Thèse De Doctorat

Ana Cristina

Martinez Maciel

Mme Laurence

Professeur M Daniel Lemordant

Émérite

M Michel

Sylvie Mme

Grugeon

M Pierre

Docteur Ingénieur

Renault Groupe

M Stéphane Examinateur

Professeur Laruelle

Tristan Lombard

François Rabuel

Emiliano Primo

Matthieu Courty

Gregory Gachot

Dominique Larcher

Christine Surcin

Serge Pilard

Sébastien Rigaud

David Mathiron

Loïc Dupont

Carine Davoisne

Arash Jamali

Laura

Mariem Caro Ronan

Julio César

Adrián Pérez

a ñL

First of all, I would like

1. Rechargeable batteries for automotive applications 1.1. Generalities of batteries

Description of a battery

A battery is a system that converts chemical energy into electricity through electrochemical reactions. It is basically constituted of two electrodes, the energy-providing, an electrolyte, ensuring the ionic conduction within the cell, and a separator, preventing contact between electrodes. Since the discovery of the electrochemical principles by Alexandro Volta in 1799, the functioning principle is the following: during discharge, at the anode occurs an oxidation (or anodic reaction) and at the cathode a reduction (or cathodic reaction), while the charge compensation is achieved by the electron flow through the external circuit 1 Batteries can be classified in primary or secondary. In secondary batteries, the redox reaction is reversible, so the battery is rechargeable; whereas in primary batteries, the redox reaction will solely proceed spontaneously in the forward direction. Precisely the rechargeability of secondary batteries has allowed them to be widely used in numerous applications such as telephones, computers and electric vehicles.

Characteristics to describe batteries

Batteries can be defined by different characteristics [1][2][3] , enlisted hereby.

• Voltage. In a battery, the nominal voltage (VO) is the average voltage that the battery outputs during cycling, at the beginning of its cycle life. The open circuit voltage (OCV) corresponds to the voltage between the battery terminals when no load is applied.

• Current. The current (i) is the rate of electric charge flow.

• C-rate. It refers to the current applied to discharge the battery per unit of time (hours). For instance, a C-rate of C/2 means that the total capacity will be delivered in 2 hours.

• Capacity. The capacity (Q), measured in ampere-hour (Ah), is the total amount of charge delivered from a battery until it is completely discharged. The theoretical capacity of an electrode material, both negative and positive, can be calculated if the quantity (in mol) of exchanged electrons (n) and the molecular weight (M) of the designed materials are known (Eq. I.1). F is the Faraday constant (96 485 C.mol -1 ). In theory, smaller molecular weight and more exchanged electrons in a material deliver higher capacity. It is measured using gravimetric specific capacity (Ah.kg -1 , mAh.g -1 ).

! = $% &.( * * [I.1]
• Energy (E). Energy is the product of the storage capacity by the average potential of the cell.

+ = ! • -. [I.2]
Energy density represents the amount of energy stored per unit mass or volume. Similar to battery capacity, it is measured in Wh.kg -1 , mWh.g -1 or Wh.L -1 , and mWh.cm -3 .

• Power (P). The power is the product of current i and potential VO (Eq. I.3). It measures the amount of current flowing at a given electric potential. Power is measured in Watts (W) and specific power in W.kg -1 , mW.g -1 . The power density of a battery refers to the energy that can be delivered per unit of time, measured in W.L -1 or mW.cm -3 .

/ = 0 • -. [I.3]
• State of charge (SOC). The SOC represents the ratio between the available capacity and the total capacity (Eq. I.4). The units of SOC are percentage points. • State of health (SoH). The SoH reflects the condition of a battery (or battery pack), compared to the battery in its original state. The units of SoH are percentage points. A battery's SoH is 100% at the time of manufacture and decreases over time.

• Coulombic efficiency (CE). It is the percentage ratio between discharge and charge capacities, given from one complete charge-discharge cycle.

• Cycle life. Cycle life corresponds to the number of charge and discharge cycles that a battery can complete before its capacity reaches a threshold that is dependent on the application. It depends on the structural stability of the electrode materials and the stability of the electrolyte and interphases during the charge/discharge process.

History of batteries in electric vehicles

Electric vehicles are older than we usually think. The concept of electric motor was first introduced by Michael Faraday in 1821 and was perfectioned from 1830s as a rotating electric motor by the physicists and inventors Ányos Jedlik, Moritz von Jacobi, Thomas Davenport, among others 4 . Initially, these motors were used in small machines but very quickly, they started to be noticeable as alternatives to horse-drawn carriages, thanks to the development of the lead-acid (Pb-acid) secondary battery by Gaston Planté in 1859 and the commercial distribution of electricity from 1880 5 . Around the 1890s, many entrepreneurs started marketing electric cars in major cities around the world, and by 1900, the number of electric cars outnumbered the gasoline cars by almost a factor of two, because they were less polluting, less noisy and did not require the rather dangerous and difficult to use hand crank starter 6 . These cars were powered by a rechargeable Pb-acid battery that exhibited low specific energy of 35-40 Wh.kg -1 and were capable to drive 50-150 km on a single charge, with top speed of 35 km.h -1 6 . A Pb-acid battery comprises a spongy lead plate as active material in the negative electrode and lead dioxide (PbO2) as active material in the positive electrode. These plates are immersed in an electrolyte of sulfuric acid (H2SO4) and distilled water, and the energy is obtained from the reaction of H2SO4 with Pb and PbO2 to produce lead sulfate (PbSO4) and water 3 (Eq. I.5):

/C + /C2 E + 2G E 12 H ↔ 2/C12 H + G E 2 [I.5]
Owing to the limited infrastructure to charge an electric vehicle in comparison to the easy access to fuel containers to refuel ICE vehicles and the decreasing price of gasoline in the following years, ICE vehicles started to regain the market, so that from 1930, the electric vehicles were virtually inexistent. From then, electric vehicles were mainly used in the aeronautic and military fields 5 .

Years later, in the 1990s, concerns about the heavy contribution of the transportation sector to greenhouse emissions, the increasing price of fossil fuels and the commitment in the Kyoto Protocol to reduce CO2 emissions 7 intensified again the development of electric vehicles.

Companies like Daimler-Chrysler, Ford, and General Motors initially used Pb-acid batteries, while Honda and Toyota used nickel metal-hydride (Ni-MH) batteries and Peugeot, Renault and Citröen nickel-cadmium batteries (Ni-Cd) 8 . A well-known electrified vehicle in the modern history of electric vehicles is the hybrid (HEV) Toyota Prius launched in 1997, powered with a Ni-MH battery pack with a capacity of 6.5 Ah. The Ni-MH battery was used as power source during acceleration and was recharged during the functioning of the vehicle in ICE-mode 9 . It exhibited specific energy between 50-70 Wh.kg -1 and used nickel oxide hydroxide (NiOOH) as positive electrode and a hydrogen-absorbing alloy (MH) as negative electrode (Eq. I.6). The manufacturing, the size, as well as the low power density limited this type of battery for a generalized usage.

MH + NiOOH « M + Ni(OH)2 [I.6]
The Ni-Cd batteries have the same positive electrode than the Ni-MH battery and metallic cadmium is used as active material in the negative electrode producing cadmium hydroxide (Cd(OH)2) during its oxidation (Eq. I.7). Through these reactions these batteries exhibited specific energy of around 40-60 Wh.kg -1 . Owing to the materials' price higher than Pb-acid and the toxicity and environmental hazard of cadmium, these batteries were not widely used 10 .

2NiOOH + Cd + 2H2O « 2Ni(OH)2 + Cd(OH)2 [I.7]

Around the same years, rechargeable lithium-ion batteries (LIBs) started heavy commercialization by Sony in 1991 11 . Quickly, most of the companies implemented LIB technology in their electric vehicles 12 . The LIBs are rechargeable devices composed of two lithium intercalation electrodes and an ionically conductor, electronically insulator porous separator soaked with non-aqueous electrolyte. The functioning principle is simple, LIBs are based in the lithium intercalation/deintercalation chemistry: during charge, lithium ions deintercalate from the positive electrode (LiMO2 as example, where M = transition metals) and intercalate into the negative electrode (graphite as example), by migrating through the electrolyte (Eq. I.8 and I.9). Inversely, in the discharge process, lithium ions deintercalate 13 .

An alternative to LIBs is the lithium solid polymer batteries commercialized by Bolloré ® in Autolib electric vehicles between 2011-2018 14 . It consisted in the use of LiFePO4 at the positive electrode and metallic lithium at the negative electrode. The lithium ions flow through a solid polymer electrolyte (typically polyethylene oxide) in which a lithium salt is dissolved 15 . The necessity to keep the battery at high temperature (60-80°C) throughout the usage limits its application in the automotive industry. HEV. On the contrary, considering its weight, size occupation and durability, rechargeable

LIBs are the best candidates for automotive applications 16 . The commercial electric vehicles containing LIBs are classified in HEVs, plug-in hybrid electric vehicles (PHEVs) and full electric vehicles (EVs). Each one differs in the driving range, the charge manner and the use. In HEVs, the main energy source is gasoline, so lowenergy batteries are enough. These vehicles are repeatedly charged and discharged on the road and thus do not need to be separately charged. PHEVs use a small combustion engine and an electric motor, with a small battery similar to the HEVs', although it gets charged externally and more energy density is needed to drive longer distance. Lastly, the EVs operate purely with an electric motor and an outer source of electricity that is stored in a battery pack with variable capacity and energy 17 . Their biggest advantage over PHEVs and HEVs is the driving range >100km within a single full charge and the zero-emission while driving.

The focus of this work is the study of LIBs for EVs applications within the framework of Renault's development plan, consequently, the Pb-acid, Ni-Cd and Ni-MH technologies will not be further discussed.

Panorama of LIBs in the electric vehicle industry

According to the International Energy Agency's (IEA) Global EV Outlook 2019 18 report, in 2018, the EV fleet exceeded 5.1 million globally, from which 45% were in China, 24% in Europe and 22% in the United States; an amazing global increase of 63% in comparison to the previous year 19 (Fig. I.4). In Norway, EVs represent 46% of the vehicles' market share, however globally the market share is still around 2.2%, mostly dominated by ICE vehicles, number that is expected to grow globally up to 30% by 2030 18 . Among the reasons behind the actual low global market share, are the limitations of an EV compared to an ICE vehicle, such as the availability of charging stations, the driving range and the price. Concerning the price, Deloitte's report of 2019 estimates that the market will reach a tipping point in 2022 when the cost to own an EV will be equivalent to an ICE vehicle 20 . Part of the price drop will be due to the decrease of the battery pack price, the most expensive element in an EV. Two different predictions estimated the price per kWh to drop to $62 18 and $95 21 (Fig. 21 .

Commercialization of EVs goes hand in hand with the requirements of the market: efficiency, safety, improved driving range and low-cost. Below, those are addressed, based on several sources 1,3,6,11,18 .

• High energy density. In EVs, the energy density is translated as driving range. The existent LIB technologies for city vehicles are limited to a driving range around 400 km p 12 , depending on the capacity of their battery, whereas the expectations of the market go above 600 km. For example, the Nissan Leaf ® can reach 252 km with 40 kWh battery pack, the Volkswagen e-Golf ® can reach 300 km with 26.5 kWh and the Renault Zoe ® 400 km with 41 kWh. To achieve longer driving ranges, Andre et al., 22 predicts that the next generations of LIBs (2025) will require positive electrode materials with higher energy densities targeting at least 800 Wh.L -1 and 300 Wh.kg -1 at the cell level.

• Low price. Important characteristic to be competitive against ICE vehicles.

• Size and weight. A battery directed to the electric vehicle market, should be as light and as small as possible.

• Safety. Ensuring safety is particularly important for the constructors of EVs. In the battery, sources for safety concern are: the volatile and flammable solvents in the electrolyte, eventual electrical short circuit between the electrodes and thermal runaway [23][24][25] ; to mention some.

• Long cycle life. Cycle life is an indicator of how long the battery can be used. The ideal cycle life of the battery for an EV is expected to exceed 15 years, compared to that of current portable devices, which span approximately five years 3 . Extreme operating conditions such as high/low temperature, deep discharge and repeated fast charge/discharge, favor early degradation of the battery components and shortens cycle life.

Fig. I.6

summarizes the target values of some of these requirements at the cell level in 2015, and the targets in 2022 and 2025 22 . In 2025, safety and power are not expected to vary so much, but the cost is expected to significantly decrease. In terms of performance, the energy density, specific energy, power at low temperature are expected to increase, as well as the charging current, as faster charging is needed. Although detailed values vary with the strategies of each EVs constructor, this figure offers a general picture of the complexity associated with future challenges. from Andre et al. 22 .

Components of the lithium-ion batteries

One of the most crucial challenges for future LIBs is the energy density, which is tightly associated with the choice of electrodes and their compatibility with the rest of the battery components. To understand the synergy between components, each component is thoroughly described below.

Electrolyte

The standard commercial non-aqueous electrolyte consists of a lithium salt dissolved in a mix of carbonate solvents and electrolyte additives. The basic requirements are the following 11,26,27 :

• High ionic conductivity in a broad temperature range. The minimum required for lithium ions migration is 10 -3 S.cm -1 at room temperature.

• Low viscosity. Viscosity, measured in cP or Pascal-second (Pa•s), is inversely related to ionic conductivity, in a sense that high viscosity is translated as low ionic conductivity.

• Safety. Since most of the organic solvents are flammable, in the case of a leakage, solvents with high ignition point and low toxicity are preferred.

• Cost. The price is a key issue to be considered, since depending on the choice of salt, solvent and additives, the price can be very elevated. For a typical liquid electrolyte composition, the price is 6% of the total cost of a battery 28 .

Regarding the requirements for the solvents, they are expected to present high dielectric constant (e), that is translated in greater ionic dissociation and higher ionic conductivity; and good chemical and electrochemical stability toward the diverse metals and polymers in the system. The physicochemical properties of the most used organic solvents are illustrated in In order to meet the standards of dielectric constant, chemical/electrochemical stability and safety; in most of the commercial electrolytes a mixture of two or three solvents is used.

Despite their high viscosity, cyclic carbonates are a must-have to increase the ionic conductivity. Additionally, EC allows the formation of high-quality solid electrolyte interphase (SEI) layer, required to prevent electrolyte degradation and graphite exfoliation along cycling 29 . On the contrary, PC is used less often because of its tendency to exfoliate the graphite 30 . Other important feature of electrolytes is the transference number, which is the fraction of the total electrical current carried in an electrolyte by a given ionic species 11 . In LIBs, the ideal Li + transference number is 1, however, in a typical LiPF6-based electrolyte, the Li + transference number is around 0.3.

Other component of the electrolyte is the lithium salt. Lithium salts usually comprise a large anion with delocalized charge to ensure good anion-cation dissociation in the solvents and to limit the ion-pair formation 28 . The most studied for LIBs are lithium perchlorate (LiClO4), lithium hexafluoroarsenate (LiAsF6), lithium tetrafluoroborate (LiBF4), lithium bis(trifluoromethanesulfonyl)imide (LiTFSI), lithium bis(fluorosulfonyl)imide (LiFSI) and lithium hexafluorophosphate (LiPF6).

LiPF6 is the industry's preferred choice owing to its high ionic conductivity in organic solvents (Table I.2), good electrochemical stability and its ability to protect the aluminum current collector from corrosion 31,32 . The other salts are not frequently used in EVs for several reasons: LiClO4 presents safety issues arising from the oxidative environment during charging, however it is used in primary batteries 29 ; LiAsF6 28 is very conductive and as such, does not present any environmental risk as its solvation shell stays intact in biological medium.

However, its reduction products, as in SEI, are very toxic; LiBF4 is not very conductive, and thus, can only be used in low-rate applications 33 ; LiTFSI allows severe aluminum corrosion near 3.6 V, due to the solubility of Al(TFSI)3 in alkyl carbonate solvents 34 ; similarly, LiFSI presents aluminum corrosion problems below 4 V vs. Li/Li + 35 , although it can be used in small quantities (< 0.33M) with LiPF6 without facing corrosion up to 4.3 V vs. Li/Li + 36 . (i) poor thermal stability above 55°C, decomposing into lithium fluoride (LiF) and phosphorus pentafluoride (PF5) (Eq. I.10). Then, PF5 is capable to react with the solvents to generate highly toxic chemicals 37 and promote solvents polymerization [38][39][40] .

LiPF6 à LiF + PF5 [I.10]

(ii) the reaction between PF5 and water traces produces toxic hydrofluoric acid (HF) and gaseous phosphoryl fluoride (POF3) 41,42 (Eq. I.11):

PF5 + H2O à POF3 + 2HF [I.11]
The last components of the electrolyte are the electrolyte additives that are usually added in small quantities (<10 wt.%). Some additives act on the electrodes' interfaces to improve the quality of the SEI or to extend the electrochemical window 23 . There are other additives that improve electrolyte properties (wettability), stabilize LiPF6 and enhance safety (overcharge protector, fire retardant) 43 .

Separator

The separator is a microporous membrane placed between the positive and negative electrode to prevent physical contact between the electrodes. The characteristics of a good separator are:

electrochemically inactive, microporous enough to let solvated lithium ions flow, possesses good mechanical strength, high wettability and to be chemically stable toward the electrolyte components 44,45 . Commercial separators are most of the times made of microporous polyolefins such as polyethylene (PE), polypropylene (PP, Fig. I.7) and polyethylene terephthalate (PET). They exhibit pores of 0.03-1 mm large and 30-50% porosity 1 .

The separators can be monolayer, multilayer and ceramic-coated to contribute to the safety of the battery. In the event of high temperatures caused by thermal runaway, the PE in a trilayer PP/PE/PP separator, melts around 140°C, shutting down the pores and interrupting ionic conduction 46,47 . Similarly, an alumina-coated separator provides thermal resistance and improves wettability due to the additional high surface area given by the alumina particles 48 . 49 .

Binder, conducting agent and current collectors

In addition to the active materials, the binder, conducting agent and current collectors are the other constituents of the electrodes (Fig I .8). The electrodes are produced from the combination of the binder and conducting agent with the active material in a solvent, usually N-metyl-2-pyrrolidone (NMP) or water. Then, this slurry is coated onto a foil of current collector that is heated to evaporate the solvent and obtain the finished electrode. The primary function of the binder is to keep the electrode's mechanical integrity and accommodate the volumetric expansion-contraction of active materials under successive lithium intercalation/deintercalation. In a negative electrode made of graphite, the structural expansion can reach 9.5% along the c-axis 50 . Thermal, chemical and electrochemical resistance are the other relevant characteristics. Suitable polymers for binders are polyvinyl difluoride (PVdF), styrene-butadiene rubber (SBR), carboxy methyl cellulose (CMC) and polytetrafluoroethylene (PTFE) 51 (Fig. I.9). For the positive electrode, PVdF remains the best option to maintain mechanical integrity 52 , using NMP as solvent. For the negative electrode a mixture of SBR/CMC is often preferred, using water as solvent. The mass-ratio of active materials and binder is usually in the range of 96:4 -88:12 11 . A conducting agent is desired to help the electronic conduction across the electrodes. The characteristics of an ideal conducting agent are: high dispersibility, high surface area and good electronic conductivity 53 . Carbon-based materials are the most commonly used, some examples are: carbon black, acetylene black and ketjen black. The difference between them is the degree of graphitization and porosity, which is highly dependent on the supplier's synthesis, that ultimately defines the microstructure and the surface area 11 .

Finally, the current collector is the intermediary passage of electrons from the active material to the external circuit, and inversely. The metals used as current collectors are selected as function of their electrochemical stability window and the typical thickness of the foils range between 6-20 µm. For the positive electrode, aluminum foils are used due to its low price, low weight and good electronic conductivity. Additionally, the use of LiPF6 salt protects aluminum from corrosion, since a passivation layer of AlF3 is formed from the native Al2O3 layer airformed. If the negative electrode has a low-operating potential, like graphite, the aluminum current collector is not used, due to the Li-Al alloying process at low potentials (~0.33 V vs.

Li/Li + ) 54 . Instead, copper foils are used, since copper does not intercalate or alloy with lithium at low operating potentials, moreover it is a good electronic conductor.

The negative electrode

An ideal active material for the negative electrode in a LIB displays low operating potential, high specific capacity and presents no significant crystal structure change during lithium intercalation/deintercalation 55 . To fulfill that, many different active materials have been tested

in the literature: phosphides, lithium alloys, oxides, silicon-based, titanium-based and carbonbased materials, to mention some. Lithium itself, is very interesting as active material due to its low density of 0.534 g.cm -3 , high theoretical specific capacity of 3850 mAh.g -1 , and low electrode potential of -3.04 V vs. (H + /H2) 56 . However, lithium metal disadvantages are the price, durability and safety. Durability because it participates in side reactions with the LiPF6based electrolyte that consume lithium ions, form thicker SEI and increase internal resistance;

and safety due to dendrites growth and consequent short circuit risk. Alternatives such as lithium alloys with Si, Ge and Sn have been investigated 57 . Despite the high discharge capacities from 500 to 4700 mAh.g -1 , they are limited by the large volume expansioncontraction that occurs during cycling, which in turn induces mechanical disintegration and early electrode failure 58 . Similarly, conversion materials made of phosphides, nitrides, sulfides and oxides have been under heavy investigation to reduce the mechanical issues and the inherent poor kinetics of these materials 59,60 .

Graphite is generally the preferred choice in almost every LIB application. Graphite is composed of graphene conductive layers with carbon atoms of the sp 2 hybrid orbital layered along a hexagonal plane (Fig. I.10). Amongst its advantages, it presents good electronic conductivity due to the delocalized π electrons in-between graphene layers, the intercalation occurs at potentials below 0.25 V vs. Li/Li + and it exhibits a theoretical specific capacity of 372 mAh.g -1 , by storing one Li + per six carbon atoms (LiC6) between its graphene layers 11 (Eq. I. On the other hand, for high-energy applications, silicon can be added to graphite thus resulting in high electrode capacity and superior performance 11 .

A notable characteristic of the graphite electrodes is the electrochemical decomposition of the alkyl carbonate-based electrolyte, that always occurs during the first charge/discharge cycles (called formation cycles), due to the electrochemical stability limits of the electrolyte. The deposit of solvent reduction products and decomposition products from LiPF6 lead to the construction of the SEI, as illustrated in Fig. I.12. The SEI formation depends on many variables: the type of graphite, temperature, electrochemical conditions and the electrolyte composition. A good SEI film improves cycle life by preventing further electrolyte decomposition, allowing the facile diffusion of lithium ions, reducing the concentration polarization and preventing graphite exfoliation 58,61 . However, the presence of the SEI film increases the internal resistance and consumes part of the available lithium ions, leading to both power and capacity loss. A commercial competitor to graphite is lithium titanium oxide, Li4Ti15O12 (LTO). During discharge, LTO can intercalate up to three Li + in the octahedral vacant sites (Eq. I.13) and achieve 175 mAh.g -1 at a mean potential of 1.55 V vs. Li/Li + 63 .

Li4Ti15O12 + 3Li + + 3e -à Li7Ti5O12 [I.13]
The advantages of LTO are the negligible structural changes induced by lithium intercalation/deintercalation and the solvents stability at its operating potential. However, the lower electronic conductivity and specific capacity than graphite, in addition to its high operating potential, results in lower available energy for the battery. On the contrary, this material is suitable for power and durability applications [START_REF] Takami | High-power and long-life lithium-ion batteries using lithium titanium oxide anode for automotive and stationary power applications[END_REF][START_REF] Wu | Investigation on Li4Ti5O12 batteries developed for hybrid electric vehicle[END_REF][START_REF] Yan | Durability comparison of four different types of high-power batteries in HEV and their degradation mechanism analysis[END_REF] .

The above-mentioned negative electrode materials are schematically summarized in

The positive electrode

Active materials for the positive electrode are commonly made of phosphates, oxides, or complex oxides of transition metals. The transition metals typically display four-or six-fold coordination with oxygen anions, anionic clusters, or ligands; and crystallize into a variety of structural arrangements. Independently of the composition, during discharge, the Li + ions intercalate into the material while an equal number of electrons enter the available d orbitals of the transition metal 11 . Higher potential is expected from compounds with transition metals with more electrons in the d orbital, like nickel or cobalt, because the electrons are more attracted to the core of the atom and more energy is required to extract them 58 .

The electrochemical properties are determined not only by the composition, but also by the structure of the materials, for instance, certain crystallographic orientations facilitate the entry of Li + more than others [START_REF] Nitta | Li-ion battery materials: present and future[END_REF] . For a variety of energy and power applications, along the years, many positive electrode materials with crystal structures of spinel [START_REF] Toprakci | Fabrication and Electrochemical Characteristics of LiFePO4 Powders for Lithium-Ion Batteries[END_REF] , olivine [START_REF] Padhi | Phospho-olivines as positiveelectrode materials for rechargeable lithium batteries[END_REF][START_REF] Lindsay | Buses and batteries: a rising sector[END_REF] and layered 71,

Generalities of layered materials

Layered materials are the most widely used active materials for the positive electrode, because they can meet the high energy density requirements, compared to spinel and olivine structures [START_REF] Myung | Nickel-Rich Layered Cathode Materials for Automotive Lithium-Ion Batteries: Achievements and Perspectives[END_REF][START_REF] Daniel | Cathode Materials Review, 1st International Freiberg Conference on Electrochemical Storage Materials[END_REF] . Layered materials present a crystal structure with R-3m crystal group, where the transition metals occupy the octahedral positions in a face-centered cubic oxygen matrix Besides high energy density, another characteristic of layered materials is their reactivity at high delithiated states accompanied by oxygen production and the loss of the layered structure [START_REF] Macneil | The reaction of charged cathodes with nonaqueous solvents and electrolytes -I. Li0.5CoO2[END_REF] . This issue will be further developed in Section 3.3.1.

Layered LCO, LNO and NCA

The oldest layered material still present in the market is the lithium cobalt oxide (LiCoO2 or LCO). It was first introduced in 1976 by Mizushima et al. [START_REF] Mizushima | LiXCoO2 (O<x<-1): A New Cathode Material For Batteries of High-Energy Density[END_REF] and heavily commercialized by Sony in the early 1990s 11 . It presents high energy density, good capacity retention up to 4.3 V vs. Li/Li + , high electric conductivity and good rate capability. Its theoretical capacity can reach 275 mAh.g -1 when fully delithiated, however, in practice only 0.5-0.7x in Li1-xCoO2 can be extracted, which corresponds to 140-180 mAh.g -1 76 . Due to the insufficient structural properties at high delithiated states [START_REF] Macneil | The reaction of charged cathodes with nonaqueous solvents and electrolytes -I. Li0.5CoO2[END_REF] , the high price of cobalt in comparison to other transition metals [START_REF] Schmuch | Performance and cost of materials for lithium-based rechargeable automotive batteries[END_REF] and its doubtful supply chains [START_REF] Nkulu | Sustainability of artisanal mining of cobalt in DR Congo[END_REF][START_REF] Olivetti | Lithium-Ion Battery Supply Chain Considerations: Analysis of Potential Bottlenecks in Critical Metals[END_REF] , LCO is not preferred for automotive applications, but is widely used in portable devices.

Alternatives to the use of cobalt emerged since the 90s, amongst them, the lithium nickel oxide (LiNiO2 or LNO). LNO presents almost the same theoretical specific capacity than LCO, 275 mAh.g -1 , and is cheaper due to the natural abundance of nickel in the earth crust [START_REF] Schmuch | Performance and cost of materials for lithium-based rechargeable automotive batteries[END_REF] .

Nonetheless, the synthesis of LNO with R-3m structure is rather difficult due to the nonstoichiometric cation ordering induced by the reduction of Ni 3+ to Ni 2+ and consequent migration of Ni 2+ to Li + sites [START_REF] Daniel | Cathode Materials Review, 1st International Freiberg Conference on Electrochemical Storage Materials[END_REF] . Furthermore, during cycling, it presents the same structural instability at high delithiation states [START_REF] Arai | Thermal behavior of Li1-yNiO2 and the decomposition mechanism[END_REF] , phenomenon that is more critical when the nickel content is high [START_REF] Myung | Nickel-Rich Layered Cathode Materials for Automotive Lithium-Ion Batteries: Achievements and Perspectives[END_REF][START_REF] Bianchini | There and Back Again-The Journey of LiNiO2 as a Cathode Active Material[END_REF] . For that reason, research efforts have been made to substitute part of the nickel by other transition metals, including Mg [START_REF] Chowdari | Cathodic behavior of (Co, Ti, Mg)-doped LiNiO2[END_REF][START_REF] Kim | Effect of (Al, Mg) substitution in LiNiO2 electrode for lithium batteries[END_REF] , Ti [START_REF] Kim | The effect of tetravalent titanium substitution in LiNi1-xTixO2 (0.025 <= x <= 0.2) system[END_REF][START_REF] Ha | Effect of titanium substitution in layered LiNiO2 cathode material prepared by molten-salt synthesis[END_REF] , Zn [START_REF] Kim | Effects of Zn or Ti substitution for Ni on the electrochemical properties of LiNiO2[END_REF][START_REF] Kwon | Effects of Zn, Al and Ti substitution on the electrochemical properties of LiNiO2 synthesized by the combustion method[END_REF] , Fe 88 , Co [START_REF] Chen | Effects of cationic substitution on structural defects in layered cathode materials LiNiO2[END_REF][START_REF] Zhecheva | Stabilization of the Layered Crystal Structure of LiNiO2 by Cosubstitution[END_REF] , Mn [START_REF] Guilmard | Effects of manganese substitution for nickel on the structural and electrochemical properties of LiNiO2[END_REF][START_REF] Stoyanova | Effetc of Mn-Substitution for Co on the Crystal.Structure and Acid Delithiation of LiMnyCo1-yO2 Solid-Solutions[END_REF] and Al [START_REF] Madhavi | Effect of aluminium doping on cathodic behaviour of LiNi0.7Co0.3O2[END_REF][START_REF] Lee | Characterization of LiNi0.85Co0.10M0.05O2 (M = Al, Fe) as a cathode material for lithium secondary batteries[END_REF] . Mn, Co and Al substitutions lead to two very stable layered forms: lithium nickel cobalt aluminum oxide (LiNi0.8Co0.15Al0.05O2 or NCA) and lithium nickel manganese cobalt oxide (LiNi1-x-yMnxCoyO2 or NMC).

In NCA, the Al 3+ atoms are electrochemically inactive but provide thermal stability to the compound, while the Co 3+ atoms keep high discharge capacity [START_REF] Zhecheva | Stabilization of the Layered Crystal Structure of LiNiO2 by Cosubstitution[END_REF] . The NCA material possesses high theoretical capacity of 280 mAh.g -1 and in practice, it can reach up to 200 mAh.g -1 with cut-off voltage of 4.5 V vs. Li/Li + 90 . NCA material has found automotive application because it shows better thermal stability at high delithiated states than LCO [START_REF] Lee | Characterization of LiNi0.85Co0.10M0.05O2 (M = Al, Fe) as a cathode material for lithium secondary batteries[END_REF] , however the high cobalt price limited its use, so far, only in the Panasonic-based battery pack that powers Tesla ® models [START_REF]Panasonic and Tesla Reach Agreement to Expand Supply of Automotive-Grade Battery Cells[END_REF] .

Layered NMC family

The NMC materials are a family of lithiated oxides of Ni 3+ , Ni 2+ , Mn 4+ and Co 3+ . These materials are capable to overcome the thermal stability and cation dissolution issues seen in LCO and LNO, and decrease the price by tuning the cobalt content.

In NMC materials the average valence state of the transition metals is 3+. The electrochemically active redox pairs are Ni 2+/3+ , Ni 3+/4+ and Co 3+/4+ , whereas the manganese atoms remain inactive but its presence improves the structural and thermal properties [START_REF] Lu | Layered LiNixCo1-2xMnxO2 cathode materials for lithium-ion batteries[END_REF] . The NMC compounds are named by the ratio of each transition metal, for instance, LiNi0.4Mn0.3Co0.3O2 is named NMC 433. Table I.3 shows a non-exhaustive list of other NMC compounds with their corresponding Ni, Mn and Co fractions. for their facile identification, and mention some of the advantages (in green) and disadvantages (in red) of getting closer to each composition. For EVs applications, it is desired to have as low cobalt as possible due to the high price (34 $/kg vs. 12-13 $/kg for nickel and 2-3 $/kg for manganese [START_REF] Paulsen | Ni-based lithium transition metal oxide[END_REF] ). Manganese is required because it retains the crystal structure and improves the thermal stability by minimizing the c-axis variation, since it is very stable at octahedral coordination in the MO2 framework at the expenses, however, of the available energy because the manganese ions are electrochemically inactive [START_REF] Ohzuku | Layered lithium insertion material of LiNi1/2Mn1/2O2: A possible alternative to LiCoO2 for advanced lithium-ion batteries[END_REF] . Nickel, meanwhile, is preferred in high amount so more energy density and decreased reversible capacity loss are obtained. Nevertheless, when the Ni 3+ content is high, production of stoichiometric Li:Ni materials becomes more difficult (see Section 3.1), the stability in air decreases (see Section 3.2) and the structural stability is compromised [START_REF] Paulsen | Process of making cathode material containing Nibased lithium transition metal oxide[END_REF] . Considering the above, the nickel content is responsible for both the advantages and disadvantages of the NMC materials. In EVs, blends of NMC with spinel LiMn2O4, or layered LCO have already been tested 22,103 .

They combine the high energy density of NMC with the high power of LiMn2O4 and the good capacity retention of LCO.

In conclusion, LCO yields good capacity and have good cycling behavior, however, the toxicity and price of cobalt makes it suitable for portable devices, but rather inconvenient for bigger applications. In an attempt to overcome these disadvantages, atomic substitutions in LCO with nickel and manganese lead to the production of LNO, NCA and NMC materials.

Even if they are still limited by the reactivity at high delithiation states, NCA and NMC materials present various advantages that makes them suitable for EV applications.

Particularly the NMC materials, in which the ratios of transition metals can be tuned to reach desired performances.

3.

Limitations of Ni-rich NMC materials associated with its reactivity In the industry, technical grade Li2CO3 is mostly the first choice because of its low price and its stability toward air atmosphere 106 . An NMC 111, for instance, can be easily prepared under flowing air atmosphere and using Li2CO3 as lithium source 106 . However, the synthesis of Nirich NMC (Nickel > 0.5) is prone to yield non-stoichiometric materials due to the Ni 3+ reduction into Ni 2+ and its migration (cation mixing) into the lithium sites 100 . To counteract this, the temperature is decreased, pure oxygen atmosphere is employed and, very often, excess of the lithium precursor is used 104,107 . The problems is that lithium precursor traces can remain even after the sintering step, which can be very deleterious for the electrochemical performance 106, 108 (detailed in the following section).

Structural degradation from environment exposure

The degradation of NMC materials from environment exposure is rooted in the chemical reactivity toward H2O and CO2, leading to the creation of surface species called soluble surface species 100, 109 (Fig. I.17). The most reported are the soluble bases lithium carbonate (Li2CO3), lithium hydroxide (LiOH) and lithium hydrogenocarbonate (LiHCO3) [START_REF] Paulsen | Ni-based lithium transition metal oxide[END_REF][110][111][112][113][114][115][116][117] . The growth of soluble bases from environment exposure is not exclusive to NMC materials, but it is intrinsic of layered materials and especially aggravated in Ni-rich NMC due to the presence of Ni 3+ 102, 112 . The consequences of the formation of soluble bases are lithium loss and increase of the cell impedance due to the obstruction of lithium migration, which leads to early capacity fading 113,115,118 , which will be addressed in Chapter III. In addition, Li2CO3 can react chemically and electrochemically inside the battery, which causes battery swelling due to gas production (see Chapter IV). The elimination of soluble bases is extremely important for the industry, along the years several actions has been taken, for instance washing 119, 120 and annealing processes prior use 102,116 .

Structural degradation upon battery functioning

During the operation of the battery, the NMC particles in the electrode are subject to degradation processes, that ultimately cause the deterioration of the electrochemical performance 109,121 . The origin is the structural instability and chemical reactivity at high delithiation degrees 121 . It consists in the production of oxygen as a consequence of nickel reduction, followed by cation migration of nickel to lithium sites, accompanied by the collapse of the layered structure to spinel and rock salt phases. In the following sections these steps are thoroughly described. At the same time that these surface processes occur, particle cracking may also take place from the surface to the bulk.

Oxygen production

The structural instability of NMC materials at high delithiation states (above 0. Adapted from Liu et al. [START_REF] Liu | Nickel-Rich Layered Lithium Transition-Metal Oxide for High-Energy Lithium-Ion Batteries[END_REF] .

During the collapse of layered NMC to spinel and to rock salt phases, the bulk structure remains intact due to the limited diffusion length of oxygen anions in the primary particle 122 .

In order to assess the transformation extent, high-resolution transmission electron microscopy The electrolyte degradation related to NMC materials can be classified into:

i. reactions with the surface species (Li2CO3 and LiOH),

ii. reactions with the lattice oxygen and iii. potential-dependent reactions.

Regarding the reactions with the surface species, it has been reported that Li2CO3 may react with the hydrolysis products of LiPF6 132 , with pure LiPF6 118, 133, 134 or with the carbonate solvents assisted by HF 135,136 , to produce a variety of degradation products. Nonetheless, none of these reactions has been proven to occur at the surface of an NMC material. On the other hand, hydroxide compounds have been reported to react with the hydrolysis products of LiPF6 to produce fluorophosphate species, LiF, H2O and CO2 40,41 or directly with linear solvents into alkyl carbonates and ROH 136 .

The degradation reactions induced by the oxidation reaction between active oxygen and the electrolyte components are directly related to the nickel content, as discussed before. The more nickel is contained in the material, the more 1 O2 can be produced. Once the reactive oxygen is produced, it either oxidizes the electrolyte solvents to produce gaseous species (CO2 and CO) and water (see Chapter IV) or two 1 O2 molecules quench together to form two molecules of O2. In addition to reactions with the alkyl carbonate solvents, other reactions between 1 O2 and the electrolyte components have been reported to produce a series of degradation products, such as MFn, LiF, LixPOyFz and Li2CO3 [START_REF] Liu | Nickel-Rich Layered Lithium Transition-Metal Oxide for High-Energy Lithium-Ion Batteries[END_REF]125,137 . The precipitation of these products onto the surface of NMC contributes to the formation of a cathode electrolyte Lastly, there are potential-dependent reactions, for instance, the electrochemical oxidation of surface species and solvents into CO2, CO and radical species. This occurs when the potential is increased beyond 4.2 V vs. Li/Li + , in order to obtain higher capacity from more extracted lithium.

In short, from the electrolyte degradation, three types of products are generated: gases, products dissolved in the electrolyte and products deposited in the SEI or CEI. Among them, gases are always produced from every degradation reaction; in fact, they are the most deleterious products in LIBs for EV applications because they provoke battery swelling and thus, safety and durability issues 23 .

Methods for the detection of degradation products

For gaseous products the simplest method is to measure the expansion of a pouch cell 139,140 . It can be done through measuring the thickness increase or in a more accurate way, by immersion of the pouch cell into a viscous liquid and the use of the Archimedes' principle 141,142 . Whilst the volume measurement indicates the extent of gas production, this method is not capable to obtain information about composition. For identification and quantification, ex-situ gas chromatography coupled with mass spectrometry [143][144][145] (GC/MS) and in-situ on-line electrochemical mass spectrometry 122, 146-148 (OEMS) have been utilized.

For dissolved products, electrospray ionization-high-resolution mass spectrometry (ESI-HRMS), Raman spectroscopy 149 and liquid nuclear magnetic resonance spectroscopy [39][40][41]150 (NMR) are some of the most common techniques. For deposited products, thermal analysis techniques like differential scanning calorimetry (DSC) 36,40 and accelerating rate calorimetry [START_REF] Macneil | The reaction of charged cathodes with nonaqueous solvents and electrolytes -I. Li0.5CoO2[END_REF]151 can be used. To study the interphase composition, Fourier Transform Infrared spectroscopy 152 (FTIR) and X-ray photon spectroscopy (XPS) are commonly employed.

Usually, a combination of many of the aforementioned techniques is necessary to understand and propose a degradation mechanism. In this thesis, special attention is given to gas analysis by GC techniques, although other techniques such as NMR, FTIR and ESI-HRMS are employed.

Surface strategies to overcome the limitations of NMC materials

Many strategies have been proposed to overcome the NMC's reactivity issue. For the industry, the selected strategy relies on the easiness, efficacy and cost. The strategies can be classified in surface or bulk. Bulk strategies such as core-shell or gradient materials are rather complex and more expensive than the traditional synthesis of Ni-rich NMC. Therefore, this thesis focusses only on relevant surface strategies.

Surface coatings

The role of the coatings is to produce a surface passivating film that hinders electrolyte degradation reactions due to HF attack, NMC reactivity or extreme cycling conditions (high potential and low/high temperature), and at the same time, is permeable to lithium ions.

Classic inorganic coatings

Coatings from post-synthesis treatments have proven to be effective to improve the capacity retention, rate capability and reduce particle cracking without sacrificing the initial capacity 131,153 . They are constituted by electrochemically inactive materials and in general their formation is easy and cheap. A vast variety of materials have been tested: oxides [154][155][156] , phosphates 157 and fluorides 158,159 ; to mention some. Among them, oxides have shown promising results to hinder NMC's reactivity. Myung et al. 154 studied the electrochemical performance of NMC 442 coated with several oxides: Al2O3, Nb2O5, Ta2O5, ZrO2 and ZnO, from which Al2O3 (alumina) was found to deliver the best electrochemical performance.

Coating originated from surface species

A relatively simple way to create a coating is to cause the existent soluble bases, Li2CO3 and LiOH from synthesis or ambient exposure to react, to generate a surface film. Jo et al. 160 showed that the soluble bases in NMC 622 react in solution with phosphoric acid (H3PO4) and after solvent evaporation and annealing at 500°C, a lithium phosphate (Li3PO4) layer is produced (Fig. I.24). In addition to reduce the quantity of soluble bases, they claim that the Li3PO4 layer scavenges HF and water.

Electrolyte additives

An easy, relatively cheap and widely used alternative to overcome the reactivity of NMC materials is the incorporation of electrolyte additives 29 . In the case of additives for the positive electrode, the objective is similar to that of the coatings: to form a CEI surface layer that limits the cation dissolution and the surface reactions at the electrode-electrolyte interface, whilst being permeable to lithium ions and electronically insulating 162 . A classical electrolyte additive decompose and precipitate during the formation cycles, forming a CEI of only few nanometers 163 .

Examples of electrolyte additives are compounds that contain nitrile, sulfur or boron groups, and EC-derived. Some are briefly described in Table I. Based on many reports about the synergic effect of mixtures of additives Moreover, since traces of water are unavoidable elements of every electrolyte component, it has been proposed by several authors, that the organo-fluorophosphates products from the reaction between water and the electrolyte components can also be used as electrolyte additives 177,178 . One example is the lithium difluorophosphate (LiPO2F2) 41 , which has been used to improve cycling performance by enhancing the CEI layer 179 and decrease cell's impedance 180,181 , cation dissolution 182 and structural transformations 179 . The advantage is that the CEI is richer in fluorine and phosphorous atoms 177,178 .

Up to date, there is no miracle additive or mixture of additives that remedy the limitations of NMC materials, still, they remain as the easiest way to protect the NMC materials from degradation.

Chapter conclusion and objectives of this thesis

Due to the increasing concern about the greenhouse gases emitted by ICE cars and the shortage of fossil fuels, in the last years, the market of EVs has been gaining terrain. The constant improvement of the battery components has certainly been a key parameter for their success. Nonetheless, there are still big goals to be met in terms of safety, cycle life and driving range. In this regard, the comparison of several positive electrode materials, in conjunction with the enlisted requirements of the automotive industry, showed the relevance of layered materials for automotive applications, particularly, of the NMC materials.

EVs containing NMC materials are already being commercialized. The next milestone to gain in energy density is to use Ni-rich (Ni > 0.5) NMC materials. Nevertheless, numerous reports have shown that the properties of these materials are compromised due to their reactivity toward air exposure and toward the electrolyte in delithiated state. The former leads to the formation of surface species and the latter to oxygen release and structural transformations, particle cracking, cation dissolution; all linked to the presence of nickel.

Inevitably, at the same time, the electrolyte endures degradation into dissolved, deposited and gaseous species that affect directly the performance of the battery. Considering that gases are almost always produced from the degradation processes, it is particularly relevant for this project to study the degradation mechanisms leading to gaseous products and to look for ways to prevent their formation.

Among the proposed solutions to hinder the reactivity of NMC, the literature shows that surface modifications are effective and cheap strategies. Among them, coating and electrolyte additives stand out. Coatings have been created from surface species LiOH and Li2CO3, however, they include a high-temperature reaction with more chemicals that most probably increases the cost and makes the industrial production harder. An easier option is the inorganic coatings made of oxides. Another approach rather practical is the use of electrolyte additives that promote the formation of a CEI. In this regard, so far, there is no single miraculous additive or mixture that solves the totality of the limitations of Ni-rich NMC materials, but instead, the choice depends on the electrolyte composition, the operating conditions and the nickel content.

In view of the reported limitations and solutions to the reactivity of Ni-rich NMC materials, this thesis intends to:

-Contribute to the understanding of the surface reactivity and degradation processes of Nirich NMC materials toward air through a stepwise study of surface species quantification, formation mechanisms and chemical/electrochemical reactivity.

-Understand the reaction mechanisms issued from the reactivity of lithiated or delithiated Ni-rich NMC that lead to gas production upon high-temperature storage.

-Propose solutions to the reactivity of Ni-rich NMC materials in terms of electrolyte additives.

Introduction

Throughout the industrial production of layered materials, precursors residues can be left from the synthesis and constitute what are called the soluble bases: Li2CO3 and LiOH 1, 2 (Eqs. II.1 -3 in Table II.1). In addition, during the handling of the materials, they may be exposed to air, also creating Li2CO3 and LiOH 3,4 , and supposedly LiHCO3 as well 5,6 , through the reaction of H2O and CO2 with the outermost surface of NMC particles. Among layered materials, the promising Ni-rich NMCs (Ni > 0.5 in LiNi1-x-yMnxCoyO2) are known to be highly reactive to H2O and CO2, thus stunting their utilization as positive electrode materials in LIBs for EVs 7 .

As a matter of fact, the presence of OH -groups induces slurry gelation by reacting with the PVDF binder during the preparation of electrodes 8 , the insulative nature of the soluble bases increases cell impedance 4 , capacity fade is often observed with air-exposed materials 5,9 and more gas is produced during battery operation [10][11][12][13] .

The formation of soluble bases from air is not exclusive to NMC materials, but is characteristic of layered materials. Despite having been identified for a long time, the formation mechanisms are still source of debate. Table II.1 summarizes all of the proposed mechanisms.

Goodenough's group [14][15][16] reported, very early, the migration of lithium ions to the surface of layered materials LiVO2 and LCO upon air exposure, to form Li2CO3, but did not show any mechanism. Years later, Matsumoto et al. 17 proposed a redox reaction to explain how layered materials LiMO2 (M = Ni, Co, Mn, Al) react with O2 and CO2 from air to form Li2CO3 (Eq. II.4 in Table II.1). Based on this study, and on XPS and XRD analysis of an as-received and a 1 year air-exposed LNO material, Liu et al. To explain soluble bases formation in NMC 111, Mijung et al. 18 and Huang et al. 19 , proposed that H2O get adsorbed on the particles' surface, followed by the adsorption of CO2 that form slightly acidic surface with H2CO3 (Eq. II.10). Then, Li2CO3 is formed from the reaction of H2CO3 with LiMO2, while the O 2-lattice ions combine with H + to produce LiOH and keep charge balance. LiOH can further react with CO2 to form Li2CO3 (Eqs. II.11 and II.12).

Moreover, works by Schkrob et al. 20 and Faenza et al. 6 , also supported the absence of a redox processes by proposing that in NMC 532 (Ni 3+ 0.2Ni 2+ 0.3) and NCA (Ni 3+ 0.8), LiOH originates first from an ion-exchange process involving partial Li + substitution by H + , when the OH - from water strongly adsorbs via coordination with lithium or with the transition metals (Eq. II.13). After LiOH is formed, it would react with CO2 to form LiHCO3 (Eq. II.14) or Li2CO3.

For years, only these lithium surface species have been considered to originate from the degradation of layered materials after air exposure, even if it was recognized very early, that there might exist other non-identified surface species 5 . In this respect, very recently, new evidence has been published regarding the existence of nickel-based surface species. From the same group, Sicklinger et al. 21 and Jung et al. 22 Acid-base reaction: In the literature, direct identification and quantification techniques of soluble bases in powder materials have been reported (Table II.2). For example, XRD has been used to identify Li2CO3 as a secondary phase when the ratio toward the bulk is high 4,5,18,25,26 , and in few cases, the Rietveld refinement of the XRD data allowed to quantify it. Liu et al. 3 showed that in a 1-year air-exposed LNO material, the amount of Li2CO3 was 13.3 wt.%. Matsumoto et al. 17 quantified 8 wt.% of Li2CO3 in exposed LiNi0.81Co0.16Al0.03O2 and Zhuang et al. 4 6 wt.% in exposed NCA. Conversely, when the ratio of soluble bases toward the bulk mass is small and does not constitute a secondary phase (below ~3 wt.%), the carbonate and hydroxide groups have been identified by Fouriertransform infrared (FTIR) 4,6,25,27,28 , Raman 22,25,29 and XPS spectroscopy 3,9,26,30 . Only Faenza et al. 6 and Zhuang et al. 4 reported the quantification of LiOH.H2O and Li2CO3 by FTIR using a standard.

)# ( + $ ( # → $ ( )# * [II.10] &'"# ( + K = $ ( )# * → &' >-K "# (-K ( ⁄ + K ( &'#$ + K = &' ( )
One less common and indirect technique that has been used to quantify Li2CO3 is the gas chromatography-mass spectrometry (GC/MS). In this respect, Jung et al. 22 quantified the CO2 produced from the chemical reaction of Li2CO3 and HCl (Eq. II.15). Through this method, they reported high experimental error between 12 -29% and ignored that CO2 can remain dissolved in the slurry. Hence, this method is not preferred.

Currently, the most reliable and widespread of these techniques to quantify CO3 2-and OH -, is the Warder's method 31 , also called acid-base titration. This technique thus emerged as the technique of choice for quantifying the soluble bases present in the as-received and exposed NMC materials.

Table II.2. Methods for the identification and quantification of soluble bases.

Identification technique Characteristics

Bulk (> 6 wt.%) XRD Li2CO3 phase identification, if it exceeds 6 wt.% quantification is possible through Rietveld refinement 3,4,17 Surface (< 3 wt.%) FTIR Mostly qualitative of CO3 2-and OH -4, 6, 25, 27, 28 Quantification is possible with internal reference 4,6 Raman Qualitative identification of Li2CO3 and LiOH 22,25,29 XPS Qualitative identification of Li2CO3 and LiOH 3,9,26,30 Indirect GC/MS Quantification through CO2 detection in:

!" # $% & + 2)$* → 2!"$* + $% # + ) # % [II.15]
Warder's titration Identification and quantification of CO3 2-and OH -in a precise way 31 Overall, these studies revealed that determining the precise surface composition in Ni-rich NMC materials is still challenging. As it is a prerequisite to further understand the whole processes occurring upon NMC storage under air and subsequent reactivity upon cycling, we undertook a quantitative study on different Ni-rich NMC materials. Beforehand, a deep characterization of the materials followed by a stepwise optimization work on the entire procedure (solubilization and titration) was required to get high measurement accuracy. Complementary techniques as FTIR, AAS, ICP-AES and TGA-MS were, in turn, indispensable to identify other surface species and to highlight a novel formation mechanism.

Characterization of as-received Ni-rich NMC materials

In addition to perform several analyses by using surface characterization and/or analytical tools, the difficulty of achieving the entire identification and accurate quantification of surface species may also rely on the precise knowledge of the post-synthesis conditioning and storage conditions of the materials as well as their physical-chemical characteristics.

As received Ni-rich NMC materials conditioning

The five studied Ni-rich NMC materials LiNi0.5Mn0.3Co0.2O2 (532), LiNi0.6Mn0.2Co0.2O2 (622), LiNi0.8Mn0.1Co0.1O2 (811), LiNi0.9Co0.1O2 (901), and alumina-coated LiNi0.6Mn0.2Co0.2O2 (622C)

were received in individual tight-closed 1 kg bags originally sealed in a dry atmosphere, from Umicore ® Korea.

Considering the air-sensitivity of Ni-rich NMC materials and the importance to maintain a clean surface state, the bags were transferred to a dry room with dew point between -58 and -63°C (~5-10 ppm of H2O). In there, the powders were reconditioned in PTFE flasks of 25 cm 3 that were filled and tapped to ensure that the least amount of dry air remains inside. Then, the closed flasks were put inside aluminum laminated bags and sealed under vacuum without further treatment.

Throughout the duration of this project, flasks were transferred, one by one, when needed, from the dry room to an argon-filled glovebox (< 0.1 ppm O2 and H2O, MBraun) without air exposure.

Henceforward, the powders coming from the glovebox are called as-received NMC.

As-received Ni-rich NMC materials characterization

The molar composition of each transition metal (Table II.3) of the five NMC materials and of aluminum in NMC 622C were determined at the Technocentre-Renault, after dissolution in aqua regia (protocol in Annex I) and through inductively coupled plasma atomic emission spectroscopy (ICP-AES). The transition metal molar percentages compare very well (≤1.8% of error) with the terminology given by the Umicore. Based on the quantity of aluminum found, the coating of Al2O3 was determined to constitute 18.5 µmol.gNMC622C -1 .

Particle sizes and BET surface areas were obtained from the data sheet supplied with the materials. The particle size distribution (PSD) values, classified in D10, D50 and D90, are equivalent to what has been reported in the literature for Ni-rich NMC materials [32][33][34] .

Table II given by the arrangement of primary particles (≤1µm). Upon increase of the nickel content, the surface of the particles becomes smoother and BET surface areas decrease, which is supposed to be due to the adaptation of the annealing temperature of precursors between 750 -950°C, to limit cation disordering between Ni 2+ and Li + sites and yield highly crystalline materials 35,36 . Noh et al. 35 , for instance, annealed during 15h at 950°C to yield NMC 111 and 532, at 900°C for 622 and at 780°C for 811. The five NMC materials were all issued from the same supplier that controlled the synthesis conditions to the highest standard. They present close structural and global morphological features, except the surface microstructure evolving from granular to relatively smooth with increase of Ni 3+ amount from NMC 532 to 901. These high-grade materials were thus considered as suitable samples for the comparative study of their reactivity toward different atmospheres exposure.

Quantification of surface species in as-received NMC materials

Inherently, as-received NMC materials contain soluble bases that originate either from the synthesis, since excess lithium precursors (LiOH or Li2CO3) are utilized to avoid getting nonstoichiometric NMC after synthesis 1,23 and may be left after reaction; or from air exposure as a consequence of NMC's structural instability 7,39 . The first part of this section is devoted to their quantification through acid-base titration after experimental conditions optimization. In the second part are implemented complementary techniques, with the aim of achieving the complete characterization of the soluble surface species. Based on the likelihood that the accurate quantification of soluble bases is affected by the NMC's instability in water 24,40 and the atmosphere CO2 solubilization 41 , the impact of the agitation time upon NMC washing and the atmosphere composition was evaluated. From the gradual but fast increase of pH when the Li2CO3 is added, it can be deduced that Li2CO3 in both solutions takes around five minutes to be solubilized. Under argon bubbling (blue curve), the pH levels off at a value of 10.7 for at least 40 min. On the other hand, the pH of the solution stirred under air (red curve) reaches only 10.45. This is explained by the simultaneous solubilization of Li2CO3 and CO2 from air which entails a pH decrease with time, according to the following reactions (Eq. II.16 and 17):

) # % + $% # → $% # . ) # % [II.16] $% & #. + $% # . ) # % → 2)$% & .
[II.17]

These experiments prove the necessity of performing the solubilization and titration steps under inert atmosphere, to avoid quantification errors given by CO2 solubilization.

3.1.1.2. Impact of the stirring time upon NMC washing procedure After soluble bases solubilization from NMC surface, the pH of the aqueous solution increases due to an ion-exchange phenomenon that takes place when a lithium ion in the material is substituted by a proton in solution 24,40 (Eq. II.18). Furthermore, this process can result in structural defects initiated by cation mixing 42,43 .

!"/% # + 0) # % → 0!" 1 + 0%) . + !" 2.3 ) 3 /% # [II.18]
The actual effect of stirring time on the pH evolution during the NMC's washing procedure was thus investigated. To this end, 5 g of each as-received material were weighed inside the glove box and taken out without air exposure. They were then transferred to an adapted flask for the washing step and magnetically stirred with 100 ml of bi-distilled water (mass ratio 1:20) for up to minutes for 532, 622 and 622C, and 8 minutes for 811 and 901. In the literature, surface transformations have been detected as a consequence of a prolonged water washing procedure, up to 20 minutes for NMC 811 28 and 1 hour for NMC 532 34 .

In order to ensure that the NMC's texture and bulk structure were not affected by the washing step, SEM and XRD analyses were performed. Through SEM, it was proven that the granular morphology of every NMC material is kept and the distribution of the coating particles (Fig.

Impact of the powder-to-water mass ratio

Driven by the patents of Paulsen et al. 1,2,24 , the NMC powder-to-water mass ratio of 1:20 was previously used for the solubilization procedure. Although the results unveiled complete soluble bases solubilization, this parameter was considered relevant to investigate with the aim of guaranteeing reliable measurements with, potentially, higher degrees of accuracy.

In the literature, Park et al. 44 stated that a mass ratio 1:1 of NMC 8 1.5 0.5/water and a washing time of 10 minutes is not enough to dissolve entirely the soluble bases, therefore proving that a bigger ratio is needed. On their side, Xiong et al. 28 , performed the solubilization step for 20 minutes with a 1:2 and a 1:10 mass ratio, but only showed quantification from the 1:10 mass ratio, probably due to the inconvenience to obtain the filtrate solution from the 1:2 mass ratio solution. In light of these works, a compromise has to be found between the complete solubilization of Li2CO3, a sufficient volume of water to be filtered, not wasting too much NMC powder and enough concentration of soluble bases to be titrated with accuracy.

For the sake of comparison with the previous 1:20 mass ratio (5 g in 100 ml), lower masses of 1, 2, 3 and 4 g of each NMC, corresponding to 1:100, 1:50, 1:33, 1:25 mass ratios, were stirred in an equivalent 100 ml volume of bi-distilled water. Thereafter, the powder and the solution are recovered through vacuum filtering for five minutes maximum. A flux of argon is implemented around the system, to minimize air exposure. The recovered powder is vacuum-dried for 12h at 120°C, in a Büchi oven for future characterizations.

Quantification of soluble bases in as-received materials through acid-base titration

After solubilization, the soluble bases contained in the filtrate solutions obtained from each asreceived NMC material were titrated. As shown in Table II.5, The OH -concentration varies between 40 and 214 µmol.gNMC -1 , increasing with the nickel content from 532 to 901. The CO3 2- quantity does not vary much and reaches a maximal value of 40 µmol.gNMC -1 in every NMC. It is worth noting that there is no significant difference between coated and uncoated 622 materials, which means that this type of alumina coating does not interfere with soluble bases solubilization, most probably because the particles' surface is unevenly covered, as shown by SEM previously.

Table II.5. Acid-base titration results of CO3

Complementary quantification

The acid-base titration technique has proven to be a powerful technique for the accurate quantification of OH -and CO3 2-basic species. AAS and ICP-AES analysis of the filtrate solutions were required as complementary techniques to correlate the amount of CO3 2-and OH - with the Li + amount in the filtrate solution and to unveil the possible presence of other soluble salts. These elemental analyses were carried out on the remaining half of the filtrate solution.

Complementary analysis of filtrate solution through AAS

The AAS analysis of the filtrate solution is based on the energy absorption at a specific wavelength of an atom in the ground state. For each diluted sample, the concentrations were determined through comparison of the amount of absorbed light, to the calibration curves constructed with the standards of known concentration (protocol in Annex I).

Typically, the atomization of a lithium-containing solution into the AAS blue flame turns to yellow. However, upon atomization of the diluted filtrate solution, the flame turned to orange color, indicating that other dissolved metals co-exist with lithium. Indeed, sodium was also Three main observations stand out from the AAS results:

i. The lithium concentration increases with the Ni amount from 532 to 901 materials. These results are in line with those reported by Noh et al. 35 , who quantified by titration 84, 171 and 806 µmol of lithium.gNMC -1 for NMC 532, 622 and 811, respectively; assuming, as above-reported, that all detected hydroxide and carbonate were lithium salts.

ii. The lithium concentration values in each material are higher than those obtained from acid-base titration. This leads to the conclusion that some soluble lithium salts are not titrated due to a weak or a non-acid-base property.

iii. Constant values of sodium are found for each NMC, in concentrations equal to 31 ± 2 µmol.gNMC -1 , meaning that sodium does not depend on the NMC composition, but it remained from the synthesis in the same quantity, in each material.

Even though AAS is a powerful technique for elementary analysis, the temperature of the flame limits the detection of elements other than metals and over-estimation of Na + in the presence of Li + is possible, since both are easily ionizable elements that enrich the flame in electrons.

Complementary analysis of filtrate solution through ICP-AES

The ICP-AES equipment is capable to atomize a wider range of elements than AAS, since it does not require lamps to be changed. Only elements with significant concentrations were compared to calibration curves constructed with the intensities emitted by standards (protocol in Annex I).

For the sake of comparison, that the lithium quantification by AAS was accurate. However, the sodium concentrations obtained by ICP-AES were found to be lower (14 ± 3 µmol.gNMC -1 ) than by AAS (31 ± 2 µmol.gNMC -1 ), confirming an over-estimation in the order of two and a half times. Moreover, the detection of sulfur suggests that some sulfur-based salts remain after synthesis at the surface of the material, in concentrations equal to 40 ± 2 µmol.gNMC -1 . Assuming that the sulfur element detected by ICP-AES originates from sulfate(s) and considering the additional lithium, sodium and sulfur concentrations detected by ICP-AES, it can be inferred that these elements are stemming from Li2SO4 and Na2SO4 compounds. The presence of sulfate species was confirmed in each as-received NMC spectrum through the n S-O band of SO4 2-(orange) centered at ca. 1125 cm -1 , confirmed by the Li2SO4 reference. Over the last two years, You et al. 45 showed a spectrum recorded on a freshly-synthesized LiNi0.939Co0.061O2, in which a non-assigned band emerged in the same wavenumber range; thereafter, Sicklinger et al. 21 attributed a peak in the same position in NCA to SO4 2-species. As expected from acid-base titration results, all NMC materials revealed bands (blue) at 878, 1425 and 1500 cm -1 assigned to the bending mode and the asymmetric and symmetric C-O stretching modes in CO3 [2][3][4]27 . Intriguingly, there is no sign of the n O-H band (green) near 3675 cm -1 of LiOH 6, 21 . This raises questions about its presence in our materials, which will be further discussed in Chapter III.

Conclusion

Thanks to a stepwise optimization of the washing and titration parameters (atmosphere composition, duration of the solubilization step and powder-to-water ratio), the soluble bases contained on the surfaces of NMC materials from synthesis or reactivity toward air atmosphere have been accurately quantified through acid-base titration. As compared to literature, their amounts were found to be relatively low with an OH -concentration between 40 and 214 µmol.gNMC -1 , increasing with the nickel content and a maximal CO3 2-concentration value of 40 µmol.gNMC -1 . The complementary techniques ICP-AES, FTIR and SEM/EDX evidenced the presence of Na2SO4 and Li2SO4 in similar concentrations (40 ± 2 µmol.gNMC -1 ) as leftovers from the synthesis, whatever the material. Only the presence of LiOH was not confirmed through FTIR in the as-received materials, which leads to the hypothesis that it might originate from a hydrolysis reaction upon washing procedure. This assumption will be confirmed later while studying their reactivity in Chapter III. This in-depth characterization of the surface species contained in the as-received materials provide us sound background data that will be meaningful for deciphering the global formation mechanism of components grown upon exposure of NMC materials to a specific H2O/CO2 atmosphere.

Characterization of surface species in NMC materials exposed to specific atmospheres

As reported in Chapter I, since the early 2000's, several investigations have been devoted to defining the impact of air exposure on NMC particles 3, 5, 10-13, 17-19, 29, 30 . However, despite these very interesting works, reaction mechanisms regarding surface species formation still lack of clarification. Therefore, to gain insight into the composition of the surface species grown in air and their formation mechanisms, storage experiments under specific conditions were carried out. • argon, in the glove box (< 0.1 ppm O2 and H2O),

• CO2/Ar, in an 8/92 molar ratio inside a closed desiccator,

• dry air, in the dry room (5 -10 ppm of H2O) and

• air, under a fume hood.

At the end of the storage time, the Petri dishes in the desiccator and dry room were closed and As was foreseeable, the powders stored under argon do not present any soluble bases content change, as observed by other authors before 30,46 and the storage under CO2/Ar and dry air entails very little variation, compared to the as-received materials. As a matter of fact, in these two atmospheres where CO2 is available, no significant additional Li2CO3 is produced. This observation had also been featured by Matsumoto et al. 17 and Faenza et al. 6 . The latter stored NCA powder in a bag containing only dry CO2 gas (<10 ppm), and noticed that Li2CO3 was not produced. On the contrary, when stored under flowing air, Li2CO3 grew fast due to the presence of H2O, which was already evidenced in Williams et al.'s work 47 . A two-step reaction with hydrated lithium hydroxide as intermediary species (Eqs. II.21 and II.22) was proposed:

2!"%) + 2) # % ↔ 2!"%). ) # % [II.21] 2!"%). ) # % + $% # ↔ !" # $% & + 3) # % [II.22]
The NMC materials stored under air reveal a large amount of CO3 2-, a part of which originates from the H2O-assisted consumption of OH -. The other part implies the formation of Li2CO3 through a delithiation process occurring at the NMC's surface, as suggested, very early, by

Goodenough's group [14][15][16] in other type of layered materials. However, as reported in Section 1 by several authors 3,5,18,20 , the degradation mechanism is a matter of controversy.

Therefore, clarifying this point was envisaged through promoting the NMC surface species growth under a controlled atmosphere of H2O/CO2/Ar (hereby simply called H2O/CO2); thus, satisfying the required simultaneous presence of CO2 and H2O. LiOH and Li2CO3, respectively 21 . Note that the amount of soluble bases increases with the Ni 3+ concentration in both as-received and exposed materials, which was also previously observed by You et al. 45 . Li2SO4.1H2O dehydrate around 100-120°C 49 , the water removal is expected to occur after the exposure, upon the drying process at 120°C.

Furthermore, from the acid-base titration results, the reacted lithium quantity was calculated (Eq. II. 23), assuming that all of the carbonate and hydroxide salts are lithium salts. The results are shown in Table II.6. Lithium quantity was found to increase with the amount of nickel in NMC.

For the exposed 532, 622 and 622C materials, the molar percentage of extracted lithium does not exceed 2 mol%. Likewise, exposed NMC 622C and 622 display very close values of 1.72 and 1.77 mol%, respectively, demonstrating that the alumina coating does not prevent soluble bases formation upon H2O/CO2 exposure. More substantially, values as high as 5 and 8 mol% were reached for 811 and 901 materials, respectively, in only six days of exposure. These elevated values were expected as a higher reactivity toward CO2 was reported in literature to depend on Ni 3+ content 26 . For the sake of comparison, from titration analysis, Bichon et al. 34 reported 0.3 mol% of extracted lithium from a 532 material exposed in air for one year and Matsumoto et al. 17 , a value of 8 mol% for a NCA material exposed for 500 hours to air. Owing to the prominent extraction of lattice lithium, these materials are expected to experience cation mixing and structural transformations in the near-surface region 3,5,46 . In the literature, estimates of the layer thickness through high-resolution transmission electron microscopy (HRTEM) have been reported. For the sake of comparison, a layer of ~6 nm thick and 3.5-5 nm thick was depicted at the surface, respectively, of an exposed NMC 622 (28 days in at 55°C and 80% RH) 25 and NMC 811 (30 days in air) 28 . The growth of Li2CO3 with average crystallite size of 50 nm into a layer of around 75 nm thick, was found in an exposed NMC 8-0.5-1.5 (2 years in air) 4 . When compared, for instance, with the layer thickness of the exposed 811, estimated to <150 nm from Fig. II.16l, it can be inferred that the specific atmosphere chosen in this study is very efficient in promoting fast degradation.

RS5T:SU !" = V WD XYZ[\XE .V WD B\]^XCXD_XE V `ab * 100 [II.23]
In order to assess the surface degradation state over an extended time, NMC 622C was stored under similar conditions, during 6, 10, 20, 36 and 120 days and the resulting soluble bases were quantified by acid-base titration. agrees well with the amount of Li2CO3 found by acid-base titration, in the sense that NMC 532 is the material with the lower amount and 901 has the highest, whereas 622C and 622 have pretty close amount. Hence, this FTIR analysis turns out to be a convenient and fast way to evaluate the degradation state of sulfate species-containing NMC surface upon air exposure. Moreover, if the normalized integrated area of the 1300 -1600 cm -1 peaks of NMC 532 is defined as x, the other compositions are 2.4x, 3.1x, 6.2x and 10.5x, in ascendant nickel content, respectively. If the same reasoning is applied to the titration results (x = [CO3 2-]532), we found 2.4x, 2.5x, 5.3x and 7.4x.

Owing to the proximity of the values, we can consider FTIR peak integration, through a sulfate peak normalization, as a semi-quantitative method.

Deciphering the degradation mechanisms through thermal analysis

The soluble bases, LiOH and Li2CO3, decompose respectively above 400°C into Li2O and H2O 51 , and above 700°C into Li2O and CO2 52 (Eqs. II.24 and II.25). Therefore, from the thermogravimetric analysis coupled with mass spectrometry (TGA-MS) of the exposed materials, only CO2 and H2O are potentially detected at those temperatures.

2!"%) cdee°g h⎯⎯⎯j !" # % + ) # % [II.24] !" # $% & ckee°g h⎯⎯⎯j !" # % + $% # [II.25]

Thermal analysis of as-received NMC materials

All the NMC materials were analyzed in a TGA-MS instrument placed inside an argon-filled glove box, the experimental conditions can be found in Annex I. TGA curves do not present noticeable weight losses between 40 -600°C, until the characteristic thermal decomposition of NMC above 600°C, associated with O2 evolution. The figures clearly demonstrate the influence of the increasing Ni 3+ content, as 532 and 901 materials start decomposition at 800 and 625°C, respectively. Since these experiments were performed under argon flow (O2 < 2ppm), the oxygen partial pressure was not sufficiently elevated to stabilize the Ni 3+ oxidation state at these high temperatures, leading to its reduction through the following structural transformation, exemplified for NMC 811 (Eq. II.26):

!"l" e.m /9 e.2 $7 e.2 % # ∆, pq h⎯⎯⎯⎯j 0.3!"l" 2 & ⁄ /9 2 & ⁄ $7 2 & ⁄ % # + 0.7l"% + 0.7 2 ⁄ !" # % + 0.7 4 ⁄ % # [II.26]
Noteworthy that H2O (m/z 18) evolution profile does not depict any noticeable weigh loss around the temperature where the thermal degradation of LiOH should take place, here again suggesting that LiOH is not contained in as-received NMC. On the other hand, the CO2 (m/z 44) profile presents a minimal increase above 700°C resulting from the low amount of Li2CO3. Note that, SO2 could not be detected as Li2SO4 and Na2SO4 decompose above 900°C 53 .

Thermal analysis of exposed NMC materials

The H2O/CO2 exposed materials present two main weight loss regions between 150-450°C and above 650°C ( First, H2O evolves from each material between 150 to 300°C, displaying an almost undetectable to a clear two-step signal from NMC 532 until 901. CO2, on the other hand, is detected between 200-550°C, far below the Li2CO3 decomposition temperature range, and the peak broadens with the Ni 3+ content. Lastly, O2 is also produced between 250-370°C.

In the literature, CO2 produced at low temperature is believed to originate from the decomposition of adsorbed carbonate-containing species. Comforted by the fact that H2O and CO2 peaks evolved simultaneously around 270°C, Liu et al. 5 proposed the decomposition of LiHCO3 into Li2O, H2O and CO2. However, the existence of LiHCO3 as a solid compound can be thermodynamically ruled out. The other reported explanation for H2O and CO2 evolution at low temperature, is the decomposition of cobalt (+II) and manganese (+II) carbonate hydrate (CoCO3.H2O and MnCO3.H2O) and of NiCO3.Ni(OH)2.xH2O, proposed by Gasteiger's group 21,22 . However, they do not provide a mechanism that considers the release of O2 between 250-370°C or that explain the reduction of Co 3+ and Mn 4+ to divalent metals.

So as to ascertain the existence of NiCO3

Proposed degradation mechanism

In light of the TGA-MS results and the quantification of soluble bases in the exposed materials, a novel degradation mechanism is proposed. Upon H2O/CO2 exposure, LiOH turns into Li2CO3, and the production of more Li2CO3 is accompanied by lithium extraction reactions, this process increasing with the Ni 3+ content. Considering that the soluble bases formation takes place in every NMC, regardless of the presence of Ni 3+ to be reduced 3,29,30 , it is plausible to consider that no redox process takes place during exposure and that the metallic elements maintain their oxidation state. As a consequence, the Li2CO3 production can be explained by the transformation of the NMC materials into nickel or cobalt oxyhydroxide (Ni1-2x-yCoy)OOH; NiO and MnO2; in accordance with the proposed following reaction (Eq. II.29):

!"l" 2.#3.y (1zzz) l" 3 (1zz) /9 3 (1z>) $7 y (1zzz) % # { ? | g| ? ⁄ h⎯⎯⎯⎯⎯j 1 2 ⁄ !" # $% & + (l" 2.#3.y $7 y ) (1zzz) %%) + 0l" (1zz) % + 0/9 (1z>) % # [II.29]
As cobalt presents an oxidation state of Co 3+ , it is highly probable that it forms also an oxyhydroxide compound. Following this proposition, NMC 901 is more prone to degrade in H2O/CO2 atmosphere than NMC 532, because it possess higher content of Ni 3+ , whose bonding energy to oxygen in the octahedral MO2 structure is smaller than for Co 3+ or Mn 4+ 56 .

Furthermore, the presence of oxyhydroxides is supported by the fact that NiOOH decomposes between 150-400°C, into NiO along with H2O and O2 release 57 ; precisely in the temperature range where O2 and H2O were detected in the TGA-MS experiments of the exposed materials. In this case, the reaction would be (Eq. II.30a and b):

}l" 2.#3.y $7 y ~%%) 2e.dee°x h⎯⎯⎯⎯⎯⎯j }l" 2.#3.y $7 y ~% + 1 2 ) # % + 1 4 % # [II.30a] }l" 2.#3.y $7 y ~%%) 2e.dee°x h⎯⎯⎯⎯⎯⎯j (1 -20 -u)l"% + u 2 $7 & % d + 1 2 ) # % + + 1 2 + 20 - 1 3 u 2 % # [II.30b]
Nonetheless, the formation of oxyhydroxide compounds does not explain the CO2 production in the 250-450°C range, far from the decomposition temperature of Li2CO3. An alternative explanation is the CO2 production from the consumption of Li2CO3 by a re-lithiation solid-state reaction occurring at low temperature 58,59 . In this regard, a reaction of this type is likely to occur in the exposed materials with the oxyhydroxide decomposition products and Li2CO3. To validate this theory and acquire more insight into the thermal decomposition process, an exposed NMC 811 sample was subjected to water washing procedure, to solubilize all Li2CO3. Afterwards, the sample was analyzed by TGA-MS. 300°C, the weight loss profile is identical for both samples and H2O and O2 gas evolve in a similar way. This result supports the proposed mechanism (Eq. II.29) involving the formation of oxyhydroxides; these compounds present a low solubility in water 60,61 and thus remain at the surface of the particles after the solubilization procedure. Beyond 300°C, the weight loss depicted for the exposed + washed material is less important (of 0.35 wt.%) than for the exposed sample. The CO2 gas is almost undetectable along the complete temperature range, indicating that nearly all Li2CO3 was removed from the surface. Considering these results, we propose the following reaction to explain the CO2 evolution below 500°C from an exposed material; the metal oxides produced by the thermal decomposition of (Ni1-2x-yCoy)OOH (Eq. II.30) reacts in-situ, upon the thermal analysis, with O2 and Li2CO3, to form layered oxides and gaseous CO2, following the chemical reaction exemplified for the mixed oxide (Eq. II.31):

}l" 2.#3.y $7 y ~% + 1 4 ⁄ % # + 1 2!" # $% & ⁄ ∆ → !"}l" 2.#3.y $7 y ~%# + 1 2 ⁄ $% # [II.31]
Remarkably, the shell-like degradation layer in the exposed NMC 811 (Fig by HRTEM 20,25,28 , that were wrongly attributed by Xiong et al. 28 and Chen et al. 25 as the result of processes involving the reduction of Ni 3+ to Ni 2+ . between 120 and 450°C, without CO2 evolution 40 , which supports our hypothesis of the in-situ reaction of Li2CO3 consumption (Eq. II.31). Unfortunately, they not propose any formation mechanism.

Attempts to characterize surface species in exposed NMC through Raman and XRD

According to the assumption that nickel hydroxicarbonate (NiCO3.2Ni(OH)2.xH2O) as a surface specie can be detected on NMC through Raman spectroscopy 22 , supplementary investigation was conducted (protocol in Annex I) to characterize the surface species that grew in the exposed 901 material. Owing to the absence of Ni To summarize, coupled TGA/MS experiments allowed for the proposition of a degradation mechanism that accounts for the formation of Li2CO3 detected through titration, Raman and XRD techniques and the gaseous compounds H2O, CO2 and O2 released upon heating. • Li and Na sulfates have been quantified in a constant amount in all Ni-rich NMC materials, and were not affected by the exposure procedure. These compounds could act as internal standard for analyzing the degradation upon exposure to air through facile FTIR-ATR measurements.

Conclusion of the chapter

• FTIR and TG-MS results suggest LiOH is not present at the surface of as-received NMC materials. It may arise from a hydrolysis reaction with lithium oxide, which will be confirmed in the next chapter.

• The thermal analysis revealed an in-situ re-lithiation process of metal oxide in presence of Li2CO3 to form a layered metal oxide LiMO2, which can explain some positive effects of heating a Ni-rich NMC material on the electrochemical performances, as already reported in literature.

Introduction

Li2CO3, LiOH, Li2SO4 and Na2SO4 soluble surface species were identified and quantified through titration, AAS and ICP after aqueous washing of five as-received Ni-rich NMC materials and it was determined that upon H2O/CO2 exposure, the uppermost surface material reacts into Li2CO3, (Ni,Co)OOH, NiO and MnO2.

It is well-known that the presence of surface species from synthesis or grown from air exposure alter the surface state and the electrochemical performance of layered materials 1,2 .

Reasonably, the most reported are Li2CO3 and LiOH, since, for years, they were the only identified surface species. They have been reported to be electronic insulators, in addition to being capable to react chemically and electrochemically with the battery components [3][4][5][6][7] .

However, a better analysis of these processes would certainly contribute to the understanding and prevention. For this reason, in a first instance, this chapter elucidates the chemical reactivity of relevant surface species in contact with a LiPF6-based electrolyte, as well as an as-received and an exposed NMC 811, focusing on the analysis by liquid NMR of the degradation products in the electrolyte. Then, as Li2CO3 can be electrochemically oxidized and produce gases 8,9 , the electrochemical consumption of surface species in as-received and exposed materials was evaluated. Finally, based on the evidence of the thermal consumption of (Ni,Co)OOH presented in Chapter II, a heat treatment was performed to as-received, washed and exposed NMC 622C and their electrochemical performance was examined.

Reactivity of surface species toward electrolyte

The chemical reactivity of surface species with the LiPF6-based electrolyte was studied through storage experiments at 25 and 55°C. First, the reactivity of reference compounds Li2CO3, LiOH, NiO and Li2SO4 was evaluated. As its presence on NMC surface was suggested by Paulsen et al. 10 , the reactivity of Li2O was also investigated. Then, the chemical reactivity of surface species on as-received and exposed NMC 811 was assessed under similar storage conditions.

Reactivity of reference compounds

Reference materials were added to 1 ml of electrolyte (1M LiPF6 in EC/DEC/EMC 1:1:1 vol.) inside an aluminum laminated bag for four days of storage at 25 or 55°C. The chemical inertness of the inner component (polyethylene) of the bag toward the electrolyte was verified prior to the experiments. After storage, 0.5 ml of the electrolyte were recovered in the glove box by centrifugation and transferred into an NMR tube for liquid 19 F NMR analysis (all experimental conditions can be found in Annex II). The corresponding reaction mechanisms leading to LDFP production with LiPF6 from the electrolyte can be proposed for the three reactants as following (Eqs. III.1-3). The first reaction (Eq. III.1) has been already proposed by Parimalam et al. 12 during the study of the solid electrolyte interphase (SEI) generated at the negative electrode surface, but has never been reported to occur with Li2CO3 contained on the positive electrode.

2!" # $% & + !"ÇÉ Ñ → !"Ç% # É # + 4!"É + 2$% # [III.1] 2!" # % + !"ÇÉ Ñ → !"Ç% # É # + 4!"É [III.2] 2!" # Ö%4 + !"ÇÉ Ñ → !"Ç% # É # + 4!"É + 2Ö%3 [III.3]
The LiPO2F2/LiPF6 molar ratios R calculated from the NMR peaks integration, ranges from almost zero, when the reaction temperature is 25°C, to 0.645 in presence of Li2CO3 at 55°C.

Considering the reaction stoichiometry and the starting molar ratio M/LiPF6 (M=Li2CO3, Li2O, Li2SO4) of 1.35, a maximum value of R = 2.08 can be expected. Note that this value could not be reached at 25°C as the LDFP solubility in carbonate solvents is less than 0.5 M 13 .

The FTIR analysis of the white precipitate recovered after the Li2CO3 storage test at 55°C did not reveal any Li2CO3 but only LDFP and LiPF6, thus, revealing that Eq. III.1 is complete with this reactant at this temperature.

Regarding the storage tests with Li2O, the R value reaches only 0.07 at 55°C. Hence, the Li2CO3 turns out to be much more reactive toward LiPF6 to produce LDFP, than Li2O. Li2SO4 is even less reactive as the test at 25°C revealed no reaction and a low R value of 0.013 is obtained at 55°C. and 55°C (b). It can be observed that LiOH is extremely reactive, even at 25°C. Based on previous studies [14][15][16][17] , most of the peaks can be ascribed to fluorophosphates such as:

• OPF(OH)2 with a doublet centered at δF = -77.3 ppm (J(P,F) = 927 Hz), labelled as 1 in

Reactivity of as-received and exposed NMC materials

In a similar way, reactivity tests were performed with the as-received and exposed NMC 811 materials. In this case, 100 mg of as-received or exposed NMC 811 powder were added to 1 ml of LiPF6-based electrolyte and subjected to the same storage conditions and NMR analysis. From the 19 F RMN analysis, LiPF6, LDFP and LiF were detected in both materials and at both temperatures. However, the absence of fluorophosphates is intriguing, as they were found to be easily produced from hydroxide species. As already mentioned in Chapter II, most probably, Li2O instead of LiOH is present at the NMC surface. Hydroxide ions analyzed by acid-base titration would be the reaction products of Li2O hydrolysis that takes place upon NMC washing procedure (Eq. III.8): It is noteworthy that for the exposed NMC 811 material, the oxyhydroxides species did not generate fluorophosphates when in contact with electrolyte.

!" # % + ) # % → 2!"%) [III.
The LiPO2F2/LiPF6 molar ratios R deduced from the different NMR analyses are reported in Table III.1. For both as-received and exposed materials, the LDFP formation was facilitated by the temperature augmentation, as it was found for the reference compounds. From these R values, the LDFP moles quantity in 1 ml of electrolyte was calculated and compared to the one expected from the complete reaction of Li2CO3, Li2O and M2SO4 (M = Li, Na) with the electrolyte, considering Li2O instead of LiOH from acid-base titration results (Eqs. III.1-3).

As each material contains 40 ± 2 µmol.gNMC -1 of M2SO4, an expected LDFP value of 2 ± 0.1 µmol per 100 mg of NMC can be calculated from sulphate reactants. Note that, in the as-received 811 material, a large quantity of the reactants has reacted to form LiPO2F2, whatever the temperature, which may be explained by the very thin layer. For the exposed NMC 811, higher quantity of LiPO2F2 was expected (20.9 vs. 8.1 µmol in asreceived), but the layer is thicker and consequently, as in case of the reference compound, the transformation process is slowed down; only a R value of 13 is attained at 55°C. The reaction yield may also be affected by the complexity of the exposed surface layer made of oxyhydroxides, oxides and soluble bases, which makes it difficult for the LiPF6 to react with the soluble bases and sulphates.

Furthermore, it is possible that other chemical reactions inside the battery with the electrolyte constituents and at any temperature can take place. For instance, it has been reported that Li2CO3 reacts with HF to produce LiF, H2O and CO2 (Eq. III.9, the production of HF is shown in Eqs. I.10 and 11) 4,5,21,22 . Hence, the content of Li2CO3 associated with moisture, including chemically adsorbed moisture in the NMC powder is highly undesirable.

!" # $% & + 2)É → 2!"É + ) # % + $% # [III.9]
To summarize, a stepwise investigation of the chemical reactivity of the surface species toward LiPF6-based electrolyte was carried out through 19 F NMR measurements, and following the observation that fluorophosphate-type compounds were not produced from the storage of as-received NMC 811 with LiPF6-based electrolyte, it was revealed that Li2O is present in NMC. Furthermore, it was clearly demonstrated that the NMC surface species, Li2CO3, Li2O and Li2SO4, produce LDFP upon storage with the electrolyte, even at room temperature for the first two.

Since it has been reported that LDFP is an efficient electrolyte additive, notably to decrease the impedance of the electrodes and to improve the capacity retention of batteries [23][24][25] , it is suggested that soluble bases and sulfates may benefit the battery's electrochemical performance. However, gas can be formed from Li2CO3, and the surface contains also insoluble bases (Ni,Co)OOH, MnO2 and NiO, which may impact negatively the electrochemical performance.

Reactivity of surface species under cycling

The above-results evidenced that, without specific Li-ion batteries formation conditions, the NMC surface species as Li2CO3, Li2O and Li2SO4 might never fully chemically react with LiPF6 before their first charge. Hence, Li2CO3 would be prone to electrochemical oxidation.

Electrochemical reactivity of powder references

Among the mentioned surface species, Li2CO3 has been reported to electrochemically oxidize in the first cycles 8,26 and also to induce charge transfer resistance increase 3,27 , decrease the rate capability and impact negatively the capacity retention 2-4, 6, 7 . However, there are some discrepancies regarding the oxidation potential of Li2CO3. Meini et al. 28 Given the wide potential range upon which Li2CO3 can oxidize, a decision was made to make our own assessment of its oxidation capability as pure compound before examining the electrochemical behavior of as-received and exposed NMC electrode films in half-cells.

With the aim of exploring the electrochemistry of Li2CO3 as close to the conditions on the surface of NMC as possible, Li2CO3 powder was mechanically ball-milled for 30 minutes (experimental details in Annex II). The particle size of pristine Li2CO3 was effectively reduced As expected, Li2SO4 does not show any additional oxidation peak upon charge as the calculated potential at which the electrochemical oxidation of Li2SO4, to produce lithium persulfate (Li2S2O8, Eq. III.12) would occur is above the working potential window of NMC (above 5 V vs. Li/Li + according to Ghauch et al. 30 ). The solvents oxidation occurs at high potential and would explain the increasing potential up to 4.5 V.

2!"

# Ö% d → !" # Ö # % m + 2S . + 2!" 1 E°>5 V vs. Li/Li + [III.12]
The pristine Li2CO3 incremental capacity curve depicts an oxidation peak starting around 4 V (max. intensity at 4.12 V vs. Li/Li + ) while the milled powder starts oxidizing earlier, around 3.6 V vs. Li/Li + . Because of the poor electrical conductivity of this compound, it remains difficult, even for milled particles and a low rate to efficiently oxidize (less than 1% was calculated in this case). The oxidation proceeds upon the subsequent cycles (Fig.

III.7b)
revealing a smaller oxidation peak. The electrochemical oxidation of Li2CO3 produces CO2 and O2 (Eq. III.10). Recently, it has been proposed through well-founded results, that this oxygen is produced as active oxygen 8 ( 1 O2, Eq. III.11) that rapidly oxidizes the electrolyte components, since O2 has not been identified by classical spectrometric methods 8,9 . Further discussion about the fate of 1 O2 is detailed in Chapter IV.

!" # $% & → 2!" 1 + $% # + 2 # % # + 2S . [III.10] !" # $% & → 2!" 1 + $% # + 2 # % 2 # + 2S . [III.11]
Through the analysis of the milled powder, it is clearly demonstrated that the oxidation potential is size-dependent, ranging within the working potential window of NMC materials.

Electrochemical reactivity of the Ni-rich NMC materials

The electrochemical performance of the five electrode films containing as-received NMC 532, 622, 622C, 811, 901; and exposed 622C was studied in half-cell configuration.

Fabrication of the electrodes

Electrodes of each NMC composition were fabricated in a dry room (5-10 ppm of H2O)

following three main steps (details can be found in Annex II):

i. Slurry preparation with 90 wt.% active material, 6 wt.% conductive carbon black C45

(BET = 45 m 2 .g -1 , 20 nm particle size, Sigma Aldrich) and 4 wt.% of PVDF binder (Sigma Aldrich).

ii. Slurry coating onto an aluminum foil with a target loading of 11-12 mg.cm -2 and overnight drying at 60ºC open in the dry room.

iii. Calendering of the electrodes to fix porosity at 35%, except for NMC 811 and 901, for the reasons mentioned hereafter.

At this point, it is necessary to evoke the increased reactivity of Ni-rich NMC materials toward H2O/CO2 that was studied in Chapter II. The NMC 811 and 901 electrode films were not calendered because the calendering machine is located in a room without controlled humidity.

Theoretical porosity values of 48 and 52% were calculated for 811 and 901, respectively.

In order to assess the extent of the degradation that such films would have endured during the calendering process in air, non-calendered as-received NMC 811 and 901 films were exposed 15 minutes to air, the typical duration of the calendering step, followed by electrochemical tests. no other studies of this type have been carried out.

To conclude, the calendering procedure is meant to fix equal porosity of the electrodes, but doing it in air is very damaging for NMC 811 and 901, since it induces bad capacity retention and reduced specific capacity. Therefore, in the following section, the non-calendered NMC 811 and 901 (48 and 52% porosity) are tested together with calendered 532, 622 and 622C (35% of porosity), in order to evaluate impact of soluble bases on the five as-received NMC. Although we can make suppositions about the higher irreversibility in NMC 811 and 901, at this point we cannot attribute it to with full certainty to any process. In the literature, the irreversible capacity of the first cycle in NMC-containing batteries has been attributed to [32][33][34] :

Galvanostatic profiles of the as-received materials

• Slow lithium kinetics above 0.7 Li inserted in the first discharge proved through GITT measurements of the lithium diffusion coefficient. Increasing the temperature to at least 45 °C decreases this kinetic limitation.

• Absence of calendering that prevents proper ion transport between the particles.

• The surface area of the particles. According to Choi et al. 35 , the irreversible capacity increases with the surface area of NMC 111, following a linear relationship. However, this is not our case, since, among the different compositions, NMC 532 has a higher BET surface area (0.27 m 2 .g -1 ) than NMC 901 (0.20 m 2 .g -1 ) and their irreversible capacities are 12.7 and 17.1%, respectively.

More information about the nature of the structural changes of as-received NMC materials can be obtained from the incremental charge capacity profile (Fig. III.12). After the overpotential observed at the beginning of the first charge, the NMC materials exhibit a reversible redox process starting at 3.75, 3.76, 3.73, 3.71 and 3.70 V vs. Li/Li + (ordered from highest to lowest nickel content) and lasting until 3.77 V. This potential region corresponds to the structural transition from 1 st hexagonal to monoclinic phase (H1àM) upon lithium deintercalation 36 .

Consecutive to the sharp peaks, small ones can be observed and are associated with transitions between states where lithium is ordered and disordered depending on the available octahedral or tetrahedral vacant sites in the crystal structure 31 . Then, above 3.9 V, NMC 811 and 901

present two additional peaks. The first, starting at 3.95 V, corresponds to the monoclinic to 2 nd 123 hexagonal phase transition (MàH2) and the second, starting at 4.12 V, corresponds to the 2 nd hexagonal to 3 rd hexagonal phase transition (H2àH3).

As afore-mentioned, the oxidation of Li2CO3 would represent less than 1% of the total capacity, thus, it could not be visible in the incremental capacity profiles. In conclusion, the oxidation of the Li2CO3 left from synthesis in the as-received materials is not detectable from the electrochemical tests and does not seems to affect the discharge capacity of the materials in the first two cycles.

Galvanostatic profiles of the exposed NMC 622C

In this section the electrochemical evaluation of the surface species contained on exposed NMC 622C is presented. Electrode films were made out of a 20-day exposed NMC 622C powder, containing no LiOH, 356 µmol.gNMC -1 of Li2CO3 (instead of 175 µmol.gNMC -1 for the 6-day exposed) and insoluble surface species ((Ni,Co)OOH, MnO2 and NiO). The electrodes were calendered to 35% of porosity and evaluated in half-cell. Given that the sulfate band is constant at 25°C, it was possible to get a rough evaluation of the percentage of Li2CO3 consumed, through the integration of peak areas. Taking the peak between 1350-1600 cm -1 of the reference material as 100% of Li2CO3, it can be calculated that "622C(exp.)/Li rest in OCV" material consumed ~27% of Li2CO3, which corresponds to 96.1 µmol.gNMC -1 of Li2CO3 chemically consumed (Eq. III.1); and the charged material ~59%, which corresponds to 210 µmol.gNMC -1 of Li2CO3 chemically and electrochemically consumed (Eq. III.1 and III.11). Obviously, it is necessary to consider that the consumption yield in the first charge may change from powder to film.

Through OEMS, Metzger et al. 29 These studies corroborate that Li2CO3 is not entirely consumed during the first cycle, but is likely to be consumed in the consecutive cycles by electrochemical and/or chemical pathways.

Conclusion

The high reactivity of Ni-rich NMC 811 and 901 materials compels to store and handle these powders under moisture-free environment. This was practically ascertained through exposure of electrode films under air for 15 min. As such a short exposure time leads to decreased electrochemical performances, the decision was made not to calender them; the calendering machine being located in the prototyping room under air environment. Hence, this step, though crucial to confer suitable electronical conductivity to films, was performed only with the less reactive 532, 622 and 622C materials, for incoming electrochemical tests.

Through the electrochemical reactivity experiments, it was determined that the decomposition of Li2CO3 during the first cycle of exposed NMC 622C at ambient temperature, is the result of chemical and electrochemical contributions. More specifically, it was found that pure Li2CO3 is decomposed from an onset of 3.6 V vs. Li/Li + , although the potential is strongly influenced by the size of the Li2CO3 particles. Because the layer of Li2CO3 grown after exposure is thin (<150 nm), it can oxidize at any potential within the functioning potential range of NMC materials.

Through FTIR analyses, it was determined that more than half of the Li2CO3 contained in the exposed NMC 622C can be eliminated during the first charge. Its chemical and electrochemical consumption may improve the electronic conduction, but at the same time, produce detrimental gases CO2 and 1 O2. As example, the as-received and 20-day exposed NMC 622C powders contain 37 ± 2 and 356 ± 8 µmol of Li2CO3 per gram., respectively, which would convert, at 20 °C under 1 atm., into 0.89 ± 0.05 mL and 8.6 ± 0.2 mL of CO2, considering CO2 as an ideal gas (24.055 L.mol -1 ).

Unlike the electrochemical oxidation, the chemical reaction of Li2CO3 with LiPF6 is supposed to reduce gas production within the cell as no reactive oxygen species are released, thus avoiding reactions with carbonate solvents. Besides, the sub-product LDFP is likely to improve performances (see Chapter V). Therefore, upon the industrial formation step, leaving the cell after partial charge (to avoid copper oxidation 37 ), under high temperature should be favorable to hinder as much as possible subsequent electrochemical oxidation of Li2CO3.

Modification of the surface species by a heat treatment

Annealing NMC materials has proven to be an efficient way to improve the electrochemical performances of air-exposed NMC materials. It is widely accepted that annealing above 400 °C eliminates LiOH and LiHCO3 2, 7, 26, 38 , and Li2CO3 above 700 ºC 2,39,40 . Because LiHCO3 does not exist, it is unlikely that this is the only explanation to the improved performances. In this sense, during the study of the elimination of soluble bases from synthesis, Park et al. 41 , proposed that the surface species (according to them, Li2CO3 and LiOH) in their as-received NMC 8-0.5-1.5 can be "re-sintered" into a layered phase at 700°C under O2 atmosphere, in a procedure analogous to the synthesis in oxidizing atmosphere 42,43 .

This is done through the reaction of O2, NiO (a rock salt phase that exists from synthesis) and Li2O (stemming from the decomposition of LiOH contained in their materials) (Eq. III.13):

!" # % + 2l"% + 2 # % # → 2!"l"% # [III.13]
However, according to this reaction, O2 would be necessarily required, which is not the case for other authors that have observed that annealing under argon atmosphere also improved the electrochemical performances 7 . In this sense, in Chapter II, it was demonstrated that between 150-400°C under Ar atmosphere, the re-formation of a layered material occurs through the thermal decomposition of MOOH species and, at the same time, the in-situ reaction between the formed (Ni,Co)O, O2 and Li2CO3. This layered phase is thus suspected to improve the electrochemical performance of exposed materials.

In an attempt to get further insight about the impact of the annealing procedure, the following section is devoted to the study of three surface states of the NMC 622C: as-received, washed and 20-day exposed, and their counterparts annealed in argon atmosphere. All of the annealing experiments were performed at 500 ºC, right after the decomposition of MOOH species and reformation of a layered phase (experimental conditions can be found in Annex II). The electrochemical performances of all these powders have been assessed in half cell after the elaboration of films, following the same procedure of Section 3.2.1.

4.1.

As-received, washed, exposed and their analogs annealed NMC 622C films

Loadings have been estimated from the weighing of several electrodes at different places in the film. almost the same loadings between 11 and 12.5 mg.cm -2 whereas both exposed NMC films present higher values of 14.2 and 14.3 mg.cm -2 . This seems to be related to the viscosity of the slurries, therefore, we undertook preliminary tests on as-received and exposed materials.

According to the literature, during the electrode films fabrication with PVDF binder, the viscosity of the slurry can increase owing to the presence of LiOH 44,45 . This behavior, known as slurry gelation, is due to the double bond formation after HF release (Eq. III. slurries behave similarly at high shear rates, however, at low shear rates, close to the stationary state, the differences are more pronounced. Our hypothesis is that the surface deterioration caused by H2O/CO2 exposure, influences the interaction between these particles and the polymer-carbon matrix during slurry preparation. These differences could not be explained by the afore-mentioned slurry gelation process as it would bring about a decrease of the viscosity at any shear rate.

Moreover, the shear rate value of 34 s -1 (calculation in Annex II) that the slurry endures when cast onto the aluminum foil can be used as a reference point to compare both slurries property more accurately. At this shear rate, viscosity values of 6.4 and 6.0 Pa.s were found for asreceived and exposed NMC 811 slurries, respectively. From this little difference, it can be ascertained that slurry gelation did not take place in any of them. According to Paulsen et al. 44 , above around 46 µmol.gNMC -1 of LiOH, the slurry is "unstable", which means that its viscosity In NMC 622C, LiOH detected in as-received powder was suggested to be originated from Li2O and the exposed powder does not contain any, which suggest similar viscosity results would be obtained. Therefore, loading differences do not seem to be caused by the viscosity change. Further in-depth investigations are required for the complete understanding, but are beyond the scope of this study. According to the assumption that Li2CO3 can react in situ with O2, from MOOH decomposition (Eq. II.31), no variation of the content after annealing supposes that the soluble bases are mainly originating from synthesis. This result was expected as all necessary precautions were taken to keep the as-received NMC materials under moisture-free environment. On the other hand, Park et al. 22 showed that the quantity of Li2CO3 in their asreceived NMC 80.5-1.5 decreased from 99 to 62 µmol.gNMC -1 when heated at 720°C under O2 atmosphere, the temperature at which Li2CO3 starts melting and thus could easily react.

The electrochemical performance of both films presents no conclusive change over 60 cycles, apart from the small difference in the initial discharge capacity (174 mAh.g -1 , compared to 172 mAh.g -1 of the as-received material) and a lower capacity retention after 40 cycles (95.4 %, compared to 98.5 % for the as-received material after 60 cycles). Larger-scale tests could confirm a real impact of the annealing process on as-received NMC.

Impact of heat treatment on washed NMC 622C

This section aims at defining the influence of an aqueous washing procedure on the electrochemical performances of NMC 622C. The washing procedure is similar to that optimized for acid-base titration analysis described in Chapter II (15 minutes of stirring time in argon atmosphere). It has been reported that poor performances are achieved from washed NMC 811 material 46 . As a matter of fact, SBs are not necessarily considered as impurities but rather as "surface properties". Hence, in Paulsen's patent 44 , it is proposed a threshold value of soluble base content, found experimentally, referred to as "equilibrium value". For instance, equilibrium values given for commercial NMC 111, 532 and 622 with about the same morphology (spherical dense particles with D50 of about 8-10 µm, with crystallite size of 500 nm) increases from 15-20 to 50-70 to 100-150 µmol.gNMC -1 , respectively. Later, Park et al. 22 determined that leaving 163 µmol.gNMC -1 of soluble bases improved the capacity retention. It has to be emphasized that these contents were obtained by acid-base titration without mentioning the origin of LiOH as potentially coming from Li2O.

Interestingly, the as-received NMC 622C values (145 and 166 µmol.gNMC -1 for R and R-A samples) are in the given equilibrium content range. As expected, washed samples depict lower values (31 and 85 µmol.gNMC -1 for W and W-A samples, respectively) and exposed samples depict higher values (356 and 342 µmol.gNMC -1 for E and E-A samples, respectively).

With the equilibrium value as reference, it is thus predicted that either washed or exposed materials will not deliver good performances as compared with as-received ones.

Noteworthy, the washing procedure decreased almost to zero the quantity of CO3 2-, and decreased the OH -by five. These values are not zero due to two possible reasons, (i) the presence of CO2 and H2O traces that may exist in the argon (99.999% purity); and (ii) a brief exposure to air during the transfer to and from the furnace used for the annealing treatment; in addition to the severe sensibility of washed materials toward air 22 .

The capacity retention obtained for the W and W-A NMC 622C are shown in Fig. III.17. For comparison, the as-received material is also displayed. The almost entire elimination of the soluble bases in W material did not improve the capacity retention, but slightly improved the initial discharge capacity from 172 to 176 mAh.g -1 . Unfortunately, the gain is rapidly lost, since at 20 cycles the capacity reaches the same value as the as-received material and continues to decrease until 157 mAh.g -1 at 60 cycles (red curve), compared to 170 mAh.g -1 from the as-received material (grey curve).

The W-A material presents 17 and 68 µmol/gAM powder of CO3 2-and OH -, respectively. The reason for the increase of SB content compared to the W material (5 and 26 µmol/gAM powder, respectively) is believed to be the transfer to and from the furnace used for the annealing treatment. Concerning the electrochemical effect, annealing the washed material makes it perform slightly better but still lower than the as-received material. In spite of the initial lower discharge capacity (172 compared to 177 mAh.g -1 ), after 60 cycles the washed + annealed retained 91 % of the original capacity, compared to 88 % of the washed material.

Up to this point, neither the washing nor the annealing procedures seem to improve the capacity retention of as-received NMC 622C. While in literature, surface species are considered in general terms as "surface property", it is reasonable to assume, in light of earlier results, that the capacity fade could be partly due to the absence of enough surface species (Li2CO3, Li2O, Li2SO4) to react with LiPF6 into LDFP and LiF acting, respectively, as electrolyte additive and passivating layer. Note that the role of the coating on the particles is not well defined; it has been proved experimentally that pure Al2O3 is prone to react with LiPF6 in a similar way to produce AlF3 and LiPO2F2 (Eq. III.15) 47 , but may also hydrate at room temperature to give Al(OH)3 48 . (-16 %), compared to the as-received material of 172 mAh.g -1 . From there, the regular increase of capacity seen during the first 20 cycles can be explained by the gradual chemical and electrochemical consumption of surface species. Given that the capacity was neither recuperated nor improved after 60 cycles, after Li2CO3 has been chemically or electrochemically consumed, it is reasonable to assume that the oxides and oxyhydroxides are responsible for the ceaseless low capacity. Thus, the bad electrochemical performance cannot be only attributed to the Li2CO3 presence, as has been proposed in the literature.

2Ü* # % & + 3!"ÇÉ Ñ → 4Ü*É & + 3!"Ç% # É # [III.
The surface layer in the exposed NMC 622C was expected to heavily evolve after annealing.

Indeed, the content of CO3 2-decreased by 160 µmol.gNMC -1 after annealing. Since this temperature is too low to decompose thermally Li2CO3, it underwent the in-situ reaction with O2 and Ni and Co oxides products from the decomposition of MOOH, according to previous TGA-MS findings. Since from 1 mole of Li2CO3, 2 moles of Li(Ni1-2x-yCoy)O2 are produced (Eq. II.31), the quantity of Li(Ni1-2x-yCoy)O2 formed would be 320 µmol.gNMC -1 . The presence of this newly formed layered phase along with the disappearance of part of Li2CO3 and metal oxides/hydroxides upon heating should lead to improved electrochemical performances. On the other hand, similarly as the other annealed materials, a relatively high amount of OH -(146 µmol.gNMC -1 ) was detected after annealing, which could not be explained only by the presence of impurities in argon, but mostly due to rapid degradation during the transfers.

Regarding the electrochemical performance of the corresponding films, initial discharge capacity is around 164 mAh.g -1 (+12 % compared to the exposed) which accounts for 95 % of the as-received material discharge capacity of 172 mAh.g -1 . Therefore, the annealing process at 500°C under argon allowed H2O/CO2 exposed NMC 622C materials to recover part of their capacity due to layered oxides reformation and decomposition products vanishing.

Nevertheless, the poor retention capacity of the exposed material is not improved after annealing, which may limit the applicability of this process as a long-term solution.

Correspondingly, Oh et al. 49 reported that annealing a 3-month exposed NMC 7-1.5-1.

Conclusion

With the objective to evaluate the role of surface species on cycling, as-received, washed and H2O/CO2 exposed NMC 622C were studied with and without annealing treatment at 500°C under flowing argon. The electrochemical performances of the electrode films made of nonannealed powders were predicted by comparing the values of soluble base content with the equilibrium value given in literature. It is suggested that these surface species could benefit the performance through reaction with acidic species (HF...) or LiPF6 in the electrolyte.

In addition, it was found that a too large amount of electronically insulating surface species formed upon excessive H2O/CO2 exposure leads to a large polarization and lower capacities.

In this case, the annealing treatment enables the recovering of the original polarization and a part of the capacity due to the decomposition of oxyhydroxides MOOH along with the in-situ reaction of Li2CO3 with O2 and metal oxides to reform active layered oxides (Eqs. II.29 -31 of Chapter II). Overall, this section helped clarify the relation between the decomposition of surface species and the electrochemical performance of Ni-rich NMC materials.

Conclusion of the chapter

The chapter was devoted to getting further insight into the fate of surface species in Ni-rich NMC materials when used as Li-ion battery positive active material, as well as their influence on the electrochemical performances. Therefore, a stepwise investigation of their chemical reactivity toward a classical electrolyte was first undertaken. It was proven that the soluble bases Li2CO3, Li2O, Li2SO4 react with LiPF6, especially at 55°C, to produce LDFP, LiF. On the other hand, the high reactivity of LiOH with carbonate solvents produces fluorophosphates along with LiF. This allowed to ascertain, through the absence of fluorophosphates and LiF in 19 F NMR spectra, that LiOH detected from acid-base titration in as-received NMC powders, was coming from Li2O. This result can be of paramount importance when considering the responsibility of LiOH in the PVDF-based slurry gelation process.

If the totality of Li2CO3 is not chemically consumed in contact with LiPF6, it may oxidize upon charge at a potential value that is strongly influenced by the particle size of Li2CO3. As a sacrificial salt, it leads to the production of additional Li + but also CO2 and 1 O2 species. Given the high reactivity of the latter toward carbonate solvents, producing further gas, it is assumed that the chemical consumption of Li2CO3 with LiPF6 is preferable over electrochemical oxidation. Moreover, it favors the production of LDFP, an electrolyte additive known to improve electrochemical performances and LiF, a stable passivating compound.

The removal of the surface species through a washing procedure was demonstrated to be detrimental to the capacity retention which steadily decreases from the first cycles whereas forming soluble bases in a too large amount through NMC exposure under H2O/CO2 atmosphere leads to polarization enlargement and limited delivered capacities. These results seem to corroborate the existence of an equilibrium value corresponding to an ideal amount of surface species that acts positively on the electrochemical stability of NMC materials. The annealing treatment at 500°C under argon of exposed materials helped to recover some capacity and reduce polarization owing to the decomposition of MOOH species that made it possible to reform electrochemically-active layered oxides.

Chapter IV

Gas evolution upon storage of batteries containing

Gas evolution associated with the storage of

Introduction

Implementing specific experiments to study the gas production in cells is a of particular interest for the battery community due to the associated safety risks and cycle life deterioration 1,2 . This was proved to be especially problematic when Ni-rich NMC materials are used as positive electrode. Indeed, in addition to the gas produced from reduction of carbonate solvents at the surface of the negative electrode to create the solid electrolyte interphase (SEI), gaseous species 3,4 may originate from chemical and electrochemical reactions involving NMC surface species and from materials processes that bring about electrolyte degradation. This section recapitulates and briefly discusses known mechanisms of gas production from each electrode and the main techniques to characterize them.

Gas production mechanisms

Gas associated with the positive electrode

The mechanisms of gas production associated with the positive electrode can be classified into:

i. Electrochemical oxidation of alkyl carbonate solvents

ii. Chemical oxidation of alkyl carbonate solvents by active oxygen ( 1 O2)

iii. Chemical and electrochemical reactions with surface species

i. Electrochemical oxidation of alkyl carbonate solvents

In commercial batteries containing Ni-rich NMC materials, no more than 0.6 of lithium is extracted at 4.2 V (referenced vs. Li/Li + throughout this work). To extract the totality of lithium, higher cut-off voltage is needed, which unfortunately provokes the electrochemical oxidation of alkyl carbonate solvents. The alkyl carbonate solvent that is more prone to oxidize electrochemically is EC, due to its tendency to solvate PF6 -most strongly and thus to reach the positive electrode more easily than linear solvents during charge 5 . Based on density functional theory calculations, it is plausible that there is electron withdrawing from EC by PF5, which weakens the C-O ether bonds of EC and consequently lowers the ring-opening activation energy (Eq. I.10) 6 . Then, a proton is extracted from EC• + to produce vinyl alcoholate radical, H + and CO2 (Eq. IV.1). Experimentally, through online electrochemical mass spectrometry (OEMS) analyses, Solchenbach et al. 7 proposed that the EC oxidation starts l" d1 + S . → l" &1

[IV.2]

% #. → 2 # á à â + 2S . [IV.3]
The 1 O2 is considered to be "active" as EC oxidation with ground state O2 is highly unlikely at ambient temperature 9 . Jung et al. 9 proposed an 1 O2 nucleophile attack on EC to produce CO2, CO and H2O (Eq. IV.4). Furthermore, they found a relationship between nickel content and the onset potential of gas production. In cells containing NMC 622, it starts at ∼4.7 V and at ∼4.3 V for NMC 811. For the authors, the explanation is that at equal cut-off voltage, NMC 811 reaches higher SOC, producing more Ni 4+ whose reduction releases more 1 O2.

[IV.4]

Besides gas production, another consequence of 1 O2 production is the depletion of oxygen, that causes surface phase transformations from layered to spinel and rock salt 10 (Eq. IV.5

exemplary shown for Li0.5MO2 (M = Ni, Mn, Co)). As a result, the cell impedance increases due to the reduction of ionic transport and the loss of active material 11 .

[IV.5]

iii. Chemical and electrochemical reactions with surface species

The storage of Ni-rich NMC with electrolyte, especially at high temperature, chemically produces CO2 from Li2CO3 (Eq. IV.6 also seen as Eq. III.1) and CO2 and 1 O2 from the electrochemical oxidation of Li2CO3 (Eq. IV.7 also seen as Eq. III.11).

2!"

# $% & + !"ÇÉ Ñ → !"Ç% # É # + 4!"É + 2äá à [IV.6] !" # $% & → 2!" 1 + äá à + 2 # á â à + 2S . [IV.7]
The latter reaction was demonstrated very recently, through isotope labelling of oxygen 12 or of carbon 13 in Li2CO3, to be the responsible for most of the CO2 production in the first charge.

Additionally, evidence of the production of 1 O2 was given by Mahne et al. 14 , through the use of the chemical probe 9,10-dimethylanthracene (DMA) that promptly forms DMA-O2 in the presence of 1 O2, which would not be possible with O2.

Note that CO2 is produced from each of these reactions. Therefore,

Gas associated with the negative electrode

During the formation procedure, carbon monoxide (CO), methane (CH4), ethylene (C2H4) and ethane (C2H6) are produced from one-or two-electron reduction reactions of linear (Eqs. IV.8

-11) and cyclic carbonates (Eqs. IV.12 -14) 1,3,15,16 . These gases are called formation gases and their concentration depends mainly on electrolyte composition and cycling conditions 16,17 .

In commercial LIBs, the gaseous products are degassed after the formation cycles. Another known parasitic reaction is the electrochemical reduction of H2O, which produces OH -and hydrogen (H2) (Eq. IV.15) 18,19 .

) # % 1ã ] h⎯⎯⎯j %) . + 2 # ) # [IV.15]
Lastly, it is also possible that gases diffuse from one electrode to the other and perform "crosstalk" reactions. A particular example is the chemical reduction of CO2 on a lithiated negative electrode that forms lithium oxalate (Eq. IV.16) 20,21 . This process evidences the complexity of the gassing reactions inside LIBs.

2$% # + 2!" → !" # $ # % d [IV.16]

Gas characterization techniques

In-situ and ex-situ methods are used to characterize emitted gases from LIBs. In-situ examples are the measure of pressure 22 and volume 23 evolution, and OEMS analysis 15,24 . Although the first two indicate the extent of swelling, they do not inform about composition. On the contrary, OEMS is able to characterize gas through the use of a mass spectrometer. The OEMS disadvantage is that the constant pumping of gases from the adapted cell to the spectrometer may change the electrolyte composition. Moreover, the cell configuration may be different from the commercial one and may induce different behaviors.

Concerning ex-situ techniques, gas chromatography (GC) is the main used (details about the technique can be found in Annex III). The difference with OEMS is that it offers punctual analysis, often post-mortem, and employs a chromatographic column to separate the gases.

Molecular identification and quantification are possible through the use of a variety of chromatographic columns and calibration curve. Among the different options of detectors, MS is the most used, but there are also the barrier ionization discharge (BID), flame ionization (FID) and thermal conductivity (TCD) detectors.

Conclusion

Table IV.1 summarizes the reaction mechanisms from each electrode described in this section. The production of CO2 is in common for almost all of the reactions associated with the positive electrode and most of these processes involve the reactivity of Ni-rich NMC at the delithiated state and at high temperature 8 . This is not the case of the negative electrode, in which mostly hydrocarbons, CO and H2 are produced. Therefore, in this chapter, we considered the positive electrode to be the main source of CO2 and intended to understand the role of Ni-rich NMC materials in the gas production upon storage at a charged state and high temperature (55°C), with a preliminary work on the optimization of the gas quantification. Discussion and comparisons to the literature of gas production are included along the chapter.

Optimization of the instrumentation for gas quantification

Comparison of the available gas chromatography instruments

A coupled GC/MS-FTIR (described in Reference 25 , Nicolet 6700) and a GC/BID (Tracera GC/BID-2010 Plus, Shimadzu) are available at the analytical platform (PFA) of Université de Picardie Jules Verne. Some of their characteristics are summarized in Table IV

Validity of calibration curves in time

Calibration curves do not last long Excellent Excellent

Molecular characterization

With database of ionized fragments

With FTIR database No

Flow control mode of GC Pressure Pressure or linear velocity MS and FTIR are both capable to characterize gaseous species, either through the comparison with a database of ionized fragments or with the NIST database of IR spectra, respectively. However, the ionization source in the MS evolves with time, making the calibration curves quickly unusable, and the accuracy of the quantification by FTIR is compromised by the low peak resolution. On the other hand, the BID detector uses a plasma to ionize the gaseous species, which preserves the validity of calibration curves, but it does not identify the detected molecules. To overcome this problem, in our case, the same column is installed in both devices.

In a GC device, the flow control mode is intended to direct the flow from the injector to the column and finally to the detector. The coupled system, GC/MS-FTIR, has only the option of constant pressure mode (measured in Pa), which means that the pressure of the injection port is fixed and, as the sample travels away into the column, it experiences less pressure. On the other hand, the GC/BID works in constant pressure mode or in constant linear velocity mode (measured in cm.s -1 ). The latter corrects the pressure along the acquisition in order to maintain constant the volume that travels a certain distance down the column.

Choice of the column

The variety of chromatographic columns suitable for gas analysis are classified in capillary and micropacked types ( The most frequently used are the capillary columns because they are capable to separate most of the permanent gases (gases at room temperature) 26,27 . Nonetheless, the classic HP-PLOT Q capillary column (used in our study) presents the disadvantage that the gases from air (N2, O2, CO), H2 and CH4 reach the detector at very similar retention times and thus appear as a single merged peak 25 . On the other hand, micropacked columns are rarely used, but some of them seem to separate hydrocarbons and inorganic gases including hydrogen 28 .

Considering this information, a HP-PLOT Q capillary column connected to a postcapillary column Rtx-1, and a ShinCarbon ST micropacked column were tested in this work with the main objective to improve the separation and quantification of gases within a single acquisition. This was done through the improvement of the GC experimental settings.

Use of a capillary column with the GC/MS-IR equipment

A mixture of 200 µl of gas from a cycled battery (whose cycling conditions are not relevant at this stage) were injected, plus 500 µl of argon were injected with a gas-tight syringe (Hamilton with PTFE plunger tip and 1 ml of volume), into the GC/MS-IR equipped with the capillary column HP-PLOT Q. The injector temperature was set at 100 ºC, helium was used as carrier gas at a flow rate of 1.3 ml.min -1 and a split value of 1/10 was used (the analyte was diluted 10 times). The oven program detailed in Annex III was set to obtain the best separation in the shortest amount of time. EMC, appear at retention times higher than 30 min. Given the limitations of the number of FTIR spectra acquisitions per minute, the peaks in a Gram-Schmidt are constructed from fewer points, which gives less precision to peak area calculation.

For the same acquisition, an MS chromatogram is obtained as well (Fig. IV.3c). The first gases are detected ten minutes later than in the Gram-Schmidt, this corresponds to the time required for the gases to flow through the coupling system. The first detectable peak corresponds to the merged peak of argon, air (N2, O2), CO and CH4; which appears very visible due to the considerable amount of argon (500 µl) that was added to the battery to help gas extraction. Immediately after, CO2, C2H4 and C2H6 appear, followed later by the alkyl carbonate solvents. As above-mentioned, main limitation of the GC/MS equipment is the gradual ageing of the ionization source that renders the calibration curves unusable after only one day. Daily construction of calibration curves is time consuming, thus this option was ruled out.

PF5, expected gaseous product from the decomposition of LiPF6 (Eq. I.10), is not detected either because it reacts rapidly with the stationary phase of the column or it reacts with water to form gaseous POF3 (Eq. I.11). In case of POF3, it is not detected because the fluorine is easily substituted and the resulting compound stay dissolved in the electrolyte To conclude, the co-elution of H2, O2, CO and N2 with Ar at early retention times seems to be an unavoidable characteristic of GC/MS when working with a HP-PLOT capillary column under these conditions. After deconvolution, the Gram-Schmidt is an adequate option for gases identification and quantification, although it can exist a weak error on the peak area.

Alternatively, to improve separation, a more retentive column could be used or the experimental conditions could be modified.

Use of a capillary column in the GC/BID equipment

GC/BID is a valuable alternative to GC/MS-IR, to overcome the problem of the short validity of calibration curves, since the plasma ionization is constant in time 29 .

First, the same HP-PLOT Q capillary column and experimental conditions were used. A mixture of 200 µl of the gas sample and 500 µl of helium were injected. Favorably, this flow mode decreased the broadening of the peaks and thus separated the CH4 peak. This is because the constant linear velocity mode compensates for the flow changes that can broaden a peak, caused by the differences in temperature between the injector and the column.

In the course of the thesis, this method improvement was achieved after that several meaningful results were obtained with the previous method. Therefore, in the remainder of this chapter, it is indicated if the "regular" (constant pressure) or the "improved" (constant linear velocity) method is used. For the quantification results, all of the calibration curves from both methods used in the GC/BID equipment are shown in Annex III.

Use of a micropacked column with the GC/BID equipment

The micropacked column was used with the aim of improving gas separation. According to the application note of the manufacturer 28 , the ShinCarbon ST micropacked column is capable to separate the totality of permanent gases by using a specific oven temperature program (Annex III) and the pressure program starting at 250 kPa and increasing at 15 kPa.min -1 until 400 kPa.

From the injection of gases from a pouch cell under these conditions to the GC/BID, no peaks were detected. We suspected that the initial pressure of 250 kPa was not high enough to push the gases through the dense stationary phase. Therefore, the pressure was set at 400 kPa throughout the acquisition. This pressure increase made the gases detectable ( Only within three minutes, the peaks of H2, the air peak (O2, N2), CO and CH4 appear perfectly separated from each other. Each peak was ascribed thanks to the injection of gaseous standards. As it can be seen on Fig. IV.6, a gaseous compound was eluted at the same time than C2H4, so its corresponding peak is less well defined than with the previous mode.

Although the GC equipment can reach the pressure of 400 kPa, we observed that the flow of carrier gas to maintain it through this column of ca. 1000 ml.min -1 is unexpectedly high, as compared to the 70 ml.min -1 value announced by the manufacturer and to the 1.3 ml.min -1 value in case of the GC/MS-IR device. Owing to unusual pressure variations, a good reproducibility cannot be obtained (an example is shown in Annex III).

In an attempt to decrease helium consumption and improve reproducibility, the split ratio was changed from 1/10, to splitless. The helium consumption was indeed reduced; however, a massive peak that impacts the detection of CO and CH4 appears at very early retention time (Fig. IV.7). This is an overloading peak characteristic of unadapted analyte concentration 29 .

The conclusion was drawn from this result that helium saving in splitless mode is not an adequate solution. Despite the fact that the separation of permanent gases was successful, unfortunately, the pressure of 400 kPa was too high for sustainable and reproducible measurements. Therefore, we conclude that the implementation of the micropacked column was not achievable at this point.

Conclusion

Following a brief comparison of the characteristics of both GC techniques available at the PFA, it was decided to use the GC/MS to identify the gaseous compounds and then the GC/BID to quantify them, both using the same capillary column. It was found that some of the permanent gases co-elute at early retention times when using the constant pressure flow mode in both GC/MS-IR and GC/BID. To solve this, first, argon was replaced by helium when extracting the gas from the pouch cell, but unfortunately this did not eliminate the co-elution problem. Then, it was discovered that by changing the flow mode to constant linear velocity in the GC/BID device, it was possible to successfully separate the CH4 gas from the co-eluted peak.

With a view to detect all the permanent gases with a single equipment and a single acquisition, a micropacked column was installed in the GC/BID equipment. This column presented the main advantage of H2, CH4 and CO separation, but the high helium flow needed to maintain the pressure at 400 kPa did not allow the acquisition of reproducible measurements. Profound optimization was left out of the scope of this thesis, although these experiments proved the big potential of this column to be used in future tests.

In the rest of the chapter, only results obtained from the capillary column with "regular" or "improved" method will be presented.

Gas evolution associated with the storage of Ni-rich NMC materials

This section focuses on studying the gases produced from the storage i) of the electrolyte in contact with NMC powder or film, with the objective of identifying the thermal degradation that the electrolyte endures in the presence of this material, then, ii) of NMC/Graphite(Gr) cells, by comparing two NMC compositions and modifying the concentration of surface species on NMC 811.

Gas production from the storage of NMC with electrolyte

This section investigates the thermal electrolyte degradation in the presence of the five Ni-rich NMC powders and of NMC 622C films at different states of charge (SOC).

Powder storage

The storage experiments of NMC powder + electrolyte were performed in 18650 stainless steel cans. Inside the argon-filled glovebox, 200 mg of NMC 622C and 0.2 ml of electrolyte were put inside the 18650 can and closed with a septum. The can was placed in an oven at 55ºC for four days, and then, cooled to 20ºC to extract 1 ml of gas for GC/MS analysis (Fig. The source of (C2H5)2O is the thermal decomposition of DEC through decarboxylation 30 (Eq. IV.17), assisted by PF5 from the decomposition of LiPF6 (Eq. I.10). Methyl-ethyl ether and dimethyl ether could be expected from the decomposition of EMC, but they were not detected.

[IV.17] The fluorinated alkanes, CH3F and C2H5F, come from the reaction of alkyl carbonate solvents with POF3. First, PF5 reacts with H2O traces (this electrolyte contains ~18 ppm) to produce POF3 and HF (Eq. I.11). Then, POF3 reacts with linear solvents (Eq. IV.18, exemplary shown for DEC) to produce CO2 and an alkyl fluorophosphate 30,31 . In the same way, POF3 may also react with cyclic solvents to form phosphoro-fluorinated species and CO2 (Eq. IV. 19).

[IV.18] [IV.19]

Note that there are two common features in reactions IV.17 -19:

-CO2 is always a product. The fact that multiple sources exist confirm why CO2 appears as an intense peak in the chromatogram.

-PF5 intervenes directly or indirectly (as POF3). In fact, the production of PF5 is facilitated by the available dead volume inside the 18650 can (total volume = 13.64 cm 3 ), as will be evidenced later.

To evaluate the impact of PF5 in CO2 production, the quantity of CO2 coming from a batch of three normal-size 18650 cans stored at 55ºC with LiPF6-based electrolyte was compared to other three batches: (i) normal-size cans with 1M lithium bis(trifluoromethanesulfonyl)imide (LiTFSI)-based electrolyte (to avoid PF5 presence), (ii) normal-size cans with only the solvents mixture and (iii) reduced-size 18650 cans (total volume = 6.5 cm 3 ) with LiPF6-based electrolyte.

Since CO2 is the only gas that exhibits a defined peak, it is the only one that was quantified. In the presence of 1M LiPF6, CO2 is heavily produced whereas only ca. 5% of this CO2 quantity is obtained in its absence. Similarly, replacing LiPF6 by LiTFSI produced around 4% of CO2. On the other hand, when the dead volume is decreased almost by half, but with the LiPF6-based electrolyte volume kept at 0.2 ml, the CO2 amount decreases by more than half.

In conclusion, CO2 production is linked to the dead volume and the presence of PF5 (and POF3) and it is the main initiator of gassing from the storage of powder NMC.

Afterwards, the five Ni-rich NMC powders were evaluated with the same protocol (Fig.

IV .10). Curiously, the CO2 quantity decreases with the increase of nickel content. This seems contradictory to the general knowledge that higher nickel content equals higher reactivity 32 .

However, in the literature the reactivity has been mostly observed in delithiated states where Ni 4+ can reduce to Ni 3+ . In this case, the materials are fully lithiated and have not experienced electrochemical processes, indicating that the differences come from a physicochemical origin. The BET surface areas ranging from 0.20 to 0.27 m 2 .g -1 are practically identical, so the CO2 difference is not related to them. In contrast, the water content varies between 0-114 ppm

(measured with a Metrohm Karl-Fisher titrimeter), but this did not show any relationship to CO2 produced. Again, there is no relation between the expected quantity of CO2 that would result from the chemical reaction of the total amount of Li2CO3 (Eq. IV.6) and experimental CO2. Note that these experimental amounts are much bigger than the expected.

Moreover, it is worth to look at the fractional composition of transition metals. Indeed, there is a direct relation between the decrease of Ni 2+ , Mn 4+ and Co 3+ , (more evident for Ni 2+ and Mn 4+ ) and the decrease of CO2. The exact reason is unclear, but we hypothesize that the production of PF5 is linked to the surface energy of each material that varies depending on the transition metals composition and surface species content. Lastly, other characteristic that changes between NMC compositions is the surface texture, the materials become smoother upon nickel content increase (see insets in Fig. IV.10). This observation makes us think that it might exist privileged planes on the multifacet surfaces that have a catalytic effect on the degassing of PF5. However, this has not been confirmed yet.

To conclude, the storage tests at 55ºC of electrolyte allowed us to identify PF5 as the main cause for gas production, mostly whenever dead volume is available, as a consequence of the thermal decomposition of the electrolyte, producing CO2. Upon the addition of NMC, surprisingly, the quantity of CO2 varies depending on the composition of NMC: the more nickel there is, the lower CO2 is produced. The reason for this behavior remains unexplained.

Film storage

The thermal electrolyte degradation at two different state of charge (SOC) was evaluated through the storage of NMC 622C films and electrolyte in Swagelok-type cells (loading = 11.4 ± 0.5 mg.cm -2 ). Films were used as the electrochemical delithiation is more practical. The entire protocol is illustrated in Fig. IV.11 and consists in:

i. preparation of the Swagelok-type half-cell in the glovebox, as described in Annex II;

ii. cycling at 30 ºC and a rate of C/10 up to a delithiated (4.3V) or lithiated state (2.8 V) (the highest delithiated state is Li0.4Ni0.6Mn0.2Co0.2O2, which by definition here corresponds to 100% of (achievable) SOC;

iii. transfer of the cell to the glovebox and relocation of the lithiated or delithiated film in a half Swagelok-type cell with a septum cap and fresh electrolyte; iv. storage during 4 days at 55°C, followed by cooling down to 20ºC, and v. collection of 1 ml of gas for GC/BID analysis.

The advantage of this system is that it does not contain formation gases and reduction products to interfere with the thermal reactions of the electrolyte. At least 130 µmol.gNMC -1 were produced from the three samples, which corresponds at least to the thermal decomposition of alkyl carbonate solvents. This result is unfortunately not comparable to the previously-shown powder storage due to the several variables: the different dead volume, the NMC-to-electrolyte mass ratio and the carbon and binder contained in the film. However, it is sure that in both cases PF5 contributed to total CO2, due to the existence of dead volume (regardless of the SOC in this case).

Furthermore, the storage of the delithiated film produced two times more CO2 than the lithiated films, and both lithiated films produced around the same quantity. The difference of 128 µmol.gNMC -1 between the delithiated and lithiated film is the evidence of the influence of the SOC in CO2 production, and also constitutes a proof that the delithiated NMC is more reactive toward the electrolyte. This is possible through the production of active oxygen ( 1 O2), that reacts chemically with EC to produce CO and CO2 (Eq. IV.4) 33,34 . CO was not identified due to the separation property of the capillary column. According to Eq. IV.4 and the amount of NMC in the film, 128 µmol.gNMC -1 of CO2 corresponds to 1.8 µmol of consumed 1 O2, and to ~0.6 mol% of oxygen from NMC. This calculation gives an idea of the reactivity of delithiated NMC 622C and can eventually be compared to other NMC compositions.

The production of 1 O2 at this temperature is supported by the work of Xiong et al. 8 , who showed through TGA/MS that delithiated NMC 442 produces oxygen from 40°C. Kim et al. 35 reported that upon storage at 60°C of delithiated NMC 532 (4.45 V), approximately 75% of the gas is produced within the first 10 hours of storage. These results imply for us that the 1 O2 production and the following reactions occur fast at the early stage of the storage at high temperature. Our result complements these works by giving a direct comparison between the lithiated and delithiated state. In upcoming work, it would be interesting to compare the other NMC materials.

Gas production from the storage of NMC/Graphite pouch cells

The complete cell is a more complex system because, in it, co-exist gases produced from both electrodes and electrolyte. In a first instance, this section deals with the optimization of a 2and 3-electrode pouch cell format exhibiting reproducible electrochemical performances and making it easy to sample gases. Then, will be described the special storage protocol that consists of formation cycles followed by a storage step at high temperature. Lastly, the parameters that influence the production of gases from the formation cycles are studied. Gases after the cycling step and the storage step are investigated with two objectives:

i. Differentiate the contribution of Li2CO3 source to total CO2.

ii. Investigate the role of NMC composition.

Optimization of the pouch cell format

For these storage experiments, pouch cells were chosen, since there are customizable.

Additionally, the dead volume can be reduced almost to zero.

Graphite was chosen as active material for the negative electrode. Since graphite does not degrade in air, this film was manufactured with an automatized coating machine that controls tightly the parameters of temperature, coating velocity and gap. The electrode composition The assembly of pouch cells is done inside a dry room (dew point ≤ -60°C ≤ 8ppm of H2O) based on a pre-established pouch cell patron developed by this research group 36 . Dimensions, sealing areas and sealing conditions were customized (Fig. IV.13), in order to meet the requirements of reproducibility, leak-proof sealing, resistance at 55°C and space to store gas.

3-electrode pouch cells contain a thin lithium foil laminated on a copper grid and placed between the electrodes. Note that the cells are laboratory-made (slurries casting, electrodes and separator cutting and assembling).

After the assembly and three-sides sealing, the pouch cells were kept in a vacuum oven for 12h at 80°C to eliminate adsorbed water. The final sealing was done after electrolyte filling (80 µl of base electrolyte: 1M LiPF6 in EC/DEC/EMC 1:1:1 vol.), with a vacuum sealer set at 98% vacuum (percentage referred to the vacuum capacity of the oil pump, details in Annex III). The sealing method and conditions were chosen due to its reproducibility, capacity to leave the minimal amount of dry air inside the pouch cell and assurance that the minimum amount of electrolyte is lost during the vacuum procedure (only 3 -4 wt.% was lost). To extract the gas within a single pierce from the reservoir, a protocol was established. It consists in sticking a rubber septum with cyanoacrylate glue to both sides of the gas reservoir to ensure that the needle do not pierce across the pouch cell. Then, to ensure that always enough gas volume is injected to the GC and that the majority of gas is collected, 500 µl of inert gas are first added to the reservoir, and then the totality of gas is collected. A very recent example of this type of gas recuperation can be found in the literature 37 .

Complete storage protocol of pouch cells

To study specifically the gas production after the storage of NMC/Gr pouch cells when varying the NMC composition and the concentration of surface species, a complete storage protocol was implemented. It includes a cycling step followed by a storage step at 55ºC.

For industrial production of batteries, the formation step varies depending on the materials and the application. Very often, it includes couple of cycles at low C-rate and a resting period at high temperature 38 . Although battery manufacturers evacuate the gases after the formation step, parasitic reduction reactions as cycling proceeds is likely to occur. In this study, a simpler formation step was used, similar to the one used with Swagelok-type cells in Section

Gas production of the formation step

With the aim of identifying which gases are produced during the formation step and the better conditions for the complete storage protocol, this section deals with temperature and the C-rate variations during the formation step.

Four batches of three pouch cells NMC 811/Gr with equivalent loadings, were cycled at 30 or 55°C and at a rate of C/5 or C/10 (Table IV The gases were extracted with the protocol previously mentioned and analyzed by GC/BID using the capillary column and the improved method (CH4 is visible). The gases that were identified from the formation cycling, regardless of the temperature or C-rate, are CH4, CO2, C2H4, C2H6, H2O, DMC, EMC and DEC. ). However, it is not ensured that the totality of Li2CO3 is oxidized electrochemically during the formation step.

For comparison, in Chapter III through FTIR analyses it was found that in exposed NMC 622C powder around 59% of Li2CO3 (from which 27% was chemically consumed) could be consumed after one charge at 25°C and a rate of C/10. Due to the use of NMC films in the pouch cell system, it was not possible to quantify the Li2CO3 consumption through FTIR (because the reference sulfate band was covered by PVDF bands). Alternatively, in the next section, an experiment was put in place to try to discern the contribution of CO2 from Li2CO3.

From these results, cycling at 30°C and C/10 were the chosen conditions to perform the formation step. Indeed, the lower current density allows to form a SEI with better properties 16 , and the lower temperature hinders the thermal degradation of the electrolyte.

Gas production of NMC 811 with modified surface composition

Since the identification of CO2 production from Li2CO3 was not possible from the formation step, the complete storage protocol was performed with an as-received and exposed NMC 811 material in complete pouch cell.

After the storage at 55°C, the pouch cells were cooled down to 20°C and the gas was collected and analyzed by GC/BID with the improved method (CH4 is visible). Table 5 shows the main characteristics of the pouch cells. Note that both positive electrodes have equivalent loading, but a very different quantity of Li2CO3, the surface of as-received NMC 811 contains only 1.04 µmol, whereas the exposed contains 20.9 µmol. As detailed in Chapter III, the electrochemical performance of exposed electrodes is greatly perturbed by the presence of surface species after H2O/CO2 exposure. The as-received 811/Gr pouch cells exhibit a 2 nd charge capacity of 10. It can be observed that the peak areas of formation gases CH4, C2H4 and C2H6 remained almost unchanged regardless of surface composition, and that the CO2 produced from 811(exposed)/Gr (0.099 ± 0.015 µmol) almost doubles compared to the as-received 811/Gr (0.052 ± 0.005 µmol). The most relevant difference between both materials was the concentration of surface species. Therefore, it is safe to assume that the difference of 0.047 µmol of CO2 came exclusively from the chemical (Eq. IV.6) and electrochemical (Eq. IV.7

and IV.4) reactions of Li2CO3. This value is far from the theoretical quantity of CO2 from the complete consumption of Li2CO3: 20.9 µmol. There are two possible explanations for this:

-Incomplete chemical and electrochemical consumption of Li2CO3.

-Partial dissolution of CO2 in the electrolyte 40,41 . As an example, the solubility of CO2 in a solvent mixture 1:1 wt. EC/DMC is 0.568 molCO2.L -1 at 20ºC 42 . This means that in 80 µl of solvents, a maximum of 45 µmol of CO2 can be solubilized. Because LiPF6 is contained in the electrolyte, the actual quantity of dissolved CO2 would be six times lower, as predicted by the Setchenov law (salting-out effect) 41 .

In fact, both explanations are likely to occur. The first was proved in Section 3.2.4 of Chapter III, whereas the latter could explain the lesser quantity detected, compared to the expected theoretical one. Deeper study of these phenomena was left out of the scope of this work.

These experiments revealed the contribution of Li2CO3 to total CO2. Unfortunately, it was not possible to differentiate the chemical from the electrochemical contribution. Note that similar quantities of gas originating from solvents reduction to form SEI are obtained, regardless of the NMC surface state, which show the reliability of the protocol.

Gas production of different NMC compositions

In order to investigate the role of NMC composition on the gases produced in storage, this section studies the gas production at different stages of the complete storage protocol of twelve pouch cells NMC 622C/Gr and twelve pouch cells NMC 811/Gr. Among them, a third of each composition performed the formation step (C/10 and 30°C), a third performed the complete storage protocol (formation + storage) and the last third performed only a storage step (4 days at 55°C just after the electrolyte impregnation 5h-period) to discriminate the amount of gas issuing from the thermal solvents degradation only.

Table IV.6 shows the loadings and current densities of each NMC/Gr system. The applied current density in the formation step was higher in pouch cells containing NMC 622C than 811 (0.48 vs. 0.31 mA.cm -2 ) due to higher positive electrode loading (17.3 vs. 13.7 mg.cm -2 ). because lower current density is applied in this case, which induces more solvents reduction 17 (negative electrodes are equivalent). This can be verified through peak integration (Fig.

IV.21), where the integral value of the SEI formation peak in pouch cells 811/Gr is higher than in 622C/Gr (8.0x10 -4 vs. 5.9x10 -4 , respectively). On the other hand, higher quantity of CO2 is detected from 811/Gr than from 622C/Gr. In both systems the CO2 is originated from several processes: i. electrochemical oxidation of EC (Eq. IV.1),

ii. oxidation of EC by 1 O2, from NMC at the delithiated state during cycling (Eq. IV.4).

iii. chemical consumption of Li2CO3 (Eq. IV.6), iv. electrochemical consumption of Li2CO3 (Eq. IV.7) and v. thermal degradation of alkyl carbonate solvents (at 30°C) (Eqs. IV.17, 18 and 10),

It was not possible to quantitatively assign a value of CO2 to each process. However, conjectures can be drawn that explain why 811/Gr produces 0.018 µmol of CO2 more than 622/Gr.

First, the electrochemical oxidation of EC due to cycling above 4.2 V (i) should not vary as much from one system to the other. In contrast, the oxidation of EC due to 1 O2 (ii) is more likely to occur in pouch cells NMC 811/Gr than in 622C/Gr. This is supported in the work of Jung et al. 9 , who showed that the production of CO2 from EC oxidation by 1 O2 starts at ∼4.7 V for NMC 622 and at ∼4.3 V for NMC 811.

Concerning the chemical and electrochemical oxidation of Li2CO3 (iii and iv), in Chapter III,

through NMR experiments, we demonstrated that the chemical consumption of Li2CO3 is not efficient at 30ºC. Instead, we believe that electrochemical oxidation is predominant over chemical, but we do not know to what extent.

Lastly, regarding the thermal degradation of solvents (v), pouch cells NMC 622C/Gr stayed for 25 h at 30ºC, whereas 811/Gr stayed 28.5 h, a difference of only 3.5 h. Although not discussed, in Section 3.2.3 Fig. IV.16, it was found that in the time of 10h only 0.003 µmol of CO2 were produced from 811/Gr at 30°C, therefore, the 0.018 µmol of difference between NMC 622C and 811 pouch cells cannot be mainly attributed to thermal degradation at 30ºC.

To summarize, during the formation step of pouch cells NMC 811/Gr, the main CO2 contributors were assumed to be the electrochemical oxidation of Li2CO3 and the EC oxidation.

In conclusion, comparing the gases of the formation step from pouch cells containing NMC 622C and 811, it was found that C2H4 and C2H6 production stemming from solvents reduction depends on current density. On the other hand, CO2 depends on cycling time and nickel content.

Discussion about gas of the complete storage protocol

The analysis of the gas resulting from the complete storage protocol gets more complex because of the combination of the gases of the formation and the storage steps. As it can be seen in Fig.

IV.20b, whatever the NMC composition, more CO2 was produced and a large part of C2H4 was consumed, whereas C2H6 remained mostly unchanged.

This led us to think that there were other intermediary oxidation reactions, as proposed in Eqs.

IV.21 and 22, forming ethylene oxide (C2H4O) or CO and H2O:

2$ # ) d + % # 2 → 2$ # ) d % [IV.21] $ # ) d + 2 % # 2 → 2$% + 2) # % [IV.22]
C2H4O and CO are in fact molecules that may be present in the battery from thermal decomposition 25 , so if detected, their attribution solely to C2H4 oxidation is impossible. Note that C2H4 is not the only hydrocarbon that can be oxidized, there are CH4 and C2H6 as well, but it was already observed that the C2H6 quantity did not change greatly in Fig. IV.20 and CH4 was not detected with the used protocol. Likely, EC could also be oxidized (Eq. IV.4).

We deduce that gaseous C2H4 would be preferentially oxidized over dissolved EC, because most of the EC molecules solvate lithium ions and are thus less prone to oxidation. However, this affirmation requires additional experiments.

-Explanation to the production of CO2

The processes that contribute to total CO2 after the complete storage protocol are the same that from the formation step (enlisted hereby as i -iv). The variant is that processes ii, iii and v are exacerbated by the storage at the delithiated state and at high temperature. Moreover, the possibility that C2H4 is oxidized by 1 O2 (vi) is added to the previous five processes:

i. electrochemical oxidation of EC (Eq. IV.1),

ii. oxidation of EC by 1 O2, from NMC during storage at the delithiated state (Eq. IV.4).

iii. consumption of Li2CO3 that was not totally consumed electrochemically (Eq. IV.6), iv. electrochemical consumption of Li2CO3 (Eq. IV.7) and v. increased thermal degradation of carbonate solvents (at 55°C) (Eqs. IV.17, 18 and 19),

vi. oxidation of C2H4 by 1 O2 (to an unknown extent, if any),

Indeed, the quantity of CO2 increased from 0.006 µmol in the formation step, to 0.034 µmol after the complete storage protocol in pouch cells 622C/Gr. Similarly, CO2 increased from 0.024 µmol to 0.060 µmol in 811/Gr (recapitulated in Table IV.8). The gap between the CO2 produced from the complete protocol and the formation step will be called "difference B" in The discrepancy among the values of "difference B" is due to the reactivity of each NMC during the storage at high temperature, involving the processes ii and vi that involve reactions with 1 O2 (more 1 O2 is produced from 811 than 622C). In this sense, "difference B" corresponds to the storage step at the delithiated state, whereas "storage only" to the lithiated state. Thereby, the difference between them ("difference C" in Table IV.8) corresponds to the quantity of CO2 from the reactions with 1 O2. The values are 0.007 and 0.016 µmol, from which we deduce that 811/Gr contributes more to total CO2 than 622C/Gr due to the reactivity of delithiated NMC 811.

Finding more CO2 after the storage of pouch cells is in fact contradictory to the work of Dahn's group 20,21 , who proposed that CO2 in the cell NMC 442/Gr diffuses to the negative electrode and reduces to lithium oxalate during storage (Eq. IV.16), based on the sole observation of CO2 decrease after storage at 40°C. As a supplementary proof, Ellis et al. 43 showed XPS analyses where the O=C-O -peak was intense for the negative electrodes from stored pouch cells NMC 442/Gr, and through GC analyses showed that CO2 was the most consumed gas. At the stage of the advancement of this project it is not possible to differentiate if some CO2 has been consumed in the negative electrode, which can be envisaged to be studied through FTIR analysis of the SEI in future work.

To conclude, these tests gave evidence about the higher reactivity of NMC 811 through the consumption of C2H4 and the production of CO2 after the complete storage protocol. The higher quantity of C2H4 consumed from cells 811/Gr than from 622C/Gr, led us to deduce that the consumption occurs through the reaction with 1 O2.

Impedance evolution after the complete storage protocol

At the same time that gas is produced from storage at high temperature, increase of the cell impedance has been reported to occur 20 , due to the formation of a rock salt phase from the structural transformation of layered NMC 11 . To the author's knowledge, there are not existing comparisons between different NMC materials in terms of impedance evolution before and after a storage period at high temperature.

Adaptation of the complete storage protocol

The impedance experiments were performed on 3-electrode NMC 622C/Gr and 811/Gr pouch cells. Potentiostatic electrochemical impedance spectroscopy (PEIS) measurements were recorded before and after the storage step of the complete storage protocol, at 30ºC, at frequencies ranging from 300 kHz to 0.05 Hz, with 7 points per decade (VMP-3 Biologic).

Since self-discharge during the storage step is expected to occur, the protocol includes intermediary PEIS measurements at 90% and 80% after the formation step, in order to have a comparison scale with the PEIS measurements after storage. At the end of the storage, the pouch was allowed to descend to 30ºC and a recharge to 100% SOC was done, followed by 15 min. rest period before performing PEIS (Fig. IV.22 summarizes the different steps).

Throughout the protocol, the pouch cells were not disconnected or moved to avoid any misinterpretation. Each of these Nyquist plots is composed of three parts, explained by several authors [44][45][46][47] :

i. a high-frequency semicircle attributed to the contact interface resistance between the current collector and particles (Rcontact), with frequency values for the positive electrode between 1 -100 kHz;

ii. a middle-frequency semicircle attributed to the charge-transfer process (Rct), with frequency values for the positive electrode between 1 -100 Hz, and

iii. a low-frequency straight line, corresponding to the Li + diffusion.

The resistance value crossing the X axis corresponds to the electrolyte resistance (Relect) and the resistance of the cables.

The full cell spectra are practically identical; however, in the spectra of the individual electrodes there are slight disparities in the process of charge transfer: the shape of the semicircle decreases in the positive electrode and increases in the negative electrode; whereas the shape of the contact interface resistance is unchanged. The variations when measuring at 100, 90 or 80% of SOC are minimal, which means that, up to 20% of discharge, the different interfaces are slightly modified.

To go further in the comprehension, the extraction of the spectra characteristics was undertaken, with the aid of an equivalent circuit model previously used in the literature 44,48 (Fig. Landesfeind et al. 48 consider that the Rct extracted from this model is overestimated because the semicircle includes the contribution of the electrolyte ionic resistance within the pores. In our case, fitting was performed without considering this contribution.

3.2.6.2. Impedance from pouch cells NMC 622C/Gr and 811/Gr After the storage period, the OCV of pouch cells NMC 622C/Gr and 811/Gr stored in the charged state gradually decreased, until a corresponding calculated self-discharge of 9-12% SOC, which appeared to be totally reversible. In the literature, it has been proposed that the layered materials have a reversible and irreversible part of self-discharge upon high temperature ageing 49 . The self-discharge reversible part has been attributed to lithium ions that re-intercalate from the electrolyte into NMC; whereas the irreversible part has been attributed

to surface transformations (hexagonal to spinel in LCO 49 ). In our pouch cells, the capacity was totally recovered after recharge to 4.3 V. Maybe the irreversible part was not detected because the amount of produced oxygen ranges at least around 0.015-0.021 mol% (Section 3.2.5.2). This implies the creation of a very thin rock salt layer which will not likely affect the capacity, although it may impact the PEIS measurements. The frequency values calculated from the contact interface resistance, between 8-18 kHz, match the frequency values reported in the literature [44][45][46][47] . This is also the case of the frequencies calculated from the charge-transfer process, which lie between 1-36 Hz. By looking at the alpha values and the shape of the Nyquist plots, it has been difficult to add a third contribution corresponding to the CEI.

On the Nyquist plot of the positive electrodes before storage, the semicircle attributed to the contact interface resistance is slightly bigger than the one corresponding to the charge transfer process (Fig. IV.25c, d). In terms of values, Rcontact was found to be 2.88 Ohm.cm 2 for 622C and 6.84 Ohm.cm 2 for 811. The higher value of the NMC 811 electrode may be due to the absence of calendering, compared to 622C electrode.

After storage, the fitted values show that Relect increased very little. This slight increase means that the dissolved species from electrolyte degradation did not change the resistance of the electrolyte significantly. Additionally, Rcontact doubled in both positive electrodes, and Rct increased by 40% in 622C and 74.3% in 811. The similar increase of Rcontact means that the contributions of the interfaces between current collector and active particles and the resistance of Li + migration through the CEI (if any) evolve in a proportional way in both NMC electrodes.

In NMC 622C/Gr pouch cell, the Rcontact increased by 53.8%, compared to +40% of the Rct; whereas in 811/Gr the contrary is observed: Rct increased by 74.3% compared to +53.8% of Rcontact. The increase of the resistances in both cases can be attributed to the formation of the rock salt structure at NMC surface. In this sense, Sallis et al. 11 determined that the formation of rock salt phase in layered NCA is reflected as a marked increase of Rct and a lesser increase of Rcontact, which is consistent with the observations in NMC 811 with similar composition. This is also consistent with our analysis of the gas evolution: more CO2 is produced and more C2H4 is consumed as a consequence of the formation of 1 O2 and rock salt phase. In this way, PEIS measurements constitute supplementary evidence to the increased reactivity of Ni-rich NMC and highlight the necessity to avoid surface transformations in Ni-rich NMC materials.

To gain more insight into the impact of rock salt formation on the impedance measured after the storage step, the same protocol was performed in a 3-electrode pouch cell 811/Gr changing the storage temperature to 10°C. This experiment was undertaken considering that the 1 O2, and

Conclusion of the chapter

This chapter evidenced the reactivity of Ni-rich NMC materials from the analysis of the gaseous products upon storage with electrolyte at 55ºC. First, among the available GC devices, columns and detectors, the best sustainable and reproducible measurements were obtained from the GC/BID device equipped with a capillary column and set in "linear velocity" control mode. A micropacked column with better gas separation properties was also tested, but it was not preferred due to a lack of reproducibility.

The CO2 quantity produced from pristine NMC and electrolyte storage at 55°C was found to be driven by PF5 and dependent on dead volume. On the other hand, it was found that the storage at delithiated state increases the CO2 production, explained by the presence of Ni 4+ that reduces to Ni 3+ producing 1 O2, which oxidizes EC.

Gas production from the storage of 2-and 3-electrode pouch cells was then evaluated. For this, their fabrications were optimized and a protocol containing a formation and a storage step was established. After the formation step, it was found that the detectable formation gases (CH4, C2H4, C2H6) are dependent on current density. Additionally, several sources of CO2

were considered: chemical and electrochemical oxidation of solvents and Li2CO3, and thermal degradation of solvents. In our case, the NMC composition was found to influence CO2 production since n(CO2)811 > n(CO2)622C. After storage, two phenomena were observed: C2H4 consumption and an increase of CO2 concentration. Both are issued from the same cause:

oxidation of C2H4 and EC by 1 O2, influenced by the composition of NMC. Indeed, in compositions richer in nickel, more Ni 4+ exist, so more 1 O2 can be produced. At the same time, a restructuration of the surface with the formation of a rock salt phase occurs, causing impedance augmentation. This was evidenced in 3-electrode pouch cell experiments.

Finally, we acknowledge that there are many things left to be investigated, for instance, CO2 dissolution, crosstalk reactions and long-term gas evolution. The latter is investigated in the next chapter by using the methods here developed.

Introduction

The use of electrolyte additives is one of the most effective and economical ways to improve the electrochemical performance of LIBs. In Chapter III, the salt, lithium difluorophosphate, weight content (usually between 0.5 and 3 wt.%), LDFP is capable to:

LiPO2F2
• decrease the impedance of the positive electrode [3][4][5] , negative electrode 2,6,7 or full cell [7][8][9][10][11] • prevent lithium plating 12 and decrease cation dissolution 8

• decrease gas production after storage at high temperature 9,12,13 • impede surface transformations from layered to rock salt structure in NMC 622 14 and

• improve capacity retention after at least 100 cycles at 1C, C/2, C/3 and 2C; in half cell or pouch cells with upper cut-off voltages as high as 4.5 V vs. Li/Li + 3, 7-11, 15 .

Usually, the improved electrochemical performances when LDFP is added is ascribed to the formation of a solid electrolyte interphase (SEI) or a cathode electrolyte interphase (CEI).

However, there are contradictory opinions about the mechanism. Through theoretical calculations or experiments, some authors propose that it gets reduced 6 , oxidized 3,8 and others, that it simply precipitates onto the electrode material 2,6,11,15 . Unfortunately, none of these reports present a mechanism and simply assume that the reduction or oxidation would form an efficient SEI or CEI. Among those who propose that no redox reaction occurs, Yang et al. 2 , explain the precipitation of LDFP into the SEI and CEI through a migration mechanism, where LDFP is the lithium carrier (Fig. V.2). During charge, LixPO2F2 (x = 3) would reach the negative electrode by potential difference and deposit, while the supplementary lithium ions would intercalate. Inversely, during discharge, LixPO2F2 (x = 3) would migrate to the positive electrode, deposit, and the addititional Li + intercalate. The fact that PO2F2 -and Li + ions associate into ion-triplets in the electrolyte, due to the low dissociation of LiPO2F2 in alkyl carbonate solvents 16 , it is less likely that they deposit reversibly on both electrodes. In 2019, two other hypotheses were postulated. Qian et al. 15 , suspected that LDFP prevents electrolyte ageing through the trapping of the initiators of carboxylate species formation, such as 2,5-dioxahexane dicarboxylate (DEDOHC or compound 11). Although the GC/MS results are remarkable, no reactions were shown to support this hypothesis, so the mechanism remains unclear. On the other hand, Zhao et al. 16 explained that LDFP forms a CEI that supposedly inhibit to some extent the decomposition of LiPF6, prevents cation dissolution, HF formation and contribute to decrease electrolyte oxidation due to the presence of phosphate ions (PO4 3-), as observed by XPS. To explain the presence of PO4 3-, they proposed the formation of Li3PO4 through the "self-decomposition" of LDFP into LiF and PO2F (Eq. V.1), followed by PO2F reaction with traces of H2O in the electrolyte then H + /Li + exchange (Eq. V.2, 3) or with Li2CO3 from the SEI (Eq. V.4). Additionally, they attributed an oxidation peak around 3.1 V vs. Li/Li + to the oxidation of PO2F into PO2 + and F2 (Eq. V.5). These reactions are highly debatable, since F2 is not likely to be produced through equation V.5.

LiPO # F # → PO # F + LiF [V.1] PO # F + 2H # O → H & PO d + HF [V.2] H & PO d + 3Li 1 → Li & PO d + 3H 1 [V.3] 2PO # F + 4Li # CO & → 2Li & PO d + 4CO #↑ + 2LiF [V.4] 2PO # F .#ã ] h⎯⎯j 2P% # 1 + É # [V.5]
None of the three described hypotheses present a concrete and sufficiently grounded explanation to what occurs to the electrolyte or the interfaces when LDFP is used. Even though the last seems very complete, it is based solely on the presumed observation of PO4 3-.

In light of the previous information, there is room for better understanding of the LDFP's action mechanism that has not been reported before. In this chapter, in a first instance, the properties of the electrolyte containing LDFP were evaluated. Then, the cycling performance at 25 and 55°C is compared for NMC 811 and at 55°C for 622C, all in full cells. Given the surprising contradictory results, a following section is devoted to the analysis of the solid, liquid and gaseous electrolyte degradation products after prolonged cycling at 55°C in the hope of providing tangible explanations.

Impact of LiPO2F2 on the electrolyte properties

This section studies the electrochemical stability, viscosity and ionic conductivity changes when LDFP is added. LDFP was used with a concentration of 0.2M (1.92 wt.%) in the base electrolyte (1M LiPF6 in EC/DEC/EMC 1:1:1 vol.). As the PO2F2 -ionic volume is lower than that of PF6 -(538.1 and 625.5 Bohr 3 16 , respectively) and the dissolution energy calculated by density functional theory (DFT) is larger (-5.790 eV for LDFP and -6.254 eV for LiPF6),

LDFP is less soluble in the alkyl carbonate solvents, so it can be used as additive not exceeding 2 wt.% 11,16 .

Cyclic voltammetry (CV) at a scan rate of 30 mV.s -1 was performed on the base and base + 0.2M LDFP electrolytes (Fig. V.3) using platinum as working (disc d=3 mm) and counter electrodes, and Li as reference electrode. From the first oxidative scan (anodic scan) from 0 to 6 V, it can be appreciated that the electrolyte's stability was not modified by the presence of LDFP and the absence of additional peaks means that LDFP does not reduce or oxidize up to 6

V. In spite of the very low measured current, normally, signals around ~0.3-0.35 V and ~0.65-0.7 V corresponding to lithium alloy formation are obtained on a platinum electrode upon the anodic scan, whereas upon the cathodic scan, a signal around ~0.55-0.6 V is assigned to bulk lithium dissolution and two peaks around ~0.95-1.0 V and ~1.3 V are related to stripping of lithium from the alloy. The reduction of trace water would take place around ~1.5 V 17 . In reality, in the cathodic scan from 6 to 0 V, a pronounced wave between 3-1.8 V is observed. According to Gnanaraj et al. 17 , it corresponds to the electrochemical reduction of some of the insoluble surface species (such as Pt-X) that passivate the electrode, formed during the anodic scan. The slightly differences around 2.7 V might be due to a different construction of the SEI or an eventual impurity in the electrolyte rather than a process related to the additive. Although not shown here, the second cycle does not show this cathodic wave since the electrode is already passivated. In the literature, Kim et al. 6 reported the reduction of LDFP around 0.6 V vs. Li/Li + through dQ/dV analysis of the first cycle of graphite/Li cells. After careful analysis, we deduced that this peak corresponds only to the solvent's reduction, shifted to higher potentials upon LDFP addition. Regarding the oxidation of LDFP, Chen et al. 3 calculated through DFT that the oxidation of LDFP would occur around 5.9 V, before the calculated oxidation of EC at 6.97 V or EMC at 6.78 V 8 . However, experimentally, they explained that LDFP is oxidized below 4.5 V, supported only by the observation of a CEI layer by TEM, formed principally of LixPFyOz species, as detected by XPS.

To summarize, from these reports, no ground proof of LDFP's reduction or oxidation can be obtained. For the following experiments in this chapter, it is considered that, based on the CV experiments, LDFP is not electrochemically active in the potential window of NMC materials.

Afterwards, the ionic conductivity and viscosity of the electrolyte were investigated (experimental conditions can be found in Annex IV). The ionic conductivity was measured in the liquid form between two platinum electrodes (liquid measurement) and inside a porous In the liquid measurement, the conductivity of the base electrolyte was found to decrease by ca. 14% and 17% at 25°C and 55°C respectively upon the addition of 0.2M LDFP. In this regard, two similar trends have been recently published with different alkyl carbonate mixtures: Wang et al. 11 found that at 25°C, the original ionic conductivity of 6.34 mS.cm -1 decreased by 9.4% with the addition of 2 wt.% of LDFP (base: 1M LiPF6 in EC/DEC 1:3 wt.%). Meanwhile, Zhao et al. 16 found a decrease from 7.31 to 6.72 mS.cm -1 (-8%) with the addition of 1 wt.% of LDFP (base: 1M LiPF6 in EC/DMC 1:1 vol.), measured at room temperature.

The decrease in ionic conductivity is in line with the increase of viscosity from 3.8 to 4. Concerning the 0.8M LiPF6 + 0.2M LDFP electrolyte, it was projected to be less deleterious to replace part of LiPF6 than exceed 1M concentration. In reality, among the three electrolytes in the liquid configuration, this electrolyte exhibits the lower ionic conductivity, especially at high temperatures. Contrarily, Liu et al. 19 reported that an electrolyte 1M LiPF6 in EC/DEC/EMC 1:1:1 wt.% shows ionic conductivity of 10 mS.cm -1 , while an electrolyte with 0.8M LiPF6 and 0.2M LDFP displays a higher conductivity of 11.6 mS.cm -1 (16% more).

Within the Celgard 2500 separator, the ionic conductivity of the three electrolytes is divided almost by 10, due to tortuosity and the restricted ion transport through pores (Fig. V.4b) 20 , and the differences when LDFP is added are very small. As in liquid configuration, the 0.8M LiPF6 + 0.2M LDFP presents slightly lower values at temperatures higher than ca. 10°C and thus will be discarded in the following experiments.

The differences in ionic conductivity of electrolyte with and without 0.2M LDFP are so small, that no real impact on the electrochemical performances is expected upon following NMC/Graphite(Gr) pouch cells cycling at C/4.

3.

Impact of LiPO2F2 on the electrochemical performance of NMC 622C and 811

Previous works that studied the cycling performance of cells containing LDFP have reported the use of 0.5 -3 wt.% of LDFP in LiPF6-based electrolytes. The general idea is that LDFP always improves the capacity retention. For instance, after 100 cycles at 1C, LDFP addition improved the capacity retention of pouch cells NMC 111/Gr (4.5 V) 11 , NMC 532/Gr (4.53 and 4.5 V) 8,10 , as well as of half-cell NMC 111/Li (4.4 V) 3 . In the same manner, after 100 cycles at C/2 and C/3 and up to 4.2 V and 4.3 V, pouch cells NMC 111/Gr 5 and NMC 532(single crystal)/Gr 9 improved greatly their capacity retention at room temperature. At lower temperature (0°C and -20°C) and C/2, better capacity retention has been reported in NMC 532/Gr 7 . Nonetheless, upon temperature increase, the capacity retention of NMC 532/Gr pouch cell was not improved at 45°C and C/2 15 or was slightly improved in NMC 532(single crystal)/Gr pouch cells at 40°C and C/3 9 , both with cut-off voltages of 4.2 and 4.3

V.

In short, with 0.5 -3 wt.% of LDFP electrolyte additive, improved performances of relatively low Ni content NMCs-containing cells can be achieved from -20 to 40°C. Note that this additive did not show any improvement at a higher temperature of 45°C in full cells.

With the aim of gaining further insight into the understanding of the impact of this additive in Ni-rich cathode cells, this section compares the electrochemical performance of NMC 622C/Gr and NMC 811/Gr pouch cells cycled at 25 and 55°C, and at C/4 (considering total delithiation, although the maximum for 622C is Li0.4Ni0.6Mn0.2Co0.2O2 and for 811 is Li0.3Ni0.8Mn0.1Co0.1O2). Note that due to the nature of the laboratory-made electrodes (little reproducibility given by the absence of calendering in NMC 811 and loading differences from a film to another), a reference cycling with base electrolyte was performed for each set of cells. Before comparing the role of the additive, it can be noticed that the initial discharge capacity when cycling at 55°C is 12% bigger than at 25°C: ~210 mAh.g -1 compared to ~185 mAh.g -1 .

This gain is due to the limited kinetics of lithium intercalation on layered materials at ambient temperature. To compare, Jalkanen et al. 21 and Zhou et al. 22 nickel content (111 11,12 and 532 [8][9][10] ).

On the contrary, at 55°C, similar cells containing 0.2M LDFP show a poor capacity retention of ca. 65% after only 110 cycles, although behaving similarly up to 30 cycles. Experiencing a lower concentration of 0.1M LDFP (0.96 wt.%) at this high temperature results in delaying the negative effect. Indeed, performances are even slightly improved up to 110 cycles, and then the capacity falls down, retaining only 67% of the original capacity after 200 cycles.

To investigate the source of the capacity fade when LDFP is used at 55°C, the differential capacity versus voltage profiles were plotted (Fig. • Whatever the composition of the electrolyte, the peak centered at 3.65 V (part of H1àM) decreases its intensity until, at 200 cycles, it has completely disappeared. This has been attributed to the loss of available lithium sites, related to the gradual nickel migration to the lithium layer upon prolonged cycling 24 .

At this point, the capacity fade of the NMC 811/Gr cells cycled with LDFP at 55°C cannot be unquestionably attributed to one of the electrodes. However, it is supposed that one or both interfaces are modified by the presence of LDFP, limiting cell polarization and that some active materials are isolated.

Impact of the calendering process

The absence of calendering can be a source of capacity fade due to increased electrode polarization 22 , no more percolating particles... Hence, a method to calender air-sensitive films in the dry room was adapted to assess its impact on cells capacity retention. It consists in the implementation of a "portable" laminating machine that uses a gas source to generate pressure between two plates. Then, the film to be calendered is positioned between two brass laminates and introduced between the two plates. After several essays, a compromise was found between the achievable pressure, which was set at 56 psi, and the heating temperature, high enough to calender the material but low enough to avoid polymer melting. This temperature was set at 190°C, although the temperature in contact with the film was obviously lower since PVDF (160 < Tm < 177°C depending on crystallization state) does not seem to have melted on SEM pictures (Fig.

V.8c and d). The calculated initial porosity of the film was 49% and the target porosity 35%. However, the maximum that could be achieved with this dry room setup was 40%. The calendering process improves all the cells' capacity retention; from 71% to 82.4% with base electrolyte after 200 cycles (blue curves) and more significantly, from 65% to 78% after 110 cycles with 0.2M LDFP. The capacity improvement is likely due to enhanced electronic percolation through the electrode thanks to the elimination of the fissures and particles rapprochement. Nevertheless, this mechanical process does not allow for the entire recovery of the capacity; the negative impact of LDFP on the NMC 811/Gr cells capacity retention at 55°C is only stunted.

Post-mortem analyses

The pouch cells NMC 811(cal.)/Gr were disassembled, at the discharged state, after 200 cycles It is known that the stress and strain generated on the particles due to the rapid volume contraction of the c-axis during structural transformation from H2 to H3 during delithiation is one of the origins of particle cracking 25,26 and capacity fading 27 . Since the mechanical failure seems exacerbated in presence of LDFP after cycling at high temperature, in order to understand the role of this additive, the electrochemical performances of NMC 622C films were compared under similar cycling conditions. 33,34 ; the two bands at 885 and 930 cm -1 correspond to P-F symmetric and asymmetric stretching vibrations, whereas the two other bands at 1150 and 1270 cm -1 originate from P-O asymmetric and symmetric stretching vibrations. As a peak at 930 cm -1 of low intensity is depicted in all the LDFP-SEI spectra, it could be suggested that a small amount of LDFP is precipitated and trapped in the SEI, or that other species phosphate-containing exist, such as Li2PO3F, less soluble than LiPO2P2.

The set of at least two broad peaks between 960 -1070 cm -1 that, according to Becker et al. 35 , corresponds to P-O-C stretching vibration, implies that when 0.2M LDFP is used, the SEI contains (or contains more) phosphate species with alkyl groups attached. These molecules are supposed to originate from a nucleophilic substitution of LDFP's fluorine with alkoxide ion RO -formed during solvent reduction (Eq. V.7), as further detailed later in Section 4.2.1. These species can also be present in base-SEI as LDFP can be formed from the reaction between lithium carbonate or alkyl carbonates with LiPF6 (Chapter III).

!"Ç% # É # + R%!" → !"Ç% # É%R + !"É [V.7]

Along with phosphate-based compounds, this reaction also yields the passivating compound, LiF, hence preventing from further degradation of electrolyte. Even more, LiF can be produced from the thermal degradation of LiPF6 producing PF5 and from the reaction between LDFP and the CMC binder (Eq. V.6). In literature, XPS results agreed that the amount of LiF increases with addition of LDFP in electrolyte. However, they also highlight some contradictions with respect to LixPFyOz species content; Kim et al. 6 , reported than upon addition of 1 wt.% of LDFP, the quantity of LiF in the SEI increases, whereas LixPFyOz changes very little. Shi et al. 5 , reported that the content of LiF increases at the expenses of LixPFyOz when 1.4 wt.% of LDFP is added. Yang et al. 7 To sum up, the investigation of the SEI composition making use of the FTIR spectroscopy technique provides fruitful information about the role of LDFP on the graphite composite powder upon cycling: it does not inhibit the reduction of solvents and the SEIs are richer in LiF and phosphate or fluorophosphate species. On electrodes, it may react with CMC binder and as a consequence, marginally limit the solvents reduction.

Study of CEI composition

Upon cycling, a cathode electrolyte interphase (CEI) may form. It is similar to the SEI in a sense that it contains organic and inorganic compounds, which may form depending on the temperature and cycling conditions 36,37 . In the case of cycling with LDFP, some authors proposed that the CEI layer (referred as LDFP-CEI) is passivating toward electrolyte oxidation thanks to the deposition of LiPF6-derived decomposition products 3 or LDFP itself 2,6 .

Regarding LiF, Lei et al. 8 showed that it almost absent in the LDFP-CEI and Zhao et al. 16 did not observe any difference in content. Finally, Qian et al. 15 asserted that, based on XPS analyses, the LDFP-CEI is richer in fluorine.

In order to characterize the CEI formed upon cycling, FTIR analyses were performed on washed NMC 811 composite powders (93.7 wt.% lithiated NMC and 6.3 wt.% C45) after 3 cycles vs. lithium at 55°C and a rate of C/20. The experiments were done in FTIR in transmission mode (Annex IV), to provide a greater chance of obtaining an exploitable signal.

In the original spectra, no additional peak other than those of the active material are visible, indicating a very thin CEI layer, if any. After highly magnifying the spectra (Fig. The spectrum of the reference film shows mainly PVDF bands (Fig. 17a). After long cycling, remarkably, spectra are quite similar. Their bands can be assigned to phosphate and/or organofluorophosphate and/or fluorophosphates such as LDFP (Fig. 17b). The presence of LDFP and possible derivatives in base-CEI would not come as a surprise as this product was found to be produced along with LiF, through the reaction between Li2CO3 or Li2O surface species and LiPF6. In literature, the band centered at 1280 cm -1 was suggested to correspond to the P-O stretching vibration of the phosphorous pentoxide (P2O5), formed by the reaction of POF3 with Li2O 36 (Eq. V.8). Unfortunately, these results do not provide any clue for explaining the different electrochemical behaviors.

2Ç%É & + 2!" # % → Ç # % + 6!"É [V.8]

Analysis of electrolyte degradation products

The analysis of SEI suggested the presence of organo-fluorophosphate compounds stemming from chemical reaction, which could have changed the nature of the degradation products accumulated in the electrolyte. In literature, only Qian et al. 15 According to previous studies of electrolyte degradation by ESI-HRMS and GC/MS 38 , these three peaks correspond to the first member of a family of ethylene oxide oligomers terminated with alkyl (methyl or ethyl) carbonate groups (also called 1n, 1'n and 1"n, whose structural formula appears as inset). The more concentrated peak corresponds to 1'1 with ethyl, ethyl terminations, which matches well with the statistical presence of this group in the electrolyte containing the EC/EMC/DEC (1:1:1 vol.) carbonate mixture.

According to Sasaki et al. 39 , the reaction mechanisms to form the 11 family consist in two steps that start with the single electron electrochemical reduction of linear solvents (Eq. V.9).

Then, the produced alkoxide anions (RO -) attack the carbonate function of EC to form the socalled intermediary A, followed by a reaction with a linear carbonate. Note that these reactions (Eq. V. The phosphate-containing compounds present in both electrolytes are the product of fluorine substitution reactions between anions (alkoxides, intermediary A.) and fluorophosphates (Eq.

V.15). The latter is originating from the degradation of LiPF6 into PF5 and LiF (Eq. I.10), followed by the reaction with water to yield POF3 (Eq. V.14, also I.11).

[V.14]

[V.15]
To sum up, in base electrolyte, anions originating from the reduction of solvents bring about the formation of ethylene oxide and ethylene carbonates oligomers as well as phosphate ending oligomers due to the reaction of PF5 with water.

In presence of LDFP electrolyte additive, the reaction of alkoxides RO -with LDFP must be kinetically favored, thus preventing the formation of the above-mentioned ethylene oxide and ethylene carbonates oligomers. The generated lithium organo-fluorophosphate salt must be precipitated in SEI, as suggested by FTIR analysis, or dissociated in electrolyte. Part of the alkoxides can also start attacking EC but all the resulting anions react with POF3 formed from the reaction of LiPF6 (PF5) with water, as in base electrolyte. Note that, with LDFP or not, the amount of phosphates detected through ESI-HRMS stems from water traces.

Analysis of gas released

The limited literature on the effect of LDFP on gas production includes the work by Song et al. 13 , who claimed that the combination of 0.5 wt.% of LDFP with 15 vol. of FEC (fluoroethylene carbonate) in a LiPF6-based electrolyte reduced the swelling ratio of a charged (4.25 V) pouch cell NMC 811/SiOx@C to 1.96% after 7 days storage at 55°C, in comparison to 5.91% when only using 15% FEC. Liu et al. 12 reported a decrease of ~78% and ~88% in the volume of pouch cells NMC 532/Gr containing 1-2% of LDFP after 500h storage at 60°C, both at the charged (4.4 V) and discharged state (2.5 V), respectively. From the same group, Ma et al. 9 reported a decrease of ~83% of the volume of pouch cells NMC 532/Gr containing 0.5 wt.% of LDFP stored in the charged state. According to Dahn's group 9,12 , the decreased gas production is due to "the suppression of parasitic reactions", caused by the passivation of the SEI by LDFP deposition, but the mechanisms are not detailed. The aforementioned works studied the volume change of stored pouch cells upon the formation cycles, but neither of them studied the elemental composition of the gases, nor their evolution after prolonged cycling. Therefore, this section reports on the gas content and composition analysis from cycled pouch cells NMC 811/Gr, in the hope of obtaining further information about the role of LDFP.

Two batches of four 811/Gr whose cycling results were previously shown (Section 3.1), were taken for gas analysis after 200 cycles at 55°C by means of GC/BID, following the preparation, gas recuperation and quantification protocols described in Chapter IV. Three possible explanations stand out:

i. The sources of CH4 and C2H6 are the linear carbonates EMC and DEC; whereas the source of C2H4 is the cyclic EC. Hence, linear carbonates would be continuously reduced along the cycling and not EC. This is not likely to occur since, according to the calculations made by Tasaki et al. 41 , EC is the carbonate solvent more prone to undergo the first electron reduction ahead of the other linear solvents.

ii. According to some works 42,43 , C2H4 is electrochemically reduced along cycling producing Li2C2H4. However, this reaction has been reported to occur only on solid argon 44,45 and thus, it cannot be envisaged in battery conditions.

iii. C2H4 may be oxidized by 1 O2. For comparison, in the last chapter, pouch cells NMC 811(non-cal.)/Gr at C/5 and 55°C produced 0.087 ± 0.019 µmol of C2H4 after the formation cycles, an almost 23 times bigger value than that found here, after the long cycling at C/4 and 55°C. It can be presumed that C2H4 has been oxidized by 1 O2 to ethylene oxide or CO2, as discussed in the previous chapter.

With respect to the quantification of CO2 released from pouch cells containing base electrolyte, it is remarkable that the quantity is almost three times bigger when NMC 811 is used compared to 622C. The thermal degradation of solvents cannot be at the origin of this difference as pouch cells spent an almost equivalent time span in the oven at 55°C (40 days for 622C/Gr and 41.5 days for 811/Gr).

Other CO2 sources are the oxidative processes of C2H4 and alkyl carbonate solvents with 1 O2.

As already discussed, the latter is supposed to be produced from layered NMC phase transformations (to spinel and rock salt) upon aging and the consumption of Li2CO3 through electrochemical oxidation or chemical reaction with LiPF6 (Chapter III); both processes being exacerbated in case of NMC 811 compared to 622.

On the other hand, the impact of LDFP is very clear: there is less CO2 produced in the order of ~50 and ~25% less in pouch cells containing NMC 811 and 622C, respectively, whenever LDFP is contained in the electrolyte. This huge decrease in case of NMC 811/Gr cells must be interpreted with caution as the cells displayed a premature capacity loss (PCL) in presence of LDFP and thus, might not have spent the same accumulated time at oxidative high voltages.

However, the cells with and without LDFP, for each NMC, spent exactly the same time span in the oven at 55°C, hence discarding the thermal degradation at the origin of the difference. It is interesting to note that the PCL seen with NMC 811 is not accompanied with an important CO2 gas release related to electrolyte oxidation. Therefore, it can be assumed that the usual NMC aging process that brings about oxygen releasing-structural transformations is not at the origin of this PCL phenomenon.

NMC 622C/Gr cells that were found to depict improved performances when cycled with LDFP at 55°C also release a less amount of CO2, which thus proves the efficiency of this additive to mitigate solvents degradation. In literature, the creation of a CEI at the surface of NMCs was suggested to explain the less reactivity of these interfaces toward electrolyte, what we were not able to confirm from FTIR analysis. Further investigations are required to determine the origin of such CO2 decrease in presence of this additive, through chemical simulation tests.

To conclude about the gas study, the similar amount of gases, C2H4, CH4 and C2H6, with and without LDFP, whatever the NMC composition, 811 or 622C, is in compliance with the previous suggestion that the additive does not hinder the solvent reduction processes. The decreased amount of C2H4 between the formation step and the 200 th cycles may be due to its oxidation with reactive singlet oxygen released from NMC transformation phases, or the oxidation of Li2CO3 from the SEI or the surface species at the surface of the raw NMC materials.

The presence of LDFP was found to inhibit the production of CO2 and maybe more importantly in case of NMC 811 than 622C, nonetheless, further investigations are required to explain this effect. Last but not least, a conclusion can be drawn that the PCL observed with NMC 811 pouch cycled at 55°C with LDFP does not involve or is not due to supplementary production of the gases hereabove studied.

Discussion and conclusion

The study of degradation products in the electrolyte was conducted with the aim of better grasping the contradictory role of the LDFP additive upon complete cells cycling at 55ºC, which is beneficial in case of NMC 622C and deleterious in case of NMC 811. To this end, a stepwise investigation of interfacial layers, electrolytes and gases has been achieved through FTIR, mass spectrometry and GC/BID or MS analysis, respectively.

It was determined that LDFP does not act as classical additives (VC, FEC…) since it does not reduce before alkyl carbonate solvents and thus, does not prevent the formation of lithium carbonate and lithium alkyl carbonate in the SEI of graphite powder. Only a marginal decrease in the reduction of solvents has been revealed by analyzing the electrochemical curves from pouch cells, but this phenomenon may result from a reaction between CMC binder and LDFP.

On the other hand, the additive is prone to scavenge alkoxides (RO -) formed from linear carbonate reduction, before they react with EC to yield ether/carbonate-based oligomers. The "scavenging" substitution reaction leads to the formation of lithium organo-fluorophosphates and LiF as potential additional salts in the SEI layer. This new SEI composition may confer higher passivating properties and thus contribute to the capacity retention improvement in presence of LDFP. This additive also decreases the CO2 gas release, which still needs conclusive explanations. Overall, all these analyses help clarify the beneficial role of LDFP.

However, they do not allow to conclude about the premature capacity loss of NMC 811/Gr pouch cells cycled at 55°C. Only few observations can be stated:

i. the positive electrode is the culprit,

ii. the cells do not generate a higher amount of gases upon cycling and thus, the classical aging processes inducing structural transformations at the surface of NMC might not be involved,

iii. the highest the LDFP concentration, the less pronounced the capacity fade and iv. the electrode calendering has a beneficial impact.

In light of all these results, the structural collapse of the NMC 811 particles evidenced by

Conclusion of the chapter

There are two main motivations to study LDFP as electrolyte additive in LIBs containing Nirich NMC materials. First, to take advantage of the in-situ reactions between surface species and LiPF6 that produce LDFP (Chapter III). Second, to test if LDFP improves the electrochemical performance and reduces gas production, as agreed in the literature.

In a first instance, it was found that LDFP does not behave as a typical additive that reduces/oxidizes to generate a SEI/CEI. In fact, virtually no difference in the CEI formed with/without LDFP was found and the only difference in the SEI is the enrichment with fluorine-and phosphor-based species when LDFP is used. The evaluation of the cycling performance revealed that LDFP improves the capacity retention after 200 cycles of cells NMC 622C/Gr cycled at 55ºC and NMC 811/Gr cycled at 25ºC, but intriguingly, PCL is observed at 55ºC. A possible explanation was given by TEM analyses that revealed the presence of an interphase in-between the primary grains of NMC 811, possibly made of soluble bases. In agreement with the study reported in Chapter III, we hypothesized that LiPF6 reacts with this interphase to produce LDFP, inducing grain decohesion and contact loss. The presence of LDFP additive in electrolyte favors its deposition and thus exacerbates this phenomenon, which ultimately causes the particle disintegration observed when the 811/Gr pouch cells were opened.

Regarding the controversy around the action mechanism of LDFP, it was found that it scavenges RO -, thus preventing electrolyte ageing. In addition, it inhibits some CO2 production in 622C or 811-containing cells, although this part merits further clarification.

These two characteristics are very interesting for the application of this additive in LIBs for the automotive industry, nonetheless, its implementation in cells containing polycrystal NMC 811 cycled at high temperature appears to be restricted by its surface chemistry. The use of monocrystal NMC can be envisaged so that the reactivity of NMC 811 is minimized.

General conclusions and perspectives of the work

This work has been dedicated to the surface reactivity of Ni-rich NMC materials and the understanding of the related degradation mechanisms. A focus was given to the reactions that lead to gas production upon high-temperature storage and to the study of the impact an electrolyte additive.

To begin with, this work introduced the main limitations of Ni-rich NMC materials associated with their reactivity: non-stoichiometric synthesis that leaves unreacted precursors in the form of soluble bases, the degradation in air that produces surface species and the electrolyte degradation upon battery functioning.

Concerning the soluble bases left from synthesis, we identified and accurately quantified, Regarding the degradation of layered materials in air, for years, only lithium-based surface species were assumed to be produced. In this work, the degradation as function of NMC composition was studied through an accelerated exposure process in specific atmospheres. As a result, a novel degradation mechanism upon exposure to H2O/CO2, that does not involve any reduction process, was proposed. We suggest that, in addition to Li2CO3, insoluble oxyhydroxide and oxide surface species of transition metals, are produced. Furthermore, the heaviest degradation was detected in the material with the highest nickel content, meaning that fresh Ni-rich NMC materials require more careful storage conditions than a typical NMC 111.

After the surface composition of our materials was exhaustively examined, their reactivity upon battery functioning was investigated. It was found that the presence of certain surface species, namely of Li2CO3, Li2O and M2SO4; could be advantageous for the electrochemical performances of the battery, since they produce LiPO2F2 (a known electrolyte additive) and LiF (a known stable passivating compound), by chemically reacting at high temperature with the electrolyte. As a perspective, these findings could be used to improve the industrial formation process of LIBs, in which by adjusting the storage temperature and time, an electrolyte additive would be generated in-situ. If not reacted, the risk is that Li2CO3 gets electrochemically oxidized into harming CO2 and 1 O2, at a potential dependent on the particle size of Li2CO3.

It is known that NMC materials that contain surface species often suffer from low electrochemical performances. In an effort to improve them, and comforted by the in-situ reactions evidenced by TG/MS analyses, a heat treatment at 500ºC under argon atmosphere was carried out on as-received, washed and exposed NMC 622C. First, it was found that the capacity retention of the NMC material was not improved by the elimination of most of the soluble bases through washing; this result could have been expected from the literature as the quantity falls below a puzzling experimentally determined equilibrium value 4 . Similarly, annealing the as-received and the washed materials did not improve them either. On the other hand, annealing helped recover some capacity of the exposed material. We suggest that this is due to the in-situ reaction between Li2CO3, oxygen and metal oxides, that decreased polarization and reformed an electrochemically-active layered phase. This would explain the literature about the improvement of the electrochemical performances after an annealing treatment at relatively low temperature.

Other consequence of the presence of surface species is slurry gelation. In our case, it was found that at a defined shear rate value, the viscosity of an as-received or of an exposed NMC remained the same, but at low shear rates it differed. With the knowledge of surface composition, a study devoted to slurries stabilities would be interesting to be undertaken.

Having found that many of the reactions of surface species produce gas, and knowing the safety risks and cycle life deterioration associated with gas production, the next step was to study the reactivity of Ni-rich NMC materials, from the analysis of gaseous products.

One of the goals of the thesis was to quantify permanent gases produced from batteries during storage and/or cycling. The GC-BID has demonstrated to be the best technique to achieve it.

The key point of the gases quantification is their separation which operates through the use a column. Unfortunately, in our case, the appropriated conditions (column, GC mode…) to separate all the permanent gases were not found, H2 and CO were missing. Although interesting results were obtained, this left an incomplete view of the total reaction mechanisms of the solvents degradation leading to gas formation, and did not allowed to estimate the total volume of produced gases. So, it seems important to pursue the work on separation optimization in the future.

As a preliminary stage to quantifying gas from pouch cells, storage tests of materials with electrolyte in specific containers were performed with the aim of assessing the impact of the NMC materials, the state of charge, the electrolyte salt and the dead volume on CO2

production. Beforehand, it was found that the CO2 produced from the electrolyte storage at 55 ºC, was driven by PF5 and thus was dependent on the dead volume of the container. Upon storage of pristine NMC powders with electrolyte, surprisingly, lower-nickel materials produced larger amount of CO2, suggesting different electrolyte absorption properties.

Afterwards, it was proven that the storage of delithiated NMC films increases CO2 production, and was dependent of the nickel content. We inferred that this occurs through the oxidation of EC by 1 O2.

The analyses in optimized 2-and 3-electrode pouch cells showed that the inside of the battery is a complex and dynamic system, with reactions of gas generation, and surprisingly, also of gas consumption. Through the use of a simplified storage protocol, it was possible to evidence that CH4, C2H4, C2H6, from the formation step, are dependent mainly on current density.

Additionally, several sources of CO2 were evidenced (but not differentiated): chemical and electrochemical oxidation of solvents and Li2CO3, and thermal degradation of solvents. After storage, two phenomena were observed: C2H4 consumption and an increase of CO2 concentration. Both are issued from the same cause: oxidation of C2H4 and EC by 1 O2, influenced by the composition of NMC, due to more 1 O2 being produced from higher nickel compositions. At the same time, the layered phase becomes rock salt phase, triggering impedance augmentation. This was evidenced in 3-electrode pouch cell experiments, but a meticulous TEM study to link the surface reactions to the suspected structural transformations would be ideal to confirm it.

These conclusions convey the key message that there is still room for understanding the crosstalk reactions and the mechanism through which 1 O2 oxidizes C2H4 and EC. The work on gas evolution constitute the bases for following research.

As stated at the beginning of this work, the use of electrolyte additives is one of the most effective and economical ways to improve the electrochemical performances of LIBs containing Ni-rich NMC materials. LiPO2F2 (LDFP) was chosen as electrolyte additive after we found that it is produced from the in-situ chemical reactions of certain surface species.

Moreover, the literature agrees that it improves the cycling performances and decreases the gas production; although its action mechanism remains a controversial point.

It was found that LDFP does not behave as a typical additive that reduces or oxidizes in the first cycles and also does not hinder the solvent reduction processes. In fact, virtually no differences in the CEI formed with/without LDFP were found and the only difference in the SEI, is the enrichment with fluorine-and phosphor-based species, when LDFP is used.

Regarding the action mechanism, electrolyte analyses revealed that LDFP scavenges RO -, thus preventing electrolyte ageing. However, determining its exact role on CO2 production in NMC 622C-and 811-containing cells, needs further investigation.

Moreover, it was shown that LDFP exhibits a contradictory role upon complete cells cycling at 55ºC: beneficial in case of NMC 622C and deleterious in case of NMC 811. Upon opening of the 811/Gr cells, broken particles are observed, mostly when LDFP is used. TEM analyses helped us to hypothesize that chemical reactions occurred on the interphase in-between the primary grains of NMC 811, possibly made of soluble bases, inducing grain decohesion and contact loss. This phenomenon is exacerbated due to the increased LDFP concentration and eventually causes particle cracking. Therefore, we conclude, that the use of LDFP in cell containing NMC 811 is restricted by its surface chemistry. In this regard, the use of monocrystal NMC can be envisaged to minimize these reactions. Xiong et al. 5 have already reported reduced reactivity from NMC 532 toward the electrolyte, although NMC with higher nickel content has not been tested yet.

Lastly, the results in this thesis not only open up lines of research that might be pursued in this research group, but also helped to enrich the understanding of Ni-rich NMC materials for its implementation in the automotive industry.

VIII

poly(dimethylsiloxane)). Helium was used as carrier gas and was maintained at a constant flow rate of 1.3 ml.min -1 . The oven program starts at 40 ºC for 6min, then ramps at 10 ºC.min -1 up to 90 ºC, increased at 5 ºC.min -1 up to 190 ºC, held for 5 min, increased at 10 ºC.min -1 up to 250 ºC, where it was held for 10 min (Fig. A.6). The total duration of GC analysis was 52 min.

The transfer line was maintained at 250 ºC in the GC/MS-IR and the ion source was set at 200 ºC. In the GC/BID the plasma temperature was set at 280 ºC and the injector port at 150 ºC. minutes of resting time in between measurements. Then, the temperature was increased to 25 ºC, with a space of 30 minutes in between, to check that the electrolyte has not changed its properties; and then the measurements were taken every 5 ºC.

For the measurements in separator, pouch cells with a Celgard 2500 separator sandwiched between two symmetrical electrodes of copper of 4cm 2 were constructed. 1 h of impregnation was allowed after adding the electrolyte and sealing under vacuum.

• Graphite composite preparation and cycling

The graphite powder (Timcal TIMREX SLS, ~1.5 m 2 .g -1 ) was grounded with conductive carbon black C45 (20nm and BET = 45m 2 .g -1 , Sigma Aldrich) in a 9:1 mass ratio to form conductive composites. Half-cells Swagelok-type were assembled following procedure in Annex II. Around 11mg of the composite were used. One cycle with potential limitation was performed using a VMP-3 potentiostat/galvanostat (Biologic SA). The potential window was 1.5 -0.01 V at C/20 and 25 ºC or 55 ºC.

• FTIR analyses in transmission mode

The composite powders were recovered following the protocol of Annex II. Then, KBr pellets were prepared inside the glovebox by grounding in a mortar the dry composite and the dry KBr, followed by analysis in a FTIR device (Nicolet iS20) with minimal air exposure.

XX

!"l" 2.#3.y (1zzz) l" 3 (1zz) /9 3 (1z>) $7 y (1zzz) % # { ? | g| ? ⁄ h⎯⎯⎯⎯⎯j 1 2 ⁄ !" # $% & + (l" 2.#3.y $7 y ) (1zzz) %%) + 0l" (1zz) % + 0/9 (1z>) % # Ayant constaté que de nombreuses réactions des espèces de surface produisent du gaz, et connaissant les risques pour la sécurité et la détérioration de la durée de vie associés à la production de gaz, l'étape suivante a consisté à étudier la réactivité des matériaux NMC riches en Ni, à partir de l'analyse des produits gazeux.

XXII

Au moyen d'un appareil GC/MS-IR et d'une colonne capillaire, les principaux gaz identifiés à partir de nos batteries étaient CH4, CO, C2H4, C2H6 et CO2. Étant donné que la quantification est limitée dans le cas de l'appareil MS par la co-élution de certains gaz et que la précision des pics obtenus avec l'appareil IR n'est pas optimale, un appareil GC/BID a été utilisé. Cela a conduit à deux résultats principaux : (i) avec la colonne capillaire, un constituant du pic coélué a été séparé par l'amélioration des conditions expérimentales (Fig. 6a). (ii) avec la colonne « micropacked », le H2 a été détecté et le pic co-élué a été complètement séparé (Fig. 6b). Malgré l'excellente séparation, celle-ci nécessite encore une amélioration avant d'être utilisable dans des conditions acceptables.

XXV Résumé

Ce travail de thèse a porté sur la compréhension des mécanismes de dégradation de surface ainsi que sur l'étude de la réactivité des matériaux NMC riches en nickel, en mettant l'accent notamment sur la compréhension des phénomènes générant des gaz.

Dans un premier temps, la quantification des espèces solubles, laissées lors la synthèse ou produites par l'exposition à une atmosphère spécifique, a été réalisée par titrage mais également par des techniques complémentaires telles que l'AAS, l'ICP-OES et la spectroscopie infrarouge. Cette étude comparative révèle la réactivité de cinq matériaux NMC commerciaux. Les analyses TG/MS de ces matériaux ont permis de proposer un nouveau mécanisme qui explique (i) la dégradation des matériaux NMC riches en nickel au contact de l'air, produisant des espèces de surface solubles et insolubles; et (ii) la récupération partielle des performances électrochimiques des matériaux exposés, après un processus de recuit. Il a été prouvé que certaines de ces espèces présentent à la surface de NMC réagissaient chimiquement avec les composants de l'électrolyte pour produire le sel LiPO2F2.

Dans le but de quantifier les produits gazeux, préalablement, la mise au point d'un protocole de séparation a été réalisé à l'aide d'un appareil GC/BID. L'analyse des gaz produits issus d'un protocole de stockage défini, a mis en évidence la consommation de C2H4 et les multiples sources de CO2; les deux ayant en commun, la réactivité du matériau NMC délithié. Enfin, étant produit à partir de réactions chimiques in situ et décrit dans la littérature comme un additif d'électrolyte efficace, le sel LiPO2F2 a été évalué comme solution pour améliorer les performances électrochimiques des cellules à base des matériaux NMC riches en nickel. L'analyse des produits de dégradation solides, liquides et gazeux après le cyclage a permis de clarifier le mécanisme d'action par lequel cet additif peut améliorer la rétention de capacité et réduire le gazage. De manière surprenante, une perte de capacité prématurée a été observée lors du cyclage des cellules contenant du NMC 811, à 55 ºC. Grâce à des analyses en microscopie électronique, l'hypothèse d'une désintégration des particules, causée par les réactions entre l'électrolyte et l'interphase des joints de grains a été émise.

Mots-clés : NMC riches en nickel, bases solubles, stockage de batteries, génération de gaz, LiPO2F2

Abstract

This work was dedicated to the understanding of the surface degradation mechanisms and reactivity of Ni-rich NMC materials, with a focus on the reactions that generate gas.

In a first approach, the exact nature and accurate quantification of the water-soluble species left from synthesis or produced from specific atmosphere-exposure were investigated through complementary techniques as titration, AAS, ICP-OES and FTIR. This comparative study unveils the reactivity of five commercial NMC materials, as function of the increasing nickel content. TG/MS analyses of these materials allowed the proposition of a novel mechanism that explains (i) the degradation of Ni-rich NMC materials toward air, producing soluble and insoluble surface species; and (ii) the partial recovery of the electrochemical performances of exposed materials, after annealing at relatively low temperature. Some of those NMC surface species were proved to chemically react with the electrolyte components into LiPO2F2.

With the aim of quantifying gaseous products, first, different separation conditions were studied by making use of a GC-BID apparatus. The analysis of gases produced from a defined storage protocol, evidenced the consumption of C2H4 and the multiple sources of CO2; both having in common, the reactivity of delithiated NMC. Lastly, being produced from in-situ chemical reactions and described in the literature as an efficient electrolyte additive, LiPO2F2 was evaluated as a solution for enhancing electrochemical performances of Ni-rich NMC materials-based cells. The analysis of solid, liquid and gaseous degradation products after cycling helps clarify the action mechanism through which it can improve the capacity retention and seemingly reduces gassing. Surprisingly, a premature capacity loss was observed when cycling NMC 811-containing cells, at 55 ºC. Through electron microscopy analyses, it was hypothesized that it was due to the particle disintegration, caused by the reactions between the electrolyte and grains boundaries interphase.

Key-words:

Ni-rich NMC, soluble bases, battery storage, gas generation, LiPO2F2

  (Fig. I.1). During the charge, this process is inversed: the anodic reaction takes place at the positive electrode and the cathodic reaction at the negative electrode. Henceforth, the negative electrode is referred as anode and the positive electrode as cathode in the discharged state.

Figure I. 1 .

 1 Figure I.1. Scheme of a battery during discharge. Based on Scrosati et al. 1 .

Figure I. 2 .

 2 Figure I.2. Scheme of a LIB. Adapted from Abraham et al. 13 .

  With the aim of comparing the battery technologies explored during the development of electric vehicles, Fig. I.3 shows these batteries evaluated by size (volumetric density) and weight (energy density). Heavy and bulky Pb-acid and Ni-Cd are not anymore adapted for the requirements of modern electric vehicles, whereas Ni-MH batteries are still used in some

Figure I. 3 .

 3 Figure I.3. Specific density vs. energy density in different batteries. Adapted from Scrosati et al. 1 and updated with references 17 and 21.

Figure I. 4 .

 4 Figure I.4. Global EV market share and sales growth from 2013 to 2018. Taken from the Global EV Outlook 2019 report of the IEA 18 .

. 5 )

 5 in 2030, compared to ~$175 in December 2019, as a result of a decrease of the prices of raw materials and manufacturing.

Figure I. 5 .

 5 Figure I.5. Prediction of the price drop of LIBs per kWh. Taken from Marinaro et al. 21 .

Figure I. 6 .

 6 Figure I.6. Targets of LIBs characteristics at the cell level in 2015, 2022 and 2025. Taken

Figure I. 7 .

 7 Figure I.7. Scanning electron microscope (SEM) image of a Celgard® PP separator. Taken from Yang et al. 49 .

Figure I. 8 .

 8 Figure I.8. Representation of electrode composition. Adapted from Park et al. 11 .

Figure I. 9 .

 9 Figure I.9. Molecular representations of a) PVdF, b) SBR, c) CMC and d) PTFE.

Figure I. 10 .

 10 Figure I.10. Side view of lithiated graphene layers. Taken from Park et al. 11 .

Figure I. 11 .

 11 Figure I.11. Schematic galvanostatic curve and staging effects during lithium intercalation between graphene layers. Taken from Park et al. 11 .

Figure I. 12 .

 12 Figure I.12. Peled's schematic representation of the SEI arrangement 62 .

  Fig. I.13.

Figure I. 13 .

 13 Figure I.13. Approximate range of average discharge capacity and potential of some of the most common active materials for the negative electrode. Taken from Chagnes et al. 56 .

Figure I. 14 .

 14 Figure I.14. Crystal structure of layered materials LiMO2. TM stands for metals or transition metals. Adapted from Liu et al. 75 .

Fig. I. 15 .

 15 Fig. I.15. positions these NMC compositions in the ternary diagram LiCoO2-LiMnO2-LiNiO2

Figure I. 15 .

 15 Figure I.15. Ternary phase diagram for the LiNi1-x-yMnxCoyO2 materials with their main advantages and disadvantages shown in green and red colors, respectively. Adapted from Schipper et al. 100 .

Figure I. 16 .

 16 Figure I.16. Comparison of capacity retention and discharge capacity of several NMC materials in the voltage window 3-4.3V and 0.1 C, after 100 cycles. Adapted from Noh et al. 101 .

Figure I. 17 .

 17 Figure I.17. Schematic representation of the degradation of layered materials in air. Taken from Cho et al. 117 .

  6 of extracted lithium) and subsequent oxygen production can be explained with their electronic diagram (Fig. I.18). In the fully lithiated state, the energy bands of the transition metal redox couples, eg of Ni 3+/4+ and t2g of Co 3+/4+ , are located near the top of the O 2-2p band, while Mn 3+/4+ lies well above all of them. When the material is delithiated and most of the nickel is Ni 4+ , the Fermi levels eg and t2g lowers, overlapping the band 2p of oxygen O 2-. Consequently, electron transfer occurs from the oxygen to the Ni 4+ , reducing it to a more stable oxidation state Ni 3+ and Ni 2+ . In regards to the oxygen, due to the deficiency of lithium and the electron transfer to Ni 4+ , the M-O bond covalency increases, decreasing the O 2-charge density and favoring oxidation of the oxygen anions from O 2-to a lower charge density. At the same time, the repulsive interactions among the negatively charged oxygen layers, increases the crystallographic parameter c and facilitates the oxygen release 122 .

Figure I. 18 .Figure I. 19 . 3 . 3 . 2 .

 1819332 Figure I.18. Schematic illustration of the redox couples positions relative to the top 2p band of the oxygen in NMC materials. Adapted from Choi et al. 123 .

Figure I. 20 .

 20 Figure I.20. Schematic representation for the phase transitions driven by cation migration.

(

  HRTEM) has been extensively used 126, 127 . As example, Fig. I.21 shows NMC 532 surfaces cycled with different cut-off voltages, 4.5 V (a) and 4.8 V (b). It can be seen that the rock salt phase is favored when the cut-off voltage is higher and that the structural transformations does not exceed 30 nm 127 .

Figure I. 21 . 3 . 3 . 4 .

 21334 Figure I.21. HRTEM examples of phase transformations in NMC 532 particles after 50 cycles between 3-4.5 V a) and 3-4.8 V b). Taken from Jung et al. 127 .

Figure I. 22 . 3 . 4 .

 2234 Figure I.22. TEM images of an NMC 622 particle before and after 150 cycles at 60 °C and 0.5C, between 3-4.45V vs. Li/Li + . Taken from Kim et al. 131 .

  interphase (CEI) (Fig. I.23) 138 .

Figure I. 23 .

 23 Figure I.23. Schematic representation of the species deposited as CEI and the structural reconstruction in the vicinity. Adapted from Li et al. 138 .

Figure I. 24 .

 24 Figure I.24. Illustration of the reaction of SB with H3PO4 to form Li3PO4. Taken from Jo et al. 160 .

  4 and their structures are shown in Fig. I.25.

Figure I. 25 .

 25 Figure I.25. Examples of electrolyte additives that have impacted positively the performance of LIBs containing NMC as positive electrode. a) SN, b) ES, c) TMS, d) LiBOB e) VC and f) FEC.

Figure II. 1 .

 1 Figure II.1. Representative SEM pictures of Ni-rich NMC materials: a) 532, b) 622C, c) 622, d) 811 and e) 901.

Figure II. 2 .

 2 Figure II.2. SEM pictures of the alumina coating in NMC 622C, at a) 25000X and b) 50000X magnification. c) EDX point analysis showing aluminum presence.

Figure II. 3 .

 3 Figure II.3. Powder XRD patterns of the five Ni-rich NMC materials.

3. 1 . 3 . 1 . 1 .

 1311 Quantification of soluble bases through acid-base titration technique Optimization of the acid-base titration protocol The protocol for the titration of soluble bases includes a solubilization step (NMC washing) which consists in stirring the powder material for a certain time in water, followed by vacuum filtration in order to recover separately the washed powder and the solution containing the soluble bases (Fig. II.4). Then, the filtrate solution containing the soluble bases is titrated with a strong acid.

Figure II. 4 .

 4 Figure II.4. Schematic diagram of the protocol for the soluble bases solubilization and analysis

3. 1 . 1 . 1 .

 111 Impact of the atmosphere composition It is known that the filtrated solution of CO3 2-and OH -resulting from the washing procedure can be acidified with CO2 from the atmosphere 41 . The following experiments show the range in which the pH of an aqueous solution containing dissolved Li2CO3 varies while being stirred under argon or air atmosphere. For this, 0.18 mmol of Li2CO3 (Sigma-Aldrich, purity > 99%) were added to 100 ml of bi-distilled water under air or under inert atmosphere (argon), and the pH was immediately recorded as function of time by means of an automatic titrimeter equipped with a pH electrode (TitroLine 7000, SI Analytics GmbH) (Fig. II.5).

Figure II. 5 .

 5 Figure II.5. Impact of air (red curve) or argon (blue curve) atmosphere on pH variation as function of time of a solution of Li2CO3 dissolved in bi-distilled water.

  30 minutes under argon atmosphere. The pH was recorded as a function of time, immediately after the introduction of the NMC powder in water (Fig. II.6).

Figure II. 6 .

 6 Figure II.6. a) Variation of pH values over time for a suspension of 5 g in 100 ml of bi-distilled water for NMC 532, 622C, 811 and 901. Amounts of OH -and CO3 2-determined by acid-base titration of filtrate solutions at different stirring times for b) 532, c) 622C, d) 811 and e) 901.

Figure II. 7 .

 7 Figure II.7. Acid-base titration curve (black) with two equivalence points evidenced by the derivative curve (red), including dosing reactions.

  II.8a) remains unchanged. Moreover, through XRD, it was confirmed that the layered F-3m bulk structure (Fig. II.8b) was not affected. To be concise, only the results from NMC 622C are shown.

Figure II. 8

 8 Figure II.8. a) Representative SEM picture of the alumina coating on the as-received NMC 622C and b) powder XRD patterns of as-received and washed 622C.

  Fig. II.9a and b show the titrated amount of OH -and CO3 2-as a function of the powder-to-water mass ratio. The error bars were constructed from the standard deviation from three repetitions of each mass ratio and each material. The linear fit of every material data indicates that all of tested mass ratios would be adequate in terms of soluble bases solubilization. However, for the mass ratios of 1:100 and 1:50, bigger standard deviations are found, so they must to be avoided to ensure reproducibility.

Figure II. 9 3 . 1 . 1 . 4 .Table II. 4 .

 931144 Figure II.9. a) OH -and b) CO3 2-amount in different powder-to-water mass ratios.

  50 ml of the filtrate solution are weighed inside the customized closed flask of the automatic titrimeter, with three entrances for the pH electrode, the titrant solution of HCl 0.05 M and the argon bubbling. The acid-base titration analyses performed in dynamic titration mode ends when the pH reaches 3. The amount of CO3 2-and OH -is calculated from Eqs. II.19 and II.20.

  detected in noticeable concentrations. The synthesis of the transition metal precursors of NMC materials (M(OH)2 and MOOH) requires the precipitation of MSO4 (M= Ni, Mn and Co) by NaOH in the presence of ammonia solution as complexing agent 2, 24 , consequently, sodium-based species might remain in the surface of the final materials. For the sake of comparison, Fig. II.10 shows the quantification of lithium and sodium obtained by AAS and acid-base titration. Assuming that all soluble bases detected by acid-base titration were lithium salts, their concentrations were converted into lithium concentrations with the respective molar equivalences of 2 and 1 for Li2CO3 and LiOH.

Figure II. 10 .

 10 Figure II.10. Acid-base titration and AAS results comparison of lithium and sodium elements contained in the filtrate solutions.

  Fig. II.11 shows the acid-base titration, AAS and ICP-AES results obtained from the filtrate solutions of each NMC composition. ICP-AES analyses unveil noticeable concentrations of lithium, sodium and sulfur elements. The lithium concentrations determined by ICP-AES, are ranged within the standard deviation of the AAS data, corroborating

Figure II. 11 . 3 . 2 . 3 .

 11323 Figure II.11. Acid-base titration (ABT), AAS and ICP-AES results comparison of lithium (gray bars), sodium (red bars) and sulfur (blue bars) contained in the filtrate solutions.

Figure II. 12 .

 12 Figure II.12. FTIR-ATR (Ge crystal) spectra of as-received 532, 622, 622C, 811, 901 NMC; and Li2SO4, Li2CO3 and LiOH reference compounds.

Figure II. 13

 13 Figure II.13. a) and e) SEM images in backscattered (BSE) mode of the as-received NMC 532 material and their corresponding Na (b and f), S (c and g) and Ni (d and j) EDX mappings.

4. 1 .

 1 Preliminary storage studies under Ar, CO2, dry air and air atmospheres 10g of the five Ni-rich NMC materials were stored in open Petri dishes, for two weeks under four different atmospheres:

  transferred without air exposure, to the glove box; whereas the powder stored in air was vacuumdried in a Büchi oven for 24h at 120°C prior to being transferred to the glove box. Each powder was gently milled in a mortar inside the glove box to homogenize the sample, then weighed for the soluble bases quantification by acid-base titration (Fig. II.14a-e), following the optimized conditions reported in Section 3.1.1.4.

Figure II. 14 .

 14 Figure II.14. Acid-base titration results of NMC materials a) 532, b) 622C, c) 622, d) 811 and e) 901; in the as-received form and after storage in different atmospheres for two weeks.

4. 2 . 4 . 2 . 1 .

 2421 Storage under specific H2O/CO2 atmosphere 10g of each NMC material were stored during six days at 35°C, inside a desiccator with a controlled atmosphere of 8/92 molar ratio of CO2/Ar and with 90% of relative humidity. After exposure, the samples were vacuum-dried in a Büchi oven for 24h at 120°C, prior to being transferred to the argon-filled glove box. These materials stored in the atmosphere H2O/CO2 are called exposed NMC materials. Acid-base titration quantification after H2O/CO2 storage The amounts of soluble bases detected by acid-base titration in the exposed NMC materials are compared to the as-received and displayed in Fig. II.15a and b. In the as-received materials, mainly OH -species are detected. After exposure, OH -is detected only in the case of the NMC materials containing highest Ni contents, 811 and 901, and its concentration does not exceed 30 µmol.gNMC -1 . It is worth noting that the 811 values are close to the ones reported in literature for a 811 material (diameter ~25µm) exposed for 7 days in air, as being 82 and 257 µmol.gNMC -1 for

Figure II. 15 .

 15 Figure II.15. Amounts of OH -and CO3 2-determined by the analysis of the filtrate solution from the a) as-received and b) exposed (H2O/CO2 atmosphere for 6 days) NMC materials.

2 .

 2 Further investigations through SEM and FTIR The prominent delithiation process that happened upon H2O/CO2 storage, certainly induces surface textural modification, clearly observed through SEM analysis (Fig. II.16a-o). In the exposed NMC 532 (Fig. II.16b, c), almost no noticeable soluble bases growth took place, in comparison to the as-received material (Fig. II.16a). This result agrees with the acid-base titration results of the exposed NMC 532, that showed only 0.63 mol% of reacted lithium. In the exposed NMC 622C and 622, where heavier delithiation occurs, the growth of flakes (up to 4 µm) can be observed at the outermost surface of the materials (Fig. II.16e, f, h, i). Astonishingly, exposed NMC 811 and 901 present substantial surface textural modifications. The flakes form a shell-like layer that covers, almost completely, the granular morphology (Fig. II.16k, l, n, o). Some particles depict a non-uniform shell (denoted with ellipses) which we suppose is the result of an uneven exposure to H2O/CO2.This SEM investigation visually shows the extent to which the surface is affected by the H2O/CO2 exposure, with a clear dependence on the Ni content. It is supposed that the degradation reactions start at the surface of the material and form a thin film, as it has been observed before25, 28, 44 , then, once the products are crystallized enough, flakes become visible. When these reactions are favored due to the high reactivity of the material, the flakes form a thick layer that tends to detach from the NMC surface, as in the case of NMC 811 and 901. Interestingly, Cho et al.50 performed time-of-flight secondary ion mass spectroscopy (ToF-SIMS) analysis in an airexposed NMC 730, and revealed that the Li2CO3 layer lies above the LiOH layer and both of them are situated above the granular morphology of NMC 730.

Figure II. 16 .

 16 Figure II.16. Representative SEM images of as-received and exposed 532, 622C, 622, 811 and 901 materials for 6 days under H2O/CO2 atmosphere. Arrows point Li2CO3 flakes growth.

  Fig. II.17 shows the evolution of Li + amount as function of the square root of the storage time. It can be observed than the production of soluble bases follows a linear relationship up to 36 days of storage, confirming the kinetics observed by Chen et al.25 , who stored an NMC 622 up to 42 days in air at 55°C and 80% RH. However, from there, the linear relationship is lost and the degradation rate decreases owing to the difficulty to extract deeper lithium. At 120 days of 622C exposure, the reacted lithium amount reaches an astonishing percentage of 7.2 mol%.

Figure II. 17 .

 17 Figure II.17. Delithiation evolution determined by acid-base titration vs. the square root of the storage time in H2O/CO2 atmosphere.

Figure II. 18 .

 18 Figure II.18. FTIR-ATR spectra of the exposed NMC materials stored for 6 days in H2O/CO2 atmosphere.

  Fig. II.19a-d display the TG and CO2 (m/z 44), O2 (m/z 32) and H2O (m/z 18) MS curves for each as-received material.

Figure II. 19

 19 Figure II.19. a) TGA curves of as-received NMC materials and mass traces of b) CO2 (m/z 44), c) O2 (m/z 32) and d) H2O (m/z 18).

  Fig. II.20a).

Figure II. 20

 20 Figure II.20. a) TGA curves of exposed (H2O/CO2 for 6 days) materials and mass traces of b) CO2 (m/z 44), c) O2 (m/z 32) and d) H2O (m/z 18).

Figure II. 21 .

 21 Figure II.21. TGA-MS curves of NiCO3•2Ni(OH)2•1H2O reference.

Fig

  Fig. II.22 shows the comparison of the 6-days exposed NMC 811 before and after washing. Up to

Figure II. 22 .

 22 Figure II.22. TGA curves along with mass traces of H2O (m/z 18), O2 (m/z 32) and CO2 (m/z 44) of an NMC 811 exposed (H2O/CO2 for 6 days) and exposed + washed.

  . II.23a) disappears after washing (Fig.II.23b). This means that this layer is composed mainly of Li2CO3 and that insoluble (Ni1-2x-yCoy)OOH, NiO and MnO2 species are located into the surface degradation layers at the extreme surface of NMC particles. This is in accordance with the presence of NiO of rock-salt structure and other crystalline structures detected at the surface of air-exposed materials

Figure 2 . 23 .

 223 Figure 2.23. Representative SEM images of a) exposed (H2O/CO2 for 6 days) and b) exposed 811 after washing.

Figure II. 24 .

 24 Figure II.24. Raman spectra of the exposed NMC 901 (H2O/CO2 for 6 days) and several references.

Figure II. 25 .

 25 Figure II.25. Powder XRD patterns of the as-received, exposed 901, and the Li2CO3 reference.

Fig. III. 1

 1 Fig. III.1 shows the 19 F NMR spectrum of Li2CO3, Li2O and Li2SO4 at 25 and 55°C. In all cases, the doublet signal corresponding to LiPF6 is centered at δF = -74.2 ppm (J(P,F) = 709 Hz). In some cases (Fig. III.1a, b, d and e), another doublet centered at a lower chemical shift (δF = -84.6 ppm, J(P,F) = 929.4 Hz) can be attributed to LiPO2F2 (LDFP) 11 . For simplicity, the spectra were cut at 95 ppm, but upon enlarging to 160 ppm, the peak of LiF is also clearly observed at δF = -155.8 ppm (exemplary shown for Li2CO3 at 55°C in Fig. III.2). Note that, unlike Li2CO3, Li2O, LiOH and Li2SO4 reference compounds, NiO did not reveal any reactivity toward the electrolyte, whatever the temperature.

Figure III. 1 .

 1 Figure III.1. Liquid 19 F NMR spectra of the electrolyte recovered after storage, for 4 days at 25 and 55°C, in presence of a, b) Li2CO3, c, d) Li2O and e) Li2SO4 at 55°C; with LiPO2F2/LiPF6 molar ratio (R) determined from the integration of the peaks.

Figure III. 2 .

 2 Figure III.2. Liquid 19 F NMR complete spectrum of the electrolyte recovered after the storage of Li2CO3 and electrolyte for 4 days at 55°C.

  LiOH, Fig. III.3 shows the 19 F NMR results of its storage with electrolyte at 25 (a)

Fig

  Fig. III.1a, b • OPF(OR)(OR') with R = CH3 or C2H5 and R'= H, CH3 or C2H5 -a doublet centered at δF = -81 ppm (J(P,F) = 944 Hz), labelled as 2 -a doublet centered at δF = -81.3 ppm (J(P,F) = 943 Hz), labelled as 3 • OPF2(OH) with a doublet centered at δF= -85.2 ppm (J(P,F) = 950 Hz), labelled as 4.

Figure III. 3 .

 3 Figure III.3. Liquid 19 F NMR spectra of the electrolyte recovered after storage for 4 days in presence LiOH at a) 25 and b) 55°C. Label 1, 2, 3 and 4 stands for the doublets 1) OPF(OH)2; 2 and 3) OPF(OR)(OR') with R = CH3 or C2H5 and R'= H or CH3 or C2H5 and 4) OPF2(OH), respectively.

Figure III. 4 .Figure III. 5 .

 45 Figure III.4. 19 F NMR spectrum of electrolyte after storage with LiOH at 55°C for 4 days.

8 ]

 8 This result explains the absence of the LiOH band in FTIR spectra (Fig. II.12, Chapter II) and the absence of H2O release around 500°C in TGA/MS profile (Fig. II.19 a and d, Chapter II).

from 10 ±

 10 3 µm to 7 ± 5 µm (Fig. III.6a and b).

Figure III. 6 .

 6 Figure III.6. SEM images of a) pristine and b) ball-milled Li2CO3.

Figure III. 7 .

 7 Figure III.7. Incremental capacity curves (5 mV increment) of a) the first galvanostatic charge of Li2SO4 and pristine and milled Li2CO3 and b) first six cycles of the milled Li2CO3, all measured in half cells Li2CO3-or Li2SO4-C45/Li at C/40 and 25°C.

First, the

  growth of carbonates was evidenced by FTIR-ATR analysis of the pristine and airexposed NMC 811 and 901 films (Fig. III.9). The spectra were normalized by the SO4 2-band centered at 1170 cm -1 , similar to what has been done in Section 4.2.2 of Chapter II. The enlargement of the CO3 2-peak located at 878 cm -1 denotes the increase of carbonate content, and thus, confirms the surface degradation of NMC materials after only 15 minutes in air.

Figure III. 9 .

 9 Figure III.9. FTIR-ATR spectra of the pristine and air-exposed film of NMC a) 811 and b) 901.

Figure III. 10 .

 10 Figure III.10. Cycling performance of as-received and 15 min air-exposed non-calendered films of NMC a) 622C, b) 811 and c) 901; between 2.8-4.3V vs. Li/Li + at C/10 (rate considering a total delithiation) and 25°C.

Figs

  Figs. III.11a and b. show the first and second galvanostatic cycle of the five as-received materials.

Figure III. 11

 11 Figure III.11. a) First and b) second charge-discharge cycle of each Ni-rich NMC material, between 2.8-4.3V vs. Li/Li + at C/10 (rate considering total delithiation) and 25°C.

Figure III. 12 .

 12 Figure III.12. Incremental capacity profile of the first charge of each material.

Fig

  Fig. III.13a shows the potential profile versus capacity curves for the first cycle of the exposed NMC 622C. For comparison purposes, the as-received material is also shown.

Figure III. 13 3 . 2 . 4 .

 13324 Figure III.13. a) Potential profile in half-cell of the first galvanostatic cycle of the as-received and exposed NMC 622C materials. b) Corresponding incremental capacity profiles.

Figure III. 14 .

 14 Figure III.14. FTIR-ATR spectra of the exposed NMC 622C powders.

  14) that renders the polymer more rigid. It is possible that PVDF also reacts with Li2O. Even if it is less soluble in NMP than LiOH, trace of water in NMP would turn it into LiOH. -(CH2-CF2)-+ xOH -à -(CH=CF)-+ xHF + xOH - [III.14] Given this reaction, a decision was made to adapt the choice of the material tested to evaluate the role of this potential surface species. Hence, NMC 811 was chosen as LiOH was detected after H2O/CO2 exposure in an amount of 25 µmol.gNMC -1 . Its potential impact on the rheological properties of slurries was studied at 25°C with a rotational rheometer Kinexus Lab+ (Malvern) employing planar plate geometry. For comparison purposes, Fig. III.15 shows the viscosity versus shear rate plot of both as-received and 6-day exposed NMC 811. Both

  increases and may cause gelation or flocculation. In the 6-day exposed NMC 811, only 25 µmol.gNMC -1 of OH -were quantified; a too low amount to induce gelation. In the as-received material, 214 µmol.gNMC -1 of OH -were found, a sufficient amount for visible viscosity increase at every shear rate, which is not observed in Fig. III.15. This result confirms the existence of Li2O instead of LiOH, as proposed in Chapter II.

Figure III. 15 .

 15 Figure III.15. Viscosity versus shear rate plot of two slurries. The dashed gray line signals the shear rate of 34 s -1 corresponding to that induced by the coating of the slurries onto an Al foil.

4. 2 .

 2 Fig. III.16 and the quantities of CO3 2-and OH -from Li2O are written on the inset table.

Figure III. 16 .

 16 Figure III.16. Discharge capacity retention curves of the R and A-R NMC 622C/Li half-cell at C/10 between 2.8-4.3 V. Inset: table with the quantities of soluble bases obtained from acid-base titration.

Figure III. 17 .

 17 Figure III.17. Discharge capacity retention curves of the as-received and as-received + annealed NMC 622C/Li half-cell at C/10 between 2.8-4.3 V. Inset: table with the quantity of soluble bases obtained from acid-base titration.

15 ] 4 . 4 .

 1544 Impact of heat treatment on exposed NMC 622C Upon 20-day exposure to specific H2O/CO2 (Chapter II), the totality of Li2O contained in asreceived NMC 622C transformed into Li2CO3 and through further degradation, almost ten times more Li2CO3 than in the as-received material was created (inset in Fig. III.18), equivalent to an astonishing value of 5.3 mol% of reacted lithium. Due to this heavy degradation, the initial discharge capacity of exposed NMC is considerably lower, 145mAh.g -1

  5 at 600°C under O2 atmosphere gave the biggest improvement among the annealing temperatures of 100, 400 and 600°C. Similarly, Liu et al. 2 reported that 98.8% of the original discharge capacity of a 6-month exposed NMC 802 can be recovered by annealing at 725°C under flowing O2. These two works attributed the capacity improvements to the reconstruction of a layered phase because of the use of O2 and the high annealing temperature. Based on the findings of this work, we think that the reconstruction is in reality due to the reaction of the MOOH species with O2 and Li2CO3.

Figure III. 18 .

 18 Figure III.18. Discharge capacity retention curves of the R, E and E-A NMC 622C/Li halfcell at C/10 between 2.8-4.3 V. Inset: table with the quantity of soluble bases obtained from acid-base titration.

Figure III. 19 .

 19 Figure III.19. Voltage profiles vs. capacity for half-cells at a rate of C/10 between 2.8-4.3 V for the as-received, exposed and exposed + annealed materials.

Figure III. 20 .

 20 Figure III.20. TGA curves along with mass traces of H2O (m/z 18), O2 (m/z 32) and CO2 (m/z 44) of an E and E-A NMC 622C.

Ni-rich NMC materials 1 .1. 1 . 1 . 1 . 1 . 1 . 1 . 2 .

 11111112 Introduction....................................................................................................................145 Gas production mechanisms ........................................................................................... Gas associated with the positive electrode .............................................................. Gas associated with the negative electrode ...........................................................

3. 1 . 3 . 1 . 1 . 3 . 1 . 2 . 2 . 3 . 2 . 1 . 3 . 2 . 2 . 3 . 2 . 3 . 3 . 2 . 4 . 2 . 5 . 2 . 5 . 1 . 2 . 5 . 2 . 2 . 6 . 2 . 6 . 1 . 3 . 2 . 6 . 2 . 4 .

 13113122321322323324252512522626132624 Ni-rich NMC materials......................159 Gas production from the storage of NMC with electrolyte............................................... Powder storage....................................................................................................... Film storage ........................................................................................................... 3.Gas production from the storage of NMC/Graphite pouch cells ....................................... Optimization of the pouch cell format .................................................................... Complete storage protocol of pouch cells ............................................................... Gas production of the formation step ...................................................................... Gas production of NMC 811 with modified surface composition ............................ 3.Gas production of different NMC compositions...................................................... 3.Discussion about gas of the formation step ......................................................... 3.Discussion about gas of the complete storage protocol ....................................... 3.Impedance evolution after the complete storage protocol ........................................ 3.Adaptation of the complete storage protocol ....................................................... Impedance from pouch cells NMC 622C/Gr and 811/Gr .................................... Conclusion of the chapter..............................................................................................188 5. References.......................................................................................................................189

at 4 . 2 V

 42 vs. Li/Li + on a carbon black electrode. This onset oxidation potential varies as a function of electrolyte concentration. [IV.1] ii. Chemical oxidation of alkyl carbonate solvents by active oxygen The chemical oxidation of solvents arises from the reaction of EC with active oxygen ( 1 O2). As explained in Chapter I, at high state of charge (SOC) a large part of the nickel is in Ni 4+ state, so electron transfer occurs from the 2p band of O 2-to the Fermi level of Ni 4+ , reducing it to Ni 3+ (Eq. IV.2). Meanwhile, due to the lithium deficiency and the presence of Ni 3+ , the M-O bond covalency increases, decreasing O 2-charge density and favoring oxygen release (Eq. IV.3) 8 .

  Fig. IV.1 summarizes the zones in which these reactions are expected to produce CO2 along the first cycles.

Figure IV. 1 .

 1 Figure IV.1. Scheme of the stages of CO2 production associated with NMC materials. "Chem." or "C" stands for chemical reaction and "Electroch." or "E" for electrochemical reaction.

  Nonetheless, parasitic reactions along the cycle life of the battery are susceptible to continue the production of these gases, especially at temperatures above 40 ºC when Lewis acid PF5 is

  Fig. IV.2). The capillary type contains a stationary phase coated in the interior walls, whereas in the micropacked type, it fully packs the interior. In general, micropacked columns have wider external diameter (ranging from 0.74 to 1.58 mm vs. 0.25 to 0.53 mm) are shorter in length (ranging from 1 to 2 m vs. 15 to 30 m) and require higher pressure to flow the carrier gas. Consequently, they are more retentive than the capillary type.

Figure IV. 2 .

 2 Figure IV.2. Schematic representation of the two types of chromatographic columns.

  Fig. IV.3 shows the results.

Figure IV. 3 .

 3 Figure IV.3. GC/MS-IR analysis of a pouch cell, using the HP-PLOT Q column. a, b) Gram-Schmidt chromatogram (from FTIR) and c) GC/MS chromatogram.

  Fig. IV.4 shows the resulting chromatogram. Not surprisingly, the chromatogram is equivalent to what has been previously observed in Fig. IV.3c. Note that even if argon has been changed to helium to avoid argon detection, the co-elution of gases issue has not been eliminated.

Figure IV. 4 .

 4 Figure IV.4. GC/BID analysis of a pouch cell, using a HP-PLOT Q column with the constant pressure mode.

Figure IV. 5 .

 5 Figure IV.5. GC/BID analysis of a pouch cell, using a HP-PLOT Q column with the constant linear velocity mode.

  Fig. IV.6).

Figure IV. 6 .

 6 Figure IV.6. GC/BID analysis of a pouch cell, using the ShinCarbon ST micropacked column.

Figure IV. 7 .

 7 Figure IV.7. GC/BID analysis from a pouch cell, using the ShinCarbon ST micropacked column.

IV. 8

 8 ).

Figure IV. 8 .

 8 Figure IV.8. GC/MS chromatogram of gases from the storage of 200mg of NMC 622C and 0.2 ml of LiPF6-based electrolyte in an 18650 can for 4 days at 55°C.

  The quantification on Fig. IV.9 is represented with respect to the total dead volume (dead volume = total volume of the can -volume occupied by the liquid). The error bars correspond to the standard deviation calculated from three samples.

Figure IV. 9 .

 9 Figure IV.9. CO2 quantification in GC/BID (regular method and capillary column), from stored 18650 cans containing different electrolytes. Inset: molecular structure of salts.

Figure IV. 10 .

 10 Figure IV.10. CO2 quantification in GC/BID (regular method and capillary column) from 18650 cans stored 4 days at 55°C. Insets: surface morphologies of NMC 532 and 901.

Figure IV. 11 .

 11 Figure IV.11. Illustration of the steps in the protocol used for the Swagelok-type storage tests at the delithiated or lithiated state (100% or 0% of achievable SOC).

Figure IV. 12

 12 Figure IV.12. a) GC/BID chromatogram (regular method and capillary column) of the gas from the storage of a half Swagelok-type cell containing an NMC 622C film at 100% SOC. b) CO2 quantification from the storage at 100 or 0% SOC.

  was 92 wt.% graphite (Timcal TIMREX SLS, ~1.5 m 2 .g -1 ), 4 wt.% CMC (ACROS Organics, Mw = 700000, degree of substitution = 0.9) and 4 wt.% conductive carbon black C45. The electrodes were calendered to 30% porosity and the electrochemical performance was evaluated through half-cell cycling (Fig. IV.13). At C/10 and C/4, specific capacities of 330 mAh.g -1 and 220 mAh.g -1 were obtained, respectively.

Figure IV. 13 .

 13 Figure IV.13. Electrochemical performance in half-cell of the graphite film at two different C-rates. The roman numbers correspond to the plateaus mentioned in Chapter I.

Figure IV. 14 .

 14 Figure IV.14. Scheme of the interior of a pouch cell after customization for gas storage. The volume of the reservoir is ~0.7 ml.

3. 1 . 2 .

 12 It consists of a five-hour rest (electrolyte impregnation) and one cycle and a half, at 30°C, between 2.8-4.3 V, to reach a delithiated state (Fig. IV.15), which corresponds to Li0.4Ni0.6Mn0.2Co0.2O2 in NMC 622C and Li0.3Ni0.8Mn0.1Co0.1O2 in NMC 811.

Figure IV. 15 .

 15 Figure IV.15. Storage protocol established for pouch cells.

  Fig. IV.16a-d reports peak areas of CH4, CO2, C2H4 and C2H6 as function of temperature and C-rate.

Figure IV. 16 . 5 (

 165 Figure IV.16. Peak areas of a) C2H4, b) C2H6, c) CH4 and d) CO2 detected by GC/BID (improved method and capillary column), from pouch cells NMC 811/Gr after the formation step as function of temperature and current density.

  6 mAh, compared to 9.2 mAh of the exposed 811 (Fig. IV.17a). This difference can be explained by the increased polarization of 250 mV (represented for the first charge in Fig. IV.17b), most probably due to the presence of Li2CO3, MOOH and MO.

Figure IV. 17

 17 Figure IV.17. a) Capacity vs. potential curves and b) time vs. potential curves of the formation step for pouch cells NMC 811/Gr or NMC 811(exposed)/Gr.

Figure IV. 18 .

 18 Figure IV.18. Gases detection by GC/BID (improved method and capillary column) after the complete storage protocol of as-received or exposed NMC 811/Gr pouch cells.

Fig. IV. 19

 19 Fig. IV.19 shows the formation step of two representative pouch cells. The pouch cells containing NMC 811 took longer time to cycle and deliver higher specific capacity: 186 versus 174 mAh.g -1 from 622C.

Figure IV. 19 .

 19 Figure IV.19. Time versus potential curves of the formation cycles of two representative pouch cells NMC 622C/Gr and 811/Gr.

Figure IV. 20 3 . 2 . 5 . 1 .

 203251 Figure IV.20. a) Gases from the formation step at C/10 and 30°C, b) gases from the complete storage protocol and c) gases only from the storage step. Error bars correspond to the standard deviation from three measurements.

Figure IV. 21 .

 21 Figure IV.21. Incremental capacity profiles of the first charge of representative pouch cells showing the peak of SEI formation.

Figure IV. 22 .

 22 Figure IV.22. Cycling step of the formation protocol established for the testing of 2 and 3electrode pouch cells NMC 622C/Gr.

Fig. IV. 23

 23 Fig. IV.23 shows the impedance spectra of 3-electrode pouch cells 622/Gr that performed only the formation step and the PEIS measurements at 100, 90 and 80% SOC. The Nyquist plots of the full cell (a), positive (b) and negative electrode (c) are shown individually.

Figure IV. 23 .

 23 Figure IV.23. Nyquist plots collected at different SOC after the formation step in 3-electrode NMC 622C/Gr pouch cell. a) Complete cell b) positive and c) negative electrode.

  .24). According to this model, Rct and Rcontact are in parallel with constant phase elements, modeling the imperfect double layer capacitance obtained across the electrolyte interphases, and in series with a Warburg element describing the diffusion tail at low frequency.

Figure IV. 24 .

 24 Figure IV.24. Equivalent circuit model used to fit PEIS positive electrode data in Z-view ® .

Fig. IV. 25

 25 Fig. IV.25 shows the Nyquist plots of 3-electrode pouch cells NMC 622C/Gr and 811/Gr before and after storage from the complete cell (a, b), the positive electrode (c, d) and the negative electrode (e, f). After storage, both electrodes modified their impedance. As selfdischarge process is observed in both cases (inferior to 20% of SOC), it implies that surface chemical processes occurred during storage. Note that the main difference between the two is that NMC 811/Gr (b) present bigger semicircles before and after storage than 622C/Gr (a).

Figure IV. 25 .

 25 Figure IV.25. Nyquist plots collected before and after 4 days of storage at 55°C, in pouch cells NMC 622C/Gr (a, c, e) and 811/Gr (b, d, f).

  consequently the rock salt production, are less kinetically favored at low temperature. The slight evolution of the middle frequency domain in Fig. IV.26 demonstrates that the rock salt formation was reduced at low temperature. Again, the main difference between the spectra before and after storage is observed for the charge transfer process, but in a lesser way than what was observed at 55ºC.

Figure IV. 26 .

 26 Figure IV.26. Nyquist plots from a pouch cell NMC 811/Gr before and after storage at 10°C.

  (LDFP, Fig. V.I) was proved to be a product of the chemical reaction of NMC surface species (Li2CO3, Li2O, Li2SO4) with LiPF6. Attracting much attention as capacity retention improver, the decision was made to evaluate its impact as electrolyte additive in batteries containing Ni-rich NMC materials.

Figure V. 1 .

 1 Figure V.1. Structural formula of LDFP molecule

Figure V. 2 .

 2 Figure V.2. LDFP precipitation mechanism from Yang et al. 2 .

Figure V. 3 .

 3 Figure V.3. Cyclic voltammetry of both electrolytes on a Pt working electrode, Pt counter electrode and Li reference at 30 mV.s -1 .

  polypropylene commercial separator (Celgard 2500 measurement) sandwiched between two copper electrodes. In addition to the base and the base + 0.2M LDFP electrolytes, an electrolyte containing 0.8M LiPF6 + 0.2M LDFP (in the same solvents mixture) was also tested. This mixture was used with the objective to keep the molar concentration of the salt at 1M, since the maximum ionic conductivity of LiPF6 in alkyl carbonate solvents is obtained at this concentration18 . The results are shown in Fig. V.4.

Figure V. 4 .

 4 Figure V.4. Ionic conductivity measurements performed in a) liquid and b) Celgard 2500 separator. Inset values in a) correspond to viscosity measurements.

  4 cP at 25°C and from 2.1 to 2.5 cP at 55°C (insets on Fig. V.4a). At the same time, it is suspected the formation of ion-pairs of LDFP, instead of a complete dissociation like with LiPF6.

3. 1 . 3 . 1 . 1 .

 1311 NMC 811 as positive electrode Cycling performance of pouch cells NMC 811/Graphite at 25 and 55°C The preparation of the pouch cells follows the procedure described in Chapter III, using noncalendered NMC 811 electrodes, due to the sensibility to air during the calendering process. The electrolytes employed were base and base + 0.2M LDFP and the cycling performance was evaluated at 25°C and 55°C and C/4 between 2.8-4.3 V vs. Li/Li + . All cycling tests were reproduced 3 times. The results are shown in Fig. V.5.

Figure V. 5 .

 5 Figure V.5. Electrochemical performances of representative pouch cells of NMC 811(noncalendered)/Gr cycled between 2.8 -4.3 V at C/4 and at a) 25 and b) 55°C without and with LDFP.

  V.6). On the first cycle in presence of LDFP, the peak corresponding to the reduction of solvents appears always lower in intensity than the base electrolyte. This possibly means that less solvents were reduced, signifying a differently formed SEI. Further characterizations of electrolyte degradation products are portrayed in the following sections.

Figure V. 6 .

 6 Figure V.6. Representative derivative dQ/dV curves of the first cycle, evidencing SEI formation differences without or with LDFP.

Figure V. 7 .•

 7 Figure V.7. Differential capacity vs. voltage profiles of pouch cells 811/Gr. a) Initial state of the three cells and b, c, d) each cell with each electrolyte, separately.

  Fig. V.8a-d shows representative SEM pictures of the NMC 811 electrodes before and after calendering.

Figure V. 8 .

 8 Figure V.8. Representative SEM pictures of the NMC 811 film a) before (49% porosity) and b) after calendering (40% porosity).

Figure V. 9 .

 9 Figure V.9. Representative pouch cells of 811(cal.)/Gr cycled between 2.8-4.3 V at C/4 and at 55°C without or with 0.2M LDFP. Shown for comparison, watermark curves correspond to the cycling with non-calendered films.

  at 55°C. SEM and electrochemical characterizations of each electrode were performed with the aim of determining the origin of the capacity retention fade in presence of LDFP. Prior to use, electrodes were washed three times with dimethyl carbonate and let dry for one hour under vacuum. It is possible that throughout the washing procedure some of the products deposited on the interfaces have been washed away. The electrodes were cut in the form of discs to be cycled in Swagelok-type cell versus lithium, at C/10 and 25°C, with base electrolyte soaking the glass fiber separator. For comparison, Fig. V.10a and b show the potential vs. specific capacity profiles of the recovered and fresh electrodes. Cycling at 55°C with LDFP provoked capacity loss on both electrodes. In particular, the NMC 811 electrode lost 40% of its original discharge capacity, compared to only 6% of the cell with base electrolyte (Fig. V.10b). Meanwhile, the graphite electrode that cycled with base electrolyte kept almost all of the original capacity, only 2% was lost, and 7% when LDFP was used (Fig. V.10a). This clearly identifies the positive electrode as the main source of the capacity fade at 55°C. However, at this point, it is not possible to know the origin of the capacity fade (structural or chemical).

Figure V. 10 .

 10 Figure V.10. Re-cycling at 25°C and C/10 of the washed a) graphite electrode between 1.5-0.01 V; and b) NMC 811(cal.) electrode between 2.8 -4.3 V vs. Li/Li + .

Figure V. 11 .

 11 Figure V.11. Representative SEM pictures at different magnifications of the washed NMC 811(cal.) electrodes after cycling with base (blue frame) and base + 0.2M LDFP electrolytes (red frame).

3. 2 . 3 . 2 . 1 .

 2321 NMC 622C as positive electrode Cycling performance of pouch cells NMC 622C/Graphite at 55°C The representative electrochemical performances of 622C(calendered)/Gr pouch cells cycled at C/4 and 55°C, with and without 0.2M LDFP, are shown in Fig. V.12. The NMC 622C films were calendered under air to 35% of porosity. Unlike the precedent results obtained with NMC 811, the graph clearly shows that NMC 622C-containing pouch cells exhibit, after 200 cycles, enhanced capacity retention: ca. 83% with 0.2M LDFP, compared to ca. 80 % without LDFP.

Figure V. 12 . 3 . 2 . 2 .

 12322 Figure V.12. Representative pouch cells of 622C(cal.)/Gr cycled between 2.8 -4.3 V at C/4 and at 55°C without or with 0.2M LDFP. Shown for comparison, watermark curves correspond to the cycling 811(cal.)/Gr.

Figure V. 13 .

 13 Figure V.13. Representative SEM pictures at different magnifications of the washed NMC 622C(cal.) electrodes after cycling with base electrolyte and base + 0.2M LDFP.

4 . 4 . 1 . 4 . 1 . 1 .Figure V. 14 .

 44141114 Figure V.14. Incremental capacity curves of the SEI formation during the first lithiation of graphite composite powder, in half-cell configuration.

Figure V. 15 .

 15 Figure V.15. FTIR analyses of a) the SEI formed on graphite composite powders at 25°C and 55°C without/with LDFP and b) LDFP reference.

  V.16), almost identical peaks appear but could not be objectively assigned to a CEI component or impurity.

Figure V. 16 .

 16 Figure V.16. FTIR analysis of LDFP and the CEI formed on washed NMC 811 composite powder after 3 cycles at 55°C without or with LDFP.

Figure V. 17 .

 17 Figure V.17. FTIR analysis of a) the surface of non-cycled and washed NMC 811 electrodes after 200 cycles at 55°C in pouch cell, without or with LDFP and b) pure LDFP.

4 . 2 . 1 . 15 Figure V. 18 .

 4211518 Fig. V.18 shows the chromatograms obtained from both base and base + LDFP electrolytes. In addition to the peaks corresponding to the alkyl carbonate solvents EMC, DEC and EC, three defined peaks marked as 1n, 1n '' and 1n ' (Fig. V.18) are clearly visible in case of the base electrolyte, but are completely absent when 0.2M LDFP is used, even in the background.These results are in good agreement with those of Qian et al.15

4 . 2 . 2 .

 422 10) regenerate ROLi. Consecutive reactions of the intermediary A with 11 compounds can lead to the elongation of the chain to form oligomers (Eq. V.11) as shown in the following section. seem to indicate that LDFP does not inhibit the solvents reduction and as a consequence the formation of alkoxides. Hence, the absence of 11-type compounds suggests that the alkoxides are rapidly scavenged before reacting with EC. They are supposed to substitute fluorine of LDFP, yielding a lithium organo-fluorophosphate salt. The latter may correspond to a phosphate-based compound detected in SEI, through FTIR analysis (Eq. V.12). It is noticeable that all of the substitution reactions result in increasing the LiF amount in interfacial layers. [V.12] Analysis of electrolytes by ESI-HRMS ESI-HRMS experiments were performed in the positive ion mode on a Q-TOF Ultima Global instrument (Waters-Micromass, Manchester, UK). Data acquisition and processing were performed with TargetLynx software and a listing method to distinguish the main compounds was employed (details in Annex IV). The results from the base and base + LDFP electrolytes recovered after cycling, shown in Fig. V.19, are organized in two families: none phosphate-and phosphate-containing compounds. The structural formulae are shown in Table V.1 and V.2. Strikingly, the base electrolyte contains both families, whereas the base + 0.2M LDFP electrolyte contains mostly the phosphate-containing compounds in a slightly lower amount.

Figure V. 19 .

 19 Figure V.19. Phosphate-and non-phosphate-containing compounds analyzed by ESI-HRMS in electrolyte after cycling without or with LDFP.

  oxide and ethylene carbonate units: -The ethylene oxide-based compounds correspond to the 1n family produced through Eqs. V.10 and 11 presented in Section 4.2.1. It comprises the molecules detected by GC/MS (m/z 215.053 and 229.068) and those with n=2 (m/z 259.079 and 273.094).-The carbonate-based molecules are produced through a transesterification reaction involving EC and a carbonate oligomer in presence of PF540 (Eq. V.13):[V.13]All of these compounds are produced from the hereabove described reactions initiated by alkoxides. Their quasi absence in case of LDFP electrolyte underpins the assumption that most alkoxides are scavenged by LDFP (as precipitated salt or dissociated salt in electrolyte, not detected here) or other fluorophosphates before reacting with EC.

Fig. V. 20 presents

 20 the results expressed in peak areas to be able to compare CH4 and C2H6 for which calibration curves are not available. The inset values are the results of the quantification of CO2 and C2H4 and the error bars correspond to standard deviation of the analyzed samples in each batch.

Figure V. 20 .

 20 Figure V.20. GC/BID analysis of gas from pouch cells 811/Gr after 200 cycles at 55°C, without or with LDFP.

Figure V. 21 .

 21 Figure V.21. GC/MS analysis of gas from pouch cells 622C/Gr, after 200 cycles at 55°C without/with LDFP.

  SEM at the surface of the electrode after the separator is taken off, was considered as the main issue that causes the capacity fade. It has to be recalled the different global morphologies. As shown in SEM images (Fig. V.22a and b), both NMC 622C and 811 particles are polycrystalline, however, only the NMC 811 particles present an inner porosity.

Figure V. 22 .

 22 Figure V.22. Illustrative SEM pictures of broken particles of NMC a) 622C and b) 811 showing their polycrystalline structure and the porous texture of NMC 811.

Fig. V.

  Fig. V.23a of the cycled electrode exhibits a large inner void with detached agglomerates of primary particles as well as a shell made of cohesive primary particles. Magnified images of this shell (Fig. V.23b) and of a particle before cycling (Fig. V.23c) show, systematically, amorphous grain boundaries of ~10 nm around the primary particles, probably made of light elements due to the color contrast. From results of Chapter II, we can hypothesize that this

Figure V. 23

 23 Figure V.23. a) TEM picture NMC 811 particles cycled with base electrolyte (100 cycles at 55°C), cut transversally by FIB. Magnified images b) after and c) before cycling.

Figure V. 24

 24 Figure V.24. a, b) Two different TEM magnifications of the particles in NMC 622C cycled with base electrolyte (100 cycles at 55°C), cut transversally by FIB.

Li2CO3,

  Li2O and M2SO4 (M= Li or Na), on five commercial Ni-rich NMC materials (532, 622, 622C, 811 and 901), by means of acid-base titration, ICP and AAS analyses after finetuning of water-solubilization parameters as duration, atmosphere and powder-to-water mass ratio. Attention must be paid to hydrolysis reaction, as it was proved to be at the origin of the LiOH soluble base. Commonly reported in the literature as being part of NMC surface species, this salt affects the electrode processing through NMP slurries gelation1 . Its absence in our materials was confirmed through FTIR, TGA/MS and NMR analyses. Beyond the scope of this work, establishing the stability limits of NMC materials in water is important for the development of water-based electrode processing, to replace the conventional NMP-based slurries, that so far has been attempted for NMC 111 2 and 532 3 .

Figure A. 6 .

 6 Figure A.6. Oven temperature program for the capillary column.

Figure A. 7 .•

 7 Figure A.7. Oven temperature program for the micropacked column.

Figure A. 8 .Figure A. 9 .Figure A. 10 .•

 8910 Figure A.8. Calibration curves obtained from the analysis of multiple acquisitions in GC/BID with a capillary column and the "regular" method.

•

  Focused-ion beam (FIB) cut and transmission electron microscopy (TEM)Samples of NMC 622C and 811 in film were shipped for TEM sample preparation by FIB cut to the Institut d'Electronique, de Microélectronique et de Nanotechnologie (IEMN) through the RENATECH platform. Each sample was first positioned in a sample holder and taken inside of a FIB dual Beam FEI Strata DB 235 equipped with both a gallium ion source and an electron source. Then, a platinum thin layer was deposited on top of an area of interest toXVIIRésumé développé de la thèse en FrançaisAu cours de la dernière décennie, le marché des batteries s'est rapidement développé avec l'adoption croissante des véhicules électriques. Pour répondre aux besoins en termes de densité d'énergie, les matériaux lamellaires LiNi1-x-yMnxCoyO2 (NMC) de composition 111, 442 et 532 ont fait leurs preuves comme matériaux actifs pour l'électrode positive. Néanmoins, la tendance actuelle s'oriente désormais vers l'utilisation de matériaux NMC riches en nickel (Ni> 0,5), tels que NMC 811, en raison de l'augmentation de la densité d'énergie donnée par la teneur en nickel (Fig.1a). Cependant, les performances électrochimiques et son utilisation étendue sont grandement limitées par la réactivité de surface vis-à-vis de l'air et de l'électrolyte à base de LiPF6 (Fig.1b).

Figure 1 .

 1 Figure 1. a) Diagramme de phase ternaire LiCoO2-LiMnO2-LiNiO2 avec les principaux avantages en vert et inconvénients en rouge. Adapté de Schipper et al. 1 b) Comparaison de la performance électrochimique des matériaux NMC. Adapté de Noh et al 2 .

Figure 2 .

 2 Figure 2. Images SEM représentatives de matériaux NMC riches en nickel : a) 532, b) 622C, c) 622, d) 811 et e) 901.

[ 1 ]

 1 Chapitre III : Réactivité chimique et électrochimique des espèces de surface dans les matériaux NMC riches en nickel. L'objectif de ce chapitre est d'évaluer l'impact des espèces de surface dans des conditions de batterie. Pour ce faire, leurs produits de dégradation d'électrolyte provenant du stockage à haute température (réactivité chimique), ou des tests de cyclage (réactivité électrochimique), sont analysés. Il a été constaté que la présence de certaines espèces de surface, à savoir de Li2CO3, Li2O et M2SO4; pouvait être avantageuse pour les performances électrochimiques de la batterie, puisque ces dernières produisent du LiPO2F2 (un additif d'électrolyte bien connu) et du LiF (un composé passivant stable connu), par réaction chimique à haute température avec l'électrolyte (Fig. 4 exemplifie le stockage du Li2CO3 à 55 ºC).

Figure 4 .

 4 Figure 4. Spectre RMN liquide 19 F de l'électrolyte récupéré après stockage de Li2CO3 et de l'électrolyte frais pendant 4 jours à 55 ºC.

Figure 5 . 2 ]

 52 Figure 5. a) Courbes de rétention de capacité de décharge en demi-pile NMC 622C/Li « asreceived », « exposed », et « exposed + annealed » à C/10 entre 2,8 et 4,3 V vs. Li/Li + . Encart : résultat de la quantification des bases solubles. b) Profils de tension vs. capacité.

Figure 6 .

 6 Figure 6. Analyses GC/BID des « pouch cells », a) en utilisant une colonne capillaire avec les conditions expérimentales améliorés et b) en utilisant la colonne « micropacked ».

Figure 7 . 3 ]

 73 Figure 7. Résultats de la quantification du CO2 en utilisant l'appareil GC/BID et une colonne capillaire, à partir des films NMC 622C à 100 ou 0% d'état de charge (« SOC »).

Figure 8 .

 8 Figure 8. Performances de cyclage des « pouch cells » représentatives 622C/Gr et 811/Gr cyclées entre 2,8 -4,3 V vs. Li/Li + à C/4 et à 55 º C; sans ou avec 0,2M de LDFP.

Table I .

 I 1. Each solvent presents advantages and disadvantages. For instance, the e values that favor lithium dissociation are >20, values that only cyclic carbonates EC and PC have. A peculiarity of EC is that it is solid at ambient temperature and requires a liquid solvent to dissolve. On the other hand, linear DMC, DEC and EMC, exhibit viscosity values lower than

	Solvent	Molecular representation	Tm (°C)	Tb (°C)	Dielectric e constant	Viscosity (cP)
	Ethylene
	carbonate		39	248	89.6	1.86 (at 40°C)
	(EC)
	Propylene
	carbonate		-49.2 241.7	64.4	2.53
	(PC)
	Dimethyl
	carbonate		0.5	90	3.11	0.59
	(DMC)
	Diethyl
	carbonate		-43	126.8	2.81	0.75
	(DEC)
	Ethylmethyl
	carbonate		-55	108	2.96	0.65
	(EMC)

1 cP, in comparison to 1.86 (at 40°C) and 2.53 of EC and PC, respectively. Table I.1. Physicochemical properties of organic solvents for LIBs. Adapted from Park et al. 11 .

Table I . 2 .

 I2 Ionic conductivity of electrolytes containing 1M LiPF6. Taken from Park et al. 11 .

		Ionic conductivity
	Solvent mixture	at 25°C (mS.cm -1 )
	EC/DMC (50/50 vol%)	11.6
	EC/EMC (50/50 vol%)	9.4
	EC/DEC (50/50 vol%)	8.2
	Nonetheless, LiPF6 salt presents two main disadvantages:

Table I . 3 .

 I3 Composition and name of NMC materials.

	Composition	Fractional composition Ni 3+ Ni 2+ , Mn 4+ Co 3+	Abbreviation
	LiNi1/3Mn1/3Co1/3O2	0	2/3	1/3	111
	LiNi0.4Mn0.3Co0.3O2	0.1	0.6	0.3	433
	LiNi0.4Mn0.4Co0.2O2	0	0.8	0.2	442
	LiNi0.5Mn0.3Co0.2O2	0.2	0.6	0.2	532
	LiNi0.6Mn0.2Co0.2O2	0.4	0.4	0.2	622
	LiNi0.8Mn0.1Co0.1O2	0.7	0.2	0.1	811

Table I . 4 .

 I4 Examples of electrolyte additives used in LIBs containing NMC as positive electrode.

	Additive	Advantages	Disadvantages
	Nitriles-containing Succinonitrile (SN)	Exceptional electrochemical and thermal stability due to the -CN group at high potentials 164	Cell impedance increase 165
	Sulfur-containing Ethylene sulfite (ES). Trimethylene sulfate (TMS)	Increases battery lifetimes 166 , decrease impedance growth 167 and suppress gas production 168	Electrolyte oxidative decomposition 169
	Boron-containing	Formation of protective CEI films in
	Lithium bis(oxalate)borate	high voltage LIBs 170 . Stabilizes PF6 -
	(LiBOB)	anion and prevent PF5 formation 171
			Poly(VC) formed via
	Vinylene carbonate (VC)	Reduces at the surface of the negative electrode during the first charge 43	oxidative decomposition of
			VC 172
	Fluoroethylene carbonate (FEC)	Improves the electrochemical performance of NMC by forming a thin CEI composed mainly of LiF 169

Table II .1.

 II Summary of the surface species proposed to exist in layered materials along with their origin, formation mechanisms and credibility

	Origin	Surface specie		Reaction mechanisms		[#]	Credibility	Refs
		Li2CO3	Residual from:	2"##$ + &' ( )# *	+,,-./,º1, 4 5 6⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯8 2&'"# ( + )# ( + $ ( #	[II.1]	High	1, 2, 7,
												23
	Synthesis	LiOH	Residual from:	"##$ + &'#$. $ ( #	+,,-./,º1, 4 5 6⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯8 &'"# ( + 2$ ( #	[II.2]	High
		Li2O	Calcination reaction:				2&'#$	∆ → &' ( # + $ ( #		[II.3]	never detected High but almost	24
			Redox reaction (1):	&'"# ( +	< =	# ( +	< (	)# ( → &' >-< "# ( +	< (	&' ( )# *	[II.4]	Low, today it is known O2 does	17
												not intervene
		Li2CO3	Redox reaction (2):	?' *@ + # (-# -+ # -→ # (-(BCDDEFG) → ?' (@ + # -(CFDEIG) + > 2&' @ + # (-(CFDEIG) + )# ( $ ( # ⁄ → &' ( )# * 2&'#$ ⁄ ( # (	[II.5] [II.7] [II.6]	Low, does not without Ni 3+ consider NMC	3, 5
	Air-grown
			suggested, through Raman spectroscopy and
	thermogravimetric/mass spectrometry (TG-MS) results interpretation, that nickel-
	hydroxycarbonate (NiCO3.2Ni(OH)2.xH2O) is the major degradation product on NMC 111
	and 811 upon air exposure. Despite the experimental support, neither of them presented
	mechanisms that explained the degradation of NMC to nickel-based species.

Table II . 6 .

 II6 Lithium content and reacted percentage in as-received and exposed (H2O/CO2 for 6 days) NMC materials.

	Lithium quantity	532	622C	622	811	901
	as-received (µmol.gNMC -1 )	80 ± 3	164 ± 6	182 ± 6	252 ± 7	292 ± 7
	exposed (µmol.gNMC -1 )	143 ± 4	341 ± 8	365 ± 9	788 ± 20	1079 ± 27

  Residual lithium salts in Ni-rich NMC materials left from the synthesis, or formed as a consequence of NMC's high reactivity toward air atmosphere, are well-identified issues with respect to the performance of Li-ion batteries in terms of gas release, interfacial resistance and decreased capacity. Given the abundant literature on this topic, first conclusion drawn was that degradation mechanisms involving a nickel reduction process and the formation of nickel hydroxicarbonate were approved by the battery community. However, after a deeper analysis of results, no tangible proof could have been brought to light. This incited us to undertake a stepwise work aiming at achieving the complete identification of surface species. Provided by a single supplier, different compositions of Ni-rich NMC with controlled morphology and structure helped assess the increasing reactivity associated with the presence of Ni

3+

. Before and after exposure to a specific and rather effective H2O/CO2 atmosphere, soluble species were accurately quantified through acid-base titration, AAS and ICP-AES techniques; Li2CO3 is the main soluble compound formed upon exposure. It is formed either from LiOH through a water-requiring twostep reaction or from a newly proposed degradation reaction enlightened through TGA-MS analysis. Contrary to all expectations, the proposed mechanism does not involve either reduction processes or the formation of a metal hydroxycarbonate, but instead, metal oxyhydroxide, and potentially oxides NiO and MnO2 if NMC contains Ni 2+ and Mn 4+ . Other outstanding observations from this chapter are:

Table III .1. Comparison

 III

		Soluble bases from ABT (µmol.gNMC -1 )	Expected n LiPO2F2 from	NMR experiments
				Li2CO3, Li2O and
	Material	Li2CO3	LiOH from Li2O	M2SO4 for 100 mg of NMC in 1 ml	Storage temperature (°C)	R	n LiPO2F2 in 1 ml (µmol)
				(µmol)
	NMC 811 as-received	19 ± 1	205 ± 14	8.1 ± 0.9	25 55	0.006 0.007	6 7
	NM 811 exposed	370 ± 10	17 ± 5	20.9 ± 1.3	25 55	0.003 0.013	3 13

of expected and experimental LDFP moles in 1 ml of LiPF6-based electrolyte after storage with as-received and exposed NMC 811.

Table III .

 III 2 shows a summary of some electrochemical performance data from the first two cycles. The 2 nd cycle discharge capacity values are in good agreement with previously reported data for these compositions31 , except for NMC 901, for which there is no published data. Irreversible capacities range between 11.6 to 18.2 % without following a trend with the nickel content. The theoretical capacity delivered by the complete electrochemical oxidation of Li2CO3 contained in these NMC materials, represents only 0.5 to 1.2% of the first charge capacity. These values are not necessarily increasing with the nickel content and cannot explain the values of irreversible capacity among the different compositions. Therefore, it seems that the oxidation of Li2CO3 does not contribute significantly either to the irreversible or to the reversible capacity.

Table III . 2 .

 III2 Summary of the electrochemical performances of the five Ni-rich NMC materials.

	Material Reversible disch. capacity	Irreversible	Theoretical capacity from
		2 nd cycle (mAh.gNMC -1 )	capacity	Li2CO3 (mAh.gNMC -1 )
	532	158 (0.57 Li)	12.7 %	0.96 (0.6%)
	622C	170 (0.61 Li)	14.2 %	1.98 (1.1%)
	622	175 (0.63 Li)	11.6 %	2.14 (1.2%)
	811	192 (0.70 Li)	18.2 %	1.02 (0.5%)
	901	200 (0.72 Li)	17.1 %	2.10 (1.0%)

  attributed 0.24 µmol of CO2 in the first charge of NMC 111/Graphite cell cycled at C/10, exclusively to the electrochemical oxidation of Li2CO3, representing 88.9% of the theoretical quantity. Strangely, ~0.20 µmol of CO2 at the same onset potential were evolved during the second cycle. Even if no explanation was given, most probably, there was a significant contribution from other degradation processes, such as solvent oxidation, that were ignored. Similarly, Renfrew et al.

8

determined in a NMC 622/Li half-cell cycled at C/10, that 0.81 µmol of CO2 were produced, 86.2% of the total Li2CO3.

Table III . 3 .

 III3 CO3 2-and OH -amounts quantified through titration in each untreated and annealed NMC 622C and loadings of each corresponding film.As shown in TableIII.3, as-received, washed and corresponding annealed NMC films depict

	Material NMC 622C Abbrev.	CO3 2-(µmol.gNMC -1 )	OH -(from Li2O) (µmol.gNMC -1 )	Loading (mg.cm -2 )
	As-received	R	37	90	11.4 ± 0.5
	As-received + Annealed	R-A	36	130	11 ± 0.3
	Washed	W	5	26	11.1 ± 0.4
	Washed + Annealed	W-A	17	68	12.5 ± 0.5
	Exposed (20 days)	E	356	0	14.2 ± 0.6
	Exposed + Annealed	E-A	196	146	14.4 ± 0.6

  Gas characterization techniques ...................................................................................... 1.3. Conclusion ..................................................................................................................... Use of a capillary column with the GC/MS-IR equipment ...................................... 2.2.2. Use of a capillary column in the GC/BID equipment .............................................. 2.2.3. Use of a micropacked column with the GC/BID equipment ....................................

	..
	1.2.

2. Optimization of the instrumentation for gas quantification......................................151 2.1. Comparison of the available gas chromatography instruments ......................................... 2.2. Choice of the column...................................................................................................... 2.2.1. 2.3. Conclusion .....................................................................................................................

Table IV . 1 .

 IV1 Summary of the mechanisms producing gas in a LIB.

		Source
	Gas	Positive electrode	Negative electrode
	1 O2	-Lattice oxygen at high SOC -Electrochemical oxidation of Li2CO3	-1 O2 or O2 can be reduced at the negative electrode
		-Chem. oxidation of EC by 1 O2	-Chemical reduction of CO2
	CO2	-Chem. and electr. oxidation of Li2CO3	(consumption of CO2) from the
		-Electrochemical oxidation of solvents	electrode's crosstalk
	CO	-Chem. oxidation of EC by 1 O2	-Electr. reduction of cyclic solvents
	CH4	X	-Electr. reduction of linear solvents
	C2H4	X	-Electr. reduction of cyclic solvents
	C2H6	X	-Electr. reduction of linear solvents
	H2	X	-Electr. reduction of water

  .2. Main characteristics of the detection modes available at the analytical platform.

	Table IV.2. Characteristic	MS	FTIR	BID
	Peak area accuracy	Excellent	Good	Excellent

Table IV . 3 .

 IV3 Relevant characteristics of Ni-rich NMCs and experimental results (last column).

	NMC	Fractional composition Ni 3+ Ni 2+ Mn 4+ Co 3+	BET (m 2 .g -1 )	Water content (ppm)	Expected quantity of CO2 from Li2CO3 (µmol.gNMC -1 )	Experimental quantity of CO2 (µmol.gNMC -1 )
	532	0.2	0.3	0.3	0.2	0.27	14	18 ± 0.5	400 ± 30
	622C	0.4	0.2	0.2	0.2	0.27	0	37 ± 1	305 ± 15
	622	0.4	0.2	0.2	0.2	0.27	114	40 ± 1	300 ± 35
	811	0.7	0.1	0.1	0.1	0.24	10	19 ± 0.5	140 ± 45
	901	0.9	0	0	0.1	0.20	18	39 ± 1	100 ± 35

  .4). The applied current density was calculated in respect to the mass of NMC. Representative cycling can be found in Annex III and the values in TableIV.4. It was observed that, at equal temperature, there is negligible difference in between the 2 nd cycle discharge capacity at a rate of C/5 or C/10; and that pouch cells cycling at 55°C deliver higher capacity than those at 25°C. This may be explained by the improvement of the lithium kinetics during re-lithiation process at temperature above 40°C39 .

	Table IV.4. Summary of experimental specifications
		C/5		C/10
	Specification	30°C	55°C	30°C	55°C
	Negative electrode loading (mg.cm -2 )			7.1 ± 0.2
	Positive electrode loading (mg.cm -2 )		13.7 ± 0.1
	Current density (with x=1) (mA.cm -2 )	0.38 at C/10 and 0.76 at C/5
	2 nd cycle discharge capacity (mAh)	10.7 ± 0.1 11.2 ± 0.1 10.8 ± 0.1	11.4 ± 0.1

Table IV . 5 .

 IV5 Summary of some characteristics of the pouch cells

	Characteristic	As-received NMC 811	Exposed NMC 811
	Negative electrode loading (mg.cm -2 )	7.3 ± 0.3
	Positive electrode loading (mg.cm -2 )	13.3 ± 0.1	13.0 ± 0.3
	Current density (with x=1) (mA.cm -2 )	0.37 ± 0.01	0.36 ± 0.01
	Quantity of Li2CO3 (µmol.gNMC -1 )	19 ± 1	380 ± 10
	Quantity of Li2CO3 in the electrode (µmol)	1.04	20.9

Table IV . 6 .

 IV6 Summary of some characteristics of the pouch cells

	Specification	622C/Gr	811/Gr
	Negative electrode loading (mg.cm -2 )	7.4 ± 0.2
	Positive electrode loading (mg.cm -2 )	17.4 ± 0.3	13.7 ± 0.3
	Current density (with x=1) (mA.cm -2 )	0.48 ± 0.01	0.31 ± 0.01

Table IV

 IV

.8. The results were 0.028 and 0.036 µmol, in pouch cells 622C/Gr and 811/Gr, respectively.

Table IV . 8 .

 IV8 Quantities of CO2 issued from each step and the calculated difference, obtained from Fig. IV.20.

		CO2 (µmol)
	Gases step
		622C/Gr	811/Gr
	Formation step (from processes i to v)	0.006 ± 0.001	0.024 ± 0.006
	Complete storage protocol (formation + storage)	0.034 ± 0.002	0.060 ± 0.001
	Storage only (storage at the lithiated state)	0.021 ± 0.002	0.020 ± 0.004
	Difference B = n(formation + storage) -n(formation) (from processes ii and vi and in lesser extent i and iv)	0.028	0.036
	Difference C = n(Difference B) -n(storage only) (the CO2 result is only due to reactions with 1 O2)	0.007	0.016

Table IV . 9 .

 IV9 TableIV.9. Fitted results of resistances from the positive electrode obtained with Z-view.

	Fitted values			NMC 622C electrode	NMC 811 electrode
	(Ohm.cm 2 )			before	after	before	after
	R "electrolyte" (Ohm.cm 2 )	2.04	2.12	1.20	1.28
	R contact (Ohm.cm 2 )	2.88	6.24 (+53.8%)	6.84	14.8 (+53.8%)
	CPE contact (F.s a-1 /cm	2 )	7x10	-6	6.7x10 -6	5x10 -6	5x10 -6
	a contact			0.93	0.87	0.93	0.88
	fmax contact (Hz)		17.8 kHz	17 kHz	10.1 kHz	7.8 kHz
	R ct (Ohm.cm 2 )		2.76	4.60 (+40%)	5.76	22.4 (+74.3%)
	CPE ct (F.s a-1 /cm	2 )		0.0062	0.007	0.0057	0.0075
	a ct			0.75	0.73	0.89	0.89
	fmax ct (Hz)			35.7 Hz	18.5 Hz	7.3 Hz	1.2 Hz

  found that both amounts of LiF and LixPFyOz increase drastically with 1wt.% of LDFP. Furthermore, Ma et al.9 reported that the SEI formed with 0.5 wt.% of LDFP contains fewer carbonate species and more fluorophosphate species, after the formation cycles. It is worth noting here that P, O, F-based compounds (LiPO2F2, organofluorophosphates…) are not directly formed from electrochemical reduction processes, hence, they can be present in electrolyte and not necessarily inside the SEI layers. This could explain the different results obtained by XPS technique regarding the amount of LixPFyOz.

Table V .

 V

1. List of m/z values, retention time (RT) and formulae of compounds without phosphates, observed in ESI-HRMS. m/z RT (

min) Brute formula Structural formula

	215.053	6.86	C7H12NaO6
	229.068	7.62	C8H14NaO6
	259.079	7.05	C9H16O7Na
	273.094	7.73	C10H18O7Na
	303.068	7.51	C10H16NaO9
	317.084	8.09	C11H18NaO9
	319.041	7.63	C10H16KO9
	405.100	8.39	C14H22NaO12

Table V . 2

 V2

. List of m/z values, retention time (RT) and formulae of compounds containing phosphates, observed in ESI-HRMS. m/

z RT (min) Brute formula Structural formula

	293.076	7.24	C9H19O7NaP
	317.032	6.45	C7H17O7FNaP2
	331.048	6.89	C8H19O7FNaP2
	355.055	7.64	C10H18FNaO9P
	367.076	7.18	C11H21O10NaP
	381.092	7.66	C12H23O10Na P
	253.025	6.48	C6H12O6FNaP
	263.028	5.74	C7H13O7NaP
	265.044	6.14	C7H15O7NaP
	267.040	7.15	C7H14O6FNaP
	279.060	6.67	C8H17O7NaP

The none-phosphate-containing compounds are low-mass oligomers (n ≤ 2) with ethylene

Chapter V: Study of LiPO2F2 as electrolyte additive in batteries containing Ni-rich NMC materials

-Explanation to the consumption of C2H4 First, to explain the consumption of C2H4, Table IV.7 recapitulates the quantities of C2H4 from the formation step, from the complete storage protocol and from the "storage only". The "difference A" corresponds to the calculated C2H4 consumed during the storage step.

Table IV.7. Quantities of C2H4 issued from each step and the calculated difference (from Fig.

IV.20).

Gases step C2H4 (µmol)

622C/Gr 811/Gr Formation step 0.093 ± 0.020 0.155 ± 0.024

Complete storage protocol (formation + storage) 0.022 ± 0.002 0.055 ± 0.010

Storage only 0 0

Difference A = n(formation) -n(formation + storage) (C2H4 consumed) 0.071 ± 0.022 0.100 ± 0.034

The explanation of C2H4 consumption is rather complicated and to the author's knowledge, there is very limited information. The most relevant work is the one of Ellis et al. 43 , who injected C2H4 into pouch bags containing lithiated graphite electrodes (without electrolyte and positive electrode) and observed that C2H4 was almost completely consumed. They proposed the reduction of C2H4 by lithium to form Li(C2H4)n (n = 1,2,3) or Li2C2H4. These species can be synthesized, but are only stable at very low temperature (T ~ -189°C) 43 , so this hypothesis is highly unlikely.

Instead, we hypothesize that C2H4 is oxidized by 1 O2 (Eq. IV. 20), knowing that the storage of delithiated NMC at high temperature facilitates the production of oxygen (Eq. IV.5) 8 . In this case, the 0.071 and 0.100 µmol of C2H4 consumed corresponds to ~0.015 and ~0.021 mol% of the oxygen content in NMC 622C and NMC 811, respectively.

Furthermore, if the reaction is complete, 0.14 and 0.20 µmol of CO2 are expected from 622C/Gr and 811/Gr, respectively. These values do not correspond to the actual increase of CO2 of only 0.028 and 0.036 µmol of CO2 in 622C/Gr and 811/Gr, respectively (Table IV.8).

Annexes

I

Annex I (Chapter II)

• Scanning electron microscopy (SEM)

SEM images of each NMC powder were carried out by means of a field emission gun scanning electron microscope fitted with an energy dispersive X-ray spectrometer (SEM-EDX, Oxford Instrument), operating at 15 kV and in secondary electrons (SE) or backscattered electrons (BSE) mode.

• Powder X-ray diffraction (XRD)

An XRD (Bruker) using Cu Ka radiation, accelerating voltage of 40 kV and current of 100 mA was used to identify the crystalline structure of each powder. The data were obtained in the 2theta range of 10-80º with a step size of 0.02º and scan rate of 2º.min -1 .

• Atomic absorption spectroscopy (AAS)

The AAS equipment (Perkin Elmer, Analyst 300) is composed of a hollow cathode lamp For the element analysis in the filtrate solution, the filtrate solution was diluted to a factor of 2 and acidified to 2% with 1M HNO3. Calibration curves were constructed from the elements that were found in significant concentration in the analyte solution. For the powders, 250 mg of each NMC powder were dissolved separately in 30 ml of aqua regia (HCl/HNO3 2:1, vol.), during four hours in a sand bath at 140 ± 10 ºC. The transition metals and aluminum were dosed from the solution diluted by a factor of 1000 and 100 in ultra-pure water, respectively.

In both ICP and AAS, every analysis was performed at least three times to ensure reproducibility and the error bars were constructed from the standard deviation of the repetitions, in addition to the error inherent to water dilutions. The amount of every element, in ppm, in each filtrate solution was converted into concentration mol.gNMC -1 , given that the amount of water and powder is known.

• Fourier transform infrared (FTIR)

FTIR was carried out in ATR mode (FTIR-ATR, Nicolet iS20) using a Ge crystal, in a dry room with dew point between -58 and -63 ºC (~5-10 ppm of H2O). Anhydrous lithium sulfate was obtained from the vacuum-drying at 130 ºC overnight of Li2SO4.1H2O (Sigma-Aldrich, 99%).

• Thermogravimetric analysis-mass spectrometry (TGA-MS)

The TGA-MS instrument (STA 449F3, NETZSCH, Germany) is placed inside an argon-filled glovebox (< 0.1 ppm O2 and H2O, Jacomex). The experiments were performed from 25 to 900 ºC at a rate of 5 K.min -1 under argon atmosphere at a flow rate of 40 ml.min -1 .

• Raman Spectroscopy

The exposed NMC powder was measured without further treatments, on a Raman DXR microscope (Thermo Fisher Scientific) with a He-Ne laser (λ = 532 nm, 4mW), an objective with 10X magnification and 50µm pinhole aperture. The Raman spectra were measured under air within no more than five minutes to avoid heavy decomposition and on different areas to ensure reproducibility.

Annexes

III

Annex II (Chapter III)

• Preparation of reference materials and electrolyte

The reference materials Li2CO3 (> 99%) and Li2O (97%) were purchased from Sigma-Aldrich.

Li2SO4 was obtained by vacuum-drying at 130 ºC of Li2SO4.1H2O (Sigma-Aldrich, 99%).

LiOH was synthesized by Li reaction with methanol, then vacuum-dried.

The electrolyte consists in 1M LiPF6 in a mixture of EC/DEC/EMC 1:1:1 vol. (99.9%, H2O < 20 ppm; Solvionics).

• Nuclear magnetic resonance (NMR)

The NMR analyses were performed in a Bruker Ultrashield 400 MHz and the fluorine (376.5 MHz) spectra were acquired with a 5-mm nitrogen cryoprobe and recorded with a relaxation delay of 5s prior to one 90º pulse excitation (13.5 µs) and an acquisition time of 0.8 s (total relaxation for relative quantification).

• Preparation for ball milling

Inside the argon-filled glovebox (< 0.1 ppm O2 and H2O, MBraun), to avoid air exposure, 5 g of Li2CO3 powder were weighed and introduced to the SPEX cell to be mechanically milled by means of a ball miller (SPEX 800) during 35 minutes, in order to reduce the particle size.

• Preparation for electrochemical analysis with powder in half-cells

The Li2CO3 (pristine or milled) and Li2SO4 powders, and the 20-days exposed NMC 622C, were grounded separately with conductive carbon black C45 (and BET = 45m

• Powder preparation for FTIR analysis

Immediately after the electrochemical evaluation, the Swagelok-type cells containing NMC powder were transferred to the glovebox and the powders were rinsed with dimethyl carbonate (DMC) and centrifuged three times to remove the electrolyte salts. Then, they were left to vacuum dry at ambient temperature overnight. The FTIR experiments were carried out in ATR mode, using a Ge crystal, in the dry room.

• Annealing of powder materials

The annealing consisted in putting 10g of the material in an alumina boat inside a tubular furnace under constant argon flow (0.5 ml.min -1 ). The temperature program was defined to heat at 5 K.min -1 up to 500 ºC, then 7 hours fixed at this temperature and finally cool at 5 K.min -1 until 40 ºC. At the end of the temperature program, the tube was isolated from both sides while keeping argon atmosphere inside, and the annealed powders were transferred from the tube to the glovebox with limited air exposure.

Annexes

V

• Calculation of shear rate (g) Shear rate is defined as the rate at which a shearing deformation is applied to the slurry (Fig. In the case of the shear rate during the coating of a film onto the aluminum film, the velocity used is 0.85 cm.s -1 and the gap 0.025 cm, which gives a shear rate of 34 s -1 .

Annexes

VI

Annex III (Chapter IV)

• Description of the gas chromatography (GC) device GC is composed of an inert gaseous mobile phase that flows through the separation column toward the detector, and a stationary phase, which is a solid or liquid coated or packed onto the column located inside a temperature-controlled oven (Fig. A.4). Gaseous and liquid samples are injected with gas-tight syringes. After the injection, the analyte sample gets vaporized in the heated injector (3) and then enters to the column (5) contained inside a temperature-controlled oven (4), where each component is retained differently depending on the interaction of the stationary phase and the analytes. Finally, the quality of the chromatogram (7) depends on the conditions defined in the method (2), the type of gas in the mobile phase (1) and the sensibility of the detector (6).

• Pouch cell sealing conditions

All of the pouch cells were prepared in the dry room, where three sides were sealed prior electrolyte filling with a Joke Registron machine at 115 ºC during 3.5 s, followed by 3 s of cooling. Then, the pouch cells were vacuum-dried overnight at 80 ºC. The pouch cell was degassed for 8 seconds and the last side was sealed after electrolyte filling (80 µl) with a vacuum sealer (Henkelman Boxer 42 XL) until 98% of the oil pump capacity. The electrolyte was 1M LiPF6 in a 1:1:1 vol. mixture of EC/DEC/EMC (99.9%, H2O < 20 ppm; Solvionics).

The separator is the Celgard 2500.

• GC/MS-IR experimental conditions

GC/MS-IR analyses were performed using a trace GC ultra-gas chromatograph (Thermo Scientific) connected to an ISQ 1100 mass spectrometer (Thermo Scientific) and to a FTIR system Nicolet 6700 with a mid-infrared source and a medium band, liquid nitrogen cooled mercury cadmium telluride detector. The GC was interfaced with a light-pipe GC/FTIR (Thermo Scientific), according to the procedure described by Gachot et al. 26 . The Aldrich vapour phase FTIR library and the National Institutes of Standards (NIST) library of ionized compounds were used.

• Capillary column for GC/MS-IR and GC/BID

The capillary column used is a HP-PLOT Q capillary column from Agilent J & W Technologies (30m, d=0.32mm, 0.020 mm ID, polystyrene-codivinylbenzene-based) connected to a postcapillary column Rtx-1 from Restek (15 m, 0.25 mm ID, 0.25 µm, 100% Annexes XII protect it from the ion irradiation. In the selected area, gallium ions were used to pulverize the matter to create a section of a hundred of nanometers thick over 3-4 µm in depth. The created section is cut on the edges and attached using platinum to a micromanipulator to be afterwards positioned at the edge of a grid holder. It is then separated from the micromanipulator and attached to the TEM grid using platinum. Finally, the sample was thinned by consequent gallium ion irradiation with lower beam energy until the desired thickness was achieved.

The samples were then shipped back and analyzed by (TEM) (FEI Tecnai G2 electron microscope operating at 200 kV) without further treatments and with minimal air exposure.

• Electrolyte recovery from pouch cells and "liquid" GC/MS conditions

The electrolyte was extracted from pouch cells by vortexing each separator in 1mL of acetonitrile for 1 minute. The resulting liquid was diluted with dry acetonitrile (HPLC grade) by a factor of 200 and then 2µL were analyzed by "liquid" GC/MS and by a factor of 2000 for ESI-HRMS analysis.

"Liquid" GC/MS analyses were performed using a trace 1300 series GC Ultra gas chromatograph interfaced with an ISQ mass spectrometer (both Thermo Scientific). The chromatographic separation was performed on a "BPX70" cyanopropyl polysilene-siloxane based capillary column (30 m x 0.25 mm i.d., 0.25 mm) from SGE.

• High resolution mass spectrometry -electrospray ionization (ESI-HRMS)

Experimental conditions ESI-HRMS experiments were performed in the positive ion mode on a Q-TOF Ultima Global instrument (Waters-Micromass, Manchester, UK) equipped with a pneumatically assisted electrospray ion source (Z-spray) and an additional sprayer for the reference compound (LockSpray). The prepared solutions were directly introduced (5 L.min -1 ) via an integrated syringe pump in the electrospray source. The source and desolvation temperatures were kept at 120 and 450 ºC, respectively. Nitrogen was used as a drying and nebulizing gas. The cone gas flow and desolvation gas flow rates are 50 and 800 L.h -1 , respectively; and the capillary voltage was 3.0 kV (POS).

Method for ion listing

-Record the maximum chromatographed ions associated with a retention time in each sample and delete duplicates.

Annexes

XIII

-Create a method for integrating peaks by TargetLynx by specifying the mass measured at the appropriate retention time.

-Use TargetLynx on all mode samples (positive or negative), including at least one blank per sample.

Method for cleaning TargetLynx results

-Replace with 0 all values less than /7uS99S òôöVõú + (10 × ûT5ü::uóS òôöVõ ).

-If variables are equal to 0 in all samples (blanks excluded), then delete the variables.

-Application of a threshold for the positive mode in order to keep only the strongly expressed variables: if the integration of a variable is greater than 1000 in at least one of the samples, this variable is kept, otherwise it is deleted.

-Format data as a matrix (column = ions; row = samples). Les résultats de cette thèse ouvrent non seulement des pistes de recherche qui pourraient être poursuivies, mais également contribue à enrichir la compréhension des matériaux NMC riches en Nickel en vue de leur utilisation dans l'industrie automobile.

Annexes

XIV