Prof Patrice Bourson

Dr David Chapron

Prof Alonso Carlos

Prof Moreno

Bravo Ivan

Prof Perez

Antonio Jose

Castro Osma

Enrique Murillo

Victor, Sergio Rosario Fernando

Prof Manuel Piñeiro

Dr Carolina Hermida

Maite, Inma, Juancar Maria

Cold-crystallization enthalpy ΔHm Melting enthalpy TABLE OF CONTENTS Chapter 1. Introduction

If you want to go fast, go alone. If you want to go far, go together" Certainly, there are many people involved in this thesis, to whom I am more grateful than I can herein express.

First, I would like to thank my supervisors for their continuous support and help during these three years. Thanks to

PREFACE

Since Berzelius coined the word polymeric in 1833, the study of the behaviour of such "long molecules" brought about new insights such as the vulcanization of caoutchouc (Goodyear, 1839) and the synthesis of nitrocellulose (Schönbein, 1845) or the celluloid (Brothers Hyatt, 1869), setting the onset of plastics revolution. Since then, plastics have burst into our daily life due to their versatility, cheapness, lightness, and excellent thermal properties, with packaging use accounting for the largest plastic sector on the market (39,7%).

Moreover, the application of synthetic polymers in medicine was firstly reported during the Second World War, when Nicholas Harold used poly(methyl methacrylate) (PMMA) as an artificial corneal substitute, which was followed by the employ of polyethylene terephthalate (PET) for vascular grafts, polydimethylsiloxane (PDMS) for breast implants, and polyethylene (PE) for hip joint replacements. The design of novel polymeric materials must fulfil the environmental and societal demands to diminish the carbon footprint required to adopt manufacturing strategies that meet the European policies such as "A European Strategy for Plastics in a Circular Economy", which was launched in 2018 to address how plastics are designed, used, and recycled. Likewise, the European Parliament has recently recognised the potential role of bioplastics and compostable plastics in the circular economy and sustainability. Consequently, bio-based plastics synthesised from renewable resources such as Poly-(lactic acid) (PLA) are promising environmentally friendly candidates for the development of biomedical and pharmaceutical applications whilst contributing to the circular economy. However, the clinical application of PLA-based materials is limited despite the research abundance on both its structural characterization and potential applications which might be due to the lack of know-how transfer between the fundamental research and clinical trials.

In this context, the University of Castilla-La Mancha (Spain) and the University of Lorraine (France), as well as different collaborators in Spain, France and the Netherlands (Figure 1), developed a global project with an interdisciplinary approach consisting of the synthesis in a controlled manner of biocompatible polymers obtained from renewable sources and the establishment of the structure-processing-properties relationship to tailor polymeric materials from the genesis for biomedical applications aiming at the end-user requirements. This thesis, focused on PLA and particularly on stereocomplex PLA (SC-PLA), is the foundation stone of this project.

The state-of-the-art of PLA as a biomaterial for its medical application is reviewed from a global approach by highlighting the interconnections among the architectural designing parameters with the desired applications. Moreover, special emphasis has been placed on the stereocomplex phase and the problems encountered so far to obtain pure stereocomplex phase (Chapter 1).

Moreover, the synthesis of the long desired-HMw stereo-diblock copolymer of PLA (PLA-SBC) is described by using a Zn-based heteroscorpionate catalyst synthesized by the Inorganic group of the University of Castilla-La Mancha (UCLM), in Spain. The living character of the catalyst selected as well as the high sterochemical control achieved enable the customisation of block copolymers of PLA (tri-, tetra-, penta-block copolymers) as well as hetero-block copolymers in combination with polymers presenting complementary physicochemical properties to PLA such as Poly-caprolactone (PCL), Poly-etilenglicol (PEG), etc. However, in this chapter, the synthesis of symmetric block copolymers exhibiting different molecular weight has been established as the main objective (Chapter 2).

In addition, the crystallization mechanism from the melt of the PLA-SBC under different thermal conditions is studied at different length scales, in collaboration with the beamline 26 of the European Synchrotron Radiation Facility (ESRF), in France. Likewise, the crystallization mechanism of the PLA-SBC blended with different nucleating agents (NA) is also analysed in collaboration with the University of Maastricht, as a strategy to overcome one of the main drawback of PLA, the low crystallization rate (Chapter 3).

The selection of a trace element-based catalyst to accomplish the synthesis of PLA stereoblock copolymers was aimed to apply the PLA-derivatives in the biomedical field. During my studies of pharmacy, I was introduced to the world of nanocarriers through the collaboration with the Neuroscience research group of the faculty of Pharmacy at the UCLM.

Moreover, there is a gap between material research, preclinical test, clinical requirements and regulatory aspects that hamper the translation of PLA nanoparticles research to the clinical level. Likewise, gelatin-based drug delivery systems (DDS) containing nanoparticles formulated from PLA-derivatives presenting different stereoregularities are designed as well as the release rate and local applicability of the bionanocomposites, in collaboration with the Universities of Vigo and Castilla-La Mancha, in Spain (Chapter 4).

ABSTRACT

The pharmaceutical industry demands constantly the development of drug delivery systems that improve therapeutic efficacy and reduce side effects. The traditional strategy of manufacturing standardised drug doses is ineffective for 70% of patients, creating the need to replace mass production with personalised medicine. In particular, nanoparticles (NP) and generally from polymeric sources were largely investigated as drug nanocarriers to improve the stability and solubility of the encapsulated molecules. In addition, NP formulations improve safety and efficacy. However, the current gap between the characterisation of the physicochemical properties of polymeric NP and their performance under physiological conditions hampers the translation from research to the clinical level. Typically, Poly-(lactic acid) (PLA) is an aliphatic and synthetic polyester obtained from renewable sources that was extensively employed to generate NPs as drug carriers due to its bioresorbable nature related to its hydrolysis in physiological media into non-toxic products.

Likewise, the current thesis aims to establish from a global approach the design of tailored PLA block-copolymers with defined nanostructures to target materials with specific structure-processing-properties relationship that enable to obtain personalised biomedical and pharmaceutical solutions from the genesis aiming at the end-user requirements.

The rich PLA nanostructure together with the polymorphic nature offers the possibility to twitch the desired physicochemical properties. Particularly, the stereocomplex (SC) phase has been largely aimed due to the exhibited promising physicochemical properties featuring higher melting temperature, mechanical performance and hydrolysis resistance compared to the enantiomeric counterparts. Likewise, biomedical and pharmaceutical applications have been extensively designed with SC-PLA such as drug delivery systems. However, the typical SC crystallization by the blend of enantiomer (PLLA and PDLA) derivatives diminishes for high molecular weight (HMw) PLA and enantiomeric homocrystals (HC) are obtained instead.

In addition, the synthesis of PLA stereo-block copolymers (PLA-SBC) exhibiting HMw and high stereocontrol requires catalysts with stereocontrol for both enantiomers that are extremely rare.

Herein, a novel series of PLA derivatives comprising different nanostructures were synthesized by Ring-Opening Polymerization through a novel biocompatible heteroscorpionate catalyst with high control of the chain stereochemistry. Particularly, the long-desired HMw PLA-SBC presenting high tacticity has been synthesized for the first time without the use of a co-initiator. Moreover, the analysis of the mechanism of crystallization of the PLA-SBC proved the prevalence of the SC-phase, avoiding the typical HC-and SC-phase separation that occurs in the enantiomeric blends. Furthermore, the PLA industrial drawback related to its slow crystallization rate has been tackled by the addition of an organic nucleating agent that increases the crystallization rate of PLA-SBC whilst preserving its biocompatibility with a significant impact on the new processing technologies. NP-PLA with different nanostructures were then formulated with 2 antitumoral compounds. However, the NP-PLA so generated exhibited an amorphous structure upon formulation, erasing the designed nanostructures and their different release rate. Therefore, a straightforward synthetic approach was designed to generate bionanocomposites of gelating hydrogels with embedded NP-PLA with the aimed crystal phase and degree of crystallinity by its post-formulation thermal annealing step. Likewise, the gelatin hydrogel network acts thereby as a mechanical sustain for both the storage and local administration of the therapeutic agent as well as avoiding the aggregation of nanoparticles upon annealing whilst offering support to complementary drugs with different profile releases.

Graphical ABSTRACT RESUME

L'industrie pharmaceutique exige constamment le développement de systèmes d'administration de médicaments qui améliorent l'efficacité thérapeutique et réduisent les effets secondaires. La stratégie traditionnelle consistant à fabriquer des médicaments à dose médicamenteuse standardisée est inefficace pour 70 % des patients, créant le besoin de substituer la production en masse par une médecine personnalisée. En plus, les nanoparticules (NP) formulées normalement à partir de polymères permettent d'améliorer la stabilité et la solubilité des molécules encapsulées, ce qui améliore la sécurité et leur efficacité. Néanmoins, il existe un écart entre la caractérisation des propriétés physico-chimiques de NP polymériques et ses performances dans des conditions physiologiques qui limite l'application au niveau clinique. L'acide polylactique (PLA) est un polyester synthétique et biosourcé qui présente des caractéristiques biorésorbables grâce à son hydrolyse en milieu physiologique en produits non toxiques.

Cette thèse vise à établir les relations structure-traitement-propriétés des matériaux polymères, et en particulier du PLA stéréocomplexe (SC) à blocs, pour concevoir des solutions biomédicales personnalisées depuis la synthèse jusqu'aux exigences du patient final.

La richesse de la nanostructure du PLA associée à sa nature polymorphe, offrent la possibilité de modifier les propriétés physicochimiques en fonction de l'application souhaitée.

En particulier, la phase SC est une solution en raison des propriétés physico-chimiques prometteuses qu'elle présente, notamment une température de fusion, des performances mécaniques et une résistance à l'hydrolyse plus élevées que celles de ses homologues énantiomères. De même, les applications biomédicales et pharmaceutiques ont été largement conçues avec des PLA SC, comme les systèmes d'administration de médicaments. Cependant, la cristallisation typique des SC par le mélange de dérivés énantiomères (PLLA et PDLA) diminue pour les PLA à haut poids moléculaire (HMw) et des homocristaux énantiomères (HC) sont obtenus à la place. En outre, la synthèse de copolymères stéréoblocs de PLA (PLA-SBC) présentant une masse moléculaire élevée et un contrôle stéréochimique élevé nécessite des catalyseurs avec un contrôle stéréochimique pour les deux énantiomères, qui sont extrêmement rares.

Dans ce travail, une nouvelle série de dérivés de PLA comprenant différentes nanostructures a été synthétisée par polymérisation par ouverture de cycle à l'aide d'un nouveau catalyseur hétéroscorpionate biocompatible avec un contrôle élevé de la stéréochimie. En particulier, le HMw PLA-SBC présentant une tacticité élevée, a été synthétisé pour la première fois sans l'utilisation d'un coinitiateur. De plus, l'analyse du mécanisme de cristallisation du PLA-SBC à différentes échelles a prouvé la prévalence de la phase SC, évitant la séparation des phases HC et SC qui se produit dans les mélanges d'énantiomères. En outre, l'inconvénient industriel du PLA lié à sa vitesse de cristallisation lente a été résolu par l'ajout d'un agent de nucléation organique qui augmente la vitesse de cristallisation du PLA-SBC tout en préservant sa biocompatibilité.

Des nanoparticules de PLA (NP-PLA) avec différentes nanostructures ont ensuite été formulées avec 2 composés anti-tumoraux. Cependant, les NP-PLA ainsi générées présentaient une structure amorphe lors de la formulation, effaçant les nanostructures conçues. Par conséquent, une approche simple a permis de générer des bionanocomposites d'hydrogels gélifiants avec des NP-PLA. Le taux de cristallinité a été ensuite contrôlé par une étape de recuit thermique après formulation. De même, la matrice d'hydrogel de gélatine agit ainsi comme un support mécanique pour le stockage et l'administration locale de l'agent thérapeutique, tout en évitant l'agrégation des NP lors du recuit et en offrant un support à des médicaments complémentaires ayant des profils de libération différents.

INTRODUCTION

In this Chapter the state of the art of the synthesis, processing, and properties of PLA is summarized, as well as its application in the biomedical and pharmaceutical field.

Particularly, the influence of PLA synthesis and processing on its nanostructure is emphasised to correlate it with its thermo-mechanical properties that in turn, will define the suitability of the generated end-product for the desired application (Scheme 1.1).

Scheme 1.1 Summary of the state-of-the-art of PLA as a material for its application in the biomedical and pharmaceutical field, considering the correlation of the synthesis-processingstructure-properties relationship.

Introduction

Poly-(lactic acid) or Polylactide (PLA), classified as an aliphatic polyester, is the biopolymer most employed in industrial fields such as fibre manufacturing, food handling, or textile industry [1] and it accounts for 24% of the global production capacity for biodegradable polymers [2]. PLA is derived from renewable sources such as corn sugar, potato, and sugar cane [3], is compostable [4], and its properties are comparable to those of petroleum-based polymers, such as polyethylene, polypropylene, polystyrene, polycarbonate and polyethylene terephthalate [5]. Furthermore, PLA hydrolyses in physiological media, yielding lactic acid, a non-toxic component that is eliminated via Krebs cycle as water and carbon dioxide [6].

Moreover, PLA is a polymorphic polyester comprising a chiral carbon in its structure that yields a wide variety of nanostructured derivatives (isotactic, atactic, syndiotactic, stereocomplexed, etc.) [7][8][9] featuring different physicochemical properties. Particularly, the stereocomplex (SC) phase, which is formed upon blending enantiomeric poly(L-lactide) (PLLA) and poly(D-lactide) (PDLA) or by synthesizing stereo-block copolymers (PLLA-b-PDLA), exhibits higher melting temperature, mechanical performance and hydrolysis resistance compared to the enantiomeric counterparts, and has been extensively investigated for biomedical and pharmaceutical applications, as well as commodity and industrial applications [7,10]. However, the SC crystallization of the blended enantiomers (PLLA and PDLA) diminishes for high molecular weight (HMw) PLA and enantiomeric homocrystals (HC) are obtained instead [11]. In addition, the synthesis of PLA stereo-block copolymers (PLA-SBC) exhibiting HMw and high stereocontrol requires catalyst extremely rare and has only been reported twice [12,13].

The biocompatibility, biodegradability, and resorbability characteristics of PLA have promoted its use in the biomedical field for a wide range of applications (suture threads, bone fixation screws, drug delivery systems, etc.), offering an alternative to conventional biocompatible materials such as metals and ceramics [14]. In addition, the ability to tailor the mechanical, thermal, and degradation properties of PLA derivatives due to the range of afforded nanostructures depending on the chemical architecture and processing conditions allows for designing personalised medical solutions. Indeed, the novel available synthetic PLA approaches to generate multiblock copolymers as well as the advent of current processing technologies broaden its suitability to advance into the customisation of the generated endproducts to minimise the side effects. In particular, several highly cutting-edge PLA-based therapeutic applications have recently attained the clinical phase, such as drug-eluting stents [15][16][17][18] as well as personalised pharmaceutical agents that were designed from an interdisciplinary approach to avoid the serendipitous progress, emphasising the beneficial interaction between the materials and the biomedical fields. Likewise, the frequent dismissed PLA nanostructure of the designed biomedical solutions hamper the systematic advance by correlating the structure-property relationship of the system with the application performance. Several reviews about PLA have been published covering a wide range of topics, such as the PLA synthesis [9,19], the physicochemical and mechanical properties of PLA [1],

the crystallisation and structure-properties relationship of PLA [20], the characteristics of the promising stereocomplex PLA phase [7,10], and PLA applications in widespread fields [3,14,[21][22][23][24][25].

Herein, the state-of-the-art of PLA-based biomedical applications is reviewed from a global approach by highlighting the interconnections among the architectural designing parameters with the desired applications, with an emphasis on the stereocomplex phase of PLA.

Poly-(lactic acid)

Poly-(lactic acid) is a poly-α-hydroxy acid synthesized from lactic acid (LA; 2hydroxypropanoic acid), which exists in two optically active stereoisomers, namely L-LA and D-LA (S and R in absolute configuration, respectively) [26]. Approximately 90% of the total lactic acid produced worldwide is made by bacterial fermentation, which offers the advantages in both utilization of a renewable source and the production of optically pure Land D-lactic acid, depending on the strain selected (the chemical synthesis of lactic acid always results in a racemic mixture). The dehydrated cyclic dimer of lactic acid is commonly called lactide (3,6-dimethyl-1,4-dioxane-2,5-dione). Lactide exists in three different forms due to the presence of two asymmetric carbon atoms in the molecule: L-lactide, D-lactide, and mesolactide (Figure 1.1). In addition, a racemate of D-lactide and L-lactide exists as rac-lactide [26].

The polymerization of optically pure L-and D-lactide yields isotactic homopolymers of Poly-(L-lactide) (PLLA) and Poly-(D-Lactide) (PDLA) respectively (Figure 1.1). Both PLLA and PDLA are semicrystalline polymers, showing a melting temperature (Tm) around 170°C [8,9] and a thermal degradation temperature around 200°C [27]. The PLA derivative crystallinity as well as their melting and glass transition temperatures (Tg) usually decrease with diminishing optical purity of the lactate units [1,28]. PLLA polymers with D-lactide content lower than 10% tend to be crystalline (or PDLA with L-lactide content), whilst homopolymers with lower optical purity are amorphous [29]. Besides, the random insertion during the polymerization of D-and L-LA units of both rac-and meso-lactide monomers generates an atactic polymer, Rac-PLA, which is completely amorphous. Moreover, syndiotactic PLA is obtained when D-and L-lactic acid units are placed alternatively along the chain whereas heterotactic chain architecture is attained when D-and L-lactide units are inserted alternatively on the polymer chain [30]. Syndiotactic PLA is a semicrystalline polymer exhibiting higher crystallization temperature (Tc) than isotactic PLA but lower Tm, whilst heterotactic PLA is amorphous [31]. Furthermore, the one-pot sequential addition polymerization method [32] of D-and L-lactide monomers yields stereo-block copolymers with blocks of opposite chirality, featuring melting temperatures 50°C higher than isotactic homopolymers (220°C) [7].

Besides, PLA displays different crystalline phases (α, α', β, γ) established by the chain architecture and the specific crystallization mechanism or thermo-mechanical history imposed during its processing, which defines the properties of the final product [33] (Figure 1.2). The α-phase is the more stable PLA homo crystal structure that corresponds to an orthorhombic unit cell in which the helices are packed in a hexagonal fashion, containing two antiparallel chains per unit cell. The α-phase is normally obtained when isotactic PLLA or PDLA are crystallized from the melt above 130°C or by crystallization from solution, characterized by a melting temperature of around 170°C [34,35]. The α'-phase (or δ-phase) is the disordered form of the α-phase that is generated either from crystallization from the melt at temperatures below 110°C or by cold-crystallization after quenching PLA to the glassy state. The α'-phase is also organized in the orthorhombic crystallographic unit cell, however, it contains two parallel helices per unit cell which increases the lattice parameters when compared to the ordered α phase. Besides, a mixture of α-and α'-phases is obtained when PLA is crystallized between 110°C and 130°C, although the α'-phase recrystallizes into α-phase when heated near the Tm (150-160°C) [34]. Moreover, the β-phase is developed by stretching PLA fibres in the α-phase at high temperature (130°C-140°C) and high draw ratios as well as by casting thin films from solution [35,36]. The chain conformation of the β-phase is a threefold helix in a trigonal unit cell containing three chains per unit cell, and its Tm is ca. 10°C lower than α-phase Tm [37] [38].

In addition, the β-phase exhibits piezoelectricity that allows the interchange of mechanical and electrical energy [39], broadening its potential applications. Finally, the γ-phase is produced by epitaxial crystallization of PLA on hexamethylbenzene, forming two antiparallel threefold helices in an orthorhombic unit cell [40].

Furthermore, a new crystal structure, the stereocomplex (SC) phase, is formed from the co-crystallization of the two stereoisomers of PLA (PLLA and PDLA) that features a trigonal unit cell comprised of six threefold helices per unit cell. The structural peculiarity of the SC phase with the nearest neighbours of any stem being of a different polymeric chain provides them easy access to the growth front for both enantiomeric species. In addition, the specific C-H…O-H hydrogen bonds within the crystal lattice that stabilize the structure [41] confirms stereocomplexes with a higher melting (220°C) as well as degradation temperature (220-260°C) [27]. The SC phase was first found by casting a mixed solution of both enantiomers [42] and since then, the SC phase has typically been obtained intentionally from the blend of both enantiomers in solution (in an appropriate solvent such as dichloromethane or chloroform at room temperature, and acetonitrile around boiling temperature [43]) or in the solid-state from the melt [10]. However, the SC crystallization of the blended enantiomers diminishes for high molecular weight (HMw) PLA, and enantiomeric homocrystals (HC) in the α-phase are obtained instead [11]. Additionally, the critical Mw to exclusively obtain SC crystallization is lower for blends obtained from the melt than from solution [44] which hampers its industrial application. The optical purities of the polymers and the mixing ratio of the isomeric chains also affect the obtained ratio of SC to HC crystallites and thus, the preparation of pure SC-PLA requires meticulous specific conditions [45]. SC crystallites can also be generated through the synthesis of block copolymers by the one-pot sequential monomer addition to a truly living polymerization catalyst, which allows retaining the SC crystallization in HMw polymers [12,13].

The diverse synthetic approaches afford a wide range of chain architectures that can be generated through the different ratios of L-and D-Lactide monomers that offer the possibility to tailor the properties of the final polymeric product depending on the intended application.

Furthermore, understanding the advantages and drawbacks of the different synthetic processing methods to obtain PLA is crucial to tailoring the foreseen applications. Moreover, since PLA still exhibits performance drawbacks such as low mechanical properties, low thermal resistance and hydrophobicity that limit its applications in some biomedical fields, novel materials with unique properties can be obtained through the blend or copolymerization of PLA with other biodegradable or non-biodegradable polymers, such as polyethylene (PE), polypropylene (PP), Polyhydroxhyalkanoates (PHA), PMMA, Poly(ethylene-co-vinyl acetate) (PEVA), etc. [5,46,47]. In addition, nanocomposites can be fabricated by mixing PLA with other complementary compounds such as silk [48], gelatin [49], collagen [50], tungsten disulfide [51][52][53], natural fibres (flax, jute, hemp,) [54], ceramics (ZnO, TiO2) [55,56], etc. to enhance their performance. However, HMw PLA by ROP on an industrial scale was only attained by the mid-1990s [26].

Synthesis of PLA

The lactic acid monomer can be converted to PLA through a polycondensation process by the reaction of the hydroxyl (-OH) and carboxylic acid (-COOH) groups with the removal of the detrimental by-products such as water. Generally, catalysts are added to polymerization to increase the reaction rate. The removal of water, enhanced under vacuum pressure, is critical to producing HMw polymers due to the increased viscosity of the reaction mixture as the reaction proceeds. However, side reactions, such as transesterification, can also occur during the polycondensation of lactic acid, resulting in the formation of ring structures of different sizes, such as lactides. Transesterification reactions lower the overall Mw and the stereocontrol over the chain architecture, decreasing the physical properties of the PLA afforded as well as reducing the reaction yield [28]. The HMw PLA is mainly synthesized by ROP due to the accurate chemical control in terms of molecular weight, polydispersity, polymer chain-ends, and tacticity. Moreover, ROP requires relatively mild conditions (130°C) when compared to polycondensation (180-200°C) [28,57]. Three reaction mechanisms have been proposed for ROP of lactide: anionic, cationic, and coordination-insertion mechanisms.

In both anionic and cationic polymerizations, a monomer-activation mechanism occurs first that permits the catalyst to be independent of the propagating polymer, and thus, can be easily removed as the polymerization finishes. However, undesirable side and racemization reactions are likely to occur due to the highly activated monomers. On the contrary, coordination-insertion polymerization attains HMw PLA with higher control over the Mw distribution [6,28]. Metal complexes of several metals have been widely employed as the catalysts for the ROP of lactides [28] is the catalyst utilized for the industrial synthesis of PLA, due to a large extent to be approved by the FDA for the use in medical (<20 ppm [58]) and food applications. Moreover, lauryl alcohol (1-dodecanol) is usually added as an initiator [28].

The polymerization is induced by a coordination-insertion three-step mechanism supported by the catalyst [57], which was first formulated in 1971 by Dittrich and Schulz [6] (Figure 1.3). Firstly, Sn(Oct)2 reacts with the lauryl alcohol to form a tin alkoxide.

Subsequently, the exocyclic carbonyl oxygen of lactide temporarily coordinates with the tin atom of the catalyst being in the alkoxide form. The formed coordination system enhances the nucleophilicity of the alkoxide part of the initiator as well as the electrophilicity of the lactide carbonyl group that enables the reaction to each other. Finally, the acyl-oxygen bond of lactide is disrupted and the generated linear chain of the lactide turns into the alkoxide part of the catalyst, promoting the coordination with a new lactide molecule and thus, the polymerization propagation [28]. Finally, the active metal-alkoxide bond is hydrolyzed as the monomer is entirely consumed and the formation of a hydroxyl end-group occurs [57]. In the last stage of the propagation step, as the monomer concentration becomes significantly lower (~80%), both intra-and inter-molecular transesterification reactions occur and molecular weight distribution increases. However, the degrees of racemization and chain scrambling achieved by metal complexes that follow coordination-insertion mechanisms, are notably lower than those obtained by anionic or cationic catalysis [28]. The microstructure of the final polymers depends both on the initial monomers added to the reaction mixture (Figure 1. [6,28,59,60]. However, several drawbacks to controlling the synthesis of HMw stereo-block PLA have emerged since then, such as the decrease of the living character of the catalyst due to the increase in the reaction heterogeneity [61], the detrimental side reactions due to multiple nuclearities exhibited by the catalysts [62] as well as the long reaction time required to achieve the desired architectures and molecular weight [6]. Recently, novel catalysts that fulfil the synthetic requirements whilst exhibiting low toxicity for the application of the PLA derivatives in the biomedical and pharmaceutical fields were just attained [13,[63][64][65] which offers the possibility to design multiblock copolymers featuring simultaneously the PLA stereo-block to attain higher physicochemical properties with other complementary blocks to enhance the PLA limitations.

PLA processing

Once synthesized, PLA is usually manipulated into its final shape by the use of different processing techniques that apply diverse thermomechanical histories (Figure 1.

2).

Melt processing is a three-step process generally used to transform PLA into different commodity products at an industrial scale [66]. Firstly, the polymer is melted to subsequently be moulded into the desired shape which is then generally cooled to stabilize its dimensions.

The industrial techniques such as extrusion and injection moulding are the two most common melt-based processes for manufacturing thermoplastic polymers. Besides, novel common additive manufacturing technologies with promising personalised biomedical applications such as fused deposition modelling follow a melt-based process [67]. Likewise, the tailoring of Tm of PLA is crucial for the performance of applications employing melt-based processing techniques as the process temperature must be above Tm to form a homogeneous melt although low enough to minimize thermal degradation [68]. In addition, the cooling rate during the third step will influence the properties of the final product which determine the crystallization conditions that dictate the crystallinity degree and crystalline phase.

Particularly, quenching PLA from the melt at a high cooling rate (>500°C/min, such as during injection moulding) will result in a highly amorphous polymer [69], whilst semicrystalline PLA is obtained when the cooling rate is reduced (<30°C/min [70]). The α-form is usually developed from PLLA or PDLA during typical melt processing, however, a mixture of α and α' phases is obtained when the cooling rate is higher than 2°C/min since the α-form has very slow crystallization kinetics (pure α-phase is obtained at slower rates) [71]. Furthermore, postproduction treatments such as annealing can be implemented to increase the thermal stability and mechanical properties of the final product. The α-phase develops completely from a meltcrystallized material exhibiting a mixture of α and α' when annealing at 140°C for 1h [71].

Additionally, obtaining only the SC phase from the melt from the equimolar blend of PLLA and PDLA enantiomers can be just achieved by restricted thermodynamic conditions such as the relatively low cooling rates (20°C/min) to avoid the phase separation or from low molecular weight enantiomers (~20kDa) [11] and the isothermal crystallization at temperatures above the Tm of homocrystals (~175°C) [44] as well as by melt-spinning under high tensile stress conditions [72] (Figure 1.2). However, PLA materials crystallized in the SC phase has not achieved the market yet [73].

Three-dimensional (3D) printing is an additive manufacturing technology, with the unique ability to produce personalized objects with complex designs at reduced costs and a high-resolution precision [74,75] that have already reached the market in the biomedical field such as the manufacturing of 3D scaffolds for studying the response of particular tissues to different stimuli [START_REF] Sun | Crystallization Behavior of Asymmetric PLLA / PDLA Blends[END_REF]. In particular, PLA is the most-used polymer for 3D printing since the slow crystallization rate compared to polyolefins (i-PP or PET) avoids warping between layers [77]. «Classical» processing methods such as particulate leaching, gas foaming or solvent-casting were used for tuning the internal architecture of 3D scaffolds due to their adequacy for replacing tissues with high regenerative capacity. However, 3D printing enables the required control over the scaffold architecture for less regenerative tissues, such as tendons or nerves [78]. Likewise, PLA 3D printed scaffolds have already been investigated for bone [79][80][81][82][83][84],

neural [85] and musculoskeletal soft [86] tissue engineering. In addition, PLA nanofibers have also been used as part of a fibrous bioink for the 3D printing of a meniscus construct to study the proliferation of human adipose-derived stem cells that prove a higher cell proliferation and metabolic activity [86]. Recently, the assessment of PLA scaffold geometry effect on the orthopaedic applications [87] revealed that the presence of hydroxyapatite (HA) in the scaffold enables efficiently the mineralization as well as induces the crystallization of PLA after being 3D printed, whilst PLA without HA remained amorphous. The presence or absence of crystalline domains within the 3D printed PLA scaffold will invariably influence the hydrolysis degradation rate, which is crucial to controlling the optimum performance of the biomaterial. However, the usual lack of structural study for most of the reported 3D printed PLA scaffolds for biomedical applications impedes determining the relationship between the processing conditions, the crystalline structure and the biomedical performance.

Furthermore, in the biomedical industry, electrospinning has been considered a promising method to fabricate polymer nanofibers due to its simplicity and cost-effectiveness of the technique. Electrospinning, unlike drying spinning which relies on mechanical extrusion, uses the electrostatic force to spin the solution into fibres [69]. The fibres thus obtained have a nanometric diameter, producing materials with a high area/volume ratio, high flexibility, and superior mechanical properties compared with other materials formats [88].

Electrospun PLA usually exhibits either an amorphous structure or a semicrystalline structure, although with a very low crystallinity (~10%), due to the rapid solidification of the fibres during the process which entails post-processing thermal treatment between the Tg and the Tm to enhance the crystallinity by cold-crystallization in the α-phase [89,[START_REF] Gómez-Pachón | Characterisation and Modelling of the Elastic Properties of Poly(Lactic Acid) Nanofibre Scaffolds[END_REF]. Recently, the straightforward fabrication of PLA electrospun fibres exhibiting β-phase without further postprocessing treatment was achieved [START_REF] Echeverría | Development of Highly Crystalline Polylactic Acid with β-Crystalline Phase from the Induced Alignment of Electrospun Fibers[END_REF], facilitating the development of PLA-based devices with piezoelectric properties for potential biomedical applications, as will be further discussed [START_REF] Ribeiro | Piezoelectric Polymers as Biomaterials for Tissue Engineering Applications[END_REF]. Furthermore, PLA fibres in the SC phase were also generated from the electrospinning of PLLA/PDLA blends resulting in more uniform fibres [START_REF] Kang | Multilevel Structural Stereocomplex Polylactic Acid/Collagen Membranes by Pattern Electrospinning for Tissue Engineering[END_REF], however, an annealing postprocessing step is usually required to obtain crystalline structures [START_REF] Feng | The Crystallization Behavior of Poly(l -Lactic Acid)/Poly(d -Lactic Acid) Electrospun Fibers: Effect of Distance of Isomeric Polymers[END_REF]. The large efforts dedicated to electrospun nanofibers in the biomedical field arise the possibility to mimic the extracellular matrix since the human tissues and organs are formed by nanofibrous scaffolds [START_REF] Al-Enizi | Polymer-Based Electrospun Nanofibers for Biomedical Applications[END_REF][START_REF] Bhattarai | Biomedical Applications of Electrospun Nanofibers: Drug and Nanoparticle Delivery[END_REF]. However, the industrial scale up production of PLA nanofibers has not been achieved yet, due to the low throughput of the technique, and the requirement of specific solvents [69].

Nevertheless, PLA nanofibers are a topic in continuous research due to their crucial role in several biomedical applications such as bone regeneration [START_REF] Siqueira | Influence of the Addition of β-TCP on the Morphology, Thermal Properties and Cell Viability of Poly (Lactic Acid) Fibers Obtained by Electrospinning[END_REF][START_REF] Holderegger | Preclinical in Vivo Performance of Novel Biodegradable, Electrospun Poly(Lactic Acid) and Poly(Lactic-Co-Glycolic Acid) Nanocomposites: A Review[END_REF], drug delivery systems [START_REF] Cui | Fabrication and Structure Analysis of Poly(Lactide-Co-Glycolic Acid)/Silk Fibroin Hybrid Scaffold for Wound Dressing Applications[END_REF] and wound dressing [100].

Furthermore, PLA can be also processed into nanoparticles to generate drug delivery systems in which the drug release rate can be controlled by varying parameters such as the processing method, the microstructure of the starting polymer or its concentration in the PLA (L/D ratio) [108,109]. Moreover, PLA can also be part of hydrogel materials as a mechanically reinforcing and/or drug-eluting component [49,110,111]. PLA-based hydrogels studies are less common than PLA-based nanoparticles. However, a deeper structural characterization is usually accomplished.

PLA properties

When designing a PLA-based material for biomedical applications, physicochemical properties such as mechanical and thermal behaviour as well as the degradation degree under physiological conditions must be considered from the chain architecture synthesis together with the subsequent impact of the processing methods to correlate the structure-property relationship with the targeted application.

Mechanical properties

The mechanical properties of PLA are mainly associated with the crystalline phase and crystallinity degree [33] which must be tailored by controlling the stereocontrol exerted during the synthesis as well as by the subjected thermal history and pre-treatment conditions. The mechanical properties of PLA range from soft and elastic (amorphous PLA) to stiff and high strength (semicrystalline PLA) [1], and thus, semicrystalline PLA is preferred when robust mechanical properties are required. For amorphous PLA, the Tg is one of the most important parameters as below Tg amorphous PLA behaves like glass and its flexibility is drastically reduced, whilst above the Tg amorphous PLA behaves like a rubbery material. For semicrystalline PLA both, Tg and Tm, are important parameters since above the Tm the semicrystalline PLA behaves as a viscous liquid [66].

Amorphous PLA is typically obtained from rac-PLA as well as from isotactic PLA quenched from the melt (Figure 1.2). Semicrystalline PLA has a tensile strength of 60 MPa, flexural strength of 100MPa, and elongation at break of about 4%, whilst amorphous PLA has lower tensile and flexural strength (40MPa and 84MPa, respectively) and higher elongation at break (7,5%) [66]. The mechanical properties are also dependent on Mw [1] in particular, semicrystalline PLA is more affected compared to amorphous PLA [66] as the chain entanglements, which influence the crystallinity and hence the mechanical properties, increase as the Mw increases [112]. Likewise, the tensile strength increases by a factor of 10 when Mw is increased only by a factor of 4 [113]. Semicrystalline PLA presents higher mechanical properties than a few relevant commodity polymers such as PS, i-PP and PET but inferior thermal resistance [114]. PLA is then more susceptible to thermal degradation during processing, lowering its Mw and hence its final mechanical properties (Mw of PLA decreased 14-40% after injection moulding [115]).

The mechanical properties of PLA as a function of crystal polymorphisms have also been investigated, although mainly for α-and α'-forms due to their abundant availability under industrial processing conditions. Likewise, heat deflection temperature of injectedmoulded PLA materials was found to increase linearly with increasing crystallinity above ca. 35% -40% as well as with the annealing temperature increase, which could be explained by the rise of the α/α' ratio [71,116]. Furthermore, PLLA materials crystallized in the α-phase by annealing after quenching feature a larger Young modulus, tensile strength and storage modulus as well as lower elongation-at-break and water vapour permeability when compared to those crystallised in the α'-phase [116,117], which, may be attributed to the tighter molecular packing of PLA chains in the α-phase [117]. However, PLA crystallized from the melt showed a decrease in tensile strength, elongation at break, and unnotched impact strength as the temperature of crystallization increased (isothermally analysed between 100°C and 130°C) [118]. Moreover, the Young modulus increased as the temperature increased except for the PLA crystallized at 130°C (α-phase), which is associated with the formation of an aggregate between the amorphous and crystal region due to the formation of fewer nuclei [118]. The discrepancy between results may be attributed to the difference in spherulite density dependent on thermal treatment (annealing after quenching vs. crystallization from the melt)

[119]. In addition, the tensile modulus and tensile strength of PLA also increase for the β phase compared to the α-phase [36]. Similarly, the tensile properties such as tensile strength, Young's modulus and elongation at break, as well as storage modulus, are also improved by stereocomplexation of PLA [120].

Degradation

Degradation is a crucial feature to be considered when designing PLA biomedical materials with tailored properties. The main factors that affect the mechanism and rate of PLA degradation depend either on the polymer characteristics such as the molecular weight, crystallinity and shape or/and size as well as physicochemical parameters related to the surrounding medium as temperature, pH or/and enzymes [121].

Commonly, PLA degradation is a multistep process that is initiated through a hydrolytic process followed by enzymatic action, since biotic attack only occurs when PLA reduces its Mw [122]. However, in the human body, only the hydrolytic process occurs, followed by the degradation of lactic acid in the Krebs cycle to finally yield water and oxygen.

Polymer factors

Hydrolytic degradation is inversely proportional to Mw below 4x10 4 g/mol, due to the higher mobility of chains and larger density of hydroxyl groups that enhance the probability of water-soluble oligomers formation, thus catalyzing PLA degradation [125]. However, the hydrolysis rate is no longer dependent on Mw above 8x10 4 g/mol since other parameters, such as total crystallinity have a higher effect on hydrolysis [126].

Besides, semicrystalline PLA is more resistant to hydrolytic degradation than amorphous polymeric chains, including the amorphous regions between crystalline parts.

[121]. Hydrolysis of amorphous PLA showed a weight decrease of about 14% after 18 weeks, whilst semicrystalline PLA lost a similar weight after 20 months (at pH 3.4 and 37°C) [124].

However, accelerated hydrolysis in the early stages has been observed with increasing the initial polymer crystallinity (Xc) at pH 7. 4 [127]. The higher density of hydrophilic terminal groups in the amorphous regions between the inter-crystalline areas accelerates the initial hydrolytic degradation of crystallized PLA regions by enhancing the diffusion of water molecules into the bulk material. Likewise, films formed by mainly amorphous PLA undergo hydrolytic degradation with a nearly constant hydrolysis rate constant (kh) throughout the process, whilst films composed of predominantly semicrystalline PLA are subjected to a twostep hydrolytic degradation, initiated by a higher kh stage, followed by a smaller kh. Moreover, the degradation rate of the SC crystalline phase is smaller than that of the α-phase [128].

In addition, degradation is faster as the thickness of the object is increased [129]. The size is related to the hydrolytic degradation as a result of the solubility of the oligomers close to the surface that are drained into the solution as the ageing time increases. Besides, the solubilized oligomers within the matrix remain entrapped and contribute to the autocatalytic effect (Eq. 1.2). Likewise, the main PLA degradation mechanism was found to be dependent on several size regimes: bulk erosion for material thickness lower than 2mm, core-accelerated erosion between 2mm and 7.4 cm, and surface erosion for thickness higher than 7.4 cm [121].

Media factors

Furthermore, the main media factors that affect PLA hydrolysis are pH and temperature. Particularly, understanding the degradation kinetics dependence of PLA on pH is crucial to generating materials with tailored properties for biomedical applications due to the pH of the human body range from strongly acidic to almost neutral. Recent efforts have been focused on elucidating the effect of pH in the PLA hydrolysis mechanism of degradation.

The pH dependence of the kh was found to be described by the general rate constant equation

(Equation 1.2) [130]: kh = ko + kH [H + ] + kOH [OH -] (1.2)
where, ko is the first-order rate constant degradation in water, and kH and kOH are the secondorder rate constants catalysed by protons and hydroxyl ions respectively. A pH regime dependence described the hydrolysis mechanism, specifically hydroxyl-catalysed above pH 5 In addition, the effect of temperature on the hydrolytic degradation (Th) was found to be differentiated into three thermal regimes related to the chain mobility: Th<Tg, Tg≤Th<Tm and Tm≤Th. The hydrolytic degradation rate increases when Th is higher than Tg as a consequence of the higher chain mobility. Likewise, if Th exceeds Tm, crystalline regions melt and disappear, and thus, the hydrolytic degradation in the melt occurs homogeneously, similarly to racemic PLA. However, when Th is lower than Tm, the hydrolysis degradation causes a heterogeneous Mw reduction as a result of the lower water diffusion. Indeed, the activation energy of the hydrolytic degradation increases from 50.9 kJ/mol for the thermal range Tm≤Th, to 83 kJ/mol for Th<Tg, whilst intermediate activation energy of 69.9 kJ/mol is found for the temperature regime Tg≤Th<Tm [121].

(
Moreover, the rate of degradation is also dependent on the magnitude and type of the applied stress that is crucial to designing biomedical materials employed in physiological conditions. Materials subjected to a static load experience a higher rate of degradation, which is further increased when a dynamic load is applied [133].

The recent trends of PLA in pharmaceutical and biomedical research will be addressed in the following section. The pharmaceutical section comprises the applications related to drug delivery such as nanoparticles whilst the biomedical section includes the applications related to tissue reparation, although the diffuse limit between both fields.

Chiral drugs

The development of chiral drugs has recently become important for the pharmaceutical industry as frequently, the enantiomers of a racemic drug have different physiological

Polymer therapeutics

Polymer therapeutics is an alternative approach to conventional drug delivery systems which comprises polymer-drug/protein conjugates as well as polymers with a therapeutic effect by themselves (polymeric drugs). The lack of therapeutic activity featured by PLA has determined its use conjugated with drugs due to its higher stability, precise drug loading efficiency and sustained drug release compared to drug encapsulation systems [196,197]. Several synthetic methods have been proposed to conjugate PLA into different drugs (Table 1.3). Particularly, paclitaxel has been employed as an initiator for the ROP of rac-PLA through the reaction of the 2'OH of paclitaxel with lactide, ensuring the presence of paclitaxel at the chain termini [198]. In addition, paclitaxel has also been conjugated to rac-PLA throughout the PLA chain in a controlled way through an azide-alkyne click reaction which was found by in vitro cytotoxic analysis a higher therapeutic effect than the free drug [199]. Besides, antihypertensive drugs such as lisinopril have also been conjugated to PLA-PEG-PLA copolymers followed by micelle formulation exhibiting a more sustained release compared to physically loaded micelles [200]. Moreover, anti-inflammatory drugs such as indomethacin have also been covalently immobilized onto PLLA films as a potential metallic implant coating to avoid biofilms and blood clot formation [201]. However, the unusual nanostructure characterization of the PLA component of the polymer-drug conjugate which is typically based only on the molecular weight, morphological aspects as well as thermal properties (Tg, Tc, Tm, crystallinity), deter the understanding of the release kinetics [202]. a new technology in the field of percutaneous coronary intervention to provide temporary mechanical support and drug delivery followed by bioresorption of the scaffold in the vessel.

The first biodegradable stent in the clinical setting was developed 30 years ago and was constituted by PLLA [15].

Bone regeneration

Bone fracture is one of the more common injuries which is associated with treatment costs exceeding billions of dollars, societal productivity loss, and individual disability.

Approximately 5 to 10% of fractured bone results in incomplete healing [226]. Currently, stainless steels and titanium-based bone plates remain dominant in bone fracture internal fixation [227]. However, metallic implants usually lead to health side complications, including bone atrophy and infection, and their subsequent removal can result in increased weakening and bone re-fracture due to the presence of the remained screw holes. Additionally, the organization and function of each bone are highly site-specific and related to its embryological origin [228].

In recent years, the use of biodegradable polymer plates for bone regeneration is gaining attention [14]. PLA and PGA are the biopolymers more commonly used due to the featured higher mechanical properties compared to other polyesters but a sufficiently low elastic modulus to avoid the stress shielding caused by metallic implants [227]. PLA was proposed as a potential substitute for titanium plates to heal bone fractures 30

Tendon regeneration

PLA, and in particular its copolymer with PGA (PLGA) has also been employed for other tissue applications such as tendon regeneration. Tendinopathies represent about 30% of tendon-related injuries [220] and spontaneous tendon healing results in problematic restitution due to the low cellular and hypo-vascular nature of tendon tissue [221]. Biopolymeric 3D scaffolds provide an implement to both heal tendon injuries as well as to understand cell behaviour and differentiation in response to defined external biochemical and mechanical stimuli [222]. Particularly, a 2 cm partial resection of the Aquiles tendon was repaired in rabbits using a composite scaffold composed of an outer part of knitted PLGA, and an inner part of unwoven PGA fibres. The generated scaffold was incubated in adiposederived stem cells culture and cell-seeded assemblies that formed a neo-tendon presenting a histological structure similar to native tendon 45 weeks after implantation [216]. PLGA has also been employed for tendon regeneration as a nanocarrier to ensure sustained and controlled delivery of human growth differentiation factor 5 [222], which is crucial for the expression of genes linked to the neo-tendon type. Similarly, stem cells with different origins were selected to understand the physiopathology of tendinopathy in which PLGA was always used as a 3D scaffold [223,224]. Recently, the influence of certain production parameters to tailor the final nanofiber morphology such as the alignment and the fibre size were probed for an electrospun PLA scaffold on the differentiation of amniotic epithelial stem cells towards tenogenic lineage, concluding that highly aligned electrospun fibres [221] and smaller fibre diameter [220] enhanced cell differentiation and immunomodulation. Moreover, the mechanical properties of the electrospun scaffold were also influenced by the fibres alignment and its diameter, particularly, the scaffold featuring a higher diameter exhibited lower Young's modulus [220] and higher fracture strain whereas the alignment of the fibres increase the stress and strain values [221].

Nerve regeneration

Peripheral nerve injuries are the most common injury types affecting the nervous system and the posterior recovery of sensibility and motor functions is estimated to be less than 3% and 25%, respectively. Peripheral nerve injuries can be repaired by surgery, however, the use of grafts turns necessary when an important loss of physical substance occurs (>5mm)

Perspectives

Piezoelectric activity

The β-phase of PLA [39] as well as the oriented α form of PLLA films fabricated by solution casting and uniaxial stretching [240] exhibit piezoelectric characteristics (Figure 1.9) that could enhance the functional complexity designed for therapeutic applications. Indeed, the piezoelectricity of PLA has already been explored for its biomedical application as a film

Shape memory polymers

3D printing (3DP) is an additive manufacturing technology that allows mimicking the complex and multicomponent structure of human tissues at reduced costs [74,75,78]. Moreover, the 4D printing processing method has emerged as an extension of 3D printing in which printed objects exhibit the capability to self-transform over time, or in response to an external stimulus such as heat, light, pH, magnetic or electric forces [246]. Likewise, shape memory polymers (SMP) are capable of recovering their initial shape upon an external stimulus after being experienced a shape deformation [247] (Figure 1.10). PLA has already been used as a bio-ink for 4D printing applications. A spiral shape scaffold of a composite composed of PLA with magnetic iron oxide nanoparticles was printed to be used as an SMP.

Iron oxide can be remotely heated and attracted by alternating magnetic fields demonstrating that it could be applied as a self-expandable stent with a minimally invasive intervention [247] as well as overcoming the lack of radial strength found for PLA stents compared to metallic

Conclusions

PLA is a bio-based polymer that has attracted large interest due to its biodegradation capability and low price. PLA-derivatives attaining different architectures are obtained depending on both the starting lactide derivative and the polymerization method. The recently employed living polymerization assisted by a single-site catalyst allows affording PLA derivatives with high stereocontrol, high Mw and narrow polydispersity including the stereoblock copolymers. The availability of high Mw PLA stereo-block copolymers which feature enhanced physicochemical properties, expand the potential biomedical and pharmaceutical applications. In addition, the rich PLA nanostructure with diverse crystalline phases featuring different thermo-mechanical properties, as well as particular physico-chemical properties such as piezoelectricity, are obtained through several processing techniques such as electrospinning, 3D printing or nanoparticle formulation. Besides, the PLA crystalline domains degrade in physiological media at different rates depending either on the crystalline PLA phase as well as physicochemical and morphological parameters related to the surrounding medium, which must be considered when designing a biomedical device for a specific application.

Remarkably, PLA based materials have already attained the clinical level in both pharmaceutical and biomedical applications such as sutures, stents, microparticles and dermal fillers. In addition, PLA-based nanoparticles and 3D scaffolds are currently developed for drug delivery and tissue regeneration applications, respectively. Besides, the piezoelectricity and chirality exhibited by PLA enlarge the functionalities for potential biomedical devices.

Likewise, the generation of bionanocomposites based on PLA affords the customisation of therapeutic materials with multifaceted capabilities. Furthermore, personalised therapeutic approaches have become feasible with the advent of 3D printing technology, designing specific pharmaceutical formulations for specific patients, pointing toward personalized medicine.

However, the current gap between the PLA structural characterisation and the performance of PLA in biomedical applications hinders the systematic conception of personalised biomedical solutions based on the correlation of the structure-property relationship with the required functionalities. The end-user requirements should be considered from an interdisciplinary approach from the genesis of PLA-based material to realistically reach the pharmaceutical/biomedical market. Likewise, the PLA-nanoparticles formulation has greatly advanced to attain drug carriers down to nanoparticles with a fair control; however, the nanostructure development during the polymer processing is frequently dismissed, which is key to tuning the drug release profile as well as diminishing the batch-tobatch variability usually obtained through conventional techniques that hamper the scalability of nanoparticle formulation to the industrial scale. Moreover, the comprehensive nanostructure control achieved for a wide range of smart PLA-based hydrogels with specific functionalities must be followed by in vitro and in vivo studies to correlate the physicochemical properties with its performance under physiological conditions. In addition, the screening focus to develop new synthetic pathways to efficiently conjugate drugs to PLA should be also associated with a targeted PLA conjugate nanostructure in the applied medium in order to be translated reliably to both the release mechanism and the molecular recognition activity required to efficiently target the damaged tissue. Moreover, the tissue regeneration progress achieved from the materials approach requires a further understanding of the particular physicochemical and mechanical properties of each tissue, the alterations occurring under pathological conditions, and its evolution upon healing for designing materials according to the real requirements. Furthermore, the performance of the PLA-based stents suffers from structural stability due to the exhibited lack of radial strength that could be tacked by iterative structure-properties studies of novel multiblock copolymers. The mastering of novel polymer processing techniques, especially 3D techniques, should be achieved to generate customised materials with performing mechanical properties to attain consistently demanding applications.

Aims and objectives of the thesis

The main objective of this thesis is to establish from a global approach the structureprocessing-properties relationship of tailored PLA-block copolymers to obtain personalised drug delivery systems.

Firstly, the synthesis of PLA-derivatives controlling the stereo-regularity of the polymeric chain architecture, particularly the long desired PLA-stereo-block copolymers will be described (Chapter 2). Moreover, the correlation of the structure-processing-properties relationship will be established, particularly the crystallization mechanism (Chapter 3) for the design of multifunctional bionanocomposites PLA-based with personalised drug release rates (Chapter 4).

The following chapters include a brief introduction of the subject, an enumeration of the specific objectives, a description and discussion of the main results as well as an enumeration of the main conclusions. The main experimental techniques won't be described in detail, but the specifications of each technique for the reproducibility of the results will be described after the conclusions (and not after the introduction), to maintain the narrative thread.

1.9

Introduction

Previously, stereo-block-PLA derivatives have been achieved by the lengthening of building blocks on pre-polymers by a stepwise fashion [1], the extension of prepared PDLA and PLLA segments by reactive end-groups [2][3][4] or by solid state polycondensation [5][6][7].

However, limitations on the preciseness and complexity of the desired polymer architecture are the main drawbacks of the prepolymer extension synthetic approaches (asymmetric blocks, high molecular weight distributions, transesterification reactions, etc). Numerous efforts have been devoted over recent decades to the development of specific catalytic systems to promote the Ring-Opening Polymerization (ROP) of lactide (LA) attaining different tacticities under mild conditions to combine both the efficiency and polymerization control of ROP over polycondensation [8,9]. However, the slow advance of metalorganic catalysts for ROP that are able to retain the living character whilst assuring stereocontrol to avoid the common side reactions during the synthesis of HMw polymers at mild experimental conditions has deterred the development of PLA derivatives with specific tacticities. Likewise, isotactic PLA synthesized from rac-LA by aluminium Schiff base catalytic systems has been obtained with a high stereoselectivity [10]. Furthermore, heterotactic and syndiotactic PLA were also afforded from rac-and meso-LA, respectively, by catalysts featuring specific ligand characteristics, such as bulky substituents or chiral components [8,[11][12][13].

Numerous efforts have also been devoted to synthesize PLA stereo block copolymers (SBC) of HMw following the one-pot sequential addition method. Particularly, the first PLA-SBC of medium molecular weight was attained in the 90 although the increase of the reaction heterogeneity was found to decrease the living character of the catalyst [14]. Besides, detrimental side reactions were later ascertained to occur during polymerizations with homoleptic catalysts due to multiple nuclearities exhibited by the catalysts, which resulted in broad molecular weight distributions and lack of stereocontrol over the polymerization [10].

PLA-SBC have also been achieved by the two-step polymerization of rac-LA combining stereoselection and chiral ligand exchange [8], but requiring long polymerization times (up to 8h).

Recently, achiral iron complexes has recently been introduced to initiate the ROP of rac-LA presenting relatively high heteroselectivity (92% isotacticity) and attaining medium molecular weight [15]. Moreover, monomeric and highly active Zn and Mg complexes have been introduced to initiate the ROP of homo-chiral lactides to synthesize by a pot sequential addition HMw stereo-n-block copolymers in the presence of alcohol as an activator. PLA-SBC characterized with a high melting temperature (Tm) and enthalpy (ΔHm) as well as very narrow polydispersities have been successfully synthesized by the one-pot sequential monomer addition method [16], suggesting the formation of desired SC crystals [17,18]. However, the development of novel robust catalysts able to attain HMw PLA-SBC is still required to ascertain the mechanism of crystallization of the SC phase as well as to correlate the semicrystalline structure to its promising physicochemical properties, since previous studies exclusively proved the presence of the SC-phase. Although, slight structural imperfections and modest chain-length distribution are tolerated when designing block polymers in benefit of the most synthetically economical approach, such as the dispensability of the activator [19].

A potential efficient metalorganic catalytic system to attain HMw PLA-SBC by ROP should follow the coordination-insertion mechanism to avoid the termination step while maintaining a high activity for both L-and D-LA monomers during the synthesis as well as featuring a heteroleptic ligand to avoid transesterification reactions. In addition, a freeinitiator catalyst will benefit the scalability of the synthesis to the industrial level since a more synthetically economical approach is achieved. Previously, a series of chiral scorpionate zinc alkyl complexes have been proved to catalyse as single-site living initiators through a truly living polymerization at mild conditions of both enantiopure L-LA and rac-LA monomers [20].

The high conversions and stereocontrol achieved during the ROP of lactide monomers assisted by the chiral scorpionate zinc alkyl complexes have encouraged us to synthesise stereodiblock-copolymers of PLA with HMw due to their promising physicochemical properties.

Aims and objectives

The main objective of Chapter 2 is the synthesis of high molecular weight stereo-block copolymers of polylactide (HMw PLA-SBC) by the ring-opening polymerization of lactide using a heteroscorpionate catalyst (Scheme 2.1). To accomplish the main objective, the following steps were followed in collaboration with the Faculty of Pharmacy, at the University of Castilla-La Mancha, in Spain:

 Synthesis and characterization of the

Initiator [Zn(Et)(κ 3 -bpzteH)] (bpzteH = 2,2- bis(3,5-dimethylpyrazol-1-yl)-1-para-tolylethoxide).
 Analysis of the polymerization mechanism followed by the Initiator.

 Analysis of the polymerization kinetics followed by the Initiator.

 Synthesis of homo-and racemic-PLA.

 Synthesis of stereo-diblock PLA copolymers attaining different molecular weight and tacticity. The L-Lactide monomer is polymerized in solution following a coordination-insertion mechanism by the addition of the Initiator. ( 2): The D-lactide monomer is added to the reaction mixture when the first monomer has reached the maximum conversion and the PLA-SBC is obtained.

Results and discussion

Characterization and activity of the Initiator

The bioinspiring structure of the previously synthesized Zn complexes with multifunctional ligands reproducing the active centre of the metalloproteins together with their low cost confer them as great candidates to generate feasible biocompatible polymers.

The acid character of the Zn complexes enhances the polymerization of cyclic esters as facilitates the coordination of the LA monomer to the metal. Likewise, the functionalization of the scorpionate ligand to tailor the electronic environment and the steric hindrance confers stability to the metallic complex as well as directs specific molecular conformation to coordinate with other ligands. Among the scorpionate derivative zinc alkyl complexes, the

Initiator [Zn(Et)(κ 3 -bpzteH)] (bpzteH = 2,2-bis(3,5-dimethylpyrazol-1-yl)-1-para- tolylethoxide) (Figure 2.1)
, featuring an ethyl group as well as a scorpianate ligand with a tolyl (p-MeC6H4) substituent, was selected as a result of its potential higher activity aroused from the lower Zn-C bond, as well as the potential absence of hetero-enrichment due to the tolyl substituent [20]. The reactivity of the catalyst is crucial to one pot monomer sequential addition controlled by a living catalyst without the requirement of a co-initiator that potentially leads to detrimental chain transfer reactions [21]. However, the activity of the Initiator [Zn(Et)(κ 3 -bpzteH)] to ROP of LA remained undescribed in the previous work [20] as well as its mechanism of reaction. solution experiment due to the existence of an asymmetric carbon in the ligand [20]. The presence in solution of only meso-enantiomer was confirmed by the appearance of only one signal for each proton. The crystalline structure manifested a tetrahedral arrangement for the metallic centre coordinated with a single pyrazol group and with a coordination κ-NN-μ-O to the ligand, whilst a distorted tetrahedral environment found for the Zn was found for a similar monomeric structure with a κ 3 -NNO coordination with the methyl group as alkyl substituent for the Zn complex [20]. The obtained 1 H NMR chemical shifts along with the crystal structure suggested a tetrahedral environment for the zinc atom with the two pyrazole rings in cis-and trans-conformation in reference to the para-tolyl group for the monomeric monalkyl Zn complex [20]. However, the 1 H NMR spectra is unable to ascertain the presence either the dimer or monomeric species in solution.

The synthesis of the ligand 2,2-bis((3,5-dimethyl-pirazol-yl)-1-para-tolylethoxide) (bpzteH) and the Initiator [Zn(Et)(κ 3 -bpzteH)] was accomplished following the previously published procedure [20], which is briefly described in the experimental section and its molecular structure was proved by 1 In addition, Densify Functional Theory (DFT) theoretical calculations were performed to provide an overall interpretation of the coexistence of monomeric and dimeric forms of the Initiator. Both the dimeric and monomeric conformations of the Initiator were optimized and the energy profile of the most stable geometries in the process of rupture of the dimeric species was obtained (Figure A2.4). The yielded Gibbs free energy showed that the monomeric form was significantly more stable than the dimeric compound for both gas and solution (toluene)

phases, affording ΔG of -5.36 and -14.51 kJ/mol, respectively, in agreement with VT-1 H NMR analysis.

The key initiator parameters are the polymerization control and mechanism, the stereocontrol and rate [22]. Therefore, an analysis of the mechanism of reaction, tacticity and polymerization kinetics was accomplished.

Polymerization mechanism

The Moreover, the degree of control exerted by an initiator is generally assessed by the fulfilment of the following criteria: linear increase of Mn with conversion, linear increase of Mn with increasing the ratio monomer/initiator, narrow PDI, the ability to undergo sequential

Polymerization stereocontrol

Furthermore, the stereocontrol exerted by Initiator was ascertained for the ROP of a racemic LA mixture by assessing the presence of either the hetero enrichment or random insertion by homonuclear decoupling 1 H NMR of poly(rac-LA) using both Bernouillan and non-Bernouillan statistics. Likewise, the cis configuration of the relative orientation of two sequential alkyls of the diade is referred as i ("isotactic") or m ("meso"), whilst the trans orientation is denoted as s ("syndiotactic") or r ("racemic"). In this case, m and r notations are used. Thereby, five possible tetrad combinations are generated during the polymerization of rac-LA (Table 2.1). The stereocontrol achieved by a particular initiator is expressed in terms of probability to generate certain stereosequences. Pm is known as the probability of isotactic enchainment while Pr is the probability of syndio/heterotactic enchainment [22]. Likewise, a random insertion is indicated when Pm = 0.5 whilst isotactic PLA is designated by Pm = 1, with the total probability expressed as Pm + Pr = 1 [30]. Stereocontrolled polymerization can be mediated by two distinct mechanisms, namely, chain-end control (CEC) and enantiomorphic site control (ESC). In the former case, the control of the chirality is known to be associated with the propagating chain end in which the transition state of the next monomer insertion defines the chirality of the next monomer unit to be incorporated into the growing chain. In contrast, for polymerizations mediated through an ESC mechanism, the chirality of the catalyst determines the chirality of the next monomer unit [30]. However, stereocontrol in the ROP of PLA is commonly achieved by CEC even when the initiator contains a chiral component. The deconvolution of the homodecoupled methane resonance of poly(rac-LA) derivative (Table A2.1, entry 14) revealed five signals in methane region (Figure 2.5), which were assigned to the five tetrads obtained from rac-LA polymerization [23] and were analysed by Bernouillan and non-Bernouillan statistics based on both CEC and ESC mechanism, respectively (Table A2.2). However, the probabilities calculated using non-Bernouillan statistics remained unresolved indicating that the stereocontrol exerted by Initiator follows a CEC mechanism. Likewise, the Pm obtained by CEC for poly (rac-LA) was 0.52, in agreement with a completely atactic polymer. The homodecoupled methine resonance employing a CEC mechanism was applied to the afforded PLA derivatives to elucidate the tacticity values.

Polymerization kinetics

Polymerization rate constants were determined by solution kinetic studies of a 2.2), with apparent rate propagation constants (kapp) of the polymerization conducted in toluene one order higher than the kapp values found for polymerizations carried out in THF with the same family of initiators [20]. In addition, identical order of magnitude of kapp was found for polymerization previously performed in dichloromethane with other Zninitiators [23]. The linearity of the plots corresponds to a first-order dependence with respect

𝐿𝑛 𝑘 𝑎𝑝𝑝 = - 𝐸 𝑎 𝑅 ( 1 𝑇 ) -𝐿𝑛 𝐴 (Eq. 2.1)
where kapp is the apparent rate propagation constant (10 2 s -1 ), Ea the activation energy (J mol -1 ), R the gas constant (R=8.314 J mol -1 K -1 ), T the temperature (K) and A the Arrhenius constant.

The Ea for the LA polymerization obtained, 61.57 kJ mol -1 , was higher than the Ea data for organoaluminium initiators (~6.5 kJ mol -1 ) [24], which explains the slower polymerization kinetics observed. Moreover, the polymerization rate constant of both homopolymers, PLLA and PDLA (Table 2.2, entries 4 and 5), were compared to confirm the absence of hetero-enrichment yielded by the heteroscorpionate complexes counterparts previously described [20], which would be evidenced by a significant higher rate constant for one of the enantiomers. Generally, a catalyst featuring a chiral carbon will invariably result in stereocontrol of polymerization exerted by an ESC mechanism. However, the CEC mechanism becomes predominant in polymerizations using a chiral catalyst under harsh reaction conditions (i.e. high temperature and monomer concentration). In addition, a slight difference in the activation energy barrier between L-and D-LA monomers to form the coordinated complex catalyst-monomer might influence certain kinetic parameters when the polymerization approaches equilibrium, such as the monomer concentration at equilibrium that potentially would control the stereochemistry of the generated polymer. Likewise, L-and D-LA monomer conversions during two independent homopolymer reactions were monitored by 1 H NMR, following the decrease in monomer concentration of both enantiomers during the PLA polymerizations.

The modelled polymerization kinetics manifested a slightly higher rate for L-LA polymerization over the D-LA (

Homochiral and Racemic Polymerization of Lactide

The ROP reaction conditions (time, temperature, and solvent) were optimized (Table 2 Likewise, the specific optical rotation (α) was evaluated from polarimetry measurements to corroborate the reaction stereocontrol (Table 2.3). The chiral purity may be estimated higher than 95% for all enantiopure derivatives (Table 2.3, Entries 1-8) based on calculations from the reported value of (α) =-173° at 598 nm for PLLA featuring perfect optical purity [25]. The specific optical rotation of poly(rac-LA) (Table 2.3, Entry 9) was close to 0°, in agreement with the previous Pm analysis by deconvolution of the homonuclear decoupled 1 H NMR spectrum.

Stereo-Block Copolymerization by Sequential Monomer Addition

Stereo-diblock PLA co-polymers exhibiting different Mw and tacticities were obtained by the one-pot sequential addition method ( represent the length of the block measured in lactide units (also the monomer/catalyst ratio).

The Mw was controlled by the monomer/catalyst ratio and the tacticity was controlled by changing the time at which the second monomer was added to the reaction, although the different between the experimental and theoretical Mn increase as the [LA]/[cat] ratio increases probably due to the increase of the reaction mixture viscosity. The 1 H NMR monitored monomer consumption revealed conversions higher than 80% for both block sequences, which was in good agreement with the Mw determined by GPC. Likewise, the kinetic experiments of the addition of the second monomer yielded an equivalent kapp 6.2 ± 0.002 s -1 (Figure A2.14)

with a first order of reaction, confirming the living character of polymerization upon the growth of the second block. to the appropriate tetrads following previous reports [37], and a Bernouillan statistical model was applied to the deconvoluted spectra. The obtained Pm values for the series of stereodiblock copolymers (Table 2.2, Entries 1-4 and 7-9) confirmed the high stereocontrol achieved in the ROP for both blocks (>0.88), evidencing the almost full consumption of the two LA enantiomers as well as the isotactic stereo-diblock microstructure nature of the polymeric derivatives. Moreover, the Pm value of two stereo-diblock copolymers (Table 2.2, Entries 5-6)

was lower than 0.81 due to the addition of the second monomer before the total consumption of the first one, increasing the number of stereoerrors in the second block, hence decreasing its tacticity. Finally, an ultra-high molecular weight asymmetric block copolymer was synthesized (Table 2.2, Entry 10) by increasing the L-LA/D-LA ratio to 6. The high tacticity of the asymmetric block copolymer was confirmed by the Pm value obtained of 0,93 and the specific optical rotation measurement confirmed the presence of 80% and 20% of L-LA and D-LA, respectively. 3).

Conclusions

PLA stereocomplexation has typically been achieved by the blend of the PLA enantiomers with detrimental phase separation at high molecular weight that diminishes the promising physicochemical properties featured by the SC phase. In Chapter 2, the synthesis of high molecular weight stereo-diblock copolymers of PLA that might potentially fully crystallize in the SC phase has successfully been achieved.

 The first organometallic Initiator able to generate high molecular weight stereodiblock copolymers of PLA without the use of a co-initiator has been reported.

 The Initiator follows a coordination-insertion polymerization mechanism.

 Kinetic experiments confirm the living character of the initiator, the first-order reaction and similar rate constant of the polymerization for each block.

 The reaction conditions have been optimized to attain homo-, racemic-and stereocomplex-PLA exhibiting high molecular weight and relatively narrow PDI.

 The PLA-derivatives tactities and specific optical rotations values confirmed the fine stereocontrol achieved and minimal transesterification reactions on both symmetric and asymmetric PLA-SBC.

In Chapter 3, the characterization of the crystallization mechanism of the PLA stereodiblock copolymers will be accomplished.

Experimental section 2.5.1 Materials

Solvents and reagents were commercially acquirable from Sigma-Aldrich (St. Louis, MO, USA) apart from L-lactide (L-LA) and D-lactide (D-LA) monomers that were purchased from Rex scientific. The lactide monomers were purified three times by sublimation and then stored in a glovebox at 4 _C. Toluene and tetrahydrofuran (THF) were pre-dried over sodium wire as well as distilled under nitrogen from sodium and subsequently stored over molecular sieves (3Å) in a glovebox.

Synthesis

All the synthetic manipulations were performed under nitrogen, using standard Schlenk techniques.

The synthesis of the Initiator followed the previously published procedure [24].

Synthesis of the Lig and 2,2-bis((3,5-dimethyl-pirazol-yl)-1-paratolylethoxide) (bpzteH)

The ligand (bpzteH) was synthesized via the following procedure: In a 250 mL Schlenk tube, Bis 3,5-(dimethyl-pirazol-yl)-methylene (bpm) (3 g, 14.7 mmol) was dissolved in dry THF (70 mL) and then cooled to -78°C. A 2.5 M solution of BunLi (6.47 mL, 16.17 mmol) in hexane was added to the bpm solution and stirred for 1h. The resulting mixture was added dropwise to a cooled (-10°C) solution of p-tolyl (1.86 g, 15.43 mmol). The reaction mixture was allowed to warm up to ambient temperature and was stirred for 1h. The obtained product was hydrolyzed with saturated aqueous NH4Cl (15 mL). The organic layer was extracted, dried over MgSO4, filtered, and the solvent was removed in vacuum to yield the product as a yellow oil, which was titrated with hexane to afford the pure product as a yellow solid (1.54 g, 75%)).

Synthesis of the Initiator [Zn(Et)]-[2,2-bis((3,5-dimethyl-pirazol-yl)-1-paratolylethoxide)] (Zn(Et)(κ3-bpzteH)

The catalyst (Zn(Et)(κ3-bpzteH) was synthesized by the following procedure: The ligand (bpzteH) was dissolved in dry toluene and Zn(Et)2 was added through a cannula to the solution placed on ice-bath (0°C). The resulted solution was stirred for 2h at room temperature and the solvent was removed in vaccum. Yield : 1,54g, 75%.

Polymerizations of LA

Ring-opening polymerizations (ROP) were performed on a Schlenk line in a dried Schlenk flask equipped with a magnetic stirrer. The Schlenk tubes were loaded in the glovebox with the required amount of LA and Initiator 1, separately, and dissolved in the appropriate amount of solvent (Table A2

Kinetic experiments

Reaction kinetic studies were conducted on the polymerizations of PLLA400 S1, entries 1-6). The polymerization procedure was followed as previously described and aliquots were collected using a glass pipette at appropriate time intervals. All the volatiles in the aliquots were removed, and the residue was subjected to monomer conversion determination that was monitored by integration of methine resonances and its subsequent monomer against polymer ratio calculation in 1H NMR (CDCl3).

Instrumentation and measurements

1 H NMR, 13 C NMR and homonuclear decoupled 1 H NMR

Spectra were recorded at room temperature on a Varian Inova FT-400/500 spectrometer and referenced to the standard in the deuterated solvent with the relaxation time fixed to 4 s.

The polymeric derivatives were dissolved in CDCl3 for the synthetic characterization as well as homonuclear decoupled 1 H NMR and Toluene-d8 for variable temperature 1 H NMR, respectively.

Homonuclear decoupled 1 H NMR spectroscopic analysis was conducted by applying the decoupling into the methyl region (δ = 1.5). The global spectrum deconvolution technique was performed when the multiplicity of the proton associated with the CH group of the polymer backbone corresponded to the tetrad signal (δ= 5.00-5.20), using both non-Bernouillan and Bernouillan statistics, further experimental calculations in Supplementary Information.

MALDI-TOF analysis

Matrix-assisted laser desorption/Ionization-time of flight (MALDI-TOF) spectra were acquired using a Bruker Autoflex II TOF/TOF spectrometer using dithranol (1,8,9trihydroxyanthracene) and NaI as the matrix and the cationisation reagent, respectively.

Samples co-crystallised with the matrix mixture in a ratio 100:1 on the probe were ionized in the positive reflector mode. External calibration was performed by using Peptide Calibration

Standard II (covered mass range: 700-3200 Da) and Protein Calibration Standard I (covered mass range: 5000-17,500 Da).

GPC analysis

Molecular weight (Mn, Mw), as well as polydispersity index (PDI) of the afforded polymer derivatives, were determined from Gel permeation chromatography (GPC) using chloroform as eluent. The molecular weight characterization was performed with a Shimadzu

Prominence-I LC-2030 equipped with a Shodex GPC KF-805L column (Shodex, Tokyo, Japan) and a Shimadzu RID-20A detector (Shimadzu, Kyoto, Japan). Analytical grade CHCl3 was used as the mobile phase at 40 _C, with a flow rate of 1 mL/min. GPC samples were prepared by dissolving ca. 5 mg of the corresponding polymer in 1.5 mL of solvent and left overnight under constant agitation. Thereafter, the samples were filtered over a 0.2 µm polytetrafluoroethylene (PTFE) syringe filter prior to its injection. Polystyrene with a molecular weight of 172 kDa was used as a reference to determine the overall molecular weight of the samples.

Optical rotation

The specific optical rotation (α) was measured in dichloromethane (DCM) at a concentration of 1mg/mL and 25 °C by a JASCO P-2000 WI using a beam wavelength of 589 nm.

Computational methodology

Computational Simulations were conducted by using Gaussian09 (Rev. C.01) software package (Gaussian 09). The theoretical calculations employed a hybrid density functional B3LYP method to minimise the geometries of all species involved. 6-31G (d) basic set was used to represent all atoms except Zinc (Zn). For Zn, LANL2DZ has successfully been applied to describe the geometry of complexes [26,27]. Besides, the solvent (toluene) was simulated using the Polarizable Continuum Model (PCM) methodology implemented in Gaussian09 package [28]. The initial estimations for the different intermediates were generated from the crystallographic structures of the dimer. Geometry at the selected level of theory was optimized for dimer structure, and then monomeric species were generated from them.

Subsequently, an initial conformation analysis was carried out for all species to obtain the lowest energy conformation. The nature of the stationary points was assessed using the normal vibration frequencies calculated from the analytical second derivatives of the energy. The firstorder saddle points -which are related to transition states-must show an imaginary value for the frequency associated with the eigenvector primarily describing the product formation step, whereas the real minima of the potential energy hypersurface -which are related to stable species-have to show positive real values for all the vibrational frequencies. Gibbs free energies used were obtained as the sum of electronic and thermal free energy at 298.15 K

Calculation of Pm values

Homodecoupled 1 H NMR spectra were obtained for racemic and stereo-diblock copolymers of PLA.

The methane region was deconvoluted and the curves were integrated using Origin software. The five tetrad dyads of the racemic spectrum were assigned according to literature [23] and probability of each dyad was obtained after normalizing by the total area. Both, Bernouillan and non-Bernouillan statistics based on chain-end control mechanism (CEC) and enantiomorphic site control mechanism (ESC), respectively, were applied. For CEC mechanism, Pr was firstly obtained from

[rmr] = 0.5 Pr 2 (1) to finally calculate Pm from

[mmm] =Pm 2 + 0.5 Pm Pr (2)
For ESC mechanism, Pm was directly obtained from

[mmm] = [Pm² + (1-Pm)² + Pm 3 + (1 -Pm) 3 ]/2 (3)
Since the probabilities calculated using non-Bernouillan statistics remained unresolved, Bernouillan statistic was applied to the afforded stereo-diblock copolymers for calculation of Pm.

CHAPTER 3

Characterization of the mechanism of crystallization of PLA stereo-di block copolymers

Introduction

The mechanism of crystallization of a semicrystalline polymer is well accepted to involve two main steps: the nucleation and subsequent growth [1]. The nucleation step can proceed homogeneously or heterogeneously. Homogeneous nucleation consists on the separation of a crystal nucleus from the amorphous phase by an interfacial surface. Below the melting temperature, a thermodynamic competition between the energy gain in the nuclei and the energy loss in the surface occurs, until a critical nucleus size is formed and the free energy barrier is exceeded. In heterogeneous nucleation an additive such as a nucleating agent (NA) is added, which already constitutes the surface where the polymer starts the crystallization [2],

reducing the energy barrier required to form a stable nucleus. Subsequently, during the growth step and accordingly to the Lauritzen-Hoffman theory [3,4], the polymer chain folds back and forth into crystal stems to form lamellas, which are ca. 10-20 nm thick platelets with the growth axis that advance perpendicularly to the polymer chain [5]. Thereafter, the crystal growth continues with the stacking of the lamellas that growth radially from the nucleus to form the spherulite crystal structure (>1000 nm) which continues advancing with larger radius until impingement occurs [6]. The nucleation (i) and growth rate (g) can be determined by the application of the Lauritzen-Hoffman theory to the spherulites growth rate (G) determined by polarized optical microscopy (POM):

𝑮 = 𝑮 𝒐 𝒆 [ -𝑼 * 𝑹(𝑻 𝒄 -𝑻 ∞ ) ] 𝒆 [-𝑲 𝒈 /(𝑻 𝒄 𝜟𝑻𝒇) (Eq. 3.1)
where G0 is a constant known as front factor, U* is the activation energy for local motion, R is the gas constant, Tc is the crystallization temperature, T∞ is the temperature at which motion associated with viscous flow ceases (typically the glass transition temperature, Tg), 𝛥𝑇 is the undercooling and f is a factor that account the change in heat of fusion with temperature. Kg is the nucleation constant, which is given by:

𝑲 𝒈 = 𝒂𝒃𝝈𝝈 𝒆 𝑻 𝒎 𝟎 𝒌∆𝑯 𝒎 (Eq. 3.2)
where a is a constant that depend on the crystallization regime (2 for regime II and 4 for regime I and III), b is the surface nucleus thickened, σ is the lateral surface free energy, σe is the fold surface free energy, T 0 m is the equilibrium melting temperature, ∆𝐻 𝑚 is the heat of fusion for 100% crystallinity and k is the Boltzmann constant [1].

Moreover, three regimes are identified in the Lauritzen-Hoffman theory depending on the limiting factor to the growth rate. At low undercooling (Regime I) the chain mobility to crystallize is high and the limiting factor is the nucleation rate. At high undercooling (Regime III) the nucleation rate is high but the chain mobility is reduced, limiting the overall growth rate. In middle undercooling (Regime II), the combination of a higher nucleation rate and still enough chain mobility usually prompt the highest growth rates. Both, nucleation and growth processes, are influenced by the temperature, molecular weight, chemical structure, thermal history, etc [1]. Moreover, crystallization can be studied from solution, that provides isolable polymer crystals that can be precisely studied, and from the melt, that is closer to the realistic industrial conditions but additional degree of difficulty is added to the fundamental structural studies, that will affect the nucleation and growth of a crystal [7].

Polylactide (PLA) is a semicrystalline polymer that can be arranged in different crystalline phases (Table 3.1). The α-phase is the more stable PLA homo crystal (HC) structure normally obtained when isotactic PLLA or PDLA are crystallized from the melt above 130°C

or by crystallization from solution, characterized by a melting temperature of around 170°C [8,9]. The α'-phase (or δ-phase) is the disordered form of the α-phase and is generated either from crystallization from the melt at temperatures below 110°C or by cold-crystallization after quenching PLA to the glassy state. Besides, a mixture of α-and α'-phases is obtained when PLA is crystallized between 110°C and 130°C, although the α'-phase recrystallizes into αphase when heated near the Tm (150-160°C) [8]. Moreover, the β-phase is developed by stretching PLA fibres in the α-phase at high temperature (130°C-140°C) and high draw ratios as well as by casting thin films from solution [9,10], and its Tm is ca. 10°C lower than α-phase Tm [11,12]. Finally, the γ-phase is produced by epitaxial crystallization of PLA on hexamethylbenzene, forming two antiparallel threefold helices in an orthorhombic unit cell [13]. Moreover, 11 lactate units are needed for PLLA to nucleate [14], the spherulite growth rate (G) and nucleation constant (Kg) of homopolymer (HP) PLLA varies from 2 to 10 µm/min and from 1.85x10 5 to 5.97x10 5 , respectively, depending on the crystallization temperature [15],

and the PLLA lamellar thickness is approximately 24 nm [16]. The main drawback of PLA to replace polyolefins as materials for commodity products is the slow crystallization rate that hampers its processability under industrial conditions [17].

Crystalline

The slow PLA crystallization rate is mainly due to the presence of an oxygen atom in the polymer backbone that increase its flexibility with a nearly free rotation around the O-C bond, which implies less extended chains and a higher entropy barrier to crystallize [18]. Different strategies have proved to increase the crystallization rate of PLA such as stretching in the melt that promotes the polymeric chains orientation along the drawn direction that accelerates the nucleation of PLA [19][20][21], the use of nucleating agents to reduce the nucleation barrier of PLA [22,23] as well as the stereocomplexation between PLLA and PDLA, that decrease the free energy of the system compared to the HC counterpart [24,25] due to specific C-H…O-H hydrogen bonds within the crystal lattice that stabilize the structure [26] and confirms stereocomplexes with a higher melting (220°C) as well as degradation temperature (220-260°C) [27].

The stereocomplexation of PLA was first reported to occur by precipitation from a solution containing both enantiomers, PLLA and PDLA [28]. Several parameters affect the stereocomplexation of PLLA and PDLA, such as the mixing ratio of enantiomers, the annealing time and temperature [29] as well as the molecular weight, which is a crucial factor since exclusive stereocomplexation from the melt is only achieved for PLA-enantiomers exhibiting Mw lower than ~20 kDa [29,30]. During the crystallization of physically blended enantiomers featuring higher molecular weight (HMw), a competition between stereocomplexation and homocrystallization occurs [30], since the stereocomplex (SC) formation suffers from a larger kinetic barrier despite its higher thermodynamic preference.

The mechanism of crystallization in the SC-phase of the enantiomeric blends has been largely investigated [31,32]. The formation of HC crystals during crystallization from the melt occurs after the SC formation instead of simultaneously [30], whilst the homocrystallization occurs before stereocomplexation during cold crystallization [33]. Moreover, the formation of helical pairs (PLLA and PDLA) below 220°C was observed by FTIR previously to the formation of a mesophase below 190°C that conclude with the stereocomplexation below 150°C [34]. However, the helical pairs are formed in the melt at relatively low cooling rates such as 2.5K/s which are suppressed at 3000K/s [35], leading to the appearance of the 103 helix conformation, characteristic of the α-phase. Although, the 103 helix undergoes a conformational transformation to the 31 helix, characteristic of the SC phase, upon isothermal crystallization above the Tg [36].

Moreover, the phase separation usually occurring in enantiomeric blends of HMw has successfully been avoided by different approaches, such as the cold crystallization above the melting temperature of HC [37], the crystallization from solution at a high evaporation rate (such as in spin coating or solution casting) [38], the decrease of the entanglements during the solution blending by decreasing the polymer concentration in solution [39], the miscible polymer blending with hydrogen bond interactions, such as poly(vinyl phenol) [40] or poly(ethylenglycol) [41], by a pickering emulsion approach followed by the reinforcing of the structure with cellulose [42] as well as the use of nucleating agents such as nanocrystals [43] and nanofibers [44] of cellulose. However, strategies to obtain full stereocomplexation from the melt with high crystallinity values of HMw PLA-derivatives in a relatively short processing time, such as the crystallization of PLA stereo-block copolymer (PLA-SBC), are still required to adapt to the processing conditions at the industrial scale.

Aims and objectives

In Chapter 2, a series of PLA-SBC with different molecular weight and tacticities were synthesized, particularly, the largely desired HMw PLA-SBC, to avoid the usually occurring phase separation when crystallizing the homoenantiomeric physical blend. In this Chapter, the crystallization mechanism from the melt of the PLA-SBC will be analysed and a summary of the samples that will be used in each section of the chapter is depicted in Table 3.2 based on both, adequacy and availability.

 Firstly, a brief analysis of the thermal degradation of all the symmetric PLA-SBC as a function of the molecular weight will be accomplished by Thermo-Gravimetric Analysis (TGA). Moreover, the methodology to verify the Mw/tacticity decrease of the PLA-SBC after the thermal treatments, particularly by DSC, will be explained using the MMw-HPm as PLA-SBC model.

 Subsequently, the effect of the Mw in the crystallization of the PLA-SBC will be discussed, since it is a crucial parameter that determines the HC/SC ratio when stereocomplexation occurs from the enantiomeric blend. A series of PLA-SBC with similar tacticity (Pm higher than 0,9) and in 44-134kDa Mw range will be used in this section.

 Then, the spherulite growth of PLA-SBC, particularly the L520:D520 as a HMw model, will be analysed and compared to both the corresponding homopolymer counterparts (L5 and D1), as well as to literature of stereocomplex formed from the enantiomeric blend.

 Moreover, an in-depth analysis of the crystallization mechanism under isothermal and non-isothermal conditions will be addressed, and a series of mathematical models will be applied to understand the kinetics of crystallizations. The effect of the Mw and tacticity on the crystallization mechanism will be addressed by selecting three PLA-SBC with different Mw and tacticity.

 Finally, the efficiency of different nucleating agents on the crystallization of PLA-SBC will be discussed to analyse if the crystallization rate of PLA can increase by both strategies, stereocomplexation and the use of NA.

Results and discussion

Thermal degradation of PLA-SBC

The understanding of thermal degradation/stability of the synthesized PLA-SBC (Table 2.3) is crucial to underpin the crystallization studied from the melt as involves the thermal treatment at high temperatures that might induce their degradation. temperature shifted to ca. 240°C, which is in the range of the reported SC-PLA degradation [27]. In addition, the lower Mw PLA-SBC (L115:D115) was completely degraded at 300°C whilst the higher Mw PLA-SBC completely degraded between 350 and 400°C, which is higher than the temperature of degradation reported for the HP counterparts [27,46] and in agreement with the degradation temperature previously reported for blended SC-PLA [47,48]. The higher degradation temperature of the SC-PLA compared to the homopolymer-PLA (HP-PLA) is ascribed to the endurance of the hydrogen bonds that stabilize the stereocomplex crystal phase during the melt at temperatures below ca. 260°C [34,48].

Moreover, PLA degrades upon thermal processing due to undesirable reactions that occur at or above the melting temperature (Tm), such as hydrolysis, transesterification reactions, etc. [49,50], that can affect the molecular weight distribution of the isotactic HP-PLA (PLLA or PDLA). However, transesterification reactions occurring in the PLA-SBC derivatives can also affect the stereochemistry of the polymeric chain, which might modify the crystallization mechanism. proportionally with the increase of the duration of the thermal step at Tm until the crystallization was impeded when the Mw decreased more than 60% after 60 minutes at 230°C, similarly to the decrease of Mw found for PLLA after thermal processing [50]. Interestingly, the Tc of both the HP-PLLA [17] and SC-PLA formed from the blend [30] increases as the Mw decrease, due to an increase of the molecular mobility. Moreover, SC-PLA presenting Mw similar to the PLA-SBC after 60 minutes at 230°C (~28.000 g/mol) is able to crystallize upon cooling at 10°C/min [30], suggesting that the observed decrease of Tc as a function of the holding time at 230°C (Figure 3.2) might be consequence of the transesterification reactions that changed the stereochemistry of the stereoblock copolymer rather than the decrease in Mw.

However, Raman or Fourier-Transformed Infrared Spectroscopy analysis are required to confirm the stereochemistry of the polymeric chain after the thermal treatment. Therefore, the thermal protocols imposed to the PLA-SBC derivatives were firstly introduced by a dwell step at 230°C for 2 minutes to erase the thermal history of the HMw PLA-SBC for the analysis of the crystallization mechanism from the melt. In addition, the Mw/tacticity stability was controlled after subjecting the polymer to a thermal treatment in the melt by evaluating the Tc in a supplementary cooling ramp to draw conclusions with unequivocal polymer architectures.

Effect of Mw in the PLA-SBC crystallization

The effect of Mw in the mechanism of the stereocomplex crystallization was studied due to its crucial impact to determine the sterecomplexation/homocrystallization (SC/HC) ratio obtained in the competitive crystallization of the PLA enantiomeric blends. The crystallization behaviour under cooling from the melt of PLA-SBC synthesized with different Mw was analysed by DSC by applying thermal protocols with cooling and heating rates of 10°C/min (the complete protocol is specified in the experimental section) (Table 3.3). Herein, only the PLA-SBC derivatives presenting Pm values higher than 0,9 were considered for further discussion and are highlighted in Table 3.3 (Figure 3.3), although the crystallization from the melt of all the PLA-SBC was analysed by DSC (Figure A3.1 and Table 3.

3).

The crystallization and melting behaviour of PLA-SBC were found to depend strongly on Mw, similarly to enantiomeric blends although with an inverse effect [41]. The Tc of the PLA- Moreover, the Tm of the PLA-SBC decreased as the Mw decreased, similarly to both HP-PLA [51] and SC-PLA formed from the enantiomeric blend [45]. In addition, the Tg of the PLA-SBC occurred between 41 and 56°C without a clear dependence on the Mw whilst the crystallinity values obtained from the melting enthalpy (Xc-m) increased from 28,02%

(L115:D115) to 46,58% (L500:D500) with increasing Mw.

The low and medium Mw enantiomeric blends typically crystallize completely in the SC phase, with a decrease of the SC/HC ratio as the Mw increases, exhibiting a clear HC melting endotherm transition [30,52]. However, the HMw PLA-SBC (L375:D375; L500:D500; L520:D520) fully crystallized in the SC phase upon cooling from the melt, as evidenced by the single endotherm at ca. 210°C and the absence of cold-crystallization upon heating (Figure

3.3B and Table 3.3).

Moreover, the crystal patterns of the PLA-SBC derivatives under study (with Pm values higher than 0,9) obtained by Wide Angle X-Ray Scattering (WAXS) upon crystallization from the melt at 10°C/min confirmed the presence of exclusively the stereocomplex phase in all the PLA-SBC derivatives, independently of the Mw, as evidenced by the reflections (110), (300)/(030) and (220) (q (nm -1 ) = 8.5, 14.7 and 16.9) (Figure 3.4).

Spherulite growth rate of stereo-block copolymers (PLA-SBC) and PLLA and PDLA enantiomers (PLA-HP)

The spherulite growth rate of the HMw stereo-block copolymers was studied by polarized optical microscopy (POM) for the PLA-SBC (L520:D520) derivative as model compound for the rest of the family due to featured higher Mw, under both non-isothermal and isothermal conditions and compared to the HP-counterparts of slightly lower molecular weight (PLLA500 (L5) and PDLA500 (D1) Table 2.2)), summarized in Table 3.4, to understand the kinetics of nucleation and growth of the PLA-SBC as well as the crystal morphology developed upon cooling from the melt.

Firstly, the thermal history was erased by applying an isothermal step at the Tm during 2 minutes (200°C and 230°C for HP and SBC, respectively) and subsequently were cooled down to 30°C at a constant cooling rate of 5°C/min while acquiring POM images periodically.

The HP crystallization of both enantiomers (PLLA and PDLA) started at ca. 135°C while the PLA-SBC crystallization started 25°C before, at ca. 160°C (Figure 3.5, yellow arrows). The Tc of the HP-PLA derivatives was similar to commercial PLA of similar Mw crystallized under the same conditions [23,53] and the PLA-SBC Tc was 30°C higher than the crystallization temperature of enantiomeric blends of similar Mw studied by POM [54].

Moreover, heterogeneity in spherulite size was found for both HP-PLA and PLA-SBC upon cooling to 30°C (Figure 3.5) that might be ascribed according to Hoffman-Lauritzen theory to the transition from Regime II to Regime III, as in regime II both nucleation and posterior growth contribute equally to the total growth rate whilst the Regime III is dominated by nucleation, since the reduction of chain mobility impedes the growth of spherulites.

The crystallization kinetics were also analysed under isothermal conditions. After erasing the thermal history, the PLA derivatives under study (PLLA, PDLA and L520:D520)

were cooled down at 30°C/min to the isothermal-crystallization temperature (Ti-c) and hold during 60 minutes. The Ti-c were selected based on the crystallization temperature (Tc) found during the non-isothermal treatment, ±10/15 °C around the Tc, (summarized in Figure 3.6 and Table 3.5). However, the comparison of crystallization among the PLA-derivatives must be accomplished carefully due to the difference in Ti-c for both HP-PLA and PLA-SBC. The three PLA-derivatives presented spherulites with obvious Maltese crosses and boundaries well defined in contrast with the observations found for PLA-enantiomeric blends of Mw higher than 32 kg mol -1 , in which the Maltese crosses were indistinctly observed and the boundaries were undefined [52]. Moreover, the presence of a periodical lamellar twisting was previously observed for PLLA and PDLA whilst disturbed lamella were found for the PLA-SC formed from enantiomeric blends [55]. However, both the HP-PLA and PLA-SBC exhibited a disturbed lamellar twisting morphology across the film area under study (Figure A3.2)

although it can be ascribed to a high twisting distance that was unseeable at the spherulite size grown in the samples.

The three PLA-derivatives under study (PLLA, PDLA and L520:D520) exhibited lower spherulites density as the Ti-c increased, as previously reported for PLLA [17,56] due to the decrease of the nucleation rate with temperature. However, the spherulites density featured by PLA-SBC was higher than for HP-PLA, considering the difference of the Ti-c analysed among the PLA derivatives, as previously determined for enantiomeric blends [57].

Likewise, the spherulite radius (r) of the three PLA-derivatives (PLLA, PDLA and L520:D520) was measured as a function of time (t) (Figure A3.3) to calculate the crystal growth rates (G) from its slope (Table 3.5). The G values obtained from the HP-PLA are equivalent to analogous PLA derivatives previously studied under the identical conditions (Ti-c and Mw) [58,59]. However, the crystal growth rate determined for the PLA-SBC was lower compared to lower Mw blended stereocomplex PLA [57] that might be due to the decrease of the chain mobility as the molecular weight increases.

The maximum G was obtained at 130°C (5,803 µm/min) and 135°C (6,3 µm/min) for PDLA and PLLA respectively (see orange and violet arrows in Figure 3.7A), which is in agreement with previously crystallization mechanism for similar PLA-derivatives [59,60]. In addition, the PLA-SBC exhibits its maximum G at Ti-c=155°C. However, a higher G at lower Tic cannot be discarded, which is experimentally unattainable as the high crystallization rate produces the impingement of spherulite to be able to measure their radius.

The spherulite growth rate data obtained from POM were further analysed using the Lauritzen-Hoffman equation (Eq 3.1), from which the nucleation factor (Kg) and front factor (G0) can be obtained from the slope and the intercept (Figure 3.7B) of the logarithmic plot (Eq.3.3 summarized in Table 3.3). The HP-PLA followed a single linear trend, whilst the PLA-SBC crystallization proceeded in two different linear regimes, although analysis at higher Ti-c should be accomplished. Hence, the crystallization of the PLA-SBC derivative might begin in the regime II for temperatures higher than 170°C and continued in regime III at temperatures lower than 170°C, in agreement with the transition temperature reported for PLA-SC formed from the enantiomeric blend (165°C [62] and 185°C [63]).

Ln G + U * R (T i-c -T ∞ ) = Ln 𝐺 0 - K 𝑔 (T i-c •ΔT•f) (Eq.
All the HP-PLA-derivatives exhibited a similar G0, and in the same order of magnitude to previously found for equivalent HP-PLA [17,64]. Interestingly, the PLA-SBC featured a slightly lower Kg, which might be related to the difference of the Ti-c selected to analyse the crystallization kinetics of the HP-PLA and PLA-SBC, since the nucleation rate decreases as the temperature increase [65].

Isothermal crystallization

The crystallization mechanism of merely the PLA-SBC was firstly probed at isothermal conditions for three PLA-SBC upon two different structural parameters such as Mw and tacticity (Pm), namely L500:D500, L375:D375 and L330:D330, which will be referred from now Table 3.5. Values of growth rate (G), nucleation factor (Kg) and front factor (G0) for PDLA500, PLLA500 and L520:D520.

on as high molecular weight and high tacticity (HMw-HPm), medium molecular weight and high tacticity (MMw-HPm) and medium molecular weight and low tacticity (MMw-LPm), respectively (represented by a black square, a red circle and a blue triangle, respectively, Table 3.

6).

The isothermal protocol (Figure 3 isothermal conditions previously reported (Figure 3.3 and Table 3.3) resulting in different Ti-c for each PLA-SBC derivative (Table 3

.6).

Likewise, the crystallization mechanism from the melt was firstly analysed by DSC at several temperatures revealing distinct difference among the PLA-SBC derivatives on the thermograms of the isothermal step (Figure 3.9). Moreover, the thermodynamic parameters extracted from the isothermal crystallization and subsequent cooling and melting ramps are presented in Tables A3.1-A3.3, in the appendix A3. However, it must be considered that the enthalpy of crystallization and the relative crystallinity (signed with a * symbol in Tables A3.1-A3.3), at certain Ti-c, might be inaccurate due to both an onset of the crystallization before reaching the Ti-c (unable experimentally to quench it fast enough) as well as an incomplete crystallization during the isothermal step, which impede the correct integration of the thermal peaks.

All the PLA-SBC exhibited a progressive increase of the crystallization time peak (the time at which the heat flow reach its maximum [66,67]) as the Ti-c increased. The crystallinity degree obtained from the enthalpy of crystallization (Xc-c) were c.a. 40%, 41% and 43% for MMw-LPm, HMw-HPm and MMw-HPm, respectively, independently of the Ti-c, whilst the enantiomeric blend of similar Mw PLA-derivatives exhibited crystallinity index lower than 10% under identical thermal conditions [30]. However, similar crystallinity values were obtained for HMw enantiometic PLA blends in which a pre-treatment to decrease the entanglement was applied [39].

The stereo-diblock copolymers were further cooled down to room temperature to verify if the total crystallinity was attained for the PLA-derivatives during the isothermal step and particularly, the PLA-SBC crystallised at high Ti-c could further crystallize upon cooling, exhibiting higher crystallization enthalpy for the highest Ti-c (Tables A3 HPm stereo-diblock copolymers exhibited a slight increase of Tm as Ti-c increased, probably ascribed to the formation of more perfect crystals/bigger crystals at higher temperatures [70,71] due to an increase of the polymer chains mobility. Furthermore, an exotherm/endotherm was observed prior to the melting peak of both HPm stereo-diblock copolymers that might be attributed to the recrystallization of defective crystals formed during the isothermal crystallization and/or subsequent cooling. The exotherm/endotherm events prior to the melting peak gradually shift to lower temperatures as the Ti-c decrease (black and red arrows), since the Tm of the defective crystals decrease as the Ti-c decrease.

Finally, the MMw-LPm exhibited a significantly lower melting temperature (~200°C) compared to the HPm PLA-SBC counterparts. In addition, the Tm of the MMw-LPm significantly changed with Ti-c. The MMw-LPm crystallized at 170°C exhibited exclusively a narrow-distributed endotherm centred at 200°C that splits in two (P1 and P2) as the Ti-c decreases, probably due to the formation of a lower size crystal and/or lower crystal perfection population [70,71]. The P2 remains constant in temperature with a further decrease of Ti-c, whilst P1 shifts to lower temperature until it splits again at Ti-c 150°C. P1 can be ascribed to the melting of crystals developed during the isothermal crystallization or subsequent cooling and P2 is ascribed to the melting of crystals formed in the recrystallization upon heating [30].

However, the melting of the MMw-LPm crystallized at Ti-c 140° is represented again with a single endotherm, centred at 195°C but with a much wider distribution.

A second isothermal protocol (Figure 3.11) was conducted to discard the crystallization during the cooling ramp after the isothermal step that will influence the melting behaviour.

The MMw-LPm derivative was crystallized isothermally at 160°C, as a model temperature, and subsequently heated to the melt.

A similar Tm and ΔHm was obtained for the second protocol (Figure 3.12 A) suggesting that the multiple melting endotherms observed in the PLA-SBC derivatives isothermally crystallized that might be ascribed to a recrystallization of imperfect crystals, might be originated during the isothermal step rather than during the subsequent cooling to room temperature. In addition, identical protocol was performed by modulated DSC (MDSC, Figure 3.12 B) to confirm that the endo-/exothermic transitions observed prior to the main melting peak correspond to the recrystallization process that accounts for 13,2 J/g (non-rev. heat flow)

whilst the total melting enthalpy (rev. heat flow) is 64,2J/g.

The nanostructure obtained during the isothermal crystallization from the melt was also investigated by simultaneous SAXS/WAXS experiments. The evolution of the intensity of the SAXS and WAXS patterns as surface plots during the isothermal step and the subsequent cooling to room temperature is included in the Apendix A3 for the sake of clearness (Figures

A3.4-A3.15).

The PLA-SBC derivatives crystallized in the SC phase and stacked in lamellae slightly with a low crystallinity value of 8.8%. However, the lamellae were unable to stack and thus, the distance between crystals was too high to be measured with the q range selected. The Lp of the HMw-HPm after the isothermal step at 190°C was ca. 24 nm and decreased to 18 upon cooling to room temperature. Moreover, the stack of lamellas of the HMw-HPm after the isothermal step at 200°C was undetectable, and featured ca. 15 nm upon cooling to room temperature. Moreover, the Lp of the MMw-HPm at the end of the isothermal step decrease from 28 nm to 24 nm as the Ti-c decrease from 195°C to 180°C, suggesting either an increase of the crystal thickness as the Ti-c increases or a bigger amorphous phase between crystals since the crystallinity decreased with increasing the Ti-c, as obtained by WAXS analysis. In addition, the Lp drastically decreased upon cooling to room temperature, to a constant values of 15-17 nm independently of the Ti-c, which is significantly higher that the Lp of SC-PLA formed from the enantiomeric blend of similar Mw PLA-derivatives (10-15 nm) [30],

and slightly lower to block copolymers with significantly lower Mw (8.000 g/mol) crystallized between 70 and 170°C (Lp = 7-21 nm) [72].

The MMw-LPm stereo-diblock copolymer possessed long spacing when crystallized at temperatures lower than 170°C that become slightly lower at the end of the isothermal step and was lower than the Lp of the MMw-HPm stereo-diblock copolymer, between 22 and 16 nm. Similarly, the Lp increased also as the Ti-c increased, probably due to the increase of the amorphous interphase between crystalline domains since the crystallinity degree decreased as the Ti-c increase, as observed by WAXS. Likewise, the Lp drastically decreased upon cooling reducing to constant values of 13-15 nm, independently of the Ti-c, which is still slightly higher that the Lp found for the SC-PLA formed from the blend [72], although the tacticity of the PLA-SBC is lower.

Avrami model

The 3.7).

In addition, the Avrami index "n" (Eq. 3.5) is considered to be composed of two terms: Furthermore, the crystallization rate constant k is significantly higher for HMw-HPm (0,040-0,064 min -1 ) followed by MMw-LPm (0,007 min -1 ). However, the fair difference in Ti-c

for the experiments conducted for the LPm derivatives compared to the HPm counterparts in the stereo-diblock copolymers study limits the study of the effect of the molecular weight on the crystallization rate to the later. Interestingly, SC-PLA formed from the enantiomeric blend of PLA-derivatives of similar Mw and crystallized under identical conditions, featured t1/2 values three-times higher (30 min vs 10 min) compared to the HPm stereo-diblock copolymers [30,39].

The Avrami model was also applied to the WAXS data obtained under the same thermal protocol than for DSC. Similarly, the Vc (Eq. 3.7) value was assumed to be the crystallinity obtained after the deconvolution and integration of the 1D WAXS patterns, calculated as follows:

𝑉 𝑐(𝑊𝐴𝑋𝑆) = 𝐴 𝑐𝑟𝑦𝑠𝑡 𝐴 𝑐𝑟𝑦𝑠𝑡 +𝐴 𝑎𝑚𝑜𝑟𝑝ℎ × 100 (3.7)
where Acryst and Aamorph are the area of the crystalline and amorphous reflections respectively.

Likewise, the Vc increase during the isothermal step at 190°C of the HMw-HPm was inminent and the Avrami model couldn't be applied. Moreover, the Vc evolution during the isothermal step of the MMw-LPm (Figure A3.17

Non-isothermal crystallization

The crystallization mechanism of the PLA-SBC under non-isothermal conditions was investigated to elucidate the effect of Mw and tacticity upon crystallization from the melt under industrial conditions. Firstly, the PLA-SBC were analysed by DSC and the protocol (Figure The obtained Tc were found higher for HMw-HPm, followed by MMw-HPm, which is probably ascribed to the higher molecular weight, in disagreement with the decrease of the Tc previously reported for similar SC-PLA formed from the enantiomeric blend, in which the HC-and SC-phases separated, decreasing the Tc [30]. Moreover, the Tc for HMw-HPm is 70°C higher compared to Tc values found for SC-PLA formed from the blend of homoenantiomers [30] and 30°C higher than Tc values of stereo-diblock copolymers exhibiting similar molecular weight and at similar cooling rates [75].

Furthermore, HMw-HPm polymer exhibited a crystallinity degree upon cooling of c.a.

44% independently of the cooling rate whilst the crystallinity of both MMw-HPm and MMw-LPm significantly decreased at 30°C and 50°C/min, (Table A3 crystallization rate for both MMw derivatives, being the slower the MMw-LPm. Besides, the Xc-c of the PLA-SBC exhibiting high tacticity (HMw-HPm and MMw-HPm) are slightly higher (5% over 100% of crystallinity) than similar SC-PLA derivatives formed from the blend (~40%) [75]. Furthermore, the HMw-HPm fully crystallized upon cooling even at high cooling rates such as 50°C/min, which to the best knowledge of the authors has never been reported, that increase its applicability as commodity product as the polymer processing at industrial scale typically imply high cooling rates.

Likewise, the PLA-SBC were subsequently heated to 230°C to analyse the melting behaviour of the crystallites formed under non-isothermal conditions (Figure 3 In addition, different theoretical models have been applied to the thermograms obtained upon cooling at different cooling rates to obtain further insights about the nonisothermal crystallization kinetics of the different PLA derivatives.

Avrami modified model

The Avrami equation modified by Jeziorny [66,67] is frequently applied to analyse the non-isothermal crystallization kinetics of polymers, according to which the relative degree of crystallinity (Xt) as a function of the crystallization time (Figure A3.4) can be expressed as eq.

3.8:

𝟏 -𝑿(𝒕) = 𝐞𝐱𝐩[ -𝒁 𝒕 𝒕 𝒏 ] (3.8)
where n denotes the Avrami exponent which is dependent on the crystals nucleation and growth, and Zt is a temperature-dependent constant related to the cooling rate with its linear form being eq. 3.9:

𝐥𝐨𝐠[ -𝐥𝐧 (𝟏 -𝑿 𝒕 )] = 𝐥𝐨𝐠 𝒁 𝒕 + 𝐧 𝐥𝐨𝐠(𝒕) (3.9)

However, during non-isothermal crystallization the temperature continues changing constantly and the relation between the crystallization time t and the temperature T needs to be applied, such as the eq. 3.10:

𝒕 = ( 𝑻 𝟎 -𝑻)/𝜱 (3.10)
where Φ is the cooling rate. Moreover, the crystallization rate constant should also be corrected by the cooling rate, and the modified crystallization rate constant (Zc) is expressed as eq. 3.11:

𝐥𝐨𝐠 𝒁 𝒄 = 𝐥𝐨𝐠 𝒁 𝒕 /𝜱 (3.11)
The linearization of the Avrami modified equation by the logarithmic (Eq. 3.9, Figure 3.25) exhibited a linear behaviour in the early stage of crystallization, followed by a deviation at longer crystallization times, ascribed to the secondary crystallization due to spherulite [77] and SC-PLA blends with similar Mw [43]. In addition, the Zc were found to increase for the PLA-SBC derivatives under study as the cooling rate increases, since the transition process from the melt to the crystalline state is faster when the cooling rate was higher (Table 3.8) [43].

Furthermore, the t1/2 decreased as the cooling rate was increased (Figure 3.26) in the PLA-SBC derivatives under study whereas at low cooling rates (5-20°C/min) the t1/2 was higher for the MMw-LPm, followed by the MMw-HPm, which can be ascribed to the lower tacticity and lower molecular weight, respectively. At higher cooling rates, 30-50°C/min, the t1/2 values of the PLA-derivatives should not be compared since the crystallization process was incomplete and errors have probably been introduced in the model, similarly to Avrami model applied to isothermal crystallization [74].

Ozawa model

Ozawa modified the Avrami equation and considered the non-isothermal process by incorporating the cooling rate factor as follows (Eq. 3.12):

𝟏 -𝑪(𝑻) = 𝐞𝐱𝐩 [- 𝑲(𝑻) 𝜱 𝒎 ] (3.12)
where C(T) is the relative crystallinity, K(T) is the Ozawa crystallization rate constant and m is the Ozawa exponent. The linearization of the Ozawa equation by applying the logarithms to the eq. 3.12 (Figure 3.27) (Eq. 3.13) should follow a linear behaviour if the model describes accurately the non-isothermal crystallization process.

𝑳𝒐𝒈 [-𝐥𝐧(𝟏 -𝑪(𝑻))] = 𝐥𝐨𝐠 𝑲(𝑻) -𝒎 𝑳𝒐𝒈 𝜱 (3.13)

However, the logarithmic plot of the Avrami modified equation appears as relatively straight lines at small temperature ranges and with a R 2 lower than 0.98, indicating that the model is not describing the crystallization kinetics of our PLA derivatives.

Avrami-Ozawa model

A new crystallization model was proposed to study the non-isothermal crystallization behaviour by combining the Avrami and Ozawa equation [66,67]. The relation between A3.7-A3.9). The increase of F(T) with increasing the relative crystallinity in the PLA-SBC under study indicates that at unit crystallization time, a higher cooling rate is required to obtain a higher degree of crystallinity.

Moreover, the trend of the obtained F(T) with the lower value for HMw-HPm followed by MMw-LPm suggests that MMw-HPm requires a higher cooling rate to approach an identical degree of crystallization and thus, a lower crystallization rate [78]. The α values are in agreement with previous obtained for the SC-PLA obtained from the enantiomeric blend [79] as well as their featured tight range between 0.8-1.5 indicates that the Avrami-Ozawa equation describes well the non-isothermal crystallization kinetics of PLA-SBC.

The Avrami-Ozawa model approach has been used to describe the non-isothermal crystallization kinetics of different polymers, such as PLA or PEEK [22,[78][79][80]. However, the Avrami-Ozawa model is still controversial since the two equations used to obtain the model are derived from the Avrami equation with different assumptions such as in the Ozawa equation, a constant cooling rate is assumed whilst in the Evans-Avrami equation, isothermal conditions are implied by a constant growth rate [81].

Crystallization of SBC with nucleating agents

One of the strategies to increase the crystallization rate of polymers and particularly of PLA consists of the addition of nucleating agents (NA) to reduce the energetic barrier to nucleate [82,83]. Likewise, the nucleation efficiency of three organic NA for the HMw PLA-SBC was assessed to tackle one of the major disadvantages of PLA to substitute polyolefins at the industrial scale and in particular, with the advent of novel processing technologies such as 3D

where the control of crystallization is key to tailor the properties of the final material [84].

Besides, melt-soluble bio-organic NA are generally preferred to conventional inorganic compounds as are designed to dissolve in the polymer melt and crystallize prior to the polymer on cooling, achieving higher nucleation efficiency due to the homogenous dispersion in the polymer matrix while preserving the biocompatibility of the afforded composites [23].

Likewise, the effect on the PLA-SBC crystallization of a series of low molecular weight NA featuring similar molecular groups with hydrogen bonding capability but differing on the structural rigidity and numbering of moieties was assessed. Therefore, the nucleation Firstly, the nucleating capacity of the organic NA was assessed to identify the optimal conditions to then investigate the kinetics and morphology of crystallization in the different PLA-SBC derivatives at fixed nucleating agent concentration.

Nucleating agent efficacy survey

BHET

The efficiency of a terephthalic acid-based NA, the N,N'-Bis (2-hydroxyethyl) terephthalamide (BHET) (Figure 3.29) on the crystallization of SHMw-HPm was assessed. BHET was previously reported [85] to increase the crystallization temperature of PLLA from 100°C to 120°C and the crystallization enthalpy from 34.0 to 48.3 J/g by increasing the concentration of BHET from 0 to 1.5% wt, respectively. Moreover, PLLA crystallized from the melt at 190, 170 and 140°C, at concentrations of 2.0, 1.5 and 1.0 % wt, respectively. The crystalization of films prepared with SHMw-HPm and different concentrations of BHET (0, 1 and 10% wt) was studied by DSC at fixed cooling and heating rates of 10°C/min (Figure 3.30, and the corresponding summary of the transitions in Table 3.9). The addition of 1 and 10%wt of BHET to the PLA-SBC decreased the crystallization temperature by 40 and 50°C, respectively.

Moreover, the crystallization enthalpy decreased by 11 and 14 J/g, respectively, whilst the melting temperature diminished 13 and 18°C as well as the melting enthalpy decreased 14 and 22 J/g respectively, in contrast to the nucleation effect observed in PLLA at the identical cooling rate [85]. The decrease in the thermal transition temperatures and enthalpies can be ascribed to the impingement of the BHET crystallization in the PLA-SBC melt, that in turn acts as a plasticizer, increasing the molecular mobility of the PLA-SBC and thus, decreasing its crystallization rate. Moreover, the heat capacity change during the glass transition decreased in presence of the BHET and shifted to a lower temperature (Table 3.8) confirming its plasticizing effect in the PLA-SBC matrix [86].

OXA

The efficiency of NA based on oxalamide to the enhancement of the crystallization rate of PLA-SBC was also tested due to melt solubility and higher crystallization temperature [87]. (142 J/g, crystallinity 1 ).

Oxalamide-based organic compounds (OBOC) (Figure 3.31) are an example of melt-soluble organic NA in which the end groups stimulate dissolution in the polymer matrix at elevated temperatures due to their structural similarity, reducing the viscosity of the polymer at melt, while the hydrogen bonding between the oxalamide moieties provides the driving force for crystallization from the polymer melt upon cooling, that could potentially act both as plasticizers at the melt and nucleating agents upon cooling. Interestingly, the nucleation efficiency of the OBOCs increases at high cooling rates whilst preserving the biocompatibility of PLA due to its harmless effect [87].

Films of PLA-SBC, MMw-HPm, were produced containing several concentrations (0.5, The nucleation efficiency difference between the OBOCs might be ascribed to the increase of the molecule flexibility by the increase of the number of methylene groups in- observed at 160°C by the reflections ((110), (300)/(030) and (220) (q (nm -1 ) = 8.5, 14.7 and 16.9)), and a low-intensity signal centred at q(nm -1 )=12 appeared at 100°C (Figure 3.34B) corresponding to the more intense peak of the α-phase of the homocrystallization ((110)/(200)), in agreement with the stereocomplexation occurring before homocrystallization in isothermal crystallization as previously found for similar SC-PLA formed from the enantiomeric blend [30].

Likewise, the development of the α-phase upon cooling is in agreement with the double melting peak observed by DSC, however, it melts at 75°C (Figure 3.33C-D) which is in contradiction to the α-phase melting temperature (ca. 175°C). The nanostructure of the OXA-2 was analysed to discard the correspondence of the q(nm -1 ) =12 signal to a diffraction peak of the NA, but the more intense peaks of the OXA-2 appeared at q higher than 13 nm -1 (Figure A3.31). Likewise, further analyses are required to understand the origin of diffraction reflection to the α-phase as well as its melting at 75°C.

Morphology of the crystallized NA/SC-PLA derivatives

The influence of the nucleating agent and the attained crystallization kinetics on the morphology of the PLA-SBC under study (Table 3.8) was also investigated to understand the mechanism of crystallization. Moreover, the crystallization from the melt of the PLA-SBC was also analysed by DSC at a cooling rate of 30°C/min to compare the crystallization in bulk and in film (Table A3.10). In addition, the tailoring of the crystallised morphology is a key parameter to customise their physicochemical properties such as the mechanical behaviour.

Likewise, the differences in morphology of thin films crystallised both isothermally and under constant cooling from the melt were probed by POM and WAXS using a microbeam to enlarge the length scale under study.

The morphology of the crystallization of the OXA-2 within the PLA matrix was initially studied with an amorphous PLA derivative, Rac-PLA, to assess the OXA-2 crystal growth and morphology without the influence of the PLA crystallization as well as its distribution across the film area (Figure 3 Moreover, the efficiency of OXA-2 as nucleating agent in the crystallization from the melt of the HMw-HPm PLA-SBC derivative was probed by DSC at 30°C/min (Table A3.10) as well as at high crystallization rates by coupling an optical microscope with a nanocalorimeter.

The addition of OXA-2 to the PLA-SBC derivative enabled the crystallization at cooling rates up to 400°C/min, whilst the crystallization of the neat polymer was impede at cooling rates higher than 300°C/min (Figure 3.43).

Likewise, the efficiency of the OXA-2 as a nucleating agent was evaluated in the crystallization of an asymmetric PLA stereoblock copolymer (A-UHMw-HPm-0%OXA) that crystallizes from the melt at lower temperature, 105°C, as evidenced by the exothermic signal monitored by DSC during a cooling ramp at 5°C/min (Figure A3.40). Moreover, the DSC thermograms suggest that the neat A-UHMw-HPm-0%OXA crystallizes in both HC-and SCphase as well as in their defective crystal forms (α'-and SC'-phases) as evidenced by the multiple and continuous melting endotherms monitored upon heating (Table A3.10 and Figure 3.40). Likewise, the asymmetric blend of PLLA/ PDLA enantiomers was found to exhibit two well-defined endothermic signals representing both the HC and SC phases [82,83].

In addition, the crystallization temperature of the A-UHMw-HPM upon cooling from the melt was revealed 55°C lower than the MMw-HPm by DSC (Figure A3.40), which could be ascribed to the asymmetry as well as the lower chain mobility due to the higher molecular weight of the asymmetric PLA-SBC. Similarly to symmetric PLA-SBC, the Tc of the A-UHMw-HPM remained similar with the addition of OXA-2 as revealed by POM (Figure A3.41).

Likewise, thin films of asymmetric PLA-SBC prepared at 0 and 2wt% of OXA-2, namely UHMw-HPm-0%OXA and UHMw-HPm-0%OXA, respectively, were isothermally crystallized from the melt at 120°C and 150°C, respectively, as these were the temperatures at which the crystals were large enough to be later analysed with the X-ray microbeam. The nucleation density of the neat asymmetric UHMw-HPm-0%OXA was found significantly higher than the symmetric PLA-SBC-0wt%, probably due to the lower crystallization temperature (Figure 3.44).

Likewise, the morphology of the UHMw-HPm-2%OXA thin film crystallized at 150°C has shown that the nucleating agent formed a homogeneously distributed network with a similar needle dimension. In addition, mainly heterogeneous nucleation was found for the asymmetric UHMw-HPm-2%OXA with a standard radial spherulitic growth across the film, in contrast to the perpendicular orientation observed for the symmetric SC-PLA derivatives A and B) that correspond to "defective" crystals, in agreement with the continuous melting endotherm revealed by DSC. In addition, the distribution of the "perfect" and "defective" crystals seemed heterogeneous across the film region analysed, where the "perfect" crystals distribution of one phase (blue region for the perfect SC-phase in Figure 3.47) overlaps with the "defective" crystals distribution of the second phase (Figure 3.47B, orange region for the phase), suggesting that stereocomplexation in a perfect crystal lattice leads homocomplexation in the α'-phase, whilst a disturbed SC-crystal lattice leads to defective HC to homocomplexation in the α-phase.

Conclusion

The investigation of the mechanism of crystallization of the novel PLA-SBC have revealed that:

 The increase of the Mw of PLA-SBC from 20kDa to 150kDa shifts the degradation temperature start from 220° to 240°C. Moreover, the increase of the elapsed time at the Tm to erase the thermal history of the PLA-SBC decreases both the Mw and the Tc, which might be ascribed to transesterification reactions.

 Both the Tc and Tm increases as the Mw increases. Importantly, the typical phase separation occurring in the SC-PLA formed from the blend as the molecular weight increases is absent in the high molecular weight PLA-SBC derivatives.

 The Tc obtained from the crystallization from the melt of the PLA-SBC was higher compared to both the enantiomeric counterparts (PLLA and PDLA) and the SC-PLA formed from the blend, as observed by POM. Moreover, the maximum spherulite growth rate of the HP-PLA and the PLA-SBC was found at ca. 130°C and 155°C, respectively, and a transition from the crystallization regime II to regime III was found at 170°C for the PLA-SBC, in line with previous experiments for both HP-PLA and SC-PLA formed from the blend.

 The crystallinity degree obtained from the PLA-SBC isothermally crystallized from the melt was ca. 10% higher than the enantiomeric blend of similar Mw PLA-derivatives. Moreover, the Tg of the PLA-SBC was slightly lower than the values reported for the SC-phase, suggesting a less rigid structure for the PLA-SBC compared to the enantiomeric blend.

 The PLA-SBC crystallize only in the SC-phase, as evidenced by WAXS analysis.

However, the decrease of tacticity of the PLA-SBC derivatives (Pm value) from ca. 0,93 to 0,7 decrease the Tm from 210°C to 200°C, which might be ascribed to the formation of a "defective" SC-phase due to the stereoerrors along the polymeric chain.

 The mechanism of crystallization at isothermal conditions of the PLA-SBC is fairly well described by the Avrami model using the crystallization parameters obtained by both DSC and WAXS experiments, although the dimensional growth of the crystals obtained by the Avrami index n are in disagreement with the 3D crystal spherulite observed by POM.

 Moreover, the crystallization rate constant of the PLA-SBC is significantly higher than the SC-PLA formed from the enantiomeric blend, as observed by the three-times lower t1/2 value obtained for the PLA-SBC.

 The crystallinity degree of the HMw and high tacticity PLA-SBC was 44% at high cooling rates such as 50°C/min, increasing is suitability for industrial processing.

 Similarly, the Avrami modified model reproduced the kinetics of the PLA-SBC under non-isothermal conditions, although the Avrami index contradicted the values previously found for the SC-PLA formed from the blend and the dimensional growth of the crystals found by POM. Moreover, the crystallization rate at low cooling rates (5-20°C/min) decreases as the molecular weight and tacticity decreases.

 The Ozawa model might not described the crystallization kinetics of the PLA-SBC and the Avrami-Ozawa model results are in agreement with the Avrami results, which claimed that crystallization rate increases with increasing molecular weight and tacticity.

 The BHET molecule acts as plasticizer in the PLA-SBC matrix, as observed by the decrease of the Tc and Tg of the PLA-SBC derivative.

 The OXA-2 molecule acts as nucleating agent in the PLA-SBC matrix by increasing the crystallization rate of the PLA-SBC derivatives, as observed by DSC and POM analysis. Moreover, the efficiency of the OXA-2 as nucleating agent is higher for the PLA-SBC derivatives than for the homopolymer counterparts, which might be ascribed to an epitaxial crystallization of the PLA-SBC from the nucleating agent. In addition, OXA-2 changes the PLA-SBC crystal morphology from spherulite-like crystals to row nucleation, which could be useful for tailoring the mechanical properties of the PLA-SBC.

Experimental section Materials

Solvents and reagents were acquired from Sigma-Aldrich. L-and D-lactide dimers were purchased from Rex scientific. The lactide dimers were purified three times by sublimation and then stored in a glovebox at 4°C. Toluene and tetrahydrofuran (THF) were pre-dried over sodium wire as well as distilled under nitrogen from sodium and subsequently stored over molecular sieves (3Å) in a glovebox. The N,N'-Bis(2-hydroxyethyl) terephthalamide (BHET) and the Oxalamide based organic compounds, OXA-4 and OXA-2, were kindly offered by the University of Maastricht [23,85].

Methods

Polymerizations

Polymerizations were performed as described previously in Chapter 2 [88].

Films preparation

Films containing PLA derivatives and the nucleating agents, BHET, OXA-4 and OXA-2, were prepared by solution cast. Briefly, individual 10wt% solutions in chloroform of the PLA derivatives and the nucleating agents were prepared by stirring during 1h. Subsequently, the nucleating agent solution was added to the PLA solution in different volumes to obtain the desired concentrations and the final solution was stirred for 1h. The solutions were solvent cast in a petri dish and the solvent was evaporated at room temperature. Finally, the films obtained were dry in vacuum at 80°C during 24h to eliminate all the solvent of the film.

Instrumentations and measurements

Thermogravimetric analysis (TGA)

The mass loss of the symmetric PLA-SBC was analysed by TGA with a Labsys evo instrument with a range temperature from 25°C to 1150°C. Briefly, 8-20 mg of the corresponding polymer was loaded in the TGA pans and heated from 30°C to 650°C at a constant heating rate of 5°C/min.

Gel permeation chromatography (GPC)

Molecular weight (Mn, Mw), as well as polydispersity index (PDI) of the afforded polymer derivatives, were determined from GPC using chloroform as eluent. The molecular weight characterization was performed with a Shimadzu Prominence-I LC-2030 equipped with a Shodex GPC KF-805L column (Shodex, Tokyo, Japan) and a Shimadzu RID-20A detector (Shimadzu, Kyoto, Japan). Analytical grade CHCl3 was used as the mobile phase at 40 °C, with a flow rate of 1 mL/min. GPC samples were prepared by dissolving ca. 5 mg of the corresponding polymer in 1,5 mL of solvent and left overnight under constant agitation.

Thereafter, the samples were filtered over a 0,2 µm polytetrafluoroethylene (PTFE) syringe filter prior to its injection. Polystyrene with a molecular weight of 172 kDa was used as a reference to determine the overall molecular weight of the samples.

Polarized Optical Microscopy (POM)

The morphological development during the crystallization of the samples was monitored using polarized optical microscopy on an Olympus BX53 (x20 or x50 objective) microscope mounted with an Olympus DP26 camera and a Linkam heating stage (between cross-polarisers and using a wavelength of 530 nm). The samples were heated to 200°C and 230°C, for Homopolymers and stereoblock copolymers, respectively, and held at this temperature for 3 minutes. Non-isothermal crystallization was monitored by cooling down to room temperature at 5°C/min. Moreover, isothermal crystallization was monitored during 60 minutes after cooling to the desired isothermal crystallization temperature (Ti-c) at 30°C/min.

Differential Scanning Calorimetry (DSC)

DSC measurements. Differential scanning calorimetry (DSC) were performed using a DSC Discovery 250 (TA Instruments, New Castle, DE, USA) under nitrogen atmosphere.

Temperature, heat flow and baseline were calibrated by using indium as a standard. DSC samples were prepared by loading around 5-7 mg of the polymer derivatives in Tzero Hermetic Aluminium pans (TA Instruments, New Castle, DE, USA). The thermal transitions and their corresponding enthalpies such as the temperature and enthalpy of crystallization upon cooling (Tc, ΔHc), temperature and enthalpy of cold crystallization upon heating (Tcc, ΔHcc), melting temperature and enthalpy of melting (Tm, ΔHm) as well as the glass transition temperature (Tg) were calculated from the obtained thermograms.

The thermal degradation protocol consisted of a first heating ramp at 10 °C/min from 25 °C to 230 °C and held for 2, 5, 15, 30 and 60 min. Samples were then cooled down to 25 °C at a constant rate of 10 °C/min, heated to 230°C at 10°C/min and cooled down again at 10°C/min. The melting behaviour of all samples was observed by a new heating and cooling at 10°C/min.

The non-isothermal protocol consisted of a first heating ramp at 10 °C/min from 25 °C to 230 °C and held for 2 min at the melt to erase thermal history. Samples were then cooled to 25°C at different cooling rates (1, 10, 15, 20, 30 and 50°C/min).

The isothermal protocol 1 consisted of a first heating ramp at 10 °C/min from 25 °C to 230 °C and held for 2 min at the melt to erase thermal history. Samples were then cooled to the isothermal crystallization temperature (Ti-c) at 50°C/min, held 60 minutes and cooled down to 25°C at 2°C/min. Subsequently, the samples were heated at 10°C/min to 230°C and cooled down again to 25°C at 10°C/min. A variation of this protocol (isothermal protocol 2) was implemented by heating to 230°C right after the isothermal crystallization at the Ti-c.

Modulated-Differential Scanning Calorimetry (M-DSC)

M-DSC analysis were performed using a DSC Discovery 250 (TA Instruments, New Castle, DE, USA) under nitrogen atmosphere. The heat capacity signals were calibrated using the sapphire standard. M-DSC samples were prepared by loading around 5-7 mg of the polymer derivatives in Tzero Hermetic Aluminium pans (TA Instruments, New Castle, DE, USA). The isothermal protocol 2 was applied and modulation of the heat flow was applied to the heating ramp after the isothermal step, with 120s period, 1°C/min heating rate. The measured intensity in the 2D detectors was reduced to 1D intensity profiles using bubble [21] as a function of the scattering vector (q = 4π/λsinθ) and reported in arbitrary units by correcting by the incident intensity and background subtraction.

Wide-Angel X-Ray Scattering (WAXS) and

The thermal protocol was performed with a Linkam stage DSC 600 with a 0,7x3 mm window.

The isothermal protocol was similar to the isothermal protocol 1 performed by DSC without the last heating and cooling ramps (due to the limited beamtime available). The nonisothermal protocol was also similar to the DSC non-isothermal protocol, with only one heating and cooling ramp.

The deconvolution of the 1D WAXS profile between 6 and 20 nm -1 was accomplished with Lorentzian functions for the crystalline plane reflections and two pseudo-Voigt functions for the amorphous halo, as previously described for stereocomplexed PLA [34]. Moreover, the crystallinity values from the 1D WAXS patterns were obtained by the equation XX:

𝑋 𝑐 = 𝐴 𝑐𝑟𝑦𝑠𝑡𝑎𝑙𝑙𝑖𝑛𝑒 𝐴 𝑐𝑟𝑦𝑠𝑡𝑎𝑙𝑙𝑖𝑛𝑒 + 𝐴 𝑎𝑚𝑜𝑟𝑝ℎ𝑜𝑢𝑠 × 100
where Acrystalline and Aamorphous are the area under the curve of the Lorentz and pseudo-voigt equations.

Furthermore, the SAXS intensity was corrected with the Lorentz correction function (q 2 I(q)) and the Long period (Lp) was obtained with the position of the correlation peak, qmax.

The Lp, defined as the distance between crystalline lamellas, can be obtained from the equation:

𝐿 𝑝 = 2𝜋 𝑞 = 𝐿 𝑐 + 𝐿 𝑎
Where Lc and La represent the thickness of the crystalline and amorphous fraction,

respectively. An approximation of the lamellar thickness can also be obtained from the equation:

𝐿 𝑐 = 𝐿 𝑝 𝑋 𝑐

Scanning X-ray micro-diffraction (micro-beam)

Scanning X-ray micro-diffraction experiments were carried out using the microbeam facility at the ID13 beamline of the European Synchrotron Radiation Facility (ESRF) (Grenoble, France). The micron sized X-ray beam was focused to a spot size of 2 x 2 μm² in along both directions, (H x V), using compound Beryllium lenses. X-ray photons with energy of 13 keV were used in the experiments. 2D diffraction data were collected using an Eiger4M. The scanning of the samples was done in transmission geometry with the sample surface normal to the X-ray beam. After identification of the region of interest in the film sample by help of an on-axis optical microscope, the sample was using a high precision XYX gantry scanned in continuous mode. In this mode the sample is moved continuously through the beam focus of the X-ray beam while after each line scan the sample was vertically offset by an increment for the next line scan. The size of the illuminated sample volume is defined by the counting time of the acquisition scheme and the vertical displacement increment. The parameters were chosen to achieve a resolution in the two dimensional scan of 5 x 5 µm². The measurement geometry including the norm of the scattering vector q, where |q| = 4π sin(θ)/λ, was calibrated using several diffraction peaks of alpha-alumina using the PyFAI package available at the ESRF including data reduction and integration. The data analysis and visualization 2D diffractograms were done using home-built routines designed in Igor Pro software (Wavemetrics Ltd.).

Nanocalorimeter

A custom-built ultrafast microelectromechanical systems (MEMS)-based nanocalorimeter instrument [89] was used to study the crystallization of PLA-SBC from the melt a high cooling rates. Briefly, 30-40 ng of the PLA-SBC were positioned within the active area of the nanocalorimeter sensor using a micromanipulator. PLA-SBC derivative samples were heated to 230°C and cooled down to room temperature at cooling rates from 1000°C/min to 40°C/min. The crystallization of the PLA-SBC derivative was evaluated visually by a polarized optical microscopy coupled to the nanocalorimeter sensor.

Introduction

Nanoparticles (NP) emerged in the 1970s as a drug delivery system (DDS) approach to improve the stability and solubility of the encapsulated molecules, promote drug delivery across cellular membranes as well as enhance drug circulation duration, improving safety and efficacy [1]. The composition of the designed nanoparticles has ranged from natural materials such as chitosan, dextran and gelatin or inorganic materials such as gold, silica and Cd/Zn as well as synthetic polymeric NP [2]. In particular, polymeric NP have recently received great interest due to the possibility to tailor the drug release rate by varying polymeric architectural parameters such as crystallinity, hydrophilicity or glass transition temperature [3]. Polylactide (PLA) is among the most commonly used polymers for nanoparticle formulation development due to its biocompatibility. As described in Chapter 1, PLA is a bio-based and polymorphic polyester comprising a chiral carbon in its structure that yields a wide variety of nanostructured derivatives (isotactic, atactic, syndiotactic, stereo-complexed, etc.) [4][5][6] featuring different physicochemical properties. In addition, the processing conditions during the formation of the NP determine the kinetic trapped or thermodynamic state of the final material that affect significantly the final polymer characteristics such as crystallinity degree and crystalline phase (Figure 1.2, Chapter 1), which influence the drug release rate [3].

Particularly, polyesters exhibiting low glass transition temperature (Tg), melting temperature (Tm), and degree of crystallinity generally feature higher drug release [7,8]. Furthermore, amorphous PLA is characterized by soft and elastic mechanical properties, whilst semicrystalline PLA is stiff and robust [9], which will invariably define the drug release.

Likewise, the drug release rate from PLA nanoparticles exhibiting stereocomplex (SC) phase is expected to decrease compared to the HP counterpart due to the stabilization by specific C-H…O-H hydrogen bonds of the crystalline phase that increases its hydrolytic resistance [10].

However, the impact of the processing parameters of the formation of the nanoparticles on the final nanoparticles release kinetics is generally found inconsistent due to the lack of mechanistic process control as well as the unusual analysis of the structure-properties relationship of the formulated nanoparticles. However, systematic studies correlating the structure-processing-properties relationship of nanoparticles formulation have been recently

accomplished [3,11].

The hydrophobic character of PLA increases the contact angle with water, which diminishes the interaction with physiological fluids and cell adhesion, which is commonly faced by copolymerizing PLA with hydrophilic monomers such as polyethyleneglycol (PEG)

or poly-glycolic acid (PGA) to afford block copolymers with enhanced biocompatibility [12][13][14]. Recently, nanoparticles generated from poly-L-lactide (PLLA)-3-ethylglycolide and poly-D-lactide (PDLA)-3-ethylglycolide were successfully formulated featuring the SC phase as well as attaining crystallinity degrees up to 50%. However, a long (7 days) stereocomplexation process from solution was required to obtain crystallized nanoparticles [15]. In addition, PLA micelles crystallized in SC-phase were also achieved from PEG-PLLA-PDLA stereoblock copolymers by the nanoprecipitation method, which exhibited a slower drug release rate than the amorphous and isotactic counterpart [16]. However, the copolymerization of PLA with PGA or PEG usually deters the crystallization of the nanoparticles, limiting the physicochemical properties of the PLA-derivatives to the characteristic features of the amorphous phase [15] that neutralize the PLA chirality effect on the molecular mobility and the relaxation process [17]. Alternative processing methods are urged to favor the hydrophilicity of the PLA-based material whilst retaining the tunability of the physicochemical properties for the nanostructured PLA derivatives. Several bio-based materials have been previously used for the fabrication of bioscaffolds such as collagen, chitosan, alginate, gelatin, silk, etc. [18]. Alginate is an attractive material for biomedical applications due to its structural similarities to polysaccharides; however, it is unable to meet concomitantly the different required parameters for tissue engineering applications such as the degradation degree, the mechanical properties and bioactivity [19]. Similarly, collagenbased scaffolds for the regeneration of tissues such as nerves, bone, cartilage or skin were also developed with the detrimental low solubility in water that turns gelatin into a better candidate for tissue engineering [20]. Gelatin is a natural biopolymer obtained from the denaturation of collagen [21] that preserves key bioactive functions of collagen for biomedical applications, such as cell adhesion and proliferation [22]. Moreover, gelatin is widely used in the food and cosmetic industries, and particularly, fish gelatin is becoming increasingly used for biomedical applications due to a likely less intense immune reaction in the human body than mammalian gelatin [23] as well as an approach to revalorize the by-products from fish industrial processes [24]. Moreover, gelatin can be processed in several formats depending on the intended applications, such as electrospun matrices for wound healing [25], microspheres for being loaded with a growth factor for bone regeneration [26], nanoparticles as nanocarriers for anticancer therapeutic agents [27], protein, gene and vaccines delivery [28], DDS for regenerative medicine or three-dimensional scaffolds for drug screening [29]. PLA/gelatinbased scaffolds were frequently used as a matrix for tissue engineering or drug delivery.

Recently, tetracycline-loaded PLLA microparticles embedded in a porous gelatin scaffold exhibited controlled release behavior, biocompatibility as well as antibacterial activity [30].

Furthermore, an electrospun composite scaffold of polycaprolactone and gelatin with curcumin-loaded PLA microparticles exhibited biocompatibility and good mechanical properties for wound-healing applications [31]. However, the design of bionanocomposites containing crystalline PLA nanoparticles was generally dismissed as an approach to tailor the drug delivery rate.

Aims and objectives

The main objective of Chapter 4 is the generation of a series of bionanocomposites composed of shark gelatin and PLA nanoparticles featuring different nanostructures to design multifunctional drug delivery systems with tailored release rates and local applicability (Scheme 4.1). To accomplish the main objective, the following steps were followed, in

Results and discussion

As state in the introduction, the discussion of the main results and the conclusion will be presented firstly and the experimental techniques will be briefly described at the end of the chapter. Moreover, some of the data herein described have already been discussed in Chapter 2 and 3, but they will be repeated for the sake of convenience and fluidity when comparing results. Furthermore, since the amount of nanoparticles obtained by the double emulsion method is limited, not all formulations could be analysed by all tests. However, triplicate measurements have been performed to ensure the reproducibility of the data. The correspondence between the abbreviated name of the samples and its compositions is described in Table 4.1. LA ratio was estimated to be higher than 98% for the HC-PLA and close to 50% for rac-PLA and SC-PLA, as analyzed by polarimetry measurements (Table 4.2). Furthermore, the Pm values, known as the probability of isotactic enchainment in which Pm = 0.5 is a random insertion whilst isotactic enchainment is designated by Pm = 1 [33], were obtained from deconvoluted homonuclear decoupling 1 H NMR spectra (Apendix A2) and confirmed the atactic character of the rac-PLA and the isotactic character of HC-PLA and SC-PLA (Table 4.2). Firstly, the nanostructure of the bulk polymer derivatives was characterized to understand the structure-properties relationship of the selected PLA derivatives to design different release rates. The thermal transitions of the bulk polymers were studied by DSC (Table 4.2 and Figure 4.1) to understand the crystallization mechanism that will determine the nanostructure generated at the formulation conditions. Rac-PLA exhibited a Tg at 24 °C, which was notably lower than values reported in the literature [36], which is likely due to the lower molecular weight [37], whilst isotactic PLA-derivatives, namely HC-and SC-PLA, exhibited a Tg c.a. 40 °C. Additionally, only the HC-and SC-PLA exhibited crystallization and melting transitions, as both the Tc and Tm of SC-PLA were 13 °C higher than those of HC-PLA (Figure 4.1A). Notably, the HC-PLA exhibited a double endothermic peak at 161 °C (PHC1) and 167 °C (PHC2), corresponding to the recrystallization of the α'-into the α-phase and the subsequent melting of the latter [38]. Moreover, the SC-PLA exhibited also a double melting peak (at 180 °C (PSC1) and 195 °C (PSC2)) likely ascribed to the melt-recrystallization process of the α-phase into the stereocomplex (SC) phase [39]. The melting enthalpy for both HC-and SC-PLA was ca. 47 J/g. However, the crystallinity was 20% higher for the HC-PLA due to the lower ΔHm° value.

Furthermore, the crystal phase of the bulk polymers was analyzed by Wide-Angle X-Ray Scattering (WAXS) (Figure 4.1B). Only an amorphous halo was found for rac-PLA in agreement with the atactic character obtained from the Pm value. The characteristic pattern of the α-phase at (010), (110)/(200), ( 203) and (015) (q (nm -1 )= 10.65, 11.9, 13.6 and 15.7) was observed for the HC-PLA, whilst the SC-PLA showed high-intensity reflections of the SC phase at (110), (300)/(030) and (220) (q (nm -1 ) = 8.5, 14.7 and 17)) and a low-intensity reflection at (110)/(200) characteristic of the α-phase. The presence of both α-and SC-phases on the SC-PLA, which is in agreement with the double-melting peak observed by DSC, is likely due to the typical crystallization competition occurring between homo-and stereocomplexation [39] when a quenching process occurs, such as the precipitation with methanol after the synthesis.

In addition, Raman spectroscopy analysis of PLA bulk polymers (full spectrum in A4.1). In particular, the vC-COO stretching vibration band in the spectral range between 850 and 1000 cm -1 characterizes the helix conformation of the PLA polymer chain (Figure 4.2B). Likewise, the vibrational band assigned typically to vC-COO stretching at 870 cm -1 was found to shift to a higher wavenumber, to 880 cm -1 , for SC-PLA as usually found for PLA in the stereocomplex phase [40]. Moreover, a second vibrational band of medium intensity, attributed to the coupling of the vC-C backbone stretching mode with the CH3 rocking mode, is observed either at 920 or 908 cm -1 for HC-PLA and SC-PLA, respectively, whilst it remained absent for the rac-PLA, reflecting that the PLA chain retains confirming its amorphous nature, whereas the HC-PLA carbonyl spectral region splits in three overlapping bands at 1755, 1766 and 1777 cm -1 (Figure 4.2D), as reported previously for isotactic PLA. In addition, the carbonyl band of SC-PLA appeared centered with a strong and sharp band at 1752 cm -1 , which arose from the symmetric stretching vibrations [15], which are due to the stabilization of the stereocomplex crystalline phase by hydrogen bonds [10], and a broad shoulder at ca. 1770 cm -1 resulting from the vibration of the amorphous component (Figure 4.2E).

Nanoparticles formulation

Polymeric nanoparticles from the selected PLA derivatives (NP-PLA) were successfully formulated with two different antitumoral compounds by the double emulsion method and their size and morphology were characterized by DLS ( An averaged unimodal hydrodynamic diameter of 254 ± 22 nm with narrow PDI (0.10 ± 0.03) was obtained for all formulations by DLS. Furthermore, fairly spherical nanoparticles with a roundness index of 0.86 (being 1 perfectly round) and an average diameter of 104 ± 4nm

were observed for the generated nanoparticle formulations by SEM. The size discrepancies observed by both techniques are related to shrinking effects during the freeze-drying and Aushadowed treatment required for the SEM characterization, whilst DLS measurements were conducted in solution, resulting in NP swelling. In addition, the hydrodynamic diameter calculated by DLS arises from the diffusion coefficient of nanoparticles in suspension, in which small amounts of aggregates or dust are known to disturb the size determination [42,43].

Additionally, DLS analysis was performed at longer times to verify (PDI values remained constant with time) their stability in solution and thus the NP storage capability prior to their A4.5). In addition, the nanostructure of the formulated NP-PLA was characterized to understand the effect of the processing conditions on the PLA crystallization mechanism and its correlation with the release profiles.

Nanoparticles characterization

The nanostructure of the formulated NP-PLA was characterized to understand the effect of the processing conditions on the PLA crystallization mechanism and its correlation conducted to generate the NP-PLA that prevents the polymer chains to nucleate and form the polymer crystallites [15]. However, the WAXS profiles of the NP-PLA/DOX exhibited crystalline reflections at high q values (q (nm -1 ) = 19. Notably, an exothermal peak is observed prior to the melting peak of NP-HC-PLA that can be attributed to the recrystallization of the α'-phase into the α-phase [38], which is similar to the results obtained for the bulk polymer. Moreover, a double-melting peak was observed for the NP-SC-PLA, which was likely ascribed to the melt-recrystallization process of the α-phase into the SC phase, as found for the bulk polymer. The crystallinity (Xc) degrees obtained for NP-HC-PLA and NP-SC-PLA were around 23 and 31%, respectively. 200) and (203) (q (nm -1 ) = 11.9 and 13.6)) and the SC-phase for the NP-SC-PLA ((110), (300)/(030) and (220) (q (nm -1 ) = 8.5, 14.7 and 17)), indicating that the double-emulsion method did not impede the nanoparticles by possible degradation to further recrystallize upon heating.

However, the annealing step to crystallize the freeze-dried nanoparticles leads to aggregation, deleting the pharmaceutical advantages of the nanoparticles for drug delivery applications. A simple method to prevent the aggregation of the NP-PLA upon heating is to disperse them in a matrix, which should enhance their hydrophilicity and ideally promote cellular growth as well as endorsing the DDS with storage capability and easy applicability.

Bionanocomposite formation

Gelatin is a natural polymer featuring inherent biocompatibility, uniformly distributed interconnected porosity and appreciable mechanical strength that is ideal for drug delivery and tissue engineering applications. Likewise, the SAXS intensity is dominated by hydrogel nanostructure as a result of the likely low scattering of NP-PLA internal structure considering its low concentration (Figure 8C-D).

The measured SAXS profiles were fitted to fractal aggregates with a radius of gyration of ca. 1 nm and a fractal dimension of 1 with a correlation length of ca. 6 nm. In addition, the fibrillar structure of the triple helix undergoes a transition upon thermal treatment to a higher compact fractal cluster with a radius of gyration of firstly ca. 0.33 nm and then 0.65 with a fractal dimension of 2.5 and 3 and a correlation length of ca. 1 nm, indicating the transition of the gelatin chains to a random coil during the isothermal step at 65 °C (Table A4. The morphology of the gelatin hydrogel as well as the corresponding bionanocomposites was analyzed by cryo-SEM to probe the effect of the annealing step on both the NP-SC-PLA and the gelatin hydrogel network (Figure 4.9). A porous network was found for both the shark gelatin hydrogel and the bionanocomposites in agreement with previous tuned gelatin hydrogel [27]. However, the size of the pores showed that the NP-SC- In addition, the so-called "syringe test" [45] into a potential biomaterial for liver tissue regeneration [46] as well as for the distribution into a complex anatomical human body area in particular after tumor resection where a combination of therapeutic agents is required to be administrated with low invasive techniques and that concomitantly promote the cellular growth.

Drug release analysis

Finally, the release of the loaded doxorubicin from the NP-PLA/DOX as well as from the GE/NP-PLA/DOX was monitored in PBS at 37 °C by UV-Vis spectrometry (Figure 4.12).

Shark gelatin absorption spectra present a UV-Vis band at 215 nm with a shoulder at 290 nm, whilst the doxorubicin and PLA main absorption bands are located at 490 nm and 231 nm, respectively [47], which permits the monitoring of the drug release profiles. All the NP-PLA/DOX (NP-HC, NP-SC and NP-rac) released completely the loaded doxorubicin within a few hours, with a significant burst release step during the first 2 h (98.9%, 60.9% and 51.6% for NP-HC, NP-Rac and NP-SC, respectively, Figure 4.12) that might be related to the formation of internal NPs microvoids due to heterogeneous structural relaxation during physical aging.

The microvoids could be then filled with water when exposed to an aqueous solution that promotes the doxorubicin dissolution from the inner NP region [48]. Moreover, the enthalpy relaxation values related to physical aging found by DSC (1.44 J/g, 1.2 J/g and 1.06 J/g for NP-HC, NP-Rac and NP-SC, respectively, Figure 4.5A) are in agreement with the burst release differences between formulations, being both the enthalpy and the burst release higher for NP-HC followed by NP-Rac (Figure 4.12).

In contrast, the doxorubicin release from the GE/NP-PLA/DOX exhibited a sustained pattern similar to all the NP formulations (≈5% in 120 h) that was slower than the doxorubicin release from the gelatin network itself, confirming the tailoring of the drug release profiles.

Moreover, GE/NP-SC-PLA/DOX exhibited slightly higher release, followed by GE/NP-HC-PLA/DOX and GE/NP-Rac-PLA/DOX, which might be related to the crystallinity difference found by DSC (Figure 4.6B and Table 4.3), since the doxorubicin molecule might interact preferentially with the amorphous phase of the nanoparticles, resulting in a higher doxorubicin/amorphous phase ratio for the more crystalline NP, namely GE/NP-SC-PLA/DOX followed by GE/NP-HC-PLA/DOX, and hence, a slightly higher release rate (

Conclusions

A straightforward methodology to anneal the PLA nanoparticles to customize the drug release rate was developed by generating bionanocomposites with shark gelatin hydrogels.

 The nanoparticles were successfully formulated with different PLA derivatives and loaded with two different antitumoral compounds but yielded amorphous PLA materials upon formulation, erasing the designed nanostructures.

 The dispersion of PLA nanoparticles in the gelatin hydrogel permits their thermal annealing and thus, their crystallization, without the detrimental agglomeration that diminishes their therapeutic effect.

 The SAXS and WAXS profiles as well as the cryo-SEM, DSC and rheological measurements of the bionanocomposites suggested the stability and reversibility of the hydrogel network upon the incorporation of the nanoparticles.

 The bionanocomposites generated were proved to feature storage capability at room temperature as well as being able to be injected by a syringe, which provides the possibility to be administrated locally.

Methods

Gelatin extraction

For the production of gelatin (GE), frozen skins of BS were cut in portions of around 5 × 5 cm and, after defrosting, were washed with water for 30 min under orbital agitation to remove impurities. Then, sequential chemical treatments were applied to skins, in all cases using a (1:4) solid: liquid ratio and continuous agitation (50 rpm). Each chemical stage was run for 30 min at 22 °C, intercalating water washings in between treatments. The first treatment of skins was performed with 0.05M NaOH, which was followed by acid processing using a solution of H2SO4 (0.02 M) and ending with another solution of citric acid (0.052 M). Thereafter, skins were soaked in water (1:2 ratio) at 45 °C for 16 h and gentle stirring to achieve the gelatin extraction in an aqueous medium. The obtained gelatin solution was then separated from skins by filtration (500 µm) and mixed with active charcoal (1.5%, w/v) for 3 h under constant agitation (100 rpm). Charcoal and impurities were removed by centrifugation (8000× g/20 min), and the gelatin was dried in an air-forced convection oven at 50 °C for 72 h.

Polymerizations

Polymerizations were performed as described previously in Chapter 2 [32].

Nanoparticles formulations

PLA nanoparticles (NP-PLA) were formulated from the three PLA derivatives, racemic-PLA (NP-rac-PLA), homopolymer-PLA (NP-HC-PLA) and stereo-diblock-PLA (NP-SC-PLA). In addition, all the PLA derivatives NP were loaded with the antitumoral doxorubicin and dasatinib as reference drugs [49][50][51]. Nanoparticles were prepared following the double emulsion methodology. Briefly, 10 mg of PLA was dissolved in 4 mL of CH2Cl2 and 1 mL of mQ water was added; then, the mixture was shaken in a Vortex and subsequently sonicated in an ice bath for 1 min to avoid emulsion overheating. The pre-emulsion generated was added to 10 mL of PVA (1%), shaken in a Vortex and sonicated for 5 min. The emulsion was allowed to evaporate organic solvent at room temperature for 45 min in a fume hood; then, it was centrifuged at 4 °C and 15k rpm for 20 min. Nanoparticles were collected and

SEM

NP-PLA were freeze-dried and Au-shadowed before SEM images were recorded on a Jeol 6490LV electron microscope to evaluate the size and morphology of the particles.

The resulting images were analyzed using Digital Micrograph™ software from Gatan.

Moreover, NP-PLA were quantified using ImageJ software [53]. Filters were applied to the images to normalize the contrast and highlight the edges of the nanoparticles. Then, a suitable threshold was used to binarize the images. A Watershed binary image segmentation was applied to separate touching particles. Particle agglomerates were generated in most cases, since sample preparation was complicated, individual particles were considered for sample characterization/quantification using image analysis, and large particle agglomerates were avoided as much as possible.

Cryo-SEM

The nanostructures of the GE/NP-HC-PLA, GE/NP-SC-PLA and GE/NP-Rac-PLA were evaluated in cryogenic mode. The hydrogels were mounted on the 13 mm gold sample holder and frozen at -200 °C with liquid nitrogen. Subsequently, the samples were transferred under liquid nitrogen in a cryogenic box to the Baltec equipment (MODEL, MED-020) and cut into 2-3 mm thickness slices. Finally, the samples were analyzed with an SEM (JEOL JSM-6700) at an acceleration voltage of 5 kV.

Cryo-SEM images at a magnification of ×2000 were selected for the quantitative analysis due to the presence of a significant number of pores as statistically representative of the system. The coral-like porous structure in the micrographs was examined using ImageJ software [53]. Filters were applied in order to highlight the edges of the cells and be able to threshold and segment them. Several closing steps were applied also to clean artefacts in the images and only pores having an equivalent diameter bigger than 0.1 μm were quantified to plot the cell size distributions.

DSC measurements

Differential scanning calorimetry (DSC) analyses were performed using a Q2000 DSC (TA Instruments, New Castle, DE, USA). The calibration was performed by using indium as a standard (Tm = 156.61 °C). DSC samples were prepared by sealing around 5 mg of the samples in Tzero Hermetic aluminum pans (TA Instruments, New Castle, DE, USA). The thermal protocol performed for bulk polymers consisted of a first heating ramp at 5 °C/min from 10 to 200 °C for Rac-PLA and HC-PLA, and from 25 to 220 °C for SC-PLA, followed by a cooling ramp at 5 °C/min to 10 °C and 25 °C, respectively. Moreover, the thermal protocol performed for NP-PLA and bionanocomposites (Gelatin with NP) consisted of a first heating ramp at 5 °C/min from 10 to 65 °C, followed by an isothermal step for 60 min. Samples were cooled down to 10 °C at 5 °C/min and further heated at 5 °C/min to 200 °C for the Rac and HC samples and to 220 °C for the SC samples. The thermal transitions (Tg, Tc, and Tm) and their corresponding enthalpies (ΔHc, ΔHm) were calculated from the obtained thermograms.

Raman spectroscopy

Raman measurements were performed with a micro-Raman spectrometer LabRAM HR from Horiba Jobin with a laser at 532 nm. The spectral resolution was around 1 cm -1 and a 50× microscope objective was used. The measured intensity in the 2D detectors was reduced to 1D intensity profiles using bubble [21] as a function of the scattering vector (q = 4π/λsinθ) and reported in arbitrary units by correcting by the incident intensity and background subtraction.

Small-angle X-ray scattering (SAXS) and wide-angle X-ray scattering (WAXS)

The reduced SAXS experimental data of the hydrogel derivatives were adjusted to a heterogeneous sphere with fractal structure defined with a radius of gyration 𝑅𝑔 related to the fractal size as well as the fractal dimension D that represent their compactness [54] by an inhouse program [55,56] connected to the minimization routine used by the MINUIT [57] program conceived in CERN, and thus, the SAXS profile intensity (I(q)) was fitted to:

𝐼(𝑞) = (𝐷-1) sin ((𝐷-1) tan-1(𝑞,ξ)) (𝐷-1)𝑞ξ (1+𝑞2ξ2) (𝐷-1) 2 with 𝜉 2 = 2𝑅𝑔 2 𝐷(𝐷+1)
𝜉 being the correlation length.

Rheological experiments

The mechanical properties were determined using a Physica MCR 101 rheometer (Anton Paar, Graz, Austria) able to control torques between 0.5 μN•m and 125 mN•m. Strain and sweep frequency sweeps for GE, GE/NP-Rac-PLA, GE/NP-SC-PLA and GE/NP-HC-PLA were performed using a cone-plate geometry (CP 25-1) with a constant space of 0.048 mm. The CHAPTER 4. APPLICATION 199 temperature was controlled using a Peltier P-PTD 200, which was placed at the lower plate [58]. The rheological experiments were repeated three times to analyze the reproducibility of obtained results, and the deviations in (complex) viscosities between repetitions were lower than 3%.

The linear viscoelastic range was previously determined by performing a strain sweep experiment in the range from 0.1 to 1000% at a constant angular frequency of 10 rad/s at 20 °C.

The storage modulus G′ and the loss modulus G″ were determined in the linear strain range using a constant 0.1% strain. Frequency sweep measurements were conducted from 0.05 to 200 rad/s applying a constant strain of 0.1% at 20 °C. The solutions were preheated to 60 °C for 30 min for all rheological experiments to homogenize the samples and ensure adequate filling of the geometry as well as avoid air bubbles.

The so-called "syringe test" experiments were performed using a plate-plate geometry (PP25/S) with a plate diameter of 25 mm and a geometric gap of 0.1 mm. A three-step experiment [45] was performed to evaluate the response of the gelatin-based bionanocomposites to mimic the injection in the human body of the DDS. Firstly, samples were subjected to a constant shear rate of 0.1 s -1 at room temperature, which simulates the resting state of the formulation under storage conditions and in the injection syringe. Subsequently, samples were heated up to 37 °C at a constant shear rate of 100 s -1 , which simulates an injection using a needle with a diameter equal to 0.9 mm, which was typically selected for the administration of hydrogel formulations. Finally, a shear rate of 0.1 s -1 was imposed at 37 °C, simulating the rest state of the GE/NP-PLA hydrogel mixture in the human body.

Release studies

For NP-PLA/DOX, 10 mg of lyophilized DOX-loaded NPs were placed in the dialysis membrane (molecular weight cut-off: 3500 kDa) and incubated in 8 mL of phosphate-buffered saline (PBS, pH 7.4). The suspension was incubated at 37 °C with continuous stirring (50 rpm) in an IKA incubator shaker KS 3000. At different intervals of incubation, 1 mL of releasing medium was removed to measure the DOX absorbance in a UV-Vis Carey 100 spectrophotometer at 470 nm, and new PBS was added to maintain the volume constant.

For GE/NP-PLA/DOX, 1 mL of gelatin loaded with 5 mg of DOX-loaded NPs was placed in a dialysis membrane and released in the same method described above for the NP-PLA/DOX counterparts. the PLA-SBC crystal morphology from spherulite-like crystals to oriented crystals due to the row nucleation on the OXA-2 crystals, which is key for tailoring the mechanical properties of the PLA-SBC.

Finally, a series of bionanocomposites composed of gelatin hydrogel and PLA nanoparticles featuring different nanostructures were designed to obtain multifunctional drug delivery systems with tailored drug release rates. PLA-derivatives nanoparticles were successfully formulated with two antitumoral drugs by the double emulsion method that yielded an amorphous structure independently of the chain architecture, erasing the designed nanostructure. However, post-formulation thermal annealing enabled the crystallization of the nanoparticles in the corresponding crystal phase, although detrimental nanoparticles agglomeration occurred that diminished their therapeutic effect. A straightforward methodology to anneal the nanoparticles and avoid aggregation was developed by the dispersion of the PLA nanoparticles in the gelatin hydrogel at 65°C. The gelatin serves as a mechanical support for the storage and local administration of the drug delivery system as well as a support to complementary drugs with different release profiles. However, the drug release profiles of the annealed PLA nanoparticles within the gelatin matrix were similar independently of the polymeric nanostructure, although long-term drug release analysis was unable due to the saturation of the UV-drug signal resulting from the disintegration of the gelatin.

The wide field covered by the herein thesis from the synthesis and the posterior full characterization to the application was evaluated from a multidisciplinary approach that limited the focus on a particular scope in return and thus, future lines of research are envisaged to enlarge the knowledge of the project.

The inorganic initiator employed for the synthesis enabled the polymerization of PLA derivative for the biomedical field due to biocompatility of Zn. However, "greener" synthetic approaches might be considered by the use of organic initiator that excludes metals, such as 4-(dimethylamino)pyridine, which were already reported to successfully polymerize PLA, although the synthesis of stereo-block copolymers have not been achieved yet. Moreover, the control of the synthesis enables the polymerization of multiblock-PLA (tri-, tetra-, penta-, etc.) which will exhibit tailored properties, as well as hetero-block copolymers with complementary properties such as PEG, PGA, PCL. Furthermore, the yield of the polymerization at the laboratory scale was of few grams limiting certain properties analysis that required higher yield, such as the mechanical test. Therefore, scaling the synthesis process to obtain a greater amount of polymer per batch is required to continue the characterization of the PLA-SBC as well as its processing through new processing techniques such as 3D printing.

The analysis of the thermal stability assessment performed was mainly focused on the detection of likely degradation occurring during the thermal treatment during the analysis of the mechanism of crystallization from the melt. However, a deeper analysis to understand the degradation process as well as to quantify the detrimental transesterification reactions occurring upon thermal treatment is necessary to establish the structure-processing-properties relationship. In addition, the data acquired to study the crystallization mechanism should be investigated further to describe the lamellar crystal thickness by the application of the correlation function to the SAXS profiles. In addition, the crystallization mechanism from the glass state or solution needs also to be addressed. Likewise, nano-beam analysis at the interface of the OXA-2 nucleating agent with PLA-SBC crystallized should be addressed to tackle the question of the higher nucleation efficiency of the PLA-SBC copolymer compared to the homopolymer counterpart. The analysis of the crystal phase orientation of the PLA-SBC on the OXA-2 crystal might clarify whether the row nucleation observed at the micro-scale is due to a geometric effect or an epitaxial growth.

Furthermore, the designed bionanocomposite drug delivery system should be studied with a different detection method to analyse the drug release at a longer time, such as highperformance liquid chromatography (HPLC), since the gelatin signal saturated the UV-drug detection. In addition, new methods to improve the crystallization of the nanoparticles upon formulation could be envisaged as well as a deeper analysis of the mechanistic governing the formulation of PLA nanoparticles from a thermodynamic point of view. Finally, the potential application of the PLA-SBC in the biomedical field includes the fabrication of stents, which have already been fabricated based on PLLA, although yielding devices with a lack of radial stress that might be overcome by the enhanced mechanical properties presented by the PLA stereocomplex phase. to warm up to ambient temperature and was stirred for 1h. The obtained product was hydrolyzed with saturated aqueous NH4Cl (15 mL). The organic layer was extracted, dried over MgSO4, filtered, and the solvent was removed in vacuum to yield the product as a yellow oil, which was titrated with hexane to afford the pure product as a yellow solid (1.54 g, 75%)).

APENDIX A2

A2.3.3 Synthesis of L3

The polymer L3 was synthesized following the procedure described for the L1 polymer, namely the homopolymer synthesis. Two schlenk tubes were charged in the glovebox with 10.4 mg and 1.8 g of catalyst and L-LA respectively, and 5 and 15 mL of solvent respectively. The oil bath was pre-heated at 90°C. A white powder was obtained after 45 min

A2.3.4 Synthesis of L4

The polymer L4 was synthesized following the procedure described for the L1 polymer, namely the homopolymer synthesis. Two schlenk tubes were charged in the glovebox with 10.4 mg and 1.8 g of catalyst and L-LA and 5 and 15 mL of solvent, respectively.

A white powder was obtained after 90 min (1.20g, 66 %).

A2.3.5 Synthesis of L5

The polymer L5 was synthesized following the procedure described for the L1 polymer, namely the homopolymer synthesis. Two schlenk tubes were charged in the glovebox with 10.4 mg and 1.8 g of catalyst and L-LA and 5 and 15 mL of solvent, respectively.

A white powder was obtained after 120 min (1.55g, 86 %).

RESULTATS ET CONCLUSIONS

Dans cette thèse, la synthèse de copolymères stéréo-diblocs de haut poids moléculaire et de haute tacticité de PLA a été faites avec succès (Figure 2).

BIBLIOGRAPHIE

Figure 1 .

 1 Figure 1. Map summarising the collaborators of the project.

3. 3 . 3

 33 Spherulite growth rate of stereo-block copolymers (PLA-SBC) and PLLA and PDLA enantiomers (PLA-HP) ................................................................................................................................................................. 3.3.4 Isothermal crystallization ...................................................................................................................... 3.3.4.1 Avrami model .......................................................................................................................... 3.3.5 Non-isothermal crystallization .............................................................................................................. 3.3.5.1 Avrami modified model ......................................................................................................... 3.3.5.2 Ozawa model ........................................................................................................................... 3.3.5.3. Avrami-Ozawa model ........................................................................................................... 3.3.6 Crystallization of SBC with nucleating agents ..................................................................................... 3.3.6.1 Nucleating agent efficacy survey .......................................................................................... 3.3.6.1.1 BHET ...................................................................................................................................... 3.3.6.1.2 OXA ........................................................................................................................................ 3.3.6.2 Morphology of the crystallized NA/SC-PLA derivatives .................................................. ................................................................................................................................................ 3.5.2 Methods ................................................................................................................................................. 3.5.2.1 Polymerizations ....................................................................................................................... 3.5.2.2 Films preparation .................................................................................................................... 3.5.3 Instrumentations and measurements .................................................................................................... 3.5.3.1 Thermogravimetric analysis (TGA) ...................................................................................... 3.5.3.2 Gel permeation chromatography (GPC) .............................................................................. 3.5.3.3 Polarized Optical Microscopy (POM) .................................................................................. 3.5.3.4 Differential Scanning Calorimetry (DSC) ............................................................................ 3.5.3.5 Modulated-Differential Scanning Calorimetry (M-DSC) .................................................. 3.5.3.6 Wide-Angel X-Ray Scattering (WAXS) and Small-Angle X-Ray Scattering (SAXS) ...... 3.5.3.7 Scanning X-ray micro-diffraction (micro-beam) ................................................................. 3.5.3.8 Nanocalorimeter ......................................................................................................................

1. 6 . 1 . 3

 613 Perspectives ................................................................................................................................ 1.6.1.3.1 Pharmaceutical design ...................................................................................................... 1.6.1.3.2 Chiral drugs ........................................................................................................................ 1.6.1.3.3 Antibacterial applications ................................................................................................. 1.6.1.3.4 Polymer therapeutics......................................................................................................... 1.6.2 Biomedical applications ........................................................................................................................... 1.6.2.1 Stents ...................................................................................................................................1.6.2.2 Tissue engineering .................................................................................................................... 1.6.2.2.1 Bone regeneration .............................................................................................................. 1.6.2.2.2 Tendon regeneration ......................................................................................................... 1.6.2.2.3 Nerve regeneration ............................................................................................................ 1.6.2.2.4 Cartilage regeneration ....................................................................................................... 1.6.2.3 Perspectives ................................................................................................................................ 1.6.2.3.1 Piezoelectric activity .......................................................................................................... 1.6.2.3.2 Shape memory polymers ..................................................................................................

Figure 1 . 1 .

 11 Figure 1.1. PLA microstructures.

Figure 1 . 2 .

 12 Figure 1.2. Overview of the morphology attained by the different PLA derivatives depending on the processing conditions. Dimer refers to starting lactide dimer for ROP polymerizations.

  PLA was first synthesized by polycondensation by Théophile-Jules Pelouze in 1845. In 1932, Wallace Hume Carothers developed a novel synthetic method based on the ring-opening polymerization (ROP) of the cycle lactide monomer to synthesize PLLA. ROP was later patented by Du Pont in 1954 to synthesize vinyl fluoride (U.S. Patent No. 2674632, 1954).

  2) and the catalyst stereocontrol. The control exerted by the catalyst over the nanostructure of the PLA, and particularly to synthesize PLA stereo-block copolymers of HMw, is essential to tailor the properties of the final product, and the synthesis of novel catalysts, particularly metal-based catalysts, for polymerization by the coordination-insertion mechanism has attracted much attention since the pioneering work of Kleine et al. in the 1950s

Figure 1 . 3 .

 13 Figure 1.3. ROP of lactide by Sn(Oct)2 by coordination-insertion mechanism. Reproduced from K. Masutani, Y. Kimura, PLA Synthesis. From the Monomer to the Polymer, in: A. Jiménez, M. Peltzer, R. Ruseckaite (Eds.), Poly(Lactic Acid) Sci. Technol. Process. Prop. Addit. Appl., The Royal Society of Chemistry, Cambridge, 2015: pp. 3-36.,with permision from The Royal Society of Chemistry.

Figure 1 . 4 .

 14 Figure 1.4. Manufacturing process of PLGA microparticles by emulsion methods and the parameters affecting the final properties of the formulation. Reprinted from K. Park, A. Otte, F. Sharifi, J. Garner, S. Skidmore, H. Park, Y.K. Jhon, B. Qin, Y. Wang, Formulation composition, manufacturing process, and characterization of poly(lactide-co-glycolide) microparticles, J. Control. Release. 329, Copyright 2021, with permision from Elsevier.

  PLA hydrolysis follows the ester moiety cleavage from the polymeric backbone chain, decreasing the Mw: -[C(RR′)COO-] + H2O → -C(RR′)COOH + -C(RR′)COH Moreover, PLA hydrolysis is an autocatalytic process, as the chain scission increases the concentration of carboxylic acid end groups in the degradation medium, which possesses a catalytic action on the process (Equation 1.1) [123]. d[COOH]/dt = k[COOH][H2O][E] (1.1) where [COOH], [H2O], and [E] are the concentrations of the carboxylic groups, water, and ester groups in the medium, respectively. The hydrolytic degradation of PLA typically occurs in a stepwise fashion. Firstly, water is diffused into the material that triggers the later hydrolysis of the chains in the amorphous region; subsequently, the diminution of Mw proceeds as a result of the hydrolytic cleavage of ester bonds and the formation of watersoluble compounds; finally, the hydrolysis of the crystalline phase occurs [124].

Figure 1 . 5 ,

 15 the slope is ~1 in the plot of log kh against pH) and specific proton catalysis occurs at pH values below pH 4 (Figure 1.5, slope ~ -1) whilst kh is minimal at pH around 4.5 [130,131]. In addition, the dependence of the hydrolysis mechanism on pH was proved and the degradation under acidic conditions occurs through the split of the last lactic acid monomer after the nucleophilic attack by water [130-132]. However, the mechanism of PLA degradation in neutral or alkaline mediums is controversial, and the decomposition route is believed to follow either an intramolecular transesterification with the formation of lactide from the two end units of the oligomer [130] or a random ester cleavage degradation mechanism [131].

Figure 1 . 5 .

 15 Figure 1.5. Dependence of hydrolytic degradation constant (kobs) on pH. Reprinted from S.J. De Jong, E.R. Arias, D.T.S. Rijkers, C.F. Van Nostrum, J.J. Kettenes-Van Den Bosch, W.E. Hennink, New insights into the hydrolytic degradation of poly(lactic acid): Participation of the alcohol terminus, Polymer (Guildf). 42 (2001) 2795-2802, copyright 2001, with permission from Elsevier.

  of biomedical applications attained by PLA materials processed by 3D technology such as the oral formulation of solid dosage forms [169] and subcutaneous implants for prolonged drug delivery [170] by fused deposition modelling envisage their use increases as the 3D techniques advance and the utilisation of 3DP is expanded in the society.

  activities and pharmacokinetic profiles (e.x. R-Thalidomide is a sedative, S-Thalidomide leads to birth deformity). However, the vast majority of drugs used clinically are still employed as racemates likely due to the difficulty in separating racemic compounds into enantiomerically pure isomers. The chirality of PLA enables its use for enantioselective drug delivery [171] as proved by the generation of chiral particles of PLLA and PDLA loaded with racemic naproxen which preferably releases S-naproxen isomer in ethanol and chloroform for both PLLA and PDLA particles. The enantioselective release of S-naproxen was found to result from a different distribution of R-and S-naproxen within the particle caused by enantioselective interaction [172] (Figure 1.7). Furthermore, the drug release of the designed drug nanocarrier

Figure 1 . 6 .

 16 Figure 1.6. Machine learning approach for predicting the 3D printability of medicines (M3DISEEN). Reprinted from M. Elbadawi, B. Muñiz Castro, F.K.H. Gavins, J.J. Ong, S. Gaisford, G. Pérez, A.W. Basit, P. Cabalar, A. Goyanes, M3DISEEN: A novel machine learning approach for predicting the 3D printability of medicines, Int. J. Pharm. Copyright 2020, with permision from Elsevier.

Figure 1 . 7 .

 17 Figure 1.7. Polylactide-based chiral particles with enantio-differentiating release ability. Reprinted from J. Liang, B. Yang, J. Deng, Polylactide-based chiral particles with enantiodifferentiating release ability, Chem. Eng. J. 344 (2018) 262-269., copyright 2018, with permision from Elsevier

Figure 1 . 8 .

 18 Figure 1.8. Schematic diagram of the release of Ag+ ions from AgNPs comprised in PLA electrospun membranes exhibiting antibacterial properties. Reprinted with permission from C. Liu, J. Shen, K.W.K. Yeung, S.C. Tjong, Development and Antibacterial Performance of Novel Polylactic Acid-Graphene Oxide-Silver Nanoparticle Hybrid Nanocomposite Mats Prepared by Electrospinning, ACS Biomater. Sci. Eng. 3 (2017) 471-486. Copyright 2017 American Chemical Society.

  years ago and several generations of PLA-based materials were developed since then, especially for the treatment of maxillofacial fractures [217]. The third generation materials composed of uncalcined hydroxyapatite and PLLA, have reached the market (FIXSORB MX or OSTEOTRANS MX) and several clinical studies have been conducted exhibiting promising results [217]. However, the titanium plates remain in daily clinical use due to the persistence of PLA properties drawbacks, in particular, the slow degradation rate which leads to a foreign body reaction, and thus inflammation, after two years of implantation. Besides, the increase in the thickness of PLLA plates to match the strength of titanium plates increases the risk of exposure and patient discomfort [229], which might be improved by the fabrication of PLA plates in the SC phase to increase the strength avoiding the thickness increase. In addition, a fourth PLA-based material generation was recently developed which includes a small amount of PGA to overcome the slow degradation issue and its associated detrimental body inflammation [217,230,231]. PGA exhibits the highest initial mechanical properties, however, lessening strength very rapidly in the body due to its hydrolytic instability resulting in incomplete bone healing and high local acid concentration, causing inflammation [232], which turns the copolymer poly (lactide-co-glycolide) (PLGA) into the best option for bone regeneration. PLGA mechanical properties and degradation lifetime can be modified based on the PLA/PGA ratio (the lower the ratio PLA/PGA, the faster PLGA is expected to degrade). Moreover, silk [214], hydroxyapatite [233] or collagen [50] can be included in the scaffold to increase hydrophilicity, bone regeneration ability and tensile strength of the composite, respectively Recently, the PLA applicability in bonne tissue repair has advanced to tackle their characteristic limitations by the enhancement of the mechanical properties, degradation rate and the optimization of PLLA osteogenic activity [229]. Moreover, optimization of 3D printing parameters of PLA/Hydroxyapatite composites for bone plates has recently been developed, correlating the processing conditions with the final mechanical properties of the composite [234].

1 . 6 . 2 . 2 . 4

 16224 [235]. Synthetic conduits of PLA/PLDLA combined with longitudinal PLA filament scaffolds showed to enhance axon myelinization in a rat sciatic nerve lesion model when compared to a silicone conduit [215]. In addition, the neural precursor cells which are an interesting source of cells for neural tissue regeneration, generate a high amount of detrimental reactive oxygen species during metabolism and the lactic acid degradation product released from PLA scaffolds has been shown to reduce the intracellular redox state that increases the proliferative capacity of the neural cell population [236] Cartilage regeneration Similarly to tendons, cartilage tissue exhibits little ability to self-repair due to the limited vascularity. The polymeric approaches were envisaged due to their versatility and the lack of current satisfactory solutions for cartilage tissue regeneration by other means [237,238]. Electrospun scaffolds based on PLA and gelatin loaded with the antioxidant resveratrol showed high specific surface area, slow drug release and thinner diameter to promote the repair of cartilage injury [239]. In addition, thermogels containing PEG-sc-PLA-Chol were evaluated as scaffolds for cartilage tissue engineering in which the cholesterol block enhanced the mechanical properties, enlarged the pore size and improved chondrocyte adhesion. The degradation cycle of the PLA scaffold was according to the regeneration of cartilage defects and in vivo results showed a good differentiation of the loaded chondrocytes in cartilage-like tissue [225]. The crystallization in the α-or SC-phase of the PLA block depending on the initial materials was evidenced by different techniques, however, only the stereocomplexed materials were analysed in vivo [225].

Figure 1 . 9 .

 19 Figure 1.9. Piezoelectric activity of PLA. (b) Schematic of the bending mode of high aspect ratio structures with aspect ratio ϕ, which results in reduced effective stiffness by a factor of ϕ -2 with respect to the bulk material. (c) Proposed mechanism by which a cell can electromechanically stimulate itself by interacting with the piezoelectric nanostructures. (d) Simulation of a PLLA nanotube with axial polymer chain orientation, showing the potential developed in response to bending. Inset shows the opposing potentials developed across the tube diameter and the orientation of the corresponding electric field. (e) An example of the PLLA nanotube arrays produced via melt-press template wetting (i) in cross-section and (ii) in plan view. Adapted with permission from M. Smith, T. Chalklen, C. Lindackers, Y. Calahorra, C. Howe, A. Tamboli, D. V. Bax, D.J. Barrett, R.E. Cameron, S.M. Best, S. Kar-Narayan, Poly-l -Lactic Acid Nanotubes as Soft Piezoelectric Interfaces for Biology: Controlling Cell Attachment via Polymer Crystallinity, ACS Appl. Bio Mater. 3 (2020) 2140-2149. Copyright 2020 American Chemical Society.

Figure 1 . 10 .

 110 Figure 1.10. 4D scaffold of c-PLA/Fe3O4 ink and its potential biomedical application. (a) optical image of the multilayer scaffold; (b) schematic diagram of the detailed structure of the scaffold; (c) top view of the optical image of the printed scaffold; (d) deformation shape of the printed scaffold; (e) schematic of the restrictive shape recovery process; (f) demonstration of the restrictive shape recovery process triggered by a 30 kHz alternating magnetic field; (g) recovery shape under restrictive conditions; (h) potential application of the 4D scaffold as an intravascular stent. Reprinted with permission from H. Wei, Q. Zhang, Y. Yao, L. Liu, Y. Liu, J. Leng, Direct-write fabrication of 4D active shape-changing structures based on a shape memory polymer and its nanocomposite, ACS Appl. Mater. Interfaces. 9 (2017) 876-883. Copyright 2017, American Chemical Society.

  stents. Similarly, porous tissue scaffolds attaining different internal structures based on PLA with embedded magnetic nanoparticles have recently been developed by 3D printing showing great cell attachment and proliferation ability [248].

Scheme 2 . 1 .

 21 Scheme 2.1. Ring-Opening polymerization of L-and D-lactide following the one-pot sequential addition method to synthesize high molecular weight stereo-diblock copolymers of polylactide. (1):

(Figure 2 . 1 )

 21 Previously, an asymmetric dimeric conformation of the Initiator [Zn(Et)(κ 3 -bpzteH)] featuring dinuclear complexes was found by X-ray diffraction analysis (XRD) in the solid state, whilst an equilibrium with the monomeric counterpart was not discarded in solution[20]. Furthermore, the existence of different enantiomers of Initiator was assessed by the addition of a chiral shift reagent (R)-(-)-(9-anthrul)-2,2,2-trifluoroethanol for 1 H NMR

Figure 2 . 1 .

 21 Figure 2.1. Chemical structure of the Initiator [Zn(Et)(κ3-bpzteH)].

  polymerization mechanism was assessed by following the ring-opening of the LA by the Initiator in a [LA]/[Zn] = 1 mixture, monitoring the key functional groups involved in the reaction by VT-1 H NMR (Figures 2.3 and A2.5). The coordination of the Zn atom to the carboxylic oxygen of LA and the subsequent insertion of the alkyl (Et) group formed an intermediate that was evidenced by the broadness of the ethyl signals (δ 1.9 and 0.95; blue and orange arrows) immediately after the addition of LA to the reaction mixture (δ 1.4 and 4.0-4; red box). Thereafter, the acyl-oxygen bond of LA is disrupted and the generated linear chain of LA turns into the alkoxide part of the catalyst (δ 1.2; green box), confirming the coordination-insertion mechanism.

Figure 2 . 2 .

 22 Figure 2.2. Variable temperature 1 H NMR (500 MHz, Tol-d8, δ) spectra of the Initiator. 1 H NMR (CDCl3, 297 K): 1.88-1.91 (t, 3H, H9); 1.79 (s, 3H, H1); 0.95-0.96 (q, 2H, H8).

Figure 2 . 3 .

 23 Figure 2.3. VT-1 H NMR analysis of a [LA]/[Zn]:1 mixture.

Figure 2 . 4 .

 24 Figure 2.4. Plot of number average molecular weight (Mn) and molecular weight distribution values (PDI=Mw/Mn) obtained from gel permeation chromatography (GPC) as a function of monomer conversion (%). [LA]0/[Zn]0 = 500, toluene, 90 °C.

Figure 2 . 5 .

 25 Figure 2.5. 1 H NMR spectra (500 MHz, 298 K, CDCl3,) of the homodecoupled CH resonance of poly(rac-lactide). r and m are the trans-and cis-configuration of the relative orientation of two sequential alkyls of the diade. [rmr] represents the probability of finding a sequence with the following stereochemistry -LDDL-or -DLLD-; [rmm]= -LDDD-or -DLLL-; [mmr]= -LLLD-or -DDDL-; [mmm]= -LLLL-or -DDDD-; [mrm]= -LLDD-or -DDLL-.

[

  rmr]= -LDDL-or -DLLD-[rmm]= -LDDD-or -DLLL-[mmr]= -LLLD-or -DDDL-[mmm]= -LLLL-or -DDDD-[mrm]= -LLDD-or -DDLL-Pm= 0.52

  [LA] 0 /[Zn] 0 =400 polymerization at three different temperatures: 90, 75 and 60 °C. The order of reaction was calculated by plotting the ln([LA]0/[LA]t) versus reaction time where [LA]0 is the initial LA monomer concentration and [LA]t is the LA concentration at a given time t (Figure 2.6A). The polymerization rate increased with increasing temperature, typical of Arrhenius systems (Table

Figure 2 . 6 .

 26 Figure 2.6. a) First-order kinetic plots of L-LA polymerization in toluene at different temperatures. [LA] 0 /[Zn] 0 =400; [LA] 0 = 400 mM; [Zn] 0 = 1mM. b) Plot of ln k app versus T -1 for Initiator in toluene with [LA] 0 /[Zn] 0 =400; [LA] 0 = 400 mM; [Zn] 0 = 1mM.

Furthermore, the

  specific optical rotation (α) measured close to 0° for all samples, evidenced an L-LA/D-LA ratio of nearly one. The stereocontrol for the stereo-diblock copolymer formation was assessed by determining the tacticity of the copolymers obtained from deconvoluted homonuclear decoupling 1 H NMR spectra using Bernouillan statistics. The deconvoluted methine spectral regions of the homonuclear decoupled 1 H NMR spectrum of L375:D375 stereo-diblock co-polymer (Figure 2.7 as a representative example) were assigned

Figure 2 . 7 .

 27 Figure 2.7. 1 H NMR spectra (500 MHz, 298 K, CDCl3) of the homodecoupled CH resonance of stereodiblock copolymer 100L-b-100D (entry 2, Table3).

(

  [LA]0/[Zn]0 = 400; [LA]0 = 400 mM; [Zn]0 = 1 mM) at 90, 75 and 60 _C, PLLA and PDLA500 ([LA]0/[Zn]0 = 500; [LA]0 = 625 mM; [Zn]0 = 1.25 mM) at 90 _C, as well as L200:D200 ([LA]0/[Zn]0 = 400; [LA]0 = 400 mM; [Zn]0 = 1 mM) at 90 _C, to establish the reaction constant rate of each homopolymer enantiomer (Table

  Thermogravimetric analysis (TGA) .................................................................................................. Gel permeation chromatography (GPC) .......................................................................................... Polarized Optical Microscopy (POM) .............................................................................................. Differential Scanning Calorimetry (DSC) ........................................................................................ Modulated-Differential Scanning Calorimetry (M-DSC) .............................................................. Wide-Angel X-Ray Scattering (WAXS) and Small-Angle X-Ray Scattering (SAXS) .................. Scanning X-ray micro-diffraction (micro-beam) ............................................................................. Nanocalorimeter ..................................................................................................................................

  The onset of the stereocomplex-PLA (SC-PLA) degradation is at ca. 220-260°C[27] and thus, most of the calorimetry analyses reported in the literature erase the thermal history of SC-PLA by heating up to 230°C-260°C[30,34,41,45]. Likewise, the thermal degradation of the PLA-SBC derivatives was analysed by following the loss mass during a 5°C/min heating ramp from room temperature to 650°C by TGA (Figure3.1). The PLA-SBC of low molecular weight (LMw) (L115:D115; L220:D220; L250:D250; L275:D275) started to degrade at ca. 220°C whilst the onset HMw PLA-SBC (L300:D300; L375:D375; L500:D500; L520:D520) degradation

Figure 3 . 1 .Figure 3 . 2 )

 3132 Figure 3.1. TGA thermograms of the PLA-SBC (middle) and extension of the start (left) and end (right) of the degradation curve.

Figure 3 . 2 .

 32 Figure 3.2. Molecular weight and crystallization temperature of the L375:D375 as a function of the holding time at 230°C.

  c ,crystallization enthalpy; Xc-c, crystallinity obtained from crystallization; T g , glass transition ; T cc, cold-crystallization temperature; ΔH c , cold-crystallization enthalpy; T m on , onset melting temperature; T m p , peak melting temperature; ΔH m , melting enthalpy; 𝑋 𝑐-𝑚 = ∆𝐻 𝑚 ∆𝐻 𝑚 0 × 100 avec ∆𝐻 𝑚 0 = (142 J/g, crystallinity 1 ); Pm, probability of finding meso-dyads. SBC derivatives increased as the Mw increased (Figure 3.3A) and further rheological analysis are needed to determine the entanglement degree before crystallization and its relationship with the PLA-SBC molecular weight since generally, the Tc of polymers decrease as Mw increase. However, the Tc decreased for the highest Mw PLA-SBC derivative (L520:D520) that might be ascribed to the decrease of chain mobility due to higher entanglement [17]. Similarly, the Tc of the low Mw PLA-SBC derivative (L115:D115) slightly increased, due to an increase of chain mobility. However, a decrease of the crystallization rate as the Mw decreased was suggested by the cold crystallization observed upon the subsequent heating of the lower Mw PLA-SBC derivatives (L115:D115; L275:D275) (Figure 3.3B, red square).

Figure 3 . 3 .

 33 Figure 3.3. (A) Cooling and (B) heating ramps at 10°C/min of the PLA-SBC under analysis monitored by DSC.

Figure 3 . 4 .Table 3 . 4 .

 3434 Figure 3.4. Wide-Angle X-Ray Scattering (WAXS) patterns of the PLA-SBC after cooling from the melt at 10°C/min

Figure 3 . 5 .

 35 Figure 3.5. Optical microscopy images of PDLA500, PLLA500 and L520:D520 obtained upon cooling form the melt at 5°C/min. The size reference found at the right bottom of the images represents 50µm.

Figure 3 . 6 .

 36 Figure 3.6. Optical images obtained by POM after 5 and 15 minutes at the isothermal crystallization temperature indicated above for HP-PLA (PLLA500 and D500) and SBC (L520:D520) respectively.

3 . 3 )

 33 by using the following parameter values: U* = 1500 cal mol -1 and T∞ = Tg and R= 1,987 cal K -1 mol -1[61].

Figure 3 . 7 .

 37 Figure 3.7. (A) Spherulite growth rate as a function of Ta and (B) Logarithmic plot of the Lauritzen-Hoffman equation for PDLA500, PLLA500 and L520:D520.

. 8 )Table 3 . 6 .

 836 consisted of a first heating ramp at 10 °C/min from 25 °C to 230 °C and held for 2 min at the melt to erase thermal history. Samples were then cooled down to the isothermal-crystallization temperature (Ti-c) at 50°C/min, hold 60 minutes and cooled down to 25°C at 2°C/min. Subsequently, the samples were heated at 10°C/min to 230°C and cooled down again to 25°C at 10°C/min to analysed the melting behaviour of the isothermally crystallized PLA-SBC and verify the Mw/tacticity stability, respectively. The crystallization temperatures for isothermal crystallization (Ti-c) were selected based on the crystallization temperature (Tc) found in the preliminary experiments under non-Abbreviations of the PLA-SBC used in the following sections. HMw: high molecular weight; MMw: medium molecular weight; HPm: High tacticity; LPm: Low tacticity; Mw; weight-averaged molecular weight; Pm: probability of finding meso-diads.

Figure 3 . 8 .

 38 Figure 3.8. Isothermal protocol conducted by DSC.

. 1 -

 1 A3.3), indicating that the crystallization was incomplete during the isothermal step. The Tg of the high tacticity stereo-diblock copolymers (HMw-HPm and MMw-HPm) was c.a. 50°C independently of the

Figure 3 . 9 .

 39 Figure 3.9. DSC plots of the isothermal steps at different annealing temperatures for L500:D500 (A), L375:D375 (B) and L330:D330 (C).

Figure 3 . 10 .

 310 Figure 3.10. DSC plots of the heating ramp after the isothermal step and subsequent cooling for (A) L500:D500, (B) L375:D375 and (C) L330:D330

Figure 3 . 11 .

 311 Figure 3.11. Second isothermal protocol conducted by DSC

Figure 3 . 13 .

 313 Figure 3.13. WAXS patterns of (A) HMw-HPm or L500:D500, (B) MMW-HPm or L375:D375 and (C) MMw-LPm or L330:D330 after 60 minutes isothermally crystallized at the temepratures indicated.

Figure 3 .

 3 Figure 3.12. (A) DSC plot of the heating ramp after isothermal protocol 1 and 2 and (B) M-DSC plot of the heating ramp of L330:D330 at Ta=160°C

Figure 3 . 14 .

 314 Figure 3.14. SAXS profiles of (A) HMw; HPm or L500:D500, (B) MMW; HPm or L375:D375 and (C) MMw; LPm or L330:D330 after 60minutes isothermally crystallized at the temperatures indicated.

Figure 3 . 14 )

 314 crystallization mechanism at isothermal conditions was further analysed by modelling the kinetics of the PLA-SBC crystallization by the Avrami model. The Avrami equation (Eq. 3.3) is expressed as: 1 -Vc (t) = exp (-kt n ) (3.3) where: Vc is the relative volumetric transformed fraction, n is the Avrami index and k the overall crystallization rate constant, that involves both, nucleation and growth. The evolution of Vc obtained from DSC with time is slower when increasing Ti-c for the stereo-diblock PLA copolymers under study (Figure A3.16). The linearization of the Avrami equation 3.4 permits to both determine the Avrami index and the crystallization rate by: of Log [-Ln [1 -Vc(t)]] vs Log (t) and the subsequent linear regression of the experimental data is conducted to extract the Avrami index n and the overall crystallization rate k from the slope and the intercept, respectively (Table

Figure 3 . 15 .Table 3 . 7 .

 31537 Figure 3.15. (A) Crystallinity values obtained from the deconvolution of the WAXS patterns and (B) long period obtained from the Lorentz-corrected SAXS profiles after the isothermal step (full symbols) and subsequent cooling to room temperature (empty symbols).

Figure 3 . 16 .

 316 Figure 3.16. Avrami plots obtained from DSC analysis for PLA-SBC isothermally crystallized at different Ti-c (A) HMw-HPm (L500:D500), (B) MMw-HPm (L375:D375) and (C) MMw-LPm (L330:D330).

  ) followed a sigmoidal function at Ti-c 150°C and 155°C, indicating that was fully crystallized during the isothermal step following an Arrhenius behaviour, whilst at 160°C followed a much more linear function, suggesting an incomplete crystallization during the isothermal step, similar to the MMw-HPm at 195, 190 and 185°C. The linear Avrami plots of the obtained from WAXS were almost parallel to each other (Figure 3.18), similar to Avrami plots obtained from DSC thermograms. The obtained n values of 2.6 for MMw-LPm are similar to n values calculated from DSC, however the t1/2 were much higher (Table 3.7). Conversely, MMw-HPm presented n values slightly lower than the obtained from DSC and slightly lower t1/2 (Table 3.7).

Figure 3 . 17 .

 317 Figure 3.17. Optical images for (A) HMw-HPm at Ti-c = 175°C, (B) MMw-HPm at Ti-c = 170°C and (C) MMw-LPm at Ti-c = 152°C

3. 19 )Figure 3 . 18 .

 19318 Figure 3.18. Avrami plots obtained from WAXS analysis for PLA-SBC isothermally crystallized at different Tc-i (A) MMw-HPm (L375:D375) and (B) MMw-LPm (L330:D330).

Figure 3 . 19 .

 319 Figure 3.19. Non-isothermal protocol accomplished by DSC.

. 4 -

 4 A3.6) mainly due to an incomplete crystallization upon cooling at high cooling rates, as evidenced by the coldcrystallization exhibited upon subsequent heating (Figure A3.19-A3.20), suggesting a lower

Figure 3 . 20

 320 Figure 3.20 Crystallization temperature (A) and crystallinity degree obtained from the crystallization enthalpy (B) as a function of cooling rate for the SBC-PLA crystallized under non-isothermal conditions from the melt.

  .21 and A3.18-A3.20). Firstly, the Tg was found to be dependent on the tacticity, shifting from 54-56°C for the HPm stereo-diblock copolymers to 50°C for the low tacticity stereo-diblock copolymers, probably due to the lower crystallinity obtained upon cooling. Moreover, the heating ramps of both MMw PLA-SBC derivatives, the HPm and the LPm, suggested a lower crystallization rate compared to the HMw stereo-diblock copolymer due to the presence of coldcrystallization after cooling at 50 as well as 30 and 50°C/min, respectively. The melting temperatures of the HPm stereo-diblock copolymers (Figure3.21) was c.a. 210°C whilst the low tacticity PLA-SBC (MMw-LPm) exhibited a melting temperature 10°C lower, similarly to the Tm after isothermal conditions. In addition, the melting endotherms of the high tacticity PLA-SBC exhibited a well mono distributed form independently of the cooling rate, whilst the low tacticity PLA-SBC showed a wide distributed melting endotherm for the crystallized derivative at 50°C/min, that becomes more defined as the cooling rate decreased, probably due

Figure 3 . 21 .

 321 Figure 3.21. Melting temperature as a function of cooling rate for the SBC-PLA crystallized under non-isothermal conditions from the melt.

Figure 3 . 22 .Figure 3 . 23 .

 322323 Figure 3.22. WAXS patterns of (A) HMw-HPm or L500:D500, (B) MMw-HPm or L375:D375 and (C) MMW-LPm or L330:D330 at room temperature crystallized from the melt at different cooling rates

Figure 3 . 24 .

 324 Figure 3.24. Crystallinity (A) and Long period (B) of the PLA-SBC derivative under study obtained from WAXS and SAXS experiments, respectively.

Figure 3 . 25 . 3 . 8 .Figure 3 . 26 .

 32538326 Figure 3.25. Logarithmic plot (Eq. 3.9) of the Avrami modified equation for (A) HMw-HPm or L500:D500, (B) MMw-HPm or L375:D375 and (C) MMw-LPm or L330:D330

Figure 3 . 27 . 3

 3273 Figure 3.27. Logarithmic plot of the Ozawa model of the SBC-PLA non-isothermally crystallized from the melt.

Figure 3 . 28 .

 328 Figure 3.28. Logarithmic plot of the Avrami-Ozawa model of the SBC-PLA nonisothermally crystallized from the melt.

Table 3 . 8 .

 38 efficiency was assessed for N,N'-Bis(2-hydroxyethyl) terephthalamide [85] as well as 5,6,11,12tetraoxo-4,7,10,13-tetraazahexadecane-1,16-dioate and diethyl 5,6,13,14-tetraoxo-4,7,12,15etraazaoctadecane-1,18-dioate [23] with a series of PLA-SBC derivatives (Table 3.8) featuring different Mw as well as an PLA-SBC derivative with an asymmetric block to evaluate structural parameters on the kinetics mechanism. PLA-SBC derivatives and the corresponding code under study with the nucleating agent. Mw: Molecular weight; Pm: Probability of finding meso-dyads; Rac-PLA: racemic Polylactide; PLLA: Poly (L-lactide); PDLA: Poly (D-lactide); SHMw: Super-high molecular weight; HPm: high tacticity; MMw: Medium molecular weight; HMw: High molecular weight; A-UHMw: Asymmetric-Ultra high molecular weight.

Figure 3 . 29 .Figure 3 .

 3293 Figure 3.29. Chemical Structure of Compound N,N′-Bis(2-hydroxyethyl)terephthalamide (BHET)

Figure 3 . 31 .Table 3 . 8 .

 33138 Figure 3.31. Chemical structure of the OBOCs used in this study, abbreviated based on the number of methylene groups in-between the oxalamide units. OXA-2: n=2 ; OXA-4: n=4.

1

  and 2%) of two OBOCs, OXA-4 and OXA-2. The dissolution and crystallization of the NA in the PLA derivative matrix were clearly observed by POM. The PLA-SBC containing OXA-4 crystallized above 140°C following homogeneous nucleation (Figure 3.32A), whilst the OXA-4 crystals appeared later than PLA spherulites, yielding also heterogeneous nucleation (Figure 3.32 B, see arrow). In addition, OXA-2 crystals were formed above 200°C allowing the onset of the nucleation of PLA-SBC around 170°C (Figure 3.32 C-D).

Figure 3 . 32 .Figure 3 . 33 .

 332333 Figure 3.32. Polarized optical microscopy images of PLA film containing 2 wt % of OXA-4 (A,B) and OXA-2 (C,D).

Figure 3 . 34 .

 334 Figure 3.34. WAXS pattern of the L375:D375 film containing 2%wt of OXA-2 upon cooling (A) at 30°C/min and heating (C) at 10°C/min. (B) and (D) are zoom.

. 35 )

 35 . The generated morphology by POM confirmed the needle-like growth that tends to form a branched network due to the likely concentration gradient provoked by the diffusion through the PLA melt as well as the depression of the crystallization temperature as previously observed [23]. Likewise, the isothermal crystallization temperatures were then selected by screening the crystallization from the melt of the PLA-SBC derivatives under study. Therefore, the generated PLA-SBC thin films with and without OXA-2 were cooled from the melt at a constant rate of 20°C/min due to the technical limitations of the heating stage coupled to POM (highest cooling rate attainable) to visually determine the crystallization temperature (as shown for the MMw-HPm derivative, Figure 3.36 and for the rest of the PLA-SBC derivatives under study in Apendix A3 Figures A3.32-A3.33 and Table3.10). The addition of OXA-2 to the PLA derivatives under study has produced the inverse effect on the crystallization temperature of the PLA derivatives compared to the OXA-2. Particularly, a general slight increase of the Tc was observed for the PLA derivatives at the applied cooling rate whilst the crystallization of OXA-2 proceeded at lower temperatures. However, the depression of the Tc

Figure 3 . 35 .

 335 Figure 3.35. POM images of the thin film of Rac-PLA containing 2 wt% of OXA-2 nonisothermally crystallized from the melt at 20°C/min.

Figure 3 . 36 .Table 3 . 10 .

 336310 Figure 3.36. POM images of the thin films of MMw-HPm containing 0 wt% and 2 wt% of OXA-2 non-isothermally crystallized from the melt at 20°C/min.

Figure 3 . 37 .Figure 3 . 38 .Figure 3 . 39 .

 337338339 Figure 3.37. POM image of the MMw-HPm film after X-Ray microbeam.

Figure 3 . 40 .Figure 3 . 41 .

 340341 Figure 3.40. Azimuthal integration of (110) SC crystallographic reflection of MMw-HPm thin films with (A) 0xt% and (B)2wt% of OXA-2 crystallized isothermally at 170°C.

Figure 3 .

 3 Figure 3.42. X-Ray micro beam images of MMw-HPm thin films with (A) 0xt% and (B)2wt% of OXA-2 crystallized isothermally at 170°C, where yellow and black colours represent high and low intensities of the (110) SC plane.

Figure 3 . 43 .

 343 Figure 3.43. POM images obtained at room temperature of the HMw-HPm PLA-SBC derivative with 0 and 2wt% of OXA-2 crystallize from the melt in a nanocalorimeter.

(

  MMw-HPm-2%OXA and HMw-HPm-2%OXA) during the isothermal step, suggesting lower nucleation along the needles like crystals formed by the nucleating agent that enable the growth of spherulites in all directions. The homocrystal and stereocomplex phase distribution suggested by DSC was also probed by microWAXS in the thin films of both A-UHMw-HPm-0%OXA and A-UHMw-HPm-2%OXA. The A-UHMw-HPm-0%OXA presented a macroscopic phase separation upon crystallization from the melt, as revealed by the intensity distribution of both the (110)SC (Figure 3.45A)and (110)HC plane (Figure 3.45B). The macroscopic phase separation might be ascribed to the survival of the stereocomplex hydrogen bonds interactions in the melt that assist the phase separation upon high undercooling. However, the nucleating agent, OXA-2, prevented the macroscopic phase separation observed in the UHMw-HPm-0%OXA and both HC-and SC-phases were homogeneously distributed across the film (Figure 3.46). Moreover, the stereocomplexation occurred firstly upon reaching 20-30µm of diameter (Figure 3.46A) followed by the homocomplexation in the

Figure 3 . 44 .

 344 Figure 3.44. POM images of the thin films of A-UHMw-HPm containing (A) 0 wt% and (B) 2wt% of OXA-2 isothermally crystallized at 120°C and 150°C, respectively.

Figure 3 . 45 .

 345 Figure 3.45. X-Ray micro beam images of A-UHMw-HPm thin films with 0xt% OXA-2 crystallized isothermally at 120°C, where yellow and black colours represent high and low intensities of the (110) SC plane (A) and the (110)HC plane.

Figure 3 .

 3 Figure 3.46. X-Ray micro beam images of A-UHMw-HPm thin films with 2wt% OXA-2 crystallized isothermally at 150°C, where yellow and black colours represent high and low intensities of the (110) SC plane (A) and the (110)HC plane (B).

Figure 3 . 47 .

 347 Figure 3.47. Peak position analysis of the X-Ray micro beam images of A-UHMw-HPm thin films with 2wt% OXA-2 crystallized isothermally at 150°C. (A) (110) SC plane and (B) (110)HC plane.

  Small-Angle X-Ray Scattering (SAXS) SAXS/WAXS experiments were conducted both at DUBBLE (bm26) ESRF in Grenoble (France) and bl11 NCD-SWEET at ALBA, Cerdanyola del Vallès (Spain) with an X-ray wavelength of 12 Kev. The WAXS patterns were recorded with a Pilatus 300 kw (1472 × 195 pixels configuration) with a pixel size of 172 µm 2 × 172 µm 2 in DUBBLE, whereas a Rayonix LX255-HS characterized by a pixel size of 40 µm 2 × 40 µm 2 and active area of 85 mm 2 × 255 mm 2 (h × v) was used at NCD-SWEET. The SAXS patterns were collected with a Pilatus 1M with a detector size (981 × 1043) with a pixel size of 172 µm × 172 µm at a sample to detector distance of ca. 2.5 m. The calibration of the scattering angular was performed with either α-Al2O3 (alumina) for the WAXS scattering angles or AgBe for the SAXS scattering angle range.

4. 5 . 2 . 1

 521 Gelatin extraction .................................................................................................................... 4.5.2.2 Polymerizations ....................................................................................................................... 4.5.2.3 Nanoparticles formulations ................................................................................................... 4.5.2.4 Bionanocomposites generation ............................................................................................. 4.5.3 Instrumentation and measurements ...................................................................................................... 4.5.3.1 1 H NMR ................................................................................................................................... 4.5.3.2 GPC analysis ............................................................................................................................ 4.5.3.3 Optical rotation ........................................................................................................................ 4.5.3.4 DLS technique .......................................................................................................................... 4.5.3.5 SEM ........................................................................................................................................... 4.5.3.6 Cryo-SEM ................................................................................................................................. 4.5.3.7 DSC measurements ................................................................................................................. 4.5.3.8 Raman spectroscopy ............................................................................................................... 4.5.3.9 Small-angle X-ray scattering (SAXS) and wide-angle X-ray scattering (WAXS) ............ 4.5.3.10 Rheological experiments ...................................................................................................... 4.5.3.11 Release studies .......................................................................................................................



  collaboration with the Universities of Vigo and Castilla-La Mancha, as well as The European Synchrotron Radiation Facility (ESRF): Selection of three PLA-derivatives described in Chapter 2 exhibiting different chain arquitecture and characterization of its nanostructure. Formulation and characterization of PLA nanoparticles with two different antitumoral compounds based on the three different PLA derivatives selected.  Blending of PLA nanoparticles with fish gelatin at 65°C to generate the multifunctional bionanocomposite.  Characterization of the storage capability and the feasibility for local administration of the bionanocomposite.  Drug release analysis from both the PLA nanoparticles, gelatin hydrogel network and the bionanocomposite.

Scheme 4 . 1 .

 41 Scheme 4.1. Summary of the steps followed along the chapter to obtain the bionanocomposites.

Figure 4 . 1 .

 41 Figure 4.1. DSC thermogram (A) and WAXS patterns (B) of bulk polymer at room temperature.

Figure 4 .

 4 Figure 4.2A) was accomplished to identify the vibrational bands associated with phase variations (Figure 4.2B-E and TableA4.1). In particular, the vC-COO stretching vibration

Figure 4 . 2 .Figure 4 .

 424 Figure 4.2. Full Raman spectrum of the bulk polymers (A), C-COO stretching Raman region of the bulk polymers (B) and carbonyl region of the Rac-PLA (C), HC-PLA (D) and SC-PLA (E)

  with the release profiles. The nanostructure of the lyophilized NP-PLA with doxorubicin (NP-PLA/DOX) probed by WAXS evidenced the kinetically trapped state in the amorphous phase for the PLA derivatives (Figure 4.4A) during the NP formation, removing the targeted crystallinity of the designed semicrystalline PLA derivatives (NP-HC-PLA and NP-SC-PLA).The amorphous structure obtained from the formulation conditions is likely related to the subjected short incubation time of the polymer solution during the water quenching

Figure 4 . 3 .

 43 Figure 4.3. SEM image of NP-HC-PLA (A) and histogram of the statistical diameter analysis (B).

Figure 4 . 4 .

 44 Figure 4.4. WAXS (A) and SAXS (B) of the Doxorubicin-loaded and no-loaded PLA nanoparticles at room temperature.

  3, 22.4, 38.8 and 45, Figure 4.4A), corresponding to doxorubicin reflections [44], indicating that doxorubicin maintained its crystalline structure upon encapsulation in PLA nanoparticles. In addition, the NP-SC-PLA/DOX featured a crystalline structure (Figure 4.4A, zoom), which suggests the nucleating effect of doxorubicin. Likewise, the absence of the long period in the SAXS profile of the loaded nanoparticles confirms the quenching mechanism during the formation of the nanoparticles (Figure 4.4B). Thermal treatment for the formulated NP-PLA was applied by annealing them at 65 °C, above the observed Tg at ca. 40 °C (Figure 4.5), for 1 h to modify the kinetically trapped nanostructure afforded upon formulation whilst selecting a temperature that will maintain the stability of the gelatin matrix. The thermal parameters obtained from the thermograms are summarized in table 4.3. The NP-PLA Tg at ca. 40°C was followed by an endothermic peak at ca. 55 °C that might correspond to an enthalpy relaxation that occurred due to physical ageing during the storage period before the thermal analysis (Figure 4.5A). Moreover, both NP-HC-PLA and NP-SC-PLA exhibited a cold crystallization during the isothermal annealing, confirming the quenched effect by the NP formation and discarding any degradation effect upon formulation (Figure 4.5B). NP-Rac-PLA did not crystallize upon heating in agreement with the bulk polymer structure. Additionally, upon further heating, NP-HC-PLA and NP-SC-PLA exhibited an endothermic peak at 165 °C and 200 °C, respectively, which is consistent with the melting of the α-and SC-phases, respectively, corroborating the crystallization of the semicrystalline PLA derivatives formulations during the annealing step (Figure A4.6).

Figure 4 . 5 .

 45 Figure 4.5. DSC thermograms of the no-loaded nanoparticles during heating step (A) and isothermal step at 65°C (B).

Table 4 . 4 .

 44 Thermal parameters obtained from DSC analysis of theDoxorubicin-loaded and non-loaded NP-PLA. Furthermore, the doxorubicin-loaded nanoparticles revealed by DSC an increase of Tg upon heating (Table 4.4 and Figure 4.6A), which might be ascribed to the interaction of PLA with doxorubicin. Moreover, the isothermal treatment at 65 °C of the doxorubicin-loaded nanoparticles revealed the crystallization of both NP-HC-PLA/DOX and NP-SC-PLA/DOX (Figure 4.6B). However, a lower crystallization enthalpy was found with respect to the Sample Tg

Figure 4 . 6 .Figure 4 . 7 .

 4647 Figure 4.6. DSC thermograms of the Doxorubicin-loaded nanoparticles during heating step (A) and isothermal step at 65°C (B).

Figure 4 .

 4 A bionanocomposite hydrogel formed by NP-HC-PLA and fish gelatin (GE/NP-HC-PLA) was formed upon cooling to room temperature of the homogenized gelatin solution by sonication at 65 °C as the gelatin hydrogel network is disrupted, allowing the NP-PLA dispersion and the concomitant cold crystallization of the PLA nanoparticles in the α-phase, as confirmed by WAXS measurements at room temperature (Figure 4.8A, red line). The development of the α-crystalline phase reflections was observed over the WAXS profile of the gelatin, which features a broad amorphous halo with the superimposed signals of the triple helix with the characteristic domain distances of 11-12.6 Å and mainly the single α-helices with repeat distance of 2.9 Å (Figure A4.7A). Furthermore, the NP-HC-PLA embedded in the gelatin hydrogels were isothermally annealed at 65 °C for 30 min, and the nanostructure development was probed by simultaneous SAXS/WAXS experiments (Figure 4.8B-4.8D). The crystallinity of the homo phase ((110)/(200) and (203) (q (nm -1 ) = 11.9 and 13.6), respectively) further increased during the isothermal treatment (8B), whilst the α-helices of the gelatin diminishes (FigureA4.7A). Furthermore, the SAXS profile of the annealed bionanocomposite was exclusively sensible to the gelatin hydrogel nanostructure in the q range under study (FigureA4.7B) due to the large size of the NP-PLA.

  In addition, the effect of the nanoparticles embedded in the hydrogel thermal stability was assessed by DSC analyses. The sol-gel transition of the gelatin hydrogel network was shifted upon the incorporation of the NP-PLA from ca. 20 °C for the shark gelatin hydrogel to ca. 45-50 °C (FigureA4.9), which confirms that NP-PLA percolation throughout the hydrogel network reinforced the physical crosslinks, suggesting storage capability.

Figure 4 . 8 .

 48 Figure 4.8. (A) WAXS pattern of NP-HC-PLA (black line), the bionanocomposite of NP-HC-PLA at room temperature (red line), 65 °C (blue line) and after 30 minutes at 65 °C (pink line). (B) WAXS pattern evolution during the isothermal at 65 °C. (C) SAXS pattern of NP-HC-PLA (black line), the gelatin (green line), the bionanocomposite of NP-HC-PLA at room temperature (red line), 65 °C (blue line) and after 30 minutes at 65 °C (pink line). (D) SAXS pattern evolution during the isothermal at 65 °C.

Figure 4 . 9 . 4 . 3 . 5

 49435 Figure 4.9. Cryo-SEM images of GE with (A-C) and without (B-D) NP-SC-PLA at x2000 (A-B) and x5000 (C-D) magnification

Figure 4 . 10 .

 410 Figure 4.10. Strain and frequency sweep for GE (green color), GE/NP-SC-PLA/DOX (blue color), GE/NP-HC-PLA/DOX (red color) and GE/NP-Rac-PLA/DOX (black color). Storage (G′; filled symbols) and loss moduli (G″; hollow symbols) depicted versus strain (A) and angular frequency (B) at 20 °C.

Figure 4 . 11 .Figure 4 .

 4114 Figure 4.11. Syringe test results of gelatin hydrogel (A), GE/NP-Rac-PLA/DOX (B), GE/NP-HC-PLA/DOX (C) and GE/NP-SC-PLA/DOX (D). Viscosity (black color), shear rate (red color), temperature (blue color).

Figure 4 . 12 )Figure 4 . 12 .

 412412 Figure 4.12. Release profiles of DOX from nanoparticles in suspension (NP-HC, NP-SC and NPrac), gelatin (Gelatin) and nanoparticles embedded in gelatin (Gelatin-NP-HC, Gelatin-NP-SC and Gelatin-NP-Rac) in PBS pH 7.4 at 37 ° C. The drug releases were tested in three replicates. Error bars are 2σ.



  The generated bionanocomposites based on gelatin and comprising PLA nanoparticles with several nanostructures might have potential in the development of DDS for local release in soft tissue treatments requiring the administration of Solvents and reagents were acquired from Sigma-Aldrich. Blue Shark (BS, Prionace glauca) skin by-products were kindly provided by Propegal S.L. L-and D-lactide dimers were purchased from Rex scientific. The lactide dimers were purified three times by sublimation and then stored in a glovebox at 4°C. Toluene and tetrahydrofuran (THF) were pre-dried over sodium wire as well as distilled under nitrogen from sodium and subsequently stored over molecular sieves (3Å) in a glovebox.

A2. 1

 1 Characterization and activity of the catalyst A2.1.1 Synthesis of the 2,2-bis((3,5-dimethyl-pirazol-yl)-1-para-tolylethoxide) (bpzteH) The ligand (bpzteH) was synthesized via the following procedure: In a 250 mL Schlenk tube, Bis 3,5-(dimethyl-pirazol-yl)-methylene (bpm) (3 g, 14.7 mmol) was dissolved in dry THF (70 mL) and then cooled to -78°C. A 2.5 M solution of BunLi (6.47 mL, 16.17 mmol) in hexane was added to the bpm solution and stirred for 1h. The resulting mixture was added dropwise to a cooled (-10°C) solution of p-tolyl (1.86 g, 15.43 mmol). The reaction mixture was allowed

1 H

 1 NMR (500 MHz, CDCl3, 297 K): δ 7.25-7.23 (d, 2H, H3), 6.91-6.88 (d, 2H, H2), 6.28-6.26 (d, 1H, H6), 5.90-5.89 (d, 1H, H6), 5.59 (s, 1H, H4), 5.44 (d, 1H, H5), 2.14-1.55 (m, 12H, H7), 1.44 (s, 3H, H1).

Figure A2. 1 : 1 HFigure A2. 3 .

 113 Figure A2.1: 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of the ligand 2,2-bis((3,5-dimethylpirazol-yl)-1-para-tolylethoxide) (bpzteH) ): δ 7.25-7.23 (d, 2H, H3), 6.91-6.88 (d, 2H, H2), 6.28-6.26 (d, 1H, H6), 5.90-5.89 (d, 1H, H6), 5.59 (s, 1H, H4), 5.44 (d, 1H, H5), 2.14-1.55 (m, 12H, H7), 1.44 (s, 3H, H1).

Figure A2. 4 .Figure A2. 6 .A2. 3

 463 Figure A2.4. Energy profile for the dimeric and monomeric species. ΔG = 0 for the dimer in gas and toluene

Figure A2. 7 : 1 H

 71 Figure A2.7: 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of L1.

( 1 .

 1 44g, 80%).1 H NMR (500 MHz, CDCl3, 297 K): δ 5.21-5.16 (q, 1H, CH(PLA)), 5.07-5.02 (q, 1H, CH(L-LA)), 1.71-1.70 (d, 3H, CH3(L-LA)), 1.61-1.59 (d, 3H, CH3(PLA)). GPC (CHCl3): Mw=115118, Mn=62564, Ð=1.84.

Figure A2. 9 :

 9 Figure A2.9: 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of L3.

  1 H NMR (500 MHz, CDCl3, 297 K): δ 5.21-5.16 (q, 1H, CH(PLA)), 5.07-5.02 (q, 1H, CH(L-LA)), 1.71-1.69 (d, 3H, CH3(L-LA)), 1.61-1.59 (d, 3H, CH3(PLA)). GPC (CHCl3): Mw=75898, Mn=42880, Ð=1.77.

Figure A2. 10 :

 10 Figure A2.10: 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of L4.

  1 H NMR (500 MHz, CDCl3, 297 K): δ 5.12-5.06 (q, 1H, CH(PLA)), 4.98-4.93 (q, 1H, CH(L-LA)), 1.62-1.60 (d, 3H, CH3(L-LA)), 1.52-1.47 (d, 3H, CH3(PLA)). GPC (CHCl3): Mw=105673, Mn=48253, Ð=2.19.

Figure A2. 11 : 1 H

 111 Figure A2.11: 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of L5.

Figure A2. 12 :

 12 Figure A2.12: 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of D1.

Figure A2. 15 : 1 H

 151 Figure A2.15: 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of L115:D115.

Figure A2. 16 :

 16 Figure A2.16: 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of the homodecoupled CH resonance of L115:D115.

Figure A2. 17 : 20 Figure A2. 18 : 1 H

 1720181 Figure A2.17: 13C NMR spectrum (500 MHz, 298 K, CDCl3) of L115:D115.

Figure A2. 19 :

 19 Figure A2.19: 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of the homodecoupled CH resonance of L220:D220.

Figure A2. 22 :

 22 Figure A2.22: 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of L275:D275.

Figure A2. 23 : 1 H

 231 Figure A2.23: 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of the homodecoupled CH resonance of L275:D275.

Figure A2. 24 :

 24 Figure A2.24: 13 C NMR spectrum (500 MHz, 298 K, CDCl3) of L275:D275.

Figure A2. 25 : 1 H

 251 Figure A2.25: 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of D275:L275.

Figure A2. 26 :

 26 Figure A2.26: 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of the homodecoupled CH resonance of D275:L275.

Figure A2. 34 : 1 H

 341 Figure A2.34: 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of L520:D520.

Figure A2. 35 . 1 H

 351 Figure A2.35. 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of the homodecoupled CH resonance of L520:D520.

Figure A2. 37 . 1 H

 371 Figure A2.37. 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of L1200:D300.

Figure A2. 38 . 1 H

 381 Figure A2.38. 1 H NMR spectrum (500 MHz, 298 K, CDCl3) of the homodecoupled CH resonance of L1200:D300.

Figure A3. 3 .

 3 Figure A3.3. Spherulite radius as a function of time at different isothermal crystallization temperatures of (A) PDLA500, (B) PLLA500 and (C)L520:D520.

Figure A3. 4 .

 4 Figure A3.4. WAXS (A-B) and SAXS (C-D) patterns of the HMw-HPm isothermally crystallized from the melt at 200°C (A-C) and the subsequent cooling to room temperature (B-C).

Figure A3. 5 .

 5 Figure A3.5. WAXS (A-B) and SAXS (C-D) patterns of the HMw-HPm isothermally crystallized from the melt at 190°C (A-C) and the subsequent cooling to room temperature (B-C).

Figure A3. 6 .

 6 Figure A3.6. WAXS (A-B) and SAXS (C-D) patterns of the MMw-HPm isothermally crystallized from the melt at 200°C (A-C) and the subsequent cooling to room temperature (B-C).

Figure A3. 7 .

 7 Figure A3.7. WAXS (A-B) and SAXS (C-D) patterns of the MMw-HPm isothermally crystallized from the melt at 198°C (A-C) and the subsequent cooling to room temperature (B-C).

Figure A3. 8 .

 8 Figure A3.8. WAXS (A-B) and SAXS (C-D) patterns of the MMw-HPm isothermally crystallized from the melt at 195°C (A-C) and the subsequent cooling to room temperature (B-C).

Figure A3. 9 .

 9 Figure A3.9. WAXS (A-B) and SAXS (C-D) patterns of the MMw-HPm isothermally crystallized from the melt at 190°C (A-C) and the subsequent cooling to room temperature (B-C).

Figure A3. 10 .

 10 Figure A3.10. WAXS (A-B) and SAXS (C-D) patterns of the MMw-LPm isothermally crystallized from the melt at 175°C (A-C) and the subsequent cooling to room temperature (B-C).

Figure A3. 11 .

 11 Figure A3.11. WAXS (A-B) and SAXS (C-D) patterns of the MMw-LPm isothermally crystallized from the melt at 170°C (A-C) and the subsequent cooling to room temperature (B-C).

Figure A3. 12 .

 12 Figure A3.12. WAXS (A-B) and SAXS (C-D) patterns of the MMw-LPm isothermally crystallized from the melt at 165°C (A-C) and the subsequent cooling to room temperature (B-C).

Figure A3. 13 .

 13 Figure A3.13. WAXS (A-B) and SAXS (C-D) patterns of the MMw-LPm isothermally crystallized from the melt at 160°C (A-C) and the subsequent cooling to room temperature (B-C).

Figure A3. 14 .

 14 Figure A3.14. WAXS (A-B) and SAXS (C-D) patterns of the MMw-LPm isothermally crystallized from the melt at 155°C (A-C) and the subsequent cooling to room temperature (B-C).

Figure A3. 15 .

 15 Figure A3.15. WAXS (A-B) and SAXS (C-D) patterns of the MMw-LPm isothermally crystallized from the melt at 150°C (A-C) and the subsequent cooling to room temperature (B-C).

Figure A3. 16 .

 16 Figure A3.16. Crystallinity evolution (Vc) of (A) HMw-HPm, (B) MMw-HPm and (C) MMw-LPm during the isothermal ramps obtained by DSC.

Figure A3. 17 .

 17 Figure A3.17. Crystallinity evolution (Xc) of (A) HMw-HPm, (B) MMw-HPm and (C) MMw-LPm during the isothermal ramps obtained by WAXS.

Figure A3. 18 .

 18 Figure A3.18. DSC thermograms of HMw-HPm upon non-isothermal crystallization from the melt at different cooling rates (A) and the subsequent heating at 10°C/min (B).

Figure A3. 19 .

 19 Figure A3.19. DSC thermograms of MMw-HPm upon non-isothermal crystallization from the melt at different cooling rates (A) and the subsequent heating at 10°C/min (B).

Figure A3. 20 .

 20 Figure A3.20. DSC thermograms of MMw-LPm upon non-isothermal crystallization from the melt at different cooling rates (A) and the subsequent heating at 10°C/min (B).

Figure A3. 22 .

 22 Figure A3.22. WAXS patterns of the HMw-HPm non-isothermally crystallized from the melt at 10°C/min (A) and the subsequent heating at 10°C/min (B).

Figure A3. 23 .

 23 Figure A3.23. WAXS (A-B) and SAXS (C-D) patterns of the HMw-HPm non-isothermally crystallized from the melt at 5°C/min (A-C) and the subsequent heating at 10°C/min (B-C).

Figure A3. 24 .

 24 Figure A3.24. WAXS (A) and SAXS (B) patterns of the MMw-HPm non-isothermally crystallized from the melt at 50°C/min.

Figure A3. 25 .

 25 Figure A3.25. WAXS (A) and SAXS (B) patterns of the MMw-HPm non-isothermally crystallized from the melt at 10°C/min.

Figure A3. 26 .

 26 Figure A3.26. WAXS (A-B) and SAXS (C-D) patterns of the MMw-HPm non-isothermally crystallized from the melt at 5°C/min (A-C) and the subsequent heating at 10°C/min (B-C).

Figure A3. 27 .

 27 Figure A3.27. WAXS (A-B) and SAXS (C-D) patterns of the MMw-HPm non-isothermally crystallized from the melt at 1°C/min (A-C) and the subsequent heating at 10°C/min (B-C).

Figure A3. 28 .

 28 Figure A3.28. WAXS (A-B) and SAXS (C-D) patterns of the MMw-LPm non-isothermally crystallized from the melt at 50°C/min (A-C) and the subsequent heating at 10°C/min (B-C).

Figure A3. 29 .

 29 Figure A3.29. WAXS (A-B) and SAXS (C-D) patterns of the MMw-LPm non-isothermally crystallized from the melt at 10°C/min (A-C) and the subsequent heating at 10°C/min (B-C).

Figure A3. 30 .Table A3. 7 .Table A3. 8 .

 3078 Figure A3.30. WAXS (A-B) and SAXS (C-D) patterns of the MMw-LPm non-isothermally crystallized from the melt at 5°C/min (A-C) and the subsequent heating at 10°C/min (B-C).

Figure A3. 32 .

 32 Figure A3.32. POM images of the thin films of HMw-HPm containing 0 wt% and 2 wt% of OXA-2 non-isothermally crystallized from the melt at 20°C/min.

Figure A3. 33 .

 33 Figure A3.33. POM images of the thin films of A-UHMw-HPm containing 0 wt% and 2 wt% of OXA-2 non-isothermally crystallized from the melt at 20°C/min.

Figure A3. 34 .

 34 Figure A3.34. POM images of the thin films of PLLA containing 0 wt% and 2 wt% of OXA-2 non-isothermally crystallized from the melt at 20°C/min.

Figure A3. 35 .

 35 Figure A3.35. POM images of the thin films of PDLA containing 0 wt% and 2 wt% of OXA-2 non-isothermally crystallized from the melt at 20°C/min.

Figure A3. 36 .

 36 Figure A3.36. POM images of the thin films of MMw-HPm containing 0 wt% of OXA-2 isothermally crystallized from the melt at 170°C.

Figure A3. 37 .

 37 Figure A3.37. POM images of the thin films of MMw-HPm containing 2 wt% of OXA-2 isothermally crystallized from the melt at 170°C.

Figure A3. 38 .

 38 Figure A3.38. POM images of the thin films of HMw-HPm containing 0 wt% of OXA-2 isothermally crystallized from the melt at 170°C.

Figure A3. 39 .

 39 Figure A3.39. POM images of the thin films of HMw-HPm containing 2 wt% of OXA-2 isothermally crystallized from the melt at 170°C.

Figure A3. 40 .

 40 Figure A3.40. A-UHMw-HPm thermogram obtained by DSC upon cooling at 5°C/min and subsequent heating at 10°C/min.

Figure A3. 41 .

 41 Figure A3.41. POM images of the thin films of A-UHMw-HPm containing 0 wt% and 2 wt% of OXA-2 non-isothermally crystallized from the melt at 20°C/min.

Figure A4. 4 .Figure A4. 5 .

 45 Figure A4.4. SEM images of NP-HC-PLA/DOX

Figure A4. 6 .

 6 Figure A4.6. Heating ramp followed by DSC of the NP-PLA after the isothermal treatment at 65°C.

Figure A4. 7 .

 7 Figure A4.7. WAXS (A) and SAXS (B) of the gelatin hydrogel

Figure A4. 9 .

 9 Figure A4.9. DSC thermograms of the Gelatin and NP-PLA bionanocomposites during the heating step from room temperature (A) and the isothermal step at 65°C (B).

Figure A4. 10 .

 10 Figure A4.10. Histograms of the size cell (A) and the size cell above a certain size (B) of the x2000 magnification images for both GE without (w/o NP) and with PLA NPs (w NP).

Figure 1 .

 1 Figure 1. Résumé graphique de la thèse

  L'amorceur organométallique employé pour la polymérisation du PLA constitue le premier amorceur capable de générer les copolymères stéréo-diblocs longtemps désirés sans l'utilisation d'un co-amorceur, ce qui représente un avantage pour l'évolutivité vers la production industrielle. L'initiateur suit un mécanisme de coordination-insertion, présente un caractère vivant ainsi qu'une réaction cinétique de premier ordre. De plus, une constante de vitesse similaire pour les deux énantiomères PLA a été trouvée, ce qui permet la synthèse de copolymères stéréoblocs en un temps relativement court, ainsi que la synthèse de PLA racémique. La tacticité des copolymères séquencés peut être adaptée en modifiant le moment auquel le second monomère est ajouté à la réaction, bien que la diminution de la pureté énantiomérique de chaque monomère puisse également être une solution pour adapter la tacticité.Par la suite, la caractérisation du mécanisme de cristallisation à partir de la masse fondue, sous différentes conditions thermiques, des copolymères stéréo-diblocs PLA a été réalisée à différentes échelles grâce à une approche multi-techniques. Il est important de noter que la séparation de phase typique se produisant dans le mélange énantiomère de PLA est

Figure 2 .Figure 3 .

 23 Figure 2. Résumé du procédé de synthèse

  Mechanical properties .............................................................................................................................. 1.5.2 Degradation ............................................................................................................................................. 1.5.2.1 Polymer factors .......................................................................................................................... 1.5.2.2 Media factors .............................................................................................................................

	Table of contents
	1.1 Introduction
	1.2 Poly-(lactic acid)
	1.3 Synthesis of PLA
	1.4 PLA processing
	1.5 PLA properties
	1.5.1

1.6 Medical applications 1.6.1 Pharmaceutical applications .................................................................................................................... 1.6.1.1 Nanoparticles ............................................................................................................................. 1.6.1.2 Hydrogels ...................................................................................................................................

1.6.1 Pharmaceutical applications 1.6.1.1 Nanoparticles

	1.6 Medical applications
	PLA and its copolymers are currently used at the clinical level in different
	pharmaceutical applications such as dermal fillers and drug delivery systems (DDS) as well as
	biomedical applications such as sutures or tissue engineering (TE) [134-137]. The first
	biomedical application of PLA at the clinical level, namely the suture Vycril ® composed of
	glycolide and lactide components, dates from 1974. Since then, several sutures composed of
	different ratios of glycolide and lactide have reached the market, such as Polysorb ®, Radik ®
	or XLG ® [134]. Moreover, Decapeptyl SR was the first pharmaceutical product based on
	PLGA microparticles approved in Europe in 1996 for the treatment of prostate cancer.
	Thereafter, more than 15 microparticles PLA-based products have been approved and
	marketed to control drug delivery [135] as well as a soft-tissue augmentation agent [136,137].
	PLA is among the biocompatible polymers the most frequently used as DDS (Table
	1.1) due to the ease to tune the drug release profile from PLA formulations by modifying
	different architectural parameters such as polymer molecular weight, stereochemical
	composition or polymer crystallinity [138]. Nowadays, several PLA-based microparticles
	formulations have successfully reached the market [135], however, microsized DDS features
	inherent drawbacks to traverse biological barriers as well as their rapid clearance after
	systemic administration, limiting microparticles application to subcutaneous and
	intramuscular injections [139]. In contrast, nanosized DDS increase systemic circulation time
	and facilitate tissue penetration as well as sterilization via filtration [139]. PLA nanoparticles
	have been designed for the treatment of several pathologies such as cancer [140,141], Chagas
	disease [142], Alzheimer disease [143] or insulin-dependent diabetes [144]. Besides, the surface
	of PLA nanoparticles has been functionalized with antibodies [145] or cell-penetrating
	peptides [146,147] as a synthetic approach to target specific tissues or cells and both increase
	the therapeutic effect and decrease side effects. Moreover, the attachment of PLA nanoparticles
	to cells exhibiting blood barrier brain migratory properties (T lymphocyte) has also been
	proposed as a novel approach to drug delivery into the brain [148].
	However, the unusual establishment of the correlation of the structure-property
	relationship of the PLA-based nanoparticles after formulation, such as the Tg, the crystallinity

Table 1 .1

 1 Summary of the PLA-derivative nanoparticles.

	or the physical ageing unable to understand the drug release mechanism from a nanoscopic
	scale. Particularly, the rearrangement of polymer chains during physical ageing, which
	typically occurs during the storage step and depends on the Tg, may increase water absorption
	leading to an initial burst release [150]. The control or the analysis of the polymeric

Material Processing Technique Polymer Structure Molecular Weight (kDa) Material Characterisation 1 Studies 2 Ref. Nanoparti cles

	physicochemical parameters of the formulations is crucial to designing a specific degradation
	rate considering the influence of the conditions on the drug-release mechanism followed by
	the formulation: drug diffusion through water-filled pores, diffusion through the polymer
	matrix, osmotic pumping or surface and bulk erosion [139].
	Double emulsion-solvent evaporation 1.6.1.2 Hydrogels	PLLA-PEG PLLA	85-160	Size, PDI, EE (%)	In vitro/in vivo [140]
	Nanoprecipitation-solvent displacement PLA-based hydrogels that are formed through different external stimuli have already Rac-PLA -Size, PDI, Z-potential, LE (%), EE (%), Morphology (SEM) In vitro/in vivo/clinical reached the clinical level. Particularly, Atridox® is a doxycycline formulation used for chronic	[141]
	Interfacial deposition-solvent displacement periodontitis containing PLA and N-methyl-2-pyrrolidone mixtures that solidifies to a wax-Rac-PLA-Peg 18-28 Size, PDI, Z-potential In vivo like consistency upon contact with gingival crevicular fluids. Furthermore, OncoGel® is a	[142]
	Emulsion-solvent evaporation paclitaxel formulation used for local tumour treatment based on PLGA-PEG-PLGA triblock Maleimide-PEG-PLA (no PLA specification) PEG3-PLA70 copolymers that exhibit a sol-gel transition at body temperature [151]. The fabrication of PLA-Size, Z-potential, In vivo hydrogels usually requires copolymerization with other polymers such as PEG, polyurethane Morphology (TEM) Methoxy-PEG-PLA (no PLA specification) and poly glycolic acid [152] as well as the blend with collagen and gelatin [49,50] to increase PEG3-PLA50 the hydrophilic character. Moreover, different physicochemical and mechanical properties,	[143]
	Emulsion-solvent evaporation such as the drug release rate, can be modulated by varying the size of PEG, polymer CE (%) complexation PLGA (75L:25G) 15 efficiency, Z-potential, In vivo concentration as well as the stereoisomery of PLA [153,154]. Likewise, a series of morphology (SEM) Nanoprecipitation-solvent displacement stereocomplexed PLA pentablock hydrogels were synthesized with tailored microstructure by Trastuzumab-PEIcoating-Rac-PLA 22 LE (%), EE (%), Morphology In vitro varying the stereostructure, crystallization and stereocomplexation and their physical (TEM) properties were thoroughly characterized (Table 1.2). The PLA hydrogels featuring symmetric	[144] [145]
	Maleimide-PEG-PLA (no pentablock copolymers exhibited higher gel-sol transition temperature, higher storage PEG3.4-
	Emulsion-solvent evaporation modulus and slower biodegradation and drug release compared to asymmetric pentablocks PLA specification) PLA34 Size, PDI, Z-potential, In vitro/In vivo [146] [155]. Recently, the synthesis of hydrophobic polymers, in particular PLA, to produce Morphology (TEM) Methoxy-PEG-PEG3.4-thermohydrogels has been achieved without the use of solvent, which enhances PLA (no PLA specification) PLA30 thermohydrogels biocompatibility [156]. Furthermore, a PLA-based hydrogel was designed
	Maleimide-that might be applied as a delivery vehicle with immediate release upon contact with reducing
	Emulsion-solvent conditions, such as tumour microenvironments [157]. The polymer solution gelled between PEG-PLA (no PLA specification) PEG3.4-PLA34 Size, PDI, Z-potential, EE (%), LC (%), Morphology 32°C and 40°C due to the decrease of polymer solubility and collapsed upon exposure to In vitro/In vivo [147] evaporation (TEM) Methoxy-PEG-strong reducing agents due to the cleavage of the disulphide bond incorporated in the polymer PEG3.4-PLA (no PLA PLA30 structure [157]. Likewise, a PLA-based hydrogel was recently designed for the treatment of specification) Nanoprecipitation-solvent displacement Rac-PLA-OH chronic inflammatory diseases, such as rheumatoid arthritis, which fulfil an on-demand drug 10.5 Size, Z-potential, In vitro [148] Rac-PLA-PEG-PLA4.5-release depending on the severity of the disease by using nitric oxide (overproduced in Morphology (TEM) NH2 PEG3.5-NH2 inflammatory environments) as a hydrogel degradation agent [158]. PLA-based injectable
	Nanoprecipitation-solvent evaporation hydrogels were also envisaged to be used for regenerative medicine requiring higher PLA -Size, Z-potential, pH, EE(%), LE (%), Morphology (AFM, SEM), Molecular structure In vitro mechanical properties. Particularly, hydrogel composites comprising electrospun-PLA	[149]
	(FTIR) nanofibers and gelatin nanoparticles lead to a 15-fold increase of the storage modulus without
	compromising the injectability [159]. Moreover, the incorporation of PLA nanoparticles into a
	gelatin hydrogel also leads up to a 100-fold increase in the viscosity, without compromising
	the injectability [110].

Table 1 .2 Summary of PLA-derivative hydrogels. 1.6.1.3 Perspectives 1.6.1.3.1 Pharmaceutical design

 1

	The application of 3D printers for the customization and personalization of
	pharmaceuticals is one of the most revolutionary and powerful tools in the last decades
	[162,163]. The traditional « one-size-fits-all » treatment approach in the healthcare industry is
	ineffective in up to 70% of patients, according to the National Health Service, due to varied
	pharmacokinetic traits of different patients, creating the need to shift from mass production to
	personalized dosage medicine [164]. 3DP technologies have been used for the production of
	pharmaceutics for pre-clinical animal models, due to the accuracy in producing dosage forms
	of appropriate geometry and size [165]. Besides, the Food and Drug Administration (FDA)
	approved in 2015 the first drug manufactured using 3DP technology, Spritam® [166,167],
	opening the scope of oral drug delivery using 3DP. Likewise, a five-in-one dose combination
	polypill was manufactured by three-dimensional extrusion with two independently
	controlled and well-defined release profiles [168], simplifying the drug therapy and
	potentially improving the adherence of patients to the prescribed treatment. Additionally,
	artificial intelligence machine learning techniques have been developed to predict key 3DP
	fabrication parameters and advance the 3DP fabrication process, reducing time, costs and
	resources normally invested in formulation development [74] (Figure 1

.6). The current variety Material Processing Technique Polymer Structure Molecular Weight (kDa) Material Characterisation 1 Studies 2 Ref.

		Temperature increase of solutions (RT --> 37 °C)	PLLA-Castor Oil Rac-PLA-Castor Oil	4.8 4.6	Viscosity, specific optical rotation, Tg, shear stress	-	[153]
			PLLA	1.5	Sol-gel transition,
		Temperature	Rac-PLA	1.5	micelles size, circular
		increase of solutions	PLLA-PEG	14.4	dichroism, morphology	-	[160]
		(RT --> 37 °C)	Rac-PLA-PEG	14.4	(TEM), nanostructure (XRD)
			PLLA-PEG-
	Hydrogels	Hydrogelation by concentration	PLLA Rac-PLA-PEG-	11.5-15.5	Storage modulus, nanostructure (WAXD)	-	[161]
			Rac-PLA
		Temperature increase of solutions (RT --> 37 °C)	OS-rac-PLA-serin) PEG-rac-PLA-OS (olygomer	~3	Sol-gel phase transition (depending on PEG Mw)	In vitro/In vivo	[154]
		"Click" reaction	DA-NOCCL + N3 + rac-PLA-PEG-N3 (NP)	~6	Aggregates size, properties morphology (Crio-SEM and TEM), mechanical	In vitro/In vivo	[158]

Deng, Polylactide-based chiral particles with enantio- differentiating release ability, Chem. Eng. J. 344 (2018) 262-269., copyright 2018, with permision from Elsevier 1.6.1.3.3 Antibacterial applications

	Bacterial infections originated from implants and medical devices, are typically treated
	with antibiotics and their misuse has promoted their resistance to drugs. Multidrug-resistant
	bacteria is one of the main threats to health and food security worldwide, according to the
	World Health Organization, and alternative treatments mainly based on bacterial proliferation
	prevention need to be addressed [184]. Polymers usually do not exhibit intrinsic antibacterial
	properties, although antibacterial additives such as silver nanoparticles [185,186], titanium
	dioxide [187] or essential oils [188] can be used to develop nanocomposites exhibiting
	antimicrobial activity. Moreover, antibacterial PLA filaments for 3DP containing cooper
	particles are already commercially available (PLACTIVE AN1 Copper3D [189], PLActive Red
	[190]) and antibacterial 3D prostheses have been developed to minimize skin disorders related
	to microbial infections [191]. PLA-based nanocomposites exhibiting antibacterial
	characteristics can be manufactured in a wide variety of architectures due to the capability of
	PLA to be processed by different techniques such as 3DP [192], electrospinning [193] (Figure

1.8), extrusion

  [194], spin coating[195], etc. that broads its applicability in the biomedical field.

Table 1 .3 Summary of PLA-derivative polymer therapeutics. Material Processing Technique Polymer Structure Molecular Weight (kDa) Material Characterisation 1 Studies 2 Ref. Polymer therapeutics

 1

	ROP mediated by	PEG-RacPLA-
	Paclitaxel ROP mediated by	Ptxl PEG-RacPLA-	~30	Structure (NMR), Mw (GPC)	-	[198]
	Docetaxel	Dtxl
	Azide-alkine click reaction	Ptxl (23%)-RacPLA-PEG (25%)	~10-15	Size (DLS), structure (GPC) (NMR), morphology (TEM), molecular weight	In vitro	[199]
	Conjugation + micelles formulation	RacPLA-PEG-RacPLA-Lisinorpil	~10	Size (DLS), morphology (AFM)	In vitro drug release	[200]
	Hot press film	PLLA-indomethacin	-	Conjugation contact angle (fluorescence, XPS),	-	[201]
	ROP	(PLLA)3-Camptothecin [3-armed]	~30	Molecular weight (GPC), structure (NMR), morphology (AFM), DSC) thermal properties (TGA,	In vitro drug degradation release and enzymatic	[202]

1.6.2 Biomedical applications 1.6.2.1 Stents

	Stents are medical devices placed in a lumen of the body, particularly blood vessels, to
	prevent its closure. Metal stents are typically used to treat coronary disease due to their
	flexibility, radial force, resistance to fracture, radiopacity, biocompatibility and low
	thrombogenicity. However, several late clinical complications including stent thrombosis,
	restenosis and neoatherosclerosis still exist [17]. Bioresorbable vascular scaffolds emerged as

Table 1 .4 Summary of PLA-derivative stents. Material Processing technique Polymer structure Molecular weight (kDa) Material characterization 1 Studies 2 Ref STENTS

 1

	Bought (REMEDY) 3	Rac-PLA	-	-	Clinical	[15]
	Desolve Cx 3	PLLA-Novolimus	-	In vitro/ in vivo lost) degradation (MW	Clinical	[16]
	Mirage 3	PLLA (4% D-LA)-sirolimus	-	-	Clinical	[17]
					Clinical trial
	ABSORB V G2 3	PLLA	-	-	(RENASCENT	[203]
					III)
	MAGNITUDE 3	PLLA-sirolimus	-	-	Clinical	[18]

1.6.2.2 Tissue engineering

	stents were found to increase as both the initial pitch angle as well as the initial diameter
	decreased[210].
	Tissue engineering (TE) is one of the main strategies of regenerative medicine that
	consists of the regeneration of neotissues by applying 3D scaffolds for cell attachment and
	growth [211]. The designed scaffolds must fulfil different criteria with ambivalent properties
	such as biocompatibility, and mechanical properties suitable for the intended application
	whilst supporting the normal functioning of cells and tissues. Polymers are the most popular
	materials for scaffold production due to their physicochemical characteristics and bioactivity
	[212]. Particularly, synthetic polymers have drawn large attention as an alternative to natural
	polymers such as peptides, nucleic acid and polysaccharides, due to their price and ease to be
	produced [213]. Among the synthetic polymers, PLA, PGA, poly-(e-caprolactone) (PCL),
	polydioxane and poly-(trimethylene carbonate) are currently the most extensively used for
	tissue engineering applications [213] and the specific polymer selection is based on the
	structure-property relationship required for the targeted application. Likewise, PCL features
	high drug permeability and less acidic byproducts compared to other polyesters but a
	relatively slow degradation rate, limiting its usage to long term applications. In contrast, PLA
	presents a faster degradation rate compared to PCL, although still relatively slow. However,
	long degradation times coupled with the high crystallinity of the remained fragments might
	cause inflammatory reactions in the body. Likewise, the PLA is used frequently as rac-PLA to
	diminish detrimental health effects, being the latter rapidly degraded without the formation
	of crystalline fragments [14]. Besides, alternative approaches to circumvent the individual
	limitations of the single polymers such as synthetic co-polymers, the combination with natural
	polymers or/and the scaffold surfaces functionalization are frequently used to overcome the
	used manifested hydrophilicity, cell attachment and biodegradability issues [212].
	However, 3D scaffolds generated from PLA and its copolymers [212] have already been
	Subsequently, several drug-eluting stents mainly based on PLLA evaluated for a wide variety of regenerative applications with different tissues such as bone
	and Rac-PLA were developed and particularly BVS Absorb is the more spread [204]. BVS [214], nerve [215], tendon [216] and cartilage which enlarge the applicability of conventional
	Absorb was approved in 2016 by the FDA based on non-inferiority test after 1 year of ceramic and metal solutions focussed on the orthopaedic field (Table 1.5). Particularly,
	implantation compared to the corresponding "gold-standard", namely the metallic polymers can be shaped into complex topologies for customised medical solutions with the
	everolimus-eluting stent [205]. However, compared with metallic stents, the BVS appears to advent of 3D printing technologies.
	be associated with both lower efficacy and higher thrombotic risk after 2 years on average
	[206] mainly due to scaffold discontinuity, malposition due to under expansion and
	neoatherosclerosis [207,208]. In addition, the current challenging generation of PLA-based
	scaffolds with equivalent radial strength and flexibility to their metallic stent counterparts
	[209] was recently faced by investigating the influence of the stent geometric parameters on
	the mechanical properties. The radial stiffness and peak compression force of the PLLA helical

Table 1 .5 Summary of PLA-derivative scaffolds. Material Processing technique Polymer structure Molecular weight (kDa) Material characterization 1 Studies 2 Ref SCAFFOLDS

 1

	Freeze-drying, salt-leaching (3D scaffold)	PLGA (75L:25G)+ Silk +HA( Hidroxiapatite)	90-126	Structure (FTIR), (%), Mechanical properties Degradation temperature (TGA), Morphology (SEM), Swelling (%), water uptake	In vitro/ In vivo	[214]
	Wet-spinning (Microfilament)	P(L-co-rac-LA)(75:25)	200	Morphology (SEM)	In vivo	[215]
	Bought (kinnet filmanets)	PGA/PLA (2:1)	-	Morphology (SEM and TEM), Biomechanical properties	In vitro/In vivo	[216]
		PLLA-PGA
	Bought (forged composite sheets)	(88:12) + HA	-	Morphology (SEM)	In vivo	[217]
		PLLA + HA
	bought			Molecular weight,
	(OSTEOTRANS	PLLA + HA	-	Crystallinity, Morphology	Clinical	[218]
	MX)			(SEM) [after surgery]
		PLLA/PGA
	Bought (Plates)		-	-	Clinical	[219]
		PLLA/HA
				Morphology (SEM),
	Electrospinning	PLGA (85/15)	285	Mechanical properties,	In vitro	[220]
				Structure (FTIR)
				Morphology (SEM),
	Electrospinng	PLGA (85/15)	285	Mechanical properties,	In vitro	[221]
				Structure (FTIR)
	Supercritical Emulsion extraction	PLGA carriers + fibrin hydrogel	38-54	Size, Morphology (SEM)	In vitro	[222]
	Electrospinning	PLGA + silk	-	Morphology (SEM), Mechanical properties	In vitro	[223]
	Knitted fibers	PLGA (10/90)	-	Mechanical properties	In vivo	[224]

Table of contents 2.1 Introduction 2.2 Aims and objectives 2.3 Results and discussion

2.3.1 Characterization and activity of the Initiator .......................................................................................... 2.3.2 Polymerization mechanism ...................................................................................................................... 2.3.3 Polymerization stereocontrol ................................................................................................................... 2.3.4 Polymerization kinetics ........................................................................................................................... 2.3.5 Homochiral and Racemic Polymerization of Lactide ............................................................................... 2.3.6 Stereo-Block Copolymerization by Sequential Monomer Addition ......................................................... 2.4 Conclusions 2.5 Experimental section 2.5.1 Materials .................................................................................................................................................. 2.5.2 Synthesis .................................................................................................................................................. 2.5.2.1 Synthesis of the Lig and 2,2-bis((3,5-dimethyl-pirazol-yl)-1-para-tolylethoxide) (bpzteH) 89 2.5.2.2 Synthesis of the Initiator [Zn(Et)]-[2,2-bis((3,5-dimethyl-pirazol-yl)-1-paratolylethoxide)] (Zn(Et)(κ3-bpzteH) ..................................................................................................... 2.5.2.3 Polymerizations of LA .............................................................................................................. 2.5.2.4 Kinetic experiments .................................................................................................................. 2.5.3 Instrumentation and measurements ........................................................................................................ 2.5.3.1 1 H NMR, 13 C NMR and homonuclear decoupled 1 H NMR ............................................... 2.5.3.2 MALDI-TOF analysis................................................................................................................ 2.5.3.3 GPC analysis .............................................................................................................................. 2.5.3.4 Optical rotation .......................................................................................................................... 2.5.4 Computational methodology .................................................................................................................... 2.5.5 Calculation of Pm values ..........................................................................................................................

Table 2 . 1 .

 21 Tetrad probabilities of ESC and CEC mechanisms based on non-Bernouillan and Bernouillan statistics.

	Tetrad	Probability of ESC (non-	Probability of CEC (Bernouillan)
		Bernouillan)
	mmm	[Pm 2 +(1-Pm) 2 +Pm 3 +(1-Pm) 3 ]/2	Pm 2 +0.5PmPr
	mmr	[Pm

2 (1-Pm)+Pm(1-Pm) 2 ]/2 0.5PmPr mmm [Pm 2 (1-Pm)+Pm(1-Pm) 2 ]/2 0.5PmPr rmr [[Pm 2 (1-Pm)+Pm(1-Pm) 2 ]/2 0.5PmPr mrm [Pm(1-Pm)] 0.5(Pm 2 +PmPr) ppm Peak Analysis [mmm]

Table 2 . 2

 22

. Rate constants. a [LA]0/[Zn]0=400; [LA]0= 400 mM; [Zn]0= 1mM. b [LA]0/[Zn]0=500; b [LA]0= 625 mM;[Zn]0=1.25mM

Table 2

 2

	Entry	Sample	T (°C)	kapp x 10^2 (s -1 )
	1	a PLLA400	90	6,59 ± 0,002
	2	a PLLA400	75	2,56 ± 0,002
	3	a PLLA400	60	1,05 ± 0,0002
	4	b PLLA500	90	5,61 ± 0,003
	5	b PDLA500	90	4,71 ± 0,001

.2, entries 4 and 5) with a difference in monomer conversion of 10% after 45 min of reaction (85 and 95% for D-and L-LA, respectively). The similar polymerization rate constants for both D-and L-LA enantiomers obtained from

  .3 and TableA2.1, Entries 6-14) from previously generated conditions to attain HMw PLAderivatives[20]. Mild reaction conditions at low [LA]/[cat] with 100 eq. of L-LA in THF at 50 °C prevented the onset of the polymerization even after 2 h of reaction (Table2.3, Entry 1).The solvent exchange by a lower polarity solvent as toluene increased the conversion only to 5%. Only under more extreme conditions was the Initiator able to promote the polymerization, converting 43 and 94% of monomer, in 1 h at 70 and 90 °C, respectively (Table2.3, Entries 3-4). HMw polymers obtained with monomer conversions of 94% were obtained after reaction

	Entry	Sample	[LA]/[cat]	Temperature (°C)	Time (min)	Conv. (%)	Mn (theor.) (Da) a	Mn (exp.) (Da) b PDI [α] c
	1	PLLA100	100	50	120	0	0	-	-
	2	PLLA100	100	50	120	5	720	-	-
	3	PLLA100 (L1)	100	70	60	43	6.192	12.824 1,36 -159,4
	4	PLLA100 (L2)	100	90	60	94	13.536	30.055 1,62 -171,4
	5	PLLA500 (L3)	500	90	45	94	67.680	62.564 1,84 -172,2
	6	PLLA500 (L4)	500	70	90	76	54.720	42.880 1,7 -162,7
	7	PLLA500 (L5)	500	70	120	94	67.680	48.253 2,19 -168,4
	8	PDLA500 (D1)	500	70	90	85	61.200	41.292 1,78 154
	9 rac-PLA600 (Rac1)	600	70	120	90	64.800	40.754 1,59 -1,1
								homopolymers L4
	and D1 present similar Mn value whilst different conversions, which might be ascribed to an
	error in the conversion due to the heterogeneity of the reaction mixture ascribed to the high
	viscosity. Molecular weight distributions (PDI) obtained by GPC were found to be narrow,
	although slightly increased for HMw PLA derivatives due to both the steric hindrance and

with 500 eq. of L-LA at 70 and 90 °C, in 120 and 45 min, respectively (Table

2

.3, Entries 5 and 7). Moreover, monomer conversions higher than 85% were achieved for D-and rac-LA in 2 h (Table

2

.3, Entries 8-9). Table 2.3. Polymerizations of L-, D-and rac-LA Polymerization conditions: 25 µmol of Initiator 1, toluene as solvent except entry 1, 5 and 20 mL of solvent for [LA]/[cat]=100 and 500, respectively. a Theoretical Mn = (monomer/Initiator 1) x (% conversion) x (Mw of LA). b Determined by GPC relative to polystyrene standards in chloroform. c Specific optical rotation. ([α]PLLA=-173°) [25]. The obtained molecular weights were in good agreement with the calculated values as the LA/Initiator ratio increased, confirming the living character of the polymerization as well as the single-site catalyst mechanism, apart from polymerizations at ratio [LA]/[cat] = 100 (L1 and L2), which are likely related to a high kprop./kinic. ratio. However, PLAviscosity rise in the reaction mixture. Furthermore, polymerization of L-LA and D-LA yielded stereochemically pure PLA (Figures A2.7-A2.13) indicating that epimerization of either the monomer or the polymeric derivative was absent under the polymerization conditions.

Table 2 .4 and Figures A2.15-A2.36), which

 2 consist on the addition of the second monomer to the reaction when the conversion of the first one has reached a certain percentage. The names of the PLA-SBC (Table2.4) present the following

structure: X123-Y123, were X and Y denote D and L, depending on the lactide chirality and 123

Table 2 .4. Stereo

 2

	Entry	Sample	[LA]

-block copolymerizations of L-LA and D-LA.

Polymerization conditions: 25µmol of Initiator, toluene as a solvent at 90°C . a 30-60 min was maintained between each monomer addition, depending on the monomer loading. b Theoretical Mn = (monomer/Initiator 1) x (% conversion) x (Mw of LA). c Determined by GPC relative to polystyrene standards in chloroform. d Molar mass distribution= Mw/Mn. e Specific optical rotation ([α]PLLA=-173°)

[25]

f Probability of finding dyads calculated from homonuclear decoupling 1 H NMR spectra after deconvolution; calculations are based on CEC statistics..

/[cat] Solvent Temperature (°C) Time (min) a Conversion (%) Mn(theor.) (Da) b Mw(exp) (Da) c Mn (exp.) (Da) c PDI d [α] e Pm f

		(L115:D115)	250	Toluene	90	50	95	34.299 32.823 17.007 1,93 -2,47 0,96
		(L220:D220)	500	Toluene	90	90	90	64.800 63.313 28.779 2,20 -5,09 0,88
		(L250:D250)	600	Toluene	75	120	81	58.320 72.508 32.958 2,20 -2,67 0,92
		(L275:D275)	600	Toluene	90	60	95	82.080 79.320 35.253 2,25 -1,07 0,9
		(D275:L275)	600	Toluene	90	60	94	81.216 79.107 33.238 2,38 1,41 0,81
		(L330:D330)	700	Toluene	90	60	84	72.576 94.181 44.636 2,11 -1,93 0,69
		(L375:D375)	800	Toluene	90	100	97	97.776 107.912 63.853 1,69 -0,69 0,95
		(L500:D500)	1100	Toluene	90	110	89	102.528 139.517 53.660 2,60 0,32 0,94
		(L520:D520)	1200	Toluene	90	120	85	146.880 150.094 64.976 2,31 -1,65 0,92
	10	(L1200:D300)	1600	Toluene	90	95	90	194.400 202.280 99.257 2,04 -87,7 0,93

  .1). Both LA and initiator Schlenk flasks were attached to the vacuum line and temperature equilibrium was ensured by stirring the solutions for 15 min in an oil bath. Monomer and catalyst solutions were poured together by a glass bent adaptor and polymerization onset times were measured from that point. Stereo-diblock copolymerizations were achieved by the addition of the second monomer upon the first monomer high conversion that was probed by monitoring reaction aliquots by 1 H NMR. Methanol was used to terminate the reaction and precipitate out the polymer synthesized. The obtained polymer was collected by filtration, dried at room temperature, exposed to vacuum over 24 h in the

Schlenk line, and stored upon characterization. The polymer derivatives experimental details and their respective characterization is detailed in the Appendix 2 (Figures A2.7-A2.

13

for homopolymers and A2.15-A2.

38

for PLA-SBC). Moreover, the names of the PLA-SBC present the following structure: X123-Y123, were X and Y are D and L, depending on the lactide chirality and 123 is the length of the block measured in lactide units.

Table of contents 3.1 Introduction 3.2 Aims and objectives 3.3 Results and discussion 3

  .3.1 Thermal degradation of SBC-PLA ......................................................................................................... 3.3.2 Effect of Mw in the PLA-SBC crystallization ....................................................................................... 3.3.3 Spherulite growth rate of stereo-block copolymers (PLA-SBC) and PLLA and PDLA enantiomers (PLA-HP) ................................................................................................................................................................. 3.3.4 Isothermal crystallization ...................................................................................................................... 3.3.4.1 Avrami model .......................................................................................................................... 3.3.5 Non-isothermal crystallization ..............................................................................................................

3.3.5.1 Avrami modified model ......................................................................................................... 3.3.5.2 Ozawa model ........................................................................................................................... 3.3.5.3. Avrami-Ozawa model ........................................................................................................... 3.3.6 Crystallization of SBC with nucleating agents ..................................................................................... 3.3.6.1 Nucleating agent efficacy survey .......................................................................................... 3.3.6.1.1 BHET ...................................................................................................................................... 3.3.6.1.2 OXA ........................................................................................................................................ 3.3.6.2 Morphology of the crystallized NA/SC-PLA derivatives .................................................. 3.4 Conclusion 3.5 Experimental section Materials ......................................................................................................................................................... Methods .......................................................................................................................................................... Polymerizations ................................................................................................................................... Films preparation ................................................................................................................................ Instrumentations and measurements .............................................................................................................

Table 3 .

 3

1. Crystalline phases of PLA.

Table 3 . 3

 33

. Thermal characterization by DSC of the PLA-SBC.

Table 3 . 9 .

 39 Thermodynamic parameters obtained from the thermal analysis by DSC of the MMw-HPm PLA-SBC with OXA-2.

Table of contents 4.1 Introduction 4.2 Aims and objectives 4.3 Results and discussion 4

  .3.1 Selection and characterization of PLA-derivatives ................................................................................ 4.3.2 Nanoparticles formulation ..................................................................................................................... 4.3.3 Nanoparticles characterization .............................................................................................................. 4.3.4 Bionanocomposite formation.................................................................................................................. 4.3.5 Rheological properties of the bionanocomposite ..................................................................................... 4.3.6 Drug release analysis .............................................................................................................................

	4.4 Conclusions
	4.5 Experimental section

4.5.1 Materials ................................................................................................................................................ 4.5.2 Methods .................................................................................................................................................

Table 4 . 1 .

 41 Correspondence between the abbreviated name and the composition.

	NAME	GELATIN	PLA	BULK POLYMER/ NANOPARTICLE	DRUG
	Rac-PLA	-	Racemic	Bulk	-
	HC-PLA	-	Homopolymer	Bulk	-
	SC-PLA	-	Stereoblock copolymer	Bulk	-
	GE	Yes	-	-	-
	GE/DOX	Yes	-	-	Doxorubicin
	NP-Rac-PLA	-	Racemic	Nanoparticle	-
	NP-HC-PLA	-	Homopolymer	Nanoparticle	-
	NP-SC-PLA	-	Stereoblock copolymer	Nanoparticle	-
	NP-Rac-PLA/DOX	-	Racemic	Nanoparticle	Doxorubicin
	NP-HC-PLA/DOX	-	Homopolymer	Nanoparticle	Doxorubicin
	NP-SC-PLA/DOX	-	Stereoblock copolymer	Nanoparticle	Doxorubicin
	NP-Rac-PLA/DAS	-	Racemic	Nanoparticle	Dasatinib
	NP-HC-PLA/DAS	-	Homopolymer	Nanoparticle	Dasatinib
	NP-SC-PLA/DAS	-	Stereoblock copolymer	Nanoparticle	Dasatinib
	GE/NP-Rac-PLA/DOX	Yes	Racemic	Nanoparticle	Doxorubicin

Table 4 . 2 .

 42 Bulk polymer characterization. Molecular weight in number average obtained by GPC relative to polystyrene standards in chloroform. b PDI=Mw/Mn; Polydispersity Index obtained by GPC. c Glass transition temperature. d Crystallization temperature. e Crystallization enthalpy. f Melting temperature. g Melting enthalpy. h Crystallinity 𝑋 𝑐 = Same polymer as Entry 4 in Table 2.3.

	∆𝐻 𝑓 ∆𝐻 𝑓 0 × 100 with ∆𝐻 𝑓𝐻𝐶 0	=

a

m ol) a PDI b Tg (°C) c Tc (°C) d ΔHc (J/g) e Tm (°C) f ΔHm (J/g) g Xc (%) h Pm i [α] j

	PDLLA	Rac-PLA 59.594 1,77	24,12	-	-	-	-	-	0,52 -1,1
	PLLA	HC-PLA 36.047 1,92	40,84	101,24 46,102	170,24 49,538 53,75 0,99 -159,4
	(*)PLLA-b-PDLA	SC-PLA 35.253 2,25	41,28	113,83 44,767	193,93	46,91	33,03 0,90 -1,07

Table 4 .3 and A4.2) and

 4 SEM (Figures 4.3 and A4.4, as representative examples). Doxorubicin and dasatinib were selected to assess the loading capacity of the NP-PLA with two different model compounds that possess distinct structures either featuring a stiff structure imposed by the aromatic groups for the doxorubicin or a flexible chain structure for dasatinib.

Table 4 .

 4 3 DLS characterization of PLA nanoparticles.

  HS characterized by a pixel size of 40 µm 2 × 40 µm 2 and active area of 85 mm 2 × 255 mm 2 (h × v) was used at NCD-SWEET. The SAXS patterns were collected with a Pilatus 1M with a detector size (981 × 1043) with a pixel size of 172 µm × 172 µm at a sample to detector distance of ca. 2.5 m. The calibration of the scattering angular was performed with either α-Al2O3 (alumina) for the WAXS scattering angles or AgBe for the SAXS scattering angle range.

	SAXS/WAXS experiments were conducted both at DUBBLE (bm26) ESRF in Grenoble
	(France) and bl11 NCD-SWEET at ALBA, Cerdanyola del Vallès (Spain) with an X-ray
	wavelength of 12 Kev. The WAXS patterns were recorded with a Pilatus 300 kw (1472 × 195
	pixels configuration) with a pixel size of 172 µm 2 × 172 µm 2 in DUBBLE, whereas a Rayonix
	LX255-

Table A2 .

 A2

	1. Polymerizations catalysed by Initiator. Polymerization conditions: 25µmol of
	initiator. a Theoretical Mn = (monomer/initiator) x (% conversion) x (Mw of LA). b Determined by
	GPC relative to polystyrene standards in chloroform. c Specific optical rotation ([α]PLLA=-173°)
	[30].

Table A4 .

 A4

	Particle	Average size (nm)	PDI	EE [%]
	NP-Rac-PLA/DOX	286.35	0.17	0.3
	NP-HC-PLA/DOX	214.8	0.18	0.24
	NP-SC-PLA/DOX	270.0	0.17	0.18
	NP-Rac-PLA/DAS	237.6	0.16	1.03
	NP-HC-PLA/DAS	282.85	0.15	2.8
	NP-SC-PLA/DAS	275.55	0.17	0.23

2. DLS analysis of drug loaded nanoparticles

Tg: glass transition temperature; Tc: crystallization temperature; Tm: melting temperature; ?: No information found. a [40] b [38] c [10] d [34] e [31].

CHAPTER 4

Multifunctional PLA/Gelatin Bionanocomposites for Tailored Drug Delivery Systems dissolved in PBS pH = 7. Drug-encapsulated nanoparticles were prepared following the abovementioned methodology. The amount of the drug (1 mg) was dissolved along with the polymer in 4 mL of CH2Cl2. The encapsulated drug in the NP-PLA is indicated at the end of the abbreviation name: NP-rac-PLA/DOX. The correspondence between the sample abbreviations and their compositions is summarized in Table 4.3.

Bionanocomposites generation

The bionanocomposites (GE/NP-Rac-PLA, GE/NP-HC-PLA, GE/NP-SC-PLA, GE/NP-Rac-PLA/DOX, GE/NP-HC-PLA/DOX, GE/NP-SC-PLA/DOX) were prepared by blending NP-PLA or NP-PLA/DOX with shark skin gelatin and tri-distilled water in a 0.4/25/74.6 w/w (NP/GE/H2O) proportion under constant stirring in an ultrasound bath (SELECTA, 40 kHz, 125 W) at 65 °C for 30 min. Furthermore, a hydrogel of gelatin (GE) was prepared at the same ratio 25/75 (GE/H2O) as a reference to compare the drug release [52]. Likewise, the corresponding sample abbreviation with its chemical composition is listed in the Supplementary Information (Table 4.3).

Instrumentation and measurements

1 H NMR

Spectra were recorded at room temperature on a Varian Inova FT-400/500 spectrometer and referenced to the standard in the deuterated solvent with the relaxation time fixed to 4 s.

The polymeric derivatives were dissolved in CDCl3 for the synthetic characterization by 1 H NMR.

GPC analysis

The molecular weight and polydispersity index of the obtained PLA derivatives were determined by gel permeation chromatography (GPC) with a Shimadzu Prominence-I LC-2030 equipped with a Shodex GPC KF-805L column (Shodex, Tokyo, Japan) and a Shimadzu RID-20A detector (Shimadzu, Kyoto, Japan). CHCl3 was used as the mobile phase at 40 °C, with 1mL/min as the flow rate. Then, 5 mg of the PLA derivatives was dissolved in 1.5 mL of solvent overnight in constant agitation. Subsequently, the PLA solutions were filtered through a 0.2 µm polytetrafluoroethylene (PTFE) syringe filter before injection. Polystyrene with a molecular weight of 172 kDa was used as a reference to determine the overall molecular weight of the samples.

Optical rotation

The specific optical rotation [α] was measured by a JASCO P-20000 WI using a beam wavelength of 598 nm, using dichloromethane as solvent at a concentration of 1 mg/mL.

DLS technique

The nanoparticle size and its polydispersity index were measured in a 0.5% v/v PBS solution by using DLS techniques in a Zetasizer Nano ZS instrument (Malvern instruments).

CONCLUSIONS AND PERSPECTIVES

The main objective of the thesis was to design novel materials for personalized pharmaceutical applications from a global approach by tailoring the PLA-block copolymers architecture to generate defined nanostructures with specific structure-processing-properties relationships required for the end-product performance. Particularly, the PLA stereocomplex phase, which is typically obtained from the physical blend of the PLA enantiomers, exhibits promising physicochemical properties such as higher melting temperature, hydrolysis resistance and mechanical performance compared to the enantiomeric counterparts. However, detrimental phase separation occurs with high molecular weight derivatives that deteriorate their physicochemical properties performance. The synthesis of high molecular weight PLAstereoblock copolymer was long time desired to overcome the phase separation, however, sophisticated catalysts with high stereocontrol for both enantiomers are required that are extremely rare.

Herein, the synthesis of high molecular weight and high tacticity stereo-diblock copolymers of PLA (PLA-SBC) has successfully been achieved. The organometallic initiator employed for the PLA polymerization constitutes the first initiator able to generate the stereodiblock copolymers without the use of a co-initiator, which represents an advantage for scalability to industrial production. The polymerization follows a coordination-insertion mechanism and the initiator features a living character with a first-order constant rate. In addition, a similar rate constant for both PLA enantiomers was found, which enables the synthesis of stereo-block copolymers in a relatively short time, as well as the racemic-PLA counterpart. The tacticity of the block-copolymers can be tailored by changing the time at which the second monomer is added to the reaction, although the decrease of the enantiomeric purity of each monomer might be also a solution to tailor the tacticity.

Subsequently, the characterization of the PLA-SBC mechanism of crystallization was studied from the melt, under different thermal conditions, at different length-scales by a multitechnique approach. Importantly, the typical phase separation occurring in the enantiomeric blend of PLA is avoided in the high molecular weight PLA-SBC derivatives even at high cooling rates such as 50°C/min, which increase the suitability of the PLA-SBC for its processing under industrial conditions. The temperature of degradation, melting and crystallization of the PLA-SBC increases as the molecular weight increases. Moreover, the decrease of tacticity of the PLA-SBC decreases the melting temperature from 210°C to 200°C which might be ascribed to the formation of a defective SC crystal phase. The crystallinity degree and the crystallization rate of the high molecular weight PLA-SBC are 10% higher and 3-times higher, respectively, than the SC-PLA formed from the enantiomeric blend. Furthermore, the addition of OXA-2 as the nucleating agent can further increase the crystallization rate enabling the crystallization at high cooling rates such as 400°C/min. Interestingly, the OXA-2 changes also

A2.2 Polymerization mechanism

A2.3.2 Synthesis of L2

The polymer L2 was synthesized following the procedure described for the L1 polymer, namely the homopolymer synthesis. Two schlenk tubes were charged in the glovebox with 10.4 mg and 360 mg of catalyst and L-LA respectively, and 2 and 3 mL of solvent, respectively. The oil bath was pre-heated at 90°C rather than 70°C. A white powder was obtained after 60 min (270mg, 75%).

A2.3.7 Synthesis of Rac1

The polymer Rac1 was synthesized following the procedure described for the L1 polymer, namely the homopolymer synthesis. Racemic-LA (L-LA and D-LA) was used rather the L-LA. Two schlenk tubes were charged in the glovebox with 10.4 mg and 1.8 g of catalyst and Rac-LA and 5 and 15 mL of solvent, respectively. A transparent and rubbery product was obtained after 120 min (650 mg, 36 %). 1 H NMR (500 MHz, CDCl3, 297 K): δ 5.25-5.17

A2.4.2 Synthesis of L275:D275

The block-copolymer L275:D275 was synthesized following the one-pot sequential addition. Three schlenk tubes were charged in the glovebox with 10.4 mg, 1,080 g and 1,080 g of catalyst, L-LA and D-LA, respectively. 5 ml, 10 ml and 10 ml were added to the schlenk flasks respectively. The second monomer solution (D-LA) was added after 30 min of L-LA polymerization. Methanol was used to terminate the reaction and precipitate the polymer synthesized after 60 min. The obtained polymer was collected by filtration and dried at room temperature exposed to vacuum over 24h, obtaining a white powder (1,826 mg, 83%). 1 H NMR (CDCl3, 297 K): δ 5,15-5,10 (q, 1H, CH(PLA)), 5,02-4,98 (q, 1H, CH(L-LA)), 1,66-1,64 (d, 3H, CH3(L-LA)), 1,56-1,54 (d, 3H, CH3(PLA)). 13 C NMR (CDCl3, 297 K): δ 169,56 (CO), 68,95 (CH), 16,60 (CH3). GPC (CHCl3): Mw=79.320, Mn=35.253, Ð=2,25.

A2.4.2 Synthesis of L330:D330

The block-copolymer L330:D330 was synthesized following the one-pot sequential addition. Three schlenk tubes were charged in the glovebox with 10.4 mg, 1,080 g and 1,080 mg of catalyst, L-LA and D-LA, respectively.

A2.4.2 Synthesis of L500:D500

The block-copolymer L500:D500 was synthesized following the one-pot sequential addition. Three schlenk tubes were charged in the glovebox with 10.4 mg, 1,8 g and 1,8 mg of catalyst, L-LA and D-LA, respectively. 5 ml, 15ml and 15ml were added to the schlenk flasks respectively. The second monomer solution (D-LA) was added after 60 min of L-LA polymerization. Methanol was used to terminate the reaction and precipitate the polymer synthesized after 110 min. The obtained polymer was collected by filtration and dried at room temperature exposed to vacuum over 24h, obtaining a white powder (2,7 g, 75%).

A2.4.2 Synthesis of L520:D520

The block-copolymer L520:D520 was synthesized following the one-pot sequential addition. Three schlenk tubes were charged in the glovebox with 10.4 mg, 2,2 g and 2,2 mg of catalyst, L-LA and D-LA, respectively. 5 ml, 15ml and 15ml were added to the schlenk flasks respectively. The second monomer solution (D-LA) was added after 60 min of L-LA polymerization. Methanol was used to terminate the reaction and precipitate the polymer synthesized after 120 min. The obtained polymer was collected by filtration and dried at room temperature exposed to vacuum over 24h, obtaining a white powder (2,880 g, 80%). 1 H NMR (500 MHz, CDCl3, 297 K): δ 5,16-5,10 (q, 1H, CH(PLA)), 5,02-4,99 (q, 1H, CH(L-LA)),