# Influence of polymer structure and structuration at different scales for improved properties towards sustainable materials

Matthieu Gervais

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Matthieu Gervais. Influence of polymer structure and structuration at different scales for improved properties towards sustainable materials. Materials. Conservatoire national des arts et metiers - CNAM, 2023. tel-04293234

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Submitted on 21 Nov 2023

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Conservatoire National des Arts et Métiers

Mémoire d'Habilitation à Diriger des Recherches

Présenté par

Matthieu Gervais

Maître de Conférences au Conservatoire National des Arts et Métiers
Laboratoire PIMM - UMR 8006 ENSAM - CNRS - CNAM
Équipe « Polymères et Composites »

Influence of polymer structure and structuration at
different scales for improved properties towards
sustainable materials

Soutenue le 16 novembre 2023

Jury:
- Laurence Rozes, Professeur, LCMCP, Sorbonne Université
- Odile Fichet, Professeur, LPPI, Cergy Université
- Sylvain Caillol, DR, ICGM, Université de Montpellier
- Jean-Marie Raquez, Professeur, LPCM, Université de Mons
- Thomas Karbowiak, Professeur, PAM, Institut Agro Dijon
- Cyrille Sollogoub, Professeur, PIMM, CNAM

Présidente
Rapportrice
Rapporteur
Rapporteur
Examinateur
Examinateur

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le cnam

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Table des matières

<table><tr><td>I. General introduction</td><td>5</td></tr><tr><td>I.1. Personal presentation</td><td>5</td></tr><tr><td>I.2. Research activities</td><td>6</td></tr><tr><td>Scientific overview</td><td>8</td></tr><tr><td>II. Multinolayer coextrusion for eco-design of food packaging films</td><td>10</td></tr><tr><td>II.1. Principle of multinolayer device</td><td>10</td></tr><tr><td>II.2. EVOH crystallization behavior</td><td>12</td></tr><tr><td>II.2.1. Morphology of crystallization</td><td>12</td></tr><tr><td>II.2.2. Kinetics of crystallization</td><td>13</td></tr><tr><td>II.2.3. EVOH crystallization under confinement</td><td>15</td></tr><tr><td>II.3. Multinolayers food packaging films made of polyethylene and EVOH</td><td>19</td></tr><tr><td>II.3.1. EVOH combined with PE: "soft" confinement</td><td>19</td></tr><tr><td>II.3.2. EVOH combined with another EVOH: "hard" confinement</td><td>23</td></tr><tr><td>II.4. Conclusion</td><td>27</td></tr><tr><td>III. High gas barrier coatings</td><td>28</td></tr><tr><td>III.1. Polymer matrixes coatings</td><td>30</td></tr><tr><td>III.1.1. Coating processing</td><td>30</td></tr><tr><td>III.1.2. Helium permeability of pure polymer coatings</td><td>32</td></tr><tr><td>III.2. Addition of nanofillers towards high gas barrier coatings</td><td>32</td></tr><tr><td>III.2.1. VMT dispersion in the coatings</td><td>33</td></tr><tr><td>III.2.2. Helium permeability of VMT nanocomposite coatings</td><td>34</td></tr><tr><td>III.3. Chemical modification of chitosan for humid gas permeability improvement</td><td>37</td></tr><tr><td>III.4. Conclusion</td><td>41</td></tr><tr><td>IV. Mechanical properties improvement by blends of polymers</td><td>42</td></tr><tr><td>IV.1. Upcycling of plastics from WEEE</td><td>42</td></tr><tr><td>IV.1.1. Control of the morphology of a ABS/PC blend</td><td>43</td></tr><tr><td>IV.1.2. Morphology characterization</td><td>44</td></tr><tr><td>IV.1.3. Mechanical properties of structured blends</td><td>46</td></tr><tr><td>IV.1.4. Chemical compatibilization of the blend</td><td>47</td></tr><tr><td>IV.1.5. Conclusion</td><td>48</td></tr><tr><td>IV.2. Improvement of PLA/PHBV blends mechanical properties</td><td>49</td></tr><tr><td>IV.2.1. Compatibilization of PLA/PHBV blends</td><td>49</td></tr><tr><td>IV.2.2. Preparation of oligomers of PLA</td><td>50</td></tr><tr><td>IV.2.3. Thermal characterization of compatibilized PLA/PHBV blends</td><td>51</td></tr><tr><td>IV.2.4. Morphology characterization</td><td>52</td></tr><tr><td>IV.2.5. Mechanical properties</td><td>53</td></tr><tr><td>IV.3. Conclusion</td><td>54</td></tr></table>
<table><tr><td>V.</td><td>Block copolymer self assembly and its combination for actuator application</td><td>55</td></tr><tr><td>V.1.</td><td>Block copolymer combination with ionic liquids</td><td>55</td></tr><tr><td>V.1.1.</td><td>Specific incorporation of IL in PMMA phase</td><td>56</td></tr><tr><td>V.1.2.</td><td>Block copolymer morphology evolution with IL content</td><td>58</td></tr><tr><td>V.2.</td><td>BCP/ILs combination mechanical properties</td><td>59</td></tr><tr><td>V.3.</td><td>Actuator application of BCP/IL combinations</td><td>60</td></tr><tr><td>V.4.</td><td>Conclusion</td><td>63</td></tr><tr><td>VI.</td><td>Conclusions and Research project</td><td>64</td></tr><tr><td></td><td>Theme 1: Effect of humidity on the barrier properties of structured polymer materials</td><td>65</td></tr><tr><td></td><td>Theme 2: Improvement of mechanical and electromechanical properties by polymer blend structuration</td><td>68</td></tr></table>
# I. General introduction

In this manuscript I will present the research work I've been carried out since 2012 as a Maître de Conférences in PIMM's laboratory and then I will explain my plans for my future research over the following years. This manuscript will try to explain my goals of merging scientific interests to societal impact and especially over the environnemental aspects that are linked to the academic field that I'm dedicated to, meaning polymeric materials science. My understanding of the academic position that I'm in is to allow our society to think of different ways of acting on environnemental aspects that our world is facing nowadays. In the field of materials, this is closely linked to industrial issues and engineering problems.

## I.1. Personal presentation

After a master's degree in polymer chemistry I joined the Laboratoire de Chimie des Polymères Organiques (LCPO) in Talence for my PhD on anionic polymerization of epoxides. After this PhD, I did several post docs. The first one was realized in LCPO as well but on a very different topic as it was my first research study really focused on materials. The aim was to study several composites and the formulation of the matrixes of those composites in terms of mechanical properties (impact test) for a helicopter pilot helmet. After this first encounter with materials, I did a post-doc in the USA, in the Lawrence Berkeley National Laboratory (LBNL) (California) on the synthesis of polymers for fuel cell and Lithium-ion batteries applications. After two years in the USA, I came back to France and started another post-doc position in the Laboratoire Analyse Modélisation Matériaux pour la Biologie et l'Environnement (LAMBE) in Evry (France). I worked on the condensation of a silane in interaction with DNA components. This involved precise NMR studies especially. Finally, I joined the CNAM in 2012 as an Assistant Professor in the Procédés et Ingénierie en Mécanique et Matériaux (PIMM) laboratory. The scientific fields studied in PIMM are linked to material science and in particular, the link between polymer processing, micro/nanostructure and properties (mostly mechanical and gas barrier). The aim of my recruitment, with a chemist profile, was to allow the understanding of the effect of the chemical structure on the different processes when nanoscale is at stake. Indeed the PIMM lab had developed since 2007 an innovative process tool that allows the creation of films composed of thousand of stacked layers. The thicknesses attained in those layers were close to the size of a polymer chain itself and so an understanding of the behaviour of the influence of the chemical structure was then needed and a chemist point of view could help understanding those phenomena. To meet this goal, however, as a chemist, I needed to develop my knowledge on material science more deeply and also on polymer processing. Furthermore the PIMM lab is not a synthesis lab and so I had to modify my way of doing research at that time.

An adaptation period was then required for me to fully develop a scientific methodology and understand the scientific locks that I can try to tackle. Fortunately, my colleagues involved me in different research projects and I could develop new skills related to polymer processing and microscopical analyses as well as a more physics-related approach of the scientific locks. It helped me change my perspective from the relation chemical structure/properties to a larger chemical structure/micro-nanostructure/properties. My research is now linked to this triptych combined with a close look at the environmental impact of the results of my research. Indeed, as I came closer to the material science and to the direct industrial uses of those materials, my environmental consciousness became more important in my scientific topics choice. For example, food packaging became one of my most important topic and I enrolled in the Réseau Mixte de Technologie (RMT) PropackFood. I also take the pedagogic responsibility of an engineer's diploma based on the materials for packaging, located in Angoulême in addition to several other teaching responsibilities in Paris.

In parallel to the research topics, I also became the referee for sustainable development in PIMM's lab and did the greenhouse gas emissions report (BGES in French) of the lab.
## I.2. Research activities

As mentioned above, the main interest guiding my research activities is the environnemental aspect of the polymeric materials. In order to carry a research that may have an impact on the environnement, I focused my activities on the most widespread application of plastics: packaging which corresponds roughly to 40wt% of the plastics production in the world. In this field, some environnemental issues are currently highly at stake. Indeed, the end-of-life of packaging is largely evolving with the increase of recycling efficiency lead especially by legal decisions which are themselves lead by population's will to get rid of plastic pollution that has been highly highlighted mediatically. In particular the plastic continent which contains 800 000 tonnes of plastics and represents an area of 1.3 millions km² retains a lot of attention.

However in this recycling evolution, some of the food packaging can not still be recycled due to packaging made of differents polymers blended together or multilayered and then can not be sorted separately which will pollute the recyclates leading to bad performances of them. This decreases the interest for those recycled polymers as the flow does not perform consistencilly. A large need for monomaterial food packaging is then at stake. But multilayer materials are used for a reason which is a need for high gas barrier performances to protect the food from oxygen and moisture to avoid contamination and spoilage. The usual way of obtaining good moisture and oxygen barrier properties is to combine two polymers having each one of them. The most widespread one is a combination of polyolefin (polyethylene (PE) or polypropylene (PP)) for moisture barrier properties and ethylene-vinyl alcohol copolymer (EVOH) for oxygen barrier properties. The amount of EVOH is usually lower than the one of the polyolefin as its cost is higher. My research activities are then focused on a way to obtain food packaging that may be recycled without any bad impact on the recyclate and that will maintain those good barrier properties to prevent the spoilage of the food which is the worst scenario environmentally speaking.

To do so, different paths have been followed which all are linked to the same idea which is to decrease the amount of the oxygen barrier polymer enough to make it negligible towards the recycling of the material. In those paths have been considered the improvement of the current polymer, i.e EVOH, through the control of its cristallization which is supposed to be an important factor for its barrier properties. To do so, multilayer coextrusion has been used to multiply and by then decrease the size of EVOH layers in order to obtain a confinement of EVOH that will induce an evolution of its cristallinity. For this kind of research, my activities then tend to study the cristallization behavior of polymer especially in confined environnement. Also, the confinement and multi-layered structures can have different effects on EVOH behaviour in general in addition to the crystalline phase. These effects have been studied, in particular under humid conditions.

Another path is to use a different polymer than EVOH that will present better barrier properties. Thus to obtain similar properties than the current packaging, smaller amount of this polymer will be needed. For example the use of pure vinyl alcohol polymer (PVOH) without copolymerization with PE leads to better properties. However, generally, those polymers possess good barrier properties which are due to a large hydrogen bonds network but this leads to a more complex processing through classical thermal techniques like extrusion. To avoid this problem, a different processing methodology has been followed. The first one will be to process the polymer in solution, through coatings mainly. One of my approach has then been to develop high barrier coatings composed of several different polymers. This approach allows also the addition of loads in the coating that may improve the barrier properties. I've followed this approach with the addition of loads of the nanometric size in large quantities compared to what is usually done in the literature (>25 wt%). Those approaches allowed me to work on viscosity problems, loads dispersion and barrier properties. This also implied to consider the impact of humidity on the properties as a coating is not protected from the external relative humidity. As the barrier properties are linked to the hydrogen bonds network, the humidity has a dramatic effect on them by plasticizing the coating. I've tried to tackle this issue through the chemical modification of the polymer used and by modifying the hydrophilicity of it.
Then, in another approach, and to improve the end-of-life of polymer materials, I also tried to improve the recycling of already used materials. To tackle this issue, in a similar approach, I focused on some of the largest flow of recycled materials which is the one emanating from the Waste of Electric and Electronic Equipments (WEEE). In this, one of the main problems of those recycled polymers flow is that the different polymers have been aged during their service lifetime and can not be used as virgin materials due to lowered performances. The idea here is to improve those properties using a structuration of the matter through the multinanolayer device already discussed for the food packaging application. The scientific lock that I want to tackle here is the link between a structuration down to the nanoscale of a blend of polymer with the mechanical properties and especially in this case with the impact properties. Two polymers largely present in the WEEE were thus used, a copolymer made of acrylonitrile – butadiene and styrene (ABS) and a polycarbonate (PC). Nanostructured pellets were realized and dogbone samples were injected from these pellets. The retention of the nanostructuration through the injection step was studied and the impact of the structure obtained on the impact properties was observed. An improvement of this property was observed with the size and orientation of the domains of PC in the ABS matrix.

Also, in order to avoid the increase in polymeric wastes, the use of biodegradable plastics have been a topic of interest for a lot of researchers and industrial companies for decades now. I also try to tackle this topic by using different kinds of bio-based polymers that present a certain level of biodegradability (polylactide (PLA) for example). Blends are usually used due to the fact that the mechanical properties, as well as the gas barrier properties for food packaging applications, of those polymers can not withstand the one of competing petrosourced polymers. To improve these, I worked on the classical addition of poly(3-hydroxybutyrate-co-3-hydroxyvalerate) (PHBV) in a PLA matrix to modify the brittle behavior of the PLA. But, to improve that blend, the addition of a compatibilizer is needed to obtain a good synergy of it. I then worked on the addition of oligomers of PLA to act as a compatibilizer and plasticizer of PLA/PHBV blend.

Finally, in an final approach of the impact of structuration on material properties, the self-assembly of block copolymers was used to create this structuration. Nanodomains of a block copolymer made of poly(methyl methacrylate) and poly(butyl acrylate) were obtained. The effect of the addition of an ionic liquid on the structuration was studied. This ionic liquid was added also to inject an ionic conductivity in the material. This ionic conductivity was used to develop an electromechanical response of the material under a voltage. The impact of the morphology on this electromechanical response is then at stake and actuators have been developed and characterized in order to understand the behaviour of the ions trapped in a nanostructured polymeric environment.

To sum up, different approaches of nanostructuration have been employed to create nanodomains of different polymeric structures. The effect of the morphologies developed were then studied regarding different properties according to the applications envisaged. The evolution of gas barrier properties regarding the multilayer structure or through a composite coating have been studied, especially in regard to the relative humidity. The evolution of mechanical properties was also investigated regarding a multilayer structure or through bio-based materials compatibilized. Finally, the electromechanical properties of a nanostructured block copolymer structure additivated with ionic liquids were characterized.

Two main issues emerged from these studies which will be my main interests in my future research : the effect of moisture on gas barrier properties of nanostructured polymer and the hydrophily/hydrophoby balance required to develop the best properties. The second will be the control of the orientation and type of morphology for mechanical and electromechanical properties and the modelling of these behaviors.
Scientific overview

As described in the introduction of this manuscript, the environmental impacts of the plastics are my main interest. In order to deal with this problem, I have chosen to deal with different topics. First, I'll present the work realized for the food packaging application. This will be divided into two different approaches with the multilayer coextrusion device and the coating approach. Then, I'll focus on the improvement of mechanical properties of plastics through blending and firstly in the view of recycling WEEE plastics and secondly on the improvement of the mechanical properties of biobased polymers. Finally I'll describe the use of electroactive functionnalities in a combination of plastics with ionic liquids. This scientific overview will be divided into four parts for each of these issues.

My first focus will be on food packaging applications as 40 wt% of the produced plastics is for the packaging industry and correspond to the part of plastics that will present the shortest life time before its disposal. This field is then the most problematic one for plastic disposal and an improvement of this issue is required. The regulation is also evolving and especially in France the AGEC law requires the food packaging to become all recyclable and single-use plastics are prohibited¹. However, food packaging needs to possess properties such as transparency, being biologically inert but one of the most prominent and complicated to achieve is the high gas barrier properties as well as moisture barrier. Indeed to prevent the spoliation of the food which is the worst case scenario environmentally speaking², the food has to be protected from moisture and oxygen that would cause rotting and oxidation of the it. Therefore both oxygen and moisture needs to be stopped by the packaging. The main issue is that usually one plastic is not capable of being a good barrier to both. The solution developed classically industrially is to combine at least two different polymeric materials that will each block one of the problematic gases: one for dioxygen and one for water vapor. The polymers used for gas barrier properties exhibit usually a dense hydrogen bonds network that will then lead to a impermeable material as no specific movement is able in this matter. It also leads to a high sensitivity to water vapor of those materials due to the hydrogen bonds that induce a large drop of the gas barrier properties in humid conditions. The polymers used for those applications are usually poly(vinyl alcohol) (PVOH), a random copolymer of ethylene and vinyl alcohol (EVOH) and some polyamides³.

To protect those high gas barrier polymers from water vapor, hydrophobic polymers are used in combination to the oxygen barrier polymers to avoid any contact between the hydrophilic one with the moisture. Those polymers are usually polyolefins such as polyethylene (PE) and polypropylene (PP) or poly(ethylene terephtalate) (PET)⁴. PP is mainly used when the packaging will have to go through a retorting process which implies to heat up the packaging at a temperature higher than 100°C which could not be endured by PE without a complete melting of the polymer as Tm (PP) ~160°C and Tm (PE) ~100°C. The retorting process involves sealing a food product in a hermetically sealed container and applying sufficient heat to render the product commercially sterile. These polymers are then often coextruded in order to obtain a multilayer structure with the hydrophilic polymer surrounded by hydrophobic ones for its protection. These structures are very efficient for food preservation and have been widely used until now. However, their multimaterial nature is a major inconvenient for their mechanical recycling as the shredding of these materials will create a flow of mixed polymers which will present poor properties compared to a flow of pristine and virgin polymer. The reuse of this flow is then complicated and it is then quickly discarded. To avoid this, the amount of gas barrier polymer should be the lowest possible that will not have any disturbing effect on the recycled flow of the main polymer. Indeed PE and PP are currently recycled polymers as their flows are large enough to allow the existence of a recycling industry of them.

The multimaterial nature of those packaging is then the problem of those products. The objective is so to develop polymers that will present similar levels of gas barrier properties but with the lowest amount of a second material inside a potentially recyclable one.

Another problem that will have to be taken care of is the influence of water vapor on the behavior of the gas barrier polymer. Indeed, during the retort process, the increase of temperature of the system decreases the water vapor barrier properties of the surrounding polymer (PP) and water may reach the central barrier
polymer. With the decrease of temperature after the process, the water remains trapped into the central layer and will affect durably the oxygen barrier properties. The development of a system that is not impacted in those harsh conditions would improve drastically the lifetime of the food.

Different approaches will then be followed: the coextrusion of EVOH with an hydrophobic polymer in a multilayer approach with the multiplication of layers in order to achieve nanometric thick layers and the use of coatings made of highly gas barrier polymers and their chemical modification to sustain their properties in humid conditions. In an environmental approach, the use of biobased polymers and/or nanoloads will also be favoured.
## II. Multinanolayer coextrusion for eco-design of food packaging films

The eco-design of food packaging films is highly at stake currently and a good combination of light materials with good properties and recyclability are mandatory. To answer these issues, the plastic food packaging has many interesting properties such as its weight and versatility of properties. However, its end of life presents many problems and a better design to facilitate its recycling is needed. To do so, the multimaterial nature of those films needs to be limited by adding the minor material in a proportion low enough to avoid any impact on the recycling of the major material. In this chapter, I'll describe a way to improve the gas barrier properties required for this application through a processing step that will allow the creation of films composed of multiple layers of the minor and major materials with thicknesses as low as several nanometers which will modify the behaviour of those polymers.

### II.1. Principle of multinanolayer device

The first way to develop high gas barrier properties «mono» material is by the development of multinanolayered structures through layer multiplying coextrusion. This process consists on two extruders (A and B in Figure 1) that will be combined into a 3 layers flow in a feedblock and this flow is then divided, spreaded and recombined in layer multiplying elements (LME) (Figure 1 b) leading to the multiplication of the number of layers as described in Figure 1 a⁵. With n LMEs, the total layer number of created layers can be predicted by a simple equation, see eq. 1. A final extruder can also bring an additional external layer that will surround the multiplier core (Extruder C) prior to the rectangular die. Finally, the film obtained in pulled using chill rolls with controlled temperature and speed of drawing that will control the final thickness of the films. Two conditions can then be obtained: a multiplied (A/B/A)ₙ film or a C/[(A/B/A)ₙ]/C film.

<figure><img src="image_4.png" /><figcaption>Figure 1. a) Multiplication of layers principle by the multilayer coextrusion process; b) Layer multiplier element (LME) representation</figcaption></figure>

For a feedblock giving three initial layers, the total number of layers N will be:

$$N = 2^{(n+1)} + 1 \qquad \text{eq. 1}$$

With n the number of LME used. Also, the thickness of each layer can be determined through eq. 2:

$$e_A = \frac{e_{\text{film}} \times \Phi_A}{N_A} \qquad \text{eq. 2}$$
Where $e_A$ is the layer thickness of the material A, $e_{film}$ is the film thickness, $\Phi_A$ is the volume proportion of material A and $N_A$ is the number of layers of the material A. By varying the number of LME from 0 to 12 in the system, multilayered films having number of layers varying from 3 to 8193 are created, with each layer thickness that can reach tens of nanometers close to lamellar dimension while total film thickness remains unchanged.

Pioneering work on this process has been performed by Eric Baer's team of the National Science Foundation center for Layered Polymeric Systems (CLiPS) with the collaboration of Dow Chemical. Starting from a microlayer system designed in the nineties, improvements have permitted to reach homogeneous and continuous multilayered films with nanolayer thickness with amorphous polymers couple$^{6-8}$. It has thus opened new horizons in the field of architECTured nanomaterials for many applications as electronic, optic or packaging domains$^{9,10}$. This team and others have then worked on the relationship between structure and properties of such composites and particularly on the crystallization under confinement of numerous polymers induced by a substrate in the multinanolayer coextrusion$^{11-15}$.

This was specifically used for the improvement of gas barrier properties through the confinement of a semicrystalline polymer, the poly(ethylene oxide) (PEO), by a random copolymer made of ethylene and acrylic acid (EAA) in a multinanolayered structure. Thickness reduction of PEO from 20 µm to 25 nm has permitted a morphology transition from 3-D spherulite to single in-plane lamella, as can be seen on Figure 2 a) and b), revealing two orders of magnitude decrease of oxygen permeability from 0.3 to 0.0052 barrier, Figure 2 c)$^{16}$. Same results were obtained with a confinement induced by PS without being affected by moisture. Oxygen diffusion coefficient follows the same trend and solubility remains constant with layer thickness while PEO crystallinity decrease from 75% to 68%$^{17}$.

<figure><img src="image_5.png" /><figcaption>Figure 2. Morphology transition of PEO induced by multinanolayer coextrusion process followed by WAXS a), AFM b). PEO layer and multilayered film oxygen permeability are plotted in function of PEO layer thickness c) adapted from ¹⁷</figcaption></figure>

This enthusiastic result has however been realized on a polymer which doesn't present high gas barrier properties intrinsically. The aim is then to adapt this methodology on the classical polymers used in food packaging such as copolymers of ethylene-vinyl alcohol (EVOH)$^{18}$.

In order to develop similar results with EVOH, the crystallinity behavior of this latter has to be perfectly understood as in-plane lamellae will be needed to improve the gas barrier properties. This behavior has then
been studied extensively, in bulk and more importantly under confinement to observe the possible
orientation of the EVOH crystalline lamellae¹⁹.

## II.2. EVOH crystallization behavior

EVOH polymers are random copolymers composed of ethylene and vinyl alcohol units, Figure 3, and arise from the hydrolysis of ethylene/vinyl acetate copolymers²⁰. Those latter are synthesized through radical polymerization that leads to a random and branched structure. Different grades of EVOH are used in the food packaging industry and are named after the weight proportion of ethylene units incorporated, i.e. an EVOH44 contains 44 molar% of ethylene units.

<figure><img src="image_6.png" /><figcaption>Figure 3. EVOH chemical structure</figcaption></figure>

### II.2.1. Morphology of crystallization

According to this complex structure, the crystallinity has been studied in bulk to observe the characteristic
temperature, crystalline morphology and size of the crystalline structures. DSC, Polarized Optical
Microscopy (POM), in-situ AFM and Small Angle Light Scattering (SALS) were used to determine these
different parameters on EVOH44²¹. A spherulitic structure was observed with typical size below 1µm using
POM, SALS and AFM. However, the nucleation and spherulite growth could not be observed due to fast
crystallization as can be seen on the in-situ AFM imaging of the crystallization in Figure 4.

<figure><img src="image_7.png" /><figcaption>Figure 4. AFM phase image of EVOH44 during annealing at 145 °C a) and b) and cooling to room temperature c), at 150 °C, two different zones d) and e) and after cooling f). The scale bar corresponds to 1 µm.</figcaption></figure>
II.2.2. Kinetics of crystallization

The kinetic of crystallization was then further investigated using in-situ X-ray diffraction and Flash DSC.
Indeed, isothermal crystallizations experiments realized on conventional DSC could not allow us to observe
a complete peak of crystallization as the phenomenon starts prior to the stabilization of the apparatus on
the specified temperature as can be observed in Figure 5b for temperature as high as 145°C where no
crystallization should occur with the cooling ramp used to reach the isothermal temperature (Figure 5a).

<figure><img src="image_8.png" /><figcaption>Figure 5. a) DSC thermogram of EVOH44 cooled at 2, 5, 10 and 50 °C/min and b) crystallized at different isothermal temperatures after a 200°C step to completely melt the crystalline phase with a 50°C/min cooling ramp to the isothermal temperature.</figcaption></figure>

To overcome this limitation, in-situ WAXS experiments were carried out in a Synchrotron apparatus, the Swing line in Soleil Synchrotron with a heating stage allowing high scanning rate, typically of 150°C/min compared to the maximum 50 °C/min obtained in conventional DSC. The evolution of different scattering peaks linked to the crystalline structure of PE and of PVOH has been identified and followed versus time at specific temperatures during isotherms. This evolution (Figure 6) exhibits the increase of the crystallinity during the isotherm, with the identification of three characteristic peaks corresponding to the following crystalline planes: (110) and (200) from the orthorhombic PE lattice and (100) from the PVOH monoclinic lattice²².

<figure><img src="image_9.png" /><figcaption>Figure 6. WAXS diffractograms of an EVOH44 film assembled during an isotherm experiment at 140 °C during 15 min</figcaption></figure>

This allowed us to plot the double logarithmic from equation 3 with α the relative crystallinity (or the conversion ratio), Figure 7a, to deduce from the slope of the curves in this domain the value of the Avrami
exponent n<sub>a</sub>, and from the intercept, the value of the crystallization rate k<sub>a</sub>, both reported in Table 1. Using those parameters, we can plot the evolution of the relative crystallinity as a function of time and compare the plots with the experimental values (Figure 7 b). The Avrami models fit well the experimental value for relative crystallinity below 0.6. However, for higher values of relative crystallinity (above 0.6), the experimental curves show a slower crystallization than the one predicted by the Avrami model, which is typical for the occurrence of a secondary crystallization (corresponding also to a bending of the double logarithm plots). The values of n<sub>a</sub> globally range between 2 and 3.

$$ \ln(-\ln(1-\alpha)) = \ln(k_a) + n_a \ln(t_0) \quad \text{eq. 3} $$

<figure><img src="image_10.png" /><figcaption>Figure 7. (a) Double logarithmic plot and (b) relative crystallinity as a function of time. The red lines correspond to the fits from Avrami model.</figcaption></figure>

<table><caption>Table 1. Avrami parameters determined from WAXS experiments.</caption><thead><tr><th>Crystallization temperature (°C)</th><th>n<sub>a</sub></th><th>k<sub>a</sub> (min<sup>-n</sup>)</th></tr></thead><tbody><tr><td>136</td><td>0.3</td><td>2.18</td></tr><tr><td>138</td><td>1.9</td><td>14.1</td></tr><tr><td>140</td><td>2.2</td><td>16.4</td></tr><tr><td>142</td><td>1.9</td><td>4.57</td></tr><tr><td>145</td><td>1.6</td><td>0.99</td></tr><tr><td>147</td><td>2.8</td><td>0.54</td></tr><tr><td>150</td><td>3.3</td><td>0.01</td></tr></tbody></table>

However, the results obtained with the WAXS must be considered with caution. As already noticed, at high undercooling, crystallization is very fast and may begin either during cooling or during the transient due to the switching from the cooling to the isothermal step. The data acquired each five seconds do not allow capturing with a high resolution the early stages of crystallization for these temperatures, crucial for the determination of the Avrami parameters. On the contrary, at low undercooling, the crystallization is incomplete during isothermal step and impingement of spherulites may not occur. In such conditions, the assumptions of Avrami model are not verified, especially the one assuming a complete transformation of the sample.

Flash DSC was then used to overcome these limits. A critical cooling rate able to suppress crystallization can be determined by melting the sample at 250 °C then cooling it down to -70 °C at different rates varying between 0.1 and 1000 K/s. When the cooling rate is equal to 1000 K/s, the subsequent heating ramp at 8000 K/s shows no melting, neither cold-crystallization signatures. Thus, both nucleation and crystallization from the melt are inhibited, proving that isothermal crystallization can be followed by Flash DSC on initially amorphous samples. Experiments could then be carried out to determine the Avrami parameters similarly
to WAXS experiments and the relative crystallization and double logarithmic plot are shown in Figure 8. The global consistency between results shows the advantages of Flash DSC in the investigation of crystallization kinetics.

<figure><img src="image_11.png" /><figcaption>Figure 8. (a) Crystallization advancement as a function of the annealing time and (b) Double logarithmic plot from Avrami equation 3. The solid lines represent the linear fittings.</figcaption></figure>

The Avrami parameters obtained from these results are then summarized in Table 2.

<table><caption>Table 2. Avrami parameters determined from Flash DSC experiments</caption><thead><tr><td>Annealing temperature</td><td>n<sub>a</sub></td><td>k<sub>a</sub> (min<sup>-n</sup>)</td></tr></thead><tbody><tr><td>125 °C</td><td>2.7</td><td>548</td></tr><tr><td>128 °C</td><td>2.5</td><td>39</td></tr><tr><td>130 °C</td><td>2.9</td><td>396</td></tr><tr><td>132 °C</td><td>2.7</td><td>33</td></tr><tr><td>134 °C</td><td>3.3</td><td>2</td></tr></tbody></table>

In the investigated temperature domain, the Avrami exponent n<sub>a</sub> is about 3 which is in good agreement with the spherulitic growth observed microscopically and by SALS experiments. The use of flash DSC was thus proven to be efficient for the understanding of the crystallization of EVOH44.

After this understanding of the crystalline structures and kinetics of crystallization of EVOH44, the most important parameter for gas barrier properties in multinolayer process is the orientation of lamellae under confinement as explained by Baer et al¹⁰. Indeed an in-plane orientation is needed to expect an improvement in barrier properties while an on-edge orientation would in contrary leads to a decrease in properties.

### II.2.3. EVOH crystallization under confinement

To develop oriented lamellae of EVOH, a confinement of the spherulites of EVOH is needed. Different ways are possible to obtain this confinement. Polymers can be trapped into multinolayers as already discussed¹⁰ but also by having thin films by spin-coating²³ or in nanodomains created by block copolymers²⁴ or in inorganic nanoporous templates²⁵. First, in order to be able to observe the lamellae and their orientation in a layer configuration, thin films of EVOH were spin-coated on silicon wafers but no direct observation of crystalline structure of EVOH was realized due to the harsh process that is spin-coating. Thus, to observe the appearance of lamellae, in-situ observation of spin-coated EVOH films by hot stage atomic force microscopy allowed us to investigate the effect of the film thickness and of the crystallization temperature on the crystal structure and orientation.
EVOH has been dissolved in a ethanol/H₂O (90/10 vol%) mixture at 60°C and spin-coated on silicon substrate. Those substrates were treated by different chemical treatments in order to add an hydrophilic (piranha water and UV beam) or hydrophobic (hydrofluoric acid and grating of benzylfunctions on the surface using benzylchlorodimethylsilane) nature to it. However the spin-coating of a regular thin EVOH film was rendered difficult because of the amphiphilic character of the copolymer, due to the presence of both hydrophilic (hydroxyl) and hydrophobic (ethylene) units. When spin coated, the thin film dewetted before complete evaporation of the solvent, leading to the formation of numerous irregular droplets on the substrate. The treatments realized on the substrate could not helped the improvement of this film quality due to the amphiphilic character also.

Observation of the formation of lamellae could nevertheless be attempted by in-situ AFM imaging. Like for the bulk film, the early stages of the crystallization were difficult to capture due to the long time required to scan an image (around 5 minutes). Crystalline lamellae were visible, both orientated in-plane and on-edge (Figure 9).

<figure><img src="image_12.png" /><figcaption>Figure 9. AFM phase images of an isothermal crystallization of EVOH (8 g/l) spincoated after 40 min a) and 70 min b). Oriented lamellae are indicated with arrows and height profile was measured with Nanoscope software c).</figcaption></figure>

In our case, it is difficult to correlate systematically an orientation of the lamellae with a specific parameter, for example the film thickness especially due to thickness heterogeneity of the droplets. Anyway, it seems that no transition due to thickness exists because both orientations coexist for coating thicknesses of the same order of magnitude. Still, it is possible thanks to height profiles, to measure the typical dimensions of a lamella, especially its thickness comprised between 15 and 20 nm. Moreover the treatment realized onto the surfaces did not make any change of behavior in orientation so no link between polarity of the surface and orientation of the lamellae could be made.

After confirmation that on-edge lamella could be obtained but without understanding on the control of the orientation, different multinolayer films have been processed to obtain films with EVOH layers in the range of the confining regime, i.e. around a few tenth of nm in the case of EVOH regarding the thickness of a lamella observed in AFM. Different systems have been processed with different confining polymers for EVOH. The first one used cyclic olefin copolymers (COC), Figure 10, which presents a glass transition temperature higher than the melting temperature of EVOH. Indeed, to allow a correct confinement of the polymer, the layers surrounding the confined layer need to become rigid before the crystallization of the polymer. If not, the crystallization may proceed as usual by pushing away the soft confining layers. This phenomenon called « hard confinement » requires then that the confining layer has for amorphous polymers
a glass transition temperature higher than the crystallization one of the confined polymer or for semi-crystalline polymers a higher crystallization temperature. In our case, the COC used presented a Tg of 178°C, measured by DSC, while the EVOH44 crystallizes as previously observed below 150°C.

<figure><img src="image_13.png" /><figcaption>**Figure 10.** Chemical structure of cyclic olefin copolymer (COC)</figcaption></figure>

Different films were produced with different numbers of LME to obtain different final EVOH layer thicknesses. Also PMMA was used as a skin layer after the multiplication of the layers, see Figure 1 in extruder C, to decrease the thickness of the layers in the total thickness of the films. The conditions obtained are summarized in Table 3. Thicknesses in the range of the confining regime could be obtained with 8 LME, a proportion of 9wt% of EVOH and a total film thickness of 15 µm.

<table><caption>**Table 3.** Characteristics of COC/EVOH multinolayered films</caption><thead><tr><th></th><th colspan="2">PMMA-(COC-EVOH44-COC)<sub>n</sub>-PMMA</th><th colspan="2">35-(10.5-9-10.5)<sub>n</sub>-35wt%</th></tr><tr><th>LME number</th><th>6</th><th>8</th><th>6</th><th>8</th></tr></thead><tbody><tr><th>Total film thickness (µm)</th><td>230 ± 11</td><td>160 ± 9</td><td>13 ± 2</td><td>15 ± 2</td></tr><tr><th>Number of layers</th><td>129</td><td>513</td><td>129</td><td>513</td></tr><tr><th>Nominal EVOH layer thickness (µm)</th><td>0.940</td><td>0.166</td><td>0.054</td><td>0.015</td></tr><tr><th>Measured EVOH layer thickness (µm)</th><td>0.696 ± 0.233</td><td>0.123 ± 0.041</td><td>0.069 ± 0.014</td><td>0.027 ± 0.06</td></tr></tbody></table>

The first important parameter is the layer continuity. This has been verified by different microscopic techniques depending on the thickness of the layers as can be seen Figure 11 along with the layer thicknesses depicted in Table 3.
<figure><img src="image_14.png" /><figcaption>Figure 11. Optical microscopy image in reflection mode of cross section of (COC-EVOH-COC)₈ film with individual EVOH layer thickness around 700 nm a). AFM phase images of cross sections of(COC-EVOH-COC)₈ films with individual EVOH layer thickness of 120 b), 70 c) and 30 nm d)</figcaption></figure>

As the layer continuity is verified, the effect of the layer thickness on the crystalline structure of the EVOH and lamellae orientation has been investigated through WAXS experiments. To study the orientation, the extruded films had to be observed in different directions as presented Figure 12. The differences observed between extrusion (ED) and transverse (TD) direction would be related to the stretching imposed by extrusion process while a difference between extrusion and normal (ND) would be related to confinement.

<figure><img src="image_15.png" /><figcaption>Figure 12. Directions of observations of a multinanolayered film: Extrusion Direction (ED), Transverse Direction (TD) and Normal Direction (ND)</figcaption></figure>

Figure 13a presents the scattering patterns of as-extruded films in the different directions for bulk and multinanolayered films. Patterns in ED and TD are similar, suggesting that an anisotropic result would mainly come from the geometrical confinement rather than extrusion flow. While for the bulk and 700 nm samples, diffraction patterns exhibit isotropic rings in all three directions, some anisotropy appears in the ED and TD patterns when the 120 nm layer thickness is reached. Indeed, some broad equatorial concentration of the (110) intensity suggests preferential orientation of lamellae. In order to confirm this
orientation, azimuthal scan of the (110) plane is plotted on Figure 13b. It appears that (110) crystalline plan orientation seems to shift from 90 ° with EVOH layer thickness reduction from 120 to 30 nm, which is possibly the result of transition from on-edge to in-plane lamellae. Still, the peaks from (110) and (200) planes become hardly discernable with the decrease of EVOH thickness and the interpretation of the orientation of the lamellae remains complicated. This observation suggests however that the confinement by a high Tg polymer, i.e COC here, may allow to induce an orientation of the crystalline structure within the layer when thin enough layers are created and the critical thickness to obtain this orientation seems to be between 120 and 30 nm for EVOH44.

<figure><img src="image_16.png" /><figcaption>Figure 13. a) WAXS 2D images of EVOH44 control and (COC-EVOH44-COC)ₙ multinanolayered films in extrusion direction (ED), transverse direction (TD) and normal direction to the film plan (ND); b) Azimuthal integration of the (110) plan for multinanolayered samples having 700, 120 and 30 nm EVOH44 layer thickness.</figcaption></figure>

Finally, as COC is not interesting for food packaging applications, no permeability measurements were carried out on this combination and more investigation were realized on more classical polymer combination in the food packaging industry.

## II.3. Multinanolayers food packaging films made of polyethylene and EVOH

### II.3.1. EVOH combined with PE: “soft” confinement

A classical combination of EVOH with a polyolefin was realized and multiplied in order to reach nanometer thick layers. The first combination attempted was the industrial combination with low density polyethylene
(LDPE) which is LDPE/tie layer/EVOH/tie layer/LDPE with the tie layer being a poly(ethylene)-grafted-maleic anhydride (PE-g-MA) which ensures the adhesion between PE and EVOH and prevents delamination of the final film. The anhydride maleic function of the PE-g-MA allows a chemical linkage with the hydroxy function of the EVOH creating interfacial copolymers EVOH-co-PE-g-MA²⁶ while the PE nature of it permits the adhesion to the LDPE phase. However, this combination is problematic in the multilayer coextrusion process as the interface created with a reactive polymer will induce strong defects that deteriorate the film quality with the apparition of white spots and thickness heterogeneities. Those defects are amplified when increasing the number of LME. Those defects have not been suppressed when replacing LDPE-g-MA by linear low density polyethylene-grafted-maleic anhydride (LLDPE-g-MA) and have been amplified when adding more LMEs. The corresponding layer structure, observed with optical microscopy, is presented Figure 14b and exhibits quite heterogenous and chaotic layer arrangements. It seems that the separation, spreading and recombination steps in the multiplying elements do not occur properly, provoking thus chaotic layer multiplication when grafted tie layer is used²⁷,²⁸. The repartition of these interfacial copolymers has been shown to be inhomogeneous at the interface and responsible for interfacial stabilities when the interface is increased, for example during the flow in a flat die²⁹. The inhomogeneity of the repartition of the interfacial copolymers (created between EVOH and the reactive tie layer) is dramatically increased by the LMEs, due to the increase of interfacial area during the flow in LME (the interfacial area is multiplied by two in each LME). This inhomogeneous repartition creates interface corrugation and is responsible for the observed defects as can be seen Figure 14.

<figure><img src="image_17.png" /><figcaption>Figure 14. a) Photos of 180 µm multilayered films with 4 layers of EVOH44 having theoretical 12 µm layer thickness (2 LME) and different types of tie layer: LLDPE-g-MA and LDPE-g-MA; b) Optical microscopy image of 180 µm (LLDPEg-EVOH38-LLDPEg)₂ cross section multilayered film with transmission mode. The black arrows indicate EVOH38 layers</figcaption></figure>

Those results show that this kind of reactive systems cannot be multiplied in LME system correctly. To overcome this limitation and retain the adhesion between the layers, a non-reactive system was employed using linear low density polyethylene (LLDPE) in place of the low density polyethylene (LDPE) previously used as the interaction between LLDPE and EVOH are higher than between LDPE and EVOH. Two sets of films were then realized: only the multiplied LLDPE/EVOH/LLDPE but also, as LDPE should remain the main polymer that will be recycled at the end, the other system developed was the following one:

$$ \text{LDPE} / (\text{LLDPE/EVOH38/LLDPE})_n / \text{LDPE} = \frac{40}{(7/6/7)/40} \text{ wt}\% $$

With LDPE used as an external layer and a multiplied core with alternated layers of LLDPE and EVOH38, with 38 being the molar proportion of ethylene in the copolymer. In Figure 1, LLDPE and EVOH are introduced in extruders A and B while LDPE is added in extruder C as the skin layer. At the end, only 6wt% of EVOH is present in the film obtained which is close to the general recommendation to maintain the polymer film recyclable³⁰. Films were coextruded with a number of LME from 0 to 6 leading to films with 3 to 129 layers in the multiplied core. The thickness of the final individual EVOH layers has been then from 3 µm to 50 nm with total thicknesses of the films around 50 µm. A higher number of LME has led for this configuration to instabilities manifested by a dramatic width reduction of the core probably due to the occurrence of slippage between layers as suggested by Zhao and Macosko³¹. Nonetheless, films with
continuous and homogeneous layers were obtained with thicknesses as low as 47 nm as can be seen on Figure 15. For (LLDPE/EVOH/LLDPE)ₙ, more LME could be used (up to 10 so 2049 layers in the end) but the total thickness couldn't be less than 180µm leading to individual layer thickness of EVOH close to 50nm without breakage of the layers so similar to the ones obtained with a LDPE skin layer and a lower number of LME.

<figure><img src="image_18.png" /><figcaption>Figure 15. AFM phase images of 60µm LDPE-(LLDPE-EVOH38-LLDPE)n-LDPE cross section films with a) 16 layers (4 LME) and b) 64 layers (6 LME) of EVOH38</figcaption></figure>

In order to investigate the evolution of the crystalline degree with the thickness of the EVOH layers, DSC measurements were performed on (LLDPE-EVOH38-LLDPE)ₙ multinanolayered films with EVOH layer thickness going from 44 µm to 50 nm and EVOH layer number from 1 to 1024. Two melting peaks in the multilayered samples can be distinguished, one corresponding to LLDPE around 123 °C and the other one to EVOH38 around 170 °C. Focusing on EVOH melting peak, it is systematically slightly lower (around 4°C) in the multilayered coextruded films, probably due to a rapid cooling at the die exit. Hence, the crystallinity of EVOH38 has been calculated for all the multinanolayered films and does not vary significantly by reducing EVOH38 thickness in the film (Figure 16). DSC was not realized on the LDPE/(LLDPE/EVOH/LLDPE)ₙ/LDPE due to the low amount of EVOH in those films leading to a poor signal. As the thicknesses are similar in both types of films, the evolution of the crystallinity of EVOH is expected to be similar as well.

<figure><img src="image_19.png" /><figcaption>Figure 16. DSC thermograms of 44 µm thickness control EVOH38, 65 µm thickness control LLDPE and 180 µm thickness (LLDPE-EVOH38-LLDPE)n multinanolayered films with different LME a). Crystallinity of EVOH38 calculated in the multinanolayered film compared to an EVOH38 control film of 44 µm thickness b). The numbers above the experimental dots are the EVOH38 layer number.</figcaption></figure>

Contrary to previous multilayered COC/EVOH system, here the LLDPE/EVOH system does not lead to a hard confinement as the T_m of LLDPE (~120°C) is lower than the one of EVOH (~170°C). No specific orientation of the crystalline structure is then expected and has not been studied. Only the oxygen permeability of the complete films has been investigated. As presented in Figure 17, the film oxygen
permeability has been measured in different conditions of relative humidity and the permeability of only the EVOH has been calculated using equation 3.

$$ \frac{1}{P_{film}} = \frac{\varphi_{EVOH}}{P_{EVOH}} + \frac{\varphi_{LLDPE}}{P_{LLDPE}} + \frac{\varphi_{LDPE}}{P_{LDPE}} \quad \text{eq. 3} $$

Since the permeabilities of LLDPE and LDPE are very low compared to EVOH38, they can be neglected, leading thus to equation 4:

$$ \frac{1}{P_{film}} = \frac{\varphi_{EVOH}}{P_{EVOH}} \quad \text{eq. 4} $$

As explained in the introduction, food packaging films may have to go through thermal treatment with highly humid conditions. It usually leads to a high loss of properties of the EVOH as moisture can penetrate the internal layer under those conditions and the water stays trapped in the EVOH layer at the end of the treatment³². This remaining humidity decreases drastically the barrier properties as those latter are linked to the creation of a hydrogen bonds network between EVOH chains and humidity will replace a part of those bonds while keeping a large mobility. To avoid this situation, the barrier properties of the system under high relative humidity conditions has to be realized to anticipate this decrease.

<figure><img src="image_20.png" /><figcaption>Figure 17. Oxygen permeability of 60 µm thick LDPE-(LLDPE-EVOH38-LLDPE)n-LDPE multinanolayered films a) and EVOH38 oxygen permeability measured in the multilayered films with equation 4 b) at 23°C, 0, 50 and 75 % RH evolution with EVOH38 layer thickness. The red, magenta and black solid lines correspond to film permeability having 1 layer of EVOH38. The dashed lines correspond to EVOH38 permeability at 0, 50 and 75 % RH provided by technical sheet at 0, 50 and 75 % RH. The numbers above the experimental dots are the EVOH38 layer number.</figcaption></figure>

Global constant values are found at 0 % RH, while a decrease of factor 2.7 and 3.3 is observed at 50 and 75 % RH respectively. Thus, no effect is observed in dry conditions as the barrier properties are stable with the number of layers and their thicknesses, as already observed by Alipour et al.³³ and Lee et al.³⁴ even if their layers didn't reach the nanometer scale. However, this seems of prime importance when the relative humidity increases. Indeed, decreasing the thickness of each layer decrease the permeability of the total film in the two RH conditions tested. This might be explained by two hypotheses. Either, the « soft » confinement induced by the LLDPE leads to a modification of the crystalline structure or of the hydrogen bonds network, or the multiplication of the interfaces allows a great decrease of the diffusion of water inside the different layer and by this decrease the effect of it on the EVOH behavior.
II.3.2. EVOH combined with another EVOH: “hard” confinement

To go further, a system based on the hard confinement of EVOH has been designed. To do so, an EVOH44 was used and confined by a different grade of EVOH, an EVOH29 with only 29 mol% of ethylene units. Indeed, the difference in melting temperature of those different grades allows a hard confinement of one by the other. T<sub>m</sub> (EVOH44) is closed to 170°C while T<sub>m</sub> (EVOH29) is around 185°C. Multinanolayered films of multiplied (EVOH29-EVOH44-EVOH29)<sub>n</sub> with the proportion 35-30-35 wt% and a total thickness of 100 µm have been manufactured with 4 (2 LME), 64 (6 LME), 256 (8 LME) and 1024 (10 LME) EVOH44 layers. Multinanolayered films with 64 layers of EVOH44 have been stretched at different draw ratios using the chill rolls during the cooling to reach thinner films. As EVOH cannot be extruded alone and to reach nanometric layers without numerous LME, 50 %wt of LDPE sacrificial layer has been used to surround the core and to avoid direct contact of EVOH with the die walls. Multinanolayered films have been successfully coextruded, without any evidence of instabilities, and coextruded films, display good transparency with EVOH44 layer thicknesses ranging from 7 µm to 30 nm. The thinnest layers have been observed by AFM which reveals the continuity of the layers, Figure 18.

<figure><img src="image_21.png" /><figcaption>Figure 18. AFM images of (EVOH29-EVOH44-EVOH29)n cross section multinanolayered films of 64 EVOH44 layers with a EVOH44 layer thickness of 455nm a) then 116 nm using a higher drawing ratio b). AFM images of (EVOH29-EVOH44-EVOH29)n multinanolayered films of 30 nm EVOH44 layer thickness using 1024 layers of EVOH44 c)</figcaption></figure>

Then the crystallinity was followed using DSC, Figure 19. It is interesting to notice that while the EVOH29,
the confining material, presents a stable crystallinity degree, the EVOH44, the confined material, displays a
decrease of its crystallinity with the decrease of the thickness of the layers from 35 % with 4 layers to 25%
with 1024 layers and an individual layer thickness of EVOH around 30 nm. This can be easily explained by
the constraint induced by the hard confining material which hinders the mobility of the confined chains
leading to a decrease of ability of the chains to organize themselves contrary to the soft confinement
previously observed where crystallinity degree of EVOH38 wasn't affected by the thickness of the layers
(Figure 16). Similar crystallinity reduction with geometrical confinement has previously been observed, for
example for HDPE confined by PS³⁵ or for PCL confined by PMMA (from 48 % to 26 %)³⁶. Very often,
this decrease in crystallinity was accompanied by a change of crystalline morphology (from a 3-D isotropic
spherulites to orientated lamellae).
<figure><img src="image_22.png" /><figcaption>Figure 19. DSC thermograms of control EVOH29, EVOH44 and multinanolayered films a). Crystallinity of EVOH29 and EVOH44, calculated by peak deconvolution, in function of EVOH thickness b). The numbers above the experimental dots are the EVOH44 layer number.</figcaption></figure>

As the objective of this confinement is to improve the gas barrier properties of the films, a decrease of the crystallinity can be considered as negative as the crystalline structures are impermeable and a higher crystallinity degree usually induces a lower permeability. So to further establish the effect of layer multiplication combined with a hard confinement, oxygen permeability was realized, in dry conditions and at 75% of RH as well, on those films, Figure 20.

<figure><img src="image_23.png" /><figcaption>Figure 20. Oxygen permeability of (EVOH29-EVOH44-EVOH29)ₙ multinanolayered film versus EVOH44 thickness. The red and black solid lines correspond to the theoretical permeabilities of 3 layers film having one layer of EVOH44 at 0 and 75 %RH. The numbers above the experimental dots are the EVOH44 layer number.</figcaption></figure>

As previously observed by other authors for micronic or submicronic layers³³,³⁴, no significant improvement of the gas barrier properties is observed in dry conditions with 1024 layers compared to only 4 layers even if the crystallinity degree has decreased. More interestingly, the permeability observed in humid conditions is decreasing with the decrease in the layer nominal size, ie the number of layers, and reaches an even lower level than the classical three-layer film in dry conditions. This decrease may again be explained by an antiplastification effect due to the strong linkage of the water to OH functions of the EVOH through H-bonding leading to a decrease in the free volume available for the gas permeability. We may hypothesize that the decrease in layer thickness may affect the possible absorption of water even in equilibrium state, which is the state where the films are for this measurement. This phenomenon may be linked to the confinement of the EVOH chains as even if a slight decrease is observed in permeability between 7 µm and
0.1 µm for the thickness of each EVOH layer, the more significant drop is observed when this thickness is around 30 nm. This drop is more significant than what has been observed for LLDPE/EVOH38/LLDPE films which were based on a “soft”confinement while here, the EVOH29/EVOH44/EVOH29 combination allows a “hard” confinement of the EVOH44 and this non-linearity of the drop is resembling the results reported by Baer’s team¹⁷.

To confirm the effect of the confinement and so of the layer thickness on the permeability, similar films were realized with a thick external layer made of LDPE+LDPE-g-MA in order to reach nanometric layers with less multiplying elements and so less final total number of layers, Figure 21.

<figure><img src="image_24.png" /><figcaption>**Figure 21.** Oxygen permeability of the total EVOH29-EVOH44 present in the two multinolayered films (LDPE+LDPEg)-(EVOH29-EVOH44-EVOH29)n-(LDPE+LDPEg) named “skin” and (EVOH29-EVOH44-EVOH29)n named “core” in function of EVOH44 thickness. The numbers above the experimental dots are the EVOH44 layer number.</figcaption></figure>

Figure 21 shows that the permeability values obtained depend only on the layer thickness, irrespectively of the number of interfaces/layers. This tends to confirm the fact that the confinement effect has a predominant effect on the permeability.

In order to understand this phenomenon, water uptake measurements has been realized and a slight increase of the water uptake has been observed for the different multi-layered systems compared to pure EVOH but without any incidence of the layer thickness or number of interfaces which could then not explained the evolution of the permeability in presence of water. Similarly, the evolution of the crystallinity at 75% RH has been studied and concluded that the crystallinity slightly decreased with the addition of moisture and with the increase in number of layers/interfaces, which then could not explain the evolution of the gas barrier properties.

Finally, DMTA measurements were carried out to understand if the water has a classical plasticizer effect or on the contrary an antiplasticizer one, Figure 22.
<figure><img src="image_25.png" /><figcaption>Figure 22. DMTA curves of (EVOH29-EVOH44-EVOH29)ₙ multinolayered film realized at 0% and 75 % RH.</figcaption></figure>

Figure 22 highlights an expected decrease around 30 °C of α transition related to the Tg for each samples due to plasticization phenomenon. The additional transitions α' corresponding to molecular orientation³⁷ or segmental motion close to lamella surface and β related to segment motion within the interfacial region are also decreased (of around 10 °C). Also, the initial storage modulus is not affected by the relative humidity whatever the samples. However, no differences have been noticed between the two multi-layered samples. The plasticizer effect of water is then confirmed for all the different films without any link to the number of layers. To understand the permeability behaviour in humid conditions, the remaining hypothesis is the hydrogen bond network which is known to play a significant role in EVOH gas barrier properties. Reducing EVOH layer thickness in multilayered film could possibly modify this network, for example reinforce it, preventing water molecules from destroying it and forcing them to fill the free volume. However, uptil now, no clear proof could be brought to support this hypothesis.

Finally, we may conclude that the use of multinolayer coextrusion process allowed the achievement of increase barrier properties in EVOH when this EVOH is hardly confined in nanometers thick layers. This increase could not be observed in dry conditions and by then that water is highly involved in this evolution. The most probable hypothesis is a reinforcement of the hydrogen bond network by interaction with the water molecules.

As food packaging might have to go through a pasteurisation process that involves water and high temperature and which leads to a non-neglectable amount of water emprisonned in the EVOH layers, this improvement in humid conditions may help improving the overall gas barrier properties of this polymer and then decreasing the thickness required in a food packaging films to prevent the food from spoliation.
This would lead to more recyclable films with less “impurities” that may affect the recycling of the main
polymer of the films, PE or PP.

II.4. Conclusion

To understand the behaviour of EVOH, its cristallization habits have been investigated in bulk and under confinement (using spin-coating techniques as well as multinanolayer coextrusion with cyclic olefin copolymer to induce a hard confinement). These analyses allowed us to display a spherulitic shape of the crystalline phase, to show the interest of using Flash DSC to study the kinetics of crystallization, the possibility of having in plane and orthogonal lamellae and finally a potential orientation of those lamellae in multinanolayer films according to the thickness of the layers under a confinement due to cyclic olefin copolymer.

In another hand, different multilayer films have been successfully produced with different EVOH grades
and LLDPE while emphasizing the fact that a chemical compatibilizer can not be used in LME coextrusion
system. The effect of the thickness and number of the layers have been investigated on EVOH crystallinity
degree and on barrier properties of the whole films. Two types of confinement have been used : a soft one
with LLDPE and a hard one by a smart combination of two different grades of EVOH possessing different
melting temperature. While no effect of the confinement could be detected on the barrier properties in dry
conditions, a decrease by a factor close to 10 has been observed for the hard confinement system and only
a factor of 3 with a soft one with the decrease of the thickness of the EVOH layers from 10 µm to 30 nm
when the permeability test was run under a 75% relative humidity. This improvement might be of great
interest to improve the protection of the EVOH layers regarding industrial processes such as pasteurisation
or retort that involves high relative humidity and high temperature.
### III. High gas barrier coatings

The second path to attain high gas barrier properties for food packaging films is to use coatings with polymers that present very low gas permeability. This property is usually linked to hydrogen bonding network as discussed with EVOH layers. The problem with polymers presenting high hydrogen bond networks is the fact that those polymers are usually complicated to process, due to this network. Indeed, the degradation temperature of those polymers are usually close or higher than their processing temperature. So, the melt-processing of those polymers are either complicated or impossible³⁸. The coextrusion with PE or PP can then not be envisaged. To be able to process those polymers, the softening and modelling of it will be realized with the use of a solvent, usually water, and the shape will be realized by coating and evaporation of this solvent. This is the way that have been attempted in this part with the use of different types of polymers. The aim here is to deposit thin layers (around 2 to 3 µm) of those high gas barrier polymers on top of a recyclable film (PET in a first approach). The low thickness of the coating will allow this coating to be considered neglectable in the end-of-life industry compared to the PET substrate.

This method is already employed in the food packaging industry but with a polymer that presents an hazardous end-of-life, the poly(vinylidene dichloride) (PVDC). This polymer possesses in its backbone two chlorines which are problematic in end-of-life treatments and especially in waste incineration where toxic products may be formed³⁹. The industry tries then to get rid of this polymer. However, the interesting part of PVDC is the fact that this polymer exhibits high gas barrier properties in dry and humid conditions. Indeed, those properties are here linked to strong van der Waals interactions (Keesom and Debye one predominantly) due to the fact that the two chlorines induce strong dipoles in the backbone of the polymer. Therefore strong interactions are created between PVDC chains without the involvement of hydrogen bonds which are strongly affected by water presence. To be able to replace PVDC, the coatings will then need to present good gas barrier properties in dry and humid conditions as well. This is of prime importance, as contrary to the multilayer films, the coatings are not protected from the moisture as they are on the surface of the film, directly in contact with the environment or the food.

Different polymers were then envisaged for this application. As high gas barrier properties are expected, polymers with strong hydrogen bond network were chosen, Figure 23. The followed polymers were then tested:

- Poly(vinyl alcohol) (PVOH)

- Chitosan (CS)⁴⁰

- Fish gelatin (FG)

<figure><img src="image_26.png" /><figcaption>Figure 23. PVOH and Chitosan chemical structure</figcaption></figure>

Finally, another type of polymer was attempted that present lower hydrophilicity: a phenoxy resin (PKHW) which is use as a dispersion (latex) in water. This latter was used to expect better properties in humid conditions.

Also in order to improve drastically the gas barrier properties, the addition of nanofillers in those polymeric matrix coatings was attempted. Indeed, the use of non-permeable particles may lower significantly the overall permeability by inducing a strong tortuosity in the coating. Different types of nanofillers have also been attempted. The different combinations of polymer matrix/nanofillers under dry and humid conditions have been realized and will be explained in the following parts.
Nanofillers are particles with at least one dimension in the nanoscale range. These fillers have the particularity of possessing a high specific area, yielding more matrix/filler interface. They can be categorized by their morphologies and shapes (disks, spheres, cylinders, see Figure 24), or by their nature and composition (mineral, organic, etc...).

<figure><img src="image_27.png" /><figcaption>**Figure 24.** Different filler morphologies (a) cylinder (b) spheres (c) platelet/disk morphology.</figcaption></figure>

Amongst the wide range of nanoparticles, the following groups can be found:

- Petrochemical nanofillers: classical carbon derived nanofillers such as black carbon are used as a reinforcing filler in rubber products. Carbon nanotubes (CNT), double-wall and multi-wall carbon nanotube (DW and MW-CNT) also present outstanding properties (very high Young modulus up to 1 TPa, heat conductivity of 3000 W.m⁻¹.K⁻¹, electric conductivity of 107 S.m⁻¹)⁴¹. Graphene is a bi-dimensional nanofiller (the theoretical thickness is below 1 nm), and the fundamental unit of CNT and graphite. Its high specific surface area, strength and stiffness (130 GPa and 1 TPa) generate great enthusiasm, however, its high cost is a limiting factor⁴².

- Mineral nanofillers: nanoclays are the most common mineral nanoparticles due to their abundance in nature and their favorable aspect ratio (up to 100). They can be separated into two big families: phyllosilicates (which includes kaolinite, smectite, talc, mica, montmorillonite, and saponite) and polysilicate, both natural and synthetic. Clays are known for their layered structure linked by ionic bonds, which makes them quite sensitive and easily penetrated by water. Among them, vermiculites are phyllosilicate minerals that are composed of hydrated sheet silicates, which contain layers of water molecules within their internal structure⁴³.

- Polysaccharides: nanoparticles can be made from polysaccharides such as alginate⁴⁴, chitosan⁴⁵, and pectin nanoparticles⁴⁶, chitin nanowhiskers⁴⁷, cellulose nanocrystals or nanowhiskers (CNC or CNW) and cellulose nanofibrils (CNF). The properties of these particular NPs are directly linked to their extraction method, and the chemical used in the process⁴⁸. The commercialization of polysaccharide nanoparticles is still relatively recent, but it is quickly expanding into a growing and promising industry.

Typical inclusion rates of nanofillers vary from 0.1 to 10 wt%. Due to the small size of nanoparticles, the surface area to volume ratio is small when compared to the same volume of material made up of bigger particles. However, for tortuosity impact, higher content value may be envisaged to increase drastically the diffusion path. Also, as the processing method used for coatings are in solution, the large increase of viscosity caused by the nanofillers is not a blocking issue. Indeed, when thermic processing such as extrusion is involved, the evolution of viscosity is usually problematic and leads to a limit in nanofillers incorporation. In solution, very high content such as 50 wt% or even pure nanofillers can be attempted to see the optimized value of nanofillers required. Different content will be then be attempted and the quality of the dispersion will be studied as well as the orientation of the nanofillers, as different patterns can be observed, as described on Figure 25.
<figure><img src="image_28.png" /><figcaption>**Figure 25.** Basic patterns of dispersion and distribution in nanocomposites (a) poor dispersion and bad distribution, (b) poor dispersion and good distribution, (c) poor dispersion and good distribution with orientation, (d) good dispersion and bad distribution, (e) good dispersion and distribution and (f) oriented fillers with both good dispersion and distribution</figcaption></figure>

In this context, nanofillers have been selected due to their availability, shape and biodegradability: vermiculites. Indeed, vermiculites have a platelet shape and will help obtaining a high diffusion path through tortuosity.

The combination of those polymer matrixes with these nanofillers was then attempted through a coating process approach and their gas permeabilities were studied.

### III.1. Polymer matrixes coatings

In a first part, the polymeric matrixes have been tested in terms of coating processing and gas barrier properties.

#### III.1.1. Coating processing

First, solutions were prepared for the different polymer matrixes. The solvent used were either pure water (PVOH, gelatin, PKHW dispersion) or a acidified aqueous solution (1vol% of acetic acid in water) for CS. Then the coating itself has been realized by bar-coating using bar-coaters of different thicknesses with a manual set-up or an automatic one, Figure 26.

<figure><img src="image_29.png" /><figcaption>**Figure 26.** Coating apparatus, manual on the left and automatic on the right</figcaption></figure>
Prior to the deposition, a corona treatment has been applied on PET substrate in order to improve the wettability of the film to obtain a homogeneous deposition. Then the solutions are coated on the substrate and dried using an infra-red radiation or in an oven (45°C) under vacuum. The final theoretical thicknesses attained were in the range of 1 to 5 µm and were characterized by scanning electron microscopy (SEM) or transmission electron microscopy (TEM). The main issue of these aqueous coatings is the drying step. Indeed, to obtain a perfectly dried film, the thinnest solution films must be realized to allow the water molecules to be released from the polymer matrix. As those matrixes are highly hydrophilic, the water removal is complicated and a large amount of it may be kept imprisoned in the matrix after the drying of only the surface of the film. To avoid this phenomenon, it was decided to coat thin films and to repeat the coating and drying steps several times to obtain the final expected thickness. SEM images of CS coatings can be seen on Figure 27 to show the homogeneity of the final coatings. No separation between the successive coatings can be observed indicating that probably the surface of the previous layer is dissolved by the following coating solution leading to a perfectly homogeneous coating. The homogeneity of thickness was also investigated and with the increase of number of layers, the thickness becomes more homogeneous. This might be explained by the fact that successive coatings may make up for the possible holes created during the process. So heterogeneities are compensated after each step, so the more steps there is, the less inhomogeneities are observed.

<figure><img src="image_30.png" /><figcaption>Figure 27. a) SEM cross-section micrographs thickness pictures of the PET films with coated layers: 1) one layer chitosan, 2) two layers chitosan and 3) three layers chitosan b) Average thickness of coated layers measured by SEM.</figcaption></figure>
### III.1.2. Helium permeability of pure polymer coatings

Helium permeabilities of the different types of coatings attempted are summarized in Figure 28. One can observe that even if the PVDC presents very high gas barrier properties, the use of chitosan allowed improving this property.

<figure><img src="image_31.png" /><figcaption>Figure 28. Helium permeabilities of coatings made of PVDC, PVOH, Fish gelatin, PKHW and Chitosan with 1, 2 and 3 successive coatings</figcaption></figure>

The different helium permeabilities obtained using the different matrixes are in the range of the PVDC even the PKHW one. The best results obtained were with chitosan but the drying step is different for this polymer. Indeed, in that case, a drying with an oven under vacuum was realized which would lead to a better drying of the coating. We may then suppose that PVOH and gelatin coatings could be improved by improving the drying step which is not necessary for PVDC coating which is less hydrophilic. Nonetheless with an optimized drying of the coating, higher barrier properties of a factor close to 10 can be achieved with at least chitosan. The sequential coatings of different layers of chitosan would also lead to similar value of permeabilities of the coatings. A slight increase is observed with the increase of number of coating steps. This might also be linked to the fact that by increasing the final thickness of the coating, the drying might lose its efficiency and slight amount of moisture may stay imprisoned in the coating, decreasing its barrier ability.

## III.2. Addition of nanofillers towards high gas barrier coatings

Intercalation of clay in several polymers has been introduced as a promising route, leading to versatile polymer composites, especially those containing a nanoscale silicate of high aspect ratio. Because of the small size of the structural unit and the high surface-to-volume ratio, the properties of nanocomposites are greatly influenced by the degree of mixing between the two phases⁴⁹,⁵⁰. In this work, we build a better barrier by adding vermiculite (VMT) to the different matrixes with various content ranging from 5 to 90 wt%. The
innovative part is to add very high amount compared to what is usually done in the literature. This is helped
by the fact that the processing is realized in solution and not by melt-processing (i.e. extrusion). Thus, VMT
has been dispersed in water and mixed with the different solutions of polymer matrixes. The homogeneity
of the dispersion was insured by mechanical stirring with an Ultraturax apparatus at 5000 rounds/min. The
final dispersion was characterized by Dynamic Light Scattering and allow the observation of 1µm particles
which is coherent with a size of a single platelet of vermiculite. Then those solutions were coated and dried
similarly to previously presented.

### III.2.1. VMT dispersion in the coatings

Dispersion of the VMT in the final coated layer after the drying step was investigated by Transmission Electron Microscopy (TEM), see Figure 29 for the mix fish gelatin with 20wt% of VMT. It can be observed that the dispersion of the VMT was not maintained through the drying step as aggregation of VMT are visible in Figure 29. Nevertheless, particles are overall oriented parallel to the surface of the PET and the aggregates form what can be assimilated as a VMT layer contained in the matrix layer. Individual VMT particles observed present a lateral dimension of 100 to 200 nm and some isolated particles are also visible in the images. The global organization of the VMT might result from a combination of the drying step with the viscosity of the solution and the affinity of the VMT with the matrix.

<figure><img src="image_32.png" /><figcaption>Figure 29. Representative TEM images of Gelatin/VMT (20wt%) coatings. a) the entire PET film with the coating can be observed, b) a zoom on the coated layer. The circular patterns observed on the images arise from the microscopy grid.</figcaption></figure>

Figure 30 presents the images of PET films coated with two layers of CS-VMT with various rates of VMT. The transparency of edible/bio-based polymer films is a key parameter to good film acceptance by users⁵¹. First of all, visually all the showed films were transparent, for either uncoated PET or PET coated with CS (Figure 30 a. b). Also, it can be noticed that nanocomposites films had a translucent surface (Figure 30 c.d. e). The brightness of all films remained fairly constant, for VMT percentage varying from 0 to 25%. However, the presence of 40, 50, and 60 wt% of VMT in the coated films increases significantly the yellow tint. To conclude, all the films were visually acceptable.
<figure><img src="image_33.png" /><figcaption>Figure 30. Images of films coated with two layers : a) uncoated PET b) PET coated with CS c) PET coated with CS-25%VMT d) PET coated with CS-40%VMT e) PET coated with CS-50%VMT</figcaption></figure>

### III.2.2. Helium permeability of VMT nanocomposite coatings

Then, the helium permeability of coated PET films were characterized. The results of measurements of the helium permeance of the PET films according to the coated layers with CS-VMT with different VMT content (25, 40 and 50 wt%) are represented in Figure 31. The coated films with CS-VMT are significantly less permeable than the uncoated PET film and the permeability decreases with the increase of VMT content. Thus the presence of VMT in the chitosan layers decreases significantly the helium permeance of coated PET films. Firstly, a decrease in the permeance of the nanocomposite coated films is observed when the number of coated layers increases. Also, it is accompanied with a decrease of permeance when the composition of the nanocomposite is increased in VMT. A improvement of a factor of around 100 was obtained when adding 50wt% of vermiculite compared to pure CS. It should be noted that even higher effect on the gas barrier properties could be expected, however, the limit of detection of the apparatus is attained and the evolution cannot be detected as seen in Figure 31, as well as the permeance becomes nearly stable from two to three layers on nanocomposite from 50% VMT fillers. From a kinetic point of view, the gas permeation is much faster for a layer and it results in a higher gas permeance. This is explained by the decrease in the permeance of the CS/VMT when the thickness of the coated layer increases. These results lead to calculate the helium permeability of coated layers (similarly to eq. 3). Helium permeability of two layers of nanocomposite versus the VMT fraction of the nanocomposite is presented in Figure 31 b. In the present work the helium permeability decreases after adding VMT fraction to the CS, also when the composition of the nanocomposite increased on VMT. It can be noted that the helium permeability of two layers of nanocomposite decreases from 24.4×10⁻²⁰ (m³.m/m².s.Pa) up to 1.1×10⁻²⁰ (m³.m/m².s.Pa) by only adding 40% of vermiculite to chitosan in dry conditions. Even higher effect on the gas barrier properties could be expected for 50 and 60 wt% vermiculite, as said previously.
a)

<figure><img src="image_34.png" /></figure>

b)

<figure><img src="image_35.png" /><figcaption>Figure 31. Helium permeance of a 130 µm PET film coated with two layers of CS-VMT (a) and helium permeability of two coated layers of CS-VMT (b) both in dry condition (23 °C and 0% RH).</figcaption></figure>

Similarly for the other matrixes, the addition of even a slight amount of VMT (~10 wt%) allows a significant decrease in helium permeance (factor 1000 for PVOH) of coated PET films, see Figure 32. However for PVOH, after 20wt%, no improvement could be observed and even a loss of barrier properties can be observed above 50wt% probably due to the fact that at those contents, the adhesion between PET and the coated layer is not correct anymore and also the brittleness of the layer increases dramatically leading to crazes in it very easily. However, as can be seen in Figure 32, the PVOH allows the best properties compared to fish gelatin and PKHW. As expected, the less polar nature of PKHW limits its ability to achieve high gas barrier properties and indicates that its affinity is probably less efficient than with PVOH or gelatin. However, as explained previously, these systems will have to endure humid environment during their lifetime and the evolution of the gas barrier properties have to be investigated in relation with the relative humidity.
<figure><img src="image_36.png" /><figcaption>**Figure 32.** Helium permeance of 12µm PET film coated with PVOH, PKHW, Fish Gelatin, and PVDC with a content of VMT varying from 0 to 90wt%</figcaption></figure>

This has been carried out on PVOH and Gelatin coated on PET with a load of 20wt% of VMT as it was observed as an optimum content. The transmission rate of oxygen in relation with relative humidity can be seen in Figure 33. As expected, the permeability of those very polar coatings are highly sensitive to moisture and a large loss of property is observed starting at 60 %RH. PVOH which is more polar than gelatin presents an even higher loss with a transmission rate similar to pure PET at 75%RH while gelatin remained slightly lower. Regarding those results and in order to develop efficient food packaging, a protection towards humidity is required for those polar polymers which can be realized through a lamination of a PE film on top of the coating.

<figure><img src="image_37.png" /><figcaption>**Figure 33.** Oxygen transmission rate for 12µm PET films coated with PVOH or Gelatin with 20wt% of VMT as a function of relative humidity at 23°C</figcaption></figure>
### III.3. Chemical modification of chitosan for humid gas permeability improvement

In order to improve the barrier properties in humid conditions of polar polymers, a modification of the hydrophilicity of the CS was attempted. A balance between a large hydrogen bond network needed for good barrier properties and a low hydrophilicity to limit the effect of water on those properties is then required. To do so, different paths may be followed and during Fatima Essabi's PhD⁵², the grafting of a fatty acid on some amine functions of the CS has been attempted. Indeed, the chemical nature of the CS allows a simple chemical modification with an amidification step. A chemical modification from the literature⁵³ was followed with an expected reaction of 20mol% of the amine function into amide with palmitic acid, as can be seen in Figure 34. This acid was chosen along its biosourcing and intermediate hydrophobicity. Indeed a longer chain would have prevented the solubilization of this acid in the CS solution. The ratio of 20mol% was also chosen in order to keep a majority of amine functions to maintain a solid hydrogen-bonded network in the amorphous phase of the CS to avoid gas permeability.

<figure><img src="image_38.png" /><figcaption>**Figure 34. Synthesis of palmitic acid-grafted chitosan**</figcaption></figure>

The solution was then dialyzed using dialysis membrane against distilled ethanol first and then against water with successive exchange of fresh solvent to remove unreacted palmitic acid and the catalyst: 1-ethyl-3-(3-dimethylaminopropyl)carbodiimide (EDC).

In order to verify the efficiency of the grafting, the material was analyzed using different devices. First, FTIR analysis was conducted on CS, palmitic acid-grafted-CS (PA-g-CS) and a blend of palmitic acid (PA) and CS, Figure 35. The signals linked to PA can be seen in the range of 1600-1300 cm⁻¹ with especially three peaks at 1558, 1375 and 1311 cm⁻¹. Those peaks are observed for PA-g-CS and PA/CS blend as expected even after the dialysis step for PA-g-CS where unreacted PA has been removed from the solution. A peak linked to the amide function is also observed at 1652 cm⁻¹. The peak is observed with all samples as amide function is present in the CS itself due to a non-complete deacetylation of the initial chitin. However, this peak is more intense for the PA-g-CS which means the grafting of PA by creation of amide functions. In order to confirm these results, a thermal analysis was realized. The thermogram of palmitic acid shows a melting peak at 65°C. In the case of PA-g-CS, the presence of PA was confirmed by FTIR, but in DSC analysis, no melting peak was observed at 65°C. On the other hand, this peak is present in the case of the mixture between CS and PA (PA/CS). If this peak has disappeared after grafting on the CS, this implies that its crystallization is prevented due to the grafting. This can be explained by the fact that, being grafted, the palmitic acid chains cannot be organized and therefore cannot crystallize. This change in crystallization behavior tends to confirm that residual palmitic acid observed in FTIR is well grafted on chitosan chains. Finally this grafting was confirmed by solid state ¹³C NMR analysis and the presence of palmitic acid peaks were observed. Unfortunately, no shift of the peak corresponding to the amine function that becomes an amide function could be observed. However, since the CS was washed after the grafting reaction effectively eliminating any residual presence of unreacted palmitic acid, the presence of palmitic acid carbons (after grafting) on solid phase ¹³C RMN spectra shows that the reaction did occur.
<figure><img src="image_39.png" /><figcaption>Figure 35. FITR spectra of CS, PA-g-CS and PA/CS mixture</figcaption></figure>

PA-g-CS and PA-g-CS/VMT coatings were prepared in a similar manner to ungrafted chitosan coatings to observe the effect of grafting on coating properties. Two VMT rates were chosen: 25 and 60%. The 25% will achieve a reasonable value, in particular in terms of mechanical properties and is a good compromise barrier properties / general quality of the coating. It can be noted that PET film coated with PA-g-CS has
a somewhat translucent surface compared to the film PET coated CS. In addition, yellow tint increases with
increasing VMT rate as previously observed for the ungrafted. Overall, all films are visually transparent.
These results are consistent with the results of CS/VMT coated PET films.

Similarly to CS without grafting, the thickness and homogeneity of the layers were analyzed by SEM imaging
and the results were in good accordance with the CS ones. However, it has been observed that coated layers
of coated CS-g-PA layers are thicker than ungrafted CS layers for the same rates of nano-fillers. The
thickness of the CS-g-PA layer is in the order of 2.5 µm. However, the thickness of the CS layer is in the
order of 2.1 µm or an increase of 0.4 µm with the grafted one. After the addition of VMT, layer thickness
increases to 3.01 and 2.9 µm for 25 and 60 wt% of VMT, respectively. This may be due to the bar coater
method (layers have been manually spread) but also to the difference in density between pure chitosan
(1.411 g/cm³) with very high cohesion via hydrogen bonds and grafted chitosan (1.317 g/cm³) where grafts
push chains back from each other. Since the coated layers with grafted chitosan have the same qualities as
the pure chitosan and thicknesses are also comparable, the permeability studies can be conducted in a similar
manner. First, it is important to know the modification of barrier behavior in the dry state. Indeed, the
addition of grafts on the amine functions of chitosan may reduce the amount of hydrogen bonds present in
the material and these are responsible for barrier properties. It is therefore necessary that these properties
remain not too affected by grafting. Figure 36 shows a comparison between the results of the oxygen
transmission rate (OTR) of the PET film coated with two layers of CS, CS/VMT, PA-g-CS and PA-g-
CS/VMT in dry state (23°C and 0% RH). According to this figure, the OTR of the grafted and non-grafted
chitosan is perfectly superimposed, regardless of the rate of nano-loads added. The latter decreases when
the VMT is added to the PA-g-CS. In addition, a PET film coated with 60 wt% VMT, where the thickness
of the coated layer is of the order of 2.9 µm, exposes an OTR below the commercial instrumentation
detection limit (0.008 cm³/m².day). It can be concluded that these results are similar to those obtained with
non-grafted chitosan and the properties are not degraded by the presence of grafting even if the thicknesses
deposited with the grafted CS are slightly higher than those of ungrafted CS. Therefore the chemical
modification of the chitosan seems to not modify the barrier properties at least at this degree of grafting. It
can be noted that the oxygen permeability of the two nanocomposite layers decreases after the addition of
25wt% VMT, this layer exposes a permeability of 6,18.10⁻²⁴ m³.m/m².s.Pa with a thickness of 3,01 µm.
Although a higher effect on the gas barrier properties could be expected when 60 wt% of VMT was added,
the detection limit of the device is reached and the evolution cannot be detected as shown in Figure 36. We
can conclude that the loss of barrier properties due to grafting does not affect the general behavior.

<figure><img src="image_40.png" /><figcaption>Figure 36. Oxygen permeability of two coated layers of CS/VMT and CS-g-PA/VMT with various VMT content both in dry conditions (23°C and 0% RH)</figcaption></figure>
Finally, the effect of relative humidity was observed on those coatings. This was realized using Helium permeance while varying the relative humidity of the gas. In Figure 37.a) at 0% RH, a decrease in helium permeance is observed between pristine chitosan and those with grafted chitosan, regardless of the amount of VMT. This result is quite surprising in comparison to what was observed in OTR where the barrier properties are similar for both cases in Figure 36. The impact of layer thickness may be a factor explaining this difference since we have a greater thickness of grafted chitosan than pristine chitosan. The behaviour at 54% RH follows the same trend between grafted and ungrafted with stable properties compared to dry conditions, as already observed. From 74% RH, the tendency of the grafted chitosan is the same as that of the ungrafted, namely, an increase in the permeance of the film. However, in the films produced, there are still lower permeance values with palmitic acid grafted chitosan coating. Finally, at 98% RH, the permeability is strongly increased for all cases. However, the grafted chitosan coating containing 60 wt% VMT still has a lower value than simple PET. An improvement of a factor 2 is observed in these extreme conditions. In order to eliminate the fact that the layers do not all have the same thickness, the permeability of the coated layers was calculated using measured thicknesses and permeance values. It is noted that the permeability of the grafted chitosan layers (Figure 37 b) remains better than the one of pristine chitosan, regardless of conditions. However, the difference in permeability is all the more clear with the incorporation of VMT into the coating. The combined effect of grafting and addition of vermiculite would appear to limit the impact of moisture on the barrier properties of the material up to 74% RH only. The improvement effect of grafting is in any case more noticeable in the wettest conditions. This is logical because the hydrophilicity remains slightly modified by the addition of the grafts and its effect will be more visible to the highest relative humidity. Prepared films based on grafted CS therefore have better results than those of CS-based films in wet conditions, as hoped. Thus, this appears to be a synergistic effect of grafting and VMT for the helium barrier behaviour of chitosan.

<figure><img src="image_41.png" /><figcaption>**Figure 37.** Helium permeance (a) and permeability (b) of two coated layers of CS, CS-g-PA, CS/VMT and CS-g-PA/VMT at different humidity rates (0, 54, 74 and 98 %RH) with different VMT content (0, 25, 50 and 60wt%) (VA: approximate values since the detection limit was exceeded)</figcaption></figure>
### III.4. Conclusion

Different nature of coatings have been attempted and their permeabilities have been measured with helium or oxygen permeance. The addition of nanofillers was also realized in order to improve drastically the barrier properties with high amount of fillers (up to 60 wt%). The effect of the chemical nature of the polymer on the moisture absorption was studied as well as its effect on the gas permeability. The thickness of the coated layers was successfully determined by SEM (from 1 to 4 µm). Helium and oxygen permeance of PET films coated with CS were measured with a drop of two decades compared to PET. After adding the VMT to the CS the helium and oxygen permeances decrease when the composition of the nanocomposite increased on VMT. An improvement factor of about 100 in helium and more than 10 in oxygen were obtained when adding 50% of VMT which make those coatings more efficient than PVDC ones in dry conditions. Also, the permeability of coated CS and CS/VMT layers remains stable up to 54% RH and from 74% RH, the layers gradually lose their barrier properties regardless of the nano-loads rate. At 98% RH, all barrier properties of coatings are almost completely lost and the permeance of the film corresponds closely to that of PET alone. To improve this critical point from an application point of view, the grafting of a fatty acid, palmitic acid, on the CS chain was successfully carried out. This grafting was confirmed by the appearance of the amide peak in the infrared spectrum as well as by the DSC thermogram of PA-g-CS, which shows no peak corresponding to the melting of the palmitic acid, which is therefore prevented from crystallizing. Finally, ¹³C NMR solid phase analysis confirmed the presence of peaks corresponding to the methyl groups of palmitic acid (CH₂) in the PA-g-CS NMR spectrum. In addition, solubility tests have also proven that grafted CS is not soluble in organic solvents or in an acidified aqueous solution, which is a good solvent of pristine CS. The grafting of palmitic acid to the chitosan chain improved the barrier properties of the CS and CS/VMT film at all RH conditions for helium permeability. At 98% RH, an improvement factor of 2 is obtained compared to pure PET with a PA-g-CS coating and VMT at 60 wt%. At the same humidity content, the helium permeability value of the same sample (PA-g-CS/60%VMT) is measured at 1.5x10⁻¹⁹ m³/m².Pa with a layer thickness of 2.9 µm.
# IV. Mechanical properties improvement by blends of polymers

Along with the gas barrier properties, I tried to tackle other environmental impacts such as the loss of mechanical properties of recycled material or the ones of biobased materials in order to compete with petrobased ones. To study these issues, I have chosen to produce blends of polymers and control their structuration at the nano or micro-scale. These studies are summarized in two parts, based on recycled thermoplastics emanating from Waste of Electrical and Electronic Equipment (WEEE) and on improvement of biobased materials: poly(lactide) (PLA)/poly(3-hydroxybutyrate-co-3-hydroxyvalerate) (PHBV) blends.

## IV.1. Upcycling of plastics from WEEE

Plastics waste issued from WEEE has dramatically increased for the last several years, triggered by a constant need for technological innovation and a fast substitution of both telecommunication and information technology (IT) materials. Considering that less than 30% of the plastic waste generated in Europe every year is collected for recycling⁵⁴, the management of plastic waste has become a major environmental and societal concern.

Mechanical recycling of plastics refers to their re-processing into new plastic products and is considered as the most economically viable and eco-friendly route for recycling plastics waste⁵⁵,⁵⁶. However, it is limited, especially for WEEE, by the strong heterogeneity of the waste (presence of a wide variety of incompatible polymers, of additives, degradation products, and contaminants) and the degradation of polymers both during reprocessing and the product lifetime. The resultant poor properties of such recycled plastics make them highly uncompetitive with virgin polymers ⁵⁷. There is therefore a constant need for new strategies to upgrade the properties of recycled polymers and make them suitable for new applications. A promising strategy, investigated in this project consists in turning their quality defects into assets, by controlling and structuring the blend morphology from micro to nanoscale to enhance final properties. The final blend morphology (size, shape and spatial distribution of one phase into the other) depends not only on polymer parameters (elasticity, viscosity, chemical affinity between the phases) but also on processing technologies and conditions (temperature, time and intensity of mixing, nature of the flow). Considering this interplay of polymer parameters and processing conditions, the control of the phase morphology in polymer blends during the melt flow in the process represents a real scientific and technological challenge and has been the purpose of many studies⁵⁸. By a fine-tuning of processing and material parameters, elongated morphologies have been obtained and stabilized, leading to improved mechanical or gas barrier properties⁵⁹,⁶⁰. More recently, new developments have led to a control of the structuration down to the nanoscale, for example through reactive blending⁶¹,⁶² or using innovative structuring processes⁶³. Such nanostructured polymer blends exhibit unique improved properties (mechanical, gas barrier, optical, etc.) due to increased interfaces and confinement effects⁶⁴.

We proposed to use the multilayer coextrusion device, also known as Layer Multiplier Element system (LMEs), to achieve micro and nano multilayers blend composites. However, despite the fact that the LME system is currently available in the plastic processing industry, it is rather sparsely used. Indeed, this system has been extensively studied for the barrier properties of polymer, especially for food packaging. But the mechanical properties are less studied even if they are greatly impacted. A recent study on PET/PEF multilayered films has shown for example a great improvement in puncture energy⁶⁵.
IV.1.1. Control of the morphology of a ABS/PC blend

Blends of immiscible polymers exhibit different morphologies (from drop, laminar, fibrillar to co-
continuous blend systems) linked to the thermal and mechanical processes endured. The control of the
morphology obtained can induce the access of different levels of properties. A terpolymer of acrylonitrile,
butadiene and styrene (ABS) is commonly used in blends with other polymers coming from the WEEE or
other sources like polycarbonate (PC). The mechanical behavior of a PC + ABS mixture has been extensively
studied with great improvement of impact strength without losing the interesting properties of each polymer
(high processing capacity of ABS with excellent mechanical properties and the high heat and shock
resistance of PC). However, the control of the morphology may lead to an even better increase. The
objective is then to understand and control the creation of morphologies in blends of polymers. The classical
way of studying the morphologies developed during the extrusion process are linked to the behavior of the
pellets in the device. Indeed, the pellets first endure high shearing stress while softening which will lead to
the formation of sheets of polymers emanating from the pellets. This sheeting mechanism has been
observed in single or twin-screw extruder⁶⁶,⁶⁷. The film formed will then be teared up and turned into
cylinders and fibrils and ultimately into nodules⁶⁰, Figure 38. The intermediate morphologies are highly
unstable during the process⁶⁸, especially the elongated ones like fibrillar or lamellar.

<figure><img src="image_42.png" /><figcaption>Figure 38. Morphology development during melt blending⁶⁰.</figcaption></figure>

The control of the morphologies is influenced by several parameters such as the geometry of the screw, the viscosity ratio between the polymers as well as the shearing intensity⁶⁹ and the residence time. Finally, the compatibility between the polymers shows also an importance on the obtained morphology⁷⁰ as well as its thermal stability⁷¹. In order to develop those intermediate morphologies different approaches have been classically followed like the chaotic mixing⁷² where two flows of polymers are introduced in a mixing element using two rods that will rotate alternately creating a chaotic mixing. The morphology is then controlled by the sequences of rotation of the rods leading to lamellar or fibrillar morphologies. Another route will be to form the morphology after the extrusion step using a stretching step towards a fibrillar morphology⁷³. A lamellar morphology can also be obtained by using a coextrusion process to produce multilayered pellets like ABS/PC/ABS) followed by an injection step that will destroy partially the layered morphology into the lamellar one, Figure 39.
<figure><img src="image_43.png" /><figcaption>**Figure 39.** LME extrusion system to produce lamellar sample</figcaption></figure>

This latter approach was followed using LME equipment with a starting (ABS/PC/ABS) three layers flow with a ratio 70/30 wt% ABS/PC to obtain multilayered structures, similarly to what has been discussed in Chapter 2. The aim here is to produce layered structures of ABS/PC with a range of individual thickness of layers going from millimeter to nanometer scale<sup>74</sup>.

### IV.1.2. Morphology characterization

In order to verify the efficiency of the control of the morphologies with different numbers of LME, strands were collected after extrusion and before pelletizing and were sliced by microtome. Those slices were observed with microscopic techniques adapted to the scale of the structures. Figure 40 presents different optical micrographs of transversal slices obtained with the microtome, and one AFM image from the 10 M strand. White areas are attributed to PC phase (smooth areas in AFM).

<figure><img src="image_44.png" /><figcaption>**Figure 40.** Blend morphologies observed in slices of the obtained strands. Optical micrographs for (a) 0 M (b) 3 M (c) 6 M (d) AFM image for 10 M</figcaption></figure>

As expected, the strands obtained with the multilayer coextrusion process are composed of alternated layers of ABS and PC (3 layers with 0 mixing element, 17 layers with 3 mixing elements, 129 layers with 6 mixing elements).
elements and 2049 with 10 mixing elements). In the 0 M image (Figure 40 (a)), the thickness of the PC layer is irregular along the sample. In all multilayered sample, there is a distortion of the layers when traveling from the square section coextrusion channel into the circular die. This distortion does not appear to break layers. These thickness variations are maximum in 0 M where the single PC layer is highly deformed. With higher amount of multipliers, only outer layers are distorted.

After this, the strands were cut into pellets and could be injected into dogbone samples in order to test the mechanical properties of those blends. The injection step will obviously degrade drastically the micro/nanostructure that has been developed during the first step and the aim is then to characterize the remaining structures after the injection process. To characterize it, the dogbones samples were cut into two parts and the center part of the samples is then observed by SEM to determine the blend morphology, as seen in Figure 41.

<figure><img src="image_45.png" /><figcaption>**Figure 41.** Dogbone samples cut in half on the left and a SEM images for a ABS/PC samples injected from structured pellets</figcaption></figure>

This analysis has been carried out in function of the number of layer multiplying elements used, see Figure 42.

<figure><img src="image_46.png" /><figcaption>**Figure 42.** SEM micrographs at higher magnification of the skin region of injected specimens (longitudinal cut) for 0, 3, 6 and 10 LME and with a reference made by twin screw extrusion (TSE)</figcaption></figure>

This investigation was carried out at high magnification with SEM and showed that the layers are discontinuous, leading to a mix of lamellar, fibrillar and nodular morphology. Figure 42 presents SEM micrographs in the skin region of injected specimens for each sample. To a certain extent, all samples presented some degradation compared to their pellets' morphology. Thus, all samples contain nodules, most contain fibrils and only 0 M, 3 M and 6 M retained some platelets. As the initial layers gets thinner, the
amount of lamellae drops in the injected sample. This breakup is complete in 10 M where only nodules can be observed, leading to a microstructure comparable to the sample obtained by twin screw extrusion (TSE) which is supposed to be the classical morphology obtained by blending polymers, i.e. a nodular one.

<table><caption>**Table 4. Description of the blend morphologies in injected specimens**</caption><thead><tr><td>Sample</td><td>Morphology</td><td>Mean aspect ratio</td></tr></thead><tbody><tr><td>TSE</td><td>Nodules</td><td>3.0 ± 0.1</td></tr><tr><td>0 LME</td><td>Platelets/Fibrils/Nodules</td><td>3.4 ± 0.2</td></tr><tr><td>3 LME</td><td>Platelets/Fibrils/Nodules</td><td>5.0 ± 0.4</td></tr><tr><td>6 LME</td><td>Platelets/Fibrils/Nodules</td><td>3.4 ± 0.1</td></tr><tr><td>10 LME</td><td>Nodules</td><td>2.2 ± 0.1</td></tr></tbody></table>

Table 4 gives an overview of the structures observed in each injected sample. All observed structures are listed and the most common ones are written in bold. To obtain the aspect ratio, cross sections of the skin region were observed at high magnitude. A large number of PC particles were analyzed and their maximum and minimum dimensions were measured. The aspect ratio of a particle is the ratio of its maximum dimension over its minimum one. A mean aspect ratio of 1 would translate to a nodular morphology. The higher the mean aspect ratio is, the more elongated the PC particles are. 3 LME retain a high aspect ratio indicating an elongated morphology, followed by 0 LME and 6 LME. TSE and 10 LME have the lowest values, as expected. It is of interest to point out that the aspect ratio in 10 LME is the lowest of all samples pointing to a morphology closer to nodular than TSE, even though it started with thin platelets. The 0 LME value is surprisingly low considering the trend observed for other samples. This is due to a high number of nodule like structure on the studied micrographs. Platelets are frequent, but the high number of nodules hinders the mean aspect ratio value.

Whichever considered sample, initial morphology undergoes stretching during the injection step, leading to a significant reduction of the thickness compared to initial features. For instance, the minimum dimension in 0 LME is never higher than 5 µm, where initial layers were 150 µm thick. Nodules from TSE were stretched leading to the observed 3.0 ± 0.1 aspect ratio. It is important to emphasize that these measurements were made in the skin region, where shear rate is the highest, hence such stretching was expected.

### IV.1.3. Mechanical properties of structured blends

The influence of morphology on impact strength was then studied and it appears obvious and maximum for the 0 LME sample, as seen in Figure 43. For this morphology, the impact strength shows a 65% increase compared to basic mixing law and double the value of the “classic” nodular morphology obtained in the TSE sample. Every sample that exhibits special morphology (non-nodular) also exhibits an impact strength value higher than mixing rule. In the case of TSE and 10 LME, the nodular morphology leads to lower impact strength. However, it should be noted that 10 LME has a significantly higher value than TSE, which suggests that part of the initial morphology may be retained.

Platelet morphology leads to increased impact strength and elongation at break compared to standard nodular morphology.
<figure><img src="image_47.png" /><figcaption>Figure 43. Impact strength of all injected sample. Pure ABS, pure PC and expected value from mixing rule are also given in dashed lines</figcaption></figure>

The control of the morphology may thus allow an improvement of the mechanical properties, especially impact strength, of blends of polymers emanating from WEEE.

#### IV.1.4. Chemical compatibilization of the blend

In order to control the morphology and its retinment after the injection step, a chemical modification of ABS was attempted to improve its compatibility with PC. A grafting of maleic anhydride (MA) was carried out through reactive extrusion with a radical mechanism using a peroxide initiation, Figure 44.

<figure><img src="image_48.png" /><figcaption>Figure 44. Grafting of maleic anhydride on ABS terpolymer</figcaption></figure>

This grafting has been realized in two steps: a first one in twin screw extrusion to create a batch of ABS with MA without initiator and a second step in twin screw extrusion as well with the introduction of the blend previously formed and the initiator. This methodology prevented the reaction of MA with itself as the dissociation of dicumyl peroxide is fast at this temperature (half-life time around 1 min at 180°C) which would have then occurred before the mixing of ABS and MA.

The efficiency of the grafting reaction was verified by dissolution of 1g of this product in acetone. The solution is precipitated in ethanol to remove any unreacted MA. The polymer is finally characterized by FTIR and the presence of a carbonyl function evidenced the grafting of MA on the ABS. Indeed, a similar system without use of initiator was analyzed and no presence of carbonyl function could be observed after the purification step. The maleic anhydride is then expected to allow van der waals interactions or even a chemical reaction with PC to create copolymers of ABS and PC allowing an improvement of the compatibility of the two phases.

Multi-layered systems ABS/PC/ABS were coextruded with the introduction of this maleic anhydride-grafted-ABS (MAgABS) additive in the ABS or PC flow with a content of 2 to 5 wt%. This addition didn't
have a noticeable change in the thickness and continuity of the layers obtained in the pellets except for the 10 LME conditions where some breaking up of the layers is observed. After the injection step however, an evolution of the shape of the minor phase is observed with more spherical particles of PC than without addition of the compatibilizer, as can be seen on Figure 45 where the maximal and minimal dimensions of the particles are closer in compatibilized sample than the one without the additive. This might be explained by the fact that the compatibilizer added will affect the surface tension around PC domain phases and will favour the more stable morphology thermodynamically, i.e. a nodular one. The elongated structures is thus more easily lost during the injection step.

<figure><img src="image_49.png" /><figcaption>Figure 45. Maximal (D<sub>max</sub>) and minimal (D<sub>min</sub>) dimensions of PC particles in 0 LME ABS/PC samples with (orange histogram) or without (blue histogram) addition of ABS-g-MA.</figcaption></figure>

In terms of mechanical properties, the same trend as what has been observed previously without additive is obtained with the best results for 0 LME pellets and a decrease in the impact strength with an increase in LME number. However, the different results are lower in every cases for additivated systems which is in good coherence with the morphology observation.

### IV.1.5. Conclusion

The link between microstructure and mechanical properties of 70/30 wt% ABS/PC blend was investigated. ABS/PC being a compatible but immiscible blend, it is ideal to study the links between morphology and properties of the blend as we can rule out the influence of incompatibility.

Structured pellets were extruded with three different techniques. A twin screw extruder was used to obtain basic nodular morphology. Coextrusion and LMEs were used to create four different types of structured pellets consisting of stacking thin ABS/PC layers. These structured pellets were then injection molded and obtained samples were studied to evaluate both the impact of morphology on mechanical properties and the degree of morphology retention during injection.

Microscopic observation revealed that the microstructure was partially preserved during injection molding. However, as the specific dimension of the domains decreases, the ability to retain the morphology during injection diminished. The morphology loss was total for the 10 LME sample (initial layer thickness lower than 1 µm). These results confirmed that nodular morphology is the most stable as platelets collapsed into fibrils and further into nodules.

It is clear from the impact strength tests that the morphology has a significant influence on final mechanical properties. Platelets is the most favorable morphology to boost impact strength. An increase of 65% in impact strength compared to mixing law was found for the 0 LME sample (platelets). This reaches twice the value of the TSE sample with conventional nodular morphology.

A PhD student is currently trying to improve this methodology by realizing plates of PC/ABS through multilayer coextrusion without the use of injection to determine the effect of the nanostructuration, without its deterioration step, on mechanical properties and impact test especially.
## IV.2. Improvement of PLA/PHBV blends mechanical properties

On the topic of mechanical properties, the use of biopolymers have received great attention both in industry and in academia due to an increased concern toward the environmental impact of plastic waste and the saving of limited fossil energy. Poly(lactide) (PLA), one of the most used biobased and biodegradable polymers, is synthesized from lactic acid obtained from renewable resources (corn, sugar, potato). Its properties make it a potential candidate to be used in various applications, such as packaging, medicine, agriculture, textile, and automotive industries⁷⁵. However, despite its good optical properties and high tensile strength, its inherent brittleness and also high gas permeability significantly limit its applications, especially in food packaging⁷⁵⁻⁷⁸. Many strategies have been recently developed and proposed to improve these properties, like the addition of modifiers, copolymerization or blending⁷⁹⁻⁸².

Blending PLA with another biobased and biodegradable polymer has attracted much attention since it offers simply the opportunity to widen the range of physical properties, without compromising environmental features of the final material. Among polyhydroxyalcanoates (PHAs), which are aliphatic polyesters obtained by microbial fermentation, Poly(3-hydroxybutyrate) (PHB) and its copolymers (poly(3-hydroxybutyrate-co-3-hydroxyvalerate) (PHBV), poly(3-hydroxybutyrate-co-3-hydroxyoctanoate (PHBHO), poly(3-hydroxybutyrate-co-3-hydroxyhexanoate) (PHBHHx)... ) have been largely used for PLA-based blends⁸³. Some of the resulting blends exhibit improved properties compared to neat PLA: higher toughness and increased elongation at break⁸⁴⁻⁹⁰, improved gas barrier properties⁹¹,⁹², increased thermal stability⁸⁹, enhanced biodegradability⁹⁰.

### IV.2.1. Compatibilization of PLA/PHBV blends

Those improvements are partly due to the fact that, although immiscible, the two polymers have a certain degree of chemical affinity. Miscibility and compatibility of PLA/PHB(V) have been thoroughly studied and have been found to depend classically on the molecular weight of both components, the blending method, the blend composition and the crystallinity of both components. Blümm and Owen⁹³ observed miscibility of low molecular weight poly(L-lactide) (PLLA) (M̅n = 1 800 g/mol) and immiscibility of high molecular weight PLLA (M̅n = 160 000 g/mol) blended with PHB (M̅n = 222 000 g/mol) over the whole composition range. Similarly, Koyama and Doi⁹⁴ investigated the miscibility of PHB (M̅w = 650 000 g/mol) blended with PLA of various molecular weights and determined that the blend is immiscible above a critical value of molecular weight of PLA around 20 000 g/mol. More generally, all literature agrees on the immiscibility of the blend as soon as the molecular weight of both components is high. Zhang et al.⁹⁵ found that melt blended samples exhibited greater compatibility than those prepared by solvent casting at room temperature, possibly due to a transesterification reaction between PHB and PLA occurring at high temperature. This transesterification, or more generally the presence of specific interactions between the two polymers, has been evidenced by many other authors⁸⁸,⁹⁰,⁹⁶,⁹⁷, especially for blends containing a low amount of PHB(V). For such PLA/PHBV blends (typically weight fractions of 90%/10%), Gerard and Budtova⁹⁶ observed a peculiar morphology, with very small PHBV nodules (about 400 nm) well dispersed in PLA matrix, with increased adhesion at the interface, presumably explaining the increase of ductility (elongation at break of 200 % when tested immediately after sample preparation) measured for this composition by the authors.

All those studies have highlighted the importance of the interface on the final properties of PLA/PHB(V) blends and some attempts have been undertaken to modify it, using different strategies. Yoon et al.⁹⁸ compatibilized a mixture of PHB and PLA (M̅n = 56 000 g/mol) using synthesized Poly(ethylene glycol) – b – PLLA (PEG-b-PLLA) (PEG is known to be compatible with PHB) and poly(vinyl acetate) (PVAc, compatible with both PLA and PHB). A small improvement in toughness was obtained with the addition
of 2 wt% of block copolymers or PVAc, but the overall mechanical properties were not improved much. One possible reason proposed by the authors could be the solubilization of the compatibilizer molecules in either PHB or PLA, or in both, instead of placing themselves at the interface between the PHB and PLA. An alternative approach was proposed by Yang et al.⁹⁹ by adding transesterification catalysts in order to induce chemical reaction between the two components. They found a significant effect on the mechanical properties by adding a small amount of zinc acetate (0.1 wt% in the blend).

Finally, some authors have formulated PLA/PHB(V) blends in order to improve their mechanical properties and/or processability, adding different components as plasticizers, either commercial or synthesized ad-hoc: PEG¹⁰⁰,¹⁰¹, ATBC¹⁰⁰, limonene¹⁰², polyester plasticizer (Lapol 108)¹⁰³, lactic acid oligomer (oLA)¹⁰⁴ and tributyrin¹⁰⁵. It has been accepted that concentrations from 10 to 20 wt% and up to 30 wt% are required to provide a substantial plasticization effect, which strongly modifies the mechanical properties: if the ductility is usually enhanced, the Young's modulus and the tensile strength are decreased. Besides, the gas barrier properties of the plasticized blends are often deteriorated as a consequence of the incorporation of the plasticizer, which induces a gain of mobility of the polymer chains. Nevertheless, some authors noticed that, when the plasticizer has good miscibility with both of the components, it can have a beneficial effect on the compatibility between them. For instance, Armentano et al. [31] observed a reduction of the size of PHB domains with the addition of oLA ($\overline{M}_n$ = 957 g/mol) in PLA/PHB (85/15 wt%), indicating improved interfacial properties. The enhanced ductility can therefore be interpreted as the consequence of a combined effect of plasticization and compatibilization. This oLA was added from 15 to 30 wt%.

### IV.2.2. Preparation of oligomers of PLA

Here, a different strategy to compatibilize PLA/PHBV, see Figure 46, blends was developed, which will preserve the tensile strength and bio-based character. It consists in introducing a much smaller amount (between 0.1 and 1 wt%) of hydrolyzed PLA having an oligomer-type molar mass ($\overline{M}_n$ = 5 000 g/mol) to the PLA/PHBV blends. It is anticipated that the low molecular weight PLA (oLA) will improve the compatibility between PLA and PHBV without a significant drop in the mechanical properties. The resulting thermal, morphological and mechanical properties of the obtained biobased PLA/PHBV blends are investigated¹⁰⁶.

The hydrolysis is realized in distilled water at neutral pH and at a temperature of 100 °C. About 40 g of PLA are introduced into 4 L of water in a mount of reflux containing a round-bottom flask of 5 L. The use of a large amount of water will keep a pH close to 7 during degradation. Indeed the production of lactic acid during the degradation tends to acidify and catalyze the hydrolysis reaction which may cause a rapid degradation of the macromolecular chain or a change in the degradation mechanism.

<figure><img src="image_50.png" /><figcaption>Figure 46. Chemical structures of a) Poly(lactic acid) and b) poly(3-hydroxybutyrate-co-3-hydroxyvalerate)</figcaption></figure>

The obtained samples are dried to a constant weight in a ventilated oven at 65 °C and their molecular weight is measured by size-exclusion chromatography (SEC). SEC analyses of the PLA after 2 to 43 hours of hydrolysis have two molecular chains distributions (one around 5 000 g/mol and one with a higher average
molar mass). After 45 hours, a single distribution around 5 000 g/mol is obtained. Therefore the oLA ($\overline{M}_n$ ≈ 5 000 g/mol) with D close to 1.25 is selected for this study.

#### IV.2.3. Thermal characterization of compatibilized PLA/PHBV blends

Films of PLA, PHBV, PLA/PHBV (90/10 wt%) and this blend with the addition of oLA with a content of 0.1, 0.5 and 1% were realized through single-screw extrusion with a flat die.

<figure><img src="image_51.png" /><figcaption>Figure 47. First heating run in DSC for PLA, PLA/PHBV and PLA/PHBV/oLA films and PHBV pellets.</figcaption></figure>

The DSC thermograms from the first heating run of these films are shown in Figure 47. The curves reveal different thermal events as follows: the glass transition (Tg), the cold-crystallization (characterized by Tcc and the enthalpy of crystallization $\Delta$Hcc), the melting point (characterized by Tm and the melting enthalpy $\Delta$Hm). All these data are summarized in Table 5. The Tg value for the PLA neat films was about 60 °C, while the exothermic peak is observed at about 118 °C and the multi-step melting endothermic peak is obtained at about 151-157 °C. The DSC thermograms for the PLA/PHBV blend display a slight shift to lower Tg values with respect to neat PLA and a similar multi-step melting with the first and second peaks (Tm1 and Tm2) due to the PLA and the third one (Tm3) corresponding to the PHBV around 170 °C. This suggests no complete miscibility between the two polymers<sup>107</sup>. The double melting peak in the range of 147–156 °C can be attributed to the formation of different crystal structures for the PLA<sup>89,108</sup>.

In the first heating run, characterizing the thermal properties of the films “as prepared”, a slight decrease of the glass transition temperature of PLA from 58 °C for the neat PLA film to 56 °C for the PLA/PHBV film and 54 °C for the PLA/PHBV/oLA 89/10/1 is observed. This gradual decrease of the Tg of PLA phase can be understood as an increase of compatibility of PLA/PHBV blends due to the presence of the oLA, miscible with both PLA and PHBV. The cold crystallization temperature of PLA is decreased from 119°C to 105°C with the addition of PHBV, as already observed by previous authors<sup>88,89,109</sup>, and is not influenced by oLA addition. The PLA degree of crystallinity also remains extremely low in every cases (<5%).

When blended with high molar mass PLA, PHBV loses its crystallinity, with a sharp decrease from 85 % to 51 %, as confirmed by Ma et al.<sup>85</sup>. But the degree of crystallinity of PHBV increases when an optimum value of oLA is added (from 51 % without any oLA to 74 % with 0.5 wt% oLA). This follows the same trend as
what has been observed for PHBV/oLA blends, with a maximum of cristallinity for 0.5 wt% oLA and can be seen as a plasticizing effect of the oLA on the dispersed PHBV phase. This effect will then hinder the crystallization at higher amount (1 wt%). Such effect is also observed in the second heating run with an increase from 18 to 31% in the PHBV degree of crystallinity with 0.5 wt% of oLA compared to pure PHBV.

Finally, it can be noted that PHBV and oLA seems to have an impact on the crystalline forms of PLA. PHBV almost prevents the appearance of the crystalline form that melts at 150°C while the addition of oLA has the opposite effect. With 1% of oLA, the 150°C form is even predominant and in higher concentration than in PLA alone.

<table><caption>Table 5. Thermal analysis of PLA, PLA/PHBV and PLA/PHBV/oLA films as a function of the amount of oLA (wt%).</caption><thead><tr><th></th><th></th><th>T<sub>g</sub><br>(°C)</th><th>T<sub>m1</sub><br>(°C)</th><th>T<sub>m2</sub><br>(°C)</th><th>T<sub>m3</sub><br>(°C)</th><th>T<sub>cc</sub><br>(°C)</th><th>ΔH<sub>m3</sub><br>(J/g)<br>(PHBV)</th><th>ΔH<sub>cc</sub><br>(J/g)</th><th>X<sub>c</sub><br>(%)<br>(PLA)</th><th>X<sub>c</sub><br>(%)<br>(PHBV)</th></tr></thead><tbody><tr><td rowspan="5">2<sup>nd</sup><br>heating<br>run</td><td>PLA</td><td>60</td><td>151</td><td>157</td><td>-</td><td>118</td><td>-</td><td>34</td><td>0.1</td><td>-</td></tr><tr><td>PLA/PHBV</td><td>57</td><td>150</td><td>156</td><td>169</td><td>111</td><td>1.96</td><td>29</td><td>3.2</td><td>18</td></tr><tr><td>PLA/PHBV<br>/oLA 0.1%</td><td>57</td><td>150</td><td>156</td><td>167</td><td>114</td><td>1.76</td><td>29</td><td>1.4</td><td>16</td></tr><tr><td>PLA/PHBV<br>/oLA 0.5%</td><td>57</td><td>151</td><td>156</td><td>167</td><td>117</td><td>3.38</td><td>31</td><td>0.6</td><td>31</td></tr><tr><td>PLA/PHBV/<br>/oLA 1%</td><td>58</td><td>151</td><td>-</td><td>168</td><td>117</td><td>1.47</td><td>31</td><td>0</td><td>13</td></tr><tr><td rowspan="5">1<sup>st</sup><br>heating<br>run</td><td>PLA</td><td>58</td><td>152</td><td>157</td><td>-</td><td>119</td><td>-</td><td>34</td><td>1.1</td><td>-</td></tr><tr><td>PLA/PHBV</td><td>56</td><td>147</td><td>156</td><td>170</td><td>105</td><td>5,53</td><td>29</td><td>2.7</td><td>51</td></tr><tr><td>PLA/PHBV/<br>/oLA 0.1%</td><td>55</td><td>147</td><td>156</td><td>170</td><td>104</td><td>5,54</td><td>29</td><td>1.5</td><td>51</td></tr><tr><td>PLA/PHBV/<br>/oLA 0.5%</td><td>55</td><td>148</td><td>156</td><td>170</td><td>105</td><td>8,08</td><td>28</td><td>1.2</td><td>74</td></tr><tr><td>PLA/PHBV/<br>/oLA 1%</td><td>54</td><td>148</td><td>156</td><td>171</td><td>105</td><td>5,92</td><td>26</td><td>0.6</td><td>54</td></tr></tbody></table>

IV.2.4. Morphology characterization

<figure><img src="image_52.png" /><figcaption>Figure 48. AFM micrographs of PLA/PHBV (90/10) (a) and PLA/PHBV/oLA (89/10/1) (b)</figcaption></figure>

Figure 48 shows the AFM micrographs of the fracture surface of PLA/PHBV (Figure 48 (a)) and PLA/PHBV/oLA blends (Figure 48 (b)). AFM images of PLA/PHBV films underlined that the dispersed PHBV phase had relatively small average diameter with a typical nodular morphology dispersed in the PLA
polymer matrix. With the addition of 1 wt% of oLA, the size of the PHBV domains is reduced but the shape remains unchanged. The mean area goes from 0.14 µm² for the PLA/PHBV film to 0.11 µm² for the PLA/PHBV/oLA 89/10/1 film while the ratio length/width is around 2 in both cases. To support this observation and explain the standard deviation of the domains area measured for PLA/PHBV/oLA 1 wt%, the global distribution of the nodules sizes of PHBV is presented in Figure 49. If the PLA/PHBV film presents a broad single distribution, the additivated film shows a double distribution, with a sharper main distribution shifted to smaller size and the presence of a small portion of remaining large domains of PHBV (>0.4 µm²) that increase the average size of the nodule observed. However a clear tendency to the decrease of PHBV size domain is observed: excluding the few remaining large domains, the average area of the compatibilized nodules is less than 0.09 µm² (while it is 0.13 µm² for the PLA/PHBV nodules), confirming the compatibilizing effect of oLA.

<figure><img src="image_53.png" /><figcaption>Figure 49. PHBV nodule size repartition in PLA/PHBV and PLA/PHBV/oLA 1%wt blends. Shapes of the distributions are eyeguides only.</figcaption></figure>

#### IV.2.5. Mechanical properties

The results of mechanical analysis for PLA/PHBV and PLA/PHBV/oLA with different amount of oLA films are summarized in Table 6.

The addition of small amounts of oLA reveals a slight improvement in the Young's modulus (E<sub>Young</sub>) over PLA/PHBV (around 10%) while no effect is observed for single PLA, which may be linked to the increase in the crystallinity of the PHBV nodules. Yield strength however, remains similar for all the samples tested. The elongation at break (ε<sub>b</sub>) of PLA/PHBV films increases with the addition of PLA short chains, showing an important enhancement in blends ductility. Films with 1 wt% oLA shows the highest elongation at break and homogeneity, since low values of standard deviations are measured.

This improved elongation at break can be interpreted by two ways: the first explanation is that oLA acts as compatibilizer so it creates an interphase between PLA and PHBV facilitating stress transmission at the interface. The second explanation is based on the fact that the oLA acts as a plasticizer. A non-significant evolution in elongation at break from 3.1 % for the neat PLA to 2.8 % for PLA/oLA (99/1) blend is observed. This result shows that the plasticizing effect is negligible for the PLA phase and the
compatibilizing effect of the oLA led to the improvement of the elongation at break for PLA/PHBV/oLA films.

<table><caption>Table 6. Results from tensile test for PLA, PLA/oLA and PLA/PHBV films as a function of the amount of oLA (wt%)</caption><thead><tr><th rowspan="2">Samples</th><th colspan="3">Mechanical properties</th></tr><tr><th>σ<sub>y</sub><br>(MPa)</th><th>ε<sub>b</sub> (%)</th><th>E<sub>Young</sub><br>(MPa)</th></tr></thead><tbody><tr><td>PLA</td><td>45 ±4</td><td>3.1±1</td><td>2050 ±481</td></tr><tr><td>PLA/oLA 1%</td><td>48 ±7</td><td>2.8±0.5</td><td>2107 ±210</td></tr><tr><td>PHBV</td><td>39</td><td>2</td><td>2800-3500[46]</td></tr><tr><td>PLA/PHBV</td><td>43 ±5</td><td>3 ±1</td><td>2507 ±252</td></tr><tr><td>PLA/PHBV/oLA 0.1%</td><td>47 ±1</td><td>8 ±2</td><td>2747 ±452</td></tr><tr><td>PLA/PHBV/oLA 0.5%</td><td>45 ±4</td><td>6 ±1</td><td>2709 ±382</td></tr><tr><td>PLA/PHBV/oLA 1%</td><td>44 ±2</td><td><b>14 ±3</b></td><td>±133 2799</td></tr></tbody></table>

IV.3. Conclusion

Two different multimaterial systems have been used to determine a link between morphology and mechanical properties and the impact of a chemical compatibilization on it.

ABS/PC multi-layered pellets were produced and final dog bone samples were injected with a retention of the previous morphology into a lamellar one. The influence of this morphology on impact resistance was observed and an improvement of this property could be measured with the more elongated shape of minor phase. A compatibilization of those two polymers was attempted through reactive extrusion by producing maleic anhydride-grafted-ABS however this additive led to a faster degradation of the produced morphology to a more thermodynamically stable nodular one. Finally, this additive couldn't help improving the mechanical properties as the needed morphology is more easily lost with this addition. The upcycling of ABS/PC emanating from WEEE can then be envisaged through the best conditions observed in this study, by using pellets simply coextruded in a three-layer configuration that eventually lead after the injection to an improvement of impact strength of 65% in comparison to blends obtained by twin-screw extrusion.

For the second system, bioplastics have been used and films composed of blends of PLA with PHBV were extruded. The addition of an oligomer of PLA was carried out and its effect on crystallinity of both PLA and PHBV, on compatibility between both and on the morphology of the blend has been studied. The size of the minor domain has then been impacted and a general decrease of these sizes has been observed from 0.13 µm² to 0.09 µm² with the addition of only 1wt% of oLA.

Finally an improvement of elongation at break of PLA/PHBV blends as well as an increase in Young's modulus have been observed with the addition of this oLA from 0.1 to 1wt%. This could be understood as an improvement of compatibility between PLA and PHBV as well as an impact of the size of the morphology on these properties.

In this chapter, it has been demonstrated that the morphology of blends has a significant importance on mechanical properties. The modification of the chemical composition of these blends was attempted through the addition of chemical functions and other molecules in order to investigate the influence of the compatibility on the morphology and in fine the mechanical properties as well.
# V. Block copolymer self assembly and its combination for actuator application

To obtain nanostructures made of two polymers, different approaches have been seen in the previous chapters. However, one of the most used methodology is to chemically link two polymers into a block copolymer (BCP). This chemical link will prevent a phase separation in the micrometer range and will keep it at a nanoscale and will lead to specific morphologies according to some parameters such as the Flory-Huggins parameter $\chi$ between the two blocks, the length of each block and the global polymerization degree (N)¹¹⁰. This chapter will deal with another kind of parameter that may be used to control the morphology created with the addition of a molecule that will be incorporated specifically in one of the phases. The chosen molecules are ionic liquids which are a combination of a cation and an anion that presents the specificity of being in the liquid state over a large range of temperature¹¹¹. This will then help modifying the obtained morphology by swelling one of the two phases as well as influencing the $\chi$ parameter but will also lead to the addition of specific properties such as ionic conductivity and will modify the mechanical properties of the block copolymer. Finally, the combination of those different properties will be used to develop actuators that will present a deformation under a voltage when placed between electrodes. The interest in these type of soft actuators has been growing recently due to higher flexibility, faster response to electrical stimulus and light compared to classical “hard” actuators made of piezoelectric ceramic or metallic materials¹¹². Different kind of soft actuators have been developing and the category aimed with the combination BCP/IL has a main advantage of being usable with low voltages (only a few volts)¹¹³.

## V.1. Block copolymer combination with ionic liquids

Blends of BCPs with functional additives have proven to be increasingly interesting since adding small molecules gives new compelling properties to the preexisting BCPs without the implementation of complex syntheses, and provides a versatile method for fabricating predictably ordered microstructures.

More recently, ionic liquids (ILs) have emerged as a class of chemicals which are room-temperature molten salts composed entirely of ions¹¹⁴,¹¹⁵. They show unique properties such as high thermal and chemical stability, non-inflammability, negligible vapor pressure, very low toxicity, and high conductivity.¹¹⁶,¹¹⁷ ILs become great candidates for the creation of functional materials due to their tunable properties such as hydrophilicity, acidity, and hydrogen bonding because of the unlimited possibilities for anions and cations combinations. These properties make them highly compatible with various synthetic polymers and good candidates as solvents or additives¹¹⁸⁻¹²⁰.

In the literature, several diblock copolymers, such as poly(styrene-b-ethylene oxide) (PS-b-PEO)¹²¹, poly(styrene-b-methyl methacrylate) (PS-b-PMMA)¹²²⁻¹²⁴, poly(styrene sulfonate-b-methyl butylene) (PSS-b-PMB)¹²⁵,¹²⁶ or poly(styrene-b-2-vinyl pyridine) (PS-b-P2VP)¹²⁷⁻¹²⁹, have been blended with ILs which appear to be miscible with the more polar block of the copolymer. For instance, for PS-b-PMMA, certain ILs have been found to preferentially associate with the polar PMMA block of copolymer¹²²⁻¹²⁴. The ILs act as plasticizers as they enhance the mobility of this block, reducing the glass transition temperature. Interestingly, they also increase the Flory-Huggins interaction parameter $\chi$ and influence the center-to-center distance by swelling the IL-containing phase¹²²⁻¹²⁴,¹²⁸,¹³⁰,¹³¹. Order-to-order lyotropic morphology transitions from lamellar to disordered spherical have been observed by studying different ILs and changing their relative amount in a same BCP¹²²⁻¹²⁴.

Materials based on ILs and BCPs have already found applications in several areas, such as in the manufacturing of polymer electrolytes for batteries, membranes, fuel cells, actuators, and photonics¹²⁵,¹³²⁻¹³⁵. The common denominator in all these applications is the need to produce a material that is not only
conductive but also solid, which are two antagonistic properties for this type of materials. Indeed, the conductivity is necessarily associated with molecular mobility, which de facto reduces the mechanical properties of the final material. Then, the use of ABA triblock copolymers, ensuring both mechanical strength and ionic conductivity, appears obvious. By selectively incorporating ILs within nanostructured morphologies of triblock copolymers, the best of both worlds could be achieved. Watanabe et al. investigated the microstructures, rheological properties, and ionic conductivity of ionic gels formed by self-assembly of PS-b-PMMA-b-PS triblock copolymers with 1-ethyl-3-methylimidazolium bis(trifluoromethyl-sulfonyl)imide ([Emim][Tf2N]).136 PS, immiscible with IL, creates spherical domains that serve as physical crosslink points. PMMA forms the continuous phase with the IL, which serves as the ionic conduction path. These polymer electrolytes were then used as an electroactive polymer actuator composed of ion gel electrolytes sandwiched between two composite carbon electrodes. By applying low voltages (<3.0 V) to the electrodes, the actuator exhibited a slow bending motion towards the anode side. The IL concentration and the length of the actuator (geometry of materials) have effects on the movement of the actuator. Various studies have confirmed the observation of greater displacement of the actuator by increasing the concentration of IL in the material and/or the length of the actuator136. The ionic conductivity measurements on these materials confirmed the relation between tortuosity of different morphologies and ionic conduction path, which is crucial for the application137,138. But the effects of ILs to obtain different possible morphologies for self-assembly and the effects of different types of morphologies on mechanical strength and ionic conductivity of the final material remain as a challenging question in the literature. For this, the gyroid morphology is targeted as it allows cocontinuous nanostructures leading to high conductivity. In classical BCP systems, N is usually decreased to obtain low xN values close to the Weak Segregation Limit, in which the gyroid morphology is present. However, in this region, small molecular weights lead to low mechanical properties, which limits their practical use.

### V.1.1. Specific incorporation of IL in PMMA phase

Here, we used a BCP with high molecular weight and low x parameter to access the Weak Segregation Limit at the disorder-order transition¹³⁹. The objective was to promote the gyroid morphology by adding an ionic liquid. Thus, we studied the effect of ionic liquids additives on the self-assembly process of ABA type triblock copolymers with an elastomeric central B block and rigid end A blocks which are well-known commercially available thermoplastic elastomers (TPEs), namely poly(methyl methacrylate-b-butyl acrylate-b-methyl methacrylate) (MAM), Figure 50 c. The two blocks have chemical similarities that provide a very low x parameter. Two different compositions of the same copolymer (MAM M52 with lamellar morphology and MAM M53 in the disordered state) were tested to study the impact of IL on the structuration (from disorder to ordered state) and improvement of the existing morphologies (from ordered state to well-ordered morphology) of BCPs. The MAM polymer films with and without ILs additives were characterized by different techniques to analyze their final morphologies. Two different kinds of IL were tested to study the influence of the chemical nature of the IL: 1-ethyl-3-methylpyridinium bis(trifluoromethyl-sulfonyl)imide ([EMPy][Tf₂N]) and 1-ethyl-3-methylimidazolium bis(trifluoromethyl-sulfonyl)imide ([Emim][Tf₂N]), see Figure 50 a and b.

The IL content in the BCP was also varied to enhance the formation of well-ordered morphologies in solvent-cast films. A lyotropic behavior was also detected.
<figure><img src="image_54.png" /><figcaption>Figure 50. Chemical structures of a) 1-ethyl-3-methylpyridinium bis(trifluoromethylsulfonyl)imide ([EMPy][Tf₂N]); b) 1-ethyl-3-methylimidazolium bis(trifluoromethylsulfonyl)imide ([Emim][Tf₂N]) and c) MAM copolymer</figcaption></figure>

<figure><img src="image_55.png" /><figcaption>Figure 51. DSC curves for MAM M53 (a) and MAM M52 (b) block copolymers blended [EMPy][Tf₂N] over different weight concentrations (φ_IL): 0%, 5%, 13%, 18%, 23%, 33%, and 44%. The arrows indicate the Tgs of the PBuA block (black) and the Tgs of the PMMA block (red for M53 and blue for M52). Insets show the evolution of the Tg of the PMMA block as a function of IL content (wt%).</figcaption></figure>

The selectivity of ILs with BCPs is of critical importance to understand the blends nanostructuration. The miscibility of the small molecule in one of the blocks of the copolymer is usually detected by glass transition temperature (Tg) shifts¹²³,¹³¹,¹⁴⁰. DSC is then the method of choice to study the impact of the presence of ILs on Tg. [EMPy][Tf₂N] was used first to study this incorporation. The heat flow thermograms of MAM block copolymers M52 and M53 with increasing ILs weight concentrations are shown in Figure 51. At the studied concentrations, the IL seems to be fully dispersed in the BCP matrix as no melting peak was observed at its melting temperature, i.e. 20 °C. The Tg of the PMMA block for the pure M53 BCPs is equal to 117 °C (glassy at room temperature), whereas the Tg of the PBuA block is close to -40 °C (rubbery at room temperature). Adding IL to the BCP causes a significant drop in the Tg of the PMMA block, decreasing to 2 °C with φ_IL = 44 wt%. The Tg of the PMMA block of the M52 BCP is 124 °C decreasing to 50 °C at the highest IL content (44 wt%). Both insets in Figure 51 show the evolution of the Tg of the PMMA block as a function of IL content. It shows a monotonous decrease, which follows the blend rule. In the same time, the Tg of the PBuA block remains fairly stable (rubbery at room temperature) for all IL concentrations, even though a very small shift towards lower temperatures can be observed at higher IL concentrations.
V.1.2. Block copolymer morphology evolution with IL content

The evolution of the self-assembly of the BCP was then followed by AFM imaging with increasing IL content. As a reminder, BCP has the ability to self-assemble into specific morphologies according to their polymerization degree, Flory-Huggins parameter χ and the BCP composition. The different morphologies that can be obtained are presented Figure 52¹⁴¹.

<figure><img src="image_56.png" /><figcaption>Figure 52. a) Phase diagram of diblock copolymer predicted by SCMF theory. b) Various microdomain organization patterns of a linear AB diblock copolymers. *f*: volume fraction of one block; χ: Flory-Huggins interaction parameter; N: degree of polymerization; L: lamellae; H: hexagonally packed cylinders; Q²³⁰: double-gyroid phase; Q²²⁹: body centered spheres; CPS: closed-packed spheres; and DIS: disordered¹⁴¹.</figcaption></figure>

Without IL addition, the BCP MAM M53 is in the disordered state as no specific morphology can be observed, Figure 53. During the evaporation of the non-selective solvent (CHCl₃), the blocks of the copolymer are not highly repulsive so that they stay in the disordered state as the concentration of CHCl₃ decreases to the point where no further mobility can be reached. The addition of IL ([EmPy][Tf₂N]) in the initial BCP solution induces a gradual disorder-order transition (DOT) from a disordered state to an ordered morphology between 0 and 33 wt%. At an IL concentration of 5 wt%, small ordered domains are seen by AFM, but most of the sample seems to remain in the disordered state, as also observed theoretically by Nakamura et al. and experimentally by several authors, in the case of BCP in the presence of Lithium salts¹⁴²,¹⁴³. As the IL concentration increases up to 13 wt%, larger ordered domains are detected. It is noteworthy that it is difficult to assign a specific morphology as both dots and stripes are seen by AFM. At an IL concentration of 18 wt% and 23 wt%, a coexistence of two different morphologies and a transition between dots and stripes were observed on AFM images. This suggests a transition from gyroid to lamellar morphology according to the composition of PMMA and PBuA blocks in BCP and the swelling of PMMA induced by the selective incorporation of IL¹²³,¹²⁴. However, it is important to analyze these films by another method to confirm the exact morphologies. As for the 33 wt% morphology, only stripes are observed, which favors the assumption of a lamellar morphology. This last AFM image also shows that the domain size in which a uniformly oriented morphology is observed, increases compared to the samples containing less IL. The 44 wt% sample could not be observed by AFM due to the plasticizing effect induced by the presence of IL, which prevented us from preparing the surface through ultramicrotomy. In Figure 53, SAXS 1D profiles of several samples consisted of BCPs mixed with different IL weight concentrations φ<sub>IL</sub> are shown. SAXS analysis results show that for MAM M53 where the morphology of the initial film without IL and of the samples with low IL concentrations (below 13 wt%) can be attributed to a disordered state with broad peaks, in good agreement with the AFM observations. However, increasing the concentration of IL in the BCP significantly narrows the width of the q* peak and at higher concentrations, higher order Bragg peaks appear, which is consistent with ordering of the BCP morphology. SAXS patterns of MAM M53 films
with 13 wt% and 23 wt% of IL reveal several positions of higher order scattering peaks, especially at $\sqrt{4/3}q^*$, $\sqrt{10/3}q^*$, $\sqrt{11/3}q^*$, $2q^*$, $3q^*$, which can be attributed to a coexisting gyroid/lamellae morphology$^{122-124}$. This result is in good agreement with the AFM images presented in Figure 53 c-e showing a coexistence of dots and lines. Lamellar morphology was also confirmed for the MAM M53 samples with higher IL concentration (33 wt% and 44 wt%) by the appearance of numerous higher order peaks at $2q^*$, $3q^*$, and $4q^*$.

<figure><img src="image_57.png" /><figcaption>**Figure 53.** Left: AFM phase images for MAM M53 block copolymer blended with ionic liquids (ILs) ([EmPy][Tf₂N]) with different weight concentrations φIL: (a) 0%, (b) 5%, (c) 13%, (d) 18%, (e) 23%, and (f) 33%. Right: SAXS profiles of MAM (a) M53 with different ionic liquid (IL) content φIL indicated on the graphs 0%, 5%, 13%, 23%, 33%, and 44%.</figcaption></figure>

## V.2. BCP/ILs combination mechanical properties

After understanding the morphology behavior in presence of ILs and in order to develop the actuator application, the mechanical properties were initially characterized and are presented in Figure 54-a. Thinking that the IL was simply going to act as a plasticizer, a lowering of the Young's modulus was expected. But the evolution presented on Figure 54-b shows that the lowering is not monotonous and that it would seem that around the IL compositions where a well-ordered morphology begins to appear (visible in SAXS and in AFM) a mechanical reinforcement is suggested due to the appearance of gyroid cocontinuous morphology. This evolution was also noticed by Cho et al.$^{144}$ who observed an influence of the addition of lithium salt on the mechanical properties which was characterized by an increase in the elastic modulus during the transition from hexagonal to gyroid morphology while this transition took place during a temperature heating. On the other hand, by continuing the heating, during the transition from gyroid to lamellar morphology, the elastic modulus then suddenly dropped. This evolution of mechanical properties is also accompanied by the appearance of a plasticity threshold presenting a necking then absent for a disordered, lamellar or cylindrical morphology, as Dair et al. have already noticed$^{145}$. This necking is visibly characteristic of the gyroid morphology and would be due to the 3D structure of the gyroid morphology which ensures continuity of the vitreous phase in the direction of stretching. We were also able to observe the appearance of this plasticity threshold with the addition of IL in a certain range of intermediate compositions where the gyroid morphology was suspected (between 13 and 23 wt%). With the highest IL concentrations studied (above 33 wt%), plasticization of the vitreous phase makes the material extremely flexible in a rubbery state where it is difficult to measure an elastic modulus in uniaxial tension.
<figure><img src="image_58.png" /><figcaption>Figure 54. a) Tensile tests with 10mm/min for MAM M53 block copolymers blended with ionic liquid (IL) ([EmPy][Tf₂N]) over different weight concentrations (φIL): 0%, 9%, 13%, 18%, 23%, 28% and 33%; b) Young's modulus of MAM M53 block copolymer blended with ionic liquid (IL) ([EmPy][Tf₂N]) over different weight concentrations (φIL): 0%, 9%, 13%, 18%, 23%, 28% and 33%</figcaption></figure>

### V.3. Actuator application of BCP/IL combinations

Then, to be able to use those materials as actuators, electrodes were needed to be implemented on both sides of the films. Gold leaves of around 100 nm thick were used as they are easily taped on a 50µm film. Those leaves were selected due to their high conductivity and very low thickness which shouldn't impact drastically the mechanical properties of the actuator and its deformation. The actuator is then sollicitated with a square tension and the deformation is followed by the displacement of the tip of the actuator by a camera. The actuator is oriented perpendicular to the ground in order to avoid the effect of gravity as described in Figure 55.

<figure><img src="image_59.png" /><figcaption>Figure 55. Scheme of the actuator analysis system</figcaption></figure>

The principle of the actuator mechanism is linked to the fact that the size of the ions composing the IL are different and when a current is applied on both side of the films, the ions will be attracted towards the different electrodes. The electrode that will attract the largest ions will then lead to the expansion of this side of the film while the other will shrink due to the smaller space occupied by the smallest ions, as described in Figure 56.
<figure><img src="image_60.png" /><figcaption>Figure 56. Scheme of the principle of actuator mechanism in IL based actuator¹²⁵</figcaption></figure>

The applied voltages vary in amplitude from 2 V to 50 V. As can be seen in Figure 57, the amplitude of displacement of the actuator is linked to the voltage applied on it. The first noticeable displacement can be detected as voltages as low as 2V. Then with an increase of the voltage, the amplitude is increasing in a similar manner. It can be noted that the displacement is almost linear with time during the entire voltage application meaning that an ion displacement is rather slow and is continuous over this period of time (120 seconds here).

<figure><img src="image_61.png" /><figcaption>Figure 57. Variation of the amplitude of displacement (right) in function of the voltage applied (left) for an actuator made of MAM M53 with 18wt% of IL ([EmPy][Tf₂N])</figcaption></figure>

Different IL content were used for this application and similar experiments were carried out at 2V. The maximum of amplitude was collected for the different content and summarized in Figure 58 a). An increase of this amplitude is observed with the increase of IL content which can be expected as higher amount of ions are present and a higher mobility is attained with higher IL content due to the plasticization of the polymer. Indeed, a decrease in the glass transition temperature of the PMMA phase is observed with the increase of IL content (Figure 51) which would increase the mobility inside the phase where the IL are miscible in. This effect is also confirmed by the increase of ionic conductivity that has been measured on different films of BCP/IL at different temperatures, see Figure 58 b). Finally, a decrease of the amplitude is observed for 38wt% of IL but it is due to a too soft material that deforms only by gravity. The mechanical properties of the blend at those compositions can't thus be properly studied. A limit of IL incorporation is then attained in this type of system for actuator application. No direct effect of the morphology can be observed at this stage even if a higher conductivity may be expected for gyroïd morphology due to the co-continuous nature of it. Indeed, the increase of IL content has several effects that may play a role in the conductivity and amplitude. At a first glance, the amount of ions present seems to be the main parameter involved for the observation of Figure 58 a), but it would be interesting to compare similar IL amount with different morphologies to highlight the effect of this latter.
<figure><img src="image_62.png" /><figcaption>Figure 58. a) Evolution of the amplitude of displacement for an actuator made of MAM M53 and [EmPy][Tf₂N] as IL with various content under 2V stimulus, b) Ionic conductivity of BCP MAM M53/IL ([EmPy][Tf₂N] at different weight concentrations: 23%, 34% and 45% versus temperature</figcaption></figure>

Finally, Figure 59 presents the deformation responses to a +/- 35V tension applied to an actuator made of MAM M53 in the presence of an IL ([EmPy][Tf₂N]) (28 wt%). At voltages higher than 15V, a damping is thus observed and could be due to an electrochemical instability of the IL. A limit of voltage is thus attained and other strategies are required to develop a higher deformation. The stability of the IL is also in general at stake and the stability and reproducibility of the displacement over time.

<figure><img src="image_63.png" /><figcaption>Figure 59. Actuator based on MAM M53 and [EmPy][Tf₂N] (28wt%) under a +/- 35 V stimulus.</figcaption></figure>

a) Typical curve representing the square function of the tension (V) versus time applied to the actuators; b) displacement of actuator versus time following the tension shown in a)

A modification of the IL was thus attempted with the use of a different IL, [EMIM][Tf₂N] at 33wt%. The displacement was followed under a lower voltage of only 5V, Figure 60. At this low voltage, no degradation of the system is observed over the course of 30 cycles of 120 sec (the first 10 are represented in Figure 60).
<figure><img src="image_64.png" /><figcaption>Figure 60. Displacement of MAM M53/IL ([EMIM][Tf2N]) with 33wt% under a square tension of +/- 5V, b) Zoom on three cycles of tension.</figcaption></figure>

It can be noted also that the amplitude of this combination attained larger displacement than [EmPy][Tf₂N] combination (0.6 mm versus 0.2 mm) which highlights the importance of the IL on the electrochemical behaviour of the actuator. Further investigation of the size of the ions and hydrophilicity of the IL on these properties.

## V.4. Conclusion

The impact of the addition of ionic liquids on block copolymers has been investigated. In the case of MAM BCP, the incorporation of IL have been shown to be specific towards the more polar phase of the BCP. A significant decrease of Tg of the PMMA phase was observed while no effect could be detected on PBuA phase until 44wt% of IL. Then, the evolution of the morphology and the mechanical properties in relation with the morphology were characterized and allowed us to obtain different morphologies for the same BCP composition. Specifically, a gyroïd morphology has been obtained within a range of 13 to 23wt% of IL ([EmPy][Tf₂N]).

Then electromechanical behaviour was studied by placing the BCP/IL films between gold electrodes. The influence of several parameters (IL nature and content, voltage) has been investigated. The electrochemical instability of the IL has been highlighted in presence of high voltages (higher than 15V) by a gradual decrease of the amplitude of deformation of the actuator. Below this voltage, the stability and reproducibility of the displacement has been shown and the influence of the IL on the amplitude of displacement was also depicted.

Finally, a more complete study related to these preliminary experiments is necessary to complete these promising results. One could think in particular of varying more exhaustively the nature of the charges of the IL, the difference in size between the anions and the cations, the nature of the electrodes, their thickness, the geometry of the actuator, etc. Also, the BCP used is problematic for actuator application as IL is incorporated in the “hard” phase of the thermoplastic elastomer (TPE) where the IL mobility is highly limited and the deformation starts to be possible as soon as the rigidity of this block is sufficiently decreased. Another type of polymer with a “soft” phase that can be more polar than the “hard” phase which will thus be the host of the IL and allows it to move more freely.
VI. Conclusions and Research project

Following different strategies, the control of the micro/nanostructure of mix of polymers was realized in order to improve different properties. This control was achieved through processing conditions (multinolayer coextrusion, coating) or through the use of block copolymers.

First, the use of multinolayer coextrusion on different polymer couples allowed us to analyse its impact on mechanical and gas barrier properties. Confinement effect on crystallization of EVOH44 was studied and an evolution of the gas barrier properties was observed in humid conditions. A hard confinement was shown to be more efficient than a soft one even if this latter condition allowed also an improvement in gas barrier properties. An improvement of a factor of 10 in oxygen permeability was obtained in a couple EVOH44/EVOH29 (the hard confinement) with an individual thickness of the EVOH44 layer close to 30 nm at 75%RH while a factor around 3 was obtained for LLDPE/EVOH44 couple in similar conditions. The explanation of this improvement is however still at stake even if the importance of the thickness of the layers was evidenced by the use of different conditions of layer multiplication.

Then this multinolayer coextrusion technique was also used to improve the mechanical properties of a polymer couple. The one studied was ABS/PC, polymers that may emanate from the WEEE. Here, the effect of the morphology on the properties was observed after an injection step that will destroy partially the morphology developed in the first coextrusion step. It was observed that when a non-spherical morphology of PC in ABS was obtained, an improvement in impact strength was obtained. In this case, starting from the largest oriented PC domains (i.e. simple three layer structure) will lead to the best properties due to the fact that this morphology will be less lost during the injection due to its initial size. A chemical modification of the ABS to improve its compatibility with PC in order to retain the morphology had the opposite effect by accelerating the return to a nodular form of the minor domain. Another way that is currently investigated is the use of the ABS/PC directly after the extrusion without the injection step in order to characterize the effect of an oriented nanostructure on the impact strength of the material.

Other processing methods have been employed: coating and monoextrusion of polymer blends. First, coatings of high gas barrier polymers have been realized and their combination with lamellar fillers, i.e. vermiculite. Very high gas barrier films were then developed with the use of bioplastics such as chitosan or PVOH with high loading rates of vermiculite (up to 60 wt%). The main drawback of these systems is their sensitivity to moisture. A chemical grafting of fatty acids on chitosan was attempted and has permitted an improvement of the gas barrier properties of a modified chitosan/vermiculite (60 wt%) coating of a factor close to 10 at 75 %RH.

For the monoextrusion of polymer blends, a mix of bioplastics was studied for the improvement of their mechanical properties. A blend of PLA with PHBV additivated with oligomers of PLA (oLA) was realized and the effect of the plasticizer was characterized. First, the effect of the additive on the size of the minor phase (PHBV) was investigated and then the impact on mechanical properties was explained with an improvement in elongation at break and Young's modulus for these blends with addition of very small amount of oLA from 0.1 to 1 wt%.

Finally, the use of block copolymer and its combination with a small molecule to modify its morphology was attempted. The evolution of the self-assembled morphologies obtained with a triblock copolymer was studied in relation with the amount of ionic liquid specifically incorporated in one of the two phases. Disorder-order and order-order transitions were observed and specific gyroïd morphology could be determined for the combination of a MAM M53 with the addition of 13 to 23 wt% of [EmPy][Tf₂N]. The evolution of the mechanical properties in tensile strength, of conductivity and of electromechanical properties were investigated with the amount of IL. Actuator were realized using gold leaves as electrodes and a deformation of the actuator was observed with voltages as low as 2V. The electrochemical stability of the IL has been identified as a limiting parameter for the range of potentially useful voltages. The chemical nature of the IL was also identified as a determining parameter for the amplitude of displacement.
In relation with the different scientific projects that I've been involved in, two main scientific questions arise and will lead my future research and projects. The first scientific lock that I'm intending to tackle is the influence of humidity on nanostructured system and the link between barrier properties and hydrophilicity/hydrophobicity of polymers and composites. The second scientific lock is focused on the link between nanostructure and mechanical or electromechanical properties. I'll explain more thoroughly the different questions and ways to answer them in the following themes.

## Theme 1: Effect of humidity on the barrier properties of structured polymer materials

During my previous work, I was able to produce a number of solutions to improve the gas barrier properties of plastic materials with a view to producing more recyclable and therefore environmentally friendly packaging. One of the strong blocking points of these systems remains their sensitivity to humidity, especially during industrial processes such as pasteurization which leads to the presence of humidity within the gas barrier polymer layers and will deteriorate this property. Nevertheless, I was able to observe during my previous work changes in these properties with humidity according to several parameters such as the structuring at the nanometric scale or the control of the hydrophilic/hydrophobic balance of the material. These observations unfortunately could not yet be quantified properly and I would like to focus part of my work on these phenomena. For this I will develop two approaches related to these two parameters. The first approach will focus on nanostructuring and multi-nanolayer films based on ethylene/vinyl alcohol (EVOH) copolymers. Achieving a nanometric thickness for the EVOH layers has a direct impact on the barrier properties in the presence of humidity whereas it does not have any in the dry state. The production of films with various proportions of EVOH, of various natures of this copolymer (Ethylene/vinyl alcohol ratio) and with other types of polymer (pure poly(vinyl alcohol) for example) will make it possible to study the impact of network of hydrogen bonds of the system and its confinement on the barrier properties in humid conditions. This network will be studied using FTIR but also Dynamic Vapor Sorption (DVS) to enhance the difference of absorption of water in presence of confined EVOH. Solid-state NMR might also be considered.

This work will be a continuation of the collaboration already established with the IMP laboratory in Lyon and in particular Eliane Espuche, Fabrice Gouanvé and René Fulchiron. Industrial collaborations are also in place with, in particular, the company Leygatech, one of the leaders in the field of polyolefin-based food packaging. A PSPC project is currently on going on the problematic of EVOH gas barrier improvement.

The second approach concerns the hydrophilic/hydrophobic balance of gas barrier materials. Indeed, the majority of polymers used in food packaging are highly hydrophilic materials, characterized by the presence of a strong network of hydrogen bonds which is the basis of these permeability properties. Indeed, this network strongly limits the potential movements within the material and the possibilities for gases to move inside it and therefore their overall permeability. However, this network is also extremely sensitive to humidity due to the nature of the interactions. In order to control the impact of humidity, playing on the hydrophilic character of the material could make it possible to obtain the right balance between barrier properties and water absorption. For this, chemical modification by grafting onto the alcohol or amine functions present in the polymers of this category will be considered. This will be particularly useful on natural polymers with a high proportion of these functions and therefore a very high sensitivity to humidity. The use of polymers derived from algae such as alginates produced by the company Algaia will be considered in particular to meet several challenges in parallel: the reduction in the use of non-recyclable materials, the impact of production of the raw material and the end of life of these materials. Moreover, the mixture of these materials with lamellar fillers could make it possible to further improve these barrier properties, but these fillers also present problems of moisture absorption. The chemical modification of these fillers could also be envisaged. Finally, it will be interesting to see the impact of this hydrophilic/hydrophobic balance
on systems used in the form of coatings (in particular natural polymers which are complicated to implement other than by going into solution) but also more coextruded systems.

In this perspective, two main scientific questions will be addressed:

- First, the coating process of aqueous solution of a composite on a highly apolar polymer could be problematic and the impact of the process (deposition and drying steps) on the dispersion of the nanofillers will be studied.

- Second, the effect of the chemical nature of the coating and the nature and the impact of the chemical modification, such as the grafting ratio of fatty acids, on the moisture sensitivity for gas barrier properties will be studied.

For the first issue, the process will then be the main case of study. The expected coatings will be made of composites with the presence of nanofillers and the impact of the process on the dispersion and the organization of the nanofillers needs to be clarified. Indeed, as presented in this manuscript, the dispersion of the nanofillers was not easily obtained even when starting with dispersed solutions (Figure 29 b)¹⁴⁶.

The control of viscosity and shear imposed by the process will therefore be studied as well as the evolution of the dispersion during the drying step. For this purpose, rheological and morphological analyses will be performed using Couette apparatus, SEM, TEM and optical microscopies as well as X-ray diffraction measurements¹⁴⁷.

Indeed, the coating process imposes shear and elongation of the aqueous polymer. The understanding and quantification of the viscoelastic stresses need to be modelled, so that the limits of flow are known both in continuous shear, oscillatory shear and elongation. Ideally, the quantification should be known for different temperatures. To accomplish this, rheological analyses will be carried out using a Couette apparatus and a Sentmanat extensional accessory in a temperature-controlled oven. The results will be analysed and modelled in order to propose the appropriate constitutive equations for possible numerical simulations.

This understanding will lead to the process conditions that will allow controlling the coating thickness and the nanofiller dispersion.

Then for the second issue, the choice of polymers for the coatings will be guided by their gas barrier properties but also their availability on an industrial scale. Different chemical structures are considered for this problematic. Firstly, PVOH is a well-known and readily available polymer and has already been studied in this manuscript however without any chemical modification. That's making it a ready-to-use material and with alcohol functions that can be easily modified chemically. Secondly, chitosan (CS) which has already been studied as well has shown interesting improvement in humid environment with a chemical modification. However solely one condition of grafting has been investigated and different percentage of modification and different fatty acids may be considered to investigate the link between hydrophilicity/hydrophobility balance and gas barrier properties in humid conditions.

Thirdly, another bio-based material will be studied. Alginate is a polysaccharide like chitosan and is produced by using biomass from French coast (through seaweed). Their structure is close to the one of the CS and similar barrier properties can be expected due to strong hydrogen bonds and ionic interactions¹⁴⁸,¹⁴⁹. They are classically used as edible coatings on fruits and vegetables¹⁵⁰ but their use as gas barrier coatings on polymer films is less observed. They are supplied by a French company named Algaia specialized in products emanating from seaweed. They supply amounts of alginate in an industrial scale with a fine physico-chemical analysis of their products which will be necessary for a large development of these coatings at a large scale in the food packaging industry.

Bonding unsaturated fatty acids, or other less polar substitutes, to this chemical structure will change the solubility parameter of these polymers, which could lead to a decrease in water sensitivity as it will decrease the overall polarity of the polymer¹⁵¹. Another way of modifying the solubility parameter will be to use chitin and poly(vinyl acetate) which are the precursors of CS and PVOH before the deacetylation/hydrolysis step. In these polymers, the hydroxyl functions observed in CS and PVOH are replaced by ester functions which
are less polar. Similarly, alginate can be chemically modified¹⁵² to reduce its sensitivity to moisture when used as a thin film coating. The adhesion properties between layers are also a key issue that will require attention and will be studied in detail, especially when the polarity is changed.

The scientific objectives of this issue will therefore be:

- Establish the link between the hydrophobicity/hydrophilicity through the solubility parameter of polymers with oxygen permeability in the presence of moisture.

- Understand the influence of the chemical structure on barrier properties beyond the solubility parameter.

- Establish the link between this hydrophilicity and the solvent used for coating process and the dispersion of the nanofillers → link with the industrial applications.

- Influence of the solubility parameters on the final quality of the coating, especially concerning the adhesion level with the apolar substrate and on the dispersion of the nanofillers in correspondence with the optimization of the parameters of the first objective.

Finally, the development of these coatings as a proper alternative will require different steps:

- The development of a laboratory scale coating line to assess the scalability of the process and the first tests on an industrial line capable of drying aqueous coating solutions.

- The ageing of the coatings will have to be studied in order to predict the lifetime of the product and particularly the evolution of the barrier properties over time

- To verify the recyclability of the coated films by mechanical recycling and characterization of the mechanical properties of the recycled products

- To ensure the suitability of these coatings for food contact

Indeed, the industrialization is particularly at stake as the aqueous coatings are usually complicated to develop. Moreover, the end-of-life of the developed solutions needs to be envisaged and will allow us to indicate a criterion of nanofillers load and thickness of coatings on the final properties of the mechanically recycled product. Finally, the suitability for food contact is always an important factor when it comes to chemicals likely to be in contact with food.

To manage this approach on hydrophobicity/hydrophilicity balance and nanostructuration on permeability behaviour under relative humidity, partnerships have already been set up with the Joint Technology Network (RMT) ProPackFood in support. A collaboration with the LNE (Catherine Loriot, Elise Delbos, Phuong-Mai Nguyen) and the PAM laboratory in Dijon (Thomas Karbowiak, Frédéric Debeaufort) is being set up. A packaging manufacturer (Wipak) is also a partner in this project. An ANR proposal has been first dismissed but further development are under way especially following the AMI "Alimentarité".
Theme 2: Improvement of mechanical and electromechanical
properties by polymer blend structuration

The second theme of my research project is also related to structuration but will focus on the impact that
this structuration can have on mechanical and electromechanical properties. The structuration, whether in
terms of size of the observed domains or in terms of orientation of these domains, brings about a variation
in the behavior of macromolecules, in particular when these chains are found confined in small spaces. The
evolution of the glass transition temperature or crystal shapes and orientations in confining situations have
been extensively studied in the literature<sup>153-155</sup>. My research project focuses on developing structuring
strategies for different types of application. The first targeted application is related to the recycling of
thermoplastic materials. Indeed, it has been demonstrated that it is possible to improve mechanical
properties, in particular impact strength, by combining polymer blends with the minority phase placed in
oriented domains, such as platelets. For this, the use of the multinolayer coextrusion technique can be
further considered. However, the resulting problem of these systems is the obtention of a mixture of
polymers which is troublesome in a perspective of recycling a posteriori these materials. The idea here is to
use recycled material, which is a mixture from the start, as the initial material. In the case of Waste Electrical
and Electronic Equipment (WEEE), the mixture of acrylonitrile-styrene-butadiene (ABS) copolymer with
polycarbonate (PC) is conventionally used and ends up after the waste sorting step in the waste stream of
pure ABS. This flow therefore has degraded properties compared to either a pure ABS or an ABS/PC
mixture whose proportions are well controlled. Multinolayer coextrusion could make it possible to
consider this recycled ABS flow as one of the two layers used and to add PC as the second layer. By greatly
increasing the number of layers to be produced, the layers of virgin PC will eventually break and give nano-
domains of PC included in the recycled ABS/PC mixture. These nano-domains will also exhibit orientation.
The impact of the size and orientation of these domains, the amount of PC in the recycled ABS stream as
well as the state of degradation of the ABS will have to be studied in order to determine the possibility of
using this methodology to improve the impact strength properties of this material and ultimately to obtain
a possibility of upcycling ABS by this methodology of structuration. In general, a mechanical modeling of
the system could make it possible to understand the impact of the size of the domains on the mechanical
properties. Indeed, the understanding of the effect of the shape and size of the minor phase on mechanical
behavior and especially the crack propagation phenomenon will be carried out. This will lead to create a
model of the system that will allow us to understand the influence of the morphology on these properties
and should allow optimizing this shape and dimensions to reach the highest impact strength of the recycled
material to enhance its recyclability.

The aim will then be to create specific blend morphologies (nodular, fibrillar or lamellar at micro and nanoscale) to get upgraded mechanical properties and study the effect of the size of the morphology of the minor phase on it. Several processing tools and approaches will be used and combined to obtain and stabilize elongated morphologies (1D-fibrils or 2D-lamellas) with the highest aspect ratio, or nodular morphologies with the smallest dimensions. Dry-blending approach in a classical single-screw extruder, with optimized processing conditions (temperature, residence time, shear rate) and various physico-chemical properties of the blend, will be for example investigated as well. As already presente the multilayer coextrusion technique is proposed to produce layered structured and nanoblends with the breakage of the layers. Indeed below a critical thickness, these layers may eventually break up⁵ and form thin elongated nano-lamellae. In this approach, multinolayer coextrusion is combined with dry-blending in single-screw extruder and is expected to act like static mixers, generating strong successive shearing and elongational flow achieving new nano-blend morphologies. The impact of layers interactions on the development and on the stabilization of the morphology will be thoroughly investigated.

The development of these methodologies requires collaborations for the mechanical modelling of the blends
behaviour. Collaborations in this direction with the MSME laboratory of Marne-la-Vallée (Fabrice Detrez)
and the LGMC of Mines d'Alès (Anne-Sophie Caro-Bretelle, Patrick Ienny and Didier Perrin) are under
construction to be able to provide this modeling of mechanical properties in function of
micro/nanostructure through the deposition of ANR proposal as well as a proposal for the DIM Materre.
The second application concerns electromechanical approaches and in particular the production of actuators based on block copolymers and ionic liquids. In this context, the structuration is provided by the self-assembly mechanism of block copolymers which leads to the creation of domains of nanometric size with an organization directed by the affinity of the two blocks, the length of each of these blocks and the total molar mass of the copolymer. Here, structuration is used to obtain a good compromise between mechanical properties and ionic conduction properties when an ionic liquid is added to the system. Indeed, it has been shown that an ionic liquid can mix in a specific way with the most polar phase of the copolymer. It is thus possible to obtain an ionic conductive phase with the phase receiving this ionic liquid and an unaffected phase. Ion conduction allows the realization of electroactive actuators with the migration of ions towards electrodes placed on each side of the material. However, depending on the mechanical properties of the polymers involved in the copolymer, the addition of the ionic liquid can collapse the overall Young's modulus of the material and therefore its rigidity, which will greatly limits the actuation force that can be applied. The choice of copolymers, their chemical nature, their respective ratio and the overall molar mass, is therefore a key parameter to obtain an actuator that can present the best electromechanical properties. The ionic liquid and in particular its incorporated amount, will also affect the morphology created in the system, a parameter which thus adds to those previously mentioned. A compromise between mechanical property and conduction property must be found to maximize the possible deformation amplitude as well as the applicable actuation force and this will be parameterized by the morphology (cylindrical, lamellar, co-continuous) created as well as by the quantity of ions and their mobility in the system.

The chemical nature of the block copolymer has also to be carefully chosen to improve the mechanical stability of the actuator while having the highest ionic conduction. An ABA terpolymer with hard A blocks non polar and a soft B block more polar to host the IL will have to be considered. The aim is to allow a high mobility of the ions even at low IL content with a rubbery polymeric phase at room temperature while the A blocks are not affected by the addition of IL and so retain the thermoplastic elastomer behaviour of the terpolymer. The chemical nature and also the global BCP composition will then influence these parameters and will be studied.

The first scientific questions that will be investigated can thus be summarized as:

- What formulations (BCP composition and molar mass, type of IL, IL content) can lead to co-continuous morphologies (gyroid or perforated lamellae (HPL))? The evolution of the BCP phase diagram in function of the IL content will be examined.

- What is the evolution of the mechanical properties of the previous formulations in flexion or tension?

- What formulation and what morphology will lead to the highest ionic conductivity value? Also, what is the influence of the incorporation of an IL in a polymeric phase on its electrochemical stability?

These questions should lead to one or more formulations which will be electromechanically characterized and then used to design actuators. Then the following problem will be tackled: how to electromechanically model these polymers and design experiments allowing the understanding of the underlying physical mechanism and accessing their relevant electromechanical constants?

To respond to this scientific obstacle, a collaboration is set up with a polymer synthesis team from the ICGM laboratory (Vincent Lademiral, Patrick Lacroix-Desmazes, Christine Joly-Duhamel, Gilles Silly), a team from the IMP laboratory dealing with conduction measurements ionics (Anatoli Serghei, Laurent David) and members of the PIMM laboratory of the DISCOh team (Marc Rébillat, Pierre Margerit) specialists in electromechanical systems and mechanical modelling. An ANR proposal is currently under review on this topic with this team.
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