 13] Si composites, along with v arious Si nanostructures,
such as nanoparticles [14] , nanowires [15], nanospheres [ 16] and nanotubes [ 17] that could serve
as building blocks for high-performance anodes. Despite these extensive efforts, Si -based anode
LI Bs still inevitabl y suffer from substantial capacity decay during lithium insertion and ex traction
and thus they are now below the requirements for practical applications.
To shed light on the underlying degradation me chanisms, a range of investigation techniques
such as scanning electron microscopy (SEM) [18] , transmission electron microscop y (TEM) [ 19] ,
X-ray diffraction ( XRD) [ 20], nuclear magnetic resonance (NMR) [ 21] and Raman spe ctroscop y
[22] have been adopted. As a result, our knowled ge of th e intrinsic b ehaviour of Si upon lithi um
insertion and extraction has been si gnificantly enhanced [23] and various possible explanations
for the performance deterioration have b een p roposed. For instance, it has been discovered that,
apart from the loss of t he lithium inventor y and electrol y te/binder d ecomposition [ 24], the
preferential volume expa nsion along Si [110] directions is as larg e as ~300% during lithium
insertion and this extraordinary volum e change has been proposed as the primary factor
contributing to the fatal ca pacit y deca y [25]. Nevertheless, p rior an al y ti cal tools are inher ently
limited. For example, most of the measurements abovementioned e xplore the degradation
mechanisms on the atomic or single nanoparticle level, overlooking th e interplay s amongst a
multitude of particles and the interactions betwee n the active material and the conductive/binder
agent. What is more, a particular battery design, such as open structure and specialized electrol y te
that does not adequately simulate the real battery operating conditions is widely employed. O n
the other hand, most conventional tools only specialize in revea ling the structur al and
compositional information without imaging capability, lackin g effec tive spat iall y resolved
information about the degradation mechanisms. In parallel, some investigations are carried out ex
situ , i.e. post-mortem, as opposed to in situ characterizations of d ynamic processes. Considering
that a realistic composite electrode includes an assembly of ensembles of active particles, an
organic pol y meric binder and a conducting a ge nt, the abilit y to probe the d y namic deterioration
mechanisms on a mul ti-particle electrode level is of technolog ical and practical importance.
Therefore, it is crucial and emergent that fundamental research techniques are highl y need ed to
further pr omote the exploration of electrode degradation .
Recently , X- ra y imaging based on either laboratory X -ra y or s y nchrot ron X-ra y sour ces has
rapidly evolved into a p owerful characterization tool in materials science [26-42] . Specifically ,
X-ray imaging has enabled researchers to obtain unprecedented insights into LI Bs non-
destructively on a l ength scale r anging from particles to entire electrodes and has contributed
markedly to our unde rstanding [29-33]. The pioneering r esearch of emplo ying s y nch rotron X -ray
tomogra ph y b y Ebner et al . features a direct observation and quantification of elec trochemical
and mechanical degradation in a SnO anode [34] . Meanwhile, Gonz alez et al . have visualized the
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

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expansion of large Si parti cles during the first lithiation step in three -dimensions (3D) b y means
of tomography base d on a laboratory X-ra y source [35].
Herein, b y emplo y ing both laboratory X-r a y tomograph y and radiograph y, we re- in spect the
underly ing mechanism of performance d egradation of Si based half cells from the pe rspective of
entire electrodes, which impli es spanning the length scales from individual active particles to the
macroscopic electrode e nsemble. On the one hand, X-r a y mi crotomography grants us th e abilit y
to track the structural evolut ion induced b y volu me cha nges in thre e di mensions . On the other
hand, in operando micro X-ray radiograph y enables us to directly obse rve the changes of the
active particles and an enti re electrode during lithium insertion and extraction . By the
combinatorial X- ray tom ograph y and radio gra ph y, for the first time we highli ght that, apart from
the huge volume ex pansion/contraction of Si particles during c ycles, the st riking inhomogeneous
li thiation/delithiation mec hanism amongst ensem ble active particles obse rved in an electrode ,
which cannot be easil y detected in conventional electrochemical measurements and, the
unprece dented electrochemical “ deactivation ” of original electrochemical active particles are
another two key factors contributing to the substanti al pe rformance d egradation. This stud y
expands the inherent ex planations for performance deterioration of Si -anode material based LIBs,
and the n ew sights wo uld open new d esign p rinciples and opportuniti es for high-capacit y
electrode materials for nex t-generation energy storage s y stems.
Experimental Section
Materials:
Silicon was re ceived from Elkem AS , Norwa y. P olyviny lidene difluoride (PVDF) binder, ca rbon
black, Celgard s eparator, CR2032 coin cells and lithium were purchased fr om MT I Cor. USA. N -
methyl p y rrolidone solvent (NMP) and 1M LiPF 6 in a volume-ratio (1:1) mix ture of eth y lene
carbona te (EC) and dim ethyl carbonate (DMC) were purchased from Sigma Aldri ch. Titanium
foil is from ANKURO Int. GmbH, German y . Th e housing of p roof- of -c oncept beamline b atter y
is made of poly amide-im ide (Torlon ), from McMaster-Carr compan y .
Characterization:
Scanning electron microscopy (SEM) image wer e taken using a Z eiss ult raplus mi croscope. The
electroc hemical cha racterizations were conducted b y using an I vium CompactStat station,
Iviumtec hnolog y .
Battery Preparation:
The composite electrode is made of slurries wit h weig ht ratios o f Si :carbon black:binder of
75:15:10 in NMP. For the tomo-cell, the slur r y was first sandwiched betw een two glasses and put
into an oven at 60 ℃ to dr y and form a block, w hich was then cut into small pieces with a razor
blade and put directly onto the top of the screws . The resultant S i composite electrode was a
small piece of 1.7 x 1.7 x 0.2 mm (length x width x height). Before the Si composit e was
assembled into the tomo-cell, it was weighe d b y a digital balance and the amount of S i particles
was determined from the origian mass ration. The Si composite electrode mass in the t omo-cell
was around 0.9 mg. For the radio -cell, the slurr y was casted onto the 5µ m thi ck tit anium (Ti) foil.
The area of the slurr y w as around 4 x 3 mm (len gth x width). To remove t he NMP, the c asted Ti
foil was also dried in an oven at 60 ℃ . Before and after the casting proces s, Ti foil was weighted
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

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to determine the weight of the electrode materials. The S i composite in the radio-cell was around
1 mg . These two proof-of-concept batteries were assembled in an argon filled glovebox with
humidit y and ox y gen lev els below 0.1 ppm. During battery ass embly, the se two cells were only
hand-assembled without exerting extra pressures . For the tom o-cell, metallic lithium (2.5 mm
diameter, 1 mm thic k) was placed on top of t he other sc rew, and served as a counter and
reference electrode. For the radio-cell, a lollipop shape copper foil without inner reg ion (out
diameter is 10 mm and inner diamete r is 6 mm) was used as a cu rrent co llector and the Ti foil
was placed on top of it. Metallic lithium (6 mm diam eter, 1 mm thick) was used as the opposite
electrode . A pol y m er separator w as plac ed betw een the lithium electrode and the Si electrode in
both cells. A fter filli ng these two cells with sufficient electrol y te, the y w ere sealed off and then
taken out of th e glovebox. For a further com parison of the electrochemical a ctivit y of Si
composite electrode in the two proof- of - concept cells, commerc ial CR2 032 coin -cell with the
same Si electrode was assembled and tested.
Micro X-ray Tomography & Radiography Charac terizations:
X-ray tomo graphy was c onducted on the tom o -cell in the original state an d after the voltage h ad
dropped to 0.03V, 0.02 V, 0.01V and 0V durin g disch arge. During the discharge pro cess, the
curre nt w as set to be 0.04 mA. The five tomographies are im p lied b y alphabet a, b, c, d and e in
Fig . 4, Fig. 5 and Fi g. 6. The particul ar parameters for the X -ra y tom ography are as follows: the
energy is 60 kV, the current is 166 µA, the sou rce detector distance (SDD ) is 500 mm , the sou rce
object dist ance (SOD ) is 58 mm. Given the resulting magnification, each detector pix el therefore
repre sented 5.76 µm/voxel in the sample. The exposure time is 2.1 s. During the 360° rotation,
800 projections were recorded. During the char ge process, the current w as switched to 0. 01mA
and the tomography was conducted after the volta ge had risen to 0.5 V, 0.7 V, 1 V, 1.5 V and 2 V.
The five tomographies are im plied b y alphabet f , g, h, i, and j in Fig. 4, Fig. 5 and Fig. 6. The
locations of the p articles shown in Fig. 3, Fig. 4 and Fig. 5 are marked in SI Fig . 4. Mor e
information is given in the Supporting I nformation (SI).
For the X- ray radiography measurements, c haracterization wa s c onducted in situ during the
whole discharge/charge process of the radio -cell. During the dischar ge p rocess, the current was
0.07 mA for 0-18h and then was switched to 0.04 mA for the rest of discharge p rocess. Durin g
the charge process, the current was set to be 0.015mA for 6.5 h and then was changed to 0.01 mA
until the voltage was 2 V. The ra dio- cell was facing the X-ray source during the characterization.
Data Proce ssing:
To keep the validity o f the measurements, onl y a median filter (1×1×1) was performed in I mageJ.
3D particle visualization was performed b y VGStudio MAX. Anal ysis of the particles is
conducted by I maje J . More information is given in the S I .

Results and Discussion
X-ray microtomograph y and radiograph y was con ducted with a laborator y X -ray CT s y stem [43].
To be full y comp atible with X -ray imaging and at the same ti me to completely repre sent working
LI Bs, we have designed and manufactured two dedicated proof -of-concept electrochemica l cells:
a cell for tomograph y (tomo-cell) and a c ell for radio gra ph y (radio -cell). Fi g. 1 displa y s
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

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photographs and th e corresponding schem atic illustrations of the two p roof-of-concept b atteries
as well as a schematic representation of the X-ra y micro CT setup [44].

Fig. 1. Im ages of the proof-of-concept batteries: a) tom o - cell and b) r adio- cel l. c) Corresponding
schematic represen tation of the tom o - cell. d) and g) Corresponding schemati c representation s of t he radio-
cell. e) E nlarg ed region of i nteres t com prising from t op to bottom : lithium (blue), separato r (grey) and
electrode material (green). g) Schematic illustration of t he radio -cell, fr om top to bot tom are, the upper
housing (brown), sea l ing ri ng (yellow), lithium p late (blue) with copper wire, separat or (wh ite),
Si/carbon/bind er composite (green ), titanium foil current colle ctor (gray), lollipop - shaped copper current
collector (copper), the lower ho using (brown). f) Schematic representation of the X - ray micro CT setup.
From l eft to right: X - ray source (red), cone X-ray beam (yellow), sample represent ing either of the two
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

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cells (green) and r otation table (grey), detector (blue). The waves in c) and d) represen t t he direction of X -
rays.

After assembling the tw o proof-of-concept batteries, cyclic voltammetr y (CV) was performed to
verify the reduction and oxidation chara cteristics of Si. The ins ets of Fi g. 2 show the CV curves
of the two batteries sc anned at rate of 1 mV s -1 in the potential window of 0-2.5 V. Although
there a re parasitic redox pe aks (a r rows in CV figure) associated with c ontamination a nd/or
electrolyte d ecomposition [45, 46] , the clearly visi ble cathodic p eak above 0 V and an anodic
peak at 0.5 V a re related to the Li allo y in g and de -allo y ing with Si, r espectively. The results are
in good a greement with previousl y reported Si/Carbon composit es [9, 47] , thus confirming the
electroc hemical reactions inside our proof-of-conc ept batteries.
Then we conduc t the X-ra y ima ging investigations b y galvanostaticall y discharge/charge these
two electrochemical batteries: for tomograph y , th e discharge/charge process was stopped during
the tomographic data acquisition and continued afterwards (i.e. the measurement is quasi in situ ).
For radiograph y , the measurement is continuousl y conducted simultaneousl y durin g the
discharge/charg e p rocess of the radio-cell (i.e. i maging is trul y in situ ). The discharge/charge
curves of both c ells, showing the t y pical ch aracteristic feature of Si are displa y ed in Fig. 2 and
they are in good agreement with previous results [9]. Note that for the tomog raphic imaging, a
newly assembled battery is used, while for radiogra phic imaging, the batter y is first used to
obtain the C V scans and then directl y used for in situ radiograph y me asurement (read mor e in
Supporting I nformation (S I)) . For a further comparison of the electrochemical activity of S i
materials in the two proof-of-conce pt cells, commerc ial coin- cell assembled with the same Si
electrode was conducted and the resultant discharge/charge curv e is sho wn in SI Fig . 1. It is
found that these three cells have a low coulomb efficiency during the first charge and the reason
may stem from the sig nificant particle pulverization of the large Si particles currentl y used [19].
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

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Fig. 2. Electrochemical characterization of the tomo-cell and radio-cell: a) the first discharge-charge curve
of the tom o- cell, the inset s hows the CV curves ; b) sam e as a) f or the r adio-cell.

Micro-computed X-ra y i maging is an emerging anal y ti cal technique th at measures v ariations in
X-ray attenuation of sampl es and it is particula rly well suited to track change s from particles
morphologica l evolution to electrode architecture changes over time as c haracteriz ations can be
conducted in situ and non -destructively [ 48] . Consequently , previousl y unp recedented
degrada tion mechanisms are highli ghted here for the first time. In the present work, a dist ribution
of micron siz ed Si particles ranging from 125 to 180 µm in diameter are used for both tom o -cell
and radio-cell because of the renewed interest in using micron siz ed particles as commercial LIBs
[49] . To begin with, 3D visualization of S i particles evolution within an operating LIB durin g the
first dischar ge/charge pr ocess is presented by mi cro X-ra y tom ography. F irst, the pristine state of
the freshly prepared tomo -cell is tomographicall y r ecorded with 800 projections coverin g an
ang ular range o f 360° wi th 2.1 s exposure time. Then the batter y is galvanostatica ll y disch arged
to 0.03V [35] . At this poi nt, the discharge process is stopped and the 2 nd tomography is conducted.
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

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After this, the discharge is continued and followe d b y a 3 rd tomography. Altogether 10
tomogra ms du ring the first discharge/charge process ar e made (5 during the disch arge and 5
during the charge). All the tomograms conducted during the first discharge/charge process are
implied b y red spots in Fig . 2 a and all the para meters are the same. Fig. 3a and Fig. 3b show
projection images of the Li elect rode|separator|Si electrode ass embly within the proof -of-concept
batteries. W hile the L i electrode and the porous plasti c separator appear nearl y invi sible due to
their low X-ray absorption c oefficients, the Si elec trode is clearl y visi ble. Fig. 3c shows a
graysca le slice aft er r econstruction of the tom o-cell and Fig. 3d an enlarged part of the radio -cell.
In o rder to qualitativel y validate our X - ra y imaging te chnique s, scannin g electron micrographs
(SEM), as sho wn in Fig. 3e, are given to compare with th e reconstruct ed r aw gra y scale sli ce
image s and the radiography im ages. The unambiguous agreement indicate s that the X -ra y
imaging tec hnique accuratel y captures the morph ology of Si electrode.

Fig. 3 X-ray i mag i ng res ults of the assem bled L i electrode| se parator| Si el ectro de : a) X-ray projection
imag e of the t om o- cell, the scale bar is 1 mm. b) 2D radiography of the radio-cel l, the scale bar is 2 mm. c)
Exam ple of a reconstructed slice im age of the tomo -cell. d) Cu tout of a radiogra phic image of the radio-
cell. e) Scanning electron micrograph of t he m ixed Si electrode material. The sc ale ba r in c -e) are all
200 µm .

We now track in detail t he morpholo gical ch anges during lithium insertion and ex traction of on e
randomly chosen particle, illustrating other electroactive particles evoluti on as a func ti on of
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

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discharge/charge . Fig. 4 shows a serie s of cross-s ectional slices taken from tomographies of the
Si particle in differe nt stages of discharge /charge. The corresponding 3D rendering of this
particle is displa yed in S I Fig. 1. Histograms of the attenuation coefficients in sub volumes
containing the particle ar e displa y ed as a fun ction of discharge/charge process see Fi g. 4k and Fig.
4l. The leftmost peak corresponds to the w eakl y absorbing carbon black, binder a nd electrol yte
surrounding the particle. F eature ch anges to th e right most peak are d irectly r elated to the
evolution of the active Si particle. Prior to t he electrochemical r eduction, the attenuation
coeff icient hist ogram for the particle consists of one peak located gra y level of 6000. During
lithium insertion, see Fig. 4k, along with the magnitude of this peak progressively decreases and
moves to the left, the particle evolves gradually i nto a weakly X - ra y atte nuating material. This
implies that the particle has completely t ransformed from the high -densit y Si phase to a low -
density Li x Si phase (1<x<4.4). Meanwhile, a notable volume expansion to around 200%
(compared with pristine state ) is observed in th e fully lithiated state. These r esults are in good
ag reement with pr evious results [19, 35] . During lithium extraction, see Fig. 4l, along with the
attenuation pe ak increases and shifts towards the original direction, we observe a homogeneous
increase in attenuation c oefficient. However, i t is worth noting t hat, neithe r the pe ak position nor
the peak shape of Si is restored to the ori ginal st ate b y the end o f the first charge proc ess. This
may stem from incomplete delithiation and/or significant trapping of lith ium b y th e electrol y te
decomposed b y products [50] , which is considered as one of the reasons of the performance d ecay
during the first cyc le.

2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

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Fig. 4. Evolution of an electrochem i cally active Si particle: a) -e), tomog r aphic slice s t hrough t he particle
during the first discharge st ep (see arrow direction); f)-j), tom ogram s of the first charge process (see arrow
direction). k) and l ), attenuation coeffic ient histogram s of the specific particle as a function of discharge
and charg e (dashed r ed line is the pristine state) . Outline of the pristine state ( red outline in a)) and
outlines of the discharged state ( yellow outline in e)) and charged state (green outline i n j)) are shown f or
visual com par ison. The sca le bar in a) is 100 µm l ong and is the sam e in all the im ages.
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

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Fig. 5. Evolution of electrochem ical deactivatio n of a Si particle : a)-e), t om ographic slices during the firs t
discharge step ( see arrow direction) ; f) - j), tom ographic slices during t he first charg e st ep (see arrow
direction). k) and l ), attenuation coefficie nt histogram s during discharge a nd charg e, res pectiv ely. Outline
of the pristine state (red outline in a)) and outline s of the discharg ed state (yellow outline in e)) and
charged state (green outline i n j)) are shown for visual com par ison. The scale bar i s 100 µm long and
applies to al l the im ages.

Surprisingly, we have also detected an unp recedented phenomenon that many elec troactive Si
particles become electrochemicall y ina ctive during the first delith iation process. This
electroc hemicall y deactivation progress is illustrated in Fig. 5 , in which, the electroactive particle
originally effectivel y par ticipates in the discharge process, from Fig. 5a to Fig. 5 e, but turns to be
in active during the subsequent charge process, contribut ing to the capacity loss during the first
charge p rocess, as w e can clearly see the un changed shape of the particle, from Fig. 5f to Fi g. 5j,
and the stationary attenuation profile during delithiation, Fig. 5l. Another unexpected
phenomenon is that several Si particles n ever under go lithi ation/delithiation during the
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

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discharge/charg e process, i.e., they are electrochemicall y in active durin g the whole macroscopic
battery level lithium insertion and ex traction process, as shown in Fig. 6. On the whole, f rom an
entire electrode-scale point of view, as displayed in S I Fi g. 4, we can unambiguousl y obse rve that
some electroc hemicall y active Si particles undergo a continuous phase t ransformation during
macroscopic batter y lev el discharge/charge, some electroc hemically a ctive Si particles undergo
an electrochemical deactiv ation process after the lithium insertion and some non -electroactive Si
particles a re completel y inactive during the first c y c le even as the voltage drops to zero against
Li + / Li. In addition, we also observe that some S i particles emerge and/or disappear in the same
slice as a function of dischar ge/charge process. These distinctive phenomena clearl y im pl y that
apart from the lar ge volume expansion during lithi ation/delithiation, there are man y other factors
that contribute to the ultimate perf ormance deterioration in S i-based LIB s.

Fig. 6. Evolution of the electrochem i cally non-active Si particles: a)- e), tomog r aphic slices of the first
discharge process (see a rrow direction); f) -j), tomographic slices of the first charg e proce ss (see arro w
direction). k) and l), attenuation coef ficient histog ram s as a function of discharg e and charge, respec tively;
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

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Outline of the pristine state (red in a)) i s duplicated to the di scha rged state e) and charged st ate j ) f or
visual com par ison. The sca le bar is 100 µm .

Moreover, in operando 2D X-ra y mic roradiograph y provides more sophisticated temporal
information within the LIB. As shown in Fig. 7, we can cl early observe the expansion/contraction
of the whole electrode material (dotted blue contours from D01 to C10), the ex pa nsion and
contraction of the electro chemically active pa rticles (green panel), the deactivation phenomenon
of originall y electroactive Si particles (red p anel) and a few Si particles that are not
electroc hemicall y active throughout the whole discharge /char ge process (y ellow outl ines from
D01 to C10). Furthermore, the “ core-shell ” m odel reaction is clearly observed durin g the
discharge/charge p rocess (see the evolution of particle from Fig. 7a to 7 j) [ 8]. The significant
volume expansion -induced pulverization (see the green panel) ma y le ad to electric disconnection
of active particles from c urrent collectors [25] . These results are in good agreement with the 3D
X-ray mi crotomograph y and the whole in opera ndo discharge/charge pr ocess is presented in a
movie in S I .

Fig. 7. In situ radiographi c charac terization of the radio -cell: D01 to C10 represent stag es in the first
discharge (D ) and charg e (C) step. The reg i on encirc led in blue is the orig inal area that cov ers the
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

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electrode material, increas ing numbers ref er to el apsing time. The r egion in green boxes represents t he
evolution of an electroche mical active Si particle from a to t (the copper lead is se en at the bottom lef t of
each f ram e). T he r egion in red boxes display s the evolution of electrochem ical deac tivation from a’ t o t’.
The region encircled in yellow contains many electrochem ica lly inactive particles . T he scale bar in D01 is
1 mm l ong, that in the sm all boxes 100 µm . More infor m ation is provided in SI .

Fundamenta ll y sp eaking, LIB s oper ate throu gh th e reversible inse rtion into or removal of lithium
from electroactive host materials, respectivel y [51] . The abilit y of th e materials to accommodate
the changes associated w ith the chemical phase t ransformations that accompan y the variations in
lithium c oncentration determines the ele ctrode’ s utility and batter y p erformance . F rom an
electroc hemical point o f view, lithium insertion or extraction process entails Li + and e -
simultaneously [52]. Contrar y to proposed idealised electrochemical mod els and simulation that
[53] electroche mically d riven phase conv ersion is homog eneous and isot ropic, such conditions
are hardl y satisfied in a real comme rcial LIB due to the complex electrode structures (ensemble
of active m aterials, conductive agents and binders) and t he complic ated morpholo g y and
conditions with respec t to porosity , tortuosit y , conductivity and pe rcolation abilit y for the
electrolyte [54]. I n fac t, increa singl y mor e attention is being paid to the investigation of lo cal
electroc hemical r eactions and their relationship with electroche mical pe rformance on a
macroscopic and electrode performance [55, 56] . For example, experimentall y a discr epanc y is
observed between electrochemical me asurements that represent the overall state of a cell and
spectroscopic data that reflect the loc al state [57, 58]. I n addition, the heterogeneous lo cal depth
of discharge (DOD) and the non -uniform local current distribution have been directl y
demonstrated b y Zhang et al. [ 59, 60] and Ng et al. [ 61] respectively , by customizing multiple
working electrode LIBs. Similarl y , si gnificant inhomogeneit y of local D OD in different locations
of a commercial LI B have been dir ectl y detected b y Cai et al . [62] and Pax ton et al . [63] through
neutron diffraction and energy-dispersive X- ray diffrac tion. Moreover, by using µm -resolved
Raman spectroscop y , Nanda et al. [64] pr esented visuall y the local “ spectroscopic ”
electroc hemical v ariations on an LI B electrode. F inally, recent mod eling investigations b y Zhao
et al. [65] show a direct corre lation between energy densit y and the no n -uniformit y in local
curre nt distributi on, demonstrating a potential gain as hi gh as 40% in energ y d ensit y th rough an
improved current distrib ution. Apparentl y, these locally non -uniform electroche mical reactions,
curre nt dist ributions, ioni c/electric conductivity within el ectrodes wi ll impact on batter y
performance in a variety of wa ys, including reduced energ y and power, unde r utiliz ation of
capacity , localized heat genera tion and overcharge or over-disch arge of active materials [56, 66].
As in our case, it a re these electrochemically active Si particles w hich are ionicall y and
electronically accessed by the electrolyte and connected to the conducting network that directly
dictate battery performance. I n other wor ds, the total electrode dis charging/charging current
concentrates onl y on the electrochemicall y active Si particles, as clearl y sh own in Fig. 4, S I Fig.
4 and Fig. 7 . Thes e el ectrochemicall y active Si p articles und ergo li thium i nsertion and extraction
to store and release energ y, evidenced b y the X-r a y attenuation coefficient change. During
lithium uptake and release, these electroactive S i particles undergo notable volume expansion
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

- 95 -

accompanied by the gene ration of subst antial compressive and tensile stresses [67] . The
mechanical forces generated will drive local displacement and re arrangement of other Si particles
within the composite electrode, leading to the phe nomena that some Si particles emerge and some
disappear in one and the same slice, as shown in SI Fig. 4 and Fig. 7. On the other h and, the
observed “ deactivation ” mecha nism of o riginall y electrochemically active S i particles prob ably
has three sourc es: First, an electric contact disconnection. Under the influence of the mechanical
force s generated, a lot of displacement/rearrangement of the particles oc curs with numerous
losses or va riations in the conta cts. Taking into account that in a stack of particles the electrons
are transferred between particles by hopping or tunneling with the contact resistivit y de pendin g
exponentially on the gap between the contacts, and that a gap of 8 -10 nm makes a contact
electronically insulating [68] , it is pl ausible to say that the deactivation mechanism is attributed to
electric disconnection. The second possi bilit y is the insulating cr ystalline L i 2 O oxide la yer. B y
using in sit u TEM, He et al . report that upon initial Si lithi ation, the formed cr ystalline L i 2 O, the
product of Li reaction with the native sil icon ox ide laye r, will partially insulate the pa rticles
during subsequent delit hiation cycle [ 69]. The third possible reason is the unshrinkable
conductive matrix . Weker et al . proposed that upon the delithi ation, when the pa rticles be gin to
contract, the conductive matrix does not necessarily shrink back to fill the space created b y the
contracting particl es, thus leading to the deactivation of , in their case, germanium particles [ 70].
Right now we cannot ide ntify a single dominant factor. However, sp ecial focus in the future work
should be placed to eliminate this phenom enon due to the resultant substan tial capacit y loss es. I t
is also worth noting th at the unex pected ele ctrochemica ll y ina ctive Si particles withi n the
electrode . From the fact t hat they are electrochemically inactive from the beginning, it is assumed
these S i particles are dis connected into locations of ionically or electronicall y insul ating islands
during the electrode preparation, for example binder redistribution during drying, calendaring,
cutting and compacting [ 52] . To get a further quantitative anal ysis of the influence of dif ferent
types of Si particles on t he obtained discharge/charge capacit y, we investigated 78 particles in the
same slice in tomo-cell with diameters lar ger than 50 µm, as shown in SI Fig. 6. I t was found that
13% (10) o f the S i partic les were ele ctrochemically in -active through the first c ycle, 87% (68) of
the Si particles experienced the first lithiation process and only 24% (1 9) of the Si pa rticles
experienced the first deli thiation process. Assuming the theoretical capacity of S i is 3500 mAh/g,
the obtaine d firstl y discharged capacity (3000 mAh/g that is around 85% of the theore tical
capacity ) and the obtai ned firstl y charged capacity (500 mAh/ g that is around 14% of the
theoretica l capacit y ) agre e with the results from the quanti tative analysis (fr action of 87% and 24%
(within error considera tion)). More information can be found in SI.
The unexpected electrochemical “ deactivation ” mechanism and the presence of man y
electroc hemical inactive Si particles are alarming because the y ca nnot be easil y characterized b y
conventional macroscopic electroanal ytic al characterization techniques and should draw
attentions from electrode engineers and simulatio n ex perts. From an electrode en gineer ’ s point of
view, the electrode architecture engineering optim ization is crucially important. On the one hand,
in addition to focusing onl y on individual particle features to develop n ext -ge neration LIBs, more
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

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and more efforts shoul d be put in optimiz ing the entire electrode a rchitecture, which can involve
all active material particles ionicall y and electronicall y connecting to e lectrolyte and electric
conducting n etwork to maximize the utilization of active materials. On the other hand, new
conductive/binder agents are highl y n eeded. Th e matrix of carbon/binder is supposed to sustain
the active m aterial ’ s ex pansion/contraction during c y cles as well as to provide an efficient
dyna mic electric/ionic conducting pathwa y even under si gnifica nt el ectrode transformation.
Currently , some ex ploratory research of dev elop ing s elf-healing pol y mer [ 49] and electric and
ionic conductive pol y mer [71] is underway. From a simulation expert ’ s point of view, in order to
develop an electrochemical model that c an be used to gain insight into int ernal processes, to
predict performance a nd opera tion and opti mize ce ll desig n, the homogeneous and flawless
idealized microstructure characteristics should be compromised with the re al complex composite
electrode s.
Conclusions
In summar y , we re-explore the mechanisms of d y namic d eterioration of Si anode LI Bs on an
electrode scale b y emplo y in g X -ray ima ging tom ograph y and radiograph y and for the first time
highlig ht that, apart from the sig nificant volume expansion -induced pulverization and electric
disconnection from current collectors, ele ctrochemical “ deactivation ” contributes significa ntl y to
the capacit y loss during the first charge process. In addition, the presence of a not able numbe r of
electroc hemicall y in active Si particles is also believed to substantiall y decrease energy density
due to the ineffic ient utilization of loaded active materials. These unexpected findings, which
cannot be obtained by macroscopic electro chemical characterizations and conventional
structural/compositional characterizations, provid e us with novel insights i nto the mechanisms of
performance degradation of S i anode LIBs. From practical point of view, c ommercially o riented
researche s into the local microscopic electrochemi cal reactions could be mot ivated and a roused
as it governs directl y the energ y densit y and c apacit y retention of a real LIB. And more attention
should be paid to the further investigation of the corre lation between macroscopic
electroc hemical performance and local b ehavior of active materials , to g uide the selection and
optimization of electrode materials and the manufacture of the electrode.
Supporting Inform ation
Supporting Inf ormation is available in the online version or from the author.
Acknowledgements
We thank Norbert Bec k for fabricating the beamline batter y and Elkem AS for providing us wit h
Si particles. This work was sponsored by the Helmholtz Association and the China Scholarship
Council.

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2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

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2.4.1 Suppor ting Inf ormat ion

Investigation of failure me chanisms in silicon based half cells during the first cyc le by micro
X-ray tomography and radiography

Fu Sun a,b, *, Henning Markötter b , Kang Dong a,b , Ingo M anke b , Andre Hilger b , Nikolay
Kardjilov b and John Banhart a, b
a I nstitute of Mater ial Scien ce and Techn ol ogies
Technical Univ er sity Berlin
10623 Berl in, Germany
b Helmholtz Cent re Berlin for M aterial s and Energ y
Hahn- Meitner- Platz 1
14109 Berl in, Germany
* Correspond ing Author : fu.sun@helmholtz- berli n.de

This section in clude s :

Data Acquis ition a nd Processing
SI Fig. 1, SI Fi g. 2, SI Fig. 3 , SI Fig. 4, SI Fig. 5 an d SI Fig. 6
Caption for the Sup plementary Movie

Data Acqu isition an d Processing:
The obtained tomography raw data are first flat-field and dark-field corrected and rearran ged into
sinogra ms. A median filter with a k ernel size of 1 ×1 voxels is applied to the ori ginal im age for
noise reduction while maintaining the validit y of the dataset. The reconstruc tion software
Octopus (8.8.2-64 bit) is used to reconstruct the obtained dataset. Because of the hu ge electrode
expansion or contraction during dis charge or char ge, the reconstructed 2D slices shown in Fig. 4,
Fig . 5 and Fig. 6 do not necessarily represent exact ly the s ame plan e from one tomography to
another. However, we compare with one slice to another to find the closest slices for each particle.
For the particles shown in Fig . 4, it is located 46 µ m from the bottom stainless screw current
collector and is marked b y green cir cle in S I Fig. 5. For the particles shown in Fig. 5, th e dist ance
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

- 100 -

is 40 µm from the bottom stainless scre w current collector and is marked by ye llow rectangular in
SI Fi g. 5. For the p articles shown in Fig . 6, the distance is 80 µm from the bottom stainless screw
curre nt colle ctor and is marked in red round rectangle in S I Fig . 5. N ote that all the locations are
measured in the pristine state.
For the cal culated volum e expansion, we assume that the Si particles are round and measure the ir
diameter in the pr istine state and then compare with that of discharg ed state. The formula is:
V= 4
3 𝜋 ( 𝐷
2 ) 3
D is the diameter measured through Image J.
For 3D pa rticle p resentation, the softw are VGStu dio MAX 2.2 is emplo yed. Different thresh old
grey values are emplo y ed to best depict the particle shape.
The obtained radiographic raw data are onl y flat- field and dark-field corrected. To minimise the
dataset, we used Grouped Z Project to compress every 20 imag es into one image. The frame rate
of the movie is 20 fps.
The figure s shown in Fi g.7 represent states as a function of the following discharge or charge
times:
State

Time

State

Time

State

time

D01

0 h

a

0 h

a ’

0 h

D02

6 h

b

3.35 h

b ’

6.78 h

D03

9.9 h

c

4.85 h

c ’

8.25 h

D04

13.6 h

d

6.35 h

d ’

9.76 h

D05

17.4 h

e

7.85 h

e ’

11.28 h

D06

22.8 h

f

9.76 h

f ’

12.85 h

g

11.28 h

g ’

14.35 h

C01

23.4 h

h

12.8 h

h ’

15.85 h

C02

27.2 h

i

14.7 h

i’

18.12 h

C03

31 h

j

23.4 h

j ’

23.4 h

C04

32 h

k

26.9 h

k ’

25.4 h

l

25.2 h

l ’

25.9

m

26.3 h

m ’

26.3 h

n

26.8 h

n ’

26.8 h

o

27.2 h

o ’

27.2 h

2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

- 101 -

p

27.7 h

p ’

27.7 h

q

28.2 h

q ’

28.2 h

r

28.65 h

r ’

28.65 h

s

29.1 h

s ’

29.1 h

t

29.5 h

t ’

29.5 h

For the quantitative ana lysis of the influence of different t ypes of Si particle on the o btained
discharge/charg e capacity , we have chosen 78 particle s in the sa me slice in tomo -cell with
diameter lar ger than 50 µm , the locations of the chosen particles are sho wn in S I Fig. 6. It shoul d
be noted that, during the discharge/charge process, the l ocations of some Si particles
(electrochemic all y active through the first c ycle, electrochemicall y active during the first
lithiation but turn int o electrochemically in -active during the first delithiation process and
electroc hemicall y in-active particles throu gh the first cy cle) will change continuously due to
lithium insertion into/ ex traction from the S i electrode. W e onl y cho ose the particles that
experienced minor movement during the discharge/charge p rocess (they stay ed almost in the
same slice) and t he Si pa rticles that experienced significant movement (they were not in the same
slice) were out of consideration. It should be also noted that, we onl y choose the S i particles
whose diameter were larger than 50 µm due to the li mited resolution of the labor ator y X -ra y
source. Considerin g the limited resolution of our laboratory X - ra y sourc e, we have not taken Si
particles with diameter lower than 50 µ m into account. Due to the same reason, we cannot
quantify the lithi ation or delithiaiton sate of individual particles. That is to sa y , the individual
lithium storage and rel ease ability o f electrochemically active Si particles durin g the first
discharge and first char ge process cannot be calculated. Nevertheless, b y calculating the fraction
of Si particles of electrochemic all y in-active, electrochemicall y active and electroc hemical
deactivate d, we get a similar result by calculating the discharged/charged capacity .

2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

- 102 -

SI Fig . 1 Electrochem i cal chara cteriza tion of the coin c el l with th e same Si e lectro de : the firs t voltage-
capacity pr ofile. The mass of th e loaded Si was 4 m g and t he discha rge/ch arge cu rrent w as 0.075 A g -1 .
Note that th e Si com posi te was cast d irectly to the coin ce ll casing without using coppe r curren t collec tor.

SI Fig . 2. 3D ev olution of the elec trochem ically act ive Si part icle as shown i n Fig . 4. a) - e), during the first
discharge proce ss ; f)-j), dur ing the first charg e process ; the sca le bar is 50 µm l ong.
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

- 103 -

SI Fig . 3. 3D ev olution of the el ectrochem ical dea ctivation of S i partic les as show n in Fig . 5. a)-e), during
the first di scharge proc ess; f) - j) , dur ing the fi rst charg e proc ess; the sc ale ba r is 50 µm long.
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

- 104 -

SI Fig . 4. 3D ev olution of the el ectrochem ically inactive Si par ticles shown in Fig . 6. a) -e), dur ing the f irst
discharge proce ss ; f)-j), dur ing the first charg e process ; the sca le bar is 50 µm l ong.
2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

- 105 -

2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

- 106 -

SI Fig . 5. Evolution of S i partic les in a larg er area of the elec t rode: A ) -E), during the first discha rge
process; F)- J), during the first cha rge process ; green circle shows region s of electrochem ically activ e
particles; y ellow rect angula r shows reg ions of ele ctrochem ical ly deact ivated par ticles; red round rectangle
shows reg ions of inac tive part icles; b lue round rec tang le shows r egions o f p article displacem ent. The slice
location is 61 µm fr om the bottom st ainless screw c ur rent col lector in the pristine state.

SI Fig . 6. Locations of the chosen 78 Si part icles for q uantitativ e a nalysis. Red num bers denote the Si
particles th at are el ectrochem ical ly in-activ e through th e f irst cycle (from 1 to 10). Blue num bers deno te
the Si partic les that ex perience t he first lithiation but tu rn into elect rochem ica lly in-active during the f irst
delithiation (from 1 to 49). Green num bers deno te the Si pa rtic les tha t are elec trochem ical ly activ e
through the fir st cy cle (fro m 1 to 19).

2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

- 107 -

Caption f or the Sup plementar y Movie:

http://www .science direct.com /sci ence /artic le/pii/S 037877 5 3163051 09

In situ 2D radiogra phy of radio-cell duri ng the fir st cycle

Note that d uring the in situ radiogra phy m e asurem ent we have st opped the m icro
CT once t o cool dow n the X-ray source f or around 30 m i nutes.

2.4 Degradation of lithium ion batteries based on ~100 μ m-sized Si particles

- 108 -

2.5 Fracture be h avior of ~20 µm S i particles

Reprinted w ith perm iss ion from DOI : 1 0.1002/c elc.20 1600219. © 2016 WILEY-VCH Ver lag Gm bH &
Co. KGaA,Wein heim .

Synchrotron X-ray tomographic study of a Silicon elec trode before and after discharge and
the effect of cavities on particle fracturing

Lukas Zielke* a , Fu Sun b , He nning Markötter b , André Hilger b , Riko Mo roni a ,
Roland Zengerle a, c , Simon Thiele a, d , John Banhart b and Ingo Mank e b

Luka s Zielke, Riko Moro ni
a Laboratory for MEMS Applications, I MTEK Department of Microsystems Engineering,
University of Freiburg, Georg es-Koehler-Allee 103, 79110 Freiburg, Germany
E-mail: Luka s.Zielke@i mtek.de, [email protected]

Prof. Dr. John Banhart, Dr. I ngo Manke, Fu Sun, Dr. André Hilger, Dr. Henning Markötter
b Helmholtz Zentrum Berlin, Hahn-Meitner-Platz 1, 14109 Berlin, Germany
E-mail: banhart@he lmholt z-berlin.de, manke@helmholtz -berlin.de, fu.sun@he lmholtz -berlin.de,
hilger@ helmholtz -berlin.de, [email protected]

Prof. Dr. Roland Ze ngerle
a Laboratory for MEMS Applications, I MTEK Department of Microsystems Engineering,
University of Freiburg, Georg es-Koehler-Allee 103, 7911 0 Freiburg , German y
c Hahn-Sc hickard, Georges-Köhler-Allee 103, 79110 Freiburg, German y
Ze [email protected]

Dr. Simon Thiele
a Laboratory for MEMS Applications, I MTEK Department of Microsystems Engineering,
University of Freiburg, Georg es-Koehler-Allee 103, 79110 Freiburg, Germany
d FIT, University of Freiburg, Georges -Köhler-Allee 105, 79110 Freiburg, Germany,
[email protected]

Abstract
Silicon (Si) has been proposed as one of the most promising anode materials for nex t -g eneration
lithium ion batteries ( LIBs). However, uns atisfactory disch arge capacity/energ y densit y and
inevitable performance worsening prevent their commercialization. Herein, an in -d epth
investigation on the same Si composite electrode before and after the first discharge by
employing in sit u s y nch rotron X-r a y tomo graphy is presented. It is found that i) on the electrode
level, the Si particles located in the central part o f the electrode pre ferentially experience crack
2.5 Fracture behavior of ~20 μ m Si particles

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formation; ii) on the individual particle level, h eterogeneous electrochemical lithiation behaviour
is observed; iii) ca vities are formed during the electrode preparation and battery operation.
Moreover, the correlation between the electrochemical activities of Si particles and their
individual electrical cont act to t he electron condu cting netwo rk is investigated. For th e first time
it is quantified that S i particles will experience l ithiation only under the condition that at least
40% of th eir surface is electricall y connected. These nov el insights are possi ble explan ations for
low discharge capacit y /energ y of Si electrode LIBs, and would open new design principles and
opportunities for hig h-capacit y electrode materials for next-generation energy storage systems.

Abstract Graph ic

Keywords lithium ion batteries, si licon particles, lit hiation mechanisms, degradation mechanisms,
Synchr otron X-ra y tomog raph y

1. Introduction
Next-generation lithi um ion batteries (LIBs) wit h improved specifi c power and energy densit y
have been proposed for a variety o f demanding a ppli cations from electric vehicles to large-scale
grid stor age facilities. [1] Worldwide efforts are therefore underwa y to find novel ele ctrode
materials that will be considered as alternative options to replace the cur rently commercialized
cathode and anode ma terials. [2] On thi s search for high-capacit y anode materials, silicon (Si) has
been identified as one of the most promising candidates to substitute graphite [3] : I n contr ast to
graphite, where full y l ithiated graphite stores one Li -ion per six c arbon atom s through
intercalation, full y lithiated S i-anodes c an theoreticall y store four Li -ions per S i atom by
2.5 Fracture behavior of ~20 μ m Si particles

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chemical bonding. This l eads to larger theoretical capacities up to 3579 mAhg -1 , which is about
10 times larg er than that of graphite anodes, 372 mAh g -1 . [4 ] However, S i anodes ex hibit an
unsatisfactory discharge capacit y /energ y densit y when discharged using high currents. The y
additionally show irreversible capacit y fading during c ycling. Previous studi es, employing
various characterization tool s such as scanning electron microscope (SEM) [5] , transmission
electron mic roscope (TEM) [6] and X-ra y tomograph y [7] , have suggested that the inevitable
capacity deca y is associated intimately with the dramatic volume cha nge of ~300% during
cy cling, with preferential de/lithiation pathway along the [110] direction of crystalline Si. [8] In
addition, it was found that a two -phase boundar y between a cr ystalline Si core and an amorphous
Li -Si shell induces stress within the anode particles, resulting in significa nt par ticle fra cture and
pulverisation. [6,7,9]
Apart from boundar y-induced fracture mechanisms, other factors also co ntribute to the decrease
in energy d ensity and the irreversible c apacit y fading. The central iss ue is related to the
weake ning or even com plete loss of electrical c ontact betwee n Si particles and the conductive
network, consisting of carbon binder domain (CBD) and current collector. [1 0] I nsuffi cient contact
to the conductive network leads to partial electroche mical inactivation of these active materials,
resulting in a reduced utiliz ation of loaded material and unsatisf actory dis/char ge capacit y
retention. Unfortunatel y, ex tensive characterizations of the evolution of a complete Si based
electrode within an operational LIB hav e not been conducted due to i) the incompatibility
between conventional i nvestigation tool s and commercial LIBs [ 11] and ii) the inabilit y to
character ize ensembles of active particles contained within a realistic multi -particle ele ctrode. [ 11]
Further more, no ex perimental fra mework of inves tigating the lower energ y/capa cit y caused from
insufficient contact between active materials and CBD has been reported, although it shows that
the rate performance of a LIB electrode composit e can be markedl y improved by providing
alternative e lectron paths. [ 12]
Here, we report the imaging of the same Si based electrode in pristine and the firstl y discharged
state in three -dimensions using in situ sy nchrotron X -ray tomo graphy. Using hi gh flux es from
sync hrotron sources, [ 13] thi s characterization technique allows us to obtain large and
repre sentative datasets t o track the evolution of the entire Si elec trode on the electrode level
(millimetres) and on th e level o f individual particles (~5µm) sim ultaneousl y . Studi es on th e
de/lithiation of Si and other Li-alloying metals (e.g. Sn and Ge ) using X -ra y
tomogra ph y /microscop y we re conducted in the past [7,11,14 – 16] , in which volume
expansions/contractions, the relation betwee n X -ray absorption and lithi ation as well as fracturing
mechanisms were inves tigated. S pecificall y , T aiwo et al. su ggested th at mechanical stress
ge nerated f rom the signi ficant volume changes can de crease the contact of single Si particles to
the conducting network . [16]
By emplo y ing in-sit u s ynchrotron X -ra y tomography, we observ ed what Taiwo et al. su ggested
and ex panded thi s theor y b y a further quantitative analy sis. Other impor tant findings were i) on
the electrode level, the Si particles located in the central part of the electrode preferentiall y
experienced crack formation compared with the Si particles located in the peripheral region. And
ii) on the individual particle level, heterogeneous electrochemical li thiation behaviour was
2.5 Fracture behavior of ~20 μ m Si particles

- 111 -

observed compared with the widely emplo yed ma cro-homo geneuos model for battery sim ulation.
[17] Finally , we quantif ied the fracturing of Si particles and related particles that remain
unfrac tured with a loss of electrical/ionic c ontact to either the electron conductive network
(current collector or to the carbon binder domain (CBD)) or the Li ion co nductive elec trol y te.
Possible reasons for the contact loss have also been proposed. Th e curre nt stud y fundamentall y
expands the inhere nt explanations for low discharge capacity/energy o f Si electrode LI B s, and
would open new design principles and opportunit ies for high-capacity electrode materials for
next-generation energy storag e systems.

2. Results

2.1 Morphology of the P ristine Electrode and Electr ochemical Characterization

X-ray tomograph y is b ased on measurin g variations in X -ra y attenuation c oefficients in a rotating
sample and thr ee-dimensional (3D) reconstructions of samples with hi gh sp atial resolution can be
obtained b y using high-flux s y nchrotron X- ra y f acilities. [18] Since all phases present (Si particles,
CBD and cavit y) exhibited ex cellent contrast d ue to different X-ra y attenuation coefficients,
morphologica l changes o f particles, local CBD destruction and cavity form ation could be studi ed
in detail in 3D.
In the present stud y, a proof- of -c oncept battery , full y compatible with a s ync hrot ron X-ra y
tomogra ph y s etup ( Figure 1 a and b) and simult aneously representative of commercial LIBs, was
used. We focused on mo rphological chan ges in th e entire Si composite electrode before and after
the first disch arge. The ex perimental setup is shown in F igure 1a. Th e custom made c ell allows
imaging all inner compo nents within the proof-of -concept b attery ( Figure 1b). Before starting the
in -situ X-ra y tomograph y we conducted s canning electron microscope (SEM) characterization of
the prepared Si composite e lectrode surf ace for a reliable int erpretation of the X -ray tomograph y
data. An exemplary image is shown in Figure S1: I t can clearly be seen, that the S i particles are
well mixed with the CBD, which agree s with the tom ogra ph y dat a (Figure 1c, inset), wher e the
spherical Si particles can clearl y be discriminated from th e CBD filled pore sp ace. The nano
pores in the CBD cannot be resolved usin g s y nchrotron X -ra y tomograph y with a pix el size of
438 nm. [19,2 0] There fore, in the following CBD denotes a mix ture of small, electrol y t e filled
pores and carbon black particles glued toge ther with bi nder.
In a first step, we p erformed cyc lic voltammetry to ve rif y the re duction and oxidation
character istics of the Si particles ( Figure S2). The observed anodic/cathodi c peaks are in good
agree ment with those of previously repo rted silicon/carbon/binder (Si/C ) composites [2,21] .
Subsequently, a tomography in the pristine state (without an y c y cling) was made for reference. It
is shown in Figur e 1c: The batter y contains a Si/C composite electrode, a separator (24 µ m thick)
and a lithium metal electrode (as counter an d reference electrode). I t can clearl y be se en that the
curre ntl y investigated Si/C composite electrode consists of an ensemble o f active p articles hold
together by the e lectron conducting CBD, as in most realistic electrodes.
2.5 Fracture behavior of ~20 μ m Si particles

- 112 -

Figure 1 a) S chem atic ill ustration of experimental setup at the BAMline at the electron s torage ring
BESSY II in Berlin, Germ any. b) Battery cell holde r design. c) Cell with lithium metal (bottom ), separato r
(24 µm, centre) and a silicon/ca rbon/b inder electrode (top). The lithium f lux during di scharge is indicated
by arr ows. An exemplary part of the electrode is show n in the inset. The scale bar r epresent s 50 µm. d)
First discha rge curv e of the battery at 0.13 Ag -1 , where the sc hematic sk etch of an intact S i particl e
represents the pristin e state and the c racked Li /Si part icle represen ts the dis charg ed state.

After the tomo gra ph y in the pristine state, the battery was discharged at 0.13 Ag - 1 (~0.03 C )
based only on th e mass of the loaded Si. The co rresponding discharge curve is shown in Fi gure
1d. The obtained specific capacity w as around 1 700 mAhg -1 , which is lower than the theoretical
specific capacit y of Si (3 579 mAhg -1 ). In order to find explanations for th e low specific capacit y
of the Si electrode , anoth er tomography of the same electrode was conducted.
2.1 Morphology of the Discharge d Electrode
2.1.1 Heterogeneous Lithiation on the Electrode Level
The firstl y discha rged st ate tomograph y is show n in Figure 2 a, b and c. Compared with the
pristine state tomo graphy (Figure 1 c), it can be un ambiguousl y discerned that Si particles located
in the central part of the electrode prefere ntially undergo crac k formation, denoted by a blue
circle in F igure 2b.
2.5 Fracture behavior of ~20 μ m Si particles

- 113 -

Figure 2 a) Three- di m ensional representation of the discharg ed Si/C composi te electrode in th rough -plane
direction. The scale bar represents 50 µm . b) T op -v iew of the half-cropped reconstruc tion, showing t he
region where Si particles fract ured preferential ly (blue ci rcle). c) Zoom in on the border of the region with
preferred fracturing. T he scale bar represents 10 µm . d) Schem atic illustration showing t he 1D Li ions
flow paths in the inn er part of the Si/C com posite electr ode.

In contr ast, particles located in the peripheral regions are kept intact during the whole discharge
process (hereafter, we denote the fractured Si pa rticles as electrochemically active particles or
lithiated particles; the unfractured Si particles as electrochemical l y ina ctive particles or un-
lithiated particles). This finding was further checked b y employing two -di mensional particle siz e
distributions, shown in Figure S3. In the region with prefe rred fracturing (blue circle Figure 2b),
the particle size is s y stemati cally smaller, confirming the su ggestion from the visual analysis of
the tomographic dataset. The observed inhomo geneity of lithiation on the entire electrode level
corre sponds well with previous reports, [22] in which, Cai et al. observe d an si gnifica nt
inhomogene ous deterioration in a commerc ial LIB b y usin g in situ ne utron diffraction. The y
found that n ear the ed ges of the batte r y , both the graphite anod e and the spinel -bas ed ca thode
showed a decreased capacity while near the central are a, both electrode s functioned properl y .
They propos e some potential factors, such as the electrol yte solution loss, separator pore clogging
and non-uniform temper ature can contribute to this phenomenon. In our case, we p ropose that an
inhomogene ous Li ion flux resulting fr om an inhomogeneous press ure subjected on Si/C
composite electrode cause the inhom ogeneous lithi ation on the electrode l evel. As schematicall y
shown in Figure 2d, owning to a conv ex shape of the Si/C electrode, an in homogeneous pressure
(a lar ger pressure in the i nner pa rt of th e electrode compared with a smaller pressure at the outer
parts) c an be generated when sc rewing the S i/C electrode and L i electrod e together using the two
curre nt collector pins (Figure 1b). This is further evidenced b y the curvature of the separator
between the two electrodes (Figure 1c and 2a). Assuming a homo geneous and electricall y well
conducting CBD [23] , the lithiation process onl y occurs in reg ions with strongl y promoted Li
2.5 Fracture behavior of ~20 μ m Si particles

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conduction pathways. [24] Considering that th e fl ow of Li ions, driven b y the electrochemical
potential between the Li electrode and S i electrod e, is f aster in the cent ral part (sho rter dist ances)
than that of the outer part of S i/C electrode (longer distances), the currentl y obs erved well -
defined one dimensional (1D) Li ions pathwa y (indicated b y arrow in Figure 2d) a grees well with
previous reports. [24] Therefore it is reasonable t hat Si particles loc ated along the 1D Li ions
pathway experience significant cra ck formation (lithiation) while o thers, located o utsi de the
pathway, ar e kept intact. A direct detrimental effect of the inhomogeneous lithi ation on the
electrode level is that t he discharge current subjected on the complete batter y (0.13 A g -1 )
concentrates only on the central part of the S i/C electrode , resulting in i) high localized current
density and current hotspots that could induce fracture and accelerate capacit y fading [ 25] and ii)
significa nt under-utilization of active materials that could decrease the capacit y/energy d ensity.
[26] We therefore conc lude that it is crucial to design sophisticated elec trode architecture that
guarantees uniform Li ion flow pathway s over the whole electrode.
2.1.2 Particle Sizes and Volume Fractions
To further investi gate c auses leading to low capacit y /energy dens it y , a rectangular region of
interest (RO I) loc ated w ithin the 1D Li-ion flow pathwa y was chosen, as defined in Fi gure S3
and the re sults are sho wn in F igure 3 . First, we quantified the fracturing of p articles in this
particular ROI fo r which we calculated the p article size distribution for th e pristine and
discharged electrode. A s a r esult, we found a significant thinning o f the distribution after
discharging the batter y , which reflects a general fracturing of particles. The results agree well
with previous report. [1 ] In our case, the fracturing of particles le ads to a d ecreased mean particle
diameter from 7 µm to 3 µm (Fig ure 3a).
2.5 Fracture behavior of ~20 μ m Si particles

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Figure 3 a) Par t icle size distribution o f th e pr istine ( beige) and d i scharged electrode (blue). b) 3D
representation of a lar ge particle, pristine and fractured. c, d) Recon structed slices of the pristine electrode
(c) and the same position after discharg e (d) of a small part of the region of interest. T he arrows indicate
intact particle s with good contact t o the conducting ne twork (yellow) and bad contact (red). The scale bar
is 40 µm l ong. The white triangle is show n for o rientat ion.

The height of the first bar in the particle size distribution of the dischar ged electrode in Figure 3 a
indicates that the size of a certain am ount of fractured pa rticles is below th e estimated resolution
limit of ~1 µm (~2 times the pixel siz e of 438 nm). To confirm this, we calculated the Si volume
per el ectrode area. The calculation y i elded 31.5 µm³/ µm² in the pristine state and 24.0 µ m³/µm²
in the discharged state, which is a difference of 24%. This difference corresponds ver y well with
another group’s value of 25%, which implies a reasonable segmentation. [16] The difference
before and after the first discharg e indicates that approximately a quarter of the orig inal Si after
the first lithi ation has f ractured int o pieces smaller than 1 µm in diameter. In Figure 3b a particle
2.5 Fracture behavior of ~20 μ m Si particles

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before and after discharge is shown for illust ration, indicating the f racturing into multiple smaller
parts while kee ping the originall y spherical shape.

2.1.3 Size and Distribution of Cavities

Apart from particle fr acture, other ph enomena in the RO I are demonstrated in Figure 3c and d.
Here, the s ame in -plane region of the electrode is shown in the pristine (Fig ure 3c) and
discharged state (Figure 3d). The white trian gle shows that there is little in -plane dist ortion, even
though most of the particles exhibit strong fracturing. The limited volume expansion during the
first discharge process suggests limi ted lithium uptak e/storage ability of Si particles. [22]
Moreover, it can be observed that some Si particles still undergo no lithiation (no crack formation,
indicated b y y ellow arro ws in Fi gure 3d) even th ough th e y are located within the 1D Li ions flow
pathway, with CBD surrounding them. Actuall y , the observed inh omogeneous lithi ation
behaviour among individual particles is in agreement with previous reports that uncharged FePO 4
phases are oft en surrounded by charged LiFePO 4 phase s [ 27] and that the statistical measure of
local “spe ctroscopic” sta te -of-charge (SOC) state in an LI B is heterogeneous. [ 28] I t has been
suggested that in a r ealistic commercial LIB electrode, conditions for a homogeneous
electroc hemical reaction are hardl y satisfied sinc e an y fluctuations on the electrol y t e exposure,
electrical contact or cr y stal defe cts can result in an inhomogeneous reaction. [29] I n the present
study we quantitativel y investigated the correlation between the electrochemical activities of Si
particles and the influence of their contact to both the ionically conducti ng electrol y te and the
electrically conducting network.
Another worth y obs ervation is that particles in the dark regions in Figure 3d show no or weak
frac turing (indicated b y red arrows). As a m atter of fac t, due to the dif ferent X-ra y absorption
coeff icients of Si particle s, C BD and large por es, we could identif y th e d ark re gions as cavities in
the CBD. It is alread y known that gas evolution, which is intimatel y related with the solid
electrolyte interphase (SE I) for mation and the e lectrolyte decomposition, can locally d estro y
CBD and induce electrolyte displ acement. [ 30,31] Goers et al. have obs erved gas form ed channels
during cy cling within a LIB b y neutron radiography (NR). [32]
In the followin g, we ana lyse the cavities in 3D in terms of siz e, distributi on and their influ ences
on Si particle lithi ation. Results are shown in Figure 4 . First, we present the quantitative anal y sis
of the cavities. The mean diameter of the cavities within the complete Si/C electrode was found
to be ~5 µm. However , diameters of up to ~18 µm were measured as shown in the size
distribution in Figure 4a .
2.5 Fracture behavior of ~20 μ m Si particles

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Figure 4 a) Cavi ty si ze distribution in the discharged state. b) Fraction of cavity ar ea per el ectrode area in
the through- plane direction , st arting at the separator. c) Exem plary t omog ra m showing inhom ogeneously
distributed cav ities. One of them is shown in the inset. The scale bars represent 25 µm (i nset) and 200 µm.
d) T hree- dim ensional represent ation of the cavi ty network. T he largest connected cluster is shown in blu e
and the second larg est in re d. Rem ai ning cavity cluster s are shown in green .

We also found that cavities were mostly loc ated in the vicinity of the sep arator (Fi gure 4b) and
were not homo geneousl y dist ributed throu ghout the el ectrode as depicted in Figure 4c
(exemplary tom ogram) a nd Fig ure 4d (3D distributions of cavities). I t has to be noted that in
Fig ure 4d, the colours represent connected cavit y clusters: the largest connected cavit y comprises
63 vol. % o f the tot al c avit y volume (Figure 4d, bl ue) and spans the whole length o f the RO I; the
next largest cavity region comprises 17 vol. % (Figure 4d, red).
2.1.1 A Correlation between Partic le Fracture and Cavities
Further more a quantitative anal y sis of the influence of cavities on the fracturing of particles in
close contac t to them wa s conducted. The results are shown in Figure 5 .
2.5 Fracture behavior of ~20 μ m Si particles

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Figure 5 a) Relat ive volume of f ractures with in silicon particles versus connected surface area, calcula ted
using two different segm e ntatio n method s (diam onds and squares) ). b, c) Structurally i ntact and poorly
connected particles : b) within the large cavity network ( shown diameter was 16 µm ) and c) with localiz ed
CBD destruction around it ( shown diam eter was 9 µm). d) Partially connected parti cle with a f ew
fractures (~48% connected surface area, d iameter was 14 µm ). e, f) Well connected Si particle e) with
fractures asid e of the cavity network ( 90% connected surfa ce area, diam et er of 27 µm ) and f ) inside of the
cavity network ( 77% connected surf ace area and a diameter of 14 µm ). The white dotted line in a) is a
guide to the eye.

Four t y p es of particles were found: i) P articles well connected to the CBD th at exhibited lar ge
frac tures (Figure 5 e, f), ii) particles with only partial contac t to the C BD but with fr actures
(Figure 5 d), iii) particles with onl y partial contact to the CBD but witho ut frac tures (Figure 5 b )
and iv) particles almost full y surrounded b y a ca vit y (Figure 5 c). I t needs to be noted that the
cavity around the particles in iv) is not part of the large cavit y n etwork b ut is strong ly loc alized
around sin gle pa rticles. Actually, this kind of particle (t y pe iv) is representative of a class of Si
particles that often undergo no crack formation w ithin the 1D L i ions flo w pat hwa y (see Figure
3d). The reason for thes e stron gly localized cavi ties lies probably in the electrode p reparation
process, e.g. in the m aterial mixture and/or in the dr y in g process. [3 3] Further information on t ype
iv particles ca n be found in the supplementar y information.
For a repr esentative qu antification of the co rrelation between particle f racturing and contact loss
to both the electrical and ionic network, we investigated 21 diff erent particles in the vicinit y or
within cavities (both gas generated cavities and localized cavities). Relative volum e of fractures
within the particles, se rving as an indicator for li thiation, versus the electrically/ionically
2.5 Fracture behavior of ~20 μ m Si particles

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connected surface area of the particles is shown in Figure 5 a. We found that , above a threshold o f
40%, the degree of f ractures intensifies with increasing connected sur face area. Since the
determination of contact areas is highl y critical, all particles and c avities were segmented
predominantly m anuall y using two different methods (diamonds an d squar es ), which are
described in d etail in the Methods section. Despite the scatter of the data a clear trend is observed:
the relation between fracturing and electrical/ioni c contac t of S i particles in the battery shows a
threshold. That is, a connected surf ace area below 40% leads to no electrochemical activit y o f Si
particles. To the authors’ knowledge, this is the first investigation on th e c orrelation between th e
electroc hemical activities of Si particles and the influence of their contact area to the condu ctin g
electronic/ionic ne twork in t he battery.
From an electrochemical point of view, lithi um insertion into Si particles during the lithiation
process entails L i-ions (from the electroly te or L i electrode) and electrons (from CBD,
conductive ne twork or adjacent active particles) simult aneously . [ 12] The dependence of the
electroc hemical pe rformance of LI Bs on the contact b etween active materials and the
electrically /ionicall y conducting network is an important electrode desi gn consideration. [34]
Actually , it has been s uggested th at the electron transference between difference ph ases is
through hopping or tunn elling mechanism with contact resistivity depending exponentiall y on the
ga p between the contacts. A ga p of 8-10 nm is larg e enough to make a contact elec tricall y
insulating. [35] Thus, the cavities (either from elect rode preparation or battery ope ration)
surrounding Si particles will block the flow of electrons into the active particles during the
lithiation process to some extent, resulting in the ob served electrochemicall y in -active
(unfractured) p articles. After an in-depth quantitative anal y sis, it is discovere d for the first time
that, Si particles will experie nce lithiation onl y under the condition that at least 40% of it s surface
is electronica ll y/ionica ll y connected.

3. Conclusion
In summa r y , we perfor med an in -situ anal ysis of a S i/C electrode in its pristine and firstl y
discharged states. F irstly, we found that the Si particles in the centre of the electrode
preferentially under go crack forma tion while particles located in th e peripheral re gions are kep t
intact during the whol e fi rst discharge process. W e propose that the on e -dimensional L i ions flow
behaviour and a non-un iform contact pressure subjected on the Si/C elec trodes could be the
reason. Secondly, we revealed the presence o f large non -uniforml y distributed cavities (either
from elec trode preparation or batter y operation) within the discharged Si/C electrode s. We further
quantified their size and distribution within the electrode in th ree-dimensions. I n addition, an in -
depth quantitative analysis was conducted to investi ga te the dependence of the electrochemical
activities of active mate rials and their contact s urfaces to the electrically / ionicall y conducting
network. W e found that sili con particles experience fracture onl y when th eir contac t area to the
conducting network is at least 40% of their total s urface area . These results show that s ync hrotron
X-ray tomography is a powerful characterization tool to quantitatively investigate the dep endence
of the electrochemical pe rformance of LIB s on the electrical contact between active materials and
conducting network. The novel insights into possible causes for the lower capac it y/energy densit y
2.5 Fracture behavior of ~20 μ m Si particles

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of LI Bs are crucial because the y cannot be easil y obtained b y con ventional macroscopic
electroa nal y tical characterization techniques. Our findings provide potential guidelines to further
enhance the perf ormance of LI Bs. On the one hand, future ele ctrode engineering shoul d
concentrate on optimi zing the whole electrode architecture, which can involve all active material
particles ionically and electronicall y connecting t o electrolyte and electric conducting network to
maximize the utili zation of active materials. Particularly, an efficient electrica ll y con ductin g
CBD network and an ionicall y well- conducting pore space are hi ghly desirable to eliminate the
observed electrochemically in -active material. On the other hand, ou r analysis shows that it is
important to develop new electrolytes with reduced g as for mation to eliminate gas induced
cavities.

4. Experimental Section
Materials: Sil icon was rece ived from Elkem AS, Norway. Conductive carbon bl ack,
Pol y vin ylidene difluoride (PVDF) binder, Cel gard separator and lithium were purchased from
MTI C or. USA. N - me thy l p y rrolidone solvent (NMP) and 1M LiPF 6 in a volume-ratio mix ture
(1:1) of ethylene carbonate (EC) and dimeth y l carbonate (DMC) were purchased from Sigma
Aldrich. The housing o f the proof -of-concept bea mline ba ttery is made of polya mide -imide
(Torlon) fr om the McMaster-Carr compan y.
Battery Preparation: The electrode wa s made of electrode slurries with weight ratios of Si:carbon
black:binder of 70:20:10 in NMP. As binder, PVDF was us ed. Subsequentl y, the slurr y was cast
onto an alumini um foil . To remove the NMP , th e cast aluminium foils were dried in an oven at
60 ℃ overnight. After drying, the completed composite electrode was diced into smaller pieces of
around 1.7 mm  1.7 m m  0.2 mm (length  width  height) with a razor blade. Before the Si
composite was assembled into the batter y , it was weighed and the amount of Si particles was
determined from the o riginal mass ratio. The mass of the Si composite electrode was 0.17 mg.
The proof- of -concept batter y was assembled i n an argon-filled glovebox with humidit y and
ox y ge n levels below 0.1 ppm. Metallic lithium was placed on the top of a screw, actin g as a
counter and ref erence electrode. The pol y mer separator was placed betw een the li thium electrode
and the Si electrode. Finall y, the housing tube w as filled wi th the liquid electroly te. Current leads
were connec ted to a potentios tat for electrochemical tests.
Electrochemical Measurement: C y clic voltamme try (CV) and ga lvanostatic charge/discharge of
the battery w ere carried out with an I viumStat form I vium Te chnologies, Ne therlerlands. A
freshly assembled battery was measured from 2.5 V to 0 V at a scan rate of 1 mVs -1 to obtain CV
curves. Durin g galvanostatic c yc ling at the beamline, the assembled batte ry was measured at a
discharge current of 0.13 Ag -1 based only on the mass of Si active material.
Settings of Tomography Measurements: S ynchrotron X-ra y tomography was carried out at the
BAMline at BESSY II of the Helmholtz -Centre Berlin, Germany . The sy nchrotron beam was
monochromatized to 20 keV usin g a double multil ayer monoch romator with an energ y resolution
of about 1.5 %. The detector system comprised a 60 -µm thi ck C dWO 4 scintillator, a mi croscopic
optic and a pco4000 ca mera with a 4008×2672 pix el 2 CCD chip that is kept out of the direct
2.5 Fracture behavior of ~20 μ m Si particles

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beam b y using a mi rror. For tomograph y , 2200 projections durin g a 180° batter y rotation, each
with 4 s exposure time were recor ded before or after discharge.
Data Processing: Th e Si/ C/B domains were s egmented using th e software F IJI and appl ying a
median filter and Otsu’s me thod for binarization, followed b y a slight opening [36,37] . The cavity
domain wa s se gmented using the statisti cal region-merging tool and followe d by individual
thresholds for subsets. The y were chosen b y visu al judgement. Por e and particle size distribut ions
were calculated usin g the Delerue method, imple mented in the so ftware packages Geo Dict (3 D
global) and M atlab (3D local). The phase distributions (2D and 3D) were calculated usin g self -
programmed M atlab functions. Segmentation of contact areas was con ducted usin g manual
segmentation, and two diff erent a pproaches. Manual segmentation of particles and a global
threshold from visual ju dgement was us ed to segment the contact area b etween Si pa rticles and
CBD. A dilation of the particle followed, resulting in an overlap of particle and cavit y . To avoid
edge effects, the se gmented particle was e roded once and subsequentl y the thresholds for the
segmentation of fractures we re chosen from visual judgment (Method 1). Based on these datasets,
the particle-cavit y contact area w as manu ally co rrected in ev er y main direction, followed b y a
slight 3D median filter (Method 2).

Supporting Inform ation
Supporting Inf ormation is available from the Wiley Online Library or from the author.

Acknowledgements
We thank Dr. Heinrich Riesemeier, the beamline scientist at BESSY II, for his valuable
assistance, Anna Manzo ni for conducting scanning electron microscope ( SEM) charac terization
and Norbert Bec k for fabricating the b eamline batter y . W e also thank Elkem AS for providing u s
with the Si particle s. This work is sponsored b y the Helmholtz Association and the China
Scholarship Council.

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2.5 Fracture behavior of ~20 μ m Si particles

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2. 5.1 Suppor ting Inf ormation

Synchrotron X-ray tomographic study of a Silicon electrode before and after discharge and
the effect of cavities on particle fracturing

Lukas Zielke* a , Fu Sun b , Henning Markötter b , André Hilger b , Riko Mo roni a ,
Roland Zengerle a, c , Simon Thiele a, d , John Banhart b and Ingo Mank e b

Fig. S1 False color scann ing electron micrograph of t he mixed Si/Carbon/b i nde r composite elect rode in
the prist ine state. The rose particle s are the silicon particles and the pore -filling material is the carbon
binder m ixture. The sca le bar repr esents 50µm . T he imag e was scanned with a voltag e of 2kV.

2.5 Fracture behavior of ~20 μ m Si particles

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Fig. S2 T he cyclic voltamm etry ( CV) cur ves scanned at 1 mV s -1 in the potential window of 0 - 2.5V of t he
assembled p roof- of-concept beam li ne ba ttery

Fig. S3 a) Exem pl ary segmented in-plane im age showing particles of th e discharged elec t rode. The
particles are g ene rally larger in the outer region s ince they k ept their original sh ape (b) and f racture d to
smaller sizes i n the inne r reg ion (c). Scale bars in b an d c are 10 µm l ong.

On the search for sufficient morphological parameters determining if single particles with
strongly loca lized cavities around them fractured or not, we ca lculated characteristic sizes of
multiple particles (volume:surfa ce area). The calculated characteristic sizes were calcula ted in the
pristine electrode reconstruction and whether the particles cracked was checked in the discharged
reconstruc tion. The calculation i s shown in Fig. S4: the fracturing of those particles did not
depend solely on the c haracteristic size. Finding a parameter solel y deter m ining whether t ype iv
particles fracture or not bears a large potential for increasing the performance of Si composite
electrode s.

Fig. S4 I nfluence of ch arac teristic size ( volum e di vided by sur face are a) on wheth er the partic le with
strongly localized cavi ties around them fractured o r not.

2.5 Fracture behavior of ~20 μ m Si particles

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2.6 Diff erent (de)l ithiation behaviors of Sn par ticles

Reprinted with perm i ssion from DOI : 10.1002/cssc.2 01600220 . © 2016 WILEY- VCH V erlag GmbH &
Co. KGaA,W einheim .

In situ radiogra phic investiga tion of lithiation and delithiatio n
mechanis m s in a Sn-elec trode lithi u m ion battery

Fu Sun,* [a, b] Henning Mar kötter, [a, b] Dong Zhou, [a, b] Saad Sa be Sulaiman Alrwashdeh , [ a, b, c]
Andre Hilger, [b] Nikola y Kardjilov, [b] I ngo Manke [ b] and John Banhart [a, b]

[a] Fu Sun, Prof. Dr. John Banhart, Dr. Henning Markötter, Dong Z hou, Saad Sabe Sulaiman
Alrwashdeh
Institute of Ma terial Science and Technologies
Technica l Universit y Berlin
10623 Berlin, Germany
E-mail: [email protected]
[email protected]
[b] Fu Sun, Dr. He nning Markötter, Dong Z hou, Saad Sabe Sulaiman Alrwashdeh, Dr. Andre
Hilger , Dr. Nikola y Kardjilov, Dr. I ngo Manke and Prof. Dr. John Banhart
Helmholtz Centre Berlin for Mate rials and Energy
Hahn-Meitner-Platz 1
14109 Berlin, Germany
[c] Saad Sabe Alrwashdeh
Mechanical Enginee ring Department
Fac ult y of E ngineering, Mu'tah Universit y
P.O Box 7, Al-Karak 61710 Jordan

Abstract
The lithiation and delithiation mechanisms of multiple -Sn particles in a customized flat
radiography cell were investigated b y in situ synchr otron radio graphy. For the first time, four
hitherto unknown de/lith iation phenomena in a Sn-electrode battery s y stem are highlighted: 1
The de/lithiation behavior varies b etween different Sn particles ; 2 The t ime required to li thiate
individual Sn particle is markedly different fro m the time needed to discharge the comple te
battery ; 3 Electrochemical deactivation of originall y electrochemicall y active particles is reported;
4 A change of electrochemical behavior of ind ividual particles during cycling is found and
explained b y d yna mic changes of de/lithiation pathwa ys among st particl es within the electrode.
2.6 Different (de)lithiation behaviors of Sn particles

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These unexpected find ings fund amentaly expand the understandin g of the und erl y ing
de/lithiation mechanisms insi de commercial lithium ion batteries (LIBs) and would open new
design principles for high-performance next-generation LIBs.

Abstract Graph ic

Key words
lithium ion b attery; Sn part icles; in-situ; Syn chrotro n X-ray; rad iography

Lithium ion batteries (LIBs) are the dominant energy carrier in man y app lications rang in g from
portable electronics to h ybrid elec tric vehi cles. [1 - 4] The development of next-generation LIBs with
superior en ergy and po wer density and lon g-term c y cling stabilit y n ecessitates a fundamental
understanding of de/lithiation mecha nisms insi de batter y s y stems. Currentl y, dire ct visualization
into de/lit hiation process has been la rgel y provided by in situ transmission electron microscop y
(TEM) investi gations. [5, 6] However, Zhong et al. argue that the “end/point contact” architecture
of active m aterials used in the in situ TEM techni que ma y deviate fr om the commercial “floodin g
ge ometr y” contacts and the y find the multiple -stri pe lithiation mechanism in a flooding geometry
by investigating SnO 2 na nowires. [7] Furthermore, Gu et al. sugge st that the open-cell
config uration and the ionic/ L i 2 O electrolyte are also inherentl y diff erent from r eal commercial
batteries and the y have deve loped a sealed oper ando TEM electrochemical c ell and found that
the lithiation of Si na nowire immersed in the liquid electroly te progresses in the core -shell
fashion. [8] P revious work carr ied out based on the in situ TEM have provided insightful
2.6 Different (de)lithiation behaviors of Sn particles

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information int o the structural and chemical evolution of electrodes during d e/lithiation.
Neverthe less, one t ypical defic iency associated with the in situ TEM studies is that they are
li mited to e xplore the de/lithiation mechanism o n atomic or single nano wire/nanoparticle level
only, missing the interplay amongst multiple particles and the interactions between active
material and conductive/binder agent composite. Realisticall y , a commerc ial LIB electrode is an
assembly o f ac tive particles, org anic pol y meric binder and conducting age nt . [ 9] Thus, direct
observation of the de/lithiation process on th e multiple -particle scale can provide additionally
realistic insights into t he operation of L IBs and guide engineers designing th e electrode
architec ture in developing more advanced next- generation LI Bs.
Herein a cell suitable for in situ radiog raphy was b uilt and the d e/lithiation process of multiple-Sn
particles w as investi gated b y s y nchrotron X-ra y radiograph y. S ome d e/lithiation behaviors o f Sn
particles obtained from this in situ radiography cell are consistent wit h previous results . [8]
However, contrar y to the widespread belief t hat the lithium ions ar e proposed to diffuse
uniformly from the separator to the c urrent collector with th e formation of an electrode-l evel
de/lithiation front, [10] we find a number of ph enomena in the b attery s ystems that cast doubts on
this belief. These unex pected findings, which cannot be obtained b y the overall m acroscopic
electroc hemical characterizations and a single-parti cle de/lithiation model, fundamentally expand
our understanding of unde rly ing de /lithiation mechanisms in pra ctical c ommercial LI Bs and
would open new de sign principles and opportunities for hi gh -pe rformance next-generation LIBs.
The desi gn of the in sit u 2D radiography cell (radio-ce ll) and the schematic il lustration of the
sync hrotron setup are illustrated in Figure 1 a,c, along with a photograph of the radio -cell, Figure
1b. Figure S I of the S I (Supportin g Inf ormati on) shows a projection image of the Li
electrode /separator/Sn el ectrode assembl y within the cell obtained after the synchrotron X - ray
measurements b y usin g a laboratory micro X -ray source. [11] A weight-ra ti o of Sn:Carbon:Binder
of 60:30:10 was used to prevent possible particle overlapping in the radiographs. More details of
the cell and the measurement procedure are described in the experimental section and S I . After
assembling the radio-cell, cyclic voltammetr y (CV) wa s performed to verify the reduction and
oxidation cha racter istics of S n, as shown in Figure S2 . The cle arl y observed anodic/cathodic
peaks shown in S I Figure 2b are in good agreement with previousl y reported Sn/Carbon
composite LI B s. [1 2] I n a next step, the cell was dis/charged while conducting s y nchrotron X -ra y
radiography. Ty pical dis/charge curves shown in SI Figure 2a also a gree with previous results. [13]
2.6 Different (de)lithiation behaviors of Sn particles

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Figure 1. (a) Schematic illustration of the radiography cell used for i n situ synchrotron radiography . From
top to bo t tom : upper housing ( orang e), sealing ring (yellow), l ithium plate ( blue) with copper w ire,
separator (light g ray), Sn/c arbon/binde r com posite (green), titanium f oil current collector ( gray ), annula r
copper current collec tor (copper ), lower housing ( orang e); (b) Photog raph of t he radio -cell; (c) Schem atic
illustration of the experim ental setup at the BAMlin e, BESSY II, Hel mholtz -Zentrum Berlin, Germ any.
The radio- cell is ar ranged coaxia lly to the beam throug hout radiog raphy.

Synchr otron X-ra y imaging is an analytical technique that maps the X-ra y attenuation
coeff icients of samples. S ince cha racterizations can be conducted in situ and non-destructivel y, [14]
it is particularly well suited to track the morphological evolution of particles and change s of an
entire electrode as a f unction of discharge/charge state. The present ed s ync hrotron X -ra y
radiography wa s c onducted at the BAM line at B ESSY II , Helmholtz-Zentrum Berlin,
Germany. [15] A beam ener gy of 17 keV was chosen for optimal beam tr ansmission and image
contrast using a double mul tilayer monochromat or with an energ y resolu tion of about 1.5%. A
PCO4000 camera with a 4008×2672 pixel 2 C CD chip was used with optics resulting in a pix el
size of 0.438 µm and a field of view (FoV) of 1.7 × 1.2 mm 2 (width×heig ht). One radio graphic
image was acquired ev ery 60 s at a relativel y lo w X -Ra y flux (sync hrotron was operated in th e
single bunch mode). In order to obtain large and representative data of th e de/lithiation
2.6 Different (de)lithiation behaviors of Sn particles

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mechanism insi de the ra dio-cell during dis/ charge pro cess, consecutively three dif ferent regions
as marked region 1-3 in Figure S1 were characterized.
All normalized projec tions acquire d from the in situ radiog raphy are shown in the Supporting
Movies (SM) and Figure S3 display s an enl arged view of the pristi ne state of the consecutively
character ized thr ee r egions. Different de/lithiation behaviors inside the multi -Sn particles batter y
during dis/char ge pro cess are shown in Figure 2 . We track in detail structural changes of one Sn
particle (green dotted circle in Fi gure 2A at time 00:00) during de/lithiati on, representing the
class of electroactive pa rticles’ evolution as a fun ction of the dis/charge process. The location of
this particle is shown in Fig ure S3. As can be clearly seen from the 1 st row of Figure 2A ,
following the first lithiation (from left to right), the sharp contour of the electrochemicall y active
Sn particle becomes pr ogressivel y blurr y during which its volume i ncreases upon further
lithiation. Finally, a sudden disintegration involving multiple c racks and fragments occurs and the
particle resembles cauliflower-t ype morpholog y . During the 1 st delithiation process, as shown in
the 2 nd row in Figure 2A (from ri ght to left), the volum e decreases gradually. T he 3 rd row in
Fig ure 2A shows the 2 nd lithi ation process of this particle. I t is worth y to n ote that compared with
the 1 st lithiated state, there is a limited volume expansion during the 2 nd lithi ation, sugge sting a
limited lithium uptake capability a nd a resultant c apacity loss during the 2 nd discharge. Additional
evidence fo r L i insertio n and extraction into or from the Sn particle are the changes of the
histogram of transmission values in reg ions containing this particle a s shown in Figur e 3 A
(includes the 1 st discharge, 1 st char ge and 2 nd disc harge). The rightmost pe ak corresponds to the
weakly absorbing carbon, binder and electrol y te, the leftmost peak corresponds to the Sn particle s
and the changes to the middle peaks are dire ctly r elated to the electroc hemical evolution of Sn
particle. Following the 1 st li thiation process, the peak moves progressivel y from le ft to right,
impl y in g that the Sn particle is graduall y transforming from the high-d ensit y Sn phase to a low -
density Li X Sn ph ase (1< x<4.4). [16] During the 1 st delithiation process, we find that the p eak shifts
gradually towards the original direction. The present demonstrations that the de/lithiation -induced
volumetric expansion/contraction b y S n lattice dislocation and plastic deformation during
dis/charge process ar e in good agreement with previous reports. [ 7, 17] However, it is worth noting
that, neither the peak pos ition nor the peak shape of S n is restored to that of the ori ginal stat e b y
the end of the 1 st char ge p rocess. This ma y stem from the incomplete delithi ation and/or a
significa nt amount of lithium trapped in the dec omposed electrol y te byproducts . [18]
2.6 Different (de)lithiation behaviors of Sn particles

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Figure 2. Dem onstration of different de/lithiation behavior i n the radio - cell: Panel A displays one
electrochem ically active and one electroch em ically inactive Sn particl e (green and yellow circles,
respectiv ely, in t he figure labeled 00:00); Panel B shows that the t ime required to l ithiate indiv idual Sn
particle is different from the t ime needed to dischar g e the com plete battery; Panel C demonstrates the
electrochem ical deactivation of two Sn par ticles after t he first cycle; Panel D shows an unexpecte d
lithiation pathway, namely fracture ins tead of volum e expans ion occurs during first l ithiation ; Panel E
shows one Sn particle becom ing electrochemically activ e after the f ir st cycle incubation period; Panel F
features one Sn par ticle t hat remains electrochem i cally i nactive after t he f irst cycle incubation period. T he
locations of all the above particles are shown in Figure S3. In all figures, green rectangles mark the 1 st
2.6 Different (de)lithiation behaviors of Sn particles

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discharge, r ed r ectangle s 1 st charge, bl ue rectangles the 2 nd discharge. Green dots encircle particles in the
pristine state, red dots refer to the 1 st delithiated state, blue t o the 2 nd li thiated st ate. Arrows i ndicat e
progress in time. All the scale b ars shown are 50 µm long.

Surprisingly, we have also observed that m an y Sn particles nev er under go de/lithiation during the
dis/charge processes, that is to sa y , the y are electrochemicall y inactive during the whol e
macroscopic electrode level dis/charge process, as, for example, shown in Figure 2A whe re the
particle encircled in yellow does not show an y visible evolution. Actuall y, the observation of
electroc hemicall y in active Sn particles corresponds to the previous report b y Wang et al on
LiFeO 4 batteries, in which it wa s found that s ome pa rticles remain intact during the whole
cy cle. [1 9] I n another stud y on NiO electrodes by He et al, it was found that not all the NiO
nanosheets can be reduced even as the voltage dropped to zero against Li + /Li. [20] These recen t
investigations of the loca ll y microscopic inhomo geneous electrochemical reaction with respect to
the macroscopicall y and electrode level electrochemical performance suggest that the
electroc hemicall y driven reac tions are hardl y homogeneous in a commercial LI B du e to the
complex electrode structure and the complicated interactions . [21] The unexpected locall y
inhomogene ous d e/lithiation behavior amongst mul tiple -Sn particles could have two side effects.
On the one hand, the electrochemicall y inactive Sn particles are believ ed to substantiall y decrease
the energy densit y due to the inefficient utilization of loaded active materials . [ 22] On the other
hand, the inhomogeneous electrochemical reactio ns suggest that the overall dis/charge current is
heteroge neousl y distributed in the electrode, that is to say , onl y the ele ctrochemicall y active Sn
particles carr y the current . [9] More specifica ll y, we have c alculated the t ime required to lithiate
the elec trochemicall y acti ve Sn particle shown in Figure 2 B in the 1 st lithiation process to be
approximately 1 h. I n contrast, the time required to discharge the whole cell is approximately 15
h. The present ti me difference differs from a previous demonstration and the reason ma y stem
from the different rates subjected to the electrochemical ce lls. [ 9]
2.6 Different (de)lithiation behaviors of Sn particles

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