study of a buffer layer based on block copolymer

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Study of a buffer layer based on block copolymerelectrolytes, between the lithium metal and a ceramic

electrolyte for aqueous Lithium-air batteryLouise Frenck

To cite this version:Louise Frenck. Study of a buffer layer based on block copolymer electrolytes, between the lithiummetal and a ceramic electrolyte for aqueous Lithium-air battery. Electric power. Université GrenobleAlpes, 2016. English. NNT : 2016GREAI041. tel-01532066

https://tel.archives-ouvertes.fr/tel-01532066

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THÈSE Pour obtenir le grade de

DOCTEUR DE LA COMMUNAUTE UNIVERSITE GRENOBLE ALPES Spécialité : Matériaux, Mécanique, Génie civile, Electrochimie Arrêté ministériel : 7 août 2006

Présentée par

Louise FRENCK Thèse dirigée par le Pr Renaud Bouchet préparée au sein du Laboratoire d'Electrochimie et Physicochimie des Matériaux et Interfaces dans l'École Doctorale IMEP2

Study of a buffer layer based on block copolymer electrolytes, between the lithium metal and a ceramic electrolyte for aqueous Lithium-air battery Thèse soutenue publiquement le «16 septembre 2016», devant le jury composé de :

Dr, Elisabeth, Siebert Directeur de recherche, LEPMI, Grenoble, Présidente

Dr, Michel, Rosso Directeur de recherche, LPMS, Palaiseau, Rapporteur

Pr, Sylvain, Franger Professeur ICMMO, Orsay, Rapporteur Pr, Nitash, Balsara Professeur UC Berkeley and LBNL, Berkeley, Examinateur

Pr, Renaud, Bouchet Professeur LEPMI, Grenoble, Directeur de thèse

Dr, Philippe, Stevens Chercheur senior, EDF, Moret sur Loing, Co-encadrant

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A la mémoire de mon grand-père

Dr Allan Gustav Ringheim et de

ma grand-mère Julie Ringheim

A mes Chers parents et à ma

Chère sœur

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"Expérimenter, c'est imaginer."

Nietzsche

Un voyage de milles lieues commence toujours par un

premier pas.

Lao Tseu

"We live on an island surrounded by a sea of ignorance. As

our island of knowledge grows, so does the shore of our

ignorance."

John Archibald Wheeler

"Not everything that can be counted counts, and not

everything that counts can be counted."

Albert Einstein

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Remerciements

This PhD was for me a unique experience made of travels in France (between Paris and

Grenoble), but also to the US (Berkeley) and I had the chance to make numerous happy encounters

during these few years. If I was able to finish this ordeal, it was through the help of many people and I

would like here to thank them.

I would like to start by expressing my gratitude to all the members of my PhD committee, who

gave me the honor to evaluate my PhD work. First of all, I would like to acknowledge the two referees,

Sylvain Franger and Michel Rosso; they have taken the time to read and correct my "small" manuscript

and I would like to thank both of them for their constructive critics and comments. Also I would like

to thank Elisabeth Siebert to have accepted to be the president of my committee and for the very

interesting and pertinent discussion that we had.

Now, I would like to thank all of my three advisors for their help during those "three years".

This research work could not exist without the EDF project, therefore I would like to thank Philippe

Stevens for giving me the opportunity to work with him and his team on the Li-Air project, but also to

introduce me on the complexity of batteries and the industrial research.

Be a French researcher for one year at the Lawrence Berkeley National Laboratory was a great

experience for me and I would like to warmly thank Nitash Balsara for hosting me and making me

feel part of his laboratory. Furthermore, it was a great pleasure to work with him and to have such

stimulating individual meetings.

Last but not least, I would like to thank my PhD director Renaud Bouchet for his help and

his support from the beginning to the final end of this project including the worst part, "la redaction".

Thank you to have guided me on the path of lithium dendrites; starting from nothing it was not easy

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but I have learn so much all along this journey. You have showed me the way to become a scientist with

persistence, determination and scientific rigor. Thank you for all your guidance.

As a matter of fact, I had the chance to be in a great lab at the LBNL and I would like to

thank the whole Balsara group to welcome me and to show me how the lab works. I would like to

thank particularly some members: Chelsea for her help with the STEM. Dula and Katherine for their

help at the tomography beamline and for the data reconstruction. Adriana and Jacob for their help with

the SAXS experiment and treatment. Mahesh for your kindness in the office. And finally, Didier and

Irune for everything that you have done for me as super nice colleagues and great researchers, but also

as incredible friends ! In addition, I have met nice persons during this year in California, therefore I

would like to thank the "Frenchies" Juliette, Paul, Thomas and Igor for all the good moments at the

terrace of the Molecular Foundry and outside work, but also for the great weekends at Yosemite and

all the hiking trips, without forgetting Mike the American, Georg and Herman the Germans. And

also I can't forget my amazing neighbors: Ani, my favorite hippie, Jeff my amazing firefighter and

paramedic neighbor and Melissa his beautiful and kind wife. You were an inspiration source of how to

live as a real Californian. I really enjoyed all the moments at our rooftop. Sheila, thank you for showing

me your beautiful Arizona State, the discovery of such amazing landscapes was extraordinary.

J'aimerais maintenant remercier les personnes que j'ai rencontrées à EDF sur le site bien connu

des Renardières caché près du village fortifié de Morêt sur Loing. Tout d'abord, je dois remercier tout

le groupe M29 pour m'avoir accueillie et pour tous les bons moments passés avec eux. J'aimerais

remercier en particulier Gwenaëlle pour son aide au labo et au MEB, Marie-Christine pour toute son

aide avec tous les papiers administratifs (et il y en a eu beaucoup...) et pour sa bonne humeur

quotidienne. Merci à Delphine pour la bonne ambiance et nos discussions. Par ailleurs, j'ai eu la chance

de partager mon bureau avec mon Cher voisin Patrick, merci pour tous les fous rires et la bonne humeur

dans notre bureau plein de voyages, et merci de m'avoir fait découvrir le Club peinture. J'ai découvert

une autre partie de moi même à travers la peinture, merci à tous les membres du Club pour tous ces

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déjeuners artistiques. En particulier, merci à Patricia notre Présidente et à Pascale notre professeur avec

qui j'ai tant appris, mais aussi à Greg et Minh, qu'est ce qu'on a bien rigolé !

Les longues heures de train accumulées pendant ces années m'ont permis de faire de belles

rencontres, merci à mes chers collègues de navette et de Transilien : Thomas, Ken, Samuel, Adrien,

Houssam, Luis, Camille, Kevin et Jose pour tous nos papotages et rigolades, qui ont transformé ces

trajets en bons souvenirs.

Et oui; il reste encore un laboratoire dont je n'ai pas parlé, le LEPMI ! Malgré ma présence

intermittente; j'ai eu la chance d'être très bien accueillie par tous ses membres que j'aimerais

chaleureusement remercier. Par ailleurs, pour tous les très bons moments inoubliables que j'ai passé,

j'aimerais spécialement remercier Filou, Guillaume, Marc, Shayenne, Seng-Kian, Juan, Clément G

pour m'avoir introduit au monde de la fanfare, Clément M pour les nombreuses parties de jeu de go,

Marine pour toutes nos interminables discussions sur le Japon, Lulu pour toutes les bonnes soirées, Toc

pour ta joyeuse folie contagieuse, Lazou pour toutes les fois où tu es venu me changer les idées dans mon

bureau quand je rédigeais, et Juju pour toutes nos discussions scientifiques, non scientifiques, pour les

papotages et les cours photos. Grace à vous je me suis sentie entourée et soutenue dans cette ville qui

m'était inconnue.

Je souhaite maintenant remercier mes amis avec qui malgré la distance et les années passées rien

n'a changé, toujours la même joie de vous revoir et de passer du temps ensemble. Merci à Giselle pour

tous les bons moments passés à Berkeley. Merci à Jérem, Virgile et Prisca, Andrea et Basilus,

Chaussong, Kevin et Arnaud pour tous les super bons moments, les soirées, les verres et les rigolades à

la colloc ! Merci à Leslie pour toutes les randos, les soirées et pour tes encouragements pendant ma

rédaction. Alex et Fabien, merci pour tous les japonais, tous les fous rires et votre soutien. My Choupy,

cela fait maintenant bien trop d'années qu'on se connait pour les compter, merci d'avoir été là. Benj et

Stannou, c'est en grande partie grâce à vous que je me suis dirigée vers la thèse pour le meilleur et pour

le pire, merci de votre soutien et de votre amitié, à très vite pour un petit verre au St Hil ! Ma Clémence

merci pour tous les inoubliables moments à Stockholm, Orléans, Paris, San Francisco et Grenoble.

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Ma Chère Manue, merci pour ta bonne humeur infaillible qui m'a souvent reboostée, je ne peux plus

compter tous les fous rires que j'ai eu avec toi ! Marie, c'est grâce à toi que je me suis sentie bien entourée

au LEPMI, merci et à très vite pour de nouvelles aventures au Canada ou ailleurs sur la planète.

Irune, my Dear friend, it was a great surprise to find a beautiful person like you in the US, thanks to

you I felt at home and I was always surrounded by your positivity, your warm attitude and your smiling

face. Thank you a lot. Ma belle Yuki, mille mercis pour ta générosité, ton soutien, tes encouragements

et tes messages plein de pâtisseries ! Ma Len-chan, nous avons partagé tellement de beaux moments

ensemble, je n'ai qu'une hâte, c'est de repartir avec toi à la découverte de nouveaux pays. Merci pour ta

belle amitié.

Il n'y a pas assez de mots pour exprimer toute ma reconnaissance et ma gratitude à Juan-

Manuel. Tu es mon soutien de tous les instants et ta présence même lointaine me réconforte. Avec toi

tout semble plus facile...

Enfin, mes profonds remerciements vont à mes parents qui m'ont depuis toujours soutenue et

encouragée à aller plus loin. En dernier lieu, je remercie tendrement ma sœur Laura, pour sa belle

énergie, notre complicité et notre grande amitié.

Tables of contents

Introduction. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1

Chapter 1. General context and battery state of the art . . . . . . . . . . . . . . . . . . . 5

1. Societal and environmental context. . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

2. The main battery technologies. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11

3. Metal-air batteries. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 18

a. General context and metal-air batteries. . . . . . . . . . . . . . . . . . . . . . . . . . . . 18

b. Non aqueous lithium-air battery. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20

c. Aqueous lithium-air battery. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 26

4. Solid lithium ionic conductor as separator between the lithium and the aqueous

electrolyte. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

30

5. Possible ways to protect the ceramic. . . . . . . . . . . . . . . . . . . . . . . . . . . 33

a. Liquid electrolytes. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 34

b. Lithium phosphorus oxynitride or LiPON. . . . . . . . . . . . . . . . . . . . . . . . . 34

c. PEO-polymer based electrolytes. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36

d. Block copolymer electrolytes. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 40

e. Single-ion electrolytes. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 43

6. Lithium metal. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 45

a. The solid Electrolyte Interphase (SEI) . . . . . . . . . . . . . . . . . . . . . . . . . . . . 45

Tables of contents

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b. Model of nucleation and growth of lithium dendrite. . . . . . . . . . . . . . . . . . . . . . . 47

c. Lithium dendrite prevention. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 55

7. Conclusion. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 63

References of Chapter 1. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 65

Chapter 2. LiPON as a protective layer for the ceramic. . . . . . . . . . . . . . . . . . . 73

1. Experimental section. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75

a. Materials. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75

b. SEM characterization. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75

c. Electrode sputtering and cell assembly. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 76

d. EIS measurement. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 76

2. Results and discussion. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 77

a. Micro-structural analysis. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 77

b. Ohara GC results. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78

c. LiPON-Ohara GC-LiPON results. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 81

3. Conclusion. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 89


Chapter 3. Chemical and physical characterization of block copolymer electrolytes 93

1. Block copolymer : presentation and preparation. . . . . . . . . . . . . . . . . . . . . . 95

2. Thermodynamical properties of block copolymer electrolytes. . . . . . . . . . . . . . . 97

3. Morphology studies. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99

a. Small angle X ray scattering. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99

b. Dark field scanning transmission electron microscopy. . . . . . . . . . . . . . . . . . . . . . 108

4. Material electrical properties. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 109

a. Cell preparation. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 109

b. Cell optimization. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 110

c. Conductivity measurements. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 111

5. Transference number. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 118

6. Conclusion. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 123


Tables of contents

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Chapter 4. Dendritic growth in lithium symmetric cells. . . . . . . . . . . . . . . . . . 127

1. Cycling experiments followed by electrochemical impedance spectroscopy. . . . . . . . 129

a. Cycling routine. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 129

b. Neutral block copolymers. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 131

c. Single -ion block copolymers. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 136

2. Dendrites morphologies studied by hard X-ray microtomography. . . . . . . . . . . . . 145

a. Hard X-ray microtomography. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 145

b. Protocol. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 146

c. Neutral block copolymer. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 147

d. Single-ion block copolymer. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 152

3. Conclusion. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 157

References of the Chapter 4. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 159

Chapter 5. The polymer-ceramic composite. . . . . . . . . . . . . . . . . . . . . . . . . 163

I. Study of the polymer-ceramic composite. . . . . . . . . . . . . . . . . . . . . . . . . . . 165

1. Experimental procedure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 165


a. Electrical properties. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 167

b. Cycling. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 175

c. Characterization by hard X-ray microtomography. . . . . . . . . . . . . . . . . . . . . . . . 185

II. Quantification of polarization loss at the polymer-ceramic interface. . . . . . . . . . . . . 190

1. State of the art. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 190

2. Experimental procedure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 191


a. Experimental results. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 192

b. Discussion. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 196

Conclusions. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 201


Conclusions and perspectives. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 207

Résumé en français 211

Tables of contents

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Abstract. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 214

Introduction

The concept of sustainable development has been established in the XXth century, it was

defined as “a development that meets the needs of the present without compromising the ability

of future generations to meet their own needs”1. In a society more concerned by its impact on

the environment, and willing to have a sustainable development, energy plays an important role

if not the main. However, everything starts form the use of sustainable energy generated from

clean and renewable sources. One of the main issues for their development since there are

intermittent technologies is the energy storage. Therefore, it becomes a crucial challenge to get

different energy storage and especially electrochemical energy storage.

Nowadays, a large panel of electrochemical energy storage technologies are developed,

among them the lithium-air (Li-air) technology is one of the most promising. In particular, the

aqueous Li-air exhibits the highest specific energy and energy density compared to the other

technologies already developed or even compared to the technologies under development like

the lithium-sulfur technnology2.

However, the aqueous Li-air battery presents some issues which need to be addressed in

order to make this technology viable. One of the main issue is the reactivity of the lithium metal

electrode. The use of an aqueous electrolyte implies the protection of the negative electrode,

Visco et al.3 have proposed a protected lithium anode (PLA) composed of a bilayer of solid

Introduction

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electrolytes composed of Lithium Phosphorus OxiNitride (LiPON), which is stable in contact

with lithium metal, and a lithium ion conductor ceramic LATP (Li1+xAlxTi2-x(PO4)3) which is

stable in contact of the alkaline electrolyte. On the other hand, these materials are hard and

fragile, and during the charge of the Li-air battery, the volume variation of lithium leads to

mechanical constrains at the interface which results in a loss of contact at the lithium- LiPON

interface and then a loss of the active surface area. This is why, the protective layer should be

replaced by a material which exhibits: at first, softness property to keep the contact with the

lithium metal and secondly, it needs to mitigate dendritic growth during cycling to protect the

ceramic.

Block copolymer electrolytes (BCE) based on poly (ethylene oxide) (PEO) are good

candidates as protective buffer layers for the ceramic. Indeed, this materials is stable versus

lithium4 and they have been recently highlighted for their high performances in lithium metal

polymer battery5,6. The aim of this work is to study the use of the polymer electrolytes in the Li-

air technology, i.e. to replace the LiPON with BCE which will have two main goals, firstly to

assure the good contact between the lithium and the ceramic and secondly to protect the ceramic

from the lithium dendrites.

Such protective layer has to fulfill several criteria: to be stable and to ensure a good interface

with both lithium and ceramic, to present a high lithium ionic conductivity, to be feasible in

order to keep the contact with lithium even with the volume variation in charge and discharge

and finally to be resistive to dendritic growth in order to protect the ceramic from the contact

with the lithium metal.

In the Chapter 1, we will first introduced the general context of the energy, in order to

understand why the world needs to develop new electrochemical energy storage systems with

increased energy density, reliability, safety, etc. A state-of-the-art of the different battery

technologies will then be given, followed by a focus on the lithium air battery technology. The

insight of the aqueous Li-air technology will be developed and more particularly, what are the

possible suitable electrolytes which can be considered as a protective layer for the lithium ion

conductor ceramic, as well with the different issues encountered with the use of the lithium

metal electrode. The different models of dendritic growth, which have been developed, will then

be introduced, and finally the different methods used to prevent or mitigate the dendrite

nucleation and growth will be discussed.

Introduction

- 3 -

The Chapter 2 is dedicated to the study of the actual solution to protect the ceramic, i.e the

thin lithium phosphorus oxinitride (LiPON) films on the ceramic. The study of the electrical

properties of the ceramic on one side and the LiPON in the sandwich on the other side will be

studied by electrochemical impedance spectroscopy (EIS). Ionic conductivities and activation

energy will be calculated.

The physico-chemical characterizations (ionic conductivity morphologic property and

transport properties) of the BCEs used in this study will be given and discussed in the chapter 3.

Several techniques, such as differential scanning calorimetry (DSC), small angle X-ray scattering

(SAXS) and EIS will be used.

A cycling study of the BCE-lithium symmetric cells along with the morphological

characterization obtained in situ by hard X-ray micro-tomography of the cell before and after

cycling (post mortem) will then be discussed in Chapter 4.

Finally, in Chapter 5, the composite ceramic-BCE/lithium will be studied (interface by EIS,

cycling with EIS and finally DC polarization analysis and hard X-ray micro-tomography).

References of the Introduction

1. WCED, 1987; Bojo et al., 1992

2. Bruce, P. G., Freunberger, S. A., Hardwick, L. J. & Tarascon, J.-M. Li-O2 and Li-S batteries with

high energy storage. Nat. Mater. 11, 19–29 (2012).

3. Visco, S. J. & Nimon, Y. S. Protected lithium electrodes having tape cast ceramic and glass-ceramic

membranes. (2015).

4. Stone, G. M. et al. Resolution of the Modulus versus Adhesion Dilemma in Solid Polymer

Electrolytes for Rechargeable Lithium Metal Batteries. J. Electrochem. Soc. 159, A222–A227 (2012).

5. Bouchet, R. et al. Single-ion BAB triblock copolymers as highly efficient electrolytes for lithium-

metal batteries. Nat. Mater. 12, 452–457 (2013).

6. Devaux, D. et al. Optimization of Block Copolymer Electrolytes for Lithium Metal Batteries. Chem.

Mater. 27, 4682–4692 (2015).

Introduction

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Chapter 1.

General context and battery state of the art

Abstract

In a society where the energy demand continuously increases and where the fossil

energy is limited by the planet resources, the development of renewable energies and

electrical vehicles is a necessity. However, known energy storage technologies do not

exhibit sufficient high capacities for the needs of today and tomorrow society.

Therefore, new electrochemical storage technologies with higher capacities are required

for example to the development and the widespread of electrical vehicles. This first

chapter will introduce the general societal and environmental context in which this

project takes place. The main electrochemical storage will be then reviewed, before to

discuss more specifically around metal-air batteries. The discussion will highlight the

advantages of the aqueous lithium-air battery developed by EDF. Nevertheless, the

negative electrode as it is conceived today presents issues which still need to be

addressed. Therefore, alternative materials will be presented. Finally, a state of the art

of lithium dendrite will be given.

Chapter 1. General context and battery state of the art

- 6 -

Table of contents

Chapter 1. General context and battery state of the art .................................................. 5

1. Societal and environmental context ............................................................................................................ 7

2. The main battery technologies ................................................................................................................... 11

3. Metal-air batteries ......................................................................................................................................... 18

a. General context and metal-air batteries ......................................................................................................... 18

b. Non aqueous lithium-air battery ................................................................................................................... 20

c. Aqueous lithium-air battery.......................................................................................................................... 26

4. Solid lithium ionic conductor as separator between lithium and aqueous electrolyte ...................... 30

5. Possible ways to protect the ceramic ........................................................................................................ 33

a. Liquid electrolytes ......................................................................................................................................... 34

b. Lithium phosphorus oxynitride or LiPON ................................................................................................... 34

c. PEO polymer based electrolytes ..................................................................................................................... 36

d. Block copolymer electrolytes ........................................................................................................................... 40

e. Single-ion electrolytes ..................................................................................................................................... 43

6. Lithium metal ................................................................................................................................................ 45

a. The Solid Electrolyte Interphase (SEI) ......................................................................................................... 45

b. Model of nucleation and growth of lithium dendrites ..................................................................................... 47

c. Lithium dendrite prevention .......................................................................................................................... 55

7. Conclusion ..................................................................................................................................................... 63

References of Chapter 1 ....................................................................................................................................... 65


- 7 -

1. Societal and environmental context

Energy is known as "the lifeblood of modern societies"1. Indeed, it is critical for human

activities such as industrial manufacturing, agriculture, transportations, and communications.

Unfortunately, the high dependence of the global economy on fossil fuel makes it vulnerable to

two types of crisis that could arise in the near future : a supply disruption and an environmental

disaster.

Indeed fossil fuel supplies are by definition finite and under growing demand; on the other

hand reserves are concentrated in a small number of regions which increases geopolitical tension.

However the main concern is the rise of atmospheric, sea and land pollutions, which lead to

dramatic consequences on human and animal health, as well as on water quality and agricultural

production. The main greenhouse gas, cause of global warming, emitted by human activities to

the atmosphere is carbon dioxide (CO2), produced by fossil fuel use.

Society is becoming more conscious of the situation, and there is an increasing demand to

massively introduce renewable energies in the energy mix in order to mitigate CO2 emissions and

their effects on climate change1. In order to address climate change, countries from all around

the world met in Paris for the 21st Conference of the Parties to the United Nations Framework

Convention on Climate change (COP21) in 2015, to negotiate an international agreement and set

a direction for combating climate change and keep global warming below 2ᵒC2.

To be able to stay on the 2ᵒC scenario, a rise in renewable energies in the global power

generation is necessary. The International Energy Agency (IEA) asks for a rise of 45% of

renewable electricity generation between 2012 and 20203. Figure 1 presents renewable power

generation by region from 2000 to 2020.


- 8 -

Figure 1. Renewable power generation by region 3.

Figure 2 shows the breakdown of global renewable energy use in 2010 and an evolution

perspective for 2030 (REmap 2030), by technology and sector. The International Renewable

Energy Agency (IRENA) is expected to have 3% of electricity coming from solar sources and

11% coming from wind sources by 2030. Production from energy sources cannot be dispatched,

since technologies such as solar or wind are intermittent energies, and they are often generated

far from where they are consumed. These aspects of renewable energies represent a real

challenge for the management and reliability of the electrical grid4. The unpredictable nature of

energy from renewable sources means that it has to be stored to be available when required by

consumers. In order to provide the flat energy production curve to the grid, it is necessary to

couple the energy source to an energy storage. Energy storage has therefore emerged as one of

the greatest issues of the 21st century. The main interest of these technologies is that they can be

placed close to the place of consumption and they can be adapted in size.

0%

10%

20%

30%

0

2 000

4 000

6 000

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10 000

12 000

14 000

2000 2005 2012 2020 2025

Sh

are

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Other non-OECD Brazil IndiaChina OECD Europe OECD Asia OceaniaOECD Americas Share of renewable generation Share of renewable generation 2025

Forecast Targets


- 9 -

Figure 2. Breakdown of global renewable energy use in 2010 and in REmap 2030, by technology and sector 5.

One potential good candidate for energy storage is electrochemical storage thanks to

batteries; the technology used for renewable energies must satisfy several criteria such as a long

lifetime, low cost, effectiveness and safety6.

Besides, stationary storage of electrical energy from renewable sources that completes the

energy mix and limits the use of fossil energies, a sustainable modern society requires to electrify

its transport by providing electric vehicles that can compete with cars powered by internal

combustion engine. Moreover, according to the International Energy Agency (IEA), in order to

keep the increase in global temperature below 2ᵒC, it is necessary to have a decrease of at least

50% of greenhouse gas emissions by 2050 compared to 2005 levels3. This requires not only a

universal climate agreement, which implies strong climate policies (see the recent international

meeting COP 21), but also the widespread use of electric and hybrid vehicles (see Figure 3),

which can contribute to a decrease of 30% of the greenhouse gas emissions.

However, the transition to mass development of these technologies is directly related to

battery improvements, which is the struggling point of today. Indeed, today none of the

conventional batteries have met all the required specifications, which are in order of priority

safety, cost, long lifetime and a high energy density6.


- 10 -

Figure 3. Annual light duty vehicle sales by technology 3.

Tomorrow's electrochemical storage has therefore to fulfill a certain number of criteria

starting with safety, sustainable components, high energy density, scalability, ease of use and cost

of maintenance. However, for each applications the order of priorities are different. Indeed, for

stationary applications such as solar, wind or even tidal, weight and space are not critical whereas

the energy stored, power, lifetime and cost are crucial. In addition, stationary applications are

expected to have a lifetime between one and two decades, which implies around 3000 cycles. The

last specification for such application is the capacity which does not necessarily need to be very

high due to the possible over sizing. Whereas for loaded applications such as EV or HEV,

weight and space are very important on the contrary. Safety in both applications are very

important but need to be specific for each applications.

Nowadays, lithium ion (Li-ion) technology is extensively widespread for consumer portable

electronics such as cell phones or computers. This technology stays expensive and more

importantly the environmental cost is high due to some components for Li-ion electrodes, such

as cobalt, which are rare, expensive and toxic, as well as the presence of organic solvents in the

electrolyte. A battery for cell phone is a 4 Wh battery, when an EV needs 40 kWh. Thus, to build

an EV battery 10 000 cell phone batteries are needed. In other words, the 5 billion cell phones

on the planet is equivalent to only 500 000 of potential EVs, which is unfortunately a drop in the

car's ocean.

In the following sections, we will first review the main battery technologies since the

discovery of the voltaic pile to lithium-sulfur batteries including lithium-ion batteries and lithium

metal polymer batteries. Metal-air batteries will then be reviewed and the difference between


- 11 -

non-aqueous and aqueous technologies will be discussed. Their advantages and disadvantages

will be addressed and finally the discussion will focus on the evolution of materials composing

the aqueous lithium-air battery. Finally, the use of lithium metal and its related issues will be

discussed. The formation of passivation layers, the different models of dendritic growth and how

to prevent or limit the nucleation and growth of dendrites will be approached.

2. The main battery technologies

In this thesis, we will use the terms cathode for the positive electrode (where the reduction

takes place with a consumption of electrons during the discharge) and anode for the negative

electrode (where an oxidation reaction takes place with a production of electrons at the

discharge), these terms are usually used by the Anglo-Saxon scientific community.

As seen in Figure 4, the historical roots of the development of rechargeable battery can be

found in the 19th century. Alessandro Volta, an Italian physicist and chemist was the first to see a

continuous and stable current with his invention "the Volta pile", in 18007. In 1803 J. W. Ritter,

a German chemist physicist and philosopher, was the first to build a secondary battery 8.

Figure 4. Historical roots of the development of secondary batteries from 1803 to 1994 8.

Lead acid battery. A few decades later, reliable rechargeable battery history started with

Gaston Planté in 18599, a French physicist who invented the lead-acid battery. Figure 5

represents Planté's lead acid battery, where a spiral roll of two sheets of pure lead were separated


- 12 -

by a linen cloth immersed in a glass jar of sulfuric acid solution. The redox couples involved at

the electrodes are PbSO4/Pb at the cathode and PbO2/PbSO4 at the anode.

Twenty years later, the first commercial rechargeable lead-acid battery was marketed10,11.

More than a century after its invention, this battery is still widely used, and it represents 70% of

the secondary battery market mainly due to its use as starter batteries (in thermal motors), vehicle

lighting, engine ignition, but also as emergency power and backup systems.

Figure 5. First lead acid battery by G. Planté 9.

Nickel-Cadmium battery. In 1899, four decades after Gaston Planté, a new type of battery,

based on nickel-cadmium (Ni-Cd), was born with the discovery by Ernst Waldemar Jungner12. In

these batteries, the positive electrode is composed of nickel hydroxide (NiO(OH)) and the redox

couple involved is Ni(OH)2/NiO(OH), whereas the negative electrode is composed of metallic

cadmium (Cd) and the redox couple is Cd(OH)2/Cd. The commercialization was achieved in the

20th century with G. Neumann and his Ni-Cd sealed cell in 194713.

The nickel metal hydride battery (Ni-MH) is a technology similar to the Ni-Cd. Indeed, both

the positive electrode and the electrolyte are similar; the main difference lies on the use of

hydrogen absorbed in a metal alloy at the anode instead of cadmium. It was in the late 1960s, at

Dutch Philips Research Laboratories, that the LaNi5 compound was found to be able to absorb

reversibly large amounts of hydrogen14. A mature and reliable technology was commercialized in

the early 1990s15. This new battery type, which is cadmium free, is considered to be more

environmentally friendly and in addition it is recyclable. The metal alloy used is a mix of rare

earth and nickel (LaNi5).

Lithium-Ion Battery. In the continuous race for an always higher specific and volumetric

energy in battery technology because of the fast nomad electronics development, lithium has

been considered as a good candidate for the negative electrode as active material. In 1979 at


- 13 -

Oxford University, an American professor John B. Goodenough created a cathode material

based on lithium cobalt and lithium-manganese spinels7,16, which set the bases to the lithium ion

(Li-ion) battery. More than a decade later in 1991, Sony was the first to commercialize a lithium

ion battery also called "rocking chair" battery 17,7.

Lithium batteries firstly employed intercalation compounds as cathode and lithium metal as

the anode. However, the anode was then replaced by lithiated carbon due to safety issues leading

to the so called "Li-ion" technology. Years of research have produced a wide choice of cathode

materials for Li-ion batteries, among them lamellar compounds such as LiCoO2, lithium nickel

manganese cobalt oxide (NMC), lithium nickel cobalt aluminum oxide (NCA), spinel

compounds like LiM2O4 (with M = Ni, Mn) or even olivine compounds such as lithium iron

phosphate (LFP). Both cathode and anode materials are insertion materials and have a structure

adapted to lithium ion intercalation during oxidation and reduction processes. The electrolyte is

an organic liquid composed of a lithium salt (LiBF4, LiPF6, LiClO4, LiBC4O8 (LiBOB)) dissolved

in a mixture of organic solvents (ethylene carbonate EC, diethylene carbonate DEC, propylene

carbonate PC or dimethyl carbonate DMC) and impregnated in a porous polymer separator

made of polyolefin (polypropylene PP, polyethylene PE) 18.

Figure 6 shows a schematic of a classical Li-ion battery during discharge. The operating

temperatures in charge range from -20ᵒC to 60ᵒC, whereas in discharge from -40ᵒC to 65ᵒC 7. In

addition, thanks to the Li-ion battery, the specific energy density has increased compared to Ni-

MH technology and range from 100 to 250 Wh.kg-1 (or 220 to 400 Wh.L-1) 17. This technology is

ideal for nomad applications and it has quickly conquered this market. However, Li-ion battery

has a high cost and safety issues due to the use of flammable electrolytes.


- 14 -

Figure 6. Scheme of a common Li-Ion battery during the discharge 18.

PLiON battery. In 1994, Bellcore (former Telcordia) patented a plastic lithium ion battery

called PLiON 8. The main difference from a classic Li-ion is that the electrolyte is a copolymer,

PVdF-HFP (polyvinylidene fluoride-hexofluoropropylene), which contains two different

domains. One amorphous domain (HFP), which is gelled by the liquid electrolyte composed of

LiPF6 and a mixture of EC-PC, and a crystalline domain (PVdF) which assures mechanical

properties of the entire electrolyte in order to obtain a free standing film19. Figure 7 represents a

schematic of the PLiON battery. Despite similar specifications (150 Wh.kg-1 and 300 Wh.l-1) to

Li-ion battery, PLiON shows an interesting size and design, indeed the assembly process of this

battery is lamination and thus can be made as thin as a credit card.

Figure 7. Schematic diagram showing the construction of a polymer Li-ion cell (PLiON) 7.

In 1995, the introduction on the market of pouch cell, which uses flexible and heat sealable

foils, simplified battery packaging. Few years later, in 1999, US Oak Ridge National Laboratories

patented the first marketable lithium ion polymer battery 8.

Nevertheless, scaling up Li-ion technology, to store efficiently energy from the intermittent

renewable power resources and to widespread EV's and HEV's onto the consumer market,


- 15 -

remains a challenge due to issues such as safety because of flammable organic solvents, cost

(higher than 250-400 $/kWh) or even material availability 18. In 2014, Tesla in cooperation with

Panasonic launched the Tesla Gigafactory project in order to reduce Li-ion battery cost by more

than 30% by 2020 20. Figure 8 represents the expected evolution of Li-ion battery cost for the

next decade.

Figure 8. Evolution of Li-ion battery pack cost ($/KWh) from 2015 to 2025 20.

The battery community realized the need to go further than Li-ion batteries. Thus a novel

fundamental research activity has been started which focuses on different kinds of battery

technologies, often referred as "beyond Li-ion". Figure 9 represents practical specific energy for

electrochemical storage systems from lead acid system to beyond Li-ion technologies such as

metal-air, including zinc air and lithium-air batteries.

It is important to note that a real breakthrough is necessary in order to have a widespread

adoption of electric vehicles into the public market. We only have two options to increase the

energy densities, which is proportional to the product of the specific capacity by the emf, :

- increase the electro motive force (that is the future Li-ion see in Figure 9)

- increase the specific capacity of the active materials, which corresponds to metal-air and

lithium-sulfur technologies


- 16 -

Figure 9. Practical specific energy perspective for some rechargeable batteries, along with estimated pack prices 21.

In the quest for high energy density for the negative electrode, lithium metal has inevitably

emerged from the others candidates as a promising active material, because its electrochemical

characteristics are unique. Indeed, metallic lithium can be considered as the ultimate negative

electrode due to its high theoretical specific capacity (3862 Ah.kg-1) and its very negative

potential (-3.05 vs SHE) 7. The idea to use lithium metal in an electrochemical device dated back

to 1949, when J.J. Halek patented a primary battery using lithium metal as an anode 22. Few years

later, in 1957 this idea was specified by D. Herbert and J. Ulam in a secondary battery 23, and in

1970, Watanabe and Fukuda for Panasonic (former Matsushita Electronic Co. Ltd) started its

production 24. At this time, lithium metal battery was assembled with a cathode of TiS2, and an

electrolyte composed of a lithium salt (lithium perchlorate (LiClO4)) dissolved in a mixture of

organic electrolytes, which is made of 70% of tetrahydrofuran (THF) and 30%

dimethylformamide (DMF) 7. However, the reactivity of the high surface area of lithium that is

formed during cycling leads to a poor cyclability and safety issues 25,26. This was due to irregular

lithium deposition during charge27,28. This heterogeneous electro deposition, also called

"dendrites", shows needle-like and mossy morphologies that can shortcut the battery by passing

through the electrolyte and potentially causing fire or explosion7. For example, in 1989 a lithium

metal battery in a cellular phone burned during operation, this was due to an internal short of the

battery 29.

Lithium metal polymer. In order to prevent these kinds of safety issues, two research axes

have been developed. The first one, which was already described above, consists in replacing


- 17 -

lithium metal by an intercalation compound, lithiated carbon, lithium titanate (Li4Ti5O12) or metal

alloys, which have led to the Li-ion battery technology. The second one consists on replacing the

organic liquid electrolyte by a dry solid polymer electrolyte, the lithium metal polymer (LMP)

technology was born. The first to suggest a polymer as a good electrolyte for lithium batteries

was M. Armand during the Second International Meeting of Solid Electrolytes in 197830. He

opened new perspectives for the development of solid polymer electrolytes which led 40 years

later to a wide variety of polymer electrolytes and new lithium salts. Polymer electrolytes are

composed of a lithium salt (typically LiClO4 or lithium bis (trifluoromethane)sulfoimide salt

(LiTFSI)) incorporated into a polymer matrix (mostly based on poly (ethylene oxide) PEO)

which is then cast into thin films. LMP presents some advantages compared to the classic Li-ion,

a good flexibility of polymers which enables the design of thin batteries to be made in a wide

variety of configurations, and a higher safety thanks to a volatile solvent free technology.

However, the ionic conductivity at room temperature for such electrolytes is low and therefore

the operating temperature needs to be increased to 80ᵒC. This technology has been developed

for electric vehicles such as the Bluecar and the Bluebus by Blue solutions-Bollore and is now

widespread in Paris as the Autolib 31,32.

Lithium-Sulfur battery. In order to renew the interest in the lithium metal battery concept,

radical changes in the approach of the fundamental electrochemical processes are necessary. In

this context, another type of battery is under intense scrutiny, the lithium sulfur battery (Li-S8).

Indeed, this technology shows appealing specifications, such as a practical energy density ranging

from 400 to 600 Wh.kg-1, a cathode active material abundance (lower cost) and the non-toxicity

of elemental sulfur (environmental friendly)21,33. However, after its discovery in 195723, this

technology sank into oblivion because of its poor performances at that time. Since 2009, after

Nazar34 reported a Li-S battery with improved cycling performance, this Li-S battery got back on

track, becoming one of the most studied technologies for next generation electrochemical

storage. Nevertheless, some important barriers still prevent the realization of a practical Li-S

battery with a high energy density and a long lifetime.

The Li-S8 battery is composed of elemental sulfur (S8) as the positive electrode, lithium metal

as the negative electrode and an electrolyte in between (which can be a liquid electrolyte, an ionic

liquid based electrolyte or a solid polymer electrolyte)35. The overall reaction involved in the cell

is 16 Li + S8 ↔ 8 Li2S. However, before obtaining the ultimate product, the lithium sulfide

(Li2S), intermediate polysulfides (Li2Sx with x = 2-8) are generated by the reduction of S8 during

the discharge process. Those molecules are dissolved in the electrolytes, which leads to an


- 18 -

irreversible capacity fade36,37. They can migrate through the electrolyte to the lithium metal

electrode in a so called, shuttle effect, and form an electrochemical insulating layer composed of

Li2S2 and Li2S leading to a deterioration of the battery performance and a poor rate capability.

Figure 10 shows the polysulfide shuttle mechanism and the deterioration of the lithium anode.

Figure 10. Schematic illustration of the polysulfide shuttle mechanism during the charge 38.

In addition, growth of dendrites from the lithium metal anode still causes internal short cut

issues. These lithium metal issues, such as dendrite growth, remain unsolved and scientists need

to dive into the complex life of interfaces and mechanisms involved at this attractive material

surface. In spite of their potential attractive advantages, Li-S8 batteries are not yet a mature

technology and need further improvements.

After introducing the evolution of battery technologies since its first stammering on the early

Nineteenth century, the following section will focus on a new kind of batteries, the metal-air

batteries.

3. Metal-air batteries

a. General context and metal-air batteries

The metal-air battery is a good candidate for EV's and large electricity storage systems due to

its high gravimetric and volumetric energy density (see Table 1). Metal-air batteries are composed

of :


- 19 -

· A metallic negative electrode; which can be made of lithium (Li), sodium (Na),

magnesium (Mg), aluminum (Al), iron (Fe) or zinc (Zn).

· An air positive electrode which uses oxygen from the ambient air as the active

material, and is composed of catalysts for O2 reduction reaction (ORR) and oxygen

evolution reaction (OER) in a porous network of an electrically conducting

supporting materials.

· An electrolyte which can be either an organic electrolyte or an aqueous electrolyte39.

One of the main advantage of such a device is, in principle, that the oxygen is supplied by

the surrounding atmosphere and does not need to be stored inside the battery40. Thus, this type

of battery has a reduced weight and more available space for energy storage. Consequently,

metal-air batteries presents high theoretical specific energy ranging from 1086 Wh.kg-1 for zinc-

air system to 3582 Wh.kg-1 for aqueous lithium-air system.

Table 1. Theoretical cell voltage with specific energy and energy density for various metal-air battery compared to Li-ion 40. a based on the volume of ZnO at the end of the discharge. b Based on the sum of the volume of Li at the beginning and Li2O2 at the end of the discharge. c Based on the sum of the volume of Li + H2O consumed and

LiOH at the end of the discharge. d Based on the sum of the volume of Li at the beginning and Li2S at the end of discharge. e Based on (Na+ and Na2O2). h Based on the sum of the volume of Mg at the beginning and Mg(OH)2

at the end of the discharge

Figure 11 is a general scheme of a metal-air battery. This schematic view includes a lithium

ion conducting membrane to protect the lithium metal for example in the case of aqueous Li-air

battery. During the discharge, the metal is oxidized and the metal ions produced migrate through

the electrolyte to the positive electrode where the oxygen is reduced.


- 20 -

Figure 11. Schematic principle of a metal-O2 battery during discharge40.

Despite their attractive specifications, metal-air batteries are still under development. Indeed,

only few realistic prototypes of metal-air batteries exist. This is due to complex challenges facing

i) the design of a rechargeable metal-air battery, ii) the difficulties to reach theoretical specific

energy due to parasitic chemistry occurring during metal-air electrochemistry and iii) safety

issues.

Lithium-air batteries have the highest potential energy density above all other metal-air

technologies, this is due to the use of light materials for the negative electrode, i.e. the lithium

metal.

Two types of Li-air batteries exist, the non-aqueous lithium-air battery and the aqueous

lithium-air battery. The principle of each type of technology will be discussed in sections

hereafter. Moreover, their different advantages and disadvantages will be considered.

b. Non aqueous lithium-air battery

Non aqueous lithium-air battery was first reported by Galbraith in 1976 41, but it was only

twenty years later in 1996 that Jiang and Abraham demonstrated the working principle in a

secondary battery42. The first Li-O2 cell was composed of a conductive organic polymer

electrolyte sandwiched between a thin Li metal foil and a thin carbon composite electrode (the

air electrode). One decade later, the interest in Li-air battery increased thanks to Bruce and

Ogasawara, who proved in 2006 that Li2O2 could potentially form a rechargeable couple 43, they

actually showed that Li2O2 is removed from the electrode during charge. Consequently, the


- 21 -

overall reaction 2 Li+ + 2e- + O2 ↔ Li2O2 may be reversible. In 2009, IBM and Energy

Laboratories started exploratory research programs on Li-air battery, since then this research

field has grown exponentially 43.

There have been controversies about the mechanism of O2 reduction in the presence of

lithium ions leading to the formation of lithium peroxide . Indeed, nowadays the consensus is

that the reduction of O2 during discharge follows a mechanism in three steps ((1) to (3))44:

O2 + e- → O2- (1)

O2- + Li+ → LiO2 (2)

2LiO2 → Li2O2 + O2 (3)

Reduction mechanism of O2 implies firstly the formation of a superoxide O2- which then

reacts with Li+ to form LiO2 on the surface of the electrode. Then, because LiO2 is unstable, it

will disproportionate in Li2O2 following equation (4).

However, other studies proposed the direct reduction of O2 into Li2O245.

2Li+ + O2 + 2e_ → Li2O2 (4)

Which may be further reduced to lithium oxide according to:

Li2O2 + 2Li+ +2e- → 2Li2O (5)

During the charge process (OER) the pathway is different to the discharge and follows

equation (6) hereafter:

Li2O2 → 2Li+ + 2e- + O2 (6)

Indeed, it was proved that the oxidation of Li2O2 is direct and does not pass through LiO2 as

an intermediate44. A result of these different pathways is the observed gap in charge and

discharge voltages (see Figure 12) resulting in an extremely low energy efficiencies, which would

limit the use of this battery in practical applications.


- 22 -

Figure 12. Voltage gap for a non-aqueous lithium-air battery 21.

A suitable electrolyte is the key component of this system and remains a real challenge.

Indeed, it has to fulfill various requirements coming from both the anode side and the cathode

side specific needs. From the anode side, the electrolyte in contact with lithium metal will react

and form a SEI like layer (solid electrolyte interphase), which has to be cohesive and flexible to

ensure the anode protection. In addition, this layer has to be sufficiently Li+ conducting in order

to ensure smooth lithium plating without dendritic growth.. Moreover, contaminants such as

oxygen, water, nitrogen and carbon dioxide, coming from the air electrode must not cross

through the electrolyte and react with the lithium surface46.

From the cathode side, the electrolyte has to present a low volatility to avoid its evaporation

at the open air electrode. It also has to present a high O2 solubility and diffusion to ensure

satisfactory rate capability, and to be able to wet the electrode surface. However, the most

challenging requirement for the electrolyte is to be stable to both O2- and its reduced species as

LiOx compounds that form during the discharge. Indeed, historically carbonate based organic

electrolytes were used, mostly because they were well known to be compatible with lithium

metal, they present a low volatility and a high oxidation stability (> 4,5V vs Li+/Li). An example

of a typical electrolyte is a lithium salt (LiPF6) in a mixture of propylene carbonate and dimethyl

carbonate (50/50). However, studies have shown that O2- in aprotic solvent reacts with organic

substrates via nucleophilic attack47. In 2010, Mizuno et al. reported that carbonates based

electrolytes are degraded by the superoxide radical O2- 48.


- 23 -

Other organic electrolytes were therefore investigated in order to find the ideal one. Ester-

based electrolytes were potential candidates but a computational study on esters by Bryantsev et

al.49 revealed that, similarly to carbonates, the superoxide radical attacks the ethereal carbon atom

in both cases of linear and cyclic ester. Another study of Bryantsev et al.50 has focused on

predicting the stability of a wide range of solvents for non-aqueous Li-air batteries. They

computed free energy barriers (∆Gact) for reactions with superoxide O2-, and showed that

solvents with a ∆Gact < 20 kcal/mol are chemically unstable against the superoxide, whereas

solvents with ∆Gact > 24 kcal/mol do not present reactivity with the superoxide, which make

them good candidates as electrolyte. Among them, tetraethylene glycol dimethyl ether

(TEGDME), presents a low vapor pressure, a high lithium salt solubility as well as a large

electrochemical window spanning up to 4.5 V. However, ethers show an auto oxidation under

oxygenated radicals. The main product of the discharge appears to be ether decomposition51,

similar behavior was found with cyclic ethers. Nitrile-based solvents have also been investigated.

In fact, acetonitrile presents a sufficient stability towards oxygen reduction species. However, it

shows a high vapor pressure at room temperature, which will leads to an evaporation of the

electrolyte at the positive electrode46. In addition, it is not stable against lithium metal. Therefore,

in order to have a low vapor pressure, alternative nitriles have to be studied. It is important to

notice that studies about nitrile-based solvents for non-aqueous Li-air battery are still on their

primary stages. In addition, the long term stability of nitrile components towards superoxide ion

are not yet proven.

Another approach has been investigated, the use of solid polymer electrolyte (SPE). Indeed,

numerous polymer systems have been studied for lithium batteries, the most used is based on a

poly(ethylene) oxide matrix hosting a lithium salt, generally lithium trifluoromethanesulfoimide

LiCF3SO3 or LiTFSI. SPE is a promising alternative for the volatile organic solvents 52. Polymers

are expected to react slowly due to the absence of convection and diffusion selective if compared

to organic solvents 53. However, the electrochemical stability of the very long molecular chain

glymes is debatable because the short chain glymes react with lithium oxides species 54. In

addition, the recharge ability of such system has not been yet proven 46. Moreover, the challenge

facing SPE is to improve their ionic conductivity and manage the huge volume change in the air

electrode between charge and discharge 55.

Since almost two decades, an intense research to find the perfect suitable electrolyte for non-

aqueous lithium-air battery has been started. However, the formation of a superoxide ion O2-


- 24 -

that is very reactive over organic compounds, have complicated the task and nowadays no

electrolytes have been found stable under long term cycling.

The theoretical specific energy and energy density are calculated from the weight of active

components of the battery. In non-aqueous Li-air battery, in a fully charged state the active

component corresponds to lithium metal alone, indeed the O2 is coming directly from the

ambient air. In a fully discharged state, the active component is Li2O2 and Li2O56. In Table 1,

specific energy and energy density are calculated if there is a stoichiometric quantity of lithium

and if the cathode is only composed by Li2O2 (no porosity, no binders and no carbon)21.

However, nowadays no practical prototype exists, therefore it is difficult to predict specific

energies or energy densities which will be achievable.

Indeed, some factors influencing the practical energy of a non-aqueous Li-air battery are to

be taken into account. On one hand, there is the excess of lithium necessary in order to

compensate the lithium loss during cycling. Figure 13 presents the evolution of specific energies

and energy densities for different excesses of lithium. On the other hand, the porosity inside the

positive electrode is important to take into account. Bruce et al. have calculated a potential

specific energy and energy density if the positive electrode is composed in volume of 20% of

carbon, 20 % of the electrolyte and 60 % of active material (Li2O2 or Li2O)21. Figure 13 b)

represents the results for Li2O2 in green and Li2O in orange, when a stoichiometric amount of

lithium is taken at the negative electrode (compared to Li-ion specifications).

It appears that compared to Li-ion, this technology stays very attractive.

a)


- 25 -

b)

Figure 13. Specific energies and energy densities of a non aqueous Li-air battery for a) different excess of lithium and b) a porous cathode 21.

Beyond the electrolyte modalities, non-aqueous Li-air battery presents other serious issues

that limits the development of practical prototypes56. Among them, the high charging potential

which limits energy efficiency, the very low discharge capacities at high currents which limits

power performances, the poor Columbic efficiency and safety issues due to the use of lithium

metal and organic liquid electrolytes, and finally parasitic reactions from components in air other

than O2 (water and carbon dioxide). Indeed, the presence of H2O and CO2 leads to the

formation of Li2CO3 and LiOH that precipitate into the pores of the electrode that becomes

clogged. This leads to the necessity of removing those gases by using an O2 selective membrane

or the use of pure O2. Zhang et al. have investigated protection membrane composed of a

hydrophobic polydimethylsiloxane and silicate on a porous metal substrate sheet57,58. They

presented encouraging results for an air-cell cycling in ambient air with 20% relative humidity.

Thus so far, even very promising non-aqueous Li-air batteries suffer from several drawbacks

(positive electrode clogging, dendritic growth, electrolyte stability, use of O2 instead of air, etc...)

that postponed their commercial application for several years.

However, very recently Liu et al.59 have reported a Li-air battery which has addressed several

of the main issues of this technology. They studied a non aqueous Li-air battery using an iodide

redox-mediator (LiI) and cycling via LiOH formation and decomposition (instead of Li2O2 as

discharge product). They have shown that this battery reversibly form and remove crystalline

LiOH with sizes larger than 15 μm and they reported high specific capacities of 93.2% with a

small voltage gap of 0.2 V (instead of 2 V21). In addition, they studied the sensitivity of the cell in

the presence of water inside the electrolyte or cycling under humid O2, and in both cases no

change was observed in the electrochemical profile.


- 26 -

If the latest breakthrough in non-aqueous Li-O2 technology can be turned from a laboratory

demonstrator into a commercial product (still 10 years approximately), it will enable the

widespread of rechargeable batteries with 5 times more energy than the Li-ion batteries.

c. Aqueous lithium-air battery

Most of the current researches on lithium-air battery are focused using organic electrolyte.

However, due to numerous unsolved problems described above, some researchers turned to the

development of the aqueous lithium-air battery. Indeed, some of the non-aqueous Li-air battery

issues can be solved by replacing the organic electrolyte by an aqueous one. This technology is

based on the best of two other technologies, i.e. lithium metal battery for the anode and the

OER/ORR from fuel cells for the positive electrode.

Unlike the aprotic electrolytes,where several electrolytes are eligible, in the aqueous Li-air cell

the choice of electrolyte is limited to only acidic or alkaline solutions. Here, we will focus on the

alkaline electrolytes. However, contrary to the non-aqueous Li-air battery, the positive fuel cell

air electrode is well known and dates back to Grove (1838) and later the NASA's Apollo

program fuel cells. In order to prevent an excessively vigorous reaction between lithium and

water contained in the electrolyte, a watertight lithium ion conducting membrane must protect

the lithium electrode. Therefore, the prerequisite to obtain a practical battery is to develop this

protective layer. In addition, this layer has to be stable in alkaline solution. The pioneer in the use

of a water-stable lithium electrode was Visco et al. in 2004 from Poly Plus Battery Company 60,61.

They used a lithium ion conducting glass-ceramic (LiC-GC), manufacturing by Ohara in Japan,

combined with a LiPON coating to prevent any reaction between the LiC-GC and lithium metal.

This ceramic allows the fast transfer of lithium ions but blocks water and therefore protect the

lithium metal from violent corrosion.

Our work is focused on a prototype of an aqueous Li-air battery developed by EDF and its

partners (Figure 14), where alkaline solutions (LiOH and KOH) are used as an aqueous

electrolyte.

This battery basically comprises a positive compartment, which is composed by the air

electrode, the oxygen evolution electrode and the aqueous electrolyte and a negative

compartment, which is composed by the lithium metal, and a composite separator. Each of the

different components of this battery will be described in the next paragraphs.


- 27 -

Figure 14. Scheme of the aqueous lithium-air battery developed by EDF 62.

Positive compartment. The air electrode is composed of a catalyst supported onto carbon

powder organized into 3D porous structure. It is the interface between the aqueous electrolyte

and the ambient air, therefore it needs to be stable to the basic electrolyte and to be open to

allow the oxygen from air to access the catalysts. A good compromise has to be found in order

to have a sufficient porosity to enable the air to penetrate the positive electrode, but also to have

sufficient mechanical stability and hydrophobicity to contain the electrolyte inside its

compartment and avoid leakage. This compromise has been obtained by the addition of an

hydrophobic agent such as PTFE (polytetrafluoroethylene) above the porous electrode63.

During the discharge, the positive electrode reduces oxygen from the air. Two types of

oxygen reduction reactions can happen; a direct chemical oxygen reduction reaction (ORR) with

four electrons (equation (7)) or an indirect ORR involving two steps : 1) the reduction of O2 into

peroxide (equations (8)), then 2) the reduction of peroxides in hydroxide (equation (9)). The

reaction pathway depends on the catalysts used. For example, gold is well known to favor the

two electrons reaction 64. Whatever the case, the final product corresponds to the formation of

lithium hydroxide (LiOH) 63.

O2 +2H2O + 4 e- → 4 OH- (aq) (7)


- 28 -

O2 + H2O +2 e- → HO2- + OH-

(aq) (8)

HO2- + H2O +2 e- → 3 OH-

(aq) (9)

However, LiOH has a limited solubility in water (5.3 M at room temperature) 62 and when

saturation is reached, it will precipitates as a monohydrate LiOH.H2O (equation (15)). An excess

of lithium hydroxide precipitate can lead to clog the pores of the air electrode, blocking the

oxygen diffusion and therefore limiting the lifetime and performance of the electrode. In fact,

lithium hydroxide also appears to precipitate preferentially onto the ionic conductor glass

ceramic surface on the negative compartment, forming a non-conducting layer. Figure 15 shows

lithium hydroxide crystals precipitated on LiSICON membrane62. Nevertheless, the precipitation

reaction is slow and can be contained on the bottom of the aqueous compartment by adding a

thin layer of cationic polymer on the surface of glass-ceramic62.

Li+(aq) + OH-

(aq) + H2O → LiOH.H2O(s) (10)

Figure 15.Photograph of LiOH crystals precipitated on LiSICON membrane 62.

It is important to notice that the ORR occurs at a triple phase interface. Indeed, the reaction

occurs between a gas (the oxygen from air), a liquid (the electrolyte) and a solid (the electrode).

Thus, the positive electrode needs to be porous to facilitate the O2 diffusion. For the ORR, many

studies have focused on cobalt oxide catalysts made by different processes (sol-gel, plasma, spray

or pulsed laser ablation). However, F. Moureaux et al. showed that in a long term experiment

these types of electrodes are not stable in saturated LiOH solution and present poor

performances and lack of reproducibility 65.

During the charge, the oxygen evolution reaction (OER) corresponds to the oxidation of

hydroxide ions into O2 according to equation (11). Contrary to the non-aqueous Li-air battery,


- 29 -

the reaction in discharge follows the same pathway as in charge. However, the oxygen gas

evolution occurring at the liquid (electrolyte) and porous solid (electrode) interface leads to an

erosion of the air electrode and can cause breakdown of the electrode63. In addition, the OER

can cause the corrosion of the carbon support and the decomposition of the catalysts because of

its high oxidation potential 66. Those two phenomena accelerate the failure of the electrode.

4 OH-(aq) → O2 (g) + 2 H2O + 4 e- (11)

Thus, in order to avoid the deterioration of the air electrode, a second independent electrode

dedicated only to the OER can be added in the aqueous compartment. Since a high surface area

electrode is necessary for the OER, it enables the use of simple and cheap metal grids based on

low cost catalysts66. Recently, an interesting OER electrode made of 316L stainless steel has been

proposed 67. This new type of OER electrode showed a better lifetime and an increase in activity

with time (up to 250 hours) and then reaches a plateau. This is due to the formation of a thin

porous layer composed of nano-crystalline nickel oxide at the surface of the electrode, which

catalyzes the OER67.

At last, the use of non-treated air leads to CO2 penetration inside the alkaline aqueous

electrolyte where it reacts to form carbonates (Li2CO3). This compound has a very low solubility

(13 g.L-1) and precipitates which induces pores clogging and a decrease of the capacity due to a

consumption of lithium ions63. Therefore, the air needs to be treated to remove carbon dioxide,

which is a severe drawback of this technology. However, a composite air electrode with a single

ion, anion conducting polymer electrolyte has been developed by P. Stevens and et al.. in 2010

62,68. The polymer electrolyte conducts OH- ions and enables the ORR to occur in the electrode

but prevents LiOH and Li2CO3 from precipitating into the pores. In addition, this anionic

membrane significantly lows down the diffusion of CO2 through the electrode and limits lithium

carbonate precipitation in the liquid electrolyte63,62.

In an aqueous Li-air battery, the capacity is related to the amount of discharge product,

LiOH, that can reversibly be stored in the cell. At the limit of solubility (5.3 M at room

temperature), this is equivalent to 138 mAh/cm3 or 40 mAh/cm2 for a 3 mm thick electrolyte

compartment. To increase the areal capacity there are two possibilities: the first one is to increase

the thickness of the aqueous compartment and the second one is to increase the amount of the

solid product in the aqueous solution 63. For example, if 80% of the aqueous compartment

volume is occupied by LiOH.H2O, always for a 3 mm thick electrolyte, the positive electrode can


- 30 -

store up to 800 mAh/cm3. Moreover, this high volume fraction of solid discharge product has

no detrimental effect on the performance of the air electrode 69.

Negative electrode. Since the negative compartment is in direct contact with the aqueous

electrolyte, lithium metal needs to be protected by a watertight and stable lithium ion conducting

solid electrolyte. Lithium polymers electrolyte are not watertight, only inorganic solid state

lithium ion conductor have the required properties. The different issues of the negative electrode

are discussed in the Chapter 1.6. on lithium metal.

4. Solid lithium ionic conductor as separator between lithium

and aqueous electrolyte

Many inorganic solid-state lithium ion conductors exist in literature, a short review of these

electrolytes is presented below. The variation of conductivities as a function of temperature of

the principal families of compounds are plotted in Figure 16.

Figure 16. Arrhenius plots for ionic conductivities for selected solid electrolytes 70.

Perovskite -type. Lithium Lanthanum Titanate, also called LLTO, has been widely studied

as a solid electrolyte. Its general formula is Li3xLa(2/3)-x(1/3)-2xTiO3 and exhibits a perovskite

(ABO3) structure70. The best room temperature ionic conductivity is obtained for x = 0.11, with

a conductivity up to 10-3 S.cm-1 for the mono crystal and 5.10-4 S.cm-1 for the ceramic. In

addition, LLTO ceramic is stable in aqueous and LiOH solutions71. However, due to the


- 31 -

presence of titanium IV, which is easily reduced to titanium III, LLTO is not stable versus

lithium71.

Garnet type. Another family of lithium ion conductor with a garnet structure and a general

formula of Li5La3M2O12 (with M=Ta, Nb), was discovered by Thangadurai and Weppner72. Later,

the Li7La3Zr2O12 also called LLZ material was developed and reported by Murugan et al. in

200773. Its high ionic conductivity of 2.4.10-4 S.cm-1 at 25ᵒC for a sintered pellet (92% of the

theoretical density) and its good stability versus lithium metal makes it a good candidate for all

solid-state batteries. However, its stability in aqueous electrolyte is disputed. Shimonishi et al.74

studied the stability of this material immersed in different aqueous electrolytes, such as LiCl

saturated solution and 1M LiOH solution, during one week at 50ᵒC. They reported no changes

in XRD (X ray diffraction) pattern and in electrical conductivity. However, they reported

changes at the surface of the LLZ after immersion in 1M LiOH74. Therefore, the main issue is its

instability in concentrated LiOH solutions.

LiSICON-type. The term LiSICON, which stands for Lithium Super Ionic Conductor, was

employed for the first time in 197875 to describe the Li14ZnGe4O16 phase. LiSICON and related

systems, (Li2+2xZn1-xGeO4) were first described by Bruce and West76. The crystalline structure of

LiSICON is related to the ɣ-Li3PO4 crystal structure. However, it presents a relatively low ionic

conductivity of about 1.10-6 S.cm-1 at room temperature. Moreover, it exhibits a high reactivity in

contact with lithium metal and atmospheric CO2.

Sulfur type. Thio-LiSICON-type lithium conductor was introduced by Kanno et al.77. They

replaced in the LiSICON structure oxide ions by sulfide ion which are more polarizable and

bigger, in order to have a better ionic conductivity. However, the ionic conductivity of Li4-

2xZnxGeS4 was not improved, e.g. the compound Li4GeS4 presents a low ionic conductivity at

25ᵒC of 2.10-7 S.cm-1. On the contrary, the introduction of lithium vacancies via phosphor

addition (P5+) leading to the Li4-xGe1-xPxS4 compound presents a high ionic conductivity at 25ᵒC

of 2.2 .10-3 S.cm-1. Recently, Kamaya et al.78 reported the Li10GeP2S12 phase, which presents a

mono crystalline structure and an ionic conductivity of 1.2.10-2 S.cm-1 at 25ᵒC. However, it is

hygroscopic and has a high cost due to the presence of germanium which might restrain its

utilization.

The thio-LiSICON's are promising because they present a good ionic conductivity at 25ᵒC

and they are found to be stable in contact with lithium78. Unfortunately for the aqueous lithium-

air technology, they are unstable in contact with water and lithium hydroxide79.


- 32 -

NaSICON type. NaSICON-type ceramics (Na+ Super Ionic Conductor) exhibit an

AM2IV(PO4)3 structure (with A = Li+, Na+ and M = Zn, Ti, Si, Ge), which are formed by

M2(PO4)3 units, where two octahedra MO6 are connected at the top to three PO4 tetrahedra.

Figure 17 shows a schematic illustration of the NaSICON framework.

Figure 17. Schematic illustration of NASICON (generally rhombohedral) and framework of general formula AxMM'(XO4)3

80.

A good sodium ionic conductivity is reported for Na1+xZr2SixP3-xO12 phase80, (especially for x

= 2). Later, this structure is also considered for lithium conduction and solid electrolytes based

on lithium titanium phosphate as it was first reported by Aono et al.81 in 1989. However, the

Li1+xZr2SixP3-xO12 phase exhibits a very low ionic conductivity. The substitution of Zr4+ by Ti4+ in

the NaSICON structure leads to an increase of the ionic conductivity. For example, Li3Ti2(PO4)3,

also called LTP, exhibits an ionic conductivity of 10-6 S.cm-1 at room temperature82. Higher ionic

conductivity is reported for compounds where several Ti4+ ions are replaced by Al3+ and Li+. Xu

et al.83 reported for the Li1,4Al0,4Ti1,6(PO4)3 compound also called LATP, an ionic conductivity up

to 1.10-3 S.cm-1 at 25ᵒC for ceramics with a density greater than 97%. The best ionic conductivity

is reported for a ceramic composed of Li1,5Al0,5Ge1,5(PO4)384, which exhibits at 27ᵒC a

conductivity of 4.22.10-3 S.cm-1.

It is worth noting that Ohara Corporation85 has commercialized a Li ion conducting

NaSICON-type ceramic which is widely used nowadays. The glass-ceramic is obtained after

sintering Li2O-Al2O3-SiO2-P2O5-TiO2-GeO2. The exact formula is proprietary. They obtained an

ionic conductivity of 1.10-4 S.cm-1 at room temperature for a sintered pellet86. The crystal

structure of the commercial ceramic is presented in Figure 18.


- 33 -

The main advantage of NaSICON-type ceramic is its stability in air and water, but due to the

presence of titanium these ceramics are not stable versus lithium metal. Furthermore, their

stability in LiOH are controversial and still discussed87–89.

Figure 18. Li conductor with a NaSICON structure crystal structure and vacancy diffusion for the Ohara commercial ceramic85.

Table 2 presents the principal properties of the main solid electrolytes, which were discussed

above. None of the high ionic conductors are both stable versus lithium and in LiOH solution.

Therefore, the ceramic used needs a protection.

Type NaSICON Perovskite

Li0.5La0.5TiO3

Garnet

Li7La3Zr2O12

Thio-LiSICON

σ (S.cm-1)

at RT 4 . 10-3 81 10-3 70 8 . 10-4 90 10-2 78

Ea (eV) 0,35 81 0,3 70 0,3 90 0,25 78

Stability in H2O Yes 74 Yes 71 Yes 74 No 79

Stability in LiOH Discussed 87–89 Yes 71 No 74 No 79

Stability vs Li No 91 No 71 Discussed 74 Yes 78

Table 2. Characteristic of the main solid electrolytes discussed above.

5. Possible ways to protect the ceramic

The protective buffer layer is the interface between lithium metal and the ceramic, thus it has

to have a good adhesion and be stable with both materials. Ideally, the protective layer for the

ceramic should have a high lithium ionic conductivity at room temperature. Moreover, the layer

needs to have high resistance to dendritic growth and should preferably be a single-ion


- 34 -

conductor to avoid concentration gradients during cycling. In the section hereafter, we will focus

on potential buffer layers to protect the ceramic to the contact with lithium, based on i) organic

liquid electrolytes, ii) inorganic solid electrolyte and finally iii) polymer electrolyte.

a. Liquid electrolytes

Organic liquid electrolytes for lithium-ion batteries have been extensively studied for the past

four decades, an excellent review written by Xu92 details the advances in liquid electrolytes. One

possibility as a protective layer for the ceramic is the use of a liquid electrolyte (that provides

good conductivity) impregnated into a porous separator (that provides mechanical properties)93.

Table 3 presents a classical list of organic solvent for liquid electrolyte, with their properties.

Their main advantage is their high ionic conductivity (>10-3 S.cm-1). However they are not

stable versus lithium metal electrode and generally the lithium growth is heterogeneous. It would

therefore be better to favor solid electrolytes.

Table 3. Organic electrolyte solvents 92.

b. Lithium phosphorus oxynitride or LiPON

Amorphous thin film lithium-ion conductors are first investigated for the purpose of an

electrolyte in thin-film rechargeable battery. Lithium phosphorus oxynitride or LiPON is an

amorphous alkali phosphate thin-film glass found in 1992 by Bates and et al. at the Oak Ridge


- 35 -

National Lab94,95. Its composition can be represented by xLi2O:yP2O3:zPON, where PON is

phosphorus oxinitride. This material is amorphous. It exhibits a conductivity up to 3.10-6 S.cm-1

at 25ᵒC. One decisive advantage of LiPON is that it can be sputtered in thin layers. Typical

thicknesses for LiPON used in an electrochemical device are between 1 and 2 µm. However,

LiPON thickness can be decreased down to 12 nm when using ion beam sputtering96 of a

lithium phosphate target (Li3PO4) under pure nitrogen atmosphere (N2). Some efforts were made

in order to increase the ionic conductivity of LiPON by optimizing the composition by varying

especially the partial pressure of nitrogen97. Another important advantage of this material is its

stability towards lithium and up to 5.5 V vs Li+/Li 95. Yu and et al. determined that LiPON

exhibits a single Li+ ion conductor between -26ᵒC and 140ᵒC 98. They also cycled this material

more than 40 000 times at 25ᵒC with lithium electrode and LiCoO2 cathode. They found that

LiPON is mechanically stable and is acting as a rigid barrier against lithium dendrites growth.

However, the LiPON is not stable against moistures present in the air. Nimisha et al. has

investigated the stability of LiPON in air. They showed morphological and chemical changes on

the LiPON surface. Firstly, the LiPON smooth surface turned into a rough surface after 24

hours exposition, and secondly the ionic conductivity decreased to 9.9 10-10 S.cm-1. This is due to

instability with H2O, O2 and CO299. Recently, Schwobel et al. studied by X-Ray photoemission

the interface between lithium and LiPON and reported evidence for different chemical reactions

at the surface of the LiPON in contact with lithium metal 100. Those reactions lead to the

decomposition of LiPON into smaller units such as Li3PO4, Li3P, Li3N, Li2O. However, due to

the high cyclability of LiPON thin films reported98, they concluded that these interface reactions

lead to the formation of a passivation layer (SEI) thin enough to permit a good lithium ion

conduction100.

Mechanical properties of the LiPON were studied by nano-indentation by Herbert et al.101.

They showed that the shear modulus of LiPON is 77 GPa and is independent of film thickness,

substrate and annealing. This result suggests that the LiPON is expected to suppress dendrite

formation at the lithium/LiPON interface, due to a shear modulus 7.3 times larger than lithium.

Indeed, the Monroe and Newman theory102,103 suggested that if the electrolyte presents a shear

modulus about twice that of lithium, dendrites growth should be suppressed.

The main properties of LiPON material are synthesized in Table 4.


- 36 -

LiPON

Conductivity at 25ᵒC Up to 3.10-6 S.cm-1

Shear modulus 77 GPa

Stability vs Li+/Li 5.5 V

Table 4. Main properties of LiPON.

In the aqueous Li-air battery, a thin and flexible current collector (CC) is sputtered onto

LiPON surface and lithium can be electrochemically deposited between the CC and LiPON62. A

majority of studies have been reported on copper current collectors (Cu CC). However,

inhomogeneous electro-deposition of lithium is frequent with this CC on LiPON layer104,105,106.

Nevertheless, dense electrodeposit of lithium has been reported by Stevens et al.62

To conclude, LiPON is an interesting protective layer since it enables a dense and uniform

lithium electrode to be produced. However, despite its ability to be sputtered as a thin film (<

1µm), its conductivity is rather low at room temperature (~ 10-6 S.cm-1) which leads to high

polarization.

c. PEO polymer based electrolytes

The discovery of the ionic conductivity of PEO complexes with alkali metal salt by Wright107

in 1973 marked the beginning of the research on PEO complexes. Indeed, he was the first to

show that ether-based polymers can dissolve inorganic salt and exhibit ion conduction at room

temperature. However, it is only in 1978 that M. Armand30 suggested PEO mixed with lithium

salt as a material of interest for the development of electrochemical devices. It is the start of the

solid polymer electrolyte (SPE) story. Solid polymer electrolytes present excellent properties,

such as mechanical strength, flexibility, lower reactivity with lithium metal than liquid electrolyte

and improved safety because they are free of organic and flammable solvents. One of the main

challenge in SPE is to develop a material which exhibits high ionic conductivity at room

temperature with a Li+ transference number equal to one (single-ion), good mechanical

properties, good interfacial properties with electrodes and finally good electrochemical stability at

high potentials. However, the major drawback of classical polymer electrolytes, such as PEO, is

that the mechanical strength and conductivity vary in opposite directions.

PEO based polymer electrolytes are the most extensively researched polymer electrolytes.

Indeed, ethylene units (EO) inside PEO exhibit a high donor number for Li+ that produces its

complexation and the salt dissociation, as well as a high chain flexibility that promotes the ion


- 37 -

transport inside. Moreover, due to its relatively high dielectric constant (ε = 8 in the amorphous

phase) and strong Li+ solvation, PEO is able to dissolve lithium salts108. However, PEO is a

semi-crystalline polymer (see Figure 19) which presents a glass transition temperature, Tg, at

around - 60ᵒC, and a melting point, Tm, at 65ᵒC for high molecular weight polymer. PEO-

lithium salt complexes usually present an acceptable ionic conductivity of 10-4 S.cm-1 only above

the melting temperature 60ᵒC. Below the Tm, the conductivity drops drastically because the ion

conduction in PEO and similar polyether based media occurs in the amorphous phases of the

polymer matrix 109,110.

Figure 19. Morphologies of semi-crystalline PEO 111 and schematic representation of ion conduction in amorphous and crystalline phase in PEO.

Lithium ions are 4 or 5 dentate coordinated by the ether oxygen within the PEO matrix, in a

similar way that to complexation by crown ether based solvents112. Figure 20 presents the

mechanism of ion transport in the polymer matrix. The ion transport occurs via an oxygen-

assisted hopping mechanism. Lithium-oxygen (Li-O) bonds are forming/breaking, which result

in an intrachain or an interchain hopping mechanism113,114. Both mechanisms present a

continuous segmental rearrangement with the gradual replacement of the counter anion92. Those

mechanisms suggest a long range net displacement of lithium ions with a long range segmental

motion of the chains 108. Therefore, the lithium ions conduction depends on chain mobility, high

chain mobility results in good ionic conductivity but poor mechanical properties.


- 38 -

Figure 20. Sketch depicting the three different cation transport mechanisms in PEO-salt electrolytes. Each mechanism is characterized by specific time scale115.

Usually, for SPE based on PEO, lithium salts are designed such as the anion presents a

highly delocalized negative charge in order to facilitate the salt dissociation and solvation. A wide

variety of lithium salts have been reported, typical lithium salts used for PEO based electrolytes

are listed with their structures and properties in Table 5.

Salt name Lithium salt

(abbreviation)

Anion structure Main characteristics Ref

Lithium perchlorate LiClO4

- Broad electrochemical stability window 116

Lithium

terafluoroborate

LiBF4

- Broad electrochemical stability window 117

Lithium

hexafluorophosphate

LiPF6

- High ionic conductivity

- Decomposes in the presence of moisture to form

HF

117,118

Lithium

bis(trifluoromethanesul

foimidate)

LiTFSI

- High solubility and high ionic conductivity

- High electrochemical stability

119

Lithium

bis(fluorosulfonyl)

imide

LiFSI

- Higher ionic conductivity compared to LiTFSI

- High electrochemical stability

120

Lithium bis(oxalato)

borate

LiBOB

- High electrochemical stability and long term

stability

- Form a highly resistive SEI films (low

conductivity compared to LiPF6 and LiTFSI)

121

Table 5. Structure and properties of commonly used lithium salts for studies on polymer electrolyte 122.


- 39 -

Lascaud et al.123,124 and Vallee et al.125 have studied properties of such PEO-salt complexes,

they studied PEO mixed with different lithium salts, such as LiClO4, LiTFSI and LiFSI and

determined a phase diagram for PEO-LiTFSI complexes. They demonstrated the existence of

three defined compounds for EO/Li ratio equal to 2, 3 and 6, when a cristallinity breach exists

between EO/Li equals to 8 to 12. In addition, they showed that complexes with EO/Li >12 are

semi crystalline 125. Besides they studied the influence of salt concentration on glass transition

temperature and showed that the Tg increases with the introduction of lithium salt due to

electrostatic interactions that limit the chain mobility. They noticed that for LiTFSI and LiFSI,

the evolution of Tg is different to that with LiClO4 showing that large and "flexible" anion acts

as a plasticizer125. Since Tg is related to the segmental motion of the macromolecular chain,

lowering the Tg will promote the conductivity at a given temperature.

Different approaches to increase the ionic conductivity at low temperature have been

proposed. One of them consists in decreasing the degree of crystallization, ɣc, or reducing the

polymer melting temperature, Tm, in order to promote amorphous phases. For example, the use

of PEO have been largely studied, because small branched PEO crystallize only at low

temperatures. However, the mechanical properties are very weak.

Another possibility is the addition of a plasticizer which will reduce the crystallization of

PEO which improved the salt solvation and molecular dynamic that in returns increase the ionic

conductivity. In the literature, molecules such as succinonitrile, glymes, carbonates etc have been

studied126.

To increase the mobility of lithium-ion, it is necessary to weak the lithium-ion/polymer chain

interactions. One solution is to add room temperature ionic liquid (RIL), such as 1-ethyl-3-

methylimidazolium127 or N-methyl-N-butylpyrrolidinium bis (trifluoromethansulfonyl) imide128.

Linear copolymers were first reported in 1984 by Watanabe et al.129 They synthesized the

block copolymer electrolyte (BCE) poly(dimethyl siloxane-co-ethylene oxide) (PDMS-PEO) in

order to obtain higher ionic conductivity compared to PEO-complexes. Later, Fonseca et al.130

studied PDMS-PEO BCE loaded with different concentration of LiClO4 salt. A maximum of

conductivity is reported for the polymer-complex with 5 wt% salt and exhibits an ionic

conductivity of 2.6.10-4 S.cm-1 at room temperature.

We have presented a non-exhaustive list of PEO based polymer electrolytes, however

progress in this area has been summarized in several recent reviews111,131.


- 40 -

d. Block copolymer electrolytes

Block copolymers have been extensively studied in material science thanks to their unique

properties resulting from a wide variety of chemical functionalities. Recently, they have raised a

new interest for the design of solid polymer electrolytes. Indeed, in spite of numerous studies

dedicated to the suppression or the reduction of PEO crystallization, which improves the ionic

conductivities, the resulting amorphous polymers presented usually very poor mechanical

properties. Block copolymers electrolyte (BCE) can be a possible alternative to combine good

mechanical properties with high ionic conductivities in the same material, one block providing

mechanical properties and the other one allowing lithium ion conduction (generally based on

PEO).

Phase separation. Block copolymers are composed of two or more chemically distinct

polymer blocks covalently bound together. When mixing two different polymers, they tend to

separate at a macroscopic scale (like oil and water emulsions) with no synergy of their properties.

However, in the case of BCE, polymer blocks are covalently bounded and they cannot separate

at such a scale. Nevertheless, in order to minimize unfavorable contacts between the different

blocks, polymer chains will tend to extend perpendicularly to a reduce interface, where the

junction points (covalent bond) will be localized132. This phenomenon leads to a phase separation

that results in self-assembly properties of the BCE and form structures at a nanoscale. In bulk, a

wide variety of morphologies have been observed including spherical micro domains (S),

hexagonally ordered cylinders (C), lamellar (L), gyroid morphologies (G) depending on the

volume fraction of the blocks (Figure 21 b)) 133. Two parameters dictate the relative immiscibility

of the blocks: an enthalpic factor, the temperature dependent Flory-Huggins factor (χ), which

reflects the interaction energy between the different blocks, and an entropic factor, the

polymerization factor (N). When the product (χ . N) is superior to a certain critic value (χ .

N)ODT, which corresponds to the order to disorder transition, the block copolymer performs a

microphase separation. The balance between the interactions of monomers of the same nature

are optimized, while the interface between the dissimilar monomers are minimized. This balance

is predominantly governed by the composition of the block copolymer and especially the volume

fraction of the respective blocks fa or fb. The different morphologies can be mapped in a

theoretical phase diagram which enables morphologies as a function of χN and fa parameters to

be predicted. Figure 21 a) presents a theoretical phase diagram for a symmetrical triblock

copolymer ABA 133 and the different associated morphologies in Figure 21 b). The challenge is to

direct the self-assembly of the BCE at a long range.


- 41 -

a) b)

Figure 21.a) Theoretical phase diagram for a triblock copolymer ABA (interaction parameter as a function of the A block volume fraction) b) Block copolymer morphologies 133.

Block copolymer electrolyte. Block copolymer electrolytes (BCE) was first reported by

Giles134 in 1987. He synthesized the poly(styrene-block-butadiene-graft PEG-block-styrene) (PS-b-

PB-graft PEG-b-PS) triblock copolymer with poly(ethylene glycol) grafted to PB block. By

covalently bonding the different blocks, the BCE exhibits two different functionalities, a PS

block conferring the mechanical properties and a lithium-ion conducting PB-g-PEG block.

However, ionic conductivities are below 7.10-6 S.cm-1 at room temperature. Mechanical

reinforcement of BCE is studied in similar systems, where the PEG is grafted onto the central

block such as PS: Poly(styrene-b-(styrene-g-oligoethylene glycol)-b-styrene)135. They confirmed by

transmission electron microscopy (TEM), that this BCE is nanostructured with PS cylinders

hexagonally arranged in a PS-g-PEG matrix. An ionic conductivity of 10-5 S.cm-1 is obtained at

room temperature.

Recently, poly(styrene)-block-Poly(ethylene oxide), (PS-b-PEO also called SEO), doped with

LiTFSI salt have been extensively studied by the Balsara's group and it is a good candidate for

high performances solid electrolytes. Singh et al.136 have shown that SEO BCE presents a lamellar

morphology for a volume fraction ranging from 0.38 to 0.55 and PEO block molecular weight,

MPEO, ranging from 16 to 98 kg.mol-1. Panday et al.137 reported that the glassy PS lamellae provide

good mechanical integrity and that the PEO lamellae ensure the good ion conduction. However,

due to PEO crystallization fast ion conduction cannot be expected at room temperature, thus,

much researches have been performed above PEO melting temperature. Nevertheless, this BCE

exhibits excellent mechanical properties even at high temperature (90ᵒC) and a shear modulus


- 42 -

around 108 Pa for high molecular weight137. However, ionic conductivity are low with 10-4 S.cm-1

reached above 90ᵒC.

As discussed previously, the maximum of ionic conductivity in bulk homopolymer can be

reached by the right combination of polymer/salt. However, in nanostructured BCE, the ionic

conductivity behavior is more complex and relies on different parameters: the volume fraction of

the conductive phase, but also on its morphology, as well as on the interface properties between

the different blocks. Surprisingly, Singh et al. have showed that the ionic conductivity in PEO

based BCE increases with increasing PEO molecular weight, which is counter intuitive. A study

by Gomez et al.138 showed that the lithium distribution in the lamellae of SEO BCE by energy-

filtered transmission electron microscopy (EF-TEM). Figure 22 presents the EF-TEM image of

the nanostructured PS-b-PEO BCE. It is clearly shown that lithium-ions are segregated in the

middle of the PEO domain, due to the PEO chains stretched at the interface PEO/PS.

Figure 22. LiTFSI distribution in nanostructured PS-b-PEO electrolyte: EF-TEM picture138.

More recently, Bouchet et al.139,140 demonstrated, by thermodynamic analysis of the melting of

confined PEO in triblock PS-b-PEO-b-PS nanostructured BCE, as well as by conductivity

analysis, that it exists at the interface PS/PEO a zone called "dead zone" where both ion

concentrations and PEO chain mobility are strongly affected (Figure 23). In addition, they

estimated that the thickness of the dead zone represent 4-5 EO units (1.6 nm) and it is not

dependent on either the PEO molecular weight or on the EO:Li ratio. They supposed that the

absence of conduction and crystallization in the dead zone can be explained by the low mobility

of PEO chains in this region. Therefore, the presence of a dead zone can explain why the ionic

conductivity increases with MPEO. In other words, when MPEO increases the proportion of the

dead zone decreases.


- 43 -

Figure 23. Schematic representation of ion segregation in nanostructured PS-b-PEO-b-PS triblock copolymer electrolyte 139.

Solid polymer electrolyte have been extensively studied and ionic conductivity up to 10-5

S.cm-1 at room temperature could be reached. However, the ionic conductivity is not the only

limiting factor for a good electrolyte. During the redox process only lithium ions are involved at

the electrodes leading to an accumulation of their counter anions at the electrode interfaces. This

is due to the low transference number (tLi+) of such electrolytes. A strong concentration gradient

is then formed at the interface electrode with deleterious effects such as promoted dendritic

growth141 and limited power delivery142. Therefore, the design of single-ion conductor, with a

transference number equal to unity should prevent those limiting effects.

e. Single-ion electrolytes

In order to combine the high ionic conductivity with a Li+ transference number of 1, several

strategies have been employed142–144. In the middle of the 90s, Benrabah et al.144 have synthesized

polyanionic salts in order to blend them with PEO and then cross-linked them.

Watanabe et al.145 reported the synthesis of poly (2-oxo-1-difuluoroethylene sulfonylimide)

(LiPI). They reported a transference number of unity. However, the single-ion electrolyte

exhibits a low ionic conductivity of 10-8 S.cm-1 at room temperature and 10-6 S.cm-1 at 110ᵒC.

Another approach is to functionalize chain ends of PEO. For example lithium sulfonate and

lithium acetate have been chemically anchored to oxide ethylene glycol. At room temperature,

ionic conductivities reported ranged between 10-9 S.cm-1 to 5.10-7 S.cm-1. To improve the lithium

transport, the addition of PEG liquid oligomers plasticizes the polymer electrolytes, and

conductivities up to 5.10-5 S.cm-1 at room temperature have been reported. However, mechanical

properties of such electrolytes are very weak146.


- 44 -

Single-ion block copolymers have also been studied as a promising solid electrolyte.

Sadoway's group147,148 synthesized different BCE with PEGMA as the conductive block and

poly(lauryl metacrylate) block for the mechanical properties. A lithium metracrylate was

incorporated as the single-ion function, inside the PEGMA blocks and outside the PEGMA

blocks. These single-ion BCE exhibit ionic conductivity below 10-7 S.cm-1 at room temperature

and up to 10-5S.cm-1 at 70ᵒC.

Recently, Bouchet et al.142 reported a new single-ion conductor triblock copolymer P(STFSI-

Li)-b-PEO-b-P(STFSI-Li) exhibiting high performances (see Figure 24 a) for the BCE structure.

LiTFSI anion is covalently bonded to the styrene moieties in PS blocks. Therefore, only the Li+

are mobile. The conductivity is provided by the PEO block, where mechanical properties are

given by the PS block. The TFSI anion is grafted to PS blocks and confers to the electrolyte a

transference number close to 1 (>0.85). The ionic conductivity reaches 1.3.10-5 S.cm-1 at 60ᵒC

(see Figure 24 b)). In addition, they reported mechanical strengths of 10 MPa at 40ᵒC and an

enlarged electrochemical stability window up to 5V versus Li+/Li.

Figure 24. P(STFSILi)-b-PEO-b-P(STFSILi) single ion electrolyte: (a) structure, (b) Arrhenius plot of the conductivities142.

Later Balsara's group149,150 studied similar single-ion conductor, but in diblock copolymers

PEO-b-PSTFSI, and they reported a lamellar morphology presented in Figure 25.


- 45 -

Figure 25. Dark field scanning electron micrograph of PEO-b-PSTFSI. The bright phase represent PEO rich lamellas149.

6. Lithium metal

After have introduce the different possible materials for the protective buffer between the

Ohara-GC and the lithium metal, it is important to focus on the last part of the negative

electrode, i.e, the lithium metal.

a. The Solid Electrolyte Interphase (SEI)

Native layer. Due to its high reactivity a passive layer called "native", is instantly formed on

the surface of lithium metal (Figure 26), when it is in contact with N2, O2, H2O, CO2 or CO.

Analysis by X-ray photoelectron spectroscopy (XPS)151 and by spectroscopy such as infra-red

and Raman152, have shown that this layer is usually composed of two layers, one merely

composed of Li2O at the surface of the lithium (with a thickness between 10 to 100nm) covered

by a thinner layer composed of Li2CO3/LiOH.

Figure 26. Schematic illustration of the native passive layers at the surface of lithium metal 151.

Formation of a SEI in contact with the electrolyte. When lithium metal is in contact with

a liquid organic electrolyte, an evolution in the native layer is observed depending of the nature

of the lithium salt and of the solvent molecules as well as the impurities153. The so called Solid


- 46 -

Electrolyte Interphase (SEI) by Peled153 forms a solid interface between the electrolyte and the

lithium metal, which presents its own chemical and physical properties. In addition, this layer

plays a protective role against further corrosion. To be efficient this passive layer should be an

electronic insulator and Li+ ion conductor. However it generally presents a higher resistivity than

the electrolyte154. The SEI strongly influences the cycling performance155, especially due to its

irregular morphology and composition, that will favor both dendritic nucleation and growth

during lithium electrodeposition156. The SEI also plays a role on the Faradic efficiency due to

another difficulty with lithium, i.e. its high reactivity with the electrolyte. Thus each times a fresh

lithium is formed or when the SEI is broken, it forms new layers that consumes both lithium and

electrolyte which leads to a low columbic efficiency.

Figure 27. SEI model157. (a) Schematic view; (b) equivalent circuit; (c) electrode impedance spectrum.

The SEI model according to Thevenin157 is presented in Figure 27, it assumes that all the

lithium metal is covered by this passive film. In order to be reduced on the surface of lithium

metal, lithium ions must follow a path in three steps; the lithium ions are firstly desolvated from

their solvation sphere in the electrolyte to be then transferred into the SEI. Then the Li+ ions

moves through the SEI via migration, and finally, they are reduced at the lithium surface at the

interface Li/SEI. Due to the ionic conductivity and typical thickness of the SEI, the kinetically

determining step corresponds to the migration of Li+ through the SEI, which depends of the

applied potential E157,158.

The SEI can be modeled by an equivalent circuit composed of a resistance RSEI in parallel

with a pseudo-capacitance CSEI (see Figure 27 b)) and the impedance spectrum in Nyquist

coordinates is a semi-circle (see Figure 27 c)).

Polymer electrolyte interphase157. In the case of a polymer electrolyte, the quality of the

contact between lithium metal and the electrolyte is a key part. Indeed, contrary to the liquid

electrolytes, which are able to wet entirely the surface of the lithium electrode, the adhesion of

polymer electrolyte can be heterogeneous. The passive layer123, schematized in Figure 28, can be

discontinuous and the active surface can be only a portion of the entire surface.


- 47 -

Figure 28. Schematic presentation of the Li/PE interphase: (A) native oxide film; (B) freshly formed SEI; (C)

void; (D) pe (or cpe); L : void height 159.

b. Model of nucleation and growth of lithium dendrites

In the field of electro-deposition of metal, "dendrites" are a well-known and a common

phenomenon. Many metals such as zinc, copper, silver, lithium, etc. are reported to exhibit

different electro-deposit morphologies including fractal-like, needle-like, tree-like, bush-like,

moss-like, snowflake-like and whiskers-like depending on the electrolyte and the current

densities. Here, we will focus on the mechanisms of lithium dendrites nucleation and growth

described in the literature.

Figure 29. Main issues related to lithium dendrites nucleation and growth160.

Heterogeneous growth of lithium was first hypothesized in 197425, and directly observed in

1980161 in a lithium/organic systems. The presence of Li-dendrites at the anode leads to many

serious problems such as low energy density, safety hazards and short cycle life (Figure 29).


- 48 -

In the last forty years, many groups, both industrial and academics, have been working on

lithium dendrite nucleation and growth processes. They have proposed fundamentals models

which will be introduced in this section.

DLA model. The first model proposed to describe growth of metallic clusters via

electrochemical way is the Diffusion-Limited Aggregation, DLA, based on Witten and Sander

model162. This simple computer model consists on particles diffusing onto a cluster by random

walk one at a time. Figure 30 shows a numerical computation for dendritic growth of 3600

particles in a square lattice, the clusters grow immediately dendritic and disorderly, moreover,

they were shown to present fractal dimension163.

.

Figure 30. Random aggregates of 3600 particles on a square lattice162.

A few years later, Brady and Ball164 showed that the electrolytic deposits of copper in

diffusion-limited conditions exhibits a fractal shape and that the Hausdorff dimensionality was in

a good agreement with the simple computer model of DLA, they found D =2.4.

In this model, ions migration was considered negligible compared to diffusion, however, this

is only valid with a small electric field. In addition, the DLA model found another limitation due

to the multitude of other morphologies than fractals obtained by lithium electro-deposition such

as needle like.

Space charge model or Chazalviel's model141. It is important to notice that this model

proposes to explain the nucleation of lithium dendrites. The proposed model is an electrostatic

model based on the creation of a positive space charge upon anion depletion in the vicinity of

the negative electrode when the current density is higher that the diffusion limited current.

By solving the equations governing the potential variation ions motion in a binary electrolyte,

Chazalviel calculated both concentration profiles of ionic species in the whole cell. Those

profiles are presented in Figure 31, the evolution of ions concentration profiles as well as the


- 49 -

potential profile across the cell according to time. In Figure 31, results obtained at the steady

state condition are given.

b)

Figure 31. a) Profile of the ion concentration Cc and Ca, and of the electrostatic potential V resulting from the numerical simulation in the hypothetical case of uniform deposition with negligible growth of the cathode. L =1

crn, Vo=1 V, Dc =Da =10-5 cm2/s, zc =za =1, ϵ=80; (a) C0=1010 cm-3 and b) Profile of the ion concentrations and electrostatic potential as a function of time; t =100 s (solid lines), t =10' s (dashed lines), t

=10 s (dotted lines). Notice the motion of the anion distribution, due to the drift in the applied electric field, and the associated rise of the space charge near the cathode.

The current density when τS is approximately equal to τd (the diffusion time) is called the

critic current density, J* and is defined according to equation (12), with C 0 the ionic

concentration, D the ambipolar diffusion coefficient, e the elementary charge, L the inter-

electrode distance, ta the anion transference number.

! "# $%&°'( " )

*+,,,,,, (12)

As soon as a current above the diffusion limited current, J*, is imposed, the vicinity of the

cathode (x=0) is depleted to form a metallic deposit and anions migrate to the anode. After a

time close to the Sand time (τs) 165, i.e. when ionic concentrations are close to zero at the cathode

and J > J*, Chazalviel showed that an area of few microns devoid of anions with an excess of

cations (zc.Cc >> za.Ca), is created. This area is called the space-charge layer and is positively

charged. Thus, two regions can be distinguished in the electrolyte, a neutral region (zone I) and a

space charge region (zone II) in the vicinity of the negative electrode. This space charge induces

a very large electric field (up to 10 kV.cm-1) in the vicinity of the cathode and therefore an abrupt


- 50 -

drop in potential localized in this region (see Figure 31) that is the main driving force for the

nucleation of dendrites.

Therefore he postulated that dendrite starts to grow when the electric field reach a critical

value, typically reached at the Sand time (see equation (13), with µa and µc anion and cation

mobility respectively, D the ambipolar diffusion coefficient, J the current density).

-s = .' " /01"%"&°1$! 2 ² /3+4313+ 2$,,,,,,, (13)

The space charge region propagates into the electrolyte with the speed of the anions

migration, thus the velocity of the front of the deposit is determined by the velocity of the

anions according to equation (14), with E the electric field applied to the cell.

5 = 6a7 8,,,,,, (14)

Dendrites have been observed by Rosso's group166,167,168. Brissot et al.169 used three

independent in situ and ex situ methods in Li/PEO-LiTFSI/Li cells and demonstrated a direct

relationship between dendritic growth and concentration gradient and especially that the onset of

dendritic growth matched very well the Sand's time.

When the current is inferior to the diffusion limited current, in principle there should be no

Sand behavior, and therefore no dendrite formation. However, experimentally dendrites

formation are observed. Rosso et al.170 and Teyssot et al.171 attributed the formation of dendrites

to the local non-uniformity of the Li/electrolyte interface, which leads to a large concentration

variations even in the depleted zone close to the conditions of Chazalviel's model.

Transverse Chazalviel. This model was proposed by Teyssot et al.171 and involved the

development of concentration instabilities due to heterogeneities at the lithium/electrolyte

interface. This model takes into account the effect of the irregular passivation surface and the

small inter-electrode distance. They reported a "transverse Sand/Chazalviel" behavior due to

variations in the local current density at the lithium/electrolyte interfaces leading to

concentration gradients along the electrodes. This model is presented in Figure 32. The

important parameter now is ∆J instead of J which explains the Sand behavior observed

experimentally even below J*.


- 51 -

Figure 32. Development of concentration instabilities in a cell where (i) the passivation layer is very non-uniform, (ii) the inter-electrode distance l is very small compared to the lateral dimension L. The cell may be regarded as the parallel arrangement of small box units with size l. In box 1, the current density at the anode is smaller than the

current density at the cathode: consequently, the concentration decreases in this box. On the opposite, the concentration increases in box 2. Transversal diffusion will eventually limit these local concentration fluctuations171.

Monroe and Newman model. A model to describe the dendritic growth and electrode

roughening due to local variation of current densities has been proposed by Monroe and

Newman102,103,172. This model is based on different previous work based on the Mullins-Sekerka

linear stability analysis173 and the Barton and Bockris dendrite-propagation model174. Two models

have been developed, one for liquid electrolyte and the second one for polymer electrolyte. The

first model172 is based on a kinetic equation that is driven by the concentration profile and the

potential and valid in the case of dendritic growth under galvanostatic conditions applicable to

liquid electrolyte and they determined an expression for the tip growth rate (equation (15)),

where Jn is the effective current density normal to the dendrite (hemispherical) tip, V is the molar

volume of Li and F is the Faraday's constant.

(15)

Dendrite growth profile at different current densities (expressed as a fraction of the limiting

current density, iL) have been calculated (see Figure 33 a)). They revealed that dendrite growth

accelerates with time and as dendrites move across the cell. In addition, lowering the current

density prolongs the linear behavior (corresponding to a slope of 1) of the growth regime.

They also demonstrated that a combination of a high diffusion coefficient (D) and high

transference number maximizes the charge passed before failure (see in Figure 33 b)).


- 52 -

a) b)

Figure 33. a) Dendrite growth profile at varying currents and b) Effect on changing D and t+ on charge passed per surface area172.

Another model102,103 for the case of polymer electrolytes has been developed. Monroe and

Newman models are based on kinetics, where the surface tension forces play a preponderant role

in local current densities. In other words, if the surface tension at the negative electrode is small,

this will favors the roughening of the interface and enables the growth of dendrites. The

originality of this model is the addition in the kinetic model of mechanical forces such as

elasticity, viscous drag and pressure, which have an effect on local exchange current densities and

potentials at the interface, that favors high flux and roughening of the interfaces. Equations (16)

represents the kinetic relationship which includes the effects of a general deformation (∆μ e-α,α')

and local deviations from macroscopic mass transfer (cβMz+).

(16)

The general deformation equation is presented in equation (17) and describes the changes in

electrochemical potential caused by mechanical forces around a isothermal roughening interface,

this equation takes into account the surface tension (2γH), elastic and plastic deformation

(∆τdα,β), viscous response of the bulk phase (∆v

α,β) and finally externally applied pressure on the

electrode or the electrolyte (∆pα,α' + ∆p

β,β').

(17)


- 53 -

Consequently, nucleation and growth of dendrite can be avoid when the electrolyte's shear

modulus is higher than metal's shear modulus. In the case of lithium metal, they have shown that

if the electrolyte exhibits a shear modulus about twice that of the lithium (~ 109 Pa), dendrite

formation can be mechanically prevented103. However, this model has limitations, indeed in this

model the dendrite growth is derived from the growth of a single dendrite without taking into

account the interactions with the neighboring dendrites and the transverse interaction, the

electrostatic field.

Coexistence of both models. Nishikawa et al.175 demonstrated that both Chazalviel's model

and Monroe-Newman's model coexist.

They followed the cell potential over time and monitored the dendrite length. Their results

are presented in Figure 34. They determined three different stages as a function of time. Stage 1,

where no dendritic growth occurred, is an incubation time. At the end of stage 1, the Sand time

is reached, the surface concentration at the cathode decreases to zero leading to a divergence in

potential (see Figure 34 stage 1).

A second stage defined by a rapid dendritic growth is then observed, the current density is

high (J > J*). Dendrites nucleates and grows at a velocity close to the anions velocity, which is

compatible with the Chazalviel's model. The potential passed by a maximum, which implies a

maximum in the dendrite velocity (see in Figure 34). Dendrites obtained in this regime exhibits

coral morphology as shown in Figure 35 a).

Figure 34. Time evolution of the cell potential (solid line) and dendrite length (open circle), for a 0.011 cm thick cell at current density of 16 mA.cm-2. The straight dashed line shows the velocity predicted from Chazalviel's

model. The electrolyte concentration C0 is 0.2 mol.dm-3.


- 54 -

The last stage observed, stage 3, it is reached when the current density becomes low (J < J*)

due to the diminution of the inter-electrode distance (L) (due to the presence of dendrites) and

thus J increases (J is inversely proportional to L). Stage 3 is defined by much reduced velocity of

dendritic growth (see Figure 34 stage 3). The potential decreases slowly and the velocity of

dendrite growth decreases leading to the incompatibility with the Chazalviel's model. However, it

is compatible with the Monroe and Newman kinetic model. Dendrites obtained in this regime

exhibits a compact structure as shown in Figure 35 b).

Figure 35. SEM images of dendrites evidencing the structural change of the deposit at the transitions. The region close to the cathode (bottom of the picture, see inset a)) corresponds to the fast velocity regime. One observes a mossy structure. Individual grains are about 1.5 µm in diameter. At the tip of the dendrites (top of the figure, see inset

(b)) after the transition between the two regimes, the deposit has a more compact and regular structure. The current

density is J#2J*. The white solid line is 200 µm long in the low magnification image, and 20 µm in the two insets.

Figure 36 summarized the concentration profile as a function of the distance of the distance

from the cathode during the three different stages described above.

Figure 36. Schematic illustration of the time evolution of the cell. Ionic concentration is plotted as a function of a distance to the cathode (on the left of the figure): a) end of stage 1, at Sand time; b) stage 2; c) the distance between

the deposit and the anode is equal to l* (beginning of stage 3)


- 55 -

Both the Chazalviel's model and the Monroe and Newman model can follow each other. It is

worth noting that the current density used is above the critical current density, J*. After stage 2,

which is ruled by the Chazalviel model, the dendritic growth induced a decrease in the inter-

electrode distance (L), leading to an increase in the critic current density up to J* = J, since J* is

inversely proportional to L. Thus a critic distance L* is reached when J* = J (see Figure 36) and

the transition to Monroe and Newman regime occurs.

c. Lithium dendrite prevention

Various approaches have been studied to suppress lithium dendrite formation and growth in

the last forty years. In this section, we will discuss the main approaches proposed.

i. Lithium dendrite prevention by in situ formed SEI layer

It is worth to note that a suitable SEI has to exhibit several physical properties, such as good

adhesion, flexibility and homogeneity, to avoid the corrosion of the lithium, a transference

number of one and a small resistance and to be able to soak the lithium permanently in order to

prevent lithium dendrites.

Additives in electrolyte solution. Various functional additives have been reported to be

effective to reduce dendrite lithium formation and growth. In fact, electrolyte additives are added

in order to enhance the SEI films on the surface of lithium. The aim is to quickly form a stable

and dense interface in order to reduce the reactivity between lithium and the liquid electrolyte.

The presence of CO2 in liquid electrolytes increases significantly the lithium cycling

efficiency117,176. In fact, in contact with lithium, CO2 will form Li2CO3 which is an efficient

passivating agent. In addition, the interfacial impedance or solutions containing CO2 are lower,

and remain constant over aging177.

The addition of small amounts of hydrogen fluoride (HF) to electrolytes increases the cycling

stability to hundreds of cycles178. This is caused by the formation of a thin and compact surface

film composed of a bilayer structure LiF/Li2O179.

In the 80's Abraham et al. reported that the presence of 2-methyl furan (2-Me-F) in liquid

electrolyte as a good additive due to its ability to form a surface film through a chemical ring-

opening reaction/polymerization180. This thin layer prevents or slows down the reaction between

the electrolyte and lithium. Columbic efficiencies are improved and other related additives

containing furan such as 2,5 dimethyl furan, 2,5-dimethyl-thiohene, 3,4-dihydrofuran, 2-methyl-


- 56 -

tetrahydrofuran and 2,5-dimethyl-tetrahydrofuran have been studied and shows similar

functionalities181. Another additive, which forms a surface layer by ring-opening polymerization,

is the vinylene carbonate (VC) and fluoroethylene carbonate (FEC). They improve the Columbic

efficiency, by forming a uniform surface film onto lithium. In addition, this film presents a lower

interfacial resistance182.

Inert additives such as ammonium chlorides with n-alkyl group, benzene and toluene have

been studied183. These additives accumulate at the electrode/electrolyte interface to form a thin

layer which efficiently prevents the formation of passivation layers at the lithium surface.

Metals ions as additives have also been studied. Mastuda et al.184 were the first to find that the

Columbic efficiency could be improved in electrolytes containing Al3+, In3+, Ga3+ and Bi3+. These

metals ions will chemically or electrochemically deposit to form thin layers of lithium alloy which

improve the surface uniformity184. In fact, the deposition of such alloy happens preferentially on

the most active points of the lithium electrode, suppressing dendrite formation and improving

the Columbic efficiency.

ii. Lithium dendrite prevention by surface coating formed ex-situ

Another approach to prevent dendrite growth is to cover the Li electrode by a protective

layer (artificial SEI layer) formed ex-situ by treating Li metal prior to its use.

One good example of an ex-situ protective layer is reported by Umeda et al.185, they shows

that the silica film formed after exposing lithium metal to tetraethoxysilane (TEOS) can suppress

dendritic growth. Indeed, the lithium is plated back into an empty region which is created

beneath the protective film during the stripping process. They reported that the impedance is

unchanged after 100 cycles of Li plating and stripping.

Other coatings such as chlorosilane derivatives186 have also been investigated. Choi et al.187

used a cross-linked gel polymer electrolyte to coat the Li metal.

Wu et al.188 exposed Li metal to N2 gas in order to form a Li3N protective film. They reported

a reduced Li/electrolyte interface resistance with the prevention of corrosion of the Li electrode

in liquid electrolyte.

These ex-situ protective films have an uniform physical contact and good adhesion with Li

electrode. In addition, these protective films suppress the reaction between lithium metal and

non-aqueous liquid electrolyte. However, while these artificial SEI are a good protection at the


- 57 -

initial stage, during cycling it is unavoidable that these films will be destroyed which leads to an

acceleration of dendritic growth.

Coating materials. Various coating materials, including glasses (such as LiPON98) or a

lithium ion conductive organic/inorganic composite protective layer189, were used to prevent and

suppress dendrite formation.

Suppression of dendrites is also possible by exerting pressure against the surface and

blocking the space that is open for further dendrite growth. Those materials presents a shear

modulus about twice that of Li, and therefore use the mechanical arguments demonstrated by

Monroe and Newman103 in order to suppress lithium dendrite.

Recent studies on thin interconnected hollow carbon nanospheres190 and multilayered

graphene coating191 has been reported to improve charge-discharge efficiency. It is important to

distinguish these new types of coatings from the one described above due to the dissimilarity in

the protection process. In other words, the previous coatings stopped dendritic growth via

mechanical properties, when this coating type separates the SEI layers, formed in contact with

the electrolyte, from dendrite growth. Figure 37 presents a schematic illustration of the modified

Li anode structure and how this coating layer creates a scaffold for stabilizing the SEI layer.

Figure 37. Schematic illustration of Li anode structure190. Modifying the Cu substrate with a hollow carbon nanosphere layer creates a scaffold for stabilizing the SEI layer. The volumetric change of the Li deposition process

is accommodated by the flexible hollow-carbon-nanosphere coating.

The major limitation of this kind of protection is that after several cycles, thin films break

due to high volume changes.

iii. Lithium dendrite prevention by mechanical blocking

Polymer electrolyte. As discussed previously, dendritic growth can be suppressed if the

shear modulus of the electrolyte is about twice that of lithium, therefore polymer electrolytes are

good candidates. As seen previously, PEO based electrolytes are the most commonly used

polymer electrolyte. However, several studies in the early 1990s have proved that PEO itself


- 58 -

cannot block dendrite growth192. This is mainly due to the necessity to work at elevated

temperature (~80ᵒC), where PEO becomes a highly viscous liquid.

A combination of classic liquid electrolyte and inert polymer network has been reported as

gel polymer electrolytes and has been studied for dendrite suppression. Tastsuma et al.193

reported that 5-10 wt% of poly(acrylonitrile) (PAN) in the electrolyte is sufficient to efficiently

suppress dendrite. Eichinger et al.194 proved that the addition of PAN in classic electrolyte

improved lifetime of the battery up to 450 cycles. However, they showed that PAN reacts with

lithium leading to an increase in bulk and interfacial resistances.

More recently, Balsara et al.195showed in the case of block copolymer electrolyte of PEO and

PS blocks that the dendritic short-circuit is mitigate when the shear modulus increases.

Polymeric single ion conductors. As seen in the section above (Chap 1.6. b)) the Sand

time is proportional to 1/ta (equation 13). For a binary electrolyte ta + tLi+ =1. Therefore, the

lithium transference number has a large impact on the electrochemical performances of Li

batteries. Indeed, if tLi+ is close to unity, ta is small leading to a large Sand time. Moreover, tLi+

close to unity means µa close to zero leading to a velocity of dendrite growth infinite small

according to Chazalviel model141. In other words, if tLi+ is equal to one, theoretically no lithium

dendrite should nucleates, meaning that Li metal could be reversibly plated and stripped.

However, in classic liquid electrolyte the transference number is less than 0.5, in the case of PEO

based polymer tLi+ < 0.2. Therefore, a new strategy to stop dendrite growth is to increase the

electrolyte transference number. As seen in a previous section (Chapter 1. 5. e)), single-ion have

been developed.

Recently, Bouchet et al.142 reported a multifunctional single-ion polymer electrolytes based on

poly-anionic block copolymers, which present impressive gains in power performances

compared to classic lithium metal polymer battery. They cycled their batteries more than 100

cycles at different current densities without any signs of dendritic growth.

Nanoporous ceramic. Recently Tu et al.196,197 proposed a novel electrolyte design where a

liquid electrolyte is hosted in the pores of a ceramic membrane with a high areal density and

nanometer-sized pores. Figure 38 presents a scheme of this new electrolyte design.


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Figure 38. Middle and left : Schematics of the structure and preparation method of PVDF-HFP/Al2O3 separator. SEM images of the PVDF-HFP/Al2O3 with 100nm nanopores: top, cross section of the composite,

right, cross section of the internal alumina layer, bottom, boundary between alumina and polymer196.

This composite electrolyte exhibits an ionic conductivity greater than 1 mS.cm-1 at room

temperature and a mechanical modulus at least of 0.5 GPa. Therefore, this new type of

electrolyte presents both high ionic conductivity and good mechanical properties. In addition,

they reported more than 1100 cycles in lithium battery and in lithium symmetric cell. This

electrolyte shows more than 1000 hours of stable operation at 0.2 mA.cm-2. Moreover, they have

observed that this electrolyte is able to contain the non-uniform lithium deposits within small

regions but does not suppress dendrites.

iv. Lithium dendrite prevention by other approaches

Effect of charge protocol. Both the Chazalviel and the Monroe-Newman models

demonstrate that the values of current density affect significantly lithium dendrite nucleation and

growth. However, a recent study by Mayers et al.198, reported that the charge protocols,

galvanostatic or pulse, has a strong impact on lithium dendrite growth. They proposed a coarse-

grained lattice model to investigate the relationship between electrode charging conditions and

deposition morphology in liquid based electrolyte.

They revealed that dendrite formation emerged from a competition between the timescales

of Li+ diffusion and their reduction at the anode. They reported an effective suppression of


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lithium dendrite up to 96% with lower over-potential and shorter electrode pulse duration.

Indeed, longer rest periods between pulses favors cations diffusion into the structure, and

therefore cation reduction during pulses will lead to high-density deposit. This approach aims for

the relaxation of concentration gradients and therefore, at the interface electrolyte/lithium, there

is permanently anions.

Effect of pressure. Gireaud et al.199 reported that lithium deposited at higher pressure is

more dense and uniform. They showed that the pressure applied on the cell confines dendritic Li

deposits to the vicinity of the negative electrode. The efficiency of lithium deposition increases

from 60% up to 90% with the pressure increasing from 0.7 kg.cm-2 to 7 kg.cm-2. However, it is

worth noting that it is not always feasible to apply pressure on the cell.

Effect of surface smoothness. The surface state of lithium electrode significantly affects

the morphology of subsequent electro-deposition. Gireaud et al.199 observed that metal

imperfections favors growth of dendritic lithium deposit. Such imperfections will locally

enhanced the current density leading to dendritic growth.

Lithium-coated polymer matrix. More recently a new anode design, exhibiting a minimum

volume change and dendrite free lithium metal anode, has been reported by Liu et al.200. They

reported an electrospun polyimide (PI) stable vs Li+/Li via a layer of zinc oxide coating, molten

lithium can then be drawn into the matrix. Figure 39 presents a schematic illustration of the

behavior of this new anode design and SEM images of the fibers after stripping and plating.

They showed that the Li stripping/plating is well confined inside the polymer matrix. Moreover,

they showed dendrite suppression after 10 cycles at 5 mA.cm-2.. This approach aims to divide the

surface to decrease current densities and thus limit dendritic growth.


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Figure 39. Well-confined stripping/plating behaviour of the Li-coated PI matrix. Top-view SEM images of (a) the exposed top fibres of the Li-coated PI electrode after stripping away 5mAh.cm-2 Li; (b) exposed top fibres partially filled with Li when plating 3mAh.cm-2 Li back and (c) completely filled PI matrix after plating an

additional 2mAh.cm-2 Li back (current density 1mA.cm-2, in EC/DEC). The polymeric matrix ensures that Li is dissolved and deposited from the underlying conductive Li substrate and, as a result, Li is effectively confined into the matrix. (d) Schematic illustrating the alternative undesirable Li stripping/plating behaviour where, after stripping, Li nucleate on the top surface, leading to volume change and dendrites shooting out of the matrix. Scale

bars, 5 mm.

Uniform Li-ion flux. Other approaches have focused on making the lithium ion flux more

uniform. This approach is based on the suppression of local transversal concentration gradients,

in other words to suppress the space charge in the vicinity of the negative electrode (Chazalviel

model). Figure 40 presents a schematic illustration of the structure of the mussel-inspired

polydopamine-coated separators201.

Figure 40. Schematic illustration of the structure of the uniform Li ion flux via a separator with good wettability and ionic liquid201.


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This approach is similar to the single ion block copolymer electrolyte approach which

suppresses the concentration gradient.

Self healing electrostatic shield mechanism. A new type of additive that functions on the

basis of electrostatic interactions with Li ions has been discovered by Ding et al.202. They reported

a self-healing electrostatic shield mechanism depending on an electrolyte additives which exhibits

an effective reduction potential lower than that of lithium ion. Figure 41 shows how the self-

healing shield mechanism can self-heal the tip of a lithium dendrite. They showed that at low

concentration (< 0.05 M) cesium and rubidium additives remain positively charged, in other

words, in the mixed electrolyte these ions will not form thin layers of alloys at the electrode

surface due to an effective reduction potential lower (for low concentration) than that of lithium.

Instead they repel incoming Li ions from sharp tips to the valley, leading to a smooth lithium

metal surface morphology even after cycling.

Figure 41. Illustration of lithium deposition process based on the SHES mechanism202.

Figure 42 presents SEM images of lithium morphologies deposited in non-aqueous liquid

electrolyte without CsPF6 (a) and with CsPF6 (b and c). It is clear that the addition of CsPF6 leads

to a regular and uniform lithium deposition. However, it only applies for low current densities.


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Figure 42. SEM images of the morphologies of Li films deposited in the 1M LiPF6/PC electrolyte with CsPF6 concentration of : a) 0 M, b) 0.005 M, c) 0.05 M, at a current density of 0.1 mA.cm-2 202.

3D patterning via micro-needle technique. A recent work by Ryou et al.203, reported an

applied micro-needle surface treatment technique for lithium metal, which enhances rate

capability and cycling stability, and reduce interfacial resistance. This technique increases the

effective surface area, by the 3D patterning, and dissipates the electron density at the given area

current density. Figure 43 shows schematic figures illustrating the micro-needle technique.

Figure 43. Schematic figures illustrating the micro-needle technique. a) An overview of the micro-needle technique creating surface patterns on the Li metal. Schematic illustration describing the Li plating mechanism on b) micro-

needle treated Li metal and c) magnified images of the part b)203.

7. Conclusion

In a society more aware of the climate change and in order to limit the increase of global

warming below 2ᵒC, the massive introduction of renewable energies in the energy mix becomes a

necessity, as well as a better mix management and the mass development of electrified transport.

In addition of stationary storage, the society needs the widespread use of electric and hybrid

vehicles which will contribute to decrease the greenhouse emissions. Therefore, the energy

storage is one of the greatest issue of the 21st century and the electrochemical energy storage with

batteries is one of the most promising candidate due to its versatility, energy efficiency, its

absence of inertia.


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Since the discovery of the Volta's pile, many different secondary batteries have been

developed from lead-acid batteries to lithium-ion batteries including nickel cadmium, nickel

metal hydride, alkaline and lithium sulfur. After the commercialization in 1991 of Li-ion batteries

that conquered the market of nomad applications, the increasing needs of always higher

performances has led to more research effort for the next battery revolution,. Nowadays, only

few battery chemistries are promising enough to fulfill the requirements for tomorrow's needs,

among them, the metal-air battery based on metallic negative electrode (Li, Na, Zn, Al, Fe) and a

positive based on the air electrode. The lithium-air battery exhibits the highest theoretical

performances with a specific energy up to 3582 Wh.kg-1 for the aqueous Li-air battery (to be

compared ti the 250 Wh.kg-1 of the best Li-ion battery).

In spite of great theoretical specifications, this technology still have very challenging issues to

be addressed: the presence of an aqueous electrolyte implies the protection of the negative

electrode, i.e. the lithium metal, by a solid lithium ionic conductor, which is watertight. The

different solid electrolytes have been reviewed. However, most of the solid electrolytes are not

stable in contact with lithium or their stability is still under discussion. Therefore, another layer

between the lithium and the ceramic has to be added: several options are available, including

classic liquid organic electrolytes, LiPON, PEO polymer based electrolytes (block copolymer

electrolytes or single-ion electrolytes).

The last issue is related to the use of lithium metal itself, which presents dendritic growth

during the recharge. In order to mitigate this heterogeneous growth, it is important to

understand the basis of dendrites nucleation and growth. We have introduced here, firstly the

SEI formed at the interface lithium-electrolyte, then the models of nucleation and growth of

dendrites, and finally, the different approaches to prevent or mitigate the dendrites growth.


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Chapter 2.

LiPON as a protective layer for the ceramic

Abstract

In this chapter, a primary part will be dedicated to the characterization

of the Ohara glass ceramic (Ohara GC) via electrochemical impedance

spectroscopy. Ionic conductivities for the bulk and the grain boundaries will be

discussed and their activation energy will be determined. Then, the Ohara

GC/LiPON sandwich will be studied. Here, the aim is to be able to

determine, by a simple method, the LiPON contribution in the Ohara

GC/LiPON sandwich. Ionic conductivities and activation energy of the

LiPON will be also determined.

Chapter 2. Ceramic protective layer alternative

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Table of contents

Chapter 2. LiPON as a protective layer for the ceramic............................................... 73

1. Experimental section .................................................................................................................................. 75

a. Materials ..................................................................................................................................................... 75

b. SEM characterization .................................................................................................................................. 75

c. Electrode sputtering and cell assembly ........................................................................................................... 76

d. EIS measurement ......................................................................................................................................... 76

2. Results and discussion ................................................................................................................................. 77

a. Micro-structural analysis .............................................................................................................................. 77

b. Ohara GC results ........................................................................................................................................ 78

c. LiPON-Ohara GC-LiPON results ............................................................................................................ 81

3. Conclusion ..................................................................................................................................................... 89

References of Chapter 2 ....................................................................................................................................... 91


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The aim of this study is to separate ceramic and LiPON contributions by an analysis of

electrochemical impedance spectroscopy. In order to investigate the different contributions

of the LiPON/ceramic/LiPON sandwich, ac impedance measurements have been

performed.

1. Experimental section

a. Materials

Plates of Li+ ion conductor NaSICON type glass-ceramic (Ohara GC) electrolyte were

purchased from Ohara Corporation1. The dimensions of the plates were 2.54 cm² with a

thickness of 160 µm or 250 µm (see Table 1). Each plate was cut into smaller pieces of

around 0,635 cm².

LiPON thin films were deposited at room temperature by rf magnetron sputtering of a

Li3PO4 target in a nitrogen atmosphere. Fleutot et al.2 investigated the sputtering parameters

in order to optimize LiPON ionic conductivity. In this study, we used the same optimized

procedure. Two different LiPON thicknesses were deposited onto the Ohara GC substrates,

cells with expected values about 0.6 µm and 1.2 µm of LiPON on each surfaces of the

ceramic were made (see Table 1).

Sample name Ohara GC

thickness (µm)

Total LiPON thickness by SEM

(µm)

Cu-LiPON1,2-Ohara-LiPON1,2-Cu 1

250 2.30 ± 0.04


250 2.30 ± 0.04


160 1.38 ± 0.04


160 1.39 ± 0.04

Table 1. Material's thicknesses for LiPON/Ohara GC/LiPON sandwich and LiPON thicknesses determined by SEM (see Figure 2).

b. SEM characterization

In order to characterize the LiPON/Ohara GC/LiPON sandwich, and determine the

thicknesses of LiPON thin films, cross sectional images were taken by a scanning electron

microscopy (SEM Hirox SH-1500 and FEG SEM Zeiss Ultra 55). Care was taken to transfer

samples from the glove box to the SEM by using a transfer airlock filled with argon in a


- 76 -

sealed recipient to prevent any reaction with ambient air. SEM images were taken at 3 kV

acceleration and a magnification of 20 k.

c. Electrode sputtering and cell assembly

In order to investigate the electrical properties of the Ohara GC, gold thin electrodes

were deposited at both surfaces at room temperature by dc magnetron (EMITECH SC760

Sputter Coater). Ceramics covered by gold were placed between two stainless steel (SS) plates

in a swagelock cell assembly (Figure 1). Care was taken in order to keep the ceramic

unbroken.

Figure 1. Schematic representation of swagelock assembly used for characterization by impedance spectroscopy.

Copper thin films were deposited by rf magnetron sputtering onto both LiPON surfaces

(Plassys MP 3005). The copper deposit thickness was about 400 nm on both side.

In order to avoid any deterioration of LiPON, the exposure to moisture and air was

minimized. All handling of materials and assembly of cells were accomplished in an argon

filled glove box with less than 1 ppm of water and less than 5 ppm of O2. The samples were

placed in a swagelock in the same manner as described before (see Figure 1) in order to

perform EIS characterizations outside the glovebox.

d. EIS measurement

The ionic conductivities of Ohara GC and LiPON/Ohara GC/LiPON were

characterized by electrochemical impedance spectroscopy (EIS) using a Solartron 1260

Ceramic

or

LiPON / ceramic / LiPON

Stainless steal

plunger

SS disk

SS disk

Spring

Stainless steal

plunger


- 77 -

frequency response analyzer, driven by Zplot Software (Scribner Associates)3. The applied

AC voltage ranged from 10 mV to 100 mV, Ac impedance analysis was performed over a

frequency span from 30 MHz to 100 mHz. In order to fit the EIS data, a complex least

square fitting program, Zview (Scribner Associates)3, was used. Swagelock cells were placed

in a climatic chamber with a thermocouple to control the temperature, and the EIS

measurements were performed from -40°C to 130°C with 5 or 10°C increments. For each

temperature, the cells were allowed to equilibrate until stabilization of the impedance.

2. Results and discussion

a. Micro-structural analysis

SEM analysis are performed on the LiPON/Ohara GC/LiPON sandwich in order to

have an accurate thickness determination of the LiPON layers for the conductivity

calculation. Figure 2 shows a typical SEM cross sectional images of the sandwich structure.

The different layers are easily distinguished due to their difference in chemical composition

for the copper current collector, LiPON thin films and lithium ion conducting Ohara GC.

The Ohara GC (bottom of SEM images in Figure 2) exhibits a polycrystalline structure in

an amorphous glass matrix, which is coherent with the literature4. In addition, it presents a

compact and dense morphology.

Figure 2 shows the two different thicknesses of LiPON thin films (a) the 600 nm and b)

1200 nm). The LiPON thin films present a highly dense, amorphous and homogeneous

structure, which is in a good agreement with litterature5–7.

Figure 2.- SEM cross-sectional image of a Ohara GC-LiPON multilayered structure with a deposited LiPON layer of a) 600nm and b) 1200 nm.


- 78 -

Several SEM images are taken and analyzed in order to have an accurate average

thickness of LiPON thin film for each of the four cells. As expected, thicknesses are found

different compared to the expected values of 1.2 and 2.4 µm. Table 1 presents the results

from SEM measurements for the four cells.

b. Ohara GC results

Figure 3 shows a typical response in electrochemical impedance spectroscopy for a Au/

Ohara GC(160μm )/Au cell at -30ᵒC in a Nyquist plot.

Figure 3. Nyquist plot of a) Au/Ohara GC/Au cell at -30°C (plot) and fit (line); and b) zoom at high frequencies.

All impedance spectra are composed of a first semi-circle at high frequency (f > 106 Hz)

which corresponds to the electrical response of the grains of the Ohara GC and thus permits

to determine the grains (g) resistance Rg. At -30ᵒC, its characteristic frequency is about 2

MHz. This contribution can be observed from -40ᵒC to 10ᵒC, at higher temperatures it

presents a characteristic frequency above the solartron frequency limits and thus could not

be identified. In the middle to high frequency (domain 106 to 104 Hz) another loop is

observed which is attributed to the electrical response of the Ohara GC grain boundaries and

presents a characteristic frequency around 16 kHz at -30ᵒC. This contribution allows the

determination of the grain boundaries resistance Rgb. This analysis is coherent with previous

study by Fu et al.4 who have reported similar contributions at - 40ᵒC. In practice, the largest

contribution is due to the grain boundary.

14x103

12

10

8

6

4

2

0

- Im

(Z)

(W)

14x103121086420

Re(Z) (W)

Ohara ceramic at - 30°C Fit result

16 kHz

a)

1000

800

600

400

200

0

- Im

(Z)

(W)

10008006004002000

Re(Z) (W)

2 MHz

Oharaceramic at - 30°C Fit result

b)


- 79 -

It is important to notice the EIS at low temperature is rarely reported and the loop

present at middle high frequency, at room temperature and higher temperature, is often

attributed to the bulk of the Ohara GC8. This is an error, this contribution corresponds to

the grain boundaries and not the grains.

Finally, at lower frequencies (f < 104 Hz), the near-vertical line observed is due to

accumulation of charges at the Ohara GC/gold interface associated with the interfacial

capacitance of a blocking electrode.

The electrical behavior of the cell is modeled using the elementary brick-layer model9,10

(BLM), which was outlined at the end of the 1970's. This model represents grains by cubes

with a certain size and they are separated by homogeneous grain boundaries with constant

thickness (generally around 5 nm according to the model of Fisher11). For polycrystalline

structures, we assume that parallel grain boundaries conduction is negligible. Therefore, the

electrical behavior of the Ohara GC is modeled via an equivalent electrical circuit showed in

Figure 4. The Au/Ohara GC/Au cell is modeled with two (R parallel CPE) circuits in series.

The first one for the grain contribution with Rg and CPEg and the second one for the grain

boundaries with a resistance Rgb and CPEgb, where CPE stands for constant phase element.

A serial resistance Rcable and an inductance Lcable is used to model the electrical behavior of

circuit cabling12, and finally a constant phase element CPEelect is used to model the blocking

electrode behavior.

Figure 4. Equivalent circuit used to model the Au/Ohara GC/Au cell.

A fit of the Au/Ohara GC/Au cell at -30ᵒC is represented by the black solid line in

Figure 3. It shows a very good agreement with the experimental results. Ionic conductivities

for both grain and grain boundaries (in this case it is an effective conductivity) are calculated

using equation (1), with resistances of grain and grain boundaries obtained by least square

fitting.

Lcable Rcable Rg

CPEg

Rgb

CPEgb

CPEelect


- 80 -

!g="lOhara"

SOhara"."Rg and !gb"=

"lOhara"

SOhara"."Rgb (1)

With:

- lOhara : Ohara GC thickness (in cm)

- SOhara : surface area of the Ohara GC (in cm2)

Finally, the effective ionic conductivity of the whole ceramic, σeff, is calculated according

to equation (2). The ionic conductivities obtained are plotted versus temperature in an

Arrhenius plot in Figure 5.

!eff="lOhara"

SOhara"."(Rg"+Rgb)" (2)

Figure 5 : Arrhenius plot of the effective Ohara GC conductivity, grain boundaries and bulk contributions from -40ᵒC to 130ᵒC.

Figure 5 shows the Arrhenius plot of the Ohara GC’s conductivity. The grain exhibits a

high ionic conductivity of 1.10-3 S.cm-1 at 25°C with an activation energy, Ea, of 0.31 eV,

which is in a very good concordance with published results for this material1,4,13. The grain

boundaries conductivity is one order of magnitude lower with 1.9.10-4 S.cm-1 at 25ᵒC and an

activation energy of 0.37 eV. Thus, the effective Ohara GC ionic conductivity is mainly due

to the contribution of the grain boundaries and we obtain 8.2.10-5 S.cm-1at 25°C with a

similar activation energy.

10-6

10-5

10-4

10-3

10-2

Co

ndu

ctivity (

S.c

m-1

)

4.54.03.53.02.5

1000/T (K-1

)

Grain boundary Grain (bulk) Effective conductivity

130 60 12 -23 -50

Ea = 0.31 eV

Ea = 0.37 eV

Temperature (°C)


- 81 -

In the case of other oxide ion conducting ceramics, such as Y2O3 doped CeO2, the higher

activation energy of the grain boundaries compared to the activation energy of the grain, is

explained by electrostatic barriers due to the formation of a space charge layer at the grain

boundaries14. However, in the case of the Ohara GC, Mariappan et al.13 reported that Mott-

Schottky-type space charge layers are not at the origin of the high grain boundary resistance.

The equivalent capacitance (C in F) of the material is calculated from the pseudo-

capacitance, Q in F.s1-n with n the phase (n=1 for a pure capacitance) obtained from the EIS

fits, according to equation (3):

C"="Q1/n"*"R(1/n)-1" (3)

The dielectric constant, εr, is then calculated according to the equation (4):

#r"=C

#0"."(S $% ) (4)

Where, ε0 is the vacuum permittivity (8.85 .10-14 F.cm-1), S and $ are the geometrical

factors.

We found for the Ohara GC an εGC = 57 which is higher than the value reported in the

literature and it is probably coming from errors due to other capacitance contribution in

parallel coming from the assembly of the cell.

c. LiPON-Ohara GC-LiPON results

A comparison of the Au/Ohara GC/Au and Cu/LiPON/Ohara GC/LiPON/Cu

spectra at -30°C is given in Figure 6 in Nyquist plots. It is clear from the zoom at high

frequencies (Figure 6 b)) that the contribution of the Ohara GC remains unchanged in the

presence of LiPON. On the contrary, we observe a strong modification of the shape of the

spectra in the middle low frequencies. We may assume that it is due to the LiPON. Finally, in

all cases, a vertical line at low frequencies is observed, which corresponds to a typical

blocking electrode’s behavior.


- 82 -

Figure 6. a) Impedance spectrum of Au/Ohara GC/Au (red) and Cu/LiPON/Ohara GC/LiPON/Cu sandwich (blue) structure at -30°C (plot) and fit (160μm) (line) and b) a zoom at high frequencies.

Before making a subjective parameterization of the spectrum with our equivalent circuit,

a difference of the spectra LiPON/Ohara GC/LiPON and Ohara GC will allow to observe

directly the evolution after adding LiPON to the Ohara GC. The result is presented in Figure

7. After the subtraction, there remains one contribution at middle frequencies, which is an

almost a perfect semi-circle. Thus, it is possible to model it by a simple resistance in parallel

with a constant phase element.

We could then calculate the dielectric constant (using the equation (3) and (4)) for the

LiPON layer and we found εLiPON =26.4, which is in good agreement with the literature7,15.

In addition, on the difference between Ohara GC and the LiPON/Ohara GC/LiPON

spectrum a capacity is still present at low frequency. This is due to the difference of

interfacial capacitance between the Ohara GC/gold and the LiPON/Cu.

40x103

30

20

10

0

- Im

(Z)

(W)

40x1033020100

Re(Z) (W)

Ohara ceramic at -30°C Fit result Ohara ceramic LiPON-ceramic sandwich at -30°C Fit result LiPON-ceramic sandwich at -30°C

a)

16 kHz 775 Hz

1000

800

600

400

200

0

- Im

(Z)

(W)

10008006004002000

Re(Z) (W)

2 MHz

b)


- 83 -

Figure 7. Result of the difference between Ohara GC EIS and sandwich EIS at 50ᵒC.

This new contribution can come from either the bulk of LiPON and/or the interface

LiPON/Ohara GC.

Finally, the electrical behavior of the Ohara GC-LiPON structure is modeled by the

equivalent circuit shown in Figure 8, where another RLiPON // CPELiPON is simply added in

series to the equivalent-circuit previously used for the Au/LiSICON/Au cells (see Figure 4)

in order to model the impact of the LiPON layer. The fit is represented in a solid blue line in

Figure 6. As expected we obtain a very good agreement with the experimental results.

Figure 8.- Equivalent circuit used to fit the response for Cu/LiPON/Ohara GC/LiPON/Cu sandwich structure.

This new contribution can be either attributed to the bulk contribution of the LiPON

(proportional to l/S) or to the ionic charge transfer at the LiPON/Ohara interface

(proportional to 1/S). Thus, in order to identify the nature of the LiPON contribution, we

can compare the two LiPON prepared. The spectra of the two samples at -30°C are given in

100

80

60

40

20

0

-Im

(Z)

(W.c

m2)

100806040200

Re(Z) (W.cm2)

0.46 MHz

Lcable Rcable Rg

CPEg

Rgb

CPEgb

RLiPON

CPELiPON

CPEelect


- 84 -

Figure 9. It appears that the contribution of the Ohara GC-LiPON 2 x 1.2 µm is almost

doubled compared to the Ohara GC LiPON 2 x 0.6 µm. Therefore, we can unequivocally

conclude that the additional contribution comes from the bulk of LiPON. By consequence,

as an interesting and original result it appears that the Li+ charge transfer at the interface

LiPON-Ohara GC is rather small and not observable.

Figure 9. a) Impedance spectrum of two Cu/LiPON/Ohara GC/LiPON/Cu sandwich structure at -30°C with two different LiPON thicknesses (plot) and fit (line) and b) zoom at high frequencies.

At low temperature, all the contributions are well defined and it is possible to obtain a

very nice fit with a good accuracy for all the different parameters. However, at higher

temperatures the characteristic frequencies of the LiPON and the Ohara GC grain

boundaries converge and only one contribution is observed. Thus, the accuracy of the

parameter values becomes low. An illustration of this effect is given in Figure 10 where we

plot the spectra obtained at different meaning temperatures. Since the time constant for a

material (τ = 1/fc = 2π/ωc = 2 π RC) converge from -30 ᵒC to 50ᵒC, it means that the

activation energy for the process associated to the bulk LiPON is higher than the one from

the Ohara GC.

100x103

80

60

40

20

0

- Im

(Z)

(W)

100x103806040200

Re(Z) (W)

775 Hz

a)

520 Hz

Ceramic-LiPON 2 x 1.2 mm

Ceramic-LiPON 2 x 0.6 mm

14 kHz

2000

1500

1000

500

0

- Im

(Z)

(W)

2000150010005000

Re(Z) (W)

2 MHz

b)


- 85 -

Figure 10 : Nyquist plot of the Ohara GC (red square) and the Ohara GC with LiPON layer (blue diamond) at a) -30°C, b) 10°C, c) 25°C and 50°C.

i. Sample characterization : Conductivity measurement

Ionic conductivities of LiPON are calculated. Results of LiPON ionic conductivity are

plotted versus the inverse of temperature in an Arrhenius plot in Figure 11. Due to the

overlap of LiPON and Ohara GC characteristic frequencies, the results are only plotted with

accuracy from -35ᵒC to 10ᵒC. We obtain for this representation a linear variation of the

conductivity values described by the Arrhenius law (see equation (5)).

40x103

30

20

10

0

- Im

(Z)

(W)

40x1033020100

Re(Z) (W)

Ohara -30°C LiPON-Ohara-LiPON -30°C

14 kHz500 kHz

800 Hz 430 Hz

a)

2000

1500

1000

500

0

- Im

(Z)

(W)

2000150010005000

Re(Z) (W)

Ohara 10°C LiPON-Ohara-LiPON 10°C

112 kHz

81 kHz

8 kHz10 kHz

b)

1000

800

600

400

200

0

- Im

(Z)

(W)

10008006004002000

Re(Z) (W)


193 kHz

300 kHz

27 kHz

15 kHz

c)

300

250

200

150

100

50

0

- Im

(Z)

(W)

300250200150100500

Re(Z) (W)


0.66 MHz

0.80 MHz

57 kHz

100 kHz

d)


- 86 -

!"="!0

Texp

-Ea

RT (5)

Where σ0 is the conductivity at infiniteT0, Ea the activation energy and R the constant of

perfect gases.

Figure 11. Arrhenius plot of LiPON conductivity on the sandwich cell.

The literal values are listed in Table 2. We obtain an activation energy of 0.56 eV and the

conductivity at 25ᵒC is 1.9.10-6 S.cm-1. These values are in very good agreement with the

results in literature for the LiPON with an activation of 0.54 ± 8 eV7 and a conductivity of (2

± 0.7).10-6 S.cm-1 2,16,17. The ionic conductivities of the Ohara GC inside the multilayered

structure are in good agreement with the results presented in Chapter 2.1.b). Therefore, the

multilayered structure shows ionic conductivity equals to the sum of pure Ohara GC and

pure LiPON conductivities.

Samples σ at 25ᵒC (S.cm-1) Ea (eV)

LiPON2,7,17 (2 ± 0.7).10-6 0.50 ± 0.06

LiPON (in this study) 1.9.10-6 0.56

Ohara4,13,18 1.0.10-4 0.36 ± 0.01

Ohara (in this study) 8.2.10-5 0.37

Table 2 : Ionic conductivity at 25ᵒC with activation energy for Ohara GC and LiPON from the literature and from this study.

10-8

10-7

10-6

Co

nd

uctivity (

S.c

m-1

)

4.24.03.83.63.4

1000/T (K-1

)

Temperature (°C)

LiPON

Ea = 0.56 eV

20 5 -10 -20 -35


- 87 -

Firstly, this shows that the interface between LiPON and Ohara GC is very good and

secondly that the ionic charge transfer at the interface LiPON/Ohara-GC is very small

which suggests that there is no significant energy barrier for this process and the charge

transfer can only give a second order contribution.

At 10ᵒC, for an Ohara GC of 250 μm and a LiPON thickness of 2 x 600nm the total

ceramic contribution is 840 Ω, whereas the LiPON contribution is 487 Ω. Thus the major

contribution is clearly the ceramic contribution.

ii. Electrochemistry a lithium plating

P. Stevens et al.19 have studied lithium electro-deposition on a stainless steel current

collector deposited on the LiPON layer. The other half cell is an aqueous electrolyte (5 M

LiOH) with a reference electrode and a counter electrode in stainless steel. They succeeded

to deposit electrochemically extremely high area capacities up to 100 mAh.cm-2 of 100%

dense lithium at a current density of 0.2 mA.cm-2. The Figure 12 shows an illustration of

such deposit obtained by SEM.

Very interestingly, columnar lithium has been observed (charged with a capacity of 80

mA.cm-2, the obtained thickness is 370 μm). In this case the lithium growth nature is one

dimensional, they did not seen lateral growth of lithium.

Figure 12.SEM micrograph of electrochemically deposited lithium from an aqueous electrolyte for an electrode charged at 80mAh.cm-2.

Once they succeeded to obtain dense lithium electrode, they performed cycling

experiment onto the cell at 0.2 mA.cm-2. Figure 13 shows their results for 100 cycles. The


- 88 -

bottom axis is split for more clarity. During the first 60 cycles, the potential is stable and the

cell exhibits good cyclability with a potential plateau reached in charge and discharge.

However, after 60 cycles the potential starts to increase.

Figure 13. Voltage versus time curve for the cycling of the LiPON-Ohara GC cell at 0.2 mA.cm-2 at room temperature, the number inside the graph correspond to the cycle number.

During cycling, some columns are disconnected from the LiPON surface20 leading to a

reduction of the active surface associated to an increase of the resistance and thus the

resulted voltage rises (Figure 13).

Figure 14 shows a SEM image after such a loss of the active surface

Figure 14. SEM image of lithium after cycling 20

-3.4

-3.2

-3.0

-2.8

Vo

lta

ge

(V

)

806040200Time (h)

800780760740720700

Ohara GC - LiPON - Lithium

1

100


- 89 -

Moreover, such delamination of the lithium from the LiPON surface has been also

observed during the lithium electro-deposition process. In fact, if the intimacy between

lithium and LiPON is lost, the corresponding surface area is no longer active. Therefore, the

lithium electro-deposition is stopped, leading to columns shorter as seen in Figure 15.

Figure 15. SEM micrograph of lithium electrochemically deposited through LiPON and ceramic at 0.2 mA.cm -2 showing the disconnection between lithium and LiPON leading to smaller column.

Even if it has been shown that this inorganic/inorganic assembly allows enables the

desired dense deposits at 0.2 mA.cm-2, the delamination seems inevitable leading to inactive

lithium.

In addition Larfaillou et al.15 reported the degradation or oxidation of Li metal at the

LiPON/Li interface after aging. This is a possible explanation of the delamination of lithium

during cycling.

3. Conclusion

In this section, firstly we studied the Ohara GC electrical properties by impedance

spectroscopy. The bulk and grain boundaries contributions have been separated (especially at

low temperature).Their resulting ionic conductivities using the geometrical factors have been

calculated, together with the activation energy (at 25ᵒC we obtained: σg = 1.0.10-3 S.cm-1 with

Ea = 0.31 and σgb = 1.9.10-4S.cm-1 with an Ea = 0.37 and an effective Ohara GC conductivity

σeff = 8.2.10-4 S.cm-1) Our results are in excellent agreement with the literature values.


- 90 -

The Ohara GC sandwiched by two LiPON layers has been then studied in order to

analyze the LiPON contribution in the sandwich structure. We have shown that we were

able to observe the LiPON electrical behavior by working at low temperature. Interestingly

in this sandwich the Li+ charge transfer contribution remains a second order contribution

that cannot be estimated. Therefore, this charge transfer must be very fast. A classic electrical

equivalent circuit has been used and the LiPON contribution could be isolated and its

conductivity and activation energy have been determined. Our values are in good coherence

with the literature data (with σLiPON = 1.9.10-6 S.cm-1 at 25ᵒC and Ea = 0.56).

The LiPON thin film is a good protective layer for the Ohara GC. The electrochemically

deposition of lithium metal through the Ohara GC/LiPON sandwich has been investigated

by Stevens et al.19 and dense and columnar lithium electro-deposit has been observed.

However, during cycling, the lithium/LiPON interface is lost resulting in inactive lithium

surface area and the increase of the polarization. Therefore, another protective layer has to

be found. An interesting alternative candidate is the solid polymer electrolyte which are

flexible and can have good mechanical properties.


- 91 -


1. www.oharacorp.com.

2. Fleutot, B., Pecquenard, B., Martinez, H., Letellier, M. & Levasseur, A. Investigation of the local

structure of LiPON thin films to better understand the role of nitrogen on their performance. Solid

State Ion. 186, 29–36 (2011).

3. ZPlot® and ZView®. www.scribner.com

4. Fu, J. Fast Li+ Ion Conduction in Li2O-Al2O3-TiO2-SiO2-P2O2 Glass-Ceramics. J. Am. Ceram.

Soc. 80, 1901–1903 (1997).

5. Bates, J. B. et al. Electrical properties of amorphous lithium electrolyte thin films. Solid State Ion.

53, 647–654 (1992).

6. Hamon, Y. et al. Influence of sputtering conditions on ionic conductivity of LiPON thin films.

Solid State Ion. 177, 257–261 (2006).

7. Y. Hamon. Nitruration de verres conducteurs ioniques en couches minces. Universite Sciences et

Technologies (2004).

8. Tenhaeff, W. E., Perry, K. A. & Dudney, N. J. Impedance Characterization of Li Ion Transport at

the Interface between Laminated Ceramic and Polymeric Electrolytes. J. Electrochem. Soc. 159,

A2118–A2123 (2012).

9. Beekmans, N. M. & Heyne, L. Correlation between impedance, microstructure and composition of

calcia-stabilized zirconia. Electrochimica Acta 21, 303–310 (1976).

10.Bouchet, R., Knauth, P. & Laugier, J.-M. Theoretical analysis of the impedance spectra of

electroceramics Part 2: isotropic grain boundaries. J. Electroceramics 16, 229–238 (2006).

11.J. Philibert. Diffusion et Transport de Matière dans les Solides. Les Editions de Physique,p 227

Les Ulis 1987

12.Bouchet, R., Lascaud, S. & Rosso, M. An EIS Study of the Anode Li/PEO-LiTFSI of a Li Polymer



- 92 -

13.Mariappan, C. R., Gellert, M., Yada, C., Rosciano, F. & Roling, B. Grain boundary resistance of

fast lithium ion conductors: Comparison between a lithium-ion conductive Li–Al–Ti–P–O-type

glass ceramic and a Li1.5Al0.5Ge1.5P3O12 ceramic. Electrochem. Commun. 14, 25–28 (2012).

14.Meyer, R., Guo, X. & Waser, R. Nonlinear Electrical Properties of Grain Boundaries in Oxygen

Ion Conductors Modeling the Varistor Behavior. Electrochem. Solid-State Lett. 8, E67–E69

(2005).

15.Larfaillou, S., Guy-Bouyssou, D., Cras, F. L. & Franger, S. Characterization of Lithium Thin Film

Batteries by Electrochemical Impedance Spectroscopy. ECS Trans. 61, 165–171 (2014).

16.Bates, J. B. et al. Electrical properties of amorphous lithium electrolyte thin films. Solid State Ion.

53–56, Part 1, 647–654 (1992).

17.Yu, X., Bates, J. B., Jellison, G. E. & Hart, F. X. A Stable Thin- Film Lithium Electrolyte:

Lithium Phosphorus Oxynitride. J. Electrochem. Soc. 144, 524–532 (1997).

18.OHARA INC. : Lithium-ion Conducting Glass-ceramics (LICGC) :Development Products.

Available at: http://www.ohara-inc.co.jp/en/product/electronics/licgc.html. (Accessed: 4th April

2016)

19.Stevens, P., Toussaint, G., Puech, L. & Vinatier, P. Very High Specific Area Lithium-Air Battery.

ECS Trans. 50, 1–11 (2013).

20.Stevens, P. et al. Development of a Lithium Air Rechargeable Battery. ECS Trans. 28, 1–12

(2010).

Chapter 3.

Chemical and physical characterizations of block

copolymer electrolytes

Abstract

This chapter is dedicated to the introduction of solid polymer electrolytes used

in this study. All polymer electrolytes are built up by polystyrene block (simple or

multifunctional) and poly(ethylene-oxide) block Three different types of block

copolymer are used, a neutral diblock copolymer laden with lithium salts, a

diblock single-ion copolymer and a triblock single-ion copolymer. Details of the

preparation of thin membranes, but also of their characterization by several

techniques such as differential scanning calorimetry, small angle X-ray scattering,

scanning electron microscopy, conductivity and transference number will be

presented and discussed hereafter.

Chapter 3. Chemical and physical characterizations of block copolymer

- 94 -

Table of contents

Chapter 3. Chemical and physical characterizations of block copolymer electrolytes ..........93

1. Block copolymer : presentation and preparation ........................................................................... 95

2. Thermodynamical properties of block copolymer electrolytes ................................................... 97

3. Morphology studies ............................................................................................................................ 99

a. Small angle X-ray scattering ................................................................................................................ 99

b. Dark field scanning transmission electron microscopy.......................................................................... 108

4. Material electrical properties ........................................................................................................... 109

a. Cell preparation ................................................................................................................................ 109

b. Cell optimization............................................................................................................................... 110

c. Conductivity measurements ................................................................................................................ 111

5. Transference number ....................................................................................................................... 118

6. Conclusion ......................................................................................................................................... 123

References of chapter 3 .............................................................................................................................. 124


- 95 -

1. Block copolymer : presentation and preparation

The most common architectures used for Block Copolymer Electrolytes (BCE) are AB

diblock and BAB triblock copolymers, where the A block is the ionic conductor block and

the B block provides other functionalities, such as good mechanical properties. In this study,

three types of BCEs are studied. The first one is a neutral AB diblock copolymer made of

poly(ethylene-oxide) (PEO) block and poly(styrene) (PS) block, laden with LiTFSI salts. The

polymers used in this study are called SEO xx-yy, where xx and yy are the molecular weights

of the PS, MPS, and PEO, MPEO, blocks in kg.mol-1, respectively. The second BCE studied is

an anionic AB diblock copolymer PEO-b-PSTFSI (called Si-EO). The A block is composed

of PEO, when the B block is composed of TFSI-Li anion covalently attached to the PS

block1. The last one is an anionic triblock copolymer PSTFSI-b-PEO-b-PSTFSI (called Si-

EO-Si). Both anionic BCEs are structurally Li+ single-ion conductors, because the counter

anion is covalently bound to the PS block, which leaves only lithium ions mobile in the

polymer.

SEO 55-52 is synthesized by the Balsara's Group as described in Singh et al 2, whereas

SEO 386-300 is purchased at Sigma Aldrich. Anionic diblock copolymer AB and triblock

copolymers BAB are synthesized by anionic polymerization as described in Bouchet et al3.

The relevant properties of the copolymers used in this study are shown in Table 1. The

calculations used to have the ratio EO/Li and the volume percentage of BCE are clarified

hereafter.

For SEO BCE :

EO

Li=!

mEOMEO"

mLiTFSIMLiTFSI"

(1)

%Vol!=!mEO

dEO+!mLiTFSI

dLiTFSImEO

dEO+!mLiTFSI

dLiTFSI+!mPS

dPS

(2)

And for single-ion BCE:

%Vol=!mEO

dEOmEO

dEO+!mPSTFSI

dPSTFSI

(3)

with a first approximation that the density of PSTFSI-Li is the average between the

density of PS and the density of LiTFSI, thus: dPSTFSILi=! dPS+dLiTFSI2.


- 96 -

Acronym Structure MPEO

(103 kg.mol-1)

MPS or MPSLiTFSI

(103 kg.mol-1)

Mntotal

(103 kg.mol-1)

EO/Li Wt% of PS or

PSLiTFSI

%vol of

PEO

SEO xx-yy

52 300

55 386

107 686

11.8 11.8

51.4 56.3

Si-EO

25

23.75

48.725

7.6

48,7

42.7

Si-EO-Si

35

2 * 6.8

48.6

18.8

28

72

Table 1. Molecular weights and compositions of the different copolymers used in this study.

The preparation of polymer-lithium salt complexes (SEO) is carried out by adding a

calculated amount of LiTFSI salt into the polymer dissolved in DMF solvent in order to

have the desired ratio EO/Li. The mixture is stirred overnight at 90°C and then, cast onto a

smooth nickel foil via a doctor blade. The films of SEO BCE are then dried overnight inside

an argon glovebox antechamber into a vacuum oven at 90°C, in order to remove any traces

of solvent and water.

The mass of lithium salt added (mLisalt) depends of the mass of EO units (mEO) which is

related to the mass of the whole polymer (mBCE).

mEO = fm . mBCE (4)

The mass of LiTFSI is expressed in function of mEO according to:


- 97 -

mLisalt!=(fm!.!mSEO!.!MLiTFSI)

(EO/Li.!MEO) (5)

With :fm, the weight percent of EO units, MEO = 44 g.mol-1 and MLiTFSI = 287,087 g.mol-1.

The salt concentration chosen in this study (EO/Li = 11.7) is found to maximize the

conductivity in SEO electrolytes4, moreover this salt concentration is found to be in the

crystallinity gap5.

Both Si-EO and Si-EO-Si BCEs are dried at 90°C in the glovebox antechamber into a

vacuum oven overnight. In order to get free standing thin films, they are then melt pressed

between two teflon films at 70°C inside the argon filled glovebox.

All films are stored in a dry-argon glove box. All sample preparation and cell assembly

are carried out inside a glove box filled out of argon (MBraun, O2 <8 ppm and H2O <

0.1ppm).

2. Thermodynamical properties of block copolymer electrolytes

The thermodynamic properties of BCEs are studied by differential scanning calorimetry

(DSC) using a Thermal Analysis Q200 calorimeter. Samples are hermetically sealed in

aluminum pans inside an argon glovebox. The measurement method relies on the fact that

the initial state and thermal history of a polymer-lithium salt complex has an impact on

thermodynamic properties measured6. Therefore, in order to obtain a similar thermal history

for all samples, the temperature is increased a first time until 150°C, i.e. above Tm(PEO) and

Tg(PS), and then it is decreased to -90°C, the next cycle of heating and cooling is then realized

to obtain an exploitable thermogram. DSC experiments are run with 10°C.min−1 heating

rates and 2°C.min−1 cooling rates. The temperature ranged from -90ᵒC to 150°C. Melting and

glass transition temperatures are obtained from the analysis of the second heating run.

The thermogram for the second heating run of the sample SEO 386-300 is presented in

Figure 1. The melting temperature (Tm) of the polymer-salt complex is taken at the

intersection between the baseline and the tangent at the inflexion point of the endothermic

melting peak (Tonset). The melting enthalpy (∆Hm) is given by the surface area under the

melting peak. The glass transition of PEO and PS are calculated according the two insets in

Figure 1. In the case of SEO BCE, both PS and PEO glass transitions are observed.


- 98 -

Figure 1. Thermogram of the second heating cycle of SEO 386-300-LiTFSI complex.

By determining the melting enthalpy, it is possible to calculate the degree of crystallinity

(Xc) 7 of the polymer-salt complex according to equation (6):

Xc=!#H(m)

fEO!.!#H$(m) (6)

With : ∆H(m) the melting enthalpy determined by DSC (in J.g-1), fEO the weight percent of

PEO in the BCE and ∆Hᵒ(m) the standard melting enthalpy for a 100% crystallized PEO 8,9,

that is to say 195 J.g-1 .

The melting temperature, the melting enthalpy, the calculated degree of crystallinity, and

the glass transition temperatures for PEO and PS are given in Table 2. Nevertheless,

determining the Tg is not trivial 8 and this parameter could not be determined for all samples.

Indeed, the intensity of the Tg signal depends of the amorphous phase percentage and in the

case of the anionic BCE, the crystalline phase is predominant.


- 99 -

Sample

Tm

°C

ΔHm

J.g-1

Degree of crystallinity

%

Tg (PEO)

°C

Tg (PSLiTSFI)

°C

SEO 386-300 33 ± 1 29.8 38 -38 ± 1 102 ± 1

SEO 55-52 24 ± 1 7.334 7 -41 ± 1 107 ± 1

Si-EO 55 ± 1 84.6 84 - -

Si-EO-Si 54 ± 1 66.1 47 - -

Table 2. Thermodynamic properties determined by DSC for the different BCE studied.

The PEO glass transition is observed at -38ᵒC and -41ᵒC for both SEO BCEs, which is

in good agreement with values obtained for homo PEO5. The Tg of PS is observed at 102ᵒC

and 107ᵒC for both SEO BCE.

3. Morphology studies

In this section, morphology studies will be focused on single-ion BCEs since SEO

morphologies have already been studied during the last two decades. It is well known that

poly(styrene) (PS) and poly(ethylene oxide) chains are highly incompatible, which leads to a

microphase separation10. Moreover, it has been shown that the tendency for microphase

separation is enhanced by the presence of ions in neutral BCE11. Singh et al 2 have studied the

morphologies of SEO BCE loaded with LiTFSI salt, by small angle X-ray scattering (SAXS)

and transmission electron microscopy (TEM). They have determined that SEO (74-98), (40-

54) and (16-16) present a simple lamellar morphology. In other words, the BCE with a PEO

volume fraction ranged from 45 to 55% present a lamellar morphology.

In this section, SAXS measurement as a function of temperature will be discussed for

both anionic single-ion BCEs. The morphological characterization of the samples is then

performed based on dark field scanning electron microscopy (STEM). The domain spacing is

determined by both techniques and are compared and discussed.

a. Small angle X-ray scattering

Small angle X-ray Scattering is a powerful technique for examining meso-structures and it

is especially adapted to analyze the self-assembly of BCEs. For single ion BCEs, such as

PSTFSI-b-PEO diblock copolymers, recent studies have shown a microphase separation12.

However, the presence of PSTFSI-Li instead of PS in the BCE suggests a better

compatibility with the PEO block. Indeed, the LiTFSI is highly compatible with PEO, thus

we can suppose that it will favor the miscibility of the PSTFSI-Li block with the PEO block.


- 100 -

In this section, phase behavior of single-ion BCE will be studied in a wide range of

temperature using SAXS.

Experimental protocol

Polymer films for X-ray scattering experiments are produced by a melt pressing method.

A small amount of BCE is melt-pressed inside two Teflon (fluoropolymer) films using a

custom-made hand-press heated up to 70°C. Finally, films are sealed into a home-made

airtight holders with Kapton® windows. The samples are then annealed at 90°C under

vacuum for at least 24 hours and they are slowly cooled down to room temperature.

Thicknesses ranges from 958 to 974 µm. A blank cell with Kapton® windows is produced

for signal treatment after the experiment. Samples are then placed into a custom-built 8-

samples heating stage. The SAXS experiment is performed from room temperature to 70ᵒC

and the samples are annealed at each temperature for 20 minutes before taking

measurements. The temperature of the heating stage is controlled using a Walthow EZ zone

controller and monitored using a thermocouple K.

Measurements are made at Beamline 7.3.3. at the Advanced Light Source at the

Lawrence Berkeley National Laboratory (LBNL)13. Silver behenate is used as a standard to

determine the beam center and sample-to-detector distance. Scattering patterns are reduced

using the Nika program for Igor Pro available from Jan Ilavsky at Argonne National

Laboratories14 in order to produce one-dimensional (1D) scattering profiles. The background

scattering from the Kapton® cell is subtracted from the reduced SAXS data. The azimuthally

averaged scattering intensity, I, is reported as function of the magnitude of the scattering

vector, q, where q = (4π/λ)sin θ, with λ the wavelength and 2θ the scattering angle.

Results and discussion

SAXS intensities, I, as a function of magnitude of the scattering vector, q, of Si-EO and

Si-EO-Si BCEs, obtained during heating run, are shown in Figure 2 a) and b), respectively.

The location of the primary scattering peak q1* is indicated in the graphs as well as q2 and q3.

The domain spacing d of the copolymers can be determined according to the equation (7):

d = 2π / q* (7)

In Figure 2 a), the primary scattering peak is difficult to locate due to the low-q up-turn.

Similar up-turns have already been observed before in other charged BCEs15.


- 101 -

a) b)

Figure 2. Scattering data are shown vertically offset for clarity. SAXS intensity versus the magnitude of the

scattering vector, q, from 25.9°C to 70°C for a) PEO-b-PSTFSI BCE and b) PSTFSI-b-PEO-b-PSTFSI BCE.

In order to locate more precisely q*, the baseline is subtracted from the scattering profile

using the TN020 procedure available for Igor Pro software. A polynomial function in 1/q2 is

employed as a baseline. Similar data treatment is performed onto the Si-EO-Si scattering

profiles. A treated scattering profile with baseline correction is given in Figure 3 in the case

of Si-EO-Si BCE. In addition, the graph shows how different parameters are determined

such as the maximum intensity, Imax and the width at half maximum.

After baseline correction, a fourth order peak is observed in the scattering profile in the

case of Si-EO-Si BCE (see in Figure 3). However, such peak could not been determined in

the case of Si-EO BCE due to the low q-up-turn which makes the baseline correction very

difficult to produce accurately at low q.


- 102 -

Figure 3. Intensity versus magnitude of the scattering vector with baseline correction for Si-EO-Si BCE at 25.9ᵒC.

In order to enhance the different peak in the scattering profile, another graph is drawn

by representing the square of the scattering vector multiply by the intensity, Iq2, versus q in

the case of the Si-EO-Si BCE (Figure 4). Due to the low q-up-turn in the case of the Si-EO

BCE, such representations did not enhance the peak position. The scattering profile are

shown vertically offset for clarity.

0.14

0.12

0.10

0.08

0.06

0.04

0.02

0.00

-0.02

Inte

nsity (

a.u

)

1.61.41.21.00.80.60.40.2

q (nm-1

)

25.9°C with baseline correctionImax

Half value width

q1*

q2

q3

q4


- 103 -

Figure 4. Intensity multiply by square of the scattering vector versus scattering vector for the Si-EO-Si BCE.

At room temperature, scattering profiles suggest the presence of a lamellar morphology

in both cases, with a primary peak located at q* = 0.24 ± 0.01 nm-1 for the Si-EO and q* =

0.23 ± 0.01 nm-1 for Si-EO-Si and higher-order peaks on the vicinity of 2q*, 3q* and 4q*. The

location of the higher-order peak are not in a perfect agreement with the expected locations

for a lamellar morphology. The domain spacing, d, defined as the center to center distance

between adjacent PEO rich lamellae is 26 ± 2 nm and 27 ± 2 nm for Si-EO and Si-EO-Si

respectively. Interestingly, the domain spacing for both BCEs are found to be very similar,

which is probably due to the small difference in the total molecular weight for both BCEs .

In Figure 2 b), peak intensities are found to be higher suggesting a higher segregation of the

two phases in the case of the Si-EO-Si BCE16.

A recent study by Rojas et al.17 has shown a domain spacing of 27.8 nm for similar single

ion BCEs composed of PEO-b-PSTFSI but with block molecular weights: Mw(PEO) = 5

kg.mol-1 and Mw (PSTFSI) = 3 and 4 kg.mol-1. This results are surprising given the difference in

Mw in our study, which are Mw (PEO)= 25 and 35 kg.mol-1 and Mw (PSTFSI) = 23.75 and 2*6.8

kg.mol for Si-EO and Si-EO-Si BCE respectively.

4

5

6

78

1

2

3

4

5

6

78

10

2

3

4

5

6

78

100

log

Iq

2(a

.u.)

1.51.00.5

q (nm-1

)

Triblock single ion 25.9°C Triblock single ion 30°C Triblock single ion 35°C Triblock single ion 40°C Triblock single ion 45°C Triblock single ion 50°C Triblock single ion 55°C Triblock single ion 60°C

q1*

q2

q3q4


- 104 -

The different scattering peak values as a function of temperature are collected for Si-EO

and Si-EO-Si.

Temperature (ᵒC) q1* (nm-1) q2/q1 q3/q1 q3/q2

25.9 0.238 2.42 3.70 1.53

30 0.235 2.45 3.71 1.51

35 0.231 2.47 3.71 1.50

40 0.226 2.49 3.71 1.49

45 0.208 2.54 3.82 1.50

50 0.187 2.74 4.42 1.61

55 0.171 2.83 4.64 1.63

60 0.183 2.82 4.75 1.68

Table 3. Primary scattering peak and scattering peak ratio in function of temperature for Si-EO BCE.

Temperature (ᵒC) q1* (nm-1) q2/q1 q3/q1 q4/q1 q3/q2 q4/q2 q4/q3

25.9 0.229 2.69 4.42 6.31 1.64 2.36 1.43

30 0.228 2.70 4.43 6.30 1.64 2.34 1.42

35 0.221 2.77 4.56 6.39 1.65 2.31 1.40

40 0.228 2.64 4.38 6.18 1.66 2.34 1.41

45 0.213 2.78 4.63 6.65 1.67 2.39 1.44

50 0.203 2.81 4.74 6.76 1.69 2.40 1.42

55 0.189 2.86 4.91 7.00 1.71 2.44 1.43

60 0.183 2.77 4.70 6.72 1.69 2.42 1.43

Table 4. Primary scattering peak and scattering peak ratio as a function of temperature for Si-EO-Si BCE.

Table 3 and Table 4 present the primary scattering peak q1* and the scattering peak ratio

for the Si-EO and Si-EO-Si respectively, as a function of temperature. The scattering peak

ratios are found to be different from the theoretical ratio of 2 and 3 for q2/q1 and q3/q1

ratios respectively in the case of lamellar morphology in both BCEs. Rojas et al17 attributed

this disparity to the complexity of order formation in the presence of crystallization18.

Besides no simple relation is found between q1* and the higher order scattering peaks.

The SAXS signal in semi-crystalline polymers is due to the contrast in electronic density

between crystalline and amorphous domains. A recent study has concluded that the ordering

formation in diblock single-ion copolymer is driven by the crystallization of PEO and that

the PSTFSI block is accommodated within the amorphous phase17. As expected, here we

obtain similar results since the scattering peaks disappear above 60ᵒC, i.e. above the PEO

melting temperature. Interestingly, the first peak seems to disappear more abruptly than the

q2 and q3 peaks.


- 105 -

An interesting phenomenon is observed in the SAXS experiment, the scattering peak

values decrease with increasing temperature. This evolution in the q value as a function of

temperature is presented in Figure 5 a) and b) for Si-EO and Si-EO-Si BCEs, respectively.

The scattering peak value evolution is found to be similar in both BCEs cases. In this

section, the three different scattering vectors q1, q2 and q3 are considered as three

independent primary peaks. For more clarity in the graph, the scattering peaks are

normalized by the qi (with i = 1, 2 or 3) obtained at 25.9ᵒC. When temperature increases q2

and q3 are found to be relatively constant at least up to 50ᵒC, and q1 is found constant only up

to at least 40ᵒC. In addition, q1 shows a higher decrease with temperature. Both q2 and q3

appeared to have a similar behavior with temperature, whereas q1 seems decorrelated from

them (see in Figure 5 a) and b)). Each q shows a primary behavior before an abrupt drop. By

drawing a tangent for each slope, we could determined at the intersection a transition

temperature for each q. It is important to notice that the transition temperature for q1, which

is found at 42ᵒC and 40ᵒC for Si-EO and Si-EO-Si BCE respectively, is distinct from the

transition temperature for q2 and q3 54ᵒC and 48ᵒC for Si-EO and Si-EO-Si respectively. In

the case of the Si-EO BCE the transition temperature for q2 and q3 is very close to the Tm.

a) b)

Figure 5. Evolution of the scattering vector value q normalized as a function of temperature for a) Si-EO BCE and b) Si-EO-Si BCE.

Three different domains spacing are calculated from the three scattering vectors and

plotted versus temperature in Figure 6 a) for the Si-EO BCE. In the case of the Si-EO-Si

BCEs, four domains spacing could be calculated from the scattering vectors and are


- 106 -

presented as a function of temperature in Figure 6 b). The d1 for the first scattering peak,

which is considered as the domain spacing for the lamellar morphology, is constant up to

40ᵒC with a value oscillating around 26 nm and 27 nm for Si-EO and Si-EO-Si BCE

respectively. The domain spacing then becomes larger up to 44.5 nm and 33 nm for Si-EO

and Si-EO-Si BCE respectively at 60ᵒC. If we suppose that this increase in d1 corresponds to

a dilatation of the lamellae, from room temperature to 60ᵒC, it would corresponds to a

dilatation of 68% and 21%, for Si-EO and Si-EO-Si respectively, which is not physically

possible. For d2 and d3, we calculated a dilation of 36% and 21% for SI-EO BCEs and 21%

and 17% for the Si-EO-Si BCEs. Therefore, the dilatation of the lamellae is not a reasonable

explanation for the evolution of the scattering peak and thus for the evolution of the domain

spacing.

a) b)

Figure 6. Evolution of the domain spacing in function of temperature for a) Si-EO BCE and b) Si-EO-Si BCE.

Heating the sample results in a decrease in peak intensities, the normalized intensity for

the primary peak, which is the peak corresponding to the lamellar morphology, is shown in

Figure 7 for both BCEs. The variations of peak intensity for both electrolytes are found to

be very similar. Indeed, in both cases, a drop of intensity is observed above 55°C with a

complete disappearance of the peak above 60ᵒC. This order to disorder transition (ODT)

seems therefore correlated to the melting of the PEO crystallites as it has been also observed

for low molecular weight Si-EO by Rojas et al17. In addition, the small intensity of the

scattering peak above the melting temperature implies an order loss.


- 107 -

Figure 7. Normalized SAXS intensity for the primary scattering peak versus temperature for Si-EO and Si-EO-Si BCEs.

In the literature, the ODT for block copolymers with and without lithium salts has been

extensively studied11,19,20. Typically, the ODT is accompanied by a strong increase in the

width of the primary scattering peak, this broad peak is representative of the large amplitude

concentration fluctuations. However, in the disordered single-ion BCEs (above 60ᵒC) the

primary SAXS peak is absent (see in Figure 2), indicating at the same time the absence of

concentration fluctuations. The same observations have been reported in the literature for

similar BCEs12.

The full widths at half maximum (FWHM) are determined for the primary peak in both

BCE cases and are presented as a function of temperature in Figure 8 a) and b). Increasing

the temperature results in a refinement of the peak with a decrease in the width at half

maximum height, which generally corresponds to a smaller distribution in the size of the

nano-structured domains and/or a better long range ordering.


- 108 -

a) b)

Figure 8. Half (value) width at half maximum for the primary scattering peak a) diblock single ion BCE and b) triblock single ion BCE.

Here, we report that both single ion BCEs exhibit a lamellar morphology for temperature

below 60ᵒC. The domain spacing calculated from the primary scattering peak are 26 ± 2 nm

and 27 ± 2 nm for Si-EO and Si-EO-Si respectively. In addition, for temperature above the

melting temperature, BCE presents a loss in the order of the micro-phase separation.

b. Dark field scanning transmission electron microscopy

In order to confirm the BCE lamellar morphology, dark field transmission microscopy

(STEM) are performed on the samples. They are stained with ruthenium tetraoxide (RuO4)

vapor for 10 minutes. STEM are performed on a Tecnai F20 UT FEG equipped with a high

angle annular dark field (HAADF) detector using 200 keV acceleration voltage. The bright

phase represents PEO-rich lamellae21.

The micrograph of PEO-b-PSTFSI is shown in Figure 9 a). Lamellar morphology is

confirmed and the domain spacing between two PEO rich lamellas is calculated to be d = 36

± 3 nm, which is in reasonable agreement with the scattering data (dSAXS = 26 ± 2 nm).


- 109 -

a) b)

Figure 9. Dark field scanning transmission electron micrograph of a) PEO-b-PSTFSI BCE and b) PSTFSI-b-PEO-b-PSTFSI BCE.

A micrograph of PSTFSI-b-PEO-b-PSTFSI is shown in Figure 9 b). Lamellar

morphology is also confirmed in this case and the domain spacing between two PEO rich

lamellae is calculated to be d = 25 ± 3 nm, which is in very good agreement with the

scattering data (dSAXS = 27 ± 2 nm). The triblock copolymer exhibits smaller dark lamellae

compared to the diblock which seems to be coherent with the Mw of the PSTFSILi blocks.

4. Material electrical properties

a. Cell preparation

Ionic conductivity measurements are performed in lithium symmetric cells. Films of

block copolymer electrolyte (BCE) are produced by a melt pressing method. A small amount

of BCE is melted inside two foils of Teflon (fluoropolymer) using a custom-made hand press

heated to 70°C. The membrane is then pressed onto a Kapton® spacer, with a 1/8 inch hole

diameter using the same method, in order to have a well-known geometry of the sample.

Typical thicknesses of the membrane ranged from 50 µm to 100 µm. Lithium metal chips of

150 µm thick electrodes and nickel tabs are used to assemble symmetric cells inside an argon

filled glove box. Cells are then vacuum sealed (Showa-Denka) in an air tight pouch material

in order to be able to carry out the experiment outside the glove box. Figure 10 shows a

typical lithium symmetric cell assembly.


- 110 -

Figure 10. Ionic conductivity assembly for lithium polymer symmetric cell.

b. Cell optimization

In order to maximize the lithium-polymer interface in lithium symmetric cells, the

optimization of the surface state of lithium metal chips is made. Lithium chips are purchased

from MTI corporation, however stripes are visible on the whole surface of both sides of

lithium chips. Those heterogeneities come from the manufacturing process of lithium and

the lithium surface state is not smooth and featureless. Therefore, different pre-treatments

are performed on the lithium chips in order to suppress those asperities and to determine the

best treatment. Three lithium symmetric cells are assembled, one with lithium chips without

any treatment, another with lithium pressed at room temperature with a press set up at 3 bar

for 5 seconds and a last one assembled with lithium pressed at 90ᵒC for 5 second at 3 bar.

Here, we will focus on both the electrolyte and the interface resistances, which will be

affected by the pre-treatment of the lithium chips.


- 111 -

Figure 11. Nyquist plots of SEO 386-300-lithium symmetric cell at 90ᵒC assemble with lithium chips with different pre-treatment.

We present the results, for three different cells of SEO386-300-lithium symmetric cells

with thickness of 75 μm, in Figure 11. The interface resistance in the case of lithium chips

pre-pressed at 90ᵒC is smaller than in the two other cases, however the electrolyte resistance

is high which corresponds to 1.54.10-4 S.cm-1 of ionic conductivity. In the case of the lithium

chips pre-pressed at RT, the interface resistance is around 900 Ω but the electrolyte

resistance corresponds to an ionic conductivity of 3.2.10-4 S.cm-1. Therefore, we choose to

pre-pressed the lithium at RT before the experiment.

c. Conductivity measurements

Conductivity experiments are carried out using a Bio-Logic VMP3 (multichannel

potentiostat). Cells are placed into a homemade heating stage, where measurements are

carried out from 30°C to 90°C. A first heating cycle is carried out from room temperature to

90°C with 10°C steps, then a cooling cycle is performed with the same temperature steps to

30°C and finally a second heating cycle finalized the experiment. Ionic conductivities are

determined from impedance spectroscopy measurements. An excitation signal of 50 mV is

applied from 1MHz to 1Hz. An impedance spectroscopy measurement is performed on the

cell at every temperature steps after temperature stabilization.

1600

1400

1200

1000

800

600

400

200

0

- Im

(Z

) (W)

160012008004000

Re (Z) (W)

Lithium unpressed Lithium pressed at RT Litihum pressed at 90°C


- 112 -

Figure 12 shows a typical spectrum obtained in a Nyquist plot representation for a

lithium/SEO 386-300/lithium cell at 90°C. All spectra are composed of a first semi-circle at

high frequency (f > 105 Hz) which corresponds to the bulk of the electrolyte that enables the

electrolyte resistance Rel to be determined and the ionic conductivity of the electrolyte (σel) is

calculated according to equation (8):

&el=!lel

S!.!Rel (8)

With: lel: electrolyte thickness (in cm) and S: surface area of the electrolyte film (in cm2).

In the middle frequency domain (from 105 to 102 Hz), another loop is observed in all

spectra. This is generally attributed to the formation of surface layers22–24, the well known

SEI25 (Solid Electrolyte Interphase) which forms at the interface between the polymer

electrolyte and lithium. This interface resistance is called here Rint. Another contribution is

present at low frequencies (f < 10 Hz), it is characteristic of the transport by diffusion in the

bulk of electrolyte and modeled by a short Warburg element (see in Figure 12)24.

Figure 12. Typical Nyquist plot obtained from a lithium/SEO 386-300/lithium cell at 90°C, the equivalent circuit used to fit the data is given in inset.

The SEO 55-52 BCE presents a similar spectrum at 90ᵒC and is given in Figure 13.

However, the main difference lies in the interface resistance. Indeed, the R int is very high

(around 7000 Ω.cm2 compared to around 120 Ω.cm2 for SEO 386-300). This isolating


- 113 -

interface could arise from the wrong orientation of the lamellae due to an increase of PS

percentage. However, this BCE presents 51.4% of PS compared to 56.3% in SEO 386-300,

therefore we should have a better interface in SEO 52-55 due to a higher amount of PEO

potentially at the interface. Thus, this is not the explanation for the high Rint. Furthermore,

the characteristic frequency of the interface lithium/polymer here is 18 Hz, which is two

orders of magnitude lower than for SEO 386-300 (2.4 kHz). Therefore, we can suppose that

the chemistry of the interface in the SEO 52-55 differs from the SEO 386-300, meanings

that impurities are still present in this BCE and form an isolating interface layer with lithium.

Figure 13. Nyquist plot obtained from a lithium/SEO 52-55/lithium cell at 90°C.

All impedance spectra are modeled via the equivalent circuit represented in the inset in

Figure 12. It is assumed that the different phenomena are in series from an electrical point of

view. An inductance (Lc) in series with a resistance (Rc) described the response due to the

electrical connectors24. For the electrolyte bulk and the lithium/polymer interface a resistance

in parallel with a constant phase element (CPE) circuit are used. In this section, the

impedance measurements are not carried out at very low frequencies (from 1Hz to 10-3 Hz),

therefore points representing the diffusion are not taking into account.

The ionic conductivities of SEO BCEs are presented in an Arrhenius plot in Figure 14.

The error bars correspond to the distribution obtained in three cells, the ionic conductivities

are measured at the second heating cycle. For SEO 386-300, ionic conductivities are found

ranging from 3.10-6 S.cm-1 at 20°C to 2.5.10-4 S.cm-1 at 90°C. Whereas SEO 55-52 showed

higher ionic conductivities ranging from 8.7.10-6 S.cm-1 at 20°C to 3.2.10-4 S.cm-1 at 90°C.

7000

6000

5000

4000

3000

2000

1000

0

- Im

(Z)

(W.c

m2)

6000400020000

Re(Z) (W.cm2)

18 Hz

100

80

60

40

20

0

- Im

(Z)

(W.c

m2)

100806040200

Re(Z) (W.cm2)

0.3 MHz


- 114 -

This result is in good agreement with literature, Panday et al.4 have reported that the ionic

conductivity increases with increased molecular weight and reaches a plateau for MPEO > 60

kg.mol-1. The ionic conductivities increased with temperature smoothly, which confirms the

low melting temperature and a low degree of crystallinity3 measured by DSC previously (see

chapter 3. 2). In fact, SEO 386-300 presented a melting temperature at 33°C and SEO 55-52

a melting temperature of 24 °C. Furthermore, it is important to remind that in neutral BCE

there is a "dead zone" at the near interface of PS block26 and that this "dead zone" represents

4-5 EO units and that it is not dependant of molecular weight or EO/Li.

Figure 14. Plots of conductivity as a function of inverse temperature for the two SEO BCEs.

Figure 15 shows a typical spectrum in Nyquist coordinates for the Si-EO BCE in lithium

symmetric cell at 90°C. Similar contributions are observed compared to SEO electrolytes.

However, as expected for a single-ion at low frequency no additional contribution for the

diffusion is visible. In addition, the characteristic frequency for the polymer-lithium interface

is found to be considerably different compared to the SEO copolymers, 2.4 kHz for SEO

386-300 against 422 Hz for Si-EO. This difference can be explained by a difference in the

chemistry of the interface for the anionic BCEs compared to SEO BCEs. Moreover, the

10-6

2

4

6

810

-5

2

4

6

810

-4

2

4

6

810

-3

Con

du

ctivity (

S.c

m-1

)

3.53.43.23.02.92.8

1000/T (K-1

)

90

SEO 386-300 SEO 55-52

7080 60 50 40 30 20

Temperature (ºC)

90


- 115 -

interface resistance is substantially high, about 2200 Ω.cm2, when for SEO 386-300

electrolytes Rint is about 120 Ω.cm2.

Figure 15. Typical Nyquist plot obtained from a lithium/PEO-b-PSTFSI/lithium cell at 90°C.

Figure 16 shows a typical spectrum in Nyquist coordinates for Si-EO-Si-lithium

symmetric cells at 90°C. Again, two contributions are observed, at high frequency (0.24

MHz) the electrolyte contribution and at middle frequency (4.2 kHz) the interface

contribution. The characteristic frequency for the interface in the case of the Si-EO BCE and

the Si-EO-Si BCE are found to be very different, which suggests a difference in the

chemistry of the SEI. Moreover, in the case of the Si-EO-Si the interface resistance is about

50 Ω.cm2 compared to 2200 Ω.cm2 for the diblock copolymer. It is worth noting that the Si-

EO BCE (diblock) exhibits an orange color, which is probably due to remaining impurities

from the synthesis. This is probably the origin of the high interface resistance observed.


- 116 -

Figure 16. Typical Nyquist plot obtained from a lithium/PSTFSI-b-PEO-b-PSTFSI/lithium cell at 90°C.

The evolution of the ionic conductivities for single ion BCEs are shown in Arrhenius

coordinates in Figure 17. For the Si-EO BCE, conductivities ranged from 3. 10-9 S.cm-1 at

40°C to 2.5 .10-5 S.cm-1 at 90°C and for the Si-EO-Si they are higher (3 times), ranging from

1.10-8 S.cm-1 at 40°C to 4.10-5 S.cm-1 at 90°C.


- 117 -

Figure 17. Plots of conductivity and the normalized SAXS intensity as a function of inverse temperature for the two anionic BCE.

Moreover, a drop in conductivity of three orders of magnitude is observed for both

anionic single-ion BCEs due to PEO crystallization. Above the Tm, the conductivity rises

strongly indicating a strong increase in mobile charges concentration.

The normalized SAXS intensity obtained in section chapter 3.3.a) is plotted in Figure 17

with the ionic conductivities. A correlation between phase separation, ordering and

conductivity is highlighted. Indeed, the high ionic conductivities values are obtained for

disordered morphology seen by SAXS, when the normalized intensity reached zero whether

for temperature above 60°C. This result is consistent with previous study 12,17, where they

have shown the same correlation between SAXS experiment and in situ ionic conductivity

measurements (presented in Figure 18).

10-9

10-8

10-7

10-6

10-5

10-4

Co

nd

uctivity (

S.c

m-1

)

3.33.23.13.02.92.82.7

1000/T (K-1

)

1.0

0.8

0.6

0.4

0.2

0.0

No

rma

lize

d S

AX

S in

ten

sity

(a.u

.)

s PSTFSI-b-PEO-b-PSTFSI

s PEO-b-PSTFSI Inormalized PSTFSI-b-PEO-b-PSTFSI

Inormalized PEO-b-PSTFSI

7090 80 60 50 40

Temperature (Cº)


- 118 -

Figure 18. Scheme to explain the order to disorder morphology17.

Thus, in this system the conductivity arises from at least a partial miscibility of the PEO

and PSTFSI-Li above the melting temperature17. In other words, below the melting

temperature, PEO and PSTFSI are well phase separated and the ionic conductivity is rather

low. Indeed, the PSTFSI-Li block is an insulator and the lithium ions cannot be solvated in

this block, this implies that the Li+ ions can only be solvated at the interface PEO/PSTFSI-

Li. Therefore, due to a small amount of lithium ions solvated below the Tm, the ionic

conductivity is very low. However, above the Tm the order decreases implying a partial

miscibility of the two blocks leading to the solvation of Li+ in the PEO blocks resulting in

the increase of ionic conductivity of three order of magnitude.

5. Transference number

Transference number determination is an old and complex problem. The lithium-

polymer-lithium symmetric cells used for transference number measurements are similar to

those described in the section above. For SEO BCEs the steady-state current method is

used, whereas for single ion BCEs low frequency electrochemical impedance spectroscopy is

used. Both methods are applied to cells at 90°C and carried out using a VMP3 potentiostat

from biologic.

However, it worth noting that the two methods are similar. Indeed, according to Mac

Donald27, who made the hypothesis of a binary electrolyte (type LiX) diluted, the

transference number can be calculated following equation (9), where Rel is the electrolyte

resistance, Rdiff is the diffusion resistance.


- 119 -

t+!=!Rel!

(Rel!+!Rdiff) (9)

However, the polarization, ∆V, used in the steady state current method can be calculated

from the initial current I0 and the different resistances Rel and Rint according to equation (10):

Rint!+!Rel!=!#V

I0 (10)

Or using the steady state current I∞ and all the resistances Rel, Rint and Rdiff according to

equation (11):

Rel!+!Rint!+!Rdiff=!#V

I' (11)

Hence, if we substitute the terms in equation (9) by the one obtained using equations

(10) and (11) the transference number can be calculated by equation (12). which is the Bruce

and Vincent equation28. The initial and steady state interfacial resistances, R0int and R∞

int

respectively, are determined by fitting the equivalent circuit (see inset in Figure 12).

t+!=!#VIo!-!Rint

#VI'!-!Rint

=!I'!(#V-I0Rint)

I0!(#V-I'Rint) (12)

The initial current is generally obtain after 1 second and is compared to (Rel + Rint).

However, along the current circulation through the interface the Rint evolutes, thus a right

equation should be:

t+!=!#VIo!-!Rint

#VI'!-!Rint

=!I'!(#V-I0Rᵒint)

I0!(#V-I'R∞int)

(13)

With Rᵒint the initial interface resistance and R∞int the interface resistance taken after the

polarization experiment.

The steady-state current method consisted in applying a constant potential of 20 mV and

the current is measured as a function of time28,29. Impedance spectra are performed every

hour in order to probe the bulk and the interfacial resistances. These values are determined

by fitting an equivalent circuit (see the inset in Figure 12) to the data using Ec-lab software.

The transference number is determined using equation (12).


- 120 -

The current density as a function of time for SEO 386-300 during a 20 mV polarization

experiment at 90°C is shown in Figure 19. The inset shows the ac impedance of the cell

before (turquoise) and after (blue) polarization. The transference number is determined to be

t+ = 0.14 which is rather consistent with literature for SEO BCEs30.

Figure 19. Transference number determination of the SEO electrolyte. Current density as a function of time for

SEO 386-300 at 90°C. The inset shows the nyquist plot of the cell before (turquoise) and after (blue) polarization.

To confirm that the two methods are rather in good agreement the Figure 20 presents

the low frequencies EIS experiment for a lithium-SEO386-300 symmetric cell (pink) and the

normal EIS (blue) at 90ᵒC.

For the low frequencies EIS method, the Mac Donald equation for binary electrolyte in

dilute approximation is used27. It consisted in measuring an electrochemical impedance

spectrum from high frequencies (1MHz) to low frequencies (1mHz) and the transference

number is determined using equation (9). For a material presenting a transference number

different from unity, an additional typical Warburg loop (1/4 lemniscate) at low frequency is

observed24. Usually, the diffusion response is modeled by a Warburg short element.

We observe that additional contributions at low frequencies are observed in the case of

SEO BCE. The transference number is calculated according to Mac Donald equation and t+


- 121 -

= 0.16, it is in good agreement with literature31 but it is slightly different from the value

obtained by the Bruce and Vincent method.

Figure 20. Nyquist plots obtained from a lithium/SEO 386-300/lithium cell at 90°C.

Transference number experiments are performed with Si-EO BCE in lithium symmetric

cell and the resulting impedance spectra are shown in Figure 21. The very first spectrum

showed no additional semi-circle at low frequency. However, during the polarization time the

interface between the polymer and the lithium changed leading to a diminution of the R int.

We attributed this to the rearrangement of the passive layer at the interface lithium-polymer.

Nevertheless, the transference number of the diblock copolymer PEO-b-PSTFSI is

calculated to be 1 due to the absence of a diffusion resistance confirming the single-ion

nature of this BCE.

300

250

200

150

100

50

0

-Im

(Z)

(W.c

m2)

300250200150100500

Re(Z) (W.cm2)

SEO 386-300 low frequencies SEO 386-300


- 122 -

Figure 21. Nyquist plots obtained from a lithium/PEO-b-PSTFSI/lithium cell at 90°C.

Same experiment is performed with Si-EO-Si BCE in lithium symmetric cell at 90ᵒC and

results are shown in Figure 22. A small additional contribution is observed at low

frequencies, therefore the transference number calculated is 0.9. However, it is highly

possible that the small diffusion observed is related to the diffusion of lithium ions inside the

passivation layers. In fact, due to the counter anion covalently bond, the transference

number of the BCE is structurally unity.

Figure 22. Nyquist plots obtained from a lithium/PSTFSI-b-PEO-b-PSTFSI/lithium cell at 90°C.

2000

1500

1000

500

0

-Im

(Z)

(W.c

m2)

2000150010005000

Re(Z) (W.cm2)

PEO-b-PSTFSI low frequency PEO-b-PSTFSI

800

600

400

200

0

-Im

(Z)

(W.c

m2)

8006004002000

Re(Z) (W.cm2)

PSTFSI-b-PEO-b-PSTFSI low frequency PSTFSI-b-PEO-b-PSTFSI


- 123 -

6. Conclusion

We have characterized the thermodynamical, morphological and transport properties of

the different BCEs used in this study. We have determined the Tm and Xc of the different

BCEs. For the neutral BCEs, we can described the decrease of the Tm and Xc by a

confinement effect. Whereas, for the Si-EO and Si-EO-Si BCEs the Tm and Xc are similar

to the one of the homo-PEO.

In addition, we presents a morphology study of the single-ion BCE. For both anionic

single-ion BCE, an ordered lamellar morphology has been confirmed by SAXS and STEM

below 60ᵒC. However, the higher order scattering peak values are not in perfect agreement

(i.e. 2q* and 3q*). The domain spacing at 25ᵒC determined by SAXS are 26 ± 2 nm and 27 ±

2 nm for Si-EO and Si-EO-Si respectively, whereas, by STEM we found 36 ± 3 nm and 25

± 3 nm respectively. Interestingly, the scattering peak values decreases with increasing

temperature, leading to an increase in domain spacing. Moreover, the disappearance of the

scattering peaks with temperature, i.e. the order to disorder transition, seems to be different

for the primary scattering peak value q1 and the higher order q2 and q3.

Ionic conductivity measurements have been performed by EIS. The results for neutral

BCEs are in good agreement with the literature. In the case of the single-ion BCE, a

correlation has been highlighted between the high ionic conductivity and the disordered

phase. Indeed, above the melting temperature, the order loss leading to a partial miscibility

between PEO and PSTFSI-Li blocks results in a strong increase of ionic conductivities.

Transference number for single-ion BCE has been confirmed to be unity, whereas for SEO

BCE, the t+ is equal to 0.15 ± 0.01.


- 124 -

References of chapter 3

1. Meziane, R., Bonnet, J.-P., Courty, M., Djellab, K. & Armand, M. Single-ion polymer electrolytes

based on a delocalized polyanion for lithium batteries. Electrochimica Acta 57, 14–19 (2011).

2. Singh, M. et al. Effect of Molecular Weight on the Mechanical and Electrical Properties of Block


3. Bouchet, R. et al. Single-ion BAB triblock copolymers as highly efficient electrolytes for lithium-metal

batteries. Nat. Mater. 12, 452–457 (2013).

4. Panday, A. et al. Effect of Molecular Weight and Salt Concentration on Conductivity of Block


5. Lascaud, S. et al. Phase Diagrams and Conductivity Behavior of Poly(ethylene oxide)-Molten Salt

Rubbery Electrolytes. Macromolecules 27, 7469–7477 (1994).

6. GRENET Jean & LEGENDRE Bernard. Analyse calorimétrique différentielle à balayage (DSC).

\iTech. Ing. Méthodes Therm. Anal. (2010).

7. FONTANILLE Michel & GNANOU Yves. Structure moléculaire et morphologie des polymères.

(1994). at http://www.techniques-ingenieur.fr

8. Beaumont, R. H. et al. Heat capacities of propylene oxide and of some polymers of ethylene and

propylene oxides. Polymer 7, 401–417 (1966).

9. Beaudoin, E. et al. Effect of Interfaces on the Melting of PEO Confined in Triblock PS-b-PEO-b-PS

Copolymers. Langmuir 29, 10874–10880 (2013).

10. Eitouni, H. B.; Balsara, N. P. In Physical Properties of Polymers Handbook, 2nd ed.; Mark, J. E., Ed

Springer: New York, 2007; Chapter 19, pp 339-356.

11. Teran, A. A. & Balsara, N. P. Thermodynamics of Block Copolymers with and without Salt. J. Phys.

Chem. B 118, 4–17 (2014).



13. Hexemer, A. et al. A SAXS/WAXS/GISAXS Beamline with Multilayer Monochromator. J. Phys. Conf.

Ser. 247, 12007 (2010).

14. Ilavsky, J. Nika: software for two-dimensional data reduction. J. Appl. Crystallogr. 45, 324–328 (2012).


- 125 -

15. Eitouni, H. B. & Balsara, N. P. Effect of Chemical Oxidation on the Self-Assembly of Organometallic

Block Copolymers. J. Am. Chem. Soc. 126, 7446–7447 (2004).

16. Hamley, I. W. & Castelletto, V. Small-angle scattering of block copolymers. Prog. Polym. Sci. 29, 909–

948 (2004).

17. Rojas, A. A. et al. Effect of Lithium-Ion Concentration on Morphology and Ion Transport in Single-

Ion-Conducting Block Copolymer Electrolytes. Macromolecules 48, 6589–6595 (2015).

18. Loo, Y.-L., Register, R. A. & Ryan, A. J. Modes of Crystallization in Block Copolymer

Microdomains:Breakout, Templated, and Confined. Macromolecules. 35, 2365–2374 (2002).

19. Leibler, L. Theory of Microphase Separation in Block Copolymers. Macromolecules 13, 1602–1617

(1980).

20. Khandpur, A. K. et al. Polyisoprene-Polystyrene Diblock Copolymer Phase Diagram near the Order-

Disorder Transition. Macromolecules 28, 8796–8806 (1995).

21. Trent, J. S., Scheinbeim, J. I. & Couchman, P. R. Ruthenium tetraoxide staining of polymers for

electron microscopy. Macromolecules 16, 589–598 (1983).

22. Fauteux, D. Lithium Electrode/PEO-Based Polymer Electrolyte Interface Behavior Between 60° and

120°C. J. Electrochem. Soc. 135, 2231–2237 (1988).

23. Peled, E., Golodnitsky, D., Ardel, G. & Eshkenazy, V. The sei model—application to lithium-polymer

electrolyte batteries. Electrochimica Acta 40, 2197–2204 (1995).

24. Bouchet, R., Lascaud, S. & Rosso, M. An EIS Study of the Anode Li/PEO-LiTFSI of a Li Polymer


25. Peled, E. The Electrochemical Behavior of Alkali and Alkaline Earth Metals in Nonaqueous Battery

Systems—The Solid Electrolyte Interphase Model. J. Electrochem. Soc. 126, 2047–2051 (1979).

26. Bouchet, R. et al. Charge Transport in Nanostructured PS–PEO–PS Triblock Copolymer Electrolytes.

Macromolecules 47, 2659–2665 (2014).

27. Macdonald, J. R. Binary electrolyte small-signal frequency response. Electroanal. Chem. Int. Electrochem.

53, 1–55 (1974).

28. Bruce, P. G. & Vincent, C. A. Steady state current flow in solid binary electrolyte cells. J. Electroanal.

Chem. Interfacial Electrochem. 225, 1–17 (1987).


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29. Doyle, M. & Newman, J. Analysis of Transference Number Measurements Based on the Potentiostatic

Polarization of Solid Polymer Electrolytes. J. Electrochem. Soc. 142, 3465–3468 (1995).

30. Devaux, D., Bouchet, R., Glé, D. & Denoyel, R. Mechanism of ion transport in PEO/LiTFSI

complexes: Effect of temperature, molecular weight and end groups. Solid State Ion. 227, 119–127

(2012).


Mater. 27, 4682–4692 (2015).

Chapter 4.

Dendritic growth in lithium symmetric cells

Abstract

In the previous chapter, the different polymers were characterized. This

chapter investigates the performance of the polymers during cycling, and the

dendritic growth through the polymer electrolyte will be discussed. Two parameters

in particular will be discussed: the electrolyte and the interface resistances, both

evolving during cycling. Finally, the BCE-lithium interface (dendrite) morphology

will be characterized by hard X-ray micro-tomography for both SEO and single-

ion electrolytes.

Chapter 4. Dendritic growth in lithium symmetric cells

- 128 -

Table of contents

Chapter 4. Dendritic growth in lithium symmetric cells ............................................. 127

1. Cycling experiments followed by electrochemical impedance spectroscopy ................................... 129

a. Cycling routine ........................................................................................................................................... 129

b. Neutral block copolymer ............................................................................................................................. 131

c. Single-ion block copolymers ......................................................................................................................... 136

2. Dendrites morphologies studied by hard X-Ray microtomography .................................................. 145

a. Hard X-Ray microtomography ................................................................................................................... 145

b. Protocol ...................................................................................................................................................... 146

c. Neutral block copolymer ............................................................................................................................. 147

d. Single-ion block copolymer .......................................................................................................................... 152

3. Conclusion ................................................................................................................................................... 157

References of Chapter 4 ..................................................................................................................................... 159


- 129 -

1. Cycling experiments followed by electrochemical impedance

spectroscopy

a. Cycling routine

Cells for dendritic growth experiment are prepared following the same assembling

routine than cells for the conductivity experiments. They are galvanostatically cycled using

either a Maccor 4000 Tester Series or a VMP3 potentiostat in either a custom made heating

stage or an oven at 90ᵒC. Cells are allowed to equilibrate at the temperature of interest during

four hours before starting the cycling experiment. Each cell is then submitted to 15 pre-

conditioning cycles that consists of a current density of 0.02 mA.cm-2 imposed in one

direction for 4 hours, followed by 45 minutes of rest, followed by the imposition of a

constant current of 0.02 mA.cm-2 in the opposite direction for 4 hours, followed by another

45 minutes period of rest. They are then cycled at a current density of 0.175 mA.cm-2 with

the same time intervals until the cell shorted or has to be removed due to time constrains.

The thickness of lithium transferred at each cycle at 0.175 mA.cm-2 is 3.13 µm. During each

rest period an EIS measurement is measured in order to follow the evolution of the different

phenomena occurring after charge and discharge.

Figure 1. Scheme representing the evolution of lithium dendrite inside a BCE-lithium symmetric cell and the voltage associated.


- 130 -

The Figure 1 presents the evolution of lithium dendrite inside a BCE-lithium symmetric

cell and the voltage associated during cycling. The stage 1 corresponds to the initial stage and

the beginning of the cycling, the voltage reaches the steady state and stays constant over

cycling. The stage 2 corresponds to the nucleation and growth of lithium dendrites. When

the lithium dendrites cross the whole BCE the stage 3 is reached, short-circuits with a

voltage drop appears. Another phenomenon could be observed, the soft short-circuit, which

is characterized by a reduce voltage compared to the initial one. The last possible

phenomenon is the fuse effect, which is characterized by the "healing" of the voltage after a

voltage drop due to a short-circuit.

A typical cycling routine is shown Figure 2, where the voltage is plotted versus time for a

lithium symmetric cell with SEO 386-300. The cycles represented in pink are the pre-

conditioning cycles at 0.02 mA.cm-2, whereas the cycles in navy blue are the one cycled at

0.175 mA.cm-2. Typical signature of short-circuit is highlighted in Figure 2, but also typical

signatures of soft short-circuits with a reduce voltage and a typical fuse effect1.

Figure 2. Typical cycling routine for a lithium polymer cell at 90ᵒC, voltage versus time. For more clarity the time axis is splitted.

-0.04

-0.02

0.00

0.02

0.04

Voltage (

V)

2001801601401201008060

Time (h)

800750700650600

Cycles at 0.02 mAcm-2 Cycles at 0.175 mA.cm

-2

Short circuitFuse effect

Lithium symmetric cell with SEO 386-300 BCE


- 131 -

b. Neutral block copolymer

During cycling two different parameters will be followed, the electrolyte resistance and

the interface resistance between lithium metal and the polymer. The first one is related to the

ionic conductivity of the electrolyte and the second is related to the Li/electrolyte interface

(SEI). Here, we will present results for two cells assembled with the neutral BCE in lithium

symmetric cells.

a) b)

Figure 3. a) Nyquist plots obtained from a lithium/SEO 386-300/lithium cell at 90°C after 1, 50, 70 and 92 cycles at 0.175mA.cm-2 and b) zoom

Four Nyquist plots after 1,50, 70 and 92 cycles at 0.175mA.cm-2 are represented in Figure

3 for the cell 1. For the first and the 50th cycles, the spectra are very similar and are

composed of three contributions (described in Chapter 3.4): the electrolyte (Rel), the interface

(Rint) and the diffusion (Rdiff) contributions. During the first 50 cycles, Rel is constant, whereas

Rint increases slightly. The spectra after 70 and 92 cycles represent a typical behavior for a

smooth short-circuit. Indeed, Rel decreases by a factor two and Rint by a factor 9. In addition,

the characteristic frequencies of the interface increases from 3 kHz to 7kHz in the case of

the soft short-circuit.

It is possible to model this soft short-circuit by adding a resistance representing the

short-circuit (RSC) in parallel to the equivalent circuit used to model the whole cell1 (see in

Chapter 3.4.c), this equivalent circuit is presented in Figure 4 a). We fix all the parameters

(which were previously determined with an unshorted cell) except the RSC. We present the fit


- 132 -

results obtained for the cycle number 70 (in red solid line) and for the cycle number 92 (in

blue solid line) in Figure 4 b) and c). We found that the short-circuit for the cycle 70 is 589

Ω, whereas the RSC for the number 92 is 186.3 Ω.

a)

Figure 4. a) Equivalent circuit used to model the soft short-circuit, b) Nyquist plots obtained from a

lithium/SEO 386-300/lithium cell at 90°C after 70 and 92 cycles at 0.175 mA.cm-2 (in open markers) and the fit result obtained from the equivalent circuit given in a) (solid line) and c) zoom on the cycle 92.

In the case of the cell 2, three EIS spectra after 1, 50 and 70 cycles, are represented in a

Nyquist plot in Figure 5. It is important to notice that only the interface contributions are

evolving during cycling, the electrolyte resistance remains constant over all the cycling. In

addition, characteristic frequencies of Rint is constant, meanings that there is no chemistry

change in the interface during cycling.

Lcable Rcable Rel

CPEel

Rinterface

CPEinterface

Rsc

400

300

200

100

0

- Im

(Z

) (W

)

4003002001000

Re (Z) (W)

Cycle 70 Fit cycle 70 Cycle 92 Fit cyle 92

b)

200

150

100

50

0

- Im

(Z

) (W)

200150100500

Re (Z) (W)

Cycle 70 Fit cycle 70 Cycle 92 Fit cyle 92

c)


- 133 -

Figure 5. Nyquist plots obtained from a lithium/SEO 386-300/lithium cell at 90°C after 1, 50 and 70 cycles at 0.175mA.cm-2.

Experimental results of the study by EIS are shown in Figure 6 and Figure 7 for the two

different cells made with the same SEO electrolyte. Electrolyte resistance versus charge

passed are shown in Figure 6 a) and in Figure 7 a). The ionic conductivity for both cells are

similar (2.5 10-4 S.cm-1) and in a good agreement with the ionic conductivity reported

previously in Chapter 3.4. During cycling, the electrolyte resistance stayed stable after charge

and discharge. However the lifetime before the short-circuit is different, the cell 1died after

around 255 C.cm-2 (50 cycles), whereas cell 2 is still running after about 382 C.cm-2 (77

cycles). Unfortunately, due to constraints of channel availability the cell 2 had to be stopped

before it dies. Nevertheless, this experiment shows clearly that SEO 386-300 is very stable

over cycling, which is an expected result due to its high Mw providing good mechanical

properties.


- 134 -

Figure 6. SEO 386-300 lithium symmetric cell 1 a) electrolyte resistance versus charge passed and b) interface resistance versus charge passed.

In our experiments, even if both cells are produced with the same routine, Rint are

different. This is possibly due to various factors such as physical interface, surface state of

lithium metal chips, SEI forming at the surface of lithium when it is in contact with polymer.

Interface resistances are shown in Figure 6 b) and Figure 7 b), in both cells the interface

resistance after charge and after discharge are not similar. For the cell 1, the difference is

about 10 Ω.cm-2, when for cell 2 the difference is up to 40 Ω.cm-2 at the beginning of the

cycling. This gap is reduced over cycling from 10 to 6 Ω .cm-2, and from 40 Ω.cm-2 at the

beginning to 16 Ω.cm-2 at the final cycle for cell 1 and cell 2 respectively.

However, the two cells present different behavior for the interface resistance over

cycling. Cell 1 presented an interface resistance which is reasonably constant during cycling

until an abrupt drop in Rint from 80 Ω.cm-2 to 3 Ω.cm-2, due to a dendritic soft short-circuit

visible after 255 C.cm-2. This smooth short-circuit (see spectra in Figure 3) is accompanied by

a decrease in potential on the chronopotentiometry of the cell (see Figure 2).

40

30

20

10

0

Rele

ctr

oly

te (W

.cm

2)

4003002001000

Charge passed (C.cm-2

)

806040200

Cycle number

'After charge SEO 386-300' 'After discharge SEO 386-300'

a)

120

100

80

60

40

20

0

Rin

terf

ace (W

.cm

2)

5004003002001000


)

806040200

Cycle number

After charge SEO 386-300 After discharge SEO 386-300

b)


- 135 -

Figure 7. SEO 386-300 lithium symmetric cell 2 a) electrolyte resistance versus charge passed and b) interface resistance versus charge passed.

In the case of cell 2, for the first 25 cycles at 0.175 mA.cm-2 Rint is relatively constant after

charge and discharge, then it started to decrease from 180 Ω.cm-2 to 125 Ω.cm-2 after charge

and from 137 Ω.cm-2 to 109 Ω.cm-2 after discharge, which corresponded to 30% and 20%

respectively. Moreover, the gap between Rint after charge and Rint after discharge is

substantially reduced, from 40 Ω.cm-2 at the beginning to 16 Ω.cm-2 at the final cycle. We

attributed the difference in interface resistance after charge and after discharge to a memory

effect. During cycling, a slight decrease of the interface resistance between the polymer and

the lithium electrode is observed due to the disruption of the SEI and also to the presence of

fresh lithium2.

It is worth noting that the charge passed in both experiments with SEO 386-300 are

good results. The Balsara's group have reported similar experiments with SEO 240-269 with

the same EO/Li ratio, i.e. a BCE with a lower molecular weight compared to the one used in

this study. They have observed for a BCE using a film with a thickness around 30 μm, that

the charge passed is 123 ± 60 C.cm-2 at 90ᵒC with the same cycling procedure3,4. Hallinan et

al.5 have reported that the charge passed is linearly dependant of the thickness in SEO BCE,

which means that the dendrites growth has a higher impact compared to the nucleation.

Indeed, if the nucleation of dendrites was the limiting factor, the charge passed should be

proportional to the inverse of the thickness6. Therefore, due to a thickness twice thicker in

30

28

26

24

22

20

18

16

Rele

ctr

oly

te (W.c

m2)

4003002001000


)

6040200

Cycle number

'After charge SEO 386-300' 'After discharge SEO 386-300'

a)

200

150

100

50

0

Rin

terf

ace

(W.c

m2)

4003002001000


)

6040200

Cycle number

After charge SEO 386-300 After discharge SEO 386-300

b)


- 136 -

our case (around 60 μm), we should expect a charge passed twice larger than the one they

obtained, i.e. about 250 C.cm-2. Here, we report a charge passed ranging from 255 C.cm-2 to

higher than 386 C.cm-2. This increase in the charge passed is probably due to the increase of

mechanical properties due to the increase of molecular weight. Stone et al.7 have shown that

the charge passed increase with the modulus of the BCE. Moreover, Hallinan et al. have

observed that the condition in lithium symmetric cells are stronger than in batteries, and they

reported about three times longer the lifetime in batteries.

c. Single-ion block copolymers

PEO-b-PSTFSI block copolymer. Similar cycling experiments are performed on

single-ion block copolymers at 90ᵒC. The voltage versus time during the cycling experiment

is presented in Figure 8 a). Here, we presents only cycles at 0.175 mA.cm-2 after the 15 pre-

conditioning cycles at 0.02 mA.cm-2. The voltage stays constant during the 37 cycles. And a

direct short-circuit with a drop to zero of the voltage is observed after 37 cycles. This

behavior is different from that observed for SEO BCE, where the voltage decreased but no

direct short-circuit has been observed.

-0.3

-0.2

-0.1

0.0

0.1

0.2

0.3

Vo

lta

ge

(V

)

3002001000Time (h)

Lithium symmetric cell with PEO-b-PSTFSI

a)


- 137 -

Figure 8. Voltage versus time graphic for a PEO-b-PSTFSI BCE-lithium symmetric cell at 90ᵒC, b) zoom on cycles 5 and 6, c) zoom on cycle number 37 and short-circuit.

For more clarity, a zoom of two cycles is given in Figure 8 b), the voltage behavior is

unexpected. Indeed, we suppose that for a single-ion BCE the voltage should reach almost

instantly the steady state and stays constant, however, here, we observe that the voltage

increases linearly during the polarization. When the polarization is stopped the voltage

dropped at 0 instantly. Finally, a zoom of the last cycle and the direct short-circuit is shown

in Figure 8 c).

Figure 9 shows four spectra obtained for a PEO-b-PSTFSI-lithium symmetric cell after 1,

15, 25 and 37 cycles at 0.175 mA.cm-2. As described previously (Chapter 3.4.) two

contributions are visible. Rel stays constant over cycling, and Rint slightly decreases. However,

the characteristic frequency did not change during cycling.

-0.3

-0.2

-0.1

0.0

0.1

0.2

0.3

Vo

lta

ge

(V

)

55504540Time (h)

b) -0.3

-0.2

-0.1

0.0

0.1

0.2

0.3

Vo

lta

ge

(V

)

370365360355Time (h)

c)


- 138 -

Figure 9. Nyquist plots obtained from a lithium/Si-EO/lithium cell at 90°C after 1, 15, 25 and 37 cycles at 0.175mA.cm-2.

The electrolyte resistances and the interface resistances versus charge passed for the

diblock copolymer PEO-b-PSTFSI are presented in Figure 10 a) and b) respectively.

The ionic conductivity is in a good agreement with the conductivity measured in Chapter

3, here it is constant at about 2.7 10-5 S.cm-1. Rel is measured after each charge and each

discharge and is relatively constant during cycling at about 240 Ω.cm2. After 37 cycles at

0.175 mA.cm-2 (charge passed of 197 C.cm-2), the cell experienced a short-circuit and dies.


- 139 -

Figure 10. PEO-b-PSTFSI lithium symmetric cell a) conductivity versus charge passed and b) interface resistance versus charge passed.

The interface resistance experienced changes during cycling. As seen for SEO electrolyte,

values after charge and after discharge are different. However, from the first cycle to the last

one, resistance decreased from 1263 Ω.cm-2 to 970 Ω.cm-2 after charge, which corresponds to

a decrease of 23%. A similar decrease is observed after the discharge, from 1170 Ω.cm-2 to

932 Ω.cm-2, which corresponds to a decrease of 20%. This large change in the interface

resistance between the polymer and lithium is already observed for this single-ion in the case

of the transference number experiment (see Chapter 3.5.). One part of this decrease in

interface resistance can be attributed to the reorganization and/or the breaking of passive

layers at the interface during polarization.

This single ion BCE present a low resistance to dendritic growth with only 37 cycles at

0.175mA.cm-2, compared to SEO BCE which exhibited for the worst 50 cycles at the same

current density. However, it is very important to remember that this BCE has a low

molecular weight (48.725 . 103 kg.mol-1) and is very weak compared to SEO (686 . 103

kg.mol-1). In addition, it is important to note that the Si-EO BCE exhibits poor mechanical

properties and more especially at 90ᵒC. We could not performed mechanical properties

characterization for this BCE, however, we observe that it becomes a viscous liquid above

60ᵒC. Therefore, the results obtained are very encouraging, if we compare this BCE to SEO

with poor mechanical properties we should obtained smaller charge passed7.

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- 140 -

PSTFSI-b-PEO-b-PSTFSI block copolymer. Similar experiments are performed with

PSTFSI-b-PEO-b-PSTFSI electrolyte. We present the voltage versus time and the current

versus time in Figure 11 a) for four preliminary cycles performed at 0.02 mA.cm-2. It is

interesting to see how the voltage follows intimately the current. When the polarization starts

the voltage reaches instantly the steady state and stays relatively constant. The voltage drop

to zero during the rest time. The Figure 11 b) shows regular cycles at 0.175 mA.cm-2, for

more clarity two zooms on three cycles are given. A first one from cycle number 20 to 23

and a second one from the 92th cycle to the 95th. We observe that the voltage is almost flat

when the cell is under polarization. In addition, the voltage is pretty stable during cycling (a

very small difference between the 20th cycle and the 95th).

-0.003

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a)


- 141 -

Figure 11. a) Preliminary cycles at 0.02 mA.cm-2 for a PSTFSI-b-PEO-b-PSTFSI BCE-lithium symmetric

cell at 90ᵒC: voltage versus time (burgundy) with the current versus time (turquoise) and b) zoom on cycles 20 to 23 and 92 to95 at 0.175 mA.cm-2.

Four EIS spectra after 20, 60 and 95 cycles at 0.175 mA.cm-2, are given in a Nyquist plot

in Figure 12. The electrolyte resistance is almost constant during cycling and the interface

resistance fluctuates around 450 Ω.cm2.

Figure 12. Nyquist plots obtained from a lithium/Si-EO-Si/lithium cell at 90°C after 1, 20, 60 and 95 cycles at 0.175mA.cm-2.

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- 142 -

Figure 13 presents the evolution of the electrolyte resistance and the interface resistance

as a function of charge passed. The absence of point for some cycles are due to the high

noise observed in EIS measurement and the impossibility to obtain the right values for the

resistances.

Figure 13. Si-EO-Si lithium symmetric cell a) electrolyte resistance versus charge passed and b) interface resistance versus charge passed.

The electrolyte resistance is in good agreement with the conductivity measurement

presented in Chapter 3.4. In addition, the electrolyte resistance is relatively constant over

cycling around a value at 400 Ω.cm2. The interface resistance is also constant over cycling

and the value is around 450 Ω.cm2. However, we observe some fluctuations of the interface

resistances. This cell has been cycled at 0.175 mA.cm-2 for 95 cycles, which corresponds to

around 500 C.cm-2, the experiment has to be stopped due to time constraints. It is worth to

note that this BCE as the Si-EO BCE is weak and exhibits poor mechanical properties.

Therefore, a charge passed of 500 C.cm-2 without short-circuit is an excellent and very

encouraging result.

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- 143 -

Figure 14. Mechanical strain versus stress for the Si-EO-Si BCE at 40ᵒC.

The tensile test has been performed on the Si-EO-Si BCE, the Figure 14 shows the result

for the BCE at 40ᵒC. The Young's modulus is calculated with the slope of the linear part

below the elastic limit and we obtain a Young modulus of 21 MPa at 40ᵒC and 0.11 MPa at

60ᵒC (above the Tm of the BCE). For the comparison, the modulus of SEO BCE (240-260)

is 51.6 MPa at 90ᵒC without lithium salt7. Thus, the Si-EO-Si presents a modulus at least

three order of magnitude lower than SEO at 90ᵒC, nevertheless, it presents exceptional

results with more than 500 C.cm-2 of charge passed. We can suppose that this result is due to

the suppression of the concentration gradient, indeed the Si-EO-Si present a Li+ transference

number of unity.

Many different electrolytes have been studied in order to mitigate lithium dendritic

growth. Here, we compare our results to a wide variety of literature studies, mostly extracted

the recent Choudhury et al.8 study. We plot the results obtained for the three BCEs in the

graphics representing the short-circuit time (Tsc) versus current densities in Figure 15. The

average result for the SEO BCE is represented by a blue square, the result for the Si-EO

BCE is represented by a turquoise circle and finally the result for the Si-EO-Si BCE is

represented by a red triangle. It is important to remember that the Si-EO-Si BCE has not

experienced a short-circuit, therefore we plotted the time value obtained after the last cycle,

but it is higher.

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- 144 -

Figure 15. Short-circuit time of the three BCEs studied (turquoise circle for SEO, blue square for Si-EO and red triangle for Si-EO-Si) compared with other state of the art battery performance. Red squares and red circles indicate the Tsc for strip-plate test and polarization test respectively in Archer et al. work8. The black filled symbols represent polarization tests done at room temperature, while the open symbols represent elevated temperature experiments. Black closed triangles represent silica tethered with imidazolium (Si-IM-IL) and piperidinium ionic liquid (Si-PP-IL) at various volume fractions of silica, indicated in parenthesis9. Black closed diamonds indicate anion tethered hybrid silica nanocomposites10. The high temperature data include crosslinked PE-PEO solid polymer with different plasticizer content given in parenthesis11. Other data points are PVdF-HFP/PEO composite12, high molecular weight polymer13, silica -polymer composite14, polymer with ionic liquid15 as well as their combination16. The blue symbols indicate neat/pristine electrolyte systems. Error bars denote deviations from different measurements.

If we compare our results to Rosso et al.13 study about high molecular weight PEO

containing lithium salt, the SEO BCE exhibits a time before short-circuit (Tsc) at least one

order of magnitude higher, whereas the Si-EO-Si BCE exhibits a Tsc two orders of

magnitude higher. The Figure 15 shows also other studies such as Liu et al.14–16, who have

studied silica-PEO based nanocomposite electrolytes, solid polymer electrolytes based on

PEO X LiTFSI repeating units and a combinations of the two. Khurana et al.11 used cross-

linked PEO-PE-PEO polymers and Sannier et al.12 have studied electrolyte using gel-polymer

membranes based on PEO and PVdF-HFP polymers. Comparison between these literature

studies and our study shows that the SEO BCE and the two single-ion BCEs exhibit as good

or better performances than these studies. The recent study from Choudhury et al.8 about

cross-linked hairy nanoparticles reports very promising results slightly higher compared to

ours. However, it is important to remember that these electrolytes present high mechanical

properties on the contrary of the Si-EO and Si-EO-Si BCEs. For comparison, results for

liquid electrolyte/separator (here noted as neat electrolyte) are also given in Figure 15, the

Tsc is smaller than our results. In addition, it is worth to note that the presence of a


- 145 -

separator increases the Tsc. To conclude, here we report very promising results with high

Tsc, especially for the Si-EO-Si BCE with more than 500 C.cm-2.

2. Dendrites morphologies studied by hard X-Ray micro-

tomography

a. Hard X-Ray micro-tomography

Hard X-ray micro-tomography is a radiographic imaging technique, it is non destructive,

can produce 3D images of a material's internal structure at a spatial resolution of about a

micrometer and it is available for a diverse range of samples. This technique is becoming

more mainstream since instruments, analysis and modeling code have been improved. The

tomography technique requires to record a series of x-ray radiographs of the sample over a

range of different angles in order to reconstruct a 3D image with this series of 2D images.

Thus, advanced mathematics are required to obtain the 3D reconstruction which is first

introduced by Radon in 1917 17. X-ray micro-tomography uses an X-ray scintillators to

convert X-rays to visible light, in addition this light is recorded by a visible light microscope

and a charge coupled area detector (CCD).

In this work, X-ray micro-tomography is performed at the Advanced Light Source (ALS)

at Lawrence Berkeley National Laboratory (LBNL) in California at the beamline 8.3.2. 18 The

X-ray source is coming from a synchroton radiation with a 1.9 GeV electron beam.

The sample mounting stage assembly and the X-ray camera detector are shown in Figure

16. The sample is placed on a rotary stage which is a circular plate magnetically coupled to

the top of a rotary air bearing. The sample is placed on a magnetic plate which enables a

rotation from 0ᵒ to 360ᵒ.

For the data acquisition, the typical protocol is the rotation of the sample from 0 to 180

degrees during which 2049 radiographs are recorded. The system operates in a continuous

sample rotation mode, which reduce the acquisition time. The number of images over the

angular range, the exposure time and the blur limit are selected by the user and can be tuned

depending of the material studied 18.


- 146 -

Figure 16. Schematic layout of the sample mount and X-Ray camera at the ALS beamline 8.3.2. 18

The tomographic reconstruction is performed with Fourrier methods using the

commercial Octopus software. For the image processing, visualization and analysis, both

Fiji19 and Avizo20 are used.

b. Protocol

In the case of BCE-lithium symmetric cells, the orientation of the cells for the hard X-ray

micro-tomography is shown in Figure 17. Due to the low molecular weight of the different

BCE and the lithium, it is possible to image the entire cell.

Figure 17. Orientation of a polymer-lithium symmetric cell on the rotating stage for hard X-ray microtomography experiment.


- 147 -

Two kinds of cells are imaged, the uncycled cells and the shorted cells.

After cells shorted, they are taken back into the glovebox to remove the nickel current

collectors in order to improve the quality contrast of the X-Ray micro-tomography images.

The new cells are vacuum sealed in a new pouch cell and transferred from the glovebox to

the micro-tomography beamline 8.3.2.

c. Neutral block copolymer

Recent studies using X-ray micro-tomography experiments have shown a new

morphology of dendrites inside the lithium-SEO symmetric cells. Indeed, the formation of

globular dendritic structures at the lithium metal/SEO interface, with a part of the dendrite

residing within the lithium electrode, have been observed and reported3,4,21.

A slice through a typical tomogram obtained from an uncycled cell is shown in Figure 18.

The image shows the cross-section of an uncycled lithium symmetric cell. We can observe

the three expected phases; one thin and bright layer corresponding to the electrolyte, the

block copolymer SEO 386-300, and two surrounding darker layers corresponding to lithium

metal. It is important to notice that the grayscale pixel value in the tomogram corresponds to

the relative X-ray linear absorption coefficients of the material at that point. Consequently,

polymer electrolyte appears brighter than lithium metal because they have higher electron

densities, which makes them more opaque to the X-rays. Besides, the electrolyte-electrode

interface is surrounded by a thin bright band on the electrolyte side and a thin dark band on

the electrode side. This is due to the Fresnel phase contrast, which appears during the

imaging of samples containing interfaces22. This effect is also apparent in the case of

dendritic structures. However, some variation in pixel brightness can possibly arise from

numerous source of noise. The resolution of this technique in our conditions is of the order

of the micrometer.

Figure 18. 2D X-ray tomography slice showing the cross-section of symmetric lithium cell.

Faceted bright particles are observed in the uncycled cell. They are lying in the lithium

electrode and they are identified as one of the main contaminants of the lithium metal. One


- 148 -

example is given in Figure 18 on the top lithium electrode. In order to show that those

impurities are inside all the lithium metal, a slice (parallel to the slice showed in Figure 18)

through the bulk of the top lithium electrode is shown in Figure 19. We observe 11 faceted

impurities (inside the white circle) on this slice and they are randomly distributed. According

to the manufacturer, the main lithium impurity is nitrogen, which is expected in the form of

lithium nitride (Li3N)23 and we can suppose that it is those impurities that we observed.

However, a recent study by energy dispersive spectroscopy (EDS) showed that such

impurities are more likely to be either lithium hydroxide (LiOH) or lithium oxide (LiO2)21.

Whatever the case, those crystalline impurities are insulating and randomly distributed within

the electrode.

Figure 19. Slice through the lithium metal showing the faceted impurities.

The uncycled lithium-SEO symmetric cell is devoid of any other noticeable features, and

thanks to our process an intimate contact between lithium and polymer is obtained.

Before the discussion about our results, it is important to note that Harry et al.4, reported

a new morphology of dendrite in SEO block copolymer electrolyte, a globular and

multiglobular morphology observed by hard X-ray microtomography. Moreover, they

showed that in the early stages of dendrite formation, a subsurface structure is formed and its

volume is larger than that occupied by the one protruding out from the electrode surface.

Figure 20 presents the evolution of dendrite growth as a function of charge passed. Then

dendrites grow through the bulk of the electrolyte.


- 149 -

Figure 20. Evolution of dendrite growth. a–d, X-ray tomography slices showing the cross-sections of symmetric lithium cells cycled to various stages. The thin, bright horizontal strip through the centre of the images is the

polystyrene-block-poly(ethylene oxide) copolymer electrolyte sandwiched between two lithium metal electrodes. The amount of charge passed, C, for each cell is: 0 C cm-2 (a), 9 C cm-2 (b) and 84 C cm-2 (c). d, Shorted cell: 296 C cm-2. Dendritic structures are evident in b–d. e–h, Magnified, 3D reconstructed volumes of cells shown in the top panel. e, An uncycled cell with no dendritic structures, C = 0 C cm-2. f, Heterogeneous structures begin to form in the bottom electrode in early stages of cycling, C = 9 C cm-2. g, Dendritic structures in both electrolyte and electrode phases are seen at the intermediate stage of cycling, C = 84 C cm-2. h, Dendritic structures that span the thickness

of the electrolyte are seen in the shorted cell, C = 296 C cm-2. The arrow indicates the colour scale for voxel brightness.

Later, they observed that all dendrites nucleate from an impurity particle located inside

lithium metal below the nucleation site24. They proposed a mechanism for the nucleation and

growth of the lithium globular structure and it is presented in Figure 21. They assumed the

presence of a SEI between the lithium metal and the electrolyte and they postulated that the

SEI is interrupted at the edge of the impurity (see Figure 21) leading to a preferential

deposition of lithium at the corner of the impurity, due to an increase in local conductivity.

However, this is possible only for impurities located close to the electrolyte.


- 150 -

Figure 21. A schematic showing a proposed mechanism by Harry et al. 21 for the nucleation of and growth of

lithium globular structure.

After, this brief reminder of the dendritic growth in SEO BCEs, we presents our results.

Figure 22 shows a cross section of a lithium-SEO shorted cell after 255 C.cm-2 of charge

passed. Numerous dendritic structures are observed in the portion of the cell imaged and a

typical example of dendritic structure is visible in Figure 22. The dendrite shown here is

composed of several agglomerated globules of lithium, which appears slightly darker. It is

interesting to notice that the globular structure observed is a mix between lithium and BCE,

i.e. a "porous" dendrite structure. This structure extends both in the lithium electrode and in

the electrolyte. In front of the dendrite, the lithium electrode forms a clear crater

complementary to the shape of the dendrite. Finally, the presence of a small crystallites at the

foot of the dendrite can also be seen. Multi-globular structure is consistent with previous

publications for a similar SEO electrolytes4,3,21. However, it is important to specify that this

type of dendrite was not commonly observed in PEO homo-polymer, for example, where

morphology of needle, tree like, or mossy-like25, with sharp tips and highly branched

structure have been discussed. In this study the dendrites are blunt and not branched.

On the top of the dendrite presented in Figure 22, a split is visible and continuous

lithium is seen which suggested a short-circuit at this point.

Moreover, electrolyte spanning structures are observed. Indeed, one large globule

structure (227 μm of diameter with 175 μm of height) spans the electrolyte, when several

others globules (with smaller sizes) do not span the electrolyte and are contained inside the

large globule (Figure 22).


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Figure 22. X-ray tomography slice showing the cross section of a symmetric lithium cell after shortage for a SEO 386-300 electrolyte. An impurity crystallite is visible at the dendrite's foot.

The multi-globular structure of the dendrite is coming from the successive cycling back

and forth of the lithium21. Indeed, polarization experiment in which lithium is deposited only

in one direction, on the bottom side of the lithium symmetric cell is presented in Figure 23

c). This experiment resulted in a single globule of lithium with electrolyte spanning as shown

in Figure 23 c), whereas, several cycles resulted in a multi-globular structure (Figure 23 b)).

Figure 23. Cross section slices through reconstructed X-ray tomograms of lithium metal, polymer electrolyte, lithium metal symmetric cells21. a) Before the passage of current, b) after cycling, lithium filled multi globular

structures form in the polymer electrolyte, and c) when lithium is passed in one direction from the top to the bottom electrode, structures containing a single lithium filled globule form.

Figure 24 shows the three-dimensional (3D) reconstructed volumes of selected regions

around the slice shown in Figure 22. Each reconstructed volume should be viewed as a 3D

array of brightness values. Voxels with a brightness below a certain threshold are rendered

transparent in order to observe only the lithium dendritic structure itself.

The three dimensional nature of the dendritic structures formed in lithium-SEO

symmetric cells is clearly observed. On the right and left side of the dendrite, we can see two

flat sheets which are the electrode/electrolyte interface. Due to a lower X-ray absorption for

lithium compared to SEO electrolyte, lithium appeared with darker voxels, which explained


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the absence of features above and below the two electrode-electrolyte interfaces. Some spots

are visible in Figure 24 and are due to noise.

a) b)

Figure 24. 3D volume rendering of a) whole dendrite and b) cross section inside the dendrite.

The dendrite presents two different parts, one that lies inside the electrolyte which is

highly expected, and another that lies inside the electrode. This imaging technique enables

the lithium-polymer interfaces to be observed, we can see that the lithium-polymer interfaces

are running through the electrolyte and the lithium electrode. The two dendrite parts are thus

filled with ramified lithium-polymer interfaces and it is clearly shown in Figure 24 b).

However, we can notice that the lithium/electrolyte interface in the upper electrode is always

intimate that means that the SEO is able to follow the evolution of the lithium morphology

by wetting the lithium, implying that the BCE has good viscous-elasticity properties.

In addition, an impurity faceted crystallite is observed at the base of the dendrite in

Figure 22 and such impurity is present for every dendritic structures observed in this cell.

Thus, this result is consistent with literature3,4,24.

d. Single-ion block copolymer

In this section we will investigate the behavior of single-ion BCE. Contrary to the SEO

electrolytes studied in the section above, Si-EO and Si-EO-Si BCEs have a Li+ transference

number of unity, which results in the absence of a concentration gradient over cycling6 and

thus lithium dendritic nucleation should be avoided according to the model of Chazalviel6.

Nevertheless, short-circuit is observed for several cells. In order to investigate what

phenomenon are implied in the failure of the cells, hard X-ray micro-tomography is

performed at the ALS.

Si-EO BCE. Similar experiment has been performed on lithium symmetric cell

containing Si-EO BCE. A slice through the lithium symmetric cell before cycling is given in

Figure 25. As seen previously, in the case of SEO BCE, the polymer electrolyte exhibits an


- 153 -

intimate and homogeneous contact with lithium. The polymer layer presents small variations

in its thickness but it is mostly homogeneous and planar.

Figure 25. X-ray tomography slice showing the cross-section of a PEO-b-PSTFSI lithium symmetric ell before cycling.

A slice through the polymer after the cell failed after 197 C.cm-2 is given in Figure 26.

The polymer-lithium interface exhibits an irregular morphology very different compared to

the one previously shown in Figure 25. Contrary to the SEO BCE, there are no clear defined

objects. Besides, the polymer layer exhibits some parts where the interface is still planar,

meanings that the lithium is homogeneously deposited and stripped, and some other parts

where the interface BCE-lithium is irregular on both sides suggesting that the lithium electro-

deposition during cycling is not homogeneous. Moreover, the irregularities of the interface

on both sides of the BCE are not similar, meanings that the local current densities are not

symmetrical. Nevertheless, the contact between the BCE and lithium is still intimate. In

addition, the lithium presents the same gray value even in the irregularities implying that the

lithium is dense. The last interesting observation is that no crystalline impurities are present

close to the polymer surface. Thus, the mechanism of nucleation and growth of lithium

"dendrites" should be different from the one implicated in SEO BCE. The lithium objects

which have grown are however dense, which suggests that the lithium deposition through

the single-ion BCE is irregular but dense. The dense objects present various sizes (from 12.5

to 62.5 μm in diameter) and shapes.

Figure 26. X-ray tomography slices showing the cross section of a symmetric lithium cell after the cell failed for a PEO-b-PSTFSI electrolyte.


- 154 -

A 3D volume rendering of the interface polymer-lithium is given in Figure 27. The cross

section of the Si-EO-lithium symmetric cell is given in Figure 27 a), we choose an area where

the interface BCE-lithium is very irregular on purpose. Here, it is even more clear that the

morphologies of the lithium deposits on both sides of the Si-EO BCE are different and

independent from the other BCE-lithium interface. The Figure 27 b) presents the top surface

of the 3D reconstruction and it is clearly observed that the polymer-lithium interface is very

irregular and rough. The increase of roughness of the interfaces may clearly explain the

decrease of interface resistance observed in the cycling experiment (Figure 10).

a) b)

Figure 27. 3D volume rendering of a PEO-b-PSTFSI lithium symmetric cell after cycling a) cross section and b) upper surface.

Moreover, on the 3D reconstruction (Figure 27 b)) we can distinguished large and small

irregularities of the interface. The large objects are easily observed even in the tomography

slice Figure 26, but the small irregularities of the interface observed in the 3D reconstruction

is hard to detect with only the tomogram. The size of these irregularities is inferior to 5μm.

These local variations of the interfaces (large and small) are probably due to local current

density variations, that may come from the variation of the SEI properties26. In addition, it is

important to remember that the Si-EO BCE exhibits a yellowish color, which suggests that

impurities are still present in the BCE. These impurities are probably responsible of local

heterogeneous composition of the BCE, which can increase the local current densities

leading to more irregular deposition of lithium.

Therefore, in the case of the single-ion BCE we do not observed porous lithium

dendrites, but instead dense lithium objects with different shapes and sizes.

Si-EO-Si BCE. Similar characterization experiment are performed on Si-EO-Si-lithium

symmetric cells. Prior to the cycling experiment, the lithium symmetric cell was imaged and

very similar tomography slices compared to Figure 25 are obtained, i.e. intimate contact

between the electrolyte and the BCE.


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Figure 28 a) shows an X-ray micro-tomography slice showing the cross section of a

Li/Si-EO-Si/Li cell after a charge passed of 500 C.cm-2 which corresponds to 95 cycles at

0.175 mA.cm-2.The first observation is that globally the interface is very smooth, but it

appears few large lithium objects. The morphology of the objects observed is completely

different to the dendritic structures observed in SEO BCE. We observe a concave semi

ellipsoid structure, which punctures the electrolyte membrane. No electrolyte spanning is

observed in the lithium electrode. Another major change is the absence of lithium impurity at

the base of the dendritic structure, which is similar to the result obtained for SI-EO BCE.

Moreover, this object presents a gray value similar to the lithium value, which implies that

this object is composed of dense lithium metal suggesting that the electrolyte presents a

transference number of unity. However, the presence of such objects suggested

heterogeneities in the electrolyte itself or in the SEI, which produced heterogeneous current

densities at the surface of the polymer and lead to heterogeneous deposition. In addition, the

size of the dense object is large, in this case we measure the diameter to be 56 μm with a

depth of 18 μm.

a) b)

Figure 28. X-ray tomography slice showing the cross section of a symmetric lithium cell after failure for a Si-EO-Si BCE electrolyte for a) a cell after 500 C.cm-2 and b) another dense lithium object.

Another example of lithium object is shown in Figure 28 b). The morphology is similar

to the first dendrite observed, dense lithium is deposited and formed a concave hole inside

the electrolyte layer. The dimension of this object is larger than the previous one, i.e. 143.7

μm of diameter with a depth of 53 μm.

A 3D volume rendering is performed on the tomogram in order to have a better idea of

the morphology of the dendritic objects, the result is given in Figure 29 for the object shown

in Figure 28 a). Two different views of the same dendritic objects are represented: Figure 29

a) the top surface where we can observe the hole and b) the bottom surface. Two large

dendritic objects are observed in this area with smaller lithium objects. The two large objects

exhibit a length of 120 for the bigger one and 90 μm for the second one, with a maximum

width of 70 and 50 μm respectively. In addition of the concave hole, the object presents a


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convex semi ellipsoid on the other side of the object (Figure 29 b) and d)) which is the

complementary deformation of the depression on the other side. Thus, we can suppose that

the current densities are almost symmetrical in this case. This result is different compared to

the Si-EO BCE, where the deformations of the interface BCE-lithium were independent

from one side to another. A schematic of the 3D volume rendering of the dense dendritic

object is represented in Figure 29 c) and d) for more clarity.

a) b)

c) d)

Figure 29. a) and b) 3D volume rendering of the same dendrite observed in the Si-EO-Si lithium symmetric cell, c) and d) schemes of the 3D volume rendering.

It is important to remember that this BCE is cycled at 90ᵒC, i.e. at a temperature higher

than its melting temperature (Tm=56ᵒC see Chapter 3.2). As previously reported in Chapter

4.1.c, the Si-EO-SI BCE exhibits poor mechanical properties above 60ᵒC. If we compare this

BCE to SEO BCE with similar mechanical properties, the charge passed before failure

should be smaller (inferior to 100 C.cm-2)7. In addition, according to the Monroe and

Newman model27–29, the dendritic growth should be mitigated by mechanical barriers,

however, here we report a very high resistance to dendrite with poor mechanical properties.

The presence of this dense lithium objects in single ion BCEs is surprising, since the Li+

transference number is equal to unity (see Chapter 3.5). Indeed, according to Chazalviel's

model6, no space charge can be formed on the vicinity of the negative electrode in single ion

BCE. And one possible mechanism for the nucleation of dendrites is the formation of a

space charge. In our materials, the concentration or activity of anions is constant because the


- 157 -

TFSI anions are covalently bond to the backbone of the BCE. In addition, according to

Monroe and Newman27–29 (a model only on the dendrite growth), when tLi+ is unity the Sand

time, i.e. the starting time to lithium growth, is infinite meanings that no dendrite should

grow. One possible explanation, for the formation of dense dendritic objects, is the current

instabilities due to the irregular SEI at the electrolyte/electrode interface26. Nevertheless, it is

worth to note that the Si-EO-Si BCE exhibits exceptional dendrites resistance performances,

despite the presence of these lithium objects.

Figure 30. a) Dendrite formed in a lithium battery using a liquid electrolyte (EC/DMC 2:1 and 1M of LiPF6 salt) after one charge at 2.2 mA.cm-2

In the Figure 30, we present the morphology of a typical lithium dendrite which have

grown in a battery assembled with an organic liquid electrolyte30. The morphology is very

different from the morphology obtained in our experiments. In fact, in both neutral and

single-ion BCEs, the lithium objects are larger and the shape are neither tree-like, nor needle-

like (the classical dendrite shapes). In addition, in the case of single-ion BCEs the lithium

objects exhibit no porosity.

3. Conclusion

In this chapter, we reported the behaviors of BCEs under galvanostatic cycling at a

current density of 0.175 mA.cm-2 during 4 hours (3.15 μm of lithium displaced at each time).

In the case of the SEO BCE laden with LiTFSI salt a classical behavior is observed, i.e. after

the large initial voltage increase, the voltage increases smoothly due to different processes.


- 158 -

The steady state is never reached in our conditions. A slow relaxation voltage is then clearly

visible before reaching 0V, due to the re-equilibration of concentration along the cell during

the rest time. However, in the case of Si-EO-Si BCEs the voltage attains the steady state

almost instantly and stays constant during the polarization time simultaneously. For both

single-ion, when the current is stopped the potential drops to zero without relaxation. This is

a clear demonstration of the single-ion nature of these electrolytes.

For the neutral SEO, we observed that the electrolyte resistance is pretty constant during

cycling (before the short-circuit), whereas the interface resistance decreases during cycling

and then stabilizes. This electrolyte, which is mechanically very strong, presents as expected a

high resistance to dendrite corresponding to more than 386 C.cm-2. In the case of single-ion

BCEs, the Rint also decreases during cycling. A very promising result is that despite their poor

mechanical properties at 90ᵒC, these materials exhibit good (for the diblock) and excellent

(for the triblock) resistance to dendrites with 197 C.cm-2 and 500 C.cm-2 (the value should be

higher because the cell did not short-circuit) of charge passed) respectively.

Post-mortem hard X-ray micro-tomography has been performed on the different cells

after cycling, in order to determine the morphology of the lithium cycled. In the case of SEO

BCE, globular structures have been observed, which is in very good agreement with

literature results for similar materials21,24. In the case of single-ion BCEs, for the first time,

dense lithium objects have been observed: which leads to a dense irregular interface. The

BCE-lithium interface in the case of Si-EO BCE presents some area where it is very planar,

meanings that the electro-deposition and stripping of lithium is very homogeneous. But, in

other parts, the BCE-lithium interface is very irregular, dense lithium objects are observed

with different sizes and shapes. But in general there are large and smooth.

This is a surprising result because no dendritic object is expected in the case of

electrolytes, which present a Li+ transference number of unity. The formation of these dense

dendritic objects are probably due to heterogeneities of the local current densities because of

heterogeneous SEI or local different BCE composition for example. Whatever the case,

these objects exhibit a morphology very different from any kind of dendritic objects already

observed and described in the literature.


- 159 -


1. Rosso, M. et al. Dendrite short-circuit and fuse effect on Li/polymer/Li cells. Electrochimica Acta 51,

5334–5340 (2006).

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4. Harry, K. J., Hallinan, D. T., Parkinson, D. Y., MacDowell, A. A. & Balsara, N. P. Detection of

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69–73 (2014).

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11. Khurana, R., Schaefer, J. L., Archer, L. A. & Coates, G. W. Suppression of Lithium Dendrite Growth

Using Cross-Linked Polyethylene/Poly(ethylene oxide) Electrolytes: A New Approach for Practical

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12. Sannier, L., Bouchet, R., Rosso, M. & Tarascon, J.-M. Evaluation of GPE performances in lithium

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lithium/polymer cells. J. Power Sources 97–98, 804–806 (2001).

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15. Liu, S. et al. Lithium Dendrite Formation in Li/Poly(ethylene oxide)–Lithium

Bis(trifluoromethanesulfonyl)imide and N-Methyl-N-propylpiperidinium

Bis(trifluoromethanesulfonyl)imide/Li Cells. J. Electrochem. Soc. 157, A1092–A1098 (2010).

16. Liu, S. et al. Effect of co-doping nano-silica filler and N-methyl-N-propylpiperidinium

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Li/poly(ethylene oxide)-Li(CF3SO2)2N/Li. J. Power Sources 196, 7681–7686 (2011).

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Stripping Behavior of Lithium Metal across a Rigid Block Copolymer Electrolyte Membrane. J.


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Stripping Behavior of Lithium Metal across a Rigid Block Copolymer Electrolyte Membrane. J.


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26. Anna Teyssot, C. B. Inter-electrode in situ concentration cartography in lithium/polymer

electrolyte/lithium cells. J. Electroanal. Chem. 584, 70–74 (2005).

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- 162 -

Chapter 5.

The polymer-ceramic composite

Abstract

This chapter will be focused on a new type of electrolyte, a composite

electrolyte composed of an inorganic ceramic (Ohara) and a protective organic

layer made of a block copolymer electrolyte. A first part will be devoted to the

study of the composite electrolyte, characterization experiments such as ionic

conductivity, cycling experiment and hard X-ray microtomography have been

performed. A second part will be dedicated to the quantification of the

polarization loss at the interface polymer-ceramic. In this purpose, a

galvanostatic steps experiment has been performed on polymer-lithium symmetric

cells and composite-lithium symmetric cells.

Chapter 5. The polymer-ceramic composite

- 164 -

Table of contents

Chapter 5. The polymer-ceramic composite ................................................................ 163

I. Study of the polymer-ceramic composite ................................................................................................ 165

1. Experimental procedure.............................................................................................................................. 165


a. Electrical properties ...................................................................................................................................... 167

b. Cycling ......................................................................................................................................................... 175

c. Characterization by X-Ray microtomography ............................................................................................... 185

II. Quantification of polarization loss at the polymer-ceramic interface ................................................. 190

1. State of the art .............................................................................................................................................. 190

2. Experimental procedure.............................................................................................................................. 191


a. Experimental results ................................................................................................................................... 192

b. Discussion .................................................................................................................................................... 196

Conclusions ............................................................................................................................................................ 201

References of chapter 5 ........................................................................................................................................ 204


- 165 -

I. Study of the polymer-ceramic composite

1. Experimental procedure

Cells preparation and assembly. Composite cells are assembled using Ohara GC,

presented in Chapter 2, SEO 55-52 and the two single-ion block copolymers presented in

Chapter 3, i.e. PEO-b-PSTFSI and PSTFSI-b-PEO-b-PSTFSI. The ceramic is sandwiched

between two thin membranes of BCE.

For the composite assembly, extra care is taken when handling the ceramic because of its

brittleness. Two membranes of BCE melt pressed onto a kapton spacer with a hole of ¼ inch

diameter are first produced (following the procedure in Chapter 3.4.). They are then placed on

both sides of the ceramic and finally slightly and gently pressed inside a custom homemade

hot press heated up at 60ᵒC, in order to optimize the contact between polymer and ceramic.

A picture of the assembly is shown in Figure 1 b).

Lithium metal chips of 150 µm thick and nickel tabs as current collectors are then added

to the cell, and finally the assembly is vacuum sealed (Showa-Denka) in an air tight pouch

material in order to carry out the experiment outside the glove box. A schematic of the

symmetric composite-lithium assembly is shown in Figure 1.

It is worth noting that the assembly of BCE onto a fragile ceramic is not trivial and

delicate. The assembly of the composite with the high Mw SEO 386-300 BCE has been

experimented. However, this polymer presents poor adherence properties onto the ceramic

surface and it was impossible to obtain a correct assembly. Therefore, unfortunately this

composite is not studied.

In the case of the two single-ion BCE, that are very soft above the Tm, the composite

cells obtained exhibit a nice and very homogeneous surface as it is shown in Figure 1 a).

However, in the case of the SEO 55-52 BCE, the composite surface presented signs of

defective adherence due to the presence of hard PS blocks.


- 166 -

a) b)

Figure 1. a) Picture of the polymer-ceramic composite cell using single-ion BCE and b) Schematic of the assembly for a composite-lithium symmetric cell.

EIS measurement. Cells ae characterized by EIS using a Bio-Logic VMP3 potentiostat

driven by EC Lab software1. The applied AC voltage was 50 mV and the analysis is

performed over a frequency spanning from 1 MHz to 1 Hz.

Ionic conductivity measurement. Cells are placed inside a custom homemade heating

stage and ionic conductivity are carried out from 30ᵒC to 90ᵒC, with 10ᵒC step. Samples are

allowed to equilibrate at each temperature for 30 min after temperature stabilization. The

same heating and cooling cycles as in chapter 3 are used, i.e. a first cycle from room

temperature to 90ᵒC, followed by a cooling cycle to 30ᵒC, and finally a second heating cycle to

90ᵒC.

Cycling. Composite-lithium symmetric cells are galvanostatically cycled using a VMP3

potentiostat in a custom homemade heating stage at 90ᵒC. Cells are cycled following the same

cycling routine as in chapter 4 for the dendritic growth experiment, in other words pre-

conditioning cycles are performed at 0.02 mA.cm-2 for 2 hours, followed by cycling performed

at 0.175 mA.cm-2 for 4 hours. Electrochemical impedance spectroscopy are also carried out

onto the cells during the rest period (after 30 minutes of rest) after each charge and discharge.


- 167 -

2. Results and discussions

a. Electrical properties

Ionic conductivity measurements ae performed on the three different materials used,

which are the Ohara ceramic (Chapter 2), block copolymer electrolytes (Chapter 3) and finally

in this section, we will focus on the ionic conductivities of the composite.

A schematic representation of half of the composite cell is presented in Figure 2 a), we

describe here the different phenomena possibly occurring in our cell and the different

contributions expected in EIS measurement. Characteristic frequencies of the different

contributions are added to the expected EIS graph (shown Figure 2 b)), based on the

frequencies observed in previous chapters.

Due to the presence of the Ohara GC (grain and grain boundaries) and the BCE, we

expect to see their contributions as three loops. However, due to time constants, which are

close for the grain boundaries and the BCE, only two discernable loops at high frequencies

should be observed. At middle high frequencies, the lithium/polymer interface, composed by

the SEI contribution and the lithium charge transfer contribution, is expected. In addition, the

polymer/Ohara GC interface, corresponding to the lithium ion charge transfer between the

BCE and the ceramic can be expected in the same range of frequencies. Finally, in the case of

SEO electrolytes, due to its low transference number (inferior to 0.22,3), a contribution due to

the diffusion is expected at low frequencies.

The theoretical impedance spectrum from high frequencies to low frequencies (black dash

line) is shown in Figure 2. However, due to the VMP3 constraint at high frequency, EIS

measurements are performed only from 1 MHz and due to the time constraint the

measurement are stopped at 1 Hz. Therefore, the spectrum expected is represented by black

solid line in Figure 2b). An electrical equivalent circuit for the expected impedance spectrum

is presented in Figure 2 c). All the phenomena are expected to be in series.


- 168 -

a)

b)

c)

Figure 2. a) Schematic representation of a half composite-lithium cell, b) theoretical electrochemical impedance spectrum for a composite-lithium symmetric cell. The different contributions are presented in color, when the expected spectrum is presented in black dash line. c) Electrical equivalent circuit for the theoretical spectrum with: Lc the cable inductance, Rc the cable resitance, Rg and CPEg, the GC grain resistance and constant phase element respectively, Rgb and CPEgb the grain boundaries resistance and CPE, Rinter BCE/GC and CPEinter BCE/GC the polymer-Ohara GC interface resistance and CPE, Rinter Li/BCE and CPEinter Li/BCE, the lithium-polymer interface resistance and CPE, and finally in the case of SEO BCE WSEO, the SEO short Warburg.


- 169 -

Typical spectra for a composite (Ohara GC-SEO 55-52) lithium symmetric cell at 30ᵒC

and 90ᵒC are shown in Figure 3 a) and c) respectively. Zooms at high frequencies are

presented for both cases in Figure 3 b) and d).

The spectrum at 30ᵒC presents a first loop at high frequency, which corresponds to the

Ohara GC contribution and a second loop at middle high frequency (42 kHz) corresponds to

the SEO BCE contribution. When the temperature is increased, the characteristic frequency

increases. At 90ᵒC, three contributions are observed. The first one at high frequencies should

represent the composite contribution, i.e. the Ohara GC and SEO contributions. The second

one at medium frequency is clearly deformed. We suppose that this loop is composed of the

SEO-Ohara GC interface and the lithium-polymer interface contributions. The last

contribution at low frequency represents the diffusion of ionic species into the SEO

electrolyte. Unfortunately, the transference number of this BCE was not measured, but we

can assume a t+ < 0.22.

a) b)

140x103

120

100

80

60

40

20

0

-Im

(Z)

(W.c

m2)

140x103120100806040200

Re(Z) (W.cm2)

5 Hz

SEO + Ohara GC at 30°C20x10

3

15

10

5

0

-Im

(Z)

(W.c

m2)

20x103151050

Re(Z) (W.cm2)

Composite contribution

42 kHz

SEO + Ohara GC at 30°C


- 170 -

c) d)

Figure 3. Nyquist plot obtained from a lithium/composite with SEO 55-52 /lithium cell a) at 30ᵒC, b) zoom at

high frequencies and b) at 90°C with d) a zoom at high frequencies at 90ᵒC.

In addition, in chapter 1 we reported that the characteristic frequency for the lithium-

polymer interface is around 20 Hz, therefore only the low frequency part of the second loop

should correspond to the Li-polymer interface.

Figure 4. Nyquist plot obtained from a) a lithium/composite with Si-EO/lithium cell at 90°C and b) a

lithium/composite with Si-EO-Si/lithium at 90°C.

A typical spectrum for a composite-lithium symmetric cells with the Si-EO and the Si-

EO-Si BCEs are shown in Figure 4 a) and b) respectively. Only two semi circles are observed.

The first one at high frequencies represents the composite electrolyte, i.e. Ohara and Si-EO or

100

80

60

40

20

0

- Im

(Z)

( W.cm2)

100806040200

Re(Z) /(W.cm2)


- 171 -

Si-EO-Si contributions, the second one at medium frequencies represents the lithium-polymer

interface contribution and Ohara GC/BCE interface. The absence of another loop which

would correspond to the Ohara GC-polymer interface implies that the charge transfer at this

interface is not dominant and small.

In the case of the Si-EO BCE (see in Figure 4 a)), the characteristic frequency of the

composite electrolyte is slightly higher than the polymer alone (0.56 MHz for the composite

and 0.18 MHz for the polymer). This increase in the characteristic frequency is due to the

contribution of the ceramic, which presented a higher characteristic frequency compared to

the anionic block copolymer at 90ᵒC. Besides, the interface characteristic frequency for both

polymer and composite are in the same range, 422 Hz and 505 Hz respectively, meanings that

the interface is similar in both cases. In fact the interface resistance in the composite cell is

due to the direct contact between lithium and Si-EO BCE, therefore the SEI should be the

same as the SEI formed in Si-EO/lithium symmetric cells. Thus, following the interface

frequency should indicate when lithium touches the ceramic.

In the case of the Si-EO-Si BCE (see in Figure 4 b)) , the characteristic frequencies for the

composite and the polymer are found to be close at 0.21 MHz and 0.24 MHz respectively. In

addition, the interface characteristic frequency are also similar in both cases, 4,2 kHz for the

Si-EO-SI/lithium symmetric cell and 1.3 kHz for the composite. However, the interface

resistance for the composite is found higher compared to the polymer, 800 Ω.cm2 and 50

Ω.cm2 respectively. This difference could be due to the lack of good contact between the

polymer and the lithium, because it is difficult to press lithium onto the composite, due to the

brittleness of the Ohara ceramic.

In Figure 2 c), a theoretical equivalent circuit have been proposed, however the different

contributions are not all observed in ours experimental spectra. Firstly, because in the

frequency domain explored, the grain contribution of the Ohara GC is presumed to be a

simple resistance at high frequencies. In addition, due to the close time constants for the grain

boundaries of the Ohara GC and BCE, we observe only one loop. In the case of the single-

ion BCE, the polymer-Ohara GC interface is not visible, therefore we consider only the

lithium-polymer interface at middle high frequencies. Thus, the composite-lithium cells can be

modeled via a simplified equivalent circuit compared to the one presented in Figure 2 c). They

are presented in Figure 5 a) and b) for the composite with single-ion BCE and SEO BCE

respectively. The composite contribution is modeled with a resistance in parallel with a

constant phase element, Rcompo and CPEcompo respectively, which represent BCE contribution


- 172 -

added to the grain boundaries of the Ohara GC contribution. The polymer-lithium interface is

modeled with another resistance in parallel with a CPE, R inter and CPEinter.. In the particular

case of the composite with SEO electrolyte, another RinterSEO/GC //CPEinterSEO/GC is added to

model SEO/Ohara GC interface. In addition, here, the diffusion contribution is not take into

account for the fit.

Figure 5. a) Equivalent circuit for lithium-composite-lithium cells with single-ion BCE and b) Equivalent circuit for lithium-composite-lithium cells with SEO BCE.

Due to the high impact of the BCE crystallization on the ionic conductivity, only the

behavior of the composite for temperatures above its melting temperature is discussed. The

ionic conductivities of the composite are calculated from the fitted values. The ionic

conductivities of the composite cell with the SEO 55-52 are presented in Figure 6. Figure 7 to

Figure 8 compare ionic conductivities performed on gold/ceramic/gold cells (Chapter 2),

BCE lithium symmetric cells (Chapter 3) and composite lithium symmetric cells., whereas the

ionic conductivities for the Si-EO and Si-EO-Si BCE are presented in Figure 7 and Figure 8

respectively. At 90ᵒC, the ionic conductivity of the ceramic is found to be almost two order of

magnitude higher than BCE, 1.05.10-3 S.cm-1. The effective ionic conductivity of the

composite follows the BCE trend, i.e. a severe drop in ionic conductivity due to PEO

crystallization. Nevertheless, the effective ionic conductivity increases by a factor three for the

Si-EO BCE at 90ᵒC, 1.04.10-4 S.cm-1 and by a factor six for the Si-EO-Si BCE 8.82.10-5 S.cm-1.

In the case of SEO 55-52 BCE (Figure 6) a very low conductivity inferior to the polymer itself

is obtained for the composite cell. This unexpectedly poor effective ionic conductivity is

possibly due to the rigid character of the SEO. During the assembly process, we noticed the

poor wettability of the SEO with the Ohara GC. Therefore, we can assume that only a part of

the surface Ohara GC/SEO is really intimate.


- 173 -

In order to confirm that the interface between the polymer and the Ohara are suitable in

the composite cell, the effective ionic conductivities are calculated by simply taking into

account the BCE and ceramic contributions in series (equation (1) and (2)). The result of the

effective conductivity has been added in the Arrhenius plot Figure 6 to Figure 8 (solid line).

Reff = Rp+RGC = l/Sgeo.σp+l*/Sgeo.σ GC = (l+l *)/Sgeo.σeff (1)

σeff = (l!+!l!*)!"p!.!"GC

l!"p!+!l!*!.!"GC! (2)

With

· Rcompo, Rp, RGC : resistance of the composite, the polymer, the Ohara GC respectively

(Ω)

· l and l * : thickness of the BCE and the Ohara GC respectively (cm)

· Sgeo : surface area (cm2)

· σcomposite, σp, σGC : ionic conductivity of the composite, the polymer and the Ohara GC

respectively (S.cm-1)

Figure 6. Comparison of conductivity of Ohara GC, SEO and the composite with SEO 55-52 BCE. The solid line corresponds to the simple series contributions of the Ohara GC and SEO BCE.

10-5

10-4

10-3

Co

ndu

ctivity (

S.c

m-2

)

3.33.23.13.02.92.82.7

1000/T (K-1

)

Temperature (°C)

'Ohara GC' 'SEO 5552' 'Composite with SEO 52-55' Model series

90 80 70 60 50 40 30


- 174 -

Figure 7. Comparison of conductivity of the Ohara GC, Si-EO and the composite. with Si-EO BCE. The solid line corresponds to the simple series contributions of the Ohara GC and Si-EO BCE

Figure 8. Comparison of conductivity of the Ohara GC, Si-EO-Si and the composite with Si-EO-Si BCE. The solid line corresponds to the simple series contributions of the Ohara GC and Si-EO-Si BCE

In the molted state, the series model for single-ion BCEs is in a very good agreement with

the experimental results, showing that the interfaces are good. On the contrary, in the case of

the SEO55-52 the series model is a factor 3 or 4 below the experimental results. This may

10-9

10-8

10-7

10-6

10-5

10-4

10-3

10-2

10-1

Co

ndu

ctivity (

S.c

m-1

)

3.43.33.23.13.02.92.82.7

1000/T (K-1)

90 80 70 60 50 40 30 20

Temperature (°C)

'Ohara GC' 'PEO-b-PSTFSI' 'PEO-b-PSTFSI + Ohara GC' Model series

10-8

10-7

10-6

10-5

10-4

10-3

10-2

10-1

Co

ndu

ctivity (

S.c

m-1

)

3.43.33.23.13.02.92.82.7

1000/T (K-1

)

90 80 70 60 50 40 30 20

Temperature (°C)

'Ohara GC' 'PSTFSI-b-PEO-b-PSTFSI' 'PSTFSI-b-PEO-b-PSTFSI + Ohara GC' Model series


- 175 -

confirm our previous observation that the true contact area between the ceramic and the BCE

is far less than the geometric contact area due to a lower polymer elasticity in this case. The

proportion of contact, α, is therefore estimated in first approximation according to equation

(6), where Rtheotot and Rexp

tot are determined according to equations (3) to (5) and where Sexp

corresponds to the surface in contact at the interface Ohara GC/SEO.

Refftheo!=!Rp!+!RGC!! ! ! ! !!!!(3)!

!! ! ! ! !!Refftheo!=!

1

Sgeo!(lp

"p!+!

lGC

"GC)!!!!!! ! ! ! !!!!(4)!

! ! ! ! Reffexp!=!1

Sexp!(lp

"p+!

lGC

"GC)!! ! ! ! !!(5)!

! ! ! ! !!!!!#!=Sexp

Sgeo!=!

Rtheotot

Rexptot! (6)

Below the melting temperature, α is very low with 23%, and it increases slightly with

temperature. Above the Tm, α stabilizes around 30% and at 90ᵒC the effective surface is 33%

of the theoretical surface.

b. Cycling

Before describing the result of the cycling experiment, it is important to note that the

composite-lithium cell does not have the same behavior as the polymer-lithium symmetric

cell. Therefore, this paragraph intends to clarify the differences which are expected in the case

of composite-lithium symmetric cells.

In the case of polymer-lithium symmetric cell, when a dendrite grows it can reach the

opposite electrode creating a short circuit and thus the voltage of the cell will drop to zero

abruptly4 (as seen in Chapter 4.1.a)) However, in the composite-lithium symmetric cells the

presence of the Ohara GC is a mechanical and physical obstacle for lithium dendrite. A

schematic representation of composite-lithium symmetric cells at different stages during

cycling is given in Figure 9.


- 176 -

Figure 9. Scheme representing the evolution of lithium dendrite inside a composite-lithium symmetric cell and the voltage associated.

The first stage corresponds to the initial conditions and the first cycles, i.e. when the cell is

pristine (step 1 in Figure 9). In the case of the single-ion BCE, the voltage is constant over

cycling. The second stage corresponds to the nucleation and growth of dendrites. The stage 3

coincides to the contact between lithium and ceramic. Indeed, when a dendrite crosses the

whole polymer layer, it won't induce a short circuit but it will first reduce the ceramic (stage 3

in Figure 9). The reaction between lithium and the Ohara GC leads to a local volume

expansion inside the ceramic implying mechanical stress, which finally leads to the break of

the ceramic (stage 3 and 4 in Figure 9). In addition, the reaction products are insulating thus

the voltage of the cell will increase. After cycling, the analysis of the post mortem cells

showed that the glass-ceramic became black on the impact points with lithium metal.

Therefore, the Ohara GC becomes an electronic conductor where it is reduced. Finally, the

stage 4 describes the moment when the cell experiences a short circuit due to the cracks right

through the ceramic, which enables an electronic pathway leading to a short circuit.

Now that the expected behavior of composite-lithium symmetric cells is clarified, the

experimental results of cycling is presented. The EIS is measured after each charge and

discharge. Each spectrum is modeled and fitted using the equivalent circuit showed in Figure

5 a) for the composite with single-ion BCE and in Figure 5 b) for the composite with SEO


- 177 -

BCE. Two parameters are mainly extracted: the electrolyte resistance, Rel, and the interface

resistance, Rint.

Composite with the SEO BCE. The voltage versus time during cycling at 0.175mA.cm-2

at 90ᵒC for the composite-lithium symmetric cell is presented in Figure 10 a). The pre-

conditioning cycles are not presented here. The voltage decreases from 0.6V to 0.4V during

the first 10 cycles after which the voltage is constant. The decrease in voltage during the first

cycles potentially comes from the breaking of the passive layers (SEI), as well as the formation

of fresh lithium at the interface polymer-lithium.

Figure 10. a) Voltage versus time graphic for a composite with SEO 55-52 lithium symmetric cell cycled at 90ᵒC at 0.175 mA.cm-2 and b) a zoom of the cycles 5 to 10.

A zoom on four cycles is given in Figure 10 b), the voltage increases abruptly when the

polarization starts but the steady state is not reach instantly. In addition, when the polarization

is stopped the voltage relax slowly, due to the relaxation of the gradient of concentration.

0.8

0.6

0.4

0.2

0.0

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Vo

lta

ge

(V

)

8006004002000

Time (h)

After charge After discharge

a)-0.6

-0.4

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0.2

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0.6

Po

ten

tia

l (V

)

80757065605550time (h)

Composite with SEO BCE

b)


- 178 -

Figure 11. a) Nyquist plots obtained from a lithium/composite with SEO 55-52/lithium cell at 90°C after 1, 17, 50 and 91 cycles at 0.175mA.cm-2 and b) zoom at high frequencies.

Four EIS spectra, for a composite cell with SEO 55-52 electrolyte, after respectively 1, 17,

50 and 91 cycles at 0.175 mA.cm-2 are represented in a Nyquist plot in Figure 11. After the

first cycle, the lithium symmetric cell presents a high interface resistance between the polymer

and lithium. As expected from the drop of polarization, during cycling, Rint decreases strongly

and stabilizes, confirming our assumption about the interface evolution. In addition, the

characteristic frequency for Rint increases over cycling, characterizing the evolution of its

chemistry. The electrolyte contribution increases over cycling resulting in a decrease in ionic

conductivity.

Figure 12 plots the electrolyte and the interface resistances, extracted from the EIS

measurements for a composite cell with SEO 55-52 electrolyte. Figure 12 a) shows the

electrolyte resistance as a function of charge passed. This cell was cycled for 92 cycles which

corresponded to 467 C.cm-2, before it was stopped (due to time constrain). The ionic

conductivity is coherent with the previous result presented in the section above. The

electrolyte resistance remains almost stable around 45 Ω.cm2 for the first 20 cycles, then Rel

increases linearly up to 65 Ω.cm2 after 92 cycles which represents an increase of 45%. This

could be due especially to a partial loss of contact during cycling (due to the bad adherence of

this BCE onto the Ohara GC).

Rint decreases rapidly during cycling (Figure 12 b)), then it stabilizes around 50 Ω.cm2. The

low frequency loop and the voltage (see in Figure 10) suggests that no short circuit has


- 179 -

occurred. However, the strong decrease in Rint until a stabilization that suggests i) the SEI are

completely modified in term of chemistry and ii) that the active surface area probably

increases due to heterogeneous lithium electro-deposit. However, the small Rint suggests that

lithium metal is not in contact with the ceramic, meanings that the "dendrites" have not cross

all the BCE protective layer after 92 cycles.

Figure 12. SEO 55-52-ceramic lithium symmetric cell a) electrolyte resistance versus charge passed and b)

interface resistance versus charge passed at 90ᵒC.

Unfortunately, due to time constraints this cell could not be imaged by X-ray micro-

tomography to verify the status of the ceramic.

Composite with the Si-EO BCE. The voltage versus time for the composite-lithium

symmetric cell with PEO-b-PSTFSI BCE is shown in Figure 13 a). The voltage is stable for

the first ten cycles, then it increases. We observe during the discharge of the 10th cycle an

abrupt decrease of voltage and a similar behavior during the charge of the 13th cycle with an

increase of voltage is also observed. For more clarity, a zoom of the voltage versus time for

these cycles is shown in Figure 13 b). This is probably the signature of the direct contact

between a lithium dendrite and the ceramic as seen in stage 3 in Figure 9. Therefore, due to

the presence of this behavior in discharge for the 10th cycle and in charge for the 13th cycle, we

can suppose that lithium has reached the ceramic on both of its side. Two other zooms are

presented in Figure 13 c) to show typical signatures of short circuits, as seen in stage 4 in

Figure 9. This abrupt drop of voltage is followed by a healing, with the voltage that comes

100

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Rel (W

.cm

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a)

350

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100

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After charge : composite SEO 55-52 After discharge : composite SEO 55-52

b)


- 180 -

back to the standard polarization. This effect has been discussed as a fuse effect in the

literature4.

Figure 13. a) Voltage versus time graphic for a composite with PEO-b-PSTFSI BCE-lithium symmetric cell at 90ᵒC, b) zoom on cycles 10 to 13, and c) zoom on cycle number 14

Four EIS spectra, for a composite cell with the Si-EO BCE electrolyte, after respectively

1, 17, 25 and 32 cycles at 0.175 mA.cm-2 are represented in a Nyquist plot in Figure 14. After

the first cycle, the lithium symmetric cell exhibits a high lithium-polymer interface resistance.

During the first cycles, Rint is stable and after 15 cycles it starts to increases. In addition, the

electrolyte contribution doubles between the 17th and 25th cycle (associated with a decrease of

the characteristic frequency), leading to a decrease of the ionic conductivity.

-0.4

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0.0

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0.4

Voltage

(V

)

300250200150100500Time (h)

Composite with Si-EO BCEa)

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0.4

Vo

lta

ge

(V

)

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b)

Cycles 10 to 13

-0.4

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lta

ge

(V

)

134132130128126124Time (h)

Cycle number 14

c)


- 181 -

Figure 14. a) Nyquist plots obtained from a lithium/composite with PEO-b-PSTFSI/lithium cell at 90°C after 1, 17, 24 and 32 cycles at 0.175mA.cm-2 and b) zoom at high frequency.

The evolution of the two parameters Rel and Rint is shown in Figure 15 as a function of the

total exchanged electricity. This cell was cycled for 32 cycles, corresponding to 170 C.cm-2,

before it was stopped for characterization at the micro-tomography beamline. The ionic

conductivity initially measured is around 2.10-4 S.cm-1, which is coherent with the previous

measurements (see previous section). The Rel remains constant after charge and discharge for

the first 100 C.cm-2 (Figure 15 a)), i.e. the first 18 cycles and the value is around 150 Ω.cm2. It

then increases in few cycles to a constant value around 350 Ω.cm2. This increase of 140% of

the initial electrolyte resistance could be due to lithium dendrites, which pass through the

single-ion protective layer and touch the Ohara GC, leading to its reduction and fractures

which result in a loss in contact. Therefore, the net result is the increase of the whole

resistance of the composite.

The interface resistance versus the charge passed is shown in Figure 15 b). At the

beginning Rint is relatively constant around 1400 Ω.cm-2 and an increase up to 1750 Ω.cm-2 is

then observed at the same cycle number. This increase corresponds to a rise of 25% of the

initial Rint, which is quite significant. In chapter 4.1.c, we saw a decrease of 20% in the

interface resistance in the case of Si-EO-Si/lithium symmetric cells during cycling. Rint is

correlated to Rel, however it is less pronounced than the Rel and can come from a reduce

section for the current to go through. The increase observed is therefore probably due to the

contact between the lithium and the ceramic. Indeed, the interface between lithium and an

insulating compound is more resistive than with the single-ion BCE.


- 182 -

Figure 15. PEO-b-PSTFSI ceramic lithium symmetric cell a) electrolyte resistance versus charge transferred and

b) interface resistance versus charge transferred at 90ᵒC.

Composite with the Si-EO-Si BCE. The voltage versus time for the cycled composite-

lithium symmetric cell with the single-ion triblock copolymer PSTFSI-b-PEO-b-PSTFSI BCE

is presented in Figure 16 a). One remarkable aspect is that the voltage reaches almost

instantaneously the value of the steady state voltage under polarization as well as a direct drop

to zero when the cell is relaxed. This feature are characteristic of single-ion systems. The

voltage decreases during the first 2 cycles and then stabilizes over cycling. A zoom of one

typical cycle (red rectangular in Figure 16) is given in Figure 16 b). We observe a dissymmetry

between the charge, where the voltage is relatively constant and the discharge, where the

voltage decreases linearly. But more importantly, no signs of dendrites have been observed in

the voltage profile.

500

400

300

200

100

0

Rel (W..cm

2)

16012080400


)

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Cycle number

'After charge : composite with diSI' 'After discharge : composite with diSI'

a)

2000

1800

1600

1400

1200

1000

Rin

t (W

.cm

2)

16012080400


)

302520151050

Cycle number

After charge : composite with diSI 'After discharge : composite with diSI'

b)


- 183 -

Figure 16. a) Voltage versus time for a cycled composite with PSTFSI-b-PEO-b-PSTFSI BCE-lithium symmetric cell at 90ᵒC and b) zoom for the characteristic cycle in the red rectangular.

Four EIS spectra, for a composite cell with PSTFSI-b-PEO-b-PSTFSI electrolyte, after

respectively 1, 17, 25 and 32 cycles at 0.175 mA.cm-2 are represented in a Nyquist plot in

Figure 17. The characteristic frequency of the composite contribution is at 0.35 MHz, which

is coherent with characteristic frequencies of the polymer and the grain boundary of the

ceramic. The lithium polymer interface resistance shows characteristic frequency at 1.2 kHz,

which is in a very good agreement with the characteristic frequency determined in chapter 3,

4.2 kHz for the Rint between the Si-EO-Si and lithium at 90ᵒC. During cycling, the composite

electrolyte resistance stays constant, whereas the interface resistance slightly decreass at first

and stays almost constant along cycling.

-0.4

-0.2

0.0

0.2

0.4

Vo

lta

ge

(V

)

300250200150100500Time (h)

Composite with PSTFI-b-PEO-b-PSTFSI

a) -0.3

-0.2

-0.1

0.0

0.1

0.2

0.3

Voltage (

V)

134132130128126124Time (h)

b)


- 184 -

Figure 17. Nyquist plots obtained from a lithium/composite with PSTFSI-b-PEO-b-PSTFSI/lithium cell at

90°C after 1, 17, 24 and 32 cycles at 0.175mA.cm-2.

Both composite cells with single-ion block copolymer (with Si-EO and Si-EO-Si BCEs)

were started at the same time and stopped together in order to image the cells, i.e. 32 cycles.

The EIS experiment results for the composite with the Si-EO-Si BCE are presented in Figure

18. The electrolyte resistance versus charge passed is shown in Figure 18 a). On the contrary

to the composite with Si-EO BCE, the electrolyte resistance remains relatively constant over

the 32 cycles. However, a slight difference in the Rel after charge and discharge is observed.

The initial Rel after charge is at 700 Ω.cm2 and decreases to 600 Ω.cm2. The fact that the

electrolyte resistance during cycling does not increase suggests that no dendrites crossed the

BCE protective layer, which is in good agreement with the high dendrite resistance already

observed for this BCE in Chapter 4.


- 185 -

Figure 18. PSTFSI-b-PEO-b-PSTFSI ceramic lithium symmetric cell a) electrolyte resistance versus charge

passed and b) interface resistance versus charge passed at 90ᵒC.

The interface resistance as a function of charge passed is plotted in Figure 18 b). Initially,

Rint is different after charge and after discharge, 940 Ω.cm2 and 1040 Ω.cm2 respectively. A

decrease in the resistance is then observed, followed by a stabilization around 800 Ω.cm2. A

similar behavior is seen in the case of the Si-EO-SI lithium symmetric cell (chapter 4). The Rel

and thus the ionic conductivity remains constant over cycling, suggesting that the ceramic is

pristine. All this results shows that the composite-lithium symmetric cell works very well

without any dendrite that touch the Ohara GC.

We were obliged due to the beamline time for the tomography experiment to stop this

cells.

c. Characterization by X-Ray microtomography

In order to observe the different parts of the lithium symmetric cells at the initial state and

after cycling, but also to confirm our hypothesis about the state of the Ohara GC, hard X-ray

micro tomography were performed on the cells.

The presence of a ceramic composed of heavy elements, such as germanium, combined

with the impossibility to cut the composite-lithium symmetric cell into smaller sections, made

the imaging more complex. Since hard X-ray micro-tomography lies on the transmitted X-ray

beam, heavy elements absorb a large quantity of the emitted signal leading to a very low

transmitted signal, which results in a very low resolution, or even no transmitted signal, if the

1000

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Rel (W

.cm

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Cycle number

'After charge : composite with triSI' 'After discharge : composite with triSI'

a)

1400

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800

600

400

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.cm

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16014012010080604020


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302520151050

Cycle number

After charge : composite with triSI After discharge : composite with triSI

b)


- 186 -

elements absorb all the emitted signal. To avoid complete extinction of the emitted signal, the

composite-lithium symmetric cells are imaged in a vertical orientation as shown in Figure 19

a). This orientation enables the unfavorable orientation of the ceramic to be reduced,

compared to when the whole length of the ceramic is in the beamline (see in Figure 19 b)).

Indeed, when the beam has to cross the whole length of the ceramic, the transmitted signal is

very poor and it is impossible to discern any materials.

a) b)

Figure 19. a) Vertical orientation of a composite-lithium symmetric cell on the rotating stage for hard X-ray micro-tomography experiment b) unfavorable orientation of the cell due to the diameter of the ceramic.

A slice through a typical tomogram obtained from an uncycled composite-lithium

symmetric cell is shown in Figure 20. The image shows the cross section of an uncycled

lithium symmetric cell and it is dominated by four phases. The two thin, bright and horizontal

phases correspond to the polymer electrolyte, here, the anionic block copolymer PEO-b-

PSTFSI, whereas the two surrounding darkest phases, on the top and the bottom, correspond

to lithium metal electrodes, and finally the dark thick phase sandwiched inside the two

polymer layers corresponds to the Ohara GC. The uncycled cell image presented a low

resolution due mainly to the presence of the pristine ceramic. However, a smooth contact

between the polymer and the ceramic is observed in addition with the nice contact between

the polymer and lithium. Bad contact between two materials would be represented by a dark

black signature in hard X-ray micro-tomography. However, here no such signature is visible.

The large semi-circles observed are due to dust on the scintillator and could not be completely

removed during the image reconstruction. Polymer layers are homogeneous on both sides of

the ceramic. The uncycled composite-lithium symmetric cell is devoid of any other noticeable

features (see Figure 20). Here, the tomogram does not reveal the presence of the crystalline


- 187 -

impurities in the lithium electrode, however it is important to remember that those impurities

are nevertheless present and are randomly dispersed.

Figure 20. X-ray tomography slice showing the cross section of a composite with PEO-b-PSTFSI-lithium symmetric cell before cycling.

The composite-lithium symmetric cell was imaged after cycling for 170 C.cm-2, a slice

showing the cross section through the cell is shown in Figure 21. The main noticeable feature

is that the ceramic is broken. It is known that the Ohara GC is not stable versus lithium, the

transition metal (Ti) is reduced leading to a volume expansion. One example of such reaction

is observed inside the white circle in Figure 21. Thus, the lithium metal is in direct contact on

this part of the Ohara GC. However, this inflation has not yet lead to local stress, which will

induce later a fracture of the ceramic. We observe several fractures on the right side of the

slice in Figure 21, meanings that lithium metal have already reduced the ceramic on another

part of the ceramic, and the fractures were widespread. This observation is in agreement with

our results obtained by EIS, i.e. lithium dendrite reached the ceramic and reduced it. It is

important to note that no impurities are observed close to the supposed dendrites. This result

is also in agreement with our previous experiment with the single-ion BCE-lithium symmetric

cells in Chapter 4.2.d). In addition, it is worth to note that the current densities, where the

ceramic is broken, are probably high implying a deterioration of the single-ion BCE4. Thus, it

is possible that the Si-EO BCE layer delaminates from the Ohara GC. Nevertheless, due to

the fractures of the ceramic the resolution is low and it is hard to distinguish such

phenomenon.


- 188 -

Figure 21. X-ray tomography slice showing the cross section of a composite with PEO-b-PSTFSI-lithium symmetric cell after cycling.

Another interesting feature is the irregularity of the interface polymer-lithium in some

parts of the cell, which is clearly observed on the left side of the cell, where the ceramic is

unbroken. The morphology of the interface BCE-lithium is in good agreement with the

previous results shown in Chapter 4.2.d). We observe here, an irregular dense electro-

deposition of the lithium metal, probably due to the local heterogeneities in the SEI or in the

Si-EO BCE. In other parts, the interface lithium-Si-EO is devoid of any features meanings

that the lithium deposition and stripping is homogeneous.

The hypothesis made in section 2 on the cycling behavior of the composite-lithium

symmetric cell is confirmed by hard X-ray micro-tomography. Indeed, the lithium dendrites

went through the Si-EO layer and touched the ceramic leading to its reduction followed by

fractures provoking the increase of the electrolyte resistance and the interface resistance. It is

interesting to note that almost 50% of the surface of the Ohara GC is broken (it has been

confirmed by the other slices), which corresponds to the increase observed for the composite

electrolyte resistance.

Despite the low resistance to dendrites over cycling, this result is very encouraging due to

the poor mechanical properties of the Si-EO BCE at 90ᵒC.

Similar imaging experiments were performed onto the composite using the triblock single-

ion BCE, the Si-EO-Si. A typical slice through the cell before cycling is presented in Figure

22. The same four different phases are observed as in the cell described above. The polymer

layers are homogeneous and the interfaces with both the ceramic and lithium metal are

smooth. No other noticeable features are observed on the uncycled cell.


- 189 -

Figure 22. X-ray tomography slice showing the cross section of a composite with PSTFSI-b-PEO-b-PSTFSI-lithium symmetric cell before cycling.

Here, after 32 cycles at 0.175 mA.cm-2 (170 C.cm-2) the cell is imaged by hard X-ray

micro-tomography, a slice through the cell is shown in Figure 23. The first important

observation is that the Ohara GC is pristine without any fractures, suggesting that no lithium

dendrite have reached the ceramic. Because the Ohara ceramic is pristine the resolution of the

tomogram slice is poor.

Figure 23. X-ray tomography slice showing the cross section of a composite with PSTFSI-b-PEO-b-PSTFSI-lithium symmetric cell after cycling.

However, we observe that the interfaces polymer-lithium are slightly deformed compared

to their initial state (Figure 22). This observation is in agreement with the previous results

presented in Chapter 4.2.d), where irregular but dense electro-deposition of lithium has been

observed. The growth of lithium is dense, but lithium objects grows preferentially in some

places, leading to a non planar growth front. In addition, we have previously reported that the

Si-EO-Si exhibits a high resistance to dendrite growth which is confirmed here.


- 190 -

II. Quantification of polarization loss at the polymer-

ceramic interface

1. State of the art

Previous studies5–7, on lithium-ion conducting ceramic with liquid electrolytes carbonate

based, have revealed the presence of an additional well resolved semi-circle at low frequency

in Nyquist plots, which is attributed to the lithium-ion charge transfer between solid and

liquid electrolytes. In addition, Abe et al.5 have shown that the activation barrier at the

interface is affected by the solvating energy of Li+ in each solvents in liquid electrolytes.

Similar interfacial resistances with frequency resolved R // Q processes in Nyquist plots have

been reported for PEO20: LiCF3SO3 laminated on La0.55Li0.35TiO3 (LLT)8, with an activation

energy of 1.02 eV and in laminated thin film electrolytes consisting of PEO-LiCl4 and

LiPON9. However, more recently Tenhaeff et al.10 have reported a laminated electrolyte

structure composed of Ohara ceramic and PEO10: LiTFSI, which presented Nyquist plots

without additional resolved semi-circle at medium or low frequencies. In addition, they

reported an interfacial resistance of 47 Ω.cm2 at 40ᵒC. The small interface resistance is

attributed to their fabrication process, which ensures good contact between the ceramic and

the polymer electrolyte.

Gondran et al.12 have calculated the PEOx-NaI-(AgI)0.25/NaSICON interfacial

contribution in a four electrode cell using EIS as a function of the inverse of temperature and

their results are presented in Figure 24. They showed that this interface contribution is high

and increase with temperature. In addition, they showed that the ceramic-polymer interface is

more influenced by the ionic conductivity compared to the recrystallization process and that

the interfaces are related to the concentration of charge carriers.


- 191 -

Figure 24. NaSICON/PEOx-NaI(AgI)0.25 interfacial resistance versus the inverse of temperature12.

A recent study by Mehrotra et al.7 has highlighted a high polarization loss at the liquid

electrolyte-ceramic interface. They have reported that the liquid electrolyte/ceramic junction

polarization could be significant around 0.5 kΩ.cm2 (in the case of EC:DEC (1:1), DMSO and

PC solvents). The ion charge transfer at the liquid electrolyte/ceramic interface has been

widely studied by Ogumi group5,6,8,11, they have been argued that this large polarization loss is

due to lithium ion desolvation/solvation process at the interface liquid electrolyte/ceramic.

2. Experimental procedure

Cells preparation. In this section, we will endeavor to quantify the polarization loss at

the polymer-ceramic interface. For this purpose, similar composite cells are assembled

following the same procedure as in the section above. The three different block copolymers

are used in this study: the neutral block copolymer SEO 55-52, and the two single-ion block

copolymers, PEO-b-PSTFSI and PSTFSI-b-PEO-b-PSTFSI.

Composite-lithium symmetric cells are produced. Their preparation and assembly have

been described in the previous chapters.

EIS measurement. Prior to the cycling experiment, cells are characterized via EIS using

a VMP3 potentiostat driven by Ec Lab software. The applied AC voltage is 20 mV, the

measurement was performed over a frequency span from 1 MHz to 1 Hz. The cells are placed

inside a climatic chamber heated at 90ᵒC and equilibrated at this temperature for 6 hours

before running the measurements.


- 192 -

Direct current experiment. Prior the DC experiment, pre-conditioning cycling is

performed at low current density (0.02 mA.cm-2) for 15 cycles. The cycle routine was 1 hour

of charge, 30 minutes rest followed by the same routine in discharge. Different current

densities ranging from 0.02 mA.cm-2 to 0.26 mA.cm-2 with 0.02 mA.cm-2 increments, are then

applied in charge and discharge for a polarization time, tp, and after each polarization the cells

are allowed to relax during a relaxation time tr . The current densities are first increased then

decreased, in order to confirm the reversibility of the results. The current densities are applied

to the cells using a MACCOR battery cycler. The voltage is recorded over time.

3. Results and discussion

a. Experimental results

The voltage variations during the galvanostatic steps are shown in Figure 25 for the

composite-Si-EO symmetric cell (turquoise), in Figure 26 for the composite-Si-EO-Si

symmetric cell (burgundy), and in Figure 27 for the neutral SEO 55-52 BCE composite

symmetric cell (navy blue). For each material, a zoom at three different current densities, i.e.

0.02, 0.04 and 0.06 mA.cm-2 is presented in figure b). In addition, in figure c) a zoom of one

or two cycles is provided, finally Figure 25 d) and Figure 26 d) the voltage obtained at a

current density of 0.02 mA.cm-2 for the 1st and the 23rd cycle is given for comparison.

-1.0

-0.5

0.0

0.5

1.0

Vo

lta

ge

(V

)

6050403020100Time (h)

Composite with PEO-b-PSTFSI BCE

a)

0.02

0.26


- 193 -

Figure 25. a)Voltage plotted over time for a composite-Si-EO BCE symmetric cell (turquoise) the values of current densities given in the graph are in mA.cm-2, b) zoom for three different current densities (0.02, 0.04 and 0.06

mA.cm-2), c) zoom on cycle 9th and 10th and d) voltage obtained at i = 0.02 mA.cm-2 for the 1st cycle and the 23rd.

The steady state voltage for the composite-lithium symmetric cell with Si-EO BCE, is

reached as expected for a single-ion almost instantly (inferior to 1 second), then a plateau in

voltage is observed during the time of polarization. When polarization is stopped, the voltage

dropped to zero abruptly. However, we observe during the discharge of the 9th cycle a strong

increase in voltage (see in Figure 25 c)). According to our previous results, this increase could

be due to the stage 3 of lithium dendritic growth shown in Figure 9, i.e. the contact between a

lithium dendrite and the ceramic. Moreover, we have previously observed that this single-ion

diblock copolymer electrolyte does not present a very high resistance to dendrite growth (see

in Chapter 4.1.c)). In the following cycles, the absolute value of the voltage is higher, as it is

clearly shown in Figure 25 d), where the 1st and the 23rd cycles (both at 0.02 mA.cm-2) are

presented. We observe that the voltage has increased by a factor two for the same current

density. (Therefore we only focused on the first nine cycles in order to avoid other interfering

phenomena.)

-0.2

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0.1

0.2

Vo

lta

ge

(V

)

76543210Time (h)

Composite with PEO-b-PSTFSI BCE

I1 = 0.02

-I1

I2 = 0.04

-I2

I3 = 0.06

-I3b)

-0.5

0.0

0.5

Voltage (

V)

252423222120Time (h)

c)

I9 = 0.18

I10 = 0.20

-0.2

-0.1

0.0

0.1

0.2

Vo

lta

ge

(V

)

2.01.51.00.50.0Time (h)

Cycle number 1 Cycle number 23

d)

I = 0.02


- 194 -

Figure 26. a) Voltage plotted over time for a composite (Si-EO-Si-ceramic) lithium symmetric cell and b) zoom for three different current densities (the values of current densities given in the graph are in mA.cm-2), c) zoom at high

current density i = 0.24 mA.cm-2 and d) voltage obtained at i = 0.02 mA.cm-2 for the 1st cycle and the 23rd.

In the case of the Si-EO-Si BCE composite-lithium symmetric cell, a similar behavior as

presented above is observed. The zooms at different current densities presented in Figure 26

b), c) and d) show that the voltage abruptly reaches the steady state voltage and remains

constant during the polarization time. In addition, no dendrite signatures or strong increase in

voltage are observed, suggesting that dendrites have not cross the BCE layer during the

galvanostatic steps and that the ceramic is still pristine. Moreover, the voltage for the same

current density is similar for both the rise and the fall of current steps: this is highlighted in

Figure 26 d), where the 1st and the 23rd cycle are represented both obtained at 0.02 mA.cm-2,

and we observe clearly the similarity in voltage.

-0.10

-0.05

0.00

0.05

0.10

Vo

lta

ge

(V

)

6050403020100Time (h)

Composite with PSTFSI-b-PEO-b-PSTFSI BCE

a)

-0.04

-0.02

0.00

0.02

0.04

Voltage (

V)

76543210Time (h)

Composite with PSTFSI-b-PEO-b-PSTFSI BCE

I1 = 0.02

-I1

I2 = 0.04

-I2

I3 = 0.06

-I3

b)-0.10

-0.05

0.00

0.05

0.10

Vo

lta

ge

(V

)

32.031.531.030.530.0Time (h)

c)

I12 = 0.24

-I12-0.03

-0.02

-0.01

0.00

0.01

0.02

0.03

Voltage (

V)

4.54.03.53.02.5Time (h)

Cycle number 1 Cycle number 23

d)

I = 0.02


- 195 -

We can notice that at high current density the plateau is slightly deformed, a zoom on the

13th cycle is given in Figure 26 b). The voltage slightly increases over the polarization time.

The voltage deviation at high current densities could be caused by electro convection effects

at the interface, due to the heterogeneous nature of the BCE.

The response of the composite-lithium symmetric cell with SEO 55-52 is shown in Figure

27. The general shape of the voltage versus time curve is very different compared to the

rectangular curves obtained for the composite with single-ion BCE. Here, the voltage

increases abruptly, then it increases slowly but never reaches the steady state during the time

of the polarization. In addition, during the relaxation time, the same behavior is observed, i.e.

the voltage drops slowly to zero. This time evolution of the voltage corresponds to the

formation of the concentration gradient during galvanostatic polarization and its relaxation

during rest period, with the interfacial concentrations that go back to their initial value.

-1.5

-1.0

-0.5

0.0

0.5

1.0

Vo

lta

ge

(V

)

1086420Time (h)

a)



- 196 -

Figure 27. a) Voltage plotted over time for a composite SEO 55-52-ceramic lithium symmetric cell and b) zoom for three different current densities (0.04, 0.06 and 0.08 mA.cm-2) and c) zoom on the cycle 5 at 0.1mA.cm-2 to

show the voltage noise of dendrites.

Interestingly, the voltage in charge and in discharge are different for the same current

density. In charge, the voltage is positive and is always higher compared to the one in

discharge, thus the dissymmetry of the cell could be due to the first direction of polarization

that can create a sort of memory effect. After the 6th cycle, we observe a voltage noise (see in

Figure 27 c)) with a drops of voltage that are characteristic of dendrite growth and soft short

circuits4. Then the current is reversed, the dendrites shorts are healed and the curves are very

smooth. Thus only the voltage in discharge is taken into account in the following section.

b. Discussion

The voltage is collected from the experimental data at the end of each polarization step to

construct a current-voltage plot (I-V plot) (filled marker).

The experimental data for the composite with Si-EO and Si-EO-Si BCEs are presented in

Figure 29 and Figure 30, respectively. The evolution of the polarization is rather linear, which

shows that there is no nonlinear behavior that would comes especially from the interfaces (see

Figure 28). The polarization resistances obtained by the linear regression of the experimental

data is 2.66 kΩ.cm2 and 1.27 kΩ.cm2 for the Si-EO and Si-EO-Si BCEs respectively.

-0.8

-0.6

-0.4

-0.2

0.0

0.2

0.4

0.6

0.8

Voltage (

V)

4.54.03.53.02.52.0Time (h)

b)


I2 = 0.04

I3 = 0.06

I4 = 0.08

-I2

-I3

-I4

1.0

0.5

0.0

-0.5

-1.0

Voltage (

V)

6.05.85.65.4Time (h)

Cycle number 5

c)

I5 = 0.10

-I5


- 197 -

Figure 28. Schematic of the evolution of the interfacial polarization versus current for different kinetic coefficients of

a simple electrochemical reaction as ox + e- → red. This has been drawn with a standard electrolyte with mass transport limitations.

We over-imposed the polarization, which is calculated by taking into account the

electrolyte resistance and interfacial resistance (Rmeasured) obtained from EIS measurement

made before the DC experiment (equation (3)). The result is represented as a solid line in

Figure 29 to Figure 31. The agreement with the experiment is very good for both single-ion

BCEs cases showing that compared to the ohmic contribution (electrolyte/SEI), the charge

transfer resistance must be very small leading to an overall fast interfacial exchange.

Utheo = Rmeasured * Iapplied (3)


- 198 -

Figure 29. I-V plot for a composite-lithium symmetric cell with PEO-b-PSTFSI BCE at 90ᵒC.

Figure 30. I-V plot for a composite-lithium symmetric cell with PSTFSI-b-PEO-b-PSTFSI BCE at 90ᵒC.

The Figure 31 presents the results of the polarization experiment on the composite with

SEO 55-52 BCE. This cell shows also a linear behavior, which is quite surprising at high

current densities given the nature of the charge carriers with a low Li+ transference number.

0.6

0.5

0.4

0.3

0.2

0.1

0.0

Vo

lta

ge

(V

)

0.150.100.050.00

Current density (mA.cm-2

)

Composite : Ohara GC-Si-EO Correlation with EIS

0.35

0.30

0.25

0.20

0.15

0.10

0.05

0.00

Vo

lta

ge

(V

)

0.200.150.100.050.00


)

Composite : Ohara GC-Si-EO-Si Correlation with EIS


- 199 -

Figure 31. I-V plot for composite with SEO 55-52 at 90ᵒC.

Mehrotra et al.7 modeled the concentration profile of a lithium symmetric cell with the

same Ohara GC we used in this study, but with 0.5 M of LiPF6 salt dissolved in EC:DEC

(1:1) as electrolyte (Figure 32). Given the similarity in transference number for this liquid

electrolyte (t+ = 0.38) and our SEO BCE (t+ = 0.152), we assume that the behavior of our

electrolyte will be close to the one presented in Figure 32. The two layers of BCE mimic each

other with the formation of a concentration gradient, where the inset, which zooms in the

region close to the Ohara GC, shows the constant concentration inside the Ohara GC as the

ionic current is carried only by migration (t+ = 1). As time progresses, the concentration

profile evolves until it reaches the steady-state (black solid line) with linear profile in the liquid

electrolyte.

1.4

1.2

1.0

0.8

0.6

0.4

0.2

0.0

Voltage (

V)

0.250.200.150.100.050.00


)

Composite : Ohara GC-SEO Correlation with EIS


- 200 -

Figure 32. Model calculations showing the time evolution of Li+ concentration profiles in a Li-Li symmetric cell with Ohara GC sandwiched between two liquid electrolyte made of 0.5M LiPF6 in EC:DEC (1:1) for I = 0.1

mA.cm-2 (Mehrotra et al.7). The inset is a zoom in the region close to the Ohara GC.

In the case of SEO BCE, in addition with the composite and the interface resistances to

calculate the polarization, we need to take into account the diffusion resistance. Two different

ways to find the diffusion resistance are used.

In the first experiment, a constant polarization at different current densities on a SEO-

lithium symmetric cell is performed. The experimental results are presented in Figure 33. The

diffusion resistance of the SEO 55-52 BCE could be calculated from this experiment. A

model prediction without taking into account the diffusion resistance is calculated and plotted

in blue solid line in Figure 33 a). The experimental curve is then subtracted from the model

and the diffusion resistance is given by the slope of the obtained curve. The diffusion

resistance obtained is 208 Ω.cm2.

We then calculated the diffusion resistance in the composite-lithium cell according to:

!Rdiff!compo!=!Rdiff!theo!!.!lSEO!compo

lSEO (7)


- 201 -

With, Rdiff compo is the diffusion resistance in the composite cell, Rdiff theo is the diffusion

resistance calculated by the DC experiment in the SEO/lithium symmetric cell and lSEO compo

the thickness of the SEO inside the composite and lSEO the thickness in the SEO cell.

A diffusion resistance for the SEO inside the composite cell is calculated to be 438 Ω.cm2. As

seen previously, the SEO BCE and the Ohara GC presents a poor adherence and the

coefficient of contact was determined as 0.33 at 90ᵒC. The diffusion resistance will therefore

be affected by the effective surface and thus Rdiff is multiply by 1/0.33. Therefore, for the

composite cell we obtain Rdiff = 4443 Ω.

Figure 33. I-V plot for a SEO-lithium symmetric cell at 90ᵒC with a) a model without take into account the diffusion.

We then calculate the theoretical voltage taking into account the three different

contributions as shown in the Figure 31 in pink solid line. It can be observed that the model

and the experiment are slightly different. The voltage obtained during the experiment is

slightly higher than the voltage expected. This difference is potentially a junction polarization

loss at the interface SEO 55-52-Ohara GC. The polarization loss is 0.87 kΩ.cm2, which is the

slope of the curve determined by subtracting the model curve to the experimental curve.

Conclusions

In the first part of Chapter 5, the EIS study of the composite-lithium symmetric cells has

been performed with three different BCEs: a neutral BCE (SEO)laden with LiTFSI salt and

100x10-3

80

60

40

20

0

Voltage (

V)

0.200.150.100.050.00


)

SEO 55-52 Umodel without diffusion


- 202 -

the two single-ion BCE (SI-EO and Si-EO-Si). With both single-ion BCE the contact with

the ceramic seems very good. The effective ionic conductivities of the composite-lithium

symmetric cells have been measured to be 1.04.10-4 S.cm-1 and 8.82.10-5 S.cm-2 for composite

using the Si-EO and the Si-EO-Si, respectively. One possible perspective is to decrease the

thickness of the polymer layers in order to decrease the whole resistance of the cell. However

due to the high content in PS for the SEO BCE the contact between the SEO and the

ceramic is quite poor. In this case, we have estimated that only 1/3 of the surface of the

interface is intimate at 90ᵒC. In the case of the single-ion BCEs, we did not observe an

additional contribution in the EIS spectra due to the ionic charge transfer at the interface

BCE-Ohara GC. Thus, this phenomenon is quite fast and not limiting in our case.

Galvanostatic cycling experiments, at 0.175 mA.cm-2 during 4 hours, have been then

performed on the composite-lithium symmetric cells. The EIS analysis have allowed to follow

both the electrolyte and the interface resistances during cycling. The composite cell using the

SEO BCE was stopped before the short circuit (due to time constrains) after around 500

C.cm-2 and it was not possible to image it. On the other side, the composite cells using the

single-ion BCEs have given for the Si-EO a short-circuit with fuse effect associated with an

increase of the average voltage after 13 cycles (100 C.cm-2). On the contrary, for the Si-EO-Si

BCE did not experienced a short circuit after 32 cycles (170 C.cm-2) and the voltage was

constant during cycling.

Composite cells using both Si-EO and Si-EO-Si BCEs have been imaged using hard X-

ray micro-tomography. In the first case, after cycling, in half of the cell the Ohara GC is

broken implying that the dense lithium dendrites have grown through the Si-EO layer, which

explains the voltage increase observed. In addition, the polymer-lithium interface is irregular,

which is in good agreement with our previous results in polymer-lithium symmetric cells. For

the composite cell using the Si-EO-Si BCE, after 170 C.cm-2 the Ohara GC is still pristine,

meanings that no lithium dendrites have crossed the polymer layer. However, the Si-EO-Si-

lithium interface presents irregularities implying again a dense but irregular electro-deposition

of lithium metal.

In the second part of the Chapter 5, an experiment using increasing galvanostatic steps

has been performed onto the composite-lithium symmetric cells. The first important

observation is that for the single-ion BCEs the steady state voltage is instantly reached when

the polarization starts, and drops to zero when the polarization is stopped, which confirms

their single-ion nature. Whereas for the SEO BCE, a classical behavior with an increasing


- 203 -

voltage over polarization and the steady state is never reached in our conditions. A slow

relaxation voltage is then clearly visible before reaching 0V, due to the re-equilibration of

concentration along the cell during the rest time.

The voltage obtained at the end of the polarization step is plotted in an I-V plot and a

linear behavior is observed for all three composite-lithium symmetric cells. In the case of the

composite cells using single-ion BCEs, the voltage calculated from the EIS measurement

before the DC experiment is compared to the experimental one and is in good agreement

with the experimental voltage. Therefore, the charge transfer at the BCE-Ohara GC should be

fast. For the composite cell using SEO 55-52, a polarization loss at the BCE-Ohara GC

interface has been calculated to be 0.87 kΩ.cm2.


- 204 -

References of chapter 5

1. EC-Lab® software. Bio-Logic - Science Instruments Available at: http://www.bio-logic.info/potentiostat-

electrochemistry-ec-lab/software/ec-lab-software/. (Accessed: 29th April 2016)

2. Devaux, D., Bouchet, R., Glé, D. & Denoyel, R. Mechanism of ion transport in PEO/LiTFSI

complexes: Effect of temperature, molecular weight and end groups. Solid State Ion. 227, 119–127

(2012).


Mater. 27, 4682–4692 (2015).

4. Rosso, M. et al. Dendrite short-circuit and fuse effect on Li/polymer/Li cells. Electrochimica Acta 51,

5334–5340 (2006).

5. Abe, T., Sagane, F., Ohtsuka, M., Iriyama, Y. & Ogumi, Z. Lithium-Ion Transfer at the Interface

Between Lithium-Ion Conductive Ceramic Electrolyte and Liquid Electrolyte-A Key to Enhancing the

Rate Capability of Lithium-Ion Batteries. J. Electrochem. Soc. 152, A2151–A2154 (2005).

6. Sagane, F., Abe, T. & Ogumi, Z. Li+-Ion Transfer through the Interface between Li+-Ion Conductive

Ceramic Electrolyte and Li+-Ion-Concentrated Propylene Carbonate Solution. J. Phys. Chem. C 113,

20135–20138 (2009).

7. Mehrotra, A., Ross, P. N. & Srinivasan, V. Quantifying Polarization Losses in an Organic Liquid

Electrolyte/Single Ion Conductor Interface. J. Electrochem. Soc. 161, A1681–A1690 (2014).

8. Abe, T., Ohtsuka, M., Sagane, F., Iriyama, Y. & Ogumi, Z. Lithium Ion Transfer at the Interface

between Lithium-Ion-Conductive Solid Crystalline Electrolyte and Polymer Electrolyte. J. Electrochem.

Soc. 151, A1950–A1953 (2004).

9. Tenhaeff, W. E., Yu, X., Hong, K., Perry, K. A. & Dudney, N. J. Ionic Transport Across Interfaces of

Solid Glass and Polymer Electrolytes for Lithium Ion Batteries. J. Electrochem. Soc. 158, A1143–A1149

(2011).

10. Tenhaeff, W. E., Perry, K. A. & Dudney, N. J. Impedance Characterization of Li Ion Transport at the

Interface between Laminated Ceramic and Polymeric Electrolytes. J. Electrochem. Soc. 159, A2118–A2123

(2012).


- 205 -

11. Yamada, Y., Abe, T. & Ogumi, Z. Lithium-ion Kinetics at Interface between Lithium-ion Conductive

Electrolyte/DMC-based Electrolyte Interfaces. ECS Trans. 16, 135–139 (2009).

12. Gondran, C., Albert, F. & Siebert, E. Kinetics of sodium and silver exchange on a

PEOx 0.25 based internal reference system. Solid State Ion. 84, 131–138 (1996).

13. Macdonald, J. R. Binary electrolyte small-signal frequency response. Electroanal. Chem. Int. Electrochem. 53,

1–55 (1974).


- 206 -

- 207 -

Conclusions and perspectives

This PhD has been focused on the replacement of the hard inorganic LiPON layer by a

solid block copolymer electrolyte, as a protection of the Ohara GC to prevent any contact

with the lithium metal negative electrode in the aqueous lithium-air batteries. Indeed, despite

its stability versus lithium metal, its good mechanical properties and the advantage of thin

layer deposition, the LiPON presents a structural rigidity, which leads to the loss of the active

surface area during cycling. Moreover, its restrictive deposition technique makes it a very

expensive material. Thus, we have proposed a protective buffer layer based on neutral block

copolymer electrolyte or single-ion block copolymer electrolyte, which exhibits good Li+ ionic

conductivity, flexibility and high resistance to lithium dendritic growth.

After introducing the general energy context, a review of the main battery technologies

has been presented (Chapter 1). The advantage of the aqueous lithium-air (Li-air) technology

has been discussed along with their limitations. A review of the different possible materials

for the protection of the ceramic has been given. Finally, the issues related to the use of the

lithium metal has been discussed, as well as the different approaches to prevent the dendritic

growth.

The first step of this PhD (Chapter 2) was to study, by electrochemical impedance

spectroscopy, the Ohara glass-ceramic, as well as the initial protective layer, i.e. the LiPON

deposited on the Ohara GC. A simple difference of spectra at 50ᵒC have allowed to show


- 208 -

only one contribution due to the LiPON layer. Thus we have demonstrated the lithium ionic

charge transfer at the interface LiPON-Ohara GC is very small and not measurable (second

order phenomenon). A simple equivalent circuit has allowed to isolate the LiPON

contribution in the sandwich LiPON-Ohara GC. We were able to discriminate with a very

good agreement the different electrical properties (conductivities and activation energy) of the

Ohara GC and the LiPON layer. However, the loss of the LiPON-lithium interface during

cycling remains the main issue, therefore the replacement of the LiPON with a flexible

material is necessary.

In this context, three block copolymer electrolytes (BCE) have been investigated: a neutral

PS-b-PEO laden with LiTFSI salt (SEO) at EO/Li = 11.7, and two single-ion BCEs, one

diblock PEO-b-PSTFSILi and one triblock PSTFSILi-b-PEO-b-PSTFSILi. A morphology

study by small angle X-ray spectroscopy (SAXS) and dark field scanning transmission electron

microscopy (STEM) has been performed for the two single-ion BCEs. A lamellar morphology

for both BCEs has been observed, and interestingly their domain spacing are very close to

previous studies1,2 performed with similar BCEs, but with much smaller molecular weight. In

addition, SAXS experiments have been performed as a function of temperature and we

observed a decrease in the scattering peak values with increasing temperatures, meanings that

the domain spacing increases. However, this increase is very large to be related to a dilatation

phenomenon. Besides, the high ionic conductivity has been correlated to the disordered phase

above the melting temperature of the PEO domain, which implies a partial miscibility of the

PSTFSI and PEO blocks. Finally, the Li+ transference number for these two materials has

been confirmed to be unity (or very near) using electrochemical methods.

It is important to remember that the electrolyte has to fulfill requirements in order to

protect the Ohara GC. Beyond the stability versus lithium metal, one of the main requirement

is a high resistance to lithium dendritic growth along cycling. Thus, the galvanostatic cycling

behavior of these BCEs has been studied in lithium symmetric cells (Chapter 4). An original

study by EIS measurements during the cycling has been reported for the three BCEs. Two

mains parameters were followed, the electrolyte and the interface resistances. The electrolyte

resistance stays constant over cycling, whereas the interface resistance decreases in most of

the cases. Excellent results have been reported for the first time and especially for the two

single-ion BCEs, which exhibits very poor mechanical properties, but very good cycling with

more than 95 cycles for the Si-EO-Si BCE, that finding definitely shows the importance of

transference number in the mitigation of the dendritic growth.


- 209 -

Hard X-ray micro-tomography has been performed on the cells at their initial stage and

after cycling. The intimacy of the interface BCE-lithium for the three BCEs has been

observed before and after cycling: no interface losses have been observed implying a good

flexibility. In addition, the morphology of cycled lithium electrode has been analyzed: for the

SEO BCE porous globular dendrites very recently reported have been also observed, whereas

for the Si-EO and Si-EO-Si BCEs dense lithium objects with different sizes and shapes have

been observed for the first time. We supposed that this dense heterogeneous electro-

deposition is due to local fluctuations of the current densities probably caused by

heterogeneous passivation layers. This result is encouraging due to the fact that an electrolyte

can present poor mechanical properties but high resistance to dendrite, which shows the

preponderance of blocking the nucleation by appropriate transport properties3 instead of

blocking the growth by mechanical barriers4,5.

After the characterization of the three potential protective materials for the Ohara GC,

composite-lithium symmetric cells using the BCEs have been assembled and studied (Chapter

5.I). The low effective ionic conductivity of the composite using SEO BCE permitted to

conclude to a poor adherence on the glass-ceramic, whereas choosing both single-ion BCEs

gives the expected values confirming the nice interfaces. The galvanostatic cycling of the

composite-lithium symmetric cells with an EIS study over cycling has been performed.

Imaging the composite cells at its initial state allowed us to observe the intimacy of the BCEs

and the Ohara GC in addition with the BCE and the lithium metal. The composite cells using

single-ion BCEs have been cycled for a small amount of cycle before to be imaged using hard

X-ray micro-tomography (due to the constraints of the beam time at the synchrotron). We

confirmed our previous results, i.e. that Si-EO-Si BCE exhibits a high resistance to dendritic

growth, the lithium could not cross the whole polymer layer and the Ohara GC is pristine. On

the contrary, the Si-EO BCE presents a smaller resistance to dendrites growth and the lithium

has reduced the glass-ceramic, thus fractures were observed. The morphology of the interface

polymer-lithium in both cases is irregular but dense in a similar manner than in the single-

ion/lithium symmetric cells analyzed in Chapter 4.

Finally, an experiment using increasing galvanostatic steps has been performed onto the

composite cells (Chapter 5.II). A first observation is that with single-ion BCEs the steady state

voltage is always instantly reached, and the voltage drop to zero when the polarization is

topped. A simple correlation between the EIS measurement and the voltage induces by the

different current steps is given and a good agreement has been reported. Those experiments


- 210 -

in steps confirmed that no additional contribution due to the interface single-ion BCE and

Ohara GC is observed. This result implies that the contribution of the ionic charge transfer at

the interface BCE-Ohara GC is at least small compared to the others phenomena (electrolyte

resistance, lithium/electrolyte interface resistance).

As a perspective: further galvanostatic steps would have been interesting with an

operando following of the EIS. The cycling of the single-ion at different current densities to

analyze the law and the morphology of the dendrites. Moreover, the study of BCE-lithium

symmetric cells and composite cells using thinner films (few microns) would have been very

interesting to decrease the working temperature. Finally, the cycling test in a composite

lithium-air battery at 65ᵒC (above the PEO melting temperature) would have been very

interesting.

References conclusions



2. Rojas, A. A. et al. Effect of Lithium-Ion Concentration on Morphology and Ion Transport in Single-

Ion-Conducting Block Copolymer Electrolytes. Macromolecules 48, 6589–6595 (2015).

3. Chazalviel, J.-N. Electrochemical aspects of the generation of ramified metallic electrodeposits. Phys.

Rev. A 42, 7355–7367 (1990).





- 211 -

Résumé en français

La technologie Lithium-air (Li-air) développée par EDF utilise une électrode à air qui

fonctionne avec un électrolyte aqueux ce qui empêche l’utilisation de lithium métal non

protégé. Une membrane céramique conductrice d’ion Li+ est utilisée pour séparer le milieu

aqueux de l’électrode négative en lithium métal. Cependant, cette céramique n'est pas stable

au contact du lithium, il est donc nécessaire d'ajouter une couche de protection. Par ailleurs,

dans l'idéal cette protection doit résister à la croissance dendritique du lithium. C'est dans ce

contexte que s'inscrit ce projet qui a pour but le développement d'une couche de protection;

à base d'électrolyte copolymère à blocs (BCE), entre le lithium métal et une céramique

conductrice d'ions lithium, pour les batteries Lithium air.

A ce jour, dans la technologie Li-air d'EDF, une couche de lithium phosphorous oxy

nitride (LiPON) est employée entre le lithium métal et la céramique. Dans une première

partie, l'étude de la céramique ainsi que de l'assemblage LiPON-céramique en température

sera réalisée. La caractérisation de ces matériaux se fera par spectroscopie d'impédance.

Apres avoir identifié les deux contributions de la céramique, soit le grain et les joints de grain,

l'assemblage LiPON-céramique sera étudié. Le but de cette étude est de pouvoir distinguer la

contribution du LiPON de celle de la céramique. Pour cela, un simple circuit équivalent est

utilisé.

Cependant, ce matériau inorganique est dure et pendant la charge de la batterie, la

croissance du lithium sur cette surface dure peut engendrer des contraintes mécaniques, qui

ont pour conséquence la perte de l'interface LiPON-lithium et donc une réduction de la

surface active. C'est pour cela qu'il est nécessaire de remplacer ce matériau par un matériau

qui est suffisamment mou pour absorber les variations dimensionnels du lithium lors de sa

croissance, et par ailleurs suffisamment dur pour résister à la croissance dendritique du

lithium. C'est dans ce contexte que les BCEs ont été choisi comme matériau pour jouer le

rôle de couche protectrice de la céramique.

Le comportement des ces BCE est tout d'abord étudié en cellule symétrique lithium-

lithium afin de déterminer leurs propriétés telles que leurs conductivités et nombre de

transport, ainsi que leur résistance face à la croissance dendritique du lithium. Plusieurs

techniques de caractérisation sont utilisées avec notamment du "small angle X-ray Scattering"

(SAXS), afin d'étudier la morphologie de ces BCEs et en particulier leur nano-séparation de

- 212 -

phase, cette technique sera couplée à de la microscopie (STEM) et les résultats obtenus par

les deux méthodes sont comparées. Le suivi par spectroscopie d'impédance pendant le

cyclage de ces cellules, permet de suivre l'évolution de la résistance d'électrolyte d'une part et

la résistance d'interface polymère-lithium d'autre part. Avant et après cyclage, des analyses

par micro-tomographie des rayons X sont réalisées pour analyser la morphologie du lithium.

Pour des électrolytes possédant un nombre de transport de l'unité, la croissance dendritique

du lithium est supposée être supprimée, cependant pour la première fois, la visualisation

d'une croissance non homogène et la formation d'objets denses de lithium après cyclage à

travers des électrolytes polymères poly anioniques de type "single-ion" est reportée.

Dans une dernière partie, le composite céramique-polymère est caractérisé par

spectroscopie d'impédance. Tout d'abord, les conductivités ioniques des composites BCE-

céramique sont étudiées. Puis le cyclage de ces composites en cellule symétrique lithium-

lithium et l'analyse des spectres d'impédance après chaque cycle permet ainsi de déterminer si

les dendrites de lithium ont atteint la céramique. Par ailleurs, la quantification de la perte de

polarisation à l'interface céramique-lithium est évaluée par des expériences en polarisation, la

contribution inter facial entre le polymère et la céramique est faible.

- 213 -

The lithium-air (Li-air) technology developed by EDF uses an air electrode which works with an

aqueous electrolyte, which prevent the use of unprotected lithium metal electrode. A Li+ ionic conductor

glass ceramic is used to separate the aqueous electrolyte compartment from the negative lithium electrode.

However, this glass-ceramic is not stable in contact with lithium, it is thus necessary to add a protective

buffer layer. In another hand, this protection should ideally resist to lithium dendritic growth. It is in this

context that this research project which has as goal the development of a protective buffer layer based on

block copolymer electrolytes (BCE) between the lithium metal and the lithium ionic conductor ceramic,

for lithium-air battery.

In a first part, the BCE is studied in lithium-lithium symmetric cells, in order to determine their

properties such as their ionic conductivities, their transference number, and their resistance to dendritic

growth. Several characterization techniques are employed and especially the hard X-ray micro-tomography

to analyze the lithium morphology before and after cycling. For single-ion BCE, we expect to suppress

dendritic growth, however, we report here for the first time, the visualization of a homogeneous growth of

lithium but the formation of dense lithium objects.

In another part, the composite BCE-ceramic is studied by electrochemical impedance

spectroscopy. The cycling of composite-lithium symmetric cells and the analysis of the EIS measurements

after each cycle permit to determine if the dendrites have cross the electrolyte and are in contact with the

ceramic. Besides, the quantification of the polarization loss at the interface polymer-ceramic is evaluated by

polarization experiments. This contributions is found to be small.

Développement d’une couche d’interface entre Lithium métal et électrolyte céramique à base de block-copolymère pour batterie Lithium air aqueuse

rechargeable

La technologie Lithium-air développée par EDF utilise une électrode à air qui fonctionne avec un

électrolyte aqueux ce qui empêche l’utilisation de lithium métal non protégé. Une membrane céramique

conductrice d’ion Li+ est utilisée pour séparer le milieu aqueux de l’électrode négative en lithium métal.

Cependant, cette céramique n'est pas stable au contact du lithium, il est donc nécessaire d'ajouter une

couche de protection. Par ailleurs, dans l'idéal cette protection doit résister à la croissance dendritique du

lithium. C'est dans ce contexte que s'inscrit ce projet qui a pour but le développement d'une couche de

protection; à base d'électrolyte copolymère à blocs (BCE), entre le lithium métal et une céramique

conductrice d'ions lithium, pour les batteries Lithium air.

Tout d'abord le comportement des ces BCE est étudié en cellule symétrique lithium-lithium afin

de déterminer leurs propriétés telles que leurs conductivités et nombre de transport, ainsi que leur

résistance face à la croissance dendritique du lithium. Plusieurs techniques de caractérisation sont utilisées

et notamment la microtomographie par rayons X pour analyser la morphologie du lithium après cyclage.

Pour des électrolytes possédant un nombre de transport de l'unité, la croissance dendritique du lithium est

supposée être supprimée, cependant pour la première fois, la visualisation d'une croissance non homogène

et de la formation d'objets denses de lithium après cyclage à travers des électrolytes polymères poly

anioniques est reportée.

Dans un second temps, le composite céramique-polymère est caractérisé par spectroscopie

d'impédance. Le cyclage du composite en cellule symétrique lithium-lithium et l'analyse des spectres

d'impédance après chaque cycle permet de déterminer si les dendrites de lithium ont atteint la céramique.

Par ailleurs, la quantification de la perte de polarisation à l'interface céramique-lithium est évaluée par des

expériences en polarisation, la contribution inter facial entre le polymère et la céramique est faible.

Keywords: Block copolymer electrolyte/Impedance spectroscopy/Dendritic growth/Single-ion

electrolyte

study of a buffer layer based on block copolymer

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