Keywords: thermoplastic elastomer, nitroxide-mediated polymerization, 𝛽-myrcene, isoprene ..…………………………………….…………........ styrene statistical copolymerization………..……………………...……………

The aim of this study was to produce a novel thermoplastic elastomer (TPE), used as a tough and stable joining between composite matrices. Having typically thermo-reversible crosslinks, TPE can be processed as thermoplastics and exhibit elastic behavior similar to that of chemically crosslinked elastomers in a certain temperature range. Nitroxide-mediated polymerization (NMP) was selected to synthesize linear styrenic block copolymer TPE. 𝛽-Myrcene (My) was first considered to manufacture the soft elastomeric segment poly(𝛽-myrcene) P(My) (glass transition temperature Tg ~ -77 o C). NMP of My at 120 o C in bulk using succinimidyl ester-functionalized BlocBuilder TM alkoxyamine (NHS-BB) was well-controlled, allowing styrene (S) chain-extension from cis-1,4-P(My) macroinitiator. The resulting P(My-b-S) diblock copolymers, exhibiting relatively low number-average molar mass (Mn < 50 kg.mol -1 ), showed limited stress-strain behavior (ultimate tensile strength σB < 1.1 MPa, elongation at break εB < 16%). Meanwhile, the introduction of functional groups into the soft segment was implemented to subsequently aid the adhesive bonding / welding process between the TPE and the considered polar thermoplastic composites.

Well-tailored epoxide functionalized P(My) were thereby synthesized by My/glycidyl methacrylate (GMA) nitroxide-mediated copolymerization. BlocBuilder-terminated poly(ethylene-co-butylene)-(SG1)2 (SG1 = chain-end nitroxide group) difunctional initiator was used to produce S-My-S triblock copolymers with Mn = 56-66 kg.mol -1 , which exhibited improved extensibility (σB < 0.8 MPa, εB < 200%). Poly(styrene) PS blocks were then substituted by higher Tg blocks, namely poly(isobornyl methacrylate) P(IBOMA) to enhance the toughness and the service temperature of the candidate TPE. Micro-phase separated IBOMA-My-IBOMA type triblocks exhibited improved mechanical properties (σB = 3.9 MPa, yield strength σY = 5.0 MPa, εB = 490%) associated with an extended upper service temperature of about 140 o C. However, My-based TPE did not satisfy the ArianeGroup mechanical requirements at room temperature, which prompted us to replace P(My) by poly(isoprene) (PI), which has much lower entanglement molar mass compared to P(My). Well-defined and active 1,4-PI-(SG1)2 macroinitiators were first chain-extended with S, leading to self-assembled S-I-S triblocks (Mn = 95-109 kg.mol -1 , molar fraction of S in the copolymer FS = 0.30-0.49, dispersity Đ = 2.11-2.29). With FS = 0.38, S-I-S showed σB = 4.1 ± 0.2 MPa and εB = 380 ± 60 %. IBOMA-I-IBOMA type copolymer was then produced (Mn = 94 kg.mol -1 , Đ = 1.76, FIBOMA = 0.35) and improved stress-strain properties were obtained at room temperature (σB = 11.4 ± 0.6 MPa and εB = 1360 ± 210 %). Lastly, hydrogenation of I-based block copolymers was performed at normal pressure and resulted mostly in an enhanced thermal stability, a greater tensile stress at break and a reduced elongation at break.

Finally, I want to address my sincerest thanks to Laura for her unwavering support. I also thank Vachette, Rose, Archange and Moune for their kind words of encouragement and for having graciously looked after me over the past years. I spare an affectionate thought for my beloved Mic and René. iv RÉSUMÉ L'objectif de ce projet fut consacré à la synthèse d'un élastomère thermoplastique (TPE) innovant, utilisable en tant que joint entre des matrices composites. La polymérisation radicalaire contrôlée par des radicaux nitroxyde (NMP) fut choisie afin de produire des TPE copolymères à blocs de type styrénique. Le 𝛽-myrcène (My) fut dans un premier temps sélectionné pour synthétiser le bloc souple poly(𝛽-myrcène) P(My). La NMP du My en masse à 120 o C en utilisant l'amorceur BlocBuilder TM fonctionnalisé avec l'ester succinimidyl (NHS-BB) fut très bien contrôlée, permettant l'extension de chaîne du macro-amorceur cis-1,4-P(My) avec du styrène (S). Les copolymères diblocs P(My-b-S) obtenus ayant une masse molaire moyenne en nombre modérée (Mn < 50 kg.mol -1 ) montrèrent un comportement fragile en test de traction uniaxiale (résistance à la rupture en traction σB < 1,1 MPa, allongement à la rupture en traction εB < 16%).

L'incorporation d'unités fonctionnelles au sein du segment souple fut réalisée en parallèle via la NMP du My avec du méthacrylate de glycidyle (GMA) afin de favoriser possiblement le processus de soudage / collage entre le TPE et les composites thermoplastiques envisagés. Un amorceur bifonctionnel, le poly(éthylène-co-butylène)-(SG1)2 (SG1 = groupe nitroxyde), permit par la suite de synthétiser des copolymères triblocs S-My-S ayant une plus haute masse molaire moyenne Mn = 56-66 kg.mol -1 et une plus grande extensibilité (σB < 0,8 MPa, εB < 200%). Les segments PS furent remplacés par des segments ayant une température de transition vitreuse (Tg) plus élevée, à savoir des blocs poly(méthacrylate d'isobornyle) P(IBOMA) afin d'augmenter la résistance mécanique et la température de service du TPE candidat. Des triblocs de type IBOMA-My-IBOMA, dont les domaines furent micro-structurés, montrèrent de meilleures propriétés mécaniques (σB = 3,9 MPa, contrainte à la limite d'élasticité σY = 5,0 MPa, εB = 490%) et une température de service maximale d'environ 140 o C. Toutefois, ces TPE à base de My ne satisfirent pas le cahier des charges industriel à température ambiante, ce qui nous poussa à substituer le bloc flexible P(My) par du poly(isoprène) PI, ayant une masse molaire d'enchevêtrement bien plus faible. Des macro-amorceurs 1,4-PI bien définis et actifs furent d'abord étendus avec du styrène, ce qui permit d'obtenir des triblocs S-I-S auto-assemblés (Mn = 95-109 kg.mol -1 , fraction molaire en styrène FS = 0,30-0,49, dispersité Đ = 2,11-2,29). σB = 4,1 ± 0,2 MPa et εB = 380 ± 60 % furent mesurés pour un S-I-S ayant FS = 0,38. Un tribloc de type IBOMA-I-IBOMA fut par la suite synthétisé (Mn = 94 kg.mol -1 , Đ = 1,76, FIBOMA = 0,35) et montra de meilleures propriétés en contrainte-déformation (σB = 11,4 ± 0,6 MPa, εB = 1360 ± 210 %). Enfin, des copolymères à blocs à base d'isoprène furent hydrogénés à pression ambiante en utilisant l'agent diimide, ce qui permit notamment d'améliorer la stabilité thermique et la résistance à la rupture de ces matériaux. Boltzmann's constant (J.K -1 ) kc Rate of recombination (L.mol -1 .s -1 ) kd Dissociation rate constant (s -1 ) kp Propagation rate constant (L.mol -1 .s -1 ) K Equilibrium constant (mol.L -1 ) KP(YY), αP(YY)

Mark-Houwink-Sakurada (MHS) parameters for the polymer P(YY) (K in dL.g -1 ) LV or FSG1

Living molar fraction of chains terminated by a SG1 unit (mol%) Me Entanglement molecular weight (kg.mol -1 ) Mn Number-average molecular weight (kg.mol -1 ) Mn,GPC Number-average molecular weight directly measured by GPC (kg.mol -1 ) Mn,MHS Corrected Mn,GPC value using MHS coefficients (kg.mol -1 ) vii Mn,theo Target number-average molecular weight at X = 100% (kg.mol -1 ) Mn,theo,X Target number-average molecular weight at conversion X (kg.mol -1 ) MYY Molecular weight of component YY (g.mol -1 ) mv Mass of the empty vial (g) mv+p Mass of the vial containing the polymer after drying (g) mv+p+s Mass of the vial containing the reaction solution (g) Mw Weight-average molecular weight (kg.mol -1 ) Na Avogadro constant (mol -1 ) Nc Number of chains per volume unit R Ideal gas constant (J.mol -1 .K -1 ) R 2 Squared linear regression coefficient rYY

Copolymer reactivity ratio for YY vulcanisation; temps de cycle de fabrication relativement longs; impossibilité de remise en oeuvre (déchets rarement réutilisés)…De plus, étant chimiquement réticulés, ces caoutchoucs ne peuvent pas être utilisés dans des procédés novateurs comme le soudage où la matrice polymère doit être fondue ou du moins ramollie (dans le domaine de l´écoulement caoutchoutique). Il est ainsi légitime de s'orienter vers de nouveaux types d'élastomères haute-performances disposant de propriétés intrinsèques aussi intéressantes que celles des caoutchoucs traditionnels, tout en étant pas réticulés chimiquement.

Entre les thermoplastiques à la mise en oeuvre aisée mais aux propriétés élastiques restreintes, et les élastomères vulcanisés aux propriétés élastiques remarquables, mais à la mise en oeuvre plus complexe, des matériaux intermédiaires ont fait leur apparition, les élastomères thermoplastiques (TPE). Ils offrent une combinaison de propriétés particulières : des propriétés mécaniques comparables aux caoutchoucs thermodurcissables mais limitées à un domaine de températures modérées, inférieures aux températures de ramollissement des domaines rigides ; facilité de mise en oeuvre des thermoplastiques, sans vulcanisation ; facilité de recyclage des déchets. Les TPE de type copolymères à blocs, constitués d'une phase souple (PI par exemple) généralement prédominante et d'une phase rigide (PS par exemple) reliées entre elles par des liaisons covalentes, peuvent être vus comme des matériaux de substitution prometteurs aux élastomères vulcanisés. Des noeuds de réticulation physiques, qui peuvent se définir comme des liaisons intermoléculaires non-covalentes et thermoréversibles, sont formés par la phase rigide, qui une fois dans le domaine de l´écoulement caoutchoutique à haute température, permet la mise en oeuvre du TPE.

Alors que les caoutchoucs vulcanisés ne peuvent être joints à d'autres surfaces polymère que par collage ou par fixation mécanique, la possibilité de joindre les TPE dans le domaine de l´écoulement caoutchoutique ou à l'état fondu par soudage à haute température apparaît avantageux. Le soudage ne fait pas appel à des matériaux tiers dans la liaison et permet à deux pièces assemblées de reconstituer une structure monolithique. On retrouve dans la liaison de deux pièces soudées des caractéristiques proches de celles des matériaux constituant ces pièces. Le soudage est un procédé rapide et contrôlé qui s'impose dans beaucoup de domaines où il est nécessaire et (g') poly(méthacrylate de glycidyle) P(GMA).

10

En parallèle de cette étude, nous nous sommes penchés sur la fonctionnalisation du segment souple P(My). En effet, la capacité du TPE à assurer la liaison entre deux matériaux composites est un objectif essentiel du projet doctoral. Si l'assemblage se réalise par collage (utilisation de composites thermodurcissables portant des groupes fonctionnels réactifs), la modification du TPE permettrait l'adhésion par liaisons covalentes de ce dernier avec la matrice thermodurcissable par exemple. Dans le cas d'un soudage (utilisation de composites thermoplastiques ayant une matrice relativement polaire), l'ajout d'unités fonctionnelles au sein du TPE candidat pourrait augmenter la miscibilité à l'interface avec les composites via la formation de liaisons intermoléculaires et favoriserait donc le processus d'inter-diffusion des chaînes macromoléculaires de part et d'autre de la surface de contact. Il fut envisagé de copolymériser My avec un monomère fonctionnel par NMP afin d'obtenir en une seule étape un segment élastomère portant des groupes polaires. La copolymérisation de My avec l'anhydride maléique (MA) ne fut pas concluante, possiblement due à une réaction secondaire de type Diels-Alder. De même, la copolymérisation par NMP de My avec l'acrylate de tert-butyle (tBuA) ne fut pas contrôlée, probablement causée par une réaction secondaire de type ``back-biting`` (réaction de transfert intramoléculaire) engendrant de nombreuses ramifications au sein des chaînes polymère. Toutefois, la synthèse contrôlée par NMP de copolymères P(My-co-MMA) (MMA = méthacrylate de méthyle) avec des valeurs Mn proches de celles prédites et Đ ≤ 1,32 nous poussa à étudier la copolymérisation de My avec un autre ester méthacrylique portant un groupe fonctionnel, à savoir le méthacrylate de glycidyle (GMA) possédant une fonction époxyde. La copolymérisation du couple My/GMA par NMP à 120 o C en masse, amorcée par NHS-BB, fut entreprise pour une fraction molaire initiale en GMA (fGMA,0) comprise entre 0,10 et 0,90 par rapport à My.

Les rapports de réactivité, déterminés par les méthodes de Fineman-Ross (rMy = 0,80 ± 0,31, rGMA = 0,71 ± 0,15), Kelen-Tudos (rMy = 0,48 ± 0,12, rGMA = 0,53 ± 0,18) et par une technique non-linéaire des moindres carrés résolvant l'équation de Mayo-Lewis (rMy = 0,49 ± 0,13, rGMA = 0,50 ± 0,13) indiquèrent la nature statistique de cette copolymérisation. D'un point de vue cinétique, quel que soit fGMA,0, la polydispersité des chaînes fut relativement faible (Đ ≤ 1,55) et Mn évolua linéairement avec la conversion X (Mn/Mn,theo ˃ 70% en moyenne) et ceci même à X ˃ 80%. L'extension de chaîne réussie de ces copolymères statistiques P(Mystat-GMA) avec My, GMA et S témoigna d'une synthèse contrôlée par NMP, permettant de conserver en bout de chaîne le groupe nitroxyde thermiquement labile (fraction molaire ˃ 70%, mesurée par 31 P NMR).

L'insertion d'unités époxyde dans le segment P(My) ouvre la porte à de multiples modifications possibles.

T ( o C) Module de conservation (MPa)

Module de conservation Facteur d'amortissement référence S-I-S hydrogéné fut supérieur, que cela soit à l'état caoutchoutique (T < Tg,PS) ou dans un état tel que la température soit superiéure à Tg,PS. De plus, la baisse du module de conservation à T ˃ 90 o C fut beaucoup plus modérée, indiquant une plus grande résistance thermique.

Enfin, trois TPE de référence furent caractérisés. Il s'agissait de deux triblocs S-I-S de Sigma-Aldrich ® (S4-I-S4, Mn = 98 kg.mol -1 , Đ = 1,29, FS ~ 0,08 et S8-I-S8, Mn = 97 kg.mol -1 , Đ = 1,25, FS ~ 0,16) et d'un tribloc SEBS de Kraton ® (Mn = 88 kg.mol -1 , Đ = 1,16, FS = 0,48), synthétisés industriellement par polymérisation anionique. L'objectif était d'évaluer le potentiel de certains triblocs commerciaux de type styrénique, différents en composition (S4-I-S4 et S8-I-S8) et en structure (S-I-S et SEBS), vis-à-vis des propriétés exigées en contrainte-taux d'extension à température ambiante. Malgré une extensibilité remarquable (εB = 4250 ± 230 %), les matériaux S4-I-S4 manquèrent de résistance à la traction pour l'application recherchée (σB = 7,5 ± 0,6 MPa), du fait d'une fraction molaire en styrène relativement faible (~ 8 mol%). Les propriétés mécaniques du S8-I-S8, contenant deux fois plus d'unités styrène, se rapprochèrent d'avantage des exigences d'ArianeGroup avec σB = 12,7 ± 0,7 MPa et εB = 3740 ± 140 %, bien que manquant encore de résistance mécanique. Cela permit de rendre compte qu'une fraction molaire en styrène supérieure à 20% doit être ciblée pour cette famille de TPE. Pour les essais mécaniques de ces S-I-S de référence, l´entrée en plasticité est suivi d'une phase de durcissement, provoquant un accroissement significatif de la résistance à la rupture. Un tel comportement, caractéristique des élastomères thermoplastiques, ne fut pas observé pour les S-I-S préparés par NMP, ce qui pourrait s'expliquer par une distribution en masses molaires plus large (Đ = 2,18-2,29, c´est-à-dire la présence de chaînes courtes fragilisant le copolymère) en comparaison à celle des produits de référence (Đ ≤ 1,29). Il peut également être avancé qu'une plus haute polydispersité des chaînes ne favorise pas une séparation nette des phases, se traduisant par des propriétés mécaniques inférieures aux attentes. Pour finir, un SEBS commercial (version hydrogénée du S-B-S) ayant une haute fraction molaire en styrène (48 mol%) En continuité avec l'étude réalisée pour cette thèse, les futurs travaux pourraient se consacrer à l'optimisation des copolymères à blocs à base d'isoprène synthétisés par NMP afin de satisfaire entièrement le cahier des charges d'ArianeGroup. Le copolymère tribloc de type IBOMA-I-IBOMA, qui montra les caractéristiques mécaniques les plus intéressantes, pourrait être utilisé en tant que référence et modifié de la manière suivante :

-Synthèse de PI en milieu dispersé afin de préparer un segment souple de plus haute masse molaire moyenne et avec une dispersité plus faible (amélioration des propriétés en traction uniaxiale attendue du fait d´une phase souple plus enchevêtrée et contenant moins de chaînes courtes).

-Remplacement de P(IBOMA) par un polymère au moins aussi rigide mais plus stable thermiquement tel que le poly(α-méthylène-γ-butyrolactone) P(MBL) (augmentation de la température de dégradation).

-Hydrogénation du bloc PI central (plus grande résistance en traction et stabilité thermique).

MAIN INTRODUCTION

Thermoplastic elastomers (TPE) have been defined by the International Institute of Synthetic Rubber Producers as: "Polymers, polymer blends or compounds which, above their melt temperatures, exhibit thermoplastic character that enables them to be shaped into fabricated articles and which, within their design temperature range, possess elastomeric behavior without crosslinking during fabrication. This process is reversible and the product can be reprocessed and remolded."

The development of thermoplastic materials having essentially elastic properties started in the early 1930's with the invention of plasticization of poly(vinyl chloride) (PVC) at the B. F. Goodrich Company 1 . This invention led to further interest in flexible plastics and eventually to the preparation of blends of PVC and butadiene-acrylonitrile rubber (NBR) 2,3 . The PVC/NBR blends, when properly formulated, have a rubber-like behavior and bridge the gap between liquid plasticized PVC and conventional cured elastomers. Thus, they can be considered as the precursors of TPE, as we know them today. A breakthrough occurred with the discovery of the basic diisocyanate polyaddition reaction 4 in 1937, which was first applied to produce polyurethane fibers and then to the development of some elastomeric polyurethanes at DuPont and ICI companies [5][6][7] . The work at DuPont focused on elastic fibers and eventually led to the invention of elastic linear copolyesters, prepared by melt-ester interchange between two melt copolymerized polymers 8 . These synthetic elastomers had higher strength than vulcanized natural rubbers and exhibited a faster elastic recovery. They were used to make fibers by extrusion or spinning from solutions and can be considered as the first types of TPE 9 .

The global consumption of TPE is expected to grow from 4.2 million tons in 2015 to nearly five and a half million tons in 2020 10 . This can be explained by the ability of TPE to be processed as thermoplastic materials while possessing the physical and mechanical properties of vulcanized elastomers.

Industries increasingly use TPE due to the significant cost savings arising from the TPE ability to be processed on machinery designed for thermoplastics. Conventional rubbers, whether natural or synthetic ones, are thermosetting materials that must undergo a chemical crosslinking reaction during molding or extrusion, denoted curing or vulcanization. Due to this reaction, they are not processable using conventional thermoplastic equipments. The thermoplastic molding and extrusion processes used for TPE, on the other hand, do not involve the crosslinking. Furthermore, TPE processing scrap, reject parts or end of life products can be easily reprocessed, whereas most thermosetting elastomers end up as land fill.

In addition to the advent of TPE as affordable substitute materials to traditional thermoset rubbers, the increased use of high-performance thermoplastic composites, especially in some sectors of the aerospace industry, is apparent. Thermoplastics provide several important advantages over thermosets when used in composite materials, most notably better impact damage resistance, high damage tolerance, good fatigue resistance, higher fracture toughness, higher operating temperatures, low storage cost and infinite shelf life. Thermoplastic composites offer several processing benefits over thermoset composites as well. They are reprocessable, repairable, and reformable, which together provide ease of fabrication and cost-effectiveness. Thermoplastic composites can also be deformed to complex shapes and then cooled to solidify in shorter times than required by thermoset composites, which have relatively lengthy manufacturing times due to slow cure kinetics of the thermoset resins. The group of thermoplastics that are mostly used in aircraft composite structures currently are called polyketones, and includes poly(etherketone) (PEK), poly(etherketoneketone) (PEKK) and, the most common poly(etheretherketone)

(PEEK).
This promising replacement of thermoset materials by thermoplastics, offering a larger process window, prompts engineers to rethink the joining between aerospace structures.

The research project proposed and funded by ArianeGroup for this doctoral thesis falls within the context discussed above, where the substitution of traditional thermoset materials by thermoplastics has become a major issue of the aerospace industry. One of the key activity of ArianeGroup is devoted to the manufacture of launchers which consist of a tank joined with skirts by a soft junction. The tank and the skirts are currently thermoset composites whereas the soft junction is a thermoset rubber. The joining is operated by adhesive bonding using a specific resin. Such a thermoset-based soft junction remains problematic, mostly from an economical point of view. The vulcanization of the rubber, the use of a resin and the adhesive bonding performed in an autoclave during long cycle processing times are, to name but a few, detrimental. Taking mostly into account the economical disadvantages of using thermosets and performing long and multi-step processes, ArianeGroup is studying the development of a thermoplasticbased tank/rubber/skirt assembly. Indeed, for the next-generation of launchers, ArianeGroup examines the production of tank and skirts in thermoplastic composites and the replacement of the traditional thermoset rubber used for the flexible junction by a thermoplastic elastomer. This complete change from thermoset composites and rubber to thermoplastic composites and elastomer will allow the formation of the final assembly by welding, a promising and cost-efficient joining technique.

Exclusive to thermoplastic materials, welding (or thermal bonding) has great potential for the joining, assembly, and repair of composite components. This process involves heating and melting the polymer on the bond surfaces of the components and then pressing these surfaces together for polymer 20 solidification and consolidation. Welding allows short processing cycles and offers additional advantages including reduced surface preparation requirements, reprocessing, recyclability and improved integrity/durability. Suitable for joining thermoplastic composites directly with each other, this technique is also appropriate for assembling thermoplastic composites with thermoplastic elastomers.

Due to the difficulties to find a suitable commercial thermoplastic elastomer, this work aimed to develop experimentally a novel thermoplastic elastomer exhibiting properties similar to these of the current thermoset rubber. Furthermore, the weldability of the synthesized thermoplastic elastomer directly with the substitute thermoplastic composite or indirectly by using an in-between thermoplastic layer was a prominent requirement. This doctoral work was, to a large majority, focused on the synthesis of a novel thermoplastic elastomer in order to replace the current vulcanized rubber. A specific controlled radical polymerization, termed nitroxide-mediated polymerization, was selected to produce styrenic triblock copolymer thermoplastic elastomers. Poly(β-myrcene), a bio-sourced elastomer, was first considered as the elastomeric mid-segment before being substituted by poly(isoprene). This manuscript is divided into three main parts: (I) State of the art; (II) a first experimental section devoted to the synthesis of β-myrcene-based copolymers by nitroxide-mediated polymerization (NMP);

(III) a second experimental section mainly focused on the synthesis by NMP, the characterization and the hydrogenation of isoprene-based triblock copolymers. The literature review (I) is first dedicated to thermoset rubbers and TPE to understand the will from industrials to take advantage of technologies based on thermoplastics. Styrenic block copolymer TPE, considered in this research, are then introduced more specifically as well as the polymerization processes applicable to synthesize them. The second part (II) includes three chapters. The first chapter presents the synthesis of β-myrcene/styrene (My/S) copolymers by NMP and concludes with the achievement of S-My-S triblock copolymers. The second chapter is devoted to the functionalization of P(My) by nitroxide-mediated copolymerization of My with epoxy-functional glycidyl methacrylate (GMA). The introduction of polar groups into the TPE aimed to possibly encourage the subsequent welding with relatively polar thermoplastic composites initially considered for this project.

The third chapter focuses on the synthesis of My/isobornyl methacrylate (IBOMA) copolymers by NMP.

The objective was to replace PS outer blocks by higher glass transition segments such as P(IBOMA) to enhance the toughness and the service temperature of the candidate TPE. The last major part (III) includes two final chapters. The fourth chapter is dedicated to the synthesis of isoprene-based triblock copolymers by NMP whereas the fifth chapter deals with the characterization of commercial styrenic block copolymer TPE. In these two last chapters, the partial hydrogenation of experimental and commercial TPE's PI segment is detailed as well.

INTRODUCTION GÉNÉRALE

Les élastomères thermoplastiques (TPE) ont été définis par l'Institut International des Producteurs de Caoutchoucs Synthétiques comme: ``Des polymères, des mélanges de polymères ou des formulations de polymères qui, au-dessus de leur température de fusion (ou de la température de transition vitreuse du domaine rigide), se comportent comme des thermoplastiques, ce qui permet de les mettre en forme aisément, et montrent, lorsqu'ils sont utilisés dans leur température de service, un comportement élastique semblable aux caoutchoucs thermodurcissables traditionnels alors qu'ils ne sont pas réticulés. N'ayant donc pas de réseau chimiquement réticulé, les élastomères thermoplastiques peuvent être remis en oeuvre

TECHNOLOGICAL OBJECTIVES • Introduction

Equally owned by Airbus and Safran, ArianeGroup is a world leader in access to space, providing innovative and competitive solutions for civil and military launch systems and space applications.

ArianeGroup is an expert in the most cutting-edge technologies, from all aspects of complete propulsion systems right down to the individual items of equipment and materials. One of the core businesses of ArianeGroup is dedicated to the production of launchers in thermoset composites, made by winding processes. In such a process, a stationary mandrel rotates, while a carriage arm travels horizontally up and down the length of the mandrel. As the mandrel turns, the rovings wrap around it to form a composite layer over the mandrel's surface. Before encountering the mandrel, the fibers are impregnated in a resin, which later solidifies with the fiber to create the final composite material.

The launchers are composed of a tank (vessel) in a composite joined with skirts by a soft junction.

Both the tank and the skirts are made of composite materials and undergo different deformations, as depicted in Figure 1. They must support the weight of the launcher. Moreover, the tank exhibits deformations in diameter (radial deformation) and in height. The tank can be deformed under pressure.

Accordingly, an elastomeric junction has to be used to transmit mostly shearing forces generated by these deformations. Such a joint is termed ``shearing ply``.

• Current tank-skirt assembly

The tank and the skirts are presently thermoset composites and a commercial thermoset is used as the elastomeric joint. The assembly tank/rubber/skirt is implemented by using a specific adhesive, as represented in Figure 2. Accordingly, the joining technique used is an adhesive bonding of thermosets: the adhesive is placed between the parts to be bonded (adherends) where it serves as the material that joins the parts and transmits the load through the joint. This technology relies on the covalent bonding at the interface via chemical reactions. This assembly requires critical steps which lead to increasing costs. First, vulcanization, in a specific press at high temperature and pressure, must be implemented to crosslink the rubber. Then, the assembly tank/adhesive/rubber/adhesive/skirt to bond is placed in a giant autoclave to proceed to the curing (hardening) of the adhesive. The use of an adhesive bonding technology to make the joint is detrimental due to the long assembly times. Moreover, the limited service life of the adhesive and the storage of the elastomer are non-negligible issues. Indeed, the rubber before vulcanization has a use-by date due to its reactivity (UV hardening / brittleness).

• First technological challenge: replacement of the thermoset rubber by a thermoplastic elastomer

The substitution of the current vulcanized rubber by a thermoplastic elastomer (TPE) would give significant benefits: no vulcanization step; reduction of the cost of the raw material (potentially ten times cheaper); versatility of the process based on that of thermoplastics and thus versatility of the joining technique; and unlimited storage.

The substitute TPE must exhibit inherent properties as good as those of the current thermoset rubber. Consequently, TPE specifications are required as indicated in Table 1. It must be noted that the TPE will be used under usual atmospheric conditions.

The TPE stress-strain properties desired by ArianeGroup were targeted after uniaxial tensile tests, and must be highly reproducible, and performed on the specific thermoset rubber currently used. The Young's modulus of the TPE at 100 % elongation will be important to determine since it will affect directly the thickness of the joint: the lower the modulus, the thinner the joint. A life time of minimum 20 years is mostly due to the very long period for assembling a launcher. Even though the major aim of replacing the thermoset rubber by a TPE is to allow the implementation of a more robust and cost-effective joining technique, it must be noted that the TPE will be first joined to the tank and skirt by adhesive bonding, as currently performed. Indeed, this joining process as well as the subsequent characterization, using a by double lap shear configuration for instance, is optimized in the ArianeGroup facilities. This means that the TPE must be robust (capability of preserving its properties) at elevated temperatures since an adhesive bonding would necessitate to heat at 160-180 o C for more than 10 hours.

Mechanical properties of the

• Second technological challenge: Joining of the assembly tank / TPE / skirts by welding

Even though the substitution of the thermoset rubber by a suitable TPE would constitute a major advance, the main technological breakthrough consists of the development of a new cost-effective and robust joining technique to assemble the tank, the TPE and the skirts. For future launchers, ArianeGroup studies manufacturing of tank and skirts in thermoplastic composites instead of thermoset composites.

Specific matrices reinforced with carbon nanotubes or carbon fibers are thus considered.

Thermoplastic composite technology intends to achieve improved properties and low-cost processes compared to composites containing a chemically crosslinked matrix. The use of thermoplasticbased composites and elastomer for this assembly represents a major advantage since processing, joining or shaping can be done using standard techniques used for thermoplastics in the melt state. Taking the advantage of handling thermoplastics, the production of welded structures is possible. The process of welding involves heating and melting the polymer on the bonded surfaces of the components and then pressing these surfaces together for polymer solidification and consolidation. This is an affordable method and the structures have typically a higher strength in comparison to adhesive bonding. The limited shelf lives and working lives of adhesives, the rigorous surface preparation and the lengthy cure times caused by adhesive bonding are eliminated by using a welding technique.

A resistive welding process is mainly studied. Resistance welding uses an electrically resistive element or implant, such as metal mesh or carbon strip, between the joint elements. An electrical current is applied to the resistive implant to increase the temperature of the implant due to resistance heating. The temperature of the implant increases at the interface, melts the polymer around the implant, and creates a weld under the application of pressure. Figure 3 shows a possible welding configuration.

Resistance welding can be used to weld most thermoplastic-based materials. The process is reasonably fast, simple, cheap, and clean. Since the implant remains in the bond, material compatibility between the implant and joint elements is important. To obtain a good quality weld, the resistance of the heating element should be determined to adjust the level of power input during the process. Also, controlling the heating/cooling rates improve the weld quality and mechanical performance of the weld. Resistance welding has great potential for welding high-performance continuous fiber thermoplastic composites especially for aerospace applications. Thermal insulation, input energy, welding time, fiber orientation, and type of heating element are the parameters that influence the quality and performance of the resistance welded parts.

Obviously, the direct welding between the TPE and the thermoplastic composites would be ideal.

This can be performed in principle by using the considered composites as electrically resistive elements since they contain carbon nanotubes or carbon fibers (network of conductive particles) having a concentration higher than the percolation threshold. Such a welding would lead to a tank / TPE / skirt assembly without any additional intermediary layers, as depicted in Figure 4a. However, the use of a thermoplastic composite exhibiting a high processing temperature imposes stringent thermal requirements for the welding, which would require temperatures that would probably degrade the TPE. The addition of another thermoplastic composite, situated between the TPE and the tank/skirts, can therefore be envisioned for the resistive welding process. This intermediate material, requiring a lower processing temperature and fully miscible with the thermoplastic composite considered initially, would impose lower thermal requirements. Such an assembly is depicted in Figure 4b.

• Objectives of the thesis

This research project aimed to synthesize a novel thermoplastic elastomer capable of providing a tough and robust joining between the tank and the skirts of the launcher. This TPE must exhibit stress-strain properties comparable to these of traditional thermoset rubbers at room temperature. Moreover, a limited decrease of its mechanical properties at high temperature is desired (continuous service temperature = 10- Whatever the joining technique considered for assembling the TPE with the tank and the skirts (adhesive bonding or welding), the assembly must be stable and robust with a high tensile strength at break as well as an elevated shear modulus at 100% elongation. The cohesive fracture of the connection is also required. In the case where adhesive bonding is implemented, the candidate TPE would be heated at 160-180 o C for more than 10 hours to be effectively assembled with the adhesive. If welding is applied, significantly shorter processing time would be needed (a few minutes possibly) but higher temperature may be applied (240-280 o C). In any case, the TPE must not degrade nor flow at the joining temperature. In addition, it must be capable of preserving its properties after the bonding step.

MAIN OBJECTIVES OF THE RESEARCH PROJECT

• Introduction h Many independent routes could have been considered for this doctoral thesis to produce a thermoplastic elastomer (TPE) satisfying ArianeGroup specifications. Due to time constraints and in order to follow a relevant route, privileged axes were selected for the following key points of the research: method of synthesis (anionic, radical, cationic…) ; type (block copolymer, rubber/plastic blend…) and nature (styrenic, polyurethane, polyester, polyamide…) of the TPE ; composition (monomers and their respective molar fraction) and morphology (spherical, cylindrical, lamellar…) of the TPE ; solutions to enhance the TPE properties (thermal stability, service temperature, mechanical behavior, weldability…). This part aims to outline the major directions chosen during this research project. Please note that no references have been added in this section purposely since each scientific point is developed and detailed more exhaustively in the next parts, particularly in the state of the art.

• Polymerization technique

Nitroxide-mediated controlled radical polymerization (NMP), a reversible-deactivation radical polymerization (RDRP), was selected. A controlled radical polymerization was thereby preferred to another polymerization technique, such as a living ionic polymerization process, since it does not require stringent process conditions. Moreover, the difficulty to polymerize functional monomers by living ionic polymerization was a hurdle for this project since the functionalization of the candidate TPE was envisioned to enhance its joining properties. In contrast to reversible addition-fragmentation chain transfer (RAFT) polymerization and atom transfer radical polymerization (ATRP), two other RDRP processes, purely organic systems can be employed when using NMP. NMP does not require the use of sometimes unpleasant sulfur chemistry that is associated with the chain transfer agents used in RAFT. NMP systems, comprising radical initiator / nitroxide or alkoxyamine with or without additional nitroxide, are also typically simpler initiating systems compared to ATRP initiating systems, which require, at minimum, activated halide, metal halide, and ligand, but often include other additives to facilitate the polymerization. An additional advantage that NMP offers over ATRP and RAFT is in the ability to control the polymerization of conjugated 1,3diene monomers, a first essential step to produce the desired TPE. For some time, NMP was the only RDRP method that could be used to prepare homopolymers and block copolymers of 1,3-dienes such as isoprene and 1,3-butadiene. It was thus possible to rely on several publications reporting isoprene and butadiene polymerizations by NMP. Lastly, being co-supervised by Prof. Maric, it also appeared crucial for me to perform NMP in order to take advantage of his expertise in this specific field. The presence of commercial nitroxides and alkoxyamines in his laboratories and the competence of his research team to synthesize welltailored monofunctional and difunctional initiators for NMP was beneficial as well. It must be added that NMP under homogeneous conditions in bulk or in organic solvents was favored. Even though promising, NMP in dispersed media (mostly water) was not implemented because of the time it might have required to find the optimal recipe.

In conclusion, NMP was essentially chosen for its simplicity to implement, its robustness for controlling the polymerization of 1,3-dienes and the knowledge of Prof. Maric's laboratories for this nitroxide-based RDRP system.

• Type and class of TPE

Block copolymer TPE were targeted in this thesis since, as mentioned previously, a controlled radical polymerization was selected, which appeared totally adequate to synthesize well-tailored block copolymers. Thermoplastic vulcanizates (TPV) using dynamic vulcanization, polyamide elastomers synthesized by esterification and polycondensation, copolyester elastomers (COPE) manufactured by transesterification or even reactor thermoplastic polyolefins (RTPO) produced by catalytic copolymerization of olefins, were for instance disregarded. More specifically, styrenic block copolymers (SBC) were chosen considering the large variety of vinyl monomers being polymerizable by NMP and leading to the achievement of flexible, hard or even functional blocks. A triblock structure was desired since it allows generally the TPE to have higher tensile strength compared to a similar diblock copolymer, in which the rubber chains are attached to a rigid chain at only one end and are therefore not physically cross-linked. Moreover, a plethora of styrenic triblock copolymers, reported in the literature or available commercially, can be used as references. Most of them satisfy the ArianeGroup mechanical requirements at room temperature, which is a good starting point. While exhibiting typically low upper service temperature limit, modifications of the SBC in the laboratory were planned to enhance its thermal robustness as discussed below. Even though investigated by RDRP, TPE with graft (hard segments are chemically bonded to the elastomer-like main chain as pendant chains), star (rigid arms and rubbery core)

or brush (glassy backbone and rubbery brushes for example) topologies were not considered.

From an experimental viewpoint, a difunctional initiator (Di-I) was favored to prepare triblock copolymers. It can be used to make the rubbery segment. This way, since the two chain ends are ideally capped by a nitroxide moiety, the flexible block synthesized can be subsequently used as a macro-initiator to polymerize a second batch of monomer(s) in order to add two outer rigid blocks.

• Composition and morphology of TPE

Two conjugated 1,3-diene monomers were selected to produce the flexible middle block, namely β-myrcene (My, 7-methyl-3-methylene-octa- -4 °C) making the polymerization more demanding and its similarity with isoprene monomer. About the subsequent addition of the two outer hard blocks, poly(styrene) PS was first selected as a reference since it is used in a large majority of styrenic block copolymers. Furthermore, the NMP of styrene was comprehensively studied by Georges, Hawker, Charleux and Gigmes amongst others.

Micro-phase separation is expected due to the immiscible polymer segments of the TPE. The preparation of TPE showing a spherical morphology (hard blocks forming discrete minority spheres embedded in a rubbery matrix) was selected due to their isotropic nature. Indeed, since the spherical domains are randomly dispersed, the stress-strain behavior exhibits isotropic mechanical property. In that case, triblock copolymers having about 10-20 mol% of hard blocks were targeted (arbitrary percentages since it mostly depends on the nature of the segments used). However, higher hard blocks content going up to 40 mol% were also envisioned (cylindrical, gyroid or lamellar morphologies assumed) to increase the hardness of the TPE if needed.

• Solutions to enhance TPE properties

After achieving optimized NMP-based styrenic triblock copolymers, modifications can be performed so that the industrial requirements can be satisfied. First, the functionalization of the TPE can be implemented since this latter is supposed to be welded to relatively polar thermoplastic composites. The goal was thus to increase the miscibility at the welding interface to favor chains inter-diffusion. Even if an adhesive bonding is first processed with thermoset composites, the incorporation of functional units into the TPE could be valuable to promote covalent bonding at the interface. Considering the distinct advantage of NMP about its robustness to readily add functional groups, it was decided to perform the TPE polar modification by copolymerizing directly the diene monomer with glycidyl methacrylate (GMA), carrying an epoxide group.

Secondly, poly(styrene) blocks (glass transition temperature, Tg ~ 100 o C) initially considered can be substituted to higher Tg hard blocks to improve the tensile stress and the service temperature of the TPE. The active feature of NMP-based polymers is based on the reversible capping of the active chainend radical with a nitroxide leaving group. At the end of the synthesis, chain-ends should thus be ideally terminated by a thermally labile nitroxide moiety. Removing the nitroxide groups after synthesis, which will prevent their liberation during subsequent heating processes (extrusion, injection, welding, adhesive bonding, hydrogenation…), can enhance thermal stability. Accordingly, this step appeared necessary for our optimized TPE. Many efficient strategies have been reported to remove the terminal nitroxide and can be directly applied in our situation. Lastly, a second post-polymerization process termed hydrogenation was envisioned regarding the soft poly(diene) middle block. Achieving a TPE having a saturated mid-segment could improve its resistance to the thermal and oxidative degradations. Moreover, higher TPE tensile strength can be expected after hydrogenation, which can be valuable for this industrial problematic. hydrogenation of the soft block (higher tensile stress and degradation temperature). Figure 6 shows schematically these various modifications. The notations from 1) to 4) correspond to the chronological order to perform the modifications.

Poly(isobornyl methacrylate) P(IBOMA) is mentioned herein as an example to replace the PS outer blocks.

THESIS OUTLINE • A collaborative project

• Structure of the PhD manuscript

A. Introduction

Thermoplastic elastomers (TPE) combine the functional properties of comparable thermoset elastomers with the manufacture advantages of thermoplastics. They are defined as ``a family of rubberlike materials that, unlike conventional vulcanized rubbers, can be processed and recycled like thermoplastic materials`` (ASTM D1566). These materials of interest are highlighted in this state of the art, organized according to the industrial problematic. Three main sections constitute the literature review.

The first section (B, pages 47-59) is dedicated to the thermoset elastomers and block copolymer TPE and highlights the industrial benefits of using thermoplastic polymers in order mainly to implement advantageous joining processes. After defining the notion of elasticity, the state of the art introduces the traditional thermoset elastomers, used massively in the industry, and their main shortcomings. A focus on nitrile rubber (NBR), a commercial vulcanized elastomer currently used by ArianeGroup to provide the tough joining in question between thermoset composites, is supplied as well. Block copolymer TPE, which can be seen as a way of substitution of the vulcanized rubbers, are then presented. This transition from thermoset to thermoplastic polymers follows the desire of ArianeGroup to replace thermoset elastomer/composite by thermoplastic materials. It should be noted that three major types of TPE exist, namely: block copolymers; rubber/plastic blends; dynamically vulcanized rubber/plastic alloys denoted as thermoplastic vulcanisates. This literature review is intentionally focused on the block copolymer TPE, to narrow down the scope of the research area and to be in line with the industrial project. The concept of physical crosslinks, the phase structure and the phase separation of these copolymers are introduced to understand their peculiar thermoreversible networks. Furthermore, the main advantages of block copolymer TPE over thermoset rubbers are given to support the desired changeover from conventional thermosets to new thermoplastics exhibiting similar properties. Eventually, joining techniques, corresponding to one of the main industrial key issue, are presented for thermosets and thermoplastics. The possibility to perform welding (or thermal bonding) for thermoplastics, a promising method offering many advantages over the other joining processes, is underlined.

The second section (C, pages 60-80) focuses on the general properties of styrenic block copolymers (SBC), the family of TPE selected for this project. It is first of importance to know how to optimize the formulation of these materials to achieve the desired properties. The influences of the nature of the poly(diene) middle block, the average chain length and the styrene fraction over the characteristics of these TPE are thus discussed. Moreover, the glass transition temperature (Tg) of PS blocks limiting the upper service temperature of styrenic block copolymers, the substitution of the typical PS segments by higher Tg hard blocks is then presented. After synthesis, post-polymerization modifications of styrenic TPE can be envisioned as well to strengthen its properties. Hydrogenation of the poly(diene) midsegment, introduced subsequently, allows the resulting TPE to be mostly more stable thermally and tougher. Considering that the candidate TPE has to support welding or adhesive bonding process at T > 160 o C, such an enhancement could be valuable. Eventually, attention is paid to directly available commercial SBC, which can be used as references for the experimental block copolymers synthesized in laboratory.

The last section (D, pages 81-88) of this literature review is devoted to the various polymerization techniques used to make styrenic block copolymer TPE. Living anionic polymerization, proceeding by formation of carbanions at the polymer chain ends, was first introduced. Applied industrially, it is the most robust and effective technique to manufacture well-defined styrenic TPE such as SBS and SIS. Living cationic polymerization, allowing the synthesis of well-defined poly(isobutylene), poly(styrene) and poly(styrene) derivatives, can also be employed to prepare some interesting block copolymers. Despite the excellent results obtained from living ionic polymerizations, several main constraints appear regarding our industrial project: the stringent and demanding experimental conditions required combined; the nonnegligible number of preliminary steps to follow prior to the polymerization; the difficulty to polymerize functional monomer. Accordingly, the use of a reversible-deactivation radical polymerization (RDRP), which does not require demanding experimental conditions and can be apply to a large range of monomers, is promising. The synthesis of block copolymer TPE via the three main RDRP processes, namely atomtransfer radical polymerization (ATRP), reversible addition-fragmentation chain transfer (RAFT) and nitroxide-mediated polymerization (NMP) are presented. Due to its capability to control effectively 1,3diene monomers and its simplicity of implementation, NMP appears to be an appropriate synthesis method for the industrial problematic.

B. From thermoset to thermoplastic elastomers (TPE) B.1. Elasticity and thermoset elastomers

Please note that the terms ''elastomer'' and ''rubber'' are mentioned herein in the exact same way as a polymer showing the property of ''viscoelasticity'' (colloquially ''elasticity'', defined below).

B.1.1. What is an elastomer?

Considered as the most versatile of engineering materials, elastomers are an important class of polymers exhibiting randomly distributed chains connected by crosslinks. They behave very differently from plastics and metals, particularly in the way they deform and recover under load. Elastomers consist of relatively long polymer chains, which by their high degree of flexibility and mobility allow a significant deformation 1 .

One of the most important property of an elastomer is its elasticity which can be defined as the ability of the rubber to return to its original shape and size after being stretched, compressed, twisted or bent 2 . Elastic deformation lasts only as long as a deforming force is applied and disappears once the force is removed. When a stress is applied, the configuration of the chains may be changed rapidly due to the high mobility of the elastomer. Thanks to the crosslinking of the network (formation of a three-dimensional structure from a linear polymer by a chemical or physical method) 3 , the macromolecules are prevented from flowing relative to each other under stress. On removal of the external force, the rubber is rapidly restored to its original dimensions, with essentially no residual or non-recoverable strain. Without the crosslinks, the applied stress would result in a permanent deformation. It is of importance to underline that the response of the rubbery polymer is intramolecular. The force is transmitted to the chains through the linkages at their extremities. The chains exhibit then a different conformation and act like an individual spring, according to Hooke's Law 4 . The elastomer undergoes immediate, linear and reversible high strain to an applied force.

The extensibility is caused by uncoiling of the random molecular coils (high entropy) into more linear conformations (lower entropy). The limit to elastic response is reached when the macromolecules are fully extended (one conformation resulting in a minimal entropy) 5 . A representation of the two different conformations is depicted in Figure 7.

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Another essential feature of an elastomer is its excellent resilience. Resilience as applied to elastomers is essentially their capacity to return rapidly to their original shape after temporary deflection.

It indicates thereby the speed of recovery, unlike compression set, which measures the degree of recovery.

When an elastomer is deformed, an energy input is involved, part of which is not returned when it regains its original shape. This part of the energy which is not returned is dissipated as heat in the elastomer. The ratio of energy returned to energy applied to produce the deformation is defined as the material's resilience 6 .

Furthermore, reversibility and immediate response are obtained with macromolecules that have flexible chains with weak intermolecular forces. The large reversible strain must be rapid which means the restraining intermolecular forces must be minimal. Consequently, polar groups and hydrogen bonding groups are not desirable if a polymer is to be an elastomer. Steric hindrance via bulky pendant groups or rigid intra-chain groups has to be minimal as well to encourage the uncoiling. The arrangement of the macromolecules and their configuration is also of importance. The amorphous nature of the elastomer results from the irregularity of the chains due to the difference in geometrical arrangement. Formation of unstrained random coils means that the elastomer must be non-crystalline since any regular crystal structures would be unable to contribute to elastomeric properties. These latter are optimized with the cis configuration although the motion of the bonded carbon atoms is limited by the double bond. The trans configuration favors a regular planar zig-zag conformation that may crystallize 5 .

The majority of elastomers have a number of other interesting features such as: low permeability to air, gases, water and steam; good electrical and thermal insulation; good mechanical properties; ability to adhere to various fibers, metals and plastics. It should be added that rubbers with improved or specific properties can be produced to meet a wide variety of service conditions by adequate selection of compounding ingredients. This combination of properties is the reason elastomers are used in a large number of engineering needs in fields dealing mainly with sealing, shock absorbing, vibration damping, and electrical and thermal insulation 5 . Historically, most elastomer components were manufactured using vulcanization, which prompts us to introduce the family of the thermoset elastomers.

B.1.2. Thermoset elastomers

Most types of elastomers are thermosets. It consists of single-phase materials, without the dual hard and soft phases, which can be polymerized as homopolymers or random copolymers/terpolymers. The structure is amorphous, exclusive of crystalline domains, which means that these rubbery materials undergo a glass transition temperature, but no melting point.

Thermoset elastomers gain most of their strength after vulcanization, an irreversible crosslinking of their polymer chains that occurs when the compound is subjected to pressure and heat in the presence of sulphur 7 . The polymer is traditionally crosslinked after molding or shaping to fix molecules into their positions. A thermoset elastomer typically consists of a rubber-like polymer containing a reinforcing filler, a plasticizing agent, one or more stabilizers, a crosslinking agent, commonly sulfur, several accelerators which increase the rate of crosslinking, and additives as needed. The crosslinking reaction is a rate limiting step in the manufacturing process to form conventional thermoset rubber articles. Much of the attention paid to forming these articles has focused on developing the formulations, known as compounds. These compounds must achieve the performance requirements and also chemically react in sufficiently controlled fashion to crosslink as quickly as possible at the time desired in the processing, but not prior 7 . The vulcanization of cis-1,4-poly(isoprene) is illustrated in Figure 8.

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applications 13 . Thermoset rubbers require mixing of curatives, stabilizers, process aids and other additives.

The part-to-part consistency is not optimal due to the compound difference that can be caused by the variations in weighing and metering of the various components. On the other hand, little to no compounding or mixing is required for thermoplastics. The thermoset rubber process includes several steps such as mixing, shaping and vulcanization, which can make the operation costly and reduce the reproducibility of the fabrication. The processing operation for a thermoset has longer cycle times due to the type of process involved (compression molding) and the need for vulcanization. Typically, thermoset process cycle times are measured in minutes whereas thermoplastic ones are in seconds. Consequently, the production rate for thermoplastic materials, whose process uses less energy, is much greater. While a thermoplastic scrap can be reground and reprocessed with no significant loss in performance, a vulcanized rubber cannot be reprocessed. It must be ground to a powder or crumb form to be recycled and then only be used as a filler material. It is traditionally treated as scrap and must be disposed of as waste since thermoset rubber process scrap has limited reuse capability. This limitation has been a source of economic and environmental concern for conventional thermoset rubber. Expended thermoset rubber goods pose an end-of-life disposal concern given their lack of recyclability. Since more scrap is produced in thermoset rubber processing, the production cost is made higher by the loss of material and disposal cost of the scrap 14 .

In direct connection with the industrial problematic, another major inconvenient can be mentioned: some thermoplastic processes, such as bloc molding, thermoforming and heat welding, are not possible with thermoset rubbers. In the manufacturing of aerospace composite structures, thermal bonding or welding has great potential for the joining and offers many advantages over other joining techniques.

However, this method cannot be applied to thermoset materials since they cannot melt. It follows that thermoset and vulcanized rubber components can only be joined using adhesive bonding or mechanical fastening methods 15 .

Taking these shortcomings into account, it is legitimate to look for high-performance rubbers which are not chemically crosslinked and exhibiting properties as good as those of vulcanized rubbers.

B.2. Block copolymer TPE

Please note that the following literature review only focuses on block copolymer thermoplastic elastomers (TPE).

B.2.1. Physical crosslinks

Polymers having a significant average chain length generate entanglements by molecular intertwining with a spacing in the bulk state depending on the particular macromolecular structure. A polymer melt with a high degree of polymerization can thus exhibit transient rubber-like behavior even in the absence of any permanent intermolecular bonds 16 . In a crosslinked elastomer, many of these entanglements are permanently locked in and, at high enough degree of crosslinking, they may be regarded fully equivalent to crosslinks, contributing to the elastic response of the material. Generally, a network is obtained by linking of polymer chains together. The chemical and permanent crosslinking of elastomers, mostly via vulcanization, was discussed above. Physical crosslinks, which can be defined as non-covalent and thermo-reversible intermolecular bonds holding polymer chains together in network under some conditions but not under others, can be generated as well. Physical crosslinks interconnect polymer chains using interactions other than covalent bonding, such as hydrogen bonding, hydrophobic interaction, ionic bonding, and coordination bonding. The key parameter is temperature. The bonds are strong at low temperatures, but the materials can flow at higher temperatures due to the disappearance of these physical crosslinks. These latter can be mostly formed in three distinct polymeric networks:

-Partially crystalline polymers for which the crystalline regions are embedded in the amorphous matrix and serve as physical crosslinks. For instance, the crystalline regions in low density polyethylene can behave in this way. The flexible chains surrounding the crystallites impart extensive deformability to the material 17 .

-Ionomer formation where ionic groups such as sodium methacrylate may be introduced along a polymer chain. At a level of about 5 mol%, a segregation takes place with the ionic groups forming small ionic domains 18 .

-Two-phase block copolymers comprising of a major proportion of a soft segment, affording the flexible elastomeric nature, and a minor proportion of a hard segment, providing both physical cross-link sites and filler reinforcement characteristics 19 .

This latter polymeric network is characteristic of a large group of polymers denoted as block copolymer thermoplastic elastomers (TPE). They can be defined as block copolymers with thermosreversible crosslinks which can be processed as thermoplastics (by melt processing) and exhibiting elastic behavior similar to that of chemically crosslinked conventional elastomers in a certain temperature range 20 .

Attention can be now paid to the peculiar structure of block copolymer TPE.

B.2.2. TPE microstructure

Block copolymer thermoplastic elastomers (TPE) display phase-separated morphologies. One phase displays hard behaviour at ambient temperature whereas the other is an elastomer exhibiting a subzero glass transition temperature (Tg). The phases are bonded chemically by block polymerization. The hard phase gives these TPE their strength and leads to physical crosslinks. Without the physical crosslinking, the elastomer phase would be free to flow under stress and the polymer would be practically unusable. On the other hand, the elastomer phase provides flexibility and elasticity to the system. When the hard phase is melted or dissolved in a solvent, the material can flow and can be processed by usual thermoplastic processing. The individual polymers constituting the respective phases retain most of their characteristics so that each phase exhibits its specific Tg or crystalline melting temperature (Tm). Tg and Tm are the temperatures at which the particular elastomer goes through transitions in its physical properties 21 . Figure 10a shows briefly the measurement of the flexural modulus versus temperature for block copolymer TPE.

There are three distinct regions:

-Below the glass transition of the elastomeric phase, the copolymer is stiff and brittle since both the phases are glassy. -Above its Tg, the elastomeric phase softens, and the material is elastic, resembling a conventional vulcanized rubber.

-As the temperature increases, the modulus stays relatively constant (a region referred to as "rubbery plateau") until the point where the hard phase softens or melts. At this point, the material becomes a viscous fluid. Above the crystalline melt temperature or the glass transition temperature of the rigid domain, the block copolymer chains are no longer locked into position and all the chains are free to flow.

Accordingly, the service temperature range lies usually between the Tg of the elastomeric phase (lower service temperature) and the Tg (or Tm) of the hard phase (upper service temperature). A simple representation of the structure of a block copolymer TPE is given in Figure 10b. The hard crystalline regions act as physical crosslinks and the soft rubbery phase provides the overall flexibility of these copolymers.

The rubbery phase does not crystallize with the rigid blocks and form a continuous domain of softer chains.

These flexible macromolecules are held together by the crystalline domains, which allow to lock the copolymer chains in specific locations. When these copolymers are deformed, the hard blocks remain crystalline and do not deform. The soft rubbery phase is easily deformed and provides elastic behavior. The recovery of these materials is good as long as the domains are not strained too greatly, and the temperatures are well below the Tg (or Tm) of the hard segments 22 .

B.2.3. Micro(nano)-phase separation

To obtain valuable mechanical properties in a two-component polymeric system, the two components must not be totally incompatible nor fully soluble (homogeneous phase). Most of systems are compatible to the extent that a minimum degree of mixing is apparent or interfacial bonding is generated directly, such as in graft or block copolymers. The polymer incompatibility is due to the very small entropy gained after mixing different sorts of macromolecules. In the limit of high average molecular weight, only polymer pairs with zero or negative heats of mixing form one phase. Generally, materials mix to form a single-phase system if the free energy of mixing (∆Gm) is negative. ∆Gm can be expressed in terms of enthalpy of mixing (∆Hm) and entropy of mixing (∆Sm):

∆Gm = ∆Hm -T∆Sm
where T is the temperature. Phase separation occurs if ∆Gm > 0, thus, ∆Hm > T∆Sm. It has to be considered that ∆Hm for hydrocarbon polymers is almost always positive because of the lack of strongly interacting 56 groups. ∆Hm increases as the structures of the two polymers forming the segments become less alike. The temperature and the entropy of mixing will always be positive, -T∆Sm will always be negative.

Nonetheless, -T∆Sm will tend to zero as the average molecular weights of the blocks become high and/or as the temperature decreases. Accordingly, it appears obvious that three main parameters favor phase separation: blocks with highly different structures; blocks with high average molecular weight; low temperature 23, 24 .

Unlike blend of polymers, where the different polymers can separate at macroscopic scale, the separation occurs at the microscopic (or nanoscopic) scale for block copolymers. This is explained by the covalent bond linking the blocks of different polymers, which forces them to regroup in smaller domains.

The two (or more) distinct and incompatible segments provide unique solid state and solution features to block copolymers, which can then be used for a variety of applications. The microphase separation in block copolymers gives rise to formation of various types of microdomains, including spherical, cylindrical, lamellar, and others. The phase behavior of undiluted A-B diblock copolymers is essentially determined by where α and β are constant depending on composition and architectural constrains of the block copolymer 25 .

At equilibrium, the block copolymer macromolecules are in the lowest free energy configuration. Increasing χAB by lowering the temperature encourages a reduction of the contact between A and B. If N is high enough, it is accomplished with some loss of translational and configurational entropy by local compositional ordering. Such local segregation is referred to as microphase separation in the block copolymer.

Alternatively, if χAB or N is low enough, the entropic factor will be predominant, resulting in a compositionally disordered phase. Since the entropic and enthalpic contributions to free energy density scale respectively as N -1 and χAB, the product χAB N, termed the reduced interaction parameter or lumped interaction parameter, imposes the block copolymer phase state. When χAB N is higher than a specific value for the considered system, the microphase separated structure evolve below which the system is in the disordered state. This phenomenon is called order-disorder transition (ODT) 26 . As the temperature is raised above a certain critical value, the microdomain structure disappears completely, giving rise to a disordered homogeneous phase. The temperature at which the microdomain structure completely disappears is referred to as the order-disorder transition temperature (TODT). This transition is depicted in Figure 11 for a symmetric A-B diblock copolymer.

For a symmetric A-B diblock copolymer (i.e., FA = 0.5), the transition occurs when χAB N ≈ 10.5.

At sufficiently large values of χAB N, different ordered structures are obtained and are represented in the form of a phase diagram (Figure 12). For χAB N << 10.5, the copolymer melt is disordered and homogeneous and the A-B interaction is sufficiently low that the individual chains assume unperturbed Gaussian statistics. At χAB N ≈ 10.5, a weak demixing between the blocks A and B is apparent. The repulsion between the blocks is strong enough that phase separation occurs. This regime is called weak segregation limit and the copolymers showing this behavior are characterized by a widened interface due to a phase mixing. As the value of χAB N is greater than 10, nearly pure A and B domains are formed. The chain conformation is no longer Gaussian but rather perturbed. This regime is termed as strong segregation limit and is caused by the higher immiscibility of the two different segments resulting in a reduced interface.

The stretching of the chains is increased 26 .

A variety of structures develop in block copolymers upon microphase separation. This phenomenon is due to two main effects. First, dissimilar blocks tend to segregate because of their inherent chemical incompatibility. The spatial extent of phase separation is, however, limited by the connectivity of the blocks dictated by the architecture of the macromolecules. As a compromise of both the effects, periodic microstructures evolve. The geometry of the microphase separated structure is, therefore, very sensitive to the chemical nature and molecular structure of the copolymer as well as its total composition 26 .

58

After investigating the phase structure and the phase separation of these particular block copolymers, it is of interest to explore their specific features compared to thermosets exhibiting similar properties in order to consider block copolymer TPE as replacement materials.

B.2.4. Advantages / disadvantages of block copolymer TPE

Block copolymer TPE offer some valuable advantages in comparison to vulcanized rubbers. First of all, the process required is easier with fewer steps. Indeed, TPE rely on processing techniques used for thermoplastics, more effective and less costly. Consequently, the final cost of the finished material is lower.

The process has also shorter fabrication times (injection molding, rapid and highly economical, is generally used for thermoplastics) which allows to reduce finished part costs as well. Since molding cycles for TPE last usually several seconds as opposed to minutes for conventional thermoset rubbers, the productivity is significantly increased. It can be thus stated that the thermoplastic process uses typically less total energy (lower energy consumption) due to shorter molding cycles and simpler processing. Secondly, since the majority of TPE is supplied fully formulated and ready for fabrication, compounding or mixing are rarely required, with no need to add reinforcing agents, stabilizers or cure systems. Hence, batch-to-batch variations in weighting and metering components are absent, leading to improved consistency in both raw materials and fabricated articles. The simple formulations used combined with straightforward processes allow a better-quality control and closer tolerances of finished parts. Furthermore, the more important reproducibility and consistency of properties of TPE resins result in reduced quality control costs. It can be noted that TPE processes are adequate for high speed automation and amenable to robotic assistance while thermoset rubber processing is usually much more laborious. The possibility of reusing scrap in the same fashion as with thermoplastics is an important benefit as well. The TPE scrap can be reused as a regrind frequently producing materials having the same properties as the virgin material (minimal loss in performance). Lastly, a paramount advantage of TPE over conventional rubbers is the possibility to use specific thermoplastic processes such as blow molding, heat welding and thermoforming 14, 20 .

Nonetheless, some detriments have to be mentioned when handling block copolymer TPE. The upper service temperature of these latter is limited due to the glass transition (or melting) of the hard phase.

This inherent property limits the use of block copolymer TPE to service temperatures well below their upper glass transition temperature or melting point. Moreover, the raw materials to synthesize the TPE can be expensive and thus the cost of TPE materials are generally higher than the thermoset rubbers. TPE have lower temperature resistance, thermal stability, chemical resistance and worse form recovery (high compression set) after being under load. Although it is a common step for thermoplastic polymers, drying prior to processing is required for TPE whereas it is mostly never performed for vulcanized materials 14 .

Although some shortcomings such as the relatively low heat resistance and thermal stability can be raised when using TPE, a major advantage is that processing, joining or shaping can be done using standard techniques used for thermoplastics in the melt state. Styrenic block copolymers are defined by a specific polymeric structure consisting of multiple blocks of poly(styrene) (PS) connected by segments that are essentially in rubbery state. They exhibit mostly A-B-A triblock structures, where A is a PS segment and B is an elastomeric segment. This latter soft segment is usually a polydiene, such as poly(butadiene) (PB) or poly(isoprene) (PI). Please note that the more convenient and more widely used nomenclatures for these copolymers are S-B-S and S-I-S.

These polymers form two separated phases as discussed previously, retaining many of the properties of the respective homopolymers. Accordingly, they exhibit high strength combined with elastomeric characteristics without the use of vulcanization chemicals. Styrenic block copolymers have two distinct glass transition temperatures (Tg). At room temperature, PS is rigid and strong whereas the flexible phase is elastic and can be extended. The PS phase is typically present in a minority and therefore consists of separate regions. These glassy domains are attached to the ends of elastomeric chains and create multifunctional junction points similar to chemical crosslinks (the elastomeric midsegments are tied together in a three-dimensional network). However, these crosslinks are of a physical nature and are therefore less stable thermally. At ambient temperature, these triblock thermoplastics behave in many ways like vulcanized rubbers. At T > Tg,PS ≈ 100 o C, PS domains soften, the network becomes weaker, and eventually the material is capable of flowing. Likewise, the domains totally lose their strength when the material is dissolved in a solvent (styrenic block copolymers can typically be dissolved in solvents in which both respective homopolymers are soluble). When the polymer is cooled down or the solvent is evaporated, the domains harden and the network regains its original integrity 20 . Most of the currently used hard and soft segments of TPE based on styrenic copolymers with their respective Tg are given in Figure 13.

It has to be added that branched styrenic block copolymers with the structure (S-E)nx (E = elastomeric segment; x = junction point with a functionality of n) have similar properties. However, structures exhibiting only one rigid segment such as I-S-I or B-S-B are not capable of forming continuous networks because only one end of each polydiene chain is terminated by a polystyrene block and the resulting materials are weak with no resemblance to conventional vulcanized rubber 22 .

Commercially, several flexible midsegments, namely PB, PI, poly(ethylene-butylene) P(EB), poly(ethylene-propylene) P(EP), poly(isobutylene) P(IB) and poly(ethylene-ethylene-propylene) P(EEP), have been used in association with the two PS outer segments. Styrenic block copolymers are the largest volume and one of the lowest priced commercial TPE 20 .

C.1.2. Effects of the elastomer type

Since styrenic TPE consist of two incompatible phases at service temperature, the final properties of these polymers are influenced by the properties of the individual phases and the interface. It is therefore necessary to consider several factors. As regards styrenic block copolymers, the constituents' nature (soft and hard segments), their molecular weight and relative molar fraction are crucial parameters to optimize in order to obtain valuable TPE characteristics.

The change in the styrenic TPE center-block results in difference in properties. Poly(butadiene) PB and poly(isoprene) PI have one double bond in the original monomer unit, which are prone to a chemical attack and limit the thermal and oxidative resistance of the SIS and SBS block copolymers. In contrast, poly(ethylene-co-butylene) P(EB) is saturated and SEBS copolymers are much more stable than the former. The modulus of these block copolymers should be inversely proportional to the molecular weight between the entanglements (Me) 27 . Some values of Me for various possible elastomeric segments 28 are given in Table 3. The Me for P(EB) can be assumed to be similar to that for poly(ethylene-co-propylene).

Poly(diene) Me (g.mol -1 )

Poly(ethylene-co-butylene) (a) In comparison, SIS copolymers have much lower viscosity than SBS copolymers since the PI chains have much higher entanglement molecular weight than the PB chains. On the other hand, SBS copolymers display a much higher elastic modulus and better low temperature properties. Thus, most SIS copolymers are traditionally used in hot-melt pressure sensitive adhesives while SBS copolymers are mainly used in footwear and asphalt modification. Due to a more entangled elastomeric continuous phase, SBS typically exhibits higher tensile stress than SIS having a similar composition. This is illustrated in Figure 14 which compares tensile properties at room temperature of SBS and SIS samples having similar Mns and styrene contents 29 . Although SBS samples shown slightly reduced elongations at break, mainly due to their greater hardness, compared to SIS triblocks, SBS tensile stresses at break were significantly higher than those of SIS polymers. The soft segment type also influences the driving force for phase separation and hence the physical properties. domains can be theoretically achieved at lower weight percentage of styrene. Hydrogenation of the diene phase reduces the solubility parameters of the elastomer phase and allows the formation of a stronger phase separation. Therefore, SEBS can be copolymerized for phase separation with less styrene blocks Mn while still obtaining a good phase separation, compared to unsaturated SIS and SBS TPE 32 .

C.1.3. Effects of average molar mass and hard blocks content

The condition for complete miscibility of two substances is that the Gibbs free energy of mixing (∆GM) be favorable, that is, negative. On the other hand, if ∆GM is positive, phase separation will occur.

In the case of block copolymers, a positive value of the free energy of mixing means that separate domains will be formed. The free energy of mixing can be expressed as ∆GM = ∆HM -T∆SM where ∆HM is the enthalpy of mixing, ∆SM is the entropy of mixing, and T is the absolute temperature. For styrene-diene block copolymers, the enthalpy of mixing is positive because there are no strongly interacting groups and it increases as the two polymers forming the segments become less alike. T and ∆SM will always be positive and consequently the product (T∆SM) will always be negative. This term will approach zero, however, as the molar masses of the segments become large and/or as T decreases 23, 24 . Thus, the domain formation will be favored by the following conditions:

-High degree of structural difference between the segments -High average molar masses of the segments -Low temperature Based on this approach, the values of critical molar masses and temperatures for domain formations have been developed and can be predicted quite well. Findings reported from studies 33,34 are:

-SBS triblock copolymer with PS end-segments having Mn of 7 kg.mol -1 changes to a one-phase system at about 150 o C.

-Similar block copolymers with PS end-segments having Mn of 10 kg.mol -1 and greater appear to be separated at temperatures up to 200 o C.

-The critical molar mass for domain formation in SBS appears to be about 7 kg.mol -1 .

Styrenic TPE must also exhibit a highly entangled elastomeric phase to be mechanically valuable.

The physical chain entanglements contribute mainly to the network density and to the stress at small elongations. The soft block should thereby have a molar mass which is at least several times higher than the molar mass between entanglements Me (Table 3 above). Highly entangled poly(diene)s show ductile fracture due to the energetically favorable shear yielding deformation mechanism. Low entanglement density of the soft phase likely leads to a brittle fracture behavior of the styrenic TPE.

SB become harder and stiffer as the proportion of the hard styrene segments is increased. The stress-strain behavior of otherwise similar block copolymers with a wide range of polystyrene contents shows a family of stress-strain curves, as it can be seen in Figure 14. As the styrene content increases, the products change from soft, weak, rubbery materials to strong elastomers, then to leathery materials, and finally to hard, glassy thermoplastics. As the content of styrene is increased, the morphology of the polystyrene changes from spheres to cylinders, both dispersed in the continuous elastomer phase. When the volume fractions of both components are about equal, the two form alternating lamellae. As the styrene content increases, a continuous polystyrene phase forms in which either cylinders or spheres of the elastomer are dispersed 35 . This latter situation is illustrated in Figure 15, showing various morphologies of S-rich P(I-b-S) diblock copolymers where PS forms (in an unusual way) the continuous domain 36 . SBC mechanical behavior are mainly governed by their microphase separated structures. Accordingly, the styrene fraction must be wisely selected. As many applications benefit from low-cost styrenic block copolymers (SBC), high-temperature applications and other advanced consumptions of SBC, such as in tire rubber, are largely limited by the relatively low glass transition temperature (Tg) of poly(styrene) PS. When the service conditions approach 95°C, softening of PS domains dramatically reduces the tensile stress of these thermoplastic elastomers (TPE). One major research interest is thereby to increase their upper service temperature. These efforts mainly explored anionic polymerization of polymers with higher glass transition temperatures 37 . Table 5 below gives some hard segments allowing to increase the upper limit service temperature of SBC.

Some examples can be detailed in order to understand the effects of adding rigid segments having a Tg > 100 o C. Attention can be first paid to the substitution of PS blocks with poly(α-methyl-p-methyl styrene) P(MMS) end blocks in styrenic TPE. P(MMS)-P(My)-P(MMS) triblock (P(My) = 1,4-poly(βmyrcene) in majority) was synthesized by sequential anionic polymerization 40 . To elucidate thermal behavior along with viscoelastic properties, this triblock was subjected to dynamic mechanical analysis (DMA) and compared to a conventional PS-PI-PS triblock, as shown in Figure 16a. The PS-PI-PS sample has an elastic modulus of approximately 1 MPa between -40 and + 90 °C, at which point the modulus drops due to softening of the PS phase, as expected (black trace). On the other hand, the rubbery plateau (also around 1 MPa) for P(MMS)-P(My)-P(MMS) extends to 160 °C before a drop in the modulus is observed at the softening point of the P(MMS) blocks. This suggests that this latter TPE may well have an attractively high service temperature. A similar observation was reported by Jérôme et al. when comparing by DMA hydrogenated or not P(MMA)-PB-P(MMA) (PB = poly(butadiene)) and P(IBOMA)-PB-P(IBOMA) triblocks (Figure 16b) made by sequential anionic polymerization 42 . P(IBOMA)-PB-P(IBOMA) and hydrogenated P(MMA)-PB-P(MMA) have a comparable G′ up to about 90 °C. At higher temperatures, G′ of hydrogenated P(MMA)-PB-P(MMA) starts to decrease rapidly, whereas it persists up to 160 °C for the sample having P(IBOMA) hard outer segments. Substitution of P(MMA) outer blocks by P(IBOMA) results in a substantial increase in the upper service temperature, in agreement with a much higher Tg for P(IBOMA) (190 °C, by DSC) compared to that of P(MMA).

Furthermore, a higher tensile stress is typically observed for styrenic TPE replacing PS hard blocks by higher Tg blocks. Stress-strain data for PS-PI-PS and P(MS)-PI-P(MS) triblocks, synthesized using 1,4dilithio-1,1,4,4-tetraphenylbutane as the initiator 39 , are shown in Figure 17. It is readily apparent that the P(MS)-PI-P(MS) block polymer exhibited a markedly higher tensile strength than the PS-PI-PS sample of equivalent composition. Thus, it would seem that end blocks with glass transition temperatures higher than polystyrene enhance the tensile strength of the material. It might be added that the interaction coefficient between P(MS) and PI is virtually the same as that of PS and PI. Hence, the amount of phase blending in these two triblock polymers studied should be nearly identical. It can be assumed that the effectiveness of the glassy domains (the sites of stress concentration) is greatly dependent on the degree of phase separation. Thus, the better the phase separation, i.e., the less the glassy domains are diluted by the rubbery component, the greater is the energy absorbed prior to the disintegration of these domains.

C.2. Hydrogenation of the elastomeric block

Please note that this section is only devoted to the hydrogenation of elastomers, such cis-1,4poly(isoprene) or cis-1,4-poly(butadiene), as well as the hydrogenation of the soft segment of styrenic block copolymers. Accordingly, hydrogenated styrene/isoprene triblock means herein that only the poly(isoprene) segment was saturated for instance. Even though the hydrogenation of hard blocks such as poly(styrene) is well-known, it is not the purpose of this review.

C.2.1. Introduction

Chemical modifications of unsaturated polymers have received a great deal of attention for many years. Today, the hydrogenation of polymers is perhaps the most widely practiced post-polymerization chemical reaction. Hydrogenation is a chemical reaction between molecular hydrogen (H2) and another compound, usually in the presence of a catalyst such as nickel (Ni), palladium (Pd) or platinum (Pt). This process is generally conducted to reduce or saturate compounds. Hydrogenation of unsaturated polymers can provide an alternate route to a variety of unique and useful elastomers and thermoplastics with specific structures and properties. Hydrogenation processes are directly affected by the nature of the catalyst, purity of substrates, temperature, pressure, contaminants in the substrate, instability of intermediates and solvents.

Many polymers with new monomer sequence distributions and compositions, difficult or impossible to prepare directly through conventional polymerizations, can be easily obtained via hydrogenation.

Moreover, hydrogenated poly(diene)s have normally enhanced resistance to oxidative and thermal degradations in comparison to the parent unsaturated polymers. Hydrogenated polymers are produced and available commercially, and have also been used as well-defined substrates, providing a basis for the study of a variety of important aspects of polymer physical behavior, including spectral analysis of polymeric hydrocarbons, the effect of local chain microstructure on rheological behavior, crystallization and melting properties as well as polymer-polymer phase mixing 45 .

C.2.2. Hydrogenation processes

The predominant means of hydrogenation of poly(diene)s has been their combination with molecular hydrogen in the presence of a soluble transition metal catalysts. CaCO3 is known to have good activity for PI-and PB-based polymers without bringing about chain cleavage 47 .

Figure 18.

Selective saturation of styrene-butadiene block copolymer using a Ziegler-type hydrogenation catalyst. This latter is formed by combination of a group 8-10 transition metal pre-catalyst and a trialkylaluminium co-catalyst in a hydrocarbon solvent such as benzene.

Industrially, the hydrogenation is efficiently performed using these two catalytic techniques.

However, these reactions remain highly hazardous. The primary hazard associated with any form of hydrogen is obviously producing a flammable mixture, leading to a fire or detonation when exposed to air, oxygen or sparks. Use of hydrogen, pyrophoric catalyst, flammable solvents and generation of in situ unstable intermediates are the main safety concerns when hydrogenations are run. Substrate impurities, the wrong choice of catalyst, contaminated hydrogen (carbon dioxide, carbon monoxide and hydrocarbons) and low concentration of catalyst are major contributing factors for the runaway reactions. Proper handling of catalysts is critical due to high reactivity before and after introducing hydrogen. Existence of any impurity or colored impurity in substrate can inhibit the catalyst activity resulting in partial reduction of the substrate and occasionally in a run-away reaction. Substrates should be completely purified prior to subjecting them to hydrogenations. Hydrogenation reactions are largely exothermic in nature and can eventually increase overall pressure inside the vessel. Moreover, on a laboratory scale, the equipment required to perform hydrogenations at moderate to high hydrogen pressure is often not readily available.

Lastly, catalytic hydrogenation reactions are expensive methods due to the use of costly H2 gas and metal catalysts and the need of a high pressure reactor 48, 49 .

Ambient pressure hydrogenations are the safest among all hydrogenations. The hydrogenation of carbon-carbon double bonds (C=C) using chemical ways has also been developed. In addition, the complete elimination of chain degradation chemistry was not apparent for PI. This effective TSH / TPA couple technique has since been used to hydrogenate a large range of alkene-containing polymers, where its convenience and tolerance of functional groups is of interest and has recently been extended to hydrogenate alkynes. With no need of high pressure, H2 gas and catalyst, this method is straightforward and safe 50, 51 .

C.2.3. Properties of hydrogenated poly(diene)s

The main enhancement of hydrogenated poly(diene)s compared to their original unsaturated polymers is undoubtedly their better thermal stability. This latter can be measured by thermogravimetric analysis (TGA) by determining for instance the temperature at which weight loss is most apparent (Tdeg,max).

Tdeg,max of hydrogenated natural rubber and hydrogenated synthetic cis-1,4-poly(isoprene) samples increased with increasing the degree of hydrogenation 52 . This can be explained in that a weak π-bond within the original elastomers was converted to a stronger C-H σ-bond by the hydrogen molecule. Similar observations were reported for cis-1,4-poly(butadiene) and trans-1,4-poly(2,4-hexadiene) for instance 53 , as summarized in Table 6.

It must be added that the glass transition temperature (Tg) is different after hydrogenation as well.

Regarding PI samples, while the degree of hydrogenation of the natural rubber increased from 0 to 85.8 %, the Tg increased from -63.8 to -60.0 •C 52 . This can be explained by the replacement of amorphous segments of poly(isoprene) by crystalline ethylene-propylene units in the polymer structure. A similar shift in Tg was observed for hydrogenated synthetic cis-1,4-poly(isoprene) 54 . A similar trend was also reported for methyl methacrylate/butadiene triblock copolymer having its poly(butadiene) midsegment Even after the glass transition (increase of mobility and deformation of the polymer chains with increasing temperature), E' remains higher for hydrogenated PI. The higher E' of hydrogenated natural rubbers can be explained in that the replacement of the ethylene/propylene segments by hydrogenation of the isoprene unit in the polymer chain tends to decrease the mobility and the free volume of the polymer 52 .

Such a trend of greater E' (or G') with hydrogenation degree, confirmed for PI, cannot be generalized for other poly(diene)s. The storage modulus for methyl methacrylate/butadiene, isobornyl methacrylate/butadiene and tert-butyl methacrylate/butadiene triblocks having their poly(butadiene) midsegment hydrogenated were lower compared to the original copolymers in the glassy region, whereas the reverse situation was observed in the rubbery plateau region 55 . A similar observation was reported by Saito and coworkers regarding hydrogenated styrene-block-styrene-co-butadiene-block-styrene (S-SB-S, hydrogenation of the butadiene units) 56 , as depicted in Figure 20b. Interestingly, all these studies showed that hydrogenated elastomers or thermoplastic elastomers were stiffer in the rubbery plateau, at T > Tg. The effect of hydrogenation on the mechanical properties of elastomers or thermoplastic elastomers is of great interest as well. Stress-strain curves of a series of hydrogenated S-SB-S triblock copolymers with various degrees of hydrogenation (DH) and equivalent PS contents are given in Figures 21a and21b.

Polymer DH (%) Tdeg,max ( o C)

The modulus of the initial slope increased with increased DH (Figure 21b), whereas the tensile elongation at break decreased with increased DH (Figure 21a). In the hydrogenated S-SB-S with high DH, a sharp increase in stress was observed in high-strain regions and in low-strain regions after increasing DH. The tensile stress at break for highly hydrogenated S-SB-S (DH > 99 %) doubled the tensile stress at break for neat S-SB-S (DH = 0 %). The sharp increase in stress and the high strength at break point are behaviors similar to the behaviors of cross-linked rubbers. These results indicate that hydrogenated S-SB-S strengthens with increasing DH, although the content of the hard PS outer block segment did not change.

A study of the phase behavior was led by the research team to understand the variation in stress-strain styrene-block-styrene-co-butadiene-block-styrene S-SB-S triblock copolymer and its hydrogenated counterparts (poly(butadiene) middle block exhibiting various degree of hydrogenation (DH)) 56 .

properties with the DH. TEM and SAXS results indicated that the microphase structure expands with increasing DH owing to an increased incompatibility between the PS and hydrogenated SB block segments.

Figures 21c and21d show schematic illustrations of phase-separated structure for hydrogenated S-SB-S with different DH. Accordingly, the enhancement of the mechanical properties can be attributed to the increased segregation of the microphase structure with increasing DH 56 .

Similar results were obtained regarding tert-butyl styrene/isoprene (TBS/I) star block copolymers.

While neat samples exhibited very high elongations up to 3 000 % combined with very low tensile strength (0.5 MPa maximum), the hydrogenation of about 90 % of the isoprene units produced samples having reduced ultimate elongations (< 700 %) but significantly greater tensile stresses at break (> 25 MPa). The change in solubility parameter that occured when PI segment was hydrogenated to poly(ethylene-copropylene) presumably caused the rejection of hard P(TBS) from the rubber phase to yield discrete domains of P(TBS). Accordingly, it can be assumed that hydrogenation allowed the thermoplastic elastomer to perfect its phase separation to provide improved elastomeric properties 57 .

C.3. Commercial styrenic block copolymer TPE C.3.1. Industrial producers and applications

The first commercial styrenic thermoplastic elastomers deriving their properties from an ABA block copolymer structure were poly(styrene)-block-poly(isoprene)-block-poly(styrene) PS-PI-PS (or SIS)

and poly(styrene)-block-poly(butadiene)-block-poly(styrene) PS-PB-PS (or SBS) introduced in 1965 at an ACS Rubber Division Meeting at Shell Development Company. Initial applications were in molded shoe soles, hot melt adhesives and injection molded rubber goods. To further extend the range of applications, these polymers were made more stable by full hydrogenation in the early 1970s. Fully hydrogenated styrenic block copolymers exhibited poor resistance to oils, leading to the introduction of polymers with selectively hydrogenated rubber blocks. Through selective hydrogenation of the flexible segment, SBS became SEBS (styrene-ethylene/butylene-styrene triblock) and SIS became SEPS (styreneethylene/propylene-styrene triblock). A more recent development, although not yet commercialized, is styrenic block copolymers with an isobutylene mid-segment (S-iB-S). A wide range of triblock, diblock and star architectures were introduced in the 1970s with styrene contents ranging typically from 10 to 50 %. In 1998, worldwide production of styrenic TPE from more than 20 manufacturers was more than 10 6

tons 58, 59 . The major current manufacturers and the trade names of their products are listed in Table 7. It is of interest to note that only anionic polymerization was applied to synthesize commercial styrenic TPE.

SBC differ from other TPE in at least two major ways. First, both hard and soft domains are amorphous, and thus pure SBC are soluble in common solvents such as toluene. Secondly, in their various end uses, these polymers are always compounded with large amounts of other ingredients. These latter can be other polymers, oils, fillers and resins. In the majority of their applications, the styrenic block copolymers comprise less than 50 % of the final product.

Injection molding and extrusion are performed to shape compounded products. Products based on SBS are generally compounded with PS, hydrocarbon oils and fillers. In those based on SEBS, poly(propylene) PP often replaces PS since PP gives better solvent resistance and increases the upper service temperature. Footwear, wire and cable insulation, automotive and pharmaceutical items can be given as typical applications Solution-polymerized styrene-butadiene rubber (SSBR).

Company Trade name Type Applications

Kraton

Kraton ® D SBS, SIS, SIBS (a) Hot melt spray, diaper adhesives, elastic films.

Kraton ® FG Functionalized SEBS (b)
Impact modifiers in nylon and polyesters.

Kraton ® G SEBS, SEP(S) (c)
Soft, strong compounds for handles and grips, elastic components in diapers, oil gels.

Cariflex ® IR I (d)
Gloves, medical products, tackifiers, paints.

Dexco

Vector ® SIS, SBS Hot melt pressure sensitive adhesives, photopolymer plate applications, hygiene.

Taipol ® SBS, SEBS Compounding, plastic modification.

INEOS Styrolution

K-Resin ® SBS Packaging, toys, medical components. Styroflex ® SBS Flexible films, medical tubes, stretch hoods.

Zeon

Quintac ® SIS Grades for hot melt pressure sensitive adhesive.

Fina

Finaprene ® SB, SBS Improve the general adhesion of the compounds, reduce the blend viscosity.

S&E Specialty Polymers

Tufprene ® SBS, SEBS Molding and over-molding, footwear, gaskets, toys.

Asahi

Asaprene ® SBS Adhesives component, asphalt modifier.

Dynasol

Calprene ® SBS, SEBS Bitumen application, footwear.

Solprene ® SBS, SSBR (e) Elastomeric modifier for asphalt mixtures for paving and roofing applications.

Enichem

Europrene Sol T ® SBS

Compounding for both footwear and technical applications, toughening agent in polymer modification and plastic recycling.

Firestone

Stearon ® SB, SBS Grafting and melt blending plastic applications, non-woven hot melt adhesives.

Kuraray

Septon ® SEP, SEPS (c) , SEBS, SEEPS Tapes, cold seal food packaging, reactive adhesives, toys and sporting goods.

J-VON Hercuprene ® PP / EPDM blends

Flexible material for appliances, hose and tubing, connectors, insulation.

Saint-Gobain C-Flex ®

Not specified Pharmaceutical and biopharmaceutical applications for fluid processing.

Teknor Apex

Elexar ® Not specified Wire and cable jacketing and insulation, electrical connectors and plugs.

Chi Mei

Kibiton ® SBS Shoe soles, toys, asphalt modifications, mechanical products, adhesive agents.

78

Styrenic TPE are extensively used in adhesives, sealants and coatings. Tackifying and reinforcing resins are used to obtain a desirable balance of properties. Oils and fillers can also be added. SBC-based adhesives and sealants can be applied either from solvents or as hot melts. A major application is in pressure sensitive hot melt adhesives. SIS triblocks are softer and stickier and so they are typically employed to formulate adhesives of this type. SBC blending with thermoplastics or other polymeric materials can be also mentioned. Styrenic TPE are compatible with a wide variety of other polymers. Blends with many other thermoplastics have enhanced impact resistance amongst others 60 .

C.3.2. Styrenic TPE

Although this doctoral thesis was mostly dedicated to the synthesis of a novel styrenic triblock copolymer, some commercial styrenic block copolymers can be explored and used as benchmark materials.

In this section, only products from Kraton ® company were investigated since Kraton ® is specialized in the production of high-performance styrenic TPE and detailed technical data sheets are provided. Nine commercial Kraton ® are detailed herein and consist of linear SIS, SBS, SEBS and functionalized SEBS triblock TPE, as summarized in Table 8.

First of all, it can be noticed that all these Kraton ® triblocks satisfied ArianeGroup stress-strain properties at room temperature (tensile strength at break > 15 MPa, Tensile elongation at break > 400 %).

However, important differences must be mentioned according to the nature of the elastomeric mid-segment and the composition of the TPE. SIS Kraton ® polymers are high performance TPE with a combination of high strength, low hardness and low viscosity for easy processing as a melt or in solution. They are the lowest viscosity and hardness of all the styrenic TPE. For samples D1161 B and D1164 P presented in the SBS Kraton ® are versatile with a combination of high strength, wide range of hardness and low viscosity. Although similar to SIS samples, SBS are typically harder, which is mainly due to the more entangled poly(butadiene) matrix (lower entanglement molar mass compared to poly(isoprene). For the industrial problematic, SBS appear more interesting due to their both greater modulus at 300 % elongation and stress at break, in comparison to SIS Kraton ® . Moreover, their melt flow index at 200 o C is significantly lower than that of SIS exhibiting a similar composition, which is valuable for our technological issue.

Table 8. Some properties of linear Kraton ® triblocks. For similar S weight fractions (29-34 wt% S), SEBS exhibit a higher modulus at 300 % elongation and an improved resistance to flow (only 5 g/10min at 230 o C) in comparison with SBS and SIS samples. While keeping similar stress-strain properties, richer-S SEBS (40-57 wt% S) have a relatively high modulus at 300 % elongation (6.4-7.9 MPa) and a highly limited flowing at high temperatures (230-260 o C).

Furthermore, although not indicated in the material data sheets, it can be assumed that SEBS Kraton ® exhibit enhanced thermal stability and service temperature since it consists of fully hydrogenated SBS (see section C.2.3).

Lastly, a functionalized SEBS Kraton ® termed FG1901 G is mentioned as well. It is of interest to find commercial styrenic TPE carrying functional units since such a polar modification could encourage adhesive bonding process or welding process with composite matrices as regards the industrial project.

FG1901 G consists of a linear SEBS triblock with maleic anhydride (MA) grafted onto the rubbery midblock (1.0-2.0 wt% MA). MA-grafting improves the adhesion to nylon, poly(ester), ethylene vinyl alcohol, aluminum, steel, glass and many other substrates. It is very efficient impact modifier in nylon and poly(ester)s for making super tough engineering thermoplastic materials. FG1901 G exhibit sufficient stress-strain properties even though it appears to flow more easily at 230 o C compared to a similar SEBS (G1652 H).

In conclusion, SIS and SBS Kraton ® appear to be decent reference materials for mechanical properties at room temperature although an apparent lack of hardness. SEBS Kraton ® look more advantageous for our project, being tougher and more robust thermally (maintenance of the mechanical performance with temperature). The major drawback of these commercial SBC is presumably their poor thermal robustness especially at T > 90 o C when PS hard blocks start softening.

D. Synthesis of styrenic block copolymer TPE D.1. Living ionic polymerization

Ionic polymerization is the most studied approach for the formation of styrenic block copolymers (SBC). It consists of a controlled/living chain polymerization that have ionic species as the active polymerization centers. Because of the differences in reactivity between carbanions and carbocations, different monomer classes are suitable for anionic and cationic polymerizations, although there is some overlap 61 .

D.1.1. Anionic polymerization

Anionic polymerization proceeds by formation of carbanions at the polymer chain ends as the active center and is a well-established method for the synthesis of tailored block copolymers. In industrial practice, this method is used to prepare several important classes of block copolymers including the SBC.

The anionic polymerization of styrene and diene monomers is the typical example of a controlled block copolymer synthesis. These polymerizations are typically initiated by an organolithium species and proceed in the absence of chain termination or transfer. Under the appropriate conditions, these polymerizations are living and can be used for the formation of a variety of block copolymers with prescribed molecular weights, narrow molecular weight distributions, and well-defined block lengths. The solvents usually used are inert hydrocarbons, such as cyclohexane or toluene. Anionic polymerization methods are used for the preparation of the commercially important S-I-S and S-B-S triblock copolymers. More generally, a large majority of the styrenic block copolymer TPE commercially available are made by anionic polymerization. This synthesis can be accomplished by either sequential addition of monomers (polymerization starts at one end of the molecule and continues to the other end) or by chain coupling (polymerization starts at each end of the molecule and then the reactive chains are joined together by a coupling or linking agent). The preferred initiator for the sequential and coupling polymerization methods is sec-butyllithium because it initiates the polymerization very readily (initiation rate high in comparison to the rate of the subsequent polymerization).

Styrenes and dienes make up the group of monomers most used for the preparation of block co polymers by anionic polymerization. Although there are only a handful of commercially available styrenes (styrene, alpha-methyl styrene…) and dienes (isoprene, butadiene…), by changing the synthetic conditions, changes in the diene polymerization regio-and stereo-selectivity can lead to changes in the polymer properties 61, 62 .

For example, the glass transition temperature of poly(butadiene) can be tuned by altering the fraction of the 1,2-content along the polymer backbone 63 . In this way, the physical properties of the corresponding block copolymers can be systematically modified.

D.1.2. Cationic polymerization

Cationic polymerization (sometimes referred to as carbocationic polymerization) is used to polymerize monomers that cannot be polymerized anionically, although it can only be employed for a limited range of monomers. The most used monomers are styrenes, vinyl ethers, and isobutylene. In comparison to the anionic approach, the initiating systems in controlled cationic polymerization are often more complex. For preparing triblock copolymer, F-R-F difunctional initiators are used, where F represents a functional group such as chlorine, hydroxyl or methoxy group and R is a hydrocarbon moiety. In addition to the initiating species, a ''co-initiator'' is often used to modify the naturally high reactivity of the propagating cationic species. By decreasing the reactivity of the propagating center in cationic polymerization, side reactions such as chain transfer to monomer and chain termination can be significantly reduced 64 .

A fairly recent development in the synthesis of SBC is the living cationic polymerization producing well-defined poly(styrene-block-isobutylene-block-styrene) (S-iB-S). The process is typically carried out at T = -80 °C in a moderately polar solvent. TiCl4 or BCl3 are used as co-initiators 65 . Even though only P(iB) can comprise the center elastomeric block via this approach, various outer hard blocks can be inserted.

In contrast to anionically prepared SBC, the styrenic block can be made up from α-methyl styrene, paramethyl styrene, para-tertiary-butyl styrene, para-chloro styrene, and other monomers 66 . The high Tg values of these substituted styrenes and others mean a higher upper service temperature than the analog S-iB-S and S-EB-S block copolymers. Cationic polymerization using multifunctional initiators opens the possibilities to synthesize a variety of SBC as well, such as multi-arm star block copolymers and arborescent block copolymers comprising a dendritic (tree-like) P(iB) core carrying multiple PS outer blocks. The multiarm starblock copolymers are alternatives to linear block copolymers because their superior combination of and processing properties, including high combined with lower viscosities at similar molecular weights 67 .

D.1.3. Drawbacks of living ionic polymerization

Without any doubt, living ionic polymerization, and more importantly anionic polymerization, is the most robust and effective technique to manufacture well-defined styrenic block copolymers. However, considering our industrial issue, two limiting factors can be mentioned.

First, the experimental conditions are stringent and demanding. Oxygen, water, or any other impurity must be completely eliminated to prevent undesirable reaction of the highly reactive propagating species. Even a concentration at the part per million level of these impurities can markedly affect the polymerization. The anionic polymerization high vacuum technique is the best method for the synthesis of model macromolecules even though it is a demanding, time-consuming method leading to a limited quantity of products. A high vacuum line (10 -5 to 10 -9 mbar) is typically composed of a high-quality mechanical vacuum pump (pressure reduced to 10 -2 -10 -3 mbar), a mercury or oil diffusion pump (pressure down to about 10 -5 mbar), a liquid nitrogen trap, an upper glass tube rig and a lower glass tube rig and stopcocks, as shown in Figure 22a. A glass cutter and a hand torch are also necessary tools for making reactors. Moreover, prior to the polymerization, numerous steps are paramount.

84

Degassing (mainly oxygen via freeze-thaw cycles), distillation of liquids (solvents, monomers…), dilution (concentrated reagent diluted with purified solvent) and ampoulization (collection of small quantities of solvent additives or alcohols) are common procedures. The initiator, such as secbutyllithium (sec-BuLi), has to be dried over CaH2 on the vacuum line, degassed thoroughly and distilled into a calibrated ampoule. Solvents and monomers have to be carefully purified as well. Benzene can be taken as an example. It is purified by stirring over concentrated sulfuric acid for a week in a conical flask inside a fume hood. Then, benzene is transferred carefully into a round-bottom flask containing finely grounded CaH2 and a magnetic stir bar, attached to the vacuum line, and degassed. It is left for reaction of CaH2 with and water present overnight. Then the benzene is again degassed and distilled into a calibrated cylinder containing polystyryllithium (PSLi) made from n-BuLi and styrene 68 . In comparison with the high vacuum techniques, the Schlenk technique is a bit simpler to perform allowing anionic polymerization reactions to be carried out under inert gas condition in a commercially available Schlenk flask. The Schlenk method is also more amenable to carrying out reactions on a larger scale, whereas polymerizations carried out under high vacuum conditions with break seal techniques are usually conducted on the scale of about 10 g or less 69 . A schematic of a Schlenk line is given in Figure 22b.

A second disadvantage of the living ionic polymerization is that functionalized monomers are difficult to polymerize because the functional moieties are rarely compatible with the ions and counterions formed during the polymerization. Living anionic polymerization of acrylates and methacrylates for example is not so straightforward because of the serious side reactions derived from the functional groups in the polar monomers. In addition, exothermic rapid propagation and partially aggregated propagating

D.2. Reversible-deactivation radical polymerizations (RDRP) D.2.1. Control of the macromolecular architecture by RDRP

Ionic polymerization is the most efficient method to produce well-defined block copolymers with generally an excellent control of the dispersity of the polymer chains lower than 1.1. Nevertheless, as discussed above, this kind of polymers is difficult to prepare (no oxygen, no moisture, high purity of monomers…) and quite a few monomers (styrene, butadiene, methyl methacrylate…) are consequently involved in industrial processes.

Such precise macromolecular synthesis employs concepts of living polymerization, in which the contribution of chain breaking reactions is minimized and the apparent simultaneous growth of all chains can be achieved via nearly instantaneous initiation. A combination of fast initiation and an absence of termination seemingly conflicts with the fundamental principles of radical polymerization, which proceeds via slow initiation and in which all chains are essentially dead at any given instant. However, the This approach relies on the persistent radical effect (PRE), which is a peculiar kinetic feature which provides a self-regulating effect in RDRP systems. Propagating radicals P• are rapidly trapped in the deactivation process (with a rate constant of deactivation, kdeact) by species X, which is typically a stable radical such as a nitroxide (by NMP). The dormant species are activated (with a rate constant, kact) either spontaneously/thermally, in the presence of light, or with an appropriate catalyst (as in ATRP) to reform the growing centers. Consequently, while all chains are terminated in the conventional radical process, the terminated chains constitute a small fraction of all the chains (~ 1-10 %) as a result of the greater number of growing chains in a RDRP process. The remaining chains are dormant species, capable of reactivation, functionalization, and chain extension to form block copolymers for instance 72 .

D.2.2. Block copolymer TPE by RAFT and ATRP

Various block copolymer TPE were achieved via RAFT. For instance, a series of well-controlled S-nBuA-S triblock (nBuA = n-butyl acrylate) copolymers (Mn = 77-338 kg.mol -1 , Đ = 1.41-3. The elasticity of these triblock copolymers was also demonstrated by the results of hysteresis experiments at around 250% deformation (residual strain of about 65% was observed after the deformation).

RAFT limitations include the lack of commercial availability and stability of many transfer agents, the removal of dithioester and some other end groups required due to their color, toxicity, and potential odor, and the difficultly associated with end-functionalization. Another shortcoming, directly linked to our industrial problematic, is that little has been reported on RAFT polymerizations of diene monomers.

Lebreton et al. reported on RAFT polymerization of butadiene using various fluorinated dithiobenzoates as RAFT agents 76 . Jitchum and Perrier reported on RAFT polymerization of isoprene with a decent control 77 .

The main difficulty lies with the high polymerization temperatures that are required to obtain acceptable rates of fragmentation and perhaps to the stability of the RAFT agents.

The synthesis of well-defined B-A-B type triblock copolymer TPE via ATRP was attempted as The main disadvantage of performing ATRP can be the use of a transition metal complex, which must often be removed from the final product, and a suitable ligand. More specifically for our project, ATRP of dienes remains a challenge since diene monomers can competitively displace ligands and generate less redox-active complexes. These competitive complex formation processes are especially important for low ppm copper processes. To overcome these issues, it is necessary to prepare more stable complexes (speciation) but with preserved activity and dynamics. Furthermore, high-molecular weight polymers are usually difficult to access by ATRP due to outer sphere electron transfer processes, involving oxidation or reduction radicals, as well as β-H elimination reactions 81 .

D.2.3. Focus on nitroxide-mediated polymerization (NMP)

In comparison to other RDRP processes, the NMP of diene monomers was well-reported. The NMP of conjugated dienes such as butadiene (B) 82, 83 and more importantly isoprene (I) 84 was studied from the early times of the technique by both the groups of Georges and Boutevin in both cases with 2,2,6,6tetramethylpiperidinyl-1-oxy (TEMPO), a stable free nitroxide. Well-defined homopolymers and copolymer architectures based on either B or I could be achieved using the 2,2,5-trimethyl-4-phenyl-3azahexane-3-oxyl nitroxide (TIPNO), demonstrating the ability of the technique to provide high molar mass PB with low polydispersity indexes. Worth mentioning are block and random copolymers based on associations of 1,3-dienes with functionalized vinyl monomers, including styrene, acrylate and methacrylate derivatives using TIPNO 85-89 . The controlled radical copolymerization of chloroprene or I with dimethyl 1,3-butadiene-1-phosphonate in the presence of TEMPO was later investigated, in a bimolecular initiating system 90 . The living polymerization of bicyclic dienes was also studied with the aim of preparing thermally cross-linkable block copolymers 91 . N-tert-butyl-N-[1-diethylphosphono-(2,2dimethylpropyl)] nitroxide (SG1) also demonstrated some great ability to control the polymerization of I 92- 94 . In particular, Nicolas and co-workers reported on the effects of temperature, initiator concentration, initiator structure and solvent on the preparation of PI from a range of SG1-based initiators containing primary, secondary or tertiary homolytic initiating groups with various functionalities 93 . Control over I polymerizations using acid-functionalized alkoxyamines was poor compared to polymerizations employing the corresponding N-hydroxysuccinimide (NHS)-ester alkoxyamines, which also opened the door to efficient functionalization. Accordingly, an important advantage that NMP offers over ATRP and RAFT polymerizations is in the ability to control effectively the polymerization of 1,3-dienes.

Moreover, it is today acknowledged that NMP is the easiest RDRP process, employing purely organic systems 71 . The latter, comprising radical initiator/nitroxide or alkoxyamine with or without additional nitroxide, are also typically simpler initiating systems compared to ATRP initiating systems, which require activated halide, metal halide, and ligand, but include traditionally additional additives to facilitate the polymerization. In comparison to RAFT polymerization, NMP does not require the use of sometimes unpleasant sulfur chemistry

Introduction

Terpenes are hydrocarbons that contain one or more carbon-carbon double bonds and share the same elementary unit as isoprene. They are a class of important bio-derived compounds 1,2 and have been known for hundreds of years as components of essential oils obtained from leaves, flowers, and fruits of many plants 3 . Among the vast members of the terpene family, β-myrcene (7-methyl-3-methylene-octa-1,6diene, abbreviated My) particularly is an interesting "renewable" monomer, as it is a component of the bay, ylang-ylang, wild thyme, lemon grass, and juniper berry essential oils 3 . It is an acyclic monoterpene with a highly active conjugated diene structure that can be polymerized by classical polymerization methods.

Based upon its relatively low cost 3 and ease of isolation 4-8 , My has been studied quite extensively as a starting material for polymers. Similar to other 1,3-dienes, such as butadiene and isoprene, My also forms rubbery polymers when polymerized. It is either used in small amounts as an additive to adjust elastomeric properties, or as a main component in poly(myrcene)s P(My)s 3 . For instance, the syntheses of styrene/βmyrcene (S/My) copolymers and My-based composites showed that the use of My improved the mechanical properties of poly(styrene)-based plastics 9 . Likewise, Cawse and coworkers incorporated hydroxylfunctionalized P(My) into a poly(urethane) leading to improved stress-strain behavior and impact properties 10 . Hence, My, either as a homopolymer or when copolymerized with another monomer, can provide a wide latitude of polymer properties. polymerization initiated by n-butyllithium (n-BuLi) 18 . This approach allows manufacturing tailor-made macromolecules with precise and pre-determined molar masses, compositions, topologies and functionalities. P(My-b-S) triblock copolymer, which could be used as a thermoplastic elastomer, was made in this way by Quirk et al 19,20 . Hoye and Hillmyer reported a similar polymer combined with bio-based αmethyl-p-methylstyrene prepared by the dehydrogenation of limonene 21 . In spite of these successful outcomes by living anionic polymerization, this technique exhibits one main disadvantage which is the requirement of stringent reaction conditions, especially the use of chemically ultrapure reagents as well as the absolute removal of air and of traces of water 22 . Polymer chemists have long desired to develop polymerization processes that combine the control of microstructure typified by ionic polymerization but with the ease of industrial implementation typified by other types of polymerization such as free radical or addition polymerization.

The advances in macromolecular synthesis in the last two decades provided a game-changing synthetic tool to easily achieve complex macromolecular architectures: reversible-deactivation radical polymerization (RDRP) [23][24][25][26][27] . In comparison to living anionic and cationic polymerizations, RDRP differs by its relative ease-of-use since only dissolved oxygen has to be eliminated, it can polymerize a wide variety of vinylic monomers by a radical mechanism that are not possible ionically, and the latitude in process conditions possible (bulk, solution, emulsion, dispersion, etc.) 28 . The three major families of RDRP are nitroxide-mediated polymerization (NMP) However, NMP has so far never been applied for My polymerization, whereas this process exhibits some advantages: NMP does not rely on sulfur-based chain transfer agents like RAFT does or metallic ligands that ATRP required, which may contaminate polymers for some applications 37, 38 . NMP is historically the first and represents perhaps the easiest RDRP technology to apply as the initiators and mediators are directly available from commercial sources. Although the synthesis of this poly(terpene) by NMP has not been studied yet, other 1,3-diene monomers were polymerized in a controlled manner. Indeed, isoprene and Given the positive results achieved previously with isoprene and butadiene using BlocBuilder TM type initiators, the NMP-based My chemistry appears very promising. Furthermore, My exhibits two valuable features in comparison to isoprene and butadiene: It is a natural renewable compound with low volatility (boiling point Tb = 167 °C) 3 , meaning that pressurized equipment is not needed to handle it.

In the present study, we report first the effects of target number-average molecular weight (Mn,theo), reaction medium, temperature and additional SG1 free nitroxide on the control of NMP of My. A BlocBuilder TM -modified initiator, the succinimidyl ester terminated NHS-BlocBuilder (NHS-BB, Figure 1,3) 54-57 was used during this study and the relevance of this choice is explained below in the discussion.

The optimized reaction conditions were used to chain-extend cleanly P(My)-SG1 macroinitiator with a fresh batch of S. Thermal and stress-strain properties of the resulting P(My-b-S) diblock copolymer were also investigated. Further, NHS-BB initiated the copolymerization of My and S via NMP was studied in terms of reactivity ratios and the resulting glass transition temperatures of the copolymers. To the best of our knowledge, this paper is the first comprehensive kinetic investigation of the nitroxide-mediated (co- The experimental conditions of the various My polymerizations performed and the characterization of the resulting poly(myrcene)s P(My)s at the end of the experiments are summarized Table 1 and Table 2 respectively.

• Choice of NHS-BlocBuilder (NHS-BB) as controlling alkoxyamine.

SG1-based initiators such as BlocBuilder TM (BB) or NHS-BB 54 have lower activation energy and much higher dissociation rate constants (kd) compared to TEMPO-based initiators 62 . As a result, they can be activated at lower temperatures and allow controlled polymerizations of a wider range of monomers, such as methacrylates, 1,3-dienes and styrenic monomers. However, the use of additional SG1 nitroxide is needed to further reduce the reaction rate and produce well-defined polymers when polymerizing nonstyrenic monomers initiated by BlocBuilder TM 50,52,63, 64 . In comparison, the higher dissociation rate (~ 15 times) and slightly lower activation energy (~ 0.9 times) of NHS-BB, both measured at 120 °C in tert-butyl benzene, lead to rapid initiation and quick release of the SG1 rate moderator 54 . Such intrinsic features permit

102

polymerizations to be done without any added free nitroxide at the onset of the polymerization. The NHS-BB was thus used to effectively control the homopolymerization of isoprene in bulk at 115 °C (Đ ≤ 1.15 and Mn = 2.4 kg.mol -1 after 16 h) 53 and styrene (S) in bulk at 120 °C (Đ ≤ 1.35 and Mn = 21.2 kg.mol -1 after 150 min) 54 without aid of additional SG1 nitroxide. Accordingly, the NHS-functionalized alkoxyamine is predicted to be effective in obtaining well-defined P(My) without additional SG1 free radical.

NHS-BB-initiated P(My) (experiment My-6) and BB-initiated P(My) (experiment My-6') were synthesized under the same experimental conditions (Table 1) in order to deem the effectiveness of BlocBuilder TM and NHS-BB to control the homopolymerization of My. As Figure 2.a indicates, the kinetic plots were linear with respect to reaction time. Interestingly, kpK increased when using BB (7.2 ± 0.4 10 -5

s -1 , Table 2) instead of NHS-BB (4.3 ± 0.7 10 -5 s -1 , Table 2). A similar evolution of Mn with conversion can be noticed for both experiments with Mn increasing linearly with conversion up to about 40% and slightly plateauing afterwards as shown in • Kinetic order and apparent rate constant of My polymerization by NMP.

The semi-logarithmic kinetic plots of ln((1-XMy) -1 ) (XMy = My conversion) versus time are illustrated in Figure 3. All My polymerizations obeyed first order kinetics as described by Equation 4and generated good fits to linear kinetic plots until 300 min polymerization. A loss of linearity was apparent at XMy > 0.3.

ln([M]0 / [M]t) = ln([M]0 / ([M]0(1-XMy))) = kp × [P • ] × time (4) ID [BB]0 (M) [NHS- BB]0 (M) [SG1]0 (M) r (a) [My]0 (M) Solvent (b) [Solvent]0 (M)
K = ([P • ] × [SG1 • ]) / [P-SG1] (5) 106 kpK ≅ (kp × [P • ] × [SG1 • ]0) / [NHS-BB]0 = kp × [P • ] × r (6)
Table 2 summarizes the experimental kpK values for the various polymerizations led. Without additional SG1, all kpK values are in the same order of magnitude, namely 10 -5 s -1 , regardless Mn,theo from 20 to 50 kg.mol -1 (DP = [My]0 / [NHS-BB]0 from 147 to 367), the reaction medium (in bulk or in solution) and the temperature (110 to 130 °C). When adding additional SG1 free nitroxide in the feed (My-8 and My-9), the polymerization has, by contrast, a kpK about an order of magnitude lower (10 -6 s -1 ). It was expected since SG1 free radical acts as a rate moderator, slowing the reaction down 65 . However, a clear increase of the polymerization rate at the final stages of the polymerization was observed for the experiment My-9 with 11.7 mol% of additional SG1 relative to NHS-BB. This may be explained by a slower initiation process in the presence of a significant amount of SG1 free nitroxide, giving rise then to a higher [P • ]. Interestingly, induction periods may be apparent for some polymerizations such as experiment My-3. Indeed, at the beginning of this reaction (t < 100 min), very low XMys, almost plateauing, were observed. This can presumably be explained by the role of the SG1 nitroxide trapping the radicals generated at the early stages of the reaction until SG1 reached its equilibrium concentration with the dormant species. At this point, the induction time is terminated and followed by the controlled NMP (at t > 100 min for experiment My-3, XMys increased linearly with time as initially expected). kpK values corresponding to NMP-based polymerization of isoprene in bulk at 115°C without additional SG1 are in the same range, from 0.4 to 1.4

x 10 -5 s -1 53 . These relatively slow NMP-based polymerizations of 1,3-diene monomers compared to styrenic and methacrylate monomers which exhibit higher kp 44,66-68 , was previously noted by Hawker et al regarding the NMP of isoprene and butadiene initiated by a TIPNO-based initiator 41 .

• Influence of alkoxyamine concentration.

Attention was first paid to the theoretical number-average molecular weight (Mn,theo = MMy × DP) in order to know if high molecular weight P(My) samples could be achieved by simply increasing the monomer/initiator initial molar ratio. Therefore, various amounts of NHS-BB alkoxyamine were used to prepare P(My)s by heating My at 110 °C for 7 h in 50 wt% toluene, allowing to target three different Mn,theo at quantitative conversion (My-1, My-2 and My-3 in Table 1). P(My)s with Mn ranging from 7.5 to 12.8 kg.mol -1 and dispersities of 1.35-1.68 (Table 2) were synthesized. Mn were on average 87%, 83% and 51%

of Mn,theo during the polymerizations (Mn/Mn,theo average values from 0 to 7 h) as regards the experiments My-1, My-2 and My-3 respectively (Figure 4.a). Accordingly, deviations from the theoretical molar masses are higher with increasing target Mn,theo. Besides, the lowest Đ were obtained when targeting the lowest Mn,theo (Đ of 1.14-1.35, Figure 4.a) confirming that irreversible termination reactions and/or some chaintransfer reactions 69 to My or P(My) or toluene solvent are more prevalent as Mn,theo increasing, generating the likely loss of control. In the ideal case of a controlled/living radical polymerization, fast initiation as compared to propagation and fast exchange between active (P • in low concentration) and dormant species (P-SG1 in our situation) allow all chains to begin to grow at the same time and the reaction to be free of irreversible termination and transfer reactions. This results in a linear increase of Mn with XMy while Đ stays narrow (< 1.3). 2.

The rate of polymerization was almost unaffected with the increase of the NHS-BB concentration, kpK staying around 2.0 10 -5 s -1 . An increase of the kpK with lower Mn,theo could have been expected as those polymerizations may have a reduced probability of irreversible termination reactions compared to polymerizations with higher target Mn 63 . Moreover, in such a situation with a lower Mn,theo, the equilibrium may not strongly favour the alkoxyamine such that the instantaneous chain radical concentration is not low enough to avoid bimolecular termination reactions. Accordingly, bimolecular termination events may be more apparent, resulting in an increase of the overall polymerization rate until the steady state predicted by the PRE is reached. Harrisson and al reported 53 an increase of the rate of isoprene polymerization in a nonlinear manner with the concentration of BlocBuilder TM , the rate of reaction being proportional to

[BlocBuilder TM ] 0.42 .
A target number-average molecular weight of 30 kg.mol -1 (DP = [My]0 / [NHS-BB]0 = 220) is thereafter chosen since it allowed achieving relatively high average chain length P(My) with a decent degree of control.

• Influence of the reaction medium.

The presence of a solvent, due to its polarity and its viscosity for instance, can induce drastic changes for both kd 70 and kc (rate of recombination) 71 of the initiator. A loss of control can be generated by chain-transfer reactions to the solvent as highlighted by Charleux and coworkers for the NMP of acrylic acid in 1,4-dioxane 64, 72 . On the other hand, the solvent can also improve the control of the NMP. Harrisson and al reported better control of isoprene polymerization by SG1-based NMP when using 1,4-dioxane and pyridine instead of bulk, due to a disruption of intramolecular hydrogen bonding within the alkoxyamine initiator and stabilization of the SG1 free nitroxide 73 . Accordingly, it is of interest to determine the influence of the reaction medium, which may play a significant role on the NMP of My.

The polymerization of My, mediated by NHS-BB, was carried out at 110 °C in toluene (50 wt%) and in bulk, and at 100 °C in 1,4-dioxane (50 wt%) due to its lower boiling point (Tb,1,4-dioxane = 101 °C), while keeping Mn,theo constant around 30 kg.mol -1 (My-2, My-5 and My-4 respectively in Table 1). No change in the overall rate of polymerization was observed as kpK values were not significantly different and the final conversion was around 0.4, regardless of the reaction medium (Table 2). The molecular weights for each experiment followed the same trend with a linear increase of up to 20% conversion and a slight plateau at higher conversions ( Accordingly, My polymerization in bulk gave satisfactory results in terms of control compared to that in solution using 1,4-dioxane and toluene. Bulk polymerization of My was thus applied for further studies.

• Determination of the appropriate polymerization temperature.

A third series of My NMPs was undertaken, varying only one experimental parameter, namely the reaction temperature, ranging from 110 °C to 130 °C (My-5, My-6 and My-7 in Table 1). This was made possible through the low volatility of this monoterpene (Tb,My = 167 °C). Increasing the temperature from 110 °C to 120 °C more than doubled the kpK value with significantly greater conversions (~ 1.7 times on average) during the experiment (Table 2). Therefore, an increase in the overall rate of polymerization, as expected by the Arrhenius equation 74 , was observed. Furthermore, contrary to what one might think, this was not accompanied by a loss of control over the polymerization. Indeed, Mn increased linearly with XMy with slight discrepancies at the final stages. Mn values were on average 90% of the theoretical ones while Đ remained below 1.3 (Figure 4.c). Unexpectedly, polymerization at 130 °C was not faster than that at 120 °C. Its overall rate of polymerization was very close to that of the polymerization performed at 110 °C (Figure 3.c). Besides, despite low Đ ≤ 1.3, significant deviations from the predicted line were observed regarding the average molecular weights (Figure 4.c). This marked loss of control occurring from the early stages of the experiment may be due to irreversible termination reactions, which were becoming more frequent at higher temperature.

Consequently, a reaction temperature of 120 °C (experiment My-6) was deemed satisfactory since it provides a polymerization at least as controlled as at 110 °C combined with a higher overall rate of reaction. Similarly, 115 °C was the optimized reaction temperature of the SG1-based NMP of isoprene for a control polymerization at an acceptable rate of reaction 53 . The GPC traces for My-6 are given in Figure 5 and exhibit a monomodal shift of relatively narrow peaks taken at different points in the polymerization. A second and minor population of longer P(My) chains can be generally detected at the early stages of the polymerization as observed at t = 1 h. This higher Mn peak, generally observed for almost every My polymerization led, appeared at the commencement of the reaction (t < 60 min), remained constant (no clear shift to lower elution volumes) to become not detectable by GPC at the last stages of the experiment.

A negligible fraction of high Mn P(My) was therefore obtained at the end of the reactions compared to the amount of well-controlled polydiene chains. This may be due to My auto-initiation increasing the amount of propagating radicals that would not necessarily be reversibly terminated by the SG1 and therefore would not be controlled (deviation from the initial stoichiometry with [P • ] > [SG1 • ]). Although autopolymerization should be moderated by the equilibrium established by the PRE, My auto-polymerization is likely significantly faster than My NMP making the mediation by SG1 more complicated. It should be noted that the presence of a high Mn peak was likewise observed at the early stages of most of the My polymerizations performed such as experiment My-10-S (Figure 6.a) and experiment My/S-50 at 1 h (Figure 9). Lastly, a GPC measurement of commercial My used (grade ≥ 90%, stored at 4-5 0 C to prevent early polymerization) indicated the absence of any polymers. The assumption that high Mn P(My) could be already present before polymerization does not appear relevant, auto-polymerization being inhibited at low T and impurities consisting mainly in monomers and dimers. Since the 1950s, My is known to polymerize spontaneously at elevated temperature as other conjugated dienes whereas its loss by auto-polymerization at Troom was estimated to one third after three months without inhibitor 75 .

• Effects of additional SG1 free nitroxide.

The effectiveness of additional SG1 over the control of non-styrenic monomer polymerizations initiated by the BlocBuilder TM initiator was demonstrated. For instance, the control of tert-butyl acrylate

Elution Time (min) 1 h 2 h 3 h 5 h 7 h

Figure 5. Normalized GPC traces of P(My) initiated by NHS-BB at 120 °C in bulk targeting Mn,theo = 30 kg.mol -1 at quantitative conversion (experiment My-6). The arrow indicates the peak/shoulder observed at t = 1 h corresponding to a minor fraction of high molecular weight P(My) (Mn = 45.3 kg.mol - 112 polymerization was improved by introducing 5-6% of free nitroxide relative to the BlocBuilder TM initiator, reducing both the overall rate of reaction and the Đ values 76 . In order to know if such a strategy could be efficient for the system My/NHS-BB, two extra experiments were undertaken: My-8 and My-9 with respectively 4.9% and 11.7% of additional SG1 mediator relative to NHS-BB and implemented under the optimized reaction conditions, namely in bulk at 120 °C targeting 30 kg.mol -1 at XMy = 100% (Table 1).

The results were compared with experiment My-6 where no additional SG1 was used. Adding about 5% of additional SG1 relative to NHS-BB in the feed did not affect the polymerization of My, exhibiting nearly identical overall rate of reaction (Figure 3 and coworkers for the system SG1/MMA (methyl methacrylate) when using large excesses of nitroxide 77 .

In our case, the percentage of free nitroxide was relatively low and a further analysis would be necessary to clearly identify the possible main chain-end forming event occurring. Another assumption explaining these results for experiment My-9 with the highest r = 0.117 may be the contribution to extra SG1 nitroxide to high Mn propagating macroradicals resulted from self-initiation of My. In that case, higher deviations of the measured Mn from Mn,theo can be expected as well without however bringing about a loss of control (linear trend of Mn versus XMy and Ɖ < 1.3 except at the end of the reaction, Figure 4.d).

• Microstructural features of NMP-based P(My).

My is expected to yield P(My) polymer with a variety of structural units. Such a substituted diene system can produce units of 1,2-, 3,4-, 1,4-cis and 1,4-trans. 1,2-, 3,4-and 1,4-isomers were detected by 1 H NMR whereas the signature peaks for 1,4-cis and 1,4-trans motifs were obtained from 13 C NMR spectra (see Supporting Information for the spectral assignments of 1 H NMR and 13 C NMR from typical experiments, Figures S2a andS3a). The analyses were in good accordance with the literature 21, 58 . The proportions of these four different types of microstructure at the end of each polymerization are given in Table 2. It can be concluded that the NMP of My initiated by NHS-BB affords a P(My) that is 1,4-regular (80 to 90%) with 1,2-(5 to 10%) and 3,4-defects (5 to 10%). This is in agreement with the predominant 1,4-microstructure (96%) of P(My) synthesized by RAFT with 2-ethyl-sulfanylthiocarbonylsulfanylpropionic acid ethyl ester chain-transfer agent 36 . More generally, the 1,4-motif is typically dominant (> 75%) by free radical polymerization 12,13,15,16 of My. Similar proportions of 1,4-addition (81%), 1,3-addition (13%) and 1,2-addition (6%) were determined for the NMP in bulk of isoprene 53 . Moreover, the NMPbased poly(terpene) exhibits a high cis-selectivity, higher than 65%. The coexistence of cis and trans isomers was qualitatively reported regarding RAFT-based 36 and emulsion-based 16 P(My)s.

Synthesis of poly(β-myrcene-block-styrene) diblock . copolymers

• Chain-extension with styrene (S) from poly(β-myrcene) P(My) macroinitiator.

The active feature of NMP-based polymers is enabled due to the nitroxide moieties capped at the chain ends. To investigate the ability of the expected SG1-terminated P(My) homopolymer to re-initiate a second batch of monomer, chain-extension experiments were performed using S monomer and three P(My) macroinitiators (My-10, My-11 and My-12) exhibiting different degrees of polymerization (13.2 kg.mol -1 < Mn < 30.4 kg.mol -1 ). The experimental conditions and results of these experiments are summarized in Table 3 and Table 4, respectively. With a fresh batch of purified styrene, P(My) macroinitiator was extended with a poly(styrene) PS block with Mn ranging approximately from 12 to 26 kg.mol -1 . Đ of the diblock copolymers (1.33, 1.83, 1.50 and 1.88) were higher than that of their macroinitiator (1.18, 1.32, 1.32 and 1.38). This broadening of Ɖ after the chain-extensions can be due to the diffusion of the mediating SG1 radical from the propagating chain end to styrene monomer. Side reactions involving radical crossovers for NMP were reported by Hawker et al 78 . This assumption is consistent with the reactivity ratios of My and S determined in part C highlighting the shiftless and unprivileged addition of S towards the propagating radical --My • . Besides, irreversible termination and some macroinitiator chains which were not initiated could also have contributed to the loss of control observed during these chain-extensions. The S molar composition of the synthesized P(My-b-S) copolymers (FS), measured via 1 H NMR spectroscopy ( 1 H NMR spectrum of recovered P(My-b-S) copolymer My-11-Sa provided in the Supporting Information, Figure S4a) varied from 27 to 62 mol%.

The GPC chromatograms of the chain-extension experiments are shown in Figure 6. For every experiment, the P(My-b-S) chain-extended block copolymer (dotted line) shifted to lower elution time compared to its P(My) macroinitiator (solid line). These marked shifts demonstrated the increase of Mn with polymerization time. Moreover, the GPC traces of the chain-extended polymers retained generally their monomodal nature, indicating a high level of chain end fidelity. For each experiment, the tails of chainextended copolymers overlap to some degree with the associated macroinitiator traces. This tailing, partly due to the presence of "dead" chains from the macroinitiator, was slightly more apparent for the chainextensions My-11-Sa and My-12-S where high Mn,theo was targeted (≥ 90 kg.mol -1 ) combined, in the case of My-11-Sa, with a long reaction time. Under such conditions, irreversible terminations are more prevalent, generating a loss of control which can be also highlighted by the increase of Đ > 1.8 at the end of these chain-extensions.

In summary, these results confirmed the ability of NMP-based P(My) macroinitiator to reinitiate a new batch of styrene and polymerize in a controlled way. Slight tails were observed, likely due to a small portion of inactive macroinitiator as well as the occurrence of irreversible termination events.

• Tensile properties of P(My-b-S) diblock copolymer.

Uniaxial tensile testing was performed on the resulting P(My-b-S) diblock copolymers, with FS = 0.27-0.62, in order to determine their mechanical behavior. Table 4 gives the tensile properties of the various copolymers and Figure 7 Regardless of the molar composition of the samples, all stress-strain curves were marked by a linear elastic region, followed by a short failure region. No plastic deformation was observed, resulting in a failure point very close to the yielding point. Every specimen, breaking after reaching their elastic limit, exhibited both brittle behavior with low ultimate tensile strength (< 1.1 MPa, Table S1a in Supporting Information) and low extensibility with poor elongation at break (< 16%, Table S1a in Supporting Information).

Materials from experiment My-10-S, even more fragile, fractured before performing the mechanical test.

As a comparison, isoprene/styrene P(I-b-S) triblock copolymer having slightly higher overall Mn (68.5 kg.mol -1 ) and FS (0.40) exhibited thermoplastic elastomeric mechanical behavior with significant strength (tensile stress at break > 20 MPa) combined with high elongation at break (> 1000%) 79 . Likewise, butadiene/styrene P(B-b-S) triblock sample exhibing Mn = 73 kg.mol -1 and FS = 0.40 had largely better stress-strain properties with tensile stress at break higher than 20 MPa and elongation at break higher than entanglement molecular weight (Me,PI = 4-6 kg.mol -1 and Me,PB = 2-4 kg.mol -1 ) 81 , allowing the matrix phase to have a high entanglement density. By contrast, the maximum Mn of the poly(terpene) segment in the P(My-b-S) copolymers was ~ 30.4 kg.mol -1 (experiment My-12-S), only slightly higher than P(My) entanglement molecular weight (Me,P(My) = 22-31 kg.mol -1 ) 81 . Although My-12-S sample may have been entangled, it is very likely that My-11-Sa and My-11-Sb samples exhibiting shorter P(My) chains (Mn,P(My) = 22.8 kg.mol -1 ) did not form an entangled elastomeric phase. This difference in mechanical behavior between P(My-b-S) diblock samples and P(I-b-S) / P(B-b-S) triblock samples was presumably due to different molecular architectures as well. When the rubbery segment (P(My), PI or PB) is in excess of the rigid segment (PS) so that the rubber forms the continuous phase, the triblock copolymers have generally higher tensile strengths compared to similar diblock copolymers 82 . The rigid segment is aggregated in both cases and both types of block copolymers have the characteristics of a rubber containing rigid fillers.

However, the rubber phase acts like it is cross-linked in triblock polymers because both ends of each elastomeric chain are attached to a rigid dispersed phase (entanglements are physically trapped). In the diblock polymer, the rubber chains are attached to a rigid chain at only one end and therefore are not crosslinked. Adhikari studied the influence of the presence of 60 wt% of P(B-b-S) diblock chains (FS = 0.5) to linear P(B-b-S) triblock copolymers (FS = 0.5) on the mechanical properties of this latter. Although the strain at break values were similar for both types of structures (~ 900%), P(B-b-S) triblock copolymer containing an important amount of P(B-b-S) diblock chains exhibited tensile stress at break (< 15 MPa) significantly lower in comparison to the pure P(B-b-S) triblock (> 30 MPa) 83 . Therefore, it can be assumed that the poor tensile properties exhibited by the P(My-b-S) diblock copolymers were also due their network architecture since only one end of the P(My) chains was anchored to a PS glassy domain. The untrapped P(My) end-blocks were thus able to relax stress by chain reptation, reducing the elastic response of the physical network.

The modulus of elasticity (E), also known as Young's modulus, is defined by the slope of the linear elastic region of the stress-strain curve 84 . As expected, synthesized P(My-b-S) Young's modulus, ranging from 10 1 to 10 2 MPa in order of magnitude, is between P(My) Young's modulus (EP(My) ~ 10 -3 MPa) 16 and PS Young's modulus (EPS ~ 10 3 MPa) 85 . This elastic modulus range measured experimentally is in agreement with that of P(I-b-S) triblock copolymers (1 MPa < E < 900 MPa) prepared by anionic polymerization and exhibiting an overall Mn of about 100 kg.mol -1 and FS = 0.19-0.58 86 . Furthermore, it is of interest to note that tensile testing for P(My-b-S) samples showed ultimate tensile strength values ranging from 0.4 to 0.9 MPa and elongations at break ranging from 1 to 14%. This suggests that mechanical properties of P(My-b-S) block copolymers can be readily tuned based on molar composition. The greater the molar fraction of My, the higher the elongation at break and the lower the Young's modulus and the ultimate tensile strength as previously reported by Bolton et al. for β-myrcene/α-methyl-p-methylstyrene triblock copolymers 21 .

Synthesis of styrene-β-myrcene-styrene triblock copolymers

Varying molecular architecture can dramatically influence the physical properties of block copolymers. It can be assumed that My-S diblock copolymers prepared previously exhibited poor mechanical properties since only one end of the rubbery block was anchored to a glassy domain. The synthesis of symmetric S-My-S triblock copolymers (all entanglements are ideally physically trapped) by NMP was thereby envisioned. To that end, a poly(ethylene-stat-butylene) statistical copolymer capped by a SG1 nitroxide moiety at each end, abbreviated PEB-(SG1)2 (Mn = 5.7 kg.mol -1 , Ð = 1.17

β-Myrcene / styrene statistical copolymerization

My/S mixtures, with initial S molar feed compositions fS,0 = 0.10-0.94 were copolymerized by NMP in bulk using NHS-BB as initiator and a target molecular weight at full conversion of Mn,theo = 30 kg.mol - 5 indicates the formulations of the various My/S copolymerizations.

• Effect of feed composition on copolymer composition.

An investigation of how feed composition controls the copolymer composition was first performed.

The compositions of the copolymers (FS and FMy = 1 -FS, Table 6 for the final compositions) were determined by 1 H NMR (see experimental section for full information). The terminal model 88 KT method. The errors associated with the experimental data were derived from the standard errors of the slopes from FR and KT plots. Although the KT method is more reliable than the FR method, it is still a linearization. Statistically, a non-linear least-squares fit to the Mayo-Lewis equation is probably the soundest method to determine the desired parameters 92 . Therefore, the reactivity ratios were also determined using a non-linear least-squares fitting of the data. Using the reactivity ratios determined by the KT method as initial guesses, the statistical fit to the data yielded reactivity ratios rMy = 2.90 ± 1.25 and rS = 0.47 ± 0.27 (see Supporting Information), which are consistent compared to those obtained by the other approaches. For all the methods, rMy > 1 and rS < 1, suggesting that My monomer is more reactive than S monomer toward both propagating species (--My • and --S • ). The macro-radicals had a marked preference for My and the copolymers formed instantaneously were always richer in My than S.

122

A compositional drift was observed as indicated by the product rMy × rS being less than unity. In the initial stages of the copolymerization, My was incorporated faster, producing a My-rich copolymer. When • Effect of feed composition on kinetics.

The initial molar feed composition of S (fS,0) not only influences the copolymer composition but also influences the polymerization kinetics. <kp><K> kinetic values for this binary system (the product of the average propagation rate constant <kp> and the average equilibrium constant <K>) are given in Table 6, which introduces the final molecular characterization of the various copolymers as well (semilogarithmic plot of overall conversion ln((1-XNMR) -1 ) versus time, Mn versus XNMR plot and Đ versus XNMR plot provided in the Supporting Information, Figures S10a, S11a and S12a respectively). All the GPC chromatograms indicated the distributions were monomodal as illustrated in Figure 9 for the experiment My/S-50.

Furthermore, dispersity versus overall conversion plot displayed a Poisson-like distribution (Đ ≈ 1 + 1/DPn) 102 for every experiment with relatively broad Đ values at the early stages at XNMR < 0.20 (XMy < 0.26 and XS < 0.17) and narrower ones during the subsequent stages of the copolymerization. Such a kinetic feature highlights a rapid initiation with respect to propagation and negligible chain-transfer and termination reactions.

For fS,0 = 0.10-0.69, the initial molar composition of the mixture did not significantly affect the copolymerization kinetics. The overall polymerization rate was relatively low with <kp><K> = 1.11-2.21 10 -5 s -1 and XNMR ≤ 0.52 (XMy ≤ 0.56 and XS ≤ 0.47) after 7 h. The first-order kinetic behavior was strongly marked as observed with the linear trend of the ln((1-XNMR) -1 ) versus time curves (R 2 ≥ 0.96 from 0 to 420 min). This kinetic behavior was closely similar to that of My homopolymerization under the same experimental conditions (experiment My-5, Table 1 and Table 2). Accordingly, for fS,0 < 0.70, the copolymerization kinetics was governed by the monoterpene kinetics. Besides, a good control of the reactions was observed with Mn increasing linearly with XNMR and slightly plateauing at higher conversions

126

(Mn on average 87% of Mn,theo at the end of the experiments) and Đ ≤ 1.43 from the onset to the end of these copolymerizations.

For fS,0 = 0.80-0.94, My/S statistical copolymerizations were faster as indicated in Table 6 The higher overall conversions achieved for fS,0 > 0.70 resulted in an apparent loss of control.

Indeed, non-linear ln((1-XNMR) -1 ) versus time curves combined with higher deviations of experimental average molar masses from the predicted line were observed. Mn were on average 58% of Mn,theo at the end of the experiments My/S-80, My/S-90 and My/S-95. It should be first noted that the linear PS standards used for calibrating the GPC might not be very accurate for these statistical copolymers containing FMy = 0.16-0.34. The differences in hydrodynamic volumes of these copolymers having flexible C6/C8 side chains and that of the PS standards may be apparent. Although corrections were made using Mark-Houwink-Sakurada equation, the only MHS coefficients found for P(My) may not be perfectly appropriate for this study.

Indeed, KP(My) and αP(My) were obtained at 30 o C for a broad average molar mass range of 1,4-P(My)s (Mn = 6.3-55.5 kg.mol -1 ) with 10 mol% of 3,4-content. However, such significant deviations cannot be only attributed to the GPC data. It is likely that irreversible termination reactions became more prevalent at high conversions resulting in a loss the SG1 functionality on the chain-end and thus more dead chains. Chain transfer side reactions to monomer, generating monomeric radicals giving rise to a population of very short chains, was well-investigated for the NMP of styrene 106,107 . Even at a temperature as low as 110 o C, NMP of S in toluene mediated by SG1 generated a significant population of new short chains presumably due to chain transfer to S 108 . For this polymerization, it was reported that the total number of chains increased linearly with conversion, resulting in a gradual increase in Ɖ combined to a plateau of Mn < Mn,theo. Despite of higher overall polymerization rates, My/S copolymerizations for fS,0 > 0.70 still exhibited relatively narrow molecular weight distributions (Đ ≤ 1.37).

• Active feature of My-rich and S-rich P(My-stat-S) copolymers.

Two different My/S statistical copolymers, namely My/S-07 (FS = 0.07, My-rich P(My-stat-S)) and

My/S-55 (FS = 0.55, S-rich P(My-stat-S)), were synthesized under the same previous experimental conditions and their chain end fidelity was studied. Table 7.A gives the molecular characterization of My/S-07 and My/S-55 copolymers. My-rich P(My-stat-S) macroinitiator was shown to have an elevated fraction of chain-end nitroxide groups (~ 94 ± 6 mol% SG1 chain-end functionality, standard deviation derived from the difference in Mn value between the GPC result and the corrected Mn using MHS equation) based upon 31 P NMR measurements (see Supporting Information, Figure S5a with diethylphosphite as an internal reference, the presence of the phosphorus atom in the SG1 molecule was used as a tracer 72,109 ). It should be noted that the calculation of the fraction of SG1 groups was done by considering similar relaxation times between diethylphosphite and the polymer. This elevated fraction of SG1 groups was mainly due to a slow copolymerization resulting in a low overall conversion and a negligible contribution of side reactions (Mn,1 was 89% of Mn,theo,X1 and Đ < 1.3). S-rich P(My-stat-S) macroinitiator contained a lower "living" chain fraction (~ 72 ± 5 mol%), due to a higher overall conversion achieved leading to a slight loss of control (Mn,1 was 81% of Mn,theo,X1).

In order to confirm the active feature of these two macroinitiators first characterized spectroscopically, two sets of tests were then performed to evaluate the effectiveness of the expected SG1capped copolymers to cleanly initiate a second batch of monomer: chain-extensions of My/S-07 and a) The indexes "1" and "2" refer respectively to the final features of the P(My-stat-S) macroinitiator and the whole chain-extended P(My-co-S) copolymer (macroinitiator + P(My) or PS segment added).

b) FSG1 corresponds to the living molar fraction of macroinitiator chains terminated by a SG1 unit and was measured by 31 P NMR in CDCl3.

c) Mn,theo,X1 corresponds to the predicted Mn,1 at the respective X1 measured experimentally and was calculated as follows:

Mn,theo,X1 = X1 Mn,theo,1 with Mn,theo,1 = 40 kg.mol -1 at X = 1.0. d) Mn,theo corresponds to the targeted number-average molecular weight of the whole chain-extended P(My-stat-S) copolymer at X = 1.0. e) Mn,theo,X2 corresponds to the predicted Mn,2 of the whole chain-extended P(My-stat-S) copolymer (macroinitiator + second block added) at the respective X2 measured experimentally and was calculated as follows: Mn,theo,X2 = X2 (Mn,theo -Mn,1) + Mn,1 (= Predicted Mn of the second block added at conversion X2 + experimental Mn of the macroinitiator). Accordingly, the living feature of both My-rich and S-rich P(My-stat-S) macroinitiators was confirmed as also observed by the monomodal nature and the clear shift to the left of the GPC traces corresponding to the chain-extended products (see Supporting Information, Figure S13a). Furthermore, both My and S monomers were able to be cleanly added to the macroinitiators with slower My chain-extension kinetics as previously highlighted in comparison to My/S copolymerization kinetics. Indeed, X = 0.15-0.16 were achieved after 3 h reaction for each chain-extension with My resulting in the addition of relatively short elastomeric block ≤ 9.3 kg.mol -1 (we did not let the polymerization occur for long enough).

Chain-extensions with My from My/S-07 macroinitiator (experiment My/S-07-My.b in Table 7.C) and from My/S-55 macroinitiator (experiment My/S-55-My.b in Table 7.C) in bulk at 120 °C for 5 h were then performed in order to increase the polymerization rate. As expected, faster chain-extensions were achieved with X = 0.31 at 3h for experiment My/S-07-My.b and X = 0.39 at 3h for experiment My/S-55-My.b compared to lower X = 0.15-0.16 at 3 h in toluene at 110°C. Furthermore, the final conversions X ≥ 0.48 at 5 h (Table 7.C) for the chain-extensions performed in bulk at 120 o C were significantly greater.

Interestingly, the number-average degree of polymerization of the bloc added for the chain-extensions Tg values were also measured for P(My-stat-S) statistical copolymers exhibiting FS = 0.08-0.96, as presented in Figure 11 (DSC traces given in the Supporting Information, Figure S16a). Evidently, as the final molar content of styrene in the copolymer increased, Tg increased due to the pendant aromatic groups of S repeat units, providing rigidity and hardness. An exponential trend was observed with liquid rubber P(My-stat-S) copolymers achieved for FS ≤ 0.48, exhibiting Tg < -40 °C. A more significant dependence of Tg on FS was apparent for S-rich copolymers. Indeed, Tgs ranged from -9 °C to + 77 °C for FS = 0.66-1.00. This may be explained by the sequence distribution of My repetitive units within the polymer chains. The propagating species has a clear preference for My, even at a low molar fraction of My in the final copolymer, and short P(My) segments can be formed, resulting in "elastomeric knots" in the matrix. These peculiar connections, exhibiting a high flexibility with long side chains, may significantly disrupt the PS network. Theoretical Tg, Tg,th, for these copolymers can be predicted by the Fox equation 115 (7) where ɯMy and ɯS are the weight fractions of My and S, and Tg,My and Tg,S are the experimental Tgs of P(My) and PS homopolymers.

1/Tg,th = ɯMy/Tg,My + ɯS/Tg,S copolymers. Interestingly, a similar dependence of Tg on FS was reported by Halasa and coworkers for P(Iran-S) random copolymers 120 as shown in Figure 11. The higher Tgs measured can be mainly attributed to the higher experimental Tg for PI (-65 °C) than that of P(My) (-77 °C) and the greater Mn > 120 kg.mol - 1 exhibited by the various P(I-ran-S) copolymers.

Conclusion

In this study, the optimized NMP of My was accomplished at 120 °C in bulk initiated by the alkoxyamine NHS-BB without requiring additional SG1 free nitroxide. Relatively low apparent rate constant (kpKMy = 4.3 10 -5 s -1 ) compared to S homopolymerization by NMP, first-order kinetic behavior (linear trend of ln((1-XMy) -1 ) versus time with R 2 > 0.99) and Mn increasing linearly with XMy (Mn on average 90% of Mn,theo) were apparent under these experimental conditions when targeting moderate Mn,theo ~ 30

Chapter 2. Functionalization of poly(β-myrcene) by β-myrcene /

glycidyl methacrylate nitroxide-mediated copolymerization

Preface

The thermoplastic elastomer desired by ArianeGroup has to provide a tough and stable join between composite matrices. The assembly can be done either by welding (use of polar thermoplastic composites)

or by adhesive bonding (use of thermoset composites bearing functional groups). Accordingly, the functionalization of the targeted styrenic block copolymer TPE appears to be important. In the case of a welding, the functional units carried by the TPE would increase the miscibility at the welding interface to favor chains inter-diffusion with the polar thermoplastic composites. If adhesive bonding is implemented with thermoset composites, the incorporation of functional units into the TPE could be valuable to encourage covalent bonding at the interface.

Abstract

Introduction

Elastomers based on 1,3-dienes exhibit unique and diverse properties. Their significant reversible elongation with low hysteresis combined with low modulus and poor abrasion make them suitable materials for automobile and aircraft tires, steam hoses and roofing membranes among others. 1 However, the incompatibility of hydrocarbon elastomers with polar entities considerably narrows their potential range of applications. Generally, the addition of functional units to non-polar polymers has been extensively explored in order to use them as coupling agents, adhesives, impact modifiers and compatibilizers, for instance, and thus enhance their service properties. 2 Post-polymerization modification was widely used to prepare functional elastomers. Functionalization of poly(3-methylenehepta-1,6-diene) by reacting the pendant vinyl groups with various polar groups via a light-mediated thiol-ene method allowed the synthesis of functional butadiene-based materials with greater hydrophilicity. 3 An epoxide functionalized poly(isobutylene-co-isoprene) derivative was also used to cleanly prepare a specific exo-diene, allowing subsequent functionalization by Diels-Alder cycloaddition and/or ring-opening of maleic anhydride adducts. This way, valuable multifunctional materials were achieved such as a graft copolymers with poly(ethylene oxide) (PEO) along with carboxylic acid moieties to name but one. 4,5

The insertion of polar groups in the elastomer can also be achieved directly during the polymerization process. Although coordination, 6 anionic 7,8 and cationic 9 polymerizations are typically performed to synthesize polydienes, their sensitivity to impurities and to the presence of functional groups limits their use for functionalization of diene-based elastomers by direct polymerization. 10 It is for instance commonly reported that functional groups can coordinate with selective catalysts and poison its activity during coordination polymerizations. 9 An intriguing tool is reversible-deactivation radical polymerization (RDRP) 11-14 which is robust for the synthesis of well-tailored complex macromolecular architectures with controlled molecular weights. This technique appears to be well-suited for the synthesis of diene-based elastomers. 15-18 Moreover, RDRP is desirable as a route towards well-defined functional polymers thanks to its great tolerance for functional units of monomers and reaction conditions. Well-defined glycidylmethacrylate-based polymers were mediated by 2-cyanoprop-2-yl 1-dithionaphthalate (CPDN) RAFT agent and initiated by 2,2'-azobisisobutyronitrile (AIBN) in bulk and in benzene. 19 Among RDRP techniques, nitroxide-mediated polymerization (NMP) 12,13,20 under homogeneous conditions in bulk is one of the easiest systems to implement. 2,2,6,6-Tetramethylpiperidine-N-oxyl (TEMPO) nitroxide as trapping agent and benzoyl peroxide as free radical initiator served as the pioneering bicomponent initiating system used for polymerizing styrenic monomers in bulk. 21 However, the high polymerization temperature needed (T ~ 120-130 o C) when using TEMPO combined with the inability to control effectively monomers other than styrenics 22-24 triggered the development of new alkoxyamine structures. The synthesis of alkoxyamine unimolecular initiators like 2,2,5-trimethyl-4-phenyl-3azahexane-3-oxyl (TIPNO) 25 and 26 increased the variety of monomers polymerizable by bulk NMP with lower temperatures accessible. TIPNO and its derivatives were used to control effectively homopolymerizations of acrylates such as tert-butylacrylate, 27-29 acrylamides such as N,N-dimethylacrylamide 29, 30 and also 1,3-dienes such as isoprene (I). 27, 30 The NMP of I in bulk using TIPNO was effectively controlled with a dispersity Đ < 1.20 and a linear relationship between the number-average molecular weight (Mn) and the conversion. Statistical copolymerizations of I in bulk with styrenics, acrylates and methacrylates were then performed in a controlled manner. 24 Random and block copolymers based on associations of 1,3-dienes with functionalized vinyl monomers using TIPNO were successfully prepared. Well-controlled 1,3-dienes polymerizations were reported for SG1-mediated NMP as well.

N-tert-butyl-N-[1-diethylphosphono-(2,2-dimethylpropyl)] nitroxide (DEPN or SG1)
Successful NMP of I in bulk at 115 o C using a range of SG1-based alkoxyamines was studied.

Poly(isoprene)s (PIs) exhibiting Đ ˂ 1.20 and Mn close to the predicted one at approximately 40% conversion were achieved. 34 We reported recently the optimized NMP of β-myrcene (My), 35 an acyclic monoterpene 1,3-diene, in bulk, initiated by the SG1-based succinimidyl ester-terminated NHS-BlocBuilder, 36-38 derived from the esterification of BlocBuilder TM initiator. 39 Moreover, SG1-type initiators led to the synthesis of well-define poly(acrylate)s 40,41 such as poly(n-butyl acrylate) and poly(methacrylate)s like poly(methyl methacrylate). 42, 43 Accordingly, these studies highlighted the ease of NMP based on the second generation of acyclic nitroxides to homopolymerize conjugated 1,3-dienes and to copolymerize them with functional monomers.

The copolymerization of ``renewable`` My with functional monomers has been rarely reported although it could provide a wide latitude of polymer properties. Trumbo studied the free radical copolymerization of My with methyl methacrylate and p-fluorostyrene at 65 o C initiated by AIBN. 44 To the best of our knowledge, the controlled/living copolymerization of My with polar monomers by ionic polymerization or RDRP has never been performed. Taken the robustness of NMP based on the second generation alkoxyamines to copolymerize 1,3-dienes with functionalized vinyl monomers into account, the synthesis of functional My-based copolymers by SG1-mediated NMP appears to be promising.

In the present study, we report on NHS-BlocBuilder-initiated copolymerization of My with glycidyl methacrylate (GMA) in a controlled manner at 120 o C in bulk. Copolymer composition and kinetics were studied according to the molar feed composition of GMA. The glass transition behavior of the resulting P(My-stat-GMA) statistical copolymers was explored. Their active feature was assessed by phosphorus nuclear magnetic resonance ( 31 P NMR) and by their ability to initiate a fresh batch of monomer. Lastly, a P(My-b-GMA) diblock was achieved and modified using morpholine in order to synthesize readily an amphiphilic poly(β-myrcene-block-2-hydroxy-3-morpholinopropyl methacrylate) P(My-b-HMPMA) diblock copolymer.

Results and discussion

Attempts to copolymerize β-myrcene with various functional monomers by NMP

The insertion of functional groups into poly(β-myrcene) P(My) elastomer was first tried by nitroxide-mediated copolymerization with various polar monomers (Figure 1). Maleic anhydride (MA), bearing the reactive acid anhydride moiety, was first attempted. Acrylic acid (AA) was not copolymerized due to the sensitivity of the SG1 free nitroxide towards decomposition by some acids. 47 Although Charleux and coworkers demonstrated that active SG1-capped low Mn P(AA)s and P(AA-co-S)s exhibiting Đ < 1.4

and Mn close to the targeted values can be produced using SG1-based alkoxyamine in 1,4-dioxane, 48, 49 Lessard et al. highlighted the possible degradation reactions by the organic acid with SG1 as well as the loss of the propagating macroradicals due to β-hydrogen transfer to SG1. The experimental conditions of the various copolymerizations performed and the characterization of the resulting copolymers at the end of the experiments are summarized in Table 1. The complete kinetic study including the semi-logarithmic plot of overall conversion ln((1-X) -1 ) versus time, Mn versus X plot and Đ versus X plot for every experiment is provided in the Supporting Information, Figures S1b, S2b and S3b. It should be added that the Mn values reported in this whole study directly came from GPC measurements, calibrated with PS linear standards, and no corrections were made by using Mark-Houwink-Sakurada coefficients. Even though trends of Mn versus X were meaningful and direct Mn comparisons between the experiments done were relevant due to the similar structures of the copolymers produced, the deviations of experimental Mn values from predicted ones have to be carefully scrutinized. They did not necessarily correspond to accurate values and were mostly discussed in comparison with deviations determined from other experiments. = 0.49) were first performed (Table 1). When adding 10 mol% of MA in the feed, it is of interest to note the early consumption of about one fifth of the monomers at the commencement of the reaction. Moreover, signification deviations of Mn from the theoretical predictions (Mn/Mn,theo ˂ 0.40 all along the experiment, strong plateauing compared to My/tBuA-90, Figure S2) suggest the disruption of the NMP process via side reactions. To provide a better insight into the influence of MA on the control of the polymerization, a second reaction was conducted with a stoichiometric initial feed composition of monomers (My/MA-50, Table 1).

No polymerization occurred as signaled by the absence of polymer peak in the GPC chromatograms and the gravimetric conversion already reached 50% at t = 0 min. The reaction medium turned dark yellow/orange after only 30 minutes and was dark brown at the end of the reaction. These observations suggest a side product was formed at the initial stages of the polymerization, possibly followed by another • My/tBuA NMP.

The polymerization of acrylates by NMP has become much more facile with second-generation alkoxyamines such as TIPNO and BlocBuilder TM . However, its main challenge is derived from selection of the proper conditions due to their elevated propagation rate constants (kp) compared to styrenics. 22,57 My/tBuA copolymerization, mediated by NHS-BB, was carried out in bulk at 115 o C and targeting about 40 kg.mol -1 at X = 100% (Table 1). An investigation of the influence of the initial feed composition over the control of the polymerization was performed with fMy,0 = 0.90, 0.70 and 0.50 (experiments My/tBuA-90, My/tBuA-70 and My/tBuA-50 respectively). Regardless of the initial feed composition, tBuA conversion was higher than My conversion (Figure S4b, Supporting Information) indicating the higher reactivity of this acrylate with a likely predominance of the propagating macro-radical to add tBuA monomer, whatever the terminal active unit (---My* or ---tBuA*). Relatively low dispersities lower than 1.40 were obtained, even at high overall conversion X ˃ 60%, for every experiment (Figure S3). However, deviations of experimental Mn from Mn,theo were greater as ftBuA,0 increased. Although the copolymerization of My with 10 mol% of tBuA led to a satisfactory control of the NMP with a linear increase of Mn with conversion, the initial addition of 30 and 50 mol% of tBuA in the feed affected significantly the kinetics of the reaction. Indeed, experiment My/tBuA-70 exhibited Mn plateauing very early and remaining constant during the whole reaction and this observation was even more marked for experiment My/tBuA-50 with higher ftBuA,0 (Figure S2). The predominance of chain transfer side reactions due to the presence of tBuA, terminating irreversibly the copolymer chains and thus lowering Mn values can be hypothesized. A secondary backbiting reaction, consisting of an intramolecular chain transfer event in which the propagating chain-end radical loops back and abstracts a hydrogen atom on its own backbone, may have occurred. It is now well-established that backbiting occurs during radical polymerization of acrylates, which can lead to propagation or chain-scission. 58, 59 More specifically, backbiting reactions in the presence of tBuA leading to the formation of mid-chains radicals was also reported. ) and a low chain end fidelity were observed, suggesting that substantial irreversible termination reactions took place as well. As ftBuA,0 increased, the molar fraction of 1,2-content increased whereas the fraction of 1,4-addition decreased (Figure S5b, Supporting Information). While the SG1 NMP of My resulted in 1,4-rich-P(My) containing a minor proportion of 1,2-and 3,4-contents (˂ 20%), 35 it appears that 1,2-addition is encouraged due to the presence of tBuA likely as a terminal or penultimate unit of the propagating species.

• My/MMA NMP.

Although some exceptions were reported using well-tailored alkoxyamines such as 2,2-diphenyl-3-phenylimino-2,3-dihydroindol-1-yloxyl nitroxide (DPAIO), 63 the simplest way to control effectively the polymerization of methacrylate monomers by SG1-mediated NMP is to use a small concentration of controlling comonomer such as styrene. 42, 64 Benoit and coworkers reported the effective control of the bulk NMP of MMA initiated by 2,2,5-trimethyl-3-(1′-phenylethoxy)-4-phenyl-3-azahexane alkoxyamine and using isoprene at about 20 mol% in the feed. 54 In order to know if methacrylates can be readily copolymerized with My by NMP, MMA monomer was selected as a reference methacrylate. A stoichiometric feed composition of monomers (fMy,0 = 0.50, experiment My/MMA-50, Table 1) was heated at 110 o C in the presence of NHS-BB initiator. As shown in Figure S1 by the linear relationship between ln((1-X) -1 ) versus reaction time t (correlation coefficient R 2 ˃ 0.98), a first-order kinetic behavior was observed, assuming a quasi-constant concentration of the active propagating species during the polymerization. Linear trends can be noted regarding individual My and MMA conversions with respect to the polymerization time as well (Figure S4b, Supporting Information). Moreover, Mn values increased linearly with the overall conversion (Figure S2) and narrow molecular weight distributions were obtained (Đ ≤ 1.32, Figure S3). Consequently, the reversible dynamic equilibrium between the dormant SG1terminated P(My-co-MMA) copolymers and active P(My-co-MMA) radicals was likely established. At X ˃ 20%, non-negligible deviations of the Mn from the predicted ones can be noted yet (Figure S2). This tendency of Mn to plateau at increasingly higher conversions is presumably due to irreversible terminations, commonly observed for copolymerizations of methacrylates using SG1-based alkoxyamine. 52,65,66 Since the copolymerization was done in bulk, it can be assumed that the side reactions resulted mainly from βhydrogen transfer to SG1 52 or intramolecular chain transfer. 67 The satisfactory results regarding the SG1 nitroxide-mediated copolymerization of My with MMA in bulk prompted us to study more exhaustively the copolymerization of My with another methacrylic ester that contains functional groups, such as the epoxy-functional glycidyl methacrylate (GMA).

My/GMA statistical copolymerization

We reported previously the optimized NMP of My (experiment My/GMA-100 herein, Tables 2 and3) performed at 120 o C in bulk and initiated by NHS-BB alkoxyamine, leading to Mn on average 90% of predicted Mn,theo, Đ ˂ 1.3 and a resulting SG1-capped P(My) able to reinitiate cleanly a second batch of styrene. 35 Accordingly, we attempted to copolymerize My and GMA, with initial My molar feed compositions fMy,0 = 0.10-0.90, under the same reaction conditions. Notably, well-controlled NMP copolymerizations of isoprene (I) with small amounts of GMA (1-10 wt%) in toluene were also performed at 120 o C in a medium pressure reactor. 68 The formulations for the various My/GMA copolymerizations are summarized in Table 2.

Table 2. My/GMA copolymerization formulations for various compositions at 120 o C in bulk initiated by NHS-BB and targeting Mn,theo = 30 kg.mol -1 at X = 100%. • Effect of feed composition on copolymer composition.

The influence of fMy,0 over the molar composition of the resulting P(My-co-GMA) copolymer was explored at relatively low overall conversion (X ≤ 23.2% for every copolymerization). The compositions of the copolymers synthesized (FMy and FGMA = 1 -FMy, Table 3 for the final compositions) during the experiments were determined by 1 H NMR (see Experimental Section for full information and Figure S6b in Supporting Information for the spectral assignments of 1 H NMR from experiment My/GMA-50). It should be noted that the terminal model was used, which assumes that the reactivity of an active center depends only on the nature of the terminal monomer unit of the growing chain. 69 Table S1b in the Supporting Information contains experimental data used to estimate the monomer reactivity ratios for low-conversion P(My-co-GMA)s via the Fineman-Ross (FR) 70 and the Kelen-Tudos (KT) 71 approaches (equations used for the calculation of the reactivity rations can be found in the Supporting Information). Figures 2a and2b show the plots of the parameters for both linear methods, yielding rMy = 0.80 ± 0.31 and rGMA = 0.71 ± 0.15 (rMyrGMA = 0.57 ± 0.27) via the FR approach and rMy = 0.48 ± 0.12 and rGMA = 0.53 ± 0.18 (rMyrGMA = 0.25 ± 0.17) using the KT formula. The errors associated with the experimental data were derived from the standard errors of the slopes from FR and KT plots.

Although monomer conversions considered were relatively low, the effect of possible composition drift during the copolymerization could not be taken into account by FR and KT approaches. Tidwell and Mortimer highlighted the defects of the linear methods and suggested the use of a non-linear least-squares (NLLS) procedure. 72 This latter method is known to avoid the compositional drift, altering the chemical composition of the copolymer over the period of conversion. Therefore, the non-linear least-squares fitting to the data was also used to determine the reactivity ratios (Supporting Information). Since a starting point needs to be provided, the reactivity ratios calculated by the KT method were used as initial guesses. NLLS procedure applied to the Mayo-Lewis equation 73 showed good agreement with the other methods employed in this study, with rMy = 0.49 ± 0.13 and rGMA = 0.50 ± 0.13 with 95% confidence bounds. These results have to be taken carefully into account since only 11 data points were used to implement the NLLS method.

Regardless of the method applied, rMy ˂ 1 and rGMA ˂ 1 showing the preference of --My • and --GMA • propagating species for adding the other monomer (GMA and My respectively) and each of these propagation steps proceed at a similar rate since rMy ≈ rGMA. Interestingly, 0.25 ≤ rMyrGMA ≤ 0.57 indicating that the behavior of My/GMA system cannot be considered as an ideal copolymerization (rMyrGMA = 1, the relative rates of incorporation of the two monomers into the copolymer are independent of the identity of the unit at the end of the propagating species). As the product of the reactivity ratios decreases from one toward zero, there is an increasing tendency toward alternation, where each of the two types of propagating species preferentially adds the other monomer. Since rMyrGMA ≥ 0.25 is relatively far from zero, the predominance of the statistical placement over alternating arrangement of My and GMA monomers along the copolymer chain can be assumed. Other diene/GMA were copolymerized by conventional free radical polymerization. P(I-stat-GMA) statistical copolymers were synthesized by benzoyl peroxide-initiated free radical polymerization in bulk at 80 o C and rI = 0.195 ± 0.005 and rGMA = 0.135 ± 0.005 were determined graphically. 74 More recently, Contreras-Lopez and coworkers reported rI = 0.119 ± 0.048 and rGMA = 0.248 ± 0.161, via the nonlinear method of Tidwell-Mortimer with 95% confidence intervals, for I/GMA copolymerization in bulk at 70 o C initiated by benzoyl peroxide. 75 Accordingly, lower reactivity ratios were calculated for I/GMA system with rIrGMA less than 0. Figure 2c shows the curve of instantaneous copolymer composition (mole fraction of GMA) obtained for a given feed monomer composition regarding the My/GMA system as well as I/GMA system studied by Contreras-Lopez et al.. 75 A similar dependence of FGMA over fGMA,0 can be noted for both systems with GMA-rich instantaneous copolymers produced at fGMA,0 ≲ 0.45 and diene-rich instantaneous copolymers formed at fGMA,0 ≳ 0.55. The Mayo-Lewis plot exhibits an azeotropic point at fGMA,0 = 0.47 for My/GMA system where the instantaneous composition of the copolymer is equal to the composition of the monomer feed. The greater deviations observed from the azeotropic composition for the I/GMA copolymerization were due to the lower reactivity ratios and thus the higher rates of cross-propagation measured.

• Effect of feed composition on kinetics.

It is of interest to know if My and GMA monomers are generally more or less reactive in copolymerization than indicated by their respective rate of homopolymerization. Even though kinetic data regarding My polymerization are scarce, isoprene (I) can be chosen as a reference conjugated 1,3-diene.

The propagation rate constant of free radical polymerization of I (kp,I) in bulk initiated by di-tert-butyl peroxide under irradiation at 5 o C was estimated at 125 ± 30 L.mol -1 .s -1 . 76 kp,I = 50 L.mol -1 .s -1 was calculated by Morton et al. for the emulsion polymerization of I at 60 °C with diisopropylbenzene monohydroperoxide / tetraethylene pentamine (DIBHP/TEPA) catalyst system. 77 On the other hand, kp,GMA for the free radical polymerization of GMA at 50 o C was estimated through pulsed laser polymerization (1 230 L.mol -1 .s-1) 78 and through quantum chemistry simulations (1 010 L.mol -1 .s -1 ). 79 Consequently, it can be argued that GMA polymerizes at a rate of about ten times faster than My via a radical process.

The semi-logarithmic kinetic plots of ln((1-X) -1 ) (X = overall conversion) versus reaction time are illustrated in Figure 3a. The linear trend seen for every My/GMA copolymerization until the end of the experiment (squared linear regression coefficient, R 2 ≥ 0.95) indicates the first-order kinetic behavior of the polymerizations, as described by Equation 2.

ln([M]0/[M]t) = ln([M]0/([M]0(1-X))) = ˂kp˃[P • ]time (2)
.

In Equation 2 f) No kinetic study led due to the very early ``caking`` (high viscosity) of the reaction medium.

g) Results from My homopolymerization performed previously under the same experimental conditions. 35 the lack of irreversible termination reactions so that [P • ] was nearly constant. 22 As expected for the NMP process, the homolytic decomposition of NHS-BB alkoxyamine into the propagating alkyl radical and the persistent radical SG1 • allowed the establishment of a dynamic equilibrium between a low concentration of active propagating chains and a predominant concentration of SG1-capped dormant chains. These latter are unable to propagate or terminate, allowing thus the controlled growth of the macro-radicals. The reversible trapping can thus be assumed effective for GMA/My copolymerizations. It should be noted that deviations from the linear trend can be observed at the last stages of copolymerization, mostly for experiments My/GMA-40, My/GMA-50 and My/GMA-60 (Figure 3a). A decrease in [P • ], which might result from irreversible termination reactions increasing the concentration of the persistent radical, could lead to the downward curvature trend seen for these reactions. ln((1-X) -1 ) versus time plot in the linear region allows the determination of the apparent rate constant, which can be defined as the product of the average propagation rate constant ˂kp˃ with the concentration of propagating macro-radicals [P • ], ˂kp˃[P • ]. The apparent rate constant can be related to the average equilibrium constant ˂K˃ as the product ˂kp˃˂K˃ by making some assumptions regarding the equilibrium between dormant and active chains. It should be noted that an average propagation rate constant and average equilibrium constant were used since a statistical copolymerization is studied, consisting of a diene and a methacrylate, which likely have very different individual equilibrium constants. The ˂K˃ is defined in terms of the [P • ], the concentration of free nitroxide [N • ] and the concentration of dormant alkoxyamine terminated species [P-N], as shown in Equation 3.

˂K˃ = ([P

• ][N • ]) / [P-N] (3)
The overall polymerization rate can be defined by the rate of chain propagation which is given by Equation 4.

Rp = -d[M]/dt = ˂kp˃[P • ][M] (4
The reported ˂kp˃˂K˃ in Table 3 were measured during the early stages of the copolymerization, at low or moderate final conversion (X ≤ 45.2%) except for experiments My/GMA-40 and My/GMA-10

where the final overall conversions were 67% and 57%, respectively. In our study, a stoichiometric amount of SG1 free nitroxide and radical initiator just after initiation could be hypothesized since no extra SG1 was added in the initial feed. In any case, a linear growth of Mn versus X was observed (Figure 3c), making Equation 5 applicable. As fGMA,0 increased, the ˂kp˃˂K˃'s increased with ˂kp˃˂K˃ = (9.1 ± 0.2) 10 -5 s -1 at fGMA,0 = 0.10 to ˂kp˃˂K˃ = (157.0 ± 0.9) 10 -5 s -1 at fGMA,0 = 0.90 (Table 3). These kinetic values are not surprising since kpK of pure My under identical reaction conditions 35 was estimated at (4.3 ± 0.7) 10 -5 s -1 while ˂kp˃˂K˃ = (66 ± 4) 10 -5 s -1 was measured for GMA/S copolymerization with fGMA,0 = 0.94 initiated by NHS-BB at 90 o C in 50 wt% 1,4-dioxane solution. 36 The presumable large difference in the kp between My (kp,I ~ 10 2 L.mol -1 .s -1 for isoprene, a structurally similar 1,3-diene) 76, 77 and GMA (kp,GMA ~ 10 3 L.mol - 1 .s -1 ) 78,79 , as discussed above, can partly explain the ˂kp˃˂K˃ increase with fGMA,0 even though the contribution of the two individual activation-deactivation equilibrium constants cannot be neglected. The higher ˂kp˃˂K˃ values determined when using methacrylate-richer initial feeds agrees with previous studies, reporting similar kinetic trends for the tert-butyl methacrylate/S copolymerization at 90 o C in bulk initiated by BlocBuilder TM , 80 the GMA/S copolymerization at 90 o C in 1,4-dioxane initiated by NHS-BB 36 and the MMA/S binary system initiated by BlocBuilder TM in bulk at 90 o C. 64 A well-controlled RDRP exhibits linear Mn versus X and low dispersity with monomodal distributions. The general increase in ˂kp˃˂K˃ as a function of increasing fGMA,0 suggested the polymerizations were becoming less controlled. This was reflected in faster polymerization rates (Table 3) and higher Đ as shown in Figure 3b. While relatively low dispersities (Đ ≤ 1.38) were measured for Myrich initial feed composition (fGMA,0 ≤ 0.29), higher Đ values mainly ranging from 1.25 to 1.50 were obtained for experiments My/GMA-40, My/GMA-50 and My/GMA-60 with fGMA,0 = 0.40-0.60. This tendency was amplified for GMA-rich feeds with broader molecular weight distributions measured, with values typically between 1.35 and 1.55 (Figure 3b). Independently of the initial molar composition, similar Đ versus X trends can be observed. Disregarding the first inaccurate Đ values for X ˂ 10% (likely underestimation of the dispersity due to the low intensity of the polymer peaks on the GPC chromatograms), Đ was quite elevated at the commencement of the copolymerization (X = 15-25%) but became gradually lower as the overall conversion increased (X = 25-70%). It can be assumed that the dynamic equilibrium between the propagating macro-radicals and the dormant SG1-terminated species was not being established sufficiently at the early stages of the polymerization for the persistent radical effect to be present. At the last stages of the copolymerization (X = 70-95%), a broadening of the molecular weight distribution was apparent for every My/GMA experiment, probably underscoring the greater prevalence of termination reactions, resulting in a loss of the SG1 functionality on the chain end. Since a solvent free process was applied, it may be hypothesized that many of the irreversible terminations occurring resulted from β-hydrogen transfer 156 (disproportionation) from propagating P(My-stat-GMA) radical to SG1 52,67 or intramolecular chain transfer to polymer. 67 Figure 3c gives Mn versus X for the various experiments. It should be first noted that Mn values have to be carefully compared to the theoretical ones since all GPC measurements were performed using linear poly(styrene) standards, which might not be very accurate for P(My-stat-GMA) samples. A similar linear growth of degree of polymerization with overall conversion for every My/GMA copolymerization with Mn = 18.4-21.9 kg.mol -1 at X ˃ 80% can be seen (Figure 3c). This can be explained by the active feature of the macro-radical chain end for a prolonged period of time, presumably due to the relatively low average equilibrium constant ˂K˃ obtained when polymerizing GMA with My. Mn values started slightly plateauing at X ~ 40% and non-negligible deviations from the predicted line can be observed at the final 2 and Table 3.

The same legend at the bottom right of the figure is used for each of the three plots.

stages of the copolymerizations (Mn/Mn,theo = 71-88% at the end of the experiments, Table 3). It may be due to the occurrence of irreversible termination reactions, which were no longer retarded at X ˃ 40%.

The GPC traces for experiment My/GMA-50 exhibiting a nearly stoichiometric initial feed composition are shown in Figure 4. A monomodal shift of the GPC chromatograms can be observed with a slight tail at longer elution times, most likely resulting from irreversible termination reactions generating low Mn P(My-stat-GMA) chains. At t = 15 min, a minor population of longer chains was detectable at relatively short elution times (23-26 min, Figure 4). As discussed in a previous study reporting the NMP of My, 35 it may be caused by My auto-initiation, 81 generating an excess of P • and a fast auto-polymerization which makes the mediation by SG1 difficult. This second peak was easily detectable at the commencement of My-rich feed copolymerizations (fMy,0 = 0.60-0.90) and hardly noticeable for experiments with higher fGMA,0. This tends to confirm that this minor fraction of polymer consisted of high Mn P(My) chains resulting from My auto-polymerization.

Lastly, the influence of fGMA,0 over the regioselectivity of My units in the final copolymer can be discussed (see Figure S6b in Supporting Information for the determination of 1,2-, 3,4-and 1,4-motif for experiment My/GMA-50). As already mentioned for My/tBuA system above, the molar fraction of 1,2addition increased and that of 1,4-addition decreased with fGMA,0 (and thus FGMA). While My NMP produced P(My) exhibiting a large predominance of 1,4-content (~ 80 mol%) with a minor proportion of 1,2-units (˂ 10 mol%, experiment My/GMA-100, Table 3), the initial addition of 50 mol% of GMA resulted in P(Mystat-GMA) more than three times richer in 1,2-addition (28.5 mol%, experiment My/GMA-50, Table 3).

For rich-GMA feeds (fGMA,0 = 0.71-0.90), the fraction of 1,2-content exhibited by the final copolymers was around 40 mol%. This change in regioselectivity for My units was likely due to the presence of GMA as a terminal or penultimate unit, which might bring about either steric hindrance or electronic effects.

158

• Active feature of My-rich and GMA-rich P(My-stat-GMA) copolymers.

A paramount feature of a NMP-based polymer is its ability to form a block copolymer via the reactivation of the dormant chains capped by the nitroxide group. Two extra statistical P(My-stat-GMA) copolymers were synthesized under the same previous experimental conditions (My/GMA-37 with FMy,1 = 0.37 and My/GMA-78 with FMy,1 = 0.78, Table 4A) in order to deem their degree of the active groups, in other words, the fraction of SG1-terminated P(My-stat-GMA) chains. Several strategies can be adopted to demonstrate whether a SG1-mediated (co)polymer is significantly active enough to re-initiate a second block.

Phosphorus nuclear magnetic resonance ( 31 P NMR) was first used to probe the chain end fidelity of the synthesized My/GMA copolymers. This quantitative approach allows to determine the living chain fraction (LF) by detecting the phosphorus-containing SG1 nitroxide end group with diethyl phosphite as an internal reference (see Experimental Section for further details). 42 LF values for My-rich (My/GMA-78) and GMA-rich (My/GMA-37) macroinitiators were 82 ± 5 % and 71 ± 7 %, respectively (Table 4A, standard deviation derived from the difference in macroinitiator Mn value obtained from PS calibration and PMMA calibration, 31 P spectra in Supporting Information, Figures S7b andS8b), indicating a satisfactory retention of the SG1 fragment. The lower ``living`` fraction exhibited by My/GMA-37 can be mainly attributed to the relatively high ˂kp˃˂K˃ value (~ 40.10 -5 s -1 for My/GMA-30 with fGMA,0 = 0.70, Table 3) and thus the more difficult mediation of the binary system, as witnessed by the broader molecular weight distribution (D1 = 1.39), compared to My-rich NMP. Generally, these spectroscopic results are a clear indication of the active character of NMP-based P(My-stat-GMA) copolymers, which may allow efficient chain-extension to produce well-defined block copolymers.

To confirm the 31 P NMR quantitative analysis, My/GMA-78 and My/GMA-37 were used as macroinitiators in the polymerization of My, GMA and styrene (S) in 50 wt% toluene solution at T = 110-120 o C (Table 4B). For every experiment, a high Mn,theo was targeted (80-85 kg.mol -1 ) to ensure a clear shift of the GPC peak. The final characterization of the chain-extended products is given in Table 4C.

Regarding My-rich macroinitiator My/GMA-78, a clean chain-extension with S was performed (experiment My/GMA-78-S) with a significant increase of Mn from 11.3 to 28.4 kg.mol -1 and Đ remaining low (˂ 1.40, Table 4C). The monomodal nature of P[(My-stat-GMA)-b-S] was confirmed by GPC as indicated in Figure 5a'. Furthermore, the GMA chain-extension from the same My/GMA-78 macroinitiator was attempted in order to deem the crossover efficiency. It has to be noted that around 10 mol% of My with respect to GMA was initially added (Table 4B). Indeed, NMP of GMA with 10 mol% of My (experiment My/GMA-10, Table 2) allowed to slow down the reaction with a pseudo-control of the polymerization (Đ ≤ 1.52 at X ˃ 20%, Mn/Mn,theo ˃ 80% for every sample, Figures 3b and3c e) Predicted Mn,2 of the whole chain-extended diblock copolymer (MI block + second block added) at X2 measured experimentally and calculated as follows: Mn,theo,X2 = (X2/100)(Mn,theo-Mn,1)+Mn,1 (= Predicted Mn of the second block added at X2 + experimental Mn of MI).

and glycidyl methacrylate monomers at 120 o C in the presence of 10-20 mol% of isoprene and initiated by 2,2,5-trimethyl-3-(1′-phenylethoxy)-4-phenyl-3-azahexane. 54 This may suggest that the addition of a small amount of isoprene or β-myrcene could improve the NMP control for certain methacrylate and acrylate monomers, presumably by reducing the average equilibrium constant of the system. Table 4C shows the final features of the chain-extended polymer (experiment My/GMA-78-GMA/My) and the GPC traces in the course of the chain-extension are given in Figure 5a. Despite a slight broadening of the molecular weight distribution (Đ from 1.27 to 1.50), the chain-extension was successful with a monomodal shift of the GPC chromatograms and the addition of a GMA-rich second block exhibiting about 18 kg.mol -1 . While the macroinitiator was poor in GMA (FGMA,1 = 0.22), the chain-extension allowed to drastically alter the composition since the final diblock copolymer exhibited FGMA,2 = 0.65 (Table 4C), as determined by 1 H NMR.

In order to deem the influence of the composition of the My/GMA statistical copolymer over its ability to re-initiate cleanly a second batch of monomer, a GMA-rich P(My-stat-GMA) (My/GMA-37, FGMA,1 = 0.63, Table 4A) was used as a macroinitiator in the polymerization of My and S as described in Table 4B. For both experiments, a complete shift of the GPC peak (Figures 5b and5b') corresponding to the addition of a P(My) segment exhibiting around 21 kg.mol -1 (experiment My/GMA-37-My) and the addition of a PS segment with Mn ~ 27 kg.mol -1 (experiment My/GMA-37-S, Table 4C) showed the effective synthesis of the diblock copolymer. The 1 H NMR spectrum of the final My/GMA-37-S diblock copolymer is given in Supporting Information, Figure S9b, illustrating the ternary composition of the chainextended product synthesized. While a well-defined P[(My-stat-GMA)-b-S] was achieved (Đ2 = 1.46, Mn,2/Mn,theo,X2 ~ 98%, Table 4C), My chain-extension from My/GMA-37 was not as well controlled, as seen by the broader molecular weight distribution (Đ2 = 1.85) and the greater Mn deviation from the predicted one (~ 70%). This marked loss of control may be explained by the more elevated temperature used (T = 120 o C to perform a faster My polymerization) 35 as well as the higher conversion reached (X2 = 45.3%).

Another reason can also be mentioned. The polymerization was run for a relatively long period of time (t = 260 min) under a mild toluene reflux. A slight but gradual evaporation of the toluene was apparent during the experiment which, along with the consumption of My, made the reaction medium more and more viscous. A non-negligible influence of the viscosity, increasing with polymerization time, over the loss of control of the chain-extension can thereby be hypothesized. Moreover, the GPC samples at t = 90 min and 200 min (not included herein) exhibited Mn = 19.8 kg.mol -1 , Đ = 1.40 and Mn = 25.9 kg.mol -1 , Đ = 1.51 respectively, which demonstrates that the loss of control predominantly occurred during the last hour of the reaction. Generally, at the end of every chain-extension experiment, no major shoulder on the low molar mass side was detected, which indicates a low fraction of unreacted macroinitiator. However, a slight tail can be observed for long elution times (25-29 min) most likely resulting from irreversible terminations.

This study allowed to assess the presence of an alkoxyamine chain end for SG1 NMP-based P(Mystat-GMA) copolymers by 31 P NMR and chain-extensions. Regardless of their composition, P(My-stat-GMA) can thus be used as efficient macroinitiators in order to form versatile and well-defined diblock copolymers.

• Glass transition behavior of My/GMA copolymers.

Glass transition temperatures (Tgs) were measured for the final P(My-stat-GMA) statistical copolymers exhibiting FGMA = 0.11-0.95 (Table 3), by means of differential scanning calorimetry (see

162

Experimental Section for further details, All the DSC traces are given in Supporting Information, Figure S10b). Tg is plotted as a function of FGMA in Figure 6. Each of the statistical copolymers showed a single Tg, which provided a concave glass transition temperature-composition curve. Tg,P(My) = -77.0 o C 35 and Tg,P(GMA) = + 85.0 o C 82 corresponding to the Tg of P(My) and P(GMA) homopolymers have also been added

in Figure 6 to serve as brackets for the copolymer Tgs. As expected, the concave trend reported indicates the non-ideal mixing of My and GMA, with the specific volume for P(My-stat-GMA) being larger than that for ideal mixing. This is likely due to the presence of the pendant flexible C6/C8 groups borne by the My units along the copolymer chain, which may alter the free volume. As it can be seen in Figure 6 for FGMA ≥ 0.66, the presence of a small fraction of My units in the chain reduced considerably P(My-stat-GMA)'s Tg, likely due to the plasticizing nature of P(My). A similar trend was observed with regards to NMP-based P(My-stat-S) copolymers. 35 Several theoretical and empirical approaches have been proposed for estimating the Tg of mixtures and random copolymers from knowledge of the properties of the pure components. [83][84][85][86] Although different in detail, the proposed relationships are all based on the additivity of basic thermosphysical properties. For the My/GMA system studied, they can be expressed by the Wood equation as given below (Equation 6): 86 WMy (Tg-Tg,My) + KWGMA (Tg-Tg,GMA) = 0

where Tg corresponds to the theoretical glass transition temperature of the copolymer containing the weight fractions WMy and WGMA of the two monomers. The Fox relationship is obtained when the parameter K = Tg,My / Tg,GMA. 84 Interestingly, the Fox equation does not predict well the Tg versus composition relationship for the statistical My/GMA copolymers, as shown in Figure 6 (black dotted line) by the overestimation of the copolymer Tgs (+ 19.4 o C on average) when using this simple bulk additive approach. Alternatively, the Gordon-Taylor equation was used, relying still on Equation 6 but with K = ΔβGMA / ΔβMy where ΔβGMA or ΔβMy is the difference between the expansion coefficients of the rubbery and glassy states of P(GMA) or P(My). 83 K = 0.186 ± 0.036 was estimated via the use of a non-linear least-squares fitting of the data and the Gordon-Taylor fit is given in Figure 6 (blue dashed line). K ˂˂ 1 indicates that the volume expansion coefficient for GMA in the rubbery phase is significantly lower than that of My, according to the Gordon-Taylor approach. This appears very consistent since, above the glass transition temperature, it can be assumed that the expansivity of My is much greater than that of GMA, notably due its long pendant side group. On the other hand, stronger intermolecular forces are generated by GMA units, which might reduce its expansivity. To the best of our knowledge, the coefficients of thermal expansions (CTEs) for P(My) and P(GMA) are not available, limiting thus the discussion. However, it is useful to compare the results obtained via the Gordon-Taylor method for My/GMA to another system such as I/MMA. The CTE for the unvulcanized natural rubber (cis-1,4-poly(isoprene)) in the rubbery state (T = 0-20 o C) was estimated to be 6.6.10 -4 K -1 whereas the simulated value of its CTE in the glassy state was lower, 2.0.10 -4 K -1 at T = -73 o C. 87, 88 Furthermore, the CTEs of atactic PMMA (Tg ~ 105 o C) in the rubbery state and in the glassy state are respectively 6.1.10 -4 K -1 (T = 160 o C) and 1.8.10 -4 K -1 (T = 40 o C). 87 Consequently, K = ΔβMMA / ΔβI = 0.93 can be estimated for MMA/I system, indicating in that case a similar expansivity of both units in the rubbery phase. For MMA/I pair, K is much closer to 1, showing a reduced expansivity of I units compared to My units in P(My-stat-GMA copolymers). This may be partly caused by the short methyl side groups borne by PI segments (100 mol% cis-1,4 motif) compared to the relatively longer C6/C8 pendant groups of P(My) segments (47-80 mol% 1,4-motif, stereoselectivity not determined).

The Tg values of the diblock copolymers synthesized via chain-extensions were explored as well.

The DSC traces for P(My-stat-GMA) macroinitiators and the resulting chain-extended products are given in Figure 7. While My-rich macroinitiator My/GMA-78 exhibited a single Tg = -63. macroinitiator (-12.0 o C, Figure 7a). 35 Accordingly, a stronger incompatibility between the two blocks may be apparent, bringing about a much more distinct two-phase microstructure. This could result from the long hydrocarbon pendant groups of My units, increasing significantly the free volume and the steric hindrance around P(My) chains and thus reinforcing the heterogeneity of the two-phase system. Regarding My/GMA-37-S, only the upper Tg was observed at + 67.5 o C (Figure 7a). The disappearance of the parent macroinitiator Tg can be due to the large difference in Mn between the two blocks (8.6 and 26.8 kg.mol -1 for My/GMA-37 and PS respectively, Table 4) combined with the relatively short average chain length of My/GMA-37.

Generally, these DSC measurements suggest the possible microphase separation of the synthesized diblocks. In such instance, the use of a homopolymer as macroinitiator instead of a statistical copolymer and blocks exhibiting a higher degree of polymerization should favor a neater separation of the phases. The Flory-Huggins interaction parameter χ measuring the differences in the strength of pairwise interaction energies between species can be estimated for the My/GMA system using Equation 7(Hildebrand and Scott method): 90,91

χMy-GMA = Vm,GMA (δMy -δGMA) 2 / RT

where Vm,GMA is the molar volume of GMA (136.4 cm 3 .mol -1 ) considered as the ``solvent`` (lower molecular weight), δMy and δGMA are the solubility parameters of My and GMA respectively and R corresponds to the ideal gas constant (8.314 J.mol -1 .K -1 ). δMy = 16.2 MPa 1/2 and δGMA = 19.3 MPa 1/2 were calculated, using the reported Hansen solubility parameters for My and GMA, which gives χMy-GMA = 0.50 at T = 25 o C (see Supporting Information for the calculations). 92, 93 In order to estimate the miscibility between the P(My) and P(GMA) blocks, a mixture of 1 g of P(My) and 1 g of P(GMA) with Mn,P(My) = Mn,P(GMA) = 15 kg.mol -1 can be arbitrarily considered as a basis. The Gibbs free energy of mixing ∆Gm for this binary P(My)/P(GMA) polymer mixture can be estimated using the equation developed by

) 8
where k is the Boltzmann's constant, V is the total volume of the mixture, Vr is the molar of a specific segment, vi is the volume fraction of polymer ''I'', z is the lattice coordination number (usually between 6 and 12) and Nc is the number of chains per volume unit. ∆GM (25 o C, z = 6) = 2.52 J was calculated when assuming that the number of lattice sites per chain is 6 (usually chosen for estimating polymer mixture miscibility). 95 Consequently, ∆GM (25 o C, z = 6) ˃ 0 indicates that P(My)/P(GMA) system with Mn,P(My) = Mn,P(GMA) = 15 kg.mol -1 is not miscible at T = 25 o C, according to the Flory-Huggins theory. The upper limiting Mn,lim number-average molecular weight value allowing the P(My)/P(GMA) system to be miscible at T = 25 o C can be calculated for ∆GM (25 o C, z = 6) = 0 as well. Mn,lim ~ 1.2 kg.mol -1 was estimated, indicating that a homogeneous mixture of 1 g of P(My) and 1 g of P(GMA) with Mn,P(My) = Mn,P(GMA) ≤ 1.2 kg.mol -1 could be achieved at room temperature (see Supporting Information for the calculations). Even though this Flory-Huggins thermodynamic model does not account for a change of volume upon mixing to name but one limitation, it supports theoretically the possible phase separation of My/GMA diblock 166 copolymer. Obviously, direct experimental characterization using techniques such as transmission electron microscopy (TEM) or small-angle X-ray scattering (SAXS) should be done to unequivocally demonstrate whether a microphase separation is apparent. Regardless, in the context of the present study, the nature of the polymerization process (using statistical copolymer segments) influenced the thermal properties, not surprisingly showing some diffuse interfacial effects on the miscibility of the block copolymer segments.

Poly(β-myrcene-block-2-hydroxy-3-morpholinopropyl . methacrylate) amphiphilic diblock copolymer

The well-controlled nitroxide-mediated copolymerization of My and GMA in bulk initiated by NHS-BB allowed the synthesis of epoxide functionalized My-based polymers. Once the polymer is produced, different chemical transformations of the epoxy groups can be performed to obtain a series of polymers with varied properties. In this study, we tried to synthesize an amphiphilic diblock copolymer in a three-step process: (1) NMP of My;

(2) GMA chain-extension from NMP-based P(My) macroinitiator;

(3) Post-polymerization treatment using morpholine to yield poly(β-myrcene-block-2-hydroxy-3- resulting P(My-b-GMA) diblock copolymer had Mn = 23.8 kg.mol -1 , Đ = 1.89 and FMy = 0.58 ( 1 H NMR spectrum in Figure 9a) and the epoxy groups, inert towards the previous polymerization process, were lastly treated with a 6-fold amount of morpholine in refluxing Me-THF for 3 h at 77 o C (complete experimental conditions and results for each step summed up in Supporting Information, Table S2b), relying on the comprehensive study led by Benaglia and coworkers. 96 This ring-opening reaction of the pendant epoxides proceeded quantitatively as shown by the 1 H NMR spectrum of the resulting P(My-b-HMPMA) diblock copolymer in Figure 9b. The final P(My-b-HMPMA) exhibited Mn = 26.5 kg.mol -1 , Đ = 1.91, FMy = 0.55 and a monomodal distribution of the GPC trace (Supporting Information, Figure S11b). The slight increase of Mn after the treatment with morpholine may be explained by the occurrence of branching reactions, as reported previously. 96 It should be added that the hydrophilicity of the P(My) first segment may have been slightly increased due to the possible reaction of morpholine with the succinimidyl ester end-group (initiating group of NHS-BB alkoxyamine) of the diblock. The assignments * and ** refer to olefinic and vinyl protons of 1,2-and 3,4-P(My) contents.

168

The behavior of this morpholine derivatized P(My-b-HMPMA) in water was then explored by dynamic light scattering (DLS) to determine if this diblock copolymer is able to self-assemble under specific conditions (Supporting Information, Figure S12b). The polymer was first dissolved in a minimum amount of THF, which is a good solvent for both P(My) and P(GMA) segments, and then added in a water solution (purified by reverse osmosis), followed by intense dispersion. The suspension was heated at 50 o C for 30 min and the THF was then eliminated by dialysis. A homogeneous clear aqueous solution was obtained. Regarding the samples used, two parameters were changed, [P(My-b-HMPMA)] = 0.4-1.4 mg.ml - 1 and the duration of the stirring (three days to ten days), as indicated in Table 5. Furthermore, the DLS diameter. It was well-reported that the behavior of a polymeric micellar system toward temperature mostly depend on the nature of the hydrophilic block. 97 For P(My-b-HMPMA) diblock copolymer, the decrease of Zave at T ≥ 45 o C may be explained by the much lower salvation of the P(GMA) segments in the corona of the micelles. Four different preparations were performed before measuring the DLS samples, from a relatively high concentration (1.4 mg.ml -1 ) with a short period of stirring (3 days, I, Table 5) to a low concentration (0.4 mg.ml -1 ) combined with a long period of stirring (10 days, IV, Table 5).

Conclusion

The possibility of synthesizing well-tailored My/methacrylate copolymers by NMP has been supported by the satisfactory copolymerization of My and MMA in bulk at 110 o C initiated by NHS-BB.

The linear relationship between Mn and X combined with Đ as low as 1.15-1.32 shown by this polymerization prompted us to study the NMP of My with GMA, a highly versatile monomer bearing an epoxy group. My/GMA NMP with fMy,0 = 0.10-0.90 was performed in bulk at 120 o C initiated by NHS-BB. rMy = 0.80 ± 0.31 and rGMA = 0.71 ± 0.15 were determined by the FR approach whereas the KT method gave rMy = 0.48 ± 0.12 and rGMA = 0.53 ± 0.18. The statistical nature of this copolymerization was confirmed using a NLLS procedure yielding rMy = 0.49 ± 0.13 and rGMA = 0.50 ± 0.13. For fMy,0 = 0.60-0.90, welldefined P(My-stat-GMA) copolymers were achieved with Mn approximately 81% of Mn,theo and Đ = 1.26- ). The substitution of the PS blocks by harder blocks, exhibiting a higher Tg, can significantly increase the tensile strength of the material as well as its upper service temperature.

Accordingly, it was decided herein to replace the PS outer segments by more rigid poly(isobornyl methacrylate) P(IBOMA) segments, displaying a Tg ~ 190 o C.

Abstract

Introduction

The community of polymer researchers has long desired to elaborate polymerization techniques that couple the control of microstructure offered by living ionic polymerization 1-3 with the simplicity of compositions, topologies and functionalities. Among them, nitroxide-mediated polymerization (NMP) 7,8,10 is conceptually and practically the simplest of these systems. Based on a reversible termination mechanism between the propagating free radical and the nitroxide used as a controlling agent, NMP establishes a thermal dynamic equilibration between dormant macro-alkoxyamines capped by the nitroxide moiety and actively propagating macroradicals 11 . At a minimum, only an alkoxyamine, available from commercial sources and easily functionalizable, and the monomer(s) are necessary to implement a NMP process initiated via a mono-component pathway (i.e. unimolecular initiator). NMP relies on the elimination of oxygen prior to polymerization as the main experimental constraint 6, 12 .

The opportunity to prepare high-performance polymers by taking advantage of NMP's simplicity appears to be highly desirable from both practical and economical points of view. The synthesis of polymers having desirable mechanical properties such as a high tensile strength, a remarkable extensibility or a combination of both seems achievable by NMP. In particular, the development of styrenic-diene block copolymers, generally composed of a dispersed poly(styrene) PS or PS derivative phase in a continuous elastomer domain, 13,14 looks specifically suited to the NMP process. Indeed, the NMP of conjugated dienes such as butadiene (B) 15, 16 and especially isoprene (I) 17 was explored from the onset of the development of NMP, using 2,2,6,6-tetramethylpiperidinyl-1-oxy (TEMPO) 7 , the original stable free nitroxide. Welltailored homopolymers and copolymers based on either B or I were produced as well In the present study, we report first the copolymerization of My and IBOMA in bulk at 100 o C, initiated by the succinimidyl ester-terminated NHS-BlocBuilder alkoxyamine. The active feature of these NMP-based My/IBOMA copolymers, ideally capped by a SG1 nitroxide moiety, was then assessed quantitatively via 31 P NMR spectroscopy and qualitatively by chain-extensions with a fresh batch of IBOMA and/or My. After guaranteeing the efficient control of the SG1-mediated copolymerization of My with IBOMA, the second part of this investigation was focused on the preparation of P(My)-(SG1)2 macroinitiators using telechelic poly(ethylene-stat-butylene) initiator terminated with SG1 nitroxide groups 42 and their subsequent chain-extension with IBOMA-rich/My or S mixtures. The phase behavior, the thermal behavior, notably along with viscoelastic properties, and the stress-strain properties of the resulting triblock copolymers were elucidated to emphasize the possibility to produce valuable diene-based block copolymers from largely sustainable feedstocks, in a simple way by NMP.

Results and discussion

Two main parts are introduced thereafter and consist of the synthesis by nitroxide-mediated polymerization (NMP), via the SG1 nitroxide rate moderator, and the characterization of β-myrcene-rich (My-rich) and isobornyl methacrylate-rich (IBOMA-rich) copolymers. First, the nature of My/IBOMA copolymerization initiated by NHS-BlocBuilder (NHS-BB) was examined as well as its kinetics, its level of control and the ability of the resulting copolymers to re-initiate a second batch of monomer(s). Glass Figure 1 depicts the various polymerizations done in this study.

β-Myrcene / isobornyl methacrylate copolymerization by NMP

• Preliminary polymerization temperature study. Accordingly, these encouraging results prompted us to study two My/IBOMA copolymerizations in bulk using NHS-BB initiator at 120 o C (experiments My/IBOMA-50-T120 with fMy,0 = 0.50 and My/IBOMA-70-T120 with fMy,0 = 0.71, Table 1). Although a linear increase of Mn,MHS with the overall conversion X was observed in both cases, relatively broad molecular weight distributions were obtained throughout the reactions (Đ = 1.45-1.71 for X = 26-94%, Figure S1c, Supporting Information). This may be due to the likely high individual equilibrium constant of IBOMA, as typically reported for methacrylates • Effect of the feed composition on copolymer composition.

It is first of interest to determine the monomer sequences of P(My-co-IBOMA) copolymers via the determination of the copolymer-monomer feed composition relationship. To that end, My/IBOMA NMP, with different feed ratios (fMy,0 = 0.10-0.90), was carried out at 100 o C in bulk, initiated by NHS-BB and targeting Mn,theo = 30 kg.mol -1 , as indicated in Table 1. The copolymer composition was determined through assignments from experiment My/IBOMA-50) for relatively low overall conversion samples (X ≤ 27.5%, Table S1c in Supporting Information gives the samples used for extracting the reactivity ratios). The terminal copolymerization model, where only the terminal monomer residue is considered, was applied 57 .

The equations used for the calculation of the reactivity ratios can be found in the Supporting Information. By using the Fineman-Ross (FR) method 58 (Figure 2a), the reactivity ratios were determined: rMy = 2.16 ± 0.34 and rIBOMA = 0.07 ± 0.04 (rMy rIBOMA = 0.15 ± 0.19). Similar values were calculated via the Kelen-Tudos (KT) approach 59 (Figure 2b): rMy = 1.90 ± 0.18 and rIBOMA = 0.02 ± 0.21 (rMy rIBOMA = 0.04 ± 0.15). The errors associated were derived from the standard errors of the slopes from FR and KT plots.

Even though the KT equation refines the linearization method by introducing an arbitrary positive constant to spread the data more evenly, it remains a rearrangement of the copolymer composition equation and a non-linear least-squares (NLLS) fit to the Mayo-Lewis equation 60 is likely one of the sounder methods to determine rMy and rIBOMA. Consequently, the reactivity ratios determined by the KT method were used as the initial guesses for a NLLS fitting of the data (Supporting Information). At 95% confidence level and with a regression coefficient R 2 = 0.91, the statistical fit to the data yielded reactivity ratios rMy = 2.07 ± 0.58 and rIBOMA = 0.05 ± 0.08. These latter results are similar from those obtained by FR and KT approaches, confirming the general trends with respect to My/IBOMA NMP at 100 o C in bulk and initiated by NHS-BB: rMy > 1.50, rIBOMA < 0.30 and rMy rIBOMA < 0.50. The large difference in reactivity between My and IBOMA can be first noted. Whatever the terminal unit of the propagating species (---My • or ---IBOMA • ), the macro-radical showed a large preference for My. The insertion of IBOMA units at low conversions, either by self-propagation or cross-propagation, is not favored: IBOMA-terminated macro-radical has a strong tendency to alternate. Accordingly, it can be assumed that My/IBOMA copolymerization under these conditions produces copolymers that possess a gradient in composition consisting of My enrichment at the start of the polymer chain followed by the incremental IBOMA enrichment at the end of the chain due to My depletion. However, the gradient nature of this copolymerization was moderate as indicated by the presence of methacrylate units at the start of the copolymer chain synthesized (X = 3-12% and FIBOMA = 13-18% at t = 0-60 min for experiment My/IBOMA-50 for instance). Even though samples exhibiting relatively low conversions were used, it can be assumed that the composition drift in the monomer feed was nonnegligible, especially taking the difference between rMy and rIBOMA into account and was the main source of error by relying on the instantaneous copolymer composition model 61 . To the best of our knowledge, no reactivity ratios were reported in the literature for diene/IBOMA systems, limiting thus the discussion.

Nonetheless, for the My/GMA pair copolymerized under similar conditions (120 o C in bulk initiated by NHS-BB) 49 , statistical copolymers were synthesized with rMy = 0.48-0.80 and rGMA = 0.50-0.71. Such a difference compared to the My/IBOMA system highlights the likely influence of the bulky and rigid bicyclic isobornyl groups over the reactivity of the methacrylate monomer (steric hindrance, hydrophobicity, reduction of polar interactions).

ID (a) [NHS-BB]0 (M)

[My]0 (M) Interestingly, while minor variations can be observed for My-rich copolymers (experiment My/IBOMA-80

with FMy ≥ 0.66 during the whole reaction), as FIBOMA increases, Mn,MHS becomes greater than Mn,GPC (My/IBOMA-30 with FIBOMA ≥ 0.44 for X = 25-47%). It can be thereby hypothesized that an underestimation of IBOMA-rich copolymer Mn values was done when calibrating with PMMA due to the specific hydrodynamic volume of this copolymer coil in THF.

Before exploring the kinetics of My/IBOMA copolymerization at various initial feed compositions, attention can be paid to the homopolymerization of My and IBOMA by NMP under the same conditions (experiments My/IBOMA-100 and My/IBOMA-0, Table 1). My NMP at 100 o C in bulk mediated by NHS-BB showed a satisfactory control with Mn,MHS increasing linearly with conversion and Đ ≤ 1.36 (Figure S3c, Supporting Information), resulting in 1,4-rich P(My) (Table 2), even though non-negligible deviations of Mn,MHS values from theoretical ones were measured. The NMP of IBOMA without neither co-monomer nor additional free nitroxide was also attempted under these conditions. Controlled polymerization of methacrylates is still challenging for NMP because of their very high equilibrium constant K resulting from slow recombination of nitroxides and sterically-hindered poly(methacrylate) radicals 63, 64 . This long-time obstacle was confirmed once again by the inability to implement the SG1-based NMP of IBOMA in bulk at 100 o C in a controlled manner: the reaction medium polymerized spontaneously, giving P(IBOMA) chains exhibiting Mn,MHS = 27.1 kg.mol -1 and Đ = 1.75 (My/IBOMA-0, Table 2). Due to the early high viscosity of the medium after only a few minutes, it was decided to repeat the experiment in 50 wt% toluene (experiment My/IBOMA-0-Tol, Table 1) in vain. The polymerization stagnated at X ~ 10% with Mn,MHS = 11.4-12.3 kg.mol -1 and Đ = 1.63-1.76 (Figure S3c, Supporting Information), illustrating presumably the disproportionation between P(IBOMA) macro-radicals and SG1 nitroxide.

In order to produce IBOMA-rich polymers by SG1-mediated NMP, the controlling co-monomer approach, developed by Charleux and co-workers 39-41 , was tried where the addition of a small concentration of a co-monomer exhibiting a much lower activation-deactivation equilibrium constant K with respect to the methacrylate enhances the control of the polymerization. It was herein decided to use My as a potential controlling co-monomer although its individual K has not been reported. 10-20 mol% of My had a marked impact on the outcome of My/GMA NMP initiated by NHS-BB in bulk. It gave a linear evolution of Mn with monomer conversion together with Đ ≤ 1.52 up to almost quantitative overall conversion 49 . Previously, 10-20 mol% of isoprene, another conjugated 1,3-diene, allowed the well-controlled bulk NMP of tert-butyl acrylate and GMA 18 . In the present study, the addition of 10 mol% of My (experiment My/IBOMA-10, Table 1) into the IBOMA feed undoubtedly slowed down the reaction even though the features of a controlled system were only partially observed due to the upward curvature of the semi-logarithmic plot and Đ = 1.51-1.60 at X = 4-42% (Figures 3a and3b). Expectations were fully met by performing the NMP of IBOMA in the presence of 20 mol% of My as a co-monomer (My/IBOMA-20, Table 1) with Đ ≤ 1.39, Mn increasing linearly with X and a quasi-linearity of ln((1-X) -1 ) versus t during the reaction with R 2 = 0.92 (Figures 3a, 3b and3c). It can be assumed that the addition of My in the IBOMA-rich feed induced a significant decrease of the average <K> and therefore enhanced the control of the polymerization.

For fMy,0 = 0.30-0.90, linear Mn,MHS increase with overall X and narrow molecular weight distributions (Đ ≤ 1.41) of the P(My-grad-IBOMA)s were observed (Figures 3b and3c). These results presumably illustrate the effective mediation of the stable SG1 free nitroxide, capping the propagating chains and thus shifting the equilibrium toward deactivation 12 . Although termination and transfer side reactions were still present, they were likely minimized due to the high number of reversibly terminated chains (dormant P(My-grad-IBOMA)-(SG1) chains). Nonetheless, as fMy,0 increased, the deviations of Mn,MHS from Mn,theo became more significant. Even though the principle of universal calibration was applied, the MHS coefficients did not match exactly the GPC conditions used (THF, The dashed line indicates the theoretical Mn versus overall conversion based on the monomer to initiator ratio (Mn,theo ~ 30 kg.mol -1 at X = 100% for every experiment). All experimental ID and characterization of experiments are listed in Table 1 and Table 2. The same legend at the bottom right of the figure is used for each of the three plots.

pair. As reported previously 28 , a minor peak, corresponding to high Mn polymers, was detectable mostly at the commencement of the copolymerization (20-25 min, Figure 4), presumably due to My auto-initiation 66 .

Attention can also be paid to the kinetic plots of ln((1-X) -1 ) versus reaction time t (Figure 3a). For fMy,0 = 0.10-0.81, all copolymerizations obeyed first order kinetics as described by Equation 3 and generated good fits to linear kinetic plots (squared linear regression coefficient R 2 ≥ 0.94 for the whole experiments).

ln

([M]0/[M]t) = ln([M]0/([M]0(1-X))) = ln((1-X) -1 ) = <kp>[P•]time (3)
In Equation 3 IBOMA and My homopolymerizations performed herein (My/IBOMA-0 and My/IBOMA-100, Table 1) confirmed the kinetic singularity of these two monomers polymerized by NMP. While a quasi-spontaneous polymerization was performed when an IBOMA solution was heated at 100 o C in the presence of NHS-BB, My polymerized slowly under identical experimental conditions and this reaction exhibited kpK = (1.4 ± 0.4) 10 -5 s -1 (Table 2) consistent with the kinetic values reported as regards the NMP of dienes 28, 26 . Very interestingly, <kp><K> was on average (2.03 ± 0.94) 10 -5 s -1 for the NMP of My with IBOMA with fIBOMA,0 = 0.10-0.90 (Table 2), which was substantially close to the <kp><K> value of the NMP of My. Accordingly, it can be assumed that the copolymerization kinetics was largely governed by the My kinetics, even when the initial mixture was rich in IBOMA. This behavior is illustrated in Figure 3a where all curves almost overlap regardless fIBOMA,0.

The molar fractions of 1,4-, 1,2-and 3,4-P(My) additions in the My/IBOMA copolymers were determined by 1 H NMR (Figure S4c in Supporting Information for the assignments). Interestingly, while the experiment My/IBOMA-100 led to a P(My) homopolymer rich in 1,4-content (86.6 mol%, Table 2), the fraction of 1,2-motif in the copolymers markedly increased when IBOMA-rich feeds were used (~ 30 mol% of 1,2-content for fIBOMA,0 = 0.41-0.80). A very similar trend was observed for the systems My/GMA and My/tert-butyl acrylate copolymerized under the same conditions at 120 o C and 115 o C respectively 49 .

Therefore, it must be caused by the presence of a methacrylate radical as a terminal unit, which might have altered the steric and/or the electronic environment of the active end of the propagating species. Relying on previous studies The glass transition behavior of P(My-grad-IBOMA) copolymer with cumulative IBOMA mole fractions ranging from 0.10 to 0.96 was also characterized using DSC heat curves (Figure S6c, Supporting Information). Tg versus IBOMA composition plot is given in Figure 5. A trend curve was included, which was obtained by fitting the Gordon-Taylor equation 70 (Equation 6) to the experimental data. where wMy and wIBOMA are the weight fractions and Tg,P(My) and Tg,P(IBOMA) the glass transition temperatures of the respective homopolymers. Tg corresponds to the glass transition temperature of the gradient copolymer and K = ΔβIBOMA / ΔβMy with ΔβIBOMA or ΔβMy being the difference between the expansion coefficients of the rubbery and glassy states of P(IBOMA) or P(My). This theory is based on two basic assumptions: the volume additivity (ideal volume of mixing) and a linear change in volume with temperature. With 95% confidence bounds, K = 0.322 ± 0.090 (R 2 = 0.94) was determined via the use of a non-linear least-squares fitting of the data and indicated the larger expansivity of My units in the rubbery state than that of IBOMA units. A similar trend was reported for the system My/GMA (K = 0.186 ± 0.036) 49 , highlighting the plasticizing nature of P(My). The Gordon-Taylor fitting approach yielded a good representation of the experimental Tg data (Figure 5). A slight concave Tg versus IBOMA composition curve was obtained, pointing out deviations from an ideal mixing. As expected, P(My-grad-IBOMA)'s Tg Thermogravimetric analysis (TGA, experimental section) has been applied to decomposition study on the final dry samples taken from experiments My/IBOMA-20 and My/IBOMA-80, exhibiting 51 and 81 mol% of My units respectively (Table 2). Figure 6 represents the thermal decomposition curve of both gradient copolymers. Interestingly, the synthesized polymers exhibited markedly two step degradation pattern. A first weight loss was observed from 145 to 305 o C with a Tdec,max of 242 o C (temperature at which the weight loss is most apparent, maximum of the derivative curve) for the sample richer in My (My/IBOMA-80, Figure 6b) and from 205 to 315 o C with Tdec,max = 307 o C for My/IBOMA-20 sample (Figure 6a). This early degradation peak was likely due to the loss of the side chains by cleavage of Oisobornyl bond, releasing camphene as a main product, as reported by Matsumoto and coworkers 37 . This first weight loss was proportional to the IBOMA fraction possessed by the samples: while about 40% of mass loss was measured for My/IBOMA-20 with FIBOMA = 0.49, only about 22 wt% of My/IBOMA-80, exhibiting FIBOMA = 0.19, was lost (Figure 6). This observation supports the contribution of the IBOMA units to the premature P(My-grad-IBOMA) thermal decomposition. The second degradation step is similar for both samples, starting at 305-315 o C with Tdec,max = 370-380 o C and decomposing totally the copolymers at about 450 o C (Figure 6), which could be mostly caused by the scission reactions of the main chain (depolymerization) resulting in the formation of monomer and oligomer radicals. We previously reported the decomposition behavior of SG1-mediated P(My), from 290 to 485 o C (Tdec,max = 385 o C) 28 , matching closely the second degradation step of the two P(My-grad-IBOMA)s characterized.

• Active feature of NMP-based P(My-grad-IBOMA).

Ideally, a majority of the P(My-grad-IBOMA) chain ends synthesized by NMP should possess the SG1 moiety. Indeed, the recombination of the propagating macro-radical with the controlling nitroxide should lead to a very low [P•], limiting the self-termination reaction. Afterwards, the final SG1-terminated P(My-grad-IBOMA) would be able to be re-activated via the thermal homolytic decomposition of this dormant species into the initiating macro-radical and the SG1 persistent radical. Two strategies were herein adopted to know if P(My-grad-IBOMA) chains were effectively capped by a SG1 group: 31 P NMR analysis (see experimental section) allowing the detection of the phosphorus atom of the nitroxide group 71,72 ;

IBOMA and/or My chain-extension from P(My-grad-IBOMA) macro-initiator by the same nitroxide mediated mechanism. Two samples were studied: My/IBOMA-82 and My/IBOMA-44 exhibiting respectively 82 and 44 mol% of My (Table 3A). 31 P NMR was first performed to evaluate the living fraction (LF) of these two gradient copolymers initiated by NHS-BB alkoxyamine (Figure S5c in Supporting Information for the spectra). The proportion of chains containing a SG1 end group was 74 ± 9 mol% and 69 ± 3 mol% for My/IBOMA-44 and My/IBOMA-82 copolymers, respectively (Table 3A). It is interesting to note that about two thirds of the e) Predicted Mn,MHS,2 of the whole chain-extended diblock copolymer (MI block + second block added) at X2, measured experimentally, and calculated as follows: Mn,theo,X2 = (X2/100)(Mn,theo,2-Mn,MHS,1)+Mn,MHS,1 (= Predicted Mn of the second block added at X2 + experimental Mn,MHS of MI).

My-rich

distribution (D1 = 1.51, Table 3A). A similar LF was measured for the IBOMA-rich gradient copolymer exhibiting a lower overall conversion (X1 = 39.1%). A moderate conversion was targeted for the synthesis of this macro-initiator since an IBOMA-rich (fIBOMA,0 = 0.70) feed was polymerized, where disproportionation side reactions likely occur. It can be assumed that the synthesis of a high conversion P(My-grad-IBOMA) rich in IBOMA would have a relatively low LF, caused by the occurrence of terminations by H-transfer notably.

To further examine the capacity of these samples to re-initiate and produce diblock copolymers, two chain-extension experiments were done using My/IBOMA-44 and My/IBOMA-82 as macroinitiators (Table 3A). The reaction conditions and results of the chain-extensions are given in Tables 3B and3C respectively. My/IBOMA-44 macroinitiator (FIBOMA,1 = 0.56, Mn,MHS,1 = 8.8 kg.mol -1 ) was cleanly chainextended at 115 o C in toluene with My to obtain a P[(My-grad-IBOMA)-b-My] diblock copolymer (FIBOMA,2 = 0.07, Mn,MHS,2 = 30.1 kg.mol -1 , Table 3C) containing predominantly My units. The dispersity of this chainextended product (D2 = 1.65) was higher than that of the parent macro-initiator (Đ1 = 1.32), which suggests that termination occurred during the reaction and/or some macroinitiator was not initiated. This latter assumption can be supported by the previous quantitative spectroscopic analysis, indicating that about 25 mol% of My/IBOMA-44 was not terminated by SG1 (Table 3A), as well as the GPC traces of the chainextension experiment (Figure 7b) showing that the tails of the final chain-extended copolymer overlapped to some degree with the associated macroinitiator trace. The asymmetrical last GPC trace at 390 min may also indicate the occurrence of irreversible terminations throughout the chain-extension, which generate ''dead'' species having lower Mn (~ 15-25 kg.mol -1 ). Generally, this successful My chain-extension from IBOMA-rich P(My-grad-IBOMA) copolymer prompted us to investigate the reversed situation where a

My-rich macroinitiator can be re-activated to add an IBOMA-rich segment. With a fresh IBOMA/My mixture (93 mol% of purified IBOMA), My/IBOMA-82 (FMy,1 = 0.82, Mn,MHS,1 = 13.6 kg.mol -1 , Đ1 = 1.51, Table 3A) was extended at 105 o C in 50 wt% toluene (Table 3B). Due to the beneficial kinetic effects observed of the addition of 10-20 mol% of My over the control of the SG1-based NMP of IBOMA (Figure 3), 7 mol% of the acyclic monoterperne was initially introduced in the feed. After 210 min, the resulting chain-extended diblock exhibited Mn,MHS,2 = 23.2 kg.mol -1 , Đ2 = 1.61 and the overall My molar composition was decreased from 82 to 49 mol% (Table 3C). Similarly to the first chain-extension, the GPC trace of the chain-extended product retained generally its mono-modal nature and had a clear shift to the left (Figure 7a), indicating a high-level of chain-end fidelity of the original macroinitiator.

These satisfactory results naturally paved the way to the preparation by NMP of well-defined 1,3diene-based triblock copolymers composed of a flexible and bio-sourced P(My) mid-segment 32 and two rigid P(IBOMA) outer segments, the latter coming from partially bio-based raw material source 73,74 .

Isobornyl methacrylateβ-myrcene -isobornyl methacrylate . triblock copolymers by NMP

The synthesis by SG1-mediated NMP of triblock copolymers consisting of a soft P(My) middle block and two hard blocks, containing mostly IBOMA units, was aimed in this second part. A two-step approach was applied for preparing these triblock copolymers: • Optimized synthesis of P(My)-(SG1)2.

First of all, the preparation of a P(My) segment exhibiting a relatively high Mn and able subsequently to re-initiate a fresh batch of monomer(s) at each chain end was desired. In order to have sufficient mechanical properties, the targeted final triblock must have an entangled elastomeric phase, the physical chain entanglements must be contributing mainly to the network density and to the stress at small elongations 76 . P(My) entanglement molecular weight (Me,P(My) = 22-31 kg.mol -1 ) 77 is significantly higher than that of comparable poly(diene)s such as poly(isoprene) PI (Me,PI = 4-6 kg.mol -1 ) 77 or poly(butadiene) PB (Me,PB = 2-4 kg.mol -1 ) 77 . Accordingly, Mn > 40 kg.mol -1 for P(My)-(SG1)2 was desired to ensure a minimal entanglement density.

PEB-(SG1)2 dialkoxyamine macroinitiator (Mn,MHS = 5.7 kg.mol -1 , Đ = 1.17) was first produced, according to the method reported by Lessard et al. 42 . 1 H NMR of the final product revealed that about 84 mol% of the chain ends were capped by a SG1 unit. It can thus be deduced that a minor fraction of the synthesized macroinitiator consisted of monofunctional alkoxyamine terminated by SG1 only at one chain end and "dead" PEB chains. To know the optimal conditions for making high average chain length P(My) in a controlled manner, NMP of My initiated by various concentrations of PEB-(SG1)2 at 120 o C in bulk was then studied. Four different theoretical Mn were thus targeted at quantitative conversion, namely 79.1, 105.0, 137.2 and 174.0 kg.mol -1 (Table S2c, Supporting Information), and kinetic studies were led (Figure S7c, Supporting Information) to assess the level of control. As Mn,theo increased, higher dispersity values and greater deviations of Mn,MHS from the theoretical molar masses were observed (Figures S7b andS7c), confirming likely the greater presence of irreversible terminations and chain transfer reactions to My and P(My) with [My]0. Nonetheless, lowering the dialkoxyamine concentration resulted in the synthesis of higher Mn,MHS P(My)s. Particularly, when targeting Mn,theo = 137.2 kg.mol -1 at XMy = 100% (experiment My-137), a polyterpene exhibiting Mn,MHS = 49.5 kg.mol -1 and Đ = 1.52 was obtained at 51% conversion (Figures S7b andS7c), which consisted of the most satisfactory balance between a relatively high Mn and a moderate dispersity. It was then attempted to produce longer P(My) chains with a relatively narrow molecular weight distribution by using a low concentration of the dialkoxyamine initiator [PEB-(SG1)2] (Mn,theo = 169-173 kg.mol -1 at X = 1.0) combined with additional free SG1 nitroxide (9-18 mol% with respect to the initiator, experiments My-173SG1,9 and My-169SG1,18, Table S2). An improved control of the NMP when using a slight excess of this potent nitroxide could have been expected since the dormant state should be favored (lower concentration of propagating macro-radicals, reduction of the overall polymerization rate) 78-80 . Despite apparent rate constants somewhat lower as [SG1 • ]0 increased (Figure S8a), similar Đ versus X and Mn,MHS versus X trends were determined (Figures S8b andS8c) indicating the ineffectiveness of adding SG1 free radical to the reaction mixture. This observation echoes a previous study showing that the addition of 4.9-11.7 mol% of the SG1 mediator with respect to the initiator did not enhance the NMP of My at 120 o C in bulk initiated by NHS-BB 28 . A last experiment in 50 wt% toluene, without extra SG1, was led, predicting Mn,theo ~ 170 kg.mol -1 again (experiment My-170Tol, Table S2). The use of a solvent was thought to help, as seen by the increase in viscosity at XMy > 30%, when performing this reaction in bulk, which might account for the occurrence of diffusional limitations 81 , bringing about a loss of control.

Interestingly, while the polymerization in toluene was slower (Figure S8a), similar molecular characteristics of the growing P(My) were obtained compared to the previous mass polymerizations (Figures S8b andS8c).

Higher Đ values were even observed (Figure S8b), which may be assigned to chain-transfer side reactions to toluene 82, 83 .

Consequently, the P(My) segment was prepared via the PEB-(SG1)2-initiated NMP of My at 120 o C in bulk for 280-400 min (XMy ~ 40-50%) targeting Mn,theo ~ 140 kg.mol -1 and without any additional free nitroxide. Two 1,4-rich P(My)s (experiments My-35 and My-52, Table 4A), exhibiting Mn,MHS = 35-52 kg.mol -1 and acceptable Đ = 1.45-1.53, were synthesized this way. Attention was now shifted to the capacity of these P(My) segments, ideally capped by two SG1 units, to form triblock copolymer with a mixture of IBOMA containing a minor fraction of co-monomer.

• IBOMA/Co chain-extension from P(My)-(SG1)2.

IBOMA/My chain-extension (7.8 mol% of My used as a co-monomer) from My-35 macroinitiator was first performed at 115 o C in toluene for 90 min (Table 4B). The shifts in the GPC traces, seen in Figure S9c (Supporting Information), clearly show the shift in elution time demonstrating the increase in Mn with time. The chain-extended product My-35-IBOMA/My had an increase in Mn from 35.2 to 51.3 kg.mol -1 with a broadening of the molecular weight distribution going from 1.45 to 1.91. The monomodal nature of the GPC chromatograms confirmed that the My-35 chain ends were reversibly terminated in majority. Based on 1 H NMR analysis, the final triblock copolymer exhibited an overall FIBOMA = 0.26 (Table 4C). More specifically, each chain-extended block was composed of a large fraction of IBOMA units (91 mol%) and a minority of My (9 mol%).

Afterwards, a second triblock copolymer, with a higher overall Mn as well as harder outer hard blocks, was targeted. To that end, My-52 macroinitiator (Mn,MHS,My = 51.7 kg.mol -1 , Đ = 1.53, Table 4A)

was reacted with a IBOMA/S mixture (8.7 mol% of S used as a controlling co-monomer, as previously reported 23,39,84 ) for 3 h in 50 wt% toluene at 115 o C (Table 4B). The resulting product My-52-IBOMA/S had Mn,MHS,CE = 78.1 kg.mol -1 and Đ = 2.41. Likewise, the final chain-extended product shifted to lower elution time compared to the macroinitiator (Figure S9b). In order to increase My-52-IBOMA/S Mn, a fractional precipitation using the non-solvent addition method was implemented, relying on the benzene (solvent) / 198 methanol (non-solvent) pair. Eventually, the fractionated triblock copolymer had Mn,MHS = 94.7 kg.mol -1 and Đ = 2.23 (Table 4C), reflecting the removal of a fraction of relatively short chains. The rigid segments were constituted of about 88 mol% of methacrylate units and 12 mol% of S, as determined by 1 H NMR. A high Đ = 1.91-2.23 was measured for these final triblock copolymers, which may be mainly caused by the difficulty to control the polymerization of a methacrylic ester-rich feed (fIBOMA,0 ≥ 0.91), although a small amount of My or S were added to possibly decrease the average activation-deactivation equilibrium constant <K>. The inactivity of a fraction of PEB-(SG1)2 and P(My)-(SG1)2 macroinitiators may also contribute to the broadening of the molecular weight distribution. Besides, the very likely presence of PEB-SG1 monofunctional chains may have resulted in the synthesis of short diblock copolymer chains, assuming a similar activity between PEB-SG1 and PEB-(SG1)2. All these hypotheses would lead to a nonnegligible population of shorter than expected chains, as observed in Figure S9c polymers is about 15 o C higher than that of 1,4-rich P(My) homopolymer (Tg,P(My) = -77 o C, Figure 5), which might be due to some partial compatibility of P(My) and IBOMA-rich regions. It is of interest to note that the upper Tg of the synthesized triblocks is similar to that of P(IBOMA) homopolymer (Tg,P(IBOMA) = 171-177 o C, Figure 5) even though the outer hard blocks contained a non-negligible fraction of S (12 mol%) or My (9 mol%) units. A decrease of the upper Tg would have been expected due to the presence of softer units in the rigid segments (IBOMA/My system illustrated in Figure 5). This unaffected upper Tg can result from the extended blocks composition. For the IBOMA/My chain-extension, it can be assumed that the majority of the My units reacted at the early stages of the polymerization, considering the marked preference of the ---My • macro-radical for My (Figure 2). This way, a My-rich segment including the middle block and the beginning of the two extended blocks may have been formed as well as two quasi neat P(IBOMA) outer blocks. The same hypothesis could be made for the IBOMA/S chain-extension, taking the higher reactivity of S for this system into account (rS = 0.57, rIBOMA = 0.20 reported with benzyl peroxide at 90 o C) 85 .

A. P(My

These DSC results suggest a two-phase system, due to the immiscibility between P(My) and IBOMA-rich segments. To further examine immiscibility, the solubility parameters δ were compared to estimate the tendency towards micro-phase separation. Using the group contribution method developed by Hoftyzer and Van Krevelen 86 , theoretical solubility parameters for P(IBOMA) and 1,4-P(My) (favored regioselectivity in our case) were determined (calculations can be found in the Supporting Information). relatively high immiscibility between P(My) and P(IBOMA) segments and would suggest a strong bulk phase separation of P(My-b-IBOMA) block copolymer. Atomic force microscopy (AFM) was applied to the study of the surface morphology of My-35-IBOMA/My prepared by spin-coating a 3-4 wt% solution in chloroform. Figure 9 shows a 2 μm x 2 μm scan of the polymeric surface (two other AFM phase images of My-35-IBOMA/My in different regions are given in Supporting Information, Figure S10c). The nature of the light and dark domains was determined by quantifying the relative contribution of the light component (ImageJ software application, binary mode). This image contained 23-24% of light areas which is consistent with the molar fraction of P(IBOMA) in the characterized triblock (FIBOMA = 0.28, Table 4). The light areas can thus be attributed to P(IBOMA) and the dark ones to the rubbery P(My) midblock material. The phase separation of the two component blocks appeared clear with the embedding of glassy P(IBOMA) aggregates (disperse phase) in the soft P(My) (continuous phase). A spherical or cylindrical morphology might be obtained. An in-depth analysis, using for instance small-angle X-ray scattering (SAXS), would be required to ascertain the block copolymer morphology. It is of importance to note that the My-35-IBOMA/My selfassembly was disordered with a random location of the rigid aggregates, which exhibited various sizes (27% of large aggregates exhibiting a surface area of 2650 ± 740 nm 2 and 73% of smaller aggregates having a surface area of 930 ± 480 nm 2 , quantification via Nanoscope Analysis software). By assuming that cylinders were observed (FIBOMA = 0.28, possibly elevated for spheres formation), the radius of the cylinders R can be estimated according to the following formula (7):

R = 1.0αKM 1/2 ( 7
)
where α is the chain expansion parameter (α = 1.25 is generally assumed), K is the experimental constant related to the unperturbed root-mean-square end-to-end distance to the molecular weight (KP(IBOMA) ≈ KP(MMA) = 0.565 Å was hypothesized) and M = Mn,P(IBOMA) ≈ 16 kg.mol -1 in this case 87, 88 . Accordingly, R ≈ 8.9 nm was estimated, suggesting that the surface area of the cylinders at the top/bottom (case where cylinders are perpendicular to AFM image, comparable to spheres) was equaled to about 250 nm 2 . This may suggest that the relatively small aggregates seen on the AFM image (average surface area of 930 nm 2 )

might not correspond to spheres but more likely to disoriented cylinders. For cylinders parallel to the AFM image having a mean length of 50 nm, a surface area of 2800 nm 2 was calculated, which is comparable to the size of the large aggregates observed experimentally. Several factors may explain the asymmetrical domain organization. First, the presence of My units in the P(IBOMA) hard blocks resulted presumably in attractive interactions at the soft/hard interfaces, increasing the miscibility between the domains. The relatively broad molecular weight distribution of the sample (Đ = 1.91) may also lead to a reduction of the segregation strength (decrease in χN, where χ is the Flory-Huggins interaction parameter and N is the 202 copolymer degree of polymerization 89 ). Lastly, the experimental approach can be discussed. The analyzed samples were not annealed prior to the AFM characterization. Indeed, due do the early thermal degradation of IBOMA-based polymers (Figure 6), it was decided to not heat the samples at T > Tg,P(IBOMA) for a few hours or days. It may be argued that a stronger phase separation would have occurred after annealing, allowing a full relaxation of the P(IBOMA) chains.

Figure 8b shows the temperature dependence of the shear storage and loss moduli (G' and G'' respectively) as well as the damping factor (tan δ = G''/G') from room temperature to 220 o C for the triiblock copolymer My-52-IBOMA/S (Table 4C). The significant decrease of G' from about 13 to 5 MPa with temperature until 70 o C can be first noted. A quasi constant G' could have been expected in this temperature range due to the glassy state of the IBOMA units embedded in the soft rubbery P(My) phase.

The presence of 12 mol% of S in the extended blocks can partly explain this behavior. Compared to P(IBOMA), The solubility parameter of PS, δPS = 18.4 MPa 1/2 (Hoftyzer and Van Krevelen approach), is closer to that of P(My), which may result in the partial homogeneity of the triblock copolymer. Moreover, short PS sequences softening a relatively low temperature were likely contained by the outer blocks. It was notably reported 90 by DSC that Tg,PS = 32-69 o C for low Mn,PS = 0.9-3.0 kg.mol Lastly, the mechanical properties at room temperature of My-52-IBOMA/S were determined by uniaxial tensile testing. The stress-strain curves are given in Figure 10 along with the mean Young's modulus E, yield stress σY, tensile strength at break σB and tensile elongation at break εB values. The linear elastic region was apparent until 13-19% elongation, at which point the yield strength was observed. This low-strain region allowed the determination of E = 2.32 ± 0.28 MPa and σY = 5.02 ± 0.23 MPa (Figure 10).

Beyond the proportionality limit, the tensile strength decreased (23% on average) until about 130% elongation. This can be presumably explained by the onset of the plastic deformation, not allowing the total recovery of the strains (permanent deformation of the material). The plastic region was then markedly observed with a plateau, characteristic of the drawing of the My-52-IBOMA/S chains, ending at the fracture point (εB = 490 ± 31 %). It has to be noted that the strain hardening phenomenon, consisting of the, was not observed. Consequently, no increase in stress was measured at the fracture orientation and alignment of polymer chains in the direction of the load which increases the strength and stiffness of the material (σB = 3.90 ± 0.22 MPa). A key-factor allowing the understanding of this behavior is presumably the morphology of the triblock copolymer. The possible disorganized structure of My-52-IBOMA/S, not containing necessarily P(IBOMA) glassy spheres can be detrimental. In this case, the expected physical crosslinking structure does not exist, bringing about a decrease of the entropic contribution to strain hardening. By assuming that the strain hardening process is associated with the debonding of the entanglement network (chain disentanglement), it may be also caused by the P(My) continuous phase, not sufficiently entangled (Mn,MHS,My = 51.7 kg.mol -1 < 3Me,P(My) 77 ), resulting in an unconstrained uncoiling of the chains. Generally, a flexible plastic was identified, having moderate elastic modulus and tensile strength at break combined with a good irreversible extensibility.

Conclusion

My/IBOMA NMP in bulk at 100 o C initiated by NHS-BB highlighted the gradient nature of this copolymerization with the largely higher reactivity of the monoterpene (rMy = 1.90-2.16) toward both propagating species (---My • and ---IBOMA • ) compared of that of the bulky methacrylate (rIBOMA = 0.02- 204 0.07). While the homopolymerization of IBOMA was not controlled, likely due to the high equilibrium constant of this methacrylate, the addition of 20 mol% of My as a co-monomer allowed the synthesis of a well-defined P(My-grad-IBOMA) copolymer with Đ ≤ 1.39 and Mn,MHS close to the predicted one. More generally, well-controlled copolymerizations were performed for fIBOMA,0 ≤ 0.80 and the resulting P(Mygrad-IBOMA)s exhibited Đ ≤ 1.41, Mn,MHS = 5.9-10.9 kg.mol -1 and monomodal GPC distributions at overall conversion X = 25-65%. The deviations of experimental Mn,MHS values from theoretical ones were greater with fMy,0, which may be caused by chain transfer side reactions to the monomers. Regardless of the initial feed composition, these copolymerizations had <kp><K> = 2.0 ± 0.9 10 -5 s -1 relatively close to that of My Furthermore, this triblock copolymer exhibited a tensile strength at break σB ~ 4 MPa and an elongation at break εB ~ 500%, despite the absence of strain hardening.

(1) M. Szwarc, Nature 1956, 178, 1168.

(2) M. Szwarc, M. Levy and R. Milkovich, J. Am. Chem. Soc. 1956, 78, 2656.

Chapter 4. Synthesis of isoprene-based triblock copolymers by nitroxide-mediated polymerization (NMP)

Preface

The triblock copolymers based on β-myrcene (My) did not satisfy the mechanical requirements of the industrial application outlined previously. While S-My-S copolymers (S = styrene) exhibited poor stress-strain properties (tensile strength at break σB < 0.8 MPa, tensile elongation at break εB < 200%) at room temperature, marked improvements were shown for IBOMA-My-IBOMA triblocks (IBOMA = isobornyl methacrylate) with σB ~ 4 MPa and εB ~ 500%. However, despite a satisfactory extensibility of these latter, their lack of hardness was detrimental when compared to the tensile properties of traditional vulcanized rubbers or commercial thermoplastic elastomers. One main reason may explain such low strength of these materials: a poorly entangled elastomeric poly(β-myrcene) P(My) continuous phase.

Indeed, for all these My-based triblock copolymers, the P(My) mid-segment exhibited Mn ~ 52 kg.mol -1 at most, which is only about the double of its entanglement molecular weight (Me,P(My) = 22-31 kg.mol -1 ). Two solutions have been thus considered: (1) increase of the P(My) degree of polymerization; (2) replacement of P(My) by an elastomeric segment exhibiting a lower entanglement molar mass. The first option was disregarded since the synthesis of very high Mn P(My) by NMP under homogeneous conditions in bulk or in organic solvents might not have been feasible due to the limitations of this controlled polymerization.

NMP in a dispersed medium or anionic polymerization would have been relevant but the time constraint of the doctoral thesis did not allow considering these alternatives. Consequently, it was decided to substitute P(My) by poly(isoprene) (PI), having a much lower entanglement molecular weight (Me,PI = 4-6 kg.mol -1 ).

The synthesis of well-defined PI by NMP in bulk or in organic solvents was envisioned since it was already studied and optimized as regards the polymerization of My in a controlled manner. Moreover, the effective NMP of I was reported in the literature.

In

Abstract

The polymerization of isoprene initiated by the poly(ethylene-stat-butylene)-( satisfactory extensibility, these latter diene-styrene block copolymers showed a relatively low tensile strength at break. Two axes were followed to increase σB, namely the hydrogenation of the PI mid-segment and the substitution of PS blocks by higher Tg rigid blocks.

The saturation of the soft mid-segment of the NMP-based PS-PI-PS polymers was attempted using the diimide-generating reagent p-toluenesulfonyl hydrazide (TSH), combined with tributylamine (TBA) to mitigate the participation of side reactions. A hydrogenation degree (HD) higher than 75% was determined via 1 H NMR and no significant chain degradation was observed after the post-polymerization modification as characterized by GPC. Generally, TGA evidenced higher thermal stability of the partially hydrogenated triblock copolymers under inert atmosphere compared to the parent PS-PI-PS. Moreover, the elastic modulus and the tensile strength at break of PS-PI-PS (FS = 0.49) markedly increased after the chemical hydrogenation (HD = 76.3%) and its elongation at break decreased.

Lastly, IBOMA-I-IBOMA type triblock copolymer (Mn = 94.0 kg.mol -1 , Ð = 1.76) containing a small fraction of S units in the outer blocks (FIBOMA = 0.35, FS = 0.06) was synthesized sequentially from telechelic PEB-(SG1)2. Despite its poor micro-phase separation and its early thermal degradation (onset of decomposition ranging from 219 to 265 o C depending on the nature of the atmosphere), this copolymer displayed an upper service temperature higher than 150 o C due to the glass transition of the IBOMA-rich segments occurring at 180-185 o C. Furthermore, it exhibited at room temperature a remarkable elongation at break εB = 1360 ± 210 % and a moderate tensile strength at break σB = 11.4 ± 0.6 MPa. The lower tensile stress value measured compared to these of commercial styrene-diene thermoplastic elastomers can be mainly explained by the very limited strain hardening of this experimental sample when highly stretched.

Results and discussion

The synthesis of isoprene-based triblock copolymers by NMP was performed in a very similar manner to S-My-S (Chapter 1, Section 1.4.3) and IBOMA-My-IBOMA (Chapter 3, Section 3.4.2) polymers produced previously. Therefore, PEB-(SG1)2 was first used to prepare telechelic PI-(SG1)2. Afterwards, chain-extensions were performed from these active PI-(SG1)2 macro-initiators to add PS or P(IBOMA) rigid blocks.

Preparation of SG1-terminated poly(isoprene) PI-(SG1)2

• Optimized synthesis of PI-(SG1)2.

It was first aimed to synthesize a high Mn PI segment (elevated entanglement density desired) while having a high degree of chain-end fidelity at both chain ends (for the subsequent chain-extension). Whereas the NMP of My was not reported in the literature before this doctoral thesis, the NMP of isoprene (I) was well studied using the stable free nitroxide 2, alkoxyamines 11 . They showed that a good control of the isoprene homopolymerization can be obtained when using the commercial BlocBuilder TM initiator as well as its derivative NHS-BlocBuilder.

Accordingly, it appeared herein consistent to use PEB-(SG1)2 as the initiating system 12 , which also showed its effectiveness to mediate the NMP of My (Chapters 1 and 3). Two first experiments (PI1, from 2 to 8 h, and PI2, from 11 to 17 h, experimental conditions in Table 1) were conducted in bulk since it corresponds to the easiest system to study, consisting solely of the monomer and the alkoxyamine initiator (mono-component initiating system) 13 . The kinetic study is given in Figure 1. Interestingly, two marked regimes can be distinguished. At a moderate isoprene conversion XI < 30% (experiment PI1), a good control of the polymerization was observed with low dispersity values (Ð < 1.26, Figure 1b) and experimental Mn values close to the theoretical ones (Figure 1a, GPC calibration curve based on linear PI standards). Nonetheless, at XI > 40% (experiment PI2, reaction time t = 15-17 h), a clear loss of control was seen with significant deviations of the experimental Mn values from the predicted line (Figure 1a) and broader molecular weight distributions (Ð > 1.73, Figure 1b). This unstable second regime was presumably due to the high viscosity of the reaction medium after the early stages of the polymerization. Since PI samples exhibiting relatively low Mn,PI < 35 kg.mol -1 were produced in the controlled regime, it was attempted to improve the control of the isoprene polymerization at XI > 30% via the use of a solvent.

Harrisson and coworkers reported an enhanced control over the SG1-based NMP of isoprene in pyridine compared to the same reaction in bulk 14 . Lower Ð values were obtained, which can be likely caused by the stabilization of the polar SG1 free nitroxide as well as the disruption of intramolecular hydrogen bonding due to the presence of pyridine. Therefore, a third experiment was performed for 25 h with the initial addition of 50 v% pyridine relative to isoprene (PI3, Table 1). Mn, from 36.7 to 63.9 kg.mol -1 increased linearly with conversion with minor deviations from the theoretical line (Figure 1a). A broadening of the molecular weight distribution was also observed with conversion (Ð ≥ 1.57 at XI ≥ 27%, Figure 1b).

However, it is of interest to note that these conditions allowed the synthesis of a PI sample having a relatively high Mn ~ 55 kg.mol -1 with Ð ~ 1.55. The preparation of such a PI sample was not possible in

ID [PEB-(SG1)2] (M) [I] (M) [Pyridine] (M)
Mn,target (a) (g.mol -1 ) To produce subsequently well-defined isoprene-based triblock copolymers, three PI samples were synthesized under these optimized experimental conditions and exhibited satisfactory features with Mn,1 = 51.0-55.7 kg.mol -1 and Ð1 ≤ 1.64 (Table 2). As previously performed for one of the NMP-based IBOMA-My-IBOMA triblock copolymers (Chapter 3, Section 3.4.2), a benzene (solvent) / methanol (non-solvent) fractional precipitation 15,16 was applied to the PI samples. It was aimed at eliminating the shorter PI chains, the majority of those being presumably irreversibly terminated (not capped by a SG1 unit). This way, longer PI chains can be obtained, exhibiting ideally a higher degree of chain-end fidelity. As expected, Mn increased (+ 3-6 kg.mol -1 , Table 2) and Ð decreased (-0.01-0.07, Table 2) after fractionation, indicating the effective removal of a population of relatively short PI chains.

T ( o C) t (b) (h) PI1
• Characterization of the synthesized NMP-based PI. 1.

The 1 H NMR spectrum of PI4 is presented in Figure 3. This polymer may have four different types of microstructure, namely cis-1,4, trans-

. The spectral assignments were in good accord with the literature 17, 18 . The olefinic protons' signal of cis-and trans-1,4-addition (C=CH) was observed at δ = 5.05-5.15 ppm. The signal at 4.55-4.75 ppm was attributed to olefinic protons of 3,4addition (C=CH2). The peaks at 4.80-5.00 ppm and 5.65-5.85 ppm indicated the olefinic protons of 1,2addition (CH=CH2). These peaks confirmed the formation of 1,4-PI (86 mol%) with the presence of minor fractions of 3,4-motif (6 mol%) and 1,2-motif (8 mol%). Furthermore, the sharp chemical shift from the methyl protons of cis-1,4 and trans-1,4 units (C-CH3) was observed at 1.59 and 1.67 ppm, respectively.

This way, the stereo-selectivity of PI4, containing about 71 mol% of cis-1,4 units, was elucidated. The predominance of the cis-1,4-microstructure (1,4-addition ≥ 83 mol%, cis-1,4-content ≥ 66 mol%) was also determined for the samples PI5 and PI6. P(My) homopolymers produced previously by SG1 NMP were also rich in cis-1,4-motif (1,4-addition ≥ 80 mol%, cis-1,4-content ≥ 67 mol%, Chapter 1, Section 1.4.1, Table 2).

Thermal analysis was also used to investigate the physical properties of the NMP-based PI as a function of temperature. Differential scanning calorimetry (DSC) was performed to determine the glass transition temperature (Tg). This latter characteristic temperature plays an important role in dictating whether a particular polymer is in the rubbery state at room temperature. Elastomers are amorphous in nature and thus should have a Tg much below the ambient temperature, preferably at sub-zero range 19 .

Figure 2a shows the DSC thermogram of the sample PI4. A distinct baseline shift at around -60 o C was observed, corresponding to the Tg of PI4. This sub-zero Tg thereby implies a rubber-like behavior of the 8). This might be due to the long and flexible C6/C8 pendant groups of the P(My) segment, generating free volume.

The maximum decomposition temperature (Tdec,max, temperature at which weight loss is most apparent) of PI4 under air and nitrogen atmospheres was determined via thermogravimetric analysis (TGA, Figure 2b) as well. PI4 showed a single degradation peak at Tdec,max = 389 o C under N2 flow, which was very similar to that of P(My) (Tdec,max = 385 o C, Chapter 1, Section 1.4.5, Table 8). The thermal stability of PI4 was reduced under air with Tdec,max = 361 o C, highlighting the likely occurrence of oxidation reactions generating chain scissions.

NMP-based styrene-isoprene-styrene triblock copolymers

Ideally, the synthesis of well-defined PI terminated at each chain end by a SG1 moiety allows the addition of two outer blocks via chain-extension. Preliminary styrene (2. high level of the chain end fidelity, a low fraction of styrene (S) was incorporated with a molar fraction of S in the final copolymer FS = 7.9-14.1%. For purposes of developing relatively tough and strong triblock copolymers for the industrial issue, the addition of longer PS segments to the PI macro-initiator was targeted. Accordingly, bulk chain-extensions were considered to "push" the conversion of S to higher levels. It should be noted that well-tailored P(I-b-S) polymers can be obtained via a solvent-free process.

For instance, anionic bulk polymerization to synthesize isoprene-styrene block copolymers (sequential polymerization) by reactive extrusion was reported 21,22, 23 .

The chain-extensions of PI4, PI5 and PI6 macro-initiators synthesized previously (Table 2) with S were performed in bulk at 115 o C (temperature kept under 120 o C to avoid S auto-initiation 24 ), at different reaction times (50-140 min) to prepare S-I-S copolymers with various FS. The experimental conditions are given in Table 3A below. A marked broadening of the molecular weight distribution was observed for every experiment from ÐPI = 1.47-1.59 to ÐPS-PI-PS = 2.19-2.32 (Table 3B). Several factors may explain this increase in dispersity: dead chains from PI macro-initiators; rapid increase in S conversion at short times and high initial concentration of S bringing about irreversible terminations and chain-transfer reactions 25- 27 . However, the addition of relatively long PS outer blocks was apparent since the resulting chain-extended products exhibited Mn = 89.2-101.8 kg.mol -1 and FS = 30-49% (Table 3B). As performed previously, the synthesized PS-PI-PS triblock copolymers were fractionated, resulting in an increase in Mn = 94.6-109.4 kg.mol -1 and a slight decrease in Ð = 2.11-2.29 (Table 3B).

The thermal stability of the fractionated PS-PI4-PS copolymer was investigated using thermogravimetric analysis (TGA, Figure 4b). A higher decomposition temperature Tdec,max was determined for this triblock compared to the parent PI, whether under nitrogen (+ 31 o C) or under air (+ 38 o C) atmospheres. Such an observation agrees with the higher PS Tdec,max = 445 o C, determined previously (Chapter 1, Section 1.4.5, Table 8), compared to that of natural rubber. Under air, the early thermo-oxidative ID XS (b) (%) FS (c) Mn,1 (d) (kg.mol -1 ) Đ1 (d) Mn,2 (d) (kg.mol -1 ) Đ2 (d) PS-PI4-PS e) Fractional precipitation using the non-solvent addition method was implemented, relying on the benzene (solvent) / methanol (non-solvent) pair (2 cycles).

220

characterization of block copolymers is commonly used to discriminate between homogeneous and phaseseparated systems. The existence of two glass transitions is then taken as evidence for a phase-separated state 29 . DSC measurements were carried out for the samples PS-PI4-PS and PS-PI6-PS (Figure 4a). A subzero Tg,PI ≤ -61 o C corresponding to the glass transition temperature of the flexible PI mid-block was clearly detected. At higher temperatures, Tg,PS = 102-104 o C was determined and corresponded to the rigid PS blocks (Figure 4a). A slight change in slope can be noted for the DSC trace of PS-PI6-PS at T = 80-120 o C, which can be due to the lower S molar fraction in this copolymer (FS = 38 mol%) compared to that of PS-PI4-PS (FS = 49 mol%). The difficult detection of the PS Tg was also reported in the case of S-B-S triblock copolymers (B = butadiene) for instance [30][31][32] . It can also be noted that the Tgs of the synthesized PS-PI-PS samples were very similar to those of the respective PI and PS homopolymers, suggesting a strong heterogeneity of the micro-structure.

Atomic force microscopy was applied to the study of the surface morphology of PS-PI5-PS (FS = 0.30, Mn = 94.6 kg.mol -1 , Ð = 2.29, Table 3B) to know if the two component blocks segregated into different phases with a typical microdomain structure.

(d)

The viscoelastic behavior of PS-PI4-PS and PS-PI6-PS in the rubbery region was investigated as well (Figure 4c). The storage modulus G' for PS-PI4-PS stayed relatively constant up to 90 o C, at which point the PS phase softened. A drastic drop of G' was observed at T ~ Tg,PS = 102 o C (Figure 4a) since the block copolymer chains were no longer locked into position and all the chains were free to flow. A much more atypical behavior was seen for the sample PS-PI6-PS since G' gradually decreased with temperature without a sharp drop at T ~ Tg,PS (Figure 4c). It might be explained by the possible specific threedimensional network morphology of PS-PI6-PS. Considering the PS content of this sample (FS = 38%), a gyroid structure could have been obtained, which might offer a higher order-disorder phase transition temperature (TODT) relative to their one-and two-dimensional counterparts (lamellae and cylinders) 37 . Such an explanation is highly hypothetical and an in-depth study should be led to understand these viscoelastic properties, starting with the elucidation of PS-PI6-PS micro-structure (AFM or TEM combined with SAXS for instance).

The tensile stress-strain plots of the synthesized PS-PI-PS polymers are presented in Figure 6. The tensile bar specimens were prepared by casting a solution of polymer in dichloromethane (~ 80 wt%) and driving off the solvent at room temperature (Figure 6d). The film samples were cut into typical dog-bone shapes using a sharp blade. As expected, tensile stress at break σB increased and elongation at break εB decreased as FS increased. While PS-PI5-PS (FS = 30 mol%) can be considered as a flexible plastic with a relatively high irreversible extensibility (εB = 449 ± 36 %, Figure 6a), a hard and brittle PS-PI4-PS plastic was characterized (FS = 49 mol%) with σB = 10.8 ± 1.2 MPa and εB = 21 ± 7 % (Figure 6c). Considering our industrial criterion, PS-PI6-PS appeared to be the most satisfactory triblock copolymer, which was as soft as PS-PI5-PS (similar Young's modulus E and εB) but tougher (σB = 4.11 ± 0.22 MPa, Figure 6b). As reported previously for the NMP-based IBOMA-My-IBOMA type triblock copolymer containing 36 mol% of IBOMA (Chapter 3, Section 3.4.2, sample My-52-IBOMA/S, Figure 10), no strain hardening (strengthening of the material during large strain deformation) was observed for these samples, limiting thus their tensile strength at break. Such a behavior may be explained by their possible lamellar microstructure, not allowing a large-scale orientation and alignment of the macromolecules in the direction of the load. It may be assumed that a PS-PI-PS sample exhibiting a spherical morphology (hard blocks forming discrete minority spheres embedded in a rubbery matrix) would have favored the strain hardening due to its isotropic nature. Generally, these isoprene-based triblock copolymers showed largely improved Despite encouraging results, the PS-PI-PS triblock copolymers prepared by NMP exhibited a tensile strength at break below the industrial requirements. Two options were considered to increase σB: hydrogenation of the PI mid-segment; replacement of PS outer blocks by higher Tg hard blocks.

Hydrogenation of NMP-based styrene-isoprene-styrene triblock . copolymers

Industrially, hydrogenation is conveniently performed with elemental hydrogen in the presence of a transition metal catalyst. On a laboratory scale, however, the equipment required to perform hydrogenations at moderate to high hydrogen pressure is often not readily available. For this reason, the inorganic reagent diimide (N2H2) has been widely used to prepare laboratory scale quantities of hydrogenated diene polymers (Literature Review, Section C.2.2 for further details) 38, 39 .

Hydrogenation of the unsaturated PI mid-segment of the triblocks synthesized by NMP was performed herein with diimide generated in-situ by thermolysis of p-toluenesulfonyl hydrazide (TSH).

Since the by-products from TSH, particularly p-toluenesulfinic acid, can attack at olefinic sites (protonation of unsaturation) and degrading macromolecules, tributylamine (TBA) was added to eliminate these side reactions. The protonation of this high-boiling tertiary amine was expected to greatly decrease the propensity of the by-products to attack at the polymer backbone 40 . The chemical hydrogenation performed for PS-PI4-PS and PS-PI6-PS is depicted in Figure 7, relying on Hillmyer's study 41 . The unsaturated polymers were first dissolved in p-xylene (40 mL per gram of polymer) at 45 o C and then, TSH, TBA and a small amount of hydrogen-atom-donating butylated hydroxytoluene (BHT) was added to protect against oxidative degradation. The reaction mixture was stirred at 125 o C for 3 h and the polymer was precipitated in methanol. Prior to hydrogenation, the thermal stabilization of the chain-ends of PS-PI-PS was ensured via removal of the SG1 end-groups using thiophenol as described elsewhere 42 .

The 1 H NMR spectra of PS-PI4-PS before and after hydrogenation with TSH is given in Figure 8.

The partial hydrogenation was evidenced by the diminution of the olefinic resonance at δ = 5.08 ppm. A hydrogenation degree (HD) of 76.3% was determined spectroscopically. The TSH/TBA combination did not give complete saturation of the PI block, which is consistent with the literature reporting the incomplete hydrogenation of PI via this specific hydrogenation method 40 . Figure 9 shows the GPC traces of the starting PS-PI4-PS and PS-PI6-PS polymers as well as the partially hydrogenated polymers. A slight broadening of the TSH-treated material molecular weight distribution was apparent, although the shape of the GPC traces did not change perceptibly. Furthermore, lower Mn were determined after hydrogenation for both isoprenebased triblock copolymers. A degradation of the PI mid-segment during the hydrogenation, which may involve TSH/TBA by-products reacting with unsaturated sites, can explain this change in Mn 40 . It can also be hypothesized that some chains decomposed due to thermal or oxidative reactions. It should be added that the molecular features obtained after hydrogenation must be carefully interpreted due to the change in the GPC response of the hydrogenated polymers, exhibiting a hydrodynamic volume different from the PS-PI-PS starting materials.

Thermal gravimetric analysis (TGA) was used to determine the thermal stability for both the original polymers and the hydrogenated ones. Regardless of the hydrogenation degree, similar TGA traces were obtained for PS-PI4-PS (FS = 0.49) under air (Figure 10a). However, after hydrogenation (HD = 76.3%), this sample displayed an enhanced thermal stability under N2 atmosphere since no weight loss occurred before 362 o C (305 o C for the parent polymer) and Tdec,max = 436 o C, slightly higher than that of the neat polymer (419 o C), was determined (Figure 10a'). With regard to PS-PI6-PS (FS = 0.38), while a similar improvement of the thermal stability after hydrogenation (HD = 77.3%) was observed under inert atmosphere (Figure 10b'), higher onset and maximal decomposition temperatures (Tdec,1 = 302 o C and Tdec,max = 416 o C respectively) of the hydrogenated sample were determined under air (Figure 10b).

Generally, the greater thermal stability of the PS-PI-PS copolymers with the partial saturation of the isoprene units can be explained by the conversion of the weak π-bond within the original PI segment to a stronger C-H σ-bond 43 .

The partially hydrogenated PS-PI4-PS triblock copolymer was tested via a stress-strain experiment at ambient conditions (Figure 11, must be discussed cautiously due to the poor reproducibility of the specimens after hydrogenation) and exhibited an elastic modulus (E = 10.3 ± 2.9 MPa) about four times higher than that of the parent PS-PI4-PS. Moreover, a marked increase of the tensile stress at break (from 10.8 to 16.5 MPa) and a reduction of the elongation at break (from 21 to 9%) were observed after the hydrogenation reaction. Accordingly, the tensile strength and the Young's modulus of this partially 4A). The monomodal molecular weight distribution of PI7 is shown in Figure 12. The chain-extension of PI7 with a IBOMA/S mixture (91/9 mol%) in 50 wt% toluene was subsequently performed at 115 o C (Table 4B). A small fraction of S was added as a controlling co-monomer to enhance the control of this bulky methacrylate chain-extension. A monomodal shift in molecular weight as a function of elution time from macro-initiator to final chain-extended copolymer was observed in the GPC chromatogram (Figure 12), indicating the successful addition of IBOMA/S segments. The resulting PI7-P(IBOMA-co-S)2 triblock copolymer (Mn = 73.5 kg.mol -1 , Ð = 2.00) was fractionated via the nonsolvent addition technique using the benzene/methanol pair. After fractionation, the tail on the low molecular weight edge was partly removed (Figure 12) and the triblock copolymer exhibited higher Mn = 94.0 kg.mol -1 and narrower molecular weight distribution (Ð = 1.76). Fractionated PI7-P(IBOMA-co-S)2 possessed two outer rigid segments containing 85 and 15 mol% of IBOMA and S units respectively, as indicated in Table 4C.

The thermal stability of the synthesized triblock copolymer was assessed by thermogravimetric analysis. Figure 13b represents the thermal degradation curves of PI7-P(IBOMA-co-S)2 under air and nitrogen atmospheres. As expected, the polymer degraded at higher temperatures under inert atmosphere. e) Predicted Mn of the whole chain-extended triblock copolymer at X, measured experimentally, and calculated as follows:

Mn,theo,X = (X/100)(Mn,theo-Mn,PI7)+Mn,PI7.

the rigid IBOMA/S backbone.

Figure 13a shows the DSC trace of PI7-P(IBOMA-co-S)2. The Tg of the soft PI block was clearly observed at -63 o C, similar to that of PI homopolymer (-60 o C, Figure 2a). Although harder to detect, the Tg of the IBOMA-rich segments was also determined at + 182 o C. It is consistent with the literature reporting the Tg of P(IBOMA) which ranged from 174 to 206 o C depending on the chain tacticity 50 . The relatively low Tg of the PI7-P(IBOMA-co-S)2 rigid segments herein may thus be caused by the high isotactic content of the P(IBOMA) sequences. A decrease of this upper Tg may also be due to the presence of S units, by assuming the formation of a homogeneous phase with IBOMA units. Generally, the Tg of these IBOMArich segments was much higher than that of the PS hard blocks (102-104 o C, Figure 4a), reflecting the reduced chain flexibility due to the steric hindrance of the large isobornyl side groups. Two questions can be raised following this DSC characterization: Did PI7-P(IBOMA-co-S)2 phase separate as suggested by the detection of two Tgs? Did PI7-P(IBOMA-co-S)2 exhibit extended service temperature compared to analogous S-I-S triblock copolymers?

The morphology of PI7-P(IBOMA-co-S)2 was visualized by atomic force microscopy (AFM).

Figure 14 presents a representative AFM topography and phase trace of this triblock copolymer. An unusual organization of the two phases at the micro-scale was observed with a possible core-shell type structure (dark domain) inserted in a continuous matrix (brighter domain). Although PI7-P(IBOMA-co-S)2 exhibited FI = 0.59, assigning the brighter domain to PI is not trivial. It could be for instance argued that, considering its elastic nature, PI, which was trapped between two rigid segments, tended to protrude out from the surface. These protrusions appeared as peaks in the 3D topography image (Figure 14). In this case, the darker domain would correspond to PI. Generally, this AFM study did not allow to distinctly determine a morphology (weakly aligned cylinders/lamellae might be present) and this system could be qualified as poorly ordered.

To further elucidate thermal behavior along with viscoelastic properties, the triblock was subjected to dynamic mechanical analysis (DMA). Figure 13c shows the dynamic elastic modulus (G') data as a function of temperature for PI7-P(IBOMA-co-S)2. This sample had a modulus of approximately 100 MPa in order of magnitude from room temperature to about 150 o C. G' in the rubbery plateau was thus significantly higher than that of the analogous IBOMA-My-IBOMA type copolymer (1-10 MPa, Chapter 3, Figure 8b) and relatively similar to that of S-I-S triblock copolymers (100-1000 MPa, Figure 4c), Tensile testing at room temperature for PI7-P(IBOMA-co-S)2 was also performed to know if this NMP-based block copolymer can mechanically compete with industrial styrene/diene thermoplastic elastomers. The stress-strain curves of four PI7-P(IBOMA-co-S)2 dog-bone specimens are given in Figure 15. The limit of elastic behavior was marked by a yield point at 20-30% elongation. It was followed by a non-reversible plastic deformation until the specimen failure at εB = 1356 ± 214 %. Such a stretch capacity is comparable to commercial S-I-S and S-B-S (B = butadiene) thermoplastic elastomers as mentioned in the literature review (Section C.3.2, Table 8, εB = 880-1300%). Interestingly, a reduced extensibility was observed for S-I-S samples previously synthesized (Figure 6) with εB ≤ 450 %, while they possessed a similar PI mid-segment (Mn ~ 50-60 kg.mol -1 , Ð = 1.45-1.60). Accordingly, it might be due to the PS outer segments of S-I-S not being able to withstand mechanical stress at high elongation unlike the IBOMA-rich blocks of PI7-P(IBOMA-co-S)2. This latter sample exhibited a tensile strength at break, σB = 11.4 ± 0.6

MPa, largely higher than these of My-based block copolymers (σB < 4 MPa, Chapter 1, Section 1.4. can explain the relatively low σB of PI7-P(IBOMA-co-S)2 compared to traditional styrene/diene thermoplastic elastomers. The poorly ordered micro-domain structure of PI7-P(IBOMA-co-S)2, as suggested by the AFM analysis (Figure 14), might have been the main cause of this stress-strain behavior, hindering the orientation and the alignment of the macromolecules in the direction of the load.

It should be noted that hydrogenation of the IBOMA-I-IBOMA type copolymer was not performed due to time constraints. This latter IBOMA-I-IBOMA type block copolymer corresponded to the most conclusive material prepared by NMP in this doctoral thesis in terms of mechanical performance at ambient temperature.

Conclusion

However, its tensile strength at break remained below the industrial requirement. Hydrogenation of the elastomeric mid-segment at normal pressure using diimide can be considered to increase as implemented for a S-I-S sample (Mn = 103.1 kg.mol -1 , Ð = 2.11, FS = 0.49). σB of the partially hydrogenated triblock copolymer (HD = 76.3%) increased of about 53% compared to the parent S-I-S, despite a reduced extensibility.

( The synthesis by nitroxide-mediated polymerization (NMP) to yield isoprene-based triblock copolymers exhibiting valuable stress-strain properties was highlighted in the previous chapter. However, the lack of toughness of these synthesized copolymers can be stressed when comparing with the mechanical features of commercial thermoplastic elastomers (TPE).

This technical chapter is dedicated to the thermal, mechanical and rheological characterization of commercial styrene-isoprene-styrene (S-I-S) and styrene-ethylene/butylene-styrene (SEBS) triblock copolymers, all synthesized industrially by living anionic polymerization. Although properties of sold products are generally reported by companies, in-house measurements were required in order to use the same testing machines throughout the doctoral thesis. The objective was not to find a marketed candidate TPE for the industrial problem. It was only aimed at targeting commercial styrenic block copolymer TPE satisfying the ArianeGroup mechanical requirements at room temperature. The group of styrenic block copolymers was selected to allow direct comparisons with the previous isoprene-based triblocks made by SG1-based NMP (Chapter 4).

Abstract

The relatively low tensile strength at break (σB = 7.5 ± 0.6 MPa) of an initially supplied S-I-S possessing about 8 mol% of styrene (S4-I-S4) prompted us to study a second S-I-S richer in PS (FS ~ 0.16, S8-I-S8, the subscripts state the mol% S in the polymer). Doubling the fraction of the rigid phase resulted in higher but not sufficient σB = 12.7 ± 0.7 MPa. The preparation of tensile specimens by a solvent casting instead of compression molding allowed this latter copolymer to be tougher (σB and elongation at break εB increased by 13% and 25% respectively). Moreover, the partial hydrogenation (67 mol%) of S8-I-S8 enhanced its mechanical strength at low-and high-elongation regions. A styrene-butadiene-styrene (S-B-S) having its soft mid-segment completely hydrogenated and a high fraction of PS (48 mol%) showed the most conclusive stress-strain properties at ambient temperature with σB = 25.0 ± 2.1 MPa and εB = 730 ± 40 %. However, its limited thermal stability would impose a welding with the tank and the skirts at T ≤ 275 o C under inert atmosphere.

Results

Commercial S-I-S from Sigma-Aldrich ®

A preliminary styrene-isoprene-styrene S-I-S provided by Sigma-Aldrich ® (product number = 432393, experimental section) containing only 14 wt% of S units, as outlined in the specification sheet, was first characterized and the following features were determined in the laboratory:

-8.4 mol% S was determined by 1 H NMR (triblock thus abbreviated S4-I-S4 for the study).

-Mn = 98.3 kg.mol -1 and Ð = 1.29 were measured via GPC (Figure 1).

-DSC analysis gave Tg,soft segment = -60 o C (Figure 2a) and the Tg of the PS blocks was not detected (presumably due to the low S molar fraction).

-As typically observed for styrene-based TPE, the storage modulus stayed relatively constant (G' = 0.1-0.3 MPa) until about 100 o C where the PS hard phase softened (Figure 3).

-Tensile stress at break σB = 7.5 ± 0.6 MPa and elongation at break εB = 4250 ± 230 % were determined via uniaxial tensile testing.

The relatively low σB of this commercial styrenic TPE was not satisfactory for the industrial application. Consequently, a similar S-I-S from Sigma-Aldrich ® possessing a higher fraction of S units (22 wt% as specified, product number = 432415, experimental section) was subsequently tested and exhibited:

-15.9 mol% S as determined by 1 H NMR (triblock abbreviated S8-I-S8).

-Mn = 97.2 kg.mol -1 and Ð = 1.25 via GPC measurement (Figure 1).

-Tg,soft segment = -61 o C (Figure 2a). -Well-ordered micro-phase separation revealing possibly a cylindrical/lamellar morphology as observed by AFM (Figure 5) despite a low FS ~ 0.16.

-Higher Young's modulus E = 0.19 ± 0.04 MPa, σB = 12.7 ± 0.7 MPa and lower εB = 3740 ± 140 % compared to the mechanical properties of S4-I-S4 at room temperature (Figure 4).

-Higher G' in the rubbery plateau (1-3 MPa) as well as at T > Tg,PS (Figure 3). Although S8-I-S8 was tougher than S4-I-S4 with notably an enhanced tensile strength at low and high deformations, a higher σB was required. To improve the properties of the commercial S8-I-S8 triblock copolymer, two strategies were undertaken: (1) preparation of specimens via film casting instead of compression molding; (2) hydrogenation of the elastomeric block.

The way to prepare the samples prior to tensile testing or DMA analysis can significantly affect the results. When molding the specimens via hot press (compression-molded specimens), the relatively fast cooling phase after molding may hinder the equilibrium morphology to occur. For films prepared from direct casting, the solvent evaporates slowly thus allowing a stable organization of the macromolecules close to the thermodynamic equilibrium. Accordingly, a stronger phase separation may be expected for S-I-S cast film. For this study, dichloromethane casting allowed the preparation of dog-bone style tensile specimens and rectangular bars for DMA. lower extensibility with εB = 3310 ± 95 %. Furthermore, the TGA results before and after hydrogenation were of interest. Whereas the thermal stability of S8-I-S8 was slightly improved under air after incomplete hydrogenation (Figure 2b), the commencement of the thermal decomposition (+ 55 o C) as well as the maximal degradation at about 50 wt% loss (+ 48 o C) occurred at much higher temperatures under nitrogen (Figure 2c). Lastly, a hydrogenated S8-I-S8 cast film bar was tested via DMA (Figure 3 The preparation of cast films as well as the hydrogenation of the flexible mid-segment allowed an enhancement of the TPE mechanical and thermal properties. Partially hydrogenated S8-I-S8 cast film exhibited conclusive properties for our industrial criterion. Since solvent casting process can be considered not adequate from an industrial point of view, attention was then paid to the characterization of a fully hydrogenated styrenic block copolymer TPE.

Commercial SEBS from Kraton ®

Styrene-ethylene/butylene-styrene (SEBS) block copolymers are produced industrially in a twostep process: (1) synthesis of styrene-butadiene-styrene (SBS) by living anionic polymerization; (2) Conversion of SBS to SEBS through selective and quantitative hydrogenation of the rubbery mid-block using high-pressure H2 and a catalyst. SEBS A1536 H (industrial specifications in the experimental section)

from Kraton ® was purchased and characterized in this study. hydrogenation of common SBS copolymers can result in a predominantly PE mid-block that crystallizes.

The damping factor (tan δ = G''/G') is also given in Figure 6b (grey curve) and a local peak appeared at 89 o C, referring likely to the glass transition temperature of the PS outer blocks.

Generally, this commercial SEBS triblock copolymer satisfied the ArianeGroup specifications in terms of stress-strain properties at room-temperature. The high styrene content associated to the complete hydrogenation of the poly(butadiene) segment allowed this styrenic block copolymer to display a high mechanical strength while having a largely acceptable extensibility.

For purposes of welding the candidate TPE with the thermoplastic composite considered for the application (constituting the tank and the skirts of the launcher), the assembly must be heated at an elevated temperature (T > 250 o C) for a few minutes. Consequently, the TPE must be stable thermally and any degradation would make the welding process inefficient. To deem the ability of the purchased SEBS to resist high temperatures, isothermal TGA were led (Figure 7). Under air, While SEBS did not decompose at 200 o C for almost 90 min, the sample weight fell after 20 min at T = 230 o C (Figure 7a), thus not allowing the welding in an O2 environment. Under nitrogen, the hydrogenated SBS was stable for 90 min at T = 230-275 o C. However, a 10% weight loss was apparent when running the TGA at 290 o C (Figure 7b).

Accordingly, the limited thermal stability of SEBS imposes the welding under inert atmosphere at T ≤ 275 o C.

Conclusion

This chapter aimed to identify reference styrenic block copolymer TPE providing conclusive mechanical features at room temperature with respect to the ArianeGroup issue. Commercial S-I-S triblock copolymers (Mn = 97-98 kg.mol -1 , Ð = 1.25-1.29) exhibiting a minor fraction of S (≤ 16 mol%) were first evaluated. Despite a remarkable extensibility with εB = 4250 ± 230 %, S-I-S having FS ~ 0.08 was not strong enough with σB = 7.5 ± 0.6 MPa. A slightly less stretchable S-I-S (εB = 3740 ± 140 %) richer in S (FS ~ 0.16), which micro-phase separated into lamellae (or cylinders), showed a higher tensile strength at break σB = 12.7 ± 0.7 MPa, although not sufficient for the considered application. Following these characterizations, the importance of the fraction of the rigid phase in the copolymer was highlighted.

Mechanical properties in this family of TPE can be significantly tuned based on composition. While TPE displaying a spherical morphology (hard blocks forming discrete minority spheres embedded in a rubbery matrix) were initially privileged for this thesis due to their isotropic nature, it can be concluded herein that S-I-S having FS > 20 mol% (with likely a cylindrical/lamellar morphology) must be targeted in order to reach a satisfactory σB.

GENERAL CONCLUSION

This doctoral thesis aimed to synthesize a tough and stable thermoplastic elastomer (TPE), capable of providing a soft junction between polar composite matrices (tank and skirts of a launcher). The TPE, joined to the composites by adhesive bonding or welding, is used to transmit mostly shearing forces occurring at the joint and has to meet specific requirements, including:

-tensile behavior similar to that of vulcanized elastomers.

-minimal drop of its mechanical properties at T ≤ 80 o C.

-no flowing nor degradation during the joining process (T ~ 160-180 o C for more than 10 h if adhesive bonding is employed, T ~ 240-280 o C for a few minutes in the case of a welding).

Styrenic triblock copolymers, consisting of a poly(diene) segment (continuous domain) and two poly(styrene) (PS) end segments (discrete domain), were initially targeted since most of the commercial ones satisfy the ArianeGroup mechanical requirements at room temperature (primary focus). Nitroxidemediated polymerization (NMP) was selected mainly due to its simplicity to implement, its effectiveness for controlling the polymerization of 1,3-dienes, which also conveniently aligns with the expertise of Prof. The table below (page 255) gives a summary of the main features of the experimental and commercial block copolymers characterized during this PhD project.

With respect to the industrial issue, two main modifications were implemented to the synthesized polymers:

-Incomplete hydrogenation (~ 75%) of the elastomeric PI mid-segment of S-I-S at normal pressure using diimide-generating reagent p-toluenesulfonyl hydrazide (TSH). A higher thermal stability of the partially saturated triblock copolymers was evidenced under inert atmosphere compared to the orgininal S-I-S. Furthermore, the tensile strength at break markedly increased after the chemical hydrogenation, despite a reduced extensibility.

254

-Functionalization of P(My): the introduction of functional groups into the soft segment was performed to encourage the joining process with the polar composite matrices. Bulk NMP of My and glycidyl methacrylate (GMA) initiated by NHS-BB was accomplished at 120 o C and led to the manufacture of well-tailored epoxide functionalized P(My). During the assembly, the presence of epoxy groups into the candidate material could either increase the miscibility at the interface to favor chains inter-diffusion (welding) or allow covalent bonding at the interface (adhesive bonding).

Generally, the various block copolymers produced by NMP in this project did not satisfy the ArianeGroup specifications, due in part to their lack of mechanical strength. This deficiency can be explained by the absence of strain hardening of these NMP-based triblock copolymers in the high-strain region, not allowing the tensile stress to be greatly increased. Such a tensile behavior may be due to the Three main modifications can be undertaken to synthesize the candidate TPE by NMP, namely:

1. Synthesis of PI in a dispersed medium: NMP under homogeneous conditions in bulk or in organic solvents displays a main challenge: the synthesis of high average molecular weight polymers. It was stressed in this project with the difficulty to produce in bulk or in pyridine a well-defined PI having a number-average molecular weight (Mn) higher than 60 kg.mol -1 . The aim to increase the average chain length of the elastomeric mid-segment is to develop a more entangled matrix and thus a tougher resulting triblock copolymer. Performing isoprene NMP in an aqueous dispersed medium (suspension of stable nanoor micro-particles) may be a solution. It was found that compartmentalization of the polymerization in particles (acting as nano-reactors) can improve control over molar mass distributions and afford a higher degree of end functionality. Low dispersity can be maintained even at high conversions, allowing the synthesis of well-tailored polymers exhibiting high Mn 1 . NMP in mini-emulsion, using either a watersoluble or a lipophilic initiating system 2 , could be for instance envisioned to prepare a "living" PI latex with Mn > 60 kg.mol -1 and Ð < 1.3. ) would be of interest since it is thermally stable under N2 (Tdec,onset ~ 340 o C) 6 . Its high Tg would also ensure an extended upper service temperature compared to traditional styrenic TPE as well as an elevated tensile strength at break σB. Furthermore, welldefined P(MBL) (Mn = 22 kg.mol -1 , Ð = 1.21) was synthesized via addition-fragmentation chain-transfer polymerization (RAFT) 6 , which indicates the propensity of MBL to be polymerized by a controlled radical process. Lastly, the addition of neat P(MBL) blocks to PI would be possible while IBOMA chain-extensions were performed with a small amount of styrene (S) to enhance the control of this bulky methacrylate NMP.

It could be argued that the presence of styrene in the rigid segments can reduce the strength and the performance at high temperature of the triblock copolymer due to the relatively low Tg of PS (~ 95-100 o C) 7 .

Please note that P(MBL) is herein mentioned as an example. P(MBL) derivatives, such as poly(γmethyl-α-methylene-γ-butyrolactone) (P(MeMBL), Tg = 210-220 o C, MeMBL copolymerized with styrene via RAFT emulsion

Hydrogenation of the PI middle block:

As performed in this project for experimental and commercial S-I-S, the saturation of the soft segment appears paramount. The improved thermal stability and the higher tensile strength at break exhibited by the isoprene-based copolymer after hydrogenation are beneficial, despite a loss of extensibility. A higher upper service temperature could also result from hydrogenation, due to the stronger phase separation of the hydrogenated block copolymer (greater incompatibility between hard and soft blocks) compared to that of the parent one 12 . The chemical hydrogenation using diimide reduction in this study led to incomplete replacement of isoprene units by ethylene/propylene units. Applied industrially, the catalytic hydrogenation of PI could be considered using a transition metal and high-pressure H2 in order to saturate quantitatively the poly(diene) segment and enhance the mechanical and thermal properties of the triblock copolymer.

• (Co-)Polymerization of β-myrcene (with styrene) by NMP.

The (co-)polymerizations were performed in a 25-mL three-necked round-bottom glass flask equipped with a reflux condenser, a thermal well and a magnetic stir bar. The flask was placed inside a heating mantle and the whole set-up mounted on top of a magnetic stirrer. Table 1 and Table 5 list the formulations studied for the various My(/S) (co-)polymerizations. For example, for the experiment My/S-20 (molar fraction of S in the initial feed fS,0 = 0.20), the reactor was sealed with the rubber septa after the addition of NHS-BB (0.204 g, 0.427 mmol) and the stir bar. My (10.514 g, 77.178 mmol) and previously purified S A thermocouple connected to the temperature controller was placed inside the thermal well and connected through one of the necks. A mixture of ethylene glycol/distilled water (20/80 vol%) at a temperature of 5 °C was circulated (Fisher Scientific Isotemp 3016D Digital Refrigerated Bath) through the condenser connected to one of the necks of the reactor to prevent any evaporation loss of the monomers. A purge of ultra-pure nitrogen was then introduced to the reactor for 30 min to deoxygenate at room temperature the reactants prior to polymerization. The purge was vented through the reflux condenser. After purging, the reactor was heated at a rate of about 10 °C.min -1 to the desired polymerization temperature (T = 120 °C for My/S-20) with continuous nitrogen purge. The time at which the reactor temperature reached 100 °C was taken arbitrarily as the commencement of the reaction (t = 0 min). Samples were then taken from the reactor periodically by a syringe until the end of the experiments or until the samples became too viscous to withdraw. Reactions were then stopped by removing the reactor from the heating mantle and letting the contents cool down to room temperature, while under continuous nitrogen purge. For each sample withdrawn during the polymerization, the crude polymer was precipitated with excess methanol. After filtration and recovery, the precipitated polymer was dried at 40 °C under vacuum in the oven overnight to remove unreacted monomers. Samples were analyzed by nuclear magnetic resonance (NMR) and gel permeation chromatography (GPC). Specifications of the GPC and the NMR are fully described in the Characterization section. At the end of the experiment, the overall NMR conversion for My/S-20 was 37.1 % (individual conversions: XMy = 40.5 % and XS = 28.9 %) with Mn = 9.9 kg.mol -1 , molecular weight distribution Đ = 1.26 and the molar composition of S in the final copolymer was FS = 0.15 according to NMR spectroscopy. The exact same procedure was followed for all My/S copolymerizations and My homopolymerizations.

• Chain-extension of poly(myrcene) P(My) homopolymer macroinitiator and poly(myrcene-stat-styrene)

P(My-stat-S) statistical copolymer macroinitiator with styrene and/or β-myrcene. P(My) homopolymers and P(My-stat-S) statistical copolymers were chain-extended with purified S and/or My at 110 °C or 120 °C in 50 wt% toluene solution or in bulk. The experimental setup and procedures were the same as the syntheses for the P(My) and P(My-stat-S) polymers described earlier. The entire set of formulations for the chain-extension experiments is shown in Table 3, from P(My) macroinitiators, and Table 7.B, from P(My-stat-S) macroinitiators. An example is given to illustrate the chain-extension of My/S statistical copolymer macroinitiator with purified S. To the reactor was added My/S-20 macroinitiator (1.00 g, Mn = 9.9 kg.mol -1 , Đ = 1.26) and toluene (7.51 g, 81.51 mmol) solvent with magnetic stirring. After sealing the reactor, previously purified S (6.49 g, 62.31 mmol) was added via syringe and a purge of nitrogen was applied for 30 minutes. The reactor was then heated to 110 °C to commence polymerization while continuing the nitrogen purge. Samples were periodically removed by syringe and polymers were precipitated using excess methanol. After removal of the supernatant, the wet cake was dried under vacuum at 50 °C to obtain the My/S-20-S block copolymer. The chain-extended products were characterized by GPC, calibrated with poly(styrene) PS standards, at 40°C in THF eluent. For the specific example cited, the final S conversion was 53 % according to 1 H NMR spectroscopy with Mn = 25.3 kg.mol -1 and dispersity Đ = 1.46 determined by GPC.

272

options, 5 fitting cycles with high resolution, proportional line width new spectrum with factor 0.30) was applied to all 13 C NMR spectra in order to improve the quality of the integrations.

Two P(My-stat-S) copolymers were analyzed by 31 P NMR spectroscopy to determine the fraction of the chains terminated with SG1. The spectra were recorded in CDCl3 using a 5 mm diameter Up NMR tube with 800 scans being performed in a 200 MHz Varian Gemini 2000 spectrometer operating at 81 MHz. The NMR tubes were carefully weighed and filled with polymer (mass of polymer = 0.0682 g and 0.1179 g for My/S-07 with Mn = 11.1 kg.mol -1 and My/S-55 with Mn = 16.6 kg.mol -1 respectively) and diethylphosphite as internal reference (0.0044 g and 0.0031 g respectively). In order to know if My-based polymers and diethylphosphite have similar relaxation times, My/S-55 was run under the exact same conditions with only one scan and no dummy scans (ss = 0). A negligible difference (< 1.9 %) in integral values was measured between this spectrum and the standard one with multiple scans.

It was thereby assumed that diethylphosphite and My-based polymers with moderate Mn < 17 kg.mol -1 relax at the same rate.

The number-average molecular weights (Mn,GPC) and the molecular weight distributions (Đ = Mw/Mn) were measured using gel permeation chromatography (GPC, Water Breeze) with HPLC grade tetrahydrofuran (THF) as the mobile phase. A mobile phase flow rate of 0.3 mL.min -1 was applied and the GPC was equipped with 3 Waters Styragel ® HR columns (HR1 with a molecular weight measurement range of 10 2 -5 × 10 3 g.mol -1 , HR2 with a molecular weight measurement range of 5 × 10 2 -2 × 10 4 g.mol -1 and HR4 with a molecular weight measurement range of 5 × 10 3 -6 × 10 5 g.mol -1 ) and a guard column was used. The columns were heated to 40 °C during the analysis. The molecular weights were determined by calibration with linear narrow molecular weight distribution PS standards (PSS Polymer Standards Service GmbH, molecular weights ranging from 682 g.mol -1 to 2,520,000 g.mol -1 )

and the GPC was equipped with a differential refractive index (RI 2414) detector. The P(My) contribution to Mn,GPC was converted using the appropriate Mark-Houwink-Sakurada (MHS) coefficients (MHS parameters determined at 40 °C with THF eluent for PS 4 : KPS = 15.8 × 10 -5 dL.g -1 and αPS = 0.706 ; MHS parameters determined at 30 °C with THF eluent for 1,4-P(My) 5 , KP(My) = 7.46 × 10 -5 dL.g -1 and αP(My) = 0.772). Therefore, Every Mn value reported in the Results & Discussion section was calculated using the following formula: Mn = FMy × Mn,MHS + FS × Mn,GPC where FMy and FS are the respective molar fraction of My and S in the final polymer and Mn,MHS = [(KPS/KP(My)) × Mn,gpc αPS1 ] (1/(αP(My)+1)) using the Mark-Houwink relationship 6 .

Thermogravimetric analysis (TGA) was carried out using Q500TM from TA Instruments under nitrogen flow at a ramp rate of 10 °C.min -1 . Samples were heated in aluminum pans. This technique allowed the decomposition temperature (Td) to be determined. Differential scanning calorimetry (DSC, Q2000TM from TA Instruments) was also performed under nitrogen atmosphere to measure the glass transition temperatures (Tgs) of P(My), P(My-b-S) and P(My-stat-S) polymers. Calibrations for temperature and heat flow were done using indium and benzoic acid standards respectively. The DSC was typically performed in the temperature range from -90 °C to + 150 °C using three scans per cycle (heat/cool/heat) at a rate of 10 °C.min -1 . To eliminate the thermal history, glass transition temperatures were reported from the second heating run. The reported Tgs were calculated using the injection method from the change in slope observed in the DSC traces.

The stress-strain behavior of the synthesized P(My-b-S) diblock copolymers was recorded using a MTS Insight TM material testing system with a 5 kN load cell at room temperature and a cross-head speed of 1 mm.min -1 .

Dog-bone style tensile specimens (ASTM D638, type V, overall length = 63.5 mm, overall width = 9.53 mm) were made using a hot press (Carver hydraulic press, model #3925) at T ~ Tg,specimen + 50 °C where Tg,specimen refers to the specimen's higher Tg, under a mean pressure of 5000 psi. The width (3.18 ± 0.5 mm) and the thickness (1.40 ± 0.3 mm) of the narrow section of each specimen were previously measured using a digital ruler. Each test was considered finished after the complete break-up of the specimen in two distinct parts. A minimum of five specimens was tested for each P(My-b-S) diblock copolymer sample and averaged results were reported. This mechanical characterization allowed the Young's modulus, the ultimate tensile strength and the elongation at break to be determined, using TestWorks 4 software. The Young's modulus was determined as the slope of the stress-strain curve at strains of 0-0.5 %.

• References of the experimental section.

(

1,4-motif (predominant)

The fraction of P(My-stat-S) chains terminated with SG1 was calculated by comparing the two integrated peaks corresponding to the phosphorus atom of the SG1 group (δ = 24.8 ppm) and the phosphorus atom of diethyl phosphite (δ = 7.3 ppm) used as an internal reference 3,4 . Using the known amounts of diethyl phosphite (0.0044 g) and P(My-stat-S) copolymer (0.0682 g with Mn = 11.1 kg.mol -1 ), the proportions of chains containing the SG1 moiety attached to the My/S-07 terminal unit was estimated to 94 % ± 6 mol% (standard deviation derived from the difference in Mn value between the GPC result and the corrected Mn using MHS equation). The calculation of the fraction of SG1 groups was done by considering similar relaxation times between diethyl phosphite and the polymer. • Use of terminal models: The experimental data used to determine the reactivity ratios rMy and rS is given in Table S2a. Calculation of rMy and rS was first done by linearization of the Mayo-Lewis equation 5 (eg. the Fineman-Ross (FR) approach 6 ) and by the Kelen-Tüdös (KT) approach 7 . The latter refines the linearization method by introducing an arbitrary positive constant α to spread the data more evenly so as to give equal weight to all data points. Accordingly, the KT approach minimizes the bias of the FR approach. The appropriate plot of the variables G and H, defined in Equation C1, will provide a linear relationship giving the reactivity ratios as the slope and intercept as shown in

Equation C2. G = (fS,0 / fMy,0) × [(2FS -1) / FS] ; H = (fS,0 / fMy,0) 2 × [(1 -FS) / FS) (C1) G = rS × H -rMy (C2)
Equation C3 shows the additional variables used for the KT plot as illustrated in Equation C4.

𝜂 = G / (α + H) ; ε = H / (α + H) ; α = √𝐻 𝑚𝑎𝑥 𝐻 𝑚𝑖𝑛 (C3) 𝜂 = (rS + rMy / α)ε -rMy / α (C4)
In Equations C3 and C4, Hmax and Hmin refer to the minimum and maximum H values determined from the experimental data. rMy = 1.88 ± 0.12 and rS = 0.25 ± 0.04 (rMy × rS = 0.47 ± 0.11) were obtained via the FR method and rMy = 2.19 ± 0.07 and rS = 0.34 ± 0.19 (rMy × rS = 0.74 ± 0.41) were calculated via the KT method. The errors associated with the experimental data were derived from the standard errors of the slopes from FR and KT plots.

• Use of a NLLS model: Although the KT method is more reliable than the FR method, it is still a linearization.

Statistically, a NLLS fit to the Mayo-Lewis equation is probably the soundest method to determine the desired parameters 8 . Therefore, the reactivity ratios were also determined using a non-linear least-squares fitting of the data.

The commercial software package Matlab R2016a was used to solve the Mayo-Lewis equation 5 for My/S copolymerization (C5). FS = (rS fS,0 2 + fS,0 fMy,0) / (rS fS,0 2 + 2 fS,0 fMy,0 + rMy fMy,0 2 ) (C5)

Using the reactivity ratios determined by the KT method as initial guesses, the statistical fit to the data yielded reactivity ratios rMy = 2.90 ± 1.25 and rS = 0.47 ± 0.27 at a confidence interval of 95 % and with a regression coefficient R 2 = 0.968. (BlocBuilder TM , BB, 99%, purchased from Arkema and used without further purification), N-hydroxy-succinimide (NHS, 98%, purchased from Aldrich and used as received) and N,N'-dicyclohexylcarbodiimide (DCC, 99%, purchased from Aldrich and used as received). Styrene (S, 99%), methyl methacrylate (MMA, 99%), glycidyl methacrylate (GMA, 97%), tert-butyl acrylate (tBuA, 99%) and maleic anhydride (MA, 99%) were obtained from Fisher Scientific and were purified by passing through a column of basic alumina mixed with 5 weight % calcium hydride and then stored in a sealed flask under a head of nitrogen in a refrigerator until needed. The deuterated chloroform (CDCl3, 99.8%) was obtained from Cambridge Isotopes Laboratory. Diethyl phosphite (98%) was purchased from Sigma-Aldrich. The dialysis tube (cutoff 12 000-14 000 Dalton) was purchased from Medicell International Ltd.

• My/GMA copolymerization by NMP. All copolymerizations were done in a 10-mL three-necked round-bottom glass flask equipped with a condenser, a thermal well and a magnetic Teflon stir bar. The flask was placed inside a heating mantle, and the whole setup mounted on top of a magnetic stirrer. All formulations for My/GMA copolymerizations are found in Table 2. For every reaction, the initial molar ratio of monomers and NHS-BB was calculated to give a My/GMA copolymer sample with target number-average molecular weight Mn,theo = (MMyfMy,0 + MGMAfGMA,0)DP = 30 kg.mol -1 at complete overall conversion (X = 1.0) with DP = ([My]0 + [GMA]0) / [NHS-BB]0, the average degree of polymerization. A specific formulation for an initial feed composition of GMA (fGMA,0) equal to 0.50 is given as an example (experiment My/GMA-50, Table 2). NHS-BB (0.144 g, 0.301 mmol), My (4.407 g, 32,350 mmol) and previously purified GMA (4.571 g, 32,156 mmol) were added to the reactor and mixing commenced with the stir bar. A thermocouple was inserted through one of the reactor ports via a thermal well. A mixture of ethylene glycol/reverse osmosis water (30/70 vol%) was circulated via a chiller (Fisher Scientific Isotemp 3016D Digital Refrigerated Bath, 4 o C) through the condenser connected to the second neck of the reactor. The reactor was sealed and a purge of ultra-pure nitrogen was then introduced to the reactor for 30 min. The purge was vented through the reflux condenser. After purging, the reactor was heated via the heating mantle at a rate of about 10 °C.min -1 to 120 °C with continuous nitrogen purge. t = 0 min was taken when the reactor temperature reached 100 o C. Samples were taken periodically until the end of the experiment or until the samples became too viscous to withdraw. For each • Characterization. The (overall) monomer conversion X was determined by proton nuclear magnetic resonance ( 1 H NMR) and calculated from formula 1: X = XMy fMy,0 (+ XYY fYY,0) YY = MMA, GMA, tBuA or MA (1)

where fMy,0 and fYY,0 are the initial molar fractions of My and YY, and XMy and XYY are the individual conversions of My and YY, respectively. XMy and XYY were determined with a Varian NMR Mercury spectrometer ( 1 H NMR, 300 MHz, 32 scans) using CDCl3. My conversion was calculated via the aliphatic protons of the monomer (δ = 2.15-2.30 ppm, 4H), the aliphatic protons of the polymer (δ = 1.90-2.15 ppm, 8H) and the protons of the two methyl groups of monomer and polymer (δ = 1.55-1.75 ppm, 6H). MMA conversion was obtained via the vinyl protons of the monomer (δ = 6.10 and 5.55 ppm, 2H) and the methoxy group of the polymer (δ = 3.63 ppm, 3H). GMA conversion was determined using the peaks of the vinyl protons of the monomer (δ = 6.15 and 5.60 ppm, 2H), the non-equivalent protons of the methylene bonded to the ester oxygen (δ = 4.20-4.50 and 3.65-4.00 ppm, 2H), the methine proton of the oxirane ring (δ = 3.05-3.30 ppm, 1H) and the protons of the methylene of the ring (δ = 2.75-2.90 and 2.55-2.70 ppm, 2H) for both monomer and polymer. tBuA conversion was calculated via the vinyl protons (δ = 5.70 and 6.00 ppm, 2H) and the methyl protons (δ = 1.47 ppm, 9H) of the monomer and the methyl protons of the polymer (δ = 1.42 ppm, 9H). MA conversion was determined by comparing the peaks of the vinylic protons of the monomer (δ = 7.04 ppm, 2H) and the methine protons of the polymer (δ = 2.90-3.55 ppm, 2H). Styrene (S) conversion was determined using the vinyl protons of the monomer (δ = 6.70-6.80, 5.70-5.80 and 5.20-5.30 ppm, 3H) and the aromatic protons of both monomer and polymer (δ = 6.90-7.50 ppm, 5H). The NMR conversion from GMA to 2-hydroxy-3morpholinopropyl methacrylate HMPMA repetitive units was calculated using the methine proton of the oxirane ring (δ = 3.05-3.30 ppm, 1H) and the non-equivalent protons corresponding to the methylene of the ring (δ = 2.75-2.90 and 2.55-2.70 ppm, 2H) for P(GMA) and the methylene protons attached to the morpholine ring (δ = 3.55-3.80 ppm, 4H and 2.55-2.70 ppm, 4H) and the aliphatic protons neighbor of the tertiary amine (δ = 2.30-2.55 ppm, 2H) for P(HMPMA).

The regioselectivity of the My repetitive units in the polymers was determined by comparing the three integrated peaks at δ = 4.70-4.80 ppm (two vinyl protons of 3,4-addition and two vinyl protons of 1,2-addition), δ = 5.00-5.25 ppm (two olefinic protons of 1,4-addition, one olefinic proton of 1,2-addition and one olefinic proton of 3,4-addition) and δ = 5.30-5.50 ppm (one olefinic proton of 1,2-addition). 2,3 A 5 mm diameter Up NMR tube with 800 scans being processed in a 200 MHz Varian Gemini 2000 spectrometer operating at 81 MHz was used to perform 31 P NMR spectroscopy. Polymer (mass of polymer = 0.0528 g with Mn = 11.3 kg.mol -1 for My/GMA-78 and 0.0405 g with Mn = 8.6 kg.mol -1 for My/GMA-37), diethylphosphite as internal reference (0.0047 g and 0.0052 g respectively) and CDCl3 were added to the NMR tube. My/GMA-37 was run under the exact same conditions with only one scan and no dummy scans (ss = 0). A negligible difference (< 2.6%) in integral values was measured between this spectrum and the standard one with multiple scans. It was thereby assumed that diethylphosphite and My-based polymers with moderate Mn < 15 kg.mol -1 relax at the same rate.

The number-average molecular weights (Mn) and the dispersities (Đ = Mw/Mn) were measured using gel permeation chromatography (GPC, Water Breeze, differential refractive index RI 2414 detector, 40 o C) with HPLC grade THF as the mobile phase (flow rate of 0.3 mL.min -1 ). The GPC was equipped with 3 Waters Styragel ® HR columns (HR1 with a molecular weight measurement range of 10 2 -5 X 10 3 g.mol -1 , HR2 with a molecular weight measurement range of 5 X 10 2 -2 X 10 4 g.mol -1 and HR4 with a molecular weight measurement range of 5 X 10 3 -6 X 10 5 g.mol -1 ) and a guard column was used. Mn values were determined by calibration with linear narrow molecular weight distribution PS standards (PSS Polymer Standards Service GmbH, molecular weights ranging from 682 g.mol - 1 to 2,520,000 g.mol -1 ).

Differential scanning calorimetry (DSC, Q2000TM from TA Instruments) was used under N2 atmosphere.

Indium was used as a standard to calibrate temperature while heat flow was calibrated via benzoic acid standard. A temperature range from -90 °C to + 100 °C using three scans per cycle (heat/cool/heat) at a rate of 10 °C.min -1 was typically set. For the Tg measurement, only the second heating run was taken into account to eliminate the thermal history. The reported Tgs were calculated using the inflection method from the change in slope observed in the DSC traces.

DLS measurements were performed with a Malvern Zetasizer Nano equipped with a 532 nm 50 mW green laser. The error is given as the standard deviation from three separate repeats. The samples were prepared by dissolving P(My-b-HMPMA) diblocks in THF and adding dropwise reverse osmosis water while stirring. The suspension was heated at 50 o C for 30 min and the THF was then eliminated by dialysis. To do this, the suspension was put in a dialysis tube which was placed in a beaker filled with reverse osmosis water under magnetic stirring for one or two days. The solution from the beaker was replaced by fresh reverse osmosis water a few times throughout the procedure. The sample in the dialysis bag was eventually isolated (removal of THF confirmed by 1 H NMR) and used for DLS measurements.

• References of the experimental section. • Use of terminal models: Table S1b gives the experimental data used to determine the reactivity ratios rMy and rGMA.

It was assumed that the selected copolymers samples exhibited sufficiently low degree of conversion that the monomer composition was essentially unchanged. Calculation of rMy and rGMA was first done by the Fineman-Ross (FR) approach 1 (rearrangement of the Mayo-Lewis copolymer equation into a linear form 2 ) and by the Kelen-Tüdös (KT) approach 3 . The latter refines the linearization method by introducing an arbitrary positive constant α to spread the data more evenly so as to give equal weight to all data points. Accordingly, the KT approach minimizes the bias of the FR approach. The appropriate plot of the variables G and H, defined in Equation B1, will provide a linear relationship giving the reactivity ratios as the slope and intercept as shown in Equation B2. In Equations B3 and B4, Hmax and Hmin stand for the highest and lowest values of the set of experimental points, respectively. rMy = 0.80 ± 0.31 and rGMA = 0.71 ± 0.15 (rMy × rGMA = 0.57 ± 0.27) were obtained via the FR method and rMy = 0.48 ± 0.12 and rGMA = 0.53 ± 0.18 (rMy × rGMA = 0.25 ± 0.17) were calculated via the KT method.

The errors associated with the experimental data were derived from the standard errors of the slopes from FR and KT plots.

• Use of a NLLS model: Still despite its improvements, the KT method does not overcome the shortcomings of the linear method: use of the differential form and errors that affect both variables. In order to handle rigorously the Mayo-Lewis equation, linearization has to be avoided. From a statistical point of view, a NLLS fit to the Mayo-Lewis equation is probably the soundest method to determine the desired parameters 4 . Therefore, the reactivity ratios were also determined using a non-linear least-squares fitting of the data. The commercial software package Matlab R2016a was used to solve the Mayo-Lewis equation 1 for My/GMA copolymerization (B5). FGMA = (rGMA fGMA,0 2 + fGMA,0 fMy,0) / (rGMA fGMA,0 2 + 2 fGMA,0 fMy,0 + rMy fMy,0 2 ) (B5)

Using the reactivity ratios determined by the KT method as initial guesses, the statistical fit to the data yielded reactivity ratios rMy = 0.49 ± 0.13 and rGMA = 0.50 ± 0.13 at 95 % confidence level and with a regression coefficient R 2 = 0.99 (SSE = 0.0042, RMSE = 0.0244). from Sigma-Aldrich and BlocBuilder TM . Styrene (S, 99%) was obtained from Fisher Scientific and isobornyl methacrylate (IBOMA, Visiomer® Terra) was obtained from Evonik, and both S and IBOMA were purified by passing through a column of basic alumina mixed with 5 wt% calcium hydride and then stored in a sealed flask under a head of nitrogen in a refrigerator until needed. The deuterated chloroform (CDCl3, 99.8%) used as a solvent for nuclear magnetic resonance (NMR) was obtained from Cambridge Isotopes Laboratory. Diethyl phosphite (98%) was purchased from Sigma-Aldrich.

• My/IBOMA Copolymerization by NMP. The copolymerizations, depicted in Figure 1a, were done in a 10 mL three-necked round-bottom glass flask equipped with a vertical reflux condenser (inserted in the middle neck), a thermal well, and a magnetic stir bar. The reactor was placed inside a heating mantle, and the whole setup mounted on top of a magnetic stirrer. Table 1 gives the formulations for the various My/IBOMA copolymerizations studied.

For example, for the experiment My/IBOMA-50 (molar fraction of IBOMA in the initial feed fIBOMA,0 = 0.50), the reactor was sealed with rubber septa after the addition of NHS-BB (0.102 g, 0.213 mmol) and the magnetic stir bar. My (2.381 g, 17.478 mmol) and previously purified IBOMA (3.897 g, 17.529 mmol) were then injected with a disposable needle into the reactor. The initial molar ratio of monomers and NHS-BB was calculated to give theoretically a copolymer sample with target number-average molecular weight Mn,theo = (MMyfMy,0 + MIBOMAfIBOMA,0)DP ~ 30 kg.mol -1 at complete overall conversion with DP = ([My]0 + [IBOMA]0)/[NHS-BB]0 = 164, the average degree of polymerization. A thermocouple connected to the temperature controller was placed inside the thermal well and connected through one of the necks. A mixture of ethylene glycol/distilled water (20/80 vol%) at a temperature of 3 °C was circulated (Fisher Scientific Isotemp 3016D digital refrigerated bath) through the condenser connected to one of the necks of the reactor to prevent any evaporative loss of the monomers. A purge of ultrapure nitrogen was then introduced to the reactor for 30 min to deoxygenate at room temperature the reactants prior to polymerization. The purge was vented through the reflux condenser. After purging, the reactor was heated at a rate of about 10 °C.min -1 to 100 o C with continuous nitrogen purge. The time at which the reactor temperature reached 90 °C was taken arbitrarily as the commencement of the reaction (t = 0 min). Samples were then taken from the reactor periodically by a syringe until the end of the experiments or until the samples became too viscous to withdraw.

Reactions were then stopped by removing the reactor from the heating mantle and letting the contents cool down to room temperature, while under continuous nitrogen purge. For each sample withdrawn during the polymerization, the crude polymer was precipitated with excess methanol. After filtration and recovery, the precipitated polymer was dried at 50 °C under vacuum in the oven overnight to remove unreacted monomers. Samples were analyzed by nuclear magnetic resonance (NMR) and gel permeation chromatography (GPC). At the end of the experiment (t = 340 the overall conversion for My/IBOMA-50 was 52.7% (individual conversions: XMy = 75.2% and XIBOMA = 30.2%) as determined by 1 H NMR, with a number-average molecular-weight Mn,MHS = 9.4 kg mol -1 and a dispersity Đ = 1.28, as determined by GPC. The molar composition of IBOMA in the final copolymer was FIBOMA = 0.43 according to 1 H NMR spectroscopy. The exact same procedure was followed for all My/IBOMA copolymerizations. My/IBOMA copolymer characteristics can be found in Table 2.

• Chain-extension of My/IBOMA copolymers with My, IBOMA and S. All chain-extensions from P(Mygrad-IBOMA) macroinitiators were performed in a very similar setup to that used for the My/IBOMA copolymerizations with the use of a 10-mL reactor. The formulations can be found in Table 3B. As a brief illustration, My/IBOMA-82-IBOMA/My was synthesized using a gradient P(My-grad-IBOMA) copolymer (My/IBOMA-82, FMy = 0.82, Mn,MHS = 13.6 kg.mol -1 , Đ = 1.51, 0.367 g, 0.027 mmol) that was added to the reactor along with toluene (3.24 g), My (0.135 g, 0.991 mmol) and previously purified IBOMA (2.787 g, 12.530 mmol). The reaction medium was mixed and bubbled with ultrapure nitrogen for at least 30 min, then heated to 105 o C and allowed to react for 210 min while maintaining a N2 purge. t = 0 min was selected once the reaction medium reached 90 o C. Samples were drawn periodically via syringe at 0, 60, 120 and 210 min. The samples and the final diblock copolymer were precipitated in excess methanol and allowed to dry overnight in a vacuum oven at 50 o C. My/IBOMA-82-IBOMA/My exhibited Mn,MHS = 23.2 kg.mol -1 , Đ = 1.61 (GPC) and FMy = 0.49 ( 1 H NMR). The results of the chain-extensions from P(Mygrad-IBOMA)s are given in Table 3C.

• Synthesis of Co/IBOMA-My-IBOMA/Co (Co = My or S co-monomer) triblock copolymers by NMP.

A two-step process was implemented in a similar reactor apparatus used for the above-mentioned copolymerizations and chain-extensions: 1) My homopolymerization initiated by PEB-(SG1)2 dialkoxyamine; 2) IBOMA/Co chain-extension from P(My)-(SG1)2 macroinitiator. The entire set of formulations is given in Table 4. For example, the synthesis of S/IBOMA-My-IBOMA/S triblock polymer (Co = S), characterized then mechanically and rheologically, can be detailed. To the 50-mL reactor was added PEB-(SG1)2 difunctional initiator (Mn,GPC = 5.7 kg.mol -1 , Đ = 1.17, 0.912 g, 0.140 mmol) and My (21.005 g, 143.671 mmol) with magnetic stirring (experiment My-52, Table 4A). After sealing the reactor, a purge of nitrogen was applied for 30 min. The reactor was then heated to 120 o C to commence polymerization while continuing the nitrogen flow. t = 0 min was selected once the reaction medium reached 100 o C. At t = 360 min, the polymerization was stopped, the polymer was precipitated in excess methanol and dried overnight in a vacuum oven at 50 o C. 1 H NMR was performed to ensure the complete elimination of methanol and unreacted My in the dry polymer, and to determine the conversion (49.1%). P(My) synthesized exhibited Mn,MHS = 51.7 kg.mol - 4B). The resulting polymer sample (Mn,MHS = 78.1 kg.mol -1 , Đ = 2.41) was precipitated in excess methanol and dried under vacuum at 50 o C overnight. Lastly, fractionation was performed using a methanol (non-solvent) / benzene (solvent) pair (2 cycles) to remove the low molecular weight species. The final fractionated triblock polymer exhibited Mn,MHS = 94.7 kg.mol -1 , Đ = 2.23, FIBOMA = 0.34 and FS = 0.07. Final Co/IBOMA-My-IBOMA/Co copolymer characteristics can be found in Table 4C.

• Characterization. Monomer conversion was determined with a Varian NMR Mercury spectrometer ( 1 H NMR, 300 MHz, 32 scans) using CDCl3. In the case of a copolymerization between two monomers A and B, the overall monomer conversion X was calculated from formula 1: X = XAfA,0 + XBfB,0 ) were characterized in CDCl3 with the addition of diethylphosphite as internal reference (0.0027 g and 0.0072 g respectively). To confirm similar relaxation rates between diethylphosphite and the copolymer samples, My/IBOMA-82 and My/IBOMA-44 were run under the exact same conditions with only one scan and no dummy scans (ss = 0). Slight differences (< 3.2%) in integral values were measured between these spectra and the standard ones with multiple scans. Moderate Mn IBOMA/My copolymers were thus assumed to relax at the same rate compared to that of diethylphosphite.

The number-average molecular weights (Mn,GPC) and the dispersities (Đ = Mw,GPC/Mn,GPC) were measured using gel permeation chromatography (GPC, Water Breeze, differential refractive index RI 2414 detector, 40 o C) with HPLC grade THF as the mobile phase (flow rate of 0.3 mL.min -1 ). The GPC was equipped with 3 Waters Styragel ® HR columns (HR1 with a molecular weight measurement range of 10 2 -5 x 10 3 g.mol -1 , HR2 with a molecular weight measurement range of 5 x 10 2 -2 x 10 4 g.mol -1 and HR4 with a molecular weight measurement range of 5 x 10 3 -6 x 10 5 g.mol -1 ) and a guard column was used. Mn,GPC values were determined by calibration with 10 linear narrow molecular weight distribution P(MMA) standards (Varian Polymer Standards, molecular weights ranging from 875 to 1,677,000 g.mol -1 ). The P(My) and P(IBOMA) contributions to Mn,GPC were converted using Mn,GPC αP(MMA) +1 ] 1/(αP(IBOMA)+1)

Thermogravimetric analysis (TGA) was carried out using a Q500TM from TA Instruments under nitrogen flow at a ramp rate of 15 °C.min -1 . Samples were heated in aluminum pans. Differential scanning calorimetry (DSC, Q2000TM from TA Instruments) was used under N2 atmosphere. Indium was used as a standard to calibrate temperature while heat flow was calibrated via a benzoic acid standard. Three scans per cycle (heat/cool/heat) at a rate of 10 °C.min -1 was set with a temperature range ranging for instance from -90 o C to + 220 o C for the characterization of the triblock copolymers. Only the second heating run was taken into account to eliminate the thermal history. The reported Tgs were calculated using the inflection method from the change in slope observed in the DSC traces.

The stress-strain features of the Co/IBOMA-My-IBOMA/Co triblock polymer (Co = S) were determined using a MTS Insight material testing system with a 5 kN load cell at room temperature and a cross-head speed of 10 mm.min -1 . Dog-bone style tensile specimens (ASTM D638 type V for reference, overall length = 63.5 mm, overall width = 9.53 mm) were prepared by solvent casting. To cast the film, the polymer sample was fully dissolved in dichloromethane (~ 80 wt%) for an hour with continuous stirring using a magnetic stir bar. The solution was cast into a level Teflon Petri dish at room temperature for approximately two days. Then, the drying was completed after placing the film in the oven at 50 o C under vacuum until constant weight was measured. The film samples were cut into typical dog-bone shapes using a sharp blade. Specimens with defects (nicked sides, cracks, air bubbles) on their surface were

• Use of terminal models: Table S1c gives the experimental data used to determine the reactivity ratios rMy and rIBOMA. It was assumed that the selected copolymers samples exhibited sufficiently low degree of conversion that the monomer composition was essentially unchanged. Calculation of rMy and rIBOMA was first done by the Fineman-Ross (FR) approach 1 (rearrangement of the Mayo-Lewis copolymer equation into a linear form 2 ) and by the Kelen-Tüdös (KT) approach 3 . The latter refines the linearization method by introducing an arbitrary positive constant α to spread the data more evenly so as to give equal weight to all data points. Accordingly, the KT approach minimizes the bias of the FR approach. The appropriate plot of the variables G and H, defined in Equation B1, will provide a linear relationship giving the reactivity ratios as the slope and intercept as shown in Equation B2. • Use of a NLLS model: Still despite its improvements, the KT method does not overcome the shortcomings of the linear method: use of the differential form and errors that affect both variables. In order to handle rigorously the Mayo-Lewis equation, linearization has to be avoided. From a statistical point of view, a NLLS fit to the Mayo-Lewis equation is probably the soundest method to determine the desired parameters 4 . Therefore, the reactivity ratios were also determined using a non-linear least-squares fitting of the data. The commercial software package Matlab R2016a was used to solve the Mayo-Lewis equation 1 for My/IBOMA copolymerization (B5).

FIBOMA = (rIBOMA fIBOMA,0 2 + fIBOMA,0 fMy,0) / (rIBOMA fIBOMA,0 2 + 2 fIBOMA,0 fMy,0 + rMy fMy,0 2 ) (B5)

Using the reactivity ratios determined by the KT method as initial guesses, the statistical fit to the data yielded reactivity ratios rMy = 2.07 ± 0.58 and rIBOMA = 0.05 ± 0.08 at 95 % confidence level and with a regression coefficient R 2 = 0.91 (SSE = 0.0174, RMSE = 0.0499). (99%, provided by Arkema). Styrene (S, 99%) was obtained from Fisher Scientific, isoprene (I, 99%) was obtained from Sigma-Aldrich and isobornyl methacrylate (IBOMA, Visiomer ® Terra) was obtained from Evonik, and they were purified by passing through a column of basic alumina mixed with 5 wt% calcium hydride and then stored in a sealed flask under a head of nitrogen in a refrigerator until needed. The deuterated chloroform (CDCl3, 99.8%) used as a solvent for nuclear magnetic resonance (NMR) was obtained from Cambridge Isotopes Laboratory.

• Synthesis of PI-(SG1)2.

In a typical procedure, PEB-(SG1)2 (0.0024 mol.L -1 , Mn = 5.7 kg.mol -1 ) was placed in a 48 mL ChemGlass high-pressure reactor containing a magnetic stir bar and fitted with a plunger valve and thermowell.

Previously purified isoprene (4.928 mol.L -1 ) and pyridine (6.17 mol.L -1 ) were added and the reactor was subjected to three cycles of freeze-thaw degassing, then backfilled with nitrogen. The reactor was placed in an oil bath at 115 o C for 11 h and then cooled to room temperature. The contents of the reactor were transferred to a pre-weighed flask and weighed. Unreacted isoprene was removed in the oven under vacuum at 30 o C and the residual poly(isoprene) and any unreacted PEB-(SG1)2 (mass neglected) were weighed to calculate isoprene conversion gravimetrically. When kinetic studies were led, identical pressure tubes were prepared, degassed and allowed to react for various periods in order to trace the evolution of Mn, Ð and isoprene conversion with time.

• S or IBOMA/S chain-extension from PI-(SG1)2.

The chain-extension of PI-(SG1)2 macroinitiator with S in bulk can be detailed (experiment PS-PI4-PS, similar protocol for IBOMA/S chain-extension). PI-(SG1)2 (0.004 mol.L - 1 ) and previously purified S (5.714 mol.L -1 ) were added in a 25-mL three-necked round-bottom glass flask equipped with a condenser, a thermal well and a magnetic Teflon stir bar. The reactor was sealed, a vigorous stirring was performed and a purge of ultra-pure nitrogen was then introduced to the reactor for 30 min. • Specification sheets of the commercial triblock copolymers.

• Hydrogenation of commercial S-I-S.

Please see the characterization section of the Chapter 4 for the detailed protocol ("Hydrogenation of the synthesized S-I-S triblock copolymers").

• Characterization. The hydrogenation degree (HD) of the TSH-treated S-I-S was determined by 1 H NMR, using the PS aromatic protons (δ = 7.50-6.90 ppm, 5H) and the PI olefin peak (δ = 5.05 ppm, 1H).

The number-average molecular weights (Mn) and the dispersities (Đ = Mw/Mn) were measured using gel permeation chromatography (GPC, Water Breeze, differential refractive index RI 2414 detector, 40 o C) with HPLC grade THF as the mobile phase (flow rate of 0.3 mL.min -1 ). The GPC was equipped with 3 Waters Styragel ® HR columns (HR1 with a molecular weight measurement range of 10 2 -5 x 10 3 g.mol -1 , HR2 with a molecular weight measurement range of 5 x 10 2 -2 x 10 4 g.mol -1 and HR4 with a molecular weight measurement range of 5 x 10 3 -6 x 10 5 g.mol -1 ) and a guard column was used. Mn values were determined by calibration with 10 linear narrow molecular weight distribution PI standards (Polymer Standards Service, molecular weights ranging from 1,030 to 1,040,000 g.mol -1 ).

The stress-strain features of the commercial triblock copolymers were determined using a MTS Insight 50 °C where Tg,specimen refers to the specimen's higher Tg, under a mean pressure of 5000 psi was set. Specimens with defects (nicked sides, cracks, air bubbles) on their surface were discarded and were not used in mechanical testing.

For each specimen, the thickness and the width were measured at five different points along the small center portion with a digital caliper (Marathon, CO 030150F128) and the average thickness and width were used for all calculations.

Each test was considered finished after the complete breakup of the specimen at the narrow section (samples that broke near the grips were discarded). The averaged results were reported using TestWorks 4 software. The Young's modulus was determined as the slope of the stress-strain curve at strains of 0-0.5%.

The mechanical response of SIS and SEBS copolymers to torsional oscillation at 0.15 Hz and 1% strain over

  would like to express my sincere gratitude to my direct research co-supervisor, Prof. Milan Marić, for his continuous and valuable guidance during the PhD project. I am grateful to him for his confidence that helped foster my professional development. I convey also my deep gratitude to my thesis director, Prof. Jean-Francois Gérard, and my thesis co-supervisor, Prof. Pascal Hubert, for their advice and support throughout the entire project. I want to thank Dr. Brigitte Defoort, my industrial supervisor, and Dr. Audrey Menochet, cosupervising the project as well. During the entire project, they have been flexible and tolerant regarding the lines of research, entrusting me a high degree of autonomy, for which I am grateful. I want to also address my thanks to Mr. Philippe Briant, director of the composite materials laboratory at ArianeGroup, who accepted my application and trusted me from the beginning of the doctoral thesis. I would like to acknowledge the generous financial support from ArianeGroup and the French Defense Procurement Agency (Direction Générale de l'Armement, DGA). I am grateful to Mr. Bruno Mortaigne, my DGA supervisor, for his guidance and his encouragement. I also thank the Association Nationale de la Recherche et de la Technologie (ANRT) for its financial contribution. I would like to thank Prof. Jason Robert Tavares, Prof. Martine Dubé and PhD candidate David Brassard as well for their relevant advice during these three years. It was a real pleasure to share our research projects. I want to thank my current and past colleagues here at McGill University, namely Omar, Faezeh, Saeid, Roya, Sharmaine, Marwan, Faraz, Kelly, Simon, Marlon and Kevin for the humorous and enlightening talks. I want to also thank Yara, Kuan, Léa, Laura and Lysandre which I have had the chance to mentor and train. Their hard work and enthusiasm resulted in contributions to valuable projects. I thank the center for Self Assembled Chemical Structures (CSACS) from McGill University for the use of their equipment and specifically Petr Fiurasek for all his help and training. I also thank Mary Sullivan and Noah Macy of Arkema for their aid in obtaining the BlocBuilder alkoxyamine initiator and SG1 used in this doctoral thesis. My thanks also go out to Lou Cusmich, Frank Caporuscio and Ranjan Roy, members of the McGill Chemical Engineering Department technical staff, and Lisa Volpato, the store-keeper, without whom the successful completion of this project would not have been possible.

  Mots clés : élastomère thermoplastique, polymérisation contrôlée par des nitroxydes, 𝛽-myrcène, isoprène. the international association for testing materials ATRP Atom transfer radical polymerization B Butadiene BB BlocBuilder TM (2-([tert-butyl[1-(diethoxyphosphoryl)-2,2-dimethyl propyl] amino]oxy)-2-methylpropionic acid) alkoxyamine tBuA tert-2-[N-tert-butyl-N-(1-diethoxyphosphoryl-2,2-dimethylpropyl)aminoxy]-N-propionyloxysuccinimide NLLS Non-linear least-squares procedure NMP Nitroxide-mediated polymerization

t 1 RÉSUMÉ 2

 12 Reaction time (min) T Temperature ( o C) Tb Boiling point ( o C) Td Decomposition temperature ( o C) Tg Glass transition temperature ( o C) Troom (or RT) Room temperature ( o C) Vm Molar volume (cm 3 .mol -1 ) υP(YY) Volume fraction of P(YY) wYY Mass fraction of YY Xgrav Gravimetric conversion (%) XNMR Conversion determined by proton nuclear magnetic resonance (%) XYY YY Individual conversion (%) z Lattice coordination number Zave Z-Average diameter (nm) Greek symbols δ Nuclear magnetic resonance signal (ppm) δD, δP and δH Hansen solubility parameters (MPa 1/2 ) δYY Solubility parameter of YY (MPa 1/2 ) ∆𝛽YY Difference between the expansion coefficients of the rubbery and glassy states of polymer P(YY) (K -1 ) ∆Gm Gibbs free energy (J) εB Tensile elongation at break (%) ρ Density (g.cm -3 ) σB Tensile strength at break (MPa) σY Yield stress (MPa) σ100 Tensile strength at 100% elongation (MPa) χAB Flory-Huggins interaction parameter for the system A-B SUBSTANTIEL DE LA THÈSE Leader mondial de l'accès à l'espace, ArianeGroup, anciennement Airbus Safran Launchers (ASL), est une coentreprise détenue à parts égales par Airbus et Safran. ArianeGroup conçoit des solutions innovantes et compétitives en matière de systèmes de lancement et d'applications spatiales civiles et militaires pour des clients institutionnels, commerciaux et industriels. Maitre d'oeuvre de la famille de lanceurs européens Ariane et des missiles de la force de dissuasion océanique française, cette société est experte dans l'ensemble du cycle de vie d'un lanceur spatial : conception, développement, production, exploitation et commercialisation via la filiale Arianespace. ArianeGroup produit et exploite le lanceur Ariane 5 et développe le futur lanceur Ariane 6, dont elle est autorité de conception. Les lanceurs fabriqués par ArianeGroup sont formés d'un réservoir de propergol solide appelé corps de propulseur et de viroles appelées jupettes. La liaison entre les jupettes et le réservoir, sollicitée majoritairement par des forces de cisaillement, doit être assurée par un joint souple. Un matériau élastomère répond à cette demande. Les jupettes et le réservoir utilisés actuellement sont des composites thermodurcissables tout comme l'élastomère commercial actuellement employé. L'assemblage entre le réservoir, l'élastomère et les jupettes est réalisé par collage, sur la base de colles polymérisant à chaud, en étuve, positionnées entre l'élastomère et les composites. Cette liaison flexible repose donc sur des matières thermodurcissables, ce qui est problématique d'un point de vue économique principalement. La vulcanisation préliminaire du caoutchouc via l'utilisation d'une presse à haute température et haute pression, l'insertion de résines adhésives et le collage réalisé dans un autoclave géant pendant des cycles de longues durées à hautes températures, constituent des étapes très coûteuses et sont directement liées à l'utilisation de matrices polymère réticulées chimiquement. Un assemblage réservoir / élastomère / jupettes entièrement thermoplastique est ainsi considéré par ArianeGroup pour la prochaine génération de lanceurs. Pour cela, la société envisage de produire le réservoir et les jupettes à base d'un thermoplastique renforcé par des fibres de carbone. De plus, la substitution du caoutchouc thermodurcissable commercial par un élastomère thermoplastique (TPE), par définition ``physiquement`` réticulé, est une priorité. Le passage d'une structure thermodurcissable à une structure thermoplastique présente plusieurs avantages pour ArianeGroup : pas d'étape de vulcanisation du TPE, vraisemblablement bien moins cher que l'élastomère commercial et pouvant être stocké indéfiniment ; réalisation de l'assemblage final par soudage, une technique efficace et rapide, permettant d'obtenir une liaison souple plus robuste que celle réalisée par collage ; nul besoin de colles qui ont une durée de conservation limitée. Alors que des projets parallèles d'ArianeGroup étudient la possibilité d'identifier un TPE commercial permettant de réaliser cette liaison flexible, cette thèse doctorale fut dédiée à la synthèse et la caractérisation d'un TPE original, satisfaisant le cahier des charges industriel. Les spécifications se basent sur les caractéristiques du caoutchouc commercial actuellement utilisé. En matière de propriétés mécaniques, un comportement en contrainte-élongation du TPE semblable aux traditionnels caoutchoucs vulcanisés est exigé à température ambiante. De plus, l'abattement (pourcentage de pertes) des propriétés mécaniques avec la température ne doit pas être conséquent. Par ailleurs, ce nouveau TPE doit être conforme à la réglementation REACH, disposer d'une longue durée de vie et ses propriétés ne doivent pas être affectées lors de contacts avec certains solvants. L'assemblage réservoir / TPE / jupettes mis en place par la suite devra être stable et robuste avec une haute résistance à la rupture en tension uni-axiale, un module de cisaillement à 100% d'allongement significatif ainsi qu'une rupture cohésive de la liaison. Une liaison par soudage étant privilégiée, le TPE doit supporter des cycles de quelques minutes à des températures de l'ordre de 240-280 o C. Toutefois, une caractérisation de la liaison par collage est également considérée, ce qui implique une stabilité de la performance mécanique du TPE à 160-180 o C pendant plusieurs heures. Quel que soit la technique d'assemblage utilisée, le TPE ne doit pas fluer ni se dégrader et avoir ses propriétés inchangées une fois inséré entre les matrices composites. Compte tenu de l'ampleur de ce domaine de recherche consacré à la synthèse de TPE, des axes furent privilégiés en amont du projet. La synthèse de TPE copolymères à blocs de type styrénique par « nitroxide-mediated polymerization » (NMP, innocuité envers l'environnement, absence de toxicité, aisée à mettre en oeuvre et forte expertise du laboratoire à l'Université McGill), polymérisation radicalaire contrôlée basée sur le piégeage réversible des chaînes en croissance par des nitroxydes, fut envisagée. Les TPE considérés correspondent idéalement à des copolymères triblocs formés de la manière suivante : un bloc central souple (poly(β-myrcène) P(My) ou poly(isoprène) PI) conférant les propriétés d'élastomère aux polymères synthétisés (déformation réversible et important allongement à la rupture), deux blocs poly(styrène) PS identiques et rigides venant s'ajouter de part et d'autre du segment souple, augmentant la tenue mécanique dont la résistance à la rupture. Bien que ces triblocs de type S-My-S et S-I-S furent ciblés initialement, des modifications furent entreprises afin d'essayer de répondre au mieux au cahier des charges d'ArianeGroup. Trois grandes parties composent l'étude bibliographique. Dans un premier temps (B), les élastomères vulcanisés traditionnels (matrice réticulée chimiquement) et les TPE de type copolymère à blocs (matrice réticulée physiquement) sont présentés et comparés. Cette section vise à répondre à la question suivante : Quels sont les avantages de remplacer l'élastomère thermodurcissable par un TPE pour l'application ciblée ? La grande majorité des élastomères exploités industriellement sont des thermodurcissables qui obtiennent leurs propriétés d'élastomère et de haute tenue mécanique après vulcanisation (réticulation irréversible des chaînes du polymère en présence de soufre à haute pression et à haute température). Malgré entre autres leur excellente résistance aux produits chimiques et à la température, les élastomères thermodurcissables présentent des inconvénients non négligeables : mélangeage préliminaire avec des agents de réticulation, des stabilisateurs et d'autres additifs;

4 La 6 8 (

 468 d'obtenir une liaison mécanique forte et durable sur des pièces en série. Alors que les temps de cycle pour un procédé par collage peuvent durer plusieur s heures au minimum, quelques minutes au maximum suffisent généralement pour obtenir une soudure robuste. deuxième partie de la revue de la littérature (C) se consacre aux propriétés des TPE copolymères à blocs linéaires de type styrénique, famille de TPE ciblée pour ce projet doctoral. Ce sont généralement des triblocs de type [bloc rigidebloc souplebloc rigide] tels que S-I-S et S-B-S (B = butadiène). Les propriétés de ces copolymères séquencés dépendent principalement de la nature et de la longueur du segment souple et des segments rigides ainsi que du rapport phase souple / phase rigide. La nature des blocs rigides est par exemple primordiale et influencera principalement les propriétés mécaniques et thermiques de ces TPE. L'une des principales faiblesses de ces TPE est que leur température de service maximale est plus ou moins équivalente à la températurede transition vitreuse (Tg) des blocs rigides. Ces derniers étant en grande majorité des segments PS, ces TPE commencent à s´écouler à 95 o C environ. Le remplacement des blocs PS par des blocs plus rigides ayant une Tg plus élevée, tels que le poly(α-méthyl-p-méthyl styrène) P(MMS) (Tg ~ 183 o C) ou le poly(méthacrylate d'isobornyle) P(IBOMA) (Tg ~ 190 o C), permet d'augmenter considérablement la température de service du TPE ainsi que leur résistance en traction aux petites (< 100%) et aux grandes déformations (˃ 100%). L'hydrogénation du segment souple des TPE est également une modification chimique d'intérêt. Il s'agit de saturer le bloc flexible en faisant réagir du dihydrogène avec les doubles liaisons C=C. a) Structure du copolymère tribloc poly(styrène-bloc-butadiène-bloc-styrène) (abrégé S-B-S). b) Représentation schématique d'une morphologie possible des copolymères triblocs de type SBS où les blocs poly(styrène) minoritaires (bleu) forment des sphères rigides au sein d'un domaine poly(butadiène) flexible et continu (noir). c) Cliché obtenu par MET (microscopie éléctronique en transmission) d'un copolymere tribloc S-B-S dans lequel les blocs PS adoptent une morphologie de sphères. L'hydrogénation des triblocs S-I-S et S-B-S mène donc respectivement au S-EP-S (EP = éthylène/ propylène) et au S-EB-S (EB = éthylène/butylène). L'utilisation à haute pression d'un catalyseur homogène (complexe à base de cobalt ou de nickel par exemple) ou hétérogène (palladium sur CaCO3 par exemple) est majoritaire bien que très dangereuse (H2 gazeux, catalyseur pyrophorique, haute pression, haute inflammabilité du mélange…). L'hydrogénation à pression ambiante en utilisant le diimide est une technique beaucoup plus sûre et tout aussi efficace. Elle fut optimisée pour la saturation des homopolymères PI et PB en présence du p-toluènesulfonyl hydrazide, en tant que générateur d'agents diimide, et de tripropylamine afin d'éliminer les réactions secondaires dégradant le polymère. Les poly(diène)s hydrogénés possèdent typiquement des températures de service (+ 10 à + 40 o C) et de dégradation (+ 30 à + 90 o C) supérieures, un module de conservation E' et une résistance en tension plus élevés que les polymères originaux insaturés. La meilleure stabilité thermique et la plus grande tenue mécanique de ces triblocs de type styrénique ayant leur segment souple hydrogéné est donc d'intérêt pour la problématique industrielle. Variations du module de conservation (E', courbe bleue) et du module de perte (E'', courbe rouge) en fonction de la température (1 Hz, 5 o C.min -1 ) pour un copolymère tribloc S-EB-S commercial (Kraton TM G1651 E), possédant 31.5 m% de styrène. À T < Tg,P(EB) (température de transition vitreuse de la séquence éthylène/butylène), état vitreux du S-EB-S ; à Tg,P(EB) < T < Tg,PS, plateau caoutchoutique (température de service de l'élastomère thermoplastique) ; à T > Tg,PS, écoulement visqueux permettant la mise en oeuvre du materiau. Bien que la thèse porte essentiellement sur la synthèse de ces TPE, il est important de connaître des produits similaires disponibles sur le marché, utilisables comme références. S-I-S, S-B-S et sa version hydrogénée S-EB-S constituent la grande majorité des TPE de type styrénique commerciaux produits entre autres par les compagnies Kraton, Dexco, S&E Specialty Polymers et Dynasol. Les S-I-S et S-B-S de Kraton, contenant une minorité de blocs rigides (15-34 m% S), présentent des propriétés en traction uniaxiale satisfaisantes (résistance à la rupture en traction σB = 28-33 MPa, allongement à la rupture en traction ƐB = 880-1300%) à température ambiante bien qu'ils ne semblent pas assez résistants pour supporter des procédés d'assemblage à T ≥ 160 o C (résistance en traction à 300% d'allongement σ300 ≤ 4,1 MPa et indice de fluidité à chaud IFC relativement élevé à 200 o C). Malgré un allongement à la rupture en traction inférieur mais acceptable (≥ 500%), Les S-EB-S de Kraton (29-57 m% S) s'avèrent plus intéressants puisqu'ils disposent d'une plus grande résistance à la traction (σ300 = 4,8-7,9 MPa) et d'une plus grande résistance au fluage (IFC ≤ 5 g/10min à 230 o C). Un S-EB-S fonctionnalisé possédant environ 2 m% d'unités anhydride maléique est également produit chez Kraton. Il pourrait ainsi être testé dans le procédé d'assemblage par soudage (ou par collage) afin de savoir si les groupes polaires encouragent l'adhésion avec les matrices composites considérées pour l'application industrielle. La troisième et dernière partie de l'état de l'art (D) s'intéresse aux différentes techniques de polymérisation utilisées pour produire des TPE copolymères à blocs de type styrénique. La polymérisation anionique fut, dans un premier temps, présentée. Il s'agit aujourd'hui de la seule méthode utilisée industriellement pour la production d'élastomères thermoplastiques à base de styrène et de diènes (S-I-S et S-B-S principalement), par addition séquentielle ou par couplage. La polymérisation cationique peut être également utilisée pour la préparation de triblocs S-iB-S (iB = isobutylene) notamment. Par rapport au procédé anionique, les blocs rigides peuvent être formés à partir de différents types de monomères (α-méthyl styrène, para-chloro styrène…). Bien que ces techniques ioniques soient très efficaces, les conditions expérimentales sont exigeantes (absence d'impuretés, d'eau et d'oxygène) et de nombreuses étapes préliminaires sont nécessaires (dégazage, distillation des solvants et monomères…). De plus, les monomères portant un groupement fonctionnel sont difficilement polymérisables du fait de leur incompatibilité avec les ions et contre-ions formés durant la polymérisation. Par conséquent, la synthèse contrôlée de copolymères à blocs à base de diènes par voie radicalaire, ne nécessitant pas des conditions expérimentales exigeantes et s'adressant à une large gamme de monomères, s'avère prometteuse. Trois techniques majeures de polymérisation radicalaire contrôlée furent développées : par transfert de chaîne réversible par addition-fragmentation (RAFT), par transfert d'atomes (ATRP) et en présence de nitroxydes (NMP). La préparation d'élastomères thermoplastiques via ces trois procédés radicalaires est alors présentée dans cette revue de la littérature, mettant en avant la prédisposition de la technique NMP à polymériser de manière contrôlée les diènes tels que l'isoprène ou le butadiène. L'étude expérimentale débuta en choisissant de synthétiser le poly(β-myrcène) P(My) en tant que segment souple du TPE candidat. Le β-myrcène (My, 7-méthyl-3-méthylène-octa-1,6-diène) est un monoterpène qui peut être extrait d'huiles essentielles de plusieurs plantes de la famille des lauracées par exemple. My fut sélectionné puisqu'il s'agit d'un diène conjugué possédant une structure voisine à celle de l'isoprène, tout en ayant l'avantage d'avoir une température d'évaporation élevée (Tb = 167 o C). Ce monomère naturellement renouvelable et non-volatil est par conséquent adapté à la polymérisation radicalaire contrôlée en utilisant un simple montage à reflux. L'optimisation de la polymérisation radicalaire contrôlée par des radicaux nitroxyde (NMP) de My fut tout d'abord nécessaire afin de synthétiser un élastomère à l'architecture maitrisée (dispersité Đ = Mw/Mn relativement faible, croissance linéaire des chaînes avec la conversion X). La nature de l'amorceur Structure des alcoxyamines de formule générale RONR1R2 (haut) possédant une liaison covalente C-O thermolabile. Son homolyse forme le radical alkyle R • , permettant d'amorcer la polymérisation, et le radical nitroxyde stable R1R2NO • , permettant de contrôler la polymérisation en désactivant de manière réversible le bout de chaîne du macroradical en croissance. Schéma cinétique simplifié de la technique NMP (bas), présentant l'équilibre dynamique entre une faible concentration d'espèces actives et une haute concentration d'espèces dormantes (M = monomère(s), kd = constante de vitesse de l'homolyse de la liaison C-O (dissociation), kc = constante de vitesse de recombinaison entre R(M)n • et le nitroxyde R1R2NO • , kp = constante de vitesse de propagation, kt = constante de vitesse de recombinaison entre radicaux R(M)n • (terminaison irréversible)). BlocBuilder TM commercial abrégé BB et BlocBuilder TM fonctionnalisé avec l'ester succinimidyl, abrégé NHS-BB), le choix du milieu réactionnel (masse, 1,4-dioxane, toluène), la température (T = 110-130 o C), la masse molaire moyenne en nombre ciblée (Mn,theo = 20-50 kg.mol -1 , correspondant au produit entre la masse molaire de My et le rapport molaire entre la concentration de monomère et celle d'amorceur) ainsi que l'influence de l'ajout supplémentaire de radicaux nitroxyde SG1 • (4,9 et 11,7 mol% par rapport à l'amorceur) furent évalués. La polymérisation de My à 120 o C en masse, amorcée par NHS-BB sans radicaux nitroxyde supplémentaires et ciblant Mn,theo ≤ 30 kg.mol -1 , fut considérée optimale. En effet, ces conditions expérimentales donnèrent les résultats les plus satisfaisants avec une évolution linéaire de Mn en fonction de X (Mn/Mn,theo ~ 90% en moyenne) et Đ relativement faible (< 1,30) jusqu'à environ 70% de conversion. Une microstructure de type cis-1,4 fut prédominante (˃ 80 mol%) avec une minorité d'unités 1,2 et 3,4 (5-10 mol% pour chacune). Ce polydiène synthétisé par NMP montra une température de transition vitreuse égale à -77 o C ainsi qu'une température maximale de dégradation (Tdec, 50% de perte de masse) de 385 o C. Afin de vérifier le caractère actif des bouts de chaînes du P(My), idéalement terminé par des unités nitroxyde thermiquement labiles, une extension de chaîne de ce macro-amorceur P(My) avec du styrène fut réalisée dans 50 m% de toluène. L'addition du bloc rigide poly(styrène) PS fut confirmée par chromatographie d'exclusion stérique (SEC, augmentation uniforme de Mn) et par résonance magnétique nucléaire du proton ( 1 H NMR, détection des pics correspondant aux protons aromatiques et aliphatiques du PS). Différents diblocs My-S furent synthétisés (Mn = 33,7-49,0 kg.mol -1 , Đ = 1,33-1,88, fraction molaire en styrène incorporé FS = 0,27-0,62) et leurs propriétés en traction-élongation furent évaluées à température ambiante. Une extensibilité très limitée (εB < 16%) accompagnée d'une grande fragilité (σB < 1,1 MPa) de ces copolymères diblocs furent observées. Deux principales raisons pouvaient expliquer ces propriétés mécaniques en deça des attentes : la longueur de chaîne moyenne relativement faible des diblocs (Mn < 50 kg.mol -1 ) et surtout du segment souple P(My) (Mn ≤ 30,4 kg.mol -1 ) ; une structure dibloc qui ne favorise pas de manière optimale la réticulation physique du réseau polymère puisque les chaînes flexibles ne sont attachées aux chaînes rigides que par un seul bout. Afin d'augmenter le Mn de ce TPE et d'obtenir une structure tribloc S-My-S, permettant à chaque segment souple d'être physiquement confiné entre deux phases rigides, un amorceur bifonctionnel fut préparé. Il s'agit du macro-amorceur téléchélique poly(éthylène-ran-butylène) terminé par des groupes nitroxyde SG1 à ses deux bouts de chaîne et abrégé PEB-(SG1)2. Ce dernier fut synthétisé par estérification de son homologue PEB-(OH)2 terminé par des groupes hydroxyle, suivi d'une addition intermoléculaire radicalaire en utilisant l'alcoxyamine commerciale BlocBuilder TM . La préparation d'un segment souple P(My)-(SG1)2 bien défini, ayant une masse molaire moyenne (Mn = 44,6-59,3 kg.mol -1 ) supérieure à celle des précédents blocs P(My)-SG1 monofonctionnels (Mn ≤ 30,4 kg.mol -1 ), fut optimisée en utilisant l'amorceur PEB-(SG1)2 en masse à 120 o C pendant 5 à 6 h. Suite à l'extension de chaîne de ces macro-amorceurs P(My)-(SG1)2 avec du styrène, des triblocs S-My-S furent obtenus (Mn = 50,0-72,2 kg.mol -1 , Đ = 1,72-2,41, FS = 0,18-0,25). Bien que ces triblocs montrèrent des allongements à la rupture en traction (εB = 164-195%) largement supérieurs à ceux mesurés pour les diblocs My-S, leur résistance mécanique (σB < 0,8 MPa) ne s'est pas améliorée. Par conséquent, ces copolymères à base de β-myrcène ne répondirent pas au cahier des charges industriel. Structures (a) de l'alcoxyamine NHS-BB, (b) du radical nitroxyde SG1 • , des monomères (c) β-myrcène (My), (d) isoprène (I), (e) styrène (S), (f) méthacrylate d'isobornyle (IBOMA) et (g) méthacrylate de glycidyle (GMA) et des polymères (c') cis-1,4-poly(β-myrcène) (cis-1,4-P(My)), (d') cis-1,4poly(isoprène) (cis-1,4-PI), (e') poly(styrène) (PS), (f') poly(méthacrylate d'isobornyle) (P(IBOMA))

L 49 (

 49 'ouverture quantitative des cycles époxyde d'un dibloc P(My-b-GMA) via de la morpholine fut dans cette étude effectuée, afin de produire un dibloc amphiphile P(My-b-HMPMA) capable de s'auto-assembler en micelles dans l'eau (taille moyenne Z, Zave = 120-130 nm à T = 45-65 o C et pour une concentration [P(Myb-HMPMA)] ≤ 0,7 mg.ml -1 ). Les propriétés mécaniques à température ambiante non concluantes des triblocs S-My-S nous poussèrent dans un premier temps à remplacer les blocs PS par des blocs P(IBOMA) (IBOMA = méthacrylate d'isobornyle). P(IBOMA) ayant une Tg (~ 190 o C) bien supérieure à celle du poly(styrène) (~ 100 o C), une plus grande tenue mécanique combinée à une plus haute température de service maximale étaient attendues. Des triblocs à base de My et de IBOMA furent ainsi synthétisés, dans des conditions très similaires à celles utilisées pour la synthèse des triblocs S-My-S (préparation du macro-amorceur P(My) via PEB-(SG1)2 et extension de chaîne avec IBOMA et une petite quantité de My ou S en tant que comonomère). Deux Tg distinctes (Tg,1 = -64 à -58 o C, Tg,2 = + 179 à + 181 o C) furent détectées par calorimétrie différentielle à balayage (DSC), correspondant respectivement aux domaines souple et rigide, ce qui laissa supposer d'une organisation des blocs P(My) et P(IBOMA) en masse. La micro-séparation de phase fut confirmée par microscopie de force atomique (AFM, préparation de films par évaporation lente de dichlorométhane), montrant la présence de potentielles nano-structures à base de sphères/cylindres de P(IBOMA) dans une matrice continue de P(My) (triblocs ayant Mn = 51 kg.mol -1 , Đ = 1,91, FIBOMA = 0,28, voir figure ci-dessous). Alors que les triblocs S-My-S se décomposèrent à T = 435-445 o C, la dégradation thermique des triblocs de type IBOMA-My-IBOMA s'effectua en deux étapes. Une première perte de masse fut visible par analyse thermogravimétrique (TGA) à T = 240-310 o C, correspondant vraisemblablement au détachement des groupes isobornyle de la macromolécule, suivi d'une seconde perte de masse à T = 370-380 o C, correspondant à la scission des chaînes principales. Une température de service maximale d'environ 140 o C fut déterminée par analyse dynamique mécanique (DMA), indiquant une plus grande stabilité de la performance mécanique de ce materiau en comparaison aux TPE à base de styrène qui perdent leurs propriétés, c´est-à-dire s´écoulent, généralement à partir de 90 o C. Enfin, un tribloc de type IBOMA-My-IBOMA (Mn = 95 kg.mol -1 , Đ = 2,23 et FIBOMA = 0,36) fut testé mécaniquement et montra de meilleures propriétés en traction-élongation que les précédents copolymères S-My-S avec σB = 3,9 ± 0,2 MPa (limite d'élasticité à 5,0 ± 0,2 MPa), εB = 490 ± 30 % et un module d'Young E = 2,32 ± 0,3 MPa (figure ci-dessous). Image de phase (2 μm x 2 μm) obtenue par microscopie à force atomique (AFM) en mode Tapping de la surface d'un copolymère tribloc expérimental de type IBOMA-My-IBOMA (Mn = 51 kg.mol -1 , Đ = 1,91, FIBOMA = 0,28). Le domaine sombre représente la phase continue P(My). 12 Bien que la substitution des segments PS par P(IBOMA), plus rigides, donna des résultats encourageants, le développement par NMP de triblocs à base de My ne permit pas de satisfaire les exigences du cahier des charges. La principale lacune fut une résistance à la rupture en traction trop faible (σB ≤ 3,9 MPa). Ces faibles propriétés mécaniques peuvent s'expliquer principalement par la masse molaire entre noeuds d´enchevêtrement élevée du P(My) (Me,P(My) = 22-31 kg.mol -1 ), causant une faible densité d'enchevêtrement de la matrice souple entre les domaines rigides. La difficulté de synthétiser par NMP des triblocs de très haute masse molaire moyenne (Mn < 110 kg.mol -1 dans notre étude) est également une raison qui peut être avancée. Il fut ainsi décidé de remplacer le segment central P(My) par un segment souple poly(isoprène) PI présentant trois avantages majeurs : masse molaire entre enchevêtrements bien plus faible (Me,PI = 4-6 kg.mol -1 << Me,P(My)) ; synthèse et caractérisation de PI et de TPE à base d'isoprène abondamment rapportées dans la littérature ; disponibilité commerciale de TPE de type styrénique à base d'isoprène, pouvant être utilisables comme références. La chimie de l'isoprène par voie radicalaire étant similaire à celle du β-myrcène, l'expérience acquise lors la synthèse du P(My) par NMP fut adoptée. Tout d'abord, la synthèse du poly(isoprène) par NMP fut optimisée pour produire un segment souple de haute masse molaire moyenne et terminé par des groupes nitroxyde en bouts de chaîne Cinq courbes contrainte-élongation (température ambiante, 10 mm.min -1 ) d'un copolymère tribloc expérimental de type IBOMA-My-IBOMA (Mn = 95 kg.mol -1 , Đ = 2,23 et FIBOMA = 0,36). σB = 3,9 ± 0,2 MPa, εB = 490 ± 30 % et E = 2,32 ± 0,3 MPa furent determinés. isoprène) nitroxy-téléchélique, permettant par la suite de pouvoir ajouter deux blocs rigides identiques). La polymérisation de l'isoprène amorcée par PEB-(SG1)2 dans un réacteur pressurisé à 115 o C dans 50 v% de pyridine pendant 13-16 h donna les résultats les plus intéressants. Dans ces conditions et en ciblant une masse molaire théorique relativement élevée à 100% de conversion (Mn,theo ~ 145 kg.mol -1 ), la polymérisation fut contrôlée pour X < 30%, se traduisant par un faible écart entre les masses molaires expérimentales et théoriques et une dispersité inférieure à 1,60. La synthèse de PI ayant Mn = 51-56 kg.mol - 1 et Ɖ = 1,48-1,64 fut effective. De plus, le fractionnement de ce segment flexible par précipitation en utilisant le couple de solvants benzène / méthanol permit d'éliminer les chaînes courtes irréversiblement terminées et conduisit principalement à une plus grande longueur de chaîne moyenne (Mn = 54-61 kg.mol - 1 , Ɖ = 1,47-1,59). Une Tg égale à -60 o C et une température de dégradation maximale Tdec sous atmosphère d'azote voisine de 390 o C furent mesurées pour ces PI, constitués d'environ 80-85 mol% d'addition 1,4 (unités cis-1,4 > 65 mol%). Le caractère « vivant » du macro-amorceur PI permit de produire différents triblocs S-I-S (Mn = 95-109 kg.mol -1 , Đ = 2,11-2,29, FS = 0,30-0,49) par extension de chaîne avec du styrène en masse à 115 o C pendant 1 à 2 h environ. Deux Tg furent détectées par DSC : Tg,1 de -61 à -65 o C (domaine PI souple) et Tg,2 de 102 à 104 o C (domaines PS rigides). Une séparation de phase possiblement cylindrique ou lamellaire à l'échelle du micromètre fut observée par AFM pour un S-I-S ayant FS = 0,30 (voir figure ci-dessous). Tdec ~ 420 o C fut déterminée sous N2 (~ 400 o C sous air) pour un de ces S-I-S, ce qui est satisfaisant vis-à-vis de la problématique industrielle. Alors que le tribloc S-I-S contenant FS = 0,Mn = 103 kg.mol -1 , Đ = 2,11) fut rigide (σB = 10,8 ± 1,2 MPa, E = 2,6 ± 0,7 MPa) mais peu extensible (εB = 21 ± 7 %), le S-I-S ayant FS = 0,38 (Mn = 108 kg.mol -1 , Đ = 2,18) montra une plus grande ductilité (σB = 4,1 ± 0,2 MPa, E = 1,2 ± 0,6 MPa et εB = 380 ± 60 %). Comme attendu, leur température de service maximale se limita généralement à 90-100 o C, comme indiqué par DMA, par la chute du module de conservation à T ˃ Tg,PS. Par conséquent, bien qu'ayant des propriétés mécaniques supérieures à leurs homologues à base de My, ces triblocs S-I-S manquèrent de résistance à la traction à T = 100-160 o C. Image de phase (2 μm x 2 μm) obtenue par microscopie à force atomique (AFM) en mode Tapping de la surface d'un copolymère tribloc expérimental S-I-S (Mn = 94.6 kg.mol -1 , Đ = 2.29, FS = 0,30). Le domaine sombre représente la phase continue PI. 14 Afin de combler ces lacunes, un tribloc de type IBOMA-I-IBOMA fut développé (Mn = 94 kg.mol -1 , Đ = 1,76 et FIBOMA = 0,35) de la même manière. Bien que sa morphologie à l'échelle microscopique ne fût pas clairement identifiée par AFM, la température de transition vitreuse de la phase souple fut clairement détectée par DSC à -63 o C ainsi que celle du domaine rigide à environ 182 o C. Comme observé préalablement dans le cas des copolymères à base de β-myrcène et de méthacrylate d´isobornyle, la décomposition de ce tribloc de type IBOMA-I-IBOMA se réalisa en deux principales étapes (Td,1 = 265-345 o C, perte du cycle isobornyle ; Td,2 = 345-500 o C, dégradation de la chaîne principale) sous atmosphère inerte. Ce tribloc fut ensuite analysé quant à ses propriétés mécaniques et rhéologiques. Une importante extensibilité (εB = 1360 ± 210 %) associée à une bonne résistance mécanique (σB = 11,4 ± 0,6 MPa) furent mesurées (figure ci-dessous). Ces résultats peuvent s'expliquer par l´existence d´enchevêtrements dans les blocs souples (Mn,PI = 53 kg.mol -1 ~ 10 Me,PI) combinée à la rigidité des unités IBOMA. La réponse mécanique dynamique de ce tribloc en torsion sur une large plage de température permit de mettre en évidence une température de service maximale de l'ordre de 160 o C. Au-delà de cette température, une chute du module de conservation intervient (figure ci-dessous). Ce tribloc à base d'isoprène et de méthacrylate d'isobornyle fut ainsi le copolymère synthétisé par NMP le plus concluant, compte tenu de son comportement en traction uni-axiale et de sa potentielle haute température de service maximale. La dégradation thermique précoce de ce dernier à partir d'environ 250 o C reste toutefois problématique et son comportement mécanique à haute température ne fut pas évalué. Quatre courbes contrainte-élongation (température ambiante, 10 mm.min -1 ) d'un copolymère tribloc expérimental de type IBOMA-I-IBOMA (Mn = 94 kg.mol -1 , Đ = 1.76 et FIBOMA = 0,35). σB = 11,4 ± 0,6 MPa, εB = 1360 ± 210 % et E = 3,5 ± 1,1 MPa furent determinés. Dans le but d'augmenter principalement la résistance en traction et la température de dégradation du TPE candidat, des essais d'hydrogénation furent effectués sur des triblocs expérimentaux et de référence à base d'isoprène et de styrène. Il fut question de saturer le segment central PI. Cette modification chimique fut réalisée dans des conditions ``douces`` à pression ambiante, sans catalyseur ni dihydrogène. Spécifiquement, les S-I-S furent dissous dans du solvant p-xylène et chauffés à 125 o C pendant 3 h avec du p-toluène sulfonylhydrazide, générateur de l'agent diimide, et du tributyle amine. La conversion des unités isoprène en unités éthylène/propylène fut de 67-77 mol% ( 1 H NMR). Alors que la stabilité thermique des copolymères hydrogénés ne fut que légèrement améliorée sous air, en comparaison aux homologues insaturés, une augmentation plus importante de la température de dégradation fut constatée sous atmosphère inerte (N2). Les propriétés mécaniques à température ambiante furent également affectées par l'hydrogénation du segment souple. σB et E furent augmentés de 30-55% et de 300-440% respectivement, indiquant une plus grande résistance en traction et dureté de ces S-I-S partiellement saturés. Toutefois, ces derniers montrèrent un allongement à la rupture en traction plus faible (par exemple, εB diminuant de 4700% à 3300% environ pour un S-I-S de référence après hydrogénation). Des analyses de spectrométrie mécanique dynamique ont été menées afin de déterminer l'influence de cette réaction chimique sur le comportement thermo-mécanique de ces matériaux. Le module de conservation E' d'un matériau de Variations du module de conservation (noir) et du facteur d'amortissement (gris) en fonction de la température (0.15 Hz, 1% d'élongation, 5 o C.min -1 ) d'un copolymère tribloc expérimental de type IBOMA-I-IBOMA (Mn = 94 kg.mol -1 , Đ = 1.76 et FIBOMA = 0,35).

  montra des caractéristiques mécaniques à température ambiante très intéressantes avec σB = 25,0 ± 2,1 MPa, σ100 = 4,6 ± 0,4 MPa, εB = 730 ± 40 % et E = 0,6 ± 0,1 MPa. Il commença à se dégrader (5% de perte de masse) à 260 o C sous air et à 340 o C sous N2. Des essais TGA en condition isotherme indiquèrent que ce tribloc ne se décompose pas à T ≤ 275 o C sous atmosphère inerte. Toutefois, une diminution importante de son module de conservation G' à T ˃ 60 o C fut constatée, laissant supposer une perte non négligeable de ses propriétés mécaniques à haute température. Différents copolymères à blocs de type styrénique furent ainsi préparés par NMP dans des conditions expérimentales simples, au cours de ce projet de recherche doctoral. L´influence de la nature du segment élastomère fut évaluée : les triblocs S-My-S montrèrent un comportement fragile en traction (σB < 1 MPa, εB < 200%) alors que de meilleures propriétés mécaniques furent déterminées pour les triblocs S-I-S constitués d´une phase souple PI bien plus enchevêtrée (σB = 2,8-4,1 MPa, εB = 380-450 %). L´influence de la nature des blocs rigides fut également étudiée : le remplacement des blocs PS traditionnels par des blocs P(IBOMA) plus rigides permit d´augmenter considérablement la résistance en traction et l´extensibilité du copolymère tribloc. εB ~ 1350% et une température de service maximale d´environ 150 o C furent ainsi mesurés pour le copolymère tribloc de type IBOMA-I-IBOMA. Son insuffisante résistance mécanique (σB ~ 11 MPa) associée à son instabilité thermique (début de la dégradation à 220 o C sous air et 265 o C sous azote) furent toutefois préjudiciables pour l´application industrielle.

( 1 )( 2 )( 3 )( 4 )( 5 )( 6 )( 7 )( 8 )

 12345678 et remoulés.``Le développement de matériaux thermoplastiques ayant des propriétés élastiques débuta dans les années 1930 avec la plastification du polychlorure de vinyle (PVC) par la compagnie B. F. Goodrich 1 . Cette invention mena par la suite à la préparation de mélanges à base de PVC et de caoutchouc du type copolymère butadiène-acrylonitrile (NBR) 2,3 . Cela permit d'obtenir un matériau intermédiaire entre le PVC plastifié et les élastomères vulcanisés traditionnels. Les mélanges PVC / NBR peuvent ainsi être considérés comme les précurseurs des TPE, comme nous les connaissons aujourd'hui. La découverte de la réaction de polyaddition en utilisant du diisocyanate en 1937 constitua une avancée majeure dans le développement des TPE4 . Cette technique de polymérisation fut employée chez Dupont et ICI afin de synthétiser des fibres de polyuréthane ayant des propriétés similaires à celles d´élastomères vulcanisés 5-7 . Cela conduit par la suite à la synthèse d'élastomères de type copolyester, ayant une résistance en traction supérieure aux caoutchoucs vulcanisés et une reprise élastique aux temps courts 8,9 . Ces matériaux furent ainsi considérés comme étant les premiers TPE.La consommation mondiale de TPE devrait progresser de 4,2 millions de tonnes en 2015 à environ cinq millions et demi de tonnes en 202010 . L'intérêt des TPE s'explique par leur capacité à être mis en oeuvre comme des matériaux thermoplastiques tout en ayant des propriétés physiques et mécaniques aussi intéressantes que celles des élastomères thermodurcissables. Ils permettent donc une grande flexibilité dans leur procédé de fabrication et de mise en oeuvre. Aujourd'hui, le secteur industriel utilise de plus en plus les TPE puisqu'ils s'adaptent directement aux machines traditionnelles mise en service pour les thermoplastiques. Qu'il soit naturel ou synthétique, l'élastomère conventionnel est un thermodurcissable qui est vulcanisé. La vulcanisation est une réaction réalisée durant le moulage ou l'extrusion et permettant la réticulation chimique et permanente du caoutchouc. Du fait de cette réaction, l'élastomère ne peut plus être remis en forme. Bien que la durée de la réaction de vulcanisation dépende de nombreux paramètres, il faut compter entre quelques minutes et plusieurs heures afin de conduire à un réseau tridimensionnel efficace. En comparaison, les procédés de moulage et d'extrusion utilisés pour les TPE, ne nécessitant pas d'étape de réticulation, peuvent être réalisés très rapidement. Une vingtaine de secondes peut être suffisant, ce qui se traduit par des profits non-négligeables. De plus, la possibilité de recycler les TPE est un atout considérable. Les déchets, rebuts ou articles en fin de vie à base de TPE peuvent être aisément remis en oeuvre alors que les élastomères thermodurcissables en fin de cycle se retrouvent bien souvent dans les décharges. C'est dans ce contexte que la problématique industrielle étudiée pour cette thèse prend toute sa place. Leader mondial de l'accès à l'espace, ArianeGroup conçoit des solutions innovantes et compétitives en matière de système de lancement et d'applications spatiales civiles et militaires. Ses activités couvrent l'ensemble du cycle de vie d'un lanceur spatial : conception, développement, production, exploitation et commercialisation. ArianeGroup produit et exploite le lanceur Ariane 5, le plus fiable du marché commercial, et développe le futur lanceur Ariane 6. Le projet de recherche proposé et financé par ArianeGroup pour cette thèse doctorale s'inscrit parfaitement dans une logique de remplacement des matériaux thermodurcissables traditionnels par des thermoplastiques. Cela représente un enjeu majeur de l'industrie aéronautique. L'une des activités clés d'ArianeGroup est consacrée à la fabrication de lanceurs. Ces derniers sont constitués d'un corps de propulseur (réservoir de propergol solide) et de viroles (appelées jupettes). La liaison entre les jupettes et le réservoir, sollicitée principalement en cisaillement, est assurée par un élastomère. Aujourd'hui, les jupettes et le réservoir sont des composites thermodurcissables. L'élastomère est également un polymère thermodurcissable. La liaison flexible entre ces structures est réalisée par collage, sur la base d'une colle spécifique polymérisant à chaud, en étuve. Un tel assemblage reposant sur des matières thermodurcissables est problématique, d'un point de vue économique principalement. La vulcanisation du caoutchouc, l'insertion de la colle et le collage réalisé dans un autoclave pendant des cycles de longues durées, constituent par exemple des étapes très coûteuses et sont directement liées à l'utilisation de matrices polymère réticulées chimiquement. ArianeGroup envisage alors le développement d'un assemblage réservoir / élastomère / jupettes entièrement thermoplastique. En effet, pour la génération future de ses lanceurs, ArianeGroup souhaiterait fabriquer le réservoir et les jupettes en composite thermoplastique et remplacer l'élastomère vulcanisé assurant la flexibilité de la liaison par un TPE. Ce changement total de composites et d'élastomères thermodurcissables à des composites et des élastomères thermoplastiques permettrait de réaliser l'assemblage final par soudage, une technique efficace et économique. À la différence des autres procédés d'assemblage, le soudage ne s'applique qu'aux thermoplastiques. Seuls ceux-ci possèdent des caractéristiques rhéologiques suffisantes, lors de la fusion (pour les semi-cristallins) ou du ramollissement (pour les amorphes), pour permettre une interpénétration des couches superficielles nécessaire à leur liaison. La réalisation d'une liaison par soudage permet d'obtenir d'un point de vue industriel une robustesse supérieure aux technologies de collage. La moindre sensibilité à l'environnement et à la pollution ainsi qu'un cycle de réalisation plus rapide sont des atouts fondamentaux. En raison des difficultés à identifier un élastomère thermoplastique commercial satisfaisant le cahier des charges fixé par ArianeGroup, cette thèse avait pour objectif de développer en laboratoire un élastomère thermoplastique (TPE) permettant de rivaliser avec l'actuel caoutchouc thermodurcissable. La capacité de ce nouvel élastomère à se souder directement aux composites thermoplastiques envisagés ou indirectement via l'insertion d'un matériau thermoplastique intermédiaire constituait une exigence industrielle importante. La polymérisation radicalaire contrôlée par des nitroxydes (NMP) fut sélectionnée afin de produire un TPE de type styrénique. Le poly(β-myrcène) P(My), un élastomère biosourcé de la famille des polyterpènes, fut d'abord envisagé en tant que segment souple du TPE avant d'être remplacé par le poly(isoprène) PI. Ce manuscrit se divise en trois grandes parties : (I) état de l'art ; (II) première partie expérimentale consacrée à la synthèse de copolymères à base de β-myrcène (My) par NMP ; (III) seconde partie expérimentale qui se concentre sur la synthèse par NMP et la caractérisation de copolymères à blocs à base d'isoprène. L'état de l'art (I) introduit dans un premier temps les élastomères thermodurcissables et thermoplastiques puis se focalise plus précisément sur les TPE copolymères à blocs de type styrénique envisagés dans cette étude. La deuxième section (II) comporte trois chapitres. Le premier chapitre évoque la synthèse de copolymères β-myrcène/styrène (My/S) par NMP, permettant notamment d'obtenir des copolymères triblocs S-My-S. Le deuxième chapitre présente la fonctionnalisation du poly(β-myrcène) P(My) par la copolymérisation radicalaire contrôlée par des radicaux nitroxyde du β-myrcène avec le méthacrylate de glycidyle, un monomère vinylique portant un groupe époxyde. L'incorporation d'unités réactives au sein du TPE avait pour but d'encourager le soudage ultérieur de l'élastomère avec les composites thermoplastiques relativement polaires. Le troisième chapitre se concentre sur la synthèse contrôlée de copolymères My/IBOMA (IBOMA = méthacrylate d'isobornyle) par NMP. L'objectif était ici de remplacer les segments poly(styrène) PS par des blocs P(IBOMA) plus rigides ayant une température de transition vitreuse plus élevée et permettant ainsi d'augmenter la résistance à la rupture et d'élargir la température de service du TPE candidat. La troisième grande partie (III) de ce manuscrit comprend deux chapitres finaux. Le quatrième chapitre est consacré à la synthèse de triblocs à base d'isoprène par NMP alors que le cinquième chapitre se focalise sur la caractérisation de TPE de référence de type styrénique. L'hydrogénation partielle du segment souple poly(isoprène) de copolymères triblocs expérimentaux et de référence est également présentée dans ces deux derniers chapitres. Semon W. L., US Patent 1 929 453, B. F. Goodrich Co., 1933. Henderson D. E., US Patent 2 330 353, B. F. Goodrich Co., 1943. Wolfe J. R. In Thermoplastic Elastomers: A Comprehensive Review; Legge N. R., Holden G., Schroeder H. E., eds.; Hanser Publishers: Munich, 1987; Chapter 6. Bayer O., Siefken W., Rinke H., Orthner R., Schild H., German Patent 738 981, I. G. Farben, 1937. Christ A. E., Hanford W. E., US Patent 2 333 639, DuPont, 1940. British Patents 580 524 (1941) and 574 134 (1942), ICI Ltd. Hanford W. E., Holmes D. F., US Patent 2 284 896, DuPont, 1942. Snyder M. D., US Patent 2 632 031, DuPont, 1952. (9) US Patent 2 629 873, DuPont, 1954. (10) Thermoplastic Elastomers Market Size, Industry Share Report, Global Market Insights, 2016.

Figure 1 .

 1 Figure 1. Shearing forces (yellow arrows) generated through the flexible junction due to tank and skirt deformations (red arrows) caused by the weight of the launcher (left) and due to tank deformations in diameter (middle) and height (right, top view).

Figure 2 .

 2 Figure 2. Schematic of the current assembly performed by ArianeGroup using a specific adhesive to implement the adhesive bonding. The illustration on the right shows rubber/adhesive and adhesive/thermoset composite interfaces where covalent bonding occurred (the black solid circles represent crosslinking points whereas the grey solid circles correspond to the reinforcement components of the thermoset composite).

  The failure of a resistance-welded joint can be dominated by one of the following failure modes: interfacial failure, cohesive failure of the heating element, tearing of the heating element, and tearing of the laminate. ArianeGroup would like a cohesive fracture at the TPE / thermoplastic composite interface where the failure occurs through the heating element.

Figure 3 .

 3 Figure 3. Schematic of a resistance welding process for a lap shear coupon.

Figure 4 .

 4 Figure 4. The new thermoplastic-based tank (grey) / TPE (blue) / skirt (grey) assembly could be welded directly using a specific thermoplastic composite as the resistance heating element (a) or indirectly via the use of in-between thermoplastic composite layers (red) used as resistance heating elements (b). The illustrations on the right show (a) TPE / thermoplastic composite interface where interdifusion of the chains occured (TPE arbitrarily represented by spherical hard domains in a continuous soft matrix; the grey solid circles correspond to the reinforcement components of the thermoplastic composite) and (b) TPE / intermediate composite and intermediate composite / thermoplastic composite interfaces.

h•

  Conclusion:We proposed to perform NMP in bulk or in organic solvents using a difunctional initiator to synthesize linear styrenic triblock copolymers, exhibiting ideally a spherical nano-phase separation. Poly(β-myrcene) or poly(isoprene) were considered as constituting the elastomeric middle block (60-90 mol% of the TPE) whereas two outer poly(styrene) blocks were envisioned to provide rigidity (10-40 mol% of the TPE), as depicted in Figure5below. Four chemical modifications were planned to improve the candidate TPE properties: functionalization of the poly(diene) segment by adding glycidyl methacrylate units (encouraging the joining process); replacement of PS blocks by higher Tg hard blocks (increasing the tensile stress and broadening the service temperature); nitroxide removal (thermal stability enhancement);

Figure 5 .

 5 Figure 5. Summary outline of the linear styrenic triblock thermoplastic elastomer envisioned for this thesis. The difunctional initiator Di-I used corresponds to a dialkoxyamine. According to NMP mechanism when all the suitable parameters are met, nitroxide groups (N) are attached to the two TPE chain ends. The schematic at the bottom represents the expected nano-phase separation of the TPE, where PS rigid spheres act as physical crosslinks.

Figure 6 .

 6 Figure 6. Summary outline of the different ways considered in this thesis to improve TPE properties.

Following-

  these introductive sections (pages 1-41), three main parts constitute this thesis. The first part (pages43-92) reviews the literature relevant to the industrial problematic with an emphasis on the thermoplastic elastomers targeted for this research project and their synthesis. The second and third parts present, in the form of chapters, the experimental work led throughout the doctoral thesis in chronological order. To alleviate the chapters, the respective appendices containing the experimental sections along with the supporting information are given at the end of the PhD manuscript (pages 269-328). Please note that at the beginning of each chapter appears a preface, specially added for the PhD manuscript, to make the transition from the previous chapter to the new chapter. It explains concisely the reasons that prompted us to lead the study in question. It must be noted that the references are gathered at the end of each chapter (or part, in the case of the literature review).The second part (pages 95-207) clusters all the experimental studies based on the use of poly(βmyrcene) P(My) as the soft block of the candidate TPE. Whatever the status of the chapters (published, submitted or not submitted), they are all written in a paper-based format:-Chapter 1 (pages 97-139), a manuscript published in the journal Macromolecules, details the optimized synthesis by nitroxide-mediated polymerization of well-defined P(My) leading to the achievement of β-myrcene / styrene (My/S) block copolymers (first year of the PhD, 2015). It sums up the preliminary attempts to achieve satisfactory My-based thermoplastic elastomers (TPE), using poly(styrene) PS hard blocks. Please note that the section 1.4.3 regarding the synthesis of S-My-S triblock copolymers was not present in the original publication and was added to the PhD manuscript.-Chapter 2 (pages 141-172), a manuscript published in the Journal of Polymer Science, Part A: Polymer Chemistry, presents the well-controlled nitroxide-mediated copolymerization of My with glycidyl methacrylate (GMA). Led in parallel with the experiments presented in Chapter 1, this study aimed to functionalize the soft middle block (herein P(My)) of the candidate TPE in order subsequently to encourage the adhesive bonding or welding process with the relatively polar thermoplastic composite. Chapter 3 (pages 175-207), a manuscript which will be soon submitted to the journal Polymer Chemistry, introduces the synthesis of β-myrcene / isobornyl methacrylate (My/IBOMA) copolymers by NMP (second year of the PhD, 2016). The goal was to produce enhanced My-based block copolymers by replacing PS segments by P(IBOMA) hard blocks exhibiting a higher glass transition temperature.Since My-based copolymers were not satisfying ArianeGroup requirements, the substitution of P(My) by poly(isoprene) (PI) as the rubbery component was undertaken. The third part (pages 209-249), corresponding to the work led during the second half of the doctoral thesis, is devoted to the synthesis of isoprene-based triblock copolymers as well as the characterization of commercial thermoplastic elastomers mostly based on isoprene. Unlike the previous chapters, which were more academic, these sections are more technical and industrially oriented. They are thereby purposely more concise and results-based. Two chapters are introduced: -Chapter 4 (pages 211-237) is dedicated to the achievement of linear isoprene-based triblock copolymers by NMP. S-I-S triblocks exhibiting various molar fractions of styrene unit were first synthesized and characterized. The incomplete hydrogenation of the PI mid-segment was performed. PS blocks were then replaced by P(IBOMA) blocks to produce isoprene-based triblocks exhibiting higher upper service temperature, tensile strength at break and elongation at break. -Chapter 5 (pages 239-249) details the characterization and the hydrogenation of some commercial styrenic block copolymers, allowing comparative study with the experimental NMP-based block copolymers synthesized. Finally, a general conclusion summing up the work conducted over this research project and the recommendations for future work are given (pages 251-267).

Figure 7 .

 7 Figure 7. Schematic of crosslinked (in red) elastomer macromolecules (a) unstretched (random coil structure) and (b) under uniaxial tensile stress.

Figure 8 .

 8 Figure 8. Brief synthetic route of the vulcanization of cis-1,4-poly(isoprene) using sulfur curing agent S8 generating disulphide covalent bonds between the flexible macromolecular chains.

Figure 10 .

 10 Figure 10. (a) General trends of flexural modulus (stiffness) of a typical block copolymer TPE in dependence of temperature and (b) schematic of the morphology of a block copolymer TPE.

  three factors: the overall degree of polymerization N; the architectural constraints and the composition of the diblock (overall volume fraction of the segment A); the A-B segment-segment Flory-Huggins interaction parameter χAB 25 . While the first two parameters are controlled by the polymerization conditions and affect the translational and configurational entropy, χAB is determined by the selection of the monomers A and B. χAB has a temperature dependence given by: χAB = α / T + β

Figure 11 .

 11 Figure 11. Schematic illustration of the thermally reversible order-disorder transition for a symmetric A-B diblock copolymer (FA = 0.50). Accordingly, a lamellar morphology represents the ordered phase at T < TODT.

Figure 12 .

 12 Figure 12. Theoretical phase diagram for a linear A-B block copolymer as a function of χABN and volume fraction of block A (FA). The schematic diagrams of the ordered lamellar, gyroid, cylindrical and spherical morphologies observed in A-B diblock copolymer systems, showing the domain occupied by the short minority block (the remaining space is filled by the majority block), are given as well.

60 C

 60

Figure 13 .

 13 Figure 13. Hard blocks (blue background) A and soft blocks (orange background) B of conventional A-B-A styrenic triblock copolymers. Their Tg is approximated since it depends on the regio-and stereoselectivity as well as the average molecular weight of the segments.

Figure 14 .

 14 Figure 14. Stress-strain curves for (a) SBS triblock copolymers (Mn = 18-64-18 kg.mol -1 ) and (b) SIS triblock copolymers (Mn = 20-60-20 kg.mol -1 ) both prepared anionically using dilithium adducts of isoprene as initiators (5 inches/min cross-head speed, room temperature). Numbers on curve are the styrene contents of the polymers 29 .

Figure 15 . 66 C. 1 . 4 .

 156614 Figure 15. Transmission electron micrographs (TEM) after compression molding of P(I-b-S) diblock copolymers exhibiting a fraction of styrene of 86 wt% (left, spherical morphology), 78 wt% (middle, cylindrical morphology) and 59 wt% (right, lamellar morphology). The bright phase corresponds to the PS phase and the dark one to the PI phase 36 .

Table 5 . 68 Figure 16 .

 56816 Figure 16. (a) Elastic moduli (G') of PS-PI-PS (Mn = 23-164-23 kg.mol -1 , black solid curve) versus P(MMS)-P(My)-P(MMS) (Mn = 36-168-36 kg.mol -1 , green solid curve) with ω = 1 rad.s -1 , an initial strain of 10 % and a 3 o C.min -1 temperature range rate 40 and (b) G' of P(IBOMA)-PB-P(IBOMA) (Mn = 20-55-20 kg.mol -1 , solid curve (-)), hydrogenated P(MMA)-PB-P(MMA) (Mn = 25-80-25 kg.mol -1 , discontinuous curve (-• -)) and P(MMA)-PB-P(MMA) (Mn = 25-80-25 kg.mol -1 , dashed curve (---)) at 1 Hz 42 .

Figure 17 .

 17 Figure 17. Stress-strain curves of PS-PI-PS (Mn = 20-87-20 kg.mol -1 , open square) and P(MS)-PI-P(MS) (Mn = 21-85-21 kg.mol -1 , open circles) cast films (tetrahydrofuran used) at 20 o C and at 5 cm.min -

  The most largely performed of these techniques relies on the chemistry of the reactive intermediate diimide (HN=NH), a reactive species that can transfer hydrogen to alkene or alkyne functional groups. These methods are in wide use in laboratory settings, because they allow for efficient saturation of olefins without the need for expensive high-pressure reactors and are tolerant of a variety of functional groups. A stoichiometric excess of the diimide must be provided to result in a complete saturation. The required excess is due to disproportion of diimide. The diimide hydrogenation is briefly showcased in Figure19. Hydrogenation of poly(diene)s with the diimide-based reagent termed p-toluenesulfonylhydrazide (TSH) has been reported, including specifically the saturation of PI and PI block copolymers and PB. It was found by the researchers that acid by-products from the TSH degradation reaction can react with incompletely hydrogenated polymers, generating chain cleavage and the incorporation of a fragment of the TSH onto the polymer chain. The use of amine tripropyl amine (TPA) associated with the TSH reagent in refluxing o-xylene was found to suppress the contribution of side reactions during the saturation of PB and its copolymers. Even though it improved the integrity of hydrogenated PI, it did not give complete saturation in one treatment.

  hydrogenated and exhibiting Tg = -52.0 o C whereas its original counterpart had a lower Tg = -61.0 o C 55 .

Figure 19 .

 19 Figure 19. Use of diimide to hydrogenate poly(diene) using typical experimental conditions. To avoid the protonation of the unsaturated poly(diene) during the diimide hydrogenation, the reaction must be run in the presence of a high boiling point tertiary amine (tri-n-propylamine and tributylamine are commonly used).

  ) for various poly(diene)s and their hydrogenated counterparts (DH = Degree of hydrogenation) 52,53 .It is of interest to note that hydrogenated styrenic TPE generally show a higher upper service temperature as seen in Figure20b. While G' of S-SB-S triblock stars to drop at T ~ 90 o C, fully hydrogenated S-SB-S still has a high G' at 140 o C. This can be explained by the much stronger phase separation of hydrogenated TPE with a greater incompatibility, a greater difference in solubility parameters between the hydrogenated soft midblock and PS hard blocks.

Figure 20 .

 20 Figure 20. Temperature dependence of storage modulus (E' or G') for (a) natural rubber (NR) and hydrogenated natural rubber (HNR) exhibiting different hydrogenation degrees (HDs) 52 and for (b)

Figure 21 .

 21 Figure 21. Stress-strain curves for hydrogenated S-SB-S in (a) the high-strain regions and (b) low-strain region, and schematic illustrations of (c) ambiguous phase separation for weakly hydrogenated S-SB-S and (d) strong phase separation for highly hydrogenated S-SB-S 56 .

Table 7 .

 7 Main producers and trade names of styrenic TPE and the respective industrial applications 58,59,60 . (a) SIBS = Styrene-isoprene/butadiene-styrene. (b) SEBS with grafted maleic anhydride. (c) SEPS = Styrene-ethylene/propylene-styrene. (d) synthetic poly(isoprene) products. (e)

  (a) Typical properties determined on film cast from toluene solution. (b) SEBS samples correspond to fully hydrogenated SBS TPE. (c) MA = Maleic anhydride.SEBS Kraton ® are the strongest, the most highly dilutive and the most compatible with polyolefins and mineral oils of all the styrenic TPE. They are second generation styrenic block copolymers with a hydrogenated midblock (ethylene/butylene copolymer) and are intended for use where UV resistance, high service temperature and processing stability are essential. Four linear SEBS triblocks are presented in the

Figure 22 .

 22 Figure 22. General representation of (a) a high vacuum line 68 and (b) a Schlenk line 69 .

  species often cause broadening of the molecular weight distribution. Recently, a number of novel polymerization systems for (meth)acrylate monomers have been developed to overcome such drawbacks in controlling macromolecular architectures by suitably selecting the initiator, counter cation, solvent, additive, and polymerization temperature. One of the most striking examples is the binary initiator system of a bulky organolithium and LiCl in THF at -78 o C. More generally, this polymerization method has been incompatible with most of useful reactive functional groups, since the propagating chain ends of anionic living polymers are highly basic and extremely nucleophilic. It has been thus very difficult to polymerize functional monomers containing acidic or electrophilic substituents, such as OH, COOH, and CHO groups, under the anionic condition. It has to be noted that anionic polymerization of functional monomers has greatly progressed in the last three decades by mostly introducing the concept of protective groups. Poly(4hydroxystyrene) was for instance produced anionically by protecting the acidic phenolic OH group with a bulky tert-butyldimethylsilyl group70 .

  development of several controlled/living radical systems, termed reversible-deactivation radical polymerizations (RDRP), utilizing an intermittent formation of active propagating species has recently been realized concurrent with similar developments in ionic polymerization systems. A number of RDRP methods have been developed and the three most promising are: nitroxide-mediated polymerization (NMP), atom transfer radical polymerization (ATRP) and degenerative transfer with alkyl iodides, methacrylate macromonomers, and dithioesters via reversible addition-fragmentation chain transfer (RAFT) polymerization. To extend the lifetime of the propagating chains, each of these methods relies on establishing a dynamic equilibrium between a low concentration of active propagating chains and a predominant amount of dormant chains that are unable to propagate or terminate71 . Radicals are reversibly trapped in a deactivation/activation process, as shown in Figure23.

Figure 23 .

 23 Figure 23. Illustration of the deactivation/activation process in RDRP 71 .

  19) were prepared by emulsion polymerization with a carefully designed amphiphilic macro-RAFT agent. The ultimate tensile strength reached 10 MPa with an elongation at break of 500% at the PS composition of 40-50 wt %73 . The RAFT of N-aryl itaconimide and dialkyl itaconate, which were generated from naturally derived itaconic acid, was also investigated74 . It allowed the production of well-defined polymers with controlled Mn and relatively low Đ (≤ 1.59) and enabled block copolymerization to realize a novel, biobased thermoplastic elastomer consisting of hard itaconimide and soft itaconate segments. The AFM images of the copolymer revealed a typical microphase-separated structure with co-continuous lamellae and/or cylinder morphologies. A last study can be mentioned. A series of poly(1-adamantyl acrylate)-bpoly(tetrahydrofurfuryl acrylate)-b-poly(1-adamantyl acrylate) triblock copolymers (Mn = 47-201 kg.mol - 1 , Đ = 1.66-3.03) was prepared by RAFT and a strong microphase separation behavior was demonstrated75 . E = 1.5-52.4 MPa, σB = 1.1-2.6 MPa and εB = 240-803 % were measured for these all-acrylic-based TPE.

  well. Three examples can be cited. Wang and coworkers developed S-B-S (Mn = 19 kg.mol -1 , Đ = 2.16) and S-E-S (E = ether, Mn ˂ 13 kg.mol -1 , Đ = 1.34-1.58) by the combination of the end-group transformation of prefabricated hydroxyl-terminated oligomers and ATRP78 . The hydroxyl-terminated oligomers, such as poly(ether) diol, poly(ester) diol and hydroxyl-terminated poly(butadiene), were reacted with chloroacetyl chloride to form α,ω-bichloroacetyl oligomers, used then as macro-initiator in the bulk polymerization of styrene. Poly(MMA-b-LMA-b-MMA) triblock (LMA = lauryl methacrylate, Mn = 84 kg.mol -1 , Đ = 1.14) by a two-stage ATRP in bulk at 35 o C using a complex as a catalyst and an initiator. The resulting triblock consists of a rubbery middle block (Tg,P(LMA) ~ -65 o C) and glassy end blocks (Tg,P(MMA) = 90-105 o C) 79 . Lastly, the ATRP of MMA initiated by a α,ω-dibromo poly(n-butyl acrylate) P(nBuA) in the presence of NiBr2(PPh3)2 led to P(MMA-b-nBuA-b-MMA) triblock copolymer (Mn = 89-94 kg.mol -1 , Đ = 1.19-1.23, FMMA = 0.22-0.23) 80 .

( 39 )( 40 )( 41 )( 42 )( 43 )( 44 )( 45 )( 46 )( 47 )( 48 )( 49 )( 50 )( 51 )

 39404142434445464748495051 Fetters, L. J.; Morton, M. Synthesis and Properties of Block Polymers. I. Poly-α-methylstyrene-polyisoprene-poly-α-methylstyrene. Macromolecules 1969, 2 (5), 453-458. Bolton, J. M.; Hillmyer, M. A.; Hoye, T. R. Sustainable Thermoplastic Elastomers from Terpene Derived Monomers. ACS Macro Letters 2014, 3 (8), 717-720. Kobayashi, S.; Kataoka, H.; Ishizone, T.; Kato, T.; Ono, T.; Kobukata, S.; Ogi, H. Synthesis and Properties of New Thermoplastic Elastomers containing Poly[4-(1-adamantyl) styrene] Hard Segments. Macromolecules 2008, 41 (14), 5502-5508. Yu, J. M.; Dubois, P.; Jérôme, R. Poly[alkyl methacrylate-b-butadiene-b-alkyl methacrylate] Triblock Copolymers: Synthesis, Morphology, and Mechanical Properties at High Temperatures. Macromolecules 1996, 29 (26), 8362-8370. Imaizumi, K.; Ono, T.; Natori, I.; Sakurai, S.; Takeda, K. Microphase-separated Structure of 1,3-Cyclohexadiene / Butadiene Triblock Copolymers and its Effect on Mechanical and Thermal Properties. Journal of Polymer Science Part B: Polymer Physics 2001, 39 (1), 13-22. Wang, W.; Schlegel, R.; White, B. T.; Williams, K.; Voyloy, D.; Steren, C. A.; Goodwin, A.; Coughlin, E. B.; Gido, S.; Beiner, M.; Hong, K. High temperature Thermoplastic Elastomers Synthesized by Living Anionic Polymerization in Hydrocarbon Solvent at Room Temperature. Macromolecules 2016, 49 (7), 2646-2655. McManus, N. T.; Rempel, G. L. Chemical Modification of Polymers: Catalytic Hydrogenation and Related Reactons. J. Macromol. Sci., Part C 1995, 35 (2), 239-285. Mohammadi, N. A.; Rempel, G. L. Homogeneous Catalytic Hydrogenation of Polybutadiene. J. Mol. Cat. 1989, 50 (3), 259-275. Esneault, C. P.; Hahn, S. F.; Meyers, G. F. Development of Elastomers Based on Fully Hydrogenated Styrene-Diene Block Copolymers. In Developments in Block Copolymer Science and Technology; Hamley, I. W., Ed.; John Wiley & Sons: Chichester, 2004; 341-362. Adkins, H. Hazards in Catalytic Hydrogenation at Elevated Temperatures and Pressures. J. Am. Chem. Soc. 1931, 53 (7), 2808-2809. Parker, D. K. Emulsion Polymers (Noncatalytic Diimide Hydrogenation). In Polymeric Materials Encyclopedia, 12 Volume Set; Salamone, J. C., Ed.; CRC Press: New York, 1996; 2048-2053. Hahn, S. F. An Improved Method for the Diimide Hydrogenation of Butadiene and Isoprene Containing Polymers. J. Polym. Sci., Part A: Polym. Chem. 1992, 30, 397-408. Wu, Z.; Papandrea, J. P.; Apple, T.; Interrante, V. Cross-Linkable Carbosilane Polymers with Imbedded Disilacyclobutane Rings Derived by Acyclic Diene Metathesis Polymerization. Macromolecules 2004, 37, 5257-5264.

Chapter 1 .

 1 Optimization of the β-myrcene nitroxide-mediated polymerization (NMP) and synthesis of well-defined β-myrcene / styrene block copolymers 1.1. Preface Pending the acquisition of high pressure reactors to polymerize isoprene (low boiling point), βmyrcene (My), a non-volatile 1,3-diene, was first considered as a promising building block for making the flexible segment of the thermoplastic elastomer. This chapter presents an article published in the journal Macromolecules (Métafiot, A.; Kanawati, Y.; Gérard, J. -F.; Defoort, B.; Maric, M. Synthesis of β-Myrcene-Based Polymers and Styrene Block and Statistical Copolymers by SG1 Nitroxide-Mediated Controlled Radical Polymerization. Macromolecules 2017, 50, 3101-3120). It details first the optimized controlled polymerization of β-myrcene by nitroxide-mediated polymerization. The subsequent synthesis of My-S diblock and S-My-S triblock copolymers (S = styrene) was made possible by the active feature of the well-defined P(My) elastomeric segment.1.2. AbstractNitroxide-mediated polymerization (NMP) of β-myrcene (My) at 120 o C in bulk using unimolecular SG1-based succinimidyl ester-functionalized BlocBuilder TM (BB) alkoxyamine resulted in low dispersity (Đ = 1.1-1.4) poly(myrcene)s P(My)s with high SG1 chain end fidelity. The polymerizations also showed the number-average molecular weights (Mn) increased almost linearly with conversion. SG1-terminated P(My) macroinitiators were cleanly chain-extended with styrene (S) and the S-rich P(My-b-S) diblock copolymers exhibited two distinct glass transition temperatures (Tgs), indicative of micro-phase separation. P(My-b-S) diblocks showed brittle stress-strain behavior, plausibly due to relatively low Mn. My/S mixtures, with initial S molar feed compositions fS,0 = 0.10-0.94 were also statistically copolymerized (Mn = 8.2-19.8 kg.mol -1 , Đ ≤ 1.37 and monomodal distributions). Copolymer reactivity ratios were rS = 0.25 ± 0.04 / 0.34 ± 0.19 and rMy = 1.88 ± 0.12 / 2.19 ± 0.07 using Fineman-Ross and Kelen-Tüdös methods. The statistical P(My-stat-S) copolymers displayed a range of Tgs (-77 to + 30 o C) depending on My molar fraction. Myrich and S-rich P(My-stat-S)s were then successfully chain-extended with both S and My.

  My's polymerization was comprehensively studied by conventional free radical polymerization 11-16 . Coordination polymerization was also applied to this acyclic conjugated diene. For example, Loughmari and Georges et al. reported stereoselective polymerization of My utilizing a lanthanide-based catalyst 17 . Living or quasi-living ionic polymerizations of My have been explored as well. For instance, Sivola et al reported the synthesis of living P(My)s via anionic Table of contents graphic. Synthesis of P(My), P(My-stat-S) and P(My-b-S) polymers by NMP.

  butadiene were successfully polymerized by NMP. The achievement of relatively broad molecular weight distribution NMP-based polyisoprenes (Đ = Mw/Mn, dispersity of 1.36-1.53) with the use of TEMPO, the first-generation stable nitroxide, was reported in the late 1990s 39,40 . Hawker and co-workers used then a TIPNO-based initiator 46 , 2,2,5-trimethyl-3-(1-phenylethoxy)-4-phenyl-3-azahexane to produce polyisoprenes and polybutadienes with relatively high number-average molar masses (Mn > 80 000 g.mol - 100 poly(isoprene) sample (Đ = 1.28) using 2,2-dimethyl-3-(1-phenylethoxy)-4-phenyl-3-azapentane alkoxyamine produced by the addition of two equivalents of 1-phenylethyl radical to 2-methyl-2nitrosopropane 42 . They also reported the NMP of isoprene using a poly(ethylene oxide)-alkoxyamine macroinitiator, resulting in diblock copolymers with narrow molecular weight distributions (Đ < 1.20) characteristic of living chain growth processes 43 . The discovery of more labile second-generation initiators exhibiting a lower energy C-ON bond homolysis 44,45 such as 2,2,5-trimethyl-4-phenyl-3-azahexane-N-oxyl (TIPNO) 46 and [tert-butyl[1-(diethoxyphosphoryl)-2,2-dimethylpropyl]amino] nitroxide, termed DEPN or SG1 47 (Figure 1,1) allowed controlled polymerizations at lower temperatures. One of the most predominant alkoxyamine unimolecular initiators based on SG1 is 2-([tert-butyl[1-(diethoxyphosphoryl)-2,2dimethylpropyl]amino]oxy)-2-methylpropionic acid termed BlocBuilder TM (BB, Figure 1,2). Initiators containing the SG1 moiety have high dissociation rate constants 45 giving excellent control over the homopolymerizations of 1,3-dienes 48 and many other monomers such as acrylates 49-52 . Harrisson and coworkers reported for the first time the NMP of isoprene from a range of SG1-based initiators 53 . Wellcontrolled poly(isoprene)s were obtained without addition of extra SG1. At higher conversions, the dispersity of all synthesized poly(isoprene)s approached 1.1.

) 1 . 4 .

 14 polymerization of My leading to well-defined P(My) homopolymers, P(My-b-S) diblock copolymers and P(My-stat-S) statistical copolymers. Results and discussion 1.4.1. NMP of β-myrcene: Determination of appropriate . polymerization experimental conditions.

Figure 1 .

 1 Figure 1. Structure of the N-tert-butyl-N-[1-diethylphosphono-(2,2-dimethylpropyl)] nitroxide (SG1, 1), of the 2-([tert-butyl[1-(diethoxyphosphoryl)-2,2-dimethylpropyl]amino]oxy)-2-methylpropionic acid alkoxyamine (BlocBuilder TM , BB, 2) and of the succinimidyl ester terminated NHS-BlocBuilder alkoxyamine (3).

Figure 2 .

 2 Figure 2. (a) Semi-logarithmic kinetic plots of ln((1-XMy) -1 ) (XMy = My conversion) versus reaction time t and (b) Mn and Đ versus XMy for polymerization My-6 (solid green triangle) and My-6' (open circle) initiated by NHS-BB and BB respectively.

4 a)

 4 Initial molar concentration ratio of SG1 free nitroxide to NHS-BB initiator = r = [SG1]0 / [NHS-BB]0. b) Solution polymerizations done in 50 wt% solvent versus initial masses of My and NHS-BB. c) Experiment My-6' has been run for 230 min.

a)

  My conversion (XMy) determined gravimetrically and via1 H NMR (average value). b) Mn and Mw determined by GPC calibrated with PS standards in tetrahydrofuran (THF) at 40 °C and converted to P(My) using MHS constants. c) kpK corresponding to the product of the propagation rate constant kp and the equilibrium constant K derived from the slopes kp[P • ] ([P • ] = concentration of propagating macro-radicals) taken from the semi-logarithmic kinetic plots of ln((1-XMy) -1 ) versus time in the linear region from 0 min to 120 min (0 min to 100 min for My-6' ; squared linear regression coefficient = R 2 ≥ 0.95 for every experiment). kpK's estimated from kp[P • ] and r = [SG1]0 / [NHS-BB]0 (Equation6). Error bars derived from the standard errors in the slope from the linear fits of ln((1-XMy) -1 ) versus time. d) Regioselectivity determined by1 H NMR in CDCl3. 3,4-content% not mentioned and calculated as follows: 3,4-content% = 100 -1,4-content% -1,2-content%. e) Stereoselectivity determined by13 C NMR in CDCl3. trans-1,4-content% not mentioned and calculated as follows: trans-1,4-content% = 100cis-1,4-content%. f) kp[P • ] = (3.2 ± 0.6) × 10 -5 s -1 for My-8 ; kp[P • ] = (2.6 ± 0.5) × 10 -5 s -1 for My-9.g) 1 H and 13 C NMR peaks not clearly detectable.In Equation4, [M]0 and [M]t are concentrations of monomer at time zero and a subsequent later time t, respectively, kp is the propagation rate constant and [P • ] is the concentration of propagating macro-radicals.It can be first noted that almost every trend line fitted to data at the early stages of the polymerization has a positive y-intercept, suggesting the presence of polymers at the commencement of the reaction. This can be explained by the arbitrary t = 0 min chosen when T = 100 o C, which is relatively high compared to the NHS-BB decomposition temperature < 70 o C. For every experiment, the straight kinetic trend at XMy < 0.3 revealed a steady [P • ] established by balancing the rates of activation and deactivation via the SG1 free nitroxide mediator. This illustrates thereby the expected persistent radical effect (PRE)22 based on the dynamic equilibrium between low [P • ] chains and a predominant concentration of dormant chains.However, deviations from the linear trends are observed after the initial stages of the polymerizations, especially when a solvent was used (My-1, My-2 and My-3 in Figure3.a and My-4 in Figure3.b). This is explained by a change in the concentration of active propagating species [P • ]22 . A downward curvature, as clearly observed for experiment My-2 after 200 min (XMy > 0.3), suggests a decrease in [P • ] which may be caused by irreversible termination reactions resulting in an increase of the concentration of the persistent radical. On the other hand, an upward curvature, as seen for experiments My-1, My-2 and My-4, suggests an increase of [P • ] which may be caused by a slower initiation. This latter assumption would suggest that NHS-BB's dissociation rate constant is slightly lower in toluene and 1,4-dioxane than in bulk. In any case, these significant kinetic deviations starting even at moderate conversions (XMy ~ 30%) suggest that other undesirable processes occurred. The thermal self-initiation of My 75 and its relatively low purity level (≤ 10 mol% of impurities including limonene and isomers and dimers of My among others) might alter the persistent radical concentration as well.The slopes, calculated from four to five sample points taken in the linear region, from the kinetic plots reveal the apparent rate constant in terms of the product of kp with [P • ], kp[P • ]. For the systems studied here, the kp[P • ] values were all obtained during the early stages of the polymerization, where the numberaverage molecular weight (Mn) increased linearly with conversion (first 120 min of the reaction). From these values, kpK can be estimated from Equation5and Equation6, where K is the equilibrium constant, [SG1 • ] is the concentration of free nitroxide ([SG1 • ] = [SG1 • ]0 is assumed in the early stages of the polymerization due to a high initial SG1 concentration), [P-SG1] is the concentration of dormant alkoxyamine terminated species which is assumed to be equal to the initial concentration of NHS-BB, [NHS-BB]0 (generally the case in the early stages of the polymerization) and r is the ratio of the initial molar concentration of SG1 free nitroxide to NHS-BB, [SG1]0/[NHS-BB]0.

Figure 3 .

 3 Figure 3. Semi-logarithmic kinetic plots of ln((1-XMy) -1 ) (XMy = My conversion) versus polymerization time t for My polymerizations initiated by NHS-BB. a) Influence of Mn,theo: My-1 with Mn,theo = 20.5 kg.mol -1 (◊), My-2 with Mn,theo = 29.6 kg.mol -1 (■) and My-3 with Mn,theo = 50.1 kg.mol -1 (▲); b) Influence of the reaction medium: My-2 in toluene (■), My-4 in 1,4-dioxane (∆) and My-5 in bulk (♦); c) Influence of temperature: My-5 at 110 °C (♦), My-6 at 120 °C (▲) and My-7 at 130 °C (□); d) Influence of additional SG1: My-6 with no additional SG1 (▲), My-8 with [SG1]0 / [NHS-BB]0 = 0.049 (■) and My-9 with [SG1]0 / [NHS-BB]0 = 0.117 (○). The straight lines indicate linear fits to the experimental data during the initial stages of the polymerization (0 to 120 min). All experimental ID and characterization of experiments are listed inTable1 and Table 2.

  108

Figure 4 .

 4 Figure 4. Mn determined by GPC relative to PS standards in THF and corrected with appropriate MHS coefficients and Đ versus conversion for the various My polymerizations initiated by NHS-BB. a) Influence of Mn,theo: My-1 with Mn,theo = 20.5 kg.mol -1 (◊), My-2 with Mn,theo = 29.6 kg.mol -1 (■) and My-3 with Mn,theo = 50.1 kg.mol -1 (▲); b) Influence of the reaction medium: My-2 in toluene (■), My-4 in 1,4-dioxane (∆) and My-5 in bulk (♦); c) Influence of temperature: My-5 at 110 °C (♦), My-6 at 120 °C (▲) and My-7 at 130 °C (□); d) Influence of additional SG1: My-6 with no additional SG1 (▲), My-8 with [SG1]0 / [NHS-BB]0 = 0.049 (■) and My-9 with [SG1]0 / [NHS-BB]0 = 0.117 (○). The straight lines indicate the theoretical Mn (Mn,theo) versus conversion based on the monomer to initiator ratio for the particular experiment. All experimental ID and characterization of experiments are listed in

  Figure 4.b) indicating presumably the occurrence of irreversible termination reactions. Interestingly, narrower molecular weight distributions were obtained when performing My polymerization in bulk (Đ ≤ 1.27, Figure 4.b). The slightly higher Đ obtained for the two 110 solution polymerizations might be explained by the extra contribution of chain-transfer side reactions to toluene and 1,4-dioxane.

  .d) and Mn and Đ values over time not significantly different (Figure 4.d). More surprising, a loss of control seems to be apparent in the late stages of the polymerization (t > 240 min, XMy > 0.46) when around 12% of SG1 free radical was added initially. A loss of linearity was highlighted in the ln((1-XMy) -1 ) versus time plot (Figure 3.d) combined with greater deviations of the experimental molar masses from the theoretical ones (Figure 4.d). It might be explained by a higher rate of irreversible termination by β-hydrogen chain-transfer (disproportionation) to SG1, as shown by McHale

  illustrates the stress-strain curves of two specimens from experiments My-11-Sa (FS = 0.47) and My-12-S (FS = 0.27) (all stress-strain curves of the specimens from My-11-Sa,

  a) Mn,theo corresponding to the targeted number-average molecular weight of the whole P(My-b-S) diblock copolymer at 100% of S conversion. b) Same P(My) macroinitiator used for these two experiments. c) Tail observed at about 24 min elution time (Figure 6.a) not integrated for the measurement of the molecular weight distribution.

My- 11 -

 11 Sb and My-12-S samples provided in the Supporting Information, FiguresS6a, S7a and S8a).

700% 80 ,Figure 6 .

 806 Figure 6. Normalized GPC traces for S chain-extensions from various P(My) macroinitiators at 110 °C in 50 wt% toluene. a) is the chain extension experiment My-10-S; b) is the chain extension experiment My-11-Sa; c) is the chain extension experiment My-11-Sb; d) is the chain extension experiment My-12-S.

Figure 7 .

 7 Figure 7. Tensile stress-strain curves of two P(My-b-S) diblock copolymer specimens from experiment My-11-Sa (specimen 5) with Mn,P(My) = 22.8 kg.mol -1 and Mn,PS = 26.2 kg.mol -1 (23-26) and from experiment My-12-S (specimen 2) with Mn,P(My) = 30.4 kg.mol -1 and Mn,PS = 12.2 kg.mol -1 (30-12) at room temperature and 1 mm.min -1 . The red crosses denote failure points.

Figure 7 '.

 7 Figure 7'. Atomic force microscopy (AFM) phase image (2 μm x 2 μm) under tapping mode of operation of the surface morphology of S-My-S triblock copolymers: (a) S16-My-S16 and (b) S10-My-S10. The dark domain represents the S component (color-coded height scale given to the right of the image).

120

  Atomic force microscopy (AFM, tapping mode, AFM Bruker Multimode 8 equipped with Nanoscope V controller, scanning speed for image acquisition was 0.5 Hz) was performed to study the micro-phase behavior of the synthesized triblock copolymers. Figure7' shows AFM phase images (Nanoscope Analysis software, version 1.5) of the samples S10-My-S10 and S16-My-S16, initially dissolved in chloroform and spin coated on a silicon wafer. Interestingly, despite a marked disorganization, two phases can be observed for both images with a continuous light domain representing P(My) soft matrix and discrete dark domains corresponding to the rigid PS aggregates. For both cases, the morphology cannot be clearly defined. While the presence of disoriented spheres or cylinders might be seen for S10-My-S10 (Figure7'b), a pseudo lamellar structure may have been obtained for the sample S10-My-S10 possessing a higher fraction of S units (Figure7'a).Lastly, uniaxial tensile testing was implemented at room temperature at a cross-head speed of 5 mm.min -1 . Figure7'' gives the representative stress-strain curves of one specimen of the samples S10-My-S10 and one specimen of the sample S16-My-S16. Generally, S10-My-S10 exhibited a tensile strength at break σB = 0.28 ± 0.05 MPa and an elongation at break εB = 183 ± 11% while S10-My-S10 showed σB = 0.73 ± 0.07 MPa and εB = 89 ± 8%. Compared to the previous My-S diblocks, a higher extensibility was shown by these NMP-based S-My-S. A more entangled P(My) matrix (higher entanglement density) may explain this greater capacity to stretch.

Figure 7 '

 7 Figure 7''. Tensile stress-strain curves of one S10-My-S10 specimen and one S10-My-S10 specimen at room temperature and 1 mm.min -1 . The red crosses denote failure points.

1 .

 1 Every experiment was performed at 110 °C in order to avoid the auto-initiation of S which is negligible below 120 °C87 . However, auto-initiation pathways involving dimerization of S and My might occur at 110 o C. Order-of-magnitude estimate of the rate constant for the dimerization of S and My could be made by inhibition experiments for example. It would allow knowing if any thermal radical generation in the spontaneous My/S copolymerization (initiation step) occurs. Table

  was used herein, postulating that the chemical reactivity of a propagating chain depended only on the chemical nature of the active monomer at the chain end.Extraction of the reactivity ratios was done by linearization of the Mayo-Lewis equation 89 (eg. the Fineman-Ross (FR) approach90 ) and by the Kelen-Tüdös (KT) approach91 (equations and experimental data used to determine the reactivity ratios can be found in the Supporting Information). FR and KT plots are shown in Figures 8.a and 8.b respectively and yielded rMy = 1.88 ± 0.12 and rS = 0.25 ± 0.04 (rMy × rS = 0.47 ± 0.11) via the FR method and rMy = 2.19 ± 0.07 and rS = 0.34 ± 0.19 (rMy × rS = 0.74 ± 0.41) via the

My

  was depleted, more S monomer was added. This specific behavior is highlighted by the difference in copolymer composition between the initial and the final stages of the reaction. For instance, the copolymer was composed of 76 mol% of My at 1 h and only of 53 mol% of My at 7 h for the experiment My/S-50 with fS,0 = 0.51 (individual My and S conversions in the course of the experiment My/S-50 provided in the

a)Table 5 .

 5 Experimental identification (ID) given by My/S-XX where the number abbreviation XX refers to the rounded % initial molar fraction of S in the mixture (fS,0). b) Experiments stopped before 420 min due to a highly viscous reaction medium. My/S copolymerization formulations for various compositions at 110 °C in bulk initiated by NHS-BB and targeting Mn,theo = 30 kg.mol -1 at X = 1.0. 123 Supporting Information, FigureS9a). The reactivity ratios for My/S free radical copolymerization in bulk initiated by azobisisobutyronitrile (AIBN) at 65 °C were also determined, rMy = 1.36 and rS = 0.27, using a non-linear least-squares method 14 . Similar behaviors were observed for isoprene/styrene I/S (rI = 1.92-

124 2 .

 2 02 and rS = 0.42-0.54) 93,94 and butadiene/styrene B/S (rB = 1.35-1.83 and rS = 0.37-0.84) 95-101 free radical copolymerizations, confirming the more reactive nature of these 1,3-diene monomers compared to S. The Mayo-Lewis plot is depicted in Figure8.c and shows that NMP-based P(My-stat-S) formed instantaneously was always richer in My, regardless the initial My/S composition, with no azeotrope observed. A similar trend curve was obtained by Trumbo although the P(My-stat-S)s, synthesized by free radical polymerization in bulk at 65 o C in this case, contained a higher fraction of S 14 . This is explained by a lower My reactivity ratio (rMy = 1.36)14 in comparison to that of the SG1-based nitroxide-mediated statistical copolymerization (rMy ≥ 1.88). The possible temperature dependence of My/S copolymerization reactivity ratios may explain the difference in P(My-stat-S) composition for similar fS,0 between Trumbo's copolymers made at 65 o C and NMP-based copolymers made at 110 o C.

Figure 8 .

 8 Figure 8. (a) FR and (b) KT plots to determine the binary reactivity ratios (solid black circles corresponding to experimental outliers not taken into account for the calculations) and (c) Mayo-Lewis plot of copolymer composition with respect to S (FS) versus initial styrene feed composition fS,0 (experimental data indicated by the solid green circles (•) while the dotted green line is the associated trend line ; open circles (○) refer to the data of Trumbo for My/S system using conventional radical copolymerization at 65 °C in bulk initiated by azobisisobutyronitrile (AIBN) 14 ; The thin straight line indicates the azeotropic composition (fS,0 = FS)) for My and S statistical copolymerizations done in bulk at 110 °C using NHS-BB. See Supporting Information for the full characterization of the samples used.

  photopolymerization of S at 60 °C with 2-azobispropane as sensitizer (kp,S = 176 L.mol -1 .s -1 ) 104 and for the emulsion polymerization of I at 60 °C with diisopropyl-benzene monohydroperoxide / tetraethylenepentamine (DIBHP/TEPA) catalyst system (kp,I = 50 L.mol -1 .s -1 ) 105 . The lower propagation rate constant reported regarding the free radical polymerization of I compared to that of S is consistent with the kinetic data obtained herein for the NMP of My and S.

Figure 9 .

 9 Figure 9. Normalized GPC traces of P(My-stat-S) copolymer initiated by NHS-BB at 110 °C in toluene targeting Mn,theo = 30 kg.mol -1 at quantitative conversion (experiment My/S-50).

1 . 4 . 5 .

 145 FigureS13a), likely due to the presence of "dead" chains from the macroinitiators and irreversible terminations.

Table 8 .

 8 ).As expected, P(My), as an amorphous elastomer, had a Tg = -77 °C much below the ambient temperature and very close to the natural rubber cis-1,4-PI, which has Tg around -73 °C112 . NMP-based PS exhibiting Mn = 15.6 kg.mol -1 had a Tg = + 77 °C instead of 100-110 °C for atactic PS with Mn > 100 kg.mol-1 113 .Beyond the critical influence of Mn on Tg, the possible presence of oligomers in the PS matrix, behaving as plasticizers, could play a minor role in lowering the glass transition temperature. S-rich diblock copolymers, namely My-10-S and My-11-Sa, showed two separated Tgs: a first one at about -70 °C corresponding to the soft P(My) block and a second one around + 60 °C corresponding to the hard PS block. This is illustrated in Figure10by the changes in slope of the DSC trace for My-11-Sa sample.This suggests the possible micro-phase separation of these two heterogeneous phases. However, an in-depth analysis of the P(My-b-S)s phase behavior, including atomic force microscopy (AFM) and smallangle X-ray scattering (SAXS) for example, would demonstrate whether or not a micro-phase separation is apparent and determine the morphology of this latter. It can be noted that the Tgs of these diblocks did not exactly correspond to the Tgs of the respective P(My) and PS homopolymers. This might be due to a partial Glass transition temperatures (Tgs) and decomposition temperatures (Tdecs) for P(My) (My-12), PS (My/S-100) and P(My-b-S) diblock copolymers (My-10-S, My-11-Sa, My-11-Sb and My-12-S). a) Tg('s) measured by DSC under nitrogen atmosphere at a rate of 10 °C.min -1 using three scans per cycle (heat/cool/heat). b) Tdec,1 (onset of decomposition), Tdec,max (temperature at which weight loss is most apparent) and Tdec,2 (end of decomposition) measure by thermogravimetric analysis (TGA) under nitrogen flow at a ramp rate of 10 °C.min -1 (TGA thermograms provided in the Supporting Information). 132 miscibility of the two blocks caused by both a relatively low Mn and a vulnerable phase separation taking the structural similarities between My and S repetitive units into account. Partial mixing of PI and PS blocks were previously observed for P(I-b-S) diblocks leading to in-between glass transition temperatures 114 .Interestingly, for My-rich diblock copolymers (My-11-Sb and My-12-S) with FS ≤ 0.39, no second glass transition was observed. The Tg for the PS segment was not visible on the DSC thermograms presumably due to a too low PS content. Similarly, as regards β-myrcene/α-methyl-p-methylstyrene My/AMMS triblocks with FAMMS < 0.35, even though the Tg for P(My) was detectable, no DSC signal corresponding to the Tg of the P(AMMS) outer segments was observed21 . By contrast, both Tgs for P(I-b-S) diblock copolymers were visible even for a large fraction of PS (FS = 0.78) and the Tg for PI was not observed only for FS = 0.92 114 .

Figure 10 .

 10 Figure 10. DSC trace (second heating run) of P(My-b-S) diblock copolymer My-11-Sa. The blue arrows indicate the changes in slope observed. DSC traces of the other P(My-b-S) diblock copolymers can be found in the Supporting Information.

Figure 11 alsoFigure 11 .

 1111 Figure 11 also includes the Fox equation predictions of the P(My-stat-S) Tgs. Non-negligible deviations can be noticed from our experimental data and explained by the limitations of the Fox equation in its prediction of the glass transition temperatures. Basically, it is a simple bulk additive relation disregarding the sequence distribution of My and S monomer units in the copolymer 116 . Moreover, the characterized P(My-stat-S)s did not exhibit the same Mn, ranging from 8.2 to 19.8 kg.mol -1 . Slight changes in Tg values can thereby be expected, as predicted by the empirical Flory-Fox equation 117,118 . My, containing two isoprene units, is a monoterpene structurally related to the readily available petro-chemical isoprene (I) 3 . Both P(My) and

  kg.mol -1 at complete conversion. Resulting cis-1,4-rich P(My) containing 1,2-and 3,4-defects (< 20%) and exhibiting Đ < 1.30 was achieved. Its Tg and decomposition temperature Tdec,max were -77 °C and + 385 °C, respectively. The active feature of the NMP-based P(My) macroinitiator was confirmed by subsequent chain-extensions with S leading to well-defined P(My-b-S) diblock copolymers, retaining their monomodal nature. S-rich P(My-b-S)s showed two separate Tgs, at about -70 °C and + 60 °C, indicating possibly their micro-phase separation. P(My-b-S) diblocks exhibited better thermal stability (Tdec,max = 435-440 °C) than P(My) and poor strain-stress properties with ultimate tensile strength < 0.9 MPa and tensile elongation at break < 14%. This brittle behavior combined with a low extensibility can be largely explained by a lack of entanglement density due to the experimental P(My) Mn = 13.2-30.4 kg.mol -1 , lower than or equal the P(My) entanglement molecular weight (Me,P(My) = 22-31 kg.mol -1 )81 . BlocBuilder-terminated poly(ethyleneco-butylene)-(SG1)2 difunctional initiator was used to produce S-My-S triblock copolymers with Mn = 56-66 kg.mol -1 , which exhibited improved extensibility (σB < 0.8 MPa, εB < 200%). Statistical copolymers from My and S were also prepared by NMP with fS,0 = 0.10-0.94 at 110 °C in bulk using NHS-BB. The Fineman-Ross approach yielded rMy = 1.88 ± 0.12 and rS = 0.25 ± 0.04 (rMy × rS = 0.47 ± 0.11) whereas rMy = 2.19 ± 0.07 and rS = 0.34 ± 0.19 (rMy × rS = 0.74 ± 0.41) were determined by Kelen-Tüdös approach.Well-controlled copolymerizations were performed with resulting P(My-stat-S)s exhibiting Mn = 8.2-19.8 kg.mol -1 , Đ < 1.40 and monomodal distributions as shown by the GPC chromatograms. P(My-stat-S)s displayed a glass transition temperature between -77 and + 30 °C depending on the molar composition of S in the statistical copolymer. The high chain-end fidelity of My-rich and S-rich P(My-stat-S)s was confirmed via successful chain-extensions with both My and S. This investigation thereby presented a robust and simple way to synthesize P(My), P(My-b-S) and P(My-stat-S) by NMP.

( 21 )

 21 Bolton, J. M.; Hillmyer, M. A.; Hoye, T. R. Sustainable Thermoplastic Elastomers from Terpene-Derived Monomers. ACS Macro Lett. 2014, 3, 717-720.(22) Nicolas, J.; Guillaneuf, Y.; Lefay, C.; Bertin, D.; Gigmes, D.; Charleux, B. Nitroxide-Mediated Polymerization. Prog. Polym. Sci. 2013, 38 (1), 63-235.

  Bulk nitroxide-mediated polymerization (NMP) of β-myrcene (My) and glycidyl methacrylate (GMA) with varying GMA molar feed fraction (fGMA,0 = 0.10-0.91) was accomplished at 120 o C, initiated by SG1-based alkoxyamine bearing a N-succinimidyl ester group (NHS-BlocBuilder). Low dispersity My/GMA copolymers (Đ ˂ 1.56) with slight number-average molecular weights (Mns) deviations from predicted values (Mn,theo with Mn/Mn,theo ˃ 70%) were obtained. The copolymerization was revealed to be statistical, confirmed via Fineman-Ross (rMy = 0.80 ± 0.31 and rGMA = 0.71 ± 0.15) and Kelen-Tüdös (rMy = 142 0.48 ± 0.12 and rGMA = 0.53 ± 0.18) approaches. Glass transition temperature (Tg) of the statistical P(Mystat-GMA)s increased from -77 °C to + 43 °C as the GMA molar fraction incorporated (FGMA) increased from 0.10 to 0.90. High SG1 chain-end fidelity for My-rich and GMA-rich P(My-stat-GMA)s was assessed by phosphorus nuclear magnetic resonance ( 31 P NMR, SG1 fraction ˃ 69 mol%) and chain-extensions in toluene with My, GMA and styrene (S) (monomodal shift in Mn). Lastly, diblock P(My-b-GMA) was made and treated with morpholine to produce amphiphilic copolymer able to self-organize into micelles.

Figure 1 .

 1 Figure 1. Copolymerization of My and YY (MA, tBuA, MMA or GMA) by NMP using NHS-BlocBuilder initiator.

Table 1 .

 1 Experimental conditions and results for NHS-BB-initiated copolymerizations of My with MA, tBuA and MMA. a) Experimental identification given by My/YY-XX where YY refers to the functional comonomer and XX refers to the rounded % initial molar fraction of My in the mixture (fMy,0). b) Initial molar concentration of YY in the mixture. c) Overall monomer conversion X = XMyfMy,0 + XYYfYY,0 determined by 1 H NMR in CDCl3. d) Mn and Mw determined by GPC calibrated with PS standards in tetrahydrofuran (THF) at 40 °C. e) Undetectable polymer peak on the GPC chromatogram.transformation that was indicated by the colour change. This incompatibility of My and MA to copolymerize might be explained by the formation of a Diels-Alder adduct. The Diels-Alder cycloaddition of isoprene with MA to yield 4-methyl-4-cyclohexene-1,2-dicarbocylic anhydride is well-known55 and a similar mechanism was also reported for the My/MA system.56 Figure S4b in the Supporting Information gives the individual My and MA conversions versus reaction time for both experiments. Interestingly, a similar trend in consumption for My and MA is noticeable regardless of the initial feed composition. This supports the assumption of the formation of a cyclic side product from a stoichiometric amount of My and MA.

  a) Experimental identification given by My/GMA-XX where XX refers to the rounded % initial molar fraction of My in the mixture (fMy,0). b) Experiment stopped before the arbitrary time t = 0 min due to a highly viscous reaction medium. c) My homopolymerization performed previously under the same experimental conditions. 35 150

  03 indicating the stronger tendency to alternate compared to My/GMA system polymerized by bulk NMP at 120 o C. Two parameters may explain the difference in constants observed: the significantly higher reaction temperature for My/GMA copolymerization (120 o C) and the nature of the conjugated 1,3-diene and thus the effect of the lateral C6/C8 side group borne by My units.

Figure 2 .

 2 Figure 2. (a) FR and (b) KT plots (solid blue circles (•) corresponding to experimental data while solid lines refer to the trend lines) and (c) Mayo-Lewis plot of copolymer composition with respect to GMA (FGMA) versus initial GMA feed composition fGMA,0 (experimental data indicated by the solid blue circles (•) while the blue dashed line is the associated trend line; open circles (○) refer to the data of Contreras-Lopez and coworkers for the I/GMA system initiated by benzoyl peroxide; 75 the straight line indicates the azeotropic composition (fGMA,0 = FGMA)) for My and GMA copolymerizations done in bulk at 120 °C using NHS-BB. See Table S1b in the Supporting Information for the characterization of the samples used.

3 .

 3 , [M]0 and [M]t are the concentrations of monomers at time zero and subsequent later time t, respectively, ˂kp˃ is the average propagation rate constant and [P • ] is the concentration of propagating macro-radicals. Regardless of the My/GMA initial feed composition, at the early stages of the copolymerization (caption (d) in Table 3), the straight kinetic trends of the semi-logarithmic plots highlight Table Molecular characterization and My selectivity at the end of the experiments and kinetic data of P(My-stat-GMA) copolymers at 120 °C in bulk initiated by NHS-BB and targeting Mn,theo = 30 kg.mol - 1 at X = 1.0. a) Molar fraction of My in the copolymer (FMy) as determined by 1 H NMR (Figure S6b in Supporting Information for the spectral assignments). b) Individual XMy and XGMA determined by 1 H NMR. Average X = XMyfMy,0 + XGMAfGMA,0 (further details in the experimental section). c) Mn and Mw determined by GPC calibrated with PS standards in THF at 40 °C. d) ˂kp˃˂K˃ corresponding to the product of the average propagation rate constant ˂kp˃ and the average equilibrium constant ˂K˃ derived from the slopes ˂kp˃ [P•] ([P•] = concentration of propagating macroradicals) taken from the semi-logarithmic kinetic plots of ln((1-X) -1 ) versus time in the linear region generally from 0 to 65 min (0 to 30 min for My/GMA-30, 0 to 15 min for My/GMA-20, 0 to 10 min for My/GMA-10, 0 to 120 min for My/GMA-100 ; squared linear regression coefficient = R 2 ≥ 0.98 for every experiment). ˂kp˃˂K˃'s estimated from ˂kp˃ [P•] and r = [SG1]0/[NHS-BB]0 (Equation 5). Error bars derived from the standard errors in the slope from the linear fits of ln((1-X) -1 ) versus time. e) Determined by 1 H NMR. 3,4-content% = 100 -1,4-content% -1,2-content% (Figure S6b in Supporting Information for further details).

)

  By assuming that the initial concentration of nitroxide [N]0 is high so that [N] = [N]0 in the early stages of polymerization and that [P-N] is approximately equal to the initial concentration of initiator ([P-N] = [NHS-BB]0), the following equation 5 can be obtained from Equation 3 with r = [N]0/[NHS-BB]0 (initial molar concentration ratio of SG1 free nitroxide to NHS-BlocBuilder initiator): ˂kp˃˂K˃ ≈ (˂kp˃[P • ][N • ]0) / [NHS-BB]0 = ˂kp˃[P • ]r

Figure 3 .

 3 Figure 3. (a) Semi-logarithmic kinetic plots of ln((1-X) -1 ) (X = overall conversion) versus polymerization time t, (b) Đ versus X and (c) Mn versus X for the various My/GMA copolymerizations in bulk at 120 o C initiated by NHS-BB. The dotted line indicates the theoretical Mn versus overall conversion based on the monomer to initiator ratio (Mn,theo = 30 kg.mol -1 at X = 100% for every experiment). All experimental ID and characterization of experiments are listed in Table2 and Table 3.

Figure 4 .

 4 Figure 4. Normalized GPC traces of P(My-stat-GMA) with fGMA,0 = 0.51, initiated by NHS-BB at 120 °C in bulk targeting Mn,theo = 30 kg.mol -1 at X = 1 (experiment My/GMA-50). Curves from right to left correspond respectively to the GPC chromatograms of samples taken at 15, 65, 120 and 300 min.

4 .

 4 ) compared to the GMA homopolymerization by NMP ineffectively controlled by the SG1 free nitroxide. This observation echoes the comprehensive work done by Benoit et al., reporting the well-controlled bulk NMP of tert-butyl acrylate Table Chain-extensions of A) My-rich and GMA-rich P(My-stat-GMA) macroinitiators (MI) with B) My, GMA and S monomers in 50 wt% toluene and C) molecular characterization of the resulting chainextended products. a) The indexes "1" and "2" refer respectively to the final features of the P(My-stat-GMA) macroinitiator (MI) and the whole chain-extended diblock copolymer (macroinitiator + P(My), P(GMA) or PS segment added). b) Living molar fraction of MI chains terminated by a SG1 unit, measured by 31 P NMR (Figures S7b and S8b in Supporting Information). c) Predicted Mn,1 at X1 measured experimentally and calculated as follows: Mn,theo,X1 = (X1/100)Mn,theo,1 with Mn,theo,1 = 30 kg.mol -1 at X = 100%. d) Targeted number-average molecular weight of the whole chain-extended diblock copolymer (MI block + extended block) at X = 100%.

Figure 5 .

 5 Figure 5. Normalized GPC traces for the chain-extensions of My/GMA-78 with (a) GMA with 10 mol% of My (experiment My/GMA-78-GMA/My) and (a') S (experiment My/GMA-78-S) and for the chainextensions of My/GMA-37 with (b) My (experiment My/GMA-37-My) and (b') S (experiment My/GMA-37-S).

  6 o C relatively close to that of P(My) as predicted by Figure 6, two distinct Tgs were exhibited by the diblock copolymers My/GMA-78-GMA/My (-52.7 and + 15.8 o C) and My/GMA-78-S (-27.1 and + 61.4 o C). Such a thermal behavior is typical of the presence of two heterogeneous phases, highlighting herein the high degree of chemical incompatibility between My/GMA-78 block and the chain-extended block. Interestingly, the lower Tg of these diblock copolymers (-52.7 and -27.1 o C) does not correspond exactly to that of the parent macroinitiator (-63.6 o C). This is particularly marked with My/GMA-78-S having a subzero Tg significantly

Figure 6 .

 6 Figure 6. FGMA effects on Tg in P(My-stat-GMA)s (solid blue circles, DSC traces given in the Supporting Information, Figure S10b). The Fox equation predictions represented by the black dotted line while the blue dashed line represents the experimental data fitted to the Gordon-Taylor equation.

  Flory and Huggins as given in Equation 8: 94 ∆GM = kT [(V/Vr) vP(GMA) vP(My) χMy-GMA (1-2/z) + Nc (vP(GMA) ln(vP(GMA)) + vP(My) ln(vP(My))] (

Figure 7 .

 7 Figure 7. DSC traces of (a) My/GMA-37, My/GMA-37-My and My/GMA-37-S and (b) My/GMA-78, My/GMA-78-GMA/My and My/GMA-78-S diblock copolymers (dashed blue lines indicating the changes in slope observed; magnifications with Tg values given below for the changes in slope harder to observe, as indicated by the blue indexes in brackets).

  morpholinopropyl methacrylate) P(My-b-HMPMA) diblock copolymer. Morpholine as a nucleophlic reagent was selected relying on the conclusive results reported by Benaglia et al. demonstrating a complete conversion of the P(GMA) epoxy groups via a 10-fold amount of morpholine in DMSO at 80 o C. 96 The synthesis route is shown in Figure 8. Firstly, NHS-BB-mediated P(My) exhibiting Mn = 14.4 kg.mol -1 and Đ = 1.51 was synthesized in bulk at 120 o C after 6 h of polymerization. The chain-extension of this NMP-based polyterpene with 90 mol% of GMA and 10 mol% of My was then performed for 50 min at 110 o C in 50 wt% of toluene. The

Figure 8 .

 8 Figure 8. Synthetic route for P(My-b-HMPMA) diblock copolymer.

Figure 9 . 1 H

 91 Figure 9. 1 H NMR spectra of (a) dry P(My-b-GMA) diblock copolymer and (b) dry P(My-b-HMPMA).

  measurements were carried out for every sample at three different temperatures, TDLS = 25, 45 and 65 o C. It should be first noted that a monomodal distribution was observed for every sample, with a DLS polydispersity index relatively high for experiments conducted at room temperature (PDI = 0.24 on average at T = 25 o C) and more satisfactory at T = 45-65 o C (PDI = 0.10 on average). The relatively low PDI values make the Z-average size reliable and was reported in this study. Regardless of the preparation of the samples, as TDLS increased from 25 to 45 o C, a marked drop of the Z-average diameter (Zave) was noticeable (105 nm drop on average). However, increasing TDLS from 45 to 65 o C did not result in a further Zave decrease. The set of experiments III (Table 5) can be taken as an example. While large aggregates were detected at room temperature (˃ 300 nm), a second population of aggregates exhibiting a smaller hydrodynamic diameter (~ 130 nm) was obtained at 45 o C and 65 o C. Accordingly, a critical temperature (Tcri) between 25 and 45 o C might favor the formation of micelles exhibiting a lower hydrodynamic

  The influence of the concentration over Zave can be assessed by comparing the DLS results between the first (I) and the second (II) preparation, where samples were stirred for 3 days. Regardless of TDLS, a clear decrease of Zave was effective when reducing [P(My-b-HMPMA)] from 1.4 to 1.0 mg.ml -1 . Generally, as the time of stirring increased and the concentration of the copolymer decreased, a clear reduction of Zave was observed. The aggregations of the micelles limited by reducing [P(My-b-HMPMA)] and/or stirring rigorously for 6 days or more. It can be assumed that stirring allowed to homogenously disperse the particles in the water solution whereas a lower [P(My-b-HMPMA)] diminished statistically the formation of large aggregates. Even though other experimental conditions could be used to characterize the self-assembly of P(My-b-HMPMA) exhibiting Mn = 26.5 kg.mol -1 and FMy = 0.55, the study led herein suggests that micelles having Zave = 120-130 nm can be formed in water at T = 45-65 o C and [P(My-b-HMPMA)] ≤ 0.7 mg.ml -1 , after a few days of stirring. It could also be argued that Zave ≥ 120 nm is too high (50-80 nm expected) to form simple micelles in water for P(My-b-HMPMA) diblock exhibiting relatively low Mn.

1 . 5 .Chapter 3 .

 153 35 at X as high as 79-89%. Even though broader molecular weight distributions (Đ = 1.35-1.55) were Table Dynamic light scattering analyses of P(My-b-HMPMA) in reverse osmosis water solution. a) Z-average diameter (cumulants mean) derived from three separate repeats. The error is given as the standard deviation from the three separate repeats. b) Polydispersity Index (PDI) corresponding to the coefficient of the squared term ''c'' when scaled as 2c/b 2 , where ''b'' and ''c'' are obtained from the cumulants analysis, which fits a polynomial to the log of the G1 correlation function. c) Time of stirring. Replacement of poly(styrene) by poly(isobornyl methacrylate) as the rigid domain of the β-myrcene-based block copolymers 3.1. Preface This chapter presents a manuscript which will be shortly submitted to the journal Polymer Chemistry (Métafiot, A.; Gagnon, L.; Pruvost, S.; Hubert, P.; Gérard, J. -F.; Defoort, B.; Maric, M. β-Myrcene / Isobornyl Methacrylate SG1 Nitroxide-Mediated Controlled Radical Polymerization: Synthesis and Characterization of Gradient, Diblock and Triblock copolymers).Despite the successful synthesis of NMP-based S-My-S triblocks (Chapter 1, section 1.4.3), poor mechanical properties were exhibited by these polymers at room temperature (tensile stress at break σB < 0.8 MPa, tensile elongation at break εB < 200%). The use of P(My) as the soft segment of the triblocks, imparting ideally the elastomeric properties, can legitimately be a key factor of the brittleness of these copolymers. Indeed, the high entanglement molar mass of P(My) (22-31 kg.mol -1 ) results in a relatively low entanglement density of the continuous flexible domain. However, another parameter which can affect the mechanical properties of the copolymer is the nature of the hard segment(s), as mentioned in the literature review (section C.1.4

  β-Myrcene / isobornyl methacrylate (My/IBOMA) mixtures were copolymerized in bulk using the SG1-based BlocBuilderTM alkoxyamine functionalized with N-succinimidyl ester group, NHS-BlocBuilder, at T = 100 o C with initial IBOMA molar feed compositions fIBOMA,0 = 0.10-0.90. Copolymer reactivity ratios were rMy = 1.90-2.16 and rIBOMA = 0.02-0.07 using Fineman-Ross, Kelen-Tudos and non-linear least-squares fitting to the Mayo-Lewis terminal model, and indicated the possibility of gradient My/IBOMA copolymers. A linear increase in molecular weight versus conversion and a low dispersity (Đ ≤ 1.41) were exhibited by 176 My/IBOMA copolymerization with fIBOMA,0 ≤ 0.80. My-rich and IBOMA-rich copolymers were shown to have a high degree of chain-end fidelity by performing subsequent chain-extensions with IBOMA and/or My, and by 31 P NMR analysis. IBOMA-My-IBOMA triblock copolymers containing a minor fraction of My or styrene (S) units in the outer hard segments (Mn = 51-95 kg.mol -1 , Đ = 1.91-2.23 and FIBOMA = 0.28-0.36) were then synthesized using SG1-terminated poly(ethylene-stat-butylene) dialkoxyamine. The micro-phase separation was suggested by the detection of two distinct Tgs at about -60 o C and + 180 o C, and confirmed by atomic force microscopy (AFM). A plastic stress-strain behavior (stress at break σB = 3.90 ± 0.22 MPa, elongation at break εB = 490 ± 31 %) associated to an upper service temperature of about 140 o C were also highlighted for these triblock polymers.

  industrial implementation made possible by other processes such as free radical polymerization 4 . Since the early 1980s, an easy-to-use synthetic technique now termed reversible-deactivation radical polymerization (RDRP) 5-9 was developed, encompassing several novel controlled/pseudo-living radical polymerizations and allowing the preparation of tailor-made macromolecules with precise and predetermined molar masses, Table of contents graphic. Synthesis of well-defined My/IBOMA gradient and diblock copolymers by NMP at 100 o C in bulk, initiated by the alkoxyamine NHS-BB. Preparation of Co/IBOMA-My-IBOMA/Co triblock copolymers by NMP (Co = My or S comonomer, < 10 mol%), using PEB-(SG1)2 difunctional initiator.

18 - 22 ,

 1822 mediated by the 2,2,5-trimethyl-4-phenyl-3-azahexane-3-oxyl nitroxide (TIPNO) 18,23 , a more labile second-generation initiator. N-tert-Butyl-N-[1-diethylphosphono-(2,2-dimethylpropyl)] nitroxide (SG1)24 showed also its effectiveness to control the polymerization of I 25-27 . Recently, we studied the well-controlled polymerization of β-myrcene28 , an acyclic monoterpene, initiated by unimolecular SG1-based succinimidyl esterfunctionalized BlocBuilder TM alkoxyamine called NHS-BlocBuilder 29 , demonstrating once again the facility of the NMP process to synthesize elastomers based on 1,3-dienes with low dispersity (Đ) and substantial chain-end fidelity.Well-tailored styrene (S)/I or B statistical, diblock and triblock polymers were thereby made by NMP 18,30,31 . Surprisingly enough and to the best of our knowledge, no NMP-based styrenic-diene block copolymers were reported in the literature as possible alternative plastics with valuable mechanical properties or as thermoplastic elastomers. This observation prompted us to synthesize NMP-based block copolymers based on an unusual but promising monomer coupling, namely β-myrcene (My) / isobornyl methacrylate (IBOMA).A naturally occurring monoterpene with a highly active diene structure, β-myrcene (My, 7-methyl-3-methylene-octa-1,6-diene), provides a classical chemistry similar to unsaturated hydrocarbons32 . Its polymerization leads to a bio-sourced P(My) elastomer exhibiting an expected sub-zero glass transition temperature (Tg) of about -75 o C and a pseudo-plastic behavior 33 . My-based thermoplastic elastomers were notably produced by anionic polymerization, making P(My) a possible elastomer substitute to well-known poly(butadiene) PB and poly(isoprene) PI rubbers. In 1983, S-My-S triblock copolymers, initiated by secbutyllithium at 30 o C in benzene solution, were made by sequential monomer addition (number-average molecular weight Mn = 165-194 kg.mol -1 , final S molar fraction FS = 0.23-0.58) and exhibited tensile strengths at break σB = 4.3-12.8 MPa and elongations at break εB = 670-1290% 34,35 . Such triblocks exhibited the classical self-assembled morphologies in this case. More recently, Bolton et al. produced AMMS-My-AMMS (AMMS = α-methyl-p-methylstyrene) triblock polymers by sequential anionic polymerization 36 , which showed ultimate tensile stress values ranging from 0.5 to 10.8 MPa and εB ranging from 525 to 1340%. Interestingly, the insertion of high Tg P(AMMS)s as outer hard blocks (Tg,P(AMMS) ~ 180 o C) allowed the extension of the upper service temperature of these triblocks compared to that of traditional PS-based TPEs, which are limited to 90-100 o C. Similarly, P(IBOMA), displaying a high Tg ~ 190 o C 37 , can be used as the rigid domain of block copolymer TPEs to provide strength, hardness and an improved performance at high temperature. This was illustrated by Yu and coworkers, who prepared IBOMA-B-IBOMA triblock copolymers initiated by the mdiisopropenylbenzene / tert-butyllithium diadduct 38 , which exhibited excellent stress-strain properties and an upper service temperature of 160 o C. With the development of alkoxyamine unimolecular initiators and SG1 nitroxide in the late 1990s, the polymerization of IBOMA, a bulky hydrophobic methacrylate, is no longer a barrier towards such block copolymers by NMP. Control of methacrylate polymerizations was accomplished by Charleux and coworkers via SG1-based initiators and a small fraction of co-monomer such as S 39,40 or acrylonitrile 41 (< 10 mol%).

  transition temperature (Tg) and decomposition temperature (Tdec) of the My/IBOMA copolymers were also investigated. Secondly, the preparation of relatively high Mn P(My)-(SG1)2 was explored by using the PEB-(SG1)2 dialkoxyamine. The possibility to synthesize triblock polymers by IBOMA/Co chain-extensions (Co = My or S co-monomer ~ 8-9 mol%, relative to IBOMA) from P(My)-(SG1)2 was then presented.

Figure 1 .

 1 Figure 1. a) My/IBOMA gradient copolymerization in bulk initiated by NHS-BB and subsequent My or IBOMA/My (~ 93/7 mol%) chain-extension in toluene and b) synthesis of P(My)-(SG1)2 from PEB-(SG1)2 difunctional initiator (n ˃˃ p, q) and subsequent IBOMA/Co chain-extension in toluene (Co = My or S co-monomer, 8-9 mol%).

  obtained, by reducing likely the average activation-deactivation equilibrium constant ˂K˃ of this binary system. Regardless of fMy,0, similar Mn,MHS values were measured for the experiments run at 100 and 120 o C (Figure S1c, Supporting Information). However, narrower molecular weight distributions were obtained (Đ ≤ 1.29 for X = 9-65%) at 100 o C, for both experiments exhibiting fMy,0 = 0.50 and 0.70. This apparent enhanced control of the NMP may be due to the reduction of disproportionation side reactions (β-hydrogen transfer from a propagating radical to SG1), commonly observed during the NMP of methacrylates 52,53 , at 100 o C. Accordingly, a copolymerization temperature of 100 o C was selected for the My/IBOMA NMP. This reduction of temperature is consistent with the SG1-mediated methacrylate polymerizations performed previously with a fairly low temperature (~ 70-90 o C), notably by Charleux et al. 39,54 and Lessard et al. 55,56 .

Figure 2 .

 2 Figure 2. (a) Fineman-Ross (FR) and (b) Kelen-Tudos (KT) plots (solid black circles (•) corresponding to experimental outliers not taken into account for the calculations while solid lines refer to the linear trend lines) to determine the binary reactivity ratios for My and IBOMA for copolymerizations done in bulk at 100 o C initiated by NHS-BB (parameters G, H, 𝜼 and ε defined in the Supporting Information). (c) Mayo-Lewis plot of My/IBOMA copolymerizations with respect to IBOMA, using the final molar composition FIBOMA, and the initial monomer feed composition fIBOMA,0. The solid straight line indicates the azeotropic composition where fIBOMA,0 = FIBOMA while the dashed line is the associated trend line of the experimental data (solid blue circles (•)). Table S1c in the Supporting Information lists the samples used for these plots. (d) Individual My (•) and IBOMA (○) conversions, determined by 1 H NMR in CDCl3, versus reaction time t for the gradient copolymerization My/IBOMA-50 exhibiting fMy,0 = 0.50.

  40 o C). The likely large differences in hydrodynamic volumes of the P(My-grad-IBOMA)s and that of the PMMA standards that was used to calibrate the GPC may partly explain the deviations. The copolymerization with fMy,0 = 59-90 had a flatter slope compared to the others (Figure3c), where the unreliability of the GPC calibration cannot be the sole reason. Chain transfer side reactions to the monomers, creating new chains and resulting in

PFigure 3 .

 3 Figure 3. (a) Semi-logarithmic kinetic plots of ln((1-X) -1 ) (X = overall conversion) versus polymerization time t, (b) Đ versus overall conversion X and (c) Mn determined by GPC relative to PMMA standards in THF at 40 o C, and corrected using the Mark-Houwink relationship, versus overall conversion X for the various My/IBOMA copolymerizations in bulk at 100 o C initiated by NHS-BB.

4 ) 5 )

 45 , [M]0 and [M]t are the concentrations of monomer at time zero and subsequent later time t, respectively, <kp> is the average propagation rate constant and [P•] is the propagating radical concentration. Marked upward curvatures (Figure 3a) can be seen for experiments My/IBOMA-10 and My/IBOMA-20 having an initial feed rich in IBOMA (fIBOMA,0 = 0.80-0.90). It suggests an increase in [P•], which can be explained by the contribution of the likely high individual K of IBOMA, resulting from slow recombination of nitroxides and poly(methacrylate) radicals (decrease in the concentration of dormant chains) and leading to a high [P•]. <kp>[P•] values were determined from the slopes of the semi-logarithmic plots of conversion versus time (figure 3a) in the linear regions where [P•] is expected to remain relatively constant (See Table 2, caption d). These slopes can be related to the equilibrium between the dormant and the active chains as shown in Equation 4. K = ([P•][SG1•]) / [P-SG1] (where [SG1•] and [P-SG1] are the concentrations of free nitroxide and dormant alkoxyamine terminated polymer, respectively. At the early stages of the copolymerization, it can be hypothesized that [P-SG1] = [NHS-BB]0 (one initiating moiety generates a dormant chain) as well as [SG1•]0 = [SG1•] (large initial amount of free nitroxide being relatively constant in the course of the reaction). Equation 4 can therefore be rewritten in terms of r = [SG1]0/[NHS-BB]0. <kp><K> ≈ (<kp>[P•][SG1]0 / [NHS-BB]0) = <kp> [P•] r (Initially, we expected significant differences in <kp><K> values between My-rich and IBOMA-rich initial feeds. Indeed, even though no kinetic constants for My are reported yet, kp,I = 125 ± 30 L.mol -1 .s -1 for the radical polymerization of isoprene, another similar conjugated 1,3-diene, initiated by di-tert-butyl peroxide at 5 o C, was reported 67 . Comparatively, kp,IBOMA ~ 280 L.mol -1 .s -1 at 5 o C can be estimated from the study led by Beuermann and coworkers using the pulsed-laser polymerization-GPC technique 68 .

Figure 4 . 190 •

 4190 Figure 4. Normalized GPC traces of P(My-grad-IBOMA) with fMy,0 = 0.50, initiated by NHS-BB at 100 °C in bulk targeting Mn,theo = 30 kg.mol -1 at X = 100% (experiment My/IBOMA-50).

6 )Figure 5 .

 65 Figure 5. FIBOMA effects on Tg in P(My-grad-IBOMA) gradient copolymers. The DSC traces can be found in the Supporting Information, Figure S6c. Tg = -77.0 o C for FIBOMA = 0 was determined previously 28 . The block dotted line represents the experimental data fitted to the Gordon-Taylor Equation.

  increased from -65 to 123 o C with FIBOMA ranging from 9.8 to 96.2 mol%. Interestingly, the presence of a minor fraction of My units considerably reduced the Tg of the copolymer (Tg = 81 o C for FMy = 0.21 in comparison to Tg,P(IBOMA) ~ 174 o C) while a moderate increase of the Tg was observed for a My-rich P(Mygrad-IBOMA) containing a minority of IBOMA units (Tg = -46 o C for FIBOMA = 0.20 and Tg,P(My) = -77 o C). Two explanations can be suggested: the long C6/C8 pendant groups of My units generating significant free volume; the gradient nature of the copolymer with a largely higher reactivity for My, resulting in the formation of short P(My) segments (elastomeric "knots") and the isolation of the isobornyl rings.

Figure 6 .

 6 Figure 6. TGA traces (N2 atmosphere, 10 °C.min -1 ) of final gradient copolymers a) My/IBOMA-20 and b) My/IBOMA-80, previously precipitated in excess methanol and dried under vacuum at 50 o C. Sample weight versus temperature is represented by the solid blue line whereas the dotted line represents the derivative of weight relative to the temperature versus temperature in order to determine precisely the temperature at which weight loss is most apparent (Tdec,max).

  My/IBOMA-82 chains remained SG1-capped despite a relatively high overall conversion (X1 = 64.2%). It can thus be deduced that about 30% of these chains underwent irreversible terminations, becoming more prevalent at high conversion, resulting notably in a relatively broad molecular weight Table3. Chain-extensions of (A) P(My-grad-IBOMA) macroinitiators (MI) with (B) IBOMA and/or My in 50 wt% toluene and (C) molecular characterization of the resulting chain-extended products.a) The index "1" is associated to the characteristics of the P(My-grad-IBOMA) macroinitiators (MI) whereas the index "2" refers to the features of the whole chain-extended diblock copolymer (MI + P(My) or P(IBOMA-co-My) segment added). b) Living molar fraction of MI chains capped by a SG1 group, measured by31 P NMR (FigureS5cin Supporting Information for the spectra). Standard deviation derived from the difference in macroinitiator Mn value between Mn,GPC and Mn,MHS. c) Predicted Mn,MHS,1 at X1 measured experimentally and calculated as follows: Mn,theo,X1 = (X1/100)Mn,theo,1 with Mn,theo,1 = 30 kg.mol - 1 at X = 100%. d) Targeted number-average molecular weight of the whole chain-extended diblock copolymer (MI block + extended block) at X = 100%.

( 1 )

 1 NMP of My initiated by the PEB-(SG1)2 difunctional initiator (telechelic poly(ethylene-stat-butylene) copolymer terminated with SG1 groups) 42 ; (2) Chain-extension of the P(My)-(SG1)2 macroinitiator with a IBOMA/Co mixture (Co = My or S comonomer, ~ 8-9 mol% relative to IBOMA). The synthesis route is depicted in Figure 1b. A triblock molecular architecture was targeted so that both ends of the rubbery mid-segment are theoretically anchored to a glassy domain 75 , mimicking thus the traditional micro-structure of styrene block copolymers such as SIS or SBS (I = isoprene, B = butadiene).

Figure 7 .

 7 Figure 7. Normalized GPC traces for the chain-extensions of a) My/IBOMA-82 with a IBOMA/My (93/7 mol%) mixture (experiment My/IBOMA-82-IBOMA/My) and b) My/IBOMA-44 with My (experiment My/IBOMA-44-My) at T = 105-115 o C in 50 wt% toluene. Schematic representations of the final chain-extended copolymers, where solid red and blue circles represent My and IBOMA units respectively, are given below the GPC chromatograms.

  (A) P(My)-(SG1)2 macroinitiators with (B) IBOMA/My and IBOMA/S mixtures at 115 o C in 50 wt% toluene and (C) features of the resulting triblock copolymers. a) The indexes "My" and "CE" refer, respectively, to the final characteristics of P(My)-(SG1)2 and the whole chain-extended triblock copolymer. b) Predicted Mn,MHS,My at XMy measured experimentally and calculated as follows: Mn,theo,XMy = (XMy/100)Mn,theo,My (Mn,theo,My = 143.4 and 139.8 kg.mol -1 for experiments My-35 and My-52, respectively). c) See caption e) in Table2. d) Targeted number-average molecular weight of the whole chain-extended triblock copolymer at XCE = 100%. e) Predicted Mn,MHS,CE of the whole chain-extended triblock copolymer at XCE, measured experimentally, and calculated as follows: Mn,theo,XCE = (XCE/100)(Mn,theo,CE-Mn,MHS,My)+Mn,MHS,My.

  (Supporting Information) by the tails on the low molecular weight edge of the GPC chromatograms (elution time ~ 23-25 min). • Characterization of the NMP-based Co/IBOMA-My-IBOMA/Co triblock copolymers. The Co/IBOMA-My-IBOMA/Co triblock copolymers investigated were characterized by two discrete transition regions indicative of each polymer segment type (Figure 8a). A lower Tg of about -60 o C corresponded to the transition of the P(My) mid-segment while a largely upper transition at around + 180 o C was assumed to indicate the softening of the IBOMA-rich segments. The lower Tg of these triblock

Figure 8

 8 Figure 8. a) DSC traces (second heating run) of triblock copolymers My-52-IBOMA/S (blue) and My-35-IBOMA/My (red). The numbers near the changes in slope correspond to the Tgs determined via the inflection method. b) Dynamic mechanical analysis of the sample My-52-IBOMA/S by torsional oscillation, yielding the storage modulus, the loss modulus and the damping factor versus temperature (0.15 Hz, 1% strain, 5 o C.min -1 , N2 atmosphere).

Figure 9 .

 9 Figure 9. Atomic force microscopy (AFM, experimental section) phase image (2 μm x 2 μm) under tapping mode of operation of the surface morphology of triblock copolymer My-35-IBOMA/My cast film. The dark domain represents the My component (color-coded height scale given to the right of the image).

- 1 .

 1 The rubbery plateau was then observed at T = 70-140 o C where G' slightly decreased from 3.6 to 2.2 MPa. At T > 140 o C, the modulus dropped due to the softening of the P(IBOMA) phase, the flexible P(My) chains being no longer held together by the rigid domains. The maximum tan δ peak appeared at 171 o C, which is consistent with the P(IBOMA)'s Tg measured previously (Tg,P(IBOMA) = 171-177 o C, Figure 5). Generally, this dynamic mechanical analysis (DMA) suggests that My-52-IBOMA/S may have an upper service temperature T ~ 140 o C with G' ≥ 2.2 MPa at T ≤ 140 o C.

Figure 10 .

 10 Figure 10. Tensile stress-strain curves of five My-52-IBOMA/S samples (same batch, color only used for differentiation) at room temperature and at a cross-head speed of 10 mm.min -1 . The average Young's modulus E, yield stress σY, tensile strength at break σB and tensile elongation at break εB, obtained from these curves are also given.

homopolymerization ( 1

 1 .4 ± 0.4 10 -5 s -1 ), stressing the strong influence of My kinetics over the copolymerization kinetics. P(My-grad-IBOMA) displayed a glass transition temperature between -65 and + 123 °C depending on the molar composition of IBOMA. A two-step thermal decomposition was observed by TGA, consisting of an early decomposition peak at Tdec,max = 242-307 o C, which results likely from the release of the isobornyl group, and a final decomposition peak at Tdec,max = 370-380 o C, due to the scission of the copolymer backbone. The chain-end fidelity of My-rich and IBOMA-rich P(My-grad-IBOMA)s was quantitatively deemed by 31 P NMR (69-74% of SG1-terminated chains) and confirmed via successful chainextensions in toluene with IBOMA and/or My (monomodal shift of the GPC traces). The satisfactory mediation of the My/IBOMA system by the SG1 nitroxide allowed subsequently the preparation of triblock copolymers, composed of a soft P(My) middle block and two outer IBOMA-rich segments. 1,4-P(My)-(SG1)2 macro-initiator (Mn,MHS = 35-52 kg.mol -1 , Đ = 1.45-1.53) was first synthesized using the PEB-(SG1)2 dialkoxyamine initiator in bulk at 120 o C, followed by the IBOMA/Co chainextension (Co = My or S co-monomer, < 9 mol%) at 115 o C in toluene. Displaying two distinct Tgs at -58 and + 181 o C, My/IBOMA-My-IBOMA/My triblock copolymer (Mn,MHS = 51 kg.mol -1 , Đ = 1.91, FIBOMA = 0.28) was studied by AFM, which revealed the micro-phase separation of the continuous P(My) domain and the disperse IBOMA-rich aggregates, despite an apparent disorganization. Lastly, rheological analysis as well as stress-strain tests were performed for the NMP-based S/IBOMA-My-IBOMA/S triblock copolymer (Mn,MHS = 95 kg.mol -1 , Đ = 2.23, FIBOMA = 0.36, FS = 0.05). In comparison to traditional styrenic block copolymers, an extended upper service temperature at around 140 o C was observed by DMA.

  comparison to the previous paper-based chapters, Chapter 4 is intentionally more technical and results-based. It first details the synthesis of the telechelic PI-(SG1)2 mid-segment using poly(ethylene-statbutylene)-(SG1)2 difunctional initiator. Then, the successful S and IBOMA chain-extensions from PI-(SG1)2 macro-initiator is presented as well as the characterization of the resulting well-tailored S-I-S and S/IBOMA-I-IBOMA/S triblock copolymers.

  SG1)2 dialkoxyamine at 115 o C in 50 vol% pyridine for 14-17 h led, after fractionation, to cis-1,4-rich PI-(SG1)2 poly(diene)s exhibiting Mn = 54-61 kg.mol -1 and Ð = 1.47-1.59. Well-defined PS-PI-PS triblock copolymers (Mn = 94.6-109.4 kg.mol -1 , Ð = 2.11-2.29, FS = 0.30-0.49) were synthesized via subsequent S chain-extensions from PI-(SG1)2 macro-initiators. Generally, these triblock copolymers displayed two glass transition temperatures (Tgs), a sub-zero one from -61 to -65 o C and an upper one from 102 to 104 o C, suggesting the segregation into two different phases of the flexible PI and rigid PS blocks. A micro-phase separation, exhibiting possibly a cylindrical or lamellar morphology, was observed by AFM for a synthesized PS-PI-PS having FS = 0.30. While a hard and brittle PS-PI-PS plastic containing 49 mol% of styrene units was mechanically characterized (σB = 10.8 ± 1.2 MPa, εB = 21 ± 7 %), softer and more extensible triblock copolymers having FS = 0.30-0.38 were obtained with σB = 2.8-4.1 MPa and εB = 376-449%. Despite a

- 1 -

 1 oxy (TEMPO) and the 2,2,5trimethyl-4-phenyl-3-azahexane-3-oxyl nitroxide (TIPNO), as mentioned in the literature review (Section D.2.3) 1-10 . Specifically, Harrisson et al. studied the NMP of isoprene initiated by various SG1-based

Ph. 1 .

 1 115 o C was selected as the reaction temperature since it was shown to provide a controlled polymerization at an acceptable polymerization rate 11 . Lastly, a high target number-average molecular weight at quantitative conversion (Mn,theo ~ 145 kg.mol -1 at XI = 100%) was chosen, allowing the synthesis of relatively long PI chains. Due to the low boiling point of isoprene (Tb,I ~ 34 o C), 15 mL high-pressure reactors were used, each fitted with plunger valve and thermowell, and placed in an oil bath, as shown in Ph. 1. Polymerization of isoprene using four capacity pressure tubes placed in an oil bath.

  o C, initiated by PEB-(SG1)2 and targeting Mn,theo ~ 145 kg.mol -1 at XI = 100%. a) Rounded values including the number-average molecular weight of PEB-(SG1)2 (Mn,PEB(SG1)2 = 5.7 kg.mol -1 ). b) Samples periodically taken at: 2 h, 4 h, 6 h and 8 h (PI1); 11 h, 13 h, 15 h and 17 h (PI2); 11 h, 13 h, 16 h, 19 h and 25 h (PI3). c) Polymerization performed in bulk. bulk. Consequently, it was decided to synthesize the PI-(SG1)2 mid-segment by relying on the following experimental conditions: initiation of I with PEB-(SG1)2 at 115 o C in 50 v% pyridine for about 13-16 h and targeting Mn,theo ~ 145 kg.mol -1 at quantitative conversion.

Figure 1 .

 1 Figure 1. (a) Mn versus isoprene conversion XI and (b) Đ versus XI for the isoprene homopolymerizations performed in bulk (experiments PI1 and PI2) and in pyridine (experiment PI3) at 115 o C and initiated by PEB-(SG1)2. The dotted black line indicates the theoretical Mn versus overall conversion based on the monomer to initiator ratio (Mn,theo ~ 145 kg.mol -1 at XI = 100% for every experiment). All experimental ID are listed in Table1.

a)

  Fractional precipitation using the non-solvent addition method, relying on the benzene (solvent) / methanol (non-solvent) pair (2cycles). b) Mn and Mw measured by GPC (Water Breeze, differential refractive index RI 2414 detector, 40 o C) with HPLC grade THF as the mobile phase (flow rate of 0.3 mL.min -1 ). Mn values determined by calibration with linear narrow molecular weight distribution PI standards.

Figure 2 .

 2 Figure 2. (a) DSC trace (second heating run) of the final dry PI4 sample (Tg,PI4 = -60.3 o C, determined via the inflection method) and (b) TGA traces (10 °C.min -1 ) of final dry PI4 sample under air (Tdec,max,PI4 = 361 o C) and nitrogen (Tdec,max,PI4 = 389 o C) atmospheres.

Figure 3 . 1 H

 31 Figure 3. 1 H NMR spectrum of the final dry PI4 polymer in CDCl3 (300 MHz, 32 scans, room temperature). *Signal at δ = 4.55-4.75 ppm corresponds to olefinic protons of 3,4-PI. ♣ Signals at δ = 4.80-5.00 ppm and 5.65-5.85 ppm correspond to olefinic protons of 1,2-PI.

  3-2.8 mol.L -1 ) chain-extensions from NMP-based PI-(SG1)2 samples (Mn = 46.8-61.0 kg.mol -1 , Ð = 1.41-1.58, 0.002-0.003 mol.L -1 ) were performed in 50 wt% toluene (4.8-5.7 mol.L -1 ) at T = 110-114 o C. Despite successful chain-extensions notably marked by the monomodal shift of the GPC chromatograms to lower elution times, indicating a

  degradation of PS-PI4-PS, starting at 271 o C, was observed. The formation of a highly reactive peroxy radical intermediate (ROO•), then abstracting a labile hydrogen which gives rise to the hydroperoxide species (ROOH), likely occurred throughout this decomposition process in an O2 environment28 . ThermalA. Styrene chain-extension from PI-(

a)

  Rounded values including the number-average molecular weight of PI-(SG1)2 macro-initiators (Mn,PI-(SG1)2 = 54.3-61.0 kg.mol -1 ). b) Styrene conversion XS determined by 1 H NMR in CDCl3. c) Molar fraction of S in the copolymer (FS) as determined by 1 H NMR in CDCl3. d) Mn and Mw determined by GPC calibrated with PI standards in THF at 40 °C.

Figure 4 .

 4 Figure 4. (a) DSC traces (second heating run) of final dry PS-PI4-PS and PS-PI6-PS. The numbers near the changes in slope correspond to the Tgs determined via the inflection method. (b) TGA traces (10 °C.min -1 ) of final dry PS-PI4-PS samples under air and N2 atmospheres. The characteristic temperatures Tdec,1 = 271-316 o C (onset of decomposition) and Tdec,max = 399-420 o C (temperature at which weight loss is most apparent) were determined. (c) Dynamic mechanical analysis of the samples PS-PI4-PS and PS-PI6-PS by torsional oscillation, yielding the storage modulus versus temperature (0.15 Hz, 1% strain, 5 o C.min -1 , N2 atmosphere, pictures of DMA test bars PS-PI4-PS (top) and PS-PI6-PS (bottom) prepared by dichloromethane casting).

  Figure 5 shows 2 μm x 2 μm scans of the surface morphology of PS-PI5-PS in two different regions. A clear phase separation was observed with the dark area corresponding to the continuous PI domain whereas the light area corresponded to the discrete PS phase. The domain structure may be cylindrical (disordered cylinder phases), although a definite assignment requires further study. Despite a relatively low PS content (30 mol%) in this triblock copolymer, a lamellar microstructure might be suggested as notably supported by a few examples in the literature 33,34 .

Figure 5 .

 5 Figure 5. Atomic force microscopy (AFM, experimental section) phase images (2 μm x 2 μm, two different regions shown for the same sample) under tapping mode of operation of the surface morphology of triblock copolymer PS-PI5-PS cast film (chloroform used). The dark domain represents the isoprene component (color-coded height scale given to the right of the image).

Figure 6 .

 6 Figure 6. Uniaxial tensile stress-strain curves at room temperature and a cross-head speed of 10 mm.min -1 of (a) four PS-PI5-PS specimens (FS = 0.30), (b) four PS-PI6-PS specimens (FS = 0.38) and (c) five PS-PI4-PS specimens (FS = 0.49), all prepared via a film casting process. (d) PS-PI4-PS cast films from dichloromethane solution, cut then into typical dog-bone shapes using a sharp blade. The average Young's modulus E, tensile stress at break σB and tensile elongation at break εB obtained from these curves are also given.

  mechanical properties at room temperature compared to the NMP-based S-My-S samples (Chapter 1, Section 1.4.3, Mn = 56.4-65.9 kg.mol -1 , Ð = 1.72-1.86, FS = 20.4-32.1 mol%), which exhibited σB < 0.8 MPa and εB < 190% (Chapter 1, Section 1.4.3, Figure 7''). This is likely due to the soft PI domain of PS-PI-PS, being much more entangled (Mn,PI = 54.3-61.0 kg.mol -1 ~ 10Me,PI with Me,PI = 4-6 kg.mol -1 , the entanglement molar mass of PI) than the flexible P(My) matrix (Mn,P(My) = 44.6 kg.mol -1 ~ 2Me,P(My) with Me,P(My) = 22-31 kg.mol -1 ) of the S-My-S triblock copolymers.

Figure 7 .

 7 Figure 7. Chemical hydrogenation of the carbon-carbon double bond of the poly(isoprene) mid-segment using p-toluenesulfonyl hydrazide (TSH, 2.4 equivalent to the double bond) with tributylamine (TBA, 2.5 equivalent to the double bond). 0.1 wt% of butylated hydroxytoluene (BHT) relative to the polymer was added as well.

Figure 8 . 1 H

 81 Figure 8. 1 H NMR spectra of the final dry PS-PI4-PS in CDCl3 (300 MHz, 32 scans, room temperature) (a) before and (b) after hydrogenation. The styrene-isoprene-styrene triblock copolymer structure is given at the top of the spectra. A hydrogenation degree (HD) of 76.3% was determined.

Figure 9 .

 9 Figure 9. Normalized GPC traces of PS-PI4-PS before (Mn = 103.1 kg.mol -1 , Ð = 2.11) and after (Mn = 97.4 kg.mol -1 , Ð = 2.27) hydrogenation and of PS-PI6-PS before (Mn = 109.4 kg.mol -1 , Ð = 2.18) and after (Mn = 101.8 kg.mol -1 , Ð = 2.23) hydrogenation.

Figure 10 .

 10 Figure 10. TGA traces (10 °C.min -1 ) of final dry PS-PI4-PS (blue curves, (a) and (a')) and PS-PI6-PS (red curves, (b) and (b')) samples before and after hydrogenation (HD = 76.3% and 77.3% respectively) under air and N2 atmospheres. The characteristic temperatures Tdec,1 = 265-364 o C (onset of decomposition) and Tdec,max = 394-436 o C (temperature at which weight loss is most apparent) were determined.

Figure 11 . 4 . 3 . 4 .

 11434 Figure 11. Uniaxial tensile stress-strain curves at room temperature and a cross-head speed of 10 mm.min -1 of (a) five PS-PI4-PS specimens (FS = 0.49) and (b) four partially hydrogenated PS-PI4-PS specimens (HD = 76.3%). The average Young's modulus E, tensile stress at break σB and tensile elongation at break εB obtained from these curves are also given in black.

AFigure 12 .

 12 Figure 2b); a last and minor decomposition step up to around 500 o C which might be due to the scission of

Table 4 .

 4 Chain-extension of (A) PI-(SG1)2 macro-initiator with (B) IBOMA/S mixture at 115 o C in 50 wt% toluene and (C) characteristics of the resulting triblock copolymer. a) Mn and Mw determined by GPC calibrated with PI standards in tetrahydrofuran (THF) at 40 °C. b) Determined by 1 H NMR in CDCl3. c) Previously purified by passing through a column of basic alumina mixed with 5 wt% calcium hydride. d) Targeted number-average molecular weight of the whole chain-extended triblock copolymer at quantitative conversion.

  previously synthesized by SG1-based NMP. This relatively high storage modulus observed at T < Tg,hard block for isoprene-based triblocks highlights the presence of a more entangled elastomeric continuous phase compared to My-based polymers. The softening of the IBOMA-rich segments was evidenced by the drop of G', starting at 155-160 o C. The maximum value of the damping factor appeared at 185 o C, which is consistent with the PI7-P(IBOMA-co-S)2 upper Tg measured previously via DSC (182 o C, Figure 13a). 232 Generally, the DMA test suggested that this IBOMA-I-IBOMA type triblock may well have an attractively high service temperature.

Figure 13 .

 13 Figure 13. (a) DSC trace (second heating run) of fractionated PI7-P(IBOMA-co-S)2. The numbers near the changes in slope correspond to the Tgs determined via the inflection method. (b) TGA traces (10 °C.min -1 ) of fractionated PI7-P(IBOMA-co-S)2 under air and N2 atmospheres. The characteristic temperatures Tdec,1 = 219-265 o C (onset of decomposition) and T ≥ 345 o C (temperatures between two degradation steps) were determined. (c) Dynamic mechanical analysis of fractionated PI7-P(IBOMAco-S)2 cast film (dichloromethane used) by torsional oscillation, yielding the storage modulus (G') and the damping factor versus temperature (0.15 Hz, 1% strain, 5 o C.min -1 , N2 atmosphere).

3 and Chapter 3 ,Figure 14 .

 314 Figure 14. Atomic force microscopy (AFM, experimental section) phase image under tapping mode of operation (left) and 3D topography (right, same region) of the surface morphology of fractionated PI7-P(IBOMA-co-S)2 cast film (chloroform used).

Figure 15 .

 15 Figure 15. Uniaxial tensile stress-strain curves at room temperature and at a cross-head speed of 10 mm.min -1 of four fractionated PI7-P(IBOMA-co-S)2 specimens (dichloromethane casting). The average Young's modulus E, tensile stress at 100% elongation σ100 and at break σB and tensile elongation at break εB obtained from these curves are also given.

( 12 )( 13 )Chapter 5 .

 12135 Lessard, B.; Aumand-Bourque, C.; Chaudury, R.; Gomez, D.; Haroon, A.; Ibrahimian, N.; Mackay, S.; Noel, M. -C.; Patel, R.; Sitaram, S.; Valla, S.; White, B.; Maric, M. Poly(ethylene-co-butylene)-b-(styrene-ran-maleic anhydride)2 compatibilizers via Nitroxide Mediated Radical Polymerization. Intern. Polymer Processing 2011, 26 (2), 197. Maric M. In Nitroxide Mediated Polymerization: From Fundamentals to Applications in Materials Science; Gigmes, D., Ed.; The Royal Society of Chemistry: Cambridge, 2016; Chapter 4, p. 153. Characterization of commercial styrenic triblock copolymers 5.1. Preface

Figure 1 .

 1 Figure 1. Normalized GPC traces (40 o C, THF as the mobile phase, PI standards) of S4-I-S4 (Mn = 98.3 kg.mol -1 , Ð = 1.29), S8-I-S8 (Mn = 97.2 kg.mol -1 , Ð = 1.25) and hydrogenated S8-I-S8 (Mn = 110.1 kg.mol -1 , Ð = 1.32).

Figure 2 .

 2 Figure 2. (a) DSC traces (second heating run) of S4-I-S4 and S8-I-S8. The numbers near the changes in slope correspond to the Tgs determined via the inflection method. TGA traces (10 °C.min -1 ) of S8-I-S8 under (b) air and (c) N2 atmospheres. The characteristic temperatures Tdec,1 = 258-371 o C (onset of decomposition) and Tdec,max = 370-436 o C (temperature at which weight loss is most apparent) were determined.

Figure 3 Figure 3 .

 33 Figure 3. Dynamic mechanical analysis of S4-I-S4 and (partially hydrogenated) S8-I-S8 by torsional oscillation, yielding the storage modulus (G') versus temperature (0.15 Hz, 1% strain, 5 o C.min -1 , N2 atmosphere). The rectangular bars were made either using hot press (HP) or from dichloromethane casting (cast film, CF).

Figure 4 .

 4 Figure 4. Uniaxial tensile stress-strain curves at room temperature and at a cross-head speed of 10 mm.min -1 of S4-I-S4 and (partially hydrogenated) S8-I-S8 specimens. The average Young's modulus E, tensile stress at break σB and tensile elongation at break εB obtained from these curves are also given in the table below the curves. The dog-bone style tensile specimens were made using either hot press (HP) or from dichloromethane casting (cast film, CF).

Figure 5 .

 5 Figure 5. Atomic force microscopy (AFM) phase image (2 μm x 2 μm) under tapping mode of operation of the surface morphology of S8-I-S8 cast film (chloroform used).

  , grey curve). A marked increase of G' was observed compared to the parent unsaturated S8-I-S8 cast film (black curve), either before or after the glass transition at T ~ 90-100 o C. Such a trend can be explained by the presence of ethylene/propylene sequences in the polymer chain after hydrogenation of the isoprene units, tending to decrease the mobility and the free volume of the triblock copolymer. Interestingly, at T > Tg,PS ~ 100 o C, no drop of G' was seen for the hydrogenated S8-I-S8 and it remained almost constant at about 4.5-6.5 MPa up to 150 o C.

Preliminary 1 HFigure 6 .

 16 Figure 6. (a) Uniaxial tensile stress-strain curves at room temperature and a cross-head speed of 10 mm.min -1 of four SEBS specimens prepared by hot press. The average Young's modulus E, tensile stress at 100% elongation σ100 and at break σB and tensile elongation at break εB obtained from these curves are also given. (b) Dynamic mechanical analysis of SEBS specimens prepared by hot press by torsional oscillation, yielding the storage modulus (G', black curve) and the damping factor (grey curve) versus temperature (0.15 Hz, 1% strain, 5 o C.min -1 , N2 atmosphere). (c) TGA traces (10 °C.min -1 ) of SEBS under air (grey) and nitrogen (black) atmospheres. The characteristic temperatures Tdec,1 = 260-341 o C (onset of decomposition) and Tdec,max = 362-433 o C (temperature at which weight loss is most apparent) were determined.

Figure 7 .

 7 Figure 7. Isothermal TGA traces of the commercial SEBS A1536 H under (a) air (T = 200-230 o C) and (b) N2 (T = 230-290 o C) atmospheres.

± 2 . 1

 21 Unlike the NMP-based S-I-S previously synthesized (Mn = 95-109 kg.mol -1 , Ð = 2.18-2.29, FS = 0.30-0.38), these commercial samples prepared anionically showed a marked strain hardening in the highstrain region (ε > 2000%), allowing the tensile stress to greatly increased. It may be explained by the narrower molecular weight distribution of these block copolymers (Ð < 1.30) where all the chains, exhibiting a similar length, can straighten. The high level of control of these triblock copolymers was also evidenced by the strong phase-separation of S-I-S possessing 16 mol% S, in comparison to the more disordered PI and PS microdomains observed for the triblocks made by NMP.Preparing commercial S-I-S specimens (FS ~ 0.16) via dichloromethane casting instead of compression molding allowed to increase their toughness (σB = 14.3 ± 2.1 MPa, εB = 4690 ± 325 %) as well as their storage modulus at T = 25-150 o C. The incomplete hydrogenation of this latter triblock (saturation of 67 mol% of the soft segment) made it stronger (σB = 18.4 ± 0.7 MPa) and more stable thermally under an inert environment, with a starting decomposition at about 370 o C. Moreover, the storage modulus of this hydrogenated S-I-S was significantly higher in the rubbery plateau (G' = 35-10 MPa) and did not dramatically drop at T > 100 o C (gradual decrease from 10 to 4 MPa), suggesting an extended upper service temperature compared to typical styrenic block copolymers.Targeting a styrenic block copolymer TPE which has a relatively high molar fraction of S (FS > 20 mol%) and its elastomeric mid-block hydrogenated appeared thereby to be relevant to meet the mechanical requirements at room temperature. A commercial and well-defined SEBS triblock copolymer (Mn = 87.5 kg.mol -1 , Ð = 1.16), originating from the complete hydrogenation of SBS and possessing almost as much styrene as ethylene/butylene units (FS = 0.48), showed a satisfactory stress-strain behavior with σB = 25.0 MPa and εB = 730 ± 40 %. Its acceptable thermal stability under inert atmosphere indicated seemingly a possible welding with the thermoplastic composite at T ≤ 275 o C.

  Maric's laboratories. Furthermore, the robustness of NMP towards polar monomers and a large variety of vinyl monomers was of importance since two main modifications were proposed: (1) functionalization of the poly(diene) via direct diene/functional monomer copolymerization (to aid the joining with the polar composite matrices); (2) substitution of PS by higher glass transition temperature (Tg) polymer (to increase the tensile stress and the upper service temperature of the triblock copolymer).Pending the acquirement of pressurized reactors to polymerize isoprene (low boiling point, Tb ~ 34 o C), β-myrcene (My), a non-volatile (Tb ~ 167 o C) and renewably-sourced 1,3-diene, was first considered as a promising building block for making the flexible segment of the TPE.Poly(β-myrcene) P(My) (number-average molecular weight Mn = 23-30 kg.mol -1 , dispersity Ð = 1.32-1.38) were first produced by NMP at 120 o C in bulk using the NHS-BlocBuilder (NHS-BB) monofunctional initiator (possessing the SG1 controlling nitroxide group). The active feature of these poly(diene)s (chain end capped by a nitroxide moiety) allowed styrene (S) chain-extensions in toluene at 110 o C. The resulting My-S diblock copolymers (Mn = 39-49 kg.mol -1 , Ð = 1.50-1.88, molar fraction of S, FS = 0.27-0.47) showed poor strain-stress properties (tensile stress at break σB < 1 MPa, elongation at break εB < 16%) with no plastic deformation. Two main reasons can explain this brittle behavior: a lack of entanglement density of the soft domain caused by the high P(My) entanglement molecular weight (Me,P(My) = 22-31 kg.mol -1 ); the diblock type molecular architecture where rubber P(My) chains are attached to a rigid PS chain at only one end (not physically cross-linked). The synthesis of S-My-S triblock copolymers exhibiting a higher P(My) Mn appeared thus essential and was made possible using the SG1-terminated poly(ethylene-co-butylene) (PEB-(SG1)2) difunctional initiator. Telechelic P(My)-(SG1)2 macro-initiator capped by two SG1 units was first synthesized using PEB-(SG1)2 in bulk at 120 o C, followed by the S chain-extension at 115 o C in toluene. Greater extensibility (εB = 90-180 %) was displayed by S-My-S (Mn = 56-66 kg.mol -1 with Mn,P(My) = 45 kg.mol -1 , Ð = 1.72-1.86 and FS = 0.20-0.32) compared to My-S, in spite of a similar and low tensile stress (σB < 0.8 MPa). Generally, these NMP-based My/S block copolymers displayed tensile properties largely below the industrial specifications.PS outer segments (Tg ~ 100 o C) were then replaced by higher Tg blocks, namely poly(isobornyl methacrylate) P(IBOMA) (Tg ~ 190 o C) to enhance the tensile strength and the service temperature of the candidate material. The conditions used to produce S-My-S were applied on a similar way to the synthesis of IBOMA-My-IBOMA type triblock copolymers (Mn = 51-95 kg.mol -1 , Đ = 1.91-2.23, FIBOMA = 0.28-0.36), containing a minor fraction of My or S in the rigid segments. Indeed, due to the difficulty to homopolymerize methacrylate by NMP, My or S was added as a controlling co-monomer during the IBOMA chain-extension. Two distinct Tgs were detected at about -60 o C and + 180 o C, suggesting a twophase system. The micro-phase separation of the continuous P(My) domain and the disperse IBOMA-rich aggregates was confirmed by atomic force microscopy (AFM). Moderate mechanical strength (σB = 3.9 ± 0.2 MPa) and relatively high irreversible extensibility (εB = 490 ± 30 %) were measured via tensile testing, which revealed a marked plastic region and the absence of strain hardening. A potential upper service temperature of these triblocks at about 140 o C was highlighted by the high and relatively constant storage modulus G' ≥ 2.2 MPa at T ≤ 140 o C. Consequently, the substitution of PS by harder P(IBOMA) segments was beneficial, allowing the enhancement of the mechanical properties of the block copolymer at room temperature and presumably at higher temperatures.Compared to the S-My-S triblocks, marked mechanical improvements were shown for IBOMA-My-IBOMA type copolymers, with notably a satisfactory extensibility (εB ~ 500%). However, the lack of tensile strength of these latter (σB ~ 4 MPa) remained detrimental for the industrial application. A poorly entangled elastomeric P(My) phase could be responsible of such a stress-strain behavior. The P(My) midsegment exhibited Mn ~ 50 kg.mol -1 at most (difficulty to produce high molar mass polymers by NMP under homogeneous conditions), which was only about the double of its entanglement molecular weight (Me,P(My) = 22-31 kg.mol -1 ). It was thus decided to substitute P(My) by poly(isoprene) (PI), having a much lower entanglement molecular weight (Me,PI = 4-6 kg.mol -1 ). The polymerization of isoprene (I) in a high-pressure reactor initiated by PEB-(SG1)2 at 115 o C in pyridine led to well-defined and active PI-(SG1)2 poly(diene)s (Mn = 54-61 kg.mol -1 , Ð < 1.60). S-I-S triblock copolymers (Mn = 95-109 kg.mol -1 , Ð = 2.11-2.29, FS = 0.30-0.49) were synthesized via subsequent S chain-extensions in bulk at 115 o C. The segregation into two different phases of PI and PS blocks was suggested by the observation of two Tgs, at about -60 o C and + 100 o C. An AFM study for S-I-S having FS = 0.30 revealed a micro-phase separation, exhibiting possibly a cylindrical or lamellar morphology. Whereas a rigid and brittle S-I-S sample containing 49 mol% S was identified (σB ~ 11 MPa, εB ~ 20%), softer and more extensible S-I-S plastics containing 30-38 mol% S were characterized with σB = 2.8-4.1 MPa and εB = 380-450%. Accordingly, substituting the flexible P(My) middle block of the candidate triblock copolymer by a more entangled PI block brought about an enhancement of its stressstrain properties. Despite an acceptable elongation at break, these NMP-based S-I-S having FS ≤ 0.38 showed a low and unsatisfactory tensile strength at break. Taking the positive results obtained for the My-based triblock copolymers after replacing PS by P(IBOMA) into account, a IBOMA-I-IBOMA type triblock copolymer (Mn = 94.0 kg.mol -1 , Ð = 1.76, FIBOMA = 0.35, FS = 0.06, S as a controlling co-monomer) was developed in a similar way as S-I-S. It exhibited a remarkable εB = 1360 ± 210 % and a moderate σB = 11.4 ± 0.6 MPa. Moreover, it displayed a potential upper service temperature higher than 150 o C due to the glass transition of the IBOMA-rich blocks occurring at 180-185 o C. Nonetheless, the presence of IBOMA units in the outer segments lowered the decomposition temperature of the triblock copolymer due to the thermal instability of the isobornyl ring (onset of decomposition at 220 o C under air and 265 o C under N2). Generally, this IBOMA-I-IBOMA type triblock was the closest to the benchmarks established in this study. However, its tensile strength at break remained below the industrial requirement. Moreover, its early thermal degradation renders it unusable for a subsequent welding process where T = 240-280 o C must be applied.

  broad molecular weight distribution of the experimental triblocks (Ð = 1.72-2.29), indicating the presence of a fraction of short chains in the characterized samples, which could act as a plasticizer and could disrupt the orientation and alignment of the chains in the direction of the load. Nonetheless, this thesis represents a series of advances related to the development of novel polymers and molecular architectures via a controlled radical polymerization. Several new contributions to the field of NMP have resulted from this project, including in particular: -Optimized NMP of My at 120 o C in bulk using the NHS-BB monofunctional initiator resulted in low dispersity (Đ = 1.1-1.4) P(My)-SG1. -Well-controlled statistical My/S (rMy = 1.88-2.90, rS = 0.25-0.47) and My/GMA (rMy = 0.48-0.80, rGMA = 0.50-0.71), and gradient My/IBOMA (rMy = 1.90-2.16, rIBOMA = 0.02-0.07) nitroxide-mediated copolymerizations. -Use of My (≤ 20 mol% in the initial feed) as a potential controlling co-monomer, allowing to enhance the control of GMA and IBOMA polymerizations by NMP. -Preparation of an amphiphilic polymer able to self-organize into micelles in water, after treating a NMP-based My/GMA diblock copolymer with morpholine. -Synthesis of telechelic P(My)-(SG1)2 (Mn = 35-52 kg.mol -1 ) and PI-(SG1)2 (Mn = 53-61 kg.mol - 1 ) poly(diene)s using the PEB-(SG1)2 dialkoxyamine. -Development of NMP-based IBOMA-My-IBOMA, S-I-S and IBOMA-I-IBOMA type triblock copolymers exhibiting valuable stress-strain properties (σB = 2.8-11.4 MPa, εB = 20-1360 %).

- 2 (

 2 project. (a) Molar fraction of hard segment (S, IBOMA or both) in the copolymer. (b) Onset of decomposition. (c) P(IBOMA) outer segments possessed a small fraction of S units. (d) Hydrogenation degree (HD) = 76 mol%. (e) HD = 67 mol%.CONCLUSION GÉNÉRALECette thèse doctorale avait pour objectif de synthétiser un élastomère thermoplastique (TPE) permettant de fournir une jonction souple entre des matrices composites polaires (réservoir et jupettes d´un lanceur). Ce TPE, assemblé aux composites par collage ou par soudage, doit être utilisé afin de transmettre principalement des forces de cisaillement qui se produisent au niveau de la jonction. Il doit répondre à un cahier des charges précis, incluant :-Un comportement en contrainte-taux d'extension semblable à celui d´élastomères vulcanisés.-Une perte minimale des propriétés mécaniques à T ≤ 80 o C. -L'absence de fluage et de dégradation lors du procédé d´assemblage (T ~ 160-180 o C pendant plus de 10 h pour un collage, T ~ 240-280 o C pendant quelques minutes pour un soudage). Les copolymères triblocs de type styrénique, constitués d´un segment poly(diène) (domaine continu) et de deux segments poly(styrène) (PS) (domaine isolé), furent ciblés. La polymérisation radicalaire contrôlée par des nitroxydes (NMP) fut sélectionnée pour sa facilité à mettre en oeuvre, son efficacité à contrôler la polymérisation des 1,3-diènes et l´expertise des laboratoires du Professeur Maric. De plus, la robustesse de la NMP vis-à-vis des monomères polaires et d´une grande variété de monomères vinyliques fut primordiale puisque deux principales modifications furent proposées : (1) fonctionnalisation du poly(diène) par copolymérisation du diène avec un monomère fonctionnel, afin d´encourager l´assemblage ultérieur avec les matrices composites polaires ; (2) remplacement du PS par un polymère ayant une plus haute température de transition vitreuse (Tg), dans le but d´augmenter la résistance en traction et la température de service maximale du copolymère tribloc. Dans l'attente de réacteurs haute pression permettant de polymériser l´isoprène (température d´évaporation Tb ~ 34 o C), le β-myrcène (My), un 1,3-diène biosourcé et non volatile (Tb ~ 167 o C), fut considéré dans un premier temps pour préparer le bloc souple du TPE. Le poly(β-myrcène) P(My) (masse molaire moyenne en nombre Mn = 23-30 kg.mol -1 , dispersité Ð = 1,32-1,38) fut tout d´abord produit par NMP en utilisant l´amorceur monofonctionnel NHS-BB (possédant un groupe nitroxyde SG1). Le caractère actif de ce poly(diène) (chaîne terminée par une unité nitroxyde) permit l´extension de chaîne avec du styrène (S). Les copolymères diblocs My-S ainsi formés (Mn = 39-49 kg.mol -1 , Ð = 1,50-1,88, fraction molaire en styrène FS = 0,27-0,47) montrèrent de faibles propriétés en traction uniaxiale (résistance à la rupture en traction σB < 1 MPa, allongement à la rupture en traction εB < 16%) avec notamment l´absence de déformation plastique. Le manque d´enchevêtrement du domaine flexible causé par la haute masse molaire d´enchevêtrement de P(My) (Me,P(My) = 22-31 kg.mol - 1 ) peut expliquer ce comportement fragile. De plus, l´architecture moléculaire de type dibloc, où les chaînes souples P(My) sont attachées aux chaînes rigides PS que d´un seul bout, pourrait empêcher la réticulation physique du réseau. La synthèse de copolymères triblocs S-My-S, ayant un segment central P(My) plus long, apparaissait alors essentielle et fut possible par l´utilisation de l´amorceur bifonctionnel poly(éthylène-ran-butylène) terminé par des groupes nitroxyde SG1 (PEB-(SG1)2). Un macro-amorceur téléchélique P(My)-(SG1)2 fut ainsi préparé via le PEB-(SG1)2, suivi d´une extension de chaîne avec du styrène. Une plus grande extensibilité (εB = 90-180 %) fut mesurée pour S-My-S (Mn = 56-66 kg.mol -1 avec Mn,P(My) = 45 kg.mol -1 , Ð = 1,72-1,86 et FS = 0,20-0,32) en comparaison à My-S. La résistance en traction de ces triblocs (σB < 0,8 MPa) fut toutefois similaire à celle des copolymères diblocs. De manière générale, ces copolymères à base de My et S présentèrent des caractéristiques en traction bien en deçà des spécifications industrielles. Les segments PS (Tg ~ 100 o C) furent par la suite remplacés par des segments poly(méthacrylate d´isobornyle) (P(IBOMA)) ayant une Tg plus haute (~ 190 o C). Le but fut d´améliorer la résistance mécanique et d´étendre la température de service du matériau candidat. Les conditions expérimentales utilisées pour la fabrication des triblocs S-My-S furent employées de la même manière pour la synthèse de copolymères triblocs de type IBOMA-My-IBOMA (Mn = 51-95 kg.mol -1 , Đ = 1,91-2,23, FIBOMA = 0,28-0,36), contenant une minorité d´unités My ou S dans ses blocs rigides. Deux Tgs distinctes furent détectées à environ -60 o C et + 180 o C, suggérant un système polymère ségrégé en deux phases. La micro-séparation de phase du domaine continu P(My) et d´agrégats isolés, riches en unité IBOMA, fut confirmée par microscopie de force atomique (AFM). σB = 3,9 ± 0,2 MPa et εB = 490 ± 30 % furent également mesurés par des tests de traction, qui montrèrent une déformation plastique permanente marquée et l´absence de durcissement à de hauts taux d´élongation. Une possible température de service maximale de l´ordre de 140 o C fut mise en évidence par la stabilité du module de conservation G´ ≥ 2,2 MPa à T ≤ 140 o C. La substitution de PS par P(IBOMA), plus dur, fut ainsi concluante, permettant une amélioration des propriétés mécaniques à température ambiante et probablement à plus haute température. Par rapport aux matériaux S-My-S, les copolymères de type IBOMA-My-IBOMA montrèrent de meilleures propriétés mécaniques, avec notamment εB ~ 500%. Néanmoins, le manque de résistance mécanique (σB ~ 4 MPa) constitua une lacune importante pour l´application industrielle. Une phase P(My) faiblement enchevêtrée pourrait être à l´origine d´un tel comportement en traction. En effet, Mn ~ 50 kg.mol - 1 fut au maximum obtenu pour le segment P(My) (difficulté de produire des polymères de hautes masses molaires par NMP en masse ou dans un solvant), ce qui ne représente que le double environ de sa masse molaire d´enchevêtrement. La substitution de P(My) par le poly(isoprène) (PI), dont la masse molaire d´enchevêtrement est bien plus faible (Me,PI = 4-6 kg.mol -1 ), fut par conséquent décidée. La polymérisation de l´isoprène dans un réacteur haute-pression, amorcée par PEB-(SG1)2 dans de la pyridine, permit d´obtenir des poly(diène)s PI-(SG1)2 bien définis et actifs (Mn = 54-61 kg.mol -1 , Ð < 1,60). Des copolymères triblocs S-I-S (Mn = 95-109 kg.mol -1 , Ð = 2,11-2,29, FS = 0,30-0,49) furent synthétisés par la suite via l´extension de chaîne de PI-(SG1)2 avec du styrène. La ségrégation des blocs PI et PS en deux phases différentes fut suggérée par la détection de deux Tgs, à environ -60 o C et + 100 o C. L´étude AFM d´un S-I-S contenant 30 mol% de styrène révéla une séparation de phase, possiblement cylindrique ou lamellaire, à l´échelle du micromètre. Alors qu´un tribloc S-I-S riche en styrène (FS = 0,49) fut dur et fragile (σB ~ 11 MPa, εB ~ 20%), des matériaux S-I-S constitués de segments PS plus courts (FS = 0,30-0,38) montrèrent une plus grande souplesse et une extensibilité plus importante avec σB = 2,8-4,1 MPa et εB = 380-450%. La substitution du segment P(My) du copolymère tribloc par un segment PI plus enchevêtré permit ainsi d'améliorer ses propriétés mécaniques en traction. Malgré une extensibilité acceptable, la résistance à la rupture en traction des triblocs S-I-S ayant FS ≤ 0,38 resta problématique. Les résultats concluants obtenus pour les polymères à base de My, après le remplacement des blocs PS par des blocs P(IBOMA) plus rigides nous poussèrent à préparer de manière similaire un copolymère tribloc de type IBOMA-I-IBOMA (Mn = 94,0 kg.mol -1 , Ð = 1,76, FIBOMA = 0,35, FS = 0,06, S utilisé en tant que co-monomère lors de l'extension de chaîne). Ce dernier se caractérisa par une grande ductilité (εB = 1360 ± 210 %) et par une résistance mécanique modérée (σB = 11,4 ± 0,6 MPa). Il présenta une température de service maximale potentielle supérieure à 150 o C, du fait de la transition vitreuse des blocs durs riches en IBOMA, se produisant à 180-185 o C. Toutefois, la température de décomposition de ce tribloc fut relativement basse (début de dégradation à 220 o C sous air et à 265 o C sous azote), à cause de l'instabilité thermique des groupes isobornyle. De manière générale, IBOMA-I-IBOMA représenta le copolymère à blocs synthétisé par NMP le plus concluant, malgré une résistance à la rupture en traction encore trop faible et une décomposition thermique précoce, ne permettant pas de réaliser un soudage pour lequel T = 240-280 o C doit être appliquée. Le tableau ci-dessus (page 255) récapitule les principales caractéristiques des copolymères à blocs expérimentaux et commerciaux étudiés lors de ce projet. Par rapport à la problématique industrielle, deux principales modifications furent réalisées sur les polymères synthétisés : -L'hydrogénation incomplète (~ 75%) du segment PI de triblocs S-I-S à pression ambiante et en utilisant l'agent diimide. Une plus grande stabilité thermique des copolymères partiellement saturés fut mise en évidence sous atmosphère inerte, en comparaison aux S-I-S originaux. De plus, σB augmenta significativement suite à cette hydrogénation chimique, en dépit d'une perte notable d'extensibilité. Fonctionnalisation de P(My) : l'incorporation de groupes fonctionnels au sein du segment souple fut entreprise pour favoriser l'assemblage avec les matrices composites polaires. La copolymérisation de My avec le méthacrylate de glycidyle, amorcée par NHS-BB à 120 o C, permit d'obtenir un polymère flexible bien défini et portant des unités époxydes. Lors du procédé d'assemblage, la présence de groupes époxydes au sein de l'élastomère pourrait soit augmenter la miscibilité à l'interface afin de favoriser l'inter-diffusion des chaînes (soudage), soit permettre la formation de liaisons covalentes à l'interface (collage). De manière générale, les différents copolymères à blocs produits par NMP lors de ce projet doctoral ne satisfirent pas le cahier des charges d'ArianeGroup, du fait notamment de leur relativement faible résistance mécanique. Cette déficience peut s'expliquer par l'absence d'une phase de durcissement de ces copolymères à de hauts taux d'extension, ne permettant pas d'accroitre significativement la résistance à la rupture en traction. La haute dispersité de ces matériaux préparés par NMP (Ð = 1,72-2,29) pourrait être un facteur clé afin de comprendre ce comportement mécanique. Elle indique la présence de chaînes courtes au sein du réseau polymère, fragilisant la structure du copolymère et empêchant possiblement l'alignement de ses chaînes dans le sens de la traction. Néanmoins, cette thèse apporta des contributions scientifiques dans le domaine de la NMP. Les avancées suivantes peuvent être citées en particulier : -Synthèse optimisée de P(My) (Đ = 1,1-1,4) en utilisant l'alcoxyamine NHS-BB. -Copolymérisation controlée des couples My/S, My/GMA et My/IBOMA par NMP. -Utilisation de My en tant que co-monomère (≤ 20 mol% dans le mélange initial) afin d'améliorer le contrôle de la polymérisation des monomères GMA et IBOMA par NMP. -Préparation de poly(diène)s téléchéliques P(My)-(SG1)2 (Mn = 35-52 kg.mol -1 ) et PI-(SG1)Mn = 53-61 kg.mol -1 ) en utilisant l´alcoxyamine PEB-(SG1)2. -Synthèse par NMP de copolymères de type IBOMA-My-IBOMA, S-I-S et IBOMA-I-IBOMA présentant des propriétés intéressantes en traction (σB = 2,8-11,4 MPa, εB = 20-1360 %). 260 PERSPECTIVES Although the various diene-based block copolymers synthesized by nitroxide-mediated polymerization (NMP) in this study did not meet the industrial requirements, marked improvements were noticed when using poly(isoprene) (PI) as the soft middle segment and hard outer segments rich in isobornyl methacrylate (IBOMA). IBOMA-I-IBOMA type triblock copolymer (Mn = 94 kg.mol -1 , Ð = 1.76, FIBOMA = 0.35 with a minor fraction of styrene (S) units in the outer blocks) showed a satisfactory elongation at break (εB = 1360 ± 210 %) combined to a relatively high upper service temperature at about 150 o C. However, its low tensile strength at break (σB = 11.4 ± 0.6 MPa) and thermal stability (onset of decomposition Tdec,onset at 220 o C and 265 o C under air and N2, respectively) were detrimental. This experimental isoprene-based copolymer can thus be used as a benchmark for the preparation of optimized triblock copolymers, with respect to the ArianeGroup specifications.

2 .

 2 Replacement of P(IBOMA) by a high glass transition temperature (Tg) and decomposition temperature (Tdec) polymer: the early thermal degradation of P(IBOMA), presumably due to the release of the isobornyl group, is problematic 3 . Its substitution by a polymer having a higher thermal stability while being at least as hard (Tg ≥ Tg,P(IBOMA) ~ 180-190 o C 3 ) would be valuable. For instance, the use of poly(αmethylene-γ-butyrolactone) P(MBL) (Tg ~ 190-200 o C, atactic 4,5,6

Accordingly, the preparation 1 .

 1 by NMP and after post-polymerization treatment of a fully hydrogenated MBL-I-MBL copolymer can be targeted in the future. A schematic representation of this latter triblock copolymer is given in the figure below (page 262). Compared to the IBOMA-I-IBOMA sample, higher Mn (> 100 kg.mol -1 ) and more importantly lower D (< 1.5) could be targeted for MBL-I-MBL due to the synthesis of the first PI mid-block by NMP in a dispersed medium followed by its chainextension with MBL, presumably easier to control than IBOMA. The preparation of a narrow molecular 264 PERSPECTIVES Les divers copolymères à blocs à base de diène synthétisés par polymérisation radicalaire contrôlée par des nitroxydes (NMP) ne satisfirent pas le cahier des charges d´ArianeGroup. Néanmoins, des améliorations importantes des propriétés de ces polymères furent constatées en utilisant le poly(isoprène) (PI) en tant que segment souple et des segments durs composés majoritairement de méthacrylate d´isobornyle (IBOMA). Le copolymère tribloc de type IBOMA-I-IBOMA (Mn = 94 kg.mol -1 , Ð = 1,76, FIBOMA = 0,35 avec une minorité d´unités styrène (S) dans les blocs rigides) montra un allongement à la rupture en traction (εB = 1360 ± 210 %) satisfaisant ainsi qu´une potentielle température de service maximale d´environ 150 o C. Sa faible résistance à la rupture en traction (σB = 11,4 ± 0,6 MPa) associée à son instabilité thermique (début de la dégradation Tdec,onset à 220 o C sous air et 265 o C sous azote) furent toutefois préjudiciables. Ce copolymère à base d´isoprène peut ainsi être utilisé comme point de départ pour la préparation de copolymères à blocs plus performants, selon les exigences industrielles.Trois principales modifications peuvent être considérées afin de produire par NMP le TPE candidat, à savoir : Synthèse du PI en milieu dispersé : L´application de la technique NMP en masse ou dans des solvants organiques se confronte à la difficulté de produire des polymères de haute masse molaire moyenne.Cela fut mis en évidence dans cette étude en ce qui concerne la synthèse de PI bien définis ayant une masse molaire moyenne en nombre (Mn) maximale comprise entre 50 et 60 kg.mol -1 lorsque la NMP en masse ou dans la pyridine fut employée. Augmenter la longueur moyenne de chaîne du bloc souple permettrait d´obtenir une matrice plus enchevêtrée et de ce fait, un copolymère tribloc plus résistant et tenace. A cette fin, la polymérisation de l´isoprène par NMP dans un milieu dispersé (suspension de nano-et microparticules) peut être envisageable. Il a été démontré que le compartimentage de ce type de polymérisation, où chaque particule agit comme un nano-réacteur, peut améliorer le contrôle par NMP avec une faible dispersité des chaînes (Ð) et un degré de fonctionnalité élevé des bouts de chaînes (groupe nitroxyde permettant le réamorçage). Une distribution étroite des masses molaires peut être maintenue à de hautes conversions, ce qui permet la synthèse de polymères bien définis ayant un Mn élevé 1 . La méthode NMP en mini-émulsion 2 , en utilisant un amorceur hydrosoluble ou lipophile, pourrait par exemple être considérée afin de préparer un latex PI « vivant » ayant Mn > 60 kg.mol -1 et Ð < 1,3.

2 . 3 .•

 23 Remplacement de P(IBOMA) par un polymère possédant de hautes températures de transition vitreuse (Tg) et de décomposition (Tdec) : La dégradation thermique précoce de P(IBOMA), vraisemblablement due au détachement du groupe isobornyle, est problématique 3 . La substitution de P(IBOMA) par un polymère ayant une haute stabilité thermique, tout en étant au moins aussi dur (Tg ≥ Tg,P(IBOMA) ~ 180-190 o C 3 ) serait bénéfique. Par exemple, le poly(α-méthylène-γ-butyrolactone) P(MBL) (Tg ~ 190-200 o C, atactique 4,5,6 ) constituerait un substitut intéressant puisqu'il est stable sous atmosphère inerte (Tdec,onset ~ 340 o C) 6 . Sa haute Tg assurerait une température de service maximale bien plus élevée que celle des TPE de référence de type styrénique ainsi qu´une importante résistance à la rupture en traction. De plus, MBL fut polymérisé de manière contrôlée par RAFT (Mn = 22 kg.mol -1 , Ð = 1,21) 6 , indiquant la prédisposition de P(MBL) à être synthétisé par polymérisation radicalaire contrôlée. Il peut être ajouté que l´addition de deux blocs P(MBL) purs serait possible alors que le monomère IBOMA fut toujours copolymériser avec une minorité de styrène afin d´améliorer le contrôle de l´extension de chaîne (homopolymérisation délicate des méthacrylates via NMP). Il se peut que la présence d´unités styrène au sein des segments rigides diminue la résistance mécanique et la performance à haute température du copolymère tribloc du fait de la relativement basse Tg du poly(styrène) PS (~ 95-100 o C) 7 . P(MBL) est cité en tant qu´exemple. Des dérivés du P(MBL), tels que le poly(γ-méthyl-αméthylène-γ-butyrolactone) (P(MeMBL), Tg = 210-220 o C, MeMBL copolymérisé avec du styrène par émulsion RAFT 8,9,10 ) ou le poly(α-méthylène-γ,γ-diméthyl-γ-butyrolactone) (P(Me2MBL), Tg ~ 210 o C, Tdec,onset ~ 340 o C sous N2, Me2MBL contrôlé par RAFT 6 ) pourraient être utilisés de la même manière. D´autres types de polymères, synthétisables par un procédé radicalaire et ayant des hautes Tg et Tdec, peuvent aussi être considérés comme les dérivés du poly(styrène). Le poly(α-méthyl-p-méthylstyrène) par exemple présente Tg = 165-180 o C et Tdec,onset ~ 320 o C sous atmosphère inerte 11 . Il doit être noté que l´incompatibilité de ces potentiels blocs durs avec le bloc souple central est assumée, offrant une séparation de phase nette du copolymère (étude approfondie nécessitant l´usage des paramètres de solubilité). Hydrogénation du segment souple PI : Comme réalisée dans cette étude pour des triblocs S-I-S expérimentaux et commerciaux, la saturation du poly(isoprène) apparait essentielle. La plus haute stabilité thermique et la plus grande résistance à la rupture en traction de ces copolymères après hydrogénation sont avantageux, malgré une perte notable de leur extensibilité. Une plus haute température de service maximale peut également être apparente, du fait d´une séparation de phase plus forte du copolymère tribloc hydrogéné (plus grande incompatibilité entre les domaines durs et flexibles) 12 . L´hydrogénation chimique employée dans ce projet et utilisant le diimide, résulta en un remplacement partiel des unités isoprène en unités éthylène/propylène. Appliquée dans l´industrie, l´hydrogénation catalytique de PI en utilisant un métal de transition et du dihydrogène H2 sous haute-pression permettrait de saturer complètement le bloc poly(diène) et d´améliorer ainsi les propriétés thermiques et mécaniques du copolymère tribloc. Par conséquent, la préparation par NMP et après modification chimique post-polymérisation d´un copolymère MBL-I-MBL entièrement hydrogéné serait une option envisageable pour un travail futur. Une représentation schématique de ce polymère en question est donnée dans la figure ci-dessous. Par rapport au tribloc IBOMA-I-IBOMA, une plus grande masse molaire moyenne (Mn > 100 kg.mol -1 ) et surtout une plus faible dispersité (Ð < 1,5) pourraient être atteintes pour MBL-I-MBL puisque le segment PI préliminaire serait synthétisé par NMP en milieu dispersé et que son extension de chaîne avec MBL serait probablement plus facilement contrôlable qu'avec IBOMA. La synthèse d´un copolymère tribloc ayant une faible dispersité pourrait être essentielle afin d´obtenir des propriétés mécaniques similaires à celles des TPE commerciaux traditionnels. De manière générale, les caractéristiques suivantes seraient attendues pour ce MBL-I-MBL bien défini : -Comportement en contrainte/déformation à température ambiante : σB > 20 MPa, εB > 600 % dont le domaine plastique est suivi d'une phase de durcissement, provoquant un accroissement significatif de la résistance à la rupture (faible Ð encourageant l´orientation et l´alignement des chaînes dans la direction de traction). -Comportement thermique : écoulement caoutchoutique à T ~ 170-180 o C, décomposition débutant à T ~ 280-300 o C sous air et T ~ 340-360 o C sous atmosphère inerte. -Performance mécanique à haute température : température de service maximale T ~ 170-180 o C (perte minimale des propriétés mécaniques à T < 170 o C). Par la suite, l´assemblage du TPE candidat avec le réservoir et les jupettes du lanceur (composites thermoplastiques ou thermodurcissables relativement polaires) serait réalisé par collage ou par soudage. Schéma récapitulatif du copolymère tribloc linéaire envisageable pour un travail futur. Di-I correspond à un amorceur bifonctionnel. Materials. β-Myrcene (My, ≥ 90 %), basic alumina (Al2O3, Brockmann, Type I, 150 mesh), calcium hydride (CaH2, 90-95 % reagent grade) and 1,4-dioxane (≥ 99 %) were purchased from Sigma-Aldrich and used as received. Toluene (≥ 99 %), methanol (MeOH, ≥ 99 %) and tetrahydrofuran (THF, 99.9 % HPLC grade) were obtained from Fisher Scientific and used as received. 2-Methyl-2-[N-tert-butyl-N-(1-diethoxyphosphoryl-2,2-dimethylpropyl)-aminoxy]-N-propionyloxy-succinimide, also known as NHS-BlocBuilder (NHS-BB), was prepared according to a published method 1 from 2-(tert-butyl[1-(diethoxyphosphoryl)-2,2-dimethylpropyl]aminooxy)-2-methylpropionic acid, also known as MAMA-SG1 (BlocBuilder TM , BB, 99 %, purchased from Arkema and used without further purification), N-hydroxy-succinimide (NHS, 98 %, purchased from Aldrich and used as received) and N,N'dicyclohexylcarbodiimide (DCC, 99 %, purchased from Aldrich and used as received). N-tert-Butyl-N-[1diethylphosphono-(2,2-dimethylpropyl)] nitroxide (SG1, > 85 %) was kindly donated by Noah Macy of Arkema and used as received. Styrene (S, 99 %) was obtained from Fisher Scientific and was purified to remove the inhibitor (ptert-butylcatechol) by passing through a column of basic alumina mixed with 5 weight % calcium hydride and then stored in a sealed flask under a head of nitrogen in a refrigerator until needed. The deuterated chloroform (CDCl3, 99.8 %) used as a solvent for proton and carbon nuclear magnetic resonance ( 1 H and 13 C NMR) was obtained from Cambridge Isotopes Laboratory.

( 2 .

 2 043 g, 19.616 mmol) were then injected with a disposable needle into the reactor. For this experiment, the initial molar ratio of monomers and NHS-BB was calculated to give theoretically a My/S copolymer sample with target number-average molecular weight Mn,theo = (MMyfMy,0 + MSfS,0) DP = 30 kg.mol -1 at complete overall conversion with DP = ([My]0 + [S]0) / [NHS-BB]0 = 231, the average degree of polymerization.

( 6 )Figure S1a .

 6S1a Figure S1a. Gravimetric conversion Xgrav and 1 H NMR conversion XNMR versus reaction time t for the polymerization My-6.

Figure S2a. 1 HFigure S3a. 13 C

 113 Figure S2a. 1 H NMR spectrum of P(My) My-2 in CDCl3 (300 MHz) at room temperature.

Figure S4a. 1 H

 1 Figure S4a.1 H NMR spectrum in CDCl3 (300 MHz) at room temperature of dried P(My-b-S) diblock copolymer My-11-Sa after recovery, using methanol as the precipitant.

Figure S5a. 81

 81 Figure S5a. 81 MHz 31 P NMR spectrum of dried P(My-stat-S)-SG1 macroinitiator My/S-07 in CDCl3 at room temperature.

Figure S7a .

 S7a Figure S7a. Tensile stress-strain curves of five P(My-b-S) diblock copolymer specimens from My-11-Sb sample with Mn,P(My) = 22.8 kg.mol -1 and Mn,PS = 15.7 kg.mol -1 at room temperature and 1 mm.min -1 . The red crosses denote failure points.

Figure S8a .

 S8a Figure S8a. Tensile stress-strain curves of five P(My-b-S) diblock copolymer specimens from My-12-S sample with Mn,P(My) = 30.4 kg.mol -1 and Mn,PS = 12.2 kg.mol -1 at room temperature and 1 mm.min -1 . The red crosses denote failure points.

Figure S9a .

 S9a Figure S9a. Individual monomer conversions, determined by 1 H NMR in CDCl3, versus reaction time t for the statistical copolymerization My/S-50.

Figure S10a .

 S10a Figure S10a. Semi-logarithmic kinetic plots of ln((1-XNMR) -1 ) (XNMR = XMyfMy,0 + XSfS,0 ) versus polymerization time t for the homopolymerizations My/S-0 and My/S-100 and the various My/S statistical copolymerizations at 110 °C in bulk initiated by NHS-BB and targeting Mn,theo = 30 kg.mol -1 at X = 1.0. Note that the scale of the yaxis is discontinuous (red dotted line) in order to display the higher values.

Figure S11a .

 S11a Figure S11a. Mn determined by GPC relative to PS standards in THF and corrected with appropriate MHS coefficients versus average NMR conversion XNMR for the homopolymerizations My/S-0 and My/S-100 and the various My/S statistical copolymerizations at 110 °C in bulk initiated by NHS-BB and targeting Mn,theo = 30 kg.mol -1 at X = 1.0. The red dashed line indicate the theoretical Mn (Mn,theo) versus overall conversion based on the monomers to initiator ratio.

Figure S12a .Figure S13a .

 S12aS13a Figure S12a. Đ average NMR conversion XNMR for the homopolymerizations My/S-0 and My/S-100 and the various My/S statistical copolymerizations at 110 °C in bulk initiated by NHS-BB and targeting Mn,theo = 30 kg.mol -1 at X = 1.0.

Figure S14a .

 S14a Figure S14a. TGA traces of a) P(My) My-12, b) PS My/S-100, c) P(My-b-S) My-10-S and d) P(My-b-S) My-11-Sa. Measurements performed under nitrogen flow at a ramp rate of 10 °C.min -1 . Sample weight versus temperature is represented by the solid line whereas the dotted line represents the derivative of weight relative to the temperature versus temperature in order to determine precisely the temperature at which weight loss is most apparent (Tdec,max).

Figure S16a .•

 S16a Figure S16a. DSC traces (second heating run) of P(My) My/S-0, PS My/S-100 and P(My-stat-S) statistical copolymers. The red crosses indicate the changes in slope observed.

( 1 ) 50 Figure S1b . 50 Figure S2b . 50 FigureFigure

 150S1b50S2b50 Figure S1b. Semi-logarithmic kinetic plots of ln((1-X) -1 ) (X = XMyfMy,0 + XYYfYY,0) versus polymerization time t for the various My/YY copolymerizations initiated by NHS-BB (YY = MA, tBuA or MMA).

G

  = (fGMA,0 / fMy,0) × [(2FGMA -1) / FGMA] ; H = (fGMA,0 / fMy,0) 2 × [(1 -FGMA) / FGMA) (B1) G = rGMA × H -rMy (B2)Equation B3 shows the additional variables used for the KT plot as illustrated in EquationB4.𝜂 = G / (α + H) ; ε = H / (α + H) ; α = √𝐻𝑚𝑎𝑥𝐻𝑚𝑖n (B3) 𝜂 = (rGMA + rMy / α)ε -rMy / α (B4)

  120 °C in bulk with NHS-BB and targeting Mn,theo = 30 kg.mol -1 at X = 100 %. a) fGMA,0 is the initial molar feed composition of glycidyl methacrylate. b) FGMA is the molar composition of glycidyl methacrylate in the P(My-stat-GMA) copolymer determined by1 H NMR in CDCl3. c) Overall monomer conversion X = XMyfMy,0 + XGMAfGMA,0 determined by 1 H NMR in CDCl3. d) Mn and Mw determined by GPC calibrated with PS standards in THF at 40 °C. P(My-stat-GMA) My/GMA-50: 1 H NMR (CDCl3, 300 MHz, RT): δ = 5.20-5.00 (s br, 2H P(My) ), 4.45-4.20 (s br, 1H P(GMA) ), 3.90-3.65 (s br, 1H P(GMA) ), 3.30-3.10 (s, 1H P(GMA) ), 2.90-2.80 (s, 1H P(GMA) ), 2.70-2.55 (s, 1H P(GMA) ), 2.20-1.90 (m br, 8H P(My) ), 2.00-1.80 (m, 2H P(GMA) ), 1.67 (s, 3H P(My) ), 1.59 (s, 3H P(My) ), 1.30-0.70 (m, 3H P(GMA) ). *Signal at δ = 5.45-5.30 ppm (t br, 1H 1,2-P(My) ) corresponds to one olefinic proton of 1,2-P(My) 5 . **Signal at δ = 4.80-4.65 ppm (d, 2H 1,2-P(My) and s, 2H 3,4-P(My) ) corresponds to two vinyl protons of 3,4-addition and two vinyl protons of 1,2-addition of My units 5 . The molar composition of GMA in the final P(My-stat-GMA) copolymer (FGMA) was calculated by comparing the two integrated areas corresponding to δ = 5.45-4.65 ppm (two protons of P(My)) and δ = 4.45-2.55 ppm (five protons of P(GMA)). FGMA = 0.58 was calculated for final My/GMA-50 sample.

Figure S6b. 1 H

 1 Figure S6b.1 H NMR spectrum in CDCl3 (300 MHz) at room temperature of dried P(My-stat-GMA) statistical copolymer My/GMA-50 after recovery, using methanol as the precipitant.

Figure S7b. 81

 81 Figure S7b. 81 MHz 31 P NMR spectra of dried P(My-stat-GMA)-SG1 macroinitiator My/GMA-78 in CDCl3 at room temperature.

Figure S8b .

 S8b Figure S8b. 81 MHz 31 P NMR spectra of dried P(My-stat-GMA)-SG1 macroinitiator My/GMA-37 in CDCl3 at room temperature.

Figure S10b .Figure S11b .

 S10bS11b Figure S10b. DSC traces (second heating run) of P(My-stat-GMA) statistical copolymers. The red crosses indicate the changes in slope observed.

Figure S12b .

 S12b Figure S12b. DLS measurements at (a) 25 o C, (b) 45 o C and (c) 65 o C of 0.4 mg.ml-1 of P(My-b-HMPMA) diblock copolymer in reversed osmosis water, stirred for 10 days before analysis. The correlation functions are also given to the right of the particle size distribution plots.

c. Appendix for Chapter 3 c. 1 .

 31 Experimental section• Materials. β-Myrcene (My, ≥ 90%), basic alumina (Al2O3, Brockmann, Type I, 150 mesh), calcium hydride (CaH2, 90-95% reagent grade), thiophenol (97%) and benzene (≥ 99%, ACS reagent) were purchased from Sigma-Aldrich and used as received. Toluene (≥ 99%), methanol (MeOH, ≥ 99%), methylene chloride (CH2Cl2, 99.9% certified ACS), chloroform (CHCl3, 99.8%) and tetrahydrofuran (THF, 99.9% HPLC grade) were obtained from Fisher Scientific and used as received.2-Methyl-2-[N-tert-butyl-N-(1-diethoxyphosphoryl-2,2-dimethylpropyl)-aminoxy]-N-propionyloxysuccinimide,also known as NHS-BlocBuilder (NHS-BB), was prepared according to a published method 1 from 2-(tert-butyl[1-(diethoxyphosphoryl)-2,2-dimethylpropyl]aminooxy)-2-methylpropionic acid, also known as MAMA SG1 (BlocBuilder-MA TM , (BB), 99%, provided by Arkema and used without further purification), N-hydroxysuccinimide (NHS, 98%, purchased from Aldrich and used as received), and N,N′dicyclohexylcarbodiimide (DCC, 99%, purchased from Aldrich and used as received). Telechelic poly(ethylene-statbutylene) difunctional initiator PEB-(SG1)2 terminated with N-tert-Butyl-N-[1-diethylphosphono-(2,2dimethylpropyl)] nitroxide (SG1) was produced relying on a published synthesis route 2 , from hydroxyl terminated poly(ethylene-stat-butylene) (PEB-(OH)2, Kraton TM D2205) obtained from Kraton, acryloyl chloride (98%) acquired

1

  and Đ = 1.53 as determined by GPC. The chain-extension of P(My)-(SG1)2 macroinitiator with a IBOMA/S mixture was then done by reacting P(My)-(SG1)2 (1.50 g, 0.029 mmol), purified IBOMA (3.32 g, 14.94 mmol) and purified S (0.15 g, 1.42 mmol) in toluene (~ 4.3 g) under continuous nitrogen flow and magnetic stirring for 3 h at 115 o C (experiment My-52-IBOMA/S, Table

where fA, 0 4

 04 and fB,0 are the initial molar fractions of A and B, and XA and XB are the individual conversions of A and B, respectively. My conversion was calculated by comparing the integrated peaks of the aliphatic protons of the monomer (δ = 2.30-2.15 ppm, 4H), the aliphatic protons of the polymer (δ = 2.15-1.90 ppm, 8H) and the protons of the two methyl groups of both monomer and polymer (δ = 1.75-1.55 ppm, 6H). IBOMA conversion was obtained using the NMR peaks of the vinyl protons of the monomer (δ = 6.05 and 5.50 ppm, 2H) and the aromatic proton of both monomer and polymer (δ = 4.65-4.20 ppm, 1H). Styrene (S) conversion was determined using the vinyl protons of the monomer (δ = 6.80-6.70, 5.80-5.70 and 5.30-5.20 ppm, 3H) and the aromatic protons of both monomer and polymer (δ = 7.50-6.90 ppm, 5H). The same 1 H NMR spectra (300 MHz Varian NMR Mercury spectrometer, CDCl3 eluent, 32 scans) allowed to determine the regioselectivity of the My repetitive units in the polymers by comparing the three integrated peaks at δ = 4.70-4.80 ppm (two vinyl protons of 3,4-addition and two vinyl protons corresponding to 1,2-addition), δ = 5.00-5.25 ppm (two olefinic protons corresponding to 1,4-addition, one olefinic proton corresponding to 1,2-addition and one olefinic proton from 3,4-addition) and δ = 5.30-5.50 ppm (one olefinic proton from 1,2-addition). 3,Phosphorus nuclear magnetic resonance ( 31 P NMR) was performed using a 5 mm diameter Up NMR tube with 840 scans being processed in a 200 MHz Varian Gemini 2000 spectrometer operating at 81 MHz. 0.0352 g of My-rich My/IBOMA-82 (Mn,MHS = 13.6 kg.mol -1 ) and 0.0676 g of IBOMA-rich My/IBOMA-44 (Mn,MHS = 8.8 kg.mol - 1

  Mark-Houwink-Sakurada (MHS) coefficients (MHS parameters determined at 35 °C with THF eluent for P(MMA)5 : KP(MMA) = 12.2 x 10 -5 dL.g -1 and αP(MMA) = 0.690 ; MHS parameters determined at 25 °C with THF eluent for P(IBOMA)6 : KP(IBOMA) = 3.8 x 10 -5 dL.g -1 and αP(IBOMA) = 0.748 ; MHS parameters determined at 30 °C with THF eluent for P(My) containing 90 and 10 mol% of 1,4-and 3.4-content respectively 7 : KP(My) = 7.5 x 10 -5 dL.g -1 and αP(My) = 0.772). Therefore, the converted Mn values were calculated according to the Mark-Houwink relationship 8 and the molar composition of the copolymer samples (determined by 1 H NMR) as given in equation 2 in the case of My/IBOMA copolymers: Mn,MHS = FMy [(KP(MMA)/KP(My)) Mn,GPC αP(MMA)+1 ] 1/(αP(My)+1) + (1-FMy) [(KP(MMA)/KP(IBOMA))

G

  = (fIBOMA,0 / fMy,0) × [(2FIBOMA -1) / FIBOMA] ; H = (fIBOMA,0 / fMy,0) 2 × [(1 -FIBOMA) / FIBOMA) (B1) G = rIBOMA × H -rMy (B2)Equation B3 shows the additional variables used for the KT plot as illustrated in EquationB4.𝜂 = G / (α + H) ; ε = H / (α + H) ; α = √𝐻𝑚𝑎𝑥𝐻𝑚𝑖n (B3) 𝜂 = (rIBOMA + rMy / α)ε -rMy / α (B4)In Equations B3 and B4, Hmax and Hmin stand for the highest and lowest values of the set of experimental points, respectively. rMy = 2.16 ± 0.34 and rIBOMA = 0.07 ± 0.04 (rMy × rIBOMA = 0.15 ± 0.19) were obtained via the FR method and rMy = 1.90 ± 0.18 and rIBOMA = 0.02 ± 0.21 (rMy × rIBOMA = 0.04 ± 0.15) were calculated via the KT method. The errors associated with the experimental data were derived from the standard errors of the slopes from FR and KT plots.

Figure S2c .

 S2c Figure S2c. Mn,GPC determined by GPC relative to PMMA standards in THF at 40 o C and Mn,MHS, obtained from Mn,GPC and corrected using the Mark-Houwink relationship, versus overall conversion X for experiments My/IBOMA-80 (fMy,0 = 0.81) and My/IBOMA-30 (fMy,0 = 0.30) at 100 o C. The dotted line indicates the theoretical Mn (Mn,theo) versus overall conversion based on the monomers to initiator ratio (Mn,theo = 30 kg.mol -1 at X = 100%).

Figure S4c. 1 HFigure S6c .

 1S6c Figure S4c.1 H NMR spectrum in CDCl3 (300 MHz, tetramethylsilane TMS as internal reference) at room temperature of the final dried P(My-grad-IBOMA) gradient copolymer from experiment My/IBOMA-50 after recovery, using methanol as the precipitant.

FigureFigure S9c .Figure

 S9c Figure S7c. (a) Semi-logarithmic kinetic plots of ln((1-X) -1 ) (X = My conversion) versus polymerization time t, (b) Đ versus X, and (c) Mn,MHS versus X for the experiments My-79 (open circles), My-105 (solid grey circles), My-137 (solid blue circles) and My-174 (solid red circles). The dotted lines indicate the theoretical Mn versus overall conversion based on the monomer-to-initiator ratio for each reaction (black, grey, blue and red predicted lines for experiments My-79, My-105, My-137 and My-174 respectively). All experimental ID are listed in TableS2cabove.

  The reactor was then heated to 115 o C and allowed to react for 140 min while maintaining a nitrogen purge. The resulting polymer sample (Mn = 94.4 kg.mol -1 , Đ = 2.19) was precipitated in excess methanol and dried under vacuum at 50 o C overnight. Lastly, fractionation was performed using a methanol (non-solvent) / benzene (solvent) pair (2 cycles) to remove the low molecular weight species. The final fractionated triblock polymer exhibited Mn = 103.1 kg.mol -1 , Đ = 2.11 and FS = 0.49.326 e. Appendix for Chapter 5: Experimental section• Materials. 2,6-Di-tert-butyl-4-methylphenol (BHT, ≥ 99%), p-toluenesulfonyl hydrazide (TSH, 97%), tributylamine (TBA, ≥ 98.5%), polystyrene-block-polyisoprene-block-polystyrene with 14 wt% styrene (abbreviated S4-I-S4, specification sheet below, product number = 432393) and 22 wt% styrene (abbreviated S8-I-S8, specification sheet below, product number = 432415) were purchased from Sigma-Aldrich and used as received. Methanol (MeOH, ≥ 99%), methylene chloride (CH2Cl2, 99.9% certified ACS), p-xylene (99.9%), chloroform (CHCl3, 99.9%) and tetrahydrofuran (THF, 99.9% HPLC grade) were obtained from Fisher Scientific and used as received. Styrene-blockethylene/butylene-block-styrene A1536 H (SEBS, specification sheet below) was obtained from Kraton. The deuterated chloroform (CDCl3, 99.8%) used as a solvent for nuclear magnetic resonance (NMR) was obtained from Cambridge Isotopes Laboratory.

  material testing system with a 5 kN load cell at room temperature and a cross-head speed of 10 mm.min -1 . Dog-bone style tensile specimens (ASTM D638 type V for reference, overall length = 63.5 mm, overall width = 9.53 mm) were prepared either by solvent casting or by hot press. To cast the film, the polymer sample was fully dissolved in dichloromethane (~ 80 wt%) for an hour with continuous stirring using a magnetic stir bar. The solution was cast into a level Teflon Petri dish at room temperature for approximately two days. Then, the drying was completed after placing the film in the oven at 50 o C under vacuum until constant weight was measured. The film samples were cut into typical dog-bone shapes using a sharp blade. When using a hot press (Carver hydraulic press, model #3925), T ~ Tg,specimen +

a

  temperature range of 20 to 180 o C was measured at a rate of 5 o C.min -1 under N2 using the CDT 450 convection heated measuring chamber mounted on the Anton Paar Modular Compact Rheometer MCR302. The solid rectangular fixture (SFR) was used, consisting of an upper and a lower holder with insets. Rectangular bars were made either by solvent casting or by hot press, following the above protocols used for preparing the tensile bars. Prior to the torsion tests, dynamic amplitude sweeps at 0.15 Hz and varying strain from 0.01 to 100% were conducted at various temperatures to ensure the material was kept within the linear viscoelastic regime. Please see the characterization section of Chapter 3 for the detailed protocols of DSC, TGA and AFM analyses (identical procedures).

  'UNIVERSITE DE LYON OPEREE AU SEIN DE L'INSA LYON NOM : MÉTAFIOT DATE de SOUTENANCE : 19/10/2018 Prénoms : Adrien TITRE : Synthesis of 1,3-Diene-based Block Copolymers by Nitroxide-Mediated Polymerization for Application as Robust Joining between Composite Matrices NATURE : Doctorat Numéro d'ordre : 2018LYSEI076 Ecole doctorale : MATÉRIAUX de Lyon (n o 34) Spécialité : Synthèse et Caracterisation des Polymères RÉSUMÉ : L'objectif de ce projet fut consacré à la synthèse d'un élastomère thermoplastique (TPE) innovant, utilisable en tant que joint entre des matrices composites. La polymérisation radicalaire contrôlée par des radicaux nitroxyde (NMP) fut choisie afin de produire des TPE copolymères à blocs de type styrénique. Le β-myrcène (My) fut dans un premier temps sélectionné pour synthétiser le bloc souple poly(β-myrcène) P(My). La NMP du My en masse à 120 o C en utilisant l'amorceur BlocBuilder TM fonctionnalisé avec l'ester succinimidyl (NHS-BB) fut très bien contrôlée, permettant l'extension de chaîne du macro-amorceur cis-1,4-P(My) avec du styrène (S). Les copolymères diblocs P(My-b-S) obtenus ayant une masse molaire moyenne en nombre modérée (Mn < 50 kg.mol -1 ) montrèrent un comportement fragile en test de traction uniaxiale (résistance à la rupture en traction σB < 1,1 MPa, allongement à la rupture en traction εB < 16%). L'incorporation d'unités fonctionnelles au sein du segment souple fut réalisée en parallèle via la NMP du My avec du méthacrylate de glycidyle (GMA) afin de favoriser possiblement le processus de soudage / collage entre le TPE et les composites thermoplastiques envisagés. Un amorceur bifonctionnel, le poly(éthylène-co-butylène)-(SG1)2 (SG1 = groupe nitroxyde), permit par la suite de synthétiser des copolymères triblocs S-My-S ayant une plus haute masse molaire moyenne Mn = 56-66 kg.mol -1 et une plus grande extensibilité (σB < 0,8 MPa, εB < 200%). Les segments PS furent remplacés par des segments ayant une température de transition vitreuse (Tg) plus élevée, à savoir des blocs poly(méthacrylate d'isobornyle) P(IBOMA) afin d'augmenter la résistance mécanique et la température de service du TPE candidat. Des triblocs de type IBOMA-My-IBOMA, dont les domaines furent microstructurés, montrèrent de meilleures propriétés mécaniques (σB = 3,9 MPa, contrainte à la limite d'élasticité σY = 5,0 MPa, εB = 490%) et une température de service maximale d'environ 140 o C. Toutefois, ces TPE à base de My ne satisfirent pas le cahier des charges industriel à température ambiante, ce qui nous poussa à substituer le bloc flexible P(My) par du poly(isoprène) PI, ayant une masse molaire d'enchevêtrement bien plus faible. Des macro-amorceurs 1,4-PI bien définis et actifs furent d'abord étendus avec du styrène, ce qui permit d'obtenir des triblocs S-I-S auto-assemblés (Mn = 95-109 kg.mol -1 , fraction molaire en styrène FS = 0,30-0,49, dispersité Đ = 2,11-2,29). σB = 4,1 ± 0.2 MPa et εB = 376 ± 64 % furent mesurés pour un S-I-S ayant FS = 0,38. Un tribloc de type IBOMA-I-IBOMA fut par la suite synthétisé (Mn = 94 kg.mol -1 , Đ =

Table 1 .

 1 Main

	TPE	Mechanical properties of the assembly	Service conditions	Health and safety
	Similar to the stress-strain properties of the thermoset rubber used	Satisfactory tensile strength at break and high shear modulus at 100% elongation	Continuous service temperature = 10 to 40 o C Maximum service temperature = 80 o C	REACH compliance
	currently	Minimal loss of the mechanical properties with temperature	Life time > 20 years
		Type of failure = Cohesive fracture	Hygrometry: Relative humidity = 35 to 65 % (short periods at 100 %)
			No mechanical damage
			with diestone and 1,1,1-
			trichloroethane

industrial requirements for the candidate TPE and the resulting assembly. REACH is the European Regulation on Registration, Evaluation, Authorization and Restriction of Chemicals, came into effect in 2007.

  1,6-diene) and isoprene (I, 2-methylbuta-1,3-diene). My was first used as an isoprene substitute building block. My exhibits two valuable features in comparison to isoprene: it is a natural renewable compound with low volatility (boiling point, Tb = 167 °C) meaning that pressurized equipment is not needed to handle it. After receiving high-pressure reactors, isoprene (Tb = 34

o C), more common for the synthesis of block copolymer TPE, was handled. Well-controlled isoprene NMP was previously studied by Hawker for instance or more recently by Harrisson, which constitutes a major advantage for our research. Butadiene was disregarded for this study due to its subzero boiling point (Tb =

  This industrial thesis (thèse CIFRE-défense), done from January 2015 to December 2017, was predominantly financed by ArianeGroup with a financial assistance from the French government via the French Defense Procurement Agency (Direction Générale de l'Armement, DGA). Dr. Brigitte Defoort, from the composite materials laboratory at ArianeGroup, coordinated the project. M. Bruno Mortaigne was

my DGA responsible. Dr. Jean-François Gérard, professor and former research director at INSA Lyon and currently chairman of the European Polymer Federation (EPF), was my thesis director. Prof. M. Maric and Prof. P. Hubert, respectively from the Departments of Chemical and Mechanical Engineering at McGill University, were my thesis co-supervisors. Even though solely registered at "Institut National des Sciences Appliquées (INSA) de Lyon", as a PhD student, I spent time at McGill University as a visiting PhD candidate under the supervision of Prof. M. Maric, to synthesize and characterize polymers. The laboratories of Prof. P. Hubert were also used to implement mechanical tests. I also worked at the laboratory ``Ingénierie des Matériaux Polymères (IMP)`` from INSA Lyon under the supervision of Prof. Gérard to review the literature and characterize the polymer materials. It must be added that this research was explored jointly with another ArianeGroup project led by PhD candidate David Brassard from Polytechnique Montreal and co-supervised by Prof. J. R. Tavares from Polytechnique Montreal and Prof. M. Dubé from ETS (École de Technologie Supérieure) Montreal. This second project falls totally within our industrial issue since David Brassard aims to develop a welded flexible junction between two filament wound thermoplastic composite parts, while focusing mostly on the welding process. This junction is composed of a thermoplastic elastomer welded to the composite components. David Brassard uses an extra electrically conductive nanocomposite heating element for the resistance welding process implemented.

Table 2 . Some physical, mechanical and thermal characteristics of NBR exhibiting an acrylonitrile content ranging from 18 to 50 mol% 10-12 (

 2 resistance due to its stable, saturated polymer backbone structure. EPDM is mainly used as a standard lining material for steam hoses, automotive weather-stripping and seals, radiator, electrical insulation and roofing membrane 5 . The structures of NR, SBR and EPDM are shown in Figure 9. industrial sources: Robinson TM Rubber Products Company; K.C. Seals Inc.). The structure of NBR is given to the right of the table.

	50

contains between 34 and 40 mol% of acrylonitrile repetitive units, as an industrial standard. This gives the most versatile compound performance for oil swell and low temperature properties. The major advantages Figure 9. Repetitive units of natural rubber (NR) homopolymer and styrene-butadiene (SBR) and ethylene-propylene-diene monomer (EPDM) copolymers. with excellent creep/set resistance are generally exhibited by these vulcanized materials. However, some non-negligible shortcomings appear when handling thermoset polymers. The permanent state brought about by crosslinking and the need to perform a crosslinking reaction on elastomers are disadvantages for their

Table 3 .

 3 Entanglement molar masses of some poly(diene)s 28 .

	(a) Random copolymer with 7 ethyl

Table 4 .

 4 Hansen' solubility parameters data for PS, PB, PI and PE 30,31 . δD = Dispersion forces

	contribution, δP = Polar forces arising from dipole moments contribution, δH = Hydrogen-bonds
	between molecules contribution, δ = (δD 2 + δP 2 + δH 2 ) 0.5 .
	Polymer	𝜹D	𝜹P	𝜹H
	Poly(styrene) PS	21.3	5.8	4.3
	Poly(butadiene) PB	17.5	2.3	3.4
	poly(isoprene) PI	16.6	1.4	-0.8

Poly(ethylene-co-butylene) P(EB) Table

?

. Entanglement molar masses of some poly(diene)s 36 . (a) Random copolymer with 7 ethyl

Table 4

 4

gathers Hansen solubility parameters for PS, PI, PB, and P(EB)

30,31

	Polymer	𝜹D	𝜹P	𝜹H	𝜹
	Poly(styrene) PS	21.3	5.8	4.3	22.5
	Poly(butadiene) PB	17.5	2.3	3.4	18.0
	poly(isoprene) PI	16.6	1.4	-0.8	16.7
	Poly(ethylene) PE	15.1	-4.2	2.0	15.8
					64

. The lower solubility parameters for isoprene versus butadiene means that better phase separation between the soft and the hard

are known in which high levels of saturation can be achieved without chain degradation. Palladium (Pd) on

	These reactions are commonly
	run in organic solvents at relatively high dilution (10-30 % polymer) in a stirred pressure reactor at moderate
	temperatures. A class of complexes known collectively as Ziegler-type catalysts, cobalt (Co) and nickel
	(Ni) complexes (particularly the hydrocarbon soluble 2-ethylhexanoate and acetylacetonate complexes)
	activated with organo-lithium (Li) or organo-aluminum (Al) compounds, are probably the most extensively
	studied and commercially utilized regarding homogeneous catalytic hydrogenation. These catalysts are
	effective in the saturation of poly(butadiene) PB and poly(isoprene) PI homopolymers (including natural
	rubber) and in block or random copolymers with styrene (S), and can be optimized for a given polymer
	with respect to the ratio of transition metal to the Li-or Al-activating species 46,47 . Figure 18 depicts the
	saturation of butadiene units in a styrenic block copolymer via Ziegler-based hydrogenation. A noticeable
	drawback with the commercial use of the Zeigler-type catalysts is that the residual metal has to be removed
	before use. Heterogeneous catalytic hydrogenation can be implemented as well. Transition metal catalysts
	using a variety of inorganic supports have found utility in the hydrogenation of the polymers containing
	isolated alkenes such as those encountered in poly(diene)s and metathesis-derived polymers. Conditions

Table 6 .

 6 Maximum degradation temperature (Tdeg,max

	Natural rubber	0 85.8	381.7 465.9
	Synthetic cis-1,4-poly(isoprene)	0 89.7	386.0 464.0
	Cis-1,4-poly(butadiene)	0 > 99	440 470
	Trans-1,4-poly(2,4-hexadiene)	0 > 98	345 435

Table ,

 ,

this is highlighted by their relatively low modulus at 300 % elongation (≤ 3.1 MPa). Moreover, their melt flow index at 200 o C ranging from 8.0-19 g/10min indicate that these materials significantly soften and flow at such a temperature. It can thus be assumed that these linear SIS should not be thermally robust enough to support welding process or adhesive bonding process at T > 160 o C.

  Table above exhibiting various weight fractions of styrene unit (29-57 wt% S). SEBS Kraton ® have a relatively limited extensibility (elongation at break = 500-660 %) and a comparable tensile stress at break (28-35 MPa) compared to SIS and SBS Kraton ® . SEBS are tougher due to the hydrogenated midsegment.

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(87) Wegrzyn, J. K.; Stephan, T.; Lau, R.; Grubbs, R. B. Preparation of Poly(ethylene oxide)-block-poly(isoprene) by Nitroxide Mediated Free Radical Polymerization from PEO Macroinitiators. J. Polym. Sci., Part A: Polym. Chem. 2005, 43, 2977-2984. (88) Sundararaman, A.; Stephan, T.; Grubbs, R. B. Reversible Restructuring of Aqueous Block Copolymer Assemblies through Stimulus Induced Changes in Amphiphilicity. J. Am. Chem. Soc. 2008, 130, 12264-12265. (89) Cai, Y.; Aubrecht, K. B.; Grubbs, R. B. Thermally Induced Changes in Amphiphilicity Drive Reversible Restructuring of Assemblies of ABC Triblock Copolymers with Statistical Polyether Blocks. J. Am. Chem. Soc. 2010, 133, 1058-1065. (90) Ajellal, N.; Thomas, C. M.; Carpentier, J. -F. Controlled Radical Polymerization of Conjugated 1,3-Dienes with Methyl 1,3-Butadiene-1phosphonate. Polymer 2008, 49, 4344-4349.w (91) Luo, K.; Rzayev, J. Living Radical Polymerization of Bicyclic Dienes: Synthesis of Thermally Cross-linkable Block Copolymers. Macromolecules 2009, 42, 9268-9274. (92) Matsuoka, H.; Suetomi, Y.; Kaewsaiha, P.; Matsumoto, K. Nanostructure of a Poly(acrylic acid) Brush and its Transition in the Amphiphilic Diblock Copolymer Monolayer on the Water Surface. Langmuir 2009, 25, 13752-13762.

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25,26,29

, atom transfer radical polymerization (ATRP)

24,30-32 and

reversible addition fragmentation transfer polymerization (RAFT)

23,33-35

.

RDRP applied to My is a great opportunity to produce high value-added polymers with a segment based on renewable monomers. Surprisingly enough, RDRP-based P(My) has not been synthesized and characterized prior to 2015. Very recently, the first RDRP of My was successfully performed by RAFT

36

.

Table 1 .

 1 Experimental conditions for My polymerizations for 7 h initiated by BB or NHS-BB.

	104

Table 2 .

 2 Molecular characterization, kinetic data and selectivity of P(My) synthesized with BB or NHS-BB.

Table

Table

1 ) t (min)

	My-10-S	13.2	1.18 (c)	0.006	3.442	4.812	73	300
	My-11-Sa (b)	22.8	1.32	0.005	3.242	4.829	90	420
	My-11-Sb (b)	22.8	1.32	0.005	3.143	4.802	88	140
	My-12-S	30.4	1.38	0.004	3.026	4.809	109	120

Table 3 .

 3 Formulations

of chain-extensions with purified S using various P(My) macro-initiators in 50 wt% toluene solution at 110 °C.

Table 4 .

 4 Molecular and mechanical characterization of diblock copolymers after chain-extension with purified S from various P(My) macro-initiators at 110 °C in 50 wt% toluene.

	Tensile elongation
	at break (d) (%)

a) S conversion (XS) determined only by 1 H NMR spectroscopy in CDCl3. b) Molar fraction of S in the copolymer (FS) as determined by 1 H NMR in CDCl3. c) Mn and Mw determined by GPC calibrated with PS standards in tetrahydrofuran (THF) at 40 °C. Mn of the P(My) segment corrected for each experiment using the appropriate MHS parameters. d) Young's modulus (E), ultimate tensile strength and tensile elongation at break measured by uniaxial tensile testing via a MTS Insight TM material testing system with a 5 kN load cell at room temperature and a cross-head speed of 1 mm.min -1 . 6, 5 and 5 specimens tested for the experiments My-11-Sa, My-11-Sb and My-12-S respectively. Averaged results given with the respective standard deviations. e) Specimens too brittle to perform tensile testing.

  Mn and Mw determined by GPC calibrated with PS standards in THF at 40 °C. Appropriate MHS coefficients and FMy = 1 -FS used to convert molecular weights based on PS to the fraction of My incorporated in the P(My-stat-S) copolymer. d) See notation c) in Table 2 (R 2 ≥ 0.94).

	My/S-0	0	0.40	0	0.40	8.2	1.19	1.6 ± 0.3
	My/S-10	0.08	0.47	0.39	0.46	12.6	1.25	2.2 ± 0.1
	My/S-20	0.15	0.44	0.33	0.42	10.0	1.26	1.5 ± 0.3
	My/S-30	0.28	0.41	0.37	0.40	8.4	1.24	1.2 ± 0.3
	My/S-40	0.38	0.50	0.45	0.48	13.7	1.25	1.8 ± 0.2
	My/S-45	0.34	0.40	0.28	0.35	10.0	1.33	1.3 ± 0.2
	My/S-50	0.48	0.50	0.45	0.47	11.7	1.19	1.7 ± 0.4
	My/S-60	0.31	0.38	0.31	0.34	9.8	1.36	1.1 ± 0.6
	My/S-70	0.45	0.54	0.45	0.48	13.9	1.24	1.2 ± 0.3
	My/S-80	0.66	0.89	0.95	0.94	13.8	1.33	2.5 ± 0.6
	My/S-90	0.83	0.84	0.78	0.79	19.8	1.31	2.0 ± 0.7
	My/S-95	0.96	0.81	0.89	0.89	11.0	1.37	4.1 ± 1.1
	My/S-100	1	0	0.88	0.88	15.6	1.20	7.5 ± 0.8

a) See notation b) in Table 4. b) Individual monomer conversions XMy and XS determined by 1 H NMR. Average monomer conversion XNMR = XMy fMy,0 + XS fS,0 (further details in the experimental section). c)

Table 6 .

 6 Molecular characterization at the end of the experiments and kinetic data of P(My-stat-S) copolymers at 110 °C in bulk initiated by NHS-BB and targeting Mn,theo = 30 kg.mol -1 at X = 1.0.

Table 7 .

 7 Chain-extensions of A) My-rich and S-rich P(My-stat-S) macro-initiators with B) My and S monomers in 50 wt% toluene or in bulk and C) molecular characterization of the resulting chain-extended products.My/S-55 with My (experiments My/S-07-My.a and My/S-55-My.a in Table7.B) and S (experiments My/S-07-S and My/S-55-S in Table7.B) monomers at 110 °C in 50 wt% toluene for 3 h. Accordingly, the influence of the composition of P(My-stat-S) macroinitiators over their ability to polymerize a second batch of monomer was discussed as well as their own versatility to chain-extend My and S. The final characterization of the chain-extended products is shown in Table7.C. My-rich macroinitiator My/S-07 cleanly chain-extended My and S as demonstrated by the low dispersities obtained (Đ ≤ 1.43) and the low deviations of the experimental Mn,2 values of the chain-extended copolymers from the theoretical ones (Mn,2/Mn,theo,X2 ≥ 88%) at the corresponding conversion X2. Likewise, successful chain-extensions with My and S were performed from S-rich macroinitiator My/S-55 even though broader molecular weight distributions (Đ = 1.44 with My and 1.60 with S) were measured than those of the first set of chainextensions (Đ = 1.35 and 1.43) and slightly higher deviations of Mn,2 ≥ 85% of Mn,theo,X2. This minor loss of control can be directly related to the intrinsic features of the S-rich macroinitiator My/S-55 having likely a greater proportion of irreversibly terminated chains relative to the My-rich macroinitiator My/S-07.

Table 8

 8

shows measured Tgs (DSC plots with changes in the slope provided in the Supporting Information, Figures

S15a and S16a

  Wilbon, P. A.; Chu, F.; Tang, C. Progress in Renewable Polymers from Natural Terpenes, Terpenoids, and Rosin. Macromol. Rapid Commun. Behr, A.; Johnen, L. Myrcene as a Natural Base Chemical in Sustainable Chemistry: A Critical Review. ChemSusChem 2009, 2 (12), 1072-1095. Savich, T. R.; Goldblatt, L. A. Process for Producing Myrcene from Beta-Pinene. U.S. Patent US2507546, 1950. Quirk, R. P.; Huang, T.-L. Alkyllithium-Initiated Polymerization of Myrcene: New Block Copolymers of Styrene and Myrcene. In New

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(105) Morton, M.; Salatiello, P. P.; Landfield, H. Absolute Propagation Rates in Emulsion Polymerization. III. Styrene and Isoprene. J. Polym. Sci., Part A: Polym. Chem. 1952, 8 (3), 279-287. (106) Greszta, D.; Matyjaszewski, K. Mechanism of Controlled/"Living" Radical Polymerization of Styrene in the Presence of Nitroxyl Radicals. Kinetics and Simulations. Macromolecules 1996, 29 (24), 7661-7670. (107) Kruse, T. M.; Souleimonova, R.; Cho, A.; Gray, M. K.; Torkelson, J. M.; Broadbelt, L. J. Limitations in the Synthesis of High Molecular Weight Polymers via Nitroxide-Mediated Controlled Radical Polymerization: Modeling Studies. Macromolecules 2003, 36 (20), 7812-7823. (108) Zetterlund, P. B.; Saka, Y.; McHale, R.; Nakamura, T.; Aldabbagh, F.; Okubo, M. Nitroxide-Mediated Radical Polymerization of Styrene: Experimental Evidence of Chain Transfer to Monomer. Polymer 2006, 47 (23), 7900-7908. (109) Nicolas, J.; Dire, C.; Mueller, L.; Belleney, J.; Charleux, B.; Marque, S. R. A.; Bertin, D.; Magnet, S.; Couvreur, L. Living Character of Polymer Chains Prepared via Nitroxide-Mediated Controlled Free-Radical Polymerization of Methyl Methacrylate in the Presence of a Small Amount of Styrene at Low Temperature. Macromolecules 2006, 39 (24), 8274-8282. (110) Petton, L.; Ciolino, A. E.; Dervaux, B.; Du Prez, F. E. From One-pot Stabilization to In Situ Functionalization in Nitroxide Mediated Polymerization: An Efficient Extension towards Atom Transfer Radical Polymerization. Polym. Chem. 2012, 3, 1867-1878. (111) Nishizaki, H.; Yoshida, K. Effect of Molecular Weight on Various TGA Methods in Polystyrene Degradation. J. Appl. Polym. Sci. 1981, 26, 3503-3504. (112) Loadman, M. J. R. The Glass Transition Temperature of Natural Rubber. J. Therm. Anal. Calorim. 1985, 30 (4), 929-941. (113) Rieger, J. The Glass Transition Temperature of Polystyrene: Results of a Round Robin Test. J. Therm. Anal. Calorim. 1996, 46 (3), 965-972. (114) Georgopanos, P.; Rangou, S.; Haenelt, T. G.; Abetz, C.; Meyer, A.; Filiz, V.; Handge, U. A.; Abetz, V. Analysis of Glass Transition and Relaxation Processes of Low Molecular Weight Polystyrene-b-Polyisoprene Diblock Copolymers. Colloid Polym. Sci. 2014, 292 (8), 1877-1891. (115) Gedde, U. W. The Glassy Amorphous State. Polymer Physics; Kluwer Academic Publishers: Dordrecht, 1999; 77-98. (116) Tonelli, A. E. Possible Molecular Origin of Sequence Distribution-Glass Transition Effects in Copolymers. Macromolecules 1974, 7 (5), 632-634. (117)

Table of

 of

contents graphic. Synthesis of well-defined My/GMA statistical and diblock copolymers by SG1 NMP. GMA units can be treated with morpholine to obtain 2-hydroxy-3-morpholinopropyl methacrylate hydrophilic groups.

28,31-33

  47,[50][51][52] To avoid this, tert-butyl acrylate (tBuA), bearing a tert-butyl protecting group, was used. The subsequent cleavage after polymerization of the protecting group with trifluoroacetic acid can be performed to yield carboxylic acid groups.

53

Lastly, methyl methacrylate (MMA) was used to see if My with a methacrylate can be done in a controlled manner as has been suggested by isoprene/MMA copolymerizations at quite high MMA content in the initial monomer mixture.

54

  to the occurrence of backbiting reactions, converting chain-end radicals to less reactive mid-chain radicals. With this assumption, the synthesis of branched P(My-co-tBuA) under these reaction conditions is probable. Interestingly, NMP of tBuA initiated by BlocBuilder TM with 4.5 mol% additional SG1 relative to initiator in bulk at 115 o C exhibited a linear increase of Mn with conversion.62 However, relatively high Đ values(1.51-2.27

[60][61]

Accordingly, uncontrol My/tBuA NMP might 148 be due

24,50,51

, leading to a high concentration of propagating radicals that can promote self-termination. Consequently, it was decided to perform the two same experiments at 100 o C (experiments My/IBOMA-50 with fMy,0 = 0.50 and My/IBOMA-70 with fMy,0 = 0.70, Table

1

) in order to know if a better control of the NMP can be

Table 1 .

 1 Formulations for My, IBOMA and My/IBOMA polymerizations performed in bulk or in toluene

	[IBOMA]0 (M)	fMy,0	T ( o C)	Solvent	t (min)

Table 2 .

 2 Molecular characterization and My selectivity at the end of the experiments and kinetic data of

	1,2-(e)
	(%)

  37,38,69 , the relatively low Tgs observed herein for these rigid blocks suggest that

P(IBOMA)s rich in isotactic triads (> 50 mol%) were produced. Indeed, Yu et al. notably showed that syndiotactic P(IBOMA) (> 50 mol%) exhibited a higher Tg = 195-206 o C 38 .

MHS,CE (kg.mol -1 ) Mn,theo,XCE (e) (kg.mol -1 ) ĐCE My-35-IBOMA/My

										t
									)	(min)
	My-35-IBOMA/My	0.004		1.681	0.142	0	4.929	133.4	115	90
	My-52-IBOMA/S	0.003		1.546	0	0.147	4.848	171.2	115	180
				C. Chain-extended triblock copolymer (a)
	ID		XCE (%) Mn,28.6 FMy FIBOMA FS 0.72 0.28 0 51.3	63.3		1.91
	My-52-IBOMA/S	59.0	0.59	0.36	0.05	94.7	122.2		2.23

Table 4 .

 4 Chain-extensions of

Table 1 .

 1 Formulations for isoprene polymerizations performed in bulk or in pyridine at T = 115

		0.0051	9.973	0 (c)	139	115	8
	PI2	0.0050	10.208	0 (c)	145	115	17
	PI3	0.0024	4.928	6.17	146	115	25

Table 2 .

 2 Molecular characterization of final NMP-based PI, initiated by PEB-(SG1)2 at 115 o C in 50 v% pyridine and targeting Mn,theo ~ 145 kg.mol -1 at XI = 100% (optimized experimental conditions).

	Before fractionation	After fractionation (a)

Table 3 .

 3 Chain-extensions of (A) PI-(SG1)2 macroinitiators with S at 115 o C in bulk and (B) features of the resulting S-I-S triblock copolymers.

		44.1	0.49	94.4	2.19	103.1	2.11
	PS-PI5-PS	31.2	0.30	89.2	2.32	94.6	2.29
	PS-PI6-PS	36.8	0.38	101.8	2.23	109.4	2.18

1 )

 1 Pradel, J. L.; Ameduri, B.; Boutevin, B. Use of Controlled Radical Polymerization of Butadiene with AIBN and Tempo for the Determination of the NMR Characteristics of Hydroxymethyl Groups. Macromol. Chem. Phy. 1999, 200, 2304-2308. Pradel, J. L.; Ameduri, B.; Boutevin, B. Controlled Radical Polymerization of 1,3-Butadiene. II. Initiation by Hydrogen Peroxide and Reversible Termination by Tempo. J. Polym. Sci., Part A: Polym. Chem. 2000, 38, 3293-3302. Cai, Y.; Aubrecht, K. B.; Grubbs, R. B. Thermally Induced Changes in Amphiphilicity Drive Reversible Restructuring of Assemblies of ABC Triblock Copolymers with Statistical Polyether Blocks. J. Am. Chem. Soc. 2010, 133, 1058-1065.

	(2) (3) Keoshkerian, B.; Georges, M.; Quinlan, M.; Veregin, R.; Goodbrand, B. Polyacrylates and Polydienes to High Conversion by a Stable Free
	Radical Polymerization Process: Use of Reducing Agents. Macromolecules 1998, 31, 7559-7561.
	(4) Benoit, D.; Harth, E.; Fox, P.; Waymouth, R. M.; Hawker, C. J. Accurate Structural Control and Block Formation in the Living Polymerization
	1,3-Dienes by Nitroxide-Mediated Procedures. Macromolecules 2000, 33, 363-370.
	(5) Cheng, C.; Qi, K.; Khoshdel, E.; Wooley, K. L. Tandem Synthesis of Core-Shell Brush Copolymers and their Transformation to Peripherally
	Cross-linked and Hollowed Nanostructures. J. Am. Chem. Soc. 2006, 128, 6808-6809.
	(6) Wegrzyn, J. K.; Stephan, T.; Lau, R.; Grubbs, R. B. Preparation of Poly(ethylene oxide)-block-poly(isoprene) by Nitroxide Mediated Free
	Radical Polymerization from PEO Macroinitiators. J. Polym. Sci., Part A: Polym. Chem. 2005, 43, 2977-2984.
	(7) Sundararaman, A.; Stephan, T.; Grubbs, R. B. Reversible Restructuring of Aqueous Block Copolymer Assemblies through Stimulus Induced
	Changes in Amphiphilicity. J. Am. Chem. Soc. 2008, 130, 12264-12265.
	(8) (9) Ajellal, N.; Thomas, C. M.; Carpentier, J. -F. Controlled Radical Polymerization of Conjugated 1,3-Dienes with Methyl 1,3-Butadiene-1-
	phosphonate. Polymer 2008, 49, 4344-4349.
	(10) Luo, K.; Rzayev, J. Living Radical Polymerization of Bicyclic Dienes: Synthesis of Thermally Cross-linkable Block Copolymers.
	Macromolecules 2009, 42, 9268-9274.
	(11) Harrisson, S.; Couvreur, P.; Nicolas, J. SG1 Nitroxide-Mediated Polymerization of Isoprene: Alkoxyamine Structure/Control Relationship and
	α,w-chain-end Functionalization. Macromolecules 2011, 44, 9230-9238.

  8,9,10 ) or poly(α-methylene-γ,γ-dimethyl-γ-butyrolactone) (P(Me2MBL), Tg ~ 210 o C, Tdec,onset ~ 340 o C under N2, Me2MBL controlled by RAFT 6 ) could be used the same way. Other groups of polymers, produceable via a radical process and exhibiting high Tg and Tdec, can be considered.

	Poly(styrene) derivatives, including for instance poly(α-methyl-p-methylstyrene) (Tg = 165-180 o C, Tdec,onset
	~ 320

o C under N2)

11

, are of interest as well. It should be added that the incompatibility of these candidate rigid segments with the soft middle block is assumed, allowing a strong phase separation of the two phases (investigation of the solubility parameters required).

Table S1a .

 S1a Young

's modulus (E), ultimate tensile strength and tensile elongation at break values for every specimen tested from experiments My-11-Sa, My-11-Sb and My-12-S.

Table S2a .

 S2a Samples used for determination of reactivity ratios for My/S statistical copolymerization done at 110 °C in bulk with NHS-BB and targeting Mn,theo = 30 kg.mol

Table S1b .

 S1b Samples used for determination of reactivity ratios for My/GMA statistical copolymerization done at

	fGMA,0 (a)	FGMA (b)	X (c) (%)	t (min)	Mn (d) (kg.mol -1 )	Đ (d)
	0.10	0.19	17.4	30	4.8	1.33
	0.20	0.26	16.6	30	3.7	1.34
	0.29	0.37	14.9	15	3.2	1.38
	0.40	0.44	16.1	15	4.0	1.44
	0.51	0.48	13.8	15	3.8	1.52
	0.60	0.56	23.2	15	6.1	1.43
	0.71	0.63	21.0	10	5.9	1.48
	0.80	0.73	19.8	5	7.0	1.49
	0.90	0.88	16.9	3	6.2	1.79

Table S1c .

 S1c Samples used for determination of reactivity ratios for My/IBOMA gradient copolymerization done at

	fIBOMA,0 (a)	FIBOMA (b)	X (c) (%)	t (min)	Mn,MHS (d) (kg.mol -1 )	Đ (d)
	0.10	0.15	27.5	60	4.1	1.27
	0.19	0.10	12.5	30	1.6	1.07
	0.30	0.13	13.2	60	1.8	1.10
	0.41	0.22	17.0	120	2.7	1.27
	0.50	0.24	20.1	100	3.7	1.28
	0.60	0.28	10.4	30	3.0	1.21
	0.70	0.42	13.2	60	3.5	1.30
	0.80	0.39	21.2	180	6.6	1.36
	0.90	0.57	12.4	60	7.0	1.52

Table 2 .

 2

	ID (a)	[PEB-(SG1)2]0 (M)	[My]0 (M)	[SG1]0 (M)	Solvent	[Solvent]0 (M)	Mn,theo (b) (kg.mol -1 )	t (min)
	My-79	0.011	5.928	0	-	0	79.1	300
	My-105	0.008	5.831	0	-	0	105.0	420
	My-137	0.006	5.791	0	-	0	137.2	420
	My-174	0.005	6.178	0	-	0	174.0	420
	My-173SG1,9	0.005	6.122	0.0005	-	0	172.5	420
	My-169SG1,18	0.005	5.975	0.0011	-	0	168.5	480
	My-170Tol	0.002	2.414	0	Toluene	4.573	170.1	420

Table S2c .

 S2c Experimental conditions for My polymerizations initiated by PEB-(SG1)2 at 120 o C in bulk or at 115

o C in toluene.

Cette thèse est accessible à l'adresse : http://theses.insa-lyon.fr/publication/2018LYSEI076/these.pdf © [A. Métafiot], [2018], INSA Lyon, tous droits réservés

(crosshead speed) 39 . Cette thèse est accessible à l'adresse : http://theses.insa-lyon.fr/publication/2018LYSEI076/these.pdf © [A. Métafiot], [2018], INSA Lyon, tous droits réservés

) and narrow molecular weight distributions (Đ < 1.3) 41 . Grubbs et al synthesized a well-defined Cette thèse est accessible à l'adresse : http://theses.insa-lyon.fr/publication/2018LYSEI076/these.pdf © [A. Métafiot], [2018], INSA Lyon, tous droits réservés

, Đ = 1.28) presumably due to My auto-initiation. Cette thèse est accessible à l'adresse : http://theses.insa-lyon.fr/publication/2018LYSEI076/these.pdf © [A. Métafiot], [2018], INSA Lyon, tous droits réservés

H NMR analysis (FMy and FIBOMA = 1 -FMy, Figure S4c in Supporting Information for the spectral Cette thèse est accessible à l'adresse : http://theses.insa-lyon.fr/publication/2018LYSEI076/these.pdf © [A. Métafiot], [2018], INSA Lyon, tous droits réservés

© [A. Métafiot], [2018], INSA Lyon, tous droits réservés

measured for GMA-rich initial feeds (fMy,0 = 0.10-0.49) during the copolymerizations, the linear increase of Mn with X and the monomodal nature of the GPC peaks indicated the satisfactory mediation of these reactions by SG1. The more apparent loss of control observed for My/GMA copolymerizations exhibiting richer GMA initial concentrations was likely due to the significantly higher propagation rate as indicated by ˂kp˃˂K˃ = (157.0 ± 0.9) 10 -5 s -1 at fGMA,0 = 0.90, about 36 times higher than that of My NMP under the same reaction conditions (˂kp˃˂K˃ = (4.3 ± 0.7) 10 -5 s -1 ). The faster kinetics presumably favored the occurrence of irreversible terminations which might be β-hydrogen transfer from propagating P(My-stat-GMA) radical to SG1 or intramolecular chain transfer to polymer. P(My-stat-GMA)s exhibiting FGMA = 0.11-0.95 displayed a range of Tgs from -77 o C to + 43 o C. My-rich and GMA-rich P(My-stat-GMA)s were cleanly chain-extended with My, S and GMA, indicating the high chain-end fidelity of these macroinitiators, which was also confirmed by the quantitative estimation of the SG1 end-group via 31 P NMR (LF = 82 ± 5 % and 71 ± 7 % respectively). Most of the resulting diblock copolymers exhibited two distinct Tgs, indicative of a possible microphase separation. Lastly, a NMP-based P(My-b-GMA) diblock copolymer was treated with morpholine, a nucleophilic reagent, to yield quantitatively P(My-b-HMPMA). This latter exhibited a higher degree of hydrophilicity, as shown by its total dissolution in water, compared to P(Myb-GMA). P(My-b-HMPMA) was able to self-organize into micelles in water, exhibiting Zave = 120-130 nm at T = 45-65 o C and [P(My-b-HMPMA)] ≤ 0.7 mg.ml -1 . Thus, well-tailored statistical and diblock epoxide functionalized P(My)s can be readily synthesized by SG1-mediated NMP, opening the door to a multitude of post-polymerization treatments resulting in the achievement of multifunctional My-based polymers. weight distribution triblock copolymer may be essential for the purposes of mimicking the mechanical properties of commercial styrenic TPE. Generally, the following properties could be expected for this welltailored MBL-I-MBL triblock copolymer:

-Stress-strain behavior at room temperature: σB > 20 MPa, εB > 600 % with a marked strain hardening (low dispersity encouraging the orientation and the alignment of the chains). Afterwards, the joining of the triblock copolymer TPE with the tank and the skirt (polar thermoplastic or thermoset composites) of the launcher, via a bonding or welding process, would presumably necessitate the introduction of functional groups into MBL-I-MBL. The copolymerization of isoprene with a small amount of a polar monomer could be envisioned, as performed in this study between β-myrcene and glycidyl methacrylate bearing an epoxide group. The functionalization of P(MBL) during the chain-extension would be another alternative. Summary outline of a potential linear triblock copolymer thermoplastic elastomer envisioned for further study. Di-I corresponds to a difunctional initiator.

• Characterization. The overall monomer conversion X was determined by gravimetry Xgrav and 1 H nuclear magnetic resonance (NMR) XNMR (average deviation = 3.12 %, standard deviation = 2.48 %, see Supporting Information for Xgrav and XNMR values in the course of experiment My-6, Figure S1a) and calculated from formula 1:

Xgrav was calculated from the following formula:

in which mv is the mass of the empty vial, mv+p+s is the mass of the vial containing the reaction solution just after the sample withdrawal (polymer, unreacted monomer(s) and solvent if used), mv+p is the mass of the vial containing the polymer after drying and wS is the mass percentage of solvent relative to the initial masses of initiator and monomer(s).

In order to prevent any overestimated Xgrav, drying of the samples was performed first under intense air flow for 24 h and then under vacuum at 50 o C for 24 h to maximize the removal of non-polymer components. Average conversion XNMR was then calculated from formula 3: XNMR = XMy fMy,0 + XS fS,0 (3) where fMy,0 and fS,0 are the initial molar fractions of My and S respectively and XMy and XS are the individual conversions of My and S respectively. XMy and XS were determined with a Varian NMR Mercury spectrometer ( 1 H NMR, 300 MHz, 32 scans) using CDCl3 deuterated solvent. The signal of the solvent (δ = 7.27 ppm) was used as reference for chemical shifts. My conversion was calculated by comparing the integrated peaks corresponding to the four aliphatic protons of the monomers (δ = 2.15-2.30 ppm), the eight aliphatic protons of the polymers (δ = 1.90-2.15 ppm) and the six protons of the two methyl groups of both monomers and polymers (δ = 1.55-1.75 ppm). S conversion was determined using the three vinyl protons (δ = 6.7-6.8, 5.7-5.8 and 5.2-5.3 ppm) of the monomers and the five aromatic protons of both monomers and polymers (δ = 6.9-7.5 ppm). Please note that X was only used for homopolymerizations. In the case of the copolymerizations, XNMR was considered since My and S have significantly different molar masses.

The regioselectivity of the My repetitive units in the various My-based polymers was also determined by using the same spectra ( 1 H NMR, 300 MHz Varian NMR Mercury spectrometer, CDCl3 eluent, 32 scans). Comparing the three integrated peaks at δ = 4.70-4.80 ppm (two vinyl protons of 3,4-addition and two vinyl protons of 1,2addition), δ = 5.00-5.25 ppm (two olefinic protons of 1,4-addition, one olefinic proton of 1,2-addition and one olefinic proton of 3,4-addition) and δ = 5.30-5.50 ppm (one olefinic proton of 1,2-addition) allowed the three different types of configurations to be quantified 2 . The two stereoisomers of 1,4-P(My), cis-1,4-P(My) and trans-1,4-P(My), were quantified using 13 C NMR (300 MHz Varian NMR Mercury spectrometer, CDCl3 eluent at δ = 77.4 ppm, 1000 scans) chemical shifts of specific nuclei: methylene carbon at δ = 37-38 ppm (trans-, 37,6 ppm; cis-, 37.1 ppm) and the quaternary carbon at δ = 131-132 ppm (trans-, 131.7 ppm; cis-, 131.3 ppm) 3 . Deconvolution (Mnova ® software, GSD sample withdrawn, the crude polymer was precipitated with excess methanol and then dried first under intense air and at 50 °C under vacuum. Nuclear magnetic resonance (NMR) and gel permeation chromatography (GPC) were used to characterize the samples. For the specific example cited, the overall conversion was 89.4% (individual conversions: XMy = 84.2% and XGMA = 94.6%) at the end of the experiment (t = 300 min) and the molar composition of GMA in the final copolymer was FGMA = 0.58 as determined by 1 H NMR. The final product exhibited Mn = 21.9 kg.mol -1 and Đ = 1.45, as determined by GPC. All final My/GMA copolymer characteristics can be found in Table 3. The exact same procedure was followed for My homopolymerization (Supporting Information, Table S2b) and copolymerizations of My with MMA, tBuA and MA. The formulations used for these latter experiments are summarized in Table 1. Figure 1 sums up the various copolymerizations performed.

• Chain-extension of P(My-stat-GMA) macroinitiator with My, GMA and S.

All chain-extensions from P(My-stat-GMA) macroinitiators were performed in a very similar setup to that used for the My/GMA copolymerizations with the use of a 25-mL reactor. All formulations can be found in Table 4B. As a brief illustration, My/GMA-78-GMA/My was synthesized using a statistical P(My-stat-GMA) copolymer (My/GMA-78, FMy = 0.78, Mn = 11.3 kg.mol -1 , Đ = 1.27, 1.011 g, 0.0895 mmol) that was added to the reactor along with toluene (7.615 g), My (0.603 g, 4.426 mmol) and previously purified GMA (5.967 g, 42.021 mmol). The contents were mixed and bubbled with nitrogen for 30 min, then heated to 110 o C and allowed to react for 60 min while maintaining a nitrogen purge.

Samples were drawn periodically via syringe. The samples and the final block copolymer were precipitated in excess methanol and allowed to dry overnight in a vacuum oven at 50 o C. My/GMA-78-GMA/My exhibited Mn = 29.1 kg.mol -1 , Đ = 1.50 (GPC) and FGMA = 0.65 ( 1 H NMR). The results of the chain-extensions from P(My-stat-GMA)s are given in Table 4C. The same procedure was applied for the synthesis of P(My-b-GMA) diblock copolymer from P(My) macroinitiator (Supporting Information, Table S2b).

• Synthesis of poly(β-myrcene-block-2-hydroxy-3-morpholinopropyl methacrylate) P(My-b-HMPMA).

A conventional reflux apparatus with a 10-mL three-necked round-bottom glass flask was used. P(My-b-GMA) diblock synthesized previously (FMy = 0.58, Mn = 23.8 kg.mol -1 , Đ = 1.89, 0.210 g, 0.009 mmol, Table S2b in Supporting Information) was dissolved in Me-THF (~ 5.3 g) and then morpholine (0.330 g, 3.791 mmol, 6.4 eq. relative to GMA repeating units) was added. The reaction medium was vigorously mixed and deoxygenated via N2 bubbling (explosive morpholine vapor/air may be formed above 35 o C) for 30 min before heating to 77 o C for three hours while maintaining a gentle nitrogen purge. The final product was precipitated in reverse osmosis water (formation of a dispersion which aggregated with time), decanted, filtered via a Büchner funnel and the resulting white polymer was dried at 50 o C under vacuum. A quantitative conversion of GMA units in HMPMA units was determined by 1 H NMR (Figure 9). 91% of P(My-b-HMPMA) was produced, exhibiting Mn = 26.5 kg.mol -1 and Đ = 1.91 (Supporting Information, Figure S11b). The complete experimental conditions and results can be found in the Supporting Information, Table S2b. Dynamic Light Scattering (DLS) was eventually used to measure the hydrodynamic diameter of P(My-b-HMPMA) particles dispersed in water (in the range of 120-130 nm under optimized experimental conditions, Table 5).

Theoretical estimation of ∆GM for a mixture of P(My) and P(GMA):

• Introduction: The Gibbs free energy (∆GM) describes the thermodynamics of the system and is a function of entropy (∆SM), enthalpy (∆HM) and temperature (T). ∆GM is represented by the following Equation E1:

When mixing two different polymers, the change in free energy will indicate whether the polymers are miscible (∆GM < 0) or a phase separation occurs (∆GM > 0). A statistical thermodynamic model was developed for the mixture of polymers and the free energy of mixing was given by Equation E2:

where k = 1.38.10 -23 J.K -1 is the Boltzmann's constant, V is the total volume of the mixture, Vr is the molar of a specific segment, vi is the volume fraction of polymer ``i``, z is the lattice coordination number (usually between 6 and 12), Nc is the number of chains per volume unit and χ is the Flory-Huggins parameter, which can be calculated from Equation E3:

where Vm,1 is the molar volume of the polymer with higher molar amount or, if equal composition, the one with smaller molecular weight monomer (for My/GMA system, Vm,1 = Vm,GMA = 136.4 cm 3 .mol -1 ), R = 8.31 J.mol -1 .K -1 is the ideal gas constant, δi is the solubility parameter of the polymer ``i``8 ,9 .

• Estimation of δMy, δGMA and ӼMy-GMA: δMy and δGMA were calculated using the Hansen solubility parameters (HSPs) δD (atomic dispersion forces), δP (molecular polar forces arising from dipole moments) and δH (hydrogen-bonds between molecules) 10 . HSPs are based on the concept that the total cohesive energy density is approximated by the sum of the energy densities required to overcome δD, δP and δH. The solubility parameters δMy and δGMA can thus be determined as described in Equation E4: • Estimation of ∆GM: Let us take a mixture of 1 g of P(My) and 1 g of P(GMA) with Mn,P(My) = Mn,P(GMA) = 15 kg.mol -1 . The total volume of the mixture V can be calculated as follows:

V = VP(GMA) + VP(My) = mP(GMA) / ρP(GMA) + mP(My) / ρP(My) = 1 / 0.805 + 1 / 0.895 = 2.36 cm -3 (m = mass, ρ = density, ρP(My) = 0.895 g.cm -3 12 )

The volume fractions of P(GMA) and P(My) are then determined: υP(GMA) = VP(GMA) / V = 0.526 and thus υP(My) = 0.474. The ratio V/Vr corresponds to the number of molecules (N) in the mixture: Where Na is the Avogadro constant and MGMA, the molecular weight of GMA monomer.

Moreover: Nc = NP(GMA) + NP(My) = Na (mP(GMA) / Mn,P(GMA) + mP(My) /Mn,P(My)) = 6.025.

Consequently, a mixture of 1 g of P(My) and 1 g of P(GMA) would be miscible at T = 25 o C if Mn,P(My) = Mn,P(GMA) < 1.2 kg.mol -1 , according to this thermodynamic model. discarded and were not used in mechanical testing. For each specimen, the thickness and the width were measured at five different points along the small center portion with a digital caliper (Marathon, CO 030150F128) and the average film thickness and width were used for all calculations. The samples tested exhibited a mean thickness from 0.55 to 0.69 mm and a mean width from 3.11 to 3.32 mm. Each test was considered finished after the complete breakup of the specimen at the narrow section (films that broke near the grips were discarded). 6 specimens were tested, and averaged results were reported using TestWorks 4 software. The Young's modulus was determined as the slope of the stress-strain curve at strains of 0-0.5%. • Solubility parameter of P(IBOMA): Addition of the group contributions for P(IBOMA) gives:

Fdi (MPa • Hydrogenation of the synthesized S-I-S triblock copolymers. PI mid-segment of the triblock copolymer was saturated using the method developed by Hahn 2 , which does not require the use of high-pressure H2. The unsaturated polymer was first dissolved in p-xylene (35-40 mL per gram of polymer) at 50 °C, and then p-toluenesulfonyl hydrazide (TSH, 2.4 equiv. to the double bond), tributylamine (TBA, 2.5 equiv. to the double bond) and small amount of 2,6-di-tert-butyl-4-methylphenol (BHT) were added. The reaction mixture was stirred at 125 °C for 3 hours under nitrogen atmosphere, and then allowed to cool to room temperature. The polymer was precipitated by adding cold methanol to the solution, and the methanol was decanted after stirring for 1 hour. The polymer was purified by repeating the precipitation using p-xylene/methanol system, and then dried under vacuum at 50 °C overnight. The dry and partially hydrognenated S-I-S was characterized by 1 H NMR, GPC, TGA and tensile testing.

• Characterization. Isoprene conversion was determined by gravimetry. IBOMA conversion was obtained using The number-average molecular weights (Mn) and the dispersities (Đ = Mw/Mn) were measured using gel permeation chromatography (GPC, Water Breeze, differential refractive index RI 2414 detector, 40 o C) with HPLC grade THF as the mobile phase (flow rate of 0.3 mL.min -1 ). The GPC was equipped with 3 Waters Styragel ® HR columns (HR1 with a molecular weight measurement range of 10 2 -5 x 10 3 g.mol -1 , HR2 with a molecular weight measurement range of 5 x 10 2 -2 x 10 4 g.mol -1 and HR4 with a molecular weight measurement range of 5 x 10 3 -6 x 10 5 g.mol -1 ) and a guard column was used. Mn values were determined by calibration with 10 linear narrow molecular weight distribution PI standards (Polymer Standards Service, molecular weights ranging from 1,030 to 1,040,000 g.mol -1 ).

Please see the characterization section of Chapter 3 for the detailed protocols of DSC, TGA, DMA, AFM analyses and tensile testing (identical procedures).

• References of the experimental section.