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Synthesis and rheology of poly(lactide)s and their lignin composites Chile, Love-Ese 2017

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SYNTHESIS AND RHEOLOGY OF POLY(LACTIDE)S AND THEIR LIGNIN COMPOSITES  by  Love-Ese Chile  B.Sc. (Hons), The University of Auckland, 2011  A THESIS SUBMITTED IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF  DOCTOR OF PHILOSOPHY in THE FACULTY OF GRADUATE AND POSTDOCTORAL STUDIES (Chemistry)  THE UNIVERSITY OF BRITISH COLUMBIA (Vancouver)  August 2017  © Love-Ese Chile, 2017 ii  Abstract Synthetic plastics were first introduced 180 years ago, but the materials we have produced are likely to persist on our earth for thousands of years. Global shifts in thinking have urged researchers to focus their attention on bio-derived and biodegradable polymers. One such polymer is poly(lactic acid) (PLA). Despite its environmental benefits, PLA has several material weaknesses which hinder it’s use as a replacement for commodity plastics.   Highly active and selective indium catalysts for the ring-opening polymerization of lactide isomers have recently been developed by the Mehrkhodavandi group. By utilizing these catalysts, modification of tacticity and end-group functionality of PLAs are possible, permitting exploration into the effect of these modifications on chain interactions in PLA. The thermal and rheological behaviours of PLAs with different microstructures were compared. The molecular weight between entanglements was greatest for the syndiotactically enriched PLAs, giving rise to the lowest zero-shear viscosity. In addition, hetero- and isotactically enriched PLA had higher flow activation energies than syndiotactic variants, implying the inclusion of transient aggregate regions within these polymers due to enhanced L- and D-interactions.   A series of aryl-capped PLAs were synthesized by living ring-opening polymerization with a chain transfer agent using a previously reported dinuclear indium catalyst, [(NNO)InCl]2(μ-Cl)(μ-OEt) (A). Thermal, rheological and mechanical techniques were employed to understand the extent and strength of association caused by arylated chain ends. It is shown that the end-group has a greater effect on the properties of low molecular weight PLAs due to the larger number density of aryl end groups; significant interactions can be induced under oscillatory shear conditions in the low frequency flow regime (terminal zone). iii   The lignocellulosic biorefinery industry has been expanding in recent years and now provides researchers access to a range of bio-based composite materials through blending and copolymerization. Lignin-graft-PLA copolymers were synthesized via different routes and the PLA products were analyzed. Polymers were found to have cyclic structures at low lignin loading and star-like structures at higher lignin loading. Rheological studies were undertaken to derive useful structure-property relationships and optimize material properties. iv  Lay Summary  Synthetic plastics were first introduced 180 years ago, but the materials we have produced are likely to persist on our earth for thousands of years. Global shifts in thinking have urged researchers to focus their attention on bio-derived and biodegradable polymers. One such polymer is poly(lactic acid) (PLA). This multidisciplinary thesis links organometallic catalysis with polymer technology. The aim is to take a broader view of a research question and use what we find to create efficient catalytic systems where modifications to PLA can be easily made, and will impart useful properties in a predictable way, allowing PLA to be used wider array of consumer applications. v  Preface The work in Chapters 2 and 3 is based on a manuscript published in the journal Macromolecules.1 All experiments in these two chapters were conducted by myself. The manuscript was written by myself with significant contributions from Prof. Parisa Mehrkhodavandi and Prof. Savvas G. Hatzikiriakos. The synthesis and solution characterization of metal complexes [(NNO)InCl]2(μ-Cl)(μ-OPhOMe) (1), [(NNO)InCl]2(μ-Cl)(μ-OPhBr) (4), and [(NNO)InCl]2(μ-Cl)(μ-OPhNO2) (5) using one methodology in Chapter 4, was carried out by undergraduate researcher Alan Wong who was supervised by me. I completed the synthesis and characterization of the remaining metal complexes in this chapter, as well as performing all the polymerization studies. Crystals suitable for single-crystal x-ray diffraction were obtained by Alan Wong (for the three complexes previously mentioned) or myself, and the crystal structures presented in Chapter 4 were run and solved by Dr. Dinesh Aluthge, Dr. Brian Patrick and Tannaz Ebrahimi. This work has been published in Dalton Transactions.2 The manuscript was written by myself with significant contributions from Prof. Parisa Mehrkhodavandi and Prof. Savvas G. Hatzikiriakos. The work in Chapter 5 has been submitted for publication. All the experiments and data analysis in this chapter was completed by myself. The manuscript was written by myself with significant contributions by Prof. Parisa Mehrkhodavandi and Prof. Savvas G. Hatzikiriakos. All the experiments in Chapter 6 were conducted by myself. This work has not been published and the chapter was entirely written by myself. vi  Table of Contents  Abstract .......................................................................................................................................... ii Lay Summary ............................................................................................................................... iv Preface .............................................................................................................................................v Table of Contents ......................................................................................................................... vi List of Tables ................................................................................................................................ xi List of Figures ............................................................................................................................. xiii List of Schemes ......................................................................................................................... xxix List of Symbols ......................................................................................................................... xxxi List of Abbreviations .............................................................................................................. xxxii Acknowledgements ................................................................................................................ xxxiv Dedication ............................................................................................................................... xxxvi Chapter 1: General introduction. .................................................................................................1 1.1 Biopolymers .....................................................................................................................1 1.1.1 Poly(lactide): Synthesis, characterization, biodegradation and application. ...............3 1.1.2 Lignin: Structure, isolation, characterization and biodegradation. ............................13 1.2 Polymer rheology ...........................................................................................................22 1.2.1 Polymer processing technologies. ..............................................................................22 1.2.2 Polymer viscoelasticity. .............................................................................................25 1.2.3 Polymer viscosity. ......................................................................................................27 1.2.4 Small amplitude oscillatory shear (SAOS) experiments. ..........................................29 1.2.5 Uniaxial extensional rheology. ..................................................................................32 vii  1.3 Scope of thesis ...............................................................................................................34 Chapter 2: Living ring-opening polymerization of meso-lactide using indium initiators. ....35 2.1 Introduction ....................................................................................................................35 2.1.1 Controlled tacticity via ring-opening polymerization of lactide monomers. .............35 2.1.2 Mechanisims for stereocontrol. ..................................................................................37 2.1.3 Determining poly(lactide) tacticity. ...........................................................................38 2.1.4 Catalysts for synthesizing PLA with controlled microstructure. ...............................43 2.2 Results and discussion ...................................................................................................47 2.2.1 Purification of meso-lactide. ......................................................................................47 2.2.2 Polymerization of meso-lactide with tetra- and tridentate indium alkoxide complexes. ....................................................................................................................................49 2.2.3 In situ studies. ............................................................................................................59 2.3 Conclusions ....................................................................................................................63 2.4 Experimental ..................................................................................................................64 Chapter 3: Comparison of the rheological properties of isotactic, syndiotactic, and heterotactic poly(lactide)s. ..........................................................................................................67 3.1 Introduction ....................................................................................................................67 3.1.1 Influence of polymer tacticity on chain packing and thermal properties ...................67 3.1.2 Polymer tacticity and its influence on rheological properties. ...................................69 3.1.3 Rheological properties of poly(lactide)s. ...................................................................70 3.2 Results and discussion ...................................................................................................71 3.2.1 Synthesis of PLAs with various microstructures. ......................................................71 3.2.2 Thermal study of stereoregular poly(lactide)s. ..........................................................75 viii  3.2.3 Solution viscosity. ......................................................................................................78 3.2.4 Linear viscoelasticity of PLAs with varying microstructures. ...................................79 3.2.5 Uniaxial extensional rheology ...................................................................................91 3.3 Conclusions ....................................................................................................................96 3.4 Experimental ..................................................................................................................98 Chapter 4: Aryl initiators for the living ring-opening polymerization of rac-lactide. ........103 4.1 Introduction ..................................................................................................................103 4.1.1 PLA synthesis via living ring-opening polymerization with a chain transfer agent. ..... ..................................................................................................................................103 4.1.2 Aryloxy initiators in lactide ring-opening polymerization. .....................................106 4.2 Results and discussion .................................................................................................107 4.2.1 Synthesis racemic phenoxy-bridged indium complexes 1-5. ..................................107 4.2.2 Solution structures of phenoxy-bridged indium complexes 1-5. .............................109 4.2.3 Solid-state structures of phenoxy-bridged indium complexes 1-5. .........................111 4.2.4 Living ring-opening polymerization (ROP) of lactide using catalysts 1–5. ............116 4.2.5 In situ polymerization studies. .................................................................................121 4.2.6 Immortal ring-opening polymerization (ROP) of lactide with catalyst A and aromatic alcohols. ...............................................................................................................................123 4.3 Conclusions ..................................................................................................................140 4.4 Experimental ................................................................................................................141 Chapter 5: Thermorheological investigation of aromatic interactions in aryl-capped polylactides. ................................................................................................................................151 5.1 Introduction ..................................................................................................................151 ix  5.1.1 Network structures in polymeric materials. .............................................................151 5.1.2 Non-ionic secondary interactions to modify polymer properties. ...........................152 5.2 Results and discussion .................................................................................................155 5.2.1 Synthesis of aryl-capped PLAs via immortal ring-opening polymerization. ..........155 5.2.2 Thermal studies of aryl-capped PLAs......................................................................163 5.2.3 Solution viscosity of aryl-capped PLAs. .................................................................166 5.2.4 Melt Rheology - Linear viscoelasticity. ...................................................................168 5.2.5 Uniaxial Extensional Rheology. ..............................................................................179 5.2.6 Mechanical properties. .............................................................................................181 5.3 Conclusions ..................................................................................................................184 5.4 Experimental ................................................................................................................185 Chapter 6: Synthesis and themorheological analysis of lignin-graft-poly(lactide) copolymers and their blends. .........................................................................................................................189 6.1 Introduction ..................................................................................................................189 6.1.1 Lignin-polymer materials.........................................................................................189 6.1.2 Lignin graft copolymers ...........................................................................................190 6.2 Results and discussion .................................................................................................194 6.2.1 Synthesis and characterization of lignin-graft-PLAs. ..............................................194 6.2.2 Effect of lignin source on copolymerization. ...........................................................203 6.2.3 Other synthetic routes to lignin-graft-PLAs. ...........................................................206 6.2.4 Characterization and comparison of polymers generated from the three synthetic methodologies. .....................................................................................................................210 6.2.5 Determining lignin-graft-PLA copolymer topology. ...............................................215 x  6.2.6 Rheological analysis of lignin-graft-poly(lactide)s. ................................................226 6.2.7 Lignin-graft-PLLA/PLLA blends. ...........................................................................232 6.3 Conclusions ..................................................................................................................236 6.4 Experimental ................................................................................................................237 Chapter 7: Conclusion and future directions ..........................................................................243 Bibliography ...............................................................................................................................248 Appendices ..................................................................................................................................256 Appendix A ..............................................................................................................................256 A.1 GC-MS analysis for meso-lactide purification ........................................................256 Appendix B ..............................................................................................................................257 B.1 Isothermal time sweep experiments .........................................................................257 B.2 Van Gurp-Palmen plots............................................................................................259 B.3 Uniaxial extension experimental plots .....................................................................260 Appendix C ..............................................................................................................................264 C.1 Characterization of complexes 1-6 in the solution state ..........................................264 C.2 Characterization of complexes 2-6 in the solid state ...............................................286 C.3 In situ living ring-opening polymerization data with catalysts 1-4. ........................288 Appendix D ..............................................................................................................................291 D.1 Master curves for aryl capped poly(lactide)s ...........................................................291 D.2 Uniaxial extension experimental plots .....................................................................296 Appendix E ..............................................................................................................................298 E.1 Calculation of dn/dc values ......................................................................................298  xi  List of Tables Table 1.1 Tetrad probabilities for lactide isomers, based on Bernoullian statistics.a, b ...................8 Table 1.2 Comparison of mechanical and thermal properties for common consumer plastics.29,31,55  ................................................................................................................................................12 Table 2.1 Polymerization of lactide with dinuclear indium complexes (±)-A and (RR/RR)-A. ...52 Table 2.2 Polymerization of meso-lactide with dinuclear indium complexes (±)-B and (RR/RR)-B.  ................................................................................................................................................57 Table 2.3 Rates of polymerization of lactide isomers with various catalysts. ..............................61 Table 3.1 Trends in thermal transitions for different microstructures of three common polyolefins (all values in °C). ...........................................................................................................................69 Table 3.2 Polymerization data for various microstructured PLAs used for themorheological experiments. ...................................................................................................................................74 Table 3.3 Thermal properties of PLAs of varying stereoregularity. .............................................77 Table 3.4 Comparison of rheological properties of PLA, PS and PP in this work and obtained from the literature. ..................................................................................................................................88 Table 3.5 Average relaxation times for various polymers studied calculated from linear viscoelastic experiments. ...............................................................................................................94 Table 4.1 Selected solid-state structural data for indium catalysts (A, 2-6). ..............................114 Table 4.2 Living ring-opening polymerization data of rac-lactide with complexes 1-5. ...........118 Table 4.3 Propagation rates for the ROP of rac-LA with complexes A and 1-4. .......................123 Table 4.4 Molecular weight data for immortal ROP of rac-LA with various aryl chain transfer agents. ..........................................................................................................................................130 xii  Table 5.1 Molecular weight data for immortal ring-opening polymerization of rac-LA with various aryl chain transfer agents. ...............................................................................................159 Table 5.2 Thermal properties for various end-capped PLAs in this study (all values in °C). ....164 Table 6.1 Polymerization data from lignin-graft-poly(lactide)s formed via ring-opening polymerization using InCl3 and NEt3. ..........................................................................................196 Table 6.2 Polymerization data from lignin-graft-poly(lactide)s formed using two commercially available lignin sources. ...............................................................................................................205 Table 6.3 Hydroxy group content for Indulin AT kraft lignin (IAK) and alkali kraft lignin (AK). .  ..............................................................................................................................................205 Table 6.4 Polymerization data from lignin-graft-poly(lactide)s formed via ring-opening polymerization using TBD. ..........................................................................................................207 Table 6.5 Polymerization data from lignin-graft-poly(lactide)s formed via a graft-to strategy. 208 Table 6.6 Polymerization data for control reactions in the absence of lignin. ............................217 Table 6.7 Polymerization data for linear PLA, cyclic PLA and star lignin-graft-PLAs used in rheological analyses. ....................................................................................................................224 Table 6.8 Results for graft-from copolymerizations of 90% L-lactide with InCl3/NEt3. ............232 Table 6.9 Lignin-graft-PLLA/PLLA blend properties. ...............................................................234  xiii  List of Figures Figure 1.1 Four categories of biodegradable polymers found in the literature.17 ...........................2 Figure 1.2 Production cycle for poly(lactide). ................................................................................4 Figure 1.3 Microstructures observed for poly(lactide). ..................................................................5 Figure 1.4 Coordination-insertion mechanism for the ROP of lactide by metal catalysts. .............7 Figure 1.5 Schematic diagrams of PLA tetrad resonances corresponding to the methine region in (a) homonuclear decoupled 1H NMR of PLA from rac-lactide; (b) 13C{1H} NMR of PLA from rac-lactide; (c) homonuclear decoupled 1H NMR of PLA from meso-lactide; (d) 13C{1H} NMR meso-lactide.48 ..................................................................................................................................9 Figure 1.6 Three main monolignols: para-coumaryl alcohol, coniferyl alcohol and sinapyl alcohol, which become para-hydroxyphenyl (H), guaiacyl (G) and syringyl (S) units within the lignin framework.64 ........................................................................................................................14 Figure 1.7 Enzymatic radical pathways for coniferyl alcohol.65...................................................15 Figure 1.8 Schematic representation of lignin after radical coupling of monolignols.66 ..............16 Figure 1.9 31P NMR spectrum (CHCl3/pyridine, 25 °C, 121 MHz) for phosphitylated lignin, derivatized with 2-chloro-4,4,5,5-tetramethyl-1,3,2-dioxaphospholane (TMDP). IS = internal standard (cyclohexanol), 145.2 ppm. Reference = phosphitylated water, 132.2 ppm. Alkyl-OH signals found 145.5 – 150 ppm; aryl-OH signals found = 138 – 144 ppm, carboxylic acid-OH signals found 134 – 136 ppm. ........................................................................................................17 Figure 1.10 Biomass extraction processes to produce the four classes of technical lignin, kraft, lignosulfonate, organosolv and soda.64,69 .......................................................................................18 Figure 1.11 Modulus vs. temperature for an amorphous polymer (black solid line) and a semi-crystalline polymer (blue dashed line). ..........................................................................................23 xiv  Figure 1.12 Schematic of a single screw-extruder. .......................................................................24 Figure 1.13 Schematic of a blow mold. ........................................................................................24 Figure 1.14 Plot of the linear viscoelastic moduli, G’ and G” of a thermoplastic polymer. At temperatures below 100 °C, the storage modulus dominates; thus, the material behaves as an elastic solid. Above 100 °C the material behaves as a viscous fluid. .......................................................26 Figure 1.15 Plot of the linear viscoelastic moduli, G’ () and G” () of a hypothetical thermoplastic polymer. At deformation frequencies lower than 1 rad s-1 the loss modulus dominates thus the material behaves as a viscous liquid. Above 10 rad s-1 the material behaves as an elastic solid. ...............................................................................................................................27 Figure 1.16 Schematic of parallel plate rheometer. ......................................................................28 Figure 1.17 Viscosity curves for ideally viscous, shear thinning and shear thickening flow behaviour........................................................................................................................................29 Figure 1.18 Vector diagram showing the relationship between the complex modulus |G*|, the loss and storage moduli, G” and G’. .....................................................................................................31 Figure 1.19 Sentmanat Extensional Rheometer Universal Testing Platform (SER).95 .................33 Figure 1.20 Plot of the elongational viscosity E+ vs. time for a hypothetical polymer. At low strain rates, no strain hardening is observed. At strain rates greater than 1 s-1, strain hardening is present. ...........................................................................................................................................33 Figure 2.1 Stereoisomers of lactide and the possible microstructures of poly(lactic acid). .........37 Figure 2.2 Examples of tetrad sequences from the ring-opening polymerization of lactide. a) isotactic PLA b) heterotactic PLA and c) syndiotactic PLA. i indicates an iso linkage where two bound lactyl units have the same arrangement s indicates a syndio linkage where the lactyl units have opposite configuration. ..........................................................................................................39 xv  Figure 2.3 1H[1H} NMR spectra (CDCl3, 25 °C, 600 MHz) of PLA samples resulting from ring-opening polymerization of different lactide isomers. The chemical shifts for the methine proton (shown in purple) depend on the tetrads (labelled in black) in which it is found. .........................40 Figure 2.4 A catalyst’s preference for a certain enantiomeric site is determined by the rate at which it ring-opens at one site over the other (a vs. b), quantified by the α-value.106 ..............................41 Figure 2.5 Possible insertion errors from ROP of rac-lactide via chain-end control (CEC) and enantiomorphic site control (ESC) (where i denotes isotactic enchainment, (relative stereochemistry of adjacent lactyl units are the same) and s denotes syndiotactic enchainment (relative stereochemistry of adjacent lactyl units are different)). ..................................................43 Figure 2.6 Possible insertion errors from ROP of meso-lactide via chain-end control (CEC) and enantiomorphic site control (ESC) (where i denotes isotactic enchainment, (relative stereochemistry of adjacent lactyl units are the same) and s denotes syndiotactic enchainment (relative stereochemistry of adjacent lactyl units are different)). ..................................................43 Figure 2.7 Catalysts reported for the ROP of meso-lactide. I–Coates 1999103 II–Coates 2001115 III– Okuda 2010116 and 2012 117 IV–Okuda 2011118 V–Carpentier 2006119 VI–Hendrick 2006.104   ................................................................................................................................................44 Figure 2.8 1H NMR spectra (CDCl3, 25 °C, 300 MHz) of meso-lactide (blue star) and diminishing rac-lactide impurity (red star) a) after drying b) after three recrystallizations from iPrOH c) after five recrystallizations from iPrOH. ................................................................................................48 Figure 2.9 Plot of observed PLA Mn () and Ð (●) as functions of added meso-LA for (±)- (RR/RR)-A (top) and for (±)- (RR/RR)-B (bottom). (Mn = number averaged molecular weight, Ð = dispersity index). The black line indicates theoretical Mn values based on the [LA]:[initiator] ratio xvi  at 100% conversion. All reactions were carried out at room temperature for 16 h (A) or 4 h (B) in DCM and polymer samples were obtained at >90% conversion. ..................................................51 Figure 2.10 1H{1H} NMR spectra (CDCl3, 25 °C, 600 MHz) of polymers generated from a) (±)-A and b) (RR/RR)-A. ......................................................................................................................53 Figure 2.11 Inverse-gated 13C{1H} NMR spectrum (CDCl3, 25 °C, 100 MHz) for a polymer generated from (RR/RR)-A. ...........................................................................................................55 Figure 2.12 1H{1H} NMR spectra (CDCl3, 25 °C, 600 MHz) of polymers generated from a) (±)-B and b) (RR/RR)-B. ......................................................................................................................58 Figure 2.13 Inverse-gated 13C{1H} NMR spectrum (CDCl3, 25 °C, 100 MHz) for a polymer generated from (±)-B. ....................................................................................................................59 Figure 2.14 Plot of ln[LA] versus time for polymerization of meso-LA catalysed with a) left, (±)-A () and (RR/RR)-A () and b) right, (±)-B () and (RR/RR)-B().  Reactions were carried out in an NMR tube at 25 °C and followed to 90% conversion. 1,3,5-trimethoxybenzene (TMB) was used as internal standard. All reactions were carried out with 200 eq. of LA in CDCl3 at 25 °C and followed to 90% conversion by 1H NMR spectroscopy [A] = 0.0023 M, [LA] = 0.46 M. [B] = 0.0018 M, [LA] = 0.36 M. The value of kobs was determined from the slope of ln[LA] vs. time, averaged from at least three experiments.......................................................................................61 Figure 2.15 Mechanism for ring-opening of L- or D-LA by (RR/RR)-A.129 ................................62 Figure 3.1 Chain packing conformations reported for syndiotactic polyolefins...........................68 Figure 3.2 (-)-Salen binathphylamine alkoxide complex (I) reported to be highly syndioselective towards meso-LA by Coates et al. in 1999.103 ...............................................................................72 xvii  Figure 3.3 1H {1H} NMR spectrum of polymer produced from reaction of (-)-salen binaphthylamine alkoxide complex (I) with meso-lactide. Using Equation 2.2, the polymer tacticity was calculated to be Ps = 0.77. .........................................................................................73 Figure 3.4 TGA heating traces for the various microstructured PLAs. Thermogravimetric analysis was performed on approximately 20 mg of material. Samples were heated to 500 °C at a rate of 20 °C/min to determine the degradation onset temperature (temperature at which there is 5% weight loss), Tonset. .....................................................................................................................................75 Figure 3.5 DSC traces for various microstructured PLAs studied. Samples heated to 170 °C at 10 °C/min and cooled to 25 °C at 5 °C/min under a nitrogen atmosphere to reduce sample degradation. Glass transition and melting temperatures calculated from second heating scans. ca. 2 mg of sample used. ...................................................................................................................................77 Figure 3.6 Log-log plot of intrinsic viscosity of various types of PLAs as a function of the molecular weight. ...........................................................................................................................78 Figure 3.7 Characteristic hydrodynamic radius (Rh) as a function of molecular weight and tacticity. The radii of the iso- PLA are the highest while those of syndio-PLA are the lowest, in agreement with the intrinsic viscosity data plotted in Figure 3.6. .................................................79 Figure 3.8 Complex modulus, |G*| vs. time for stereoregular PLAs. ...........................................81 Figure 3.9 Master curve of the linear viscoelastic moduli, G’ and G” and complex viscosity, |*|, (Tref = 150 °C) for (a) 182-het-96 (Table 3.2, entry 6) and (b) 148-syn-72 (Table 3.2, entry 10). 83 Figure 3.10 Master curves of the linear viscoelastic moduli for het-96 polymers (Tref = 150 °C). (a) Loss modulus vs. angular frequency (b) storage modulus vs. angular frequency (c) complex viscosity vs. angular frequency. .....................................................................................................84 xviii  Figure 3.11 Master curves of the linear viscoelastic moduli for syn-72 polymers (Tref = 150 °C). (a) Loss modulus vs. angular frequency (b) storage modulus vs. angular frequency (c) complex viscosity vs. angular frequency. .....................................................................................................85 Figure 3.12 Van Gurp-Palmen plots of the various microstructured PLAs. .................................87 Figure 3.13 The horizontal shift factors, aT for PLAs listed in Table 3.4 at the reference temperature of 150 °C. Lines represents single fitting of the data used to calculate flow activation energy, Ea,flow. .................................................................................................................................87 Figure 3.14 Complex viscosity vs. angular frequency master curves for the different microstructure PLAs at approximately a fixed molecular weight value (Tref = 150 °C). ..............89 Figure 3.15 Zero-shear viscosity, 0, versus weight averaged molecular weight for PLAs with different microstructures. ...............................................................................................................90 Figure 3.16 Illustration of micro-crystallites forming transient cross-links between polymer chains. ............................................................................................................................................92 Figure 3.17 Elongational viscosity as a function of time at Hencky strain rate of 10 s-1 for differently microstructured poly(lactide)s at (a) 70 °C, (b) 90 °C and (c) 110 °C. Uniaxial elongation experiments on iso-72 were unable to be performed at 70 °C. ....................................95 Figure 3.18 Elongational viscosity as a function of time at Hencky strain rates from 0.01 to 10 s-1 for iso-enriched PLA (Table 3, entry 4).  It is noted that strain-hardening is present at all Hencky strain rates (not the case for all other PLA types) indicating the stronger interaction of L- and D- regions. ...........................................................................................................................................96 Figure 4.1 Coordination-insertion mechanism for the living ROP of lactide with dinuclear indium catalyst (A). ..................................................................................................................................104 xix  Figure 4.2 Coordination-insertion mechanism for the immortal ROP of lactide with dinuclear indium catalyst (A). .....................................................................................................................106 Figure 4.3 1H NMR (CDCl3, 25 °C, 400 MHz) spectrum of crude reaction mixture containing (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhOMe) (1) (blue stars). Extra peaks in the aromatic region (6-8 ppm) are assigned to the bis-phenoxy complex (red stars). ........................................................................109 Figure 4.4 1H NMR spectra (CDCl3, 25 °C, 400 MHz) of diastereotopic ligand NH-CH2-Ar protons for (a) [(NNO)InCl]2(μ-OEt)(μ-Cl) (A), (b) [(NNO)InCl]2(μ-OPhOMe)(μ-Cl) (1), (c) [(NNO)InCl]2(μ-OPhMe)(μ-Cl) (2), (d) [(NNO)InCl]2(μ- OPhH)(μ-Cl) (3), (e) [(NNO)InCl]2(μ-OPhBr)(μ-Cl) (4), and (f) [(NNO)InCl]2(μ-OPhNO2)(μ-Cl) (5). ....................................................110 Figure 4.5 1H NMR spectra (CDCl3, 25 °C, 400 MHz) bridging phenolic protons for (a) [(NNO)InCl]2)(μ-Cl)(μ-OEt) (A), (b) [(NNO)InCl]2)(μ-Cl)(μ-OPhOMe) (1), (c) [(NNO)InCl]2)(μ-Cl)(μ-OPhMe) (2), (d) [(NNO)InCl]2)(μ-Cl)(μ- OPhH) (3), (e) [(NNO)InCl]2)(μ-Cl)(μ-OPhBr) (4), and (f) [(NNO)InCl]2)(μ-Cl)(μ-OPhNO2) (5). ...............................................................................111 Figure 4.6 Molecular structure of [(NNO)InCl]2(μ-Cl)(μ-OPhMe) (2) depicted with ellipsoids at 50% probability (H atoms and all solvent molecules omitted for clarity). ..................................112 Figure 4.7 Molecular structure of [(NNO)InCl]2(μ-Cl)(μ-OPhH) (3) depicted with ellipsoids at 50% probability (H atoms and all solvent molecules omitted for clarity). ..................................112 Figure 4.8 Molecular structure of [(NNO)InCl]2(μ-Cl)(μ-OPhBr) (4) depicted with ellipsoids at 50% probability (H atoms and all solvent molecules omitted for clarity). ..................................113 Figure 4.9 Molecular structure of [(NNO)InCl]2(μ-Cl)(μ-OPhNO2) (5) depicted with ellipsoids at 50% probability (H atoms and all solvent molecules omitted for clarity). ..................................113 Figure 4.10 Plots of observed PLA Mn (closed symbols) and dispersity (open symbols) as functions of added rac-LA for a) [(NNO)InCl]2(μ-OPhOMe)(μ-Cl) (1), b) [(NNO)InCl]2(μ-xx  OPhMe)(μ-Cl) (2), c) [(NNO)InCl]2(μ-OPhH)(μ-Cl) (3), d) [(NNO)InCl]2(μ-OPhBr)(μ-Cl) (4), and e) [(NNO)InCl]2(μ-OPhNO2)(μ-Cl) (5). (Mn = number averaged molecular weight).  The line represents theoretical Mn values based on the monomer:initiator ratio at 100% conversion.  All reactions were carried out room temperature in DCM. ...............................................................117 Figure 4.11 1H NMR spectrum (CDCl3, 25 °C, 400 MHz) of free phenol (red spectrum) overlaid with 1H NMR spectrum (CDCl3, 25 °C, 400 MHz) of [(NNO)InCl]2(μ-Cl)(μ-OPhH) (3) (green) and 1H NMR spectrum (CDCl3, 25 °C, 400 MHz) of the polymerization reaction mixture (blue, Table 4.2. entry 10). .....................................................................................................................119 Figure 4.12 MALDI-ToF spectrum from polymerization of [LA]:[3] = 49:1 (Table 4.2, entry 10). An = [144.13LA]n + 94.11HOPh + 23Na+ ....................................................................................120 Figure 4.13 1H{1H} NMR (CDCl3, 25 °C) spectrum of the methylene region of PLAs generated from a) (±)-[(NNO)InCl]2(μ-Cl) (μ-OEt) (A), b) (±)-[(NNO)InCl]2(μ-Cl) (μ-OPhOMe) (1) (±)-[(NNO)InCl]2(μ-Cl) (μ-OPh) (3). ................................................................................................121 Figure 4.14 Plot of ln[LA] versus time for polymerization of rac-LA catalyzed with 1 () 2 (), 3 () and 4 ().  Reactions were carried out in an NMR tube at 25 °C. 1,3,5-trimethoxybenzene (TMB) was used as internal standard. All reactions were carried out with 50 eq. of LA in CD2Cl2 at 25 °C and followed by 1H NMR spectroscopy. [1] = 0.0039 M, [LA] = 0.19 M. [2] = 0.0043 M, [LA] = 0.21 M. [3] = 0.0037 M, [LA] = 0.22 M. [4] = 0.0041 M, [LA] = 0.22 M. The value of kobs was determined from the slope of ln[LA] vs. time, averaged from at least three experiments. ..122 Figure 4.15 1H NMR (CDCl3, 25 °C) spectrum of the alkoxy-phenoxy exchange reaction with complex A and para-methoxy phenol. The loss of the bridging ethoxy peaks (denoted by red stars) and the emergence of bridging phenoxy signals (denoted by green stars) were monitored and used to calculate the rate of exchange, kexchange. ...................................................................................125 xxi  Figure 4.16 Plot of ln[A] versus time for exchange between catalyst A, HOC6H4OMe () and HOC6H4Br () and HOC6H5 (). Reactions carried out in an NMR tube (CDCl3, 25 °C). 1,3,5-trimethoxybenzene (TMB) was used as internal standard. [A] = 0.0046 M, [HOC6H4OMe] = 0.0050 M, [HOC6H4Br] = 0.0059 M, [HOC6H5] = 0.00411 M. The value of kexchange was determined from the slope of ln[A] vs. time, averaged from three experiments. ...............................................................................................................................126 Figure 4.17 GPC traces with respect to conversion for the immortal ROP of lactide with A and phenol as the CTA ([LA]:[A]:[phenol] = 10500:1:20). ...............................................................127 Figure 4.18 Plot of molecular weight and dispersity with respect to conversion for the immortal ROP of lactide with A with phenol ([LA]:[A]:[phenol] = 10500:1:20). .....................................127 Figure 4.19 Plots of observed PLA Mn (closed symbols) and dispersity (open symbols) as functions of [LA]:([A]+[CTA]) for (a) CTA = ethanol (PLA-Et),  (b) CTA = phenol (PLA-Ph), (c) CTA = 1,5-naphthalenediol (PLA-1,5-Nap) ................................................................................129 Figure 4.20 1H NMR (CDCl3, 25 °C, 400 MHz) spectrum of PLA isolated from polymerization with [LA]:[A]:[diClPhOH] ratios of 237:1:2 (Table 4.4, entry 2). ..............................................131 Figure 4.21 1H NMR (CDCl3, 25 °C) spectrum of PLA isolated from polymerization with [LA]:[A]:[1,8-NapOH] ratios of 78:1:2 (Table 4.4, entry 5). ......................................................132 Figure 4.22 1H NMR (CDCl3, 25 °C) spectrum of PLA isolated from polymerization with [LA]:[A]:[ PhOH] ratios of 1050:1:20 (Table 4.4, entry 9). ........................................................133 Figure 4.23 1H NMR (CDCl3, 25 °C) spectrum of unconverted lactide isolated from polymerization with [LA]:[A]:[biPhenOH] ratios of 583:1:10 (Table 4.4, entry 3). ..................134 Figure 4.24 MALDI-ToF mass spectrum for a selected PLA-1,5-Nap sample. An = [72.07 LA]n + 160.17 1,5-Nap  + 23Na+ ..............................................................................................................134 xxii  Figure 4.25 1H NMR (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(2,μ-diClPh) (6) ...................................................................................................................................136 Figure 4.26 Molecular structure of [(NNO)InCl]2(μ-Cl)(2,μ-diClPh) (6) depicted with ellipsoids at 50% probability (H atoms and all solvent molecules omitted for clarity). ...137 Figure 4.27 1H (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(2,μ-biPh). ....138 Figure 4.28 1H NMR (CDCl3, 25 °C, 300 MHz) spectrum of polymerization [LA]:[6] = 110 in THF over 120 hours. Red star indicates the methylene protons on growing polymer. ...............139 Figure 4.29 1H NMR (CDCl3, 25 °C, 300 MHz) spectrum of polymerization [LA]:[6] = 110 in DCM over 120 hours. ..................................................................................................................140 Figure 5.1 Self complimentary hydrogen-bonding unit, ureidopyrimidinone (UPy). ................153 Figure 5.2 Bis(diimide) capable of forming complimentary - interactions with -electron rich species such as pyrene.206,253 ........................................................................................................154 Figure 5.3 Plots of observed PLA number averaged molecular weight, Mn, (closed symbols) and dispersity, Đ, (open symbols) as functions of [LA]:([A]+[CTA]) for (a) CTA = ethanol (PLA-EtOH), (b) CTA = phenol (PLA-Ph), (c) CTA = naphthol (PLA-Nap), (d) CTA = hydroxymethylpyrene (PLA-MePyr) and (e) hydroxymethylanthracene (PLA-MeAnth). The black lines represent theoretical Mn values based on the [LA]:([A]+[CTA]) ratio at 100% conversion.  All reactions were carried out room temperature in DCM. .........................................................157 Figure 5.4 1H NMR (CDCl3, 25 °C, 300 MHz) spectrum of PLA isolated from polymerization of [LA]:[A]:[NapOH] ratios of 1050:1:20 (Table 5.1, entry 5). ......................................................160 Figure 5.5 1H NMR (CDCl3, 25 °C, 300 MHz) spectrum of PLA isolated from polymerization of [LA]:[A]:[PyrMeOH] ratios of 1050:1:20 (Table 5.1, entry 10). ................................................160 xxiii  Figure 5.6 1H NMR (CDCl3, 25 °C, 300 MHz) spectrum of PLA isolated from polymerization of [LA]:[A]:[AnthMeOH] ratios of 1050:1:20 (Table 5.1, entry 14)...............................................161 Figure 5.7 Solution UV-Vis spectra for phenol (purple line), PLA-Ph9 (blue line), PLA-Ph38 (red line) and PLA-Ph121 (black line). Polymer solutions were in CHCl3 with a concentration of 1×10-4 M.  ..............................................................................................................................................162 Figure 5.8 Solution UV-Vis spectra for naphthol (purple line), PLA-Nap11 (black line), PLA-Nap36 (red line) and PLA-Nap101 (blue line). Polymer solutions were in CHCl3 with a concentration of 1×10-4 M. .................................................................................................................................162 Figure 5.9 Solution UV-Vis spectra for PLA-MePyr17 (black line), PLA-MePyr39 (red line) and PLA-MePyr106 (blue line). Polymer solutions were in CHCl3 with a concentration of 1×10-5 M. ....  ..............................................................................................................................................163 Figure 5.10 Solution UV-Vis spectra for PLA-MeAnth11 (black line), PLA-MeAnth37 (red line) and PLA-MeAnth124 (blue line). Polymer solutions were in CHCl3 with a concentration of 1×10-4 M.  ..............................................................................................................................................163 Figure 5.11 TGA heating traces for the various aryl-capped PLAs. Thermogravimetric analysis was performed on approximately 15 mg of material. Samples were heated to 500 °C at a rate of 20 °C/min. .........................................................................................................................................165 Figure 5.12 DSC traces for the various end-capped PLAs studied. Samples heated to 250 °C at 10 °C/min and cooled to 25 °C at 5 °C/min under a nitrogen atmosphere to reduce sample degradation. Glass transition and melting temperatures calculated from second cooling scans. ca. 5 mg of sample used. Plots shifted vertically for clarity. ..........................................................................165 Figure 5.13 A log−log plot of intrinsic viscosity of various aryl-capped PLAs as a function of the molecular weight. .........................................................................................................................167 xxiv  Figure 5.14 Characteristic hydrodynamic radius (Rh) as a function of molecular weight. .........168 Figure 5.15 a) Normalized complex modulus, G*(t)/G*(t=0), vs. time at 180 °C for end-capped PLAs in this study. b) Normalized complex modulus, G*(t)/G*(t=0), vs. time at 70, 120 and 180 °C for PLA-Nap36. c) Normalized complex modulus, G*(t)/G*(t=0) vs. time at 3.14, 31.4 and 314 rad/s for PLA-Nap36. ....................................................................................................................170 Figure 5.16 Thermorheologcial characterization of PLA-MePyrene polymers. a) Master curve of the linear viscoelastic moduli, G’ and G” and |*| complex viscosity for PLA-MePyr72 at Tref = 150 °C. b) Master curve of the linear viscoelastic moduli, G’ and G” and |*| complex viscosity for PLA-MePyr135 at Tref = 150 °C. ..............................................................................................172 Figure 5.17 Master curve of the complex viscosity vs. angular frequency (Tref = 150 °C) for naphthol-capped PLAs (Nap-PLA) to show the molecular weight dependence on the upturn in complex viscosity.........................................................................................................................172 Figure 5.18 Successive frequency sweep experiments for a) PLA-Nap36 and b) PLA-Et34 polymers. T = 120 °C,  = 0.063-630 rad/s, strain = 2%. ...........................................................173 Figure 5.19 Successive time sweep experiments for PLA-Nap36. T = 120 °C,  = 3.14 rad/s, strain = 2%. ............................................................................................................................................174 Figure 5.20 a) Storage modulus vs. frequency, b) loss modulus vs. frequency for the isothermal frequency sweep experiments for PLA-MePyr72, to show the failure of the time-Temperature superposition principle. c) Horizontal shift factors, aT at 150 °C. d) Plot of G’ vs. G” at various temperatures for PLA-MePyr72 to check for possible thermal transitions. ..................................175 Figure 5.21 Master curve of the linear viscoelastic moduli, G′ and G″, and complex viscosity, |η*| (Tref = 150 °C), for a) PLA-Nap36 b) PLA-Nap63 polymers. ........................................................176 Figure 5.22 Index of association vs. molecular weight for aryl-capped PLAs in this study. .....178 xxv  Figure 5.23 Reptation time versus molecular weight for aryl-capped PLAs in this study. ........179 Figure 5.24 Comparison of tensile stress growth coefficients for selected aryl-capped PLAs (measure of elongational viscosity) at 90 °C as a function of time at Hencky strain rate of a) 10 s-1 and b) 1.0 s-1. .............................................................................................................................181 Figure 5.25 Tensile strength for aryl-capped PLAs in this study. Low molecular weight polymers had Mn between 25 and 35 kg mol-1. High molecular weight polymers had Mn between 110 and 130 kg mol-1. “All” combines data for all polymers of the same family. ....................................182 Figure 5.26 Elastic Modulus for aryl-capped PLAs in this study. Low molecular weight polymers had Mn between 25 and 35 kg mol-1. High molecular weight polymers had Mn between 110 and 130 kg mol-1“All” combines data for all polymers of the same family. ......................................183 Figure 5.27 Elongation at break (%) for aryl-capped PLAs in this study. Low molecular weight polymers had Mn between 25 and 35 kg mol-1. High molecular weight polymers had Mn between110 and 130 kg mol-1. “All” combines data for all polymers of the same family. .........184 Figure 6.1 1H NMR spectrum (CDCl3, 25 °C, 400 MHz) for polymer generated from metal-catalyzed system (arm length = 3.2 kg mol-l). .............................................................................197 Figure 6.2 1H-1H COSY NMR spectrum (CDCl3, 25 °C, 600 MHz) for lignin-graft-PLA generated from metal-catalysed system (arm length = 13 kg mol-1). ...........................................................198 Figure 6.3 1H NMR spectra (CDCl3, 25 °C, 400 MHz) for polymers with different wt% of lignin, generated from metal-catalyzed system. ......................................................................................199 Figure 6.4 31P NMR spectrum (CHCl3/pyridine, 25 °C, 121 MHz) of polymer containing unreacted lignin (arm length = 5 kg mol-1). Derivatized by 2-chloro-4,4,5,5-tetramethyl-1,3,2-dioxaphospholane. Internal standard at 145.16 ppm (cyclohexanol). .........................................200 xxvi  Figure 6.5 31P NMR spectrum (CHCl3/pyridine, 25 °C, 121 MHz) of polymer generated from metal catalyzed system (arm length = 5 kg mol-1). Derivatized by 2-chloro-4,4,5,5-tetramethyl-1,3,2-dioxaphospholane. Internal standard at 145.16 ppm (cyclohexanol). ................................201 Figure 6.6 GPC traces for a series of lignin-graft-PLAs generated from the metal catalyzed graft-from approach. .............................................................................................................................202 Figure 6.7 GPC traces for linear PLA (red lines) and lignin-graft-PLAs (black lines). .............202 Figure 6.8 IR spectra of lignin-graft-PLA (long blue dashes), lignin (black line) and PLA (short red dashes). ..................................................................................................................................203 Figure 6.9 Lignin sludge obtained from industrial sources was boiled at 140 °C for 4 h and further dried under high vacuum for 24 h. Gravimetric analysis showed 77% decrease in H2O content with a final H2O content of 3.8%. ........................................................................................................206 Figure 6.10 1H{1H} NMR spectra (CDCl3, 25 °C, 600 MHz) for polymer generated from the metal catalyst. Pr = 0.67. ........................................................................................................................211 Figure 6.11 1H{1H} NMR spectra (CDCl3, 25 °C, 600 MHz) for polymer generated from the organocatalyst. Pm = 0.56. ...........................................................................................................211 Figure 6.12 1H{1H} NMR spectra (CDCl3, 25 °C, 600 MHz) for polymer generated by a graft-to synthesis. Pm = 0.60. ....................................................................................................................212 Figure 6.13 Thermogravimetric analysis traces for lignin-graft-PLA copolymers from the various synthetic routes, after removal of excess lignin. ..........................................................................213 Figure 6.14 31P NMR spectrum (CHCl3/pyridine, 25 °C, 121 MHz) of Indulin AT kraft lignin. Derivatized by 2-chloro-4,4,5,5-tetramethyl-1,3,2-dioxaphospholane. Internal standard at 145.16 ppm (cyclohexanol). ....................................................................................................................214 xxvii  Figure 6.15 Thermogravimetric analysis traces for pure lignin-graft-PLA and that contaminated with lignin. ...................................................................................................................................214 Figure 6.16 Complex viscosity, *| vs. angular frequency at 80 °C for pure lignin-graft-PLA copolymers and that lignin contaminated copolymers.................................................................215 Figure 6.17 Intrinsic viscosity vs. molecular weight for products from the three synthetic strategies routes to the graft copolymers. .....................................................................................................218 Figure 6.18 GPC traces (THF 4 mg mL–1, flow rate = 0.5 mL min–1, dn/dc = 0.040 mL g –1 ) for linear PLA (blue dashed lines), cyclic PLA (black dash-dot line) and lignin-graft-PLAs (black line). .............................................................................................................................................219 Figure 6.19 MALDI-ToF mass spectrum for polymer generated from a control graft-from polymerization (Mw,GPC = 63 kg mol-1). ......................................................................................220 Figure 6.20 MALDI-ToF mass spectrum for polymer generated from a graft-from polymerization with low [OH]lig (Mw,GPC = 78 kg mol-1). ....................................................................................221 Figure 6.21 MALDI-ToF mass spectrum for polymer generated from a graft-from polymerization with high [OH]lig (Mw,GPC = 18 kg mol-1). ...................................................................................221 Figure 6.22 Complex viscosity, |*| (Pa.s) vs. angular frequency, , (rad s-1) for linear PLA (Mw = 19 kg mol-1), star lignin-graft-PLA (Mw = 36 kg mol-1) and cyclic PLA (Mw = 63 kg mol-1). 223 Figure 6.23 Zero-shear viscosity,  vs. weight-averaged molecular weight, Mw, for linear PLAs (red triangles), cyclic PLAs (black circles) and star lignin-graft-PLA copolymers (blue squares). .  ..............................................................................................................................................223 Figure 6.24. GPC traces for a) cyclic PLAs and b) star graft copolymers generated in this study. .  ..............................................................................................................................................225 xxviii  Figure 6.25. Viscoelastic moduli (G’-filled symbols, G”-open symbols) vs. angular frequency,  (Pa.s) for linear PLA (red triangles, Table 6.7, entry 1 ), cyclic PLAs (black circles, Table 6.7, entry 7) and star lignin-graft-PLA copolymers (blue squares, Table 6.7, entry 9). .....................227 Figure 6.26. Viscoelastic moduli (G’-filled symbols, G”-open symbols) vs. angular frequency,  (Pa.s) for linear PLA (red triangles, Table 6.7, entry 2 ), cyclic PLAs (black circles, Table 6.7, entry 8) and star lignin-graft-PLA copolymers (blue squares, Table 6.7entry 11). .....................227 Figure 6.27. Viscoelastic moduli (G’- filled symbols, G”- open symbols) vs. angular frequency,  (Pa.s) for cyclic PLAs in this study. .........................................................................................229 Figure 6.28. Viscoelastic moduli (G’-filled symbols, G”-open symbols) vs. angular frequency,  (Pa.s) for star lignin-graft-PLA copolymers in this study. ..........................................................230 Figure 6.29 Temperature ramp experiments from 30 – 130 °C for a) cyclic PLA (Table 6, entry 4) and b) lignin-graft-PLA copolymers (Table 6, entry 10). ...........................................................231 Figure 6.30 Complex modulus, |G*|  vs. time at 180 °C for 60 minutes for lignin-graft-PLLA/PLLA blends. ....................................................................................................................234 Figure 6.31 Tensile strength of lignin-graft-PLLA/PLLA blends. Blends with graft arm length = 14 kg mol-1 () and Blends with graft arm length = 12 kg mol-1(). ..........................................235 Figure 6.32 Elastic moduli of lignin-graft-PLLA/PLLA blends. Blends with graft arm length = 14 kg mol-1 () and Blends with graft arm length = 12 kg mol-1(). ...............................................235 Figure 6.33 Percent elongation at break for lignin-graft-PLLA/PLLA blends. Blends with graft arm length = 14 kg mol-1 () and Blends with graft arm length = 12 kg mol-1(). ....................236  xxix  List of Schemes Scheme 1.1 Phosphitylation of lignin using 2-chloro-4,4,5,5-tetramethyl-1,3,2-dioxaphospholane (TMDP). .........................................................................................................................................17 Scheme 1.2 Oxidative cleavage of lignin structure by lignin peroxidase (LiP). ...........................20 Scheme 2.1 Synthesis of RR/RR-[(NNO)InCl]2(μ-Cl)(μ-OEt)] (RR/RR)-A. ................................49 Scheme 2.2 Synthesis of RR/RR-[(ONNO)In(-OEt)]2 (RR/RR)-B. .............................................50 Scheme 4.1 Synthesis of dinuclear indium complexes [(NNO)InCl]2(μ-Cl)(μ-OPhR) (R = OMe, Me, H, Br, NO2). ..........................................................................................................................108 Scheme 4.2 Ring-opening polymerization of rac-lactide with complexes 1-5. ..........................116 Scheme 4.3 Alkoxy-phenoxy exchange reaction with complex A and para-substituted phenol. .....  ..............................................................................................................................................124 Scheme 4.4 Synthesis of [(NNO)InCl]2(μ-Cl)(2,μ-diClPh) (6). ................................................135 Scheme 5.1 Synthesis of end-functionalized PLA via immortal ring-opening polymerization with dinuclear indium catalyst [(NNO)InCl]2(μ-Cl)(μ-OEt) (A). ........................................................156 Scheme 6.1 Graft-to synthesis for lignin graft copolymers. Distinct polymers are synthesized and through a second step are covalently bound to lignin. .................................................................191 Scheme 6.2 Graft-from synthesis for lignin graft copolymers where lignin is used as a macro-initiator for polymerization. .........................................................................................................191 Scheme 6.3 Synthesis of lignin-graft-poly(lactide)s via ring-opening polymerization using InCl3 and NEt3. ......................................................................................................................................195 Scheme 6.4 Synthesis of lignin-graft-poly(lactide)s via ring-opening polymerization using TBD. .  ..............................................................................................................................................207 Scheme 6.5 Graft-to synthesis of lignin-graft-poly(lactide)s employed in this study. ...............209 xxx  Scheme 6.6 Formation of cyclic-PLAs from the self-condensation of chloro-terminated PLA (PLA-Cl). .....................................................................................................................................216 Scheme 6.7 Proposed mechanism for the formation of cyclic PLAs in the absence lignin.282 ...226 Scheme 7.1 Proposed synthesis of thermoresponsive lignin/PLA composites. ..........................247  xxxi  List of Symbols [OH]lig hydroxy concentration in lignin (mmol g-1) [] intrinsic viscosity aT horizontal shift factor Đ dispersity  phase angle Ea,flow activation of flow energy (kJ mol-1) H Hencky shear rate  strain   strain rate G' storage modulus (Pa) G" loss modulus (Pa) |G*|  complex modulus (Pa) GN0 plateau modulus (Pa)  viscosity  * complex viscosity (Pa.s) 0 zero-shear viscosity (Pa.s) E+ extensional viscosity (Pa.s) ave average relaxation time (s) Me entanglement molecular weight (g mol-1 or kg mol-1) Mn number-average molecular weight (g mol-1 or kg mol-1) Mn,GPC molecular weight by GPC (g mol-1 or kg mol-1) Mw weight-average molecular weight (g mol-1 or kg mol-1) Mn,theo theoretical molecular weight (g mol-1 or kg mol-1)  density (g cm-3) Pi probability of a iso linkage within a PLA chain Pm probability of a meso linkage within a PLA chain ppm parts per million Pr probability of a racemic linkage within a PLA chain Ps probability of a syndio linkage within a PLA chain R gas constant (J K-1 mol-1) Rh hydrodynamic radius (nm)  stress T temperature (°C or K) Tg glass transistion temperature (°C) Tm melting temperature (°C) Tonset degradation onset temperature (°C) Tref reference temperature (°C) Ttotal complete degradation temperature (°C)  angular frequency (rad s-1 or s-1) wt% weight percent xxxii  List of Abbreviations  1,5-NapOH 1,5-naphthalenediol 1,8-NapOH 1,8-naphthalenediol a- atactic ATRP atom transfer radical polymerization biPhOH biphenol BnOH benzyl alcohol CEC chain-end control CTA chain transfer agent DACH diaminocyclohexane DCM dichloromethane diClPhOH dichlorophene DMF dimethylformamide eq. equivalents ESC enantiomorphic site control G guaiacyl lignin residue GPC gel permentation chromatography H para-hydroxyphenol lignin residue h- heterotactic HOMeAnth hydrxymethylanthracene HOMePyr hydroxymethylpyrene i- isotactic IAL indulin AT kraft lignin iROP immortal ring-opening polymerization KL kraft lignin LA lactide LDPE low density poly(ethylene) LiP lignin peroxidase LVE linear viscoelastic MALDI-ToF MS matrix assisted laser desorption ionization time of flight mass spectroscopy NapOH naphthol Oct octanoate PCL poly(caprolactone) PDLA poly(D-lactide) PDLLA poly(D,L- lactide) PE poly(ethylene) PET poly(ethylene terephthalate) PHB poly(hydroxybutyrate) PhOH phenol xxxiii  PLA poly(lactide) PLA-Cl chloro end-capped PLA PLLA poly(L-lactide) PMMA poly(methyl methacrylate) PP poly(propylene) PS poly(styrene) RIU refractive index unit ROP ring-opening polymerization S syringyl lignin residue s- syndiotactic SAOS small amplitude oscillatory shear SER Sentmanat Extensional Rheometer TBD triazabicyclodecene  TGA thermal gravimetric analysis THF tetrahydrofuran TMB 1,3,5-trimethoxybenzene TMDP 2-chloro-4,4,5,5-tetramethyl-1,3,2-dioxaphospholane Upy ureidopyrimidinone  xxxiv  Acknowledgements  My time spent completing this doctoral degree has been the most challenging and enriching experience of my life so far. Through the ups-and-downs I have discovered my strengths, my weaknesses and my passions. I am so happy to have this chance to express my gratitude that I did not have to do it alone.  First to my research supervisor Prof. Parisa Mehrkhodavandi, who took me into her group and through her example and guidance I learnt how to succeed in this academic setting. I am so grateful not only for her patience with my youthful idealism, but also for her constant support of my ideas for new projects and initiatives. Without her backing, I would not have had the courage to attempt a multi-disciplinary research project or had the many opportunities to engage with our community about the causes I believe in.   I would like to express deep gratitude to my co-supervisor Prof. Savvas G. Hatzikiriakos who exposed me to a whole new field of inquiry that has given me a deeper understanding of polymeric materials. With his support, I have gained an array of skills which will help me continue to connect communities and bridge the gap between fundamental and applied scientists.  A huge thank you to the members of my research groups; the old guard, Dr. Insun Yu, Dr. Kim Osten, Dr. Dinesh Aluthge, Dr. Mahmoud Ansari, and Molly Sung, who welcomed me into the group and helped me develop my fundamental laboratory techniques; as well as the avant-garde, Emiliya Mamleeva, Alex Kremer, Steve Chang, Tanja Tomkovic, Carlos Diaz Lopez, Xiaofang Zhai, and Hyuk Joon Jung, who were great colleagues and even better friends. I would like to recognize Tannaz Ebrahimi with whom I shared the experience of stepping into a completely new scientific discipline. I would like to thank her for her support in teaching me polymer rheology and her encouragement when the new subject got particularly confusing. I also xxxv  had the opportunity to work with some very bright undergraduate researchers, Garion Hicks, Sam Kaser, Thales Oliveira and Alan Wong who taught me both patience and the joy of mentoring.  Thank you to my parents Rebecca and Love Chile, who supported me not only with their generosity but also with their love and experience. And to my brothers, Denen and Junior Chile, and Slade Young, in them I see reflections of myself and who I strive to be.  Being so far away from home, I came to find a new family of friends. With Elise Caron, Maria Cleveland, Nick Hein, Jan Kozicky, Kaitlyn Lovering, Brenna Milne, and Janet Ochola, I commiserated my tribulations and celebrated my triumphs. Thank you for sharing with me both tears and laughter.  xxxvi  Dedication        “Only love can hurt like this.”  To my brother, my nephew, my uncle, and all those we honour by creating a sustainable future.1  Chapter 1: General introduction.  1.1 Biopolymers  Modern society has become inextricably entwined with plastics.3 Since 1839, with the development of vulcanized rubber by Goodyear, and Eduard Simons’ discovery of poly(styrene), synthetic polymeric materials have been integrated into every technological advancement across all industries. Designed to have strength, elasticity and durability, synthetic polymers have solved more problems than they have caused.4 However, the issues associated with their use, such as environmental persistence,5,6 non-renewable feedstocks,7 lack of disposal methods8 and toxicity,9,10 have long been said to outweigh the benefits of their on-going use.3  The new paradigm of biorefineries11,12 has emerged from a wide-spread shift in focus to develop circular economies13 and materials made from renewable resources.14 Based on the conversion of biomass to chemical building blocks, there is potential to empower local and regional economies by creating new industries, through which materials can be synthesized and manufactured within a recovery cycle.  This re-thinking has led to the investigation of renewable polymers15 (i.e. polymers derived mostly or entirely from biomass)15,16 and biodegradable polymers (i.e. polymers which are capable of being decomposed by bacteria or other living organisms).17 A review of the literature reveals four major categories of exploration: biomass-derived polymers, monomers from microbial production, monomers from renewable sources and monomers from non-renewable sources. Each of these categories have been extensively studied, but for this thesis two of these “biopolymers” will be discussed, poly(lactide) and lignin (Figure 1.1).  2   Figure 1.1 Four categories of biodegradable polymers found in the literature.17   3  1.1.1 Poly(lactide): Synthesis, characterization, biodegradation and application.  Poly(lactide), or PLA, is a biodegradable polymer whose ester linkages are inherently susceptible to hydrolysis and subsequent degradation.18 The use of PLA has various advantages including, production from renewable agricultural sources, recyclability via hydrolysis or alcoholysis,19 reduction of land-fill levels and ability to tailor physical properties by material modification.20-26  The linear monomer, lactic acid (2-hydroxy propionic acid) is produced industrially via homofermentation of carbohydrates such as glucose or lactose, found in beet sugar and corn starch.27,28 High production yields of lactic acid have been attained when using bacteria from the Lactobacillus genus, which have become the industry standard.29,30 Fermentation occurs under low oxygen concentration, at pH and temperatures of around 5.4 - 6.4 and 38 - 42 °C, respectively.28,29  PLA is commonly synthesized via two routes: a) polycondensation of lactic acid or b) ring-opening polymerization (ROP) of the cyclic dimer, lactide (LA) (Figure 1.2).28 Direct polycondensation of lactic acid gives low molecular weight PLA, which is prone to hydrolysis by water generated in the reaction. Due to its low molecular weight, this material is not useful for many commercial applications.29 Azeotropic distillation processes remove water from the reaction as it is generated, and have been used industrially to achieve high molecular weight PLA from lactic acid. However, this process requires the use of solvents and is energy intensive, making it less economically viable.18   4    Figure 1.2 Production cycle for poly(lactide).   Natureworks LLC, the leading producer of lactic acid polymer technology, synthesizes 8 × 106 tons/year of PLA using a proprietary synthesis conducted in the melt phase rather than in solution.31 Lactic acid is first polymerized to low molecular weight PLA via continuous condensation. Depolymerization of oligomeric PLA with tin(II) carboxylates or alkoxides produces lactide, which can undergo ROP with tin octanoate (SnOct2) to form high molecular weight PLA.18,29 Three stereoisomers of lactide exist: D- and L-LA, where both stereocentres have 5  the same configuration, and meso-LA, where the two stereocentres have opposite configurations. Racemic-lactide (rac-LA) is a 1:1 mixture of D- and L-LA (Figure 1.2).   Different reactivity among the three isomers has been observed (further explored in Chapter 2),1,32,33 and the disparity in reactivity was attributed to dissimilar structural configurations adopted by the isomers. L- and D-LA have twisted boat geometries34 whereas meso-LA assumes a more twisted planar geometry that stabilizes its ground state.32 Thermal properties reflect the observed differences in solid-state conformation; in that melting temperatures for L- and D-LA are ~50 °C higher than that for meso-LA (94 °C vs. 47.5 °C).35  ROP of these three isomers can form PLA chains with a variety of microstructures (or different chain tacticity): isotactic PLA where all the stereocentres along the chain have the same configuration, syndiotactic PLA where stereocentres alternate throughout the chain, heterotactic PLA which has stereocentres that are doubly alternating, and atactic PLA with stereocentres which are randomly distributed along the polymer chain (Figure 1.3).   Figure 1.3 Microstructures observed for poly(lactide).  6   Industrial synthesis of PLA is more commonly achieved via the ROP of LA using homoleptic metal alkoxide complexes such as tin(II) octanoate and aluminum(III) isopropoxide. These complexes are very active and produce high molecular weight PLA but transesterification side reactions break down the growing chains causing high polydispersity in the final polymer (Đ > 2.0). Using enantiopure feedstocks of L-LA allows the formation of highly isotactic PLA using these simple catalysts; however, they are not stereoselective towards rac-LA, giving mostly atactic PLA.   In the last few decades, a number of examples of well-defined coordination complexes that effectively control the molecular weight distribution and stereochemistry of PLA have been reported.36-38 The first notable example came from Spassky and co-workers who, in 1996, reported the synthesis of a chiral aluminum salen binaphthalene complex that was highly stereoselective for the ROP of rac-lactide, and the resulting polymer was shown to have 88% optical purity.39  Generally, polymerization of LA with metal catalysts occurs through a coordination-insertion mechanism (Figure 1.4). Lactide coordinates to the metal centre and becomes activated to nucleophilic attack by either an internal or external alkoxide. A ring-opening event occurs which generates a polymeryl propagating species which can continue polymerization.   Stereocontrol in coordination complex systems arises through two mechanisms, enantiomorphic site control, where chiral ancillary ligands impart stereoselectivity, and chain-end control, where stereochemistry of the last enchained monomer selects for the next incoming monomer.40-42   7   Figure 1.4 Coordination-insertion mechanism for the ROP of lactide by metal catalysts.   As a variety of PLA microstructures can be generated from the three LA isomers, a method to distinguish between them and quantify PLA tacticity was developed. Work done in 1992 by Kricheldorf et al.33 and continued by Thakur and co-workers,43-46 used 1H, 13C{1H} and heteronuclear correlation NMR techniques to characterize PLA chains of various chain tacticity. Vert et al. employed Bernoullian statistics to investigate the different possible chain sequences based on the stereochemistry of the starting monomer.34 These chain sequences are defined as tetrad sequences, where a tetrad is a hypothetical four-unit oligomer, each of which has a certain probability of occurring within a PLA chain (Table 1.1). The level of stereocomplexity, or tacticity, occurring within a polymer chain can be determined by combining the probabilities of observing the various tetrad sequences with integration data obtained from the 1H{1H} NMR or 13C{1H} NMR spectra (Figure 1.5).47  8   Bernoullian statistics operate under the assumption that enchainment of one monomer does not affect enchainment of the next; this assumption does not always hold true, and depends on the stereocontrol mechanism operating in the system.   Table 1.1 Tetrad probabilities for lactide isomers, based on Bernoullian statistics.a, b Possible tetrads from rac-lactideb Bernoullian probability Possible tetrads from meso-lactidec Bernoullian probability  𝑃𝑟𝑃𝑚2+ 𝑃𝑚2  𝑃𝑠𝑃𝑖2+ 𝑃𝑖2  𝑃𝑟𝑃𝑚2  𝑃𝑠𝑃𝑖2   𝑃𝑟𝑃𝑚2  𝑃𝑠𝑃𝑖2   𝑃𝑟22  𝑃𝑠22   𝑃𝑚𝑃𝑟 +  𝑃𝑟22  𝑃𝑖𝑃𝑠 +  𝑃𝑠22 a Bernoullian probability equations from Coates 2002.47 b When discussing rac-lactide, literature convention uses m/r notation, m denotes a meso linkage where the lactyl units have same configuration and r denotes a racemic linkage where two bound units have the opposite configuration. c When discussing meso-LA, literature conventions use s/i notation, i denotes an iso linkage where the lactyl units have same configuration and s denotes a syndio linkage where two bound units have the opposite configuration.    9   Figure 1.5 Schematic diagrams of PLA tetrad resonances corresponding to the methine region in (a) homonuclear decoupled 1H NMR of PLA from rac-lactide; (b) 13C{1H} NMR of PLA from rac-lactide; (c) homonuclear decoupled 1H NMR of PLA from meso-lactide; (d) 13C{1H} NMR meso-lactide.48   Most consumer plastics are resistant to microbial attack, however, the ester linkages within PLAs chain structure make it much more susceptible to biodegradation.28 Degradation and compostability of PLA has been explored under a variety of conditions. In animal models, hydrolysis breaks PLA down to soluble oligomers which are metabolized by cells.49,50 When placed in landfills or in waste water streams, PLA sheets were still present after 15 months, though they were partially degraded, displaying lower molecular weight and obvious visual physical break-down.51,52 Industrial composting conditions were found to be ideal for PLA degradation. One study investigated PLA with a molecular weight of 152 kg mol-1 which was placed in a vegetable compost with a temperature, moisture content, and pH of 45-70 °C, 40-55% and 4-8, respectively. The polymer had been broken down to chains of 4.5 kg mol-1 in 17 days, and total degradation was observed after 34 days.52 Under compost conditions, decomposition begins with random non-enzymic chain scission which breaks down the main polymer chain and can be 10  accelerated in acidic or basic environments. Low molecular weight oligomers generated then diffuse from the bulk polymer and are consumed by micro-organisms.35   Solid PLA can be either amorphous or semi-crystalline, as a rule, PLAs with optical purity greater than 78% are usually crystallizable.53 PLA crystals exist as polymer helices in orthorhombic unit cells. Stereodefects introduce ‘kinks’ into the helical confirmation which reduces crystallinity and glass transition temperatures. Physical, mechanical and degradation properties of PLA are highly dependent on the transition temperatures, and thus are impacted by PLA stereoregularity.18 By controlling properties such as branching, D-isomer content and dispersity,53 commercial producers have formulated PLA resins that can be transformed into films and fibers using current processing technologies such as injection molding, extrusion and fiber spinning.28 The physical properties of PLA are often compared to poly(styrene) (PS) and poly(ethylene terephthalate) (PET) as it is strong with high modulus but is brittle with low toughness (plastic deformation) (Table 1.2).   Due to its inherent hydrolysablility, coupled with its high strength, PLA has found application as consumer articles (i.e. feminine hygiene products and nappies), in packaging (i.e. compost bags and cold water bottles), medicine (i.e. bone sutures and drug delivery micelles)54 and agricultural applications (as matrices for controlled herbicide delivery).28 However, there have been drawbacks to the widespread acceptance of PLA because of its inferior moisture barrier properties (for long-term use), low thermal stability and brittle nature.31,55-58 Research in this field is currently focused on the rubber toughening of PLA to improve thermal stability and mechanical properties. These challenges are being targeted from two perspectives: molecular chemists tune polymer properties by adjusting PLA crystallinity, functionality and composition using concepts and tools from 11  organometallic and coordination chemistry; applied material scientists achieve PLA plasticization through the use of additives and melt-blending.  A variety of highly active (high turnover frequency), productive (high turnover number) and selective (high tacticity and molecular weight control) single-site metal initiators have been developed to achieve stereoregular PLAs.36,37,59 Varying stereoregularity and crystallinity have had a marginal effect on improving the rubber toughness of PLA and most organometallic systems are not industrially viable due to low water tolerance, catalyst recovery, and melt phase selectivity.15 Regardless, these systems give access to easily tunable, well-defined polymers, allowing researchers to investigate novel PLAs and their structure-property relationships.11   Melt blending with small molecules, PLA oligomers and flexible polymers is an economic and convenient method to improve the toughness of PLA and has been successful in forming a variety of commercially available PLA resins.28 However, the formation of useful blends is complicated by the size, volume fraction, properties and interfacial adhesion of the dispersed phase.56 Small molecules and oligomers have also been shown to migrate to the surface of the melt during processing, leading to diminished properties of the final material, 56 and generally increases in toughness are linked with decreases in strength and modulus compared to the native polymer. 12  Table 1.2 Comparison of mechanical and thermal properties for common consumer plastics.29,31,55   PLLA PDLLA PS LDPE PET Mechanical properties Yield strength (MPa) 59 44 35 20 31 Tensile strength (MPa) 70 53 46 8 – 20 47 Elongation at break (%) 7 5.4 3 – 4 100 – 1000 30 – 300 Thermal properties Glass transition (°C) 63 51 95 -100 75 Melting temperature (°C) 178 -- 240 180 250 Processing temperature (°C) 210 190 230 280 255 13  1.1.2 Lignin: Structure, isolation, characterization and biodegradation.  Lignin is the second most abundant polymer available from bio-mass.60 A constituent of most plants, lignin binds cellulose and hemicellulose while imparting strength and rigidity to cell walls. Its abundance has provoked researchers to investigate the use of lignin as a component in polymeric materials.61  Lignin is comprised from phenyl propane units originating from three aromatic alcohol precursors or monolignols: para-coumaryl alcohol, coniferyl alcohol and sinapyl alcohol.60,62 These alcohols end up as three lignin residues (Figure 1.6). Biologically mediated radical coupling reactions63 (Figure 1.7) bond the monolignol units to form a highly branched, 3-dimensional macromolecule held together by abundant hydrogen-bonding interactions between methoxy and hydroxy groups (Figure 1.8).62   The composition of lignin (Figure 1.6), specifically the ratio of para-hydroxyphenyl (H), guaiacyl (G) and syringyl (S) residues, depend on the species and environment of the plant source. Hardwood* lignin contains mostly G and S units with trace amounts of H units. Guaiacyl units make up most of softwood† lignin, although low levels of H units are also incorporated. Grasses generally consist of primarily G and S units.62,64                                                  * Flowering plants, eg. cherry trees. † Seed bearing plants, eg. oak trees. 14   Figure 1.6 Three main monolignols: para-coumaryl alcohol, coniferyl alcohol and sinapyl alcohol, which become para-hydroxyphenyl (H), guaiacyl (G) and syringyl (S) units within the lignin framework.64  15   Figure 1.7 Enzymatic radical pathways for coniferyl alcohol.65   16   Figure 1.8 Schematic representation of lignin after radical coupling of monolignols.66   Lignin has an irregular and complex structure that makes it insoluble in organic solvents. Historically, analysis of its chemical structure has been difficult to determine. Argyropoulos et al. developed a novel technique which utilizes 31P NMR spectroscopy to characterize the hydroxyl groups in lignin (Scheme 1.1).67 Phosphitylation of hydroxy groups using 2-chloro-4,4,5,5-tetramethyl-1,3,2-dioxaphospholane (TMDP) is achieved in the presence of a base such as pyridine.68 Phosphitylated hydroxyls are compared to an internal standard (usually cyclohexanol) to quantify the concentration of hydroxy species, in mmol g-1 (Figure 1.9).17  Scheme 1.1 Phosphitylation of lignin using 2-chloro-4,4,5,5-tetramethyl-1,3,2-dioxaphospholane (TMDP).      Figure 1.9 31P NMR spectrum (CHCl3/pyridine, 25 °C, 121 MHz) for phosphitylated lignin, derivatized with 2-chloro-4,4,5,5-tetramethyl-1,3,2-dioxaphospholane (TMDP). IS = internal standard (cyclohexanol), 145.2 ppm. Reference = phosphitylated water, 132.2 ppm. Alkyl-OH signals found 145.5 – 150 ppm; aryl-OH signals found = 138 – 144 ppm, carboxylic acid-OH signals found 134 – 136 ppm.145.5 – 150 ppm 138 – 144 ppm 134 – 136 ppm Chemical shift (ppm) 18   Lignocellulosic biomass is a composite of lignin, hemi-cellulose and cellulose and is a feed-stock for the pulp and paper industry.64 Pulping (delignification) processes are used to remove lignin allowing fibrous cellulose to be separated for the production of paper and its related products. Pulping relies on breaking lignin-cellulose bonds via cleavage of ester and ether bonds to release the lignin from other plant cell components.    Figure 1.10 Biomass extraction processes to produce the four classes of technical lignin, kraft, lignosulfonate, organosolv and soda.64,69  19   Two methods exist for extraction of lignin from biomass: sulfur and sulfur-free processes. These processes generate four classes of technical lignin used in academia and industry (Figure 1.10).  Sulfur processes are used primarily on wood sources, and are the main delignification techniques used by the pulp and paper industry. In these processes, lignin is cooked in the presence of ‘white liqour’, a basic solution containing NaOH and Na2S. Following acidification, kraft lignin can be isolated; however, this type of lignin is contaminated with residual levels of sulfur (1-2 %). Cooking can also be done in an acidic white liquor containing aqueous SO2 (HSO3-), and produces a class of lignins called lignosulfonates. This lignin has higher sulfur content due to the presence of sulfonate groups on aliphatic side chains. The abundance of sulfonate groups make lignosulfonates soluble in water making it a useful dispersant and binder for industrial applications. These cooking processes depolymerize lignin so the final structure has significantly lower molecular weight and is chemically different to that of natural lignin.64,65,69   Sulfur-free processes are more often used for grasses and occasionally hardwoods. The structure of lignin produced from these methods are more similar to that of the plant-bound polymer.64 Ethanol/water dissolves lignocellulosic biomass at high temperature and after precipitation with a non-solvent, organosolv lignin is isolated. This class of lignin has the highest purity and shows high solubility in organic solvents. Under alkaline conditions, cooking produces the final class of technical lignin, soda lignin, which have structures most similar to native lignin. Both these classes have the potential to produce useful building blocks for chemical synthesis, which has spurred investigation into their selective degradation.6620   One role of lignin in plant cells is to protect against degradative enzymes, which means it is inherently recalcitrant towards chemo- and biodegradation.70 Lignin is biosynthesized via a radical mechanism and as such is sensitive to radiative degradation. In biological systems, lignases utilize peroxide radical pathways to degrade lignin (Scheme 1.2).71 Thermophillic ligninolytic microfungi excrete peroxidases that generate radicals and initiate the break-down of lignin macrostructure. The white-rot fungus (Pycnoporus cinnabarinus) breaks down wood and lignin in natural environments.72 Studies of the activity of this fungus under industrial composting conditions showed this microbe cannot withstand the ideal composting temperatures (40-50 °C).73,74  Scheme 1.2 Oxidative cleavage of lignin structure by lignin peroxidase (LiP).     Currently most lignin produced in the pulp and paper industry is burned to off-set energy costs,61 with only 1-2% being used for specialty products. Recently there has been a surge of interest to increase the value of this cheap and renewable source of polymeric material.69,75,76 Direct use of lignin in materials is hindered by its highly disordered aromatic structure which makes it thermally strong and brittle and thus difficult to process.77 However, these structural properties have led to lignin being used as a filler in a variety of green composites,16,65 in particular 21  to improve the properties of other bio-derived and biodegradable polymers such as PLA.26,78 The effect of unmodified lignin incorporation into PLA blends has been investigated with a variety of lignin sources and though greater thermal stability of the composites has been shown,79-82 the overall mechanical properties were diminished or remain unchanged compared to neat PLA due to poor stress transfer as a result of insufficient compatibilization between lignin and the polymer matrix.61,64,76,80,83-8622  1.2 Polymer rheology  Rheology is the study of the deformation and flow of materials when external forces are applied.87 An important aspect of polymer rheology is the study of how viscoelastic behaviour changes with respect to temperature, molecular weight and polymer microstructure. These parameters can give an indication of how the polymer will behave during industrial processing as well as shedding light on potential end-use behaviour.  1.2.1 Polymer processing technologies.  Thermoplastics have been used for consumer products and packaging because they have useful properties such as being reformable, light-weight, heat sealable and variable in their barrier properties.88 Thermoplastic materials display temperature and time dependent behaviour. At low temperatures, polymer chains have little mobility and show glass-like response. Once stretched, polymer chains cannot fully relax back to their equilibrium positions making the material brittle. As the temperature is increased, polymer chains have more free volume  allowing them to move cooperatively.89 Polymer behaviour moves into a glass transition region where the material softens. Above the glass transition temperature, Tg, amorphous polymers show rubbery elastic behaviour and at high temperature the chains are highly mobile and the material behaves as a low viscosity liquid (Figure 1.11). Semi-crystalline polymers show further softening above their Tgs but chains only become fully mobile past their melting temperature, Tm, when crystalline regions are broken and chains move freely (Figure 1.11).90 The variable behaviour of thermoplastics means they can be easily processed at temperatures well above the glass transition where they behave as viscous melts.  23   Figure 1.11 Modulus vs. temperature for an amorphous polymer (black solid line) and a semi-crystalline polymer (blue dashed line).   Extrusion is an important technique for continuous melt phase processing and is often coupled to other forming machine systems such as injection or blow molders.91 Screw extruders are commonly used for industrial processing (Figure 1.12). Polymer pellets are fed through a hopper into a heated barrel where a screw rotates, forcing pellets to move forward. The action of the screw works to mix the pellets into a homogenous melt flow. Molten polymer leaves the screw out the front of the barrel and moves through a die and into a forming machine.92  24   Figure 1.12 Schematic of a single screw-extruder.   Injection and blow molding are processing techniques commonly used to produce consumer plastics. Molten polymer flows directly into a mold where it cools and hardens to form injection molded products. In blow molding, a preform is first filled with molten polymer before air flow is used to axially stretch the polymer to fit the preform (Figure 1.13).92   Figure 1.13 Schematic of a blow mold.  25   To avoid flow instabilities and non-uniform melts, polymers require a level of thermal stability and melt strength at processing temperatures.93 Rheological properties are useful when comparing and evaluating the performance of polymers during processing. Sensitive to microstructure changes, rheological analyses can help researchers predict and control polymer properties.91   1.2.2 Polymer viscoelasticity.  Polymer melts display complex flow behaviour; acting as viscous liquids by dissipating energy when a force is applied, as well as acting as elastic solids by storing applied energy. This dual behaviour identifies them as viscoelastic materials.87 Viscoelastic properties are often described in terms of the elastic component, or storage modulus G’, and the viscous component, or loss modulus G”. The behaviour of a material depends on which of the two moduli dominate at a given temperature or deformation frequency.90 At low temperatures G’ often dominates over G”, meaning that the material shows solid response. Increasing the temperature takes the material response through a cross-over transition where G” begins to dominate and the polymer shows more fluid response (Figure 1.14).   26  temperature (oC)60 80 100 120 140 160 180G',G" moduli102103104105106107G' G" cross-overtransition Figure 1.14 Plot of the linear viscoelastic moduli, G’ and G” of a thermoplastic polymer. At temperatures below 100 °C, the storage modulus dominates; thus, the material behaves as an elastic solid. Above 100 °C the material behaves as a viscous fluid.   The viscoelastic behaviour of polymer melts also depends on the deformation time and frequency.90 Long deformation times allow chains to respond to applied stresses and chains relax to dissipate the energy. As such, G” is observed to dominate and the material shows viscous liquid response. As deformations get shorter (frequencies increase), polymer behaviour moves through a cross-over region c. After this point deformations are faster than relaxation and energy is stored within chain segments, thus solid response where G’ dominates is observed (Figure 1.15).  27   Figure 1.15 Plot of the linear viscoelastic moduli, G’ () and G” () of a hypothetical thermoplastic polymer. At deformation frequencies lower than 1 rad s-1 the loss modulus dominates thus the material behaves as a viscous liquid. Above 10 rad s-1 the material behaves as an elastic solid.   1.2.3 Polymer viscosity.  When put in motion, liquid molecules slide past each other causing internal friction and flow resistance called viscosity. Viscosity is a property that involves resistance to continuous deformation. Polymeric materials exhibit both viscous resistance to deformation and elastic energy storage.94 The viscosity of polymeric materials can be determined by inducing a shearing rate,   and measuring the resulting stress,  in the sample using a parallel plate rheometer (Figure 1.16). In this configuration, two circular plates are arranged in the same vertical axis with a polymeric sample placed in the space between them. The upper plate is rotated at a specific angular velocity, resulting in a shear force being applied to the sample. The shear stress can be related to the torque 28  required to rotate the upper plate and the shear rate to the rotational speed. The relationship between viscosity, shear stress and shear rate is described in Equation 1.1 below: /    Equation 1.1 where η is the viscosity,  is the shear stress and   is the strain rate. 90   Figure 1.16 Schematic of parallel plate rheometer.   Viscosity values are not constant and depend on conditions such as temperature and time. Ideal viscous or Newtonian flow behaviour, occurs when viscosity is independent of shear rate (i.e. water). Conversely, non-Newtonian fluids display viscous response which is dependent on the shear rate. Pseudo-plastic or shear thinning materials show a decrease in viscosity with increasing rate of shear whereas dilatant or shear-thickening materials show an increase in their viscosity with shear rate (Figure 1.17).94 29    Figure 1.17 Viscosity curves for ideally viscous, shear thinning and shear thickening flow behaviour.   1.2.4 Small amplitude oscillatory shear (SAOS) experiments.  Deformations which do not perturb a polymer past its equilibrium position fall into the linear viscoelastic regime. These deformations are usually of low strains (< 20%) and at low oscillatory frequency (< 100 rad s-1). Oscillatory tests probe chain-chain interactions and are most important when characterizing molecules and correlating their viscoelastic properties with molecular structure.9030   Consider the parallel plate rheometer presented in Figure 1.16. If a sinusoidal oscillatory strain,  is applied (Equation 1.2), the resultant stress,  can be recorded as a function of time (Equation 1.3).  (t) = 0 sin (t) Equation 1.2 where t is the strain as a function of time, 0 is the initial strain, is the deformation frequency and t is time. (t) = sint +  Equation 1.3 where  (t) is the stress as a function of time, 0 is the initial stress, is the deformation frequency, t is time and  is the phase shift or mechanical loss angle.   An entirely elastic material will not show a time lag between the applied strain and the resultant stress (, whereas in Newtonian fluids, the stress lags the strain by 90° (. As mentioned above, polymeric materials exhibit elastic storage which retards viscous resistance and as such, viscoelastic materials exhibit a phase lag between deformation and response (Equation 1.4). (t) = 0 G'( sint) + G”( cost Equation 1.4 where  (t) is the stress as a function of time, 0 is the initial strain, G'( is the elastic response (storage modulus), G”( is the viscous response (loss modulus), is the deformation frequency, t is time and  is the phase shift or mechanical loss angle.   From SAOS experiments, linear viscoelastic (LVE) moduli, G’ and G”, can be obtained. These two terms can be used to acquire various other useful parameters which give broader descriptions of linear viscoelastic behaviour.  Complex modulus, |G*|, considers both the viscous and elastic response to describe the entire viscoelastic behaviour (Figure 1.18). Similarily, the damping factor, tan( is the ratio of the loss 31  to storage moduli and describes the two parts of viscoelastic behaviour.94 Other parameters often used to describe viscoelastic behaviour are the complex viscosity η*, a deformation frequency dependent viscosity term, and the zero-shear viscosity 0, the viscosity when no stress is applied, determined by the limits of the linear viscoelastic region.90   Figure 1.18 Vector diagram showing the relationship between the complex modulus |G*|, the loss and storage moduli, G” and G’.   As rheological properties are dependent on temperature, as such, experiments performed over a range of temperatures can be used to attain a complete map of how a material behaves. These isothermal data can be superimposed by introducing a horizontal aT, and vertical shift factors bT, required to bring the data into master curves at a reference temperature. aT describes the temperature dependence of the LVE properties and follows an Arrhenius trend, from which the activation of flow energy, Ea,flow can be calculated (Equation 1.5).9032  aT= exp{Ea,flow/R(1/T − 1/Tref} Equation 1.5. where Ea,flow is the activation energy for flow, R is the universal constant of the ideal gas law, and Tref is the reference temperature.   1.2.5 Uniaxial extensional rheology.  Extensional deformations align polymer chains and if deformation time exceeds the rate of molecular relaxation, strain hardening can occur. This effect comes about when polymer chains resist deformation, causing a sharp increase in viscosity before failure. Strain hardening is useful because a polymer can harden upon being stretched into a mold, improving strength and barrier properties.88 Normally, strain hardening is observed when long-chain branching or network structures are present within the polymer. These structures increase relaxation times by reducing the number of relaxation modes,21 meaning that there are fewer ways in which the polymer can relax to its equilibrium configuration.  A Sentmanat Extensional Rheometer Universal Testing Platform (SER)95 can be used to explore strain hardening in polymer melts. A polymer is loaded onto the instrument by placing it between two drums. Rotating one drum applies a uniaxial elongation to the polymer (Figure 1.19). Elongational viscosity or tensile strength growth coefficient, E+, can be plotted against time to reveal strain hardening, which is seen as an upturn in viscosity (Figure 1.20). This effect is influenced by both extensional rate and temperature. Fast extensions do not allow chains time to respond to deformation and are more likely to show strain hardening. Likewise, at low temperature polymer chains have less mobility and can also show strain hardening. 33   Figure 1.19 Sentmanat Extensional Rheometer Universal Testing Platform (SER).95    Figure 1.20 Plot of the elongational viscosity E+ vs. time for a hypothetical polymer. At low strain rates, no strain hardening is observed. At strain rates greater than 1 s-1, strain hardening is present. Torque shaft Leading drum Securing clamp Sample Secondary drum 34  1.3 Scope of thesis   This multidisciplinary thesis links organometallic catalysis with polymer rheology. The aim is to take a broader view of a research question and use what I find to create efficient catalytic systems where modifications to PLA can be easily made, and will impart useful properties in a predictable way, allowing PLA to be used wider array of consumer applications. In particular, I aim to use molecular chemistry to probe catalysis and polymer synthesis, in order to tune PLA properties.  The thesis begins with an analysis of the ring-opening polymerization of lactide isomers with dinuclear indium catalysts. Chapter 2 describes how catalysis operates in order to design better systems that can attain highly stereoregular PLA. This analysis of stereocontrol and stereochemistry, is followed by a comparison of the structure-property relationships for stereoregular PLAs (Chapter 3). Seeing the influence of interchain interactions on polymer properties, I targeted a series of functionalized polymers. Aryl initiators were synthesized to gain better understanding of how polymerization behaviour changes when using phenolic catalysts and initiators (Chapter 4). By exploiting a facile synthesis, functionality was installed as polymer end-groups. The chain-chain interactions imparted by end-group aggregation is explored in Chapter 5. Using insight gained from examining aryl initiators for the ROP of LA, an effort to incorporate lignin into PLA composite materials is described in Chapter 6.    35  Chapter 2: Living ring-opening polymerization of meso-lactide using indium initiators.‡  2.1  Introduction 2.1.1 Controlled tacticity via ring-opening polymerization of lactide monomers.  L-(+)-lactic acid can be synthesized industrially by the bacterial fermentation of D-glucose found in corn sugar (see section 1.1.1). High temperature polycondensation followed by metal-catalyzed depolymerization of oligomers, yields lactide (LA). Racemization during the depolymerization process forms three isomers of lactide: D-LA, L-LA and meso-LA (Figure 2.1).88 Catalytic ring-opening polymerization (ROP) of the different lactide isomers produces poly(lactide) (PLA) whose microstructure or tacticity depend on the catalysts’ preference for one enantiomer over the other.34,36   Homopolymerization of enantiopure D- or L-LA forms isotactic PLA where all chiral centres along the chain have the same stereochemistry. Highly ordered isotactic stereo-block PLA can be synthesized through the sequential copolymerization of two isotactic units. Ring-opening polymerization of the racemic mixture of enantiomers, rac-LA, gives varying degrees of stereo-gradient PLA depending on the stereo-control mechanisms acting during polymerization.96 Lack of stereo-control leads to random incorporation of D- and L-LA from rac-LA forming largely disordered amorphous atactic PLA where configurations of stereocentres are randomly distributed                                                  ‡ This work has been published in the journal Macromolecules: Chile, L.-E.; Mehrkhodavandi, P.; Hatzikiriakos, S. G.; Macromolecules 2016, 49, 909. (doi: 10.1021/acs.macromol.5b02568). 36  along the polymer chain. Heterotactic microstructures can be obtained from rac-LA through the alternating inclusion of D- and L-LA into the growing polymer chain (Figure 2.1).   Isotactic PLA has been reported to show high crystallinity, indicating the presence of significant chain-chain interactions which give rise to regions of order and aggregation of polymer chains.96-101 By tailoring the microstructure, desirable properties may be imparted to the resulting polymer (discussed further in Chapter 3).88 Consequently, reports of catalysts achieving well-ordered microstructures are very common in the literature.36,59 Following these advances, and with the associated increase in commercial PLA production, the use of meso-LA as a monomer has become more prevalent.37 Researchers have been investigating the ring-opening polymerization of meso-LA to determine if polymers with useful properties can be obtained from this abundant starting material.  Ring-opening polymerization of meso-LA has the possibly of forming polymers within three categories of microstructures: atactic, heterotactic and syndiotactic PLA (Figure 2.1). Syndiotactic PLA has a complex alternating arrangement of stereocentres along the polymer chain making this type of microstructure more difficult to synthesize and thus less common.102,103 To achieve purely syndiotactic PLA from meso-LA, ring-opening should occur predominately at one of the two enantiotopic sites; conversely, heterotactic PLA can be obtained from meso-LA if ring-opening alternates between the two enantiotopic acyl oxygen sites (Figure 2.1, sites a and b).104 If there is no preference for either enantiomorphic site during polymerization, then atactic PLA is formed. Due to their stereoregularity, syndiotactic polymers are typically crystalline and have enhanced mechanical properties compared to amorphous polymers.102  37   Figure 2.1 Stereoisomers of lactide and the possible microstructures of poly(lactic acid).   2.1.2 Mechanisims for stereocontrol.  Stereocontrol in metal-catalyzed ring-opening polymerization of lactide arises through two mechanisms: enantiomorphic site control (ESC) and chain-end control (CEC). In ESC, selectivity in the ring-opening event is promoted by utilizing ancillary ligands that prefer one enantiotopic site over the other. This differs from CEC where the stereocentre of the last monomer in a propagating chain selects for the next incoming monomer.  38   Due to the range of ordered and random microstructures reported for both chiral and achiral catalysts, it can be said that these two mechanisms are not independent and both may play an important role at different points in the polymerization process.32 As both mechanisms contribute to the final microstructure, stereodefects can be incorporated into the polymer chain due to a mismatch in stereocentres when an incoming monomer is enchained. Even small degrees of stereo-irregularity can dramatically decrease the physical properties of the resulting polymer.29 Therefore, controlling and determining the number of stereodefects along the polymer chain is crucial for synthesizing PLA with desirable properties.  2.1.3 Determining poly(lactide) tacticity.   Historically optical polarimetry has been used to determine the optical rotation of the polymer in solution, this gives the enantiomeric purity and an estimation of the level of stereoregularity within a PLA sample.32,39 Enantiomeric purity does not fully capture the configurational structures occurring within PLA, and early researchers determined that more information about chain tacticity and the stereo-control mechanism can be gained through inspection of the NMR spectra of the polymers.34   The level of stereocomplexity occurring within a polymer chain is expressed by two parameters: Pi defined as the probability of having an iso linkage where two adjacent lactyl units have the same stereochemistry (often reported as Pm when discussing rac-LA), and Ps defined as the probability of having a syndio linkage where the lactyl units have opposite configuration (often reported as Pr when discussing rac-LA). A probability value close to one implies high levels of 39  order, while values close to 0.5 imply atacticity. Bernoullian statistics§ can be used to describe the various possible tetrad chain sequences** dependent on the stereochemistry of the starting monomer (Figure 2.2).43-46 Tacticity values can be calculated from analysis of the methyl region of the homonuclear decoupled 1H{1H} NMR spectra, where resonances for each possible tetrad sequence are observable after removal of scalar coupling between methine and methyl protons (Figure 2.3).105 The relative integrations can be substituted into Equation 2.1 and Equation 2.2 (derived from Bernoullian probabilities in Table 1.1) to determine the Pi and Ps values.   Figure 2.2 Examples of tetrad sequences from the ring-opening polymerization of lactide. a) isotactic PLA b) heterotactic PLA and c) syndiotactic PLA. i indicates an iso linkage where two bound lactyl units have the same arrangement s indicates a syndio linkage where the lactyl units have opposite configuration.                                                 § Bernoullian statistics assume that the configuration of each newly enchained monomer is independent of the previous monomer, an assumption which does not always hold true (see section 1.1.1, page 8) ** A tetrad sequence is four bound lactyl units (Figure 2.2). 40   Figure 2.3 1H[1H} NMR spectra (CDCl3, 25 °C, 600 MHz) of PLA samples resulting from ring-opening polymerization of different lactide isomers. The chemical shifts for the methine proton (shown in purple) depend on the tetrads (labelled in black) in which it is found.   Pi = [iii] a = x + y + z  Pi = [sis] a = x + z Ps = [sis] + [isi] Ps = (2𝑥𝑎)12  Ps = [sss] +  [ssi/iss/isi] Pi = √2𝑥𝑎  Pi = (2𝑦𝑎)÷Pi   Ps = 2(𝑧𝑎−𝑃𝑖2)𝑃𝑖 Equation 2.1 Equation to calculate tacticity for polymers formed from rac-LA.   Equation 2.2 Equation to calculate tacticity for polymers formed from meso-LA.    41   As stated above, purely syndiotactic PLA is obtained by ring-opening meso-LA predominately at one of the two enantiotopic acyl-oxygen sites. A catalyst’s preference for one site can be determined by the rate at which the catalyst ring-opens at one position over the other (Figure 2.4, a vs. b).106 This enantiotopic selectivity ratio can be quantified by the α-value (the probability of ring-opening at one site over the other) and is determined experimentally by comparing the ratio of the integrations for the [sss] and [sis] tetrads or by examining the rate of polymerization for L- and D-LA.   Figure 2.4 A catalyst’s preference for a certain enantiomeric site is determined by the rate at which it ring-opens at one site over the other (a vs. b), quantified by the α-value.106   The chain tacticity arising from catalytic systems with high enantiomorphic site control, can be approximated by substituting the α-value into Equation 2.3 below. These equations, however, underestimate the influence of defect tetrad peaks and, if used with catalyst systems which operate through any degree of chain-end control, the calculated polymer tacticity will be artificially inflated.42  Pi = [sis] α = z(x +  z) Ps = [sss] + [ssi/iss/isi] [sss]  =  ½[α3  +  (1 − α)3  +  α2  +  (1 − α)2]  [sis] =  α(1 − α)  [ssi/iss/isi] =  3(½[α2(1 − α)  +  α(1 − α)2]) Equation 2.3  Equation to calculate tacticity for polymers formed from meso-LA using the enantiotopic selectivity ratio, α-value.103   Insertion errors are corrected differently depending on the stereocontrol mechanism. Examination of defect tetrad peaks in the methine region of the 13C{1H} NMR spectrum can give an indication of which control mechanism dominates during polymerization. In systems operating under CEC, defects are propagated until the next insertion error occurs, giving a ratio of 1:1:1 for the defect tetrads iis:isi:sii (Figure 2.5). When considering CEC systems polymerizing meso-lactide, the same ratio applies to the defect tetrads ssi:sis:iss (Figure 2.6).  In ESC, because the tacticity is determined by the chirality at the metal centre, mistakes are corrected with the next insertion, giving a 1:2:1:1 ratio for the defect tetrads iis:isi:sis:sii when polymerizing rac-lactide (Figure 2.5), and ssi:sis:isi:iss  when polymerizing meso-lactide (Figure 2.6).   To wholly characterize and compare polymers of different tacticity, it is important to supplement NMR analysis with other techniques such as differential scanning calorimetry, viscometry or rheology (see Chapter 3).  43   Figure 2.5 Possible insertion errors from ROP of rac-lactide via chain-end control (CEC) and enantiomorphic site control (ESC) (where i denotes isotactic enchainment, (relative stereochemistry of adjacent lactyl units are the same) and s denotes syndiotactic enchainment (relative stereochemistry of adjacent lactyl units are different)).    Figure 2.6 Possible insertion errors from ROP of meso-lactide via chain-end control (CEC) and enantiomorphic site control (ESC) (where i denotes isotactic enchainment, (relative stereochemistry of adjacent lactyl units are the same) and s denotes syndiotactic enchainment (relative stereochemistry of adjacent lactyl units are different)).   2.1.4 Catalysts for synthesizing PLA with controlled microstructure.  In the last few decades a number of examples of coordination complexes which effectively control the molecular weight distribution and stereochemistry of lactide isomers have been reported (major contributions are summarized in Figure 2.7).27 The first notable example comes from Spassky and co-workers who in 1996 reported the synthesis of a chiral aluminum salen binaphthalene complex which was highly stereoselective for the ROP of rac-lactide.39 The resulting polymer had 88% optical purity and showed good molecular weight control due to suppression of transesterification reactions. Following this landmark report, many researchers 44  have made important contributions to this field. Their reports convey a variety of ligand motifs to support a range of metals from the main group,48,107-109 transition,110-113 and lanthanide114  metals, achieving highly isotactic or heterotactic PLA with good control over molecular weights while using mild reaction conditions.27   Despite the vast amount of work going into the polymerization of rac-LA, there are far fewer examples of stereoselective catalysts for the ROP of meso-LA.37 The landmark catalyst in this field was reported by Ovitt and Coates in 1999. In this report, a chiral aluminum salen binaphthalene complex (Figure 2.7 I39) catalyzed the ROP of meso-LA, affording highly syndiotactic PLA in 14 h at 70 °C (Ps = 0.96, Đ ≈ 1.1).103,106 Two years after their pioneering report, Coates et al. described a series of zinc(II) alkoxides bearing easily functionalized bulky β-diiminate (BDI) groups which were shown to catalyze the ROP meso-LA, giving syndiotactically enriched polymer at 0 °C in under 4 hours (Ps = 0.76, Đ ≈ 1.1).115    Figure 2.7 Catalysts reported for the ROP of meso-lactide. I–Coates 1999103 II–Coates 2001115 III– Okuda 2010116 and 2012 117 IV–Okuda 2011118 V–Carpentier 2006119 VI–Hendrick 2006.104  45   In 2010, Okuda et al. communicated the synthesis of Group 3 metal complexes bearing (OSSO)-bis(phenolate) ligands with varying levels of rigidity in their backbones. These catalysts polymerized meso-LA to form moderately to highly syndiotactic PLA in under an hour at room temperature (Ps = 0.71-0.93); however, they did not show good control of molecular weight (Đ = 1.3-2.2).116 By switching to Group 4 metals, the OSSO-type catalysts showed improved molecular weight control (Đ = 1.0-1.2) and high heteroselectivity towards meso-LA at 100 °C in toluene (Pi = 0.90).117 Further work from Okuda and co-workers demonstrated how scandium complexes supported by sterically dominating tetradentate cyclen-derived (NNNN) macrocycles were highly active towards meso-LA giving moderately syndiotactic PLA at room temperature (Ps = 0.74, Đ = 1.7-2.1).118   Other systems have shown varied success for the polymerization of meso-LA. In 2006, Carpentier et al. described the synthesis of yttrium complexes supported by flexible dianionic alkoxy-amino-bisphenolate ligands which catalyzed the polymerization of meso-LA to give modest syndiotactic enrichment (Ps = 0.75) under ambient conditions (20 °C) in less than an hour, albeit with decreased control over molecular weight (Đ ≈ 1.2).114  Hillmyer and Tolman diverged from the use of bulky supporting ligands in 2009 and reported a simple salt system using InCl3, benzyl alcohol and triethylamine which formed a catalytic species in situ that subsequently polymerized rac-lactide to give highly heterotactic PLA (Pi = 0.96); however only atactic PLA was achieved from the polymerization of meso-LA (Ps = 0.56).120,121 By exploiting a metal-free system, high levels of heterotacticity (Pi = 0.83, Đ ≈ 1.3) at −40 °C were achieved by Hendrick and co-workers who used an organocatalyst comprising of N-heterocyclic carbenes.104   46   The origin of selectivity in these systems is often unclear. Systems reported by Coates generate differently ordered PLA microstructures when using an enantiopure complex compared to a racemic version of the same catalyst (Figure 2.7 I).47,103 For these systems it has been proposed that stereoselectivity for the catalysts arises mainly through enantiomorphic site control. Shown in Figure 2.7 II and V are two achiral coordination complexes, a bulky β-diiminate (BDI) zinc complex and a yttrium complex supported by a dianionic alkoxy-amino-bisphenolate ligand, which show a preference for ordered microstructures most likely occurring through chain-end control. 115,122   Even with the concerted effort to control the polymerization of meso-LA, the systems discussed above either show insufficient molecular weight control or require elevated temperatures and long reaction times to achieve high monomer conversions. The Mehrkhodavandi group has reported two families of indium catalysts active for the ROP of rac-LA (Chart 2.1) The first-generation catalyst is a dinuclear asymmetrically bridged diaminophenolate system [(NNO)InCl]2(μ-Cl)(μ-OEt) (A) (Chart 2.1), which exerts moderate isoselectivity (Pm ~ 0.6, Đ ≈ 1.0) for the ROP of rac-LA.123,124 The second generation catalysts [(ONNO)In(OEt)]2 (B) are more active and show greater stereocontrol towards rac-LA giving isotactically enriched polymers (Pm ~ 0.75, Đ ≈ 1.5).125  In the following sections the ring-opening polymerization behaviour of chiral catalysts A and B towards meso-LA and an examination of the resulting syndiotactically enriched polymers will be discussed.   47  Chart 2.1 Chiral indium based catalysts (RR/RR)-[(NNO)InCl]2(μ-Cl)(μ-OEt) (A) and (RR/RR)-[(ONNO)In(-OEt)]2 (B).     2.2 Results and discussion 2.2.1 Purification of meso-lactide.    Meso-LA is a by-product of lactide production from L-lactic acid. During monomer processing, substantial amounts of meso- and D-LA are produced. When meso-LA is isolated, it is contaminated with significant amounts of D- and L-LA. These impurities can cause the incorporation of rac-LA units into the growing polymer causing stereodefects in the chain. Stereodefects interrupt the chain sequence, decreasing the crystallinity and diminishing polymer properties. To adequately analyse how the catalysts behave towards meso-LA, rigorous purification of meso-LA is necessary.   Meso-LA, obtained from Purac Inc., was analysed by 1H NMR spectroscopy and was shown to contain on average 12% rac-LA impurities (Figure 2.8). Sublimation of meso-LA was successful in removing rac-LA impurities, but this technique was slow and low yielding. A more 48  efficient method of purification was achieved by drying a solution of meso-LA in DCM over MgSO4 followed by re-isolation via vacuum filtration. After at least five recrystallizations from warm dry isopropanol, meso-LA with >95% purity (by GC-MS) can be obtained (see appendix A).    Figure 2.8 1H NMR spectra (CDCl3, 25 °C, 300 MHz) of meso-lactide (blue star) and diminishing rac-lactide impurity (red star) a) after drying b) after three recrystallizations from iPrOH c) after five recrystallizations from iPrOH.   49  2.2.2 Polymerization of meso-lactide with tetra- and tridentate indium alkoxide complexes.  The racemic and enantiopure versions of proligands: asymmetrically substituted trans-diaminocyclohexane (±)- or (RR)-H(NNO) and a tetradentate diimino-diol (±)- or (RR)-H2(ONNO), can be prepared according to known methods.40,123,125 Deprotonation of these proligands with benzyl potassium followed by addition of one equivalent of indium trichloride yields the respective indium chloride species. Racemic and enantiopure alkoxy complexes, [(NNO)InCl]2(μ-Cl)(μ-OEt)] (±)- and (RR/RR)-A and [(ONNO)In(μ-OEt)]2 (±)- and (RR/RR)-B can then be obtained by reaction with potassium or sodium ethoxide, respectively (Scheme 2.1 and Scheme 2.2, only the RR/RR enantiomers are shown).   Scheme 2.1 Synthesis of RR/RR-[(NNO)InCl]2(μ-Cl)(μ-OEt)] (RR/RR)-A.  50  Scheme 2.2 Synthesis of RR/RR-[(ONNO)In(-OEt)]2 (RR/RR)-B.     Polymerization reactions with the appropriate volumes of A and B (2.8 mM) and meso-LA (2.78 M) were carried out in DCM at room temperature for 16 and 4 h, respectively (Table 2.1 and Table 2.2). Conversions and polymer tacticity were determined by 1H and 1H{1H} NMR spectroscopy (CDCl3, room temperature). Molecular weight (Mn) and dispersity (Đ) are obtained via gel permeation chromatography (GPC).   Polymers formed from the reaction of (±)- or (RR/RR)-A with meso-LA show good agreement between the calculated and experimental molecular weights (Figure 2.9). Low molecular weight distributions (Đ < 1.10) indicate that the polymerization is living as previously observed for rac-LA.40,123 Polymerizations with B show less control over molecular weight at higher monomer loadings than those carried out with A due to previously described depolymerization reactions catalysed with B as observed by transesterification detected MALDI-ToF mass spectrometry data and dispersity values greater than 1.10 (Figure 2.9).126  51         Figure 2.9 Plot of observed PLA Mn () and Ð (●) as functions of added meso-LA for (±)- (RR/RR)-A (top) and for (±)- (RR/RR)-B (bottom). (Mn = number averaged molecular weight, Ð = dispersity index). The black line indicates theoretical Mn values based on the [LA]:[initiator] ratio at 100% conversion. All reactions were carried out at room temperature for 16 h (A) or 4 h (B) in DCM and polymer samples were obtained at >90% conversion.   Analysis of the methine region of polymers resulting from catalyst A by 1H{1H} NMR spectroscopy elucidates the microstructures of isolated material (Figure 2.10). Over the range of monomer to initiator ratios, the PLA produced by (RR/RR)-A shows syndiotactic enrichment (0.71 > Ps > 0.73). PLAs isolated from (±)-A display little syndiotacticity, with Ps values of 0.54-0.56 (Table 2.1, entries 8-14). This trend is opposite to that seen with rac-LA where (±)-A shows greatest isoselectivity.    1.01.21.41.61.82.00501001502002503003504000 500 1000 1500 2000 2500Ð ()Mn(kg mol-1) (■)[LA]:[initiator][(NNO)InCl]2(μ-Cl)(μ-OEt) (±)-A 1.01.21.41.61.82.00501001502002503003504000 500 1000 1500 2000 2500Ð ()Mn(kg mol-1) (■)[LA]:[initiator][(NNO)InCl]2(μ-Cl)(μ-OEt) (RR/RR)-A 1.01.21.41.61.82.00501001502002503003504000 500 1000 1500 2000 2500Ð()Mn(kg mol-1) (■)[LA]:[initiator][(ONNO)In(-OEt)]2 (±)-B1.01.21.41.61.82.00501001502002503003504000 500 1000 1500 2000 2500Ð()Mn(kg mol-1) (■)[LA]:[initiator][(ONNO)In(-OEt)]2 (RR/RR)-B52  Table 2.1 Polymerization of lactide with dinuclear indium complexes (±)-A and (RR/RR)-A.††  Entry Catalyst [LA]:[initiator] Mn,theo (g mol-1)a Mn,eGPC (g mol-1)b Psc Psd Ð 1 (RR/RR)-A 250 35400 30200 0.71 0.83 1.02 2 (RR/RR)-A 470 66400 50300 0.71 0.83 1.04 3 (RR/RR)-A 870 79800 73800 0.72 0.83 1.00 4 e (RR/RR)-A 1000 146000 136000 0.76 0.86 1.00 5 (RR/RR)-A 1500 21800 204000 0.72 0.82 1.04 6 (RR/RR)-A 2000 284000 275000 0.73 0.84 1.03 7 (RR/RR)-A 2400 344000 299000 0.73 0.84 1.07 8 (±)-A 250 35300 44300 0.55 0.81 1.02 9 (±)-A 490 69600 72700 0.55 0.81 1.01 10 (±)-A 990 141000 127000 0.56 0.80 1.01 11 (±)-A 1006 144000 112000 0.58 0.81 1.10 12 (±)-A 1500 213000 198000 0.56 0.82 1.02 13 (±)-A 2000 281000 255000 0.56 0.82 1.03 14 (±)-A 2500 350000 219000 0.53 0.82 1.09 Reactions were carried out in DCM at 25 °C to greater than 90% conversion, [A]o ≈ 2.8 mM; a Calculated from Mn,theo = (144 g mol–1 × conversion × [LA]/[initiator]). b Absolute molecular weights were determined by triple detector GPC (gel permeation chromatography) via Universal Calibration (THF 4 mg mL–1, flow rate = 0.5 mL min–1,  dn/dc = 0.044 mL g –1 ).  c Calculated from 1H{1H} NMR spectra using Equation 2.2.46 d Calculated from 1H{1H} NMR spectra using Equation 2.3.103 e Experiment carried out at 0 °C for 48 hours.                                                    †† Tacticity values in this table differ from those published in Chile, L.-E.; Mehrkhodavandi, P.; Hatzikiriakos, S. G.; Macromolecules 2016, 49, 909. 53    Figure 2.10 1H{1H} NMR spectra (CDCl3, 25 °C, 600 MHz) of polymers generated from a) (±)-A and b) (RR/RR)-A. 54   There have been cases reported where greater stereoselectivity has been observed at lower temperatures. Hendrick et al. observed a marked increase in syndioselectivity (Pi = 0.62 to 0.83) by reducing polymerization temperatures from 25 °C to −40 °C.104 A zinc β-diiminate (BDI) complex investigated by Coates imparts moderate syndiotacticity at 0 °C (Ps = 0.76), but no room temperature experiments with meso-LA were reported.115 Polymerizations with (RR/RR)-A exhibit modest syndiocontrol. So, low temperature experiments at 0 °C were conducted for 48 hours in an attempt to enhance selectivity. Good agreement with calculated molecular weights and high conversions are seen, but a significant increase in syndiotacticity was not observed (Table 2.1, entries 4 and 11).   Inspection of the 13C{1H} NMR spectrum of polymers generated from (RR/RR)-A, confirms that the sss tetrad is the major linkage present in the polymer chain, giving more evidence towards this catalyst being syndioselective. The presence of all four other defect peaks indicates both CEC and ESC stereocontrol mechanisms are in action during polymerization, but the higher ratio of the sis peak compared to the isi peak implies that ESC is the major mechanism acting in this system (Figure 2.11).  Included in Table 2.1 are Ps values calculated using Equation 2.3 for comparison. Syndioselective catalyst (RR/RR)-A, shows Ps values which differ by 13%, whereas in the less selective system, (±)-A, the variance is close to 30%, illustrating the limitations of the using enantiotopic selectivity to approximate these values.  55   Figure 2.11 Inverse-gated 13C{1H} NMR spectrum (CDCl3, 25 °C, 100 MHz) for a polymer generated from (RR/RR)-A.   Both enantiopure and racemic versions of catalyst B are isoselective for the ring-opening polymerization rac-LA, forming polymers with Pm values of ~ 0.73.  These isorich polymers have good molecular weight control for [LA]:[initiator] ratios of up to 1000 albeit with high dispersities (Đ ~1.5).125 Catalyst B displays nominal syndiocontrol towards the polymerization of meso-LA compared with catalyst A, giving moderately heterotactic polymers with Ps values ranging from 0.35 to 0.39 for (±)-B and 0.42 to 0.43 for (RR/RR)-B (Table 2.2 and Figure 2.12).   Analysis of the defect tetrads in the 13C{1H} NMR spectra for polymers generated from (±)-B confirms that the isi and sis tetrads are most prominent, giving rise to the heterotactic 56  microstructure (Figure 2.13). The presence of all four defect tetrads is an indication that both CEC and ESC are acting during polymerization, though from the defect peaks alone, it is difficult to postulate which mechanism is dominant. Using the enantiotopic selectivity approximation only holds true for systems acting through ESC, and that this approximation completely fails with catalyst B, predicating moderate syndiotacticity (Table 2.2). These results suggest that CEC is the major control mechanism acting during polymerization of meso-LA.57  Table 2.2 Polymerization of meso-lactide with dinuclear indium complexes (±)-B and (RR/RR)-B.‡‡  entry Catalyst [LA]:[initiator] Mn,theo (g mol-1)a Mn,GPC (g mol-1)b Psc Psd Ð 1 (RR/RR)-B 250 35300 30200 0.51 0.77 1.02 2 (RR/RR)-B 490 69900 48900 0.43 0.77 1.22 3 (RR/RR)-B 980 141000 74700 0.43 0.76 1.25 4 (RR/RR)-B 1400 201000 103000 0.43 0.78 1.22 5 (RR/RR)-B 1800 260000 105000 0.42 0.78 1.18 6 (RR/RR)-B 2300 328000 121000 0.42 0.78 1.26 7 (±)-B 250 35300 31600 0.46 0.77 1.22 8 (±)-B 490 70600 52500 0.39 0.76 1.17 9 (±)-B 800 115000 52100 0.35 0.77 1.14 10 (±)-B 1300 190000 110000 0.34 0.77 1.13 11 (±)-B 2000 283000 143000 0.35 0.77 1.17 12 (±)-B 2500 353000 146000 0.35 0.77 1.14 Reactions were carried out in DCM at 25 °C to greater than 90% conversion, [B]o ≈ 2.7 mM; a Calculated from Mn,theo = (144 g mol–1 × conversion × [LA]/[initiator]). b Absolute molecular weights were determined by triple detector GPC (gel permeation chromatography) via Universal Calibration (THF 4 mg mL–1, flow rate = 0.5 mL min–1,  dn/dc = 0.044 mL g –1 ).  c Calculated from 1H{1H} NMR spectra using Equation 2.2.46 d Calculated from 1H{1H} NMR spectra using Equation 2.3.103                                                 ‡‡ Tacticity values in this table differ from those published in Chile, L.-E.; Mehrkhodavandi, P.; Hatzikiriakos, S. G.; Macromolecules 2016, 49, 909. 58   Figure 2.12 1H{1H} NMR spectra (CDCl3, 25 °C, 600 MHz) of polymers generated from a) (±)-B and b) (RR/RR)-B. 59  .  Figure 2.13 Inverse-gated 13C{1H} NMR spectrum (CDCl3, 25 °C, 100 MHz) for a polymer generated from (±)-B.   2.2.3 In situ studies.    Polymerization of meso-LA using (±)- and (RR/RR)- analogues of complexes A and B can be monitored in situ using 1H NMR spectroscopy. To determine the propagation rate, kobs, experiments were performed at room temperature in CDCl3 with a monomer to initiator ratio of 200 and monitored to at least 90% conversion.  The plots of ln[LA] vs. time for all catalysts are linear after a short initiation period for catalyst A, and a longer period for B (Figure 2.14, Table 2.3). The long initiation period for B is due to dissociation of the dimer prior to polymerization.127 The linear relationship between concentration and time implies that first order kinetics (with 60  respect to the monomer) is obeyed for these systems. Table 2.3 compares rates of reaction for all catalysts and the difference LA isomers. Generally, meso-LA is polymerized at a slower rate compared to the other isomers.33 In the solid state, L- and D-LA adapt twisted boat geometries to minimize the interaction between the two carbonyl oxygens.128 This distorted structure increases ring strain within the monomer which drives ring-opening polymerization. Meso-LA, with two different configurations at the stereocentres, adopts a more twisted planar geometry, stabilizing its ground state.32 The ~4 kcalmol-1 difference in ground-state energy translates to a slower rate of living ring-polymerization for meso-LA compared to the other two isomers (Table 2.3).  Polymerization rate kinetics were explored to elucidate the mechanisms through which catalyst A acts (Figure 2.4). (±)-A was previously shown to be a racemic mixture of (RR/RR)-A + (SS/SS)-A and not a heterochiral dimer and thus, if acting through ESC, one enantiomer of the catalyst should favour the coordination of only one isomer of LA.40 Confirming this finding, a high selectivity factor for the polymerization of L-LA vs. D-LA with (RR/RR)-A (kL/kD ~ 14) was shown, but this only resulted in atactic PLA from rac-LA.40 Figure 2.15 shows the energies for species in the polymerization of rac-LA with (RR/RR)-A as determined by DFT calculations by Maron et al. in 2013.129 The rate-determining step was shown to be the nucleophilic addition of the propagating alkoxide into the carbonyl C-O bond (k1). The energy of a matched pair is stabilized compared to a mismatch (NAL < NAD). However, similar energies were calculated for the ring-opening pathways for both a match and a mismatch (k2) which would result in the generation of atactic PLA from rac-LA despite the high selectivity factor.  61   Figure 2.14 Plot of ln[LA] versus time for polymerization of meso-LA catalysed with a) left, (±)-A () and (RR/RR)-A () and b) right, (±)-B () and (RR/RR)-B().  Reactions were carried out in an NMR tube at 25 °C and followed to 90% conversion. 1,3,5-trimethoxybenzene (TMB) was used as internal standard. All reactions were carried out with 200 eq. of LA in CDCl3 at 25 °C and followed to 90% conversion by 1H NMR spectroscopy [A] = 0.0023 M, [LA] = 0.46 M. [B] = 0.0018 M, [LA] = 0.36 M. The value of kobs was determined from the slope of ln[LA] vs. time, averaged from at least three experiments.   Table 2.3 Rates of polymerization of lactide isomers with various catalysts. Entry Catalyst kobs ( 10−3 s−1) for Psa Pmb meso-LA rac-LA L-LA 1 (±)-A 1.61(0.03)c 1.72(0.16)d 2.98(0.09)d 0.52 0.61 2 (RR/RR)-A 0.35(0.5)c 0.62(0.16)d 3.4(0.6)d 0.73 0.48 3 (SS/SS)-A 0.31(0.03)c 0.70(0.05)d 0.27(0.7)d 0.71 0.49 4 (±)-B 2.52(0.02)e 2.3(0.5)f 2.6(0.5)e 0.36 0.74 5 (RR/RR)-B 2.24(0.01)e 0.46(0.09)f 2.2(0.5)e 0.42 0.77 Values shown in parentheses are standard errors. a Ps values shown for meso-LA. b Pm values shown for rac-LA. c[LA] = 0.46 M. [catalyst] = 0.0023 M. d [LA] = 0.46 M. [catalyst] = 0.0024 M;40 e [LA] = 0.36 M. [catalyst] = 0.0018 M.125 f[LA] = 0.45 M. [catalyst] = 0.0023 M.122 c-f Reactions were carried out in an NMR tube at 25 °C and followed to 90% conversion. 1,3,5-trimethoxybenzene (TMB) was used as internal standard.  The value of kobs was determined from the slope of ln[LA] vs. time, averaged from three experiments.   -2-1.5-1-0.500.511.50 2000 4000 6000ln[LA]time (s)(±)-A(RR/RR)-A-2-1.5-1-0.500.511.520 500 1000 1500 2000 2500ln[LA]time (s)±-B(RR/RR)-B62   Figure 2.15 Mechanism for ring-opening of L- or D-LA by (RR/RR)-A.129 63   The kobs value for polymerization of meso-LA with (RR/RR)-A is over five times slower than its racemic analogue, (±)-A (Figure 2.14, Table 2.3, entries 1-2). The smaller kobs value observed for (RR/RR)-A translates to a greater syndiotacticity for the enantiopure catalyst, contrasting the observed behaviour towards rac-LA. Based on the analysis of the theoretical calculations discussed above, this observation implies that: a) kinetic preference for the different acyl-oxygen sites in meso-LA is either more significant that than between D- and L-LA, giving rise to different energies for nucleophilic addition products or b) meso-LA has a significantly different energy pathway for ring-opening compared to the other isomers. Either of these deductions could account for the difference in stereoselectivity between lactide isomers.  Catalysts (±)- and (RR/RR)-B exhibit higher rates for the ROP of meso-LA than (±)- and (RR/RR)-A, which is consistent with their relative rates of polymerization of rac-LA (Figure 2.14). (RR/RR)-B shows a kL/kD of 5, much lower than that of (RR/RR)-A, nonetheless this krel translates to moderate isoselectivites and heteroselectivities (Table 2.3).40,125 DFT studies of this catalyst system are not available but the similarities in stereoselectivity between isomers may be due to the dissociation of the pre-catalyst dimer in the presence of monomer coupled with having a flexible ligand which may promote the CEC mechanism.127  2.3 Conclusions  Two series of dinuclear indium coordination complexes were employed to investigate the catalytic ring-opening polymerization of meso-lactide. Catalyst A showed good molecular weight control for the ring-opening polymerization of meso-LA. Polymerization of meso-LA with (RR/RR)-A generated PLA with the highest syndiotacticity of all complexes in this study. Defect peak analysis showed that enantiomeric site control is the main mechanism acting during ring-64  opening polymerization. Catalyst B showed less control over molecular weight and displayed similar levels of iso- and heterotactic control, likely due to dissociation of the pre-catalyst dimer during initiation. Correlation of polymerization rate to syndioselectivity was attempted, but the interplay between enantiomorphic site control, chain-end control, kinetic and thermodynamic influences during polymerization make this relationship non-trivial.  2.4 Experimental  General considerations. All air and moisture sensitive manipulations were carried out in an MBraun glovebox or using standard Schlenk line techniques.  A Bruker Avance 300 or 400 MHz spectrometer was used to record 1H NMR spectra. A Bruker Avance 600 MHz spectrometer was used to acquire homonuclear decoupled 1H{1H} NMR spectra of PLA. 1H NMR chemical shifts are given in ppm versus residual protons in deuterated CDCl3 (δ 7.27).  Molecular weights, hydrodynamic radii and intrinsic viscosities were determined by GPC-MALS-RI-Viscometer using an Agilent liquid chromatograph equipped with a Agilant 1200 series pump and autosampler, three Phenogel 5 μm Narrow Bore columns (4.6 × 300 mm with 500 Å, 103 Å and 104 Å pore size), a Wyatt Optilab differential refractometer, Wyatt tristar miniDAWN (laser light scattering detector) and a Wyatt ViscoStar viscometer. The column temperature was set at 40 °C. A flow rate of 0.5 mL/min was used and samples were dissolved in THF (ca. 4 mg/mL). The measurements were carried out at laser wavelength of 690 nm at 25 °C. Data were processed by ASTRA software (Wyatt Technology).  Materials. THF was taken from an IT Inc. solvent purification system with activated alumina columns and degassed before use. HPLC grade DCM was purchased from Fisher Chemicals and 65  was dried over CaH2, transferred under vacuum, and degassed before use. CDCl3 was purchased from Cambridge Isotope Laboratories Inc. and dried over CaH2, transferred under vacuum, and degassed via three freeze−pump−thaw cycles before use. 2-Propanol was purchased from Fisher Chemicals and was dried over 3 Å molecular sieves. Indium(III)trichloride was purchased from Sigma-Aldrich and was used as received. Rac-LA and meso-LA were a gift from PURAC America Inc. Rac-LA was recrystallized thrice from hot dried toluene. Racemic and enantiopure complexes A and B were prepared according previously reported methods.40,125   Purification of meso-lactide. Meso-LA (50 g) was fully dissolved in DCM; magnesium sulphate (anhydrous) was added to the solution, which was then stirred for 10 min. The excess solid was removed from white suspension via vacuum filtration and the solvent removed from the filtrate under reduced pressure to yield a white crystalline powder.  This powder was recrystallized from dry 2-propanol five to seven times, to yield meso-LA with >95% purity.   Polymerization of meso-LA with catalyst A. A 20 mL scintillation vial was equipped with a magnetic stir bar. Meso-lactide (290 mg, 2.7 mmol) was dissolved in 2 mL of DCM and to this, a solution of the catalyst (2.0 mg, 0.002 mmol) in 2 mL of DCM was added. The solution was stirred at room temperature for 16 h.  Solvent was removed under reduced pressure and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in minimal DCM and precipitated using excess of cold methanol at least three times to remove residual catalyst. The polymer sample was then dried under vacuum for 48 h.   Polymerization of meso-LA with catalyst B. A 20 mL scintillation vial was equipped with a magnetic stir bar. Meso-lactide (620 mg, 4.3 mmol) was dissolved in 2 mL of DCM and to this, a solution of the catalyst (2.0 mg, 0.001 mmol) in 2 mL of DCM was added.  The solution was 66  stirred at room temperature for 4 h to minimise transesterification and back-biting reactions. Solvent was removed under reduced pressure and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in minimal DCM and precipitated using excess of cold methanol at least three times to remove residual catalyst. The polymer sample was then dried under vacuum for 48 h.   In situ observation of ROP of meso-LA. All samples for NMR scale polymerization were prepared in Young’s’ cap sealed NMR tubes under an N2 atmosphere. The NMR tube was charged with a stock solution of catalyst in CDCl3 (0.25 mL, 0.001 mmol) and frozen in liq. N2. CDCl3 (0.25 mL) was added and again frozen to create a barrier between the catalyst and the monomer layers. Finally, a stock solution of meso-LA (0.50 mL, 0.45 mmol) and the internal standard 1,3,5-trimethoxybenzene (5 mg, 0.03 mmol per 0.50 mL) in CDCl3 was added and frozen.  The sealed and evacuated NMR tube was immediately taken to the NMR spectrometer (400 MHz Avance Bruker Spectrometer). The sealed tubes were thawed and the polymerization monitored at 25 °C until greater than 90% conversion.  67  Chapter 3: Comparison of the rheological properties of isotactic, syndiotactic, and heterotactic poly(lactide)s.§§  3.1  Introduction  In the previous chapter, tri- and tetradentate indium complexes were used for the stereocontrolled ring-opening polymerization of lactide isomers to achieve stereoregular poly(lactide)s. The enantiopure tridentate catalyst showed moderate syndioselectivity towards meso-lactide, acting primarily through enantiomorphic site control. On the other hand, the tetradentate catalysts showed heteroselectivity by acting through chain-end control. The effect of stereoregularity on chain-interactions and polymer properties will be further probed in the following sections.  3.1.1 Influence of polymer tacticity on chain packing and thermal properties  Stereoselective catalyst systems with good molecular weight control have allowed researchers to the investigate the effect of chain-packing and crystallinity on the properties of stereoregular polymers.130 Four polymers that form microstructures similar to those exhibited by poly(lactide) are: poly(propylene) (PP), poly(styrene) (PS), poly(methyl methacrylate) (PMMA) and poly(hydroxybutyrate) (PHB) (Chart 3.1).                                                  §§ This work has been published in the journal Macromolecules: Chile, L.-E.; Mehrkhodavandi, P.; Hatzikiriakos, S. G.; Macromolecules 2016, 49, 909. (doi: 10.1021/acs.macromol.5b02568). 68  Chart 3.1 Common polyolefins that form polymers with various microstructures.     Syndiotactic polyolefins are reported to have two major chain conformations, helical and zig-zag planar (Figure 3.1).131-134 This chain packing forms crystalline regions within the polymer matrix which impart greater strength135 and elasticity136,137 with respect to their isotactic counterparts. Thermal transitions such as melting temperature, Tm and glass transition temperature, Tg are dependent on the presence of regions of order and thus are strongly influenced by microstructure.       Figure 3.1 Chain packing conformations reported for syndiotactic polyolefins.   Higher thermal transition temperatures for s-PMMA138 and s-PS135 compared to i-PMMA and i-PS were ascribed to strong intermolecular bonding from favourable electronic interactions between side groups on adjacent chains (Table 3.1).13669  Table 3.1 Trends in thermal transitions for different microstructures of three common polyolefins (all values in °C).   PP PSc PMMAd PHB e  Isotactic Syndiotactic Isotactic Syndiotactic Isotactic Syndiotactic Isotactic Syndiotactic Tg 4.5a 6.7a 99 100 61 131 -- -- Tm 151b 139b 240 270 -- -- 183 155 a Mülhaupt 1997139 b Ward 1995136 c Malanga 1998102 d Schultz 1998140 e Ebrahimi 2016137   Reports of syndiotactic PLA (s-PLA) are scarce in the literature (see 2.1.4 above) with only two groups reporting melting temperatures for highly syndiotactic polymers.48,103 Although highly heteroselective catalysts have been investigated,107,114,115,117,120,121,141-156 there has only been one report of the thermal properties for heterotactic PLAs (h-PLA).150   3.1.2 Polymer tacticity and its influence on rheological properties.  Rheological measurements on polymer melts are useful for characterization due to the strong correlation between viscoelastic behaviour and molecular structure. Stereoregularity is expected to influence the entanglement of polymer chains and consequently rheological parameters such as zero-shear viscosity, 0, plateau modulus, GN0 and activation of flow energy (Ea,flow).  An early report by Weese et al. assessed the melt rheology of stereoregular PMMA and showed that increasing the syn-content of the polymer enhanced the zero-shear viscosity and plateau modulus compared to their isotactic counterparts.157 Similar enhancements were shown for s-PP by Eckstien et al. in 1997139 and 1998158 whereby s-PP had an increased Ea,flow compared to i-PP (55 kJ mol-1 vs. 15 kJ mol-1) and showed zero-shear viscosity 10 times higher than i-PP. 70   A more recent report comparing rheological properties of syndiotactic and isotactic poly(styrene) (PS)159 showed that entanglement molecular weight, Me is dependent on tacticity, with s-PS having the lowest Me (14500 g mol-1), therefore giving rise to the highest zero-shear viscosity and energy barrier to flow. These results were in line with earlier studies exploring the viscoelastic properties of poly(propylene) homo-160 and copolymers.161 These accounts concluded that syndiotactic polymers exhibit higher activation energies of flow, Ea,flow compared to their isotactic and atactic counterparts, implying a higher sensitivity of rheological properties on temperature. Tacticity dependence on rheological properties was also found in solutions of poly(vinyl alcohol),162 indicating that increasing degrees of syndiotacticity enhances chain stiffness as well as the entanglement density of the polymer chain.  Poly(hydroxybutyrate) (PHB) is another polyester capable of forming polymers with various microstructures. Recently Ebrahimi et al. reported the effect of tacticity on the thermorheological properties of PHBs, showing that the large Me for i-PHBs translates to a low melt strength while syndiotactically enriched PHBs have a much smaller Me and larger entanglement density, which imparts the material with strength and elasticity.137  3.1.3 Rheological properties of poly(lactide)s.  Comprehensive studies of the viscoelastic behaviours and mechanical properties of isotactic, isotactic stereocomplex, and atactic PLA have been reported.21,22,163-166 Isotactic PLA (i-PLA) has a highly-ordered microstructure, with GN0, ranging around 0.9 ± 0.2 MPa, and Me of approximately 4400 g mol-1. These properties impart i-PLAs with chain stiffness and strength but leave the material brittle, with an elongation at break of just 3%. 71   There are no rheological and mechanistic studies of the useful processing properties of syndiotactic or heterotactic PLAs. In the following sections, stereo-selective systems are used to produce PLAs with varying levels of stereoregular microstructures with the aim of correlating tacticity with the thermorheological behaviour of the materials.   3.2 Results and discussion 3.2.1 Synthesis of PLAs with various microstructures.  Families of atactic, syndiotactic, isotactic and heterotactic PLAs with varying molecular weights were synthesized and rheological studies carried out to generate a comprehensive correlation between polymer microstructures determined by in situ NMR spectroscopic methods (Pm / Ps)*** to bulk polymer properties (Table 3.2). The rheological properties of isotactic and isotactically enriched PLA have been thoroughly investigated,15,16,159-171 thus these samples were only used for comparisons.    The current study compares: a) syndiorich PLAs (s-PLA) generated from polymerization of meso-LA with (RR/RR)-[(NNO)InCl]2(μ-Cl)(μ-OEt)] ((RR/RR)-A); b) isorich gradient PLAs (i-PLA) generated from polymerization of rac-LA with catalyst (RR/RR)-[(ONNO)In(μ-OEt)]2 ((RR/RR)-B); c) essentially atactic PLA (a-PLA) (with a slight heterotactic bias, Pr = 0.60) generated from polymerization of rac-LA with tin octanoate; and d) highly heterotactic PLA (h-PLA) generated from polymerization of rac-LA with an InCl3/amine/alcohol system (Table 3.2).120,121 The polymers are assigned labels based on their molecular weight and tacticity; for                                                  *** See Chapter 2. 72  example, 155-het-60 denotes heterotactic polymers with Mn of 155 kg mol–1 and Pr of 0.60. All polymers were isolated by precipitating thrice from DCM with cold MeOH before being dried under vacuum for two days to remove any trace amounts of solvent. Subsequently, polymers were compression molded to a width of 0.4-0.6 mm at 150 °C for rheological measurements.    Synthesis of highly syndiotactic PLA with an enantiopure aluminum salen binaphthylamine isopropoxide complex (I) was attempted (Figure 3.2).39,103 These catalysts form highly syndiotactic PLA at low molecular weights (Mn ~ 12 kg mol-1). To achieve high molecular weights required for rheological measurements (Mn > 100 kg mol-1), higher monomer loading and longer reaction times were required. Under these conditions a decrease in syndio-enrichment was observed (Figure 3.3, Ps = 0.77). We attribute this decrease in syndiotactic control to the greater extent of polymer exchange which has been reported to occur at long reaction times.106   Figure 3.2 (-)-Salen binathphylamine alkoxide complex (I) reported to be highly syndioselective towards meso-LA by Coates et al. in 1999.103    73   Figure 3.3 1H {1H} NMR spectrum of polymer produced from reaction of (-)-salen binaphthylamine alkoxide complex (I) with meso-lactide. Using Equation 2.2, the polymer tacticity was calculated to be Ps = 0.77. 74  Table 3.2 Polymerization data for various microstructured PLAs used for themorheological experiments.  entry catalyst sample Mn,GPC (g mol-1)a tacticityb Ða  1 c Sn(Oct)2 50-het-60 49900 Pr = 0.56 1.11 Atactic PLAs (het-60) 2 c Sn(Oct)2 155-het-60 155000 Pr = 0.60 1.12 3 c Sn(Oct)2 249-het-60 249000 Pr = 0.57 1.14 4 c Sn(Oct)2 352-het-60 353000 Pr = 0.55 1.10 5 c InCl3/NEt3/BnOH 148-het-96 148000 Pr = 0.95 1.04 Heterotactic PLAs (het-96) 6 c InCl3/NEt3/BnOH  182-het-96 182000 Pr = 0.96 1.05 7 c  InCl3/NEt3/BnOH  255-het-96 255000 Pr = 0.95 1.02 8 c InCl3/NEt3/BnOH  295-het-96 295000 Pr = 0.94 1.05 9 d (RR/RR)-A 130-syn-72 130000 Ps = 0.73 1.01 Syndiotactic PLAs (syn-72) 10 d  (RR/RR)-A 147-syn-72 147000 Ps = 0.72 1.02 11 d  (RR/RR)-A 224-syn-72 224000 Ps = 0.72 1.02 12 d (RR/RR)-A 279-syn-72 279000 Ps = 0.71 1.01 13 d (RR/RR)-B 67-iso-72 66700 Pm = 0.73 1.06 Isotactic PLAs (iso-72) 14 d  (RR/RR)-B 100-iso-72 100000 Pm = 0.72 1.02 15 d (RR/RR)-B 104-iso-72 104000 Pm = 0.72 1.08 16 d (RR/RR)-B 119-iso-72 118600 Pm = 0.71 1.04 a Absolute molecular weights were determined by triple detector GPC (gel permeation chromatography) via Universal Calibration (THF 4 mg mL–1, flow rate = 0.5 mL min–1, dn/dc = 0.044 mL g –1 ). b Calculated from 1H{1H} NMR spectra. c Reactions were carried out in toluene at 70 °C. d Reactions were carried out in DCM at 25 °C.  e 155-het-60 denotes heterotactic polymers with a molecular weight of 155 kg mol–1 and a Pr value of 0.60. 75  3.2.2 Thermal study of stereoregular poly(lactide)s.   Thermogravimetric analysis (TGA) and differential scanning calorimetery (DSC) were used to determine the thermal properties for PLA samples with different degrees of tacticity. Studies on poly(styrene) revealed no evidence that tacticity effects the degradation pathway as these mechanisms are dependent on repeat unit rather than microstructure. However, the authors did note a weak tacticity effect on the thermal stability of PS, where atactic PS displayed the low stability while isotactic PS showed the highest thermal stability.167 Contrary to the studies with PS, the syn-72 polymers show the highest degradation onset temperature, Tonset thus better thermal stability compared to the other microstructures (Table 3.3 and Figure 3.4).    Figure 3.4 TGA heating traces for the various microstructured PLAs. Thermogravimetric analysis was performed on approximately 20 mg of material. Samples were heated to 500 °C at a rate of 20 °C/min to determine the degradation onset temperature (temperature at which there is 5% weight loss), Tonset.  76   Glass transition temperatures, Tg are highly dependent on the presence of strong inter-chain interactions and are thus influenced by tacticity.89,138,140 Based on previous observations of syndiotactic PS and PMMA, it was expected that extended blocks of syndiotactic units would allow the formation of aggregates, increasing the Tg of s-PLA compared to other PLA microstructures. In the current study, all the polymers studied are amorphous and only show glass transitions in their DSC traces (Figure 3.5). In contrast to what was seen for other syndiotactic polymers, a significant effect of tacticity on glass transition temperatures was observed (Table 3.3). Isotactically enriched PLA, iso-72 exhibited the highest Tg of the polymers studied, likely due to the strong interactions between PLLA and PDLA domains within the stereo-gradient polymer backbone (Table 3.3, entry 5).96,99,100 As a comparison, values were obtained for purely isotactic PLA which also displays a high glass transition temperature (Table 3.3, entry 6).29 Syn-72 PLA displayed the lowest Tg (comparable to atactic PLA), indicating lasting chain aggregates are scarce within the polymer matrix. Highly syndiotactic PLA reported by Coates et al. in 1999 shows a Tg lower than that of heterotactic and isotactic PLAs,103 implying that even large syndiotactic aggregates do not impart as much strength as isotactic domains.   Although heterotactic PLAs are prevalent in the literature,36,37 there is only one report of a melting temperature after annealing (Pr = 0.95, Tg = ~46 °C, Tm = ~120 °C).150 Heterotactic PLA, het-96, exhibits the second highest Tg (37.9 °C), indicating that although this microstructure enhances chain interactions, it cannot readily form highly crystalline regions. This may be due to the structure of aggregate domains within h-PLA. However, this microstructure has not been thoroughly explored so there is no information to indicate how these aggregate structures arise.77  Table 3.3 Thermal properties of PLAs of varying stereoregularity.  Entry  Tonset (°C)a Tg (°C)b Tm (°C)b 1 het-60 258(11) 31.9 -- 2 het-96 290(5) 37.9 -- 3 syn-72 323(2)  30.8 -- 4 syn-96c  -- 34 151 5 iso-72 319(6) 46.2 -- 6 iso-100d  310 55 175 Values shown in parentheses are standard deviations. aThermogravimetric analysis was performed on approximately 20 mg of material. Samples were heated to 500 °C at a rate of 20 °C/min to determine the degradation onset temperature (temperature at which there is 5% weight loss, Tonset) bThermal analysis of samples was performed by using a differential scanning calorimeter (DSC) with ca. 2 mg of sample. Samples heated to 170 °C at 10 °C/min and cooled to 25 °C at 5 °C/min to determine Tg and Tm. Glass transition and melting temperatures calculated from second heating scans. cCoates 1999.103  dGarlotta 2001.29  50 100 150-202Heat FlowTemperature (oC) het-60 het-96 syn-72 iso-720 2 4 6 8 100246810 Figure 3.5 DSC traces for various microstructured PLAs studied. Samples heated to 170 °C at 10 °C/min and cooled to 25 °C at 5 °C/min under a nitrogen atmosphere to reduce sample degradation. Glass transition and melting temperatures calculated from second heating scans. ca. 2 mg of sample used.  78  3.2.3 Solution viscosity.   The average values of the intrinsic viscosities, [and the weight-average molecular weight, Mw, obtained from light-scattering GPC analysis are plotted in Figure 3.6. The slope of the linear regression line (exponent of the Mark-Houwink equation) is 0.73, which agrees with the values reported by Dorgan et al. for PLAs in THF.22,163,164 For linear polymers with random coil conformation, the exponent of the Mark-Houwink equation has a value from 0.5 in a poor solvent to 0.8 in a good solvent.138,168 It can also be seen that the intrinsic viscosities of the isotactic PLAs are higher compared to heterotactic PLAs which in turn are higher than those of syndiotactic counterparts with the latter possessing the lowest intrinsic viscosity of all different types of PLAs. This is attributed to the different chain conformations, although the slopes of 0.75-0.76 for all the different families imply good solvent conditions and random coil conformation.138  Mw, gmol-1104 105 106[], mL/g101102103het-60het-96 syn-72 iso-72[]syn-72 = 0.008Mw0.76]iso-72 = 0.019Mw0.75 []het-60 = 0.012Mw0.75 []het-96 = 0.011Mw0.75 Figure 3.6 Log-log plot of intrinsic viscosity of various types of PLAs as a function of the molecular weight.  79   The hydrodynamic radius dependence on the weight-averaged molecular weight for the range of polymers studied is shown in Figure 3.7. The scaling relation for heterotactic (het-96) and atactic (het-60) PLAs is Rh = 0.017Mw0.55, in agreement with trends for amorphous PLAs reported by Othman and co-workers.22 The radii of iso-72 polymers are higher than those expected by this relation while the radii of syn-72 polymers are lower. This result correlates well with the calculated mass dependence of the intrinsic viscosity plotted in Figure 3.6.  Mw (g mol-1)104 105 106Rh100101102iso-72het-60het-96syn-72Othman (2011)Rh = 0.017Mw0.56Current workRh = 0.016Mw0.56 Figure 3.7 Characteristic hydrodynamic radius (Rh) as a function of molecular weight and tacticity. The radii of the iso- PLA are the highest while those of syndio-PLA are the lowest, in agreement with the intrinsic viscosity data plotted in Figure 3.6.   3.2.4 Linear viscoelasticity of PLAs with varying microstructures.    An important aspect of polymer rheology is the study of how viscoelastic behaviour changes with respect to temperature, molecular weight and polymer microstructure. Of particular significance is 80  establishing structure-property relationships i.e. zero-shear viscosity vs. molecular weight, useful also in assessing the processing properties of materials. Currently there are no reports of these properties for heterotactic or syndiotactic PLA.  The thermal stability of stereoregular PLAs under shear stress was probed in isothermal time sweep experiments at 180 °C over 60 minutes. A parallel plate rheometer was used at a constant frequency of 0.5 rad s-1 and strain amplitude of 2%. The complex modulus, |G*| , a measure of resistance to deformation sensitive to structural changes, is plotted against time in Figure 3.8.   Heterotactic (het-96) and atactic (het-60) PLAs show a steady decrease in complex modulus and associated melt strength over the experimental time frame. Subsequent GPC analysis conducted on the collected material showed a 20% decrease in molecular weight consistent with the fundamental relationship of  for linear macrostructures. Isotactically enriched PLA (iso-72) and syndiotactic (syn-72) samples showed stable melt strength over the experimental time frame, however, a small decrease in molecular weight and a broadening of dispersity values was observed after the test.†††                                                   ††† Linear viscoelastic (LVE) measurements last about 20 minutes, during which time the decrease in molecular weight is estimated to be 6%. A molecular weight decrease of this size has a minimal impact on the reported LVE properties. 3.40 wM 81  time (s)101 102 103 104Complex modulus, |G*| (Pa)100101103104105155-het-60182-het-96147-syn-72119-iso-72 Figure 3.8 Complex modulus, |G*| vs. time for stereoregular PLAs.   Dynamic linear viscoelastic measurements were focussed on the comprehensive characterization of syndiotactic and heterotactic PLAs and comparison with isotactic and atactic PLAs.20-22,163-165,169 These measurements were conducted within the linear viscoelastic regime at temperatures in the range of 70-190 °C and at angular frequencies ranging from 0.01 to 100 rad s-1 with a constant strain of 2% and a gap of 0.5 mm was used to minimize edge effects.   The isothermal frequency sweep measurements for each microstructure show an expected decrease in complex viscosity with temperature and the materials display shear-thinning behaviour.  All stereoregular PLAs studied were found to be thermorheologically simple, allowing for the application of the time-Temperature superposition to generate master curves. Figure 3.9 shows a representative master curve for the linear viscoelastic moduli, G’ and G”, as well and the complex viscosity at the reference temperature of 150 °C for het-96 (Figure 3.9a) and syn-72 (Figure 3.9b) polymers. Deformation frequency trends for the loss modulus are very similar for both heterotactic 82  and syndiotactically enriched PLAs where the loss moduli reach a plateau value (plateau modulus) at high frequencies (101 – 106 s-1). Similar trends are observed for the storage modulus which increases with increase of frequency, reaching a maximum and minimum value at higher frequencies typical of linear monodisperse polymers. It can be seen from the characteristic slopes of 1 and 2 for G’ and G”, respectively, that the terminal zone at small deformation frequencies (and long deformation times) has been reached.  The linear viscoelastic moduli of syn-72 and het-96 polymers with varying molecular weights are shown in Figure 3.10 and Figure 3.11, respectively. Trends observed in these plots are in line with those reported for isotactic PLAs,22 whereby, with increasing molecular weight plateau regions in both storage and loss moduli become more defined, occurring over a larger frequency ranges, in addition terminal zones shift to lower frequencies. (Figure 3.10 and Figure 3.11). All these observations are typical of linear monodisperse polymers. 83  Angular frequency (rad s-1)10-2 10-1 100 101 102 103 104 105 106 107 108Storage, loss Modulus, G',G" (Pa)102103104105106107Complex viscosity, |*| (Pa.s)10-1100101102103104105G' G"|*|182-het-96 (a)Slope 1Slope 2Angular frequency,  (rad s-1)10-2 10-1 100 101 102 103 104 105 106 107 108Storage, loss Modulus, G',G" (Pa)102103104105106Complex viscosity, |*| (Pa.s)10-210-1100101102103104105106G'G"|*|147-syn-72 (b)Slope 1Slope 2 Figure 3.9 Master curve of the linear viscoelastic moduli, G’ and G” and complex viscosity, |*|, (Tref = 150 °C) for (a) 182-het-96 (Table 3.2, entry 6) and (b) 148-syn-72 (Table 3.2, entry 10). 84  Angular frequency,  (rad s-1)10-2 10-1 100 101 102 103 104 105 106 107 108Storage modulus, G' (Pa)102103104105106107182 kg mol-1148 kg mol-1255 kg mol-1295 kg mol-1het-96 (a)MwMw  Angular frequency,  (rads-1)10-2 10-1 100 101 102 103 104 105 106 107 108Loss modulus, G" (Pa)103104105106107182 kgmol-1148 kgmol-1255 kgmol-1295 kgmol-1het-96(b)MwMw Angular frequency,  (rad s-1)10-2 10-1 100 101 102 103 104 105 106 107 108Complex viscosity, |*| (Pa.s)10-1100101102103104105106107182 kg mol-1148 kg mol-1255 kg mol-1295 kg mol-1het-96 (c)MwMw Figure 3.10 Master curves of the linear viscoelastic moduli for het-96 polymers (Tref = 150 °C). (a) Loss modulus vs. angular frequency (b) storage modulus vs. angular frequency (c) complex viscosity vs. angular frequency. 85  Angular frequency,  (rad s-1)10-2 10-1 100 101 102 103 104 105 106 107 108Storage Modulus, G' (Pa)101102103104105106107147  kg mol-1130  kg mol-1224  kg mol-1280  kg mol-1syn-72 (a)MwMw  Angular frequency,  (rad s-1)10-2 10-1 100 101 102 103 104 105 106 107 108Loss Modulus, G" (Pa)102103104105106107147 kg mol-1130 kg mol-1224 kg mol-1280 kg mol-1syn-72 (b)MwMw Angular frequency,  (rad s-1)10-3 10-2 10-1 100 101 102 103 104 105 106 107 108Complex viscosity, |*| (Pa.s)10-210-1100101102103104105106147 kg mol-1130 kg mol-1224 kg mol-1280 kg mol-1syn-72 (c)MwMw Figure 3.11 Master curves of the linear viscoelastic moduli for syn-72 polymers (Tref = 150 °C). (a) Loss modulus vs. angular frequency (b) storage modulus vs. angular frequency (c) complex viscosity vs. angular frequency. 86   Reports comparing the viscoelastic behaviours of syndiotactic and isotactic poly(propylene) homo- and copolymers139,161,170 showed that increasing the degree of syndiotacticity increases chain stiffness as well entanglement density, imparting strength to these polymers, observed by an increase in the plateau modulus, GN0.136  The plateau moduli, GN0 for the various microstructures studied were obtained from the minima of van Gurp-Palmen plots (Figure 3.12). From the GN0, molecular weight between entanglements, Me, can be calculated using: 𝑀𝑒 =  𝜌𝑅𝑇𝐺𝑁0         (1) where  is the melt density (g/cm3) at the chosen reference temperature, obtained using the relationship:22  𝜌(𝑡) =  1.2836 𝑒(−7.7104𝑇)     (2)  Plateau moduli for PLAs with varying levels of isotacticity calculated by Dorgan and co-workers fell within the range of 1.0 MPa ± 0.2.164 Moderately isotactic PLA, iso-72, agrees with these values (GN0= 0.97 MPa). In contrast to work on s-PS, moderately syndiotactic PLA displaying the lowest plateau modulus and largest Me exhibits lower chain stiffness compared to heterotactic and atactic polymers.159 Heterotactic and atactic polymers display similar plateau moduli and Me thus are expected to have comparable chain stiffness (Table 3.4). 87  Complex modulus, |G*| (Pa)103 104 105 106 107phase angle,( o )020406080155-het-60182-het-96147-syn-72119-iso-72 Figure 3.12 Van Gurp-Palmen plots of the various microstructured PLAs.   Figure 3.13 The horizontal shift factors, aT for PLAs listed in Table 3.4 at the reference temperature of 150 °C. Lines represents single fitting of the data used to calculate flow activation energy, Ea,flow.   1.00E-021.00E-011.00E+001.00E+011.00E+021.00E+031.00E+041.00E+051.00E+061.00E+070.0020 0.0025 0.0030Horizaontal shift factor, aT1/Temperature (K-1)het-60het-96syn-72iso-7288  Table 3.4 Comparison of rheological properties of PLA, PS and PP in this work and obtained from the literature.   entry polymer 𝑮𝑵𝟎 (105 Pa)a Me (g mol-1)b Ea (kJ mol-1)c 1 het-60 0.58 6900 152.7 (0.5) 2 het-96 0.52 7700 141.0 (0.9) 3 syn-72 0.34 11800 130.2 (0.6) 4 iso-72 0.97 4100 167.9 (0.6) 5d a-PP 0.42 7050 -- 6 s-PP 0.87 e 3400 e 50.6 d 7d i-PP 0.43 6900 38.7 8f a-PS 0.24 17900 97(9) 9f s-PS 0.30 14500 53 (5) 10f i-PS 0.16 27200 99 Values shown in parentheses are standard errors. a Calculated from minimum of van Gurp-Palmen plot (Figure 3.12). b Obtained from equation 1 and 2 c Calculated from the slope of the horizontal shift factor, aT vs. 1/Temperature plot (Figure 3.13) d Jing 2004.160 e Hadjichristidis 1992.171 f  Wang 2011.159 To obtain the GN0 and Me the reference used for the PS and PP samples are 280 °C and 190 °C, respectively.    Complex viscosities for the various tacticity polymers studied at the reference temperature of 150 °C are shown in Figure 3.14. This plot shows that the terminal zone has been reached for all different microstructures and thus zero shear viscosities can be directly determined from the experimental data.   Previous work in our group on linear PLAs with varying levels of isotacticity showed that the zero-shear viscosities (0) of PLAs scale to 3.4 power with Mw.22 In a subsequent report, it was also shown that PLA enantiomeric diblock copolymers also follow this power law;166 both of these results are in agreement with pioneering work carried out by Dorgan et al. in 2005.164 Figure 3.15 shows a plot of zero-shear viscosity, 0 against weight averaged molecular weight, Mw for the various microstructures studied. The trend lines drawn in Figure 3.15 for the iso- and syndio- PLAs 89  are best fit with a fixed exponent of 3.4 that is well established for the atactic PLAs. Although the effect of tacticity on the zero-shear viscosity might be smaller than the experimental error and thus difficult to conclude from the graph, the isotactic and heterotactic polymers appear to have a higher zero-shear viscosity compared to syndiotactic polymers, in agreement with the intrinsic viscosity comparison presented in Figure 3.6.  Angular frequency, , (rads-1)10-2 10-1 100 101 102 103 104 105 106 107 108 109Complex viscosity, |*| (Pa.s)10-210-1100101102103104105155-het-60182-het-96147-syn-72119-iso-72 Figure 3.14 Complex viscosity vs. angular frequency master curves for the different microstructure PLAs at approximately a fixed molecular weight value (Tref = 150 °C).   90  Mw (g mol-1) 104 105 106103104105106107het-60het-96syn-72iso-72Othman 2011log(0) = -13.2 + 3.4 log(Mw)Current worklog(0) = -15 + 3.6 log(Mw) Figure 3.15 Zero-shear viscosity, 0, versus weight averaged molecular weight for PLAs with different microstructures.   Horizontal shift factors, aT determined from the time-temperature superposition were found to agree with those reported in the literature166 and followed an Arrhenius trend, aT  exp{Ea/R(1/T – 1/Tref)}, where Ea is the activation energy for flow, R is the universal constant of the ideal gas law, and Tref is the reference temperature (Figure 3.13). The obtained aT values were plotted against the reciprocal of temperature and this was used to calculate the flow activation energy for all the microstructures studied (Table 3.4).  The activation barrier of flow is highest for the isotactically enriched PLA (Table 3.4, entry 4). This is evidence that strong chain interactions between regions of L- and D-units form aggregates, which need more energy to break apart and allow the material to flow. The small flow activation energy displayed by the syndiotactically enriched material (Table 3.4, entry 3) can be an indication that lasting aggregates are not formed between the syndiotactic units within the 91  polymer chains. The high energy barrier observed for atactic PLA can be explained by the statistical distribution of the different tacticities within its structure (Table 3.4, entry 1). Depending on the synthetic conditions used to make this material, it may contain enough regions of iso- or hetero-enrichment to form transient aggregates, resulting in higher activation energy compared to the heterotactic PLA (Table 3.4, entry 2).   Relevant properties for PS and PP (similar polymers to PLA in application scope) obtained from the literature (Table 3.4) are shown in an attempt to compare to the present results, particularly for the effect of tacticity on entanglement molecular weight, molecular stiffness and plateau modulus. Poly(propylene) has a small methyl side group, similar to PLA, and the reported materials display comparable molecular weights between entanglements than the various PLAs apart from s-PLA. However, they (PPs) display significantly lower activation flow energies due to the absence of interactions i.e. PLLA with PLDA. Conversely, PS with a bulky phenyl side ring is shown to exhibit entanglement molecular weights comparable only with s-PLA, yet lower activation of flow energies are also observed. Dipole-dipole and hydrogen-bonding are the dominating interactions which govern chain interactions within PLA. These results indicate the large effect of hydrogen-bonding on the viscosity of PLA compared to commodity polyolefins.  3.2.5 Uniaxial extensional rheology    Poor melt strength is one of the factors limiting the application of PLA as a replacement for commodity plastics like poly(ethylene terephthalate) and poly(styrene). Early work by Green and Tobolosky describes the motion of highly entangled polymers as a system of transient interconnections extending throughout the polymer.172 Poly(lactide) with its ability for strong 92  hydrogen-bonding interactions between regions of stereoregularity within polymer chains is a good example of this type of system (Figure 3.16). The extensional viscosity of polymeric materials is greatly influenced by the presence of chain interactions at different chain scales, thus can be used as a measure to probe the effect of microstructure on the melt properties of PLA.    Figure 3.16 Illustration of micro-crystallites forming transient cross-links between polymer chains.   Strain-hardening for short linear polymers is unusual. Palade and Dorgan however, showed that high molecular weight isotactic PLA does exhibit strain-hardening at relatively low strain rates173 and in 2004, Yamane reported that the formation of stereocomplex micro-crystallites in PLLA/PDLA blends results in temporary cross-links between polymer chains allowing for strain-hardening.24 In addition, the time scales of the deformation compared with the terminal relaxation times of the polymers might also favor the occurrence of strain-hardening. Extensional tests at low temperatures are useful in illustrating the effect of microstructure on the melt strength of PLA as any strain-hardening will be amplified. Region of stereoregularity forming micro-crystallite Free chain Tie chain 93   Uniaxial extensional tests were carried out at 70, 90 and 110 °C at Hencky strain rates of 0.01 - 10 s-1. Strain-hardening was most significantly observed at strain rates of 10 s-1. Figure 3.17 shows the log-log plots of extensional viscosity vs. time at the three temperatures.‡‡‡ Only syn-72 PLAs do not show strain-hardening at 70 °C, indicating a lack of significant chain interactions to increase chain entanglement within the material. Both atactic and h-PLAs have enough internal order for transient aggregates to form due to intermolecular interactions and thus exhibit strain-hardening (Figure 3.17a). As the temperature is increased to 90 °C, extra thermal energy affords polymer chains more free volume in which to move. At this temperature atactic PLAs no longer exhibit strain-hardening but, isotactic and heterotactic-PLAs retain enough of their aggregate and these chain interactions manifest as strain-hardening (Figure 3.17b). At 110 °C only the strong interactions between L- and D-regions within the isotactically enriched materials remain; all other PLAs have lost their internal order (Figure 3.17c).   Strain-hardening behaviour can be observed when the molecular relaxation time exceeds the characteristic time of deformation.21,165 For a Hencky strain rate of 10 s−1, the characteristic deformation time is 0.1 s. The average relaxation times for each of the polymers studied were calculated from the linear viscoelastic experiments (LVE) experiments and are shown in Table 3.5. It is seen that the average relaxation times are high at 70 °C and 90 °C (5-10 s) explaining the presence of strain-hardening at 10 s-1 (with characteristic time of deformation of 0.1 s). Another contributing factor to strain-hardening is the interactions between L-and D-regions with formations of stereocomplex aggregates. It seems that these interactions are stronger in the case of isotactic                                                  ‡‡‡ Uniaxial extensional experiments were unable to be performed on iso-72 at 70 ˚C due to material constraints. 94  PLAs, and thus, strain-hardening persists even at temperatures of 110 °C at all extensional rates (Figure 3.18). This is not the case for the other types of stereoregular PLAs where at 110 °C there is no observable strain-hardening (see appendix A).   Table 3.5 Average relaxation times for various polymers studied calculated from linear viscoelastic experiments.  entry  ave (s)  70 °C 90 °C 110 °C 1 het-60 9.97 4.72 2.79 2 het-96 7.29 4.65 1.48 3 syn-72 9.46 7.26 2.70 4 iso-72 -- 7.22 2.97  95  time(s)10-2 10-1 100E+ (Pa.s)104105106107het-60het-96 syn-723+(a)H = 10 s-170 °C.time(s)10-2 10-1 100E+ (Pa.s)105106107het-60het-96 syn-72iso-723+(b)H = 10 s-190 °Ctime(s)10-2 10-1 100E+ (Pa.s)104105106107het-60het-96 syn-72iso-723+(c)H = 10 s-1110 °C Figure 3.17 Elongational viscosity as a function of time at Hencky strain rate of 10 s-1 for differently microstructured poly(lactide)s at (a) 70 °C, (b) 90 °C and (c) 110 °C. Uniaxial elongation experiments on iso-72 were unable to be performed at 70 °C. 96  10-210-1100101102103104105106107108119-iso-72110 oCH = 0.01 s-1 H = 0.1 s-1 H = 1.0 s-1 H = 10 s-1 3++ E (Pa.s)time (s)0.01 0.1 1 10 100 1000104105106107108 Figure 3.18 Elongational viscosity as a function of time at Hencky strain rates from 0.01 to 10 s-1 for iso-enriched PLA (Table 3, entry 4).  It is noted that strain-hardening is present at all Hencky strain rates (not the case for all other PLA types) indicating the stronger interaction of L- and D- regions.   3.3 Conclusions  Known stereoselective systems were used to produce PLAs with varying stereoregularity and the relationship between tacticity and the thermal and rheological behaviour of these materials was investigated. Comparisons of thermal properties showed that to achieve large aggregate regions, chain packing from isotactic and heterotactic domains are superior to those of syndiotactic domains, as evidenced by the different glass transition temperatures. The linear viscoelastic properties of PLA showed greater dependence on the level of isotactic rather than syndiotactic linkages present in the polymer chain, with isotactically enriched polymers having higher flow activation energy than syndiotactically enriched polymers. Heterotactic PLAs (h-PLA) exhibited 97  nascent aggregate domains which were observed by an increase in flow activation energy compared to syndiotactic PLA (s-PLA).    Plateau moduli were determined from van Gurp-Palmen plots and the influence on molecular weight between entanglements, Me, was discussed. Our results showed that the measured Me is lowest for i-PLA and highest for s-PLA. Syndiotactic PLAs possess the lowest entanglement density for a given Mw (highest Me), which gives rise to the lowest zero-shear viscosity. In addition, h-PLA and i-PLA have higher Ea,flow, than s-PLA. This is consistent with h- and i-PLA having larger aggregate regions than s-PLA. These results are further confirmed by the solution and melt viscosity data.   Furthermore, it was determined that strain-hardening was possible for h- and i-PLA, showing that transient aggregates due to enhanced L- and D- interactions, more common in isotactically and heterotactically enriched polymers compared to syndiotactic polymers, can increase the relaxation times. As a result, strain-hardening was present for i–PLA even at temperatures as high as 110 °C.  Academic researchers in the organometallic chemistry literature have been basing their expectations of the properties of PLA on studies of polyolefins capable of forming different microstructures. This study on PLAs with similar iso- and syndiotactic enrichment show that this assumption is incorrect and that syndiotactic PLAs do not have enhanced properties compared to their isotactic counterparts.   There may be an enhancement in the properties of highly syndiotactic PLAs, but considering the minute number of reports of these materials, achieving highly syndiotactic PLAs from meso-lactide does not seem favourable compared to forming heterotactic PLAs. 98  3.4 Experimental  General Methods. All air and moisture sensitive manipulations were carried out in an MBraun glovebox or using standard Schlenk line techniques. A Bruker Avance 300 or 400 MHz spectrometer was used to record 1H NMR spectra.  A Bruker Avance 600 MHz spectrometer was used to acquire homonuclear decoupled 1H{1H} NMR spectra of PLA. 1H NMR chemical shifts are given in ppm versus residual protons in deuterated CDCl3 (δ 7.27). Molecular weights, hydrodynamic radii and intrinsic viscosities were determined by GPC-MALS-RI-Viscometer using an Agilent liquid chromatograph equipped with a Agilant 1200 series pump and autosampler, three Phenogel 5 μm Narrow Bore columns (4.6 × 300 mm with 500 Å, 103 Å and 104 Å pore size), a Wyatt Optilab differential refractometer, Wyatt tristar miniDAWN (laser light scattering detector) and a Wyatt ViscoStar viscometer. The column temperature was set at 40 °C. A flow rate of 0.5 mL/min was used and samples were dissolved in THF (ca. 4 mg/mL). The measurements were carried out at laser wavelength of 690 nm at 25 °C. Data were processed by ASTRA software (Wyatt Technology). Molecular masses were determined using a Bruker Autoflex time-of-flight mass (TOF) spectrometer equipped with MALDI ion source. A differential scanning calorimeter (DSC) Q1000 (TA Instruments) was employed to measure the glass transition (Tg) and melting (Tm) temperatures. Thermogravimetric analysis (TGA) traces were collected on a PerkinElmer Pyris 6 TGA with a nitrogen flow rate of 20 mL/min.  Shear measurements were performed using a MCR 501 rheometer equipped with 8 mm parallel plates. Uniaxial extensional measurements were performed using the SER-2 extensional fixture attached to an Anton Paar MCR 502 rheometer.  Materials. THF was taken from an IT Inc. solvent purification system with activated alumina columns and degassed before use. HPLC grade DCM was purchased from Fisher Chemicals and 99  was dried over CaH2, transferred under vacuum and degassed before use. CDCl3 was purchased from Cambridge Isotope Laboratories Inc. dried over CaH2, transferred under vacuum and degassed through three freeze−pump−thaw cycles before use. 2-Propanol was purchased from Fisher Chemicals and was dried over 3Å molecular sieves before use.  Indium (III) trichloride was purchased from Sigma-Aldrich and was used as received. Triethylamine and benzylalcohol were purchased from Sigma-Aldrich and were dried over CaH2, transferred under vacuum and degassed before use. Rac-LA and meso-LA were a gift from PURAC America Inc. Rac-LA was recrystallized twice from hot dried toluene prior to use. Racemic and enantiopure complexes A and B were prepared according previously reported methods.40,125   Purification of meso-lactide. Meso-LA (50 g) was fully dissolved in DCM; magnesium sulphate (anhydrous) was added to the solution which was then stirred for 10 min. The excess solid was removed from white suspension via vacuum filtration and the solvent removed from the filtrate under reduced pressure to yield a white crystalline powder.  This powder was recrystallized from dry 2-propanol five to seven times, to yield meso-LA with >95% purity (by GC-MS).  Polymerization of meso-LA with catalyst A. A 20 mL scintillation vial was equipped with a magnetic stir bar. Meso-lactide (290 mg, 2.7 mmol) was dissolved in 2 mL of DCM and to this a solution of the catalyst (2.0 mg, 0.002 mmol) in 2 mL of DCM was added. The solution was stirred at room temperature for 16 h. Solvent was removed under reduced pressure and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in minimal DCM and precipitated using excess of cold methanol at least three times to remove residual catalyst. The polymer sample was then dried under vacuum for 48 h. No stabilizers were added. 100   Polymerization of rac-LA with catalyst B. A 20 mL scintillation vial was equipped with a magnetic stir bar. Rac-lactide (620 mg, 4.3 mmol) was dissolved in 2 mL of DCM and to this, a solution of the catalyst B (2.0 mg, 0.001 mmol) and in 2 mL of DCM was added.  The solution was stirred at room temperature for 4 h to minimise transesterification and back-biting reactions. Solvent was removed under reduced pressure and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in minimal DCM and precipitated using excess of cold methanol at least three times to remove residual catalyst. The polymer sample was then dried under vacuum for 48 h. No stabilizers were added.  Polymerization of meso-LA with aluminum salen binaphthylamine alkoxide catalyst (I). A vacuum adapted flask was equipped with a magnetic stir bar and charged with meso-lactide (1190 mg, 8.2 mmol) and stirred in 4 mL of toluene. To this, a solution of the catalyst (4.4 mg, 0.008 mmol) in 4 mL toluene was added. The reaction mixture was stirred at 80 °C for 20 h under N2. Solvent was removed under reduced pressure and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in minimal DCM and precipitated using excess of cold methanol at least three times to remove residual catalyst. The polymer sample was then dried under vacuum for 48 h.  Synthesis of heterotactic PLA. A vacuum adapted flask was equipped with a magnetic stir bar and was charged with rac-lactide (2.6 g, 0.02 mol) and stirred in 5 mL of toluene.  InCl3 (2 mg, 0.009 mmol), benzyl alcohol (0.9 μL, 0.009 mmol) and triethylamine (2.5 μL, 0.020 mmol) were then added and the volume made up to 15 mL with toluene. The reaction mixture was stirred at 80 °C for 72 h under N2.  Solvent was removed under reduced pressure and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in minimal DCM and 101  precipitated using excess of cold methanol at least three times to remove residual catalyst. The polymer sample was then dried under vacuum for 48 h. No stabilizers were added.  Synthesis of atactic PLA. A vacuum adapted flask was equipped with a magnetic stir bar and charged with rac-lactide (2.0 g, 0.01 mol) and stirred in 5 mL of toluene. Sn(oct)2 (4.5 μL, 0.04 mmol) was then added and the volume made up to 15 mL. The reaction was then stirred at 80 °C for 72 h under N2. Solvent was removed under reduced pressure and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in minimal DCM and precipitated using an excess of cold methanol at least three times to remove residual catalyst. The polymer sample was then dried under vacuum for 48 h. No stabilizers were added.  DSC measurement of polymers. Approximately 2-3 mg of polymer was weighed and sealed in an aluminum pan. Experiments were carried out under a nitrogen atmosphere. The samples were heated at a rate of 10 °C/min from 25 to 170 °C and held isothermally for 5 min to destroy any residual nuclei before cooling at 5 °C/min. The glass transition and melting temperatures were obtained from the second heating sequence, performed at 10 °C/min.  TGA measurement of polymers. Approximately 20 mg of polymer was weighed into a ceramic crucible. Experiments were carried out under a nitrogen atmosphere at a flow rate of 20 mL/min. The samples were heated at a rate of 20 °C/min from 25 to 500 °C.  Linear viscosity measurements. All polymer samples were compression molded at 150 °C into discs with diameters of 16-50 mm and thickness 0.4–0.6 mm. The dynamic linear viscoelastic measurements were carried out within the linear viscoelastic regime at temperatures in the range from 70 to 190 °C. The dynamic measurements were conducted in the range of 0.01–100 rad s-1 at a strain of 2%. A gap of 0.5 mm was used to minimize edge effects and ensure a reasonable aspect 102  ratio of plate radius and gap. Dynamic time sweep measurements were carried out at an angular frequency of 0.5 rad s-1 at 180 °C to examine the thermal stability of the samples. The rheological measurements were performed under nitrogen atmosphere to minimize degradation of the polymer samples during testing.  Uniaxial extensional measurements. Samples with diameters of 16-50 mm and thickness 0.4–0.6 mm were prepared by the same procedure used for shear measurements. Individual polymer specimens were then cut to a width of 1.5 – 3.5 mm. Measurements were conducted at 70, 90 and 110 °C at Hencky shear rates of 0.01 s-1, 0.1 s-1, 1.0 s-1 and 10 s-1.    103  Chapter 4: Aryl initiators for the living ring-opening polymerization of rac-lactide.§§§  4.1 Introduction  The preceding chapters examined the polymerization behaviour of dinuclear indium complexes capable of stereoselective living ring-opening polymerization of cyclic esters (Chapter 2). This was followed by description of structure-property relationships for these materials, established by investigating the effect of tacticity on the rheological properties of poly(lactide) (Chapter 3). The next chapters aim to expand this study by exploring the role of aryloxy initiators on the living and immortal ring-opening polymerization of lactide with the ultimate aim of incorporating complex arylated bioproducts in this process (see Chapter 6).    4.1.1 PLA synthesis via living ring-opening polymerization with a chain transfer agent.  Metal-mediated living ring-opening polymerization (ROP) is characterized by linear growth in molecular weight with respect to conversion.174,175 Figure 4.1 shows the coordination-insertion mechanism for the ROP of lactide with dinuclear indium catalyst [(NNO)InCl]2(μ-Cl)(μ-OEt) (A). In this mechanism, lactide coordinates to one of the metal centres activating the monomer for nucleophilic attack by the bridging alkoxide which inserts into the carbonyl C-O bond. A subsequent ring-opening event generates a propagating species which continues the polymerization. Propagation continues until all monomer is depleted and the catalyst is only                                                  §§§ This work has been published in the journal Dalton Transactions: Chile, L.-E.; Ebrahimi, T.; Wong, A.; Aluthge, D. C.; Hatzikiriakos, S. G.; Mehrkhodavandi, P.; Dalton Trans. 2017, doi:10.1039/c7dt00990a.  104  quenched with the addition of excess alcohol. Systems acting through this mechanism are useful as they generate polymers with low molecular weight distributions and chain ends which remain active, revived with addition of more monomer. Through this route, block co-polymers176 with a variety of structures can be easily synthesized.    Figure 4.1 Coordination-insertion mechanism for the living ROP of lactide with dinuclear indium catalyst (A).   As mentioned previously, poly(lactide) or PLA has emerged as a useful bio-sourced and biodegradable polymer with applications in a wide variety of fields,11,15,17,18,177-181 packaging (films), 3-D printing (filaments), and biomedical (tissue, implants and drug delivery). However, PLA has the drawback of lacking sites to impart extra functionality which could further improve  105  PLA properties and application scope. One efficient method to functionalize PLA that circumvents this issue, first reported by Inoue et al.,182 is living ring-opening polymerization (ROP) with a chain transfer agent (CTA) (also known as immortal ROP (iROP)). In the modified mechanism, a propagating species can undergo reversible chain transfer with a CTA in solution. If this transfer reaction is faster than propagation, the altered propagating species can continue polymerization, incorporating the CTA onto the polymer chain-end (Figure 4.2). 182,183 Along with highly tuneable catalysts,184,185 immortal polymerization has been a powerful methodology for generating multi-functionalized biodegradable polymers with varied architectures.183,186-190  One way to generate architectural density in immortal ROP systems is to include lignocellulosic alcohols. The lignocellulosic biorefinery industry has been expanding75,191,192 and now provides researchers access to a large number of bio-based composite materials through blending and co-polymerization.77,83,193 One interesting case is the use of lignin functionalized PLA194-198 as an alternative to the widely-used phenol formaldehyde resins and adhesives.199 Understanding and controlling the reactivity of phenolic200-202 bioproducts as CTAs in immortal polymerization can expand their application scope (see Chapter 6).200-203   106   Figure 4.2 Coordination-insertion mechanism for the immortal ROP of lactide with dinuclear indium catalyst (A).   4.1.2 Aryloxy initiators in lactide ring-opening polymerization.  The use of aromatic alcohols as CTAs197,204-209 and phenoxy derivatives as initiators112,210-231 for the ROP of lactones has been less explored due to their lower nucleophilicity.211 Byers et al.232 have reported the use both of internal and external phenoxy derivatives with their bis(imido)pyridine iron complexes and found a decreasing trend in activity with the electron  107  withdrawing character of the phenols (Chart 4.1C).  There are few examples of main group complexes with aryloxy-initiator ligands that are active for the ROP of lactones.233-238   The following sections explore the role of aryloxy initiators on the living and immortal polymerization of lactide with the ultimate aim of incorporating complex arylated products in this process. A family of indium complexes with para-substituted phenoxy initiators was synthesized and the effect of these substitutions on the living ring-opening polymerization of lactide investigated.  Immortal ROP of lactide with a range of arylated alcohols as CTAs is explored as are the possible limitations of the system in incorporating more complex species such as diols.   Chart 4.1 Dinuclear indium complexes [(NNO)InCl]2(μ-Cl)(μ-OEt) (A)40,123,239 and bis(imido)pyridine iron complex (C).232   4.2 Results and discussion 4.2.1 Synthesis racemic phenoxy-bridged indium complexes 1-5.  A series of phenoxy bridged complexes, analogous to complex A, was synthesized to investigate the role of phenyl initiators on the ROP of lactide. Reaction of ()-(NNO)InCl2 and KOPhR (PhR = para-C6H4R; R = OMe, Me, H, Br, NO2) forms the asymmetrically bridged complexes [(NNO)InCl]2(μ-Cl)(μ-OPhR) (R = OMe (1), Me (2), H (3), Br (4), NO2 (5)) (Scheme 4.1, Route 1). Alternatively, these complexes can be synthesized by reaction of [(NNO)InCl]2(μ- 108  Cl)(μ-OEt) (A) with a phenol through an exchange reaction (Scheme 4.1, Route 2). In Route 1, reactions involving electron poor phenoxy substituents favour the formation of the mono-phenoxy complexes 4 and 5, which are obtained in higher yields than more electron rich complexes 1 and 2, which readily form bis-phenoxy bridged dinuclear complexes as a by-product . The formation of bis-phenoxy species is evidenced by the presence of extra peaks in the aromatic region (6-8 ppm) of the 1H NMR spectra of the crude products (Figure 4.3). Complexes 1-5 were purified via recrystallization from solutions of CHCl3:pentane (2:3).   Scheme 4.1 Synthesis of dinuclear indium complexes [(NNO)InCl]2(μ-Cl)(μ-OPhR) (R = OMe, Me, H, Br, NO2).   109   Figure 4.3 1H NMR (CDCl3, 25 °C, 400 MHz) spectrum of crude reaction mixture containing (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhOMe) (1) (blue stars). Extra peaks in the aromatic region (6-8 ppm) are assigned to the bis-phenoxy complex (red stars).   4.2.2 Solution structures of phenoxy-bridged indium complexes 1-5.  Complexes 1–5 have been fully characterized by 1H, 13C{1H}, 1H–1H COSY, 1H–13C HSQC NMR spectroscopy and show similar solution structures to complex A and its related compounds77-79 (see appendix B). The 1H NMR spectra of complexes 1–5 are consistent with a mono-phenoxy bridged species with 1 : 1 ratios of the backbone NH-CH2-Ar protons (doublets at approximately 3.7 and 4.7 ppm) and the bridging phenolic protons Ph-H (multiplets at 6-8 ppm). Compared to the parent complex (A), the loss of electron density with the substitution of the ethoxy bridge for  110  a phenoxy bridge causes an upfield shift of resonances on the 1H NMR spectra of complexes 1–5. Diastereotopic splitting of the two doublets at 3-5 ppm for the NH-CH2-Ar protons increases with the electron-donating ability of the bridging phenol (Figure 4.4).    Figure 4.4 1H NMR spectra (CDCl3, 25 °C, 400 MHz) of diastereotopic ligand NH-CH2-Ar protons for (a) [(NNO)InCl]2(μ-OEt)(μ-Cl) (A), (b) [(NNO)InCl]2(μ-OPhOMe)(μ-Cl) (1), (c) [(NNO)InCl]2(μ-OPhMe)(μ-Cl) (2), (d) [(NNO)InCl]2(μ- OPhH)(μ-Cl) (3), (e) [(NNO)InCl]2(μ-OPhBr)(μ-Cl) (4), and (f) [(NNO)InCl]2(μ-OPhNO2)(μ-Cl) (5).    Another notable feature of the 1H NMR spectra of complexes 1-5 is the AB aryl coupling constants (3JH-H) for the Ar-H protons of the ligand which narrow when shifting from alkyl to aryl bridging moieties (2.51 Hz for complex A versus 2.35 Hz for complexes 1–5) (Figure 4.5).   111   Figure 4.5 1H NMR spectra (CDCl3, 25 °C, 400 MHz) bridging phenolic protons for (a) [(NNO)InCl]2)(μ-Cl)(μ-OEt) (A), (b) [(NNO)InCl]2)(μ-Cl)(μ-OPhOMe) (1), (c) [(NNO)InCl]2)(μ-Cl)(μ-OPhMe) (2), (d) [(NNO)InCl]2)(μ-Cl)(μ- OPhH) (3), (e) [(NNO)InCl]2)(μ-Cl)(μ-OPhBr) (4), and (f) [(NNO)InCl]2)(μ-Cl)(μ-OPhNO2) (5).   4.2.3 Solid-state structures of phenoxy-bridged indium complexes 1-5.  Single-crystal x-ray diffraction analysis was conducted to support the solution characterization. Crystals of complexes 2–5 synthesized from racemic ligands were obtained from slow evaporation of toluene solutions at room temperature. All are asymmetric homochiral dinuclear indium complex with (R,R) ligand configuration at each distorted octahedral indium centre. There is a “cis” relationship between the phenoxy groups of the ligand, where they are on  112  the same side of the dimeric structure (Figure 4.6 - Figure 4.9). This arrangement is observed for all related compounds generated from racemic ligands in this family.77-79   Figure 4.6 Molecular structure of [(NNO)InCl]2(μ-Cl)(μ-OPhMe) (2) depicted with ellipsoids at 50% probability (H atoms and all solvent molecules omitted for clarity).    Figure 4.7 Molecular structure of [(NNO)InCl]2(μ-Cl)(μ-OPhH) (3) depicted with ellipsoids at 50% probability (H atoms and all solvent molecules omitted for clarity).    113   Figure 4.8 Molecular structure of [(NNO)InCl]2(μ-Cl)(μ-OPhBr) (4) depicted with ellipsoids at 50% probability (H atoms and all solvent molecules omitted for clarity).   Figure 4.9 Molecular structure of [(NNO)InCl]2(μ-Cl)(μ-OPhNO2) (5) depicted with ellipsoids at 50% probability (H atoms and all solvent molecules omitted for clarity).   114  Table 4.1 Selected solid-state structural data for indium catalysts (A, 2-6).  Complex R= In-O3(Å) In-Cl3 (Å) H-Cl (Å) O1-In-Cl3 (deg) O3-In-O1 (deg) N1-H-Cl2 (deg) A C2H5 2.112(8), 2.129(8) 2.667(3), 2.636(4) 2.445, 2.788 166.8(2), 167.0(2) 93.0, (3), 93.0(3) 163.5, 131.8 2 C6H4Me 2.198(1), 2.174(1) 2.165(4), 2.423(3) 2.657, 2.914 171.7(3), 172.0(3) 94.9(4), 94.7(4) 160.6, 153.0 3 C6H5 2.186(2), 2.174(2) 2.602(9), 2.645(9) 2.729, 2.819 170.6(6), 169.2(6) 93.7(8), 94.4(8) 155.7, 153.8 4 C6H4Br 2.191(3), 2.186(3) 2.620(1), 2.630(1) 2.594, 2.907 170.7(1), 168.6(1) 93.3(1), 92.5(1) 158.2, 154.4 5 C6H4NO2 2.193(4), 2.203(4) 2.613(1), 2.625(1) 2.484, 3,752 165.4(1), 167.7(1) 89.7(1), 77.2(1) 143.9, 139.2 6 C13H11O2Cl2 2.218(2), 2.247(2) 2.633(7), 2.594(7) 2.474, -- 169.3(6), 173.8(6) 95.1(8), 98.9(8) 144.9, -- Errors are shown in parentheses.  Data shown for both halves of non-symmetric structures. a Two molecules are present in the unit cell. Both show the same geometry. 115   Table 4.1 compares selected bond lengths and angles obtained from the X-ray crystal structures, the number in brackets are for the second half of unsymmetrical molecules. The phenoxy-bridged complexes 2–5 show longer In-O3 bonds than the alkoxy-bridged complex A, increasing with electron deficiency of the bridging phenol (2.174 Å (2) vs. 2.203 Å (5)). The para-substituted phenols are weaker nucleophiles compared to ethanol, which leads to longer and weaker bonds to the indium centre. An associated shortening of the In-Cl3 bond is also observed (2.423 Å (2) vs. 2.625 Å (5)) implying that this bond strengthens to alleviate the electron deficiency of the indium metal centre. In addition, with increasing electron deficiency of the phenol, the deviation from octahedral geometry increases: the O1-In-Cl3 angle becoming less linear (171.7° (2) vs. 165.4° (5)) and the O2-In-O3 bond becoming more acute (94.9° (2) vs. 89.7° (5)). This effect culminates in complex 5 having the most distorted octahedral geometry, and showing different reactivity to the other complexes in the series (Figure 4.9).  In previous work from our group,41 it was of interest to probe the effect of internal hydrogen bonding within the complexes on the ring-opening polymerization of lactide. Table 4.1 attempts a comparison of the H1-Cl2 bond lengths and related angles. Parent complex A has a H1-Cl2 distance of 2.445 Å, comparable to that of complex 5 (2.484 Å) and the N1-H1-Cl2 bond angle for complex A is comparable to that of complex 2 (163.5° and 160.6°, respectively) suggesting that hydrogen-bonding is present in these systems. However, there does not appear to be any trend with the electron-withdrawing character of the bridging phenol. The reactivity towards rac-LA for these complexes does not correlate to the trends in hydrogen bonding strength, further supporting an earlier finding from our group showing that although these hydrogen-bonding interactions are important for the stability of the dimeric species, they are unlikely a factor contributing to the reactivity of these systems.41  116  4.2.4 Living ring-opening polymerization (ROP) of lactide using catalysts 1–5.  The efficiency of the different initiators 1-5 for the ring-opening polymerization of up to 2000 equivalents of racemic lactide (rac-LA) was examined at room temperature (Scheme 4.2 and Table 4.2). Phenoxy-bridged complexes display less control over the polymerization of rac-LA compared to catalyst A, giving polymers with higher-than-expected molecular weights at low monomer loadings, which is indicative of a slow initiation step (Figure 4.10). Catalysts with electron rich bridging phenols (1-2) show the best control over molecular weight; conversely, catalysts 4 and 5 with electron deficient bridging phenols give polymers with molecular weights which largely deviate from the theoretical molecular weights. In addition, as expected for the least nucleophilic initiator, the para-nitrophenol bridged complex 5 is significantly less active than the other four phenoxy catalysts, only reaching 10-20% conversion after 3-4 days (Table 4.2, entries 18-22). These observations are in line with those reported for complex B and its analogues, where a similar change in activity was shown as the electron-withdrawing ability of the para substituent was altered.232   Scheme 4.2 Ring-opening polymerization of rac-lactide with complexes 1-5.    117  [LA]:[I]0 200 400 600 800 1000 1200 1400Mn, kg mol-10100200300dispersity1.01.11.21.31.41.5Mn, experimentalMn, theodispersity[LA]:[I]0 200 400 600 800 1000 1200 1400Mn, kg mol-10100200300dispersity1.01.11.21.31.41.5Mn, experimentalMn, theodispersity[LA]:[I]0 200 400 600 800 1000 1200 1400Mn, kg mol-10100200300dispersity1.01.11.21.31.41.5Mn, experimentalMn, theodispersity[LA]:[I]0 200 400 600 800 1000 1200 1400Mn, kg mol-10100200300dispersity1.01.11.21.31.41.5Mn, experimentalMn, theo dispersity[LA]:[I]0 200 400 600 800 1000 1200 1400Mn, kg mol-10100200300dispersity1.01.11.21.31.41.5Mn, experimentalMn, theodispersity[(NNO)InCl]2(-Cl)(-OPhOMe) (1)  a)c)b)d)e)[(NNO)InCl]2(-Cl)(-OPh) (3) [(NNO)InCl]2(-Cl)(-OPhMe) (2) [(NNO)InCl]2(-Cl)(-OPhBr) (4) [(NNO)InCl]2(-Cl)(-OPhNO2) (5)  Figure 4.10 Plots of observed PLA Mn (closed symbols) and dispersity (open symbols) as functions of added rac-LA for a) [(NNO)InCl]2(μ-OPhOMe)(μ-Cl) (1), b) [(NNO)InCl]2(μ-OPhMe)(μ-Cl) (2), c) [(NNO)InCl]2(μ-OPhH)(μ-Cl) (3), d) [(NNO)InCl]2(μ-OPhBr)(μ-Cl) (4), and e) [(NNO)InCl]2(μ-OPhNO2)(μ-Cl) (5). (Mn = number averaged molecular weight).  The line represents theoretical Mn values based on the monomer:initiator ratio at 100% conversion.  All reactions were carried out room temperature in DCM.   118  Table 4.2 Living ring-opening polymerization data of rac-lactide with complexes 1-5.  entry Cat Time (h) Conv. (%)a [M]/[I] Mn,theob Mn,GPCc Đc 1 1 18 78 58 6550 54600 1.02 2 1 18 93 260 35000 80500 1.01 3 1 18 95 560 76400 98700 1.01 4 1 18 97 1100 152000 139000 1.03 5 2 18 86 54 6690 52800 1.04 6 2 18 89 270 34300 87200 1.06 7 2 18 97 500 71200 118000 1.11 8 2 18 97 990 139000 153000 1.08 9 2 18 95 2070 284000 225000 1.01 10 3 4 21 49 1480 6180d na 11 3 4 95 260 35200 119000 1.01 12 3 4 98 510 72200 142000 1.02 13 3 4 98 1020 145000 172000 1.01 14 4 18 12 56 990 23700 1.04 15 4 18 93 270 36700 167000 1.02 16 4 18 97 540 75900 231000 1.01 17 4 18 98 1080 153000 288000 1.01 18 5 72 12 53 930 10100 1.01 19 5 72 8.6 240 3020 13500 1.22 20 5 72 11 490 7600 42800 1.14 21 5 72 17 1000 25000 80900 1.27 22 5 96 23 1010 33400 71600 1.18 Reactions were carried out in DCM, 25 °C. a Monomer conversions determined by 1H NMR spectroscopy. b Calculated from Mn,theo = (144 g mol-1 × conversion × [LA]/[initiator]) + MROPh. c Absolute molecular weights were determined by triple detector GPC (gel permeation chromatography) via Universal Calibration (THF 4 mg mL–1, flow rate = 0.5 mL min–1, dn/dc = 0.044 mL g –1 ). d Molecular weight determined from MALDI-ToF analysis.   To confirm that the phenoxy initiators remain as the polymer chain ends, polymer samples were analysed using 1H NMR spectroscopy and MALDI-ToF mass spectrometry. Chain end analysis of polymer samples generated with catalyst 3 (50 equiv. rac-LA) by 1H NMR spectroscopy shows aryl signals that do not belong to free phenol or to 3 (Figure 4.11). MALDI- 119  ToF analysis revealed a major series of peaks separated by 144 mass units with phenol at the end-group. The minor series is separated by 72 mass units, implying there is some transesterification occurring during the polymerization (Figure 4.12).     Figure 4.11 1H NMR spectrum (CDCl3, 25 °C, 400 MHz) of free phenol (red spectrum) overlaid with 1H NMR spectrum (CDCl3, 25 °C, 400 MHz) of [(NNO)InCl]2(μ-Cl)(μ-OPhH) (3) (green) and 1H NMR spectrum (CDCl3, 25 °C, 400 MHz) of the polymerization reaction mixture (blue, Table 4.2. entry 10).    120  .  Figure 4.12 MALDI-ToF spectrum from polymerization of [LA]:[3] = 49:1 (Table 4.2, entry 10). An = [144.13LA]n + 94.11HOPh + 23Na+    121   Figure 4.13 1H{1H} NMR (CDCl3, 25 °C) spectrum of the methylene region of PLAs generated from a) (±)-[(NNO)InCl]2(μ-Cl) (μ-OEt) (A), b) (±)-[(NNO)InCl]2(μ-Cl) (μ-OPhOMe) (1) (±)-[(NNO)InCl]2(μ-Cl) (μ-OPh) (3).   The tacticity of polymers formed from the aryl initiators can be determined by inspecting the methylene region (3-5 ppm) of the 1H{1H} NMR spectra. From analysis of defect tetrad patterns, polymers generated from both alkyl and aryl initiators showed the same distribution of tetrad peaks implying the same stereocontrol mechanism is acting in all the systems (Figure 4.13).  4.2.5 In situ polymerization studies.   Polymerization of rac-LA using catalysts 1-4 ([LA]:[I] = 50:1) was monitored using 1H NMR spectroscopy at room temperature.  The plots of ln[LA] vs. time for all catalysts show an initiation period followed by a linear propagation period. Initiation times correlate with the electron-deficiency of the bridging phenoxide, thus catalyst 1 displays a shorter initiation period than catalysts 2-4 (Figure 4.14).    122   time (s)0 2000 4000 6000 8000 10000 12000 14000ln[LA]-2.0-1.5-1.0-0.50.00.51.0[(NNO)InCl]2(-Cl)(-OPhOMe) (1) [(NNO)InCl]2(-Cl)(-OPhMe) (2) [(NNO)InCl]2(-Cl)(-OPhH) (3) [(NNO)InCl]2(-Cl)(-OPhBr) (4)  Figure 4.14 Plot of ln[LA] versus time for polymerization of rac-LA catalyzed with 1 () 2 (), 3 () and 4 ().  Reactions were carried out in an NMR tube at 25 °C. 1,3,5-trimethoxybenzene (TMB) was used as internal standard. All reactions were carried out with 50 eq. of LA in CD2Cl2 at 25 °C and followed by 1H NMR spectroscopy. [1] = 0.0039 M, [LA] = 0.19 M. [2] = 0.0043 M, [LA] = 0.21 M. [3] = 0.0037 M, [LA] = 0.22 M. [4] = 0.0041 M, [LA] = 0.22 M. The value of kobs was determined from the slope of ln[LA] vs. time, averaged from at least three experiments.   The estimated rates of propagation, kobs, were determined from the slope of ln[LA] vs. time (averaged from at least three experiments). The observed rates for 1-4 are one order of magnitude slower than that for catalyst A (Table 4.3). The higher-than-expected molecular weights observed in the polymers generated with phenoxy-bridged catalysts is indicative of their incomplete activation.  Previously reported computational studies on the polymerization of lactide isomers with catalyst A showed that the rate determining step is the insertion into the carbonyl C-O bond by the bridging ethoxy moiety.129 The longer initiation period of the phenoxy bridged catalysts can be attributed to the lower nucleophilicity of the phenoxy compared to the ethoxy moieties. This lower  123  nucleophilicity may mean that not all the complex molecules are being activated leading to a lower concentration of the active catalytic species in the reaction medium.   After the initial insertion of the aryloxide, similar propagation species exist in solution. The difference in the observed propagation rates is likely due to different concentrations of active catalyst present that also causes the observed higher-than-expected molecular weights.  Table 4.3 Propagation rates for the ROP of rac-LA with complexes A and 1-4.  entry Catalyst kobs (×10−4 s−1) 1 A 18.9(5) 2 1 8.5(8) 3 2 5.3(2) 4 3 3.3(2) 5 4 1.1(1) Values shown in parentheses are standard errors. Reactions carried out in an NMR tube at 25 °C with 1,3,5-trimethoxybenzene (TMB) as internal standard with 50 equivalents of LA in CD2Cl2 at 25 °C. [A] = 0.0046 M, [LA] = 0.22 M.  [1] = 0.0039 M, [LA] = 0.19 M. [2] = 0.0043 M, [LA] = 0.21 M. [3] = 0.0037 M, [LA] = 0.22 M. [4] = 0.0041 M, [LA] = 0.22 M. The value of kobs was determined from the slope of ln[LA] vs. time, averaged from at least 3 experiments (see appendix B).   4.2.6 Immortal ring-opening polymerization (ROP) of lactide with catalyst A and aromatic alcohols.  One of the main goals of this study was to investigate the use of aromatic alcohols in the immortal ROP of rac-LA. Shifting from alkyl alcohols to using aryl alcohols does have an effect on the initiation step of the living ROP (as shown in 4.2.5). Therefore, it was of interest to examine the influence of phenols on the chain transfer reaction needed to achieve immortal ROP. To probe the kinetics of exchange, the substitution reaction between catalyst A and phenol (HOC6H5), para- 124  methoxyphenol (HOC6H4OMe) and para-bromophenol (HOC6H4Br) was followed by in situ 1H NMR spectroscopy (Scheme 4.3).  Scheme 4.3 Alkoxy-phenoxy exchange reaction with complex A and para-substituted phenol.     A 1:1 mixture of [A]:[aryl alcohol] was dissolved in CDCl3 and the loss of the bridging ethoxy moiety for A was monitored at room temperature for 60 minutes (Figure 4.15). Plots of ln[A] versus time show an initial decrease in the [A] followed by a steady state (Figure 4.16).  The slope of this graph gives an indication of the chain-transfer step (6.8(4) × 10−4  s−1 ([A] + phenol)) which is on the same order as the propagation rate kobs, for catalysts 1-4.    125   Figure 4.15 1H NMR (CDCl3, 25 °C) spectrum of the alkoxy-phenoxy exchange reaction with complex A and para-methoxy phenol. The loss of the bridging ethoxy peaks (denoted by red stars) and the emergence of bridging phenoxy signals (denoted by green stars) were monitored and used to calculate the rate of exchange, kexchange.   126   Figure 4.16 Plot of ln[A] versus time for exchange between catalyst A, HOC6H4OMe () and HOC6H4Br () and HOC6H5 (). Reactions carried out in an NMR tube (CDCl3, 25 °C). 1,3,5-trimethoxybenzene (TMB) was used as internal standard. [A] = 0.0046 M, [HOC6H4OMe] = 0.0050 M, [HOC6H4Br] = 0.0059 M, [HOC6H5] = 0.00411 M. The value of kexchange was determined from the slope of ln[A] vs. time, averaged from three experiments.   To probe the immortal ROP reaction further, a polymerization with a ratio of 10500:1:20 ([LA]:[A]:[phenol]) in DCM at room temperature was performed and aliquots were taken at regular intervals over 8 hours. The extracted aliquots were quenched and precipitated with wet methanol and analysed using GPC analysis (Figure 4.17 and Figure 4.18). GPC peaks shift to short elution times which correlates to an increase in molecular weight. This linear growth with respect to conversion is expected of a living process.   time (s)0 1000 2000 3000ln[ethoxy bridge]-4-3-2-10[A] + para-HOC6H4OMekexchange = -5.4 (6) x 10-4s-1[A] + para-HOC6H4Brkexchange= -15.6 (4) x 10-4s-1[A] + HOC6H5kexchange =  -6.8 (4) x 10-4 s-1	 127  elution time (min)10 12 14 16differential refractive index (RIU)0.00.20.40.60.81.01.216% 42%52% 56% 60% 66% 73%conversion Mn Figure 4.17 GPC traces with respect to conversion for the immortal ROP of lactide with A and phenol as the CTA ([LA]:[A]:[phenol] = 10500:1:20).   conversion (%)0 20 40 60 80molecular weight (kg mol-1)0102030405060dispersity1.01.11.21.31.41.5Mndispersity Figure 4.18 Plot of molecular weight and dispersity with respect to conversion for the immortal ROP of lactide with A with phenol ([LA]:[A]:[phenol] = 10500:1:20).    128   The immortal ring-opening polymerization of rac-LA was attempted with four different aromatic diols as chain transfer agents (CTAs) at room temperature (Chart 4.2). Under an inert atmosphere, solutions of rac-LA and CTA in DCM were prepared. After a homogenous solution had been obtained, appropriate amounts of complex A were added and the resulting reaction mixtures were stirred at room temperature. Catalyst to CTA ratios were kept constant at 1:20, while catalyst to monomer ratios were varied from 1050 to 21000 to generate polymer chains 50 to 1000 units in length. The resulting polymers using phenol (PhOH), showed excellent control over molecular weight and dispersity, compared to the living polymerization with aryl initiators (Table 4.4 and Figure 4.19). The initiating species in the immortal ROP reaction is complex A with a bridging ethoxy moiety. As the initiator is an alkoxide, it is expected that all the catalyst in solution is activated unlike with complexes 1-5 leading to well-defined polymers.  Chart 4.2 Aromatic diols used as chain transfer agents (CTAs) for the immortal ROP of rac-LA.      129  [LA]/([A]+[CTA])0 200 400 600 800 1000 1200Mn, kg mol-1050100150200250dispersity1.01.52.02.53.0MnMn,theodispersity[LA]/([A]+[CTA])0 200 400 600 800 1000 1200Mn, kg mol-1050100150200250dispersity1.01.52.02.53.0MnMn,theodispersitya) PLA-Et b) PLA-Ph [LA]/([A]+[CTA])0 200 400 600 800 1000 1200Mn, kg mol-1050100150200250dispersity1.01.52.02.53.0Mn,theoMndispersityPLA-1,5-biNapc) Figure 4.19 Plots of observed PLA Mn (closed symbols) and dispersity (open symbols) as functions of [LA]:([A]+[CTA]) for (a) CTA = ethanol (PLA-Et),  (b) CTA = phenol (PLA-Ph), (c) CTA = 1,5-naphthalenediol (PLA-1,5-Nap) 130  Table 4.4 Molecular weight data for immortal ROP of rac-LA with various aryl chain transfer agents.  Entry CTA Time (h) Conv. (%)a [LA]/([A]+[CTA]) Mn,theo (g mol-1)b Mn,GPC  (g mol-1)c Đc 1  168 0 71 -- -- -- 2d 65 44 79 5310 3160e -- 3  168 0 53 -- -- -- 4  168 0 45 -- -- -- 5i 65 33 26 1400 8750e -- 6  18 93 91 12200 14400 1.04 7 18 99 875 125000 109000 1.04 8 18 98 1730 245000 197800 1.01 9  18 99 49.8 7080 8750 1.01 10 18 92 255 33800 30400 1.00 11 18 98 993 140000 117000 1.01 12 18 97 503 70300 54400 1.00 Reactions carried out in DCM at room temperature. a Monomer conversions determined by 1H NMR spectroscopy. b Calculated from Mn,theo = (144 g mol-1 × conversion × [LA]/([A]+[CTA])+ MOR. c Absolute molecular weights were determined by GPC−MALS−RI (gel permeation chromatography−multi-angle light scattering−refractive index) via Universal Calibration (THF 4 mg mL–1, flow rate = 0.5 mL min–1, dn/dc = 0.044 mL g–1). d Reactions carried out in THF. e Molecular weights obtained from MALDI-ToF mass spectrum.   End-group fidelity was confirmed by 1H NMR spectroscopy and MALDI-ToF mass spectrometry. The 1H NMR spectra show peaks in the aromatic region corresponding to the aryl-chain ends (Figure 4.20– Figure 4.22). The integrations of the aromatic peaks and the methylene  131  proton signals (5.10-5.25 ppm) are in agreement with the calculated molar ratios of [LA]:([A]+[CTA]).  The MALDI-ToF mass spectrum for a selected 1,5-Nap-PLA sample is shown in Figure 4.24. This spectrum shows one major series of peaks separated by 72 mass units, indicative of chain transfer occurring during reaction. From analysis of the peak masses, this series is assigned to PLA with a 1,5-naphthalenediol chain-end (Figure 4.24).   Figure 4.20 1H NMR (CDCl3, 25 °C, 400 MHz) spectrum of PLA isolated from polymerization with [LA]:[A]:[diClPhOH] ratios of 237:1:2 (Table 4.4, entry 2).   132   Figure 4.21 1H NMR (CDCl3, 25 °C) spectrum of PLA isolated from polymerization with [LA]:[A]:[1,8-NapOH] ratios of 78:1:2 (Table 4.4, entry 5).   133   Figure 4.22 1H NMR (CDCl3, 25 °C) spectrum of PLA isolated from polymerization with [LA]:[A]:[ PhOH] ratios of 1050:1:20 (Table 4.4, entry 9).    134   Figure 4.23 1H NMR (CDCl3, 25 °C) spectrum of unconverted lactide isolated from polymerization with [LA]:[A]:[biPhenOH] ratios of 583:1:10 (Table 4.4, entry 3).   Figure 4.24 MALDI-ToF mass spectrum for a selected PLA-1,5-Nap sample. An = [72.07 LA]n + 160.17 1,5-Nap  + 23Na+   135   Of the diols used as CTAs, only 1,5- naphthalenediol was incorporated into the polymer chain under the initial reaction conditions; the biphenol, dichlorophene and 1,8-naphthalenediol completely hindered the polymerization activity in DCM. Previously reported computational studies showed that both metal centres are active during polymerization with complex A.129 Chelation of the diols to one or both metals would be expected to hinder polymerization.    One such species was isolated. Reaction of A with 2 equivalents of dichlorophene in toluene at room temperature for 16 h forms the dinuclear complex [(NNO)InCl]2(μ-Cl)(2,μ-diClPh) (6) (Scheme 4.4). The 1H NMR spectrum of complex 6 shows separate signals for each ligand on the two metal centres, indicating a breakdown in symmetry compared to the previous complexes in the series (Figure 4.25).   Scheme 4.4 Synthesis of [(NNO)InCl]2(μ-Cl)(2,μ-diClPh) (6).     136   Figure 4.25 1H NMR (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(2,μ-diClPh) (6)   Single crystals suitable for X-ray crystallography were obtained from the slow evaporation of a toluene solution. The solid-state structure of 6 shows a deprotonated dichlorophene molecule chelated to one indium centre with one of the oxygen atoms bridging both indium centres (Figure 4.26). Compound 6 displays the same asymmetric homochiral dinuclear structure with distorted octahedral indium centres as other complexes in this family and the “cis” relationship between the phenoxy groups of the ligand is maintained (Table 4.1).   137   Figure 4.26 Molecular structure of [(NNO)InCl]2(μ-Cl)(2,μ-diClPh) (6) depicted with ellipsoids at 50% probability (H atoms and all solvent molecules omitted for clarity).    Related complex [(NNO)InCl]2(μ-Cl)(2,μ-biPh) can be generated from reaction of complex A with biphenol. Although [(NNO)InCl]2(μ-Cl)(2,μ-biPh) was not further characterized, a similar structure to that of 6 is assumed, based on the 1H NMR spectrum (Figure 4.27). The chelating initiator explains the lack of reactivity in this system. In contrast, 1,5-naphthalenediol with hydroxyl groups on opposite sides of the naphthalene ring is active as a chain transfer agent. (Table 4.4, Figure 4.19c and Figure 4.24).   138   Figure 4.27 1H (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(2,μ-biPh).   To further confirm that these stable chelate species were shutting down reactivity, immortal ROP using A and CTAs dichlorophene and 1,8-naphthalenediol was conducted in THF. These conditions did allow for the generation of polymeric materials (Table 4.4, entries 2 and 5). We attribute this effect to the blocking of a coordination site by solvent molecules, disallowing the formation of the chelating complex. Additionally, reaction of the stable chelate complex 6 with rac-LA in THF also generated polymer, albeit with only 30% conversion after 5 days, showing that coordinating solvents can disrupt the chelate leading to polymer growth (Figure 4.28 and Figure 4.29).    139   Figure 4.28 1H NMR (CDCl3, 25 °C, 300 MHz) spectrum of polymerization [LA]:[6] = 110 in THF over 120 hours. Red star indicates the methylene protons on growing polymer.   140   Figure 4.29 1H NMR (CDCl3, 25 °C, 300 MHz) spectrum of polymerization [LA]:[6] = 110 in DCM over 120 hours.   4.3 Conclusions  We were interested in utilizing phenol rich bioproducts as chain transfer agents in catalytic reactions.  In this chapter the role of phenols as initiators and as chain transfer agents in the polymerization of lactide catalyzed by diaminophenolate-supported indium complexes was explored. The correlation between electron deficiency of the bridging phenol for the aryloxide-bridged complexes 1-5, and catalytic activity towards the living ring-opening polymerization of rac-lactide, was investigated. In situ studies showed that increasing electron deficiency of the bridging phenol leads to an associated increase in the initiation period and decrease in the rate of propagation. This effect is attributed to the lower nucleophilicity of the phenolic oxygen, which hinders the initial insertion into the monomer ester C-O bond and reduces the concentration of  141  active catalyst in solution, leading to higher-than-expected molecular weights. Although this effect is not surprising, the electron rich aryloxy species showed reaction rates only half that of the ethoxy catalysts, showing that these moieties have the potential to be active in catalysis.  Phenols were also used in tandem with ethoxy-bridged complex A to investigate the use of aromatic alcohols as chain transfer agents in the immortal ring-opening polymerization of lactide. Aromatic diols shut down polymerization by chelating one indium centre to form a stable metal complex; however, in coordinating solvents the formation chelating metal complexes is disrupted and polymerization can proceed. Immortal ROP was successful when using phenol and 1,5-naphthalenediol. Polymers with well-controlled molecular weights and molecular weight distributions were generated, and the aryl chain ends were confirmed using 1H NMR spectroscopy, MALDI-ToF mass spectrometry and UV-Vis spectroscopy.  These results show that both alkyl and electron rich aryl alcohols can be active chain transfer agents with this family of dinuclear indium complexes. Future efforts will be focused on expanding this fundamental work to complex arylated molecules such as lignin.   4.4 Experimental  General Methods. All the air and moisture sensitive manipulations were carried out in an MBraun glovebox or using standard Schlenk line techniques.  A Bruker Avance 300 or 400 MHz spectrometer was used to record 1H spectra. 1H NMR chemical shifts are given in ppm versus residual protons in deuterated solvents as follows: δ 7.27 CDCl3, δ 5.31 CD2Cl2. Molecular weights, hydrodynamic radii and intrinsic viscosities were determined by GPC-MALS-RI-Viscometer using an Agilent liquid chromatograph equipped with a Agilant 1200 series pump and  142  autosampler, three Phenogel 5 μm Narrow Bore columns (4.6 × 300 mm with 500 Å, 103 Å and 104 Å pore size), a Wyatt Optilab differential refractometer, Wyatt tristar miniDAWN (laser light scattering detector) and a Wyatt ViscoStar viscometer. The column temperature was set at 40 °C. A flow rate of 0.5 mL/min was used and samples were dissolved in THF (ca. 4 mg/mL). The measurements were carried out at laser wavelength of 690 nm at 25 °C. Data were processed by ASTRA software (Wyatt Technology). Molecular masses were determined using a Bruker Autoflex time-of-flight mass (TOF) spectrometer equipped with MALDI ion source.  Materials. THF, toluene and diethyl ether was taken from an IT Inc. solvent purification system with activated alumina columns and degassed before use. HPLC grade DCM was purchased from Fisher Chemicals and was dried over CaH2, transferred under vacuum and degassed before use. CDCl3 was purchased from Cambridge Isotope Laboratories Inc. and dried over CaH2, transferred under vacuum and degassed through three freeze−pump−thaw cycles before use. Pentane, acetonitrile and, CD2Cl2 were purchased from Sigma-Aldrich and dried over CaH2, transferred under vacuum and degassed before use. Indium (III) trichloride, 1,5-naphthalenediol, 1,8-naphthalenediol, para-methoxyphenol, para-bromophenol, para-nitrophenol, and para-methylphenol were purchased from Sigma-Aldrich and were used as received. Phenol was purchased from Fisher Chemicals and used as received. 2,2-Methylenebis(4-chlorophenol) (95%) (dichlorophene) was purchased from Alfa Aesar and recrystallized from hot toluene before use. 2,2'-biphenol was purchased from ACOS and used as received. Racemic lactide was purchased from PURAC America Inc. and was recrystallized three times from hot dried toluene prior to use. The racemic proligand 2,4-di(t-butyl)-6-(((2-(dimethylamino)cyclohexyl)amino)methyl)phenol (±)-H(NNO), dichloro complex (NNO)InCl2  143  and [(NNO)InCl]2(-Cl)(-OEt) (A) were synthesized according to previously reported methods.40   Representative synthesis of phenolic potassium salts. Inside a glovebox, a suspension of potassium tert-butoxide (7.9 mg, 0.070 mmol) in 3 mL of Et2O was added to a solution of para-nitrophenol in 3 mL of Et2O (97.5 mg, 0.701 mmol). This was stirred at room temperature for one hour. After this time, a yellow solid precipitated and was filtered off, washed with acetonitrile and dried under vacuum.   Synthesis of [(NNO)InCl]2(μ-Cl)(μ-OPhOMe) (1) via salt metathesis. To a suspension of (NNO)InCl2 (15.0 mg, 0.278 mmol) in toluene (5 mL) was added a suspension of potassium para-methoxyphenoxide (57.3 mg, 0.353 mmol) in toluene (2 mL) under N2. The solution was made up to 8 mL and stirred at ambient temperature for 18 h. After this time the solution was a cloudy yellow colour and was concentrated to give a pale yellow solid. The solid was the dissolved in 2 mL of CHCl3 and filtered through Celite. The solution was concentrated to give an off-white solid which was then stirred in acetonitrile. The crude product was then isolated by filtration and finally recrystallized twice from a 2:3 mixture of CHCl3:pentane to give a white solid (45.0 mg, 28%). 1H NMR (400 MHz, CDCl3): δ 7.45 (d, 3JH-H = 8.8 Hz, 2H, Ph-H), 7.08 (d, 4JH-H = 2.4 Hz, 2H, Ar-H), 6.60 (d, 4JH-H = 2.4 Hz, 2H, Ar-H), 6.47 (d, 3JH-H = 8.8 Hz, 2H, Ph-H), 4.68 (d, 2JH-H = 11.7 Hz, 2H, Ar-CH2-N), 3.54 (s, 3H, Ph-OCH3), 3.54 (dd, 2H, 3JH-H = 1.76, Ar-CH2-N), 3.17 (d, 4JH-H = 7.6 Hz, 2H, NH), 2.98 (td, 3JH-H = 10.0, 3JH-H = 1.8 Hz, 2H, DACH), 2.69 (s, 6H, N(CH3)2), 2.59-2.38 (m, 4H, DACH), 1.93 – 1.85 (br m, 2H, DACH), 1.813(s, 6H, N(CH3)2), 1.34 (s, 9H, C(CH3)3), 1.22 (s, 9H, C(CH3)3), 1.17 – 0.86 (m, 2H, DACH). 13C{1H} NMR (151 MHz, CDCl3): δ 162.4, 153.9, 152.8, 138.2, 135.9, 129.0, 128.2, 125.9, 125.2, 123.9, 118.8, 113.3, 64.4, 55.1, 53.0, 44.2, 37.8, 35.1, 35.1, 33.7, 30.7, 29.9, 29.7, 24.9, 24.8, 21.8. Elemental analysis calculated  144  for C53H85Cl3In2N4O4: C, 54.03; H, 7.27; Cl, 9.03; In, 19.49; N, 4.76; O, 5.43; found: C, 53.95, H, 7.09.  Synthesis of [(NNO)InCl]2(μ-Cl)(μ-OPhMe) (2) via salt metathesis. To a suspension of (NNO)InCl2 (153 mg, 0.280 mmol) in toluene (5 mL) was added a suspension of potassium para-methoxyphenoxide (61.0 mg, 0.417 mmol) in toluene (2 mL) under N2. The solution was made up to 8 mL and stirred at ambient temperature for 18 h. After this time the solution was a cloudy yellow colour and was concentrated to give a pale yellow solid. The solid was the dissolved in 2 mL of CHCl3 and filtered through Celite. The solution was concentrated to give an off-white solid which was then stirred in acetonitrile. The crude product was then isolated by filtration and finally recrystallized from a 2:3 mixture of CHCl3:pentane to give a white solid (34.1 mg, 21%). X-ray quality crystals were grown via slow evaporation of a toluene solution. 1H NMR (400 MHz, CDCl3): δ 7.40 (d, 3JH-H = 8.2 Hz, 2H, Ph-H), 7.10 (d, 4JH-H = 2.4 Hz, 2H, Ar-H), 6.69 (d, 3JH-H = 8.2 Hz, 2H, Ph-H), 6.60 (d, 4JH-H = 2.4 Hz, 2H, Ph-H), 4.66 (d, 2JH-H = 12.3 Hz, 2H, Ar-CH2-N), 3.54 (d, 3H, 2JH-H = 12.3 Hz, 2H, Ar-CH2-N), 3.20 (d, 4JH-H = 10.6 Hz, 2H, NH), 3.00 (td, 3JH-H = 11.4, 3JH-H = 3.5 Hz, 2H, DACH), 2.68 (s, 6H, N(CH3)2), 2.58-2.40 (m, 4H, DACH), 1.92 – 1.80 (br m, 2H, DACH), 1.80 (s, 6H N(CH3)2), 1.34 (s, 9H, C(CH3)3), 1.22 (s, 9H, C(CH3)3), 1.17 – 0.86 (m, 2H, DACH). 13C{1H} NMR (151 MHz, CDCl3): δ 162.5, 157.6, 138.3, 135.8, 128.8, 125.9, 123.8, 122.4, 118.9, 64.4, 53.2., 50.7, 44.3, 37.9, 35.1, 33.7, 31.7, 30.8, 29.9, 24.9, 24.8, 21.9, 20.6. Elemental analysis calculated for C53H85Cl3In2N4O3: C, 54.77; H, 7.37; Cl, 9.15; In, 19.76; N, 4.82; O, 4.13; found: C, 55.05; H, 7.34.  Synthesis of [(NNO)InCl]2(μ-Cl)(μ-OPhH) (3) via salt methathesis. To a suspension of (NNO)InCl2 (111 mg, 0.204 mmol) in toluene (5 mL) was added a suspension of potassium phenoxide (39.6 mg, 0.300 mmol) in toluene (2 mL) under N2. The solution was made up to 8 mL  145  and stirred at ambient temperature for 18 h. After this time the solution was a cloudy yellow colour and was concentrated to give a pale yellow solid. The solid was the dissolved in 2 mL of CHCl3 and filtered through Celite. The solution was concentrated to give an off-white solid which was then stirred in acetonitrile. The crude product was then isolated by filtration and finally recrystallized from a 2:3 mixture of CHCl3:pentane to give a white solid (66.6 mg, 57%). X-ray quality crystals were grown via slow evaporation of a toluene solution. 1H NMR (400 MHz, CDCl3): δ 7.54 (d, 3JH-H = 7.6 Hz, 2H, Ph-H), 7.08 (d, 4JH-H = 2.4 Hz, 2H, Ar-H), 6.93 (t, 3JH-H = 8.2, Ph-H) 6.59 (d, 4JH-H = 2.4 Hz, 2H, Ar-H), 6.54 (t, 3JH-H = 7.6 Hz, 2H, Ph-H), 4.66 (d, 2JH-H = 11.7 Hz, 2H, Ar-CH2-N), 3.55 (d, 3H, 2JH-H = 10.0 Hz, 2H, Ar-CH2-N), 3.20 (d, 4JH-H = 10.6 Hz, 2H, NH), 3.00 (td, 3JH-H = 11.1, 3JH-H = 2.4 Hz, 2H, DACH), 2.70 (s, 6H, N(CH3)2), 2.59-2.39 (m, 4H, DACH), 1.93 – 1.84 (br m, 2H, DACH), 1.83 (s, 6H N(CH3)2), 1.34 (s, 9H, C(CH3)3), 1.20 (s, 9H, C(CH3)3), 1.16 – 0.86 (m, 2H, DACH). 13C{1H} NMR (100 MHz, CDCl3): δ 162.4, 160.0, 138.2, 137.9, 129.0, 128.2, 125.9, 125.2, 123.9, 122.9, 120.2, 118.8, 64.4, 53.1, 50.6, 44.2, 37.8, 35.1, 33.7, 31.7, 30.7, 29.9, 24.9, 24.8, 12.8, 21.4.   Synthesis of [(NNO)InCl]2(μ-Cl)(μ-OPhBr) (4) via alkoxide-phenoxide exchange.  To a solution of A (115 mg, 0.112 mmol) in toluene (5 mL) was added p-bromophenol (59.2 mg, 0.342 mmol) under N2. The solution was stirred for 18h. White precipitate slowly appeared as the reaction progressed. After stirring, the solution was concentrated to yield a white solid, which was washed with ACN (3  2 mL) and consequently dried under vacuum. The crude white product was recrystallized from 3:2 pentane/CH3Cl at –30°C to afford product as long white crystalline needles (60.2 mg, 47%). Suitable crystals for x-ray crystallography were grown from slow evaporation of toluene. 1H NMR (600 MHz, CDCl3): δ 7.43 (d, 3JH-H = 8.7 Hz, 2H, Ph-H), 7.11 (d, 4JH-H = 2.6 Hz, 2H, Ar-H), 7.01 (t, 3JH-H = 8.7, Ph-H) 6.61 (d, 4JH-H = 2.6 Hz, 2H, Ar-H), 4.60 (d, 2JH-H = 12.8  146  Hz, 2H, Ar-CH2-N), 3.56 (d, 3H, 2JH-H = 12.8 Hz, 2H, Ar-CH2-N), 3.17 (d, 4JH-H = 10.2 Hz, 2H, NH), 2.98 (td, 3JH-H = 11.8, 3JH-H = 3.1 Hz, 2H, DACH), 2.84 (s, 6H, N(CH3)2), 2.57-2.43 (m, 4H, DACH), 1.92 – 1.83 (br m, 2H, DACH), 1.82 (s, 6H N(CH3)2), 1.34 (s, 9H, C(CH3)3), 1.23 (s, 9H, C(CH3)3), 1.16 – 0.92 (m, 2H, DACH). 13C{1H} NMR (151 MHz, CDCl3): δ 162.2, 159.3, 138.2, 136.2, 131.0, 126.0, 124.8, 124.0, 118.7, 112.6, 64.4, 53.1, 50.7, 44.2, 37.8, 35.1, 33.7, 31.7, 30.8, 29.9, 24.9, 24.8, 21.8. Elemental analysis calculated for C52H82BrCl3In2N4O3: C, 50.90; H, 6.74; Br, 6.51; Cl, 8.67; In, 18.71; N, 4.57; O, 3.91; found: C, 51.07, H, 6.70.    Synthesis of [(NNO)InCl]2(μ-Cl)(μ-OPhBr) (4) via salt methathesis. To a suspension of (NNO)InCl2 (148 mg, 0.272 mmol) in toluene (5 mL) was added a suspension of potassium para-methoxyphenoxide (107 mg, 0.6200 mmol) in toluene (2 mL) under N2. The solution was made up to 8 mL and stirred at ambient temperature for 18 h. After this time the solution was a cloudy yellow colour and was concentrated to give a pale yellow solid. The solid was the dissolved in 2 mL of CHCl3 and filtered through Celite. The solution was concentrated to give an off-white solid which was then stirred in acetonitrile. The crude product was then isolated by filtration and finally recrystallized from a 2:3 mixture of CHCl3:pentane to give a white solid. All characterization data was consistent with the previous description.  Synthesis of [(NNO)InCl]2(μ-Cl)(μ-OPhNO2) (5) via alkoxide-phenoxide exchange. To a solution of A (50.0 mg, 0.049 mmol) in toluene (4 mL) was added p-nitrophenol (24.1 mg, 0.173 mmol) under N2. The solution was stirred for 18 h. Yellow precipitate appeared in the solution as the reaction progressed. The solution was concentrated to yield a dark yellow solid, which was washed with ACN (3  2 mL) and consequently dried under vacuum. The crude product was recrystallized from 3:2 pentane/CH3Cl at –30 °C to yield product as short yellow crystalline needles (23.2 mg, 43%). X-ray quality crystals were grown via slow diffusion of pentane into  147  CHCl3. 1H NMR (400 MHz, CDCl3) δ 7.85 (d, 3J H-H= 8.9 Hz, 2H, Ph-H), 7.60 (d, 3JH-H = 9.0 Hz, 2H, Ph-H), 7.09 (s, 2H, Ar-H), 6.59 (s, 2H, Ar-H), 4.53 (d, 2JH-H = 12.6 Hz, 2H, Ar-CH2-N), 3.56 (d, 2JH-H = 12.4 Hz, 2H, Ar-CH2-N), 3.17 (d, 4JH-H = 9.4 Hz, 2H, NH), 2.96 (t, 3JH-H = 10.6 Hz, 2H, DACH), 2.69 (s, 6H, N(CH3)2), 2.58 – 2.38 (m, 6H, N(CH3)2), 1.35 – 1.27 (m, 22H), 1.20 (s, 18H, C(CH3)3). 13C{1H} NMR (151 MHz, CDCl3): δ 166.9, 161.9, 138.3, 136.7, 129.0, 128.2, 126.0, 124.7, 124.2, 123.5, 118.5, 64.6, 53.2, 50.7, 44.3, 37.9, 35.1, 34.6, 33.8, 31.7, 30.8, 30.0, 29.9, 24.7, 21.8.  Elemental analysis calculated for C52H82Cl3In2N5O5: C, 52.34; H, 6.93; Cl, 8.91; In, 19.24; N, 5.87; O, 6.70; found: C, 52.66, H, 6.97.   Synthesis of [(NNO)InCl]2(μ-Cl)(μ-OPhNO2) (5) via salt metathesis. To a suspension of (NNO)InCl2 (148 mg, 0.272 mmol) in toluene (5 mL) was added a suspension of potassium para-methoxyphenoxide (73.8 mg, 0.417 mmol) in toluene (2 mL) under N2. The solution was made up to 8 mL and stirred at ambient temperature for 18 h. After this time the solution was a cloudy yellow colour and was concentrated to give a pale yellow solid. The solid was the dissolved in 2 mL of CHCl3 and filtered through Celite. The solution was concentrated to give a yellow solid which was then stirred in acetonitrile. The crude product was then isolated by filtration and finally recrystallized from a 2:3 mixture of CHCl3:pentane to give a yellow solid (90.9 mg, 56%). All characterization data was consistent with the previous description.  Synthesis of [(NNO)InCl]2(μ-Cl)(2,μ-diClPh) (6). To a solution of A (83.3 mg, 0.076 mmol) in toluene (4 mL) was added dichlorophene (40.8 mg, 0.151 mmol) under N2. The solution was stirred for 18 h. White precipitate appeared in the solution as the reaction progressed. The solution was concentrated to yield an off-white solid, which dissolved in CHCl3 and filtered through Celite. The solvent was again removed and the white solid washed with ACN (3  2 mL) and consequently dried under vacuum. The crude product was recrystallized from 3:2  148  pentane/CH3Cl at –30 °C to yield product as a white powder needles (50.8 mg, 52%). X-ray quality crystals were grown via slow evaporation of toluene solution. 1H NMR (400 MHz, CDCl3) δ 7.58 (d, 3JH-H = 8.8 Hz, 1H, diClPh-H), 7.21 (d, 4JH-H = 2.3 Hz, 1H, Ar-H), 7.12 (d, 4JH-H = 2.3 Hz, 1H, Ar-H),  7.02 (d, 3JH-H = 2.9 Hz, 1H, diClPh-H), 7.01 (d, 3JH-H = 2.9 Hz, 1H, diClPh-H), 6.97 (d, 3JH-H = 8.5 Hz, 1H, diClPh-H), 6.83 (d, 3JH-H = 2.7 Hz, 1H, diClPh-H), 6.76 (d, 3JH-H = 2.6 Hz, 1H, diClPh-H), 6.74 (d, 3JH-H = 2.6 Hz, 1H, diClPh-H), 6.70 (d, 4JH-H = 2.3 Hz, 1H, Ar-H), 6.54 (d, 4JH-H = 2.3 Hz, 1H, Ar-H), 5.41 (d, 2JH-H = 15.2 Hz, 1H, diClPh-CH2-PhCl), 4.63 (d, 2JH-H = 11.6 Hz, 2H, Ar-CH2-N), 3.65 (dd, 3JH-H = 13.2 Hz, 3JH-H =2.9 Hz, 2H, CH), 3.49 (d, 2JH-H = 12.6 Hz, 2H, Ar-CH2-N), 3.37 (d, 2JH-H = 15.8 Hz, 1H, diClPh-CH2-PhCl), 3.13 (dd, 3JH-H = 13.2 Hz, 3JH-H =2.6 Hz, 2H, CH), 2.91 (s, 3H, N(CH3)), 2.89 (m, 1H,), 2.62 (s, 3H, N(CH3)), 2.61(m, 1H), 2.59 (m, 3H), 2.29 (m, 3H), 1.93 (s, 3H, N(CH3)), 1.88 (m, 5H), 1.72 (s, 3H, N(CH3)),  1.41 (s, 9H, C(CH3)3), 1.32 (s, 9H, C(CH3)3), 1.25 (s, 9H, C(CH3)3), 1.22 (s, 9H, C(CH3)3), 1.17 – 0.85 (m, 5H), 0.7 (d, J = 9.7 Hz, 1H). 13C{1H} NMR (151 MHz, CDCl3): δ 162.63, 157.8, 137.9, 135.4, 133.2, 132.7, 131.2, 130.6, 127.2, 126.8, 126.0, 124.7, 124.1, 121.7, 120.5, 99.9, 65.0, 64.2, 53.3, 44.8, 44.2, 37.7, 35.9, 33.6, 31.7, 30.1, 29.7, 25.1, 24.9.   Representative living ring-opening polymerization. Inside a glovebox under nitrogen, appropriate volumes of a standard solution of catalyst in DCM (20 mg into 1 mL) were added to a solution of rac-lactide in DCM. The solutions were stirred at room temperature for at least 18 hours before the solvent was evaporated. A small sample of crude polymer was taken for 1H NMR spectroscopic analysis, the remaining polymer was dissolved in minimal DCM (~2 mL) and precipitated with cold methanol. The methanol was decanted off and the precipitation repeated twice. The isolated polymer was then dried under vacuum for at least 24 hours before further analysis.   149   Representative immortal ring-opening polymerization. Inside a glovebox under nitrogen, appropriate amounts of rac-lactide were dissolved in DCM. Appropriate volumes of a standard solution of chain transfer agent in DCM were added and the solution stirred until homogenous. A standard solution of catalyst A in DCM (10 mg into 1 mL) was the added the reaction mixtures were stirred at room temperature for at least 18 hours before the solvent was evaporated. A small sample of crude polymer was taken for 1H NMR spectroscopic analysis, the remaining polymer was dissolved in minimal DCM (~2 mL) and precipitated with cold MeOH. The MeOH was decanted off and the precipitation repeated twice. The isolated polymer was then dried under vacuum for at least 24 hours before further analysis.  Representative in situ living ring-opening polymerization. Inside a glovebox under nitrogen, 0.25 mL of a standard solution of catalyst in CD2Cl2 (20 mg, 0.004 M) was added to a J-Young tube. This was frozen in liquid nitrogen before 0.25 mL of CD2Cl2 was added then frozen. Finally, 0.5 mL of a standard solution of rac-lactide and TMB (internal standard) in CD2Cl2 (123 mg, 0.21 M) was added and frozen. The J-Young tube was evacuated and thawed before insertion into the NMR probe.  Representative in situ phenoxy-ethoxy exchange. Inside a glovebox under nitrogen, 0.50 mL of a standard solution of catalyst in CD2Cl2 (20 mg, 0.005 M) was added to a J-Young tube. This was frozen in liquid nitrogen before 0.25 mL of CD2Cl2was added then frozen. Finally, 0.25 mL of a standard solution of phenol and TMB (internal standard) in CD2Cl2 (2.5 mg, 0.005 M) was added and frozen. The J-Young tube was evacuated and thawed before insertion into the NMR probe.  Living ring-opening polymerization with complex 6. Inside a glovebox under nitrogen, rac-lactide (31 mg, 0.215 mmol) was dissolved in THF (~2 mL). Catalyst 6 (2.5 mg, 0.002 mmol) in  150  1 mL of THF was then added. The reaction mixtures were stirred at room temperature. At regular intervals, a 0.5 mL aliquot was taken to check conversion.    151  Chapter 5: Thermorheological investigation of aromatic interactions in aryl-capped polylactides.****  5.1 Introduction  The previous chapter described the use of immortal ring-opening polymerization (iROP) as a trivial route to end-functionalized poly(lactide)s. This synthetic methodology is further utilized in the following sections to yield a variety of aryl-capped poly(lactide)s. Structure-property relationships are then studied with the goal of probing the properties of these potentially useful materials.  5.1.1 Network structures in polymeric materials.  Polymer networks can be divided into three main classes, characterized by their specific rheological properties.240,241 The mechanical response of thermoset rubbery solids depends on crosslink density and molecular mass and due to the strong covalent bonds between chains, when a stress is applied, these materials display elastic response over a large deformation frequency range (G’ > G”).   Thermoplastic polymers have mechanical response which depends on entanglement molecular weight, chain alignment and degree of crystallinity. The melt rheological behaviour of thermoplastics are very distinctive whereby at very low deformation frequencies in the terminal zone (< 10-2 s-1), thermoplastics behave as highly viscous liquids.  As the frequency increases                                                  **** This work is in revision for publication: Chile, L.-E.; Hatzikiriakos, S. G.; Mehrkhodavandi, P; J. Rheol, in revision.  152  viscoelastic response is given by the following relationships: G’  2 and G”  1. At a certain cross-over frequency, the G’ surpasses G” and the material begins to display more elastic response.  Elastomers can contain non-covalent interactions and display intermediate behaviour between that of a thermoset and a thermoplastic. These materials are often healable because after deformation or breakage the secondary interactions between polymer chains can reform the network structure, restoring the material’s mechanical properties.  5.1.2 Non-ionic secondary interactions to modify polymer properties.  Deliberate inclusion of moieties capable of secondary interactions in polymeric structures has become established in the field of functional and stimuli responsive materials.242,243 The influence of non-covalent bonds, specifically hydrogen bonding and aromatic-aromatic electrostatic interactions, have been applied to achieve a variety of supramolecular structures including hydrogels and supramolecular polymers. The non-covalent bonds in these materials can lead to enhanced properties (elasticity and strength) as well as imparting self-healing capabilities244 by altering the way in which polymer chains interact.245,246  The literature holds only a few examples of researchers using these secondary interactions to alter the properties of biodegradable polyesters.188,207,247,248 Sijbesma and co-workers synthesized functionalized poly(caprolactone) (PCL), end-capped with a self-complimentary quadrupole hydrogen-bonding group, ureidopyrimidinone (UPy) (Figure 5.1).249 The UPy groups were found to form aggregate dimers which remained present in the polymer melt (at T > Tm), giving rise to a physical network. The presence of this network influenced the melt rheological properties of the polymers and rheological analysis in the linear viscoelastic regime showed a plateau in the storage  153  modulus, G’, at low deformation frequencies similar to that observed in covalently cross-linked systems.   Recently, the same UPy hydrogen-bonding unit was incorporated in PLA (UPy-PLA) using iROP by Brzeziński and Biela who showed that stereocomplexation in between PLLA and PDLA in dilute solutions was promoted by UPy aggregation.250 To further these results, Pan et al. studied the crystallization rate of UPy-PLA polymers.251 The UPy aggregation led to an increase in the melting temperature of the end-functionalized PLLAs and homocrystallization was faster when compared to the non-functionalized counterpart. In addition, when UPy-PDLA was blended with its enantiomer, the stereocomplex PLA formed showed significantly elevated melting temperatures, indicating the strength and efficacy of tailored hydrogen-bonding interactions on modifying polymer properties.   Figure 5.1 Self complimentary hydrogen-bonding unit, ureidopyrimidinone (UPy).   A healable polymer blend was reported by Colquhoun and Hayes in 2010 which utilized an electron deficient bis(diimide) capable of forming complimentary - interactions with -electron rich species such as pyrene (Figure 5.2). These blends showed improved elasticity and toughness  154  (plastic deformation) and could maintain at least 90% of these properties after breaking and healing. The pyrene-capped polymer controls in this study did not display any re-healing character.252   Dubois and co-workers utilized iROP to synthesize pyrene-capped PLAs (pyr-PLA) and explored the effect on stereocomplexation of PDLA and PLLA units. When investigating the transport properties of pyr-PLA, slower diffusion of the end-capped polymers was observed, attributed to aromatic-aromatic interactions between pyrene chain-ends.206 With the addition of carbon nanotubes, this aggregation was prevented and diffusion of the polymer chains was promoted leading to another report in 2015 on the use of pyr-PLAs as dispersion agents in PLA/multi-wall nanotube composites.253   Figure 5.2 Bis(diimide) capable of forming complimentary - interactions with -electron rich species such as pyrene.206,253   The following sections further the exploration of aromatic-aromatic electronic interactions as a method to modify and control polymer properties. In particular, the synthesis of a family of aryl-terminated PLAs via immortal ROP using aromatic alcohols is described and thermorheological techniques are used to understand the effects of aromatic interactions (i.e. extent of molecular associations) on bulk PLA properties. 155  5.2 Results and discussion 5.2.1 Synthesis of aryl-capped PLAs via immortal ring-opening polymerization.  A series of PLAs with bulky aromatic chain ends were synthesized through the immortal ring-opening polymerization (iROP) of rac-LA with dinuclear indium catalyst [(NNO)InCl]2(μ-Cl)(μ-OEt) (A) using phenol (HOPh), 1-naphthol (Nap), hydroxymethylpyrene (HOMePyr), and hydroxymethylanthracene (HOMeAnth) as chain transfer agents (CTA) (Scheme 5.1 and Chart 5.1). Under an N2 atmosphere in a glove box, solutions of rac-LA and CTA in DCM were prepared. After a homogenous solution had been obtained, appropriate amounts of complex A were added and the resulting reaction mixtures were stirred at room temperature. Catalyst to CTA ratios were kept constant at 1:20, while catalyst to monomer ratios were varied from 1050 to 21000 to generate polymer chains from 50 to 1000 units in length. The polymer samples were designated labels based on their end group and GPC molecular weight. For example, PLA-Ph34 is a polymer with a phenol chain-end and Mn = 34 kg mol-1 (Table 5.1). 156  Scheme 5.1 Synthesis of end-functionalized PLA via immortal ring-opening polymerization with dinuclear indium catalyst [(NNO)InCl]2(μ-Cl)(μ-OEt) (A).     Chart 5.1 Alcohols used as chain transfer agents (CTAs) for the immortal ROP of rac-LA.    157  [LA]/([A]+[CTA])0 200 400 600 800 1000 1200Mn, kg mol-1050100150200250dispersity1.01.52.02.53.0MnMn,theodispersity[LA]/([A]+[CTA])0 200 400 600 800 1000 1200Mn, kg mol-1050100150200250dispersity1.01.52.02.53.0MnMn,theodispersity[LA]/([A]+[CTA])0 200 400 600 800 1000 1200Mn, kg mol-1050100150200250dispersity1.01.52.02.53.0MnMn,theodispersity[LA]/([A]+[CTA])0 200 400 600 800 1000 1200Mn, kg mol-1050100150200250dispersity1.01.52.02.53.0MnMn,theodispersity[LA]/([A]+[CTA])0 200 400 600 800 1000 1200Mn, kg mol-1050100150200250dispersity1.01.52.02.53.0MnMn,theodispersitya) PLA-Et a) PLA-Phd) PLA-MePyre) PLA-MeAnthc) PLA-Nap Figure 5.3 Plots of observed PLA number averaged molecular weight, Mn, (closed symbols) and dispersity, Đ, (open symbols) as functions of [LA]:([A]+[CTA]) for (a) CTA = ethanol (PLA-EtOH), (b) CTA = phenol (PLA-Ph), (c) CTA = naphthol (PLA-Nap), (d) CTA = hydroxymethylpyrene (PLA-MePyr) and (e) hydroxymethylanthracene (PLA-MeAnth). The black lines represent theoretical Mn values based on the [LA]:([A]+[CTA]) ratio at 100% conversion.  All reactions were carried out room temperature in DCM.     Figure 5.3 shows the excellent agreement between the expected (black line on plots) and experimental molecular weights and low dispersity values achieved from iROP. The 1H NMR spectra of low molecular weight PLAs (3 – 5 kg mol-1) show peaks in the aromatic region corresponding to the aryl-chain ends (Figure 5.4 - Figure 5.6). In addition, the integrations of the  158  aromatic peaks and the polymer methylene proton signals (5.10-5.25 ppm) are in agreement with the calculated molar ratios of [LA]:([A]+[CTA]).   UV-Vis spectroscopic analysis of PLA solutions in CHCl3 showed absorbance patterns similar to those of the free alcohols as well as showing the PLA absorbance signal at 240 nm (Figure 5.7 – Figure 5.10). Examination of the 1H{1H} NMR spectra of the aryl-capped PLA were all comparable to that of polymers produced by A alone, with Pm values around 0.6 implying low levels of isotacticity.    159  Table 5.1 Molecular weight data for immortal ring-opening polymerization of rac-LA with various aryl chain transfer agents.  Entry CTA samplea Conv. (%)b [LA]/([A]+[CTA]) Mn,theo (g mol-1)c Mn,GPC  (g mol-1)d Đd 1 PhOH PLA-Ph9 99 49.8 7080 8750 1.01 2 PLA-Ph30 92 255 33800 30400 1.00 3 PLA-Ph54 97 503 70300 54400 1.00 4 PLA-Ph117 98 993 140000 117000 1.01 5 NapOH PLA-Nap11 94 52.0 7050 11200 1.01 6 PLA-Nap63 99 497 70900 63100 1.03 7 PLA-Nap68 98 518 73300 67600 1.03 8 PLA-Nap105 99 993 142000 105000 1.01 9 PLA-Nap185 92 1020 136000 185000 1.01 10 HOMePyr PLA-MePyr17 99 96 13700 16500 1.00 11 PLA-MePyr73 99 496 70800 72800 1.00 12 PLA-MePyr137 98 970 137000 137000 1.00 13 PLA-MePyr153 70 1530 154000 153000 1.00 14 HOMeAnth PLA-MeAnth11 98 49.5 6990 10500 1.02 15 PLA-MeAnth44 99 249 35500 43700 1.05 16 PLA-MeAnth79 99 492 70200 79200 1.00 17 PLA-MeAnth128 99 979 140000 129000 1.02 18 PLA-MeAnth144 93 978 131000 144000 1.00 19 PLA-MeAnth181 98 1400 197000 181000 1.01 20 EtOH PLA-Et8 97 51.5 7220 8200 1.02 21 PLA-Et38 99 251 35800 38500 1.14 22 PLA-Et68 98 500 70600 68000 1.07 23 PLA-Et121 98 991 140000 121000 1.00 24 PLA-Et124 99 1010 144000 124000 1.03 Reactions were carried out in DCM at 25 °C for 18 h. a PLA-MePyr73 denotes a polymer with a molecular weight of 73 kg mol–1 and a hydroxymethylpyrene chain-end (Chart 5.1). bMonomer conversions determined by 1H NMR spectroscopy. c Calculated from Mn,theo = (144 g mol-1 × conversion × [LA]/([A]+[CTA])) + MOR. Absolute molecular weights were determined by triple detector GPC (gel permeation chromatography) via Universal Calibration (THF 4 mg mL–1, flow rate = 0.5 mL min–1, dn/dc = 0.044 mL g –1 ).  160   Figure 5.4 1H NMR (CDCl3, 25 °C, 300 MHz) spectrum of PLA isolated from polymerization of [LA]:[A]:[NapOH] ratios of 1050:1:20 (Table 5.1, entry 5).   Figure 5.5 1H NMR (CDCl3, 25 °C, 300 MHz) spectrum of PLA isolated from polymerization of [LA]:[A]:[PyrMeOH] ratios of 1050:1:20 (Table 5.1, entry 10).   161   Figure 5.6 1H NMR (CDCl3, 25 °C, 300 MHz) spectrum of PLA isolated from polymerization of [LA]:[A]:[AnthMeOH] ratios of 1050:1:20 (Table 5.1, entry 14).   162  PLA-Phwavelength (nm)250 300 350 400 450 500abs0.00.51.01.52.0121 kg mol-138 kg mol-18.8 kg mol-1Free PhOH  Figure 5.7 Solution UV-Vis spectra for phenol (purple line), PLA-Ph9 (blue line), PLA-Ph38 (red line) and PLA-Ph121 (black line). Polymer solutions were in CHCl3 with a concentration of 1×10-4 M. PLA-Napwavelength (nm)250 300 350 400 450 500abs0.00.51.01.52.011 kg mol-136 kg mol-1101 kg mol-1Free NapOH Figure 5.8 Solution UV-Vis spectra for naphthol (purple line), PLA-Nap11 (black line), PLA-Nap36 (red line) and PLA-Nap101 (blue line). Polymer solutions were in CHCl3 with a concentration of 1×10-4 M.  163  PLA-MePyrwavelength (nm)250 300 350 400 450 500abs0.00.20.40.60.81.017 kg mol-139 kg mol-1106 kg mol-1 Figure 5.9 Solution UV-Vis spectra for PLA-MePyr17 (black line), PLA-MePyr39 (red line) and PLA-MePyr106 (blue line). Polymer solutions were in CHCl3 with a concentration of 1×10-5 M. PLA-MeAnthwavelength (nm)250 300 350 400 450 500abs0123411 kg mol-137 kg mol-1124 kg mol-1 Figure 5.10 Solution UV-Vis spectra for PLA-MeAnth11 (black line), PLA-MeAnth37 (red line) and PLA-MeAnth124 (blue line). Polymer solutions were in CHCl3 with a concentration of 1×10-4 M.   5.2.2 Thermal studies of aryl-capped PLAs.  To identify an end-group effect on the thermal transitions, thermogravimetric analysis (TGA) and differential scanning calorimetery (DSC) were used to determine the thermal properties of the various aryl-capped PLAs in this study (Table 5.2).   164   Polymeric samples were heated at 20 °C/min from 25 °C to 500 °C, and the onset and total degradation temperatures, were calculated from the first derivative of the resulting thermographs (Figure 5.11). The aryl-capped polymers were thermally stable up to ~300 °C where they began their degradation. All the polymers including the native ethoxy-capped PLA, showed total degradation at ~ 395 °C. These observations indicate that the chain-end does not affect the overall thermal degradation pathways of PLA (Table 5.2). The most advantageous property of PLA is its biodegradability. The most advantageous property of PLA is its degradability. These results show that these potentially useful chain-end modifications will only affect melt or bulk properties while allowing for the same thermal degradation pathway as the corresponding native polymer.  Table 5.2 Thermal properties for various end-capped PLAs in this study (all values in °C). Entry Material Tg a Tonset b Ttotal b 1 PLA-Et 49.2 (6.7) 345(12) 400 (5) 2 PLA-Ph 47.9 (2.1) 340 (3) 391 (7) 3 PLA-Nap 43.3 (4.4) 341 (9) 396 (5) 4 PLA-MePyr 48.0 (5.4) 351 (8) 398 (5) 5 PLA-MeAnth 48.1 (7.5) 349 (8) 398 (9) a Thermal analysis of samples was performed by using a differential scanning calorimeter (DSC) with ca. 5 mg of sample. Samples heated to 275 °C at 10 °C/min and cooled to 25 °C at 5 °C/min to determine Tg and Tm. Glass transitions and melting temperatures calculated from the second heating scans. bThermogravimetric analysis was performed on approximately 20 mg of material. Samples were heated to 500 °C at a rate of 20 °C/min to determine the degradation onset temperature (temperature at which there is 5% weight loss, Tonset). Calculated standard errors given in parentheses.    165  temperature,oC0 100 200 300 400Normalized mass 0246810121416PLA-EtPLA-PhPLA-NapPLA-MePyr PLA-MeAnth  Figure 5.11 TGA heating traces for the various aryl-capped PLAs. Thermogravimetric analysis was performed on approximately 15 mg of material. Samples were heated to 500 °C at a rate of 20 °C/min.  50 100 150 200-0.10.10.0 0.0Heat flowTemperature (oC) PLA-Et PLA-Ph PLA-Nap PLA-MePyr PLA-MeAnth Figure 5.12 DSC traces for the various end-capped PLAs studied. Samples heated to 250 °C at 10 °C/min and cooled to 25 °C at 5 °C/min under a nitrogen atmosphere to reduce sample degradation. Glass transition and melting temperatures calculated from second cooling scans. ca. 5 mg of sample used. Plots shifted vertically for clarity.   166   In the DSC analysis of the aryl-capped PLAs, samples were heated from 25 °C to 275 °C at 10 °C/min and then cooled at 5 °C/min (Figure 5.12). To eliminate the influence of thermal history, glass transition temperatures were calculated from the second heating scans and averaged over all the molecular weights of polymers in the same family. All the PLAs were amorphous, displaying no crystallization peaks. The glass transition temperatures were comparable for all the end-capped PLAs and the enhancement in crystallization seen in studies by Yang207 and Dubois206,253 on isotactic PLA homopolymers capable of forming of  aggregates was not observed. The absence of this effect could be due to insufficient chain packing from the mostly atactic microstructure of the PLAs (Pm ~ 0.6), and thus the chain-end interactions are not sufficient to drive chain ordering between the small regions of isotacticity on the polymer chains.   5.2.3 Solution viscosity of aryl-capped PLAs.  Chain aggregation is expected to form physical networks which change the linear structure of the polymer chains thus influencing mobility of the polymer chains and their viscosity.250 To investigate this effect in solution intrinsic viscosities of the aryl-capped PLAs were obtained from the GPC analysis (in THF) and are plotted versus weight-averaged molecular weight, Mw in Figure 5.13. The data for aryl-capped PLAs fall on a single line (consistent with the Mark-Houwink equation with slopes between 0.76 and 0.80). This is in agreement with the relationship determined by Dorgan163 and Othman,166 indicating random coil behaviour of linear polymers in a good solvent. More specifically, the intrinsic viscosities for all aryl-capped polymers follow the same behaviour suggesting the absence of any particular interaction to form aggregates through the various aryl groups in dilute solution. Similarly, the calculated hydrodynamic radii for the aryl-capped polymers also show comparable trends with respect to molecular weight as previously  167  reported166 for other PLAs and the exponent implies linear microstructure and thus minimal chain aggregation in dilute solution (Figure 5.14).   Mw, (gmol-1)104 105[] (mL/g)101102103PLA-EtPLA-PhPLA-NApPLA-MePyrPLA-MeAnth[]all = 0.014Mw0.77Othman 2011[] = 0.014Mw0.75Dorgan 2005[] = 0.017Mw0.74 Figure 5.13 A log−log plot of intrinsic viscosity of various aryl-capped PLAs as a function of the molecular weight.   168  Mw, (gmol-1)104 105Hydrodynamic radius, Rh, (nm)100101PLA-EtPLA-PhPLA-NapPLA-MePyrPLA-MeAnthRh = 0.016Mw0.57Othman 2011Rh = 0.017Mw0.56 Figure 5.14 Characteristic hydrodynamic radius (Rh) as a function of molecular weight.   5.2.4 Melt Rheology - Linear viscoelasticity.    To determine the thermal stabilities of aryl-capped PLAs, time sweep experiments were conducted using a parallel plate rotational rheometer at 180 °C with constant angular frequency of 3.14 rad/s over 120 minutes. Data collection was started immediately after polymer samples were loaded onto the rheometer. The complex modulus | G*|, a measure of resistance to deformation sensitive to structural changes, was monitored over the experimental time frame. To accurately compare the structural changes occurring within the aryl-capped polymer samples, the normalized complex modulus, G*(t)/G*(t=0), is plotted against time in Figure 5.15a. Native ethoxy-capped polymers showed a steady complex modulus for much of the experiment, decreasing by ~10% as a result of thermal degradation.  In contrast, aryl-capped polymers display a significant increase in  169  their complex moduli over the experimental time frame, with the modulus for pyrene capped PLAs showing the largest increase of ~450%. To further examine how these structural changes can be induced, similar experiments were conducted on PLA-Nap36 at 70 and 120 °C (Figure 5.15b) indicating similar increases. The increase in the normalized complex modulus appears to be temperature dependent as this effect is nearly absent at low temperatures (Figure 5.15b). Chains have greater mobility at higher temperatures, and as such end associations tend to occur more readily. Figure 5.15c shows the frequency dependence of these associations. At high deformation frequencies, segmental motion is prevalent and chain-ends cannot move to form aggregates therefore no enhancement of the complex modulus is seen, indicating that the observed increase in modulus can be attributed to the gradual aromatic electrostatic association of the aryl-groups under an oscillatory shearing force. 252,254 At lower temperatures, the change in complex modulus diminishes and is absent at 70 °C. The upturn in the complex modulus is attributed to chain association through the aryl groups. At high temperature, due to higher chain mobility, the associations occur faster. Therefore, the observed upturn in modulus can be ascribed to the gradual  aromatic electrostatic association of the aryl-groups under a constant shearing force. shows the frequency dependence of these associations. Increasing the frequency of oscillation causes a reduction in chain aggregation due to the application of a higher shearing force that can disrupt these weak associations. These results further demonstrate that increased chain mobility is needed for these associations to occur which is the case at high temperatures.   170  time (s)0 2000 4000 6000 8000G*(t)/G*(t=0) (Pa)0123456PLA-Et121PLA-Ph54PLA-Nap54PLA-MePyr73PLA-MeAnth80a)time (s)0 2000 4000 6000 8000G*(t)/G*(t=0) (Pa)0.91.01.11.21.31.41.51.670 oC120 oC180 oCPLA-Nap36 b)PLA-Nap36time (s)0 2000 4000 6000 8000G*(t)/G*(t=0) (Pa)0.81.01.21.41.61.83.14 rad/s31.4 rad/s314 rad/sc) Figure 5.15 a) Normalized complex modulus, G*(t)/G*(t=0), vs. time at 180 °C for end-capped PLAs in this study. b) Normalized complex modulus, G*(t)/G*(t=0), vs. time at 70, 120 and 180 °C for PLA-Nap36. c) Normalized complex modulus, G*(t)/G*(t=0) vs. time at 3.14, 31.4 and 314 rad/s for PLA-Nap36. 171   Small amplitude oscillatory shear (SAOS) experiments in the linear viscoelastic regime were performed to study the structure-property relationships of the aryl-capped materials in order to identify subtle differences between the microstructure of the various polymers. A parallel plate rotational rheometer was used at a constant strain amplitude of 2% over frequencies ranging from 0.01 to 100 Hz. Time-temperature superposition was applied to isothermal experiments in the temperature range of 70-170 °C to generate master curves of the viscoelastic moduli, G’ and G” as well as the complex viscosity at the reference temperature of 150 °C. Figure 5.20a-b plots the master curves of the viscoelastic moduli as well as its complex viscosity at the reference temperature of 150 °C of two PLA-MePyr polymers of different molecular weights, 72 and 135 kg mol-1. The high molecular weight aryl-capped PLA behaves similarly to the ethoxy-capped PLA whereby the storage moduli exhibit a clear plateau at high frequencies and the loss moduli increase with frequency, going through a maximum and reaching a minimum value at higher frequencies.20 Deviations from the expected rheological behaviour of monodisperse linear polymers is observed as the molecular weight decreases where the aromatic groups begin to have a relatively larger effect on the rheological properties (note that there is only one aryl group per chain) (Figure 5.17). In the terminal region, at small deformation frequencies, the storage and loss moduli usually display characteristic slopes of 1 and 2, respectively. However, in the low molecular weight aryl-capped PLAs the slope of G’ becomes closer to 1 showing increase in elasticity due to associations at small frequencies. As a result, the complex viscosity also exhibits an upturn at small frequencies, deviating from its Newtonian steady value (Figure 5.17). (Additional SAOS data for all aryl-capped PLAs are presented in appendix C and show similar results.)  The plots for the lower molecular weight polymer PLA-MePyr72 (Figure 5.16a), show good superposition at medium to high frequencies (10-106 rad/s) due to the absence of associations at  172  these high frequencies, however, it clearly fails in the low frequency range. This indicates that the microstructure of the material is different at each temperature due to the presence of associations. These associations depend on both temperature and frequency as stated in the discussion of the thermal stability experiments above (Figure 5.15). PLA-MePyr73Tref = 150 oCangular frequency, (rads-1)10-1 100 101 102 103 104 105 106 107G', G" (Pa)101102103104105106107complex viscosity, |*| (Pa.s)10-1100101102103104105G'G"|*|PLA-MePyr137Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106 107G', G" (Pa)101102103104105106107108complex viscosity, |*| (Pa.s)10-1100101102103104105106G'G"|*|a) b) Figure 5.16 Thermorheologcial characterization of PLA-MePyrene polymers. a) Master curve of the linear viscoelastic moduli, G’ and G” and |*| complex viscosity for PLA-MePyr72 at Tref = 150 °C. b) Master curve of the linear viscoelastic moduli, G’ and G” and |*| complex viscosity for PLA-MePyr135 at Tref = 150 °C.   angular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106 107complex viscosity, |*| (Pa.s)10-110010110210310410563 kg mol-136 kg mol-195 kg mol-1190 kg mol-1PLA-NapTref = 150 oCMw Figure 5.17 Master curve of the complex viscosity vs. angular frequency (Tref = 150 °C) for naphthol-capped PLAs (Nap-PLA) to show the molecular weight dependence on the upturn in complex viscosity.   173   To further show that these aggregates are induced under low frequency oscillations, successive frequency sweep experiments were conducted on PLA-Nap36 and PLA-Et34 polymers (Figure 5.18). The G’ increases in the low frequency region successively after each experiment, however, it remains the same at the high frequency region (Figure 5.18a). The native polymer, on the other hand, shows no change in the storage modulus over the time period of the three tests (Figure 5.18b). These results indicate that only under low shear do the aromatic chain-ends have enough time to associate and these associations do not seem to be disrupted at high shear rates.  angular frequency,  (rad/s)10-2 10-1 100 101 102 103storage modulus, G' (Pa)1001011021031041051061st test2nd test3rd test PLA-Nap36T = 120 oCangular frequency,  (rad/s)10-2 10-1 100 101 102 103storage modulus, G' (Pa)1001011021031041051061st test2nd test3rd test PLA-Et34T = 120 oCa) b) Figure 5.18 Successive frequency sweep experiments for a) PLA-Nap36 and b) PLA-Et34 polymers. T = 120 °C,  = 0.063-630 rad/s, strain = 2%.  Time sweep experiments were conducted to show if the aromatic-aromatic associations can be disrupted (break) by high frequency deformations. The time sweep test started with an angular frequency of 3.14 rad/s for 60 minutes at 120 °C, and after this time the angular frequency was increased to 314 rad/s for another 60 minutes. After this a final test at 3.14 rad/s was then conducted for another 60 min. Figure 5.19 depicts the changes in the complex modulus over time. The G* at the high frequency is not shown on the plot in order to keep the vertical scale linear and to show the changes at small frequencies shown more clearly.  It is noted that at high frequency only a very  174  minor increase in G* was observed within experimental error (see also Figure 5.15c for the effect of high frequency on G*). However, the onset of aggregation seems to have been restored when the lower frequency has been resumed and continues nearly from the point left at the end of in the first low frequency test (a small increase of 3% from the end of low frequency test one to the beginning of the low frequency test two after 60 min). shift data to the true times120oCshear for a longer time to fully break the associationstime (s)0 2000 40006000 8000 10000G* (Pa)2000220024002600280030003200 = 3.14 rad/s1st test=  3.14 rad/s2nd testPLA-Nap36T = 120 oC Figure 5.19 Successive time sweep experiments for PLA-Nap36. T = 120 °C,  = 3.14 rad/s, strain = 2%.   This temperature dependence can be further seen in the plots of the horizontal shift factor, aT vs. the inverse temperature that we determined by shifting the data in Figure 5.20a-b to obtain the master curves plotted in Figure 5.16a-b. The shift factors in Figure 5.20c show two different slopes for the high and low temperature data. Similarly, the plot of the storage vs. loss modulus for the different isothermal experiments in Figure 5.20d. For linear PLAs isothermal experiments all collapse onto the same curve, however, the PLA-MePyr polymer plots show a temperature dependence on phase transitions in the aryl-capped polymers. This explains the failure of the time- 175  temperature superposition at low deformation frequencies where low shear cannot easily break the chain-end associations.    Loss modulus, G" (Pa)100 101 102 103 104 105 106 107Storage modulus, G' (Pa)100101102103104105106107170 oC150 oC130 oC110 oC90 oC70 oC1/Temperature (1/K)0.0022 0.0024 0.0026 0.0028 0.0030aT10-1100101102103104105106angular frequency,  (rads-1)10-2 10-1 100 101 102 103 104 105 106 107 108Storage modulus, G' (Pa)100101102103104105106107170 oC150 oC130 oC110 oC90 oC70 oCPLA-MePyr72c)angular frequency,  (rads-1)10-2 10-1 100 101 102 103 104 105 106 107 108Loss modulus, G" (Pa)100101102103104105106107170 oC150 oC130 oC110 oC90 oC70 oCPLA-MePyr72d)PLA-MePyr72Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106 107G', G" (Pa)101102103104105106107complex viscosity, |*| (Pa.s)10-1100101102103104105G'G"|*|PLA-MePyr135Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106 107G', G" (Pa)101102103104105106107108complex viscosity, |*| (Pa.s)10-1100101102103104105106G'G"|h*|a) b)PLA-MePyr7270oC90oC110oC130oC170oCPLA-MePyr72150oC Figure 5.20 a) Storage modulus vs. frequency, b) loss modulus vs. frequency for the isothermal frequency sweep experiments for PLA-MePyr72, to show the failure of the time-Temperature superposition principle. c) Horizontal shift factors, aT at 150 °C. d) Plot of G’ vs. G” at various temperatures for PLA-MePyr72 to check for possible thermal transitions. 176   The observed deviations from expected melt rheology for linear polymers are comparable to those seen in associating polymers which utilize secondary interactions to form network structures.249,255-259 In this case, the network structure is developed under oscillatory shear conditions through the formation of extended aggregate structures, though because these aggregates are transient and slow in their onset, the exact morphology of the structures is difficult to probe. Similar master curves for two lower molecular weight PLA-Nap polymers are shown in Figure 5.21. Analogous to the pyrene capped PLAs, the moduli at different temperatures do not form smooth master curves at low frequencies (displaying thermorheological complexity) mainly due to the temperature dependence of the aromatic associations in the low frequency regime. However, because these aggregates are transient and fast in their onset, the exact morphology of the structures are difficult to probe. PLA-Nap63Tref = 150 oCangular frequency, (rads-1)100 101 102 103 104 105G', G" (Pa)101102103104105106complex viscosity, |*| (Pa.s)100101102103104G'G"|*|PLA-Nap36Tref = 150 oCangular frequency, (rads-1)100 101 102 103 104 105G', G" (Pa)100101102103104105106complex viscosity, |*| (Pa.s)100101102103104G'G"|*|a) b) Figure 5.21 Master curve of the linear viscoelastic moduli, G′ and G″, and complex viscosity, |η*| (Tref = 150 °C), for a) PLA-Nap36 b) PLA-Nap63 polymers.   As the formation of these network structures is observable by an increase (upturn) in the complex viscosity in the low frequency region (Figure 5.17), the strength of association can be  177  probed by defining an “index of association” based on their increase in viscosity at low frequencies (Equation 5.1).   Index of association] × 100 Equation 5.1 where is the plateau (Newtonian) complex viscosity or zero-shear viscosity, and is the complex viscosity at 0.01 Hz (chosen arbitrarily).   The index of association is plotted against molecular weight in Figure 5.22. At high molecular weight, all the polymers show very low association factors, implying that the associations at this length scale do not affect the overall behaviours of the polymers (note that there is one aryl-group per chain). As the molecular weight decreases all the aryl-capped polymers show an increase in their aggregation due to assembling of their chain-ends. There is a rough correlation between the index of association with molecular weight as expected. Important to note is that the native polymer does not show this increase in association over all the molecular weights studied.   The favourable aggregation of aromatic systems usually occurs through off centre parallel or edge-to-face stacking.260 The index of association may be related to the bulkiness of the end-group moiety and thus how well it can move through the polymer matrix, as well as its ability to form one of the preferred stacked structures. However, it is difficult to know how significant these trends are as we have no information about the initial aggregation present within the polymer samples.    178  molecular weight, kg mol-1104 105 106index of association (%)0100200300400PLA-EtPLA-PhPLA-Nap PLA-MePyr PLA-MeAnth Figure 5.22 Index of association vs. molecular weight for aryl-capped PLAs in this study.   The reptation time, c, is defined as the time it takes for the polymer chains to respond to a sudden applied force in order to fully relax after completing its diffusion out of its constrained tube. This parameter can be roughly calculated from the cross-over point between the storage and loss moduli in the frequency sweep experiments and is plotted against molecular weight in Figure 5.23. Reptation behaviour of linear microstructured PLA chains follow the relation, c (PLA-Et) = 1.512 x 10-20 Mw3.4 (indicated by the black line on the plot). Higher molecular weights polymers follow the expected trend for linear PLAs but as the molecular weight decreases the reptation times tend to be higher than expected (Figure 5.23). The presence of chain associations and the formation of a weak network is expected to increase polymer reptation time, as the “apparent” molecular weight increases, therefore, chains will need more time to respond to applied strains. As discussed  179  above, these weak associations are only significant in low molecular weight polymers, which start to show deviation from linear behaviour, the threshold being around 100 kg mol-1.   molecular weight (kg mol-1)104 105 106reptation time, c, (s)10-410-310-210-1100101PLA-PhPLA-NapPLA-MePyrPLA-MeAnthc(PLA-Et) = 1.512 x 10-20 Mw3.4Tref = 150oC Figure 5.23 Reptation time versus molecular weight for aryl-capped PLAs in this study.   5.2.5 Uniaxial Extensional Rheology.   Chain interactions can be probed by measuring the extensional viscosity of polymeric materials. Strain hardening occurs when the deformation time exceeds the terminal relaxation time for relaxation of the polymer chains. Recently our group reported that high molecular weight PLAs do show this strain hardening effect at temperatures below 110 °C.1,22 This technique was further utilized to explore the influence of the aryl end-groups on chain interactions within PLA. Uniaxial extensional tests were carried out at 90 °C at Hencky strain rates of 0.01 to 10 s–1 (appendix C). Figure 5.24a shows that all the materials studied display strain hardening at 10 s–1 strain rates, but this effect is lost at lower strains (Figure 5.24b). Reports of polymers with moieties capable of  180  secondary interactions show that these associations lead to extensive strain hardening.261,262 The functionality in those reported polymers were dispersed throughout the polymer backbone leading to a large number of possible associations. The aryl-capped polymers in this study however, have functionality at only the chain end, and the observed strain hardening is more likely a product of PLA microstructure than the interactions of the chain-ends. To enhance the strain hardening, higher temperatures would impart more chain mobility to allow extensional stress to align the chains. Unfortunately, tests at high temperature were not possible as the materials became too viscous under these conditions.  181  time (s)10-2 10-1 100extenstional viscosity, E+ (Pa.s)103104105106PLA-Et62PLA-Nap65PLA-MeAnth63PLA-MePyr693+a)T = 90 oCH = 10 s-1time (s)10-2 10-1 100 101extenstional viscosity, E+ (Pa.s)103104105106PLA-Et62PLA-Nap65PLA-MeAnth63PLA-MePyr693+b)T = 90 oCH = 1.0 s-1 Figure 5.24 Comparison of tensile stress growth coefficients for selected aryl-capped PLAs (measure of elongational viscosity) at 90 °C as a function of time at Hencky strain rate of a) 10 s-1 and b) 1.0 s-1.   5.2.6 Mechanical properties.   Tensile strength tests were performed on compression moulded samples with a thickness of 0.3-0.5 mm. The results are shown in Figure 5.25 to Figure 5.27. These aromatic electrostatic  182  interactions are very weak (section 5.2.4) thus we did not expect a large difference in the mechanical properties of the aryl-capped PLAs and within error there does not seem to be an enhancement or depreciation of the tensile properties with the incorporation of aryl groups to the chain end of the polymers. These results give more evidence that chain mobility is the determining factor for the onset of aggregation, therefore, at room temperature all the PLAs display comparable mechanical properties. This could be a beneficial property of these materials as the onset of aggregation during extrusion could improve the melt strength of PLA during processing.  14x100 15x100 16x100Tensile strength (MPa)0510152025PLA-EtPLA-PhPLA-NapPLA-MePyrPLA-MeAnthlow Mnhigh Mn All Figure 5.25 Tensile strength for aryl-capped PLAs in this study. Low molecular weight polymers had Mn between 25 and 35 kg mol-1. High molecular weight polymers had Mn between 110 and 130 kg mol-1. “All” combines data for all polymers of the same family.   183  4x100 5x100 6x100Elastic Modulus (MPa)0102030405060PLA-EtPLA-PhPLA-NapPLA-MePyrPLA-MeAnthlow Mnhigh Mn All Figure 5.26 Elastic Modulus for aryl-capped PLAs in this study. Low molecular weight polymers had Mn between 25 and 35 kg mol-1. High molecular weight polymers had Mn between 110 and 130 kg mol-1“All” combines data for all polymers of the same family.   184  22x100 23x100 24x100elongation-at-break (%)0246810121416PLA-EtPLA-PhPLA-NapPLA-MePyrPLA-MeAnth Figure 5.27 Elongation at break (%) for aryl-capped PLAs in this study. Low molecular weight polymers had Mn between 25 and 35 kg mol-1. High molecular weight polymers had Mn between110 and 130 kg mol-1. “All” combines data for all polymers of the same family.   5.3 Conclusions  Immortal ROP was successful in generating aryl-capped PLAs with well-defined molecular weights and low molecular weight distributions. Aryl chain ends fidelity was confirmed using 1H NMR, MALDI-ToF and UV-Vis spectroscopy. Thermal analysis of these PLAs show that end-group association via aromatic-aromatic - interactions are not strong enough to induce crystallization at room temperature. All polymers are amorphous, displaying similar degradation pathways. Rheological results presented give evidence that aggregation of the chain-ends manifest  185  their presence by an increase of viscosity at small frequencies in oscillatory shear due to an increase on the “apparent” molecular weight through chain association. A temperature and frequency dependence on the formation of these associated structures was shown as well as a dependence on the molecular weight of the polymer. There was no indication of a difference between the thermal and mechanical properties of the aryl-capped polymers compared to the native PLA-Et indicating that chain mobility is necessary when forming the aggregated structures within the polymer melt.   5.4 Experimental  General Methods. All the air and moisture sensitive manipulations were carried out in an MBraun glovebox or using standard Schlenk line techniques. A Bruker Avance 300 or 400 MHz spectrometer was used to record 1H NMR spectra. 1H NMR chemical shifts are given in ppm versus residual protons in CDCl3 ( 7.27). Molecular weights, hydrodynamic radii and intrinsic viscosities were determined by GPC-RI using an Agilent liquid chromatograph equipped with an Agilant 1200 series pump and autosampler, three Phenogel 5 μm Narrow Bore columns (4.6 × 300 mm with 500 Å, 103 Å and 104 Å pore size), a Wyatt Optilab differential refractometer, Wyatt tristar miniDAWN (laser light scattering detector) and a Wyatt ViscoStar viscometer. The column temperature was set at 40 °C. A flow rate of 0.5 mL/min was used and samples were dissolved in THF (ca. 4 mg/mL). The measurements were carried out at laser wavelength of 690 nm, at 25 °C. The data were processed using the Astra software provided by Wyatt Technology Corp. Molecular masses were determined using a Bruker Autoflex time-of-flight mass (ToF) spectrometer equipped with MALDI ion source. A differential scanning calorimeter (DSC) Q1000 (TA Instruments) was employed to measure the glass transition (Tg) and melting (Tm) temperatures of the synthesized  186  samples. Approximately 2-3 mg of the samples were weighed and sealed in an aluminium pan. The samples were heated at a rate of 10 °C/min from 25 to 275 °C, then held isothermally for 5 min to destroy any residual nuclei before cooling at 5 °C/min. The transition and melting temperatures were obtained from a second heating sequence, performed at 10 °C/min. Thermogravimetric analysis (TGA) traces were collected on a PerkinElmer Pyris 6 TGA with a nitrogen flow rate of 20 mL/min. Approximately 15 mg of the samples were weighed into a ceramic crucible. The samples were heated at a rate of 20 °C/min from 30 to 500 °C, and the total degradation and degradation onset temperatures were directly determined from the thermographs.  Shear measurements were performed using the MCR 501 rheometer (Anton Paar), equipped with parallel plates, 8 mm in diameter.  The dynamic linear viscoelastic measurements were carried out within the linear viscoelastic regime at temperatures in the range from 70 to 190 °C. The dynamic measurements were conducted in the range of 0.01-100 Hz at a strain of 2%. A gap of 0.5 mm was used to minimize edge effects and ensure a reasonable aspect ratio of plate radius and gap. The samples were melted at 150 °C for at least 3 min to eliminate the residual thermal histories. Uniaxial extensional measurements were performed using the SER-2 extensional fixture attached to an Anton Paar MCR 502 rheometer. Samples with diameters of 16−50 mm and thickness 0.4−0.6 mm were prepared by the same procedure used for shear samples. Individual polymer specimens were then cut to a width of 1.5−3.5 mm. Measurements were conducted at 90 °C at Hencky shear rates of 0.01, 0.1, 1.0, and 10 s-1. Tensile tests were performed using COM-TEN 95 series tensile testing equipment (COM-TEN Industries) at ambient conditions. The films were compressed in a hot press at 160 °C for 15 min before being cooled. Tensile specimens were cut from compression-moulded films. Specimens were cut from the middle portion of the compressed films to avoid edge effects and edge imperfections. A gage length of 40 mm, crosshead speed of  187  25 mm/min and a 40 pound (178 N) capacity of load cell was used for testing all samples. To eliminate specimen slippage from the grips, double adhesive masking tape was used to wrap around the top and bottom portions of the sample. Five tests were run for each sample. The average modulus, tensile stress and elongation at break were calculated from the resultant stress-strain measurements and these are reported below along with standard deviations shown by the plotted error bars.  Materials. THF, toluene, and diethyl ether was taken from an IT Inc. solvent purification system with activated alumina columns and degassed before use. HPLC grade DCM was purchased from Fisher Chemicals and was dried over CaH2, transferred under vacuum and degassed before use. Indium (III) trichloride, 1-naphthol, hydroxymethylanthracene and 1-pyrenemethlaldehyde, were purchased from Sigma-Aldrich and were used as received. Phenol was purchased from Fisher Chemicals and used as received. Racemic lactide was purchased from PURAC America Inc. and was recrystallized thrice from hot dried toluene prior to use. The catalyst [(NNO)InCl]2(-Cl)(-OEt) (A) can synthesized according to previously reported methods.40   Synthesis of hydroxymethylpyrene. 1-Pyrenemethlaldehyde (0.25 g, 0.001 mol) was stirred in 5 mL of ethanol at 80 °C. Sodium borohydride (0.22g, 0.006 mol) was added slowly. This off-white solution was then stirred at 80 °C for one hour. The reaction mixture was then cooled to room temperature and extracted three times with DCM, washed with brine and dried over magnesium sulphate. After concentration under vacuum the product was recrystallized from hot ethanol (0.18 g, 71%). %). 1H NMR (400 MHz, CDCl3): δ 8.42 (d, 1H, Pyr-H), 8.20 (m, 4H, Pyr-H), 8.06 (m 4H, Pyr-H), 5.44 (s, 2H, CH2).  188   Representative immortal ring-opening polymerization. Inside a N2 filled glovebox, appropriate amounts of rac-lactide were dissolved in DCM. Appropriate volumes of a standard solution of chain transfer agent in DCM was added and the solution stirred until homogenous. A standard solution of catalyst A in DCM (10 mg into 1 mL) was the added the reaction mixtures were stirred at room temperature for at least 18 hours before the solvent was evaporated. A small sample of crude polymer was taken for 1H NMR spectroscopic analysis, the remaining polymer was dissolved in minimal DCM (~2 mL) and precipitated with cold methanol. The methanol was decanted off and the precipitation repeated twice. The isolated polymer was then dried under vacuum for at least 24 hours before further analysis.  189  Chapter 6: Synthesis and themorheological analysis of lignin-graft-poly(lactide) copolymers and their blends.  6.1 Introduction  After cellulose, lignin is the second most abundant polymer found in biomass,  and is produced on the millions of tons scale as a by-product of the pulp and paper industry.84,263 The surge of interest to increase the value of this cheap and renewable source of polymeric material69,75,76 has prompted the use of lignin as a filler in a variety of green composites.16,65 This chapter further explores potential modifications to poly(lactide) with the aim of improve processing properties by investigating lignin-graft-poly(lactide) copolymers and their blends; first, by assessing the various synthetic routes to these materials, followed by an analysis of their structure-property relationships.  6.1.1 Lignin-polymer materials  Unmodified lignin/polymer blends and their polymer properties have been a significant field of investigation.81,82,85,193,264,265 In 2000 Alexy et al. prepared a family of blends comprised from low-density poly(ethylene) (LDPE) and up to 30 wt% of lignin. Mechanical tests of these materials showed a 55% decrease in tensile strength as lignin incorporation increased to 30 wt%.83 A similar effect was noted by Doherty et al. who investigated poly(hydroxybutyrate) (PHB)/lignin blends. At lower than 30 wt%, lignin behaved as a plasticizer forming a single phase with PHB and increasing the fluidity of the materials under shears stress. At higher lignin loading, phase separation was observed which decreased the ability to dissipate applied stress leading to an increase in storage modulus, G’, and complex viscosity, |*|.86 Torkelson and co-workers  190  improved lignin dispersion in poly(ethylene) (PE) by using a solid-state shear pulveriser to synthesize their blends. These materials exhibited a greater elongation-at-break and tensile strength compared to previous lignin/PE blends.266  In general, lignin-polymer blends have shown greater thermal stability at high lignin loading, however the overall mechanical properties are often diminished or remain unchanged compared to neat PLA. This depreciation of composite properties has been attributed to poor stress transfer as a result of insufficient compatibilization between lignin and the polymer matrix.65,75-77,263  6.1.2 Lignin graft copolymers  Graft copolymerization is one technique to improve the dispersion of two immiscible polymers. Covalently bonding two polymers improves adhesion of the incompatible phases which can be confirmed by the presence of an intermediate glass transition temperature, Tg. Researchers have successfully grafted many synthetic201,203,267-271 and bio-derived polymers to lignin.198,272   There are two major strategies to synthesize graft copolymers from lignin: grafting-to and grafting-from. In graft-to processes, discrete polymers are synthesized and are covalently bound to lignin in a second step, often utilizing the abundant hydroxyl functionality in the lignin structure (Scheme 6.1). Grafting-from requires the use of lignin (or modified lignin) as a macro-initiator for the polymerization of monomer to form the copolymer in one step (Scheme 6.2).  191  Scheme 6.1 Graft-to synthesis for lignin graft copolymers. Distinct polymers are synthesized and through a second step are covalently bound to lignin.    Scheme 6.2 Graft-from synthesis for lignin graft copolymers where lignin is used as a macro-initiator for polymerization.     Using a grafting-to strategy to modify lignin with low molecular weight PLA, Park et al. showed increased dispersion of lignin within the PLA matrix. However, composites displayed diminished mechanical properties compared to neat PLA.273 In a similar approach, Kim and co-workers synthesized lignin-graft-PLA using methanol-soluble lignin and low molecular weight chloro-terminated PLAs (PLA-Cl). The resulting graft copolymers displayed the highest molecular  192  weights reported for this type of polymer, ranging from 160 – 900 kg mol-1. Analysis of the ash content after thermogravimetric experiments showed high incorporation of lignin and mechanical tests indicated that the grafts had increased tensile strength and modulus compared to the neat polymer.274  This method is advantageous as synthesis begins with well-defined polymers whose properties can be optimized before formation of copolymers. Grafting-to reactions, however, often require harsh reagents which could hinder the widespread industrial application. A more atom-economic synthesis has been adopted by many researchers in this field.   Grafting-from approaches were used by Kai and Loh who polymerized methyl methacrylate with ATRP-modified lignin, generating lignin-graft-PMMA copolymers. These were blended with poly(-caprolactone) (PCL) and electrospun into threads that displayed improved tensile strength and storage modulus compared to native polymers.275   In a leading report, Sattely and co-workers describe the synthesis of well-defined lignin-graft-PLA copolymers through organocatalyzed ROP of lactide. They show that the graft arm length could be controlled by pre-acetylation of lignin and by varying the lignin content in the reaction mixture.194 Other organocatalytic systems were explored to synthesize PCL267 and PLA-lignin276 composites with UV-resistant and antioxidant properties. Most notable was a recent report by Liu et al. who recognized that the high concentration of hydroxy groups constrained the chain lengths of grafted PLA. Selectively alkylating 100% of phenolic hydroxy (OH) and 70% of carboxylic OH (COOH) groups allowed for the synthesis of graft copolymers with the longest reported PLA arm lengths (up to 28 kg mol-1).277   193   In 2013, Averous and co-workers synthesized lignin-graft-PCL copolymers using a simple metal complex, SnOct2. By controlling the CL/OH ratio a variety of copolymers with varying arm lengths were achieved. The copolymer with the shortest arm length contained 50 wt% lignin (Mn = 1.1 kg mol-1).203 High dispersities were observed in the copolymers (Đ = 6.1 - 41.9) and were attributed to highly polydisperse lignin at the core of the grafts. The authors went on to investigate the grafts chain topology using differential scanning calorimetry (DSC) and melt rheological analyses. They concluded that the grafts were highly branched amorphous copolymers from the observed high Tg and G’ values for copolymers with high lignin content.270 Melt rheological analysis of the grafts highlighted the correlation between arm length and rheological response, whereby copolymers with shorter arms and high lignin content displayed elastomeric response and copolymers with longer arms showed thermoplastic response.203 Yu and co-workers also reported structural analysis for a class of lignin-graft-poly(methyl methacrylate-co-butyl acrylate) (lignin-graft-P(MMA-co-BA)) composites. Although thermogravimetric analysis (TGA) results indicated that only 0.2 wt% of lignin had been incorporated into the copolymer, the lower slope for the graft copolymers in the Mark-Houwink plots ([] vs. Mw) implied that the grafts were forming star-like shapes in solution. The grafts also displayed a lower zero-shear viscosity, 0 (2.4 MPa.s vs. 3.4 MPa.s), supporting the solution viscosity results. Mechanical tests demonstrated a 2-fold enhancement in the elongation-at-break compared to the native copolymer.267  He et al. fabricated (PLA)–lignin composites by blending lignin-graft-(PCL-co-PLA)-g-poly(D-lactide) copolymer particles with commercial poly(L-lactide) (PLLA). Light scattering and small angle x-ray scattering (SAXS) studies showed enhanced dispersion of the graft copolymer in the PLLA matrix through stereo-complexation interactions between PLLA and PDLA units.  194  These composites exhibited a staggering 700% increase in their plastic deformation (toughness) as well as associated increases in tensile strength and modulus.198  Though there has been a large focus on this field in the recent years, and an exciting variety of bioderived and compostable materials have been synthesized, many of the recent reports of these lignin graft copolymers focus on determining mechanical response of composites 194,275 or on characterizing specific behaviour (e.g. wettability269 or anti-oxidant behaviour).198 There are many key questions which still need to be addressed, including the influence of unreacted lignin on polymer properties, accurate determination of lignin integration and the structure of graft copolymers. The following sections describe a systematic study of lignin-graft-PLA copolymers, their structure and viscoelastic properties. We focus on fundamental studies of the polymerization behaviour under various conditions by comparing metal catalyzed graft-from, organocatalyzed graft-from and graft-to synthetic routes to these copolymers; and fully characterize the polymers generated as star or cyclic PLAs. The lignin incorporation and melt rheology of the graft copolymers are explored, and finally, a variety of blends are prepared to further understand how the interaction between lignin and the matrix impact blend properties.  6.2 Results and discussion 6.2.1 Synthesis and characterization of lignin-graft-PLAs.  A series of lignin-graft-PLA copolymers were synthesized via a graft-from approach using the simple salt system reported by Hillmyer and Tolman.120,121 Rac-lactide was stirred with 0.2 mol% InCl3, 0.5 mol% NEt3 and 1-35 mol% of lignin†††† in toluene. The resulting mixture was                                                  †††† [OH]lig = 23 mmol g-1 determined via 31P NMR spectroscopy (see section 6.4).   195  heated to 120 °C for 48 hours (Scheme 6.3). Graft copolymers were further purified to elucidate the way in which unreacted lignin influences polymer characterization and properties. Crude reaction mixtures were diluted with DCM and subjected to multiple centrifugation cycles. Polymer containing supernant was collected and the solvent removed under reduced pressure. Copolymers were isolated by precipitation from DCM with MeOH before drying under vacuum for 24 hours (Table 6.1). The polymers were characterized by 1H, 1H-1H COSY (CDCl3, 25 °C) and 31P NMR spectroscopy (CHCl3:pyridine, 25 °C) as well as by IR spectroscopy and gel permeation chromatography (GPC).  Scheme 6.3 Synthesis of lignin-graft-poly(lactide)s via ring-opening polymerization using InCl3 and NEt3.     The formation of PLA is confirmed by the appearance of peaks in the 1H NMR spectrum assigned to the methine proton (~5.20 ppm) and the methyl protons (~1.5 ppm) of the polymer (Figure 6.1). Peaks associated with the polymer chain end are distinguishable from the 1H-1H COSY NMR spectrum. A quartet at ~4.20 ppm was assigned to the methine proton closest to the  196  chain end, as this correlated to the signal for methyl protons of the polymer (Figure 6.2). Arm lengths were calculated from the ratio of polymer to chain-end peaks and we confirmed the literature supposition that changing the LA:OH ratio gives variable arm lengths (Figure 6.3).194,203,277 It should be noted that the size of the PLA arms do not match the theoretical length calculated from the LA:OH ratio (Table 6.1), the observed arm lengths were longer than expected indicating incomplete activation of OH moieties on lignin (see Chapter 4). The PLA arms solubilize grafted lignin, as such more features belonging to lignin can be seen on the 1H NMR spectra. Specifically, the broad aryl peaks around 7 ppm and the methoxy peaks at ~4 ppm (Figure 6.3).194,267,268,273  Table 6.1 Polymerization data from lignin-graft-poly(lactide)s formed via ring-opening polymerization using InCl3 and NEt3.  entry mol% lignin-OH [LA]:[OH]liga conversion (%)b arm length (g mol-1) c Mn, GPC (g mol-1) d Mw, GPC (g mol-1) d Đd 1 0 0 98 31500 115000 133000 1.15 2 5 7.7 96 7370 46600 56300 1.21 3 9 3.9 94 3390 444000 701500 1.66 4 21 1.5 90 1370 199000 450000 2.11 5 33 0.78 90 613 85200 110000 1.30 Reactions were carried out in toluene at 120 °C for 48 hours. a[OH]lig = 23 mmol g-1 determined via 31P NMR spectroscopy.67,194 b Monomer conversions determined by 1H NMR spectroscopy. c Calculated from integration of the polymer and chain end methine protons multiplied by the molecular weight of lactide. d Absolute molecular weights were determined by GPC−MALS−RI (gel permeation chromatography−multi-angle light scattering−refractive index) via Universal Calibration (THF 4 mg mL–1, flow rate = 0.5 mL min–1, dn/dc = 0.040 mL g –1 (calculated from 100% mass recovery, see appendix E)).    197   Figure 6.1 1H NMR spectrum (CDCl3, 25 °C, 400 MHz) for polymer generated from metal-catalyzed system (arm length = 3.2 kg mol-l).   198   Figure 6.2 1H-1H COSY NMR spectrum (CDCl3, 25 °C, 600 MHz) for lignin-graft-PLA generated from metal-catalysed system (arm length = 13 kg mol-1).    199   Figure 6.3 1H NMR spectra (CDCl3, 25 °C, 400 MHz) for polymers with different wt% of lignin, generated from metal-catalyzed system.   31P NMR spectroscopy was employed to further elucidate the types of hydroxy (OH) functionality present in the polymer.67,194 Phosphitylation of the polymer produces various phosphate moieties with distinctive signals on the 31P NMR spectrum (see section 1.1.2). Spectral analysis of polymer samples containing unreacted lignin showed the presence of alkyl and carboxylic OH groups indicative of the formation of PLA (Figure 6.4). A significant concentration of phenolic OH groups, likely from the unreacted lignin, is also observed. The 31P NMR spectra obtained from graft copolymers is absent of phenolic moieties, implying little lignin is actually being incorporated into the copolymers (Figure 6.5).    200   Figure 6.4 31P NMR spectrum (CHCl3/pyridine, 25 °C, 121 MHz) of polymer containing unreacted lignin (arm length = 5 kg mol-1). Derivatized by 2-chloro-4,4,5,5-tetramethyl-1,3,2-dioxaphospholane. Internal standard at 145.16 ppm (cyclohexanol).  201   Figure 6.5 31P NMR spectrum (CHCl3/pyridine, 25 °C, 121 MHz) of polymer generated from metal catalyzed system (arm length = 5 kg mol-1). Derivatized by 2-chloro-4,4,5,5-tetramethyl-1,3,2-dioxaphospholane. Internal standard at 145.16 ppm (cyclohexanol).   Polymers were also characterized by GPC in THF at room temperature. Broadly, increasing the [LA]:[OH]lig gives polymers whose peaks elute at longer retention times indicating lower molecular weight polymers (Figure 6.6). Samples with low lignin loading (high [LA]:[OH]lig) gave peaks with long low molecular weight tails, whereas polymerizations at higher lignin loading (low [LA]:[OH]lig) gave peaks with high molecular weight shoulders. Similar multimodal peak shapes have been reported for both lignin-graft-PMMA23,267 and lignin-graft-PCL203 copolymers. The graft copolymers were also seen to elute out slower than linear chains of similar Mw indicative of polymers with different chain architectures (Figure 6.7).   202  InCl3/NEt3time (min)10 12 14 16 18 20RIU0.00.20.40.60.81.00 wt%1 wt% 5 wt% 10 wt%25 wt%  Figure 6.6 GPC traces for a series of lignin-graft-PLAs generated from the metal catalyzed graft-from approach.  elution time (min)10 12 14 16 18 20differential refractive index  (RIU)01e-52e-53e-54e-55e-56e-58 kg mol-113 kg mol-111 kg mol-113 kg mol-1 Figure 6.7 GPC traces for linear PLA (red lines) and lignin-graft-PLAs (black lines).   The IR spectra for lignin, PLA, and a representative lignin-graft-PLA copolymer are shown in Figure 6.8. The most diagnostic peak in lignin is the broad OH stretch at around 3300 cm-1 resulting from the large number of phenolic and alkyl OH groups within the lignin framework.  203  PLA has a characteristic carbonyl stretch at 1738 cm-1 as well as C-H stretches at 2959 and 3001 cm-1. The spectrum for the graft copolymer shows a carbonyl peak at 1744 cm-1 which has shifted to higher wavenumber compared to native PLA. Averous and co-workers also observed this bathochromic shift which they attributed to hydrogen-bonding between PCL and phenolic moieties on lignin.203 The small OH band at 3542 cm-1 implies the loss of the hydroxyl groups through the grafting process and gives further evidence that a low percentage of lignin is incorporated into the copolymer.   wavenumber (cm-1)1000200030004000Transmittance (%)405060708090100ligninlignin-graft-PLAPLA Figure 6.8 IR spectra of lignin-graft-PLA (long blue dashes), lignin (black line) and PLA (short red dashes).   6.2.2 Effect of lignin source on copolymerization.  A vast array of lignin is currently available on the market, each with different chemical compositions and the potential to generate polymers with different architectures (see section 1.1.2).  204  To determine the robustness of this system, a series of polymerizations were attempted with two commercially available lignins: Indulin AT kraft lignin (IAK) from ingevity™, and alkali kraft lignin (AK) from Sigma Aldrich.   Graft copolymerization was successful with both lignin sources (Table 6.2). Polymers generated with IAK lignin had higher dispersities implying higher lignin incorporation.203 These polymers also followed the trends discussed above where changing the [LA]:[OH]lig ratio generated polymers with varying arm length (Table 6.2, entries 5 to 9). On the other hand, arm lengths for copolymers generated with AK seemed to have less dependence on [LA]:[OH]lig ratio, though this lignin source gave much longer arms then the IAK variant (Table 6.2, entries 3 vs. 9). These observations can be rationalized by inspecting the 31P NMR spectra of the lignin after phosphitylation (Table 6.3). IAK has a higher concentration of hydroxy groups than AK suggesting there are more active sites from which PLA can grow, leading to higher lignin content and shorter PLA arms. The lower OH concentration in AK has the complimentary effect of giving longer polymer chains due to fewer reactive groups. This is an interesting observation, which indicates that these materials can be further optimized and modified by tailoring the specific lignin source. 205  Table 6.2 Polymerization data from lignin-graft-poly(lactide)s formed using two commercially available lignin sources.  entry Lignin  mol% lignin-OH [LA]:[OH]lig a conversion (%)b arm length  (g mol-1)  c Mn, GPC (g mol-1) d Mw, GPC (g mol-1) d Đd 1 IAKe 5 7.7 96 7370 46600 56300 1.21 2 IAK 9 3.9 94 3390 444000 701500 1.66 3 IAK 21 1.5 90 1370 199000 450000 2.11 4 IAK 33 0.78 90 613 85200 110000 1.30 5 AKf 1 32 86 38200 74000 91200 1.26 6 AK 5 7.5 78 17700 52200 60400 1.16 7 AK 9 4.0 73 13900 9090 10100 1.12 8 AK 11 3.1 82 12700 30500 38500 1.28 9 AK 21 1.4 70 5900 32300 36300 1.11 Reactions were carried out using InCl3/NEt3 in toluene at 120 °C for 48 hours. a[OH]lig in mmol g-1 determined via 31P NMR spectroscopy.67,194 b Monomer conversions determined by 1H NMR spectroscopy. cCalculated from integration of the polymer and chain end methine protons multiplied by the molecular weight of lactide.  d Absolute molecular weights were determined by triple detector GPC (gel permeation chromatography) via Universal Calibration (THF 4 mg mL–1, flow rate = 0.5 mL min–1,  dn/dc = 0.040 mL g –1 (calculated from 100% mass recovery, see appendix E)). e IAK Indulin AT kraft lignin f AK = alkali kraft lignin.  Table 6.3 Hydroxy group content for Indulin AT kraft lignin (IAK) and alkali kraft lignin (AK).  Type of OH Indulin AT kraft lignin (IAK) alkali kraft lignin (AK) alkyl OH (150 -146 ppm)a 7.5(0.9) 5.0(0.9) phenolic OH (144 -138 ppm)a 14(1) 9.1(0.7) COOH (136 - 134 ppm)a 1.5(2) 2.3(0.2) Total OHa 23(3) 16(1) Calculated from 31P NMR spectra of 2-chloro-4,4,5,5-tetramethyl-1,3,2-dioxaphospholane derivatized lignin in CHCl3/pyridine. Internal standard at 145.16 ppm (cyclohexanol). aHydroxyl (OH) group content in units of mmol g-1.   Polymerization was also attempted using lignin sludge obtained from an industrial source. Lignin with a final H2O content of 3.8% was obtained after drying at high temperatures under vacuum (Figure 6.9). This lignin was successfully used as a macro-initiator for the ROP of lactide, showing the versatility of this catalyst system to form industrially relevant polymers.    206   Figure 6.9 Lignin sludge obtained from industrial sources was boiled at 140 °C for 4 h and further dried under high vacuum for 24 h. Gravimetric analysis showed 77% decrease in H2O content with a final H2O content of 3.8%.   6.2.3 Other synthetic routes to lignin-graft-PLAs.  A review of the literature does not reveal any reports of the chain architecture of lignin-graft-PLA copolymers. The literature also contains a variety of synthetic routes to these copolymers, each potentially producing materials with different chain topologies. It was thus of interest to generate more families of these copolymers and to characterize and compare their structures and properties.   Triazabicyclodecene (TBD) was used as an organocatalyst to produce a second family of copolymers using a graft-from approach.194 A reaction flask was charged with PLA, TBD and 1-30 mol% of lignin (Scheme 6.4). The reaction mixtures were stirred at 130 °C for 3-4 hours before quenching with 5% acetic acid in DCM, precipitating from cold MeOH and drying under vacuum for 24 hours (Table 6.4).  The polymers generated from this reaction showed the expected arm-length dependence on [LA]:[OH]lig ratio, whereby at low lignin loading long arms were observed (Table 6.4, entry 2 vs. 5). All the polymers showed high dispersities, indicative of lignin incorporation.   207  Scheme 6.4 Synthesis of lignin-graft-poly(lactide)s via ring-opening polymerization using TBD.      Table 6.4 Polymerization data from lignin-graft-poly(lactide)s formed via ring-opening polymerization using TBD.  entry mol% lignin-OH [LA]:[OH]lig a conversion (%)b arm length (g mol-1) c Mn, GPC (g mol-1) d Mw, GPC (g mol-1) d Đd 1 0 0.0 83 19300 38300 51700 1.32 2 4.8 7.6 97 31900 45700 55500 1.22 3 9.0 3.9 70 19500 87000 173000 2.33 4 20 1.5 74 8160 88100 131000 1.48 5 33 0.8 72 9390 70700 96600 1.33 Reactions were carried out in the melt at 130 °C for 3-4 h. a  [OH]lig = 23 mmol g-1 determined via 31P NMR spectroscopy.67,194 b Monomer conversions determined by 1H NMR spectroscopy. c Calculated from integration of the polymer and chain end methine protons multiplied by the molecular weight of lactide. d Absolute molecular weights were determined by triple detector GPC (gel permeation chromatography) via Universal Calibration (THF 4 mg mL–1, flow rate = 0.5 mL min–1, dn/dc = 0.040 mL g –1 (calculated from 100% mass recovery, see appendix E)).   208   A graft-to synthetic strategy, adapted from a 2016 report from Kim and co-workers,274 was used to synthesize a final family of lignin-graft-PLA copolymers (Scheme 6.5). Linear PLAs were synthesized via the immortal ROP of lactide using dinuclear indium complex [(NNO)InCl]2(-Cl)(-OEt) (A) and MeOH. These polymers were reacted with oxalyl chloride producing chloro-terminated PLA (PLA-Cl). Finally, lignin was stirred with K2CO3 in dry DMF for 30 minutes before a solution of PLA-Cl in DMF was added dropwise. The resulting mixture was heated to 50 °C and stirred for one hour before being cooled to room temperature and precipitated from DCM and MeOH (Table 6.5).  Table 6.5 Polymerization data from lignin-graft-poly(lactide)s formed via a graft-to strategy.  entry PLA-Cl Mw (g mol-1) [K2CO3]:[OH]liga mol% lignin arm length  (g mol-1)b Mn, GPC (g mol-1)c Mw, GPC (g mol-1) c Đ c 1 18200 0.158 19 7273 27700 30880 1.11 2 18200 1.01 21 6303 19300 21178 1.10 3 46400 0.988 20 290 30200 33400 1.11 4 46400 0.158 4 268 32700 113000 3.44 5 46400 0.158 8 2460 82500 483000 5.86 6 154000 <0.005 5 NA d 152000 184000 1.21 7 154000 <0.005 9 NA d 153000 172000 1.12 8 154000 0.095 5 NA d 99900 112000 1.19 9 32300 0.768 5 22300 116000 1100000 9.45 Grafting reactions were carried out in DMF, at 50 °C for 60 mins. a[OH]lig in mmol g-1 determined via 31P NMR spectroscopy.67,194 b Calculated from integration of the polymer and chain end methine protons multiplied by the molecular weight of lactide. c Absolute molecular weights were determined by triple detector GPC (gel permeation chromatography) via Universal Calibration (THF 4 mg mL–1, flow rate = 0.5 mL min–1, dn/dc = 0.040 mL g –1 (calculated from 100% mass recovery, see appendix E)). dNo chain end peaks were detected in the 1H NMR spectra.   209  Scheme 6.5 Graft-to synthesis of lignin-graft-poly(lactide)s employed in this study.     The K2CO3:lignin-OH ratio does not appear to have an impact on the formation of graft copolymers. Reactions with 10 fold difference in [K2CO3]:[OH]lig  gave polymers with comparable arm lengths (Table 6.5, entries 1 and 2). The molecular weight of the pre-polymer (PLA-Cl) may have more of a significant impact on graft formation, where the ideal molecular weight of the pre-polymer is between 20 and 50 kg mol-1 (Table 6.5, entries 2 and 3). Graft copolymerization attempted with higher molecular weight pre-polymers do not show a significant difference in their  210  Mw,GPC and display low dispersity values before and after reaction implying unsuccessful grafting to lignin (Table 6.5, entries 6 to 8). Graft-to reactions utilizing an excess K2CO3 compared to [OH]lig, and 5-10 mol% lignin-OH gave the best lignin incorporation (Table 6.5, entries 4, 5 and 9).  6.2.4 Characterization and comparison of polymers generated from the three synthetic methodologies.  The three families of copolymer were further characterized by 1H{1H} NMR spectroscopy and thermogravimetric analysis (TGA). Chain tacticity for the various graft copolymers were determined by analysis of the 1H{1H} NMR spectra. Polymers generated from the InCl3/NEt3 system displayed diminished heterotactic enrichment than that reported for linear chains, Pr = 0.67 (Figure 6.10).1,120 The organocatalyst gave polymers which were mostly atactic, Pm = 0.56 (Figure 6.11). The tacticity was determined for the two pre-polymers and final copolymer in the graft-to synthesis. There was no observed change in tacticity going from PLA to the graft copolymers, each being isotactically enriched with a Pm value of 0.60 (Figure 6.12). By tailoring the synthetic route, researchers may be able to access multiple chain microstructures giving more avenues to applications for these materials.    211   Figure 6.10 1H{1H} NMR spectra (CDCl3, 25 °C, 600 MHz) for polymer generated from the metal catalyst. Pr = 0.67.   Figure 6.11 1H{1H} NMR spectra (CDCl3, 25 °C, 600 MHz) for polymer generated from the organocatalyst. Pm = 0.56.   212   Figure 6.12 1H{1H} NMR spectra (CDCl3, 25 °C, 600 MHz) for polymer generated by a graft-to synthesis. Pm = 0.60.   Lignin has a distinctive thermal degradation profile but does not fully degrade below 500 °C. Thus, the lignin content of the copolymers can be estimated by examining the ash content after thermal exposure.274 Polymers generated from the graft-to synthesis show the highest lignin incorporation, averaging 24% lignin content across the whole family. Of the graft-from catalysts, TBD gave the highest lignin incorporation albeit only averaging 2.9% (Figure 6.13). This was an expected difficulty using lignin as a macro-initiator for ROP of lactide. Chapter 4 described the living and immortal ROP of lactide using aryl initiators, which are less active than alkyl initiators due to lower nucleophilicity. 31P NMR spectroscopic analysis of phosphorus-derivatized lignin shows an abundance of phenolic OH groups (Figure 6.14). The previous study of aryl initiators showed that not all the phenolic OH groups are likely to be activated and thus once a few have initiated polymerization, ring-opening is expected to occur more predominantly in those positions, giving copolymers with low lignin incorporation and fewer arms than expected. Alkylation197 and  213  pre-acylation194 have been used to improve polymer growth however extra steps in the synthesis reduce industrial feasibility of the materials.   TGA was performed on copolymer samples that had been purified from unreacted lignin and those which had not. Figure 6.15 clearly shows that the unreacted lignin erroneously shows higher lignin content in the copolymer, when actually a copolymer/lignin blend is being formed. The melt rheology of lignin-contaminated samples had higher storage moduli and complex viscosity compared to the pure copolymer indicating that lignin is acting as filler and further modifying the properties of the copolymer (Figure 6.16). Awareness of this contamination is important during preliminary studies of these materials as they have a strong impact on flow properties.   Temperature (oC)100 200 300 400 500normalized mass (mg)051015ligninInCl3/NEt31% ash contentTBD2.9 % ash contentGraft-to24 % ash contentlinear PLAlignin-graft-PLA Figure 6.13 Thermogravimetric analysis traces for lignin-graft-PLA copolymers from the various synthetic routes, after removal of excess lignin.     214   Figure 6.14 31P NMR spectrum (CHCl3/pyridine, 25 °C, 121 MHz) of Indulin AT kraft lignin. Derivatized by 2-chloro-4,4,5,5-tetramethyl-1,3,2-dioxaphospholane. Internal standard at 145.16 ppm (cyclohexanol).  Temperature (oC)100 200 300 400 500normalized mass (mg)0246810121416ligninlignin contaminated graftpure graftlignin-graft-PLA14% lignin8% lignin  Figure 6.15 Thermogravimetric analysis traces for pure lignin-graft-PLA and that contaminated with lignin.  215  lignin-graft-PLA80 oCangular frequency, (rad s-1)10-1 100 101 102complex viscostiy, |*| (Pa.s)103104105106pure graft (8 wt% lignin)lignin contaminated graft (14 wt% lignin) Figure 6.16 Complex viscosity, *| vs. angular frequency at 80 °C for pure lignin-graft-PLA copolymers and that lignin contaminated copolymers.   6.2.5 Determining lignin-graft-PLA copolymer topology.  Lignin graft copolymers have been reported to have different chain structures. Glasser278 and Yu267 both showed evidence for the formation of star-like copolymers, whereas Averous reported their lignin-graft-PCL polymers to be branched, forming 3-dimensional hydrogen-bonding networks at high lignin incorporation.203 An overlooked aspect of these graft copolymerizations is that under normal conditions, both InCl3/NEt3 and TBD catalyst systems require an external alcohol additive to generate linear polymers in a controlled manner (usually benzyl alcohol (BnOH)).120,121,279 Under alcohol-free conditions, analogous systems to those used in this study have been shown to form cyclic PLAs.280-282 Graft-to synthesis can also generate cyclic polymers via the self-condensation of PLA-Cl pre-polymers under basic conditions (Scheme 6.6). As such, there are a variety of possible chain structures which could be present in the final polymer, impacting polymer properties. Thus, to determine the extent of cyclic-PLA formation during graft  216  copolymerization, lignin-free control reactions were performed under the same conditions as described in sections 6.2.1 and 6.2.3 above (Table 6.6).   Only PLAs generated from well-defined dinuclear indium complex, [(NNO)InCl]2(μ-Cl)(μ-OEt) (A), showed good agreement between theoretical and GPC molecular weights (Table 6.6, entries 11 to 14). Polymers generated from lignin-free reactions with InCl3/NEt3 and TBD were less controlled, giving higher-than-expected molecular weights in most cases (Table 6.6, entries 1 to 10).  Scheme 6.6 Formation of cyclic-PLAs from the self-condensation of chloro-terminated PLA (PLA-Cl).    217  Table 6.6 Polymerization data for control reactions in the absence of lignin.  entry catalyst conversion (%)a Mn,theo (g mol-1) b arm length (g mol-1)c Mn, GPC (g mol-1) d Mw, GPC (g mol-1)d Đd 1e InCl3/NEt3 93 118000 20300 44500 62600 1.41 2e InCl3/NEt3 91 111000 NAh 27600 28900 1.05 3e InCl3/NEt3 96 76800 NAh 166000 181000 1.09 4e InCl3/NEt3 98 81100 21500 71300 80600 1.13 5e InCl3/NEt3 97 77800 41600 158000 186000 1.18 6f TBD 88 13700 NAh 56100 66500 1.19 7f TBD 76 13300 21400 31800 50200 1.58 8f TBD 83 12800 22900 35900 57400 1.60 9f TBD 89 12800 14400 19200 20300 1.05 10g TBD 84 12700 20600 59900 77600 1.29 11g A 99 7200 7300 10300 10400 1.01 12g A 98 14400 20000 18500 18600 1.01 13g A 99 37400 NAh 46400 47000 1.01 14g A 98 73400 NAh 81700 82800 1.01 a Monomer conversions determined by 1H NMR spectroscopy. b Calculated from: [LA]/[cat] × Mn,LA. c Calculated from integration of the polymer and chain end methine protons multiplied by the molecular weight of lactide. d Absolute molecular weights were determined by triple detector GPC (gel permeation chromatography) via Universal Calibration (THF 4 mg mL–1, flow rate = 0.5 mL min–1, dn/dc = 0.040 mL g –1 (calculated from 100% mass recovery, see appendix E)). e Reactions were carried out in toluene at 120 °C for 48 h. f Reactions were carried out in the melt at 130 °C for 3-4 h. g Reactions were carried out in the DCM at room temperature for 16 h. h No chain end peaks were detected in the 1H NMR spectra.   Lower intrinsic viscosity compared to linear polymers has been observed in studies of both star-like267,278,283 and cyclic284,285 polymers. The intrinsic viscosities for PLAs generated from lignin-free control polymerizations are plotted in Figure 6.17. Polymers from control polymerizations with complex A are seen to fall on a trend line higher than that of all the polymers studied, and have a Mark-Houwink slope of 0.75, which is in agreement with reports of other linear PLAs.1,163,166,286 Lignin-free reactions using InCl3/NEt3 and TBD, on the other hand, produce polymers which fall into second trend of PLAs with lower intrinsic viscosity than that of linear PLAs (Figure 6.17, red circles).   218  Mw (g mol-1)103 104 105 106 107intrinsic viscosity, |*|10-1100101102103linear PLAs from A0 wt% lignin1 wt% lignin5 wt% lignin10 wt% lignin25 wt% ligningraft-to polymerslinear PLA|*| = 0.0184Mw0.749cyclic PLA|*| = 0.0011Mw0.941lignin-graft-PLA|*| = 3.6378Mw0.166 Figure 6.17 Intrinsic viscosity vs. molecular weight for products from the three synthetic strategies routes to the graft copolymers.   Figure 6.17 shows that there are two species generated from graft co-polymerization reactions, each with lower intrinsic viscosity than linear PLAs. These results support our hypothesis that the copolymerization conditions can generate PLAs with two possible structures, stars and cycles; however, their formation seems to be independent of synthetic route. Also evident from the data is that cyclic PLAs are preferentially generated at low lignin loading (≤ 5wt% lignin) and star-graft copolymers generated at high lignin loading (≥ 10 wt%). At intermediate lignin loading, it is assumed that an indistinguishable mixture of cyclic PLAs and star-graft copolymers are being generated. Analysis of the GPC traces for cyclic and star-graft copolymers show both species elute at longer times than linear polymers (Figure 6.18).  219  time (min)10 12 14 16 18 20RIU0.00.20.40.60.81.0linear18.5 kgmol-1star graft16.3 kgmol-1cycle15.8 kgmol-1time (min)10 12 14 16 18 20RIU0.00.20.40.60.81.0linear45.5 kgmol-1star graft32.1 kgmol-1cycle36.4 kgmol-1 Figure 6.18 GPC traces (THF 4 mg mL–1, flow rate = 0.5 mL min–1, dn/dc = 0.040 mL g –1 ) for linear PLA (blue dashed lines), cyclic PLA (black dash-dot line) and lignin-graft-PLAs (black line).   To further explore the cause of the variable chain structures, MALDI-ToF mass spectral analysis of the polymers generated from control polymerizations with no lignin was conducted (Figure 6.19). A major set of peaks separated by 72 mass units with no calculated chain end was observed, indicative of the formation of cyclic species. A second distribution is calculated to have a methoxy (OMe) chain end, possibly resulting from ring-opening of the cyclic polymers when precipitating with MeOH during work-up. Inspection of the MALDI-ToF spectra of polymers produced at low [OH]lig reveal similar cyclic polymer distributions (Figure 6.20). Whereas mass spectra for polymers generated from reactions using higher [OH]lig, show distributions that correspond to polymers with chain-end masses of greater than 59 m/z, which could be from fragmentation of the lignin core during ionization (Figure 6.21). However, because lignin is highly disordered and fragments in a variety of ways,287 it is difficult to determine exactly what these ends groups are. Regardless, these results allowed us to assign the three trends seen on the intrinsic viscosity vs. Mw plots to linear PLAs, cyclic PLAs and star-shaped lignin-graft-PLAs.  220     Figure 6.19 MALDI-ToF mass spectrum for polymer generated from a control graft-from polymerization (Mw,GPC = 63 kg mol-1).  221   Figure 6.20 MALDI-ToF mass spectrum for polymer generated from a graft-from polymerization with low [OH]lig (Mw,GPC = 78 kg mol-1).   Figure 6.21 MALDI-ToF mass spectrum for polymer generated from a graft-from polymerization with high [OH]lig (Mw,GPC = 18 kg mol-1).   222   Rheological analysis was employed to compare of the linear viscoelastic moduli for the chain topologies generated from these reactions. Isothermal small-angle oscillatory shear (SAOS) experiments at angular frequencies from 0.01 to 100 Hz were conducted on the three classes of polymers. Both cyclic285,288-290 and star283,291 polymers have access to more relaxation modes giving them higher fluidity compared to linear polymers. The deviation of the slopes in the high frequency region of the complex viscosity, |*| vs. angular frequency,  plot indicate more shear thinning behaviour for the graft copolymers and cycles compared to linear PLA (Figure 6.22). Recent simulation studies by Li et al. showed that although both loops and the dense core of stars can retard deformation, cycles have a larger restoring force, making them less deformable than stars.292 This effect is also seen in the melt rheology of PLA cycles and star grafts whereby cycles have a higher viscosity compared to the star-graft copolymers. Zero-shear viscosities for linear PLAs are one order of magnitude higher those that of cyclic polymers and two orders of magnitude higher than those of graft copolymers (Table 6.7 and Figure 6.23). These results are in line with a report by Schäler and Saawächter, which showed that cyclic PCLs have 2-fold lower viscosity than linear PCLs,284 and Yu et al. who show a 5-fold decrease in the intrinsic viscosities of their lignin-graft-PMMA star copolymers.267   223  angular frequency,  (rad s-1)10-1 100 101 102complex viscosity, | (Pa.s)103104105106linear PLA, 19 kg mol-1star lignin-graft-PLA, 36 kg mol-1cyclic PLA, 63 kg mol-180 oC Figure 6.22 Complex viscosity, |*| (Pa.s) vs. angular frequency, , (rad s-1) for linear PLA (Mw = 19 kg mol-1), star lignin-graft-PLA (Mw = 36 kg mol-1) and cyclic PLA (Mw = 63 kg mol-1).  Mw (g mol-1)102 103 104 105 1060, (Pa.s)103104105106107linear PLA cyclic PLA star lignin-graft-PLA []0,linear = 1 x10-10Mw3.4[]0,cyclic = 8 x10-7Mw2.4[]0,graft = 2 x10-3Mw1.7 Figure 6.23 Zero-shear viscosity,  vs. weight-averaged molecular weight, Mw, for linear PLAs (red triangles), cyclic PLAs (black circles) and star lignin-graft-PLA copolymers (blue squares). 224  Table 6.7 Polymerization data for linear PLA, cyclic PLA and star lignin-graft-PLAs used in rheological analyses. entry type catalyst lignin mol% lignin-OH ash content a arm length (g mol–1)b Mw, GPC (g mol–1) c Đ c 0  (× 105 Pa.s)d 1e linear A -- -- -- 20000 18600 1.01 2.3 2e linear A -- -- -- NAh 47000 1.01 11 3e linear A -- -- -- NAh 82800 1.01 57 4 f cyclic InCl3/NEt3 -- 0 0.9 20300 62600 1.41 3.0 5 g cyclic TBD -- 0 0.6 26900 35400 1.32 0.42 6f cyclic InCl3/NEt3 AK 1 1.4 15200 76000 1.44 5.0 7 g cyclic TBD AK 5 2.3 19500 15800 1.26 0.12 8 g cyclic TBD AK 9 1.8 27300 36400 1.68 0.85 9 f star graft InCl3/NEt3 AK 4 2.7 17300 16300 1.13 0.025 10 g star graft TBD IAK 5 4.7 20300 28300 1.46 0.70 11g star graft InCl3/NEt3 AK 13 5.5 1750 32200 1.17 0.040 12g star graft TBD IAK 26 3.0 6270 65200 1.30 0.15 a  Ash content calculated from mass of residue after thermal treatment to 450 °C. b Calculated from integration of the polymer and chain end methine protons multiplied by the molecular weight of lactide. c Absolute molecular weights were determined by triple detection GPC (gel permeation chromatography) via Universal Calibration (THF 4 mg mL–1, flow rate = 0.5 mL min–1, dn/dc = 0.040 mL g–1 ). d Zero-shear viscosities obtained from the limits of master curves generated from isothermal frequency sweeps. e Reactions were carried out in the DCM at room temperature for 16 h. f Reactions were carried out in toluene at 120 °C for 48 h. g Reactions were carried out in the melt at 130 °C for 3-4 h.. h No chain end peaks were detected in the 1H NMR spectra.   The results in this section confirm that the three polymerization methods do in fact generate both cyclic polymers and star-like graft copolymers. The different chain topologies can be determined from GPC analysis. The GPC traces of star grafts show that they seem to have high molecular weight shoulders which could be linear contaminates and cyclic PLAs have very broad traces with long tails. It is possible that the cycles could also be contaminated with linear chains (Figure 6.24). Polymerizations with little to no lignin loading generate cyclic PLAs. Under these conditions, the combined effect of low OH concentration, inactive phenolic moieties and added  225  base leads to the self-polymerization of lactide to form cyclic polymers.281,282 The proposed mechanism for the formation of these cyclic PLAs is shown in Scheme 6.7 below. In the metal-catalyzed synthesis, InCl3 activates the monomer to nucleophilic attack by the base additive which performs the initial ring opening event. This generates a zwitterion, which at low [OH], can go on to ring-open more lactide monomers. On the other hand, at higher lignin loading, there are enough OH groups present for transesterification to occur giving lignin-graft-PLAs which can go on to propagate ring-opening. cyclic PLAtime (min)10 12 14 16 18 20 22RIU0.00.20.40.60.81.062.6 kg mol-176.0 kg mol-115.8 kg mol-1 36.4 kg mol-1star graft PLAtime (min)10 12 14 16 18 20 22RIU0.00.20.40.60.81.032.1 kg mol-116.3 kg mol-1b)a) Figure 6.24. GPC traces for a) cyclic PLAs and b) star graft copolymers generated in this study.  226  Scheme 6.7 Proposed mechanism for the formation of cyclic PLAs in the absence lignin.282     6.2.6 Rheological analysis of lignin-graft-poly(lactide)s.  Having confirmed the formation of various chain structures, further SAOS experiments were conducted on the cyclic and star-like graft copolymers. An 8-mm parallel plate rotational rheometer was used at a constant strain amplitude of 2% over frequencies ranging from 0.01 to 100 Hz. The time-temperature superposition principle was applied to isothermal experiments in the temperature range of 50-110 °C generating master curves of linear viscoelastic (LVE) moduli, G’ and G”, as well as the complex viscosity at the reference temperature of 80 °C.   227  angular freguency, , (s-1)10-2 10-1 100 101 102 103 104 105 106G', G" (Pa)101102103104105106107G'G" G'G"G'G"cyclic PLA 15.8 kg mol-1linear PLA 18.6 kg mol-1lignin-graft-PLA 16.3 kg mol-1 Figure 6.25. Viscoelastic moduli (G’-filled symbols, G”-open symbols) vs. angular frequency,  (Pa.s) for linear PLA (red triangles, Table 6.7, entry 1 ), cyclic PLAs (black circles, Table 6.7, entry 7) and star lignin-graft-PLA copolymers (blue squares, Table 6.7, entry 9). angular freguency, , (s-1)10-3 10-2 10-1 100 101 102 103 104 105 106G', G" (Pa)101102103104105106G'G" G'G"G'G"cyclic PLA 36.4 kg mol-1linear PLA 47.0 kg mol-1lignin-graft-PLA 32.2 kg mol-1 Figure 6.26. Viscoelastic moduli (G’-filled symbols, G”-open symbols) vs. angular frequency,  (Pa.s) for linear PLA (red triangles, Table 6.7, entry 2 ), cyclic PLAs (black circles, Table 6.7, entry 8) and star lignin-graft-PLA copolymers (blue squares, Table 6.7entry 11).   228  Figure 6.25 and Figure 6.26 compare the LVE moduli for linear PLA, cyclic PLA and star lignin-graft-PLA copolymers in this study. Linear PLAs exhibit characteristic thermoplastic behaviour for a linear polymer above its entanglement molecular weight, whereby G’ increases to an elastic plateau at intermediate deformation frequencies to then transition into the glassy zone at high frequencies. Comparing the LVE plots of cyclic and star graft PLAs with similar GPC molecular weights, it is clear that the frequency dependence on the loss and storage moduli are different to that of linear PLAs. Both cyclic PLAs and star graft PLAs show faster terminal relaxation for the whole molecular weight range studied. As discussed above cyclic polymers have a large restoring force, making them less deformable than stars.292 Studies of ultra-pure cyclic poly(styrene)s (PS) also reported faster relaxation compared to linear counterparts and an extended relaxation regime.293 Recently Matsushita and co-workers showed that even high molecular weight ring chains form few intermolecular entanglements leading to a lack of a rubbery plateau.294 The cyclic PLAs in this study are likely to be contaminated with linear chains, either from thermal ring-opening during analysis295 or as a by-product of polymerization. Linear chains have been shown to retard the relaxation of cyclic polymers through threading interactions leading to an enhancement in the rubbery plateau. As the molecular weight of the cyclic copolymers increases, the plateau becomes more defined, which may be attributed to more entanglement between the rings or the presence of higher molecular weight linear chain contaminates (Figure 6.27).   229  cyclic PLAangular freguency, , (s-1)10-3 10-2 10-1 100 101 102 103G', G" (Pa)100101102103104105106Table 6.7, entry 715.8 kg mol-1Table 6.7, entry 462.6 kg mol-1Table 6.7, entry 535.4 kg mol-1 Figure 6.27. Viscoelastic moduli (G’- filled symbols, G”- open symbols) vs. angular frequency,  (Pa.s) for cyclic PLAs in this study.  Master curves for star lignin-graft-PLA with GPC molecular weights of 16.2, 32.2 and 65.2 kg mol–1 all showed loss modulus dominating over the whole frequency range, G” > G’ (Figure 6.28). At medium to high frequencies, the storage and loss moduli become almost equivalent, G’ ≈ G”, implying gel-like behaviour. The terminal zone was reached at very low frequencies, where LVE moduli display the characteristic slopes of G′ ∝ ω2 and G″ ∝ ω. Another interesting feature of the master curves are the lack of a rubber plateau in the storage modulus at high deformation frequencies. A similar result was observed for 3- and 6-arm PHBs which showed that increasing the arm length from 3 to 6 reduced the intensity of the rubber plateau.283 The loss of the rubber plateau in the graft copolymers is likely due to the large number of shorts arms growing from the various hydroxyl groups in the lignin structure, which are densely packed with few chain  230  entanglements, thus reducing the ability of the polymer to form a rubbery network. It is hypothesized that by increasing the PLA graft arm length would allow for thermoplastic behaviour, however there are currently very few reports of lignin-graft-PLA copolymers with arm lengths greater than 30 kg mol-1.277   star lignin-graft-PLAangular freguency, , (s-1)10-3 10-2 10-1 100 101 102 103 104 105 106 107G', G" (Pa)100101102103104105106107108Table 6.7, entry 916.3 kg mol-1Table 6.7, entry 1265.2 kg mol-1Table 6.7, entry 1132.2 kg mol-1 Figure 6.28. Viscoelastic moduli (G’-filled symbols, G”-open symbols) vs. angular frequency,  (Pa.s) for star lignin-graft-PLA copolymers in this study.   The dynamics of star and ring blends are currently under-explored in the literature., though some assumptions can be made from studies of cyclic/linear284,293,294,296 and star/linear blends.23,291,292,295,297-301 If the graft copolymer has short chains then it is likely to have lower viscosity if contaminated with cyclic PLAs as they will not entangle with each other. Graft copolymers with long chain stairs could have retardation from threading but this could be balanced  231  by arm retraction relaxation modes. From the result presented in this study we estimate these graft copolymers to be likely have dense cores with short PLA arms.   Thermal transitions were probed using temperature ramp experiments at constant deformation strain (2%) and frequency (0.5 Hz). Both cyclic and star lignin-graft-PLAs showed glass transitions at ~45 °C. At lower temperatures, the G’ dominates and thus copolymers display glassy behaviour. Above 45 °C the polymers display gel-like behaviours up to ~ 65 °C after which fluid response dominates (Figure 6.29).   cyclic PLATable 6.7, entry 462.6 kg mol-1temperature (oC)40 60 80 100 120G', G" (Pa)101102103104105106107108109G'G"star lignin-graft-PLATable 6.7, entry 916.3 kg mol-1temperature (oC)40 60 80 100 120G', G" (Pa)10-1100101102103104105106107108109G'G"a) b) Figure 6.29 Temperature ramp experiments from 30 – 130 °C for a) cyclic PLA (Table 6, entry 4) and b) lignin-graft-PLA copolymers (Table 6, entry 10).    232  6.2.7 Lignin-graft-PLLA/PLLA blends.‡‡‡‡  A graft-from strategy was used to synthesize lignin-graft-PLLA with arm lengths of 12 kg mol-1 and 14 kg mol-1 (section 6.2.1). High molecular weight PLLA was also synthesized to use in blends. LLA (20 g) was stirred with InCl3, NEt3 and BnOH in toluene at 80 °C for 48 hours. After this time the reaction mixture was concentrated under reduced pressure and the polymer precipitated three times from cold MeOH before being dried for 48 hours under vacuum (Table 6.8).  Blends with varying compositions were made by dissolving the appropriate amounts of PLLA and lignin-graft-PLLA in DCM. Following solution casting and drying under vacuum at 40 °C for 8 h, the blends were compression molded at 165 °C to make discs for rheological and mechanical measurements (Table 6.8). Table 6.8 Results for graft-from copolymerizations of 90% L-lactide with InCl3/NEt3.  entry conversiona (%) arm length  (g mol-1) b Mn, GPC (g mol-1) c Mw, GPC (g mol-1) c Đc 1d 80 NAe  210000  214000 1.02  2 92 14200 33900  40300 1.19 3 77 12300 18900  22700 1.20  Standard conditions: InCl3/NEt3 and 5 wt% lignin stirred with 90% L-LA at 120 °C.  a Monomer conversion determined by 1H NMR spectroscopy. b Calculated from integration of the polymer and chain end methine protons multiplied by the molecular weight of lactide. c Absolute molecular weights were determined by triple detector GPC (gel permeation chromatography) via Universal Calibration (THF 4 mg mL–1, flow rate = 0.5 mL min–1,  dn/dc = 0.040 mL g –1  (calculated from 100% mass recovery)).  d 2 eq. BnOH used as co-initiator, reaction stirred at 80 °C for 2 days. e No chain end peaks were detected in the 1H NMR spectra.                                                   ‡‡‡‡ To model commercially available PLAs while maintaining solubility for characterization purposes, 90% L-LA was used as a monomer for the following experiments.  233   To determine the thermal stability of the lignin-graft-PLLA/PLLA blends under shearing conditions, experiments were conducted over 60 minutes at a constant temperature and angular frequency of 180 °C and 0.5 Hz, respectively. The complex modulus, |G*| is plotted against time in Figure 6.30. Blends show similar or increased |G*| with up to 10 wt% incorporation of graft copolymer. Further increase of graft copolymer concentration in the blend causes a marked decrease in thermal stability. Blends containing graft copolymers with shorter arm lengths (12 kg mol-1) also showed a decrease in thermal stability.   The time-Temperature superposition principle was applied to isothermal SAOS experiments performed at temperatures ranging from 80-180 °C to generate master curves of the complex viscosity at a reference temperature of 160 °C. Zero-shear viscosities obtained from these master curves and are shown in Table 6.9. The viscosity of blends was observed to decrease with increasing graft copolymer content, though this decrease is less distinct with the 10 wt% blend. Interestingly, decreasing the copolymer arm length does not seem to impact viscosity compared to the neat polymer. Under oscillatory-shear flow conditions, star-shaped lignin-graft-PLLA copolymers act mainly as a plasticizer to the PLLA matrix, decreasing the modulus and viscosity of the blends.   234  500 1000 1500 2000 2500 3000 3500 40001031041051060 wt%1 wt%10 wt%25 wt%10 wt%, 12 kgmol-1Complex modulus, |G*|, (Pa)time (s)lignin-graft-PLLA/PLLA Figure 6.30 Complex modulus, |G*|  vs. time at 180 °C for 60 minutes for lignin-graft-PLLA/PLLA blends.  Table 6.9 Lignin-graft-PLLA/PLLA blend properties. entry wt% graft graft arm length (g mol-1)a  0  (160 °C) (kPa.s)b Tensile strength (MPa) Elastic modulus (MPa) Elongation at break (%) 1 0 NAe  133 10.9 ±1.9 305 ±50 8.0  ± 0.5 2 1 14200 42.1 19.7 ±1.0 328±50 9.1±1.5 3 10 14200 80.0 36.4 ±6.7 754 ±99 6.6  ± 0.7 4 25 14200 39.0 13.8±2.4 291±42 7.5±1.8 5 10 12300 139 17.6±2.2 357±61 8.5±1.3 a Calculated from integration of the polymer and chain end methine protons multiplied by the molecular weight of lactide. b Determined from isothermal SAOS experiments at 160 °C.    Mechanical tests were performed on compression-molded samples, and the results are shown in Figure 6.31 to Figure 6.33. Similar to melt rheology results, incorporation of graft copolymer up 10 wt% improves the tensile strength and elastic modulus of the resulting blends. Undesirably, increases in tensile strength are associated with a decrease in elongation-at-break. Blend of grafts  235  with shorter arms do not show significant improvement to the mechanical properties of the resulting blends.  Figure 6.31 Tensile strength of lignin-graft-PLLA/PLLA blends. Blends with graft arm length = 14 kg mol-1 () and Blends with graft arm length = 12 kg mol-1().    Figure 6.32 Elastic moduli of lignin-graft-PLLA/PLLA blends. Blends with graft arm length = 14 kg mol-1 () and Blends with graft arm length = 12 kg mol-1().    051015202530354045500% 10% 20% 30%Tensile Strength (MPa)wt% graft copolymer01002003004005006007008009000% 5% 10% 15% 20% 25% 30%Elastic Modulus (MPa)wt% graft copolymer 236   Figure 6.33 Percent elongation at break for lignin-graft-PLLA/PLLA blends. Blends with graft arm length = 14 kg mol-1 () and Blends with graft arm length = 12 kg mol-1().    6.3 Conclusions  Lignin-graft-PLAs were successfully synthesized via two synthetic strategies. The highest lignin incorporation was achieved using a graft-to approach. The organocatalyst, TBD, gave the highest lignin incorporation of the two graft-from syntheses. Graft arm length was found to depend on a variety of factors such as: lignin source and catalysis route and showed that determining [OH]lig is important when predicting the arm length. Avoiding very low [OH]lig was found to be particularly important, as cyclic PLAs can be formed as an unwanted by-product of polymerization. Cyclic PLAs can act as plasticizers, impacting polymer properties. The influence of unreacted lignin on copolymer rheology was also shown to cause errors when predicting melt properties.  0246810120% 10% 20% 30%% Elongation at breakwt% graft copolymer 237   We investigated the melt rheology of cyclic PLAs and star lignin-graft-PLAs. Both of these materials showed viscous behaviour at low deformation frequency followed by a region of gel-like behaviour at intermediate deformation frequencies. It is predicted that by further increasing the graft arm length, thermoplastic behaviour could be imparted to the material. However, this requires careful control of [OH]lig, potentially through pre-alkylation of the lignin. A family of lignin-graft-PLLA/PLA blends was formulated. It was shown that increasing the amount of graft copolymer up to 10 wt% in PLLA blends improves tensile strength and elastic modulus; however, increasing graft copolymer content past this point does not further improve these properties. Unfortunately, an associated decrease in percent elongation at break is also observed for the 10 wt% blend.  6.4 Experimental  General Methods. All the air and moisture sensitive manipulations were carried out in an MBraun glovebox or using standard Schlenk line techniques. A Bruker Avance 300 or 400 MHz spectrometer was used to record 1H NMR spectra. 1H NMR chemical shifts are given in ppm versus residual protons in deuterated CDCl3 (δ 7.27). Molecular weights, hydrodynamic radii and intrinsic viscosities were determined by GPC-MALS-RI-Viscometer using an Agilent liquid chromatograph equipped with a Agilant 1200 series pump and autosampler, three Phenogel 5 μm Narrow Bore columns (4.6 × 300 mm with 500 Å, 103 Å and 104 Å pore size), a Wyatt Optilab differential refractometer, Wyatt tristar miniDAWN (laser light scattering detector) and a Wyatt ViscoStar viscometer. The column temperature was set at 40 °C. A flow rate of 0.5 mL/min was used and samples were dissolved in THF (ca. 4 mg/mL). The measurements were carried out at laser wavelength of 690 nm at 25 °C. The data were processed using the Astra™ software provided  238  by Wyatt Technology Corp. dn/dc values for cyclic and lignin-graft-PLA copolymers were calculated using 100% mass recovery program on the Astra 6.0™ software and an average value of 0.04 was used for molecular weight determination (see appendix E). Molecular masses were determined using a Bruker Autoflex time-of-flight (TOF) mass spectrometer equipped with MALDI ion source. Thermogravimetric analysis (TGA) traces were collected on a PerkinElmer Pyris 6 TGA with a nitrogen flow rate of 20 mL/min. Approximately 15 mg of the samples were weighed into a ceramic crucible. The samples were heated at a rate of 20 °C/min from 30 to 500 °C, and the degradation onset temperatures were directly determined from the thermographs. Shear measurements were performed using the MCR 501 rheometer (Anton Paar), equipped with parallel plates, 8 mm in diameter. The dynamic linear viscoelastic measurements were carried out within the linear viscoelastic regime at temperatures in the range from 50 to 190 °C. The dynamic measurements were conducted in the range of 0.01-100 Hz at a strain of 2%. A gap of 0.5 mm was used to minimize edge effects and ensure a reasonable aspect ratio of plate radius and gap. The samples were melted at 150 °C for at least 5 min to eliminate the residual thermal histories. Tensile tests were performed using COM-TEN 95 series tensile testing equipment (COM-TEN Industries) at ambient conditions. Tensile specimens were cut from compression-moulded films. Specimens were cut from the middle portion of the compressed films to avoid edge effects and edge imperfections. A gage length of 40 mm, crosshead speed of 25 mm/min and a 40 pound (178 N) capacity of load cell was used for testing all samples. To eliminate specimen slippage from the grips, double adhesive masking tape was used to wrap around the top and bottom portions of the sample. For each sample five tests were run. The average modulus, tensile stress and elongation at break were calculated from the resultant stress-strain measurements and these are reported below along with standard deviations shown by the plotted error bars.   239   Materials. Toluene was taken from an IT Inc. solvent purification system with activated alumina columns and degassed before use. DCM and DMF were purchased from Fisher Chemical and were dried over CaH2, transferred under vacuum and degassed before use. CDCl3 was purchased from Cambridge Isotope Laboratories Inc. and dried over CaH2, transferred under vacuum and degassed through three freeze−pump−thaw cycles before use. MeOH and K2CO3 were purchased from Fisher Chemicals and were used as received. Alkali kraft lignin was purchased from Sigma-Aldrich and was dried at 60 °C under vacuum for 48 hours before use. Indulin AT kraft lignin was a gift from ingevity™ and was dried at 60 °C under vacuum for 48 hours before use. Indium (III) trichloride, NEt3, and 2-chloro-4,4,5,5-tetramethyl-1,3,2-dioxaphospholane (TMDP) were purchased from Sigma-Aldrich and were used as received. Oxalyl chloride was purchased from Alfa Aesar and was used as received. Racemic and L-lactide were purchased from PURAC America Inc. and were recrystallized three times from hot dry toluene prior to use. [(NNO)InCl]2(-Cl)(-OEt) (A) was synthesized according to previously reported methods.40  Phosphitylation of lignin for 31P NMR spectroscopic analysis. Phosphitylation of lignin and lignin-graft-PLA copolymers were performed based on modification of the method described by Sattely et al.67,194 A solvent mixture composed of chloroform (CHCl3) and pyridine (1.6:1.0 v/v) was prepared and dried over molecular sieves prior to use. This solution was used for the preparation of relaxation reagent and internal standard solution (chromium(III) acetylacetonate, 5.0 mg/mL and cyclohexanol, 10.0 mg/mL). A total of 30 mg of sample was dissolved in a CHCl3-pyridine solvent (800 μL) followed by the addition of the relaxation reagent and internal standard solution (100 μL) and TMDP (100 μL). The resulting mixtures were left to stir for 24 hours to ensure complete derivatization prior to 31P NMR analysis.   240   Representative graft-from copolymer synthesis using InCl3/NEt3. Under an N2 atmosphere, a vacuum adapted flask was equipped with a magnetic stir bar and was charged with rac-lactide (1000 mg, 6.94 mmol) and stirred in 5 mL of toluene. Lignin ([OH]lig = 23 mmol g-1), InCl3 (2.8 mg, 0.013 mmol), and triethylamine (3.4 μL, 0.025 mmol) were then added and the volume made up to 15 mL. The reaction mixture was stirred at 120 °C for 48 h under N2.  Solvent was removed under reduced pressure and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in 15 mL of DCM and subjected to centrifugation. The supernant was collected and another 15 mL aliquot of DCM added to the solid. This was repeated at least three times. The supernant was combined and concentrated under reduced pressure. Polymers were then precipitated using an excess of cold methanol. The light brown copolymer was then dried under vacuum for 48 h.  Representative graft-from copolymer synthesis using TBD. Under an N2 atmosphere, a vacuum adapted flask was equipped with a magnetic stir bar and was charged with rac-lactide (1000 mg, 6.94 mmol).  Lignin ([OH]lig = 23 mmol g-1) and TBD (10.3 mg, 0.078 mmol) were then added and the reaction mixture were heated to 130 °C for 3-4 h under N2.  Solvent was removed under reduced pressure and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in 15 mL of DCM and subjected to centrifuge. The supernant was collected and another 15 mL aliquot of DCM added to the solid. This was repeated at least three times. The supernant was combined and concentrated under reduced pressure. Polymers were then precipitated using an excess of cold methanol. The light brown copolymer sample was then dried under vacuum for 48 h.  241   Synthesis of linear PLA for graft-to copolymer formation. A 20 mL scintillation vial was equipped with a magnetic stir bar. Rac-lactide (1000 mg, 6.94 mmol) was dissolved in 2 mL of DCM and to this a solution of [(NNO)InCl]2(-Cl)(-OEt) (A) (1.3 mg, 0.001 mmol) in 2 mL of DCM was added. The solution was stirred at room temperature for 16 h.  Solvent was removed under reduced pressure and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in minimal DCM and precipitated using excess of cold MeOH at least three times to remove residual catalyst. The polymer sample was then dried under vacuum overnight.   Synthesis of PLA-Cl for graft-to copolymer formation. PLA (500 mg, Mn = 81700 kg mol-1, Đ = 1.01) was dissolved in 5 mL of DCM. To this, a solution of oxalyl chloride in DCM (0.204 M) was added and the mixture stirred at room temperature overnight. After this time the reaction mixture was concentrated under reduced pressure and the polymer precipitated with cold MeOH.   Representative graft-to copolymer synthesis. Lignin (5 wt% with respect to PLA-Cl) and K2CO3 (1 eq. with respect to PLA-Cl) were stirred in DMF for 30minutes. After this time, a solution of PLA-Cl (200 mg, Mn = 80900 kg mol-1, Đ = 1.01) in 5 mL DMF was added dropwise. The resulting mixture was heated to 50 °C for 1 hour before being cooled to room temperature. The reaction mixture was then poured into a beaker containing 1:1 MeOH:H2O to precipitate the graft copolymer. This was dissolved in minimal DCM and precipitated again with cold MeOH. The final brown copolymer was then dried under vacuum for 48 hours.  Synthesis of high molecular weight PLLA for blends.  A vacuum adapted flask was equipped with a magnetic stir bar and was charged with L-lactide (20 000 mg, 139 mmol) and stirred in 50 mL of toluene.  InCl3 (17 mg, 0.077 mmol), benzyl alcohol (7.4 μL, 0.069 mmol) and  242  triethylamine (17.7 μL, 0.138 mmol) were then added and the volume made up to 60 mL. The reaction mixture was stirred at 80 °C for 48 h under N2.  Solvent was removed under reduced pressure and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in minimal DCM and precipitated using excess of cold methanol at least three times to remove residual catalyst. The polymer was then dried under vacuum for 48 h.  Representative preparation of lignin-graft-PLLA/PLLA blend. To a homogeneous solution of PLLA (2700 mg) in 45 mL of DCM, was added lignin-graft-PLLA (300 mg in 15 mL of DCM) and the solution stirred for 30 mins. The solution was then poured into a petri dish and the solvent evaporated overnight. After this time, the films were removed from their dishes and dried under vacuum at 40 °C for 18 h. The blends were compression molded at 165 °C to make discs for rheological and mechanical measurements.  243  Chapter 7: Conclusion and future directions  The aim of this thesis was to probe catalysis and polymer synthesis in order to gain control over polymer properties. The polymerization behaviour of well-defined dinuclear indium complexes towards meso-LA was investigated in Chapter 2. From these studies, I gained insight into which stereocontrol mechanisms are acting in these systems. Attention was focused on catalyst family A and its related compounds, 1-6. Defect peak analysis of polymers generated from both alkoxide and aryloxide initiators showed that ESC is the major mechanism affording stereoselectivity in this system. Interestingly, regardless of the nucleophilicity of the bridging species or the isomer of lactide, this family of complexes retains their dimeric structure during polymerization. Analysis of solid-state structures showed that these metal complexes are held together by an important hydrogen bonding interaction between the secondary amine of the ligand and the terminal chlorine on the metal centre. This result has continued to present itself throughout investigation of this catalyst system by the Mehrkhodavandi group and has prompted the exploration of new ligand architectures which can bind two metal centres. The formation of chelated complexes discussed in Chapter 4 indicates the limitations of using diols as co-initiators for polymerization; a surprising result as triols such as, 1,3,5-tris(hydroxymethyl)-benzene (THMB), have been successful co-initiators for this reaction.283,300 The reactivity of aryl initiators also provides useful information for ligand design. Complexes with electron deficient initiators (4 and 5) were the least active towards lactide polymerization, giving an indication of the vitality of avoiding the use of electron rich phenols as ligands to ensure that the ligands do not participate or compete in polymerization with lactide.   To gain control of polymer properties, I targeted efficient catalyst systems which would yield PLAs with controllable Mw and Đ so we would be able to precisely relate structure to properties.  244  By exploiting the systems described in Chapters 2 and 4, the controlled synthesis of both stereoregular and end-functionalized PLAs was achieved. The field was lacking an accurate comparison of stereoregular PLAs. Most of the analysis of the material properties of PLAs was focused on highly isotactic or isotactic stereocomplex PLAs. In Chapter 3, the assumption that PLA tacticity has a similar effect on properties as observed with other polymers such as PP and PS was challenged and found to be false. PLA is unique in its physical properties due to strong dipole-dipole interactions between polymer chains. These interactions are magnified between regions of high isotacticity (and to a lesser extent heterotacticity). Moderately syndiotactic PLAs displayed diminished melt properties compared to PLA of other microstructures, albeit with greater thermal stability. There is a lot of potential for researchers to apply this information to the development of methods to further control and quantify PLAs during industrial production. Moreover, this information allows researchers to explore more rewarding directions when applying this material and opens the door to using combinations of stereoregular PLAs to act as internal fillers and plasticizers within the PLA matrix, generating interesting composites comprised of just one polymer family.  End-group aggregation of aryl-capped PLAs was shown to be induced under shearing conditions in Chapter 5. Though the onset of aggregation was slow, this effect could prove to be useful under creeping conditions. Experiments in the non-linear viscoelastic regime could further elucidate the benefits of these end-group associations. Another under-explored feature of these polymers was their absorption spectra. Even high molecular weight polymers with relatively low chromophore concentrations (as polymer chain-ends) were strongly absorbing. There is potential for these materials to be used as compostable coatings for novelty applications or devices. Further investigation of the spin- or drop-casting of these PLAs is needed as well as how annealing could  245  further align the chain-ends to enhance their photochemical properties. It would also be of importance to consider the bio-toxicity of the chromophore chain-ends.  It became of interest to extend my research and use other bio-sourced materials in order to bolster developing green industries by creating new materials with appealing properties. The last seven years has seen the emergence of researchers exploring lignin/PLA materials. From the reactivity studies described in Chapters 2 and 4, I recognized the value of understanding how catalytic systems are behaving in order to impart control of the materials generated. In Chapter 6, is it shown that proper analysis of polymerization products is vital, otherwise the resultant material properties will not be predictable. In the specific graft-from processes investigated, I show that under certain conditions, cyclic PLAs are generated instead of the desired graft copolymers. The impact of these cycles on material response is an area still in need of investigation. However, the melt rheological properties of lignin-graft-PLAs were analyzed and depend on the PLA arm length. Below 30 kg mol-1 viscous flow or gel-like behaviour dominates. Studies on the barrier properties of these materials of lignin-graft-PLAs have recently been reported.267,276 Lignin composites have also been electrospun into fibers.268,275 As these grafts can be processed at relatively low temperatures (< 60 °C), these composite materials may make valuable fibers, sheets or coatings for use in agricultural, biomedical or packaging applications. Thus, it would be of interest to explore the ability of these materials to be solution cast or spin coated onto different substrates. Depolymerization and pre-alkylation of lignin in situ could be performed to further increase lignin incorporation and potentially enhance barrier properties.  Preliminary studies were conducted in an attempt to synthesize thermally responsive lignin/PLA composites using Diels-Alder chemistry (Scheme 7.1). 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Figure A.2 a) GC-MS trace of meso-lactide after purification (<1% rac-lactide) b) mass spectral data for peak eluting at 5.75 minutes.   a) b) b) a)  257  Appendix B   B.1 Isothermal time sweep experiments time (s)101 102 103 104Complex Modulus, |G*| (Pa)102103104105106350-het-60350-het-60155-het-60155-het-60 Figure B.1 Complex modulus, |G*| , vs. time for het-60 polymers of various molecular weights.  time (s)101 102 103 104Complex Modulus, |G*| (Pa)103104105106295-het-96295-het-96228-het-96148-het-96 Figure B.2 Complex modulus, |G*| , vs. time for het-96 polymers of various molecular weights.   258  time (s)101 102 103 104Complex Modulus, |G*|, (Pa)103104105224-syn-72224-syn-72279-syn-72 Figure B.3Complex modulus, |G*| , vs. time for syn-72 polymers of various molecular weights.  time (s)101 102 103 104Complex Modulus, |G*|, (Pa)102103104105119-iso-72119-iso-72 67-iso-72103-iso-72 Figure B.4Complex modulus, |G*| , vs. time for iso-72 polymers of various molecular weights  259  B.2 Van Gurp-Palmen plots  103104105106020406080 182-het-96 295-het-96 147-het-96phase angle,  Complex modulus, |G*| (Pa)  Figure B.5Van Gurp-Palmen plots of het-96 polymers with varying molecular weights.  103104105106020406080 279-syn-72 147-syn-72 224-syn-72phase angle,  Complex modulus, |G*| (Pa)  Figure B.6Van Gurp-Palmen plots of syn-72 polymers with varying molecular weights.  260  B.3 Uniaxial extension experimental plots 10-210-1100101102103104105106107108109 H = 0.01 s-1 H = 0.1 s-1 H = 1.0 s-1 H= 10 s-1 3++ E (Pa.s)time (s)155-het-6070 oC0.01 0.1 1 10 100 1000 10-210-1100101102103104105106107108155-het-6090 oC H = 0.01 s-1 H = 0.1 s-1 H = 1.0 s-1 H= 10 s-1 3++ E (Pa.s)time (s)0.01 0.1 1 10 100 1000 10-210-1100101102103104105106107155-het-60110 oC H = 0.01 s-1 H = 0.1 s-1 H = 1.0 s-1 H= 10 s-1 3++ E (Pa.s)time (s)0.01 0.1 1 10 100 1000 Figure B.7 Elongational viscosity as a function of time at Hencky strain rates from 0.01 to 10s-1 for atactic PLA.   261  10-210-1100101102103104105106107108109182-het-9670 oC H = 0.01 s-1 H = 0.1 s-1 H = 1.0 s-1 H= 10 s-1 3++ E (Pa.s)time (s)0.01 0.1 1 10 100 100010-210-1100101102103104105106107108182-het-9690 oC H = 0.01 s-1 H = 0.1 s-1 H = 1.0 s-1 H= 10 s-1 3++ E (Pa.s)time (s)0.01 0.1 1 10 100 100010-210-1100101102103104105106107108182-het-96110 oC H = 0.01 s-1 H = 0.1 s-1 H = 1.0 s-1 H= 10 s-1 3++ E (Pa.s)time (s)0.01 0.1 1 10 100 1000 Figure B.8 Elongational viscosity as a function of time at Hencky strain rates from 0.01 to 10s-1 for heterotactic PLA.   262  10-210-1100101102103103104105106107108147-syn-7270 oC H = 0.01 s-1 H = 0.1 s-1 H = 1.0 s-1 H= 10 s-1 3++ E (Pa.s)time (s)0.01 0.1 1 10 100 100010-210-1100101102103103104105106107147-syn-7290 oC H = 0.01 s-1 H = 0.1 s-1 H = 1.0 s-1 H= 10 s-1 3++ E (Pa.s)time (s)10-210-1100101102103104105106147-syn-7270 oC H = 0.01 s-1 H = 0.1 s-1 H = 1.0 s-1 H= 10 s-1 3++ E (Pa.s)time (s)  Figure B.9 Elongational viscosity as a function of time at Hencky strain rates from 0.01 to 10s-1 for syndio-enriched PLA.   263  10-210-1100101102103102103104105106107108109119-iso-7190 oC H = 0.01 s-1 H = 0.1 s-1 H = 1.0 s-1 H= 10 s-1 3++ E (Pa.s)time (s) Figure B.10 Elongational viscosity as a function of time at Hencky strain rates from 0.01 to 10s-1 for isotactic stereo-gradient PLA. 264  Appendix C   C.1 Characterization of complexes 1-6 in the solution state   Figure C.1 1H NMR (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhOMe) (1).   265   Figure C.213C{1H} NMR (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhOMe) (1).   266   Figure C.3 1H – 1H COSY (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhOMe) (1).   267   Figure C.4 1H NMR (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhMe) (2).  268   Figure C.513C{1H} NMR (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhMe) (2).  269   Figure C.6 1H – 1H COSY (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhMe) (2).  270   Figure C.71H – 13C HMBC (CDCl3, 25 °C, 40 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhMe) (2).  271   Figure C.8 1H NMR (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhH) (3).  272   Figure C.9 13C{1H} NMR (CDCl3, 25 °C, 400 MHz) spectrum (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhH) (3).  273   Figure C.101H – 1H COSY (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhH) (3).   274   Figure C.11 1H – 13C HMBC (CDCl3, 25 °C, 400 MHz) spectrum (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhH) (3).   275   Figure C.12 1H NMR (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhBr) (4).   276   Figure C.13 13C{1H} NMR (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhBr) (4).   277   Figure C.14 1H – 1H COSY (CDCl3, 25 °C, 400MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhBr) (4).  278   Figure C.15 1H – 13C HMBC (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhBr) (4).   279   Figure C.16 1H NMR (CDCl3, 25 °C, 400MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhNO2) (5).   280   Figure C.17 13C{1H} NMR (CDCl3, 25 °C, 400MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhNO2) (5).   281   Figure C.18 1H – 1H COSY (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhNO2) (5).  282   Figure C.19 1H – 13C HMBC (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(μ-OPhNO2) (5).  283   Figure C.20 13C{1H} NMR (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2(μ-Cl)(2,μ-diClPh) (6).  284   Figure C.21 1H – 1H COSY (CDCl3, 25 °C, 400MHz) spectrum of (±)-[(NNO)InCl]2)(μ-Cl)(2,μ-diClPh) (6).  285   Figure C.221H – 1H NOESY (CDCl3, 25 °C, 400 MHz) spectrum of (±)-[(NNO)InCl]2)(μ-Cl)(2,μ-diClPh) (6).  286  C.2 Characterization of complexes 2-6 in the solid state  Table C.1 Selected crystallographic parameters.   2 3 4 5 6 Empirical formula C60H93Cl3In2N4O3  C69.5H103Cl3In2N4O3  C69.5H102BrCl3In2N4O3 C52H82Cl3In2N5O5 C73H102Cl4In2N4O4 Formula weight  1254.37  1378.54  1457.44  1252.90 1471.02 T (K) 90(2) 90(2)   90(2)  90(2) 90(2) a (Å) 18.461(5)  17.272(2)  18.451(2) 17.447(2)  10.5118(4) b (Å)  20.589(5)  21.834(3)  21.673(3)  18.137(2)  24.3030(10) c (Å) 19.119(5)  18.317(2)  18.847(2)  21.872(3)  30.4559(11) α (deg)  90  90  90  111.589(2)  90 β (deg) 110.389(4)  95.278(2)  111.583(3)  102.894(3)  93.687(2) γ (deg) 90  90  90  98.066(3)  90 Volume (Å3) 6811(3)  6878.2(15)  7008.2(15)  6081.1(13)  7764.4(5) Z  4  4 4  4  4 Crystal system  monoclinic  monoclinic  monoclinic triclinic monoclinic Space group  P21/c P21/c   P21/c   P-1 P21/c dcalc (g/cm3) 1.223 1.331    1.381  1.369 1.258 μ/mm-1  0.835  0.834 1.389 1.002  0.777 F(000)  2616.0  2884.0  3020.0   2588.0   3064.0   Crystal size/mm3  0.339 × 0.042 × 0.032  0.804 × 0.014 × 0.012   0.91 × 0.15 × 0.13   0.42 × 0.12 × 0.11   0.382 × 0.289 × 0.088 Radiation  Mo Kα (λ = 0.71073)  MoKα (λ = 0.71073)  MoKα (λ = 0.71073) MoKα (λ = 0.71073)  MoKα (λ = 0.71073) 2Θ range for data collection/°  3.02 to 45.16  4.84 to 54.946  2.374 to 51.664 2.374 to 51.664  2.146 to 60.184  287   2 3 4 5 6 Index ranges  -19 ≤ h ≤ 19, -22 ≤ k ≤ 22, -20 ≤ l ≤ 20  -22 ≤ h ≤ 22, -28 ≤ k ≤ 27, -23 ≤ l ≤ 23  -22 ≤ h ≤ 21, -26 ≤ k ≤ 26, -17 ≤ l ≤ 23 -21 ≤ h ≤ 21, -22 ≤ k ≤ 22, -26 ≤ l ≤ 26  -14 ≤ h ≤ 14, -34 ≤ k ≤ 29, -42 ≤ l ≤ 39 Total no. reflections  30942  68285  62150 104569  70649 No. independent reflections (Rint) 8863 [Rint = 0.1479, Rsigma = 0.1418]  15706 [Rint = 0.0742, Rsigma = 0.0632]  13290 [Rint = 0.0548, Rsigma = 0.0749] 23169 [Rint = 0.1043, Rsigma = 0.1138]  22667 [Rint = 0.0416, Rsigma = 0.0494] Data/restraints/parameters  8863/4/773  15706/0/721  13290/0/730  23169/0/1275  22667/0/833 GOF  1.045 1.003  1.041  1.008  1.089  Final R indexes [I>=2σ (I)]  R1 = 0.0882, wR2 = 0.2034 R1 = 0.0410, wR2 = 0.0899 R1 = 0.0497, wR2 = 0.1015 R1 = 0.0546, wR2 = 0.0952 R1 = 0.0490, wR2 = 0.1246 Final R indexes [all data]  R1 = 0.1846, wR2 = 0.2556 R1 = 0.0690, wR2 = 0.1036 R1 = 0.0932, wR2 = 0.1196 R1 = 0.1182, wR2 = 0.1138 R1 = 0.0618, wR2 = 0.1350 Largest diff. peak/hole / e Å-3  1.57/-1.17 1.26/-1.21 3.06/-2.07 1.39/-1.75 1.63/-0.76 a R1 = Σ ||Fo| - |Fc|| / Σ |Fo|; b wR2 = [Σ(w(Fo2 - Fc2)2)/Σ w(Fo2)2]1/2.   288  C.3 In situ living ring-opening polymerization data with catalysts 1-4.    Figure C.23 Reactions were carried out in an NMR tube at 25 °C. 1,3,5-trimethoxybenzene (TMB) was used as internal standard. All reactions were carried out with 50 equivalents of LA in CD2Cl2 at 25 °C and followed by 1H NMR spectroscopy. [A] = 0.0046 M, [LA] = 0.22 M The value of kobs was determined from the slope of ln[LA] vs. time (shown in red), averaged from at least three experiments.   Figure C.24 Reactions were carried out in an NMR tube at 25 °C. 1,3,5-trimethoxybenzene (TMB) was used as internal standard. All reactions were carried out with 50 equivalents of LA in CD2Cl2 at 25 °C and followed by 1H NMR spectroscopy.  [1] = 0.0039 M, [LA] = 0.19 M. The value of kobs was determined from the slope of ln[LA] vs. time (shown in red), averaged from at least three experiments.   y = -1.9(1)E-03x + 8.4E-01R² = 9.9E-01-2.5-2-1.5-1-0.500.510 500 1000 1500 2000 2500 3000ln[LA]time (s)[(NNO)InCl]2(μ-Cl)(μ-OEt) (A)y = -8.5(8)E-04x + 1.6E+00R² = 9.9E-01-2-1.5-1-0.500.511.50 1000 2000 3000 4000 5000 6000ln[LA]time (s)[(NNO)InCl]2(μ-Cl)(μ-OPhOMe) (1)  289   Figure C.25 Reactions were carried out in an NMR tube at 25 °C. 1,3,5-trimethoxybenzene (TMB) was used as internal standard. All reactions were carried out with 50 equivalents of LA in CD2Cl2 at 25 °C and followed by 1H NMR spectroscopy [2] = 0.0043 M, [LA] = 0.21 M. The value of kobs was determined from the slope of ln[LA] vs. time (shown in red), averaged from at least three experiments.    Figure C.26 Reactions were carried out in an NMR tube at 25 °C. 1,3,5-trimethoxybenzene (TMB) was used as internal standard. All reactions were carried out with 50 equivalents of LA in CD2Cl2 at 25 °C and followed by 1H NMR spectroscopy [3] = 0.0037 M, [LA] = 0.22 M. The value of kobs was determined from the slope of ln[LA] vs. time (shown in red), averaged from at least three experiments.   y = -5.3(2)E-04 + 1.2E+00R² = 9.9E-01-1.5-1-0.500.510 1000 2000 3000 4000 5000 6000ln[LA]time (s)[(NNO)InCl]2(μ-Cl)(μ-OPhMe) (2) y = -3.3(2)E-04x + 1.6E+00R² = 9.9E-0100.20.40.60.811.20 1000 2000 3000 4000 5000 6000ln[LA]time(s)[(NNO)InCl]2(μ-Cl)(μ-OPhH) (3) 290   Figure C.27 Reactions were carried out in an NMR tube at 25 °C. 1,3,5-trimethoxybenzene (TMB) was used as internal standard. All reactions were carried out with 50 equivalents of LA in CD2Cl2 at 25 °C and followed by 1H NMR spectroscopy. [4] = 0.0041 M, [LA] = 0.22 M. The value of kobs was determined from the slope of ln[LA] vs. time (shown in red), averaged from at least three experiments.y = -1.1(5)E-04x + 1.8E+00R² = 9.9E-0100.20.40.60.811.20 2000 4000 6000 8000 10000 12000 14000ln[LA]time (s)e)        [(NNO)InCl]2(μ-Cl)(μ-OPhBr) (4) 291  Appendix D   D.1 Master curves for aryl capped poly(lactide)s   angular frequency, , (rads-1)10-2 10-1 100 101 102 103 104 105 106 107G', G" (Pa)101102103104105106107complex viscosity, |*| (Pa.s)10-1100101102103104105106G' G"|*|angular frequency, , (rads-1)10-2 10-1 100 101 102 103 104 105 106 107G', G" (Pa)100101102103104105106107complex viscosity, |*| (Pa.s)10-1100101102103104105G'G"|*|PLA-Et121Tref = 150 oCPLA-Et68Tref = 150 oCb)a) Figure D.1 Master curve of the linear viscoelastic moduli, G′ and G″, and complex viscosity, |η*| (Tref = 150 °C), for a) PLA-Et68 b) PLA-Et121 polymers.   292  Phen-PLA121Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106 107G', G" (Pa)102103104105106complex viscosity, |*| (Pa.s)10-210-1100101102103104105106G'G"|*|Phen-PLA37Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106 107G', G" (Pa)100101102103104105106complex viscosity, |*| (Pa.s)10-1100101102103104G'G"|*|Phen-PLA61Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106 107 108G', G" (Pa)101102103104105106107complex viscosity, |*| (Pa.s)10-210-1100101102103104105G'G"|*|Phen-PLA57Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106 107G', G" (Pa)101102103104105106complex viscosity, |*| (Pa.s)10-1100101102103104G'G"|*|a) b)c) d) Figure D.2 Master curve of the linear viscoelastic moduli, G′ and G″, and complex viscosity, |η*| (Tref = 150 °C), for a) PLA-Ph37 b) PLA-Ph57 c) PLA-Ph61 and (d) PLA-Ph121 polymers.  293  Nap-PLA63Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106G', G" (Pa)101102103104105106complex viscosity, |*| (Pa.s)100101102103104G'G"|*|PLA-Nap36Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105G', G" (Pa)100101102103104105106complex viscosity, |*| (Pa.s)100101102103G'G"|*|angular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106G', G" (Pa)101102103104105106complex viscosity, |*| (Pa.s)10-210-1100101102103104105G'G"|*|PLA-Nap190Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106 107G', G" (Pa)102103104105106complex viscosity, |*| (Pa.s)10-210-1100101102103104105106G'G"|*|a)b)PLA-Nap155Tref = 150 oCb)d) d) Figure D.3 Master curve of the linear viscoelastic moduli, G′ and G″, and complex viscosity, |η*| (Tref = 150 °C), for a) PLA-Nap36 b) PLA-Nap63 c) PLA-Nap155 and (d) PLA-Nap190 polymers.   294  angular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106 107G', G" (Pa)100101102103104105106complex viscosity, |*| (Pa.s)10-310-210-1100101102103104G'G"|*|MePyr-PLA72Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106 107G', G" (Pa)101102103104105106107complex viscosity, |*| (Pa.s)10-1100101102103104G'G"|*|MePyr -PLA135Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106 107G', G" (Pa)101102103104105106107complex viscosity, |*| (Pa.s)10-1100101102103104105106G'G"|*|MePyr-PLA155Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106G', G" (Pa)101102103104105106complex viscosity, |*| (Pa.s)10-1100101102103104105106G'G"|*|MePyr-PLA39Tref = 150 oCa) b)c) d) Figure D.4 Master curve of the linear viscoelastic moduli, G′ and G″, and complex viscosity, |η*| (Tref = 150 °C), for a) PLA-MePyr39 b) PLA-MePyr72 c) PLA-MePyr135 and (d) PLA-MePyr155 polymers.  295  angular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105 106G', G" (Pa)100101102103104105complex viscosity, |*| (Pa.s)10-1100101102103104G'G"|*|MeAnth-PLA80Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105G', G" (Pa)101102103104105106complex viscosity, |*| (Pa.s)101102103104G'G"|*|MeAnth-PLA149Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105G', G" (Pa)102103104105106complex viscosity, |*| (Pa.s)100101102103104105106G'G"|*|MeAnth-PLA188Tref = 150 oCangular frequency, (rads-1)10-2 10-1 100 101 102 103 104 105G', G" (Pa)102103104105106complex viscosity, |*| (Pa.s)101102103104105106G'G"|*|MeAnth-PLA44Tref = 150 oCa) b)c) d) Figure D.5 Master curve of the linear viscoelastic moduli, G′ and G″, and complex viscosity, |η*| (Tref = 150 °C), for a) PLA-MeAnth44 b) PLA-MeAnth80 c) PLA-MeAnth149 and (d) PLA-MeAnth188 polymers.   296  D.2 Uniaxial extension experimental plots time (s)10-2 10-1 100 101 102E+ (Pa.s)103104105106107H = 0.01 s-1H = 0.1 s-1H = 1.0 s-1H = 10 s-13+Et-PLA62T = 90oC Figure D.6 Tensile stress growth coefficient (measure of elongational viscosity) as a function of time at Hencky strain rate of 0.01 to 10 s−1 at 90 °C for Et-PLA62 time (s)10-2 10-1 100 101 102E+ (Pa.s)103104105106107H = 0.01 s-1H = 0.1 s-1H = 1.0 s-1H = 10 s-13+PLA-Nap65T = 90oC Figure D.7 Tensile stress growth coefficient (measure of elongational viscosity) as a function of time at Hencky strain rate of 0.01 to 10 s−1 at 90 °C for Nap-PLA65  297  time (s)10-2 10-1 100 101 102E+ (Pa.s)103104105106107H = 0.01 s-1H = 0.1 s-1H = 1.0 s-1H = 10 s-13+PLA-MeAnth63T = 90oC Figure D.8 Tensile stress growth coefficient (measure of elongational viscosity) as a function of time at Hencky strain rate of 0.01 to 10 s−1 at 90 °C for MeAnth-PLA63  time (s)10-2 10-1 100 101 102E+ (Pa.s)103104105106107H = 0.01 s-1H = 0.1 s-1H = 1.0 s-1H = 10 s-13+time (s)10-2 10-1 100 101 102E+ (Pa.s)103104105106107H = 0.01 s-1H = 0.1 s-1H = 1.0 s-1H = 10 s-13+PLA-MePyr69T = 90oCEt-PLA62T = 90oCa) Figure D.9 Tensile stress growth coefficient (measure of elongational viscosity) as a function of time at Hencky strain rate of 0.01 to 10 s−1 at 90 °C for MePyr-PLA62.     298  Appendix E   E.1  Calculation of dn/dc values  lignin-graft-PLAexperimental time (min)0 10 20 30 40differential refractive index uni (RIU)-0.20.00.20.40.60.81.01.21.99 g/mL4.98 g/mL7.96 g/mL9.95 g/mL Figure E. 1 GPC-RI traces for THF solutions of a lignin-graft-PLA copolymer (arm length = 17 kg mol-1).   concentration (g/mL)0.000 0.002 0.004 0.006 0.008 0.010 0.012refractive index0.00000.00010.00020.00030.0004lignin-graft-PLAdn/dc = 0.041(0.003) Figure E. 2 Plot of refractive index vs. concentration for THF solutions of a lignin-graft-PLA copolymer (arm length = 17 kg mol-1).   299  cyclic PLAexperimental time (min)0 10 20 30 40 50differential refractive index units (RIU)-0.20.00.20.40.60.81.01.21.99 g/mL4.98 g/mL7.96 g/mL9.95 g/mL Figure E. 3 GPC-RI traces for THF solutions of a cyclic PLA (Mn = 61 kg mol-1).   concentration (g/mL)0.000 0.002 0.004 0.006 0.008 0.010 0.012refractive index-0.00010.00000.00010.00020.00030.00040.0005cyclic PLAdn/dc = 0.039(0.003) Figure E. 4 Plot of refractive index vs. concentration for THF solutions of a cyclic PLA (Mn = 61 kg mol-1).  

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