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
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
References of Chapter 2. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 91
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
References of Chapter 3. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 124
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
2. Results and discussion. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 167
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
3. Results and discussion. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 192
a. Experimental results. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 192
b. Discussion. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 196
Conclusions. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 201
References of Chapter 5. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 204
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
- 2 -
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
- 4 -
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
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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%
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2000 2005 2012 2020 2025
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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
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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,
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
- 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 :
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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-
Chapter 1. General context and battery state of the art
- 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)
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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)
Chapter 1. General context and battery state of the art
- 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,
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
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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
Chapter 1. General context and battery state of the art
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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.
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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).
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
- 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.
! "# $%&°'( " )
*+,,,,,, (12)
As soon as a current above the diffusion limited current, J*, is imposed, the vicinity of the
cathode (x=0) is depleted to form a metallic deposit and anions migrate to the anode. After a
time close to the Sand time (τs) 165, i.e. when ionic concentrations are close to zero at the cathode
and J > J*, Chazalviel showed that an area of few microns devoid of anions with an excess of
cations (zc.Cc >> za.Ca), is created. This area is called the space-charge layer and is positively
charged. Thus, two regions can be distinguished in the electrolyte, a neutral region (zone I) and a
space charge region (zone II) in the vicinity of the negative electrode. This space charge induces
a very large electric field (up to 10 kV.cm-1) in the vicinity of the cathode and therefore an abrupt
Chapter 1. General context and battery state of the art
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drop in potential localized in this region (see Figure 31) that is the main driving force for the
nucleation of dendrites.
Therefore he postulated that dendrite starts to grow when the electric field reach a critical
value, typically reached at the Sand time (see equation (13), with µa and µc anion and cation
mobility respectively, D the ambipolar diffusion coefficient, J the current density).
-s = .' " /01"%"&°1$! 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*.
Chapter 1. General context and battery state of the art
- 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)).
Chapter 1. General context and battery state of the art
- 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)
Chapter 1. General context and battery state of the art
- 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.
Chapter 1. General context and battery state of the art
- 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)
Chapter 1. General context and battery state of the art
- 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-
Chapter 1. General context and battery state of the art
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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
Chapter 1. General context and battery state of the art
- 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
Chapter 1. General context and battery state of the art
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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.
Chapter 1. General context and battery state of the art
- 59 -
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
Chapter 1. General context and battery state of the art
- 60 -
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.
Chapter 1. General context and battery state of the art
- 61 -
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.
Chapter 1. General context and battery state of the art
- 62 -
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.
Chapter 1. General context and battery state of the art
- 63 -
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.
Chapter 1. General context and battery state of the art
- 64 -
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.
Chapter 1. General context and battery state of the art
- 65 -
References of Chapter 1
1. Tarascon, J. M. in Encyclopedia of Electrochemical Power Sources (ed. Garche, J.) xxvii–xxviii (Elsevier, 2009).
2. 21st Conference of the Parties http://www.cop21paris.org/.
3. International energy agency (IEA) http://www.iea.org/.
4. Yang, Z. et al. Electrochemical Energy Storage for Green Grid. Chem. Rev. 111, 3577–3613 (2011).
5. International Renewable Energy Agency http://www.irena.org/.
6. Batteries for transportation now and in the future. (2011).
7. Thomas B. Reddy, D. L. Handbook of batteries. (McGraw-Hill).
8. Kurzweil, P. in Encyclopedia of Electrochemical Power Sources (ed. Garche, J.) 565–578 (Elsevier, 2009).
9. Gaston Plante. Storage of electriccal energy. (Paul Bedford, 1859).
10. Ruetschi, P. Review on the lead—acid battery science and technology. J. Power Sources 2, 3–120 (1977).
11. Rand, D. A. J., Pavlov, D. & Garche, J. Proceedings of the International Conference on Lead/Acid
Batteries: LABAT ’89On the historical development of the lead/acid battery, especially in Europe. J. Power
Sources 31, 401–406 (1990).
12. Jungner. (1899).
13. Neumann Georg. Gastight storage battery and method of manufacturing same. (1949).
14. Notten, P. H. L. & Latroche, M. in Encyclopedia of Electrochemical Power Sources (ed. Garche, J.) 502–521
(Elsevier, 2009).
15. Hariprakash, B., Shukla, A. K. & Venugoplan, S. in Encyclopedia of Electrochemical Power Sources (ed. Garche, J.)
494–501 (Elsevier, 2009).
16. Goodenough, J. B. in Encyclopedia of Electrochemical Power Sources (ed. Garche, J.) 243–248 (Elsevier, 2009).
17. Nishi, Y. Lithium ion secondary batteries; past 10 years and the future. J. Power Sources 100, 101–106 (2001).
18. Scrosati, B. & Garche, J. Lithium batteries: Status, prospects and future. J. Power Sources 195, 2419–2430
(2010).
19. Tarascon, J.-M., Gozdz, A. S., Schmutz, C., Shokoohi, F. & Warren, P. C. Performance of Bellcore’s plastic
rechargeable Li-ion batteries. Solid State Ion. 86–88, Part 1, 49–54 (1996).
20. www.teslamotors.com/gigafactory
21. 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).
22. J.J. Halek. Brevet Francais.
23. D. Herbert and J. Ulam. Brevet Francais.
24. Watanabe, N. & Fukuda, M. Primary cell for electric batteries. (1970).
25. Selim, R. & Bro, P. Some Observations on Rechargeable Lithium Electrodes in a Propylene Carbonate
Electrolyte. J. Electrochem. Soc. 121, 1457–1459 (1974).
26. Rauh, R. D. & Brummer, S. B. The effect of additives on lithium cycling in propylene carbonate.
Electrochimica Acta 22, 75–83 (1977).
27. Yoshimatsu, I., Hirai, T. & Yamaki, J. Lithium Electrode Morphology during Cycling in Lithium Cells. J.
Electrochem. Soc. 135, 2422–2427 (1988).
28. Goldman, J. L., Mank, R. M., Young, J. H. & Koch, V. R. Structure9Reactivity Relationships of Methylated
Tetrahydrofurans with Lithium. J. Electrochem. Soc. 127, 1461–1467 (1980).
29. Yamaki, J. in Encyclopedia of Electrochemical Power Sources (ed. Garche, J.) 183–191 (Elsevier, 2009).
30. M. Armand, J. Chabagno, M. and M. Duclo. in 2nd int. Meeting on Solid Electrolytes, St. Andrews
Scotland (1978).
31. Bluecar | Blue Solutions.
32. Bluebus | Blue Solutions.
33. Dunn, B., Kamath, H. & Tarascon, J.-M. Electrical Energy Storage for the Grid: A Battery of Choices.
Science 334, 928–935 (2011).
Chapter 1. General context and battery state of the art
- 66 -
34. Ji, X., Lee, K. T. & Nazar, L. F. A highly ordered nanostructured carbon–sulphur cathode for lithium–
sulphur batteries. Nat. Mater. 8, 500–506 (2009).
35. Yin, Y.-X., Xin, S., Guo, Y.-G. & Wan, L.-J. Lithium–Sulfur Batteries: Electrochemistry, Materials, and
Prospects. Angew. Chem. Int. Ed. 52, 13186–13200 (2013).
36. Ji, X. & Nazar, L. F. Advances in Li–S batteries. J. Mater. Chem. 20, 9821–9826 (2010).
37. Manthiram, A., Fu, Y. & Su, Y.-S. Challenges and Prospects of Lithium–Sulfur Batteries. Acc. Chem. Res.
46, 1125–1134 (2013).
38. Bresser, D., Passerini, S. & Scrosati, B. Recent progress and remaining challenges in sulfur-based lithium
secondary batteries – a review. Chem. Commun. 49, 10545–10562 (2013).
39. Arai, H. & Hayashi, M. in Encyclopedia of Electrochemical Power Sources (ed. Garche, J.) 347–355 (Elsevier,
2009).
40. Kim, H. et al. Metallic anodes for next generation secondary batteries. Chem. Soc. Rev. 42, 9011–9034 (2013).
41. Galbraith, A. D. The lithium-water-air battery for automotive propulsion. in (1976).
42. Abraham, K. M. & Jiang, Z. A Polymer Electrolyte9Based Rechargeable Lithium/Oxygen Battery. J.
Electrochem. Soc. 143, 1–5 (1996).
43. Ogasawara, T., Débart, A., Holzapfel, M., Novák, P. & Bruce, P. G. Rechargeable Li2O2 Electrode for
Lithium Batteries. J. Am. Chem. Soc. 128, 1390–1393 (2006).
44. Peng, Z. et al. Oxygen Reactions in a Non-Aqueous Li+ Electrolyte. Angew. Chem. Int. Ed. 50, 6351–6355
(2011).
45. Laoire, C. O., Mukerjee, S., Abraham, K. M., Plichta, E. J. & Hendrickson, M. A. Elucidating the
Mechanism of Oxygen Reduction for Lithium-Air Battery Applications. J. Phys. Chem. C 113, 20127–20134
(2009).
46. Balaish, M., Kraytsberg, A. & Ein-Eli, Y. A critical review on lithium–air battery electrolytes. Phys. Chem.
Chem. Phys. 16, 2801–2822 (2014).
47. Frimer, A. A. & Rosenthal, I. Chemical Reactions of Superoxide Anion Radical in Aprotic Solvents.
Photochem. Photobiol. 28, 711–717 (1978).
48. Mizuno, F., Nakanishi, S., Kotani, Y., Yokoishi, S. & Iba, H. Rechargeable Li-Air Batteries with Carbonate-
Based Liquid Electrolytes. Electrochemistry 78, 403–405 (2010).
49. Bryantsev, V. S. et al. Predicting solvent stability in aprotic electrolyte Li-air batteries: nucleophilic
substitution by the superoxide anion radical (O2(•-)). J. Phys. Chem. A 115, 12399–12409 (2011).
50. Bryantsev, V. S. et al. Predicting Solvent Stability in Aprotic Electrolyte Li–Air Batteries: Nucleophilic
Substitution by the Superoxide Anion Radical (O2•–). J. Phys. Chem. A 115, 12399–12409 (2011).
51. Freunberger, S. A. et al. The Lithium–Oxygen Battery with Ether-Based Electrolytes. Angew. Chem. Int. Ed.
50, 8609–8613 (2011).
52. Hassoun, J., Croce, F., Armand, M. & Scrosati, B. Investigation of the O2 Electrochemistry in a Polymer
Electrolyte Solid-State Cell. Angew. Chem. Int. Ed. 50, 2999–3002 (2011).
53. Armand, M. B. Polymer Electrolytes. Annu. Rev. Mater. Sci. 16, 245–261 (1986).
54. Bryantsev, V. S. & Faglioni, F. Predicting Autoxidation Stability of Ether- and Amide-Based Electrolyte
Solvents for Li–Air Batteries. J. Phys. Chem. A 116, 7128–7138 (2012).
55. Lu, J., Lau, K. C., Sun, Y.-K., Curtiss, L. A. & Amine, K. Review—Understanding and Mitigating Some of
the Key Factors that Limit Non-Aqueous Lithium-Air Battery Performance. J. Electrochem. Soc. 162, A2439–
A2446 (2015).
56. Luntz, A. C. & McCloskey, B. D. Nonaqueous Li–Air Batteries: A Status Report. Chem. Rev. 114, 11721–
11750 (2014).
57. Zhang, J., Xu, W. & Liu, W. Oxygen-selective immobilized liquid membranes for operation of lithium-air
batteries in ambient air. J. Power Sources 195, 7438–7444 (2010).
58. Zhang, J., Xu, W., Li, X. & Liu, W. Air Dehydration Membranes for Nonaqueous Lithium–Air Batteries. J.
Electrochem. Soc. 157, A940–A946 (2010).
59. Liu, T. et al. Cycling Li-O2 batteries via LiOH formation and decomposition. Science 350, 530–533 (2015).
Chapter 1. General context and battery state of the art
- 67 -
60. Visco, S. J. & Nimon, Y. S. Active metal/aqueous electrochemical cells and systems. (2010).
61. E. Nimon, B. Katz, L.C.D. Jonghe and M.Y. Chu, S. J. V. 12th Int Meeting on Lithium Batteries. (2004).
62. Stevens, P. et al. Development of a Lithium Air Rechargeable Battery. ECS Trans. 28, 1–12 (2010).
63. The Lithium Air Battery. (Springer New York, 2014).
64. Wiley: Electrochemical Oxygen Technology - Kim Kinoshita. Available at:
http://eu.wiley.com/WileyCDA/WileyTitle/productCd-0471570435.html. (Accessed: 17th January 2016)
65. F. Moureaux. Etude des reactions mettant en jeu l’oxygene dans un syteme electrochimique litihum air
auqeux rechargeable electriquement. (Universite de Grenoble, 2011).
66. Toussaint, G., Stevens, P., Akrour, L., Rouget, R. & Fourgeot, F. Development of a Rechargeable Zinc-Air
Battery. ECS Trans. 28, 25–34 (2010).
67. Moureaux, F., Stevens, P., Toussaint, G. & Chatenet, M. Development of an oxygen-evolution electrode
from 316L stainless steel: Application to the oxygen evolution reaction in aqueous lithium–air batteries. J.
Power Sources 229, 123–132 (2013).
68. Stevens, P., Ghamouss, F., Fichet, O. & Sarrazin, C. Interpenetrating network of anion-exchange polymers,
production method thereof and use of same. (2015).
69. Stevens, P., Toussaint, G., Vinatier, P. & Puëch, L. Very High Specific Surface Area Capacity Lithium-Air
Battery. Meet. Abstr. MA2012-02, 1100–1100 (2012).
70. Knauth, P. Inorganic solid Li ion conductors: An overview. Solid State Ion. 180, 911–916 (2009).
71. Inaguma, Y. & Nakashima, M. A rechargeable lithium–air battery using a lithium ion-conducting
lanthanum lithium titanate ceramics as an electrolyte separator. J. Power Sources 228, 250–255 (2013).
72. Thangadurai, V. & Weppner, W. Li6ALa2Nb2O12 (A=Ca, Sr, Ba): A New Class of Fast Lithium Ion
Conductors with Garnet-Like Structure. J. Am. Ceram. Soc. 88, 411–418 (2005).
73. Murugan, R., Thangadurai, V. & Weppner, W. Fast Lithium Ion Conduction in Garnet-Type
Li7La3Zr2O12. Angew. Chem. Int. Ed. 46, 7778–7781 (2007).
74. Shimonishi, Y. et al. Synthesis of garnet-type Li7 − xLa3Zr2O12 − 1/2x and its stability in aqueous
solutions. Solid State Ion. 183, 48–53 (2011).
75. Hong, H. .-P. Crystal structure and ionic conductivity of Li14Zn(GeO4)4 and other new Li+ superionic
conductors. Mater. Res. Bull. 13, 117–124 (1978).
76. Bruce, P. G. & West, A. R. The A9C Conductivity of Polycrystalline LISICON, Li2 + 2x Zn1 − x GeO4,
and a Model for Intergranular Constriction Resistances. J. Electrochem. Soc. 130, 662–669 (1983).
77. Kanno, R., Hata, T., Kawamoto, Y. & Irie, M. Synthesis of a new lithium ionic conductor, thio-LISICON–
lithium germanium sulfide system. Solid State Ion. 130, 97–104 (2000).
78. Kamaya, N. et al. A lithium superionic conductor. Nat. Mater. 10, 682–686 (2011).
79. Chen, H. M., Maohua, C. & Adams, S. Stability and ionic mobility in argyrodite-related lithium-ion solid
electrolytes. Phys Chem Chem Phys 17, 16494–16506 (2015).
80. Masquelier, C. & Croguennec, L. Polyanionic (Phosphates, Silicates, Sulfates) Frameworks as Electrode
Materials for Rechargeable Li (or Na) Batteries. Chem. Rev. 113, 6552–6591 (2013).
81. Aono, H., Sugimoto, E., Sadaoka, Y., Imanaka, N. & Adachi, G. Ionic Conductivity of Solid Electrolytes
Based on Lithium Titanium Phosphate. J. Electrochem. Soc. 137, 1023–1027 (1990).
82. Aatiq, A., Ménétrier, M., Croguennec, L., Suard, E. & Delmas, C. On the structure of Li3Ti2(PO4)3. J.
Mater. Chem. 12, 2971–2978 (2002).
83. Xu, X., Wen, Z., Yang, X. & Chen, L. Dense nanostructured solid electrolyte with high Li-ion conductivity
by spark plasma sintering technique. Mater. Res. Bull. 43, 2334–2341 (2008).
84. Thokchom, J. S. & Kumar, B. The effects of crystallization parameters on the ionic conductivity of a
lithium aluminum germanium phosphate glass–ceramic. J. Power Sources 195, 2870–2876 (2010).
85. www.oharacorp.com.
86. 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)
Chapter 1. General context and battery state of the art
- 68 -
87. Hasegawa, S. et al. Study on lithium/air secondary batteries—Stability of NASICON-type lithium ion
conducting glass–ceramics with water. J. Power Sources 189, 371–377 (2009).
88. Jackman, S. D. & Cutler, R. A. Stability of NaSICON-type Li1.3Al0.3Ti1.7P3O12 in aqueous solutions. J.
Power Sources 230, 251–260 (2013).
89. Ding, F. et al. H+ diffusion and electrochemical stability of Li1+x+yAlxTi2−xSiyP3−yO12 glass in
aqueous Li/air battery electrolytes. J. Power Sources 214, 292–297 (2012).
90. Geiger, C. A. et al. Crystal Chemistry and Stability of ‘Li7La3Zr2O12’ Garnet: A Fast Lithium-Ion
Conductor. Inorg. Chem. 50, 1089–1097 (2011).
91. Aono, H., Imanaka, N. & Adachi, G. High Li+ Conducting Ceramics. Acc. Chem. Res. 27, 265–270 (1994).
92. Xu, K. Nonaqueous Liquid Electrolytes for Lithium-Based Rechargeable Batteries. Chem. Rev. 104, 4303–
4418 (2004).
93. Marcinek, M. et al. Electrolytes for Li-ion transport – Review. Solid State Ion. 276, 107–126 (2015).
94. Bates, J. B. et al. Electrical properties of amorphous lithium electrolyte thin films. Solid State Ion. 53–56,
Part 1, 647–654 (1992).
95. Bates, J. B. et al. Fabrication and characterization of amorphous lithium electrolyte thin films and
rechargeable thin-film batteries. J. Power Sources 43, 103–110 (1993).
96. Nowak, S., Berkemeier, F. & Schmitz, G. Ultra-thin LiPON films – Fundamental properties and
application in solid state thin film model batteries. J. Power Sources 275, 144–150 (2015).
97. 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).
98. Yu, X., Bates, J. B., Jellison, G. E. & Hart, F. X. A Stable Thin9Film Lithium Electrolyte: Lithium
Phosphorus Oxynitride. J. Electrochem. Soc. 144, 524–532 (1997).
99. Nimisha, C. S., Rao, G. M., Munichandraiah, N., Natarajan, G. & Cameron, D. C. Chemical and
microstructural modifications in LiPON thin films exposed to atmospheric humidity. Solid State Ion. 185,
47–51 (2011).
100. Schwöbel, A., Hausbrand, R. & Jaegermann, W. Interface reactions between LiPON and lithium studied by
in-situ X-ray photoemission. Solid State Ion. 273, 51–54 (2015).
101. Herbert, E. G., Tenhaeff, W. E., Dudney, N. J. & Pharr, G. M. Mechanical characterization of LiPON
films using nanoindentation. Thin Solid Films 520, 413–418 (2011).
102. Monroe, C. & Newman, J. The Effect of Interfacial Deformation on Electrodeposition Kinetics. J.
Electrochem. Soc. 151, A880–A886 (2004).
103. Monroe, C. & Newman, J. The Impact of Elastic Deformation on Deposition Kinetics at
Lithium/Polymer Interfaces. J. Electrochem. Soc. 152, A396–A404 (2005).
104. Motoyama, M., Ejiri, M. & Iriyama, Y. Modeling the Nucleation and Growth of Li at Metal Current
Collector/LiPON Interfaces. J. Electrochem. Soc. 162, A7067–A7071 (2015).
105. Sagane, F. et al. Effects of current densities on the lithium plating morphology at a lithium phosphorus
oxynitride glass electrolyte/copper thin film interface. J. Power Sources 233, 34–42 (2013).
106. Sagane, F., Shimokawa, R., Sano, H., Sakaebe, H. & Iriyama, Y. In-situ scanning electron microscopy
observations of Li plating and stripping reactions at the lithium phosphorus oxynitride glass electrolyte/Cu
interface. J. Power Sources 225, 245–250 (2013).
107. Fenton, D. E., Parker, J. M. & Wright, P. V. Complexes of alkali metal ions with poly(ethylene oxide).
Polymer 14, 589 (1973).
108. Quartarone, E. & Mustarelli, P. Electrolytes for solid-state lithium rechargeable batteries: recent advances
and perspectives. Chem. Soc. Rev. 40, 2525–2540 (2011).
109. Berthier, C. et al. Microscopic investigation of ionic conductivity in alkali metal salts-poly(ethylene oxide)
adducts. Solid State Ion. 11, 91–95 (1983).
110. 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).
Chapter 1. General context and battery state of the art
- 69 -
111. Xue, Z., He, D. & Xie, X. Poly(ethylene oxide)-based electrolytes for lithium-ion batteries. J. Mater. Chem.
A 3, 19218–19253 (2015).
112. Johansson, P., Tegenfeldt, J. & Lindgren, J. Modelling lithium ion transport in helical PEO by ab initio
calculations. Polymer 42, 6573–6577 (2001).
113. Maitra, A. & Heuer, A. Understanding Segmental Dynamics in Polymer Electrolytes: A Computer Study.
Macromol. Chem. Phys. 208, 2215–2221 (2007).
114. Maitra, A. & Heuer, A. Understanding Correlation Effects for Ion Conduction in Polymer Electrolytes. J.
Phys. Chem. B 112, 9641–9651 (2008).
115. Diddens, D. & Heuer, A. Simulation Study of the Lithium Ion Transport Mechanism in Ternary Polymer
Electrolytes: The Critical Role of the Segmental Mobility. J. Phys. Chem. B 118, 1113–1125 (2014).
116. Dukhanin, G. P., Dumler, S. A., Sablin, A. N. & Novakov, I. A. Solid polymeric electrolyte based on
poly(ethylene carbonate)-lithium perchlorate system. Russ. J. Appl. Chem. 82, 243–246 (2009).
117. Aurbach, D. et al. The Study of Electrolyte Solutions Based on Ethylene and Diethyl Carbonates for
Rechargeable Li Batteries I . Li Metal Anodes. J. Electrochem. Soc. 142, 2873–2882 (1995).
118. Plakhotnyk, A. V., Ernst, L. & Schmutzler, R. Hydrolysis in the system LiPF6—propylene carbonate—
dimethyl carbonate—H2O. J. Fluor. Chem. 126, 27–31 (2005).
119. Alloin, F., Sanchez, J. Y. & Armand, M. B. Conductivity measurements of LiTFSI triblock copolymers
with a central POE sequence. Electrochimica Acta 37, 1729–1731 (1992).
120. Han, H.-B. et al. Lithium bis(fluorosulfonyl)imide (LiFSI) as conducting salt for nonaqueous liquid
electrolytes for lithium-ion batteries: Physicochemical and electrochemical properties. J. Power Sources 196,
3623–3632 (2011).
121. Xu, K., Zhang, S., Jow, T. R., Xu, W. & Angell, C. A. LiBOB as Salt for Lithium-Ion Batteries:A Possible
Solution for High Temperature Operation. Electrochem. Solid-State Lett. 5, A26–A29 (2002).
122. Grünebaum, M. et al. Synthesis and electrochemistry of polymer based electrolytes for lithium batteries.
Prog. Solid State Chem. 42, 85–105 (2014).
123. Lascaud, S. et al. Phase Diagrams and Conductivity Behavior of Poly(ethylene oxide)-Molten Salt Rubbery
Electrolytes. Macromolecules 27, 7469–7477 (1994).
124. Lascaud, S. et al. Evidence for ion pairs and/or triple ions from transport measurements in mixed-alkali
polyether electrolytes. Electrochimica Acta 43, 1407–1414 (1998).
125. Vallée, A., Besner, S. & Prud’Homme, J. Comparative study of poly(ethylene oxide) electrolytes made with
LiN(CF3SO2)2, LiCF3SO3 and LiClO4: Thermal properties and conductivity behaviour. Electrochimica Acta
37, 1579–1583 (1992).
126. Fan, L.-Z., Wang, X.-L., Long, F. & Wang, X. Enhanced ionic conductivities in composite polymer
electrolytes by using succinonitrile as a plasticizer. Solid State Ion. 179, 1772–1775 (2008).
127. Zhu, C., Cheng, H. & Yang, Y. Electrochemical Characterization of Two Types of PEO-Based Polymer
Electrolytes with Room-Temperature Ionic Liquids. J. Electrochem. Soc. 155, A569–A575 (2008).
128. Kim, G. T. et al. UV cross-linked, lithium-conducting ternary polymer electrolytes containing ionic liquids.
J. Power Sources 195, 6130–6137 (2010).
129. Nagaoka, K., Naruse, H., Shinohara, I. & Watanabe, M. High ionic conductivity in poly(dimethyl siloxane-
co-ethylene oxide) dissolving lithium perchlorate. J. Polym. Sci. Polym. Lett. Ed. 22, 659–663 (1984).
130. Fonseca, C. P. & Neves, S. Characterization of polymer electrolytes based on poly(dimethyl siloxane-co-
ethylene oxide). J. Power Sources 104, 85–89 (2002).
131. Fergus, J. W. Ceramic and polymeric solid electrolytes for lithium-ion batteries. J. Power Sources 195, 4554–
4569 (2010).
132. Sarah Querelle. Synthese et utilisation de copolymeres triblocs ABA pour l’elaboration de membranes
poreuses a morphologies et performance controlees. (Universite de Montpellier, 2008).
133. Matsen, M. W. & Thompson, R. B. Equilibrium behavior of symmetric ABA triblock copolymer melts. J.
Chem. Phys. 111, 7139–7146 (1999).
134. Giles, J. R. M., Gray, F. M., MacCallum, J. R. & Vincent, C. A. Synthesis and characterization of ABA
block copolymer-based polymer electrolytes. Polymer 28, 1977–1981 (1987).
Chapter 1. General context and battery state of the art
- 70 -
135. Wang, C., Sakai, T., Watanabe, O., Hirahara, K. & Nakanishi, T. All Solid-State Lithium-Polymer Battery
Using a Self-Cross-Linking Polymer Electrolyte. J. Electrochem. Soc. 150, A1166–A1170 (2003).
136. Singh, M. et al. Effect of Molecular Weight on the Mechanical and Electrical Properties of Block
Copolymer Electrolytes. Macromolecules 40, 4578–4585 (2007).
137. Panday, A. et al. Effect of Molecular Weight and Salt Concentration on Conductivity of Block Copolymer
Electrolytes. Macromolecules 42, 4632–4637 (2009).
138. Gomez, E. D. et al. Effect of Ion Distribution on Conductivity of Block Copolymer Electrolytes. Nano
Lett. 9, 1212–1216 (2009).
139. Bouchet, R. et al. Charge Transport in Nanostructured PS–PEO–PS Triblock Copolymer Electrolytes.
Macromolecules 47, 2659–2665 (2014).
140. 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).
141. Chazalviel, J.-N. Electrochemical aspects of the generation of ramified metallic electrodeposits. Phys. Rev. A
42, 7355–7367 (1990).
142. Bouchet, R. et al. Single-ion BAB triblock copolymers as highly efficient electrolytes for lithium-metal
batteries. Nat. Mater. 12, 452–457 (2013).
143. Fujinami, T., Tokimune, A., Mehta, M. A., Shriver, D. F. & Rawsky, G. C. Siloxyaluminate Polymers with
High Li+ Ion Conductivity. Chem. Mater. 9, 2236–2239 (1997).
144. Benrabah, D., Sylla, S., Alloin, F., Sanchez, J.-Y. & Armand, M. Perfluorosulfonate-polyether based single
ion conductors. Electrochimica Acta 40, 2259–2264 (1995).
145. Watanabe, M., Suzuki, Y. & Nishimoto, A. Single ion conduction in polyether electrolytes alloyed with
lithium salt of a perfluorinated polyimide. Electrochimica Acta 45, 1187–1192 (2000).
146. Ito, K. & Ohno, H. Ionic conductivity of poly(ethylene oxide) having charges on the chain end. Solid State
Ion. 79, 300–305 (1995).
147. Sadoway, D. R. et al. Self-doped block copolymer electrolytes for solid-state, rechargeable lithium batteries.
J. Power Sources 97–98, 621–623 (2001).
148. Ryu, S.-W. et al. Effect of Counter Ion Placement on Conductivity in Single-Ion Conducting Block
Copolymer Electrolytes. J. Electrochem. Soc. 152, A158–A163 (2005).
149. Inceoglu, S. et al. Morphology–Conductivity Relationship of Single-Ion-Conducting Block Copolymer
Electrolytes for Lithium Batteries. ACS Macro Lett. 3, 510–514 (2014).
150. 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).
151. Ismail, I., Noda, A., Nishimoto, A. & Watanabe, M. XPS study of lithium surface after contact with
lithium-salt doped polymer electrolytes. Electrochimica Acta 46, 1595–1603 (2001).
152. Naudin, C. et al. Characterization of the lithium surface by infrared and Raman spectroscopies. J. Power
Sources 124, 518–525 (2003).
153. 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).
154. Fouache9Ayoub, S., Garreau, M., Prabhu, P. V. S. S. & Thevenin, J. Mass9Transport Properties of Lithium
Surface Layers Formed in Sulfolane9Based Electrolytes. J. Electrochem. Soc. 137, 1659–1665 (1990).
155. Tarascon, J.-M. & Armand, M. Issues and challenges facing rechargeable lithium batteries. Nature 414, 359–
367 (2001).
156. Fauteux, D., Massucco, A., McLin, M., Van Buren, M. & Shi, J. Lithium polymer electrolyte rechargeable
battery. Electrochimica Acta 40, 2185–2190 (1995).
157. Thevenin, J. Passivating films on lithium electrodes. An approach by means of electrode impedance
spectroscopy. J. Power Sources 14, 45–52 (1985).
158. Bouchet, R., Lascaud, S. & Rosso, M. An EIS Study of the Anode Li/PEO-LiTFSI of a Li Polymer
Battery. J. Electrochem. Soc. 150, A1385–A1389 (2003).
Chapter 1. General context and battery state of the art
- 71 -
159. Peled, E., Golodnitsky, D., Ardel, G. & Eshkenazy, V. The sei model—application to lithium-polymer
electrolyte batteries. Electrochimica Acta 40, 2197–2204 (1995).
160. Xu, W. et al. Lithium metal anodes for rechargeable batteries. Energy Environ. Sci. 7, 513–537 (2014).
161. Epelboin, I., Froment, M., Garreau, M., Thevenin, J. & Warin, D. Behavior of Secondary Lithium and
Aluminum9Lithium Electrodes in Propylene Carbonate. J. Electrochem. Soc. 127, 2100–2104 (1980).
162. Witten, T. A. & Sander, L. M. Diffusion-Limited Aggregation, a Kinetic Critical Phenomenon. Phys. Rev.
Lett. 47, 1400–1403 (1981).
163. Kirkby, M. J. The fractal geometry of nature. Benoit B. Mandelbrot. W. H. Freeman and co., San
Francisco, 1982. No. of pages: 460. Price: £22.75 (hardback). Earth Surf. Process. Landf. 8, 406–406 (1983).
164. Brady, R. M. & Ball, R. C. Fractal growth of copper electrodeposits. Nature 309, 225–229 (1984).
165. Wiley: Electrochemical Methods: Fundamentals and Applications, 2nd Edition - Allen J. Bard, Larry R.
Faulkner. Available at: http://eu.wiley.com/WileyCDA/WileyTitle/productCd-0471043729.html.
(Accessed: 19th April 2016)
166. Brissot, C., Rosso, M., Chazalviel, J.-N., Baudry, P. & Lascaud, S. In situ study of dendritic growth
inlithium/PEO-salt/lithium cells. Electrochimica Acta 43, 1569–1574 (1998).
167. C. Brissot. Etude des mecanismes de croissance d’agregats dendritiques, produits au cours de cyclage dans
les batteries au lithium a electrolyte polymere. (1998).
168. Fleury, V., Chazalviel, J.-N., Rosso, M. & Sapoval, B. The role of the anions in the growth speed of fractal
electrodeposits. J. Electroanal. Chem. Interfacial Electrochem. 290, 249–255 (1990).
169. Brissot, C., Rosso, M., Chazalviel, J.-N. & Lascaud, S. In Situ Concentration Cartography in the
Neighborhood of Dendrites Growing in Lithium/Polymer9Electrolyte/Lithium Cells. J. Electrochem. Soc.
146, 4393–4400 (1999).
170. Rosso, M., Gobron, T., Brissot, C., Chazalviel, J.-N. & Lascaud, S. Onset of dendritic growth in
lithium/polymer cells. J. Power Sources 97–98, 804–806 (2001).
171. Anna Teyssot, C. B. Inter-electrode in situ concentration cartography in lithium/polymer
electrolyte/lithium cells. J. Electroanal. Chem. 584, 70–74 (2005).
172. Monroe, C. & Newman, J. Dendrite Growth in Lithium/Polymer Systems A Propagation Model for Liquid
Electrolytes under Galvanostatic Conditions. J. Electrochem. Soc. 150, A1377–A1384 (2003).
173. Mullins, W. W. & Sekerka, R. F. Morphological Stability of a Particle Growing by Diffusion or Heat Flow.
J. Appl. Phys. 34, 323–329 (1963).
174. Barton, J. L. & Bockris, J. O. The Electrolytic Growth of Dendrites from Ionic Solutions. Proc. R. Soc.
Lond. Math. Phys. Eng. Sci. 268, 485–505 (1962).
175. Nishikawa, K., Chassaing, E. & Rosso, M. In situ concentration measurements around the transition
between two dendritic growth regimes. Electrochimica Acta 56, 5464–5471 (2011).
176. Aurbach, D., Gofer, Y., Ben-Zion, M. & Aped, P. An International Journal Devoted to all Aspects of
Electrode Kinetics, Interfacial Structure, Properties of Electrolytes, Colloid and Biological
ElectrochemistryThe behaviour of lithium electrodes in propylene and ethylene carbonate: Te major
factors that influence Li cycling efficiency. J. Electroanal. Chem. 339, 451–471 (1992).
177. Aurbach, D. & Zaban, A. Impedance spectroscope of lithium electrodes. J. Electroanal. Chem. 367, 15–25
(1994).
178. Takehara, Z. Future prospects of the lithium metal anode. J. Power Sources 68, 82–86 (1997).
179. Shiraishi, S., Kanamura, K. & Takehara, Z. Surface Condition Changes in Lithium Metal Deposited in
Nonaqueous Electrolyte Containing HF by Dissolution9Deposition Cycles. J. Electrochem. Soc. 146, 1633–
1639 (1999).
180. Abraham, K. M., Foos, J. S. & Goldman, J. L. Long Cycle9Life Secondary Lithium Cells Utilizing
Tetrahydrofuran. J. Electrochem. Soc. 131, 2197–2199 (1984).
181. Abraham, K. M. Recent developments in secondary lithium battery technology. J. Power Sources 14, 179–191
(1985).
Chapter 1. General context and battery state of the art
- 72 -
182. Mogi, R. et al. Effects of Some Organic Additives on Lithium Deposition in Propylene Carbonate. J.
Electrochem. Soc. 149, A1578–A1583 (2002).
183. Choi, J.-W. et al. Rechargeable lithium/sulfur battery with liquid electrolytes containing toluene as additive.
J. Power Sources 183, 441–445 (2008).
184. Ishikawa, M., Yoshitake, S., Morita, M. & Matsuda, Y. In Situ Scanning Vibrating Electrode Technique for
the Characterization of Interface Between Lithium Electrode and Electrolytes Containing Additives. J.
Electrochem. Soc. 141, L159–L161 (1994).
185. Umeda, G. A. et al. Protection of lithium metal surfaces using tetraethoxysilane. J. Mater. Chem. 21, 1593–
1599 (2011).
186. Marchioni, F. et al. Protection of Lithium Metal Surfaces Using Chlorosilanes. Langmuir 23, 11597–11602
(2007).
187. Choi, N.-S., Lee, Y. M., Seol, W., Lee, J. A. & Park, J.-K. Protective coating of lithium metal electrode for
interfacial enhancement with gel polymer electrolyte. Solid State Ion. 172, 19–24 (2004).
188. Wu, M., Wen, Z., Liu, Y., Wang, X. & Huang, L. Electrochemical behaviors of a Li3N modified Li metal
electrode in secondary lithium batteries. J. Power Sources 196, 8091–8097 (2011).
189. Hongkyung Lee, D. J. L. A simple composite protective layer coating that enhances the cycling stability of
lithium metal batteries. J. Power Sources 284, 103–108 (2015).
190. Zheng, G. et al. Interconnected hollow carbon nanospheres for stable lithium metal anodes. Nat.
Nanotechnol. 9, 618–623 (2014).
191. Kim, J.-S., Kim, D. W., Jung, H. T. & Choi, J. W. Controlled Lithium Dendrite Growth by a Synergistic
Effect of Multilayered Graphene Coating and an Electrolyte Additive. Chem. Mater. 27, 2780–2787 (2015).
192. Dollé, M., Sannier, L., Beaudoin, B., Trentin, M. & Tarascon, J.-M. Live Scanning Electron Microscope
Observations of Dendritic Growth in Lithium/Polymer Cells. Electrochem. Solid-State Lett. 5, A286–A289
(2002).
193. Tatsuma, T., Taguchi, M. & Oyama, N. Inhibition effect of covalently cross-linked gel electrolytes on
lithium dendrite formation. Electrochimica Acta 46, 1201–1205 (2001).
194. Eichinger, G. & Fabian, M. Comparison of organic and inorganic gelation agents in ethylene carbonate
based electrolytes for lithium-ion batteries. J. Power Sources 68, 381–386 (1997).
195. 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).
196. Tu, Z., Kambe, Y., Lu, Y. & Archer, L. A. Nanoporous Polymer-Ceramic Composite Electrolytes for
Lithium Metal Batteries. Adv. Energy Mater. 4, n/a-n/a (2014).
197. Lu, Y., Tu, Z. & Archer, L. A. Stable lithium electrodeposition in liquid and nanoporous solid electrolytes.
Nat. Mater. 13, 961–969 (2014).
198. Mayers, M. Z., Kaminski, J. W. & Miller, T. F. Suppression of Dendrite Formation via Pulse Charging in
Rechargeable Lithium Metal Batteries. J. Phys. Chem. C 116, 26214–26221 (2012).
199. Gireaud, L., Grugeon, S., Laruelle, S., Yrieix, B. & Tarascon, J.-M. Lithium metal stripping/plating
mechanisms studies: A metallurgical approach. Electrochem. Commun. 8, 1639–1649 (2006).
200. Liu, Y. et al. Lithium-coated polymeric matrix as a minimum volume-change and dendrite-free lithium
metal anode. Nat. Commun. 7, 10992 (2016).
201. Ryou, M.-H. et al. Excellent Cycle Life of Lithium-Metal Anodes in Lithium-Ion Batteries with Mussel-
Inspired Polydopamine-Coated Separators. Adv. Energy Mater. 2, 645–650 (2012).
202. Ding, F. et al. Dendrite-Free Lithium Deposition via Self-Healing Electrostatic Shield Mechanism. J. Am.
Chem. Soc. 135, 4450–4456 (2013).
203. Ryou, M.-H., Lee, Y. M., Lee, Y., Winter, M. & Bieker, P. Mechanical Surface Modification of Lithium
Metal: Towards Improved Li Metal Anode Performance by Directed Li Plating. Adv. Funct. Mater. 25, 834–
841 (2015).
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
Chapter 2. Ceramic protective layer alternative
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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
Cu-LiPON1,2-Ohara-LiPON1,2-Cu 2
250 2.30 ± 0.04
Cu-LiPON0,6-Ohara-LiPON0,6-Cu 1
160 1.38 ± 0.04
Cu-LiPON0,6-Ohara-LiPON0,6-Cu 2
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
Chapter 2. Ceramic protective layer alternative
- 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
Chapter 2. Ceramic protective layer alternative
- 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.
Chapter 2. Ceramic protective layer alternative
- 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)
Chapter 2. Ceramic protective layer alternative
- 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
Chapter 2. Ceramic protective layer alternative
- 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)
Chapter 2. Ceramic protective layer alternative
- 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.
Chapter 2. Ceramic protective layer alternative
- 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)
Chapter 2. Ceramic protective layer alternative
- 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
Chapter 2. Ceramic protective layer alternative
- 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)
Chapter 2. Ceramic protective layer alternative
- 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)
Ohara 25°C LiPON-Ohara-LiPON 25°C
193 kHz
300 kHz
27 kHz
15 kHz
c)
300
250
200
150
100
50
0
- Im
(Z)
(W)
300250200150100500
Re(Z) (W)
Ohara 50°C LiPON-Ohara-LiPON 50°C
0.66 MHz
0.80 MHz
57 kHz
100 kHz
d)
Chapter 2. Ceramic protective layer alternative
- 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
Chapter 2. Ceramic protective layer alternative
- 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
Chapter 2. Ceramic protective layer alternative
- 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
Chapter 2. Ceramic protective layer alternative
- 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.
Chapter 2. Ceramic protective layer alternative
- 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.
Chapter 2. Ceramic protective layer alternative
- 91 -
References of Chapter 2
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
Battery. J. Electrochem. Soc. 150, A1385–A1389 (2003).
Chapter 2. Ceramic protective layer alternative
- 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
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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:
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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
Chapter 3. Chemical and physical characterizations of block copolymer
- 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
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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).
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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
Chapter 3. Chemical and physical characterizations of block copolymer
- 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
Chapter 3. Chemical and physical characterizations of block copolymer
- 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
Chapter 3. Chemical and physical characterizations of block copolymer
- 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
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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)
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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).
Chapter 3. Chemical and physical characterizations of block copolymer
- 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+
Chapter 3. Chemical and physical characterizations of block copolymer
- 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
Chapter 3. Chemical and physical characterizations of block copolymer
- 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
Chapter 3. Chemical and physical characterizations of block copolymer
- 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.
Chapter 3. Chemical and physical characterizations of block copolymer
- 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
Copolymer Electrolytes. Macromolecules 40, 4578–4585 (2007).
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
Copolymer Electrolytes. Macromolecules 42, 4632–4637 (2009).
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).
12. Inceoglu, S. et al. Morphology–Conductivity Relationship of Single-Ion-Conducting Block Copolymer
Electrolytes for Lithium Batteries. ACS Macro Lett. 3, 510–514 (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).
Chapter 3. Chemical and physical characterizations of block copolymer
- 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
Battery. J. Electrochem. Soc. 150, A1385–A1389 (2003).
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).
Chapter 3. Chemical and physical characterizations of block copolymer
- 126 -
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).
31. Devaux, D. et al. Optimization of Block Copolymer Electrolytes for Lithium Metal Batteries. Chem.
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
Chapter 4. Dendritic growth in lithium symmetric cells
- 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.
Chapter 4. Dendritic growth in lithium symmetric cells
- 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
Chapter 4. Dendritic growth in lithium symmetric cells
- 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
Chapter 4. Dendritic growth in lithium symmetric cells
- 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)
Chapter 4. Dendritic growth in lithium symmetric cells
- 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.
Chapter 4. Dendritic growth in lithium symmetric cells
- 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).
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Chapter 4. Dendritic growth in lithium symmetric cells
- 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
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Chapter 4. Dendritic growth in lithium symmetric cells
- 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.
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Chapter 4. Dendritic growth in lithium symmetric cells
- 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.
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Chapter 4. Dendritic growth in lithium symmetric cells
- 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.
Chapter 4. Dendritic growth in lithium symmetric cells
- 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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Chapter 4. Dendritic growth in lithium symmetric cells
- 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).
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a)
Chapter 4. Dendritic growth in lithium symmetric cells
- 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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Chapter 4. Dendritic growth in lithium symmetric cells
- 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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Chapter 4. Dendritic growth in lithium symmetric cells
- 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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Chapter 4. Dendritic growth in lithium symmetric cells
- 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
Chapter 4. Dendritic growth in lithium symmetric cells
- 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.
Chapter 4. Dendritic growth in lithium symmetric cells
- 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.
Chapter 4. Dendritic growth in lithium symmetric cells
- 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
Chapter 4. Dendritic growth in lithium symmetric cells
- 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.
Chapter 4. Dendritic growth in lithium symmetric cells
- 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.
Chapter 4. Dendritic growth in lithium symmetric cells
- 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).
Chapter 4. Dendritic growth in lithium symmetric cells
- 151 -
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
Chapter 4. Dendritic growth in lithium symmetric cells
- 152 -
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
Chapter 4. Dendritic growth in lithium symmetric cells
- 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.
Chapter 4. Dendritic growth in lithium symmetric cells
- 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.
Chapter 4. Dendritic growth in lithium symmetric cells
- 155 -
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
Chapter 4. Dendritic growth in lithium symmetric cells
- 156 -
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
Chapter 4. Dendritic growth in lithium symmetric cells
- 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.
Chapter 4. Dendritic growth in lithium symmetric cells
- 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.
Chapter 4. Dendritic growth in lithium symmetric cells
- 159 -
References of Chapter 4
1. Rosso, M. et al. Dendrite short-circuit and fuse effect on Li/polymer/Li cells. Electrochimica Acta 51,
5334–5340 (2006).
2. Sannier, L., Bouchet, R., Santinacci, L., Grugeon, S. & Tarascon, J.-M. Lithium Metal Batteries
Operating at Room Temperature Based on Different PEO-PVdF Separator Configurations. J.
Electrochem. Soc. 151, A873–A879 (2004).
3. Schauser, N. S., Harry, K. J., Parkinson, D. Y., Watanabe, H. & Balsara, N. P. Lithium Dendrite
Growth in Glassy and Rubbery Nanostructured Block Copolymer Electrolytes. J. Electrochem. Soc. 162,
A398–A405 (2015).
4. Harry, K. J., Hallinan, D. T., Parkinson, D. Y., MacDowell, A. A. & Balsara, N. P. Detection of
subsurface structures underneath dendrites formed on cycled lithium metal electrodes. Nat. Mater. 13,
69–73 (2014).
5. Hallinan, D. T., Mullin, S. A., Stone, G. M. & Balsara, N. P. Lithium Metal Stability in Batteries with
Block Copolymer Electrolytes. J. Electrochem. Soc. 160, A464–A470 (2013).
6. Chazalviel, J.-N. Electrochemical aspects of the generation of ramified metallic electrodeposits. Phys.
Rev. A 42, 7355–7367 (1990).
7. 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).
8. Choudhury, S., Mangal, R., Agrawal, A. & Archer, L. A. A highly reversible room-temperature lithium
metal battery based on crosslinked hairy nanoparticles. Nat. Commun. 6, 10101 (2015).
9. Lu, Y., Korf, K., Kambe, Y., Tu, Z. & Archer, L. A. Ionic-Liquid–Nanoparticle Hybrid Electrolytes:
Applications in Lithium Metal Batteries. Angew. Chem. 126, 498–502 (2014).
10. Schaefer, J. L., Yanga, D. A. & Archer, L. A. High Lithium Transference Number Electrolytes via
Creation of 3-Dimensional, Charged, Nanoporous Networks from Dense Functionalized Nanoparticle
Composites. Chem. Mater. 25, 834–839 (2013).
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
Lithium-Metal Polymer Batteries. J. Am. Chem. Soc. 136, 7395–7402 (2014).
Chapter 4. Dendritic growth in lithium symmetric cells
- 160 -
12. Sannier, L., Bouchet, R., Rosso, M. & Tarascon, J.-M. Evaluation of GPE performances in lithium
metal battery technology by means of simple polarization tests. J. Power Sources 158, 564–570 (2006).
13. Rosso, M., Gobron, T., Brissot, C., Chazalviel, J.-N. & Lascaud, S. Onset of dendritic growth in
lithium/polymer cells. J. Power Sources 97–98, 804–806 (2001).
14. Liu, S. et al. Effect of nano-silica filler in polymer electrolyte on Li dendrite formation in
Li/poly(ethylene oxide)–Li(CF3SO2)2N/Li. J. Power Sources 195, 6847–6853 (2010).
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
bis(trifluoromethanesulfonyl)imide into polymer electrolyte on Li dendrite formation in
Li/poly(ethylene oxide)-Li(CF3SO2)2N/Li. J. Power Sources 196, 7681–7686 (2011).
17. Radon, J. On the determination of functions from their integral values along certain manifolds. IEEE
Trans. Med. Imaging 5, 170–176 (1986).
18. A. A. MacDowell, D. Y. P. X-ray micro-Tomography at the Advanced Light Source. Proc. SPIE 8506,
(2012).
19. Fiji Is Just ImageJ. Available at: http://fiji.sc/wiki/index.php/Fiji. (Accessed: 10th February 2016)
20. http://www.fei.com/software/avizo3d/.
21. Harry, K. J., Liao, X., Parkinson, D. Y., Minor, A. M. & Balsara, N. P. Electrochemical Deposition and
Stripping Behavior of Lithium Metal across a Rigid Block Copolymer Electrolyte Membrane. J.
Electrochem. Soc. 162, A2699–A2706 (2015).
22. Maia, F. et al. Compressive phase contrast tomography. in 7800, 78000F–78000F–5 (2010).
23. Frianeza-Kullberg, T. C. & Salmon, D. J. Removal of lithium nitride from lithium metal. (1988).
24. Harry, K. J., Liao, X., Parkinson, D. Y., Minor, A. M. & Balsara, N. P. Electrochemical Deposition and
Stripping Behavior of Lithium Metal across a Rigid Block Copolymer Electrolyte Membrane. J.
Electrochem. Soc. 162, A2699–A2706 (2015).
25. Tatsuma, T., Taguchi, M. & Oyama, N. Inhibition effect of covalently cross-linked gel electrolytes on
lithium dendrite formation. Electrochimica Acta 46, 1201–1205 (2001).
Chapter 4. Dendritic growth in lithium symmetric cells
- 161 -
26. Anna Teyssot, C. B. Inter-electrode in situ concentration cartography in lithium/polymer
electrolyte/lithium cells. J. Electroanal. Chem. 584, 70–74 (2005).
27. Monroe, C. & Newman, J. Dendrite Growth in Lithium/Polymer Systems A Propagation Model for
Liquid Electrolytes under Galvanostatic Conditions. J. Electrochem. Soc. 150, A1377–A1384 (2003).
28. Monroe, C. & Newman, J. The Effect of Interfacial Deformation on Electrodeposition Kinetics. J.
Electrochem. Soc. 151, A880–A886 (2004).
29. Monroe, C. & Newman, J. The Impact of Elastic Deformation on Deposition Kinetics at
Lithium/Polymer Interfaces. J. Electrochem. Soc. 152, A396–A404 (2005).
30. Orsini, F. et al. In situ Scanning Electron Microscopy (SEM) observation of interfaces within plastic
lithium batteries. J. Power Sources 76, 19–29 (1998).
Chapter 4. Dendritic growth in lithium symmetric cells
- 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
2. Results and discussion ................................................................................................................................. 167
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
3. Results and discussion ................................................................................................................................. 192
a. Experimental results ................................................................................................................................... 192
b. Discussion .................................................................................................................................................... 196
Conclusions ............................................................................................................................................................ 201
References of chapter 5 ........................................................................................................................................ 204
Chapter 5. The polymer-ceramic composite
- 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.
Chapter 5. The polymer-ceramic composite
- 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.
Chapter 5. The polymer-ceramic composite
- 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.
Chapter 5. The polymer-ceramic composite
- 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.
Chapter 5. The polymer-ceramic composite
- 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
Chapter 5. The polymer-ceramic composite
- 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)
Chapter 5. The polymer-ceramic composite
- 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
Chapter 5. The polymer-ceramic composite
- 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.
Chapter 5. The polymer-ceramic composite
- 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
Chapter 5. The polymer-ceramic composite
- 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
Chapter 5. The polymer-ceramic composite
- 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.
Chapter 5. The polymer-ceramic composite
- 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
Chapter 5. The polymer-ceramic composite
- 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
-0.2
-0.4
-0.6
Vo
lta
ge
(V
)
8006004002000
Time (h)
After charge After discharge
a)-0.6
-0.4
-0.2
0.0
0.2
0.4
0.6
Po
ten
tia
l (V
)
80757065605550time (h)
Composite with SEO BCE
b)
Chapter 5. The polymer-ceramic composite
- 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
Chapter 5. The polymer-ceramic composite
- 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
80
60
40
20
0
Rel (W
.cm
2)
4003002001000
Charge passed (C.cm-2
)
806040200
Cycle number
'After discharge : composite SEO 55-52' 'After discharge : composite SEO 55-52'
a)
350
300
250
200
150
100
50
0
Rin
t (W
.cm
2)
4003002001000
Charge passed (C.cm-2
)
806040200
Cycle number
After charge : composite SEO 55-52 After discharge : composite SEO 55-52
b)
Chapter 5. The polymer-ceramic composite
- 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
-0.2
0.0
0.2
0.4
Voltage
(V
)
300250200150100500Time (h)
Composite with Si-EO BCEa)
-0.4
-0.2
0.0
0.2
0.4
Vo
lta
ge
(V
)
1201151101051009590Time (h)
b)
Cycles 10 to 13
-0.4
-0.2
0.0
0.2
0.4
Vo
lta
ge
(V
)
134132130128126124Time (h)
Cycle number 14
c)
Chapter 5. The polymer-ceramic composite
- 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.
Chapter 5. The polymer-ceramic composite
- 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
Charge passed (C.cm-2
)
302520151050
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
Charge passed (C.cm-2
)
302520151050
Cycle number
After charge : composite with diSI 'After discharge : composite with diSI'
b)
Chapter 5. The polymer-ceramic composite
- 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)
Chapter 5. The polymer-ceramic composite
- 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.
Chapter 5. The polymer-ceramic composite
- 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
800
600
400
200
0
Rel (W
.cm
2)
1601208040
Charge passed (C.cm-2
)
302520151050
Cycle number
'After charge : composite with triSI' 'After discharge : composite with triSI'
a)
1400
1200
1000
800
600
400
200
0
Rin
t (W
.cm
2)
16014012010080604020
Charge passed (C.cm-2
)
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Cycle number
After charge : composite with triSI After discharge : composite with triSI
b)
Chapter 5. The polymer-ceramic composite
- 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
Chapter 5. The polymer-ceramic composite
- 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.
Chapter 5. The polymer-ceramic composite
- 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.
Chapter 5. The polymer-ceramic composite
- 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.
Chapter 5. The polymer-ceramic composite
- 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.
Chapter 5. The polymer-ceramic composite
- 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.
Chapter 5. The polymer-ceramic composite
- 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
Chapter 5. The polymer-ceramic composite
- 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
-0.1
0.0
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
Chapter 5. The polymer-ceramic composite
- 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
Chapter 5. The polymer-ceramic composite
- 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)
Composite with SEO BCE
Chapter 5. The polymer-ceramic composite
- 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)
Composite with SEO BCE
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
Chapter 5. The polymer-ceramic composite
- 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)
Chapter 5. The polymer-ceramic composite
- 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
Current density (mA.cm-2
)
Composite : Ohara GC-Si-EO-Si Correlation with EIS
Chapter 5. The polymer-ceramic composite
- 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
Current density (mA.cm-2
)
Composite : Ohara GC-SEO Correlation with EIS
Chapter 5. The polymer-ceramic composite
- 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)
Chapter 5. The polymer-ceramic composite
- 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
Current density (mA.cm-2
)
SEO 55-52 Umodel without diffusion
Chapter 5. The polymer-ceramic composite
- 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
Chapter 5. The polymer-ceramic composite
- 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.
Chapter 5. The polymer-ceramic composite
- 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).
3. Devaux, D. et al. Optimization of Block Copolymer Electrolytes for Lithium Metal Batteries. Chem.
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).
Chapter 5. The polymer-ceramic composite
- 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).
Chapter 5. The polymer-ceramic composite
- 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
Conclusions and perspectives
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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.
Conclusions and perspectives
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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
Conclusions and perspectives
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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
1. Inceoglu, S. et al. Morphology–Conductivity Relationship of Single-Ion-Conducting Block Copolymer
Electrolytes for Lithium Batteries. ACS Macro Lett. 3, 510–514 (2014).
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).
4. Monroe, C. & Newman, J. The Effect of Interfacial Deformation on Electrodeposition Kinetics. J.
Electrochem. Soc. 151, A880–A886 (2004).
5. Monroe, C. & Newman, J. The Impact of Elastic Deformation on Deposition Kinetics at
Lithium/Polymer Interfaces. J. Electrochem. Soc. 152, A396–A404 (2005).
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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