Journal of Energy Chemistry - Xiamen University · rector of Research Institute of Electrochemistry...

Journal of Energy Chemistry 27 (2018) 1566–1583

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Journal of Energy Chemistry

journal homepage: www.elsevier.com/locate/jechem

Review

The application of synchrotron X-ray techniques to the study of

rechargeable batteries

Zhengliang Gong

a , Yong Yang

a , b , ∗

a College of Energy, Xiamen University, Xiamen 361005, Fujian, China b State Key Laboratory for Physical Chemistry of Solid Surface, Department of Chemistry, College of Chemistry and Chemical Engineering, Xiamen University,

Xiamen 361005, Fujian, China

a r t i c l e i n f o

Article history:

Received 29 December 2017

Revised 20 March 2018

Accepted 29 March 2018

Available online 4 April 2018

Keywords:

Rechargeable battery

Synchrotron X-ray techniques

X-ray diffraction

X-ray absorption spectroscopy

Pair Distribution Function

X-ray photoelectron spectroscopy

a b s t r a c t

The increased use of rechargeable batteries in portable electronic devices and the continuous develop-

ment of novel applications (e.g. transportation and large scale energy storage), have raised a strong de-

mand for high performance batteries with increased energy density, cycle and calendar life, safety and

lower costs. This triggers significant efforts to reveal the fundamental mechanism determining battery

performance with the use of advanced analytical techniques. However, the inherently complex character-

istics of battery systems make the mechanism analysis sophisticated and difficult. Synchrotron radiation

is an advanced collimated light source with high intensity and tunable energies. It has particular ad-

vantages in electronic structure and geometric structure (both the short-range and long-range structure)

analysis of materials on different length and time scales. In the past decades, synchrotron X-ray tech-

niques have been widely used to understand the fundamental mechanism and guide the technological

optimization of batteries. In particular, in situ and operando techniques with high spatial and temporal

resolution, enable the nondestructive, real time dynamic investigation of the electrochemical reaction,

and lead to significant deep insights into the battery operation mechanism.

This review gives a brief introduction of the application of synchrotron X-ray techniques to the inves-

tigation of battery systems. The five widely implicated techniques, including X-ray diffraction (XRD), Pair

Distribution Function (PDF), Hard and Soft X-ray absorption spectroscopy (XAS) and X-ray photoelectron

spectroscopy (XPS) will be reviewed, with the emphasis on their in situ studies of battery systems during

cycling.

© 2018 Published by Elsevier B.V. and Science Press on behalf of Science Press and Dalian Institute of

Chemical Physics, Chinese Academy of Sciences

Zhengliang Gong received his Ph.D. in physical chemistry from Xiamen University in 2007. After a post-doctoral fel-

lowship at National University of Singapore, in 2010, he has been working at Xiamen University. Currently he is

an associate professor at Xiamen University. His main re-

search interests are materials for rechargeable batteries and the electrochemical processes in these systems.

∗ Corresponding author at: College of Energy and College of Chemistry and Chem-

ical Engineering, Xiamen University, Xiamen 361005, Fujian, China.

E-mail address: [email protected] (Y. Yang).

Yong Yang obtained his Ph.D. in Physical Chemistry from Xiamen University in1992. Except for a one-year

(1997,1998) academic visit at Oxford University, he has been working in the State Key lab for Physical Chem-

istry of Solid Surface at Xiamen University since 1992.

Now he is a distinguished professor in Chemistry and Di- rector of Research Institute of Electrochemistry and Elec-

trochemical Engineering over there. His main research interests are new electrode/electrolyte materials for

Li/Na-ion batteries, in-situ spectroscopic techniques, and interfacial reaction mechanism study in electrochemical

energy storage and conversion system. He has published over 200 papers in many international journals such as

Nature Energy, Energy & Environmental Science, Advanced Materials and Chem.

Mater., etc with citation > 60 0 0 (H-index = 41). He is also one of the Editors for Journal of Power Sources (IF = 6.3). He has obtained several national/international

research awards, e.g. Contribution award given by Chinese Electrochemical Society in 2017, Technology Award given by IBA (International Battery Materials Associa-

tion) in 2014.

https://doi.org/10.1016/j.jechem.2018.03.020

2095-4956/© 2018 Published by Elsevier B.V. and Science Press on behalf of Science Press
and Dalian Institute of Chemical Physics, Chinese Academy of Sciences

http://www.ScienceDirect.com

http://www.elsevier.com/locate/jechem

mailto:[email protected]


Z. Gong, Y. Yang / Journal of Energy Chemistry 27 (2018) 1566–1583 1567

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A

. Introduction

The use of rechargeable batteries (mainly lithium and sodium

ased rechargeable batteries) to electrical vehicles and grid energy

torage in recent years calls for high energy density, greater bat-

ery cycle life and safety characteristics [1–3] . To meet the above

equirements, significant efforts are being focused on materials

odification, developing promising new materials and new chem-

stry for the next generation of rechargeable batteries. The per-

ormance (energy density, cycle/calendar life and safety) and cost

f rechargeable batteries is directly linked to electrode materials

composition, structure and morphology, etc.) and their structural

volution during cycling. It is very important to better understand-

ng various composition-structure-performance relationships and

he electrochemical reaction mechanisms for battery systems in or-

er to improve their performance and guide the development and

pplication of high-performance new materials [4,5] . However, the

attery systems are inherently complex and difficult to understand.

his requires systematical and deep investigations from atomic

evel to macro level in electron structures, crystal structures, mi-

rostructures and morphologies, chemical compositions and phys-

cal properties of battery materials and their evolution during the

harge–discharge processes. Advanced ex situ and in situ character-

zation techniques have been used widely to clarify scientific and

echnological problems in rechargeable batteries.

Since the early work by Mcbreen et al. around 25 years ago

6] , synchrotron X-ray techniques have been widely used to un-

erstand the fundamental mechanism and guide the technologi-

al optimization of batteries [7–11] . The merits of synchrotron X-

ig. 1. (a) Common synchrotron X-ray techniques and their applications in battery res

iagram of the electrochemical in situ XAS. Adapted from Ref. [90] and [120] with perm

merican Chemical Society.

ays, such as high brightness (10 5 –10 12 more intense than that

rom the laboratory sources), highly collimated and energy tun-

ble, align itself perfectly to applications in battery science and

echnology. The tunability of X-ray energies allows the conduct of

xperiments which require a scan of the beam energy (e.g. XAS).

t also allows the optimization of the experiments to improve the

uality of data via optimizing the beam energy (e.g., eliminating

uorescence artifacts during XRD analysis). The highly-collimated

nd variable focus synchrotron beam allows high spatial resolution

pectroscopic mapping (point-by-point measurement) and imag-

ng of electrodes. With increasing sophisticated synchrotron tech-

iques, it also allows the conduct of high temporal resolution (up

o milliseconds) studies of the electrochemical/chemical reaction.

n particular, the ultrahigh intense and penetration ability of syn-

hrotron X-rays make the in situ and operando investigation of bat-

ery systems possible and easier to realize. Compared to ex situ

easurements, in situ and operando techniques can provide direct

nformation on the system in a nonequilibrium state, allowing a

ruer visualization of what is happening during the electrochemi-

al reduction and oxidation processes. Fig. 1 illustrates the appli-

ations of common synchrotron X-ray techniques to the study of

echargeable batteries.

This paper reviews the latest developments in the application

f synchrotron X-ray techniques, especially their electrochemical in

itu techniques to the studies of battery materials and related elec-

rochemical reaction mechanisms. Since a comprehensive review

f the field has recently been provided by Lin et al. [10] , this pa-

er will focus on the unique advantages and achievements of syn-

hrotron X-ray techniques for effectively providing structural and

earches, (b) schematic diagram of the electrochemical in situ XRD, (c) schematic

ission from The Royal Society of Chemistry, and Ref. [153] with permission from

1568 Z. Gong, Y. Yang / Journal of Energy Chemistry 27 (2018) 1566–1583

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chemical evolution information about commercially and/or funda-

mentally important electrode materials and their interfaces, and

for addressing key challenges in developing breakthroughs with

rechargeable batteries.

2. Synchrotron X-ray and designing in situ electrochemical

cells

Synchrotron X-ray was first observed at synchrotrons, which are

large and complex machines built for high-energy physics experi-

ments. Initially synchrotron X-rays were regard as a nuisance by

particle physicists, because they steal energy from their particle

beams, but in the 1960s they have been found to have extremely

beneficial properties for X-ray experiments [12] . Synchrotron radia-

tion can produce extremely bright X-rays, which is many orders of

magnitude greater than X-ray tubes. The high-particle energies of

third-generation high-energy synchrotron radiation allow for un-

dulator X-rays into the high-energy X-ray range (100 keV). The ex-

tremely bright and high-energy range of synchrotron X-ray bring

many advantages, such as easy penetration into sample environ-

ments necessary for in situ studies in batteries. Synchrotron X-

ray research in batteries, especially in situ techniques have evolved

tremendously over the last decade.

Compared with conventional ex situ approaches, in situ char-

acterization of battery materials during the operation of batteries

provides significant advantages, such as (1) can effectively avoid

the potential of short circuits during cell disassembly, (2) elim-

inate spontaneous relaxation or contamination of highly reactive

and transient species, also contamination/side reaction with atmo-

spheric species (O 2 , CO 2 , H 2 O), (3) makes the identifying of short-

lived intermediates and subtle nonlinearities in behavior possible

via efficiently probing battery materials at fine reaction intervals,

(4) allow time-resolved investigation of electrochemical reaction,

(5) improve the consistency, precision and reliability of measure-

ments by eliminating sample variance come from the use of mul-

tiple independent samples, (6) also reduce the effort of assembling

and disassembling of multiple cells [7,13] .

In situ cell design is of great importance for electrochemical

in situ study, which is one of the key factors affecting the qual-

ity of experimental data. In situ electrochemical cell should min-

imize the X-ray adsorption of no active components and window

material while maintaining reliable electrochemical performance.

Currently, several electrochemical cell designs have been reported

for electrochemical in situ XRD or XAS studies, including modi-

fied coin cell, ‘coffee bags’, pouch cells, Swagelok type cells, and

other self designed cells [14–23] . For example, Balasubramanian et

al. of Brookhaven National Laboratory design an in situ cell which

housed between two machined blocks of aluminum with windows

for X-ray penetration and holes for bolts. It uses sheets of 250 μm

Mylar as windows and a rubber gasket to make a hermetic seal

[16] . Borkiewicz et al. of Argonne National Laboratory design a

multi-purpose in situ X-ray (AMPIX) cell, which consists of a cup-

shaped body, bottom and top two electrodes, X-ray transmissive

windows ( × 2), and a flat annular gasket sandwiched between

the electrodes [24] . The AMPIX cell is suitable for a broad range

of synchrotron-based X-ray scattering and spectroscopic measure-

ments. Transmission synchrotron XRD and XAS geometry is widely

employed due to convenient setup.

While synchrotron X-ray techniques show unique advantages

for the study of rechargeable batteries, its main limitation lies in

the availability and access. Since the synchrotron facilities around

the world are limited and precious, researchers usually must ap-

ply for time to use a technique (normally limited time is avail-

able for any single user), and travel to a remote lab to carry out

their experiments. Another limitation of synchrotron X-ray tech-

niques is the difficulty to distinguish elements with low X-ray

ttenuation coefficients, such as lithium which is the most im-

ortant element in rechargeable lithium batteries, because X-rays

nteract with matter through electromagnetic interactions with the

lectron cloud of atoms, thus the interactions of X-rays with low-Z

lements are weaker than that of high-Z elements. Moreover, pos-

ible degradation results from the exposure to extremely high in-

ense synchrotron X-rays should also be carefully considered.

. The use of synchrotron X-ray techniques to recharge

atteries

.1. The application of synchrotron X-ray diffraction (XRD)

.1.1. Electrochemical in situ characterization of phase transformation

In the electrode materials, the insertion/extraction of guest ions

s often accompanied by phase transformations during cycling, in-

luding single-phase solid solution transformations, two-phase re-

ctions, order–disorder transitions and crystallographic changes.

he nature of phase transitions during charge–discharge processes

as significant effects on electrode performance in battery appli-

ations. Understanding the phase transition of electrode materials

olds promising keys to further improve the performance of bat-

eries, which is also critical to address the deformation and stress

ssues encountered during cycling [25] . Synchrotron based in situ

-ray diffraction is one of the most powerful tools to study the

tructural evolutions and structure-function relationship of elec-

rode materials during cycling.

Layered oxides for lithium ion batteries :

Layered transition metal oxides, LiCoO 2 and Li(Ni x Co y Mn z )O 2

NCM, with x + y + z = 1, defined as NCM xyz , such as x:y : z = 3:3:3

NCM333), 5:2:3 (NCM523), 8:1:1 (NCM811)) have been exten-

ively studied as important cathode materials for Lithium ion bat-

eries (LIBs). In order to meet the growing demand for higher

nergy density both in portable electronic devices and electric ve-

icles (EVs), significant research efforts have been made to im-

rove the specific capacity of layered oxide cathodes. There are

wo main approaches: (1) increase the charge cutoff voltage (such

s 4.55 V for LiCoO 2 ) to extract more Li; and (2) develop Ni-

ich NCM materials (i.e., x > 0.6, such as LiNi 0.8 Co 0.1 Mn 0.1 O 2 and

iNi 0.15 Co 0.80 Al 0.15 O 2 ), since the capacity of layered NCM materials

ncrease linearly with an increase in Ni content. Both approaches

an increase the capacity of layered oxides from 140 mAh/g to

round 200 mAh/g. However, increasing the charge cutoff voltage

nd/or increasing the Ni content lead to severe capacity fading and

afety concerns due to the accelerated material degradation in-

uced by a phase transition, decomposition of the electrolyte, and

ransition metal dissolution at the high potential [26,27] . Thus, the

hase transition and structural stability of the layered oxide mate-

ials are of great importance both fundamentally and technologi-

ally.

The two-phase and single-phase reactions accompanying with

hase transition from H1 to O1 via H2 and O1a have been

erified by many in situ and ex situ XRD investigations when

harging LiCoO 2 to higher voltage ( > 4.8 V) [28–31] . The mech-

nism of capacity fading of LiCoO 2 and the effects of Al 2 O 3

r ZrO 2 coating when charged to higher voltage ( > 4.5 V) have

een studies by Yang et al. using synchrotron based in situ XRD

32–34] . Similar structural change behavior was found for the

rO 2 -coated and Al 2 O 3 -coated LiCoO 2 . It shows that the bulk

tructural damage is minimal and the capacity fading of LiCoO 2

s mainly attributed to the increased polarization resulted from

he impedance increase due to electrode/electrolyte side reac-

ion when cycled with a higher voltage limit. It was proposed

hat the effect of ZrO 2 /Al 2 O 3 -coating layer on improving the

apacity retention during high voltage cycling is on the protec-

ion of the surface of LiCoO particles and reducing the elec-
2


Fig. 2. (a) Evolution of selected peaks from operando XRD data recorded on NCA during the first cycle, (b) a and c lattice parameters obtained from Rietveld refinements,

and (c) galvanostatic charge–discharge (4.1–2.7 V, C/15, RT) curves. The marker size of the lattice parameters determined for the NCA sample stored under ambient conditions

is dependent on the phase fraction obtained from the refinement. The estimated standard deviations are within the data markers. Reprinted from Ref. [40] with permission

of American Chemical Society.

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rode/electrolyte side reaction, thus minimize the impedance in-

rease. Unlike the LiNiO 2 or LiCoO 2 systems, no new H2 phase

as observed when LiNi 0.6 Co 0.25 Mn 0.15 O 2 cycled between 3 to

.5 V, a single-phase reaction dominated the lithiation/delithiation

eaction with both the a and c axes varying with different Li

ontents [35] . Compared with LiNiO 2 , LiCoO 2 , and LiMnO 2 elec-

rode materials, a much smaller volumetric change (4%) was ob-

erved after the first charge, which indicates a better structure

tability. However, a structural phase transition from H1 to H2

exagonal phase from the beginning of charge is also reported

or number of layered oxides including LiNi 0.15 Co 0.80 Al 0.15 O 2 (NCA)

36,37] , LiMn 0.5 Ni 0.5 O 2 , [38] L iNi 1/3 Mn 1/3 Co 1/3 O 2 , [37,39] etc. Yoon

t al. investigated the structural changes of LiCo 1/3 Ni 1/3 Mn 1/3 O 2

nd LiNi 0.8 Co 0.15 Al 0.05 O 2 during first charge to 5.2 V [37] . Similar

ith LiNiO 2 , a coexistence of H1 and H2 hexagonal phase from

he very early state of charge and the formation of H3 hexago-

al phase at the end of charge are observed, accompanying with

contraction along the a and b axis and a simultaneous ex-

ansion along the c -axis during the early stage of charge and a

light expansion along the a - and b -axis and a major contrac-

ion along the c -axis near the end of charge. A structural stabil-

ty of Li 1- x Co 1/3 Ni 1/3 Mn 1/ 3 O 2 > Li 1- x Ni 0.8 Co 0.15 Al 0. 05 O 2 > LiNiO 2 at

vercharged state is proposed according to the significantly sup-

ressed H3 phase formation with collapsed c -axis at the end

f charge. Robert et al . observed different lithiation/delithiation

eaction pathways for NCA between the first and second cy-

le/following cycles [36] . An irreversible two-phase transition plus

reversible solid solution reaction process was revealed for the

rst charge. While a solid solution process dominated during the

econd charge. They ascribe the irreversible two-phase transi-

ion to an activation of the electrode, which involving irreversible

hanges in NCA’s electrical conductivity. Recently, Grenier et al .

40] noticed the widely varying Li x compositional ranges in which

1 and H2 coexist for NCA and suggested that the earlier observed

two-phase” behavior during the early stage of charge may not be

ntrinsic to the NCA material. Through the studies of identically

repared and cycled electrodes of NCA subjected to different envi-

onments by synchrotron based in situ XRD, they demonstrate that

he intrinsic reaction mechanism for NCA during the first charge is

olid-solution rather than two-phase as proposed previously. The

reviously observed “two-phase” behavior is aroused form reac-

ion heterogeneity between secondary electrode particles induced

y nonuniform erosion of a Li 2 CO 3 surface layer ( Fig. 2 ).

Polyanionic compounds for Li-ion batteries :

Polyanionic compounds have attracted extensive interest since

he first report on the electrochemical performance of LiFePO 4

y Padhi et al . , due to their high safety, long cycle life, low cost

nd environment friendly [41,42] . Among them, LiFePO 4 has be-

ome a commercially important cathode material for LIBs, with

ide applications in transportation, power tools and large scale en-

rgy storage. Excellent rate capability can be obtained for LiFePO 4

hen carbon coated and nanosized, which is contrary to its ini-

ially reported first order two phase delithiation/lithiatioin reac-

ion mechanism, a kinetically limited process involving a volume


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change of 6.8% [41,43] . This contradictory has stimulated numer-

ous both theoretical and experimental studies that attempt to ex-

plain the charge–discharge mechanism in LiFePO 4 [43–55] . Impor-

tant new insights have been revealed with the use of advanced

synchrotron based in situ XRD. These demonstrated a very compli-

cated and fascinating phase diagram with both solid solution and

two phase transition for LiFePO 4 , which depends on both parti-

cle size and charge–discharge current rate. A first-order two phase

transition with very limited Li solubility in the two end-member

phases (lithium-rich triphilyte Li 1 −αFePO 4 and lithium-poor het-

erosite Li βFePO 4 ) is observed at or near to thermodynamic equi-

librium [47,56–58] . Thermodynamic Li-miscibility gap results in a

large lattice misfit of 6.8 vol % between LiFePO 4 and FePO 4 end

members [59,60] . The miscibility gap depends on particle size

and doping [61–65] . Nonequilibrium metastable Li x FePO 4 phases

were observed at intermediate charge–discharge current rates

[43–45,48] . High current rates and high overpotential are sug-

gested to promote the formation of intermediate phase. A dramat-

ically increase in the solubility limits in both phase is observed at

high rates (10 up to 60 C), with a major portion of the electrode

exhibited diffraction of intermediate phases at 60 C ( Fig. 3 ) [47] .

Asymmetric phase transition behavior between LiFePO 4 and

FePO 4 are also observed at high rates charge–discharging

[52,53,66] . This explains the irreversible charge–discharge prop-

erty at high rate cycling. The phase transformation from heterosite

Li βFePO 4 to triphylite Li 1 −αFePO 4 was found to be less likely to

proceed during discharge compared with the previous charge. This

results in a decrease in the discharge capacity and the increase ir-

reversible capacity between charge and discharge at high-rate con-

ditions. Using in situ energy-dispersive X-ray diffraction (EDXRD),

Paxton et al . investigated the heterogeneous reaction behavior of

LiFePO 4 electrode of an actual commercial battery [50,67] . Inho-

mogeneous behavior was observed with severely delayed discharge

occurred at part of the measured locations, suggesting special care

should be taken when comparing microscopic and macroscopic

measurements. Using in situ microbeam XRD, Zhang et al . investi-

gated the phase transformation of individual LiFePO 4 particles in-

side an electrode at different current rates [55] . They revealed a

current rate dependent phase transition mechanism within indi-

vidual electrode grains. Very slow and concurrent individual elec-

trode grains transformation via coexisting platelet-shaped phase

domains was observed at low current rates, which is contrary

to commonly assumed fast particle-by-particle or mosaic trans-

formation. While phase boundaries become diffuse speeding up

the transformation rates of individual grains was observed at high

current rates. This proves a distribution of instable intermediate

phases within a single electrode particle ( Fig. 4 ).

Ravnsbæk et al. systematically investigated the dynamic trans-

formation strain behavior in LiMn y Fe 1–y PO 4 (with average particles

size ∼50 nm) system using synchrotron based in situ XRD [49,54] .

It shows a strong correlation between the dynamic transforma-

tion strain and y in LiMn y Fe 1–y PO 4 system. A significant deviation

from what is expected from the equilibrium phase diagrams is ob-

served for the phase evolution sequence of all high rate compo-

sitions (0 < y < 0.8). They found that the capacity realized at high

C-rates from 20 to 50 C varies in inverse proportion to the mea-

sured transformation strain. This demonstrates the importance of

minimizing transformation strain for maximizing power capability

of batteries.

Layered oxides for Na-ion batteries :

S odium-ion batteries (SIBs) have attracted increasing attention

as a favorable alternative to LIBs particularly in large-scale elec-

tric energy storage, due to the abundant sodium resources and its

consequent cost advantages. Three main classes of cathode ma-

terials: layered oxides, complex polyanion frameworks and prus-

sian blue are widely explored for NIBs. Layered transition metal

xides Na x MO 2 (M is a first-row transition metal or mixture of

–3 elements) are extensively investigated due to their relatively

igh reversible capacity and open 2D diffusion paths for the larger

a ions. However, Na x MO 2 normally exhibits poor cycling stability,

hich hinders their commercial application. Compared with their

i-counterparts, Na x MO 2 exhibits rich and complex phase transi-

ions during electrochemical (de)intercalation marked by multiple

nterslab ordering arrangements of Na + ions located in octahe-

ral and prismatic coordination environments or MO 2 slab glid-

ng [68–72] . These phases are typically classified into P2, P3 and

3 type, according to Delmas’ notation [73] . The O or P nomi-

ation refers to Na + occupy octahedral or prismatic sites, while

he numerical nominations indicate the repeat period of the tran-

ition metal stacking perpendicular to Na layers. To understand-

ng the capacity fade mechanisms, the structural phase transition

ehaviors of Na x MO 2 were investigated by synchrotron based in

itu XRD. It shows that the large cell-volume changes associated

ith complex multiple phase transitions and irreversible phase

ransitions are likely to be responsible for the poor cycling sta-

ility [71,74–76] . Both solid-solution and two-phase reactions are

nvolved in the (de)sodiation process. For a majority of O3-type

aMO 2 (M = Cr, Mn, Co, Ni), complex phase-evolution is observed,

nvolving transitions from O3 to O’3 type structure in the early

tage of charge due to a distortion of the oxygen lattice induced

y a small shift of MO 2 layers, and to P3 or P’3 type structure

ith further extracting of Na + due to the gliding of the MO 2 slabs

77–79] . Boisse et al. [70] investigated the phase-evolution of P2-

ype Na 2/3 Fe 1/2 Mn 1/2 O 2 . The reaction from the open-circuit voltage

an be described as a two-phase P’2-to-P2 reaction in the early

harge step followed by a solid-solution reaction of the P2 phase

0.35 < x < 0.82), and a two-phase P2-to-“Z” reaction to the end of

harge. Pang et al. [71] revealed a multiple phase transitions re-

ction process for P2-type Na 2/3 (Fe 1/2 Mn 1/2 )O 2 , which involve the

ransitions from P 63/ mmc (P2-type at the open-circuit voltage) to

63 (OP4-type when fully charged) to P 63/ mmc (P2-type at 3.4–

.0 V) to Cmcm (P2-type at 2.0–1.5 V) symmetry structures during

esodiation/sodiation. The diverse phase transition behaviors be-

ween Na x MO 2 materials indicate that the composition and initial

tructure have significant effects on the structural evolutions and

lectrochemical properties. Considerable efforts have been devoted

o tailor both the composition and initial structure of Na x MO 2 ,

hus to tune the structural evolution and electrochemical perfor-

ance [74,80–82] . Dose et al. [82] reported a single-phase reaction

ominated desodiation/sodiation process with only a short, subtle

exagonal P2 to hexagonal P2 two-phase region early in the first

harge at C/10 for a Mn-rich P2-phase Na 2/3 Fe 0.2 Mn 0.8 O 2 cathode.

u et al. [76] investigated the effects of zinc doping on struc-

ural phase transition of P2-Type Na 0.66 Ni 0.33 −x Zn x Mn 0.67 O 2 cath-

des. Undoped Na 0.66 Ni 0.33 Mn 0.67 O 2 exhibits an irreversible P2-to-

’2 phase transition during cycling, resulting in a rapid capac-

ty and voltage fade. While Zn

2 + doped Na 0.66 Ni 0.26 Zn 0.07 Mn 0.67 O 2

hows more reversible P2 to O2 phase transition, thus better ca-

acity and voltage retention. Introduction of a small amount of

i + into NaNi 0.5 Mn 0.5 O 2 has been proven to be an effective ap-

roach to improve the electrochemical properties, especially the

ycling stability in a widened voltage range [83,84] . Zheng et

l . [75] investigated the working mechanism of Li + in O3-type

aLi 0.1 Ni 0.35 Mn 0.55 O 2 material. Reversible O3–P3–O3 phase transi-

ions during charge/discharge processes were indicated by ex situ

RD. While irreversible O3–P3 phase transition (P3 phase cannot

evert to the O3 phase upon discharge) was revealed by in situ

RD, which indicates the O3 structure is thermodynamically sta-

le and the P3 sturcture is a metastable phase. The introduction of

i + is proposed to eases the structure change during O3–P3 phase

ransition via assisting the formation of an intermediate O’3 phase,

hich is beneficial for the electrochemical performance ( Fig. 5 ).


Fig. 3. (a–d) Evolution of the {200} reflection during C/5, 5C, 10C, and 60C charging. The vertical lines indicate the equilibrium {200} position, reflecting the equilibrium

unit cell parameters a in both the Li-rich (LFP) and the Li-poor (FP) phases. (e) Diffraction pattern at approximately 50% state of charge during C/5 charging illustrating

excellent fit quality with the equilibrium LFP and FP a-unit cell parameter. (f) Diffraction pattern at approximately 5% state of charge during 5C charging displaying a weak

intermediate reflection with the a-lattice parameter equal to 10.03 ̊A representing a metastable composition between the FP and the LFP phases. (g, h) Diffraction patterns

at approximately 50% state of charge during 10 C and 60 C charging including two fits. Allowing the peak width to vary results in a better agreement between the fit and

the observed intensity, which is sandwiched between the two {200} reflections, but results in a large disagreement between the fit and observed data at the left and right

wings of the {20 0} LFP and {20 0} FP reflections, respectively. For the blue fit the C/5 peak width (no broadening) is used, which reveals the rate-induced intensity between

the two {200} reflections representing intermediate solid solutions. The Rietveld refinement was performed over the complete patterns (0.03 °–13.3 ° for the 2 θ range at

λ= 0.30996 ̊A). Reprinted from Ref. [47] with permission of American Chemical Society.


Fig. 4. (a) Schematic representation of in situ synchrotron X-ray diffraction exper-

imental setup. During the exposure the sample was continuously rotated around

the vertical axis. (b) Charging voltage curve (C/5) including the evolution of a 2D

(200) LFP and (200) FP peak showing the progressive FP formation and LFP dis-

appearance. (c) Active fraction and current density of the particles resulting from

the average transformation times. (d) Sketch of the rate-dependent transformation

upon charge as follows from the microbeam diffraction experiments. Adapted from

Ref. [55] with permission of Springer Nature.

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Polyanionic compounds for Na-ion batteries:

P olyanionic framework materials are another type of widely in-

vestigated cathode materials for SIBs, due to their stable frame-

work structure and higher operating voltages aroused by the in-

ductive effect associated with polyanion groups (PO 4 3–, SO 4

2–,

BO 3 3–, etc.). However, the theoretical capacities of polyanion ma-

terials are typically lower than layered oxides, due to the heavy

complex polyanion groups result in low mass/charge ratio. Differ-

ent phase evolution behaviors are expected for the intercalation of

Na + into polyanion hosts compared with their Li-counterparts, due

to the larger ionic radii of Na + (102 pm) compared with Li + (76

pm). For NaFePO 4 , the thermodynamically stable phase is maricite

phase rather than olivine phase in LiFePO 4 . The maricite phase is

electrochemical inactive as no cationic channels present. The elec-

trochemical active olivine phase NaFePO 4 is obtained by cation ex-

change from LiFePO 4 through the electrochemical/chemical inser-

tion of sodium in FePO 4 [85,86] . A large non-stoichiometric do-

mains and different phase evolution behaviors are observed for

olivine NaFePO 4 during cycling [87,88] . Gaubicher et al. [87] sug-

gest that the continuous variations in the metrics that mirror the

entry of Na occupancy values into thermodynamically forbidden

regions indicates the corresponding extended limits of solubility.

This smoothed phase transformation is beneficial to significantly

mitigate the lattice volume mismatch. The divergence of lattice

olume mismatch on charge and discharge is proposed to explain

he asymmetry of the electrochemical curve.

The sodium (de)intercalation reaction in Na 3 V 2 (PO 4 ) 2 F 3 is re-

orted to be a simple solid solution process in early studies.

hrough a more detailed study on the structural evolution of

a 3 V 2 (PO 4 ) 2 F 3 upon Na + extraction using high angular and inten-

ity resolution synchrotron in situ XRD, Bianchini et al. reveals an

xtremely complicated phase evolution behavior [89] . Four inter-

ediate phases with three biphasic reactions one solid solution re-

ction in the interval between 1.8 and 1.3 Na per formula unit were

bserved. And the structure of the end member NaV 2 (PO 4 ) 2 F 3 was

etermined directly from in situ measurements, which is assigned

o Cmc 21 space group with two vanadium environments: V

3 + and

5 + . Liu et al . realize 1.5 electron exchange in Na 3 VCr(PO 4 ) 3 for

odium storage with the utilization of V

3 + /V

4 + and V

4 + /V

5 + re-

ox couples [90] . An irreversible structural phase transformation

ith the formation of a metastable phase was revealed by in situ

RD during the first cycle at the room temperature, which results

n severe capacity fade. An interesting result was also shown that

he irreversible structural transition could be largely suppressed

y decreasing the temperature to −15 °C, thus leading to a much

mproved cycling stability. A solid solution characteristic with ob-

ious local lattice distortion was observed for the sodium inser-

ion/extraction reaction of Na 4 Fe 3 (PO 4 ) 2 (P 2 O 7 ) by in situ XRD [91] .

he local structural transitions were further studied by solid-state

uclear magnetic resonance (NMR). It shows that the sodium ex-

raction from the Na1, Na3, and Na4 sites causes an obvious change

n local structure. While the unchanged Na2 local structure may be

esponsible for the stability of the host structure.

.1.2. Thermal stability and reaction kinetics

Safety is one of the most important desirable performance char-

cteristics of batteries, especially for large scale automotive appli-

ations. Thermal stability of electrode materials plays an important

ole in the safety characteristics of batteries. The cathode material

ay become unstable at highly delithiated state (i.e. charged) and

ecompose through exothermic or endothermic phase transitions

92–95] . The decomposition of charged cathode materials may

ause the release of highly reactive oxygen-containing species

e.g., O

2–, O

–, O2 2–, and O

2 ), which can react with the flammable

lectrolyte and release a large amount heat. Thus the degradation

f electrode materials in charged batteries may trigger highly

xothermic reactions, which results in severe thermal runaway

nd fatal failure [96,97] . The systematic study on the ther-

al stability of electrode materials will undoubtedly provide

mportant information on understanding the composition-

tructure-thermal stability relationship, and insight into the

ational design of high capacity materials with reasonably good

hermal stability. In situ temperature dependent XRD has been

roven to be a powerful tool to study the thermal stability and

ecomposition mechanism of electrode materials, especially when

oupled with other in situ characterization techniques, such as

ass spectroscopy (MS) [93,98–105] . Bak et al . systematically

nvestigated the thermal stability of charged LiNi 0.8 Co 0.15 Al 0.05 O 2

nd LiNi x Co y Mn z O 2 (NCM433, NCM532, NCM622 and NCM811)

athode materials upon heating using combined in situ XRD and

S techniques [103,105] . They revealed that the state of charge

ffects both the structural changes as well as the evolution of

2 and CO 2 gases from Li x Ni 0.8 Co 0.15 Al 0.05 O 2 cathode materials

uring thermal decomposition. The evolution of both O 2 and

O 2 gases from charged Li x Ni 0.8 Co 0.15 Al 0.05 O 2 cathode materials

re closely correlated with phase transitions that occur during

hermal decomposition. For LiNi x Co y Mn z O 2 , a specific path of

hase transitions from layered ( R 3 ̅m ) to spinel ( Fd 3 ̅m ), and then to

ock-salt ( Fm 3 ̅m ) upon thermal decomposition during heating was

evealed. The structural changes and the oxygen release features


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uring heating are strong related to the content of Ni, Co, and Mn.

he onset temperature of the phase transition (i.e., thermal de-

omposition) decrease and the amount of oxygen release increase

ith increasing the Ni content.

In situ XRD has also been shown to be a powerful tool to

racking preparation reactions in synthesizing electrode materials

nder real reaction conditions [106–109] . In recent efforts, it has

een used to studies of phase transition behaviors in intermedi-

tes during solid-state synthesis of Ni-rich NCM [108,109] . Wang

t al. [109] investigated the kinetic reaction pathway and cationic

rdering in the intermediates on synthesis reactions for prepar-

ng LiNi 0.7 Co 0.15 Mn 0.15 O 2 (NCM71515) under ambient atmosphere.

complex cationic ordering and disordering processes is observed,

hich concurrently occur throughout the heat treatment process.

t reveals that the kinetics of cationic ordering in NCM71515 are

trongly related to temperature, arising from concurrently occurred

hermal-driven oxidation of transition metals and lithium/oxygen

oss during heat treatment. Guided by insights from in situ stud-

es, a layered NCM71515 with low cationic disordering and a high

eversible capacity (up to 200 mAh/g) and excellent retention is

repared in air through synthetic control of the kinetic reaction

athway.

.2. The application of synchrotron Pair Distribution Function (PDF)

In the last decade, as a promising tool to probe the local atomic

tructure of materials, PDF analysis of high-energy X-ray scattering

ata has been widely used to the investigation of materials that do

ot show a long-range lattice periodicity, i.e. nanoscale or amor-

hous materials, liquid, or gaseous phases [110,111] . The PDF is

btained through the Fourier Transform of the total scattering

owder diffraction pattern [111,112] . Since the total scattering

iffraction pattern includes both Bragg and diffuse scattering con-

ributions, it provides a wide range of structure information, in-

luding local, medium range and long range structural information

112] . PDFs are generated from the powder diffraction data through

sine Fourier transformation of the normalized scattering function

( Q ) [111] :

( r ) = 4 π [ ρ( r ) − ρ0 ]

2

π

∫ ∞

0

Q [ S ( Q ) − 1 ] sin ( Qr ) d Q

here ρ( r ), ρ0 and Q are the microscopic pair density, average

umber density and magnitude of the scattering vector, respec-

ively. For elastic scattering Q = 4 π sin( θ )/ λ with 2 θ and λ being

he scattering angle and wavelength of the X-rays. By plotting the

( r ) (i.e. the PDF) gives the probability of finding an atom at a

iven distance ‘ r ’ from another atom. In other words it can be de-

cribed as a distribution of bond length. While the crystallographic

pproach (XRD) can only be applied to crystalline materials and

rovides the average structure of the material, as it is based on

he analysis of Bragg intensities. The PDF analysis can be applied

o both crystalline and amorphous phases. For crystalline materi-

ls, the information of deviation from average structure can be ob-

ained. The synchrotron PDF analysis shows promising application

n rechargeable batteries, it can yield detailed insights into elec-

rode processes, including local atomic structure, phase progres-

ion and partial size/ordering, etc. [110,113–115] . Since nanomate-

ials are widely pursued to enhance electrode performance, also

morphization, disordering and/or nanoparticle formation accom-

any with pronounced structural rearrangements are observed for

onversion-based electrodes during the initial conversion reaction.

ombining the PDF analysis with NMR and XRD, both the short-

ange and long-range structure of battery materials can be revealed

n detail [113,116–122] .

Lithium thiophosphates (LPS) are extensively studied as solid

lectrolytes for solid-state lithium batteries, due to their inher-

ntly high ionic conductivity and mechanically soft nature [123–

25] . Both crystalline and glass compounds with multitude of dif-

erent compositions, Li 3 PS 4 , Li 7 P 3 S 11 , and Li 10 GeP 2 S 12 , etc., have

een proven to be very promising candidates [125–127] . Crystal-

ization from their glass phases forming glass ceramic composites

s an important approach to prepare these compounds. Revealing

he principle structure-property relationship will provide new in-

ights on the search for optimized glass compositions for superior

olid electrolytes.

PDF has been shown to be a suitable tool to identify the local

tructure of LPS glasses [115,120,121] . Using in situ PDF analysis,

ietrich et al. monitored the local environments and phase evo-

ution during the crystallization of superionic conductors Li 3 PS 4 ,

i 7 P 3 S 11 and Li 4 P 2 S 7 , which is correlated with the observed ionic

onductivity and temperature dependent changes [120] . A new lo-

al structural stability diagram is proposed for the LPS glasses,

hich highlights the importance of optimizing the thermal treat-

ent in the formation of different building units and maximiz-

ng the ionic conductivity ( Fig. 6 ). The investigation of combining

DF analysis and neutron powder diffraction shows that poor ionic

onductor and high activation barriers of Li 2 P 2 S 6 may attribute to

he rather low structural symmetry of the lithium positions, which

eads to their spatial separation in a highly distorted lithium coor-

ination polyhedron [121] .

Grenier et al . investigated the electrochemical reactions mech-

nism of rechargeable Fluoride-Ion Batteries with a Bi–BiF 3 –

BF–C composite electrode by using synchrotron XRD and PDF

nalyses [128] . Oxygen does not migrate through the elec-

rolyte is proven by quantitative PDF analysis, which sug-

ests that the fluoride ion is the sole charge transfer an-

on. Using PFD analysis, Idemoto et al. studied the local struc-

ural changes in layered 0.5Li 2 MnO 3 –0.5LiMn 1/3 Ni 1/3 Co 1/3 O 2 and

.4Li 2 MnO 3 –0.6LiMn 1/3 Ni 1/3 Co 1/3 O 2 solid solutions during cycling

129,130] . It shows that local structural changes in 0.5Li 2 MnO 3 –

.5LiMn 1/3 Ni 1/3 Co 1/3 O 2 are strong correlated with the charging

ates [129] . Before charging, a smaller distortion was observed for

nO 6 octahedra compared with that for NiO 6 and CoO 6 octahe-

ra. During charging at 1 C, the distortion of MnO 6 increase, while

t 3 C the distortion of CoO 6 octahedra is more severe. Moreover,

hen charging at 3 C, the values of the bond-length strain ( λ)

nd the bond-angle strain ( σ 2 ) increased for NiO 6 octahedra that

ad entered the Li layer as a result of cation mixing. PDF anal-

sis indicated that rearrangement of Mn and Co in 0.4Li 2 MnO 3 –

.6LiMn 1/3 Ni 1/3 Co 1/3 O 2 took place from 3.3 to 4.6 V during the first

ischarge process [130] . This suggests that a stable reversible phase

ormed at around 3.3 V during this process.

Using synchrotron XRD and PDF analysis, Xiang et al. inves-

igated the strain accommodation mechanism in NaFePO 4 , which

ossesses the largest known transformation strain ( ∼17 vol%) in

he olivine compounds [131] . A new strain-accommodation mech-

nism was revealed, that the large lattice mismatch between the

wo end members NaFePO 4 and FePO 4 is alleviated by formation of

third, amorphous phase. The amorphous phase can’t be identified

y powder diffraction alone. Pourpoint et al. [117] observed the

ormation of V

3 + –V

3 + –V

3 + trimers in Li 1 + x V 1–x O 2 , indicating the

agnetically-induced distortion of the V sublattice. Wiaderek et al.

116] investigated the structural and chemical changes in mixed-

nion FeOF electrodes. In situ PDF analysis suggests that anion par-

itioning occurs during discharge and charge, with the rock salt

hase being O-rich and the rutile phase being F-rich. The F- and

-rich phases react sequentially; Fe in a F-rich environment reacts

referentially during both discharge and charge. This unexpected

referential reaction mechanism may contribute to the attractive

lectrochemical performance of these compounds.


Fig. 5. In situ XRD patterns collected during the first charge/discharge process of the O3-type NaNi 0.5 Mn 0.5 O 2 (a) and NaLi 0.1 Ni 0.35 Mn 0.55 O 2 , (b) electrode, cycled at a current

rate of 36 mA/g ( λ= 0.6887 ̊A). Reproduced from Ref. [75] with permission from The Royal Society of Chemistry.

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3.3. The application of X-ray absorption spectroscopy (XAS)

XAS is a widely used technique for determining the electronic

structure and/or local geometric of materials. It is based on the

X-ray photoelectric effect, in which X-ray photons that have suffi-

cient energy incident on an atom within a sample is absorbed and

ejects a core electron (e.g. 1 s). XAS is element-specific technique

that measures the X-ray absorption coefficient μ( E ) of a material as

a function of X-ray energy E in an energy range that is below and

above the absorption edge of the measured element in the material

[12,132] . When an energy-tunable monochromatic X-ray beam inci-

dent on a sample is scanned through the binding energy of a core

shell, an abrupt increase in the measured μ( E ) can be observed,

which is called “absorption edge” of the element. The observed

absorption edges are correlated with photoelectric absorption of

the shell electron of an atom, with each edge (K, L, M… edges)

representing a well-defined core-electron binding energy. X rays

with high photon energies from a few keV up to about 100 keV

are called hard X-rays, while those with lower energy from tens of
r
V to a few keV are called soft X-rays. The corresponding XAS is

alled “hard” or “soft” X-ray absorption spectroscopy according to

he energies of interest.

With XAS, the oxidation state and the local environment of

he interested element within the studied samples can be inves-

igated selectively, with plenty of information, including the oxida-

ion states, bond length and coordination numbers can be obtained

133,134] . Moreover, the XAS technique can be used to almost all

amples, which including gas, liquid, or solid states in both crys-

alline and amorphous phases. Due to its distinguishing features,

AS has been widely used to the investigation of battery materi-

ls, especially the use of in situ XAS to monitor the oxidation state

nd local structure evolution of electrodes during electrochemical

eaction.

.3.1. Hard X-ray absorption spectroscopy

Hard X-ray absorption spectroscopy (Hard XAS) also known as

-ray absorption fine structure (XAFS), is generally divided into X-

ay absorption near-edge structure (XANES) and extended X-ray


Fig. 6. The observed synchrotron diffraction Bragg data of 75:25 LPS glass at room temperature only show a diffuse pattern, without any long-range order (a). No intense

peaks for r > 4.7 ̊A are observed in the corresponding Pair Distribution Functions (b). The short-range order is determined by the PS 4 -tetrahedral first coordination sphere,

which is very similar for all LPS glasses (c). Reflections at 245 °C could be assigned to β-Li 3 PS 4 (d), which was used for the simulation of the PDF profile (e). The resulting

profile difference (green line, e) was compared with the room temperature PDF data of the initial glass in the range of 0.5 ̊A < r < 5 ̊A, which corroborates the coexistence of

amorphous and crystalline phases, typical for LPS glass-ceramics (f). Reproduced from Ref. [120] with permission from The Royal Society of Chemistry. (For interpretation of

the references to color in this figure legend, the reader is referred to the web version of this article.)

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bsorption fine structure (EXAFS) two parts [12,132] . XANES and

XAFS regions contain different information, and are generally an-

lyzed separately. XANES usually refers to the part of the spectrum

ear the absorption edge (within ∼30–50 eV), which providing in-

ormation about oxidation state and molecular geometry, but is

ormally analyzed qualitatively. EXAFS refers to oscillations above

he absorption edge, which can extend for 1,0 0 0 eV or more start-

ng at 20–30 eV above the edge. EXAFS is sensitive to the radial

istribution of electron density around the absorbing atom, which

roviding structural information about the central absorbing atom

ite ligation i.e. type of neighboring atoms, coordination number

nd bond distances, is normally analyzed via a quantitative com-

arison between theoretical modeling and experiments. Hard XAS

s an ideal method for in situ investigations of battery systems, due

o the high penetrability of hard X-ray. The K-edge energy of the

ost common first row transition metals (3d TMs) in cathode ma-

erials lies in energy ranges from 4.5 to 10 keV, which can be easily

easured in in situ experiments using transmission mode.

The element specific characteristic of XAS makes it highly

uitable for the investigation of electrodes containing multiple

ransition metals, e.g. layered Li(Ni x Co y Mn z )O 2 , Li-rich x Li 2 MnO 3 ·1–x )LiMO 2 (M = Ni, Co, Mn, etc.), and olivine Li(Ni x Co y Mn z )PO 4

135–148] . For Li(Ni x Co y Mn z )O 2 , the metal K-edge XANES results

eveal that the charge compensation during delithiation is mainly

chieved by the oxidation of Ni 2 + to Ni 4 + , indicated by the shift

o the higher energy of the absorption edge in Ni K-edge XANES

pectra during charging [137,140] . While the manganese and cobalt

ons remain mostly unchanged in the tetravalent and trivalent

tate. EXAFS fitting results show a significant change in the bond

ength and Debye −Waller factor of the second shell M–M contri-

ution, which results from the generation and reduction of the

ahn −Teller active Ni 3 + concentration on the charge and discharge

ycle. Significant and similar variation trends of the bond length

nd Debye −Waller factor (indicating the distortion of the local

tructure) are observed for the Mn −M and Ni −M contributions,

hile totally different with those of Co–M. This suggests the short-

ange ordering between Ni 2 + and Mn

4 + may exist in the com-

ound, which can’t be detected by XRD.

Li-rich x Li 2 MnO 3 ·(1–x )LiMO 2 materials are very promising cath-

de candidates for next high energy density LIBs, due to is high

eversible capacity ( > 250 mAh/g). However, they suffer from se-

ere capacity fade and voltage decay, which severely hinder its

ractical application. The mechanisms of capacity fade and volt-

ge decay associated with x Li 2 MnO 3 ·(1 − x )LiMO 2 have been stud-

ed by using XAS and XRD combing with NMR, neutron diffraction

nd HR-TEM, etc. The metal K-edge XAS results reveal an inherent

tructural reorganization during the electrochemical activation of

he Li 2 MnO 3 component, which contributes to the hysteresis ob-

erved in this system [136] . It shows that the average local struc-

ure and oxidation of Ni is different in the cathode even at the

ame state of charge (i.e., the same lithium content, SOCs) between

harge and discharge. In situ metal K-edge XAS results revealed

hat the three voltage plateaus at ∼3.6, 4.2 and 4.7 V vs. Li/Li +

n LiFe 1/4 Mn 1/4 Co 1/4 Ni 1/4 PO 4 cathode are correlated with the re-

ox couples of Fe 2 + /Fe 3 + , Mn

2 + /Mn

3 + and Co 2 + /Co 3 + , respectively

138] . The Ni K-edge XANES spectra remain mostly unchanged dur-

ng charging, suggesting the apparent voltage plateau above 4.9 V

s very likely originated from the electrolyte decomposition, rather

han the Ni 2 + /Ni 3 + redox.

Electrode materials based on conversion-type reactions (typi-

ally oxides and fluorides, MF x and M x O y ) have attracted signif-

cant attention, due to their high specific capacity via multiple-

lectron transfers per redox centre. During electrochemical cycling,

he conversion electrode could reversibly react with lithium via a

onversion reaction, which forming a composite structure compris-

ng metal nanoparticles embedded in a lithium salt (LiF, Li 2 O, etc.)

149] . The conversion reaction results in dramatic electronic and

tructure changes, which makes the elucidation of electrochemical


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mechanism very complicated. Integration of multiple advanced an-

alytical techniques, such as synchrotron based in situ XRD, XAS

and solid-state NMR etc., will undoubtedly facilitate the analysis

of this complex process. Combining in situ XAS with solid-state

NMR and PDF analysis, Hu et al . studied the mechanism for gener-

ating the additional capacity in conversion material RuO 2 [150] . In

situ XAS results indicate the lithiation process can be divided into

three stages. During stage I (0–1.3 Li), a shift to the lower energy

of the absorption edge in the XANES spectra is observed, indicat-

ing the reduction of ruthenium with the insertion of lithium. The

fitting results of EXAFS spectra show a change in intensity of the

Ru–Ru correlations, consistent with the insertion of lithium into

the RuO 2 structure with a solid solution reaction. XAS results in-

dicate a two-phase conversion reaction in stage II. While no ob-

vious change is observed in XANES spectra in stage III, indicating

the capacity observed here is not related to the redox reaction of

Ru-containing phases. Combing XAS, PDF and NMR results reveal

that the extra capacity in RuO 2 mainly come from the contribution

the generation of LiOH and its subsequent reversible reaction with

Li to form Li 2 O and LiH. The in situ XAS studied on electrochem-

ical reaction mechanism of FeF 3 reveals a three stages reaction

mechanism: a two-phase intercalation reaction in the first stage

(0 to 0.46 Li), a single-phase intercalation reaction in the second

stage (0.46 to 0.92 Li), and a conversion reaction in the third stage

(0.92–2.78 Li) [151] . EXAFS fitting results show splitting trends of

the Fe −F bond lengths and the Fe −F CNs, which supports the pro-

posed phase transformation from R3c structured FeF 3 to R3c struc-

tured Li 0.92 FeF 3 during the intercalation stage. After fully conver-

sion, small Fe −Fe CN (ca. 2.2) is observed for the conversion prod-

ucts, indicating small Fe nanoparticles ( < 1 nm) are formed during

conversion stage. The conversion reaction mechanisms of copper

phosphate for lithium and sodium storage have been investigated

by combing XAS, NMR, and HR-TEM, etc. [152,153] . It reveals a

very complicated conversion mechanism with several intermediate

phases, including Li x Cu(I) y PO 4 , Li x Cu 3 (PO 4 ) 2 , Li x Cu(II) y PO 4 formed

[153] . Interestingly, it shows that Cu and Li 3 PO 4 can conversion

to Li x Cu y PO 4 and copper phosphate during the charging process,

opening a new venue to explore polyanion-type conversion reac-

tion ( Fig. 7 ).

By an appropriate combination of three XAS beamlines with

various time and space resolutions, Ouvrard et al. investigated

the inhomogeneous electrochemical reaction distribution in the

LiFePO 4 electrode [154] . They observed an obvious heterogeneous

reaction behavior within the electrode during cycling, with some

parts being advanced and others delayed, compared to the mean

charge state of the electrode. They proposed that the observed het-

erogeneous behavior is associated with the quality of grain connec-

tivity to ionic and electronic percolating networks. Using in situ

XAFS imaging technique, Katayama et al . studied the spatial distri-

bution of electrochemical reaction for the planar LiFePO 4 cathode

[155] . It also revealed the reaction distribution of active materi-

als is inhomogeneous. The electrode reaction is found to occur in

the reaction channels and expands radially, thus they concluded

that the inhomogeneous electrode reaction can attributed to the

difference in electrical conductivity, the electrochemical reaction

proceed through reaction channels with low electrical resistance

in the cathode ( Fig. 8 ).

XAS is also proven to be a powerful tool for the study of dy-

namic process in battery with high temporal resolution (up to mil-

liseconds) [156–160] . Using in situ time-resolved XAS, Yu et al . per-

formed a dynamically study on the two-phase transition behavior

between LiFePO 4 and FePO 4 during chemical and high rate electro-

chemical delithiation [156] . Their results suggest that the delithia-

tion are dominated by a two-phase reaction mode at both low and

high charge rates (from 1 C to 30 C). Using in situ time-resolved

XAS and XRD, Arai et al . monitored the phase transition processes

f Li x Ni 0.5 Mn 1.5 O 4 [157] . Their results indicate a two first order

wo phase transitions reaction mechanism, with the coexistence of

iNi 0.5 Mn 1.5 O 4 (Li1) and Li 1/2 Ni 0.5 Mn 1.5 O 4 (Li0.5) in the low poten-

ial plateau, and Li 1/2 Ni 0.5 Mn 1.5 O 4 (Li0.5) and Ni 0.5 Mn 1.5 O 4 (Li0) in

he high potential plateau. The transition between the Li1 and Li0.5

hases is faster than that between the Li0.5 and Li0 phases during

ycling at 1 C, which results in thermodynamically reversible but

inetically asymmetric behavior of this electrode.

.3.2. Soft X-ray XAS

The energies (between 50 eV and 2 keV) of soft XAS cover the K-

dge of light elements (such as Li, C, O, F, Na, Mg) and the l -edge

f 3d TMs; both are prevalent in battery materials. Typically X-

ay absorption coefficient is measured by detecting the total num-

er of either the fluorescence (photon) or Auger (electron) decays

uring photoelectric process, giving rise to the total electron yield

TEY) and the total fluorescence yield (TFY), respectively. Naively

oth TEY and TFY can be considered as proportional to the absorp-

ion cross section for dilute and thin samples, so can be used for

ecording X-ray absorption spectra [161] . The TEY is surface sensi-

ive with a probe depth of several nanometers due to the shallow

scape depth of electrons. TFY is sub-surface and bulk sensitive

ith probe depth of about 150 nm as fluorescence photons have

onger escape depth [162] . While TFY signals is interfered by the

ntrinsic saturation effect, partial fluorescence yield (PFY) and in-

erse partial fluorescence yield (iPFY) are used to settle this prob-

em [163] . The contrast between the TEY and TFY signals often pro-

ides some qualitative but valuable comparison between the ma-

erial surface and bulk [162] .

For light elements, soft XAS allows to probe core to valance

1 s –2 p ) transitions through K-edges, which is directly correlated

o the physical and chemical characteristics of battery materials.

arbon and oxygen are the constituents of conductive additives

nd binders, their K-edges provide important information about

onductivity and molecular interactions. The quality of electrode

rains connectivity to conductive additive and binder significantly

ffects the electrode performance, soft XAS was thus used to detect

he interface interaction between electrode grains and conductive

dditives and binders [164–166] . Using C K-edge XAS, Ji et al. in-

estigated the mechanism of anchoring S in graphene oxide [167] .

he spectra of GO and GO-S nanocomposites indicate that the in-

orporated S enhances the ordering of the sp 2 hybridized carbon

tructure, and a strong chemical interaction between S and the

unctional group of GO does exist. The solid electrolyte interphase

SEI) plays a very important role in battery operation. Understand-

ng the formation, properties and evolution of the interphase layer

s of vital importance to both fundamental research and practical

pplication. Soft XAS is a powerful technique to investigate the SEI

ith a tunable probing depth, due to its distinct characteristics of

oth surface and bulk sensitive [168–172] . The main components of

EI are decomposition products of electrolytes and additives, also

issolution or contamination products from electrodes, which con-

ist of mainly light elements (i.e. Li, C, O, F, B, S, P from the decom-

osition of both electrolyte and additives) and 3d TMs. The K-edges

f most of light elements and l -edge of 3d TMs can be measured

y soft XAS. Yogi et al . investigated the formation of the SEI on

pulsed laser deposition LiCoO2 electrode [169] . The spectra for

-edges of Li, B, C, O, and l -edges of Co were collected using PEY,

EY and FY three detection methods with different probing depths.

hey revealed that SEI film containing lithium carbonate was in-

tantly formed just after the contact with electrolyte and that it

ecomposed during the repeated charge–discharge reactions. They

lso showed that the additive of lithium bis(oxalate) borate (Li-

OB) can effectively suppress the decomposition of SEI and the

eduction of Co ions at the electrode surface to Co(II), thus im-

rove the cycle performance. Delacourt et al . studied the effect of


Fig. 7. (a) The first derivative curves of in situ XANES spectra (black) and the linear combination fitting curves (red), obtained using the first derivative XANES curves of

Cu 3 (PO 4 ) 2 ( x = 0) and Li 6 Cu 3 (PO 4 ) 2 as standards. The vertical axis ( x ) indicates the corresponding value in Li x Cu 3 (PO 4 ) 2 . (b) Discharge/charge profiles and fitting results

(coordination numbers and bond lengths) for Cu–O and Cu–Cu from the in situ EXAFS data. Reprinted from Ref. [153] with permission of American Chemical Society. (For

interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

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anganese contamination on the SEI properties at the inert model

node. The presence of Mn

2 + ions in the SEI films is proven by the

n l -edge XAS [170] . They proposed a multiphase transformation

rocess, which Mn

2 + from the electrolyte reduces to Mn

0 at the

lectrode surface and further reoxidizes back to Mn

2 + by reacting

ith solvent molecules.

Different with the K-edge of 3d TMs, which does not provide di-

ect information on the metal valence charge and spin density due

o the probed 1s–np transitions do not contain the metal-d orbitals

ost relevant to chemical bonding; l -edge of 3d TMs fingerprints

he formal valence, spin state and chemical bond configuration be-

ause it directly probes unoccupied valence orbitals of 3d TMs via

tilizing dipole allowed 2 p → 3 d transitions [162] . The details of the

edox reaction mechanism in Li-rich x Li 2 MnO 3 ·(1–x )LiMO 2 cathode

s still ambiguous and under intense discussion, although it is gen-

rally accepted that the anionic (O

2 −) redox processes plays an im-

ortant role. Combing oxygen K-edge and transition metal l -edges

oft XAS and in situ hard XAS, the charge compensation mecha-

ism in this system has been extensively investigated [173–178] .

ypically reversibly shifts between the charged/discharged states

re observed for Ni and Co l -edge XAS spectra, indicating that

he Ni and Co ions reversibly participate in the charge compen-

ation. While the Mn l -edge spectra remain scarcely changed dur-

ng charging, indicating that the Mn ions stay mostly unchanged

s the tetravalent states. The changes in the O K-edge spectra are

roposed to be the evidence of participation of the oxygen an-

on (O

2 −) redox in the charge-compensation mechanism ( Fig. 9 )

176,179] . Using soft XAS, Qiao et al . performed a quantitative anal-

sis of the Mn oxidation states in Na x MnO 2 electrodes at different

OC [180] . A quantitative analysis of Mn l -edge soft XAS showed

hat, during discharged from 4 to 2.6 V, the Mn

4 + concentration

ecreased monotonically and the Mn

3 + concentration increase. If

he cell was further discharged below 2.6 V, the Mn

2 + concentra-

ion increased rapidly. The comparison between bulk-sensitive hard

-ray Raman spectroscopy and soft XAS, further reveal that the

n

2 + formation occurs only on the surface of Na 0.44 MnO 2 parti-

les. A portion of the surface Mn

2 + compounds become electro-

hemically inactive after extended cycles, resulting in the capacity

ading. The formation of Mn

2 + can be suppressed by regulating the

ischarge cut-off voltage to 2.6 V or higher.

.4. The application of XPS

The electrode surface and electrode/electrolyte interphase play

n important role in the performance and lifetime of recharge-

ble batteries. Side reactions between electrode and electrolyte

esult in the formation of solid electrolyte interphase (SEI).

he nature of SEI layers and its formation depend mainly on

he electrode surface chemistry and electrode potential, i.e.,

n possible (chemisorption) dissociative, catalytic decomposition

nd electrochemical/chemical redox. A detailed understanding of

uch processes as adsorption/desorption, bond breaking/formation,

harge transfer and redox transformations occurring at elec-

rode/electrolyte interfaces is essential for the effective control of

EI by surface coating and electrolyte additives. XPS is one of the

ost versatile techniques for the study of surfaces on the atomic

cale [181] . It provides qualitative and quantitative information

bout the elemental composition and chemical specificity (e.g., ox-

dation state) of the surface and interfaces. Most of the XPS re-

earches to date have used traditional in-house laboratory Al K αr Mg K α radiation. Here, we will mainly focus on the investi-

ation into battery electrode surfaces and interfaces by XPS with

ynchrotron radiation used over a broad energy ranging from the

oft X-ray regime of a few hundred eV to the harder keV ener-

ies. Synchrotron based XPS using soft and hard X-ray are referred

s Soft X-ray Photoelectron Spectroscopy (SOXPES) and Hard X-ray

hotoelectron Spectroscopy (HAXPES), respectively. SOXPES is near

urface sensitive and HAXPES is more bulk sensitive. Compared


Fig. 8. (a) XANES spectra and (b–d) chemical state maps for the LiFePO 4 cathode during charging obtained by ex situ XAFS imaging measurements. The successive maps

obtained by in situ XAFS imaging are shown in (e). Reprinted from Ref. [155] with permission of Elsevier.

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with characteristic lab sources such as Al K α, faster measurements

and better resolution can be realized with the use of extremely

bright synchrotron radiation. Foremost, tunable depth sensitivity in

a wider range of the order of 2–50 nm can be achieved via tuning

photon energies within wide ranges, which provides the opportu-

nity for nondestructive depth-resolved analysis. Furthermore, syn-

chrotron based XPS (SXPS) offers new opportunities allowing for

surface and interface investigations under in situ conditions and

ambient or high pressure.

Philippe et al. systemically investigated interfacial mechanisms

(reaction of surface oxide, Li −Si alloying process, and passivation

layer formation) of Si/C/CMC composite electrodes by combing

Hard and Soft X-ray Photoelectron Spectroscopy [182–185] . Infor-

mation on interfacial phase transitions at the surface of silicon par-

ticles as well as the composition and thickness of the SEI was ob-

tained via variation of the analysis depth. A schematic view of the

mechanisms occurring at the surface of the silicon nanoparticles

upon discharge and charge is provided. The SEI is largely formed

at the very beginning of the first discharge, and it has similar com-

position to the SEI formed at the surface of carbonaceous elec-

trodes or Sn or Sb-based intermetallic electrodes. The thickness of

SEI layer increases with lowering of the electrode potential. Simul-

taneously, lithium reacts with the silicon nanoparticles. Li 2 O and

lithium silicate Li 4 SiO 4 are formed via reaction with the surface

SiO 2 layer, and Li −Si alloy is formed via reaction with bulk sili-

con. During the first charging, the thickness of SEI layer decreases

slightly, and Li O has disappeared. It was also found that LiPF salt
2 6
lays an important role on the stability of the silicon electrode dur-

ng cycling. The formation of partially fluorinated silicon species

SiO x F y ) is detected upon cycling at the outermost surface of the

ilicon nanoparticles due to the reaction of the materials toward

he electrolyte. A similar species is also observed by simple con-

act of the pristine electrode with electrolyte. They also showed

hat LiFSI salt is able to avoid the fluorination process, which mod-

fy the composition in silicon particles surface, and preserve the

avorable interactions between the binder and the active mate-

ial surface. The reduction products of LiFSI salt deposited at the

urface of the electrode act as a passivation layer which enhance

he performance of silicon electrode. Study on the mechanisms of

lectrolyte additive fluoroethylene carbonate (FEC) show that de-

uorination and ring-opening are the two most possible degrada-

ion routes for FEC on silicon electrodes. It influenced the decom-

osition reaction of LiPF 6 salt and may have suppressed further

alt degradation. LiF and −CHF −OCO 2 -type compounds are found

o be the main reduction products with FEC additive, while car-

on and oxygen containing species dominate the SEI layer with-

ut the presence of FEC. In the presence of FEC, a conformal SEI

as instantaneously formed on the silicon electrode, which limited

he appearance of large cracks and maintained the original surface

orphology.

Younesi et al . studied the surface compositions of a MnO 2 cat-

lyst containing carbon cathode and a Li anode of a Li −O 2 bat-

ery using LiClO 4 or LiBOB Salts [186] . Carbonate and ether con-

aining compounds is formed upon cycling on the surface of the


Fig. 9. O K-edge XAFS spectra of Li 1.16 Ni 0.15 Co 0.19 Mn 0.50 O 2 electrode during the reversible 1st and 2nd cycle processes. They are obtained in (a) TEY and (b) PFY mode. The

pre-edge structures are expanded in (c) TEY and (d) PFY mode. Reprinted from Ref. [179] with permission of Elsevier.

c

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arbon cathode as a result of the decomposition of PC and EC/DEC

olvents. However, less decomposition products were detected for

C/DEC compared to PC. Decomposition products from LiClO 4 or

iBOB were also detected and contribute to the formation of sur-

ace layers on the cathode surface when cycled in PC. While, no

egradation of LiClO 4 was observed when cycled in EC/DEC. It was

lso found that a surface layer forms and remains on the MnO 2

atalyst at the end of the charged state. The SEI on the Li anode

s made of PEO, carboxylates, carbonates, and LiClO 4 salt in the

harged state. Malmgren et al . preformed a depth profiling on lithi-

ted graphite and delithiated lithium iron phosphate electrodes in

full cell configuration (graphite vs. LiFePO 4 ) to study anode and

athode electrode/electrolyte interface using SOXPES and HAXPES

187] . The deeper parts of the anode SEI was found to mainly con-

ist of lithium oxide and alkoxides, which are not observed in the

athode/electrolyte interface (solid permeable interface, SPI). And

–H and C–O containing compounds are more common in the out-

rmost anode SEI than in the innermost. While no obvious com- g

osition gradients were detected in the SPI. Significant amounts

f decomposition products consisted of C–O and P–F compounds

ere found both in the cathode SPI and the anode SEI during cy-

ling. The active material can be detected even at lower photon

nergies indicates that the SPI is obviously thinner than the SEI.

he thickness of SEI layer on lithiated graphite was estimated to be

round two tens of nanometers, while the thickness of cathode SPI

as estimated to be a few nanometers only. Lu et al . studied the

nfluence of surface adventitious hydrocarbons and carbon dioxide

CO 2 ) on the reversibility of the Li −O 2 redox chemistry using in

itu synchrotron-based ambient pressure XPS [188] . The formation

f carboxylate and carbonate-based species was observed as a re-

ult of the irreversibly reactions between surface hydrocarbons and

O 2 with Li −O 2 reaction intermediates/products such as Li 2 O 2 and

i 2 O, which cannot be removed fully upon recharge. The increas-

ng coverage of surface carbonate/carboxylate species contributes

o the slower Li 2 O 2 oxidation kinetics. Schwöbel et al . investi-

ated the valence band structure of LiPON using SXPS with variable


m

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2

excitation energies [189] . The top of the valence band of LiPON was

found to be due to nitrogen 2p states, which suggests that the ni-

trogen content of the produced LiPON films might affect the en-

ergy band alignment between cathode and the LiPON electrolyte.

Sachs et al . studied the surface films on Li 4 Ti 5 O 12 (LTO) coated

LiNi 0.5 Mn 1.5 O 4 high-voltage cathode upon cycling using HAXPES

with photon energy between 2 keV and 6 keV [190] . A thin surface

layer (a few nanometers) was found to cover on the Li 4 Ti 5 O 12 coat-

ing as a result of electrolyte decomposition. Organic polymers as

well as metal fluorides and fluorophosphates are the main compo-

nents of the surface layer. It was found that the Li 4 Ti 5 O 12 coat-

ing plays a positive role on the size and stability of the sur-

face layer. The XPS spectra collected with conventional and syn-

chrotron X-rays at different excitation energies revealed the forma-

tion of a 5–10 nm surface layer containing organic and Li x PF y O z -

type species on cycled LiNi 0.5 Mn 0.5 O 4 as result of electrolyte de-

composition [191] . And a comparatively thinner ( < 4 nm) layer

composed of transition metal fluorides and LiF was formed close

to the LiNi 0.5 Mn 0.5 O 4 electrode surface. Nordh et al . studied the

SEI on Li 4 Ti 5 O 12 (LTO) using SXPS [192] . Surface layer formation

without resembling the typical SEI on a graphite negative electrode

was clearly observed even when the sample was only exposed to

electrolyte. Entirely or partially to disintegrate was also detected

for this independently formed surface layer, which is replaced

with a surface layer of slightly different composition and thickness

when the cell was cycled. The surface layer is mainly composed

of LiP x F y O z , C–O containing species and P–O compounds, which

prefers to form on the carbon additive in the samples rather than

on the LTO particles. The effects of manganese in the SEI layer of

Li 4 Ti 5 O 12 in LTO/LMO cell was also studied by combined NEXAFS

and HAXPES techniques [193] . It shows that a solid surface layer

containing deposited manganese from the cathode was formed on

the LTO electrode. The deposited Mn is present in the ionic state

( + II and + III oxidation states) with a dominance toward + II, in-

dependent of the lithiation and delithiation of the LTO electrode.

The chemical environment of the deposited Mn is complex and dif-

fers from that of graphite. The chemical composition of the formed

SEI in LTO electrode is not significantly affected by Mn deposition,

indicating a rather passive role of manganese during cell cycling.

Hausbrand et al . systemically investigated the interface formation

between the thin film LiCoO 2 cathode and the solvent adsorbate

of DEC, dimethyl sulfoxide (DMSO) and EC [194–197] . The results

demonstrate that the decomposition of adsorbing solvents in con-

tact with LiCoO 2 materials as observed in adsorption experiments

is complex, including the reduction of solvent, subsequent decom-

position reactions, and also catalytic effects. The spectra indicate

that partial electron transfer coupled to covalent interaction lead-

ing to a reduction of organic solvents and the formation of inor-

ganic lithium-containing reaction products.

4. Conclusions and perspectives

It is obvious that the applications of synchrotron based tech-

niques to the investigation of battery system is vast and still

rapidly developing, this review have briefly introduced some re-

cently progresses of several widely used techniques, such as XRD,

PDF and Soft and Hard XAS. Although it is inevitable incomplete,

since new techniques and results are being reported every day, it

does highlight the great potential for the future implicating and

developing of these powerful techniques.

This review illustrates the usefulness of synchrotron techniques

for the study of a wide range of phenomenon occurring in battery

systems. Moreover, it demonstrates the unique capabilities of syn-

chrotron based in situ techniques for directly monitoring the elec-

trochemical processes occurring in a battery in operando, and the

chemical reaction processes for the synthesis and thermal treat-

ent of battery materials, under a variety of conditions. With its

igh temporal resolution, it makes the real-time monitoring of dy-

amic processes that rule electrochemical energy storage in batter-

es possible, and allows direct detection of short-lived intermedi-

tes in batteries during cycling. It high spatial resolution allows in-

estigating the spatial distribution of the electrochemical reaction

n electrodes spanning a wide length scale. Using scanning mode

r imaging technique, it is able to provide spatially resolved maps

f the distribution of phases, chemical states, and elements in the

lectrode. With the rational design of experiment combining tech-

iques with different probing depths, it is possible to determine

he chemical and electronic properties of the interface, and analyze

epth profiles non-destructively on the electrode interface. These

undamental researches on batteries have provided critical insights

nto their functional mechanisms, degradation and failure mecha-

isms, thus allowed the development of innovative approaches to

mprove their performance.

We thus believe that in the future synchrotron techniques will

lay a significant role in the investigation of rechargeable batter-

es by X-rays. Its remarkable ability along with rapid development

f synchrotron techniques, will provide a precise understanding

f the nature and the evolution of the battery chemistries dur-

ng cycling. Thus shed new light on enhancing the properties of

attery materials and open up new design principles and oppor-

unities to fundamentally develop new materials for current and

ext generation rechargeable batteries with improved energy den-

ity, reliability and safety. Finally, speed up the innovations in bat-

ery science and technology. On the other hand, the rapid devel-

pments of rechargeable batteries are stimulating the explosive

rowth in interest in developing advanced analytical techniques

o explore the fundamental mechanisms of batteries. This will un-

oubtedly further push the limits for the use of synchrotron tech-

iques in rechargeable batteries research, and stimulate the devel-

pment of more sophisticated synchrotron techniques and reliable

n situ cells. To fulfill the wide requirements of batteries charac-

erization, it will be an very important direction to develop syn-

hrotron techniques with excellent temporal and spatial resolution

o investigate batteries over multiple time scale under equilibrium

nd non-equilibrium conditions, and multiple length scales from

he microscopic (such as electronic structure of materials, atomic,

olecular and crystal structure of materials) to the macroscopic

batteries and systems). Synchrotron techniques with high preci-

ion and trace sensitivity chemical identification will also be very

mpotent for the accurate characterization of charge and discharge

roducts.

In recent years, significant attentions have been paid to the

ombination of complementary techniques with different functions

n order to obtain an integrated picture of battery systems, al-

hough synchrotron based techniques are powerful enough to be

sed solely with electrochemical systems. We emphasize here that

he combination of not only different synchrotron based tech-

iques but also with no synchrotron based techniques (such as

MR, TEM and MS) will play more and more important role in the

ully understand battery systems. It has become increasingly ob-

ious that synchrotron based techniques will play a more impor-

ant role to reveal the unknown in battery science, with the con-

inuously improvement of synchrotron facilities and experimental

echniques.

cknowledgments

We are grateful to the sponsors of the research in our group

n the last decades: the National Natural Science Foundation of

hina (Grant nos. 21233004 , 21303147 and 21473148 , etc.), and

he National Key Research and Development Program (Grant no.

016YFB0901500 ).

http://dx.doi.org/10.13039/501100001809


R

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