BATTERIES FOR ADVANCED

TRANSPORTATION

TECHNOLOGIES

(HIGH-ENERGY BATTERY)

2001

ANNUAL

PROGRESS

REPORT

U.S. Department of Energy

Energy Efficiency and Renewable Energy

Office of Transportation Technologies
ACKNOWLEDGEMENT

We would like to express our sincere appreciation to Lawrence

Berkeley National Laboratory, to Argonne National Laboratory,

and to Sentech, Inc., for their artistic and technical contributions

in preparing and publishing this report.

In addition, we would like to thank all our program participants

for their contributions to the programs and all the authors who

prepared the project abstracts that comprise this report.
U.S. Department of Energy

Office of Advanced Automotive Technologies

1000 Independence Avenue, S.W.

Washington, D.C. 20585-0121

FY 2001

Progress Report for the

Batteries for Advanced Transportation Technologies

(High-Energy Battery)

Energy Efficiency and Renewable Energy

Office of Transportation Technologies

Office of Advanced Automotive Technologies

Energy Management Team

Raymond A. Sutula Energy Management Team Leader

December 2001

FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

iii

CONTENTS

INTRODUCTION............................................................................................................................ 1

TASK AREA 1: CELL DEVELOPMENT..................................................................................... 3

TASK AREA 2: Anodes ................................................................................................................. 7

TASK AREA 3: Electrolytes .......................................................................................................... 17

TASK AREA 4: CATHODES........................................................................................................ 37

TASK AREA 5: DIAGNOSTICS................................................................................................... 47

TASK AREA 6: MODELING........................................................................................................ 63

LIST OF ACRONYMS.................................................................................................................... 71

ANNUAL REPORTS ...................................................................................................................... 73
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

iv
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

1

INTRODUCTION

The Batteries for Advanced Transportation Technologies (BATT) Program is supported by the

U.S. Department of Energy Office of Advanced Automotive Technologies (DOE/OAAT) to help

develop high-performance rechargeable batteries for use in electric vehicles (EVs) and hybridelectric vehicles (HEVs). The work is carried out by the Lawrence Berkeley National Laboratory

(LBNL) and several other organizations, and is organized into six separate research tasks.

Background and Program Context

The development of an advanced battery for automotive applications is a difficult undertaking.

There is a strong need to identify and understand performance and lifetime limitations to help guide

battery scale-up and development activities. High cell potentials and demanding cycling

requirements lead to chemical and mechanical instabilities, which are important issues that must be

addressed. The BATT Program addresses fundamental issues of chemistries and materials that face

all lithium battery candidates for DOE EV and HEV applications. The Program emphasizes

synthesis of components into battery cells with determination of failure modes, coupled with strong

efforts in materials synthesis and evaluation, advanced diagnostics, and improved electrochemical

models. The selected battery chemistries are monitored continuously with periodic substitution of

more-promising components. This is done with advice from within the BATT Program, from

outside experts, and from assessments of world-wide battery R&D. The BATT Program also

educates battery and electrochemical scientists who move on to work for battery developers.

Task Descriptions

The six primary BATT Program task areas are: (1) Cell Development, (2) Anodes, (3)

Electrolytes, (4) Cathodes, (5) Diagnostics, and (6) Modeling. Task 1 comprises cell fabrication,

testing and characterization, Tasks 2-4 are aimed at identifying new materials, and Tasks 5-6 support

all BATT Program work. Brief summary descriptions of each task follow.

The Cell Development task has identified three ìbaselineî rechargeable Li cell chemistries. The

polymer-electrolyte cell chemistry includes a Li negative electrode, Li(CF3SO2)2N + cross-linked

poly(ethylene) oxide (PEO)-based polymer electrolyte, and V6O13 or a tunnel-structure LixMnO2

positive electrolyte. The gel-electrolyte cell chemistry includes a natural graphite negative electrode

or a high-capacity tin-based electrode with acceptable stability, LiBF4 + cross-linked gel electrolyte,

and a LiFePO4 or Li1.02Al0.25Mn1.75O3.92S0.03 positive electrode. The baseline Li-ion chemistry is

based on the DOE Advanced Technology Development (ATD) Program ìGeneration 2î chemistry:

a negative electrode composed of graphite + poly(vinylidene) fluoride (PVDF) binder, a LiPF6 +

ethylene carbonate-ethyl methyl carbonate (EC-EMC) electrolyte, and a positive electrode

composed of LiAl0.05Ni0.80Co0.15O2 + graphite + acetylene black + PVDF binder. A web page is

under development for reporting cell test data.

The Anodes task seeks to characterize and improve graphitic and other carbon materials, as well

as conduct exploratory research on non-carbonaceous anode materials. Every commercial Li-ion

battery uses some form of carbon as its anode material, and these batteries suffer from safety, cycle

life, and storage-life problems. It is for these reasons that either improved carbons or non-
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

2

carbonaceous anodes must be developed as possible alternatives. Low-cost metal alloys with

acceptable capacity, rate, cyclability, and calendar life are under investigation.

Polymer Electrolyte research aims to understand performance characteristics by studies of the

transport properties of the electrolyte as a function of polymer and salt structure, polymer structural

changes as a function of temperature, and interactions at the electrode/electrolyte interface related to

transport and chemical/mechanical stability. A multi-pronged approach involving chemical

synthesis, advanced diagnostic tools, and coordinated modeling studies is being used. The

development of non-flammable electrolytes (NFEs) and novel electrolyte additives is recognized as

a critical need for Li-ion battery technology. We seek to identify advanced NFEs and electrolyte

additives to determine their effectiveness for liquid-based Li-ion batteries.

The identification and development of novel Cathodes is critical because of the fundamental

limitations of cobalt-based and vanadium-based materials used in present-day rechargeable Li

batteries. For example, Co is priced at $140/kg, which is inconsistent with EV and HEV cost

targets. The focus of this effort is to develop a high-rate and stable MnO2 cathode. Although

manganese is a low-cost constituent, MnO2 cathodes tend to lose capacity at an unacceptable rate.

Research is directed at understanding the reasons for the capacity fade and developing methods to

stabilize this material, as well as the evaluation of novel forms of MnO2 cathodes.

Advanced Diagnostics are essential to investigate life-limiting and performance-limiting

processes in batteries. We use post-test analyses and enhanced spectroscopic and microscopic

techniques to investigate morphology, structure, and compositional changes of electrode materials.

Examples include providing better understanding of the solid electrolyte interphase (SEI) layers

formed on carbon through ellipsometric techniques, and a detailed investigation of the Li/polymer

interface via advanced microscopies and spectroscopies.

Sophisticated Modeling is required to support BATT Program Tasks 1-5. This effort brings

physical understanding to complex interactions through the development of comprehensive

phenomenological models. Models are being advanced to elucidate the failure mechanisms of Li

battery components and to understand the mechanisms for thermal runaway.

This report (originally published in June 2001 as the Final Report for CY 2000) summarizes the

research activities and accomplishments of the BATT Program in FY 2001. The BATT Program is

a reorganized version of the Exploratory Technology Research (ETR) Program. ETR Program

reports for prior years have been published; and they are listed at the end of this Program Summary.

A website for the BATT Program, which provides Internet links to recent quarterly and annual

reports, is found at http://www.berc.lbl.gov/BATT/BATT.html.

Tien Q. Duong

Program Manager

Office of Advanced Automotive Technologies

Office of Transportation Technologies

Department of Energy.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

3

TASK AREA 1: CELL DEVELOPMENT

Cell Fabrication and Testing

Kathryn A. Striebel

Lawrence Berkeley National Laboratory, 70-108B, Berkeley CA 94720

(510) 486-4385, fax: (510) 486-7303; email: kastriebel@lbl.gov

Objective

• Evaluate baseline new materials for high-power and low-cost Li-ion cells.

Approach

• Prepare uniform electrodes from novel battery materials, assemble them into baseline cells with proven

components and a reference electrode, and test them with a standard protocol.

• Disassemble the cells, after testing, for analysis by the diagnostic tasks within the BATT Program.

Accomplishments

• Visited Argonne National Laboratory (ANL), Quallion Corp. and Mine Safety Appliances to evaluate

techniques and materials required for pouch cell manufacture.

• Developed and used a test protocol for the evaluation of 640 mAh commercial Li-ion cells provided by

DooYoun Corp. in Korea.

Future Directions

• Prepare pouch cells from commercial cell components and LBNL-prepared electrodes in liquid and gel

electrolyte formulations for testing with set protocols.

• Receive and test novel materials from other projects within the BATT Program.

The testing of novel materials in a standard cell

with preset protocols will provide the necessary link

between the invention of novel battery components

and the diagnostic evaluation of failure modes, and

will accelerate the development of a battery-powered

EV. This task is a collaborative effort, which

includes the development of an appropriate test

vehicle (cell geometry), electrochemical cycling, and

cell disassembly. Novel components will be

developed in BATT Program Tasks 2, 3, and 4

(anodes, electrolytes, and cathodes) for baseline

liquid and gel electrolyte chemistries. These

components are incorporated into a standardized cell

with appropriate ratios of components (i.e., active

materials/electrolyte volume) and tested with a

consistent protocol to determine capacity, energy,

power, and lifetime characteristics. Components are

then delivered to the scientists who carry out BATT

Program diagnostics. Fabrication and testing of the

third baseline cell, Li metal/polymer/V6O13, will be

carried out in the electrolyte Task.

The test protocol was developed for 640 mAh

cells received from Doo Youn Corp., Korea. The

protocol involves characterization tests to establish

actual vs. rated capacity, open circuit voltage (OCV)

vs. depth of discharge (DoD), and power capability.

The cells are then cycled at different temperatures

with constant current (C/2) charge and discharge

until less than 60% of the rated capacity remains,

with periodic interruptions to monitor power

performance.

A decision was made to pursue the pouch cell as

a test vehicle for both Li-ion liquid and gel
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

4

electrolyte cells. In the case of the liquid electrolyte

cells, external compression will be applied. This

vehicle lends itself to the preparation of cells of

varying sizes and precise determination of

electrolyte volumes. In addition, there should be

good flexibility on the location of reference

electrodes and cell disassembly is straightforward.

This is a new project (started October 2000) and

the equipment necessary for the production of

laboratory pouch cells is being assembled/ordered.

Site visits were made to Quallion Corp. in Los

Angeles, Mine Safety Appliances in Maryland, and

ANL in Chicago to gather input into this process.

A preliminary web site has been established at

http://Battdata.lbl.gov/ for the posting of data

generated in this task. The data will be contained in

a database for ease of maintenance.

Active Materials Characterization Using X-Ray Diffraction and Chemical Analysis

Thomas J. Richardson

Lawrence Berkeley National Laboratory, 62-203, Berkeley, CA 94720

(510) 486-8619, fax: (510) 486-8619, e-mail: tjrichardson@lbl.gov

Objectives

• Determine the products and mechanisms of deterioration of active electrode materials that result in failure or lost

capacity.

• Accelerate evaluation of electrode materials for in-house cell development by structural characterization of active

components.

Approach

• Analyze cycled electrodes removed from BATT Program cells to ascertain phase transformations in the bulk of the

active material and accumulated inactive decomposition products.

Accomplishments

• Recorded and analyzed baseline X-ray diffraction (XRD) patterns for commercial cells received from LBNL cell

fabrication and testing projects; no new phases were detected in cells that failed short-term, high-temperature

cycling tests.

• Design and fabrication of hardware for in situ XRD studies were completed.

• Prepared and characterized some alternative cathode materials.

Future Directions

• Examine BATT and commercial battery electrodes subjected to long-term cycling.

• Establish presence or absence of inactive phases and decomposition products in cycled cells.

• Carry out exploratory studies on potentially useful alternative electrode materials in support of cell fabrication and

cathode development tasks.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

5

A small number of electrodes from LiCoO2

commercial cells that have completed cycling tests

were received and analyzed by XRD. No new

phases were detected in the cycled cathodes, as

expected from the nature of the cycling tests

performed to date.

Sampling hardware for in situ XRD experiments

has been designed and built. Planar cells with

beryllium windows will be studied in reflection

mode.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

6
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

7

TASK AREA 2: ANODES

Carbon Electrochemistry

Kim Kinoshita

Lawrence Berkeley National Laboratory, 90-1142, Berkeley CA 94720

(510) 486-7389, fax: (510) 486-4260; email: k_kinoshita@lbl.gov

ïObjective

Identify a low-cost carbon (<$10/kg) for negative electrodes in Li-ion batteries for Evs and HEVs.

Approach

• Explore chemical procedures to develop low-cost carbons for Li-ion batteries.

• Use analytical and spectroscopic techniques such as high-resolution electron microscopy, scanning electron

microscopy, Raman spectroscopy, X-ray diffraction analysis, and thermal analysis to determine the physicochemical

properties of carbon.

Accomplishments

• Investigated the intercalation rate of Li+-ions in flake natural graphite with particle sizes that ranged from 2 to 40

μm.

• Investigated how flake natural graphite with a particle size of 12 μm may provide the optimum combination of

reversible capacity and irreversible capacity loss in the electrolyte and discharge rates used in this study.

Future Directions

• Continue collaboration with Superior Graphite and HydroQuÈbec to identify a low-cost carbon with improved rate

capability in Li-ion batteries.

• Identify a chemical procedure that is less expensive than thermal processes to produce purified natural graphite for

Li-ion batteries..

The intercalation rate of Li+-ions in flake natural

graphite with particle sizes that ranged from 2 to 40

μm was investigated. The amount of Li+

 ions that

intercalated at different rates was determined by

measurement of the reversible capacity during deintercalation in 1 M LiClO4/1:1 (volume ratio)

ethylene carbonate-dimethyl carbonate (EC-DMC).

The key issues in this study were the roles of particle

size and the fraction of edge sites on the rate of

intercalation and de-intercalation of Li+

 ions.

The working electrode was fabricated from a

mixture of natural graphite and PVDF dissolved in

1-methyl-2-pyrrolidinone (NMP). The slurry was

spread onto a copper grid (Exmet Corp., Naugatuck,

CT) and dried under vacuum at 95°C for 24 h. A

typical electrode thickness was about 200 μm, and

the electrode area was 12 cm2

 (6 cm2 per side). The

working electrodes were evaluated in a threeelectrode cell that contained metallic Li as both the

counter and reference electrodes. The cells were

built in a glove box under an Ar atmosphere

containing less than 5-ppm water. Electrochemical

measurements of the charge/discharge of graphite

were conducted in cells cycled between 2.5 and 0 V

vs. Li/Li+

.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

8

Figure 1 shows representative charge-discharge

profiles for NG2 and NG30 at a discharge rate of

46.5 mA/g (C/8 rate) and a charge rate of 31 mA/g

(C/12 rate). The profiles show that NG2 discharged

for 8 h before the potential reached 0 V, and the

corresponding time for NG30 was 6 h. The time for

charge of NG2 was equal to the rate (C/12), whereas

the time for NG30 was less (8 h) because the Li+

-ion

content in the intercalated graphite was less. At the

higher discharge rates, the potential of the samples

reached 0 V much sooner than the time indicated by

the C-rate. Because of the slow transport rate of Li+

ions in the graphene layer planes, the time is too

short for the ions to completely intercalate the larger

graphite particles to the theoretical limit. Slow solidstate diffusion of Li+

 ions produces concentration

polarization in graphite that drives the potential

towards 0 V, and this cut-off potential is reached

before intercalation is complete.

0

0.5

1

1.5

2

2.5

3

0 5 10 15 20 25

NG30

NG2

Potential (V vs Li +/Li)

Time (h)

Figure 1. Charge-discharge profiles for NG2 and NG30

at a discharge rate of 46.5 mA/g (C/8 rate) and a charge

rate of 31 mA/g (C/12 rate).

The composition of intercalated graphite, x in

LixC6, which is obtained with natural graphite of

different particle size is presented in Fig. 2 as a

function of discharge rate. At low specific currents,

the composition of lithiated graphite approaches the

theoretical value, x = 1, except for natural graphite

with the largest particle size. However, x decreases

with an increase in specific current for all particle

sizes. This trend suggests that solid-state diffusion

of Li+

 ions is limiting the intercalation capacity in

graphite. At this time, we have no good explanation

for the higher reversible capacity of NG12 compared

to NG7. The particle size distribution of the

particles in an electrode has an influence on the

electrode capacity at a given intercalation rate. The

presence of a small fraction of large particles in a

composite electrode with a majority of monodisperse

small particles lowers the overall intercalation

capacity. The particle size distribution of the natural

graphite was not considered in the present study, but

this factor could play a role in the scatter observed in

the data.

0

0.2

0.4

0.6

0.8

1

0 50 100 150 200 250

NG2

NG7

NG12

NG20

NG30

NG40

x in Li xC6

Specific current (mA/g)

Figure 2. Composition of intercalated graphites, x in

LixC6, vs. discharge rate of natural graphite. Parameter is

particle size.

The flake natural graphite with a particle size of

12 μm may provide the optimum combination of

reversible capacity and irreversible capacity loss in

the electrolyte and discharge rates used in this study.

PUBLICATIONS

W. Jiang, T. Tran, X. Song and K. Kinoshita,

ìThermal and Electrochemical Studies of Carbons

for Li-Ion Batteries. 1. Thermal Analysis of

Petroleum and Pitch Cokes,î J. Power Sources,

85, 261 (2000).

T. Tran, B. Yebka, X. Song, G. Nazri, K. Kinoshita

and D. Curtis, ìThermal and Electrochemical

Studies of Carbons for Li-Ion Batteries. 2.

Correlation of Active Sites and Irreversible

Capacity Loss,î J. Power Sources, 85, 269 (2000).
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

9

W. Jiang, G. Nadeau, K. Zaghib and K. Kinoshita,

ìThermal Analysis of the Oxidation of Natural

Graphite ñ Effect of Particle Size,î Thermochim.

Acta, 351, 85 (2000).

K. Zaghib, G. Nadeau and K. Kinoshita, ìEffects of

Graphite Particle Size on Irreversible Capacity

Loss,î J. Electrochem. Soc., 147, 2110 (2000).

S.E. Sloop, J.K. Pugh, S. Wang, J.B. Kerr and K.

Kinoshita, ìChemical Reactivity of PF5 and LiPF6

in Ethylene Carbonate/Dimethyl Carbonate

Solutions,î Electrochem. and Solid-State Lett., 4,

A42 (2001).

T. Tran and K. Kinoshita, ìSurface Modifications

for Carbon Lithium Intercalation Anodes,î U.S.

Patent 6,096,454, August 1, 2000.

.

Optimization of Anodes for Li-Ion Batteries

M. David Curtis and Gholam-Abbas Nazri

University of Michigan, Department of Chemistry, Ann Arbor MI 48109-1055

(734) 763-2132; fax: (734) 763-2307; e-mail: mdcurtis@umich.edu

Objectives

• Develop new surface coatings for carbonaceous anodes in order to minimize electrolyte decomposition and enhance

Li-ion conduction for application in advanced Li batteries.

• Improve safety aspects of Li-ion technology by incorporating electron- and ion-conducting polymers with fireretardant properties in the construction of carbonaceous anodes.

Approach

• Establish a baseline for thermal behavior of carbonaceous anodes at various states of charge, and compare thermal

and electrochemical properties of the modified electrodes with the baseline data.

• Simulate lithiated graphite anodes by chemical intercalation and evaluate thermal behavior of the modified surfacecoated graphitic anodes.

• Evaluate reactivity of modified surface-coated anodes toward electrolyte decomposition, gas generation, and surface

film formation.

• Improve surface coating and binder technology for large-surface-area electrodes.

Accomplishments

• Completed the determination of thermal stability and reactivity of various graphitic anodes as a function of the

electrolyte composition and degree of lithiation. These data constitute the baseline behavior against which further

modifications of anode surface behavior will be compared.

• Synthesized and evaluated new lithium phosphonate materials for use as artificial SEI coatings on carbon anodes.

• Developing new polymers that are both ionic and electronic conductors and that can serve simultaneously as anode

surface passivators and binder.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

10

Future Directions

• Use the new polymer systems as described above to construct test cells for comparison against our baseline data.

• Optimize the overall performance of graphitic anodes for high-power and high-energy applications.

Carbonaceous anodes are currently the material

of choice for use in Li-ion batteries. However, these

materials react with the electrolyte to produce

flammable organic gases during initial charge

cycles. This reaction also produces a gelatinous

surface coating (SEI layer) that eventually passivates

the carbon surface towards reaction with the

electrolyte. Although this coating has the beneficial

property of passivation, the coating also increases

the internal impedance of the cell, leads to

irreversible capacity loss in the first cycles, and can

cause mechanical instability and loss of electrical

contact between the carbon particles.

We have established a baseline data set against

which future modifications of the carbon anode can

be compared. This baseline includes the nature

and amounts of the various gases that are released

during the initial passivation reactions, as well as

the thermal stability and chemical compositions of

the SEI, and the nature of the sites on the carbon

particles that lead to deleterious electrolyte

decompositions. The thermal stability of

chemically prepared LiCx electrodes was

determined as function of Li content.

Most of the work in the past year has been

focused on the synthesis and testing of new Liphosphonates for use as artificial SEI coatings

and/or binders for carbon anodes. It is our working

hypothesis that the active sites on the carbon

particles can be blocked from reaction with

electrolyte, thus diminishing or eliminating the

release of flammable gases. These coatings must

conduct Li ions in order that the internal impedance

not be increased unduly. Furthermore, if the coating

can act as binder, a step in the anode manufacturing

process could be avoided. For high-power

applications, the binder itself should be

electronically conductive in order for all carbon

particles to be in electrical contact with each other

and with the current collector. We are therefore now

focusing our efforts on the development of new

binder-SEI coatings that will simultaneously

passivate the carbon surface, allow facile transport

of Li ions, and serve as an electron-conducting

binder.

The types of materials that are under

development are conjugated polymers (electrically

conducting in the doped or charged state) that have

pendant ionic groups to facilitate Li-ion transport.

Examples of these polymers are shown in the

following figures.

O O

x y

SO3Li

S S

N

SO3Li

O O

x N y

Figure 3.

x y

N

S S

N

SO3Li

S

N N

S

O SO3Li

O

O

O

Figure 4.

PUBLICATION

T. Tran, B. Yebka, X. Song, G. Nazri, K. Kinoshita

and D. Curtis, ìThermal and Electrochemical

Studies of Carbons for Li-Ion Batteries. 2.

Correlation of Active Sites and Irreversible

Capacity Loss,î J. Power Sources, 85, 269

(2000).
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

11

Non-Carbonaceous Anode Materials

Michael M. Thackeray

Chemical Technology Division, Argonne National Laboratory, Argonne IL 60439

(630)-252-9183; fax: (630)-252-4176; email: thackeray@cmt.anl.gov

Objective

• Replace carbon with an alternative anode material that is inexpensive and that will improve the safety of Li-ion

cells.

Approach

• Search for, characterize, and develop inexpensive intermetallic electrodes that provide an electrochemical potential a

few hundred millivolts above that of metallic Li with capacities >400 mAh/g and 1000 mAh/ml.

• Focus on nickel-arsenide and zinc-blende-type structures, and study their structural and electrochemical behavior

during discharge and charge in Li cells.

Accomplishments

• Improved the performance of copper-tin electrodes by the use of metal-substituted compounds Cu6-xMxSn5 (M=Fe,

Ni, Zn); achieved a rechargeable capacity in excess of 300 mAh/g for the early cycles.

• Studied InSb electrodes in detail and established the importance of invariant face-centered-cubic (fcc) intermetallic

electrode structures.

Future Directions

• Improve the processing and optimize the compositions of intermetallic electrodes, particularly those containing

copper and/or tin.

• Interact with industry to evaluate the intermetallic electrodes in customized cells.

• Continue to explore intermetallic compounds in which there is a strong crystallographic relationship between parent

and lithiated structures.

Considerable progress was made during the past

calendar year in understanding the electrochemical

behavior of intermetallic electrodes and the

structural changes that occur upon reaction with Li.

Progress has also been made in improving the

electrochemical properties of these materials, with

major focus on NiAs-type and zinc-blende-type

structures. In particular, our research was directed

toward the exploitation of the Cu-Sn system by

investigating the substituted derivative structures

Cu6-xMxSn5, in which M=Mn, Fe, Co, Ni, and Zn.

We demonstrated that electrodes with substituted

compositions, such as Cu5NiSn5 and Cu5FeSn5,

showed superior capacities and cycling stability

compared to the unsubstituted parent electrode

Cu6Sn5, when cycled between 1.2 and 0.2 V vs.

metallic Li. For example, a rechargeable capacity

>350 mAh/g was obtained during the early cycling

of Cu5FeSn5 electrodes between 1.2 and 0.01 V vs.

Li. When we restricted the lower voltage of Li/Cu6-

xMxSn5 cells to 0.2 V, the structure of the lithiated

product (lithiated zinc-blende-type) exhibited a

strong crystallographic relationship to that of the

parent electrode (nickel arsenide-type). This

relationship is believed to contribute to the good

electrochemical cycling of these cells.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

12

The realization that the zinc-blende framework

of a lithiated copper-tin electrode (Li2CuSn)

provided a three-dimensional interstitial space for Li

led to a detailed investigation of the family of III-V

semiconductors, MSb (M=Al, Ga, In), as possible

negative electrodes for Li-ion batteries. Preliminary

investigations showed that InSb attains significantly

superior electrochemical behavior over AlSb and

GaSb. Our research has thus focused on the InSb

system, which can deliver a rechargeable capacity of

~300 mAh/g at a C/18 rate.

Two ternary phases in the Li-In-Sb system were

identified previously at 400°C by Sitte and

Weppner; they lie on the InSb-Li3Sb tie line [LizIn1-

z/3Sb (0≤z≤3)] in the Li-In-Sb phase diagram (Fig. 5)

and have compositions close to Li1.5In0.5Sb (z=1.5)

and Li2In0.33Sb (z=2).

The electrochemical reaction of Li with InSb is

analogous to the reversible high-temperature

reaction of Li with Fe3O4 (spinel-type structure)

during which Li is inserted into, and iron extruded

from, an invariant fcc oxygen array to generate a

series of ternary Li-Fe-O compounds. InSb has a

zinc-blende structure in which both the In and Sb

atoms are arranged in independent fcc arrays that are

interlinked to form a diamond-like lattice, each In

atom being tetrahedrally coordinated to a Sb atom,

and vice versa. There are two crystallographically

independent interstitial tetrahedral sites for Li: one

site is coordinated to four Sb atoms, and the second

to four In atoms. During the discharge of Li/InSb

cells, in situ X-ray absorption fine structure data

showed that 1) the reaction occurs by Li insertion

into, and In extrusion from, InSb to yield electrode

compositions LixIn1-ySb (0<x≤3, 0<y≤1), and 2) not

all the Li is reincorporated back into the fcc Sb array

during charge. In the Li3Sb structure, the Li atoms

occupy the In sites and all the tetrahedral interstitial

sites of the InSb structure. Therefore, a strong

structural relationship exists between the parent InSb

electrode and its fully lithiated product, i.e., the

Li3Sb structure can be regarded as having a lithiated

zinc-blende structure Li2(LiSb) in which (LiSb)

refers to the zinc-blende framework. What is

remarkable about this structural relationship is that

the unit cell parameter of pure Li3Sb (Fm3m,

a=6.572 Å) is only 1.4% larger than that of InSb

(Fd3m, a=6.478 Å), reflecting only a 4.4% increase

in the volume of the fcc Sb lattice, thereby

contributing to the reversibility of the electrode.

When the volume of the extruded In is taken into

account, the Li3Sb/In electrode is 46.5% larger in

volume than InSb. Lithiation of the extruded In

metal occurs below ~0.55 V to yield a series of LixIn

compounds (0<x≤4).

The approach we adopted to identify

intermetallic electrodes by considering structural

relationships between parent and lithiated

compounds has opened up new avenues for our

research, and it became the underlying theme of our

project in 2000; it will continue to be so during

2001.

PUBLICATIONS

J.T. Vaughey, J. O'Hara and M.M. Thackeray,

ìIntermetallic Insertion Electrodes with a ZincBlende-Type Structure for Li Batteries: A Study

of LixInSb (0≤x≤3),î Electrochem. and Solid State

Lett., 3(1) 13 (2000).

J.T. Vaughey, K.D. Kepler, C.S. Johnson,

T.Sarakonsri, R. Benedek, J.O. O'Hara, S.

Hackney and M.M. Thackeray, ìIntermetallic

Insertion Electrodes for Lithium Batteries,î The

Electrochem. Soc. Inc., Proc. 99-24, 280 (2000).

M.M. Thackeray, J.O. Thomas and M.S.

Whittingham, ìScience and Applications of Mixed

Conductors for Lithium Batteries,î Materials

Research Society Bulletin, 39 (March, 2000).

C.S. Johnson, J.T. Vaughey, M.M. Thackeray, T.

Sarakonsri, S.A. Hackney, L. Fransson, K.K.

Edstrˆm and J.O. Thomas, ìElectrochemistry and

In-Situ X-ray Diffraction of InSb in Lithium

Batteries,î Electrochem. Comm., 2 (8), 595

(2000).

C.H. Chen, J.T. Vaughey, A.N. Jansen, D.W. Dees,

A.J. Kahaian, T. Goacher and M.M. Thackeray,

ìStudies of Mg-substituted Li4+xMgxTi5O12 Spinel

Electrodes (0≤x≤1) for Lithium Batteries,î J.

Electrochem. Soc., 148, A102 (2001).

PRESENTATIONS

C.S. Johnson, J.T. Vaughey and M.M. Thackeray,

ìCycling Performance of Intermetallic Insertion

Electrodes in Lithium Batteries,î 4th Chicago
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

13

Battery Workshop, Illinois Institute of

Technology, Chicago, IL, April 13-14, 2000.

J.T. Vaughey, C.S. Johnson, A.J. Kropf, R.

Benedek, M.M. Thackeray, H. Tostmann, T.

Sarakonsri, S. Hackney, L. Fransson, K. Edstrˆm

and J.O. Thomas, ìStructural and Mechanistic

Features of Intermetallic Materials for Lithium

Batteries,î 10th International Meeting on Lithium

Batteries, Lake Como, Italy, 28 May - 2 June,

2000.

C.S. Johnson, J.T. Vaughey, D.W. Dees, A.N.

Jansen, M.M. Thackeray, T. Sarakonsri and S.A.

Hackney, ìElectrochemistry and Impedance

Analysis of Intermetallic Electrodes for Lithium

Batteries,î 10th International Meeting on Lithium

Batteries, Lake Como, Italy, 28 May - 2 June,

2000.

J.T. Vaughey, C.S. Johnson, R. Benedek, M.M.

Thackeray, A.J. Kropf, H. Tostmann, L. Fransson,

K. Edstrˆm, J.O. Thomas, T. Sarakonsri and S.A.

Hackney, ìA New Class of Intermetallic Anodes

for Lithium Batteries,î Gordon Conference on

Solid State Chemistry, NH, July 31-August 4,

2000.

J.T. Vaughey, C.S. Johnson, R. Benedek, M.M.

Thackeray, L. Fransson, K. Edstrˆm and J.

Thomas, ìMechanistic and In-situ X-ray Studies

of Intermetallic Electrodes for Lithium-Ion

Batteries,î 198th Meeting of the Electrochemical

Society, Phoenix, AZ, October 22-27, 2000.

M.M. Thackeray, J.T. Vaughey and C.S. Johnson,

ìRecent Developments in Anode and Cathode

Materials for Lithium-Ion Electrochemical Cells,î,

Materials Research Society Meeting, Boston, MA,

November 27-December 1, 2000.

M.M. Thackeray, J.T. Vaughey and C.S. Johnson,

ìIntermetallic Electrodes for Lithium Batteries,î,

DOE Workshop on Interfaces, Phenomena and

Nanostructures, in Lithium Batteries, Argonne, IL,

December 11-13, 2000.

Li

In Sb InSb

Li4In

LiIn

ìLixInî 

phases

Li3Sb

Li2Sb

Line of compensation 

LizIn1-z/3Sb

Line of reaction 

representing the 

overall composition 

ìLixInSbî > 650 mV

< 650 mV ìLi2In0.33Sbî

ìLi1.5In0.5Sbî

Figure 5. The Li-In-Sb phase diagram.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

14

Materials: Novel Anodes

M. Stanley Whittingham

Chemistry and Materials Research Center, State University of New York at Binghamton, Binghamton, NY 13902-6000

(607) 777-4623, fax: (607) 777-4623; e-mail; stanwhit@binghamton.edu

Objective

ï Replace the presently used carbon anodes with safer materials that will be compatible with manganese oxide

cathodes and associated electrolytes.

Approach

• Explore, synthesize, characterize, and develop inexpensive materials that have a potential ~500 mV above that of

pure Li and higher volumetric energy densities than carbon.

• Evaluate all materials electrochemically in a variety of cell configurations, and for thermal and kinetic stability.

Accomplishments

• Determined that vanadium and manganese oxides, in their highest oxidation states, are not prime candidates for

anode materials, despite reports from France and Japan.

• Found pure aluminum to have a high capacity and to react readily with Li, but its capacity faded rapidly upon

cycling.

• Identified some simple binary alloys of Al and of Sn.

Future Directions

• Improve the electrochemical performance of the materials identified, and complete a survey of all Al-based binary

systems and their reactivity with Li.

• Identify several additional non-Al binary alloys.

• Evaluate Mn-containing alloys, as they should be immune from poisoning by Mn migration from the cathode.

The goal of this project is to identify low-cost,

low-weight anode materials that are safer than the

presently used carbonaceous materials and are

compatible with next-generation cathode materials.

Although carbonaceous materials are being

successfully used in Sony Li-ion cells, there are

safety issues as the capacity increases. This project

is taking two approaches: use of oxides giving

essentially a simple symmetric cell:

Li1+yMnO2 |electrolyte| Li1-yMnO2

or use of simple metal alloys, such as those based on

Al:

LiAlMy |electrolyte| Li1-yMnO2

We studied several manganese oxides including

the spinel and layered phases. In each case the

material could be reduced to Li2O and Mn based on

the Li utilization. However, little capacity was

retained after recharge to 1 V vs. Li, so the effort

was switched to mixed manganese-vanadium oxides

and to vanadium oxides. French and Japanese

workers (Solid State Ionics 139, 57, 2001) have

studied MnV2O6 and found positive results but with

a relatively low storage capacity, 400Ah/kg, in the 0

to 1.0 V range. MnV2O5, a material first synthesized

by us, was expected to have a higher storage

capacity because of its lower oxygen content.

However, we found very little capacity even after

extensive mechanical grinding. This low capacity

combined with the need to find a means of initially

charging the material without wasting the active

cathode material makes such oxides unattractive as
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

15

anodes. We are therefore discontinuing work on

these manganese-vanadium oxides as anodes.

However, in the process of this work, a new lowtemperature route for the synthesis of LiMn2O4 was

found. We can now prepare both the spinels and

layered phases of LixMnO2 below 200°C.

Aluminum is the ideal anode material, being low

cost, readily available, and forming a simple alloy

with lithium, LiAl. This reaction occurs at about 0.3

V above pure Li on an essentially flat plateau as

shown in Fig. 6. However, the material expands

significantly upon Li incorporation causing break-up

of the anode. This combined with the brittleness of

the LiAl causes contact problems during cycling,

resulting in relatively poor cycling behavior as

shown in Fig. 7. Even so it was the first anode used

in rechargeable Li batteries, being coupled with TiS2

(Exxon-1978).

The data shown in Fig. 6 were obtained in a

clamped Swagelokô design cell wherein pressure is

applied continuously to the electrodes. When a

higher-capacity flooded-electrolyte bag cell is used

the capacity falls even more rapidly. Clearly cell

construction is critical for the behavior of anode

materials. This is due to the need to maintain

pressure on the particulate matter so that optimum

contact is maintained as the volume changes. We

have obtained a button cell set-up, which will allow

for more reproducible cell systems.

0.0

0.1

0.2

0.3

0.4

0.5

-0.2 0 0.2 0.4 0.6 0.8 1 1.2

Formation of LiAl at 1 ma/cm2

Cell Voltage

x inLi xAl

Figure 6. Incorporation of Li into Al at 1 mA/cm2 in the

cell Li ||electrolyte| Al.

0.0

0.2

0.4

0.6

0.8

1.0

-0.2 0 0.2 0.4 0.6 0.8 1 1.2

Cell Voltage

x in Li xAl

123 4 Cycle

Figure 7. Cycling of Al metal in a clamped Swagelok

cell.

A number of Al alloys were prepared and their

ability to take up Li is being evaluated. Preliminary

data for two Ni and Si samples are shown in Fig. 8.

Both show some capacity above 1 V suggesting

reaction with oxide films or possible electrolyte

reaction. Such reactions are not observed with Al

itself. The Ni reduced the capacity of Li at room

temperature and therefore did not merit further

study. The capacity of the Si-containing sample was

not significantly better than Al itself and its

rechargeability was worse. The addition of Cu or

Mg to Al also resulted in lower capacities.

We investigated two other binary Al

compounds, aluminum carbide Al4C3 and aluminum

boride AlB2, for their ability to react with Li. We

hope that by investigating such compounds the large

volumetric changes occurring during the reaction

between pure Al and Li can be ameliorated.

0.0

0.2

0.4

0.6

0.8

1.0

1.2

1.4

-0.2 0 0.2 0.4 0.6 0.8 1 1.2 1.4

Cell voltage

x in Li xAl

1-y

My

Al

Al0.88Si0.12

Al0.75Ni0.25

Figure 8. Electrochemical insertion of Li into Al,

Al1-ySiy, and Al1-yNiy.

Figure 9 shows the behavior of aluminum

carbide, Al4C3, a combination of two known anodeactive materials. The reaction potential of 1 V is
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

16

higher than either single component. The carbide

reacts with 3 Li per formula unit during discharge

between 0.8 and 0.05 V vs. pure Li giving an initial

capacity significantly higher than that of graphitic

carbon, but slightly less than that of pure Al. The

rechargeability of this material was poor with less

than 1 Li recovered upon charge, just as for the

aluminum silicon 'alloy'. We will test in the button

cell configuration so these cells can be kept under

compression.

0.0

0.5

1.0

1.5

2.0

2.5

-0.5 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5

Cell Volts

x in Li xAl4C3

Aluminum Carbide, Al 4C

3

Figure 9. Electrochemical insertion of Li into Al4C3.

Figure 10 shows the behavior of aluminum

boride, AlB2, in a low-capacity Swagelokô cell. It

reacts with 0.8 Li per formula unit, and up to half of

this could be recovered when charging the cell. The

initial capacity translates to 430 Ah/kg, i.e., slightly

higher than that of graphite, 340 Ah/kg. We will

evaluate this system in a button cell. Little change

in the XRD diffraction pattern was observed after

reaction.

0.0

0.5

1.0

1.5

2.0

2.5

3.0

0 100 200 300 400 500 600

Aluminum Boride

Cell Voltage

Capacity, mAh/g, of AlB 2

Figure 10. Electrochemical insertion of Li into

aluminum boride.

We began an investigation of some simple Snbased materials, and we are confirming the much

better behavior of Sn2Mn over at least the initial

cycles. Such a material is particularly attractive as it

should not be poisoned by Mn transported from the

cathode.

PUBLICATION

S. Yang, P.Y. Zavalij and M.S. Whittingham,

ìSynthesis and Crystal Structure of Li2Sn(OH)6,î

Acta Cryst., C57, 228-229 (2001).

PRESENTATION

S. Yang and M.S. Whittingham, ìAluminum Based

Anode Materials for Lithium Batteries,î Materials

Research Society Meeting, Boston MA, November

27-30, 2000.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

17

TASK AREA 3: ELECTROLYTES

Electrolytes: R&D for Advanced Lithium Batteries

John B. Kerr

Lawrence Berkeley National Laboratory, 62-203, Berkeley, CA 94720

(510)-486-6279, fax (510)-486-4995; email: jbkerr@lbl.gov

Objectives

• Determine the feasibility of the Li metal electrode with organic electrolytes and provide operating conditions that

prevent dendrite growth.

• Determine the limitations on Li-ion transport in polymer electrolytes and composite electrodes and develop new

materials capable of ambient-temperature operation with Li metal electrodes.

• Determine the limits of stability of organic electrolytes at high-voltage cathode materials (4 V) and develop

materials and methods to increase stability.

• Provide cross-linkable polymers for use in the Li/polymer and gel polymer test systems and scale up synthesis and

processing to 100 cm2

 pouch cells.

Approach

ï Apply a combination of synthesis, analysis, modeling and testing, i.e., a physical organic chemistry approach to

electrolyte design, thereby ensuring that not only are the sources of poor performance and failure pinpointed but

the problem can be corrected through as-developed materials design and synthesis capabilities.

Accomplishments

• Conducted experiments which demonstrated that the mechanical properties of the electrolyte play a significant role

in dendrite growth at Li metal electrodes.

• Developed methods to separate the mechanical properties from the ion-transport properties to provide high-rate

Li/polymer batteries.

• Prepared polymer electrolytes that contain ion-solvating ether groups which are different from the commonly used

ethylene oxide units, and observed significant increases in conductivity, that show promise for room-temperature

performance of Li/polymer batteries.

Future Directions

• Carry out detailed quantitative measurements of the transport and mechanical properties of polymer electrolytes to

optimize transport vs. interfacial behaviors to minimize dendrite growth.

• Prepare and optimize polymers with new ion-solvating groups and appropriate architectures to achieve ambienttemperature operation with solid polymer electrolytes.

• Develop polymer electrolyte systems capable of operation with 4-V cathodes and Li metal electrodes.

• Develop quantitative measures to characterize electrolyte side reactions.

• Develop synthetic capabilities to provide materials and processing for larger-scale Li/polymer and polymer gel test

systems.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

18

Links between dendrite growth and

mechanical and transport properties. Lithium

cycling experiments in symmetrical cells with linear

polymer electrolytes based on PEO have clearly

demonstrated that dendrites are formed rapidly if the

limiting current is exceeded, as expected for

electrode position of metals. The limiting current

for the Li electrode is calculated from knowledge of

the transport properties of the electrolyte (diffusion

coefficient, transference number) and the dimensions

of the cell (area, thickness). However, even with

current densities well below the calculated limiting

current, rapid dendrite growth is observed with

linear polymer electrolytes. These polymers are

essentially viscous liquids with a linear chain

structure. Cross-linked linear and comb-branch

polymer structures have been shown to possess

properties of solids (see 1999 Annual Report).

Significant improvements in the cycling behavior

were observed with increasing cross-link density and

hence higher elastic moduli. Figure 11 shows the

behavior of a comb-branch polyepoxide ether

electrolyte which has been cross-linked with a 10%

density. Cycling is continued for over 100 cycles

and a total of 150 coulombs of charge is passed

before the cell suddenly fails due to a short-circuit

arising from a dendrite. With a polymer that has a

cross-link density of only 5% the cycle life is only

15 cycles (22 coulombs). Without cross-linking,

dendrite growth is observed very quickly.

Quantitative study of these effects in addition to

study of side reactions is under way in order to allow

prediction of safe operating conditions that prevent

growth of dendrites.

Polymers with improved ion-transport

properties. New polymer structures have been

prepared with conductivity that is superior to what

has been achieved to date with PEO-based polymers.

In collaboration with modeling groups (M. Ratner,

L. Curtis and J. Halley), who provided theoretical

guidance, the polymer structures have evolved from

linear PEO-like materials to comb-branch or

dendritic polymers with side chains that contain

ether side groups that are more labile than those

based on ethylene oxide units. Among these groups

are 1,3-propanediol units. The effect of

incorporating this unit into the polymers is shown in

Fig. 12. These results show that the theoretical

understanding of ion transport through the polymer

hosts has advanced greatly in recent years and that

room-temperature solid polymer electrolytes are

within reach.

-0.2

-0.1

0

0.1

0.2

0.3

0.4

0.5

-100 0 100 200 300 400 500

Li/Li cycling of PEPE2-10%xl - 30TFSI

estimated thickness 40um

cycle number 1-2

cycle number 55-56

cycle number 102-103

Potential (V)

Time(Min)

Current Density : 200uA/cm 2

Charging Time: 2hr

Measurement Temperature : 85o

C

Figure 11. Dendrite growth during cycling of symmetrical Li/Li cells.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

19

-9

-8

-7

-6

-5

-4

-3

-2

0.0025 0.003 0.0035 0.004 0.0045

Conductivities of Comb-Branch and linear Polymers with LiTFSI salt. 

PEPExPx denotes polymers that contain 1, 3-propanediol links.

PEMO-20

PEA5-20

PVE5-20

PEPE5-20

HQ data -30 TFSI

PEPExPx-LiTFSI(30:1)#2

PEPE1P1-LiTFSI (20:1)

PEPE1P1-LiTFSI(10:1)

PEPExPx #1

Log Conductivity (S/cm)

1/T

Linear polymers

with EO units Comb polymers

with EO units

Comb polymers

with PDO units

Figure 12. Conductivity measurements for linear and comb branch polymers

based on ether solvating groups.

PUBLICATIONS

L. Edman, A. Ferry and M.M. Doeff, ìSlow

Recrystallization in the Polymer Electrolyte

System P(EO)nLiTFSI,î J. Mater. Res., 15, 1950

(2000).

L. Edman, M.M. Doeff, A. Ferry, J. Kerr and L.C.

De Jonghe, ìTransport Properties of the Solid

Polymer Electrolyte System P(EO)nLiTFSI,î J.

Phys. Chem. B., 104, 3476 (2000).

M.M. Doeff, L. Edman, S.E. Sloop, J. Kerr and L.C.

De Jonghe, ìTransport Properties of Binary Salt

Polymer Electrolytes,î J. Power Sources, 89, 227

(2000).

K.E. Thomas, S.E. Sloop, J.B. Kerr and J. Newman,

ìComparison of Lithium-Polymer Cell Performance with Unity and Nonunity Transference

Numbers,î J. Power Sources, 89, 132-138 (2000).

O. Buriez, Y.B. Han, J. Hou, J.B. Kerr, J. Qiao, S.E.

Sloop, M.M. Tian and S.G. Wang, ìPerformance

Limitations of Polymer Electrolytes Based on

Ethylene Oxide Polymers,î J. Power Sources, 89,

149-155 (2000).
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

20

Composite Polymer Electrolytes for Use in Lithium and Lithium-Ion Batteries

Saad A. Khan*, Peter S. Fedkiw and Gregory L. Baker+

Department of Chemical Engineering, North Carolina State University, P.O. Box 7905, Raleigh NC 27695-7905;

+

Department of Chemistry, Michigan State University, East Lansing, MI 48824-1322

*Phone: (919) 515-4519; Fax: (919) 515-3465; e-mail: khan@eos.ncsu.edu

Objective

• Develop composite polymer electrolytes (CPEs) that are low-cost, have high conductivities, impart electrodeelectrolyte interfacial stability, and yield long cycle life.

Approach

• Use surface-functionalized fumed silica fillers in BATT-baseline systems to determine the effects of filler type and

concentration on interfacial stability and cell cycling.

• Utilize a combination of electrochemical characterization, rheological techniques, and chemical syntheses to

correlate electrochemical characteristics with mechanical properties and materials chemistry (e.g., silica-type or

PEO-type).

Accomplishments

• Determined that the presence of fumed silica in CPEs increases cell capacity and decreases capacity fading in fullcell cycling of Li/LixCoO2 and Li/LixMn2O4 couples.

• Established effects of fumed silica surface groups on Li/LixCoO2 cell capacity and cycling.

• Synthesized and measured the conductivity of fumed silica with dual functionalities containing different ratios of

alkyl to cross-linkable groups.

• Studied the effects of cross-linkable CPE constituents on the interfacial resistance between Li and CPE.

Future Directions

• Determine how fumed silicas (hydrophobic R805 and hydrophilic A200) affect full-cell cycling using vanadiumbased, 3-V cathodes and baseline PEO-type materials in coin cells.

• Establish the relationships between interfacial stability and mechanical properties.

The objective of this research is to develop a

new range of CPEs for use in rechargeable Li and

Li-ion batteries. In particular, our goal is to develop

highly conductive electrolytes that exhibit good

mechanical properties, and at the same time show

good compatibility with typical electrode materials.

The unique feature of our approach is the use of

surface-functionalized fumed silica fillers to control

the mechanical properties of the electrolytes and

enhance electrode-electrolyte interfacial stability.

Full-cell cycling studies of Li/LixCoO2.

We have previously reported that fumed silica-based

CPEs with low-molecular-weight PEOs exhibit

conductivities exceeding 10-3 S/cm at 25°C, have

electrochemical properties (Li transference number,

conductivity) decoupled from mechanical properties,

and thus provide a range of mechanical moduli (as

high as 106

 Pa). Furthermore, the presence of fumed

silica enhances Li-electrolyte interfacial stability.

Beside the above-mentioned electrochemical and

mechanical properties, a long cycle life with a high

capacity is also crucial for the electrolytes of

rechargeable batteries. Full-cell cycling with fumed

silica-based CPEs using Li anodes and metal oxide

cathodes has been carried out.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

21

The effect of fumed silica and its surface

chemistry on full-cell cycling was studied using a

standard LiCoO2 composition (91 wt% LiCoO2, 6

wt% graphite SFG15, and 3 wt% PVDF), shown in

Fig. 13. Three types of fumed silica were used:

native fumed silica A200 with hydroxyl groups,

octyl-modified R805, and PEO-modified FS-EG3.

Cell capacity quickly fades after the first few cycles

in the absence of fumed silica, but the addition of 10

wt% particulates diminishes the capacity fade, with

the effect dependent upon the silica surface

chemistry. The native fumed silica (A200) exhibited

the best improvement of capacity fading, whereas

the PEO-modified fumed silica (FS-EG3) showed

the least improvement.

Cycle Number

0 2 4 6 8 10 12

Discharge Capacity (mAh/g LiCo

O2 )

0

20

40

60

80

100

120

140

160 0% FS

10% R805

10% A200

10% FS-EG3

 LiCoO2 + 6% Graphite SFG15 + 

Solvent: PEG-dm (Mw 

Salt: Li Imide 

C rate: C/40

Cathode: 91% 

Figure 13. Effect of fumed silica and its surface

chemistry on cell capacity of Li/electrolyte/LiCoO2

cycling.

Full-cell cycling studies of Li/LixMn2O4.

Figure 14 shows the effect of octyl-modified fumed

silica on the discharge capacity of Li/LixMn2O4

cycling. The beneficial effect of fumed silica is

independent of cathode materials; as with LixCoO2,

fumed silica increases the discharge capacity and

lowers capacity fading relative to the base

electrolyte. Such independence gives us the

flexibility to extend the application of our

composites to a variety of rechargeable Li and Li-ion

batteries with diverse cathode materials.

Conductivity of cross-linkable CPEs. The

conductivities of crosslinked and uncrosslinked

CPEs with trichlorosilyl octyl methacrylate (TOM)

1:1, TOM 4:1, and TOM 1:4 fumed silica [X:Y

refers to X moles of methacrylate groups and Y

moles of octyl group in Fig. 15] are shown as a

function of temperature and fumed silica surface

group. The conductivity of the cross-linked CPEs is

independent of the fumed silica surface group and

decreases by a factor of two after cross-linking. This

suggests that the crosslinking occurs between

adjacent fumed silica particles and does not affect

the bulk electrochemical properties of the

electrolyte. Consequently, the selection of salt and

polymer matrix can be varied to tune the

electrochemical properties. Finally, the roomtemperature conductivity of the PEGdm(250) CPEs

is close to 10-3 S/cm.

Cycle Number

0 2 4 6 8 10 12

Specific Capacity (mAh/g LiMn

O2 4)

0

20

40

60

80

100

120

140

160 10% R805 (C/40) 

w/o fumed silica (C/40)

Solvent: PEG-dm (250)

Salt: Li Imide (Li:O=1:20)

Cut-off Voltage: 4.2 V (charge)

 3.0 V (discharge)

Cathode: 91% LiMn2O4 + 6% graphite SFG15 + 3% PVDF

Theoretical capacity:

148 mAh/g

Figure 14. Effect of octyl-modified fumed silica (R805)

on Li/LixMn2O4 full-cell cycling.

1000/T (K-1)

2.5 3.0 3.5

Conductivity (S cm-1) 10-4

10-3

10-2

T (°C)

100 82 67 52 37 25 12 0

10% TOM FS

PEGdm (250) 

LiImide (Li:O = 1:20)

10% TOM FS +

20% BMA + 1% AIBN X-linked

TOM 4:1

TOM 1:1

TOM 1:4

Figure 15. Conductivity of cross-linkable CPEs

containing PEGdm(250) and various dual-functionalized

fumed silicas (TOM 1:1, TOM 4:1, and TOM 1:4).

Interfacial stability studies of Li/CPE

interface for cross-linkable CPE components.

An important concern in CPE development is the

formation of a stable interface with Li. One of our
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

22

concerns is the possible interaction of the monomer

and initiator with Li metal and developing possible

avenues for mitigating this problem. We have

previously reported the remarkable success of

interface stabilization by addition of fumed silica to

the polyethylene glycol-dimethy ether (PEG-dme)-

based electrolytes. However, the presence of

reactive species such as butyl methacrylate (BMA)

and the initiator azo-bis-isobutyronitrile (AIBN)

may pose a problem. Figure 16 shows the effect of

the cross-linkable CPE constituents on interfacial

resistance. The addition of TOM 1:1 stabilizes the

interface. The interfacial stabilizing effect of TOM

1:1 is less than that of R805 FS (data not shown).

Recall that TOM 1:1 has 25% of its surface covered

with methacrylates. Although the TOM 1:1 is less

able to stabilize the interface, the interfacial

resistance is still better than that of the base

electrolyte. The addition of BMA significantly

increases the interfacial resistance. AIBN did not

statistically increase the interfacial resistance

relative to the base electrolyte. This issue of

constituents affecting the interface merits further

study along with development of techniques to

mitigate the unfavorable interaction of BMA with

Li.

Time (Days)

0 20 40 60 80

Rint (Ω cm2

) 102

103

104

105

PEGdm(500)+LiImide(1:20) [Base Electrolyte]

" + 10% TOM 1:1

" + 10% TOM 1:1 + 1% AIBN

" + 10% BMA

Figure 16. Time-dependent stability of Li/CPE interface

for various uncross-linked CPE compositions.

Current Research Focus. We have been in a

transition period since October 2000, and we are

now focusing on the BATT-baseline system. We

have initiated our research on full-cell cycling using

vanadium-based, 3-V cathodes. V6O13 from KerrMcGee, graphite SFG15 from Timcal, and PVDF

binder from Kynar are used as our cathode materials.

We have been working on establishing a standard

procedure and optimizing compositions of vanadium

oxide cathodes. We initiated full-cell cycling at

0.08 mA/cm2

 using vanadium oxide cathodes with

liquid electrolyte, 10 wt% R805, and 10 wt% A200

composites.

The mechanism of ion transport and how to tune

the electrochemistry of an electrolyte remains an

area rich in opportunity and poor in consensus

among researchers. We continued our study of the

fundamentals of ion transport using techniques in

nuclear magnetic resonance (NMR), rheology, and

electrochemistry. Our research agrees with the

overall view that the more motion a polymer chain

has the better the total charge transport. However, at

extremely high concentrations (those corresponding

to about one polymer chain per LiTFSI) the charge

transport is not optimum but the Li transference

number approaches 0.5 due to a restriction of anion

mobility. The lower conductivity may be acceptable

under certain conditions given the higher Li

transference number.

PUBLICATIONS

S.R. Raghavan, J. Hou, G.L. Baker and S.A. Khan,

ìColloidal Interactions between Particles with

Tethered Nonpolar Chains Dispersed in Polar

Media: Direct Correlation between Dynamic

Rheology and Interaction Parameters,î Langmuir,

16(3), 1066-1077 (2000).

S.R. Raghavan, H.J. Walls and S.A. Khan, ìSilica

Dispersions in Organic Liquids: Evidence of

Solvation Forces Dictated by Hydrogen-bonding,î

Langmuir, 16 (21), 7920-7930 (2000).

H.J. Walls, J. Zhou, J.A. Yerian, P.S. Fedkiw, S.A.

Khan, M.K. Stowe and G.L. Baker, ìFumed

Silica-Based Composite Polymer Electrolytes:

Synthesis, Rheology, and Electrochemistry,î J.

Power Sources, 89, 156-162 (2000).

H.J. Walls and T.A. Zawodzinski, Jr, ìAnion and

Cation Transference Numbers Determined by

Electrophoretic NMR of Polymer Electrolytes

Sum to Unity,î Electrochem. Solid-State Lett.,

3(7), 321-324 (2000).

J.A. Yerian, S.A. Khan and P.S. Fedkiw, ìCrosslinkable Fumed Silica-Based Composite

Electrolytes for Rechargeable Lithium,î Proc.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

23

198th Meeting of the Electrochem. Soc., Phoenix,

AZ, October 2000.

J. Zhou, P.S. Fedkiw and S.A. Khan, ìInterfacial

Stability of Fumed-Silica Based Composite

Polymer Electrolytes for Rechargeable Lithium

Batteries,î Proc. 198th Mtg. of the Electrochem.

Soc., Phoenix, AZ, October 2000.

M.K. Stowe, J.Hou and G.L. Baker, ìCrosslinkable

Composite Polymer Electrolytes: Ionic

Conductivity and Polymerization Behavior Using

Methacrylate Monomers,î Proc. 198th Mtg. of the

Electrochem. Soc., Phoenix, AZ, October 2000.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

24

Highly Conductive Polyelectrolyte-Containing Rigid Polymers

Duward F. Shriver* and Semyon Vaynman**

*Chemistry Department, Northwestern University, Evanston, IL 60208

**Department of Materials Science and Engineering, Northwestern University, Evanston, IL 60208

(847) 491-4475; fax: (847) 491-1022; e-mail: svaynman@northwestern.edu

Objectives

• Synthesize a new class of rigid polymer electrolytes.

• Test rigid polymer electrolytes in rechargeable Li batteries.

Approach

• Synthesize new types of polymer electrolytes that contain a rigid polymer rather than the flexible low-Tg polymers

used in conventional polymer electrolytes.

• Fabricate electrochemical cells with these electrolytes and evaluate their performance.

• Correlate the performance of electrolytes in the cells with their chemical structure and their reactivity toward

components of the electrochemical cell.

Accomplishments

• Synthesized highly conductive rigid polymer electrolytes that contain functional groups such as carboxy and

sulfone, and discovered complexes containing carboxy groups have higher ionic conductivity than a similar

polymer-salt complex containing sulfone (~10-4 S/cm vs. 5x10-6 S/cm at room temperature).

• Conducted tests on symmetrical Li/polymer electrolyte/Li cells that indicated the formation of a very resistive

interface when a carboxy-containing rigid polymer electrolyte was used, but a much less resistive layer was formed

between Li metal and sulfone-containing rigid polymer electrolyte.

• Incorporated rigid polymer electrolytes into Li cells, and found that the capacity of the cell with carboxy-containing

rigid polymer electrolyte was not satisfactory; the capacity of the cell with sulfone-containing rigid polymer

electrolyte was more than 100 mAh per gram of active cathode material at current density of 15 μA/cm2

, but the

capacity of the cell was reduced significantly with an increase in current density due to high resistance of that

electrolyte.

• Synthesized new sulfone-containing rigid polymer electrolytes with a higher density of cation-coordinating sites.

Future Directions

• Modify rigid polymer electrolyte to increase its ionic conductivity and to reduce its reaction with anode.

• Test modified electrolyte in Li cells.

During 1999 two rigid polymer systems that

contain carboxy groups were developed in our

laboratory: poly(vinylene carbonate) (PVIC) and

poly(1,3-dioxolan-2-one-4,5-diyl oxalate)

(PVICOX), which displayed both favorable

conductivity and mechanical properties. These

polymers contain a high density of coordinating

sites, which is necessary to dissolve salts. Also, the

high density of polar groups reduces the activation

energy for ion hopping from one polar site to the

next. These properties lead to high ionic

conductivity of the salt-polymer system. PVICOX

has ionic conductivity of 10-4 S/cm at room

temperature. A polymer electrolyte containing

PVICOX was incorporated into Li cells. The

capacity of the cell was low due to the formation of

a very resistive interface between the Li anode and

electrolyte.

Because PVICOX reacted with the Li anode, we

synthesized polymers that contain sulfone groups.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

25

Sulfone or sulfoxide-functionalized organic

molecules have some of the highest dipole moments

and dielectric constants known and display very

good cation-complexing properties. Therefore, it

was expected that polymers containing sulfuroxygen bonds should readily dissolve significant

amounts of Li salts and provide rapid ion transport.

Furthermore, due to the strongly electronwithdrawing character of these groups, their

polymers should have very good anodic stability.

The synthesis of the polyarylsulfone (a) was

carried out by poly-condensation of Na-4,4ísulfonyldiphenoxide with 3,4-dichlorosulfolane (Fig.

17).

A phase transfer catalysis (PTC) which involves

a two-phase system of mutually insoluble aqueous

and organic layers (liquid-liquid PTC) was

employed. The transport of the ionic reactant from

the aqueous layer into the organic phase takes place

via a lipophilic transfer agent, tetrabutyl ammonium

bromide.

The synthesized polymer was mixed with

lithium triflate in 1:1 and 4:1 molar ratios. The

mixtures were dried for 72 h in vacuum (6x10-6 Torr)

and the ionic conductivity was measured. The

results presented in Fig. 18 indicate lower ionic

conductivity than we expected by analogy with our

highly conductive rigid polymer electrolyte

PVICOX. As with PVICOX, a higher lithium

triflate concentration leads to higher conductivity.

This trend in conductivity is opposite of that for

conventional polymer-salt complex electrolytes.

Unlike typical polymer electrolytes these new

materials are rigid at room temperature.

As reported previously, PVICOX is not

chemically stable when in contact with Li metal; the

resistance of the cell increases by a couple of orders

of magnitude in the first few hours due to formation

of a layer between PVICOX and Li. According to

Fig. 19, a polysulfone (a)-containing polymer

electrolyte is more stable toward Li metal than

PVICOX. After 64 h, the resistance of the cell only

triples.

Cells that contained Li metal, cathode material,

and the rigid polymer electrolyte polysulfone (a)

mixed with lithium triflate (1:1), were assembled

and tested. The cathode consisted of 55% LiMn2O4

(from Kerr-McGee), 35% polymer electrolyte

(MEEP-lithium triflate (4:1)), and 10% carbon (from

Superior Graphite). A capacity of up to 100 mAh/g

of active cathode material was achieved at current

density of 15 μA/cm2 (Fig. 20). However, the

capacity dropped significantly with an increase in

current density to 40 μA/cm2

, due to high cell

resistance.

Another sulfone-containing polymer

(polysulfone (b)) is being synthesized in our

laboratory at the present time using 1,4-dithiane-2,5-

diol and dibromosulfolane as starting materials. We

explore different synthetic routes to this polymer

with emphasis on high yield. We anticipate that

when doped with salt, this polysulfone (b) will

display higher ionic conductivity than that of

polysulfone (a), because the polymer (b) has a

higher density of cation-coordinating sites than (a).

We also expect high stability of polymer (b) toward

Li.

The synthesis of this polymer as shown in Fig.

21 is underway now.

NaOH

NaO S ONa

O

HO S OH

O

O

O

NaO S ON a

O

O S

Cl Cl

O O

OSO

O

O

S n

O O

DMSO

-NaCl +

polysulfone (a)

Figure 17. Synthesis of polysulfone (a).
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

26

1.E-10

1.E-08

1.E-06

1.E-04

2.8 2.9 3 3.1 3.2 3.3 3.4 3.5

1000/T, K

Conductivity, 

S/cm

1 polymer-1 triflate 4 polymers-1 triflate

Figure 18. Ionic conductivity of the polysulfone (a) -

containing electrolyte.

0

5000

10000

15000

20000

25000

0 20000 40000 60000

Real, ohm s

-Im, ohms

0 hrs

1 hr

2 hrs

4 hrs

16 hrs

38 hrs

66 hrs

Figure 19. The impedance of Li/polysulfone (a) electrolyte

system.

0

1

2

3

4

5

0 20 40 60 80 100 120

CAPACITY, mAh/g active material

POTENTIAL, volt

15 microamp/sq. cm

25 microamp/sq. cm

40 microamp/sq. cm

Figure 20. Discharge of of Li/(PVICOX)/ LiMn2O4;C;

MEEP-lithium triflate cell at 25 ̊C.

S

S

HO OH

S

S

NaO ONa +

S

O O

Br Br

S

S

O O

S

O O

n

NaH

S

S

NaO ONa

H2O2

S

S

O O

S

O O

n S

O O

n S

S

O O

O

O

O

O

polysulfone (b)

Figure 21. Synthesis of polysulfone (b).
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

27

Nonflammable Electrolytes

K. Kinoshita

Lawrence Berkeley National Laboratory, 90-1142, Berkeley CA 94720

(510) 486-7389, fax: (510) 486-4260; email: k_kinoshita@lbl.gov

Objective

ï Identify chemical additives that improve the safety of nonaqueous electrolytes for Li-ion batteries by stabilizing

the SEI layer on carbon.

Approach

• Identify species that are incorporated in the SEI layer to improve its stability.

• Conduct electrochemical evaluation of additives to determine the reversible and irreversible capacity loss using the

baseline liquid or gel electrolytes.

Accomplishment

• Studied the reaction of PF5 gas and EC/DMC suggesting that decomposition of LiPF6 to PF5 facilitates the

thermal decomposition of EC/DMC containing LiPF6.

Future Direction

ï Evaluate a chemical additive that enhances the thermal stability of the SEI layer on carbon in a liquid electrolyte

to determine its electrochemical performance compared to baseline systems.

Our experimental effort focused on the thermal

stability of electrolytes and additives, including

flame retardants such as those identified by the

Illinois Institute of Technology (IIT). The thermal

decomposition of 1 M LiPF6 in 1:1 mol ratio

EC/DMC at 85oC was compared to the reaction

between EC/DMC and PF5 gas. Both produce a

brown solution and gaseous products. The

decomposition products from PF5 + EC/DMC and 1

M LiPF6 in 1:1 EC/DMC, which were detected by

gas chromatography (GC), are similar. Proton NMR

provided evidence for the presence of oligoether

carbonates in the reaction products. The reaction of

PF5 gas and EC/DMC suggests that decomposition

of LiPF6 to PF5 facilitates thermal decomposition of

EC/DMC containing LiPF6. Limiting the

decomposition of LiPF6 should improve the safety of

Li-ion batteries. Based on our studies and the results

obtained by IIT, decomposition products from the

electrolyte and flame retardants may improve the

thermal stability of the SEI layer on carbon anodes.

A sample of hexa-methoxy-tri-aza-phosphazene

(HMTAP) was received from IIT, and thermal

stability tests were conducted by heating a baseline

electrolyte [1 M LiPF6 in a mixture of 1:1 EC/DMC]

containing 3% HMTAP at 85oC. After three days at

85oC, the electrolyte remained relatively clear (slight

brown coloration) and no gas was evolved when the

sample vial was opened. In contrast, the baseline

electrolyte without HMTAP showed a visible change

to a dark brown color, and gas evolution was

observed under the same experimental conditions.

The GC results obtained for the thermal

decomposition of the baseline electrolyte (LP30) is

compared with the electrolyte containing 3 wt%

HMTAP (LP30 + HMTAP) in Fig. 22. The plot

shows the GC peak ratios for the reaction product

detected at 14.5 min to that for EC vs. the time at
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

28

which the samples were held at 85oC. Without

HMTAP in the electrolyte (LP30), a significant

increase is observed in the amount of reaction

product detected by GC. On the other hand, when

HMTAP is present (LP30 + HMTAP), EC

decomposition is inhibited. This is due to the

increased basicity of HMTAP over EC; HMTAP

will react preferentially with PF5, which is a Lewis

acid. HMTAP reacts through ring-opening

polymerization (HMTAP-ROP), and the products

are believed to be associated with the GC peak at 33

min. The increase in the concentration of the

decomposition product with time is also shown in

Fig. 22. This analysis suggests that HMTAP inhibits

EC decomposition in LP30 at 85oC. The similar

decomposition rates of HMTAP and EC in

LP30+HMTAP imply that the reaction is limited by

generation of PF5 in the electrolyte; the additive does

not change this rate.

-0.2

0

0.2

0.4

0.6

0.8

1

-50 0 50 100 150 200 250 300

EC and HMTP decomposition products vs. time 

comparison of LP30 and LP30 + HMTP

EC product @ 14.5min : EC (LP30)

EC prod @14.5 min : EC (LP30 + HMTP)

HMPT prod. @ 33min :EC (LP30 + HTMP)

Time / hr

Figure 22. The concentration of reaction products as a

function of time for LP30 and LP30+HMTAP at 85oC.

Proton NMR studies of the thermal reaction

products from EC/DMC/LiPF6 showed peaks due to

DMC, EC and polyether carbonates

(CH2CH2O)mCOO)n as reported in the literature

[ìThe polymerization of ethylene carbonateî L.

Vogdanis & W. Heitz Macromol. Chem., Rapid

Commun. 1986, 7, 543]. It appears that PF5 causes

polymerization reactions in the electrolyte. Based

on this result, a sample of PE3C (a

polyethercarbonate) is proposed for evaluation as an

additive for Li-ion cells. This material should

simulate the end product that may be produced

during the formation of the SEI layer (electrolyte

decomposition) at the anode. Upon electron transfer

to cyclic carbonates, rearrangement and loss of CO2

can occur, thereby forming a lithium alkoxide. The

alkoxide can react with the carbonyl group and lead

to polyethercarbonate (PEC) products.

A sample of PE3C was given to IIT and

evaluated at IIT as an additive in 2016 coin cells

(Li/carbon) containing 1 M LiPF6 and EC+DEC. A

Gen-1 ATD Program anode (MCMB+ 6% SFG) was

used. The charge (Li+

-ion deintercalation)/discharge (Li+

-ion intercalation) cycles

of two cells (K2 and K3) that contained 7% PE3C

are compared to a cell that contained no PE3C (Fig.

23). Charge/discharge were carried out at the C/20

rate, and the first-cycle irreversible capacity loss of

the cell without additive was 44 mAh/g and those for

K2 and K3 were slightly higher, 57 mAh/g and 55

mAh/g, respectively. The corresponding first-cycle

reversible capacities were 314, 312 and 315 mAh/g,

respectively. Based on these data, it appears that the

presence of the additive in the electrolyte has a very

small effect on the reversible capacity but produces a

higher irreversible capacity loss. The specific

capacity is slightly higher in the cells with additives

after several cycles but this result is not conclusive.

Additional data are being collected. The role of

PE3C in enhancing the thermal stability of the SEI

layer has not been substantiated.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

29

300

310

320

330

340

350

360

370

380

0 2 4 6 8 10

Discharge - 0% additive

Charge - 0% additive

Discharge - K2/7% additive

Charge - K2/7% additive

Discharge - K3/7% additive

Charge - K3/7% additive

Specific capacity (mAh/g)

Cycle No.

Figure 23. Charge/discharge cycling of 2016 coin cell containing baseline electrolyte and PE3C additive.

Development of Nonflammable Electrolytes

J. Prakash

Illinois Institute of Technology, Department of Chemical and Environmental Engineering,

10 W 33rd Street, Chicago, IL 60616

(312) 567-3639, fax: (312) 567-8874; e-mail: prakash@iit.edu

Objective

ï Develop NFEs with high flash point (>100°C), high ionic conductivity (10-3 S/cm), and wider voltage window

(0-5 V vs. Li) in an effort to provide better thermal stability and fire safety.

Approach

• Modify existing electrolytes by using novel flame-retardant (FR) additives that are compatible with active electrode

materials and the environment.

• Use chemical, electrochemical, and thermal techniques to investigate the stability and performance of electrolytes

modified with FR additives

Accomplishments

• Completed an extensive study on the thermal and electrochemical characterization of the FR additive, HMTAP

N3P3[OCH3]6, in Li-ion cells.

• Completed the scale-up synthesis of the flame retardant HMTAP.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

30

• Supplied the HMTAP material to the ATD Program (ANL, PolyStor and LBNL) and Mitsubishi Chemical

Company, Japan, for further testing.

• Demonstrated the feasibility of using this material as a FR additive in PNGV Li-ion cells.

Future Directions

• Develop new FR additives hexa-ethoxy-tri-aza-phosphazene (HETAP) and urea-diphosphonate (UDP) and new Lisalts with high conductivity and high thermal stability to reduce self-heat rate (SHR) especially at higher

temperatures during thermal runaway.

• Perform thermal diagnostic studies of BATT Program baseline Li-polymer and low-cost Li-ion cells with and

without FR additive in order to understand the degradation, failure, and safety mechanisms.

Investigation of N3P3[OCH3]6 FR additive

in Li-ion cells. In 2000, we successfully

demonstrated the scale-up synthesis of the FR

additive (N3P3[OCH3]6) with improved yield and

purity. We synthesized 10 to 15 g batches of highpurity FR additive N3P3[OCH3]6, characterized it for

chemical and electrochemical stability, and supplied

samples to the ATD group (ANL, LBNL and

Polystor Co.) for further investigations.

Optimization of the amount of FR

additive in Li-ion cells. The thermal properties

of the HMTAP in Li-ion cells were investigated in

collaboration with ANL using differential scanning

calorimetry (DSC) and Accelerated Rate

Calorimetry (ARC) methods. Figure 24 shows a

typical self-heat rate profile of the electrolyte with

and without the FR additive in the ARC in presence

of metallic Li. It is evident that the addition of 5-10

wt% FR additive to the electrolyte shifts the

exothermic peaks beyond 260°C. The shift of the

exothermic peak with 5-10 wt% FR additive in

comparison with those for electrolyte without FR

may be attributed to a passivation layer that is

formed on the surface of the Li metal. In addition,

investigations of the pressure development within an

ARC bomb also showed that 5-10 wt% FR additives

significantly reduce the pressure buildup in the

bomb compared with those for electrolyte without

FR under similar conditions.

The effect of FR additive on the thermal

behavior of the fully charged Li-ion cell electrodes

was also investigated. Pre-cycling was carried out in

sealed graphite/LiNi0.8Co0.2 32 cm2

 cells. LiPF6 in

50:50 wt% EC-DMC was used as electrolyte. These

cells were galvanostatically cycled at a C/10 chargedischarge rate. After a few cycles, the cells were

charged at C/10 rate to 4.15 V and held at 4.15 V.

These charged cells were then opened in a dry glove

box filled under Ar atmosphere and the electrode

materials were recovered. The recovered materials

were sealed in ANL-made ARC bombs for the ARC

studies. A typical thermal profile of a fully charged

anode is shown in Fig. 25. It is clear that the

presence of the FR additive in these electrodes shifts

the onset of the exothermic peak to higher

temperature. Thermal investigations suggest that 5-

10 wt% HMTAP additive is necessary to reduce the

SHR and internal pressure.

Figure 24. Temperature profile of the Li-ion electrolyte

with and without HMTAP FR additive.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

31

Figure 25. Effect of the FR additive on the ARC profiles

of the fully charged graphite anode.

In order to investigate the effect of HMTAP

concentration on the performance of the Li-ion cells,

we cycled the sealed Li-ion cells with PNGV Gen-1

chemistry containing 1-10 w% HMTAP. The results

of 5-wt% FR on the cell performance are shown in

Fig. 26 and indicate there is no detrimental effect of

the amount of the FR additive on the Li-ion cell

performance. The results of the present

investigations show that the HMTAP (N3P3[OCH3]6)

is a viable FR for Li-ion cells and the optimum

amount of the FR additive is in the range 5-10 wt%

of the electrolyte weight.

0

0.5

1

1.5

2

2.5

3

3.5

4

4.5

0 25 50 75 100 125 150 175 200

Specific Discharge Capacity, mAh/g

Voltage, V vs. Li/Li+

5 wt% FR_F74_dis

5 wt% FR_F73_dis

5 wt% FR_F72_dis

No FR_dis

Figure 26. Electrochemical performances of a PNGV Liion cell with 5-wt% HMTAP in 1 M LiPF6 with

EC:DEC(1:1).

The shift of about 25°C in the onset temperature

of the thermal runaway (Figs. 24 and 25) is very

encouraging for the safety of Li-ion cells. However,

the SHR at temperatures above 265°C was found to

increase in presence of HMTAP. In order to address

the issue of SHR at higher temperatures and to meet

the milestone (FR additives for Li-ion batteries), we

also initiated studies on other FR additives such as

HETAP and UDP with reduced SHR at higher

temperatures as a possible replacement for HMTAP.

HETAP (N3P3[OCH2CH3]6). We also

successfully synthesized HETAP and characterized

its electrochemical stability in Li-ion electrolyte

using cyclic voltammetry (CV) in the voltage

window of 0-5 V. The new FR additive HETAP

showed good electrochemical stability up to 5 V vs.

Li. Further electrochemical and thermal studies on

this material are currently underway.

PUBLICATIONS

J. Prakash, C. Lee and K. Amine, “A Novel FlameRetardant Additive for Li-ion Batteries,” U.S.

patent application filed 2000.

C. Lee, R. Venkatachalapathy and J. Prakash,

“Development of Non-flammable Electrolytes for

Li-ion Batteries,” Electrochem. and Solid-State

Letters, 3, 63 (2000).

C.W. Lee, R. Venkatachalapathy and J. Prakash,

“Flame-Retardant Materials for Lithium

Batteries,” 197th Meeting of the Electrochemical

Society, Toronto, Canada, Abst. No.120, May

2000.

C.W. Lee, H. Joachin and J. Prakash, “Studies on

Electrochemical and Thermal Properties of

Lithium-Ion Battery Electrolytes with Different

Salts,” 198th Meeting of the Electrochemical

Society, Phoenix, AZ, Abst. No 143, October

2000.

J. Prakash, “Investigations of Nonflammable

Phosphozene Based Compounds for Li-ion

Batteries,” 10th International Meeting on Lithium

Batteries, Como, Italy, Abstr. 347, June 2000.

K. Kinoshita and J. Prakash, “Prospects for

Nonflammable Electrolytes in Li-ion Batteries,”

Proceedings of the 4th Chicago Battery Workshop,

Illinois Institute of Technology, Chicago, IL, April

2000.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

32

Nonflammable Electrolytes for Li-Ion Batteries

Alan B. McEwen

Covalent Associates, Inc., 10 State Street, Woburn, MA 01801

(781) 938-1140; fax: (781) 938-1364; e-mail: mcewen@covalentassociates.com

Objectives

• Develop nonflammable electrolytes for Li-ion batteries.

• Optimize the electrolytes for BATT Program objectives regarding liquid electrolytes and to meet DOE safety goals.

Approach

• Investigate ionic liquids as nonflammable and nonvolatile electrolyte solvents.

• Synthesize new ionic liquids and formulate these into Li-ion electrolytes.

• Determine flammability and flashpoint of the electrolytes.

• Determine the performance characteristics of these new electrolytes by measuring conductivity, Li intercalation, and

cycle life of half-cell and full-cell batteries.

Accomplishments

• Developed NFEs for Li-ion batteries using ionic liquids.

• Demonstrated reversible Li intercalation in anode and cathode materials.

• Constructed coin cell batteries demonstrating cycling performance.

Future Directions

• Synthesize new low-cost ionic liquids and formulate these into Li-ion electrolytes with the addition of Li salts and

common organic solvents.

• Determine electrolyte performance with LiFePO4 and Li1.02Al0.25Mn1.75O3.92S0.03 positive electrodes and Sn-based

negative electrodes.

• Assess the performance of the relatively inexpensive Li salt LiBF4 and LiTFSI in the ionic liquid electrolyte

formulations.

The application of fluorinated ethyl acetates as

electrolyte solvents for nonflammable Li-ion

electrolytes was the original task of this project. The

suitability of these solvents as electrolyte solvents

for Li-ion batteries was demonstrated. Good

conductivity (7.5 mS/cm) and cycling performance

on graphite and LiCoO2 was observed. We have

found, however, that the flashpoint of the fluorinated

ethyl acetate solvents and electrolyte mixtures based

on these solvents is below room temperature (Table

1). Because this fails the criteria of nonflammability

we changed the focus of our research and

development efforts.

We initiated an investigation into using

Covalentís proprietary ionic liquids (U.S. Patent

5,827,602) as Li-ion electrolytes due to the inability

of fluorinated esters to meet the nonflammable

criteria. We shifted the focus of the project to

investigate these novel materials that have a wide

liquid range, high thermal stability, and lack a vapor

pressure. In addition, they are nonflammable and

non-corrosive, giving them a thermal application

range from below room temperature to approaching

400o

C.

Using a NFE, formulated with an ionic liquid,

carbonate solvent, and 1 M LiPF6 salt, we obtained
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

33

good cycling behavior (Fig. 27) in coin cells using

graphite anodes and LiCoO2 cathodes. We cycled

between 3.0 and 4.3 V. These imidazolium based

ionic liquid Li-ion electrolytes have good

conductivity (10 mS/cm) and form stable SEIs on

graphite with the small addition of carbonate

solvents. Although by themselves carbonate

solvents are flammable, in these ionic liquid

formulations there is no flash point or burning in

flammability tests.

Table 1. Flashpoint temperatures (o

C) for

fluorinated acetate electrolytes.

Solvent 1M

LiPF6

1M

LiBeti

1M

LiMethide

Ethyl Acetate < 26 < 26 < 26

TFEAa < 27 < 27 < 27

EA:EC (1:1) < 25 < 26 < 26

TFEA:EC(1:1) < 28 < 27 < 27

TFEA: trifluoroethyl acetate

We demonstrated the reversible intercalation of

Li into graphite (Fig. 28) using an ionic liquid

formulated electrolyte. One can see the three stages

of Li intercalation into graphite. Good reversibility

was also demonstrated.

We also have obtained promising results with a

Li metal/LiCoO2 cell with the same ionic liquid

electrolyte used previously in a Li-ion cell (Fig. 29).

0.00

1.00

2.00

3.00

4.00

5.00

6.00

7.00

1 11 21 31 41 51 61 71 81

Cycle Number

Capacity (mAh)

Charge

Capacity

Discharge

Capacity

Figure 27. Cycling capacity for a coin cell with a

nonflammable Li-ion electrolyte.

-0.1

0.1

0.3

0.5

0.7

0.9

1.1

1.3

1.5

1.7

-0.2

5

0.25 0.75 1.25 1.75 2.25 2.75 3.25 3.75

capacity (mAh/cm2)

E(V)

1st cycle

6th cycle

10th cycle

Figure 28. Cycling results for an ionic liquid based

electrolyte in a Li/graphite cell.

3.00

3.50

4.00

4.50

0 2 4 6 8 10 12

Time (Hr)

Voltage (v)

Figure 29. Voltage profile of a Li/LiCoO2 cell using an

ionic liquid electrolyte.

Excellent results were obtained in a

graphite/LiCoO2 coin cell using an ionic liquid

based Li-ion electrolyte (Fig. 30). As depicted

in Fig. 27 we continue to obtain good cycling

efficiency for over 60 cycles.

2.8

3.3

3.8

4.3

-25 175 375 575 775

capacity (mAh/g)

E(V)

Figure 30. Performance of a coin cell battery with an

ionic liquid electrolyte. First cycle.

Lithium/LiCoO2

Graphite/LiCoO2
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

34

Solvent additive study. The influence that

electrolyte composition has on the capacity and

cycle life in graphite/LiCoO2 coin cells was studied.

We focused on developing a stable SEI with these

ionic liquid NFEs that have good cycle life and

excellent high-temperature stability. At constant Li

salt and ionic liquid composition a correlation

between additive identity and performance was

determined (Fig. 31).

Lithium salt study. We determined the

influence Li salt identity (e.g., LiBeti, LiImide,

LiMethide, or LiPF6) has on cell performance. We

investigated using the LiBeti [LiN(SO2C2F5)2] salt

due to its improved thermal stability relative to

LiPF6 and improved corrosion stability relative to

the LiImide [LiN(SO2CF3)2] salt. We have observed

that the capacity and cycle life is dependent on Li

salt composition (Fig. 32).

0.00

1.00

2.00

3.00

4.00

5.00

6.00

7.00

0 5 10 15 20 25 30 35 40 45

Cycle Number

Capacity (mAh)

A

B

C

D

E

Figure 31. Ionic liquid electrolyte: graphite/LiCoO2 coin

cell cycling dependence on additive formulation.

0.00

1.00

2.00

3.00

4.00

5.00

6.00

7.00

8.00

0 5 10 15 20 25 30 35 40 45

Cycle Number

Capacity (mAh)

1

3

2

4

Figure 32. Ionic liquid electrolyte: graphite/LiCoO2 coin cell cycling dependence on Li salt identity.

Development of Novel Electrolytes

Kraig A. Wheeler

Delaware State University, Department of Chemistry, Dover DE 19901

(302) 739-4934, fax: (302) 739-3979; e-mail: kwheeler@dsc.edu

Objectives

• Develop improved electrolytes for Li-ion batteries.

• Investigate the electrochemical and physical properties of polysulfone cross-linked polyethylene oxide (PEO)

electrolytes.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

35

Approach

• Prepare polysulfone cross-linked PEOís using preparative synthetic organic methods and a homologous series of

polyethylene glycol (PEG) precursors (molecular weight (MW) 400, 2000, and 4600).

• Determine polymer structure and purity using spectroscopic techniques.

• Assess dissolution potential of Li-ion electrolytes with polysulfones.

• Evaluate electrochemical performance of sulfone electrolytes in Li-ion cell applications.

Accomplishments

• Prepared and characterized a homologous series of sulfone cross-linked PEO polymers.

• Achieved effective Li-ion dissolution of LiN(CF3SO2)2 with polysulfone-chloroform mixtures.

• Detected onset of polymer decomposition at >4.0 V in Li-ion cell studies of polysulfones.

Future Direction

ï Complete electrochemical characterization of sulfone materials.

The goals of this project center on developing

polysulfone-based electrolytes for Li-ion cell

applications. Essential components of this research

contribution include the fabrication, characterization,

and structure-property assessment of a homologous

series of Li-ion polysulfone electrolytic materials.

Understanding the influence of the polymeric

skeletal framework and sulfone functionality to

electrochemical stability is of primary importance to

this project.

The fabrication of the target polymers (Fig. 33)

followed a rational design that incorporates both

anionic and radical initiated processes. Selective

deprotonation of appropriate PEG precursors (MW

400, 2000, 4600) under mild basic conditions

followed by subsequent treatment with 3-chloro-2-

(chloromethyl)propene gave the expected

polycondensates as confirmed by spectroscopic

methods. Extended reaction times, elevated

temperatures, and radical initiation processes

resulted in unwanted cross-linking of the polyethers.

The development of alternative methods using in

situ reaction conditions provided highly viscous

materials with minimal cross-linking for each PEG

employed. The radical initiated reaction of

divinylsulfone and these olefinic polyethers formed

controlled cross-linked molecular frameworks that

lack hydroxyl and olefinic functionalities

characteristic of the PEG and dichloropropene

starting materials. Purification of these functional

materials gave a self-standing polymer support with

excellent mechanical properties.

 O PE O O O PE O O O 

S

O

S O

O 

O 

O PE O O O 

Figure 33. Structural topology of sulfone cross-linked

PEO polymers.

Most polymer electrolytic systems with

relatively high crystalline content are widely

recognized as being restricted in the area of iontransport process. Elevated temperatures and

rational structure modifications provide, when

effective, tools to fabricate amorphous polymeric

substrates. The structural design criteria employed

in this study produced polysulfone architectures void

of any significant crystalline content at room

temperature due to the construction of a random

molecular framework produced from cross-linking

the PEO polycondensates and divinyl sulfone.

Our efforts to understand the electrochemical

stability of these sulfone cross-linked polymers build

upon our prior CV studies of similar arylsulfone

systems. CVís were performed at 25°C on each

polysulfone sample using a Pt microdisc (working)

and Li strip (reference and counter) electrodes.

Electrolytic materials were prepared by either

swelling the polycondensate in a 0.1 M

bis(trifluorosulfonyl) imide (LiIm) solution or direct
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

36

solvation of LiIm with the polysulfone. Inspection

of the voltammograms acquired from these studies

revealed the onset of polymer decomposition at ~4.0

V vs. Li and showed stability under reversible

electrochemical sweeps (Fig. 34). The

electrochemical behavior of polysulfone/LiIm

systems in Li-cells shows favorable responses

supported by the lack of severe electrolytic

degradation at elevated voltage potentials.

Figure 34. Voltammogram of MW 400 Polysulfone/LiIm.

PUBLICATION

J.-J. Lee, T. Bae, D.A. Scherson, B. Miller and K.A.

Wheeler, ìUnderpotential Deposition of

Aluminum and Alloy Formation on

Polycrystalline Gold Electrodes from AlCl3/MEIC

Room-Temperature Molten Salts,î J. Electrochem.

Soc. 147, 562 (2000).
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

37

TASK AREA 4: CATHODES

Novel Cathode Materials

Michael M. Thackeray

Argonne National Laboratory, Chemical Technology Division, Argonne IL 60439

(630)-252-9183; fax: (630)-252-4176; e-mail: thackeray@cmt.anl.gov

Objective

• Develop low-cost manganese oxide cathodes to replace vanadium oxide electrodes in Li-polymer cells and

cobalt/nickel oxide electrodes in Li-ion cells.

Approach

• Search for, characterize, and develop low-cost manganese oxides for Li-ion and Li-polymer cells.

• Focus on layered lithium-manganese oxides derived from Li2MnO3 for Li-ion cells, and stabilized α-MnO2 for Lipolymer cells.

Accomplishments

• Investigated the concept of exploiting a layered Li2MnO3 rock salt structure to stabilize layered LiMO2 structures.

• Synthesized and evaluated ìsolid-solutionî or composite electrodes having the general formula

[x(Li2MnO3)•y(LiMO2)]; M = Mn, Ni and Co in Li half-cells.

• Developed optimized compositions which yielded specific capacities of ~150 mAh/g between 4.5 and 3.0 V in Li

cells at 50°C, but for less than 100 cycles.

• Achieved a rechargeable capacity of more than 200 mAh/g for 40 cycles with stabilized α-MnO2 electrodes in

liquid-electrolyte Li cells.

Future Directions

• Investigate alternative (non-aqueous) processing routes to synthesize [x(Li2MnO3)•y(LiMO2)] composite electrodes

to improve their capacity and cycling stability with a performance target of 150 mAh/g for 100 cycles at 50°C in

Li/[x(Li2MnO3)•y(LiMO2)] cells.

• Stabilized α-MnO2 electrodes in Li/polymer cells in collaboration with J. Kerr at LBNL.

A major focus of our research on novel or

improved cathodes for Li-ion and Li-polymer cells

has been the exploitation of a concept to use a

layered Li2MnO3 (rock salt) component to stabilize a

layered LiMnO2 structure during electrochemical

cycling of Li/LiMnO2 cells. In this respect, it is well

known that layered LiMnO2 electrodes transform

irreversibly to a spinel-type structure; the

transformation leads to poor capacity retention and

cycle life. Promising electrochemical behavior has

been achieved with layered lithium manganese oxide

electrodes derived by acid treatment of Li2MnO3.

Such treatment removes lithia (Li2O) from Li2MnO3

and yields a layered Li2-2xMnO3-x phase that shows a

resistance to the transformation to spinel. The Li2-

2xMnO3-x structure can be represented, alternatively,

as (1-x)Li2MnO3•MnO2, in which the MnO2

component retains a layered configuration.

Electrochemical lithiation of this compound in a

non-aqueous Li cell results in the formation of a
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

38

layered product, (1-x)Li2MnO3•LiMnO2, in which

the electrochemically inactive Li2MnO3 can be

viewed as a stabilizing agent for the

electrochemically active MnO2 component. This

finding has immediate implications for using

Li2MnO3 to stabilize other layered LiMO2 structures

in order to improve electrode capacity and cycle life.

The approach outlined above has been used to

evaluate the performance of layered

xLi2MnO3•yLiMO2 (M= Mn, Ni, Co, Cr) electrodes.

In one example, an experiment was conducted to

make Li2MnO3•4LiNi0.8Co0.2O2 in which the

Li2MnO3 component comprises 20% of the

composite structure. Figure 35 shows the voltage

profile obtained from a Li/Li2MnO3•4LiNi0.8Co0.2O2

coin cell at 50°C; Fig. 36 shows the cycling stability

of the cell over the first ten charge and discharge

cycles. The voltage profile suggests that Li is

inserted into, and extracted from, the metal oxide

host structure during discharge in a single-phase

process. Apart from an approximate 25% capacity

loss during the first cycle, the capacity remains

steadily above 150 mAh/g. The extent to which the

Li2MnO3 component plays a role in stabilizing such

layered electrodes, particularly at the top of charge,

is unknown. It is speculated that the presence of the

Li2MnO3 domains, which are structurally compatible

with the LiMO2 domains, are a contributing factor to

electrode stability.

As part of the BATT program, we have recently

initiated work on stabilized α-MnO2 electrodes for

Li-polymer cells. This development effort follows a

DOE/BES-supported research project in which

stabilized α-MnO2 electrodes were identified. Our

early electrochemical data obtained from liquidelectrolyte Li cells show that capacities in excess of

200 mAh/g can be obtained at an average voltage of

2.8 V vs. metallic Li and a C/12 rate.

PRESENTATIONS

C.S. Johnson and M.M. Thackeray, ìCycling

Performance of Stabilized Alpha Manganese

Dioxide Electrode Materials,î 198th Meeting of the

Electrochemical Society, Phoenix, AZ, October

22-27, 2000.

C.S. Johnson and M.M. Thackeray, ìRecent

Developments and Prospects for Layered

Manganese Oxide Electrodes for Lithium

Batteries,î 198th Meeting of the Electrochemical

Society, Phoenix, AZ, October 22-27, 2000.

C.S. Johnson and M.M. Thackeray,

ìElectrochemical Processes of Manganese Oxide

Materials,î DOE Workshop on Interfaces,

Phenomena and Nanostructures, in Lithium

Batteries, Argonne, IL, December 11-13 2000.

1

1.5

2

2.5

3

3.5

4

4.5

80 85 90 95 100

Time in Test (h)

Cell Voltage (V)

D5 = 158 mAh/g

C5 = 164 mAh/g

Li/EC:DEC 1:1, 1 M LiPF6/Li2-xMx'M"O3 (x<0.2; Li+M'+M"=3)

4.5 to 2.0 V, I = 0.1 mA/cm2

, 50 °C 

Figure 35. The voltage profile of a Li/Li2MnO3•4LiCo0.2Ni0.8O2 coin cell.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

39

0

5 0

100

150

200

250

300

0 1 2 3 4 5 6 7 8 9 10 11

C yc le N u m b e r

Specific Capacity (mAh/g)

C h a rg e c a p a c ity

D is c h a rg e c a p a c ity

Figure 36. Specific capacities delivered by a Li2MnO3•4LiCo0.2Ni0.8O2

positive electrode in a Li coin cell over 10 cycles at 50°C (~C/8 rate).

New Cathode Materials Based on Layered Structures

M. Stanley Whittingham

State University of New York at Binghamton, Chemistry and Materials Research Center, Binghamton, NY 13902-6000

(607) 777-4623, fax: (607) 777-4623; e-mail: stanwhit@binghamton.edu

Objective

ï Find lower-cost and higher-capacity cathodes, exceeding 200 Ah/kg, that are based on benign materials.

Approach

• Place emphasis on manganese dioxides, both pure and modified with other transition metals, using

predominantly low-temperature synthesis approaches.

• Synthesize these materials and characterize them structurally for thermal and chemical stability.

• Evaluate stabilized vanadium oxides for Li/polymer cells.

Accomplishments

• Determined that layered manganese dioxides can be structurally stabilized, that their stability is a function of

current density and/or cut-off voltages, that their electronic conductivity can be significantly enhanced, that

their cell cycling can be substantially improved by addition of other transition metals, and that

hydrothermally synthesized manganese oxides cycle as well as high-temperature materials.

• Showed that vanadium oxides can also be stabilized by the addition of Mn ions, and can achieve capacities

over 220 Ah/kg.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

40

Future Directions

• Identify changes in LiMnO2 structure as a function of current density during cell cycling, determine the

structure and composition of the vanadium-stabilized LiMnO2, and increase its electrochemical capacity, to

understand better the behavior of the Li1.2Mn0.4Cr0.4O2 cathode and to determine if there is an iron analog.

• Complete, for Li-polymer cells, the evaluation of the Mn-stabilized δ-vanadium oxides and compare them to

the iron phosphates.

The goal of this project is to identify transition

metal oxides with properties superior to LiCoO2,

including higher capacity, lower cost and being

benign to the environment.

Stabilized manganese oxide cathodes.

Layered LiMnO2, which has the potential of cycling

1 Li per Mn ion, is unstable relative to the spinel

LiMn2O4 during cycling. Two approaches were

used to stabilize the layer structure of manganese

dioxide: i) the incorporation of pinning ions to hold

the layers apart thus minimizing the diffusion of Mn

which is required for spinel formation, and ii) the

partial substitution of some of the Mn ions by later

transition metals such as Fe, Co or Ni.

We have completed a preliminary

electrochemical evaluation of the use of alkali

metals to pin the MnO2 layers apart, combined with

partial substitution of part of the Mn ions by Fe, Co,

or Ni. Last year we reported the enhancement of the

electronic conductivity of these materials by partial

Mn substitution.

Surprisingly we found the stability of these

layered manganese dioxides is strongly affected by

the rate of cycling. Figure 37 shows the cycling of a

hydrothermally formed potassium manganese oxide

at 0.1 and 1 mA/cm2

. At the lower current density

the expected behavior for the layered material is

found. However, at 1 mA/cm2

 the 4 V plateau

typical of spinel is observed even for the first charge.

Thus, it appears that these alkali pillars are being

removed during cycling at high rates/high potentials.

We formed the corresponding lithium

manganese oxides and found the same spinel-like

phase formation occurring upon the first charge as

shown in Fig. 38.

However, we have shown that vanadium oxygen

pillars can be incorporated between the manganese

oxide layers [Pub. no. 5]. These pillars are expected

to be relatively immobile at ambient temperatures.

Figure 39 shows the corresponding cycling data for

these pillared compounds. Clearly in this case the

manganese oxide is maintaining its layered behavior

and is not showing any of the 4-V behavior typical

of spinel-like phases. Thus, the principle of pillaring

to prevent Mn diffusion appears to be correct. The

capacity of the vanadium-oxide pillared manganese

oxide is 150 Ah/kg but drops below 100 Ah/kg upon

increasing the discharge rate to 1 mA/cm2

. Future

efforts will be targeted at improving the capacity of

these pillared compounds at higher rates, comparing

them with the recently reported high lithium

chromium manganese compounds, and the Mnstabilized V2O5.

Stabilized vanadium zxide and

comparison to iron phosphate cathodes. We

earlier reported, [Pub. no. 1], the high-capacity, 220

Ah/kg, of the double-sheet V2O5 structures, and their

structure retention during cycling. Our original Mnstabilized δ-V2O5 compound was a mixture of two

phases. In an attempt to identify the

electrochemically active phase(s) we have succeeded

in preparing a pure phase of one of the two

components. This material, [(CH3)4N]0.2MnyV2O5,

has a lattice repeat distance of 13Å, and contains

only 0.06 Mn per V2O5; it is thermally stable to over

250°C with no structural changes. Its initial capacity

is 220 Ah/kg with a mean discharge potential of 3 V

and very little polarization. However, increasing the

Mn content from 0.06 to 0.1 drops the capacity to

180 Ah/kg as shown in Fig. 37. Researchers at

DuPont are also working on this class of material

and have confirmed the high capacity (198th ECS

Meeting, Phoenix AZ, October 2000).

We initiated studies of the iron phosphate class

of cathodes, which are being actively studied by

Armand et al at the University of Montreal and

HydroQuÈbec, following the initial discovery by the

Goodenough group. We made the Li3Fe2(PO4)3

phase using a hydrothermal approach. Its

electrochemical behavior is shown in Fig. 41, and

although the capacity is not large it is retained well

during cycling. We have also formed LiFePO4

hydrothermally, and if coated with carbonaceous

material will cycle well and at above 150 Ah/kg

capacity (discussions with Armand). We will

actively pursue the hydrothermal approach to

evaluate this class of material as the base case for
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

41

comparison to our cathode materials. LiFePO4 is

reported to have a better rate capability and higher

capacity than LiCoO2, and to improve the

performance of LiCoO2 when admixed with it

(Armand at IBA meeting, March 2001).

Figure 37. Cycling of KyMn0.9Co0.1MnO2 at 0.1 mA/cm2 (left) and 1 mA/cm2 (right).

Figure 38. Electrochemical cycling of LiyMn0.9Co0.1MnO2 and LiyMn0.9Ni0.1MnO2 at 1 mA/cm2.

 

Figure 39. Electrochemical cycling of (VOz)0.1MnO2 at 0.1 mA/cm2 (left) and 1 mA/cm2 (right).

Figure 40. Capacity of Mn-stabilized vanadium oxides, (left) Mn=0.06, and (right) Mn=0.10.

Figure 41. Electrochemical cycling of lithium iron

phosphate, Li3Fe2(PO4)3 (contains Fe2O3 impurity).

PUBLICATIONS

F. Zhang and M.S. Whittingham, ìHydrothermal

Synthesis and Electrochemistry of a δ-Type

Manganese Vanadium Oxide,î Electrochem

Commun., 2, 69-71 (2000).

M.S. Whittingham. P. Zavalij, F. Zhang, P. Sharma

and G. Moore, ìThe Stabilization of Layered

Manganese Oxides for Use in Rechargeable

Lithium Batteries,î Proc. Mater. Res. Soc., 575,

77-82 (2000).
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

42

M.M. Thackeray, J.O. Thomas and M.S.

Whittingham, ìScience and Applications of Mixed

Conductors for Lithium Batteries,î Mater. Res.

Soc. Bulletin, 39-46 (2000).

M.S. Whittingham, ì25 Years of Intercalation

Chemistry for Battery Materials,î Electrochem.

Soc. Proc., 99-24, 7 (2000).

F. Zhang, K. Ngala and M.S. Whittingham,

ìSynthesis and Electrochemistry of a VanadiumPillared Manganese Oxide, Electrochem

Commun., 2, 445-447 (2000).

F. Zhang and M.S. Whittingham, ìElectrochemistry

of the Layered Manganese Dioxides

MxMn1-y(Co,Fe,Ni)yO2, M= Li, K,î Electrochem.

and Solid-State Lett., 3, 309-311 (2000).

M.S. Whittingham and P.Y. Zavalij, ìManganese

Dioxides as Cathodes for Lithium Rechargeable

Cells: The Stability Challenge,î Solid State

Ionics, 131, 109-115 (2000).

M.S. Whittingham, ìInsertion Electrodes as SMART

Materials: The First 25 Years and Future

Promises, Solid State Ionics, 134, 169-178 (2000).

PRESENTATIONS

M.S. Whittingham, ìLithium Intercalation into

Vanadium and Manganese Oxides,î Max-Planck

Worshop, Schlofl Ringberg, Germany, March 17,

2000, .

M.S. Whittingham, ìSoft Chemistry Synthesis of

Vanadium and Manganese Oxide Hosts for

Intercalation Reactions, Chemistry Department,

University of Michigan, Ann Arbor, MI, March

21, 2000,

M.S. Whittingham, ìHydrothermal Synthesis of

Manganese Vanadium Oxides,î American

Chemical Society, San Francisco, CA, March 26,

2000.

M.S. Whittingham, ìThe Stabilization of Layered

MnO2 Relative to Spinel Formation,î IBA 2000,

Argonne National Laboratory, Argonne, IL, May

11, 2000.

M.S. Whittingham, ìThe Control of Structure and

Properties of Vanadium and Manganese Oxides

through Tailored Soft Synthesis,î European

Materials Research Society, Strasbourg, France,

June 1, 2000,

M.S. Whittingham, ìCathode Materials: Manganese

Oxides ñ Overview and Results,î DOE/LBNL

Review, Berkeley, CA, June 6, 2000.

K. Ngala and M.S. Whittingham, ìSynthesis and

Characterization of Manganese Vanadium

Oxides,î Gordon Research Conference, New

London, NH, July 2000.

P.Y. Zavalij and M.S. Whittingham, ìVanadium

Oxide Framework Structures,î Intl. Symposium on

the Reactivity of Solids, Budapest, Hungary,

August 28, 2000.

M.S. Whittingham, ìSoft Chemistry Synthesis of

Vanadium and Manganese Oxide Hosts for

Intercalation Reactions,î Solid State Chemistry

2000, Prague, Czech, September 4-8, 2000.

K. Ngala, P. Y. Zavalij and M. S. Whittingham,

ìPillared Oxides of Manganese and Vanadium for

Lithium Batteries - Geometric and Electronic

Factors,î 198th Electrochemical Society Meeting,

Phoenix, AZ, October 24, 2000.

M.S. Whittingham, ìManganese Dioxides for

Lithium Batteries: The Stability Challenge,î Japan

National Battery Meeting, Nagoya, Japan,

November 20ñ22, 2000.

P.Y. Zavalij and M.S. Whittingham, ìVanadium

Oxide Frameworks Modified with Transition

Metals,î Materials Research Society, Boston, MA,

November 27-30, 2000.

K. Ngala and M.S. Whittingham, ìSynthesis and

Characterization of Manganese Vanadium
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

43

Oxides,î Materials Research Society, Boston,

MA, November 27-30, 2000.

M.S. Whittingham, ìThe Control of Structure and

Properties of Vanadium and Manganese Oxides

through Tailored Soft Synthesis,î Intl. Symposium

on Soft Solution-Processing, Tokyo, Japan,

December 8-11, 2000.

M.S. Whittingham, ìManganese and Vanadium

Oxides as Battery Cathodes,î NTT Research

Labs., Ibaraki, Japan, December 12, 2000.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

44

Synthesis and Characterization of Cathode Materials

Marca M. Doeff

Lawrence Berkeley National Laboratory, 62-203, Berkeley CA 94720

(510) 486-5821, fax: (510) 486-4881; e-mail: mmdoeff@lbl.gov

Objective

• Develop low-cost cathodes based on benign materials (e.g., manganese oxides, lithium iron phosphate) having

electrochemical characteristics (cycle life, energy, and power densities) consistent with the goals of the USABC

and/or PNGV.

Approach

• Use novel synthesis methods to produce pure materials with optimum particle size and morphology.

• Carry out systematic substitutions to improve electrochemical properties.

• Test materials in several different cell configurations under a variety of conditions, including normal cycling and

high-rate discharge at both room and elevated temperatures.

Accomplishments

• Demonstrated that Li/1 M LiPF6 EC-DMC/tunnel-LixMnO2 cells can be cycled over 250 times with no capacity

fading, and can be discharged at up to the 5C rate repeatedly without damage.

• demonstrated excellent cycling behavior at 85°C for polymer cells containing tunnel-LixMnO2 (less than 0.02%

average capacity loss per cycle over 140 discharges at the 1.5 C rate).

• Increased reversible capacity 20% by substitution of 22% of the Mn in tunnel-MnO2 with Ti.

Future Directions

• Investigate LiFePO4 as electrode material; overcome rate limitations caused by poor electronic conductivity by

using sub-micron particles and coating with carbon during synthesis.

• Expand investigation of the effects of substitution on the electrochemical behavior of tunnel-containing manganese

oxides.

We have identified a remarkably stable

manganese oxide phase with an unusual doubletunnel structure (Na0.44MnO2). The theoretical

capacity is estimated to be at least 200 mAh/g. It

undergoes ion-exchange and reductive intercalation

processes readily and reversibly and does not

convert to spinel below about 400°C. Overcharge/over-discharge studies have also indicated

that this structure is highly resistant to damage from

abuse.

In 2000, we showed that the rate behavior and

cycling characteristics of LixMnO2 made from

Na0.44MnO2 depended strongly on the particle

morphology. Best results were obtained when

Na0.44MnO2 made by a glycine-nitrate process

(GNP) was used as the precursor. This combustion

synthesis method results in a drastic reduction in the

particle dimension along which ion diffusion takes

place. Lithium cells containing GNP-LixMnO2

electrodes show excellent rate capability (Fig. 42)

and cycle without fading. This material also

performs well in a polymer cell configuration at

elevated temperatures (Fig. 43), in sharp contrast to

spinel materials, which lose capacity rapidly.

These characteristics indicate that LixMnO2 with

the Na0.44MnO2 structure should perform well in

HEV batteries, where high power, abuse-tolerance,

and long cycle life are necessary, but high energy
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

45

density is not required. For EV applications,

however, higher capacity is needed. At present, it is

not possible to extract all of the Li ions from

LixMnO2 without oxidizing the electrolyte and

compromising cycle life. This limits the practical

capacity to about half of theoretical.

Partial substitution of Mn with other metals

changes the electrochemical characteristics, and, in

some cases, flattens the voltage profile, allowing

more capacity to be delivered. For example,

LixTi0.22Mn0.78O2 delivers 20% more capacity than

LixMnO2 in the same cell configuration.Substitution

of >22-at% level does not result in further

improvement, however, because too much electroactive Mn is replaced with inert Ti.

Although higher capacity is obtained for Li cells

with LixTi0.22Mn0.78O2 electrodes, the rate capability

and cycling behavior are inferior to those with GNPLixMnO2, negating the advantage except at low

current densities (Fig. 44). A sol-gel synthesized

product had very small average particle size, but had

markedly inferior electrochemical properties,

because the distribution of Ti differed from that of

the solid-state synthesized material.

2

2.5

3

3.5

4

4.5

0 20 40 60 80 100 120 140

Cell Potential, V

Capacity, mAh/g

Figure 42. Discharge number 118 of a Li/1 M LiPF6 ECDMC/GNP-LixMnO2 cell at 25 μA/cm2

 (___), discharge

number 98 at 0.14 mA/cm2 (-----), discharge number 121

at 0.5 mA/cm2

 (ÖÖ), and discharge number 108 at 2.5

mA/cm2 ( ) or 5C rate.

0

20

40

60

80

100

120

0 20 40 60 80 100 120 140

capacity, mAh/g

Cycle number

Figure 43. Cycling data for a Li/P(EO)8LiTFSI/GNPLixMnO2 cell discharged at 0.5 mA/cm2

 or 1.5 C rate

between 3.6 and 2.5 V at 85°C.

We have also been successful in preparing

LixCu0.11Mn0.89O2 with the tunnel structure, but its

electrochemical characteristics appear to be inferior

to those of LixMnO2. Also, attempts to substitute

Mn with V, Fe, Co, Ni, Zr, or 22-at% Cu result in

layered compounds or mixtures of phases rather than

the desired tunnel compound.

Future work will continue on the promising

tunnel manganese oxide compounds. Our goal is to

increase reversible capacity without sacrificing the

power capability, and excellent stability

demonstrated by GNP-LixMnO2. To do this, a better

understanding of intercalation into these materials is

necessary, as well as knowledge about the

substitution behavior and chemistry. We will also

begin investigating LiFePO4, another promising

electrode material. At present, LiFePO4 shows

severe rate limitations due to its poor electronic

conductivity. To ameliorate this, we will optimize

particle size and coat particles with carbon.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

46

0

20

40

60

80

100

120

140

0 50 100 150 200 250 300

Capacity, mAh/g

Discharge number

Figure 44. Cycling data for Li/P(EO)8LiTFSI/

LixTi0.22Mn0.78O2 cells at 85°C, discharged at 0.1

mA/cm2

 between 3.6 and 2.5 V (x), 0.1 mA/cm2

between 3.6 and 2.4 V (!), and 0.5 mA/cm2

between 3.6 and 2.5 V (!).

PUBLICATIONS

L. Edman, A. Ferry and M.M. Doeff, ìSlow

Recrystallization in the Polymer Electrolyte

System P(EO)nLiTFSI,î J. Mater. Res., 15, 1950-4

(2000).

L. Edman, M.M. Doeff, A. Ferry, J. Kerr and L.C.

De Jonghe, ìTransport Properties of the Solid

Polymer Electrolyte System P(EO)nLiTFSI,î J.

Phys. Chem. B., 104, 3476 (2000).

M.M. Doeff, L. Edman, S.E. Sloop, J. Kerr and L.C.

De Jonghe, ìTransport Properties of Binary Salt

Polymer Electrolytes,î J. Power Sources, 89, 227

(2000).

M.M. Doeff, T.J. Richardson, K.T. Hwang and A.

Anapolsky, ìFactors Influencing the Discharge

Characteristics of Na0.44MnO2-based Positive

Electrode Materials for Rechargeable Lithium

Batteries,î in Proceedings of the International

Meeting of The Electrochemical Society,

Honolulu, HI, October 1999, G-A. Nazri, M.M.

Thackeray, and T. Ohzuku, eds., The

Electrochemical Society, Inc., 99-24, 48 (2000).

PRESENTATIONS

M.M. Doeff, ìNovel Tunnel-containing Manganese

Oxides with Excellent Reversibility,î DOE

Workshop on Interfaces, Phenomena, and

Nanostructures in Lithium Batteries, Argonne, IL,

December 11-13, 2000.

M.M. Doeff, T.J. Richardson, A. Anapolsky and M.

Gonzales, ìStable Manganese Oxides for Lithium

Battery Applications,î invited lecture presented by

M.M. Doeff at the Pacific Conference on Ceramic

Materials, San Francisco, CA September, 2000.

M.M. Doeff, ìSynthesis and Electrochemical

Characterization of Tunnel-containing Doped

Manganese Oxides,î Solid State Chemistry I

Gordon Conference, New London, NH, July 2000.

M.M. Doeff, T.J. Richardson, K.-T. Hwang and A.

Anapolsky, ìImproved Discharge Characteristics

of Tunnel-containing Manganese Oxide Electrodes

for Rechargeable Lithium Battery Applications,î

presented by M.M. Doeff at the IBA-2000

Manganese Oxide Symposium, Chicago, IL, May,

2000.

M.M. Doeff, ìLithium-Polymer Batteries on a

Laboratory Scale,î invited lecture, Nordic Energy

Research Program Workshop, Lund, Sweden,

May 2000.

M.M. Doeff, ìElectrochemical Characterization

of Li-Polymer Batteries,î invited lecture,

Nordic Energy Research Program Workshop,

Lund, Sweden, May 2000.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

47

TASK AREA 5: DIAGNOSTICS

Diagnostics: Electrode Surface Layers

Frank R. McLarnon

Lawrence Berkeley National Laboratory, 90-1142, Berkeley CA 94720

(510) 486-4636; fax: (510) 486-4260; e-mail: frmclarnon@lbl.gov

Objectives

• Apply advanced in situ and ex situ techniques to characterize the structure, composition, formation, and growth of

surface layers on BATT Program electrodes.

• Establish direct correlations between electrode surface changes, interfacial phenomena, and cell capacity decline.

Approach

• Use ellipsometry, Raman spectroscopy, and advanced microscopic techniques to characterize electrodes taken from

baseline BATT Program cells, as well as thin-film electrodes in model cells.

• Measure changes in electrode surface morphology, electrode surface chemistry, and SEI thickness and composition

that accompany cell cycling.

Accomplishments

• Determined how the storage of model thin-film LiMn2O4 electrodes in pure DMC solvent and in LiPF6-EC-DMC

electrolyte at elevated temperatures changes the electrode structure, morphology, SEI layer, and surface chemistry.

• Measured the thickness and composition of SEI layers formed during cycling of model thin-film carbon electrodes

in LiClO4-EC-DMC electrolyte.

Future Directions

• Identify changes in LiAl0.05Ni0.80Co0.15O2 and Li1.02Al0.25Mn1.75O3.92S0.03 cathode surface morphology and chemistry

which accompany cycling in LiPF6-EC-EMC, Li(CF3SO2)2N +cross-linked PEO, and LiBF4 + cross-linked gel

electrolytes.

• Characterize SEI formation and growth on Li and carbonaceous anodes in LiPF6-EC-EMC, Li(CF3SO2)2N +crosslinked PEO, and LiBF4 + cross-linked gel electrolytes.

We extended our study of model thin-film spinel

LiMn2O4 electrodes to determine the effects of

elevated temperature on interfacial phenomena. We

used XRD, atomic force microscopy (AFM), and

Raman spectroscopy to characterize changes to the

electrode structure, morphology, and surface

chemistry caused by one-day exposure to pure DMC

solvent or 1 M LiPF6/EC-DMC electrolyte at 25, 60,

and 70°C. A CV of a fresh LiMn2O4 electrode in

LiPF6-EC-DMC electrolyte displayed two sets of

anodic and cathodic peaks, as expected for a spinel

LiMn2O4 electrode, and 43 CVs at 25°C resulted in

~9% capacity decline.

The capacity of the electrode exposed to pure

DMC at 70°C vanished completely, even though

there were no accompanying changes in its XRD

pattern, suggesting that the electrode capacity

decline was caused by a surface process, rather than

by changes in the bulk electrode structure. Currentsensing AFM (CSAFM) images of a fresh LiMn2O4

thin-film electrode revealed densely packed 100-250

nm particles with average surface roughness ~24 nm

and non-uniform surface conductance, i.e., the
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

48

conductance was lower on the upper surfaces of

individual LiMn2O4 grains than on their sides or at

grain boundaries. The surface morphology of the

electrode exposed to DMC at 60o

C was similar to

that of a fresh electrode, however its surface

conductance decreased significantly. CSAFM

images of the electrode exposed to DMC at 70°C

showed no surface electronic conductance and

moderate changes of surface morphology. The

electrode granular structure was preserved, but its

average surface roughness dropped to ~9 nm. These

observations suggest the formation of a thin SEI

layer and/or surface chemical modification of

LiMn2O4.

Our thin-film LiMn2O4 electrodes lost 3 and 4%

of their capacities upon exposure to LiPF6-EC-DMC

electrolyte at 60 and 70°C, respectively, in sharp

contrast to the complete loss of capacity upon

exposure to pure DMC at 70°C. The XRD peaks of

these electrodes shifted toward higher angles and

their respective lattice parameters decreased to 8.13

and 8.10 Å, suggesting that Li was extracted to form

LixMn2O4 with x<1. Also, the electrolyte became

brown at 70°C, indicating Mn dissolution.

AFM images revealed the formation of pits 150-

250 nm deep, clearly indicating Mn dissolution.

The surface conductance dropped significantly

at 60°C and almost completely disappeared at

70°C. However, in this case SEI layer formation did

not cause significant capacity loss, as was the case

for electrodes exposed to pure DMC.

Both Raman and surface-enhanced Raman

(SER) spectroscopies were used to determine the

compositions of SEI layers formed at elevated

temperatures. The ordinary Raman spectrum of a

fresh LiMn2O4 electrode exhibited peaks at 624, 580,

and 480 cm-1, which are assigned to A1g, T2g, and T2g

phonons, respectively. Upon exposure to pure DMC

at 60 or 70°C, a new peak appeared at 656 cm-1,

consistent with spectra of hausmanite Mn3O4 and

bixbyite Mn2O3. Because the peak intensities at 624

and 580 cm-1 remained unchanged, the 556 cm-1

peak clearly reflects Mn2O3 character and may be a

signature of surface conversion of LiMn2O4 into

Mn(III) oxide in DMC at elevated temperatures.

The Raman spectrum of the electrode stored in

LiPF6-EC-DMC electrolyte at 70°C showed even

more dramatic changes: the peak at 624 cm-1

vanished, the peak at 580 cm-1 became prominent,

and a strong resonance enhancement was observed

at 514 nm excitation. These complex changes can

be explained in terms of at least partial delithiation

and conversion to λ-MnO2, consistent with our XRD

data. SER spectra showed new peaks at 240, 889,

918, and 985 cm-1; broad multi-peak maxima at

1300 and 1600 cm-1; and peaks between 2100 and

2160 cm-1. These new peaks (except for that at 918

cm-1) are invisible in the ordinary Raman and SER

spectra of fresh electrodes, and therefore must be

associated with the SEI layer. The breadth and poor

definition of these peaks suggest a complex SEI

containing multiple organic compounds, along with

derivatives of carbonates and carboxylic acids of

complicated molecular symmetry. The broad peak

at 240 cm-1 can be assigned to the lithium-oxygen

stretch in lithium alkoxides Li-O-R (R=alkyl) and/or

Li coordinated to carboxylic Li-(CO2)-R or

carbonate (Li-O)2-C=O groups. The broad maxima

at 1300 and 1600 cm-1 correspond to symmetric and

asymmetric vibrations, respectively, of carboxyl and

carbonate groups. However, the features observed

in reference Raman and SER spectra of (COONa)2,

HCOONa, and CH3COOLi cannot account for all of

the peaks we observed. The peaks at 918/889 and

985 cm-1 are probably C-C stretching and

deformation vibrations of the carboxyl groups (-

COO). None of the reference spectra showed

intense bands in the 2100-2160 cm-1 region. These

features often accompany SER spectra and may arise

from adsorbed molecules on Ag, i.e., so-called

ìimpurity lines.î These peaks are located above and

below -C≡O (2135 cm-1), which suggest the

presence of polynuclear carbonyl complexes.

We began efforts to characterize in detail the

surface phase segregation of LiNi0.8Co0.2O2 cathodes

taken from DOE-ATD Program Li-ion cells. This

work serves three purposes: (i) to enhance our

capabilities for the study of BATT Program

electrodes, e.g., identifying the role of sulfur

additives on electrode surface chemistry; (ii) to

provide better understanding of LiNi0.8Co0.2O2

interfacial processes; and (iii) to support ATD

Program diagnostics. A ìreference libraryî was

constructed by recording Raman spectra of

LiNi0.8Co0.2O2 powders and commercially available

materials, and collecting literature data. Raman

microscopy was used not only to determine the

micron-scale chemical and structural uniformity of

these materials, but also to compare their spectral

characteristics to those of LiNiO2 and LiCoO2,

which exhibit peak pairs at 468, 546 and 487, 597

cm-1, respectively. In contrast, the micro-Raman

spectrum of LiNi0.8Co0.2O2 powder exhibits a broad
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

49

peak at ~520 cm-1 and a sharp peak at ~555 cm-1,

which indicate the presence of a solid solution of

Co-substituted nickel oxides rather than a mixture of

separate crystalline phases.

Micro-Raman spectra of LiNi0.8Co0.2O2 cathodes

taken from Li-ion cells stored or cycled at elevated

temperatures (40°C) displayed significant spectral

changes. The broad 520 cm-1 peak split into separate

peaks at 479, 553, and 565 cm-1 and a weak shoulder

at ~600 cm-1, due to thermally induced surface phase

segregation. Neither substitution of Co for Ni nor Li

intercalation-deintercalation changed the LiNiyCo1-

yO2 R3m space group symmetry. However, the

substitution of Co for Ni(III) oxides increased the

crystal lattice parameter, as reflected by a 3-4 cm-1

shift of two characteristic 474/554 cm-1 NiO2 peaks

toward lower frequencies and an increased 474/554

peak ratio. Lithium intercalation into NiO2 had the

opposite effect: the peaks shifted toward higher

frequencies and the 474/554 peak ratio decreased.

The presence of separate peaks at both sides of the

original NiO2 peaks and strong relative peak

intensity variations in the spectra of LiNi0.8Co0.2O2

cathodes suggest the formation of stable crystalline

phases of Ni(III) or Ni(IV) oxides with different

amounts of Co and Li on the cathode surface.

In collaboration with the research group of J.W.

Evans (p. 53), we carried out in situ spectroscopic

ellipsometry (SE) and electrochemical quartz crystal

microbalance (EQCM) experiments to characterize

the formation and growth of SEI layers on carbon

electrodes. We recorded ellipsometric spectra

before and after cycling of an electron-beamdeposited thin-film carbon electrode between 0 and

3 V vs. Li/Li+

 at 2 mV/s in a half cell with 1 M

LiClO4/EC-DMC electrolyte. We considered two

optical models to describe the electrode: (i) a

compact and homogeneous carbon layer, and (ii) an

effective-medium approximation of a rough, porous

carbon layer. The latter model showed better fitting

results and is consistent with the electrode surface

roughness measured by AFM. We found that the

SEI layer thickness was about 60 nm after the 1st

cycle and increased to about 90 nm after the 20th

cycle. The SEI refractive index varied from 1.41 to

1.52, and its extinction coefficient was close to zero,

i.e., the SEI layer was more optically compact than

common polymers and exhibited almost no

electronic conductivity. The EQCM frequency shift

during long-term cycling correlated well with our SE

results. Detailed analysis of our SE and EQCM data

indicated that the SEI density was ~1.5 g/cm3

, which

is larger than the densities (1.0-1.2 g/cm3

) of

expected polymer-like products of alkyl carbonate

reduction. We conclude that the SEI must contain

denser inorganic Li compounds (e.g., Li2O, Li2CO3,

LiClO4) in addition to alkyl carbonates.

PUBLICATION

F. Kong and F. McLarnon, ìSpectroscopic

Ellipsometry of Lithium/Polymer-Electrolyte

Interfaces,î J. Power Sources., 89, 180-89 (2000).

PRESENTATIONS

F. Kong, R. Kostecki, G. Nadeau, X. Song, K.

Zaghib, K. Kinoshita and F. McLarnon, ìIn Situ

Studies of SEI Formation,î 10th International

Meeting on Lithium Batteries, Como, Italy, MayJune 2000.

R. Kostecki, Y. Matsuo and F. McLarnon,

ìInterfacial Phenomena on selected Cathode

Materials,î Workshop on Interfaces, Phenomena,

and Nanostructures in Lithium Batteries, Argonne,

IL, December 2000.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

50

Battery Materials: Structure and Characterization

James McBreen

Brookhaven National Laboratory, ES&T-Bldg. 555, P.O. Box 5000, Upton, NY 11973-5000

(631)344-4513, fax. (631)344-5815; email jmcbreen@bnl.gov

Objective

• Establish direct correlations between electrode material changes, interfacial phenomena, and cell capacity decline.

Approach

• Apply a combination of in situ and ex situ synchrotron techniques to characterize electrode materials and electrodes

taken from baseline BATT Program cells.

• Develop and apply X-ray absorption spectroscopy (XAS) techniques that are sensitive to both surface and bulk

processes in electrodes.

Accomplishments

• Completed an extensive study of the effect of Li and O stoichiometry on the phase behavior of spinel LiMn2O4

cathodes during cycling at ambient and low temperatures.

• Completed a study of the effects of electrode history on the phase behavior of LiMn2O4 cathodes.

• Completed in situ XAS studies of Li-Cr-Mn-O cathodes.

• Demonstrated the feasibility of collecting high-quality in situ XAS data on very dilute dopants in cathode materials.

Future Directions

• Develop a combination of electron yield and fluorescent XAS at transition metal K- and L-edges as a tool for

identification of inorganic surface degradation products in cathodes. These will be used in conjunction with K-edge

XAS of low Z elements such as C, O and F.

• Identify the changes in LiNixMn2-xO2 and Li1.02Al0.25Mn1.75O3.92S0.03 when cycled in LiPF6-EC-DMC electrolyte.

Our objectives are to develop and apply

advanced diagnostic techniques to elucidate the

molecular aspects of materials and electrode

processes in Li-ion batteries, and to use this

information to develop electrode and electrolyte

structures with low cost, good performance, and

long life. Work in 2000 was focused on in situ highresolution XRD and XAS studies of lithium

manganese oxide materials. Also methods were

devised to collect high-quality in situ XAS data on

very dilute dopants in cathode materials.

High-resolution XRD studies of LiMn2O4.

LiMn2O4 compounds with various oxygen

stoichiometries were prepared by heat treating

Li1.07Mn1.93O4 at 600°C under Ar. The O

stoichiometry was determined from separate

thermographic analysis (TGA) measurements. The

DSC results in Fig. 45 show that the temperature for

the phase transition and kinetics of the conversion

from the cubic to the tetragonal phase increased with

decreasing oxygen stoichiometry. In the case of

stoichiometric materials the conversion occurs in the

vicinity of 0°C. For the longer heat treatments the

conversion occurred at ambient temperature or

above.

Figure 46 shows that the O stoichiometry also

had major effects on the phase transformations at

low temperature. The results indicate that the

existence of an oxygen deficiency in LiMn2O4 is a

critical condition for formation of the lowtemperature phase. Also, the greater the oxygen

deficiency the higher the transition temperature. The

XRD results at low temperature are more consistent

with a mixture of a tetragonal and a cubic phase,
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

51

rather than a single phase with an orthorhombic

structure. XRD results, at 300 K, showed a mixed

phase composition for a sample treated at 600°C

under Ar for 12 h. After treatment of the sample

under O2 at 700°C for 10 h only a single cubic phase

was seen at 300 K. This strongly supports a

mechanism based on oxygen deficiency.

T e m p e ra tu re o C

-6 0 -5 0 -4 0 -3 0 -2 0 -1 0 0 1 0 2 0 3 0 4 0 5 0

Heat Flow (W/g)

-0 .1 5

-0 .1 0

-0 .0 5

0 .0 0

0 .0 5

0 .1 0

0 .1 5

0 h tre a tm e n t

1 h tre a tm e n t

2 h tre a tm e n t

3 h tre a tm e n t

Figure 45. DSC trances between -60 and 50°C for the

untreated (0h), 1 hour (1h), 2 hour (2h), and 3 hour (3h) heat

treatment of Li1.07Mn1.93O4 at 600°C under Ar.

2θ

43 44 45 46 47 48 49 50 51 52 53

Intensity

2h 

3h 

1h 

0 

h 

Figure 46. XRD patterns (λ=1.195Å) at 240 K for the

untreated (0h), 1 hour (1h), 2 hour (2h), and 3 hour (3h)

heat treatment of Li1.07Mn1.93O4 at 600°C under Ar.

LiMn2O4 with a low oxygen stoichiometry

clearly displayed three cubic phases during charge.

Stoichiometric materials and materials with excess

Li stoichiometry display pseudo single-phase

behavior in the early stages of charge. This is

followed by two-phase behavior toward the end of

charge. Preliminary cycling studies indicate that the

oxygen stoichiometry also affects capacity

maintenance and phase behavior during cycling.

Effect of Al substitution in LiMn2O4. An

Al-doped LixMn2O4 spinel material with the

composition Li1.02Al0.09Mn0.91O4 was prepared from

a γ-MnOOH precursor. This material had excellent

stability when cycled over both the 4.1 and 3 V

plateaus. After a few initial cycles the capacity

stabilized at 200 mAh/g. In situ XRD scans were

taken several times over the first fifty cycles. The

tetragonal phase was present on the 20th cycle after

discharge to 2.2 V. This indicates that Al

substitution does not eliminate formation of the

tetragonal phase on the 3 V plateau, but rather

facilitates the process and mitigates the bad effects

of tetragonal phase formation.

Li-Cr-Mn-O cathodes. In situ XAS and

XRD studies were carried out on a new cathode

material, Li[LixCryMn1-x-y]O2, developed at NRC,

Canada. This is a layered material and replacement

of ~20% of the transition metal atoms with Li

improves the electrochemical characteristics of the

material. A careful in situ XAS study at the Cr and

Mn K-edges indicated that Mn was Mn(IV) and Cr

was Cr(III) in the as-prepared material. During

charge no further oxidation of the Mn occurs. Both

the near edge data (XANES) and the extended X-ray

absorption fine structure (EXAFS) at the Cr K-edge,

during the first charge, indicate that about 40% of

the Cr(III) is oxidized to Cr(VI) The data in Fig. 47

indicates that enormous changes occur in the Cr

environment during charge. Analysis of the EXAFS

shows that the oxidized Cr migrates from octahedral

to tetrahederal sites. The reverse occurs during

discharge, with a small fraction of the Cr(VI)

remaining on tetrahedral sites. In subsequent cycles

there is no further increase in the residual of Cr(VI)

at the end of discharge. The results also indicate

that Mn and Cr are not randomly distributed in the

lattice, but exist in separate domains; and that the

excess Li is associated with Mn vacancies around

the Mn(IV).
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

52

R [Å]

012345678

k3 transform magnitude

0

2

4

6

8

10

12

14

16

Scan 1

Scan 5

Scan 15

Scan 21

Scan 30

Figure 47. Fourier transforms of Cr K-edge EXAFS

taken during the first charge from 3.5 to 4.5 V. 30 scans

were taken during charge at the C/13 rate.

Ga XAS studies of Li0.908Co0.088Ga0.003O2:

In situ Ga K-edge XAS was done on

Li0.908Co0.088Ga0.003O2 in an effort to develop

techniques to do XAS on dilute dopants. By doing

the measurements in the fluorescence mode it was

possible to get excellent data. As with the Cr in the

previous section, enormous changes were seen in

both the Ga K-edge XANES and EXAFS. The

results indicate that during charge Ga migrates from

octahederal to tetrahederal sites. Its efficacy as a

dopant may be due to this migration preventing Ni

from dropping into Ni sites. The Cr and Ga results

demonstrate the ability of local probes such as XAS

to deliver inportant structural information on

complex systems.

PUBLICATIONS

A.N. Mansour, S. Mukerjee, X.Q. Yang and J.

McBreen, ìIn Situ X-Ray Absorption and

Diffraction Study of The Li Reaction With a Tin

Composite Oxide Glass,î J. Electrochem. Soc.,

147, 869 (2000).

Y.J. Lee, F. Wang, S. Mukerjee J. McBreen and C.P.

Grey, ì6

Li and 7Li Magic-Angle Spinning Nuclear

Magnetic Resonance and In Situ X-Ray

Diffraction Studies of the Charging and

Disacharging of LixMn2O4 at 4 V,î J.

Electrochem. Soc., 147, 803 (2000).

X.Q. Yang, X. Sun and J. McBreen, ìStructural

Changes and Thermal Stability: In Situ X-Ray

Diffraction Studies of a New Cathode Material

LiMg0.125Ti0.125Ni0.75O2,î Electrochem. Commun.,

2, 733 (2000).

Diagnostics: Interfacial and Reactivity Studies

Philip N. Ross, Jr.

Lawrence Berkeley National Laboratory, 2-100, Berkeley CA 94720

(510) 486-6226, fax: (510) 486-5530; e-mail: pnross@lbl.gov

Objective

• Establish direct correlations between electrode surface changes, interfacial phenomena, and cell failure.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

53

Approach

• Use Fourier transform infrared (FTIR) spectroscopy and X-ray photoelectron spectroscopy (XPS) to study

model electrode/electrolyte combinations, e.g., using glassy carbon electrodes and BATT Program electrolytes,

to provide the basis to interpret more-complex spectra recorded for ATD Program cell materials.

Accomplishments

• Completed the experiments designed to determine the near-surface structure in lithiated highly ordered pyrolytic

graphite (HOPG) and/or LiC6(0001).

• Completed experimental measurements of the electrochemical reduction potentials of all carbonate solvents of

interest in Li-ion batteries.

• Completed computational studies of the electrochemical reduction and oxidation potentials of all organic solvents of

interest in Li batteries.

Future Directions

• Establish thermal stability of the SEI layer on graphite anodes in ATD Program Gen 2 electrolyte.

• Identify some routes to improved stability via electrolyte additives and/or graphite pre-treatment.

We have completed experiments designed to

determine the near-surface structure in lithiated

HOPG and/or LiC6(0001), i.e., which is the

equilibrium surface plane, the graphene plane, or the

Li-ion plane. We completed this determination by

analysis of the take-off angle dependence of the Li

1s and C 1s photoelectron intensities and detailed

modeling of the photoelectron emission/diffraction

from the different surface terminations. The data

clearly indicate termination in the graphene plane, in

agreement with intuition.

The study of the reaction of diethyl carbonate

(DEC) with lithiated C(0001) in ultra high vacuum

(UHV) continues. The data is being compared in

detail with that for DEC on metallic Li. Because we

now know that the surface of lithiated C(0001) is

terminated in the graphene plane, an interesting

picture of the reaction with DEC emerges. The first

step is electron transfer from the graphene sheet to

the DEC molecule, which dissociates to form two

anions, ethyl carbonate and ethyl carbanion. These

anions are bound electrostatically to Li ions beneath

the intervening graphene sheet. Since the binding is

relatively weak, and no direct bonding to Li+

 occurs,

for the most part the anions thermally decompose

and/or recombine and desorb from the surface upon

heating to 470 K. However, some Li+

 ions are also

drawn out of the graphite lattice at 470 K to form an

alkoxide-like species. This process is postulated to

occur at step-edges in the HOPG, whose

concentration can be varied by Ar-ion etching pretreatment of the crystal. To date, CO2 is the only

gaseous reaction product clearly identified; the fate

of the ethyl carbanion fragment from the initial

dissociation is unclear.

FTIR analysis of the reduction product of EC on

glassy carbon was conducted ex situ. The reduction

potential was varied from 0.05 to 1.5 V (vs. Li/Li+

)

at 0.25 V intervals. Unfortunately the results from

these first experiments failed to show a systematic

effect of reduction potential on the FTIR spectra.

The only functional group observed, other than those

of the bulk electrolyte, was the epoxy group,

characterized by a strong, sharp C=O stretching

band at 838 cm-1, as shown in Fig. 48. This band

has never been reported in other FTIR studies of Li

or carbon anodes in any carbonate-based

electrolytes. We also observed this band on some

carbon anodes removed from Li-ion cells used in the

ATD Program. In principle, ethylene oxide is the

simple decomposition product of EC, according to

EC " H2COCH2 + CO2

But one rarely sees this reaction listed as a

possibility in the literature.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

54

700 1200 1700 2200 2700 3200

cm-1

Absorbance (arb. units) 

CO2

838

1100 C-H

a

b

c

Figure 48. Diffuse reflectance mode FTIR spectra from:

a) bulk electrolyte THF/LiClO4 + 10 mM EC; b) glassy

carbon electrode after reduction for 5 min. at 0.05 V

(Li/Li+

); c) propylene oxide.

CV experiments were carried out to determine

experimental values to compare against the

calculated solvent reduction potential that we

previously reported. The target solvent was

dissolved in a tetrahydrofuran (THF)/LiClO4 (0.1 M)

electrolyte in a concentration of 1 vol%. A gold

milli-electrode (1 mm radius) was employed with a

slow scan rate of 2 mV/s. THF was selected because

of its stability against reduction above Li metal

potential (Fig. 49). The reduction at 1.8 V is

generally ascribed to the reduction of H2O, an

impurity in the electrolyte. A summary of both

calculated reduction potential and experimental

results for four solvents is shown in the table.

Figure 49. CV for THF-LiClO4 electrolyte with 1 vol%

solvent dissolved within.

Solvent Calculated Experimental

EC 1.46 1.36

DEC 1.33 1.24

DMC 0.86 1.32

VC 0.25 1.40

VC = vinylene carbonate

Both EC and DEC show good agreements

between calculated and experimental results with the

experimental values slightly lower by about 0.1 V.

When calculating the reduction of linear carbonate

solvents (DEC and DMC), it was assumed to be a

two-electron-transfer process. An alkyl carbonate

anion and alkyl radical were formed upon the first

electron transfer while the 2nd electron was

transferred to the alkyl radical to form an alkyl

anion. This assumption seems to work well for DEC

but not DMC. It is not energetically favorable for a

methyl radical to accept an electron to form a methyl

anion, which results in a lower calculated reduction

potential of DMC. The experimental result,

however, implies the reduction of DMC follows a

different reaction pathway.

Vinylene carbonate (VC) was found to reduce at

1.40 V, much higher than the calculated value of

0.25 V, which suggests that the reduction process

does not stop at a the non-bond-broken form of

radical anion as we expected. Rather, a more

energetically favorable product should be formed.

Investigations were carried out to study the

reduction products of VC. The relatively positive

reduction potential for VC supports the postulated

role of VC as an additive to propylene carbonate

(PC)-based electrolyte for Li-ion cells. In the role,

preferential reduction of VC forms an interfacial

layer that prevents PC co-intercalation and graphite

exfoliation.

A Li-ion battery composite carbon electrode

(ATD Program Gen 1) electrode was used for the

cycling experiments with VC as an additive to PC

electrolyte. The composition of the anode was 75%

MCMB carbon, 17% SFG-6 graphite, and 8%

polyvinylidene difluoride (Kureha C) with a loading

of 5.5 mg/cm2

 applied to both sides of a Cu foil

current collector. Figure 50a shows constant-current

(0.5 mA/cm2

) cycling of the composite carbon anode

in PC electrolyte with 1 M LiClO4. The long plateau

around 0.4 V indicates significant solvent reduction.

In consequence, the reversible capacity for the first
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

55

cycle was only 12%. Severe graphite exfoliation

was observed in this electrolyte after just 10 cycles,

and concurrent loss of capacity. Figure 50b shows

the first constant-current charge and discharge of a

nominally identical composite carbon anode in the

same PC with addition of 5 vol% VC. The

reversible capacity on the first cycle is significantly

improved to 67%, and subsequent cycles had a

cycling efficiency of 93ñ95%. The electrolyte

turned slightly brown after 10 cycles. However,

there was no observable exfoliation of the graphite.

The effects of VC on PC reduction reported here are

qualitatively consistent with those reported

previously by Jehoulet et al.

0

0.2

0.4

0.6

0.8

1

1.2

1.4

1.6

1.8

2

0 20000 40000 60000 80000 100000

time (s)

V

a

b 

Figure 50. Constant-current (0.5 mA/cm2) cycling of a

composite carbon anode in (a) PC/LiClO4 electrolyte, (b)

PC/LiClO4 electrolyte with 5% VC added.

The role of VC as additive was further

investigated by in situ AFM. A solution with

1 vol% of VC in THF with 0.1 M LiClO4 was used

as the electrolyte. A polished glassy carbon (1.2 x

1.2 cm) was employed as the working electrode with

Li metal as both the reference and counter

electrodes. The potential was stepped down from

the 2 to 0.5 V with 100 mV interval. Surface

profiles were scanned at each step. The data indicate

that a surface film was starting to develop on the

electrode surface at 1.2 V. The AFM images at open

circuit potential and 1.2 V are shown respectively in

Figure 51 a and b. This finding is consistent with

the reduction of VC in THF reported previously by

our group (Fig. 49). The thickness of the surface

film was measured to be ~10 nm. This observation

confirms our previous conclusion that VC is forming

a SEI layer at high potential before Li+

 intercalation

into graphite and therefore prevents solvent cointercalation.

We simulated the Li/PEO interface by

deposition of Li onto a Pt electrode from the

electrolyte THF/LiClO4 + 35 mM PEGDME, with

the PEGDME MW 400. After the electrochemical

experiment, the electrode was removed from the

cell, rinsed with THF and inserted into an airtight

infrared cell fitted with a KBr window. Ex situ

Infrared Reflection Absorption Spectroscopic

(IRRAS) analysis of the interface was conducted and

shown in Fig. 52b along with spectrum of the

THF/PEGDME electrolyte on a Pt substrate (Fig.

52a). The two spectra are similar and no new

feature is detected on the electrode interface, i.e., no

SEI is detected. IR is therefore, not useful in

detecting SEI layer on the Li/PEO interface.

(a) (b)

Figure 51. AFM image of glassy carbon electrode surface at (a) open circuit potential 2.0 V, (b) 1.2 V.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

56

0.15

0.25

0.35

0.45

0.55

0.65

700 1200 1700 2200 2700 3200

cm-1

Absorbance (Arb. units)

a

b

Figure 52. Ex-situ IR of Pt electrode surface in electrolyte

THF/LiClO4 + 35 mM PEGDME, a) before electrochemical

reduction; b) after Li metal deposition.

Synthesis and Characterization of Electrodes

Elton J. Cairns

Lawrence Berkeley National Laboratory, 70-108B, Berkeley CA 94720

(510) 486-5028; fax (510) 486-7303; e-mail ejcairns@lbl.gov

Objectives

• Characterize the local atomic and electronic states of Li in BATT Program cathode materials, and determine how

these states change during cell cycling.

• Develop intermetallic electrodes that have capacities higher than carbon electrodes and acceptable stability during

cell cycling.

Approach

• Use 7Li MAS-NMR to characterize BATT Program electrodes before and after cycling, and after being subjected to

model failure processes.

• Determine the NMR isotropic chemical shift, linewidth, lineshape, and relaxation times for each Li species. Extend

this to other important nuclei.

• Synthesize intermetallic electrodes for use as negative electrodes in BATT Program cells, and characterize the

unique microstructures that result from synthesis methods such as high-energy milling.

Accomplishments

• Obtained preliminary 7Li MAS-NMR spectra for several tunnel-structure lithium manganese oxides.

• Synthesized Mg-Si and Mg-Sn materials with various microstructures and demonstrated that their Li storage

capacities exceed that of a typical carbon.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

57

• Determined that conversion of Mg2Sn and Mg2Si into their constituent elements contributes to rapid capacity fade

when cycled over a wide voltage window.

Future Directions

• Continue NMR work using a 100 MHz magnet to improve resolution.

• Compare results for tunnel-structure cathodes at different field strengths to assign isotropic peaks.

• Determine NMR linewidths and relaxation times for tunnel-structure cathodes.

• Obtain spectra for fresh and cycled BATT Program cathodes.

• Determine effects of intermetallic electrode microstructure on electrochemical performance by correlating cycling

results with material properties.

• Use 119Sn NMR to study the local environment of the Mg-Sn intermetallic electrodes.

• Use 7Li NMR to determine sites for Li occupancy in intermetallic electrodes.

Cathodes. Several BATT Program tunnelstructure lithium manganese oxides were obtained

from M. Doeff (p. 38 of this report), including Cuand Ti-substituted manganese oxides.

Representative spectra are shown in Fig. 53. These

spectra were obtained using a 181 MHz magnet with

MAS at 11 kHz. Although this is a rapid spinning

rate, many MAS sidebands are observed to each side

of the isotropic peaks. The resulting MAS envelopes

are quite wide due to coupling between the 7

Li

nucleus and unpaired electrons in the Mn d-orbitals.

Even so, several chemically distinct types of Li are

indicated. The Ti-substituted sample gives rise to

significant spectral overlap of the isotropic peaks

arising from a wide variety of Li sites. It is expected

that future work conducted using a 100 MHz magnet

will show greatly improved resolution of the peaks

arising from chemically distinct Li sites due to the

reduction of the bulk magnetization of the Mn delectrons.

Intermetallic electrodes. There is a strong

incentive to develop and characterize noncarbonaceous materials for use as negative

electrodes that have higher capacities than carbon

and that retain high capacities during cycling. We

have shown that the Mg-Si and Mg-Sn alloys can

store Li with capacities well in excess of typical

carbonaceous electrodes, but more work is needed to

characterize these materials and prevent rapid

capacity decline during cycling.

Electrode materials were synthesized by

mechanically alloying elemental powders in a

planetary ball mill. Mechanical alloying is a useful

method for producing battery materials because the

products have a fine particle size, the product

microstructure can be controlled by the milling

conditions and heat treatments, and interesting

metastable phases can be produced in large

quantities. Mg2Si, which is difficult to prepare by

casting because of the large disparity in melting

points between Mg and Si, was synthesized directly

by mechanical alloying. In the Mg-Sn system, the

stable, cubic Mg2Sn phase could be converted to a

metastable, rhombohedral phase with further

milling, as shown in Fig. 54.

-0.01

0

0.01

0.02

0.03

3000 2000 1000 0 -1000 -2000

Intensity (per mg per scan)

Shift vs. 1M LiCl(aq) (ppm)

Li0.47MnO2

-0.002

0

0.002

0.004

0.006

0.008

0.01

3000 2000 1000 0 -1000 -2000

Intensity (per mg per scan)

Shift vs. 1M LiCl(aq) (ppm)

Li0.47Ti0.33Mn0.67O2

Figure 53. 7Li MAS NMR spectra of tunnel-structure

lithium manganese oxides. Spinning speed was 11 kHz.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

58

Cycling these materials over a wide voltage

window from 0 to 1 V vs. Li/Li+

 produced maximum

discharge capacities greater than 600 mAh/g and 800

mAh/g for cubic Mg2Sn and Mg2Si, respectively. A

typical charge and discharge curve for Mg2Si is

shown in Fig. 55, however cycling over this wide

window led to a rapid capacity fade to capacities

>100 mAh/g within 10 cycles. XRD measurements

showed that this capacity fade was at least partly due

to conversion of the alloys into their elemental

starting materials, which presented the opportunity

for reaction of Li with Si and Sn. The separation of

intermetallic alloys into their constituents during Li

cycling is a commonly-observed phenomenon that

we wish to suppress.

0

25

50

75

100

0 5 10 15 20 25

Milling Time (hr)

Mass Percent

Sn

Mg

Cubic Mg2Sn Rhombohedral Mg2Sn

Figure 54. Batch composition vs. milling time at 2.1 W.

0.00

0.20

0.40

0.60

0.80

1.00

0 100 200 300 400 500 600 700 800 900

Capacity (mA h /g)

E vs Li/Li+ (mV)

Figure 55. Charge/discharge cycle for Mg2Si.

Lithium/sulfur cells. Our work on Li/S cells

was completed in 2000, and included an

investigation of the behavior of thin sulfur cathodes

containing 70 wt% elemental sulfur, 20 wt% carbon

black, and 10 wt% PEO (M.W. 5,000,000). These

materials were mixed in acetonitrile for several days,

and then cast onto a stainless-steel disk current

collector. As-prepared electrodes were dried at

60°C for 24 h in vacuum before introduction into a

glove box where cell assembly was carried out. The

electrolyte was LiTFSI in PEGDME (MW 500), and

the separator was Celgard 2400. Cell cycling was

carried out at room temperature. Figure 56 shows

that the cell discharge capacity increased for the first

20 cycles, a phenomenon that is probably related to

slow permeation of the high-viscosity electrolyte

into the electrode. The capacity slowly decreased

over almost 400 cycles, with a capacity loss rate of

0.46 mAh/g per cycle. At 0.02 mA/cm2

, the specific

capacity of the sulfur electrode was about 470

mAh/g, which corresponds to almost 30% utilization

of active material. The capacity decreased to only

20 mAh/g of electrode at high rate (0.5 mA/cm2

),

but it fully recovered when the current density was

returned to 0.05 or 0.02 mA/cm2

.

Figure 56. Capacity of sulfur electrode during cycling.

Cut-off 1.5-3.0 V vs. Li/Li+

. Electrode area 1 cm2.

Electrode weight 1.6 mg/cm2

. Charge/discharge current

0.02-0.5 mA/cm2

. Electrolyte LiTFSI in PEGDME 500.

These results show good promise for the

continuing development of the Li/S cell as an

inexpensive, environmentally-benign cell with high

specific energy. Modifications to the structure of the

sulfur electrode should yield even higher sulfur

utilization.

PUBLICATION

D. Marmorstein, T.H. Yu, K.A. Striebel, F.R.

McLarnon, J. Hou and E.J. Cairns,

ìElectrochemical Performance of Lithium/Sulfur

Cells with Three Different Polymer Electrolytes,î

J. Power Sources, 89, 219 (2000).
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

59

Diagnostics: Electrochemical Quartz Crystal Microbalance (EQCM) Studies of the

SEI on Carbon Anodes

James W. Evans (Lawrence Berkeley National Laboratory)

University of California, 585 Evans Hall, MC 1760, Berkeley CA 94720

(510) 642-3807, fax: (510) 642-9164; e-mail: evans@socrates.berkeley.edu

Objective

• Apply EQCM to characterize the SEI at the carbon/electrolyte interface in Li-ion batteries.

Approach

• Use in situ EQCM with combined CV to understand SEI formation and growth in Li-ion half cells.

• Characterize the system with other supplementary techniques such as ellipsometry (joint work with McLarnon

group).

• Develop a model for SEI formation and growth on carbon anodes.

Accomplishments

• Estimated the SEI density on disordered carbon thin-film electrodes in 1 M LiClO4/EC+DMC electrolyte.

• Observed mass changes corresponding to Li de-intercalation in CV with EQCM.

• Developed a model for interpreting measured EQCM data on a theoretical basis.

Future Directions

• Correlate SEI formation and growth with LiPF6 or LiTFSI/EC+DMC electrolyte and impurities (e.g., H2O).

• Complete modeling of EQCM data according to each electrolyte and develop a general

model for SEI formation and growth.

This program focuses on the characterization of

the SEI on carbon electrodes in secondary Li-ion

batteries. Along with Li de-intercalation in

secondary Li batteries, other electrochemical/

chemical reactions are known to occur during the

first few cycles and form an electrode surface layer,

often referred to as the SEI. These reactions

involving in SEI formation are associated with socalled irreversible capacity loss.

Because EQCM enables us to simultaneously

measure, with high sensitivity, the mass change of

the carbon electrode and the charge transfer

associated with electrochemical reactions occurring

on the carbon electrode, this technique is our major

experimental tool for this study. Further, it allows

us to determine the mass change per mole of

electron transfer (mpe) for the species formed on the

carbon electrode.

To fabricate a thin-film carbon electrode, a 100-

nm thick carbon film was deposited by an electronbeam deposition technique onto one side of a

smooth quartz substrate with two transition layers

(500-nm Cu and 100-nm Cr), and Raman

spectroscopy was carried out to characterize the

thin-film carbon electrode. There is no visible

dominant peak which is characteristic of graphite,

diamond, glassy carbon, or any other form of

ordered carbon but rather the superposition of all

possible vibrations of various carbon structures.

This is typical for a highly disordered/amorphous

phase.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

60

A galvanostatic cycling test was conducted with

a constant current density of 10 μA/cm2

. There is a

large irreversible capacity in the first cycle and the

reversible capacity drops to about half of the

theoretical value of graphite at later cycles. As

expected from the fact that our carbon sample is

disordered/amorphous, there is no distinct plateau

that would imply staging phenomena in the graphitic

structure.

CV and EQCM were applied simultaneously to a

carbon electrode in 1 M LiClO4/EC+DMC

electrolyte. The measured quantities (current

density, i, and resonant frequency change, ∆f) for

both techniques are plotted in Fig. 57. Because the

mass increase corresponds to a decrease in the

resonant frequency, the negative of ∆f is plotted vs.

potential to facilitate direct interpretation. Both the

current density and frequency changes for the first

cycle are larger than those of subsequent cycles.

From the second to the 20th cycle, the shapes of the

cyclic voltammograms and electrogravimetric curves

are preserved. An irreversible peak appears at about

0.8 V in the first cathodic scan, and Li intercalation

prevails as the potential approaches 0 V. For the

anodic scan, Li de-intercalation is significant in the

potential range above 0.5 V. This Li de-intercalation

behavior was also reported in papers dealing with

disordered carbon. The frequency change displays a

trend similar to the CV. That is, the frequency

shows a monotonic decrease, which implies

accumulating mass on the working electrode, during

the first cathodic scan. Furthermore, the frequency

increases with increasing anodic current and

decreases with increasing cathodic current. 

Although this frequency response coincides

qualitatively with the current behavior, there exists a

quantitative discrepancy between them. While the

charge stabilizes in later cycles, becoming almost

reversible, the frequency continues to decrease with

a decreasing amplitude of oscillation if we plot

frequency and charge as a function of time during

CV. In order to identify the source of the continuous

frequency decrease, the transient frequency change

at OCV was measured. The frequency decreases

with the square root of time, and this dependency is

the same regardless of the kind of electrodes

(carbon, Al, and Ni tested) and solvents (EC+DMC

and PC tested).

0.0 0.5 1.0 1.5 2.0 2.5

-100

-80

-60

-40

-20

0

20

OCV

number of cycle increases 

V vs. Li/Li+

i, 

μA/cm2

0.0 0.5 1.0 1.5 2.0 2.5

0

200

400

600

800

1000

1200

1400

number of cycle increases 

OCV

- ∆f, Hz

V vs. Li/Li+

(a) 

(b)

Figure 57. (a) CV and (b) resonant frequency change of

the carbon electrode during the first 20 cycles in 1 M

LiClO4/EC+DMC (scan rate = 1mV/s).

With the correction using this frequency change

at OCV as a base-line, mpe values were obtained for

the first cathodic scan. The mpe values can be

classified into three parts. The first region

corresponds to the potential range from OCV to 1.1

V where mpe values are the highest (more than 30).

The second region coincides with the irreversible

peak at about 0.8 V, and the corresponding mpe is

about 21. The last region could be termed an

intercalation region where mpe values gradually

approach 7, the theoretical value for Li intercalation.

These results imply that organic species whose

molecular weights are higher than those of inorganic

species and Li ions form first and then inorganic

species form. However, we cannot exclude the

possibility of each formation potential being

overlapped over the tested potential range.

Because most irreversible mass gain occurs in

the first cathodic scan, ellipsometry was conducted

to estimate the density of the SEI formed after the

first cathodic scan under conditions identical to the

EQCM experiment. The SEI density is estimated to
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

61

be about 1.3 g/cm3

 using the SEI thickness of about

60 nm as determined by ellipsometry. However, if

the frequency change at OCV is also due to

electrochemical/chemical reactions participating in

film formation, the density would then be 1.6 g/cm3

.

The estimated density of 1.3 to 1.6 g/cm3 lies

between those of inorganic and polymeric species

proposed to be SEI components by several

researchers. This supports the current understanding

that the SEI is a composite of organic and inorganic

species.

The resonant frequency is plotted with time

during CV to investigate the SEI growth beyond the

first cycle in Fig. 58. The measured frequency can

be resolved into three parts based on a model that the

total frequency change is due to Li de-intercalation,

SEI formation, and a time-dependent term. The

time-dependent term is estimated from the frequency

change at OCV and is assumed not to entail charge

transfer. The frequency change for SEI formation is

separated by subtracting that for Li de-intercalation,

which is estimated from the charge change, and the

time-dependent term from the total measured

frequency. This process is displayed in Fig. 58. The

remaining SEI term shows that the SEI growth is

arrested after 60000 s (equivalent to the 15th cycle

under our experimental condition).

0 20000 40000 60000 80000

0

200

400

600

800

1000

1200

1400

measured

Li-subtracted

time-dependant

SEI

time, sec

- ∆f, Hz

Figure 58. Resonant frequency changes (measured, Li

de-intercalation-removed, time-dependent, and SEIrelated) with time during CV for the first 20 cycles in 1 M

LiClO4/EC+DMC.

PUBLICATION

H. Yang, K. Kwon, T.M. Devine and J.W. Evans,

ìAluminum Corrosion in Lithium Batteries ñ An

Investigation Using the Electrochemical Quartz

Crystal Microbalance,î J. Electrochem. Soc., 147,

4399 (2000).
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

62
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

63

TASK AREA 6: MODELING

Improved Electrochemical Models

John Newman (Lawrence Berkeley National Laboratory)

University of California, Department of Chemical Engineering, 201 Gilman Hall, MC 1462, Berkeley CA 94720

(510) 642-4063, fax: (510) 642-4778; e-mail: newman@newman.cchem.berkeley.edu

Objectives

• Develop experimental and computational methods for measuring and predicting transport, kinetic, and

thermodynamic properties.

• Model the behavior of electrochemical systems to optimize performance, identify limiting factors, and mitigate

failure mechanisms.

Approach

• Refine the galvanostatic polarization technique to measure transport properties.

• Develop a molecular dynamics program to predict diffusion coefficients in multicomponent solutions.

• Develop a thermal model that accounts for concentration effects in insertion electrodes.

• Use potentiometric techniques to measure entropy.

• Develop a model of factors affecting dendrite formation on Li metal.

• Use Raman spectroscopy to measure in situ concentration profile and dendrite growth.

Accomplishments

• Completed measurement of the transference number and diffusion coefficient of polyethylene-methylene oxide

(PEMO)-LiTFSI at 85o

C and compared to transition-time experiments.

• Calculated the uncertainty in the molecular dynamics simulations of diffusivities in aqueous electrolytes.

• Measured entropy of reaction in LiMn2O4, LiAl0.2Mn1.8O4F0.2, LiNi0.8Co0.2O2, and graphite MCMB 25-28, and

compared Monte Carlo simulations of entropy to measurements.

• Derived equations for determining the heat of mixing in electrochemical systems, and compared simulations of

heat generation with calorimetry experiments.

• Developed a model of the effect of surface tension on dendrite growth.

Future Directions

• Develop a model of how side reactions affect measurement of transport properties.

• Complete molecular dynamics simulation of diffusion in multicomponent nonaqueous electrolytes.

• Measure the entropy of reaction of tunnel Li0.44MnO2 and LiFePO4.

• Develop a model and conduct experiments to determine the influence of the SEI layer and of an elastic

conducting matrix on the rate of capacity fade in lithium-tin electrodes.

• Conduct in situ confocal Raman spectroscopy to measure concentration profiles across polymer electrolytes and

observe dendrite growth on Li electrodes.

• Simulate the effect of side reactions on the performance of ìrocking chairî dual-insertion cells.

In order to develop improved polymer

electrolytes, one must be able to measure the

transport properties that one seeks to improve. We

have developed and refined the galvanostatic

polarization method for measuring a complete set of

transport properties in polymer electrolytes, and, in
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

64

collaboration with John Kerr (LBNL), have

completed measurements for LiTFSI in

oxymethylene-linked PEMO at 85 ̊C and

concentrations ranging from 0.03 to 3.1 molar.

Unlike previous reports, negative values of the

transference number were calculated at low

concentrations, possibly due to side reactions and/or

ion association. Transition-time experiments, which

measure the time until depletion of salt at the

cathode during passage of a current density higher

than the limiting current, also showed evidence of

side reactions. A model is being developed to

understand how the occurrence of side reactions

would affect galvanostatic polarization and

transition-time measurements.

Measurement of the complete set of transport

properties in multicomponent solutions is extremely

difficult experimentally. We are developing a

molecular dynamics simulation program to predict

diffusion coefficients. The program has been

validated for aqueous salt solutions, and is now

being modified to simulate the nonaqueous liquid

electrolytes of interest for Li batteries. Difficulties

with the simulation algorithm have been encountered

due to the planar structure of many of the

carbonates.

Understanding heat generation is important both

for designing thermal management systems to

prevent overheating and for understanding

calorimetry experiments. Our thermal research has

three parts: derivation of an energy balance,

measurement, and prediction of thermodynamic

properties. An energy balance has been derived that

includes heat of mixing within the electrolyte and

the electrodes. The heat of mixing will be small for

electrolytes and electrodes that possess transport

properties high enough to be considered for

commercial batteries. That heat which goes into

formation of concentration gradients during passage

of current will be released during relaxation of

gradients after the current is interrupted.

Simulating heat generation requires knowledge

of thermodynamic, transport, and kinetic properties.

Measurements of entropy have been published in the

literature for only a few electrodes. Assuming that

the entropy of reaction is either constant or

negligible is inappropriate for insertion compounds

in which Li undergoes ordering on the lattice. Our

calorimetry experiments for LiAl0.2Mn1.8O4F0.2 have

demonstrated that entropic heat generation can be of

the same order of magnitude as resistive heat

generation, and that the entropic heat generation can

vary in sign and magnitude with state of charge.

Therefore, we have developed an experimental

method for measuring the entropy of reaction as a

function of state of charge, using LabVIEW to

automatically control a water bath, potentiostat, and

voltmeter. We plan to measure the entropy of

reaction for all of the electrodes being considered for

the BATT program. In addition, we have developed

a Monte Carlo simulation of LiMn2O4 spinel to

predict the entropy of reaction. The simulations

closely match the measured values.

Alloy electrodes are being considered as

alternatives to carbon or Li metal. We have

conducted a literature review to determine which

design options appear to successfully address the

issue of loss of contact in alloy electrodes. We have

learned that there are three simultaneous factors

contributing to loss of contact: alloy expansion and

contraction leading to cracking and particle

contraction away from a continuous electrical path

to the current collector, SEI formation around

freshly exposed surface area, and particle

agglomeration causing increased stresses during

lithiation and delithiation. Therefore any design

option must i) either involve no volume change or

create a driving force to push particles together

during contraction, ii) prevent resistive compounds

from forming between particles during volume

change, or iii) prevent particle agglomeration.

Lithium metal has the highest energy density of

any electrode. However, it will never be safe to use

commercially until one can figure out how to

prevent the risk of dendrites shorting the cell. We

are examining the effects of different physical

properties, such as surface energy, transport

properties in the electrolyte, and mechanical

properties of the electrolyte and Li metal, on the

growth rate of dendrites. A preliminary model of the

effects of surface energy has been completed. A

preliminary analysis using elasticity theory has

shown that the mechanical properties of the polymer

separator affect dendrite growth more significantly

than surface forces; furthermore, the plasticity of the

Li surface needs to be accounted for in the

mechanical analysis.

In situ confocal Raman spectroscopy can be

used to measure concentration gradients with high

spatial resolution across the electrolyte and to

observe dendrite growth. These experiments, to be

performed in collaboration with Robert Kostecki
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

65

(LBNL), will allow measurement of transport

properties as well as analysis of the effect of

concentration gradients on dendrite growth. Work is

still underway to construct a cell container that is

airtight, inert, transparent, can be thermostated, and

fits inside the confocal Raman microscope.

PUBLICATIONS

M. Doyle, J.P. Meyers and J. Newman, ìComputer

Simulations of the Impedance Response of

Lithium Rechargeable Batteries,î J. Electrochem.

Soc., 147, 99 (2000).

C. Fellner and J. Newman, ìHigh-Power Batteries

for Use in Hybrid Vehicles,î J. Power Sources,

85, 229 (2000).

K.E. Thomas, S.E. Sloop, J.B. Kerr and J. Newman,

ìComparison of Lithium-Polymer Cell

Performance with Unity and Nonunity

Transference Numbers, J. Power Sources, 89, 132

(2000).

H. Hafezi and J. Newman, ìVerification and

Analysis of Transference Number Measurements

by the Galvanostatic Polarization Method,î J.

Electrochem. Soc., 147, 3036-3042 (2000).

PRESENTATION

H. Hafezi, S.E. Sloop, J.B. Kerr and J. Newman,

ìMeasurement and Analysis of the Transport

Properties of Lithium Trifluoro-methane

sulfonylimide (LiTFSI) in Oxymethylene-linked

Poly(ethylene oxide) (PEMO),î 198th Meeting of

the Electrochemical Society, Phoenix, AZ,

October 22-27, 2000.

Microstructural Modeling of Highly Porous Fibrous and Particulate Electrodes

Ann Marie Sastry

The University of Michigan, Department of Mechanical Engineering and Applied Mechanics, Ann Arbor, MI 48109-2125

(313) 764-3061; fax: (313) 747-3170; e-mail: amsastry@engin.umich.edu

Objectives

• Perform AFM imaging of Li-ion electrodes to better define simulation parameters for carbon morphologies and

connectivities.

• Refine and expand simulations on conduction and mechanical load transfer in particle networks, and couple

modeling of damage and mechanics phenomena.

• Elucidate the effects of small-scale particle-particle interconnects in materials.

• Develop common platforms for numerical implementation, suitable for use in analyzing a wide range of materials of

high interest to DOE and industry.

Approach

• Develop 3D techniques for representation of particulate/fibrous networks, including models for intersections.

• Develop closed-form approximations for materials behavior dependence upon particle/fiber junctions.

• Characterize damage progression through continued simulations, and test the validity of numerical approaches.

• Expand experimental investigations to include testing of Li-ion cells.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

66

Accomplishments

• Solved mechanics problems of closed-form small-scale composite particles (fiber-fiber).

• Investigated particulate geometries; developed algorithms for defining particle junctions.

• Completed 3D→2D correlations for small-scale mechanics problems.

• Continued development of techniques for coupled transport/mechanics for fibrous/spheroidal morphologies.

Future Directions

• Develop 3D→2D simplifications for particle/fiber morphologies.

• Continue full-scale cell testing of Li-ion cells to benchmark theoretical results from damage analysis.

• Package codes and combination with electrochemical modeling.

Degradation of Li-ion anode materials is

currently a persistent problem, with damage

presenting as blistering and/or delamination. The

loads involved in these complex networks are

multiaxial, and degradation affects not only the

networksí ability to withstand subsequent loads, but

also negatively affects conductivity. A range of

methods for representing the arrays of particles

within these materials have been explored, as shown

in Fig. 59.

Figure 59. Four possible representations for porous

materials comprised of fibers and particles, including (a)

beams representing fibers, with spherical particles, (b) 2D

beams representing fibers, (c) 2D ellipses representing

fibers and (d) ellipsoidal particles of varying aspect ratio,

representing both fibers and particles.

The connectivity of the network and the bond

densities play a key role in assessing scale effects in

simulation. Moreover, anode materials are

heterogeneous mixtures of particles which are

imperfectly bonded and non-straight. In light of the

significance of bond type, we have concentrated on

modeling the 3D interconnects in fiber/particle

networks. We characterized the interpenetration of

fibers or particles as degree of intersect (d.o.i.), as

shown in Fig. 60.

Normalized effective moduli and peak stresses

have been generated over several orders of

magnitude in torsion spring constant representing

bond stiffness, and for a wide range representative

cell sizes. The reduction in both modulus and

strength are more dramatic for compliant, rather than

stiff bonds, as the scale of simulation size increases,

in all cases. For the range of simulation sizes

studied, there were no plateaux in these values,

suggesting that unlike the conductivity problem, the

mechanics problem requires a statistical approach

detailed enough to quantify bond effects. For rigid

bonds, the effect of on moduli and strengths over

three orders of magnitude of simulation lengthscale

were moderately small; few networks studied,

however, appeared from image analysis to contain

well-defined fiber-fiber bonds. These results

suggest that for multiphase networks, the statistics of

connectivity of even unpercolated phases are

warranted.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

67

(a)

(b)

Figure 60. 3D modeling with fillet. Detail of the joint of

a finite element model of two intersecting fibers with (a)

d.o.i.=0.571 and (b) d.o.i.=0.667.

Conduction modeling has similarly been focused

on heterogeneous arrays, and the importance of nonpercolated phases embedded within a percolated

phase (Fig. 61, with material parameters chosen to

be consistent with commonly used geometries of

particles in Li-ion anodes) has been established,

mirroring results for the mechanical simulations.

The implications of our findings include the

importance of consideration of all phases in the

material, concentration of efforts on representation

of junctures rather than further attempts at resolving

individual particle shape, and the importance of

connective phases (e.g., binders) in Li-ion cells in

assuring connectivity and mechanical and

conductive performance.

PUBLICATIONS

A.M. Sastry, C.W. Wang and L. Berhan,

ìDeformation and Failure in Stochastic Fibrous

Networks: Scale, Dimension and Application,î

Probabilistic Approaches to Fracture and

Fatigue, W. Soboyejo, Ed., Trans Tech

Publications, Ltd. (2000).

C.W. Wang, L. Berhan and A.M. Sastry,

ìStructure, Mechanics and Failure of Stochastic

Fibrous Networks: Part I. Microscale

Considerations,î ASME J. Eng. Mater. Technol.,

122, 450-459 (2000).

C.W. Wang and A.M. Sastry, ìStructure, Mechanics

and Failure of Stochastic Fibrous Networks: Part

II. Network Simulation and Application,î ASME J.

Eng. Mater. Technol., 122, 460-468 (2000).

PRESENTATIONS

A.M. Sastry, C.W. Wang and L. Berhan, ìLocalized

Strain and Damage in Fiber/Particulate Networks:

Damage Mechanisms in Li-ion Carbonaceous

Substrates,î Workshop on Interfaces, Phenomena,

and Nanostructures in Lithium Batteries, Argonne

National Laboratory, Chicago IL, December 2000.

A.M. Sastry, C.W. Wang and L. Berhan,

ìDeformation and Failure in Stochastic Fibrous

Networks: Scale Dimension and Application,î

The 2000 ASME International Mechanical

Engineering Congress & Exposition, Orlando, FL,

November 2000.

C.W. Wang, L. Berhan and A.M. Sastry, ìDamage

Tolerance of Fiber-Particle Networks with

Imperfect Bonding: Applications for Battery

Technologies,î The 2000 ASME International

Mechanical Engineering Congress & Exposition,

Orlando, FL, November 2000.

A.M. Sastry, Invited Presentation, ìStochastic

Materials: Design and Selectivity,î ONR/DARPA

Synthetic Multifunctional Materials Workshop,

Sedona, AZ, September 2000.

A.M. Sastry, Invited Presentation, ìIntegration of

Battery Power Supply in Ultralight Metal

Structures,î Ultralight Metals Workshop, ONR

Ultralight MURI, Brewster, MA, August 2000.

A.M. Sastry, Invited Seminar, Princeton Materials

Institute, Princeton University, Princeton, N.J.,

April 2000.

A.M. Sastry, Invited Seminar, Department of

Mechanical Engineering, Rutgers University,

Piscataway, N.J., April 2000.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

68

Figure 61. Networks in a unit cell (original volume

fraction ≈ 50%) comprised of (a) ellipse, and (b) circles.

Images (c) and (d) are voltage contours for fields (a) and

(b). Resulting normalized effective conductivities for

there arrays arrays, Keffective/Kfiber are (a) 0.0708, and

(b) 0.0 236, respectively

Thermal Modeling/Thermal Management

James W. Evans (Lawrence Berkeley National Laboratory)

University of California, Department of Materials Science and Mineral Engineering, 585 Evans Hall, MC 1760,

Berkeley CA 94720

(510) 642-3807, fax: (510) 642-9164; e-mail: evans@socrates.berkeley.edu

Objective

ï Study heat generation, heat transfer, and thermal management of large-scale batteries for EV applications.

Approach

• Develop simulation to express the heat generation rate based upon electrical and thermal state of battery.

• Experimental investigation of heat generation rates in small laboratory sized batteries.
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

69

Accomplishments

• Prepared calorimeter for measurements.

• Prepared galvanostat program for cyclic discharge.

• Compared 2D coupled results to original ìthermal-onlyî model.

Future Directions

• Continue work to expand simulation to 3D.

• Measure thermal properties of Li/polymer batteries as a function of cyclic discharge.

• Reanalyze previous perturbation comparisons.

The primary focus of this research, for 2000, has

been the development of the coupled model, and the

extension of the model from 2D into 3D. Part and

parcel of this modeling is a comparison of the

coupled model to results predicted by the previous

ìthermal onlyî model and the perturbation analysis

presented by Song and Evans.1

In previous, ìthermal-onlyî models, classical

thermal conductivity was used to determine the

thermal distribution. The heat generation was taken

from the electrochemical parameters via simplified

relations, such as the Shepherd equations.2

However, it is well known that there is a direct

relationship between the local temperature inside a

battery stack and the local conductivity. In fact it is

this relationship which provides the opportunity for

thermal run-away of a battery system. A coupled

model was developed by Song and Evans to relate

the local thermal parameters of a stack to the ionic

conductivity.3

A comparison was made of the older ìthermalonlyî model and the coupled model. Figures 62 and

63 show a comparison of these two models for a

perturbation simulation. In these figures it is seen

that upon perturbation the coupled model returns to

thermal stability much slower than the thermal-only

model. This demonstrates that the coupled model

behaves as expected as an increase in conductivity

(hence heat generation) is induced locally at the site

of the perturbation, therefore return to equilibrium is

delayed.

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1

-0.25

0

0.25

0.5

0.75

1

Temperature difference as a function of the 

state of discharge: thermal only model,

constant current 0.5mA/cm^2

Omega=10

Omega=20

Omega=30

Omega=40

Omega=50

Utilization

(Tmax-Tmin)/Tpert

Figure 62.

0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1

0

0.2

0.4

0.6

0.8

1

Temperature difference as a function of the state 

of discharge: coupled model, constant current 

0.5mA/cm^2, electrolyte 120(10^-6)m

Omega=10

Omega=20

Omega=30

Omega=40

Omega=50

State of discharge

(Tmax-Tmin)/Tperturbation

Figure 63.

During the course of this comparative study

several questions were raised regarding the methods

used in previous studies on perturbation analysis

(using the thermal-only model). A startling

discovery was that the parameters selected for the

analytic solution showed the model becoming

unstable only when the perturbations were the same

size as the cell itself, which is a violation of the

assumptions of perturbation theory. Figure 64

shows a plot of the system's response for a

perturbation in the limit of stability as given by Song

and Evans. For each data set there was no change in
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

70

the system, except the amplitude of the perturbation.

According to an analytic solution, the response

should be independent of the amplitude, and depend

only upon the wavenumber. This indicates that there

was a violation of the assumptions and the resulting

numerical comparison, in the high wave number

limit, is invalid.

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1

0

0.25

0.5

0.75

1

1.25

1.5

Edge Effect: Results for Various 

Perturbation Amplidutes for Thermal Model,

 omega=10, 5.0mA/cm^2

6.9

14.2

32.45

39.75

77.32

Utilization

(Tmax-Tmin)/Tpert

Figure 64.

FY 2001 will include a reanalysis of the

previous 2D perturbation analysis within appropriate

limits and with accurate thermal and electrochemical

parameters.

Also in FY 2000 work began to expand the 2D

coupled model to 3D. This is important as in a

prismatic cell, there are two directions of easy

conductivity, and one direction of hard thermal

conductivity, as shown in Fig. 65. The current 2D

model only uses one of the easy direction the hard

direction. Expanding the cell into 3D will give a

more accurate interpretation of the battery. It is

hoped that results from this 3D model will be

available by the end of FY2001.

Easy

Easy

Hard

Prismatic Stack

Type Battery

Figure 65.

REFERENCES

1. Song and Evans, J. Electrochem. Soc. 145, 2328

(1998).

2. Chen and Evans, J. Electrochem. Soc. 141, 2947

(1994).

3. Song and Evans, J. Electrochem. Soc. 147, 2086

(2000).
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

71

LIST OF ACRONYMS

AFM atomic force microscopy

AIBN azo-bis-isobutyronitrile

ANL Argonne National Laboratory

ARC accelerating rate calorimeter

ATD Advanced Technology Development

BATT Batteries for Advanced Transportation Technologies

BES Basic Energy Sciences

BMA butyl methacrylate

BNL Brookhaven National Laboratory

CB comb-branch

ccp cubic close packed

CPE composite polymer electrolyte

CSAFM current-sensing atomic force microscopy

CV cyclic voltammetry

DEC diethyl carbonate

DMC dimethylcarbonate

DME dimethoxyethane

DoD depth of discharge

DOE Department of Energy

d.o.i. degree of intersect

DSC differential scanning calorimetry

EC ethylene carbonate

EMC ethyl methyl carbonate

EQCM electrochemical quartz crystal microbalance

ETR Exploratory Technology Research

EV electric vehicle

EXAFS extended X-ray absorption fine structure

fcc face-centered-cubic

FR flame retardant

FTIR Fourier transform infrared

GC gas chromatography

GNP glycine-nitrate process

HETAP hexa-ethoxy-tri-aza-phosphazene

HEV hybrid electric vehicle

HMTAP hexa-methoxy-tri-aza-phosphazene

HOPG highly oriented pyrolytic graphite

ICL irreversible capacity loss

IIT Illinois Institute of Technology

IRRAS infrared reflection absorption spectroscopy

LBNL Lawrence Berkeley National Laboratory
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

72

LiTFSI Lithium bis(trifluoromethane-sulfonyl)imide

MAS Magic Angle Spinning

MCMB mesocarbon microbead

MEEP methoxyethoxyethoxide phosphazene

mpe mass change per mole of electron transfer

MSA Mine Safety Appliances

MW molecular weight

NFE nonflammable electrolyte

NMP 1-methyl-2-pyrrolidinone

NMR nuclear magnetic resonance

NRC National Research Council

OAAT Office of Advanced Automotive Technologies

OCV open circuit voltage

PC propylene carbonate

PEC polyethercarbonate

PEG-DME polyethylene glycol - dimethyl ether

PEMO polyethylene-methylene oxide

PEO poly(ethylene oxide)

PNGV Partnership for a New Generation of Vehicles

PTC phase transfer catalysis

PVDF poly(vinylidene fluoride)

PVIC poly(vinylene carbonate)

PVICOX poly(1,3-dioxolan-2-one-4,5-diyl oxalate)

ROP ring-opening polymerization

SE spectroscopic ellipsometry

SEI solid electrolyte interphase

SERS surface-enhanced Raman spectroscopy

SHR self heat rate

TFEA trifluoroethyl acetate

TGA thermogravimetric analysis

THF tetrahydrofuran

TOM trichlorosilyl octyl methacrylate

UDP urea-diphosphonate

UHV ultra high vacuum

USABC United States Advanced Battery Consortium

VC vinylene carbonate

XANES X-ray absorption near-edge spectroscopy

XAS X-ray absorption spectroscopy

XPS X-ray photoelectron spectroscopy

XRD X-ray diffraction

XRDA X-ray diffraction analysis
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

73

ANNUAL REPORTS

1. Exploratory Technology Research Program for Electrochemical Energy Storage – Annual Report for 1999,

LBNL-45767 (May 2000).

2. Exploratory Technology Research Program for Electrochemical Energy Storage – Annual Report for 1998,

LBNL-42694 (June 1999).

3. Exploratory Technology Research Program for Electrochemical Energy Storage – Annual Report for 1997,

LBNL-41950 (June 1998).

4. Exploratory Technology Research Program for Electrochemical Energy Storage – Annual Report for 1996,

LBNL-40267 (June 1997).

5. Exploratory Technology Research Program for Electrochemical Energy Storage – Annual Report for 1995,

LBNL-338842 (June 1996).

6. Exploratory Technology Research Program for Electrochemical Energy Storage – Annual Report for 1994,

LBL-37665 (September 1995).

7. Exploratory Technology Research Program for Electrochemical Energy Storage – Annual Report for 1993,

LBL-35567 (September 1994).

8. Exploratory Technology Research Program for Electrochemical Energy Storage – Annual Report for 1992,

LBL-34081 (October 1993).

9. Exploratory Technology Research Program for Electrochemical Energy Storage – Annual Report for 1991,

LBL-32212 (June 1992).

10. Technology Base Research Project for Electrochemical Energy Storage – Annual Report for 1990, LBL30846 (June 1991).

11. Technology Base Research Project for Electrochemical Energy Storage – Annual Report for 1989, LBL29155 (May 1990).

12. Technology Base Research Project for Electrochemical Energy Storage – Annual Report for 1988, LBL27037 (May 1989).

13. Technology Base Research Project for Electrochemical Energy Storage – Annual Report for 1987, LBL25507 (July 1988).

14. Technology Base Research Project for Electrochemical Energy Storage – Annual Report for 1986, LBL23495 (July 1987).

15. Technology Base Research Project for Electrochemical Energy Storage – Annual Report for 1985, LBL21342 (July 1986).

16. Technology Base Research Project for Electrochemical Energy Storage – Annual Report for 1984, LBL19545 (May 1985).

17. Annual Report for 1983 – Technology Base Research Project for Electrochemical Energy Storage, LBL17742 (May 1984).

18. Technology Base Research Project for Electrochemical Energy Storage – Annual Report for 1982, LBL15992 (May 1983).
FY 2001 Annual Progress Report Batteries For Advanced Transportation Technologies

74

19. Technology Base Research Project for Electrochemical Energy Storage – Report for 1981, LBL-14305

(June 1982).

20. Applied Battery and Electrochemical Research Program Report for 1981, LBL-14304 (June 1982).

21. Applied Battery and Electrochemical Research Program Report for Fiscal Year 1980, LBL-12514 (April

1981).
This document highlights work sponsored by agencies of the U.S. Government. Neither the U.S.

Government nor any agency thereof, nor any of their employees, makes any warranty, express or implied,

or assumes any legal liability or responsibility for the accuracy, completeness, or usefulness of any information, apparatus, product, or process disclosed, or represents that its use would not infringe privately

owned rights. Reference herein to any specific commercial product, process, or service by trade name,

trademark, manufacturer, or otherwise does not necessarily constitute or imply its endorsement, recommendation, or favoring by the U.S. Government or any agency thereof. The views and opinions of authors

expressed herein do not necessarily state or reflect those of the U.S. Government or any agency thereof.

Printed on recycled paper
www.cartech.doe.gov

DOE/EERE/OTT/OAAT - 2001/010

Office of Transportation Technologies

Series of 2001 Annual Progress Reports

• Office of Advanced Automotive 

Technologies FY 2001 Program Highlights

• Vehicle Propulsion and Ancillary Subsystems

• Automotive Lightweighting Materials

• Automotive Propulsion Materials

• Fuels for Advanced CIDI Engines and Fuel Cells

• Spark Ignition, Direct Injection Engine R&D 

• Combustion and Emission Control for 

Advanced CIDI Engines

• Fuel Cells for Transportation

• Advanced Technology Development 

(High-Power Battery) 

• Batteries for Advanced Transportation 

Technologies (High-Energy Battery) 

• Vehicle Power Electronics and Electric Machines

• Vehicle High-Power Energy Storage

• Electric Vehicle Batteries R&D