Bioplastic Blends Incorporating Polyamide-11

BIOPLASTIC BLENDS INCORPORATING POLYAMIDE-11

by David Ruehle

A thesis submitted to the Faculty and the Board of Trustees of the Colorado School of

Mines in partial fulfillment of the requirements for the degree of Master of Science

(Chemical Engineering).

Golden, Colorado Date ______

Signed:______David A. Ruehle

Signed: ______Dr. John R. Dorgan Thesis Advisor

Golden, Colorado Date ______

Signed: ______Dr. David Marr Professor and Head Department of Chemical and Biological Engineering

ABSTRACT

Biorenewable polyamide-11 (PA11/ Nylon-11) is melt blended with partially biorenewable polyamide-6,10 (PA 6,10/ Nylon 6,10) to produce thermoplastic blends of varying renewable content. Mechanical and thermal properties of these blends are characterized by differential scanning calorimetry (DSC), thermo gravimetric analysis (TGA), dynamic mechanical thermal analysis (DMTA), tensile testing, and impact testing. Thermal properties provide insight into the structure of the blends. DSC thermograms show melting point depression for the PA 6,10 crystals with increasing PA 11 composition. This observed melting point depression indicates the two polyamides are fully miscible in the melt. There is a minimum in overall crystallinity at the 75wt%/25wt% PA11/PA6,10 blend TGA shows onset degradation temperature changes monotonically for the blends from 419 ˚C for PA 11 to 437 ˚C for PA 6,10. Mechanical properties of the blends show intermediate values compared to the homopolymers. The room temperature storage modulus of the homopolymers are 0.828 GPa for PA 11 and 1.173 GPa for PA 6,10. Blends have storage moduli of 0.847 GPa, 0.974 GPa, and 1.042 GPa for the 75/25, 50/50, and 25/75 blend compositions respectively. The heat distortion temperature (HDT) is 52 ˚C for PA 11 and 76 ˚C for PA 6,10; blends have HDTs of 57 ˚C, 60 ˚C, and 54 ˚C for the 75/25, 50/50, and 25/75 blend compositions, respectively. Young’s moduli show a modified monotonic behavior with blend composition. Young’s modulus for PA 11 is 1.3 GPa and 2.5 GPa for PA 6,10. Moduli of the blends were measured as 1.9 GPa, 2.1 GPa, and 2.5 GPa for the 75/25, 50/50, 25/75 blends. Impact strength of the blends increased with increasing PA 11 content from 43 J/m for the PA 11 homopolymer, to 51 J/m, 66 J/m, 72 J/m, and 71 J/m for the PA 6,10. Morphological properties are investigated using microscopy and x-ray techniques. Wide angle x-ray scattering (WAXS) measurements show that both crystals of PA 11 and PA 6,10 form for all blend compositions. Small angles x-ray scattering (SAXS) shows an increase in the long spacing for the polymer lamellae when samples are annealed at

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185 ˚C. These results indicate that phase separation in blends of PA 11 and PA 6,10 is driven by crystallization. In a separate study, polyamide-11 (PA 11/ Nylon-11) is melt blended with polylactide (PLA) in the presence of 0.00 to 0.10 wt% titanium isopropoxide catalyst to investigate the possible occurrence of compatibilizing transreactions between the polymers. Mechanical and thermal properties of these blends are also characterized by differential scanning calorimetry (DSC), thermo gravimetric analysis (TGA), dynamic mechanical thermal analysis (DMTA), tensile testing, and impact testing. The storage moduli of the blends increase monotonically from 0.58 GPa for neat melt blended PA 11 to 1.52 GPa for neat PLA. Similarly, tensile moduli for neat PLA and PA 11 are 3.8 GPa and 1.4 GPa respectively and blend values are 2.2, 2.9 and 3.0 GPa for the 25/75, 50/50, and 75/25 compositions, respectively. As PA 11 composition increases, the failure mode changes from the brittle fracture of neat PLA to ductile fracture of neat PA 11. Impact strength of neat PLA is 41.8 J/m, and that of neat PA 11 is 248.1 J/m; values of the blends are 49.9, 30.6 and 22.3 J/m for the 25/75, 50/50, and 75/25 compositions, respectively. DSC thermograms show two distinct melting peaks in the blends, one corresponding to the melting peak PLA at about 170-173˚C and other corresponding to that of PA 11 at about 190-194˚C. DSC also shows two separate glass transition temperatures thus indicating only partial miscibility. TGA is used to determine degradation temperatures of 348.1˚C for PLA and 429.8 ˚C for P A11. The initial onset degradation temperatures for the 25/75, 50/50, and 75/25 blends are 341.6, 344.4, and 346.2 °C respectively. The heat distortion temperatures (HDT) change monotonically from the value for PA 11 to the value for PLA; values for PLA and PA 11 are 75.5 ˚C and 58.5 ˚C and values for the blends are 62.3, 69.0, 72.1 for the 25/75, 50/50, and 75/25 blends, respectively. Base etching to remove PLA domains followed by morphological examination using field emission scanning electron microscopy (FE-SEM) confirms the two phase nature of the blends. Interchange reactions during reactive mixing were investigated by 13 C-NMR spectroscopy but the analysis shows little evidence of interchange reactions. This is true irrespective of catalyst level and mixing time over the temperature range from 185 ˚C to 225 ˚C. At the upper end of the temperature range investigated, significant degradation is observed. Therefore, the results indicate that

degradation reactions will dominate over compatibilizing interchange reactions at the higher temperatures required for sufficient reactions to compatibilize the blend. The inability of the PLA/PA 11 blend to compatibilize in-situ provided the need to produce a copolymer to strengthen the interface and improve mechanical properties. For this study, PLA was hydrolyzed to lower the molecular weight in preparation for the copolymer. Kinetics of the auto-catalyzed hydrolysis were analyzed and applied to obtain a lower molecular weight PLA chain. The successful polymerization of the block poly(11-aminoundecanoic acid-block-L-lactic acid) was shown through the analysis of DSC and FT-IR.

TABLE OF CONTENTS

ABSTRACT ...... iii LIST OF FIGURES ...... viii LIST OF TABLES ...... xi ACKNOWLEDGEMENTS ...... xiii CHAPTER 1 RENEWABLE RESOURCE BASED BIOPLASTIC BLENDS: POLYAMIDE-11 / POLYAMIDE-6,10 BINARY SYSTEM ...... 1 1.1 Materials ...... 3 1.2 Methods...... 4 1.3 Results and Discussion ...... 7 1.3.1 Differential Scanning Calorimetry ...... 7 1.3.2 Thermogravimetric Analysis ...... 17 1.3.3 Dynamic Mechanical Thermal Analysis ...... 18 1.3.4 Impact Strength ...... 21 1.3.5 Tensile Test ...... 22 1.3.6 Microscopy ...... 24 1.3.7 Small Angle X-ray Scattering ...... 28 1.3.8 Wide Angle X-ray Scattering...... 30 1.4 Conclusion ...... 33 CHAPTER 2 RENEWABLE RESOURCE BASED BIOPLASTIC BLENDS: POLYAMIDE-11 / POLYLACTIDE BINARY SYSTEM ...... 35 2.1 Materials ...... 37 2.2 Methods...... 37 2.3 Results and Discussion ...... 40 2.3.1 Extractions ...... 40 2.3.2 Nuclear Magnetic Resonance on Carbon Isotope 13 ...... 42 2.3.3 Differential Scanning Calorimetry ...... 43 2.3.4 Thermogravimetric Analysis ...... 46 2.3.5 Scanning Electron Microscope ...... 47 2.3.6 Dynamic Mechanical Thermal Analysis ...... 49 2.3.7 Tensile Test ...... 51 2.3.8 Impact Test...... 53

2.4 Conclusion ...... 54 CHAPTER 3 SYNTHESIS OF POLY(11-AMINOUNDECANOIC ACID-BLOCK- L-LACTIC ACID) FOR USE AS COMPATIBILIZING AGENT IN POLYAMIDE-11 / POLY(LACTIC ACID) BLENDS ...... 57 3.1 Materials ...... 58 3.2 Methods...... 58 3.3 Results ...... 61 3.3.1 Hydrolysis ...... 61 3.3.2 Synthesis ...... 65 3.4 Conclusion ...... 68 FUTURE WORK ...... 69 REFERENCES ...... 71 CD ...... Pocket

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LIST OF FIGURES

Figure 1.1: DSC scans showing melting peaks of all compositions; significant melting point depression of PA 6,10 is observed, suggesting miscibility across all compositions of the blend...... 8 Figure 1.2: The Hoffman-Weeks plot typically used to find equilibrium melting temperatures of semi-crystalline blends. PA 6,10 shows morphological changes in addition to the thermodynamic changes making the technique ineffective for this particular blend...... 10 Figure 1.3: Percent crystallinity of as prepared polymer blends upon second heating in the DSC. Open squares represent the percent by volume of the PA 11 that is in the crystalline state, circles represent PA 6,10, and Xs represents the overall percent crystallinity of the blend...... 15 Figure 1.4: Typical melt peak of the 50/50 PA 11 / PA 6,10 blend annealed at: the high crystallization temperature only (H, T c=181), the high then the low crystallization temperatures in sequence (H-L), and the low crystallization temperature only (L, T c=167)...... 16 Figure 1.5: The typical melt peaks of all compositions (PA 11 / PA 6,10) using the maximum crystallinity annealing sequence found above (T c1 = 181, Tc2 = 167 for 30 minutes each)...... 16 Figure 1.6: Representative TG data for PA 11 / PA 6,10 blends under an air atmosphere...... 18 Figure 1.7: Storage (G') and Loss (G") Modulus of all compositions (PA 11 / PA 6,10) with respect to temperature...... 19 Figure 1.8: tan ( δ) for all compositions (PA 11 / PA 6,10) as a function of temperature...... 20 Figure 1.9: Measured impact strength of blends without annealing. This data shows a higher than a linear mixing rule would indicate...... 21 Figure 1.10: Young's modulus and strain at break measured for as molded and thermally treated tensile bars...... 23 Figure 1.11: Tensile strength at yield for all five compositions of the polyamide blends...... 23 Figure 1.12: Cross polarized microscopy image of PA 6,10 crystallized at 185 ˚C then 162 ˚C depicting the size of spherulites in the sample...... 25 Figure 1.13: Cross polarized microscopy image of 25/75 PA11/PA6,10 blend crystallized at 185 ˚C then 162 ˚C depicting the size of spherulites in the sample...... 26 Figure 1.14: Cross polarized microscopy image of 50/50 PA11/PA6,10 blend crystallized at 185 ˚C then 162 ˚C depicting the size of spherulites in the sample...... 26

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Figure 1.15: Cross polarized microscopy image of 75/25 PA11/PA6,10 blend crystallized at 185 ˚C then 162 ˚C depicting the size of spherulites in the sample...... 27 Figure 1.16: Cross polarized microscopy image of PA 11 crystallized at 185 ˚C then 162 ˚C depicting the size of spherulites in the sample...... 27 Figure 1.17: 1-D SAXS scattering of blends annealed at the crystallization temperature of 162 ˚C showing no change between compositions (PA 11 / PA 6,10)...... 28 Figure 1.18: 1-D SAXS scattering of blends annealed at the crystallization temperature of 185 ˚C. The decreasing scatter intensity (q) correlates with increasing long spacing and suggests PA 11 chains are “trapped” as the PA 6,10 crystal structures grow...... 29 Figure 1.19: Representative lattice structure for α-triclinic geometry where a, b, and c are lattice dimensions and α, β, and γ lattice angles...... 30 Figure 1.20: 1-D WAXS scattering of blends annealed at the crystallization temperature of 162 ˚C...... 31 Figure 1.21: 1-D WAXS scattering of blends annealed at the crystallization temperature of 185 ˚C...... 32 Figure 2.1: Proposed interchange reactions in polyester-polyamide blends...... 37 Figure 2.2: Chloroform insoluble weight percent as a function of catalyst loading level at 185 ˚C...... 41 Figure 2.3: C13 NMR spectra for a) pure PLA, b) chloroform soluble fraction, c) chloroform insoluble fraction, and d) pure PA 11...... 42 Figure 2.4: DSC heat flow over the glass transition regions of PLA/PA 11 showing two glass transitions for all compositions suggesting immiscibility...... 44 Figure 2.5: DSC heat flow over the range of melting temperatures for PLA/PA 11 blends prepared at different mixing temperatures...... 45 Figure 2.6: TG data for PLA/PA 11 blends prepared at 205 ˚C for 20 minutes of melt-blending...... 46 Figure 2.7: FESEM images for 25/75 wt% PLA/PA 11 blends prepared at 205 ˚C for 20 minutes of melt-blending with the PLA phase removed through etching...... 47 Figure 2.8: FESEM images for 50/50 wt% PLA/PA 11 blends prepared at 205 ˚C for 20 minutes of melt-blending with the PLA phase removed through etching...... 48 Figure 2.9: FESEM images for 75/25 wt% PLA/PA 11 blend prepared at 205 ˚C for 20 minutes of melt-blending with the PLA phase removed through etching...... 48

Figure 2.10: Storage modulus (G') versus temperature for various PLA/PA11 blends prepared at 205 ˚C with 20 minutes of melt blending time...... 49 Figure 2.11: Tan δ vs. temperature plot for all PLA/PA 11 blends prepared at 205 ˚C for 20 minutes of melt blending...... 50 Figure 2.12: HDT values for different PLA/PA 11 blends prepared at 205 ˚C for 20 minutes of melt blending...... 51 Figure 2.13: Stress at break and tensile modulus of PLA/PA 11 blends as a function of blend composition...... 52 Figure 2.14: Toughness and strain at break of PLA/PA 11 blends as a function of blend composition...... 52 Figure 2.15: Notched Izod impact strength of PLA/PA 11 blends as a function of blend composition ...... 54 Figure 3.1: Refractive index calibration for lactic acid in water...... 60 Figure 3.2: Proposed hydrolysis reaction in PLA ...... 62 Figure 3.3: Molecular mass (weight averaged) of PLA as a function of hydrolysis time...... 63 Figure 3.4: Weight of Lactic Acid in solution as a function of hydrolysis time...... 64 Figure 3.5: Typical DSC endotherms for commercial PLA and hydrolyzed PLA...... 64 Figure 3.6: DSC experiment showing the first (11-aminoundecanoic acid) and the second run (low molecular weight PA11) of sample consisting of pure 11-aminoundecanoic acid...... 65 Figure 3.7: DSC experiment showing the first run for 11-aminoundecanoic acid and samples taken every 10 minutes from the Haake reaction...... 66 Figure 3.8: Proposed reaction of PLA and 11-aminoundecanoic acid to form copolymer (first addition of monomer 11-aminoundecanoic acid shown) ...... 67 Figure 3.9: FT-IR spectra of PA 11 in red, soluble extract of copolymerization in green , adn hydrolyzed PLA in blue for a) full spectra, b) N-H fingerprint, c) C-H fingerprint and d) amide fingerprint...... 67

LIST OF TABLES

Table 1.1: Observed melting temperatures from the average of four DSC runs for all compositions...... 8 Table 1.2: Equilibrium melting temperature and stability parameter calculated from the Hoffman-Weeks plot...... 11 Table 1.3: Calculated polymer-polymer interaction parameter of PA 11 on PA 6,10 show miscibility for all compositions...... 12 Table 1.4: Calculated solubility parameters through two different methods of contribution calculations...... 13

Table 1.5: Glass transition temperature (T g), crystallinity, and maximum crystallization temperatures averaged over four different DSC runs with one standard deviation reported as error...... 14 Table 1.6: Calculated overall crystallinity for 50/50 PA 11 / PA 6,10 blend for different annealing sequences...... 15 Table 1.7: Overall crystallinity for the annealing sequence that maximizes crystallinity for all compositions...... 17 Table 1.8: Comparison of room temperature storage (G') and loss moduli (G") for blends from values obtained with a temperature sweep and from room temperature only tests...... 19 Table 1.9: Calculated Poisson's ratio from thermally treated samples of Young's modulus and complex modulus showing a discontinuity between measured modulus...... 24 Table 1.10: Calculated index of refraction from two different contribution methods at a wavelength of 589 nm...... 25 Table 1.11: Table showing the long spacing calculated based on the maximum for all blend compositions at the two different annealing temperatures ...... 29 Table 1.12: Lattice dimensions and angles from literature, calculated d-spasings, and measured d-spacings for PA 11 and PA 6,10...... 30 Table 1.13: Associated peaks for PA 11 ( ô¿) and PA 6,10 ( õ¿) taken from the peaks of the 1-D WAXS data for the samples annealed at crystallization temperature of 162 ˚C...... 32 Table 1.14: Associated peaks for PA 11 ( ô¿) and PA 6,10 ( õ¿) taken from the peaks of the 1-D WAXS data for the samples annealed at crystallization temperature of 185 ˚C...... 33 Table 2.1: Weight percent of insoluble material for different catalyst loading levels, reaction temperatures and reacting times...... 40 Table 2.2: DSC parameters obtained from the first and second heating scans for various PLA/PA 11 blend compositions...... 45

Table 3.1: Results of Solubility Tests of reactants, polymers, and product in m-cresol and chloroform...... 67

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ACKNOWLEDGEMENTS

Chapter 1 The research was supported by the National Science Foundation through grant CMMI-0700869. I would like to thank Interface Global for providing the polyamide 11 from Arkema and Dr. Scherzer from BASF for providing the PA 6,10 used in the study.

Chapter 2 This research was supported by the National Science Foundation through grant CMMI-0700869 and by the National Research Initiative Competitive Grant 2006-35504- 16618 from the USDA Cooperative State Research, Education, and Extension Service. I would like to thank Rahki Patel for her work on this project, much of the analysis and lab work. I would also like to thank Dr. C.H. Zah from the Interface Corporation for graciously providing the Polyamide-11 used in the Polyamide-11/Polylactic Acid blends.

Chapter 3 The research was supported by the National Science Foundation through grant CMMI-0700869 and the Colorado Center of Biorefining and Biofuels (C2B2) for funding man hours through it summer internship program. I would like to thank Daniel Harrison for his work on this project, in the lab and with the analysis.

Finally, I would like to thank Dr. John Dorgan for being my remarkable advisor and guiding the studies presented in the following thesis. Lastly, I would like to thank Dr. Andrew Herring and Dr. Dan Knauss for being thesis committee members for the presented Masters of Science Thesis.

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CHAPTER 1 1 RENEWABLE RESOURCE BASED BIOPLASTIC BLENDS: POLYAMIDE-11 / POLYAMIDE-6,10 BINARY SYSTEM

Market demand for renewably based polymers is expected to increase by 18% annually through the year 2018 with non-biodegradable bioplastics growing from around 8% to 47% of the bioplastic market over the same period. This projection comes from the range of applications for renewably based polymers and consumer desire for increased sustainability and energy independence [1-4]. An established and cost- effective way to develop new polymeric materials with diversified and desirable properties is to blend polymers together [5-10]. In this study the fully renewable polyamide-11 (PA 11) is blended with a polyamide 6,10 (PA 6,10) in which the “10” structural unit is derived from renewable resources. The resulting blends have renewable carbon contents ranging from 100% to 63%[11]. The full suite of thermal and mechanical properties is reported here for the first time. Polymer blending science and blend applications have grown recently [6, 7]. There are numerous literature sources that describe theoretical and experimental techniques for determining if a polymer blend is miscible or immiscible [12-40]. Polymers typically form immiscible blends. This is because the large molar mass of polymer chains limits the entropic contribution to free energy of mixing. Accordingly very small unfavorable enthalpy of mixing values lead to phase separation. Therefore, a large portion of polymer blending research is conducted on immiscible polymer systems; the effects of compatibilization agents and blending techniques on the properties and morphology of immiscible blends are widely reported [5, 7, 8, 12, 26, 34, 41-44]. Miscible polymer blends are relatively rare but are of high interest because they do not need to be compatabilized. There is limited literature concerned with renewably based polymer blends for durable polymeric material applications [5-7, 45, 46]. However, there are a number of studies on biodegradable blends both for disposable packaging applications and for biomedical use; some of these blends are composed of both renewable and non-

renewable components [5, 47-55]. This study investigates the aliphatic polyamides, polyamide-11 (PA 11) and polyamide-6,10 (PA 6,10); both biorenewable structural units made from castor oil derived from the beans of the castor plant. [56-58]. Polyamides (PAs / Nylons) are very commonly used in fabrics, carpet manufacturing, and in other durable goods applications including: metal replacement in train cars, radiator tanks, and the bracket for the battery pack in the Chevrolet Volt [59- 62]. The high-performance polymers [57, 62] blended in this article, polyamide-11 (Nylon 11) and polyamide-6,10 (Nylon 6,10), are already used separately in many high performance applications. Accordingly, blending is considered a useful enterprise that enables tuning of the thermophysical properties for specific applications. Polyamide-11 (PA 11) is entirely derived from castor oil; it is a 100% biorenewable material. PA 11 is a high-performance, semi-crystalline, thermoplastic. When compared to petroleum based nylons and other conventional plastics. PA 11 has

low net CO 2 emissions and global warming potential. Some of the outstanding properties of PA 11 include high impact and abrasion resistance, low specific gravity (density of 1.03 g/cm 3), excellent chemical resistance, low water absorption, high thermal stability (melting point of 180-190 ˚C), and ease of processing over a wide range of processing temperatures. That is, PA 11 has excellent dimensional stability, and maintains physical properties over a wide range of temperatures and environments. PA 11 thus finds numerous applications including: crude oil and gas pipelines, hydraulic and pneumatic hoses, powder coatings, skis, snowboards and optical fiber sheathing [63]. The polyamide-6,10 (PA 6,10) used in this study is partially derived from castor oil. Specifically, the ten carbon di-acid comes from castor oil. This means the polymer is 63% biorenewable. PA 6,10 is a high-performance, semi-crystalline polymer. Outstanding properties of PA 6,10 include reduced water uptake, hydrolysis resistance, stress cracking resistance, resistance to fuels, oils, greases, most solvents, aqueous solutions and alkalis. PA 6,10 has excellent flexural stiffness, dimensional stability, heat deflection temperature and high thermal stability (melting point of about 220 ˚C). PA 6,10 has numerous applications including automotive metal replacement and electrical insulation.

PA blend literature consists mainly of polyamide-6 or polyamide-6,6 blended with styrene, polypropylene, ABS, other polymers, or rubber toughening additives. There are several papers considering aliphatic PAs blended with aromatic PAs [15-20, 22, 39], but limited number of publications on aliphatic only PA blends [21, 64-66]. It has been shown that miscible aliphatic nylon blends, are rare, but do exist. Some of these miscible blends are polyamide 4,8 / polyamide 6,6 [67] and polyamide 6,6 / polyamide 6 [21, 65]. The blends in this study have not been reported in literature. There are some sources on the interchange reactions that occur between the polyamide end groups to form block co-polymer in-situ . The blending time required to form significant interchange reactions is upwards of 3 hours at temperatures of 300 ˚C; therefore, this effect will not have any large role in the results presented in this study [68-70]. This paper shows the blends of PA 11 and PA 6,10 are miscible and therefore represents new and potentially significant findings. A recent review on all known miscible crystalline/crystalline polymer blends was published recently by B. J. Jungnickel in Journal of Polymer Science: Part B: Polymer Physics [27]. This review paper discusses the few crystalline / crystalline polymer blends and notes the strange kinetic and structural phenomena observed. It is reported that the crystals do not usually grow simultaneously, making the PA 11 / PA 6,10 blend even more scientifically interesting as there is crystal/crystal induced phase separation Crystallinity induced phase separation has been treated theoretically based on derivatives of a theory originally proposed by Dorgan [71, 72].

1.1 Materials

Rislan Rigid Grade PA 11 was originally purchased from Arkema by Interface Global (Atlanta, GA) and donated to the Colorado School of Mines. Ultramid Balance PA 6,10 was supplied by Dr. Scherzer from BASF SE (Luwigshaven, DE). The thermal stabilizers, Irganox 1076 and Irgafos 168, were donated by CIBA Specialty Chemicals Inc. and used as received. Reagent grade toluene and m-cresol were purchased from Sigma-Aldrich.

1.2 Methods

Samples of PA 11 and PA 6,10 were prepared by melt mixing in a Haake Rheomix 600 at 220 ˚C for 10 minutes at 50 RPM using roller blades at compositions of 0/100, 25/75, 50/50, 75/25, and 100/0 wt% of PA 11 / PA 6,10. To minimize hydrolysis during mixing, the polymers were dried under vacuum (23 inch Hg) at 80 ˚C for 20 hours. Pellets of PA 6,10 were introduced into the mixer first, followed by PA 11. After complete melting was achieved, the stabilizing agents were added in toluene solution to achieve an overall composition of 0.25 wt% each upon evaporation of the toluene. Blends were mixed for 10 minutes. The blended material was removed from the mixer, allowed to cool, and then ground in a Foremost granulator to a pellet size of approximately 2-3 mm in diameter. After grinding, samples were taken for differential scanning calorimetry (DSC), thermal gravimetric analysis (TGA), microscopy, and x-ray scattering. The remaining material was dried under vacuum and injection molded into bars for dynamical mechanical thermal analysis (DMTA), impact testing, and tensile testing. A Morgan Press Injection Molding instrument with an aluminum mold was use for injection molding test bars at the settings of 10 tons for the clamp force, 8.2 bars for the piston pressure, and nozzle temperatures ranging from 244 ˚C to 260 ˚C. Thermal properties of the blends and homopolymers were investigated using DSC (Perkin-Elmer DSC-7 calibrated against an Indium standard). Aluminum pans were used to seal samples with masses ranging from 10-20 mg for DSC analysis. The DSC protocol to find an observed melting point (T m), glass transition temperature (T g), and the maximum crystallinity temperature (T c,max ) was: 1) hold at 5 ˚C for 5 minutes, heat from 5 ˚C to 260 ˚C at 10 ˚C/min., 3) hold at 260 ˚C for 5 min., 4) cool from 260 ˚C to 5 ˚C at 10 ˚C/min., 5) hold at 5 ˚C for 5 min., and 6) heat from 5 ˚C to 260 ˚C at 10 ˚C/min. The

Tc,max was recorded on the cooling run (step 4); Tg, T m, and crystallinity were recorded from the second heating run (step 6). The midpoint of the heat capacity change was used as the T g. The DSC protocol for creating the Hoffman-Weeks plot was: 1) hold at 5 ˚C for 5 minutes, 2) heat from 5 ˚C to 260 ˚C at 10 ˚C/min., 3) hold at 260 ˚C for 5 min., 4) cool from 260 ˚C to a temperature in the range 170-195 ˚C at 200 ˚C/min., 5) hold at that temperature for 30 min., 6) cool from the crystallization temperature to 5 ˚C at 200 ˚C/min., and 7) heat from 5 ˚C to 260 ˚C at 10 ˚C/min. The melting temperature was

taken from the second heat (step 7). The DSC protocol for finding maximum crystallinity was: 1) hold at 5 ˚C for 5 minutes, 2) heat from 5 ˚C to 260 ˚C at 10 ˚C/min., 3) hold at 260 ˚C for 5 min., 4) cool from 260 ˚C to a maximum crystallinity temperature at 200 ˚C/min. (i.e. 195 ˚C, 181 ˚C, 163 ˚C), 5) perform hold sequence of different temperatures and combinations of the temperatures, 6) cool from the crystallization temperature to 5 ˚C at 200 ˚C/min., and 7) heat from 5 ˚C to 260 ˚C at 10 ˚C /min. The hold sequences used were for PA 11 at 163 ˚C, PA 6,10 at 195 ˚C, the 50/50 blend at 181 ˚C, at 163 ˚C, or 181 ˚C then 163 ˚C. All data was collected on the second heat (step 7). The blends thermal stability was characterized by TGA using a Seiko TG/DTA 220 instrument. Samples of 15-25 mg of blend were placed in an alumina pan. The heating program used was: 1) heat from 40 ˚C to 800 ˚C at 10 ˚C /min, 2) cool from 800 ˚C to 40 ˚C at 10 ˚C /min and were carried out under air flow. The onset degradation temperature is calculated as the temperature at the intersection of a straight line from the weight vs. temperature data before decomposition started and a second straight line fit to the decomposition step. DMTA was carried out in an ARES-LS rheometer with torsional rectangular fixtures and used to find the shear modulus as a function of temperature. Testing was carried out at 0.05 % strain and a frequency of 1 Hz at temperatures from -10 ˚C to 180 ˚C at 5 ˚C /min. The instrument was calibrated for normal force and torque before tests were performed. Glass transition temperature associated α-transition and heat distortion temperature (HDT) [73] were also found from DMTA data. Impact properties of the polymer blends were measured according to ASTM standard D256 method for notched Izod impact testing. ASTM D4066 standards were followed in the preparation of the nylon bars; therefore they were not thermally treated before the test was performed. Impact testing bars (dimensions 12.7x63.5x3.2 mm) from the Morgan Press were set on a plate for 24 hours after injection molding to control free volume relaxation and provide for amorphous phase densification. The bars were then notched using a manual RJW LTD Charpy notch instrument with a type H “V” broach. A minimum of 5 samples were conducted for every composition on a TMI electronic Izod impact tester with a 5 ft-lb swing arm to measure the impact strength after an

average of three thicknesses were measured near the notch. Error cited is one standard deviation. Tensile Tests of injection molded bars were performed according to ASTM D638. Samples were allowed to densify after injection molding for 24 hours before testing or thermal treatment. The thermal treatment sought to control the degree of crystallinity and was performed in a convection oven at the maximum crystallinity temperatures of the two homopolymers (first at 185 ˚C, second at 162 ˚C) for three hours each. The thickness and width are measured using an electronic caliper in three different places within the test region and averaged. A minimum of five samples for every composition was tested at room temperature on a NTS InstruMET Corporation load frame using a 2000 lb load cell. The crosshead was used to measure strain, but corrected against extensometer data collected for small deformations. The correction factor was to take into account the observed difference in the crosshead strain and actual strain measured by extensometer. Error sited is one standard deviation. Microscopy was performed to determine the miscibility of the polymer blend and change in size of spherulites. Films were initially made through melt pressing between Teflon sheets using a Carver Press equipped with heated plates. Samples were quenched to 0 C by placing them between aluminum block through which ice water was flowing. Films were liberated from the Teflon and placed onto glass slides. Most of the films were then thermally treated using several annealing steps in different order to observe the differences in spherulite size. Pictures were taken using an Imaging Planet USB 2.0 3.3 MPX camera mounted on a Nikon microscope at 22x magnification. Small angle x-ray scattering (SAXS) and wide angle x-ray scattering (WAXS) were performed to study changes in crystalline structure with blend composition and thermal annealing. Films were prepared through melt pressing on a Carver Press between Teflon sheets. Thermal treatment in a convection oven consisted of treating at 185 ˚C or 162 ˚C for three hours on a preheated metal plate. The SAXS and WAXS tests were performed under vacuum on a Rigaku S-Max 3000 system using a wavelength of 1.5405 Å calibrated with silver behenate. Data was collected on a multi wire photon detector (MPANT detector) for SAXS and reusable image plates for WAXS and analyzed

using SAXS Gui V2.05.02 software. All SAXS samples were run for 3 hours and all WAXS samples were run for 90 minutes.

1.3 Results and Discussion

In this study mechanical and thermal properties of these novel miscible but semi- crystalline blends are characterized through the use of: differential scanning calorimeter (DSC), thermo-gravimetric analysis (TGA), dynamic mechanical thermal analysis (DMTA), tensile testing, and impact testing. The blends show intermediate properties between the two homopolymers, expanding the application portfolio of renewably derived polymers. Morphological understanding was developed using polarized light microscopy, small angle X-ray scattering (SAXS), and wide angle X-ray scattering (WAXS). The blends show definitive evidence of amorphous phase miscibility even though a phase separation must occur below melting temperatures to enable formation of polymer crystals. This unique crystalline induced phase separation provides morphologically unique polymer systems comprised of a miscible amorphous matrix with two distinct types of crystals being present.

1.3.1 Differential Scanning Calorimetry

Thermal properties of PA 11 and PA 6,10 blends were comprehensively investigated to find miscibility. As shown in Figure 1.1, melting point depression of 15 ˚C is observed. As discussed below, such depression indicates blend miscibility. A theory for thermodynamic mixing of two polymers was developed by Scott using the Flory-Huggins approximation [38]. The theory expresses the chemical potential of a polymer in a binary blend relative to the chemical potential of the polymer in the melt state. The chemical potential of polymer species 2 in a binary blend is given by Equation 1.1.

ln 1 1 (1.1) − = + − 1 − + 1 −

0% PA 11 55 25% PA 11 50% PA 11 75% PA 11 50 100% PA 11

35 Endo Up (mW) Up Endo

20 160 170 180 190 200 210 220 230 o Temperture ( C) Figure 1.1: DSC scans showing melting peaks of all compositions; significant melting point depression of PA 6,10 is observed, suggesting miscibility across all compositions of the blend.

Table 1.1: Observed melting temperatures from the average of four DSC runs for all compositions. PA 11 / PA 11 Melting PA 6,10 Melting PA 6,10 Temperature Temperature wt% ˚C ˚C 100/0 189 ± 0.6

75/25 186 ± 0.4 210 ± 2.1 50/50 189 ± 1.1 220 ± 1.4 25/75 189 ± 0.6 223 ± 0.8 0/100 226 ± 0.6

In Equation 1.1, subscript 1 identifies an amorphous polymer and subscript 2 identifies a semi-crystalline polymer, is the chemical potential per repeat unit of polymer 2 in a reference state, is the chemical potential per repeat unit of polymer 2 in the blend, is the volume fraction, Viu is the molar volume of the repeating unit, m i is degree of polymerization, R is the gas constant, T is the absolute temperature, and is the polymer-polymer interaction parameter of polymer 1 on polymer 2. Nishi and Wang [30] derived a means to calculate the polymer-polymer interaction parameter from melting point depression. Equation 1.2 describes the chemical potential difference between the

polymer in the crystalline state relative to the same melt state of Equation 1.1.

(1.2) − = −Δ − Δ In Eqtuation 1.2 is the chemical potential per repeat unit of polymer 2 in the crystalline state and ∆ and ∆ are the enthalpy and entropy of fusion per repeating unit. If ∆ ∆ is assumed to be independent of temperature, Equation 1.2 can be / rearranged to:

. (1.3) − = −Δ 1 − In Equation 1.3 is the equilibrium melting temperature of the pure homopolymer. Subtracting Equation 1.3 from Equation 1.1 and setting the chemical potentials of the semi-crystalline polymer in the crystalline and blend states to be equal at the melting temperature yields Equation 1.4.

1 1 − ln 1 1 (1.4) − = + − 1 − + 1 − Δ In Equation 1.4 is the equilibrium melting temperature of the blend and is the equilibrium melting temperature of the pure homopolymer. Equation 1.4 can be reduced further because the system being studied consists of large polymer chains, therefore, and are very large relative to other quantities. Through setting the reciprocal of these large quantities to zero results in Equation 1.5a, which can then be rearranged to explicitly calculate the polymer-polymer interaction parameter as seen in Equation 1.5b.

1 1 − (1.5a) − = 1 − Δ

Δ 1 1 1 (1.5b) = − 1 − Several issues arise in obtaining the correct experimental data to apply to Equation 1.5b in obtaining a meaningful polymer-polymer interaction parameter [36, 37]. The main concern is with finding the true equilibrium melting temperature. Hoffman and Weeks suggest that an equilibrium melting temperature can be found through isothermally crystallizing the polymer at several different temperatures (T c) and

measuring a corresponding observed melting temperature (T m). This allows an extrapolation of the data to the theoretical limit of crystallization at the true melting temperature (i.e. T m = T c) [74, 75]. This analysis assumes surface effects of the

240

235

230 C) o 225

220

215 0/100 25/75 210 50/50

Melting Temperature ( Temperature Melting 75/25

205 T m=T c

200 170 180 190 200 210 220 230 240 Crystallization Temperature ( oC)

Figure 1.2: The Hoffman-Weeks plot typically used to find equilibrium melting temperatures of semi-crystalline blends. PA 6,10 shows morphological changes in addition to the thermodynamic changes making the technique ineffective for this particular blend.

formation of crystals are negligible, crystals at the melting temperature have equilibrium crystal perfection, and that the crystal thickening process is independent of temperature. The Hoffman-Weeks plot can be seen in Figure 1.2. In the derivation by Hoffman and Weeks Equation 1.6 is established.

(1.6) In Equation 1.6 is the crystallization − temperature = −of isothermal annealing and is a stability parameter that depends heavily on the crystal thickness, where is the most = 0 stable crystal (i.e. ) and is unstable (i.e ). The calculated = = 1 = equilibrium melting temperatures and the stability parameter, , for the blends and homopolymer can be found in Table 1.2. Table 1.2 shows that, the stability parameter is increasing with increasing PA 6,10 content rendering the Hoffman-Weeks approach unsuitable for finding the equilibrium melting temperature in the present system. Rim and Runt state that variation in

Table 1.2: Equilibrium melting temperature and stability parameter calculated from the Hoffman-Weeks plot. Equilibrium Melting Stability Parameter PA11/PA6,10 Temperature ( ) wt% ˚C - 0/100 228 ± 6.6 0.00 ± 0.016 25/75 233 ± 3.4 0.14 ± 0.007 50/50 238 ± 43.1 0.50 ± 0.047 75/25 - 2.4 morphology will be superimposed on the thermodynamic melting temperature [36] as is seen by the changing stability parameter. It has also been noted in literature that nylons do not show a sign of α-relaxation [75-77] and may exhibit lamellar thinning under isothermal crystallization [75, 78, 79]. This results in an issue of finding meaningful equilibrium melting temperatures to calculate the polymer-polymer interaction parameter. Runt et. al. has shown the difficulties with this method. This author makes the assertion that there is a dependence of measured melting temperature on the rate of heating in the DSC [36, 37] and can be seen in Zhang et. al. [40]. It should be noted that the crystal melt peak for PA 6,10 completely disappears at lower crystallinity temperatures (T c = 170-180 ˚C). This is evidence of miscibility and illustrates the favorable enthalpy of mixing, enough to cause the formation of crystals to be unfavorable in the presence of PA 11 polymer. This is an indication that a eutectic system is present. After this observation a simple analysis using Gibbs phase rule (Equation 1.7) shows that the equilibrium condition will be between the crystals of the two polymers. (1.7) In Equation 1.7 the F is degrees of freedom,F = C − CP is+ the2 number of components, and P is the number of phases present in the system. The number of components of an monodisperse polymer binary blend is two and the number of phases present in the system is three (amorphous phase, crystal phase of PA 11, and crystal phase of PA 6,10), Gibbs phase rule would calculate the degrees of freedom to be one. This would indicate that by setting either temperature or pressure the whole system would be defined. This is not possible as the temperature and pressure are independent of one another. This shows the system is not in equilibrium and kinetically limited from reaching the equilibrium state

that can contain only two phases (crystal 1-crystal 2, crystal 1-amorphous, or amorphous-

crystal 2 where the middle is realized for the 75/25 blend at some temperatures tested). Polyamides crystallize quickly and during isothermal crystallization of blends in the DSC little to no crystallization peak during the hold step was observed. Therefore the calculations for the polymer-polymer interaction parameter were carried out using the observed melting temperatures as equilibrium temperatures. The observed melting temperatures from Table 1.1 were equal to or higher than the isothermal crystallization sample for all blends, allowing this assumption to be used with some confidence. The results of the calculations are shown in Table 1.3. The negative polymer-polymer interaction parameters calculated for the blends are all less than zero indicating miscibility and favorable enthalpy from mixing.

Table 1.3: Calculated polymer-polymer interaction parameter of PA 11 on PA 6,10 show miscibility for all compositions. PA 11/PA 6,10 wt% C - , - 0/100 226 1.00 - 25/75 223 0.74 -0.817 50/50 220 0.49 -0.406 75/25 210 0.24 -0.544

Group contribution calculations can be used to estimate the miscibility for the PA 11 / PA 6,10 blend to find plausibility of the interaction parameter calculations above. Using the geometric average of the individual contributions the difference in the total interaction parameters are ∆δ Hoftyzer = 0.223 J and ∆δ Hoy = 0.166 J It is cm cm commonly accepted that if the difference of the total interaction parameter is less than 5, the mixture has a good chance of being miscible [80]. Calculations of the Hildebrand solubility parameters agree with the findings of the polymer-polymer interaction parameter above; the system is miscible. Calculations from the Hildebrand solubility parameters can be related to the Flory-Huggins interaction parameter through Equation 1.8. The calculations give a Florry-Huggins interaction parameter of χ = 0.0054 for Hoftyzer’s method and χ = 0.0144 for Hoy’s method. For a 50/50 blend by volume, the Flory Huggins theory

Table 1.4: Calculated solubility parameters through two different methods of contribution calculations. δh δp (Hydrogen δd δt (Polar Bonding (Dispersive (Total Contribution Contribution Contribution Interaction Interaction Interaction Interaction Parameter) Parameter) Parameter) Parameter)

Hoftyzer’s PA 11 17.99 4.23 5.20 16.69 Method PA 6,10 18.27 4.04 6.04 16.76 Hoy’s PA 11 21.67 9.58 7.92 17.75 Method PA 6,10 22.13 9.46 9.16 17.79 predicts phase separation for χ ( - degree of polymerization). Therefore, when > 2 the value is χ = 0.0054 the number of segments in a symmetric polymer blend could be as large as m=370; this can be compared to the values obtained by dividing the polymer molecular weight (g/mol) by the repeat unit molecular weights for the two polymers. In the present case this yields m = 97 for PA 11 and m = 61 for PA 6,10. This estimation also indicates the blends are miscible (it must be noted that the approximation of Equation 1.8 precludes any negative values for the interaction parameter, a serious limitation and at odds with the experimental facts for several polymer-polymer blends.

(1.8) = − In Equation 1.8 is the Flory-Huggins interaction parameter, is the molar volume of the repeat unit, and and are the Hildebrand solubility parameters of polymer 1 and polymer 2.

The DSC was used to find the glass transition temperature (T g), un-annealed

crystallinity, and the maximum crystallization temperature (T c,max ) of the blends and homopolymers. All values were determined using the data from the second heating scan to avoid artifacts associated with the initial heating. The glass transition temperature varies slightly in the blends with only one T g being observed. This was true for all samples indicating miscibility. However, the T g value for the homopolymers are quite close; PA 11 is 34.9 ˚C and PA 6,10 is 38.2 ˚C. The thermal signals are also weak making it difficult to evaluate even in the pure state. Accordingly, no definitive

statement can be made about the miscibility based on DSC determination of the glass transition temperatures. The crystallization endotherm peak during the cooling scan is used to determine the maximum crystallization temperature. The 75/25 blend does not have a value for the

Tc,max for the PA 6,10 component as the crystallization peak does not occur during the cooling step. This can be interpreted as another indication that the polymers have energetically favorable interactions. If the polymers were phase separated, the PA 6,10 domains would still crystallize. Instead, an annealing step is required. Percent crystallinity is reported in the table on a per polymer basis (i.e. crystalline polymer i / total polymer i). Values are reported in Table 1.5 and Figure 1.3.

Maximum Maximum Blend T Crystallinity Crystallization Crystallinity Crystallization Composition g Temperature Temperature wt% ˚C % ˚C % ˚C 100/0 34.9 ± 0.6 16.8 ± 4.2 163 ± 0.3 - - 75/25 32.1 ± 0.4 15.2 ± 4.7 161 ± 0.6 5.8 ± 1.5 - 50/50 - 17.2 ± 4.1 166 ± 0.9 13.3 ± 3.5 180 ± 2.1 25/75 - 13.6 ± 6.3 175 ± 14.2 16.6 ± 2.7 188 ± 0.7 0/100 38.2 - - 19.7 ± 0.9 195 ± 0.6

Annealing studies were performed to find maximum crystallinity of the 50/50 blend and the homopolymers by holding the samples at the two maximum crystallization temperatures of 162 ˚C (Tc,max of PA 11) and 185 ˚C (T c,max of PA 6,10 in the 50/50 blend). Table 1.6 shows that the overall crystallinity is maximized for the sample that is

annealed at both Tc,max temperatures. However, the difference is small compared to annealing only at the lower temperature. Once determined, the annealing sequence that returns the maximum crystallinity was applied as the pretreatment prior to mechanical testing. Figure 1.4 presents the complex melting peaks for the samples annealed under different conditions. The observed behavior is complex making analysis difficult. The

10 Percent Percent Crystallinity (%) 5 PA 11 Crystallinity PA 6,10 Crystallinity Overall Crystallinity 0 0 25 50 75 100 PA 11 (wt %)

Figure 1.3: Percent crystallinity of as prepared polymer blends upon second heating in the DSC. Open squares represent the percent by volume of the PA 11 that is in the crystalline state, circles represent PA 6,10, and Xs represents the overall percent crystallinity of the blend.

Table 1.6: Calculated overall crystallinity for 50/50 PA 11 / PA 6,10 blend for different annealing sequences. Annealling Overall % Temperature Crystallinity 167˚C 26.5 %

167˚C 181˚C 26.8 % 181˚C 22.6 %

two melting peaks are most clearly separated in the case of annealing at the high temperature first followed by the lower temperature. The high-low (HL) annealing procedure was applied to all samples and DSC was used to determine the overall percent crystallinity present during mechanical testing. Figure 1.5 and Table 1.7 present the results. The 25/75 and 75/25 blends both show single but broad melting peaks each of which is depressed relative to the homopolymer constituting the majority component. This is further evidence that the polymer system is miscible in the amorphous region and of the eutectic system discussed previously. If the polymers were immiscible in the melt state there would be a PA 11 rich phase and a

20 50/50 H 18 50/50 H-L 16 50/50 L

Unsubtracted Heat FLow (mW) 2

0 100 120 140 160 180 200 220 240 Temperature ( oC)

Figure 1.4: Typical melt peak of the 50/50 PA 11 / PA 6,10 blend annealed at: the high crystallization temperature only (H, T c=181), the high then the low crystallization temperatures in sequence (H-L), and the low crystallization temperature only (L, Tc=167).

40 38 36 0/100 34 25/75 32 50/50 30 75/25 28 26 100/0 24 22 20 18 16 14 12 10 8 6 Unsubtracted Heat FLow (mW) FLow Heat Unsubtracted 4 2 0 100 120 140 160 180 200 220 240 Temperature ( oC)

Figure 1.5: The typical melt peaks of all compositions (PA 11 / PA 6,10) using the maximum crystallinity annealing sequence found above (T c1 = 181, T c2 = 167 for 30 minutes each).

Table 1.7: Overall crystallinity for the annealing sequence that maximizes crystallinity for all compositions. Annealing Overall % Composition Temperature Crystallinity 100/0 163˚C 18.0 %

75/25 167˚C 181˚C 7.8 % 50/50 167˚C 181˚C 26.8 % 25/75 167˚C 181˚C 27.5 % 0/100 195˚C 26.5 %

PA 6,10 rich phase. Upon cooling it would be likely that each of the phases could produce semi-crystalline domains. This line of reasoning does the phases could produce semi-crystalline domains. This line of reasoning does open the possibility that the 50/50 blend may be immiscible in the melt. Flory-Huggins theory does predict that a blend of 50% by volume needs only a small positive interaction parameter to produce phase separation. Changes due to the annealling sequence are Highly significant. All blends except the PA 11 sample changed by greater than a factor of 35% when compared to the crystallinity of a sample without isothermal anealing. Crystallinity for PA 11 increases from 16.8 % to 18.0 %. The overall crystallinity of the blends changed from 12.8% to 7.8%, 15.3% to 26.8%, and 15.8% to 27.5% for the 75/25, 50/50, and 25/75 blends, respectively. PA 6,10 changed from 19.7% to 26.5%. The decrease in crystallinity for

the 75/25 is explained by the lack of any higher T m melting peak discussed above and plays a large role in the mechanical properties.

1.3.2 Thermogravimetric Analysis

TGA was used to find the onset degradation temperature of the homopolymers and their blends. Figure 1.6 shows PA 11 has the lowest onset degradation temperature at 419 ˚C and PA 6,10 has the highest at 437 ˚C. Onset degradation temperate for the blends have a monotonic change between the two homopolymers at 424 ˚C, 426 ˚C, and 431 ˚C for the 75/25, 50/50, 25/75 blends, respectively. These data show that at the maximum temperature of 260 ˚C used in the DSC experiments and at the lower

temperature used for thermal annealing of sample bars, that it is unlikely there is significant thermal degradation.

100 0/100 25/75 50/50 75/25 100/0

TG (%) TG

0 250 300 350 400 450 500 550 600 650 Temperature ( oC)

Figure 1.6: Representative TG data for PA 11 / PA 6,10 blends under an air atmosphere.

1.3.3 Dynamic Mechanical Thermal Analysis

DMTA is used to characterize the blends mechanical properties by measuring the storage (G’) and loss (G”) moduli at fixed frequency as a function of temperature as shown in Figure 1.7. Room temperature tests over longer sampling times were performed to find accurate storage and loss moduli reported in Table 1.8. Table 1.8 shows the storage modulus changes monotonically between the homopolymers. At room temperature PA 11 has a storage modulus of 0.828 GPa and PA 6,10 has a storage modulus of 1.173 GPa. The blends have a storage modulus of 0.847 GPa, 0.974 GPa, and 1.042 GPa for the 75/25, 50/50, and 25/75 blend compositions, respectively. DMTA is also used to characterize the blends thermal properties (HDT and α- relaxation) from the storage modulus and loss tangent as a function of temperature. The

1E9 100/0 25/75 50/50 75/25 1E8 100/0

1E8 G' G' (Pa) G'' (Pa) G''

1E7

1E7 0 50 100 150 Temp (C)

Figure 1.7: Storage (G') and Loss (G") Modulus of all compositions (PA 11 / PA 6,10) with respect to temperature.

Table 1.8: Comparison of room temperature storage (G') and loss moduli (G") for blends from values obtained with a temperature sweep and from room temperature only tests. G' G"

Room Temperature Percent Room Temperature Percent Composition Temperature Sweep Difference Temperature Sweep Difference PA11/PA610 GPa GPa % MPa MPa % 100/0 0.828 0.642 -22 1.8 7.1 304 75/25 0.847 0.806 -5 3.9 8.6 118 50/50 0.974 0.884 -9 6.8 9.4 38 25/75 1.042 0.722 -31 2.0 5.3 162 0/100 1.173 0.906 -23 5.0 8.9 78 heat distortion temperature (HDT) is the temperature at which a sample bar deflects 0.25 mm during a bending under load test described according to ASTM method D648. Takemori has established a correlation between the HDT and the temperature at which the shear modulus is equal to 0.28 GPa [73]. Using the Takemori method, the HDT was

measured as 52±0.6 ˚C and 76±6.7 ˚C for PA 11 and PA 6,10 respectively. The HDT of the blends were measured at 57±15 ˚C, 60±8.9 ˚C, and 54±1.4 ˚C for the 75/25, 50/50, and 25/75 blend compositions respectively. The HDT temperature for an amorphous polymer should be near the T g, however, for semi-crystalline polymers, the HDT is

somewhere between the T g and T m [73]. The loss tangent (tan ( δ) = G”/G’) was used to find the α-relaxation associated

with the T g. The α-relaxation occurs at the peak of the tan ( δ) as a function of temperature and is measured at 56±0.8 ˚C and 55±3.7 ˚C for PA 11 and PA 6,10 homopolymers respectively. The blends α-relaxation is measured at 55±3.6 ˚C, 56±3.3 ˚C, and 55±2.2 ˚C for the 75/25, 50/50, and 25/75 blend compositions respectively. The measured α-relaxation of all compositions remain constant at about 55 ˚C compared to literature values of 47 ˚C for PA 11 and 52 ˚C for PA 6,10 [81]. There is only a single peak as a function of temperature which, like DSC, indicates a single glass transition temperature. However, the width of the tan ( δ) peak is narrowest for the homopolymers and broadest for the 50/50 blend. Breadth of this peak can be associated with the molecular heterogeneity of the amorphous phase leading to additional supporting data for a miscible amorphous phase.

0.2 100/0 25/75 50/50 75/25 0/100

0.1 Intesnsity

0.0

0 50 100 150 Temp (C)

Figure 1.8: tan ( δ) for all compositions (PA 11 / PA 6,10) as a function of temperature.

1.3.4 Impact Strength

Figure 1.9 shows that the measured impact strength of the blends ranges from 40- 75 J/m and shows positive deviation from a simple linear mixing rule between the two homopolymers. It has been shown for other aliphatic polymer blends that mechanical properties show positive deviation from a simple linear mixing rule between two homopolymers [81]. The homopolymers impact strength are 71±13 J/m and 43±8 J/m for PA 11 and PA 6,10 homopolymer, respectively. The impact strengths 72±13, 66±5, 51±4 correspond with the 75 wt%/25wt%, 50wt%/50wt%, and 25wt%/75wt% PA 11/PA 6,10 blends, respectively. The measured impact strength for both homopolymers is consistent with literature values; this is within the uncertainty of the sample preparation and measurement procedures.

50 Impact Strength (J/m) Strength Impact

40 Measured Literature Data

30 0 25 50 75 100 PA 11 (wt %)

Figure 1.9: Measured impact strength of blends without annealing. This data shows a higher than a linear mixing rule would indicate.

1.3.5 Tensile Test

Tensile testing was used to determine Young’s Modulus (tensile modulus), strain at break, and tensile strength at yield. Young’s Modulus was measured from 1.3 to 2.5 GPa for the as molded samples and 1.5 to 2.0 GPa for the re-crystallized samples; these results are shown in Figure 1.10. A crossover in modulus values occurs as the composition of PA 11 increases; in most cases the as-molded samples have higher moduli but for the PA 11 homopolymer this is reversed. The low at the 75/25 blend can be attributed to crystallinity changes as measured by DSC. It is typical for the Young’s Modulus to increase with increasing crystallinity. This was not obsereved for blends with a major component of PA 6,10. This is the result of annealing which has been documented to increase void areas from where polymer chains have been incorporated into the crystal structure. This results from a decreased chain mobility and increases the time required to densify the amorphous regions (procedure used 40 hours at room temperature). This would increase the stress on the fewer chains in th amorphous regions, leading to more streching and a lower Young’s moduli [82]. Strain at Break was measured from 133% to 285% for as molded samples and 6% to 47% for thermally treated samples as seen in Figure 1.10. Thermally treated samples have higher crystallinity and therefore the amorphous phase is prevented from flowing into the necking regions. This decreases the elongation ability of the polymer samples. The as molded samples are observed to be slightly higher than monotonic increase between the two homopolymers. The thermally treated samples reach a maximum at the 50/50 composition. Tensile strength at yield was measured from 50-78 MPa for as molded samples and 45-60 MPa for thermally treated samples with both showing a maximum at the 25/75 blend composition as seen in Figure 1.11. It should be noted that the calculations are for engineering tensile strength at break. These ranges corraspond with the aproximate literature values found for PA 11 and PA 6,10. Tensile test derived Young’s modulus (E) and DMTA derived complex modulus

( ∗ ) can be used to find the Poison’s ratio of the blends. The equation = ′ + " relating Young’s modulus to the shear modulus through poison’s ratio is . = 21 +

450 Modulus, As Molded Modulus, Thermal Treatment Modulus Literature 400 2.5 Strain, As Molded Strain, Thermal Treatment Strain Literature 350

300

2.0 250

200

150

1.5 Strain (%) Breakat Young'sModulus(GPa) 100

1.0 0 0 20 40 60 80 100 PA11 (wt%)

Figure 1.10: Young's modulus and strain at break measured for as molded and thermally treated tensile bars.

Tensile Strength (MPa) As Molded Thermal Treatment Literature

20 0 25 50 75 100 PA11 (wt%)

Figure 1.11: Tensile strength at yield for all five compositions of the polyamide blends.

Poisson’s ratio was calculated using the shear modulus and the Young’s modulus of crystallized samples and the results are shown in Table 1.9. Poisson’s ratio for rigid polymers is typically between 0.3-0.5 with 0.5 being the maximum and a perfectly incompressible material. Literature has the Poisson’s ratio of PA 11 and PA 6,10 within the range of 0.3-0.4. The calculated Poisson’s ratio of the blends shows a poor agreement between the shear modulus and Young’s modulus. As the shear modulus was measured at room temperature with extra care to insure the samples were crystallized and dry, the Young’s modulus for the thermally treated samples must be under measured. This was likely the result of the correction factor changing with composition and crystallinity and possibly the water content within the polyamide samples. The Young’s modulus would need to be 15-50% higher in order for the calculated Poisson’s ratio to agree with literature ranges.

Table 1.9: Calculated Poisson's ratio from thermally treated samples of Young's modulus and complex modulus showing a discontinuity between measured modulus. G E PA11/PA610 (Complex Modulus) (Young’s Modulus) (Poisson’s Ratio) wt % GPa GPa - 100/0 0.83 1.88 0.13 75/25 0.85 1.55 -0.08 50/50 0.97 1.76 -0.10 25/75 1.04 1.99 -0.04 0/100 1.17 2.06 -0.12

1.3.6 Microscopy

Samples were viewed under light microscopy in order to determine if separate domains were present in the melt phase (above both melting temperatures) and in the amorphous phase (below both melting temperatures). There was no observed difference, indicating miscibility. Calculations based on contribution method were performed to find approximate values for index of refraction in the two polymers as shown in Table 1.10. The values calculated are shown in the table below and indicate that there may not be a large enough difference in index of refraction to observe separate domains of the polyamide blends. The calculations were performed following both the

electromagnetism and an empirical basis of calculations based on the contribution method [80]. It was also found that the literature reports isotropic index of refraction as 1.52 for both homopolymers (an insufficient value to determine observability between phases).

Table 1.10: Calculated index of refraction from two different contribution methods at a wavelength of 589 nm. n (Refractive Index with n (Refractive Index with

Electromagnetism basis) Empirical basis) Nylon 11 1.4812 1.4952 Nylon 6,10 1.4816 1.4990

Cross polarization was used on identically annealed polymer samples of different compositions to see the difference in the spherulitic size or shape in the blends. Figure 1.12-Figure 1.16 show all homopolymers and blends under cross polarized light microscopy at 22x magnification. It can be qualitatively observed that the size of the spherulites is slightly larger in the blends when compared to the homopolymers with the 25/75 blend being the largest.

0/100

Figure 1.12: Cross polarized microscopy image of PA 6,10 crystallized at 185 ˚C then 162 ˚C depicting the size of spherulites in the sample.

25/75

Figure 1.13: Cross polarized microscopy image of 25/75 PA11/PA6,10 blend crystallized at 185 ˚C then 162 ˚C depicting the size of spherulites in the sample.

50/50

Figure 1.14: Cross polarized microscopy image of 50/50 PA11/PA6,10 blend crystallized at 185 ˚C then 162 ˚C depicting the size of spherulites in the sample.

75/25

Figure 1.15: Cross polarized microscopy image of 75/25 PA11/PA6,10 blend crystallized at 185 ˚C then 162 ˚C depicting the size of spherulites in the sample.

100/0

Figure 1.16: Cross polarized microscopy image of PA 11 crystallized at 185 ˚C then 162 ˚C depicting the size of spherulites in the sample.

1.3.7 Small Angle X-ray Scattering

The size of the lamellar long spacing as a function of blend composition and annealing strategy were examined by SAXS. The polymers are not oriented in anyway, so the diffraction pattern is isotropic around the center of the beam block. There were two sets of data collected, one crystallized only at 162 ˚C as seen in Figure 1.17 and the other crystallized only at 185 ˚C as seen in Figure 1.18, corresponding with maximum crystallization temperature of PA 11 and PA 6,10, respectively, in the 50/50 blend. The lamellar long spacing is shown in Table 1.11. For the samples crystallized at the PA 11 maximum crystallinity temperature, it can be seen that the long spacing of the lamellar structure decreases slightly from the homopolymer. This is due to the bulk averaging limitations of SAXS. The decrease can be explained by a weighted average of the two homopolymer long spacing. The lamellar long spacing for samples crystallized at the PA 6,10 maximum crystallization temperature increases as more PA 11 is added to the system. This results from PA 11 being trapped within the crystal superstructure as the

0.8

0.7 25/75 Tc=162 50/50 Tc=162 0.6 75/25 Tc=162 100/0 Tc=162 0.5

0.4

Intesnsity 0.3

0.2

0.1

0.0 0.00 0.02 0.04 0.06 0.08 0.10 0.12 0.14 0.16 0.18 q (Å -1 )

Figure 1.17: 1-D SAXS scattering of blends annealed at the crystallization temperature of 162 ˚C showing no change between compositions (PA 11 / PA 6,10).

0.8

0.7 0/100 Tc=185 25/75 Tc=185 0.6 50/50 Tc=185 75/25 Tc=185 0.5

0.4

Intesnsity 0.3

0.2

0.1

0.0 0.00 0.02 0.04 0.06 0.08 0.10 0.12 0.14 0.16 0.18 q (Å -1 )

Figure 1.18: 1-D SAXS scattering of blends annealed at the crystallization temperature of 185 ˚C. The decreasing scatter intensity (q) correlates with increasing long spacing and suggests PA 11 chains are “trapped” as the PA 6,10 crystal structures grow.

Table 1.11: Table showing the long spacing calculated based on the maximum for all blend compositions at the two different annealing temperatures Long Spacing (Å)

PA11/PA6,10 Crystallization Crystallization wt% Temperature = 162 ˚C Temperature = 185 ˚C 100/0 90 - 75/25 90 99 50/50 88 94 25/75 88 90 0/100 - 86

PA 6,10 lamella are forming. As a result, the long spacing, which includes lamellar and inter-lamellar amorphous regions, gets larger due to the increase in the amorphous regions dimensions. This has been shown to occur for crystalline/amorphous polymer mixtures [83-86].

1.3.8 Wide Angle X-ray Scattering

The crystalline dimensions of the polyamide blends were examined as a function of blend composition and annealing strategy by WAXS. Literature reports PA 11 and PA 6,10 as having crystalline lattice structures of the α-triclinic shape which can be seen in Figure 1.19. The literature values for the lattice structures are given in Table 1.12 along with the associated calculated and measured d-spacings that will be present for the crystals of PA 11 and PA 6,10. There is correlation from the literature to the measured samples. The slight degree of systematic error is due to the long dimension (c) being along the polymer chain, which can easily stretch to a longer length and short dimensions (a and b) are pulled closer together as the c-dimension elongates.

Figure 1.19: Representative lattice structure for α-triclinic geometry where a, b, and c are lattice dimensions and α, β, and γ lattice angles.

Table 1.12: Lattice dimensions and angles from literature, calculated d-spasings, and measured d-spacings for PA 11 and PA 6,10. PA 11 PA 6,10

Lattice Calculated Measured Lattice Calculated Measured Constants d-spacing d-spacing Constants d-spacing d-spacing a (Å) 4.9 3.4 3.8 4.9 3.9 3.7 b (Å) 5.4 4.3 4.2 5.4 4.4 4.4 c (Å) 14.9 9.6 11.8 22.4 16.9 17.6 40˚ 49˚

77˚ 76.5˚

63˚ 63.5˚

There were two sets of data collected, one crystallized only at 162 ˚C shown in Figure 1.20 and the other crystallized only at 185 ˚C shown in Figure 1.21, corresponding with maximum crystallization temperature of PA 11 and PA 6,10 respectively in the 50/50 blend. The measured crystalline dimensions are given in Table 1.13 and Table 1.14 for the samples crystallized at 162 ˚C and 185 ˚C respectively. The agreement for the peaks that are known to be associated with only one of the crystal structures are very consistent (i.e. and those that are shared vary slightly as is expected. The - peak changes monotonically the from the dimension of PA 11 to that of PA 6,10. WAXS is a bulk analysis tool, meaning that the measurement is an average of everything in the sample, therefore, when two peaks overlap closely the dimension being measured show a weighted average based on the number of instances of each crystal dimension. Qualitatively, the WAXS pattern appears to shift quite significantly toward the PA 6,10 crystal as the composition increases. This is evidence

1.5 25/75 Tc=162 50/50 Tc=162 ? 75/25 Tc=162 100/0 Tc=162

1.0

? ?

Intesnsity 0.5

0.0 0.0 0.5 1.0 1.5 2.0 2.5 3.0 q (Å -1 )

Figure 1.20: 1-D WAXS scattering of blends annealed at the crystallization temperature of 162 ˚C.

1.5 0/100 Tc=185 ? 25/75 Tc=185 50/50 Tc=185 75/25 Tc=185

1.0

? ?

Intesnsity 0.5

0.0 0.0 0.5 1.0 1.5 2.0 2.5 3.0 q (A -1 )

Figure 1.21: 1-D WAXS scattering of blends annealed at the crystallization temperature of 185 ˚C.

Table 1.13: Associated peaks for PA 11 ( ) and PA 6,10 ( ) taken from the peaks of the 1-D WAXS data for the samples annealed at crystallization temperature of 162 ˚C. Associated 100/0 75/25 50/50 25/75 Peak (Å) 11.85 11.85 11.8 11.75 (Å) - 8.435 8.578 8.583 -? (Å) 4.237 4.325 4.324 4.364 (Å) - - - 4.07 -? (Å) 3.821 3.808 3.847 3.826 (Å) 2.354 2.363 2.37 2.381 ? that the crystals are simultaneously crystallizing. The mechanism is most likely consistent with crystal phase induced separation. This data shows consistent correlation to the PA 6,10 crystal dimensions. The peak starts very small, and gets lost in noise even at the lowest composition of the PA11. Qualitatively, the WAXS pattern for the samples annealed at 185 ˚C appears to remain consistent despite compositional changes. This is due to the melting temperature of the

Table 1.14: Associated peaks for PA 11 ( ) and PA 6,10 ( ) taken from the peaks of the 1-D WAXS data for the samples annealed at crystallization temperature of 185 ˚C. Associated Peak 75/25 50/50 25/75 0/100 (Å) - - - 17.58 (Å) 11.73 11.78 11.7 - (Å) 8.39 8.53 8.568 8.531 -? (Å) 4.306 4.321 4.329 4.372 (Å) 4.026 - 4.094 - -? (Å) 3.709 3.761 3.78 3.743 (Å) 2.347 2.357 2.364 2.37 ? PA 11 being lower than the crystallization temperature. There is still PA 11 crystals present as PA 11 crystallizes quickly, however, the amount measured by WAXS is lower when compared to the same composition of the samples annealed at 162 ˚C.

1.4 Conclusion

Aliphatic polyamides made from renewable resources are of increasing interest. In this study, PA 11 and PA 6,10 blends were investigated for the first time and found to represent a scientifically rich system. The thermal and mechanical properties as well as the morphology of the blends between PA 11 and PA 6,10 were investigated over the entire composition range. The DSC was used to show a preferential interaction of the PA 11 / PA 6,10 molecules in the melt and the amorphous region. Melting temperature depression was used to show miscibility in the blends through negative polymer-polymer interaction parameters. Crystallinity levels explain mechanical property changes for annealed samples. Mechanical properties of the blended polymers were found to be intermediate of the homopolymer. A higher than monotonic change between the PA 11 and PA 6,10 was observed for the impact strength. This is evidence of only one amorphous phase, as it has been shown that multiple incompatible phases decrease mechanical properties due to low interfacial energy. The tensile testing of the as molded samples showed higher than monotonic changes between the homopolymer properties for Young’s modulus, strain at break, and tensile strength at yield. Thermally treated samples showed a Young’s

modulus that changed monotonically between pure polymer samples, a strain at break that was better than monotonic and peaked at the 50/50 blend, and a tensile strength at yield that was relatively insensitive to composition. Morphology was analyzed through the use of microscopy and x-ray scattering. The miscibility of the PA 11 and PA 6,10 in the melt and the amorphous region was observed through microscopy, as well as the spherulitic size differences through the compositions. X-ray scattering was used to find the differences in the blends on the scale of crystal lattice structure and the lamellar long spacing. The full suite of thermal and mechanical properties reported for the biorenewable PA 11 / PA 6,10 for the first time had great mixing of mechanical properties and where found to be miscible in the melt and the amorphous state.

CHAPTER 2 2 RENEWABLE RESOURCE BASED BIOPLASTIC BLENDS: POLYAMIDE-11 / POLYLACTIDE BINARY SYSTEM

Polylactide (PLA) is an aliphatic, crystalline, thermoplastic polyester derived from renewable resources [87, 88]. In composting environments PLA undergoes degradation via lactic acid, which is a non-toxic, naturally occurring metabolite compound [89-92]. PLA is now produced on a large scale by Natureworks LLC, Shimadzu Corporation and Mitsui chemicals in Japan, Apack AG in Germany, and Bio Invigor in Taiwan [93-95]. So far, the major applications for PLA are in packaging, polyester fibers, and in surgical sutures [90, 96, 97]. However, PLA is also being used as material for components in laptops, cell phones, various electronic goods, as well as automotive parts, such as floor mats, headliners, and spare-tire covers. PLA is thus attracting increasing industrial attention as it becomes price competitive with plastics made from fossil fuels [98, 99]. The mechanical property window of PLA is comparable to that of some conventional commodity plastics [63, 100]. For example, PLA closely resembles polystyrene in terms of high modulus and low elongation to break and poly(ethylene terephthalate) in terms of stiffness and tensile strength [101]. These properties are favorable, but in combination with a low heat distortion temperature and moderate gas barrier properties of PLA, the range of applications is somewhat limited [98, 102]. Polyamide-11 (PA 11) is derived from castor oil, a renewable resource extracted from the castor been. PA 11 is a high-performance, semi-crystalline polymer with low

net CO 2 emission and global warming potential, when compared to petroleum based nylons and other conventional plastics. Some of the outstanding properties of PA 11 include high impact and abrasion resistance, low specific gravity (about 1.03 g/cm 3), excellent chemical resistance, high thermal stability (melting point of about 180-190 ˚C), and ease of processing over a wide range of processing temperatures (-40 ˚C to 130 ˚C) [63]. PA 11 has excellent dimensional stability, and maintains physical properties over a wide range of temperatures and environments. PA 11 thus finds numerous applications

in fields including crude oil and gas pipelines, hydraulic and pneumatic hoses, powder coatings, skis, snowboards and optical fiber sheathing. However, it is relatively expensive compared to other plastic materials including petroleum based Nylon-6 [63, 103-105]. Polymer blending is an established and cost-effective way of developing new polymeric materials with diversified and desirable properties [106]. Blending allows tailoring of material properties but immiscibility can lead to inferior properties and compatibilization is often required. Reactive compatibilization can be accomplished through the generation of graft or block copolymers in-situ during melt blending. Such copolymers are formed by chemical reactions between immiscible polymers at the interface where the compatibilizer is needed [107, 108]. In the case of condensation polymers, chemical reactions between functional end groups are common during reactive blending and are generally known as “exchange” or “interchange” transreactions. The physical, thermal, mechanical, and morphological properties of the polymer system are affected by the type and extent of these transreactions [107]. Generally, it is found that, independent of the chemical mechanism of exchange, the interchange reactions initially lead to the formation of block copolymers, followed by random copolymers [107, 109- 113]. Interchange reactions involving condensation polymers such as polyesters and polyamides have been reviewed by Kotliar in detail [114]. Possible transreactions for polyesters and polyamides are depicted in Figure 2.1 and include alcoholysis and aminolysis. Reactively compatibilized melt blends of PLA have been studied with many different polymers [98, 115-118]. However, there are only few reports on

R NH CO R R NH 1 2 Aminolysis 1 2 OCR2

HN R3 R3 NH2

R1 O COR2 R1 OH OCR Alcoholysis 2

O R3 R3 OH

Figure 2.1: Proposed interchange reactions in polyester-polyamide blends. polyamide-polyester blends utilizing a catalyst [119-122] (and none on PLA-PA 11) to enhance interchange reactions. Titanium alkoxides are widely reported as effective catalysts for interchange reactions [109, 111, 123, 124]. Accordingly, the present study seeks to use TIP to compatibilize PLA with PA 11.

2.1 Materials

Commercial-grade PLA (2002D, melt-flow index 4–8 g/10 min, < 4% D-lactide) was obtained from Cargill Dow Polymers (now Natureworks LLC, Minnetonka, MN). Polyamide-11 (PA 11) was purchased from Rilsan Arkema in pellet form. Titanium(IV)isopropoxide (TIP), was obtained from Sigma Aldrich. Irganox 1076 and Irgafos 168, added as thermal stabilizers [125], were obtained from Ciba Specialty Chemicals Inc. and used as received. Chloroform, toluene and m-cresol were reagent grade, and obtained from Sigma-Aldrich.

2.2 Methods

Blends of PLA/PA 11 were prepared by melt mixing in a Haake Rheomix 3000 at 205 ˚C for 20 min at 50 RPM using roller blades. Compositions studied were 0/100, 25/75, 50/50, 75/25 and 100/0 wt% of PLA/PA 11. To minimize the possibility of hydrolysis during the mixing process, PLA and PA 11 were first dried at 80 ˚C under vacuum (23 inch Hg) for at least 20 hours. Pellets of PLA were introduced first in the preheated mixing bowl at 205 ˚C, followed by PA 11 pellets. After the temperature stabilized at 205 ˚C the stabilizing agent was added to give an overall composition of 0.25 wt% each, followed by addition of TIP at varying compositions, both in separate toluene solutions. The polymers were then mixed thoroughly for 20 minutes. Upon cooling, the blended material was ground in a Foremost A2 grinder to a maximum pellet size of about 5 mm with the vast majority of particles being 2-3 mm in diameter. Sample bars for dynamic mechanical thermal analysis (DMTA) were prepared

by a combination of vacuum and compression molding. The sample material was first melted at 185 ˚C under vacuum (about 23 inch Hg) until degassing ceased. Afterwards the material was compression molded at 185 ˚C under a pressure of 5000 psi for 3 minutes and quenched between cooled aluminum plates. Prior to DMTA testing, sample bars were crystallized at 110 ˚C for 3 hours, and physically aged for 24 hours at ambient temperature. Tensile and impact bars were prepared by injection molding at 185 ˚C under a clamp force of 10 tons and barrel pressure of 8 bars. Chloroform is a solvent for PLA and a non-solvent for PA 11. Selective solvent extraction was achieved by adding ground pellets to chloroform at about 10 wt%. The resulting mixture was stirred for 48 hours before being filtered using a Buchner funnel and Whatman No.2 filter paper. The insoluble fraction was extensively washed with chloroform to remove any residual PLA and then dried in a vacuum oven at 80 ˚C for two days. This insoluble dried material was then used for NMR analysis. High-resolution 13 C Nuclear Magnetic Resonance (NMR) spectroscopy was used to examine possible ester-amide interchange reactions between PLA and PA 11. NMR spectra were recorded on a QE-300 NMR spectrometer with field strength of 300 MHz. The samples for NMR analysis were obtained from the chloroform insoluble fractions prepared as stated above and dissolved in m-cresol to obtain around 15 wt% solutions. Field-emission scanning electron microscopy (FE-SEM) was used to investigate the phase morphology and fracture surfaces of the blends. A JEOL JSM-7000F microscope was used with an acceleration voltage of 1.5 KeV, small probe current (2 kV), and working distance of 6 mm. Compression molded bars were freeze fractured using liquid nitrogen. Fractured bars were then treated with 0.1 M sodium hydroxide (NaOH) solution at 80 ˚C for 10 hours to etch out the PLA phase and then dried at 80 ˚C under vacuum (about 23 inch Hg ) overnight. Samples were then sputter coated with gold under vacuum at 15 mA. The thermal properties of the pure components and their blends were investigated using differential scanning calorimetry (Perkin-Elmer DSC-7 calibrated against an Indium standard). Sample masses of 10-15 mg were sealed in aluminum pans. The DSC protocol was: 1) heat from 5 ˚C to 200 ˚C at 10 ˚C/min, 2) hold at 200 ˚C for 5 min, 3) cool from 200 ˚C to 110 ˚C at 50 ˚C/min, 4) hold at 110 ˚C for 30 min, 5) cool from

110 ˚C to 5 ˚C at 10 ˚C/min, and 6) heat from 5 ˚C to 200 ˚C at 10 ˚C/min. Melting points (T m) and crystallinities were recorded based on data from the second heating run; glass transitions (T g) were recorded from the first heating run. The T g was calculated as the midpoint of the heat capacity change. Thermal stability of the blends was characterized by thermogravimetric analysis (TGA) using a Seiko TG/DTA 220 instrument calibrated against an indium standard. 10-15 mg of sample was placed in an alumina pan, and the experiment was carried out under air or nitrogen gas flow. The heating program used was heating from 30 ˚C to 800 ˚C at 10 ˚C/min. The degradation temperature is calculated as the temperature at which the maximum rates of decomposition occurred. Shear modulus as a function of temperature was determined by Dynamic Mechanical Thermal Analysis (DMTA) using compression molded bars in an ARES-LS rheometer with torsional rectangular fixtures. The machine was calibrated before testing for both normal force and torque. The testing was carried out at 0.05 % strain and a frequency of 1 Hz; the temperature was ramped from 30 ˚C to 160 ˚C at 5 ˚C/min. Impact properties of polymer blends were measured according to the ASTM standard D256 method for notched Izod impact testing. Impact testing bars (dimensions 12.7×63.5×3.2 mm) were injection molded using a Morgan press and notched using a manual Charpy broach. Impact bars were crystallized at 110 ˚C for 3 hours, and physically aged for 24 hours at ambient temperature, before testing. Notched Izod impact testing was carried out on a TMI Izod impact instrument (Model 43-02-01) at room temperature with a pendulum hammer of 2 lb and 5 lb (for 0/100 wt% PLA/PA11 blend). An average value was obtained by testing a minimum of five replicates. Tensile testing of injection molded bars was performed at room temperature on an MTS Alliance RT /100 testing machine according to ASTM D638. The tests were conducted at a crosshead speed of 1 in/min using a 1 inch extensometer. A minimum of five samples of each composition were tested.

2.3 Results and Discussion

Mechanical and thermal properties of the blends are characterized by differential scanning calorimetry (DSC), thermo gravimetric analysis (TGA), dynamic mechanical thermal analysis (DMTA), tensile testing, and impact testing. Phase morphology of the blends is examined directly using field emission scanning electron microscopy (FE- SEM). It is determined that the blends are immiscible and the occurrence of compatibilizing reactions, if any, is largely obscured by degradation reactions. Nonetheless the blends do show intermediate properties between the homopolymers that may be useful in various applications.

2.3.1 Extractions

The results of all extraction tests are summarized in Table 2.1. Blends were prepared as described above with varying levels of catalyst solution only at 185˚C. The loading levels of catalyst were changed gradually from 0.01 wt% to 0.1 wt% of total polymer mass.

Table 2.1: Weight percent of insoluble material for different catalyst loading levels, reaction temperatures and reacting times. Blend Melt Weight % of Melt blending Catalyst composition Mixing Insoluble Temperature Level (PLA/PA11) Time Material wt% ˚C Minutes wt% wt% 50/50 * - - - 52.8 25/75 205 20 0.05 67.4 50/50 185 20 0.0 55.1 50/50 185 20 0.01 56.1 50/50 185 20 0.04 56.6 50/50 185 20 0.05 57.3 50/50 185 20 0.07 57.5 50/50 185 20 0.10 57.2 50/50 205 20 0.05 60.3 50/50 225 20 0.05 58.2 50/50 225 40 0.05 56.7 *PLA and PA 11 polymer pellets were measured in 50/50 wt% composition and added to a concentration of 10 wt% chloroform at room temperature

Figure 2.2 shows the wt% of chloroform insoluble material as a function of catalyst loading level for 50/50 wt% blends of PLA and PA 11 mixing at 185 ˚C. The amount of insoluble material increases with catalyst level up to 0.07 wt% then decreases at the 0.1wt% catalyst level. The control experiment of 0 wt% catalyst demonstrates that the insoluble mass is greater than the initial PA 11 mass. This is attributed to the solution coating of PA 11 domains by PLA or by inclusion of PLA inside a shell of PA 11. Nonetheless, when catalyst is added the insoluble material obtained consistently exceeds the control experiment. If the amount of chloroform insoluble material is indicative of compatibilization, the results show that there is little difference in the catalyst loading levels from 0.05 – 0.10 wt%. Accordingly, the minimal effective amount of 0.05 wt% is used in all other experiments.

58.0

57.5

57.0

56.5

56.0

55.5 Wt.material % Insoluble 55.0

0.00 0.02 0.04 0.06 0.08 0.10 Amount of TIP (wt%)

Figure 2.2: Chloroform insoluble weight percent as a function of catalyst loading level at 185 ˚C.

Similarly, the amount of insoluble material increases as the blending temperature was increased from 185 ˚C to 205 ˚C. However, a further increase to 225 ˚C leads to a decrease in the insoluble fraction. The likely cause for this observation is that at 225 ˚C significant PLA degradation occurs; significant color formation is visible with the naked eye. Thus, 205˚C is used for all further experimentation.

2.3.2 Nuclear Magnetic Resonance on Carbon Isotope 13

13 C-NMR spectra and chemical structures are summarized in Figure 2.3; peak assignments are labeled with letters. The NMR spectra obtained from the insoluble fraction (rich in PA 11) of the selective solvent extraction experiments are simply a superimposition of the NMR spectra for pure PLA and pure PA 11. That is, they show

a) a' b' c'

e-j d d) a k b c

180 160 140 120 100 80 60 40 20 PPM O b' CH C O a' n CH3 c' O Polylactide e-j k C b d a c NH n

Polyamide-11 Figure 2.3: C 13 NMR spectra for a) pure PLA, b) chloroform soluble fraction, c) chloroform insoluble fraction, and d) pure PA 11.

only peaks corresponding to those of the homopolymers. Similarly, spectra obtained from the soluble fraction show only peaks corresponding to those of pure PLA. This behavior is observed for all PLA/PA 11 blends irrespective of the blend composition, melt blending time and blending temperature. If interchange reactions between PLA and PA 11 occur, new peaks corresponding to the copolymer should appear in the 13 C-NMR spectra. However, the absence of any new peak indicates that either no interchange reaction occurred or, if reaction occurred, the extent of this reaction was not significant enough to be detected by NMR spectroscopy. It should be noted though that PLA is being observed in the chloroform insoluble faction; some researchers take this as direct evidence of interchange reactions occurring. The literature suggests that prolonged exposure to higher temperatures may be necessary for the occurrence of interchange reactions [111, 126]. The limited interchange reaction in the present binary system is attributable to the difficulty of processing these polymers at higher temperatures for longer periods of time. At elevated temperatures, degradation reactions dominate over transreactions between polymer chains.

2.3.3 Differential Scanning Calorimetry

Figure 2.4 shows the glass transition behavior of PLA/PA 11 blends for various compositions. It can be observed in Figure 2.4 that the T g of PLA is at 61.2 ˚C, and that

of PA 11 is at 45.6 ˚C. All blends are characterized by the presence of two distinct glass transition temperatures: the first at about 45 ˚C, and the second at about 57 ˚C. It is well established that a single, compositionally dependent, glass transition is associated with miscibility and molecular level mixing of the component polymers [127].

Thus, the T g data presented in Figure 2.4 indicates that blends of PLA and PA 11 are largely immiscible. However, the T g of the PLA fraction is increasingly depressed with the addition of PA 11; this indicates some compatibility between the polymers, however, the degree of compatibility is clearly very small.

T 0/100 g,PA11

Tg,PLA Tg,PA11 ) 25/75 → Tg,PLA

Tg,PA11

50/50 Tg,PLA T 75/25 g,PA11 Tg,PLA Heat Flow (Endo 100/0

40 45 50 55 60 65 70 ° Temperature ( C) Figure 2.4: DSC heat flow over the glass transition regions of PLA/PA 11 showing two glass transitions for all compositions suggesting immiscibility.

Melting curves for different PLA/PA 11 blends are shown in Figure 2.5 and the corresponding data are summarized in Table 2.2. Neat PLA (taken as the α-phase) shows α α a melting peak at T m1 = 162.2 ˚C followed by a larger peak at T m2 = 167.9 ˚C; the bimodal melting peaks are due to a distribution in lamellar thicknesses. The true melting α point at which all of the various sized crystallites are melted occurs at T m = 174.4˚C. The double melting peaks of PLA shift towards each other as the composition decreases in PLA/PA 11 blend and a single melting peak for PLA component is observed for 25/75 wt% PLA/PA 11 blend. Double melting peaks are also observed for PA 11 (taken as the β-phase). The PLA/PA 11 blends exhibit two different melting peaks, corresponding to those of the homopolymers, which do not change with composition and therefore show no miscibility. This behavior is independent of the blend composition, melt temperature, and mixing time.

T α Tm1 α m2

Tmα PLA, 205 °C

75/25, 205 °C

50/50, 225 °C ) →

50/50, 205 °C

50/50, 185 °C Heat Heat Flow (Endo 25/75, 205 °C β Tm2 β Tm1 β Tm PA11, 205 °C

140 150 160 170 180 190 200 Temperature (°C)

Figure 2.5: DSC heat flow over the range of melting temperatures for PLA/PA 11 blends prepared at different mixing temperatures.

Table 2.2: DSC parameters obtained from the first and second heating scans for various PLA/PA 11 blend compositions. PLA component ( α) PA 11 component ( β) PLA/PA1 α α α α α β β β β β 1 (wt%) Tg Tm1 Tm2 Tm ∆Hm Tg Tm1 Tm2 Tm ∆Hm (˚C) (˚C) (˚C) (˚C) (J/g) (˚C) (˚C) (˚C) (˚C) (J/g) 0/100 - - - - - 45.6 181.5 190.1 194.9 40.17 25/75 56.1 - 164.4 170.5 5.51 44.8 179.2 189.1 194.7 32.62 50/50 56.9 160.7 166.1 172.2 13.33 43.1 179.0 188.9 194.5 23.23 75/25 58.7 161.1 167.7 173.5 29.15 42.5 - 188.1 193.6 13.89 100/0 61.2 162.2 167.9 174.4 37.44 - - - - -

2.3.4 Thermogravimetric Analysis

Figure 2.6 shows the TGA curves for the homopolymers and their blends. The neat PA 11 shows the highest degradation temperature at 429.8 ˚C, whereas the neat PLA shows the lowest degradation temperature at 348.1 ˚C. For each blend a two-stage degradation curve is observed, which shifts to a slightly higher temperature with increasing PA 11 content. Onset degradation temperatures for the 25/75, 50/50, and 75/25 blends are 341.6, 344.4, and 346.2 ˚C respectively. The presence of two distinct weight-loss steps indicates separate degradation of each polymer. In addition, the onset of the second degradation step (between 350 ˚C and 475 ˚C) corresponding to PA 11 occurs earlier in the blends than for pure PA 11. Onset degradation temperatures for the second stage of the 25/75, 50/50, and 75/25 blends are 423.7, 422.1, and 419.3 °C respectively. Clearly, the PLA/PA 11 blends have lower thermal stability than the PA 11 homopolymer.

100 PLA/PA11 (wt%) 0/100 25/75 80 50/50 75/25 60 100/0

40 Weight (%)

0 250 300 350 400 450 500 550 600 Temperature ( °C)

Figure 2.6: TG data for PLA/PA 11 blends prepared at 205 ˚C for 20 minutes of melt- blending.

2.3.5 Scanning Electron Microscope

Figure 2.7-Figure 2.9shows the phase morphology of different wt% compositions of PLA/PA 11 blends obtained by FE-SEM after removing the PLA phase via base etching. Two phases are observed in each of the blends, confirming the basic immiscible nature of the PLA/PA 11 blends. A “droplet morphology” is observed in all the blends with a continuous matrix or “major” phase and a separate dispersed or “minor” phase. For example, in the 25/75 wt% PLA/PA 11 blend, the PLA is the minor phase dispersed in the PA 11 major phase matrix. The PLA phase appears as small holes because of the base etching procedure. Similarly, the 50/50 blend shows a minor PLA phase due to the greater density, and therefore smaller volume, of the PLA. When PLA is the dispersed phase, the average size of the droplets increased with concentration from 25 wt% to 50 wt%. The approximate diameters of the dispersed phase varies from 0.3-0.7 µm for 25/75 wt% PLA/PA 11 blend to 2-3 µm for 50/50 wt% PLA/PA 11 blend. At 75 wt% PLA, the morphology is very different; now the etching procedure removes the major phase and distinct droplets of the PA 11 of approximate size 0.5-1 µm are observable. Accordingly, Figure 2.7-Figure 2.9 reveals the transition from a PA 11 major phase to a PLA major phase as the concentration of PA 11 is increased.

Figure 2.7: FESEM images for 25/75 wt% PLA/PA 11 blends prepared at 205 ˚C for 20 minutes of melt-blending with the PLA phase removed through etching.

Figure 2.8: FESEM images for 50/50 wt% PLA/PA 11 blends prepared at 205 ˚C for 20 minutes of melt-blending with the PLA phase removed through etching.

Figure 2.9: FESEM images for 75/25 wt% PLA/PA 11 blend prepared at 205 ˚C for 20 minutes of melt-blending with the PLA phase removed through etching.

2.3.6 Dynamic Mechanical Thermal Analysis

DMTA is used for polymer characterization to study both the molecular relaxation processes and to determine inherent mechanical properties [128]. DMTA allows the measurement of the storage ( G’ ) and loss ( G”) moduli. The storage modulus as a function of temperature for the PLA/PA 11 blends is plotted in Figure 2.10. The blends show intermediate behavior between the two homopolymers. The glassy modulus increases monotonically with blend composition from pure PA 11 (0.58 GPa) to pure PLA (1.52 GPa). The glassy moduli for the blends are 0.75, 0.91, and 0.97 GPa for the 25/75, 50/50, and 75/25 compositions, respectively.

PLA/PA11 (wt%) 9 10 0/100 25/75 50/50 75/25 100/0

10 8 G' (Pa)

10 7 20 40 60 80 100 120 140 160 Temperature ( °C)

Figure 2.10: Storage modulus (G') versus temperature for various PLA/PA11 blends prepared at 205 ˚C with 20 minutes of melt blending time.

The loss tangent ( ) is plotted as a function of temperature for all =” ’ compositions of PLA/PA 11 blends in Figure 2.11. The α-relaxation phenomenon associated with the glass transition of the polymer is typically characterized by a sudden drop in the modulus and sharp increase in the value of tan δ. For pure PLA and PA 11,

the tan δ shows a single peak at temperatures of 72˚C and 61˚C, respectively; however, these values should not be confused with the actual glass transition temperature as measured by DSC which more closely corresponds to the peak in G” as a function of temperature. A shoulder is observed for PLA/PA 11 blends in the region of the glass transition for PA 11, indicating the presence of two glass-rubber relaxation peaks. The intensity of the shoulder appearing in the tan δ peak is generally an indication of the mass fraction of that phase in a polymer blend [129]. The lower temperature relaxation is offset relative to that observed for PA 11 and exhibits a decrease in magnitude with increasing PLA content. The high temperature relaxation is more pronounced and corresponds to the glass transition of PLA; this peak shifts to higher temperatures with increasing PLA composition. The two tan δ peaks of varying intensity are fully consistent with the observed two phase morphology.

0.30 PLA/PA11 (wt%) 0/100 0.25 25/75 50/50 0.20 75/25 100/0

δ 0.15

Tan 0.10

0.05

0.00 20 40 60 80 100 120 140 160 ° Temperature ( C) Figure 2.11: Tan δ vs. temperature plot for all PLA/PA 11 blends prepared at 205 ˚C for 20 minutes of melt blending.

DMTA can also be used to determine the heat distortion temperature (HDT) of a material. HDT is the temperature at which a sample bar deflects 0.25 mm during a

bending under load test described according to ASTM method D648. Takemori established a correlation between the HDT and the temperature at which the complex modulus equals 0.28 GPa [73]. This allows the determination of the HDT from DMTA data. Figure 2.12 shows the variation of HDT as a function of blend composition. HDT values for melt blended PLA and PA 11 are 75.5 ˚C and 58.5 ˚C, respectively. Values for the blends are 62.3 ˚C, 69.0 ˚C, 72.1 ˚C for the 25/75, 50/50, and 75/25 blends. That is, the HDT values increase monotonically from the value for PA 11 to the value for PLA as the blend composition is varied.

70 C) ° 60

HDT (

40 0/100 25/75 50/50 75/25 100/0 PLA/PA11 (wt%)

Figure 2.12: HDT values for different PLA/PA 11 blends prepared at 205 ˚C for 20 minutes of melt blending.

2.3.7 Tensile Test

Mechanical properties (tensile strength and modulus, stress and elongation at break and toughness) are plotted in Figure 2.13 and Figure 2.14 as a function of blend composition. Tensile stress-strain curves displayed a transition from brittle fracture of neat PLA to ductile fracture for pure PA 11. PLA/PA 11 blends exhibit a lower tensile

70 Stress at break 65 Young's modulus 4

60 (GPa) modulus Young's

55 3

45 2

40 Stress at break (MPa) 35 1

30 0 25 50 75 100 PLA (wt%)

Figure 2.13: Stress at break and tensile modulus of PLA/PA 11 blends as a function of blend composition.

60 18 Toughness Strain at break 16 50 14 ) 3 40 12 J/m 5 10 10

× 30 8

6 20 4 Strainbreak at (%)

Toughness ( Toughness 10 2

0 0 0 25 50 75 100 PLA (wt%)

Figure 2.14: Toughness and strain at break of PLA/PA 11 blends as a function of blend composition.

strength compared to the homopolymers which indicates poor interfacial adhesion between two phases. Tensile modulus is calculated from the initial stress-strain slope near zero strain. A stiff material is characterized by high modulus; the melt treated PLA retains a relatively high tensile modulus of 3.8 GPa. In comparison, the modulus of PA 11 is 1.4 GPa. Values for tensile modulus of the blends are 2.2, 2.9 and 3.0 GPa for the 25/75, 50/50, and 75/25 compositions, respectively. That is, the moduli values are in between those of the homopolymers and decrease with the increasing PA 11 composition. The toughness value, calculated as the area under stress-strain curve, decreased for 75/25 wt% PLA/PA 11 blend as compared to that of pure PLA. But, it increased significantly for 50/50 and 25/75 wt% PLA/PA 11 blends. Toughness values are 4056, 727, 308 and 666 KJ/m 3 for 25/75, 50/50, 75/25 and 100/0 wt% PLA/PA 11 blends, respectively. Both toughness and breaking elongation increase with the increase in PA 11 composition but are accompanied by a decrease in tensile modulus. Values of the elongation to break are 11.6, 2.4, 1.4 and 1.9 % for 25/75, 50/50, 75/25 and 100/0 wt% PLA/PA 11 blend compositions, respectively. Values of toughness and elongation to break were highest for 25/75 wt% PLA/PA 11 blend, but it had lower tensile strength as compared to other blends. Toughness and elongation to break could not be calculated for pure PA 11 due to extensive necking which produced deformation that exceeded the travel of the extensometer.

2.3.8 Impact Test

Figure 2.15 shows the notched Izod impact strength values for PLA/PA 11 blends as a function of blend composition. The measured impact strength of the neat PLA is 42 J/m, and that of neat PA 11 is 248 J/m. Values for the PLA/PA 11 blends are 50, 31 and 22 J/m for the 25/75, 50/50, and 75/25 compositions, respectively. These values are lower than that of individual homopolymers, but increase gradually with increasing PA 11 content. Impact strength of immiscible blends is strongly determined by the blend morphology and degree and efficiency of interfacial compatibilization. For the blends studied, the morphology of PLA/PA 11 blends appears too coarse to impart the

toughening by slippage and localized plastic deformation. The lack of effective compatibilization results in decreased impact strength.

350

300

250

200

150

100

Notched Izod Impact Strength (J/m) 0 0 25 50 75 100

PLA wt% composition Figure 2.15: Notched Izod impact strength of PLA/PA 11 blends as a function of blend composition

2.4 Conclusion

The first detailed study of blends of PLA and PA 11, two bioplastics based on 100% renewable resources is presented. Reactive compatibilization between PLA and PA 11 was attempted to examine the possibility of promoting interchange reactions. The blends were investigated with respect to phase morphology, thermal, and mechanical properties over the entire composition range. NMR spectroscopy is used to follow the extent of interchange reaction between polymers but shows peaks corresponding only to the homopolymers. No new peak corresponding to the formation of copolymers is observable. Accordingly, no direct evidence of interchange reactions is detectable even when using relatively high amounts of catalyst, reaction times, and increased temperatures. There is however, some indirect evidence that PLA is retained along with

PA 11 when repeatedly washed with chloroform. Difficulties caused by degradation of PLA at the higher processing temperatures and longer reaction times preclude observations of transreactions previously reported between polyesters and polyamides. The DSC thermograms show two distinct glass transition and melting temperatures indicating the immiscible nature of the blends. DMTA spectra show distinct double peaks ascribed to the glass transition of each component, irrespective of catalyst level, mixing time and temperature. FESEM clearly shows phase-separated structure of PLA/PA 11 blends. Mechanical properties of the blends are found to be in between the values for those of pure polymers. Values of shear and elongation modulus as well as tensile strength show close to monotonic behavior with blend composition. The HDT is also monotonic in composition. Only toughness and impact strength are relatively insensitive to composition. These later trends are attributed to a poorly compatibilized two-phase morphology. The intermediate properties of the PLA/PA 11 blends indicate that the system is promising for applications in which such properties are acceptable and cost reductions relative to PA 11 are desirable. These cost reductions could be achieved while retaining 100% renewable content.

CHAPTER 3 3 SYNTHESIS OF POLY(11-AMINOUNDECANOIC ACID-BLOCK-L-LACTIC ACID) FOR USE AS COMPATIBILIZING AGENT IN POLYAMIDE-11 / POLY(LACTIC ACID) BLENDS

Polymer blending has been shown to be an established and cost-effective way to develop polymeric materials with a diversified and desirable property portfolio [5-10]. Polymers typically form immiscible blends because of the high molar mass of polymer chains. This limits the entropic contribution to the free energy of a mixture and requires a high extent of favorable polymer-polymer interactions. Immiscibility in polymer blends tends to lead to inferior mechanical properties. Compatibilizing agents or copolymers can be used to increase interfacial adhesion between phases, decreasing the unfavorable mixing of certain properties. Compatibilizing agents can include but are not limited to: reactive compatibilizing agents, in-situ generated copolymer, or synthetically created block or graft copolymers [7]. The objective of the compatibilization process is to reduce the interfacial tension and facilitate dispersion, to stabilize morphology against high stress and strain processing, and enhance adhesion between phases in the solid state [7, 130]. The effect of block copolymers on immiscible binary blends have been extensively researched [7, 130-135]. The blends of polyamide-11 (PA 11) and poly(lactic acid) (PLA) are interesting as a 100% renewable alternative to petroleum derived polymers, specifically, in use for green carpeting. It has been shown in Chapter 2 and by Stoclet [134] that PLA and PA 11 blends are not miscible and have very weak interfacial properties between the separated homopolymer phases. In -situ reactive compatibilization through interchange reactions was also shown to be ineffective in Chapter 2. Therefore, the synthesis of block copolymers of PA 11 and PLA through melt polymerization of 11-aminoundecanoic acid in the presence of low molecular weight PLA was investigated as a potential compatibilization agent in the blends. Poly(lactic acid) is a thermoplastic polyester derived from lactic acid, a fermentation product. PLA is biodegradable through natural pathways of hydrolysis and

enzymatic activity. It is also price competitive with petroleum based resins and has physical properties comparable to Poly(ethylene terephthalate). However, the application profile is limited due to the low heat distortion temperature and moderate gas barrier [136]. Polyamide-11 (Nylon 11) is a well developed, high performance, semi-crystalline bioplastic derived from castor oil [57]. PA 11 is relatively expensive compared to petroleum based Nylons, however it has high impact and abrasion resistance, excellent chemical resistance, and high thermal stability [63]. Copolymers of PLA and PA 11 can be formed by the condensation reaction of 11- aminoundecanoic acid in the presence PLA. It is expected that lower molecular weight PLA will be needed to increase the number of reactive chain ends and increasing the speed of diffusion, both during the reaction and to the interface of PLA/PA 11 blends. Therefore, lower temperatures will be needed for copolymer formation in comparison to the in-situ interchange reactions. To form low molecular weight PLA, a detailed study of PLA degradation through hydrolysis of commercial PLA was investigated. Copolymerization in the melt was studied to determine if copolymer in measureable amounts can be synthesized.

3.1 Materials

Commercial grade PLA (2002, <4% D-lactide) was obtained from Cargill Dow Polymer (now Natureworks LLC, Minnetonka, MN). Irganox 1076 and Irgafos 168 were used as thermal stabilizers in toluene. Chloroform, toluene, m-cresol, 11- aminoundecanoic acid were reagent grade and obtained from Sigma-Aldrich.

3.2 Methods

Commercial PLA was ground in a Foremost A2 grinder to a pellet size less than 2 mm in diameter. PLA and DI water were prepared at equal mass in Erlenmeyer flasks. No acid catalyst was used in this study. Flasks were sealed with a rubber stopper and parafilm. Hydrolysis was completed with 21 separate flasks on a shaker table at 50 ˚C.

A sample of equal water and PLA was taken every day for the first 15 days and every third day subsequently up to day 38. To prevent further hydrolysis samples were dried at 65 ˚C for 24 hours under vacuum and stored with the sample solution at 10 ˚C in sealed glass containers. Molecular weight was determined by dilute solution viscometry using the Schulz- Blaschke and Mark-Houwink parameters found by Dorgan et al. [136]. Dilute solution viscometry was performed using the following procedure. Samples were created at concentrations of 0.002g/ml. The concentration was increased to maintain dilute solution properties as the molecular weight decreased. The samples were filtered to a clean vial through a 0.45 µm Teflon filter. Solutions were transferred to the viscometer tube held at 40 ˚C and pulled through the capillary into a reservoir glass ball. The time for the solution level to drain from a line on the top of the reservoir glass ball to a line on the bottom was recorded for every solution a minimum of 3 times. The Schultz-Blaschke relationship, as seen in Equation 3.1, is used to calculate the intrinsic viscosity of the solutions.

(3.1) = 1 + + In Eqaution 1, is the specific viscosity, is the solution concentration in g/ml, and is the Schultz-Blaschke coefficient. The specific viscosity is defined by the measured time of the solution over the measured time of pure solvent minus one. The Mark-Houwink equation was then utilized to determine the viscosity average molecular weights as follows:

(3.2) = where is the viscosity averaged molecular weight and and are the Mark-Houwink parameters specified for PLA in chloroform. Refractive indexes were measured using a refractometer held at 30 ˚C for all sample solutions (lactic acid in water). The calibration curve for the lactic acid in water systems was carried out at compositions of 0 wt% lactic acid to 85 wt% lactic acid as seen in Figure 3.1. The resulting baseline was determined to be:

(3.3) where RI is the refractive index and= 0 LA.0011 is the wt% + 1of.3331 lactic acid in water. This was used to find the lactic acid composition in the samples removed from hydrolysis.

1.45

1.40

1.35 RI=0.0011*LA+1.3331 Refractive Index Refractive

1.30 0 20 40 60 80 100 Lactic Acid Composition (wt%)

Figure 3.1: Refractive index calibration for lactic acid in water.

Polymerization of 11-aminoundecanoic acid with low molecular weight PLA prepared by hydrolysis were conducted in a Haake Rheomix 6000 at 185 ˚C and 10 RPM for 10 minutes. PLA was prepared by hydrolysis for 38 days to a viscosity average molecular weight of 25,000 g/mol. Hydrolyzed PLA and 11-aminoundecanoic acid were dried under vacuum at 80 ˚C for 20 hours. 11-Aminoundecanoic acid was added in equal mass amounts to hydrolyzed PLA to fill 70 % volume in the Haake bowl. After full melt occurred, Irgonox and Irgafos stabilizing solution in toluene was added to the mixture. The Haake bowl top remained open to allow water vapor, produced by the condensation reaction, to escape. Extractions were performed on copolymer synthesized in chloroform. Based on the solubility of commercially available PLA and PA 11 in chloroform, the following cases can be observed: 1) PLA homopolymer is dissolved in solvent, 2) Copolymer with

main component consisting of PLA is dissolved in solvent, 3) Copolymer with main component of PA 11 remains in insoluble fraction, 4) PA 11 homopolymer remains in insoluble fraction. Differential scanning calorimetery (DSC) was used to analyze thermal properties of hydrolyzed PLA, 11-aminoundecanoic acid, and copolymer. A Perkin-Elmer DSC-7 under nitrogen atmosphere and calibrated with an indium standard used the following protocol: 1) heat from 5 ˚C to 200 ˚C at 10 ˚C/min., 2) hold for 5 minutes at 200 ˚C, 3) cool from 200 ˚C to 5 ˚C at 10 ˚C/min, 4) heat from 5 ˚C to 200 ˚C at 10 ˚C/min. Melting points and glass transition temperatures were calculated on the first heat so as to measure the samples as is, without further reaction. The extracted solution and insoluble fraction were analyzed through Fourier Transform-Infrared Spectroscopy (FT-IR). KBr pellets were prepared and a solution containing dilute sample was solution cast onto the pellets. Pellets were then dried under vacuum (23 inch Hg) for more than an hour and stored in the presence of desiccant. Spectra were recorded on a Nexus 670 for hydrolyzed PLA, melt blend 11- aminoundecanoic acid product, copolymerization product, and soluble and insoluble fractions from extraction in chloroform.

3.3 Results

In this study, the hydrolysis of PLA in water was studied using dilute solution viscometry and used to create a block copolymer of PLA-PA 11. The resultant product was analyzed using differential scanning calorimetry (DSC) and Fourier transform- infrared spectroscopy (FT-IR).

3.3.1 Hydrolysis

PLA chains are degraded through an acid catalyzed mechanism. This condition results in an auto-catalyzing reaction from the increase in Carboxylic Acid functional

groups as seen in Figure 3.2. Pitt et al. derived the equation describing associated kinetics of auto-catalyzed hydrolytic degradation as shown in Equation 3.4.

Figure 3.2: Proposed hydrolysis reaction in PLA

(3.4) = − = − where [E] represents ester molar concentration in the polymer matrix, [COOH] represents carboxyl end group molar concentration, and [H 2O] represents the molar concentration of water in the polymer matrix [137]. The solution to the differential equation can be represented in terms of degree of polymerization or molecular weight as shown in Equation 3.5.

∗ (3.5) In Equation 3.5, is the current molecular = weight,, is the starting molecular , weight of the PLA, t is time in days, and k * is the kinetic rate constant. Figure 3.3 shows the results of molecular weight as a function of hydrolysis time from experimental results. After fitting the data to the relationship above, it is found that = 174±4 , kg/mol (corresponding to the average measured molecular weight) and ∗ = 0.051±0.003 1/day. From these results a desired molecular weight can be achieved through the hydrolysis of commercial PLA, an effective means to obtain low molecular weight PLA. The initial goal of PLA molecular weight is 25,000 g/mol for copolymer synthesis. Refractive index was used to show the hydrolysis of ester bonds occurs randomly in PLA. It was found that from the 80 g of PLA start weight; there was 0.12 g of lactic

200 Equation y = a*exp(b*x) Adj. R-Square 0.97465 Value Standard Error Molecular Weight a 174.73034 3.59114 Molecular Weight b -0.05189 0.00231

150

100

50 Molecular Weight (kg/mol) Weight Molecular

0 10 20 30 40 50 60 70 Day (days)

Figure 3.3: Molecular mass (weight averaged) of PLA as a function of hydrolysis time.

acid formed from hydrolysis for 36 days. This is less than a quarter of a percent of the original PLA mass. Figure 3.4 shows the weight of lactic acid in solution as a function of time. DSC results for hydrolyzed PLA provide information on the thermal properties of

PLA with molecular weight of 25,000 g/mol. Glass transition temperature (T g) was observed to decrease with molecular weight as expected from 57.2 to 46.5 for the commercial PLA and the hydrolyzed PLA, respectively. As can be seen in Figure 3.5, the area under the melt peak increased substantially, indicating that the crystallinity increase in the hydrolyzed polymer. The melting temperature (T m) as described by the end of the melting peak for hydrolyzed PLA was 160.7 ˚C as opposed to 168.3 ˚C for

commercial PLA. This decrease in melting temperature is expected for the 25,000 M w because the chain ends act as impurities in polymer matrix, decreasing the stability of lamellar formation.

0.14

0.12

0.10

0.08

0.06

0.04

0.02

Total Wieght of Lactic Acid in Solution (g) Solution in Acid Lactic of Wieght Total 0.00 0 5 10 15 20 25 30 35 40 Day (day)

Figure 3.4: Weight of Lactic Acid in solution as a function of hydrolysis time .

Figure 3.5: Typica l DSC endotherms for commercial PLA and hydrolyzed PLA .

3.3.2 Synthesis

11-Aminoundecanoic acid undergoes a condensation polymerization to a small degree of polymerization when melt of the monomer is achieved. This can be seen in Figure 3.6 as the second run of a DSC temperature sweep of 11-aminoundecanoic acid after holding at 200 ˚C for only 5 minutes changes drastically from the first run. 11- Aminoundecanoic acid was polymerized in the Haake to show polymerization occurs at 195 ˚C. Samples were removed every 10 minutes with a total time of reaction at 30 minutes. A representative DSC for each sample is shown for the samples removed in Figure 3.7. It appears that after 10 minutes the sample remains predominantly monomer, but as the reaction progressed, the melting peak became that of low molecular weight PA 11 confirming the presence of condensation reaction in the Haake.

Low molecular weight PLA was added with 11-aminoundecanoic acid in equal mass proportions to the Haake to form copolymer. The proposed reaction of the first

2st Run 30 1st Run

Endo (mW) Therm 10

0 40 60 80 100 120 140 160 180 200 Temperature ( oC)

Figure 3.6: DSC experiment showing the first (11-aminoundecanoic acid) and the second run (low molecular weight PA11) of sample consisting of pure 11-aminoundecanoic acid.

60 Haake Melt 30 min. Haake Melt 20 min. 50 Haake Melt 10 min. 11-Aminoundecanoic acid

20 Endo Therm (mW) Therm Endo

0 40 60 80 100 120 140 160 180 200 Temperature ( oC)

Figure 3.7: DSC experiment showing the first run for 11-aminoundecanoic acid and samples taken every 10 minutes from the Haake reaction.

addition of 11-aminoundecanoic acid monomer to the PLA is shown in Figure 3.8. The reaction was run for 10 minutes and 20 minutes after the addition of stabilizing solution. Qualitatively, the polymerization product was darker brown, less viscous, and solidified slower than the PA 11 homopolymer synthesis. Extraction was performed in chloroform after recording mass of copolymer sample. After extraction, 22% of sample was soluble and 45% insoluble. The soluble fraction was analyzed as the insoluble portion may have PLA remaining from incomplete extraction. PA 11 homopolymer is not miscible in chloroform, resulting in the presence of PA 11 characteristic peaks being present in the soluble extract only if there is a larger PLA chain attached. Solubility was tested for the reactants, polymer species, and product in chloroform and m-cresol with the findings in Table 3.1. FT-IR and DSC are used to characterize the product of reaction and both fractions of the extraction.

Figure 3.8: Proposed reaction of PLA and 11-aminoundecanoic acid to form copolymer (first addition of monomer 11-aminoundecanoic acid shown)

Table 3.1: Results of Solubility Tests of reactants, polymers, and product in m-cresol and chloroform. PA11 11-Aminoundecanoic PLA Solvent homopolymer Acid homopolymer Product Partially Chloroform Not soluble Not soluble soluble soluble m-cresol soluble Soluble soluble Soluble

4000 3000 2000 1000 0 0.81 1.5 extraction soluble extraction soluble PLA 0.54 PLA PA11 PA11

0.27 Absorbance 0.00 1.0 2.34

1.56

0.78 0.5 Absorbance 0.00 Absorbance

1.98

1.32

0.0 0.66 Absorbance 0.00 3600 3500 3400 3300 3200 3100 4000 3000 2000 1000 0 Wavenumber (cm -1 ) -1 Wavenumber (cm ) b a 2.0 2.5 extraction soluble extraction soluble PLA PLA PA11 PA11 2.0 1.5

1.5 1.0

1.0 Absorbance Absorbance 0.5 0.5

0.0 0.0

3000 2900 2800 1650 1600 1550 1500 Wavenumber (cm -1 ) Wavenumber (cm -1 ) c d Figure 3.9: FT-IR spectra of PA 11 in red, soluble extract of copolymerization in green , adn hydrolyzed PLA in blue for a) full spectra, b) N-H fingerprint, c) C-H fingerprint and d) amide fingerprint.

3.4 Conclusion

Hydrolysis of commercial PLA was carried out and the kinetic parameters of the auto catalyzed reaction were reported. After derivation of the kinetic equation -1 ∗ it was found that =174 kg/mol and ∗=0.05 day . A reaction of = , , 11-aminoundecanoic acid and low molecular weight PLA was confirmed after extraction through the use of FT-IR spectroscopy and DSC to result in the creation of the poly(11- aminoundecanoic acid-block-L-lactic acid).

FUTURE WORK

The first chapter of this document has shown the miscible nature of the polyamide-11/polyamide-6,10 blends. Future work should concentrate on the analysis of the crystal-induced phase separation. This can be done through a combined differential scanning calorimetry and x-ray scattering. The phase diagram of the blend should also be studied to find the Eutectic points in DSC. The analysis should include more blend compositions and more annealing temperatures then presented in the previous discussion. It would also be interesting to test mechanical properties of the 50/50 wt% blend as a function of time after annealing to see if amorphous region can densify and improve mechanical properties while retaining crystalline compositions. The second and third chapter show that the poly(lactic acid)/polyamide-11 blends are not miscible and that a block copolymer of poly(lactic acid)-co-polyamide-11 can be synthesized. Further work on this blend should attempt other reactions to produce higher molecular weight copolymer as a compatibilizing agent. This can include performing the low molecular weight synthesis of 11-aminoundecanoic acid in the Haake and then adding commercial PLA (un-hydrolyzed), more thourough removal of water from the reaction, or a synthesis in solvent. After a copolymer of sufficient molecular weight is produced, blends of poly(lactic acid)/polyamide-11/copolymer should be prepared at several different compositions to reexamine the mechanical properties of the compatibilized blends.

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