die were cooled in a water trough and palletized using Berlyn pelletizer (Worcester, MA,. USA). ...... T. H. Solomon, S. Tomas, J. L. Warner, Phys. Fluids, 10 ..... 238. 239. D. J. Shaw, Introduction to Colloid and Surface Chemistry lh Ed., Reed.
Melt Processing of Thermoplastic/Clay Nanocomposites
by Nitin Borse
A thesis submitted to the faculty of Graduate Studies and Research in partial fulfillment of the requirement of the degree of Doctor of Philosophy
Department of Chemical Engineering McGill University Montreal, Quebec, Canada
@ August 2005
1+1
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Dedicated to the memory of my father who is no more to see what can be achieved with his encouragement.
ABSTRACT Polymer/clay nanocomposites are materials composed of a polymer matrix and nanometer size clay particles. They exhibit significant improvements in tensile modulus and strength, reduced permeability to gases and liquids, compared to the pure polymer. These property improvements can be realized while retaining clarity of the polymer without a significant increase in density (typical clay loading is 2-5%). The aim of this research was to achieve maximum exfoliation of the clay and property enhancement in the nanocomposite by melt processing. The components of the Hansen solubility parameters and the Hamaker constants of the constituents were used to establish the compatibility between the polymer matrix and nanoclay. A mathematical model was formulated to study the mechanism of exfoliation of clay platelets in the polymer matrix. The mechanism of size reduction of clay particles was shown to be by erosion or surface peeling. A novel continuous process (System A) was developed and implemented to produce nanocomposites. The process incorporates chaotic mixing, efficient shearing and stretching of the melt, frequent changes in flow direction and higher residence time, in order to enhance exfoliation. The efficiency of the process was demonstrated, using transmission electron microscopy, wide angle X-ray diffraction, viscosity measurements, tensile and highspeed
impact
testing,
and
oxygen
permeability
measurements.
Polyamide
nanocomposites prepared with System A show a higher degree of exfoliation and better enhancements in mechanical and barrier properties than those prepared using conventional twin-screw extrusion (System C). Polystyrene nanocomposites show higher increase in the d-spacing of organoclay and improvement in barrier properties for System A than for System C, indicating higher specific surface area for the filler particles. The Pukanszky parameter was negative for PS nanocomposites, indicating poor polymer/clay adhesion. The crystallization kinetics of PA-6 nanocomposites were studied using high pressure dilatometry. The nanoclay seems to act as nucleating agent, increasing the rate of crystallization of PA-6 in the nanocomposite. The effect of the clayon the kinetics of formation of different PA-6 crystalline phases was evaluated and explained.
RÉSUMÉ Le nanocomposite polymère-argile est un matériau constitué d'une matrice de
polymère et de nanoparticules d'argile. Il présente d'importantes améliorations sur le polymère pur en termes de module d'élasticité et de résistance à la traction, et une perméabilité réduite aux gaz et aux liquides. Ces améliorations peuvent être réalisées tout en conservant la clarté du polymère, sans augmentation significative de densité (la charge d'argile typique est de 2 à 5 %). Le but de cette recherche était d'obtenir par fusion le degré maximum d'exfoliation de l'argile et des propriétés améliorées du nanocomposite. Avec les paramètres de solubilité de Hansen et les constantes de Hamaker des composants, on a vérifié la compatibilité entre la matrice de polymère et le nanoargile. Un modèle mathématique, conçu pour étudier le mécanisme d'exfoliation des plaquettes d'argile dans la matrice de polymère, a fait voir que le mécanisme de réduction des particules d'argile est l'érosion, ou dé stratification de la surface. Un nouveau procédé continu a été mis au point, appelé système A, pour produire les nanocomposites. TI consiste à optimiser l'exfoliation avec un mélange aléatoire, une gestion efficace du cisaillement et de la torsion de la coulée, et une alternance fréquente du sens d'écoulement et un temps de rétention accru. L'efficacité de ce procédé a été démontrée par microscopie électronique à transmission, diffraction de rayons-x aux grands angles, mesures de viscosité et de perméabilité à l'oxygène, essais de traction et d'impact à haute vitesse. Par rapport au procédé courant d'extrusion à deux vis (système C), les nanocomposites de polyamide préparés avec le système A affichent une plus grande exfoliation et de meilleures propriétés mécaniques et barrières. Les nanocomposites de polystyrène préparées avec le système A plutôt qu'avec le système C ont un d-spacing accru de l'argile organique et de meilleures propriétés barrières, indiquant une plus grande surface de contact spécifique pour les particules de la charge. Le paramètre de Puzansky, négatif pour les nanocomposites de polystyrène, a démontré la faible adhésion polymère-argile. La cinétique de cristallisation des nanocomposites du PA6 a été étudiée avec la dilatométrie à haute pression. TI semble que le nanoargile agisse comme agent de
nuc1éation pour accélérer la cristallisation du PA6 dans le nanocomposite. On a évalué et expliqué l'effet de l'argile dans la cinétique de formation des différentes étapes de cristallisation du PA6.
ACKNOWLDGEMENTS
1 would lik:e to express my gratitude and sincere thanks to my research supervisor, Professor M. R. Kamal for for bis guidance and support throughout the course of this research. 1 am grateful to Dr. Yong Cho for the valuable discussions and suggestions and to Mr. Vincent Mollet for the efforts in characterizing TEM micrographs. 1 would like to thank Mr. Slawomir Poplawski for X-ray diffractograms of my samples. Mr. Alain Gagnon and Mr. Csaba Szalacsi from chemical engineering workshop fabricated the designed parts to my satisfaction. 1 express my gratitude towards them. 1 am thankful to Mr. L. Cusmich who helped me setting up instrumentations during this research project. 1 am thankful my colleagues in polymer research group for their continuaI support. 1 would like to express my appreciation to the National Sciences and Engineering Research Council of Canada and McGill University for financial support. Nova Chemicals and DuPont Canada supplied the polymers used in this research. It would have never been possible for me to engage into this research without support from my parents, wife and daughter. 1 express my deep gratitude towards them.
TABLE OF CONTENTS ABSTRACT ACKNOLEDGEMENTS TABLE OF CONTENTS .....•............................•........•................•.........................•...................... 1 LIST OF FIGURES ..................................................................................................................... V LIST OF TABLES .............. ~ ..................................................•.............•.......•.......•.•.•.••.•.•.•••.•.•. XV NOMENCLATURE .............................................................................................................. XVII DEFINITIONS OF SOME TERMS IN ................................................................................ XXI
POLYMERIC NANOCOMPOSITE LITERATURE ......................................................... XXI 1 INTRODUCTION ................................................................................................................... 1 2 LITERATURE REVIEW ....................................................................................................... 4 2.1
LAYERED SILICATE STRUCTURE .............................................................................. 5
2.1.1
THEHAMAKERAPPROXIMATION ............................................................................. 7
2.1.2
THE LIFSHITZ APPROACH ............................................................................................ 8
2.1.3
POLAR FORCES ............................................................................................................... 8
2.2
NANOCOMPOSITE FORMATION AND STRUCTURE .............................................. 9
2.3
THERMODYMAMICS OF NANOCOMPOSITE FORMATION .............................. 12
2.4
KINETICS OF POLYMERMELT INTERCALATION .............................................. 18
2.5
HYDRODYNAMICS OF DISPERSION AND DISTRIBUTION OF CLAY IN
POLYMER MATRIX ................................................................................................................ 21 2.5.1
MODELS FOR DISPERSION PROCESS ....................................................................... 22
2.5.2
DISTRIBUTIVE MIXING ............................................................................................... 24
2.5.3
MELT PROCESSING NANOCOMPOSITES ................................................................. 28
II 2.6
PHYSICAL PROPERTIES AND CHARACTERIZATION OF NANOCOMPOSITES ............................................................................................................................................. 32
2.6.1
MECHANICAL PROPERTIES ....................................................................................... 32
2.6.2
BARRIER PROPERTIES ................................................................................................ 34
2.6.3
RHEOLOGY .................................................................................................................... 36
2.6.4
CRYSTALLIZATION ..................................................................................................... 37
2.6.5
CHARACTERIZATION METHODS ............................................................................. 39
3 OBJECTIVES ........................................................................................................................ 41 4 MODELING AND PROCESS DESIGN ............................................................................. 43 4.1
INTERACTION FORCES BETWEEN CLAY PLATELETS ...................................... 43
4.1.1
ADHESION BETWEEN PLATELETS IN A CLAY PARTICLE .................................. .43
4.1.2
BREAKING CLAYPARTICLES INTO TACTOIDS AND PLATELETS ................... .49
4.2
MODEL APPLICATION AND RESULTS ..................................................................... 57
4.3
PROCESS DESIGN .......................................................................................................... 63
4.3.1
5
SELECTION OF THE S fATIC MIXER ......................................................................... 65
EXPERIMENTAL ............................................................................................................... 76
5.1
PROCESSING EQUIPMENT ......................................................................................... 76
5.1.1
TWIN-SCREWEXTRUDER .......................................................................................... 76
5.1.2
STATICMIXER .............................................................................................................. 78
5.1.3
DIE UNIT ......................................................................................................................... 83
5.1.4
COMPRESSION MOLDING .......................................................................................... 83
5.1.5
INJECTION MOLDING .................................................................................................. 83
5.2 5.2.1
CHARACTERIZATION TECHNIQUES ...................................................................... 84 X-RAY DIFFRACTION .................................................................................................. 84
III 5.2.2
TRANSMISSION ELECTRON MICROSCOPY ............................................................ 84
5.2.3
WATERABSORPTION .................................................................................................. 85
5.2.4
THERMOGRAVIMETRY .............................................................................................. 86
5.2.5
MECHANICAL PROPERTIES ....................................................................................... 86
5.2.6
BARRIERPROPERTIES ................................................................................................ 90
5.2.7
RHEOLOGY ..................................................................................................................... 92
5.2.8
INFRARED SPECTROSCOPY ....................................................................................... 93
5.2.9
HIGH PRESSURE DILATOMETRY .............................................................................. 93
5.3
MATERIALS ..................................................................................................................... 95
5.4
EXPERIMENTAL PROCED'URES ................................................................................ 98
5.4.1
RESIDENCE TIME DISTRIDUTION ............................................................................. 98
5.4.2
EXTRUSION ................................................................................................................... 99
6
RESULTS AND DISCUSIONS ......................................................................................... 100
6.1
RESIDENCE TIME DISTRIBUTION ..................•..•.••.•..............•.•.•.•.•...•.....•......•.•.•... 100
6.2
X-RAY DIFFRACTION ........ , '?
~
••••••••••••••••••••••••••••••••••••••••••••••••••• •••••••••••••••••••••••••••••••••
105
6.2.1
PA-6 NANOCCMPOSITES .......................................................................................... 105
6.2.2
POLYSTYRENE NANOCOMPOSITES ...................................................................... 11 0
6.2.3
FACTORS INFLUENCING EXFOLIATION ............................................................... 119
6.2.4
FACTORS INFLUENCING INTERCALATION ......................................................... 122
6.3
TRANSMISSION ELECTRON MICROSCOPY ........................................................ 130
6.3.1
TEM MICROGARPHS OF PA-6 NANOCOMPOSITES .............................................. 130
6.3.2
TEM MICROGRAPHS OF POLYSTYRENE NANOCOMPOSITES .......................... 139
6.4
WATERABSORPTION................................................................................................. 143
6.4.1
PA-6NANOCOMPOSITES .......................................................................................... 143
6.4.2
POLYSTYRENE NANOCOMPOSITES ...................................................................... 149
IV 6.5
THERMOGRAVIMETRIC ANAL YSIS ...................................................................... 150
6.6
MECBANICAL PROPERTIES .................................................................................... 153
6.6.1
MECHANICAL PROPERTIES OF POLYAMIDE-6 NANOCOMPOSITES .............. 153
6.6.2
MECHANICAL PROPERTIES OF POLYSTYRENE NANOCOMPOSITES ............. 178
6.6.3
MODELING TENSILE PROPERTIES OF POLYAMIDE AND POLYSTYRENE
NANOCOMPOSITES ................................................................................................................ 184
6.7
BARRIER PROPERTIES OF NANOCOMPOSITES •.•.•.•.............•.•.............•...•••...... 200
6.8
RHEOLOGY OF NANOCOMPOSITES .•.•.••.........•.•.•.•.•.•...............•.............•.•.•.•.•..... 203
6.8.1
RHEOLOGY OF PA-6 NANOCOMPOSITES .............................................................. 204
6.8.2
RHEOLOGY OF POLYSTYRENE NANOCOMPOSITES .......................................... 207
6.8.3
ESTIMATION OF SHEAR STRESS IN THE STATIC MIXER ................................... 209
6.9
CRYSTALLIZATION KINETICS OF PA-6 AND ITS NANOCOMPOSITES .•.•.•.• 213
7
CONCLUSIONS ................................................................................................................. 222
8
RECOMMENDATIONS FOR FUTURE WORK •......................•.................................. 225
9
ORIGINAL CONTRIBUTIONS TO KNOWLEDGE.................................................... 226
10
REFEREN CES ................................................................................................................. 22 7
APPENDICES
v LIST OF FIGURES Figure 2.1.1: Structure of2:1layered silicate (l,12) ......................................................... 6 Figure 2.1.2: Clay particles consisting of stacks of identical units (14,15) ....................... 6 Figure 2.2.1: Schematic of the basic steps in nanocomposite formation (19) ................. 10 Figure 2.2.2: Schematic of microstructures oflayered silicate nanocomposites (19) ..... 11 Figure 2.2.3: Schematic of types of interfacial interactions occurring in polymerorganoclay nanocomposites (30) ............................................................................. 11 Figure 2.3.1: Schematic representation of the system components before and after the intercalation (36) ...................................................................................................... 14 Figure 2.3.2: The change of entropy per area versus the change in gallery height, for the polymer and the surfactant functionalized surface (36) ........................................... 14 Figure 2.3.3: Sketch of the nonfunctionalized P-chains and the different conformations of the end-functionalized N-chains in the polymer melt confined between two clay sheets, equilibrated at a distance H (40) .................................................................. 17 Figure 2.3.4: Free energy per unit area, MIA, as a function of surface separation, H .... 18 Figure 2.5.1: A schematic illustration of the difference between dispersion and distribution ............................................................................................................... 22 Figure 2.5.2: Representation of Smale horseshoe function .............................................. 25 Figure 2.5.3: Elliptic and hyperbolic points ..................................................................... 26 Figure 2.5.4: Enhanced Mixing Simulator (100) ............................................................. 27 Figure 2.5.5: Schematic representation of chaotic mixing apparatus and loci of elliptie and hyperbolic points in chaotic mixing (101, 102, 103) ........................................ 27 Figure 2.5.6: Degree of dispersion (by Transmission Electron Microscopy) of various nanocomposites of PA-6/Cloisite15A plotted against mean residence time in sec 29
VI Figure 2.5.7: Mechanism of clay platelet exfoliation in the melt compounding of nanocomposites (124, 126) ...................................................................................... 29 Figure 2.6.1: Bragg's law for constructive interference ................................................. .40 Figure 4.1.1: Model Clay Particle .................................................................................... 44 Figure 4.1.2: Van der Waals forces between solids (13) ................. :.............................. .45 Figure 4.1.3: Attractive van der Waals interaction energy between platelets of unmodified and organically modified clay platelets versus gallery spacing .......... .48 Figure 4.1.4: Attractive van der Waals forces per unit area between platelets of unmodified and organically modified clay............................................................. .48 Figure 4.1.5: Schematic representation of exfoliation process ........................................ 51 Figure 4.1.6: Peeling Model ............................................................................................. 52 Figure 4.1.7: Lap Shearing model ................................................................. :.................. 54 Figure 4.1.8: Dislocation in lap joint ............................................................................... 54 Figure 4.2.1: Schematic representation of a clay particle consisting of layers of stacked platelets .................................................................................................................... 57 Figure 4.2.2: Mechanisms of breaking clay particle into smaller tactoids ....................... 58 Figure 4.2.3: Shear stress required to break 1000 nm thick clay particle into two halves . .................................................................................................................................. 59 Figure 4.2.4: Shear stresses required for peeling of tactoids of variable thickness from the surface of the clay particle with surface area 200x200 nm2 , 500x500 nm2 and 1000xl000 nm2 ........................................................................................................ 60 Figure 4.2.5: Shear stresses required for peeling clay tactoids of different thickness and area from the clay surface ........................................................................................ 61
VII Figure 4.2.6: Shear stress required for peeling of 1 nm thick platelet from the surface of a clay particle at various peeling angles ................................................................... 62 Figure 4.3.1: Types of static mixers (208) ....................................................................... 67 Figure 4.3.2: Capillary rheometry data for the viscosity of PA-6 at 240°C .................... 70 Figure 4.3.3: Shear stresses encountered during processing polymer melt with different static mixers ............................................................................................................. 70 Figure 4.3.4: Extensional efficiencies of static mixers (208) ........................................... 72 Figure 4.3.5: Stretching distribution for SMX and ISO static mixers (208) .................... 72 Figure 4.3.6: Schematic representation ofbaker's transformation in a static mixer and chaotic mixing (116) ................................................................................................ 74 Figure 4.3.7: Distributive mixing in the ISO static mixer ............................................... 74 Figure 4.3.8: Schematic representation of the proposed melt processing system for nanocomposite production ....................................................................................... 75 Figure 5.1.1: Configuration of screw elements ................................................................ 77 Figure 5.1.2: Nomenclature of the screw elements used ................................................. 77 Figure 5.1.3: Pressure drop for ISO mixer in laminar flow with viscosity 10000 cps ..... 79 Figure 5.1.4: ISO static mixer elements ........................................................................... 81 Figure 5.1.5: Static mixer housing barrel and accessories (all dimensions in mm) ......... 82 Figure 5.2.1: Typical tensile designations (ASTM D638) ............................................... 87 Figure 5.2.2: Typicalload versus displacement plot for flexural test.. ............................ 88 Figure 5.2.3: Typicalload-deflection plot for high speed impact testing using RIT-8000 .............................................,..................................................................................... 90 Figure 5.2.4: Schematic of oxygen permeability test apparatus ...................................... 92 Figure 6.1.1: Residence time distribution for System C ................................................ 10 1
vrn Figure 6.1.2: Residence time distribution for System A ................................................ 101 Figure 6.1.3: Comparison of System A and System C with the same mean residence times but different feed flow rates ......................................................................... 103 Figure 6.1.4: Comparison of System A and System C for same feed flow rate ............ 103 Figure 6.1.5: Cumulative residence time distribution curves for System A and System C under different feed flow conditions ...................................................................... 104 Figure 6.2.1: XRD plots ofPA-6/Cloisite 30B (2.1 % wt) for System A and System C. XRD plot of in-situ polymerized Ube PNC with 2 % wt Clay is shown for the comparison............................................................................................................. 106 Figure 6.2.2: XRD plots of PA-6/Cloisite 30B nanocomposite samples made using System A with different clay contents ................................................................... 107 Figure 6.2.3: XRD Plots of PA-6/Cloisite 30B nanocomposites made using System C with different clay contents ........................"........................................................... 107 Figure 6.2.4: XRD plots ofPA-6/Cloisite 15A (2.5% wt) nanocomposites made using System A and System C ......................................................................................... 109 Figure 6.2.5: XRD plots ofPA-6/Cloisite 15A (4.1 % wt) nanocomposites made using System A and System C ................................................................ ~.: ...................... 109 Figure 6.2.6: XRD plots ofPA-6/Cloisite Na+ (4.2% wt) nanocomposites made using System A and System C ......................................................................................... 11 0 Figure 6.2.7: XRD plots ofPS1301lCloisite lOA (3.2%wt) nanocomposites made using System A and System C at 210°C .......................................................................... 112 Figure 6.2.8: XRD plots of polystyrene nanocomposites with Dylark compatibilizer.. 112
IX
Figure 6.2.9: Effect of compatibilizer on the formation of PS nanocomposites. PS1301lCloisite lOA (3.2% wt) nanocomposites prepared with and without Dylark 332 compatibilizer at 210°C using System A. ....................................................... 114 Figure 6.2.10: Formation ofnanocomposite with Dylark 332 and Cloisite lOA (3.2%wt) . ................................................................................................................................ 114 Figure 6.2.11: Effect of processing temperature; PS1301lCloisite lOA (3.2% wt) Extruded using System A at 210°C and 240°C. ..................................................... 115 Figure 6.2.12: Effect of processing temperature; PS 1301 +2% wt Dylark/Cloisite lOA (3.2% wt) Extruded using System A at 210°C and 240°C. .................................... 115 Figure 6.2.13: PS1301lCloisite 15A (3% wt) nanocomposites preparedusing System A and System C at 210°C .......................................................................................... 117 Figure 6.2.14: Effect of compatibilizer; PS1301lCloisite 15A nanocomposites with and without Dylark 332 compatibilizer made using System A at 210°C ..................... 117 Figure 6.2.15: Using different organoclays with PS1301 matrix; Samples prepared using System A at 210°C ................................................................................................. 118 Figure 6.2.16: PS3900/Cloisite lOA (3.2% wt) nanocomposites with and without compatibilizer using System A and System C at 210°C ........................................ 118 Figure 6.2.17: Orientation of n-alkylammonium ions in the interlayers of layer silicates . ................................................................................................................................ 124 Figure 6.2.18: Schematic of two-dimensional spreading pressure versus area per molecule isotherm. At sufficiently high pressures TIc. the solid monolayer film collapses ................................................................................................................. 126 Figure 6.2.19: Plots of compressibility factor versus spreading pressure (239) ............ 126 Figure 6.2.20: Collapse of monolayer film under pressure (240) .................................. 127
X Figure 6.3.1: TEM micrographs of PA-6/Cloisite 30B (3.1 %wt) nanocomposites ....... 133 Figure 6.3.2: TEM micrographs of PA-6/Cloisite 30B (5.0%wt) nanocomposites ....... 134 Figure 6.3.3: TEM micrographs ofPA-6/Cloisite 15A (2.5%wt) nanocomposites ....... 135 Figure 6.3.4: TEM micrographs ofPA-6/Cloisite Na+ (4.2%wt) nanocomposites ....... 136 Figure 6.3.5: Distribution of clay particle size in PA-6/Cloisite 30B(5%wt) nanocomposite (244) .............................................................................................. 138 Figure 6.3.6: TEM micrographs ofPS1301lCloisite lOA(2%wt) at low magnification140 Figure 6.3.7: TEM micrographs ofPS1301lCloisite lOA(2%wt) with Dylark compatibilizer(2%wt) at low magnification........................................................... 140 Figure 6.3.8: TEM micrographs ofPS1301lCloisite lOA(2%wt) nanocomposites ....... 141 Figure 6.3.9: TEM micrographs ofPS1301lCloisite lOA(2%wt) with Dylark compatibilizer produced with System A at different magnifications ..................... 142 Figure 6.4.1: Water absorption of Slit Die extruded PA-6 nanocomposites after 1 day145 Figure 6.4.2: Water absorption of compression molded PA-6 nanocomposites (1 day)145 Figure 6.4.3: Water absorption of injection molded PA-6/Cloisite 30B (5%wt) nanocomposite prepared using System A .............................................................. 146 Figure 6.4.4: Water diffusion in slit die extruded PA-6 and PA-6/Cloisite 30B nanocomposites prepared using System A ............................................................. 147 Figure 6.4.5: Water absorption of PS1301/Cloisite lOA nanocomposites after 4 days .149 Figure 6.5.1: Thermogravimetric analysis of extruded PA-6 and PA-6/Cloisite 30B (5%wt) nanocomposite ........................................................................................... 151 Figure 6.5.2: Thermogravimetric analysis ofPA-6 and PA-6/Cloisite 30B nanocomposite; injection molded and extruded ..................................................... 151
XI Figure 6.5.3: Thermogravimetric analysis ofPS1301 and PS130l/Cloisite lOA (4.5%wt) nanocomposites ...................................................................................................... 152 Figure 6.6.1: Tensile Modulus ofPA-6/Cloisite 30B nanocomposites ......................... 156 Figure 6.6.2: Tensile strength of PA-6/Cloisite 30B nanocomposites ........................... 156 Figure 6.6.3: Tensile modulus ofPA-6/Cloisite 15A nanocomposites .......................... 157 Figure 6.6.4: Tensile strength ofPA-6/Cloisite 15A nanocomposites .......................... 157 Figure 6.6.5: Tensile modulus ofPA-6/Cloisite Na+ nanocomposites .......................... 158 Figure 6.6.6: Tensile strength ofPA-6/Cloisite Na+ nanocomposites ........................... 158 Figure 6.6.7: Tensile moduli of PA-6 nanocomposites produced with different clays using System A ...................................................................................................... 160 Figure 6.6.8: Tensile moduli of PA-6 nanocomposites produced with different clays using System C ....................................................................................................... 160 Figure 6.6.9: Tensile strength of the PA-6 nanocomposites with different clays using System A ................................................................................................................ 161 Figure 6.6.10: Tensile strength of the PA-6 nanocomposites with different clays using Systenl C ................................................................................................................ 161 Figure 6.6.11: Tensile moduli of extruded and compression molded
PA-
6/Cloisite 30B nanocomposites .............................................................................. 164 Figure 6.6.12: Tensile strengths of extruded and compression molded
PA-
6/Cloisite 30B nanocomposites .............................................................................. 164 Figure 6.6.13: Tensile moduli of extruded and compression molded PA-6/Cloisite 15A nanocomposites ...................................................................................................... 165 Figure 6.6.14: Tensile strength of extruded and compression molded PA-6/Cloisite 15A nanocomposites ...................................................................................................... 166
xn Figure 6.6.15: Tensile stiength of extruded and compression molded PA-6/Cloisite Na+ nanocomposites ...................................................................................................... 167 Figure 6.6.16: Tensile strength of extruded and compression molded PA-6/Cloisite Na+ nanocomposites ...................................................................................................... 167 Figure 6.6.17: Effect of extruder RPM on the tensile strength of PA-6/Cloisite 30B (3.1 %wt) nanocomposites ...................................................................................... 171 Figure 6.6.18: Effect of extruder RPM on the tensile strength of PA-6/Cloisite 30B (5.0%wt) nanocomposites ...................................................................................... 171 Figure 6.6.19: Impact strength ofPA-6/Cloisite 30B nanocomposites ......................... 173 Figure 6.6.20: Impact strength ofPA-6/Cloisite 15A nanocomposites ......................... 173 Figure 6.6.21: Impact strength ofPA-6/Cloisite Na+ nanocomposites .......................... 174 Figure 6.6.22: Tensile modulus of PA-6/Cloisite 30B (5%wt) nanocomposites processed using System A followed by compression molding or injection molding ............. 177 Figur~
6.6.23: Tensile strength ofPA-6/Cloisite 30B (5%wt) nanocomposites
processed using System A followed by compression molding or injection molding ................................................................................................................................ 177 Figure 6.6.24: Tensile modulus ofPS1301lCIoisite lOA nanocomposites .................... 179 Figure 6.6.25: Tensile strength ofPS1301lCIoisite lOA nanocomposites .................... 179 Figure 6.6.26: Tensile modulus ofPS1301/Dylark (2%wt)/Cloisite lOA nanocomposites ................................................................................................................................ 181 Figure 6.6.27: Tensile strength ofPS1301/Dylark (2%wt)/Cloisite IOA nanocomposites ................................................................................................................................ 181 Figure 6.6.28: Flexural modulus of PS1301lCloisite lOA nanocomposites .................. 182 Figure 6.6.29: Flexural strength of PS1301/Cloisite lOA nanocomposites ................... 182
xm Figure 6.6.30: Experimental data on tensile modulus and model predictions ............... 189 Figure 6.6.31: Experimental data on tensile modulus and model predictions ............... 189 Figure 6.6.32: Experimental data on tensile modulus and model predictions .............. 190 Figure 6.6.33: Experimental data on tensile modulus and model predictions ............... 190 Figure 6.6.34: Schematic representation of adhesion between a clay platelet and the polymer matrix: (a) adhesion between untreated clay, (b) adhesion between organically modified clay and the polymer matrix ................................................ 193 Figure 6.6.35: Schematic illustration of formation of hydrogen bonds in PA-6/MMT nanocomposites (265) ............................................................................................ 194 Figure 6.6.36: Experimental data on tensile strength and model predictions ................ 197 Figure 6.6.37: Experimental data on tensile strength and model predictions ................ 197 Figure 6.6.38: Experimental data on tensile strength and model predictions ~ ............... 198 Figure 6.6.39: Experimental data on tensile strength and model predictions ................ 198 Figure 6.6.40: Experimental data on tensile strength and model predictions ................ 199 Figure 6.7.1: Oxygen permeability coefficient of PA-6/Cloisite 30B nanocomposites 201 Figure 6.7.2: Oxygen permeability coefficient ofPS1301lCIoisite 10A nanocomposites ................................................................................................................................ 201 Figure 6.7.3: Comparison of the oxygen permeability coefficients ofPS/Cloisite lOA and PSlDylark (2% )/Cloisite lOA nanocomposites produced using System A ..... 202 Figure 6.8.1: Schematic representation of flow pattern of polymer and polymerie nanocomposite melt: (a) flow of pure polymer melt near stationary wall, (b) flow of polymer melt containing exfoliated platelets randomly oriented under low shear, (c) flow of polymer melt containing exfoliated platelets oriented under high shear. ................................................................................................................................ 204
XIV Figure 6.8.2: Steady shear viscosities of PA-6 nanocomposites at 240°C: Effect of using clays with different surface modifiers on the viscosity .......................................... 205 Figure 6.8.3: Steady shear viscosities ofPA-6/Cloisite 30B nanocomposites at 240°C: Effect of clay content on the viscosity ................................................................... 206 Figure 6.8.4: Steady shear viscosities of PA-6/Cloisite 15A nanocomposites at 240°C: Effect ofprocessing with System A and System C on the viscosity ..................... 206 Figure 6.8.5: Steady shear viscosities ofPS1301 nanocomposites at 210°C: Effect of using clays with different modifiers on the viscosity ............................................. 207 Figure 6.8.6: Steady shear viscosity of PS130l/Cloisite lOA nanocomposites at 210°C: Effect of clay concentration on the viscosity......................................................... 208 Figure 6.8.7: Steady shear viscosity of PSl301/Dylark (2%wt)/Cloisite lOA nanocomposites at 210°C: Effect of clay concentration on the viscosity.............. 208 Figure 6.8.8: Relation between shear stress in ISG static mixer and clay content in PA6/Cloisite 30B nanocomposite ............................................................................... 211 Figure 6.9.1: Rate of change of specifie volume of PA-6 and PA-6 nanocomposite during Isobaric heating ........................................................................................... 215 Figure 6.9.2: Rate of change of specifie volume of PA-6 and PA-6 nanocomposite during Isobaric cooling .......................................................................................... 215 Figure 6.9.3: Melting and crysta1lization temperatures ofPA-6 and nanocomposite .... 215 Figure 6.9.4: Crysta1lization kinetics of PA-6 and PA-6 nanocomposite at 50 MPa .... 217 Figure 6.9.5: Crystallization kinetics ofPA-6 and PA-6 nanocomposite at 150 MPa .. 218 Figure 6.9.6: FfIR spectra ofPA-6 and PA-6 nanocomposites: ................................... 221
xv LIST OF TABLES Table 4.3.1: Characteristic values for different static mixers (208) ................................. 69 Table 4.3.2: Computed characteristic values for static mixers (208) ............................... 69 Table 5.3.1 Matrix polymers used in this study ............................................................... 95 Table 6.1.1 Mean residence time and variance for System A and System C ................ 104 Table 6.2.1: Hansen Solubility Parameters (J/cm3)112
................................................... 120
Table 6.2.2: Values of Hamaker Constant ..................................................................... 122 Table 6.3.1: Surface density, specific surface area and the degree of exfoliation of PA-6 nanocomposites (244) ............................................................................................ 137 Table 6.3.2: Aspect ratios of clay particles in PA-6 nanocomposites (244) .................. 138 Table 6.4.1: Moisture Diffusion coefficients ofPA-6 nanocomposites ........................ 148 Table 6.6.1: Comparison of tensile moduli and tensile strengths of different nanocomposites produced using System A and System C .................................... 162 Table 6.6.2: Tensile properties of extruded and compression molded PA-6 nanocomposites with different clays ...................................................................... 168 Table 6.6.3: Percentage elongation at break for PA-6 nanocomposites ........................ 169 Table 6.6.4: Mechanical properties of injection molded PA-6 nanoconiposite ............. 175 Table 6.6.5: Tensile properties of PS nanocomposites prepared using System A ......... 183 Table 6.6.6: Tensile properties ofPS nanocomposites prepared using System C ......... 183 Table 6.6.7: Aspect ratios of the clay fillers for polyamide-6 nanocomposites calculated using different composite models .......................................................................... 191 Table 6.6.8: Aspect ratios of the clay fillers for polystyrene nanocomposites calculated using different composite models .......................................................................... 191 Table 6.6.9: Work of adhesion for different clay/polymer systems ............................... 195
XVI Table 6.6.10: Values ofparameter B for different nanocomposites .............................. 199 Table 6.7.1: Aspect ratios of different nanocomposites predicted from oxygen permeability data .................................................................................................... 202 Table 6.8.1: Pressure drop in ISG static mixer for PA-6/Cloisite 30B nanocomposite.210 Table 6.8.2: Pressure drop in ISG static mixer for different nanocomposites at extruder speed of 200 RPM and the available shear stresses ............................................... 212 Table 6.9.1: Avrami parameters and crystallization half-time for y-form and a,-form crystallization ofPA-6 and PNC ............................................................................ 219 Table 6.9.2: Characteristic IR frequencies (cm-1) for PA-6 ........................................... 220
xvn NOMENCLATURE A - Aspect ratio
Ap - Specimen area for permeability measurement Aj - Specifie Surface area of filler
Aii - Hamaker constant between similar species Aij - Hamaker constant between different species
ao - Effective area per molecule of surfactant film li - Mean silicate surface size
am - Root of zeroth order of Bassel function B - Pukanszky parameter
b - Width of platelet bj
-
Width of sample for Flexural testing
C - London constant C(t) - Function of tracer concentration
D - Diffusivity or Diffusion coefficient Di - Inside diameter of static mixer housing d - Gallery spacing between clay platelets dp
-
Diameter of particle
E - Internal energy E(t) - Residence time distribution
Eo - Zero voltage level during permeability measurement Eb - Flexural modulus Ee - Young' s modulus of composite Ee - Steady-state voltage during permeability measurement Ej-
Young's modulus of filler
Em - Young's modulus ofmatrix Er - Ratio of platelet to matrix modulus F-Force AF - Change in Free energy F(t) - Cumulative residence time function
Je - Calibration factor for permeability measurement
xvrn f - Friction factor G - Adhesive fracture energy Gm - Shear modulus h - Plank's constant
hf - Thickness of sample for Flexural testing
ho - Gallery height of unintercalated clay 1 - Moment of inertia
K - A vrami rate parameter for crystallization Kp - Power constant k - Boltzmann constant ks - Number of channels in static mixer
kj - Rate constant of intercalation L - Length of clay particle
Lt- Specimen length for Flexural testing l - Length of static mixer element
1- Peeled length of platelet Mf - Bending moment
MRF - Modulus Reduction Factor Ne
~
Newton's number
N s - Number of amphiphilic molecules
n - A vrami exponent for crystallization 02GTR - Oxygen Gas Transmission Rate Pf - Maximum load during Flexural testing p - Permeability coefficient
P-Pressure Pc - Permeability coefficient of composite Pp -
Permeability coefficient of polymer
Qp - Calibration constant for permeability measurement
Q - Volumetric flow rate R - Radius of agglomerate Re - Reynolds riumber
XIX R L - 1..oad resistor for permeability measurement S -Entropy Sf - Flexural strength SB - Order parameter in Bharadwaj model s - Standard deviation T - Temperature Tf - Transverse force tp
-
Thickness of films used for permeability measurement
t-Time t1/2 -
Crystallization half time
U - Interaction energy
V
00 -
Melt volume at the end of primary crystallization
Vo - Initial melt volume Vt - Melt volume at time t v- Velocity
vp
-
Vii -
Voltage recorded in permeability measurement Volumetrie interaction parameter in Sirnha Lattice-cell hole Theory
w - Mass flow rate W - Width of clay particle Wa - Work of adhesion
W/ - Dispersion forces Wah - Forces due to hydrogen bonding X - Relative crystallinity Z - Ratio of pressure drop through static mixer to that of empty tube
ae - Extensional efficiency a - Polarizability Z - Flory-Huggins interaction energy parameter
bL -
Lyapunov exponent
b - Thickness of clay tactoid or clay platelet bd - Dispersive component of Hansen solubility parameter bd - Hydrogen bonding component of Hansen solubility parameter
xx 4 - Polar component of Hansen solubility parameter E-
Dielectric susceptibility Energetic parameter in Simha Lattice-ceU hole Theory
Ejj -
Volume fraction
fjJ -
r- Mean shear rate y- Surface tension qJ-Angle  - Wavelength
Âs - Stretching distribution f.1 -
Viscosity
f.lc - Geometrie factor for flUer particles in Cussler model f.1D -
Dipole moment
v - Main dispersion frequency B-Angle
p- Density C1'c - Composite yield stress 0; -
Yield stress of interphase in composite
G;n - Matrix yield stress 0;. -
Rupture stress
a/ - Variance of residence time distribution r - Tortuosity factor lilt - Frequency
TIs - Spreading pressure TIc - Critical spreading pressure
XXI
DEFINITIONS OF SOME TERMS IN POLYMERIC NANOCOMPOSITE LlTERATURE
Exfoliated layered material: Individual platelets of an Intercalated Layered material dispersed in a carrier material or a matrix polymer with the distance between individual platelets larger than about 8.8 nm. Exfoliation: Forming an exfoliate from an Intercalate. Intercalant: Material sorbed between platelets of the layered material that binds with the platelet surfaces to form an intercalate. Often a part of intercalant is a cation able to form ionic bond with platelet's anion. Intercalated: Layered material with inserted between platelets organic or inorganic molecules that increase the interlayer spacing between them to at least 1.5 nm. Intercalating Carrier: A carrier comprising water with or without an organic solvent used to form an intercalating composition capable of achieving intercalation of the layered material. Intercalating Composition: A composition comprising an intercalant, an intercalating carrier for the intercalant pol ymer and layered material. Intercalation: Forming an intercalate. Layered Material: Inorganic, such as smectite clay, which is in the form of adjacent layers with a thickness, for each layer of 0.3-1 nm. Matrix Polymer: Thermoplastics, thermosetting or elastomeric polymer in which the intercalate or exfoliate is dispersed to form polymeric nanocomposite (PNC). MMT: Montmorillonite;
XXII
Na-MMT: Sodium montmorillonite; H-MMT: Protonated MMT. Nanocomposite (Ne): An oligomer, polymer or copolymer having dispersed exfoliated individual platelets obtained from an intercalated layered material.
Platelets: lndividuallayers of the layered material. Tactoid: lntercalated or exfoliated clay platelets aligned parallel to each other. Spacing: Two measures of spacing are used; the interlayer spacing (d-spacing, d(OO 1), or basal spacing) and interlamel/ar spacing. The former comprises the latter plus the platelet thickness. For example for MMT d(OOl) = lnterlamellar spacing + 0.96 (nm).
1 INTRODUCTION
1
1 INTRODUCTION Commercial polymer clay composites have existed for several decades. They consist of a polymer matrix and a dispersed phase acting as reinforcement or filler. The nature of the polymer matrix can be very diverse. Thermoplastics such as nylon or polyolefin have been used extensively. The dispersed phase can be in the form of individual fibers, woven or non-woven fabrics, particles made of carbon, glass, metal, engineering thermoplastics etc. The traditional composites are therefore diverse because they allow one to tailor the properties of the material to specific needs. However improving one property often means compromising other important properties, e.g. improving the mechanical properties by incorporating glass fibers increases density significantly. Conventional polymer composites cannot be processed into films or fibers (l). These are sorne of the serious drawbacks of traditional polymer composites. The field of nanocomposite materials has drawn the attention, imagination, and close scrutiny of scientists and engineers in recent years. This is based on the premise that using building blocks with dimensions in the nanosize range makes it possible to create new materials with unprecedented improvements in physical properties. Nanocomposites are a special class of composites, in which the dispersed phase is made of nanoparticles .. Nanoparticles have one dimension of the order of 1 nanometer (10-9 m). The term nanocomposite refers to the combination of two or more component phases where at least one phase dimension is in the nanometer range (2). These materials exhibit behavior different from conventional composite materials with microscale structure, due to the small size of the reinforcing unit and the high surface-to-volume ratio. Nanoparticles are invisible to the naked eye; hence, they may be used to engender reinforced, but transparent composites. On the molecular level, the surface energy or the solvation force of clay particle is high. As a result, adsorbed matrix molecules have a tendency to lay in layers next to the clay surface (3). In consequence of these considerations, nanocomposites normally require 2-3 vol % (5-6 wt %) of
1 INTRODUCTION
2
nanoparticles. Thus, nanocomposites appear to behave as single phase and single component materials. Composites that exhibit a change in composition and structure over a nanometer length scale have been shown to afford remarkable property enhancements relative to conventional composites (4). Although the high aspect ratio of silicate nanolayers is ideal for reinforcement, the nanolayers are not easily dispersed in most polymers, due to their preferred face-to-face stacking in agglomerated tactoids. From the fundamental point of view, the reinforcing effect of nanoparticles is related to the aspect ratio and to the particle-matrix interactions. Dispersion of tactoids into discrete monolayers is further hindered by the intrinsic incompatibility of the hydrophilic layered silicates and hydrophobie polymers. This necessitates the compatibilization of the clay surface. Toyota researchers (5) demonstrated that the replacement of the inorganic exchange cations in the galleries of native clay by alkylammonium surfactants could compatibilize the surface chemistry of the clay and the hydrophobie polymer matrix. This finding was a breakthrough in the development of polymerie nanocomposites. It has been shown that at a loading of only 4.2 wt% clay, the tensile modulus of
nylon-6 was doubled, the strength increased more than 50%, and the heat distortion temperature increased by 80 oC compared to the pristine polymer (6,7). It was further demonstrated that organoclays exfoliated in a nylon-6 polymer matrix greatly improved the dimensional stability, the barrier properties and even tlame-retardant properties. The use of organoclays as precursors to nanocomposite formation has been extended into various polymer systems (4), which include epoxies, polyurethanes, polyimides, nitrile rubber, polyesters, polypropylene, polystyrene, and polysiloxanes. For true nanocomposites, the clay nanolayers must be uniformly dispersed in the polymer matrix, as opposed to being aggregated as tactoids or simply intercalated. The complete dispersion of clay nanolayers in a pol ymer maximizes the number of available reinforcing elements for carrying an applied load and detlecting cracks. The coupling between the tremendous surface area of the clay (-760 m 2 jg) and the polymer matrix facilitates stress transfer to the reinforcement phase, allowing for tensile and toughening
1 INTRODUCTION
3
improvements. Conventional polymer-clay composites containing aggregates ordinarily improve rigidity, but they often sacrifice strength, elongation and toughness. However, exfoliated clay nanocomposites of nylon-6 and epoxy systems have shown improvements in aU aspects of their mechanical performance. Nylon-6 was the first polymer to be used in the development of nanocomposites over a decade ago. Development activities have spread to all regions of the world. Sorne of the programs are focused on polypropylene, polyester, polyvinyl chloride, acrylics, polystyrene and a range of elastomers as weU as traditional thermosets. As estimated by Principia Partners, Exton PA, USA, the global market for nanocomposites is 3 million pounds, of which 2 million pounds were nanoclay-reinforced products. Market projections show that demand will grow at comparable rates from 2004 through 2009. The market will reach nearly 1.2 billion pounds in 2009, of which 1 billion pound will be nanoclay-reinforced compound. Major applications will be automotive components, packaging, appliances, electricallelectronic parts, and building and construction products (http://www.principiaconsulting.comlreleases nanocompreachl b.cfm). Layered-silicate based polymer nanocomposites are attractive not only for their obvious potential as technological materials, but also for providing a convenient macroscopic system to study fundamental scientific issues concerning confined and tethered polymers. Studying the formation, structure and dynamics of nanocomposites can lead to better understanding of organic-inorganic hybrids and polymers in a confined environment or at a solid interface. In this research project, we studied sorne of the key issues involved in the exfoliation/intercalation of nanoclay in nanocomposites, including the requirements for process/equipment design. We designed a process, based on these considerations, for producing nanocomposites. The nanocomposites produced using this process show significant improvements in physical properties over the products from conventional processing methods.
2 LITERATURE REVIEW
4
2 LITERATURE REVIEW Nanocomposites are materials that comprise a dispersion of nanometer-size particles in a matrix. The matrix may be single phase or multiphase. Depending on the nature of the matrix, nanocomposites may be classified into polymeric, ceramic or metallic nanocomposites. The filler nanoparticles may be composed of metals, metal oxides, carbon, as weIl as inorganic mineraIs with lamellar, fibriIlar, shell-like or spherical shape. For the enhancement of mechanical and barrier properties, anisometric particles such as lamellar or fibrillar are preferred. The present study focuses on inorganic clayreinforced polymeric nanocomposites. There are several reasons for using clays, e.g., availability, cost and aspect ratio. The layered structure of clay particles (8) makes it feasible to produce nanoparticles with high aspect ratio during processing. The smectite clays, in particular, are unique in the fact that they exist in nature in turbostratic units that are hydrophilic and can be broken down into one nanometer thick platelets (1). Polymer-clay nanocomposites have their origin in the pioneering research conducted at Toyota Central Research Laboratories (5,6,7,9,10,11), where these two divergent organic and mineraI materials were successfully integrated. The first practical application of nanocomposites was iü the use of nylon-montmorillonite clay nanocomposite as a timing belt coyer on a Toyota Camry automobile. This nanocomposite exhibited large increases in tensile strength, modulus, and heat distortion temperature without 10ss of impact resistance. The composite also had lower water sensitivity, permeability to gases, and thermal coefficient of expansion. In this review, clay structure, formation of hybrid nanocomposite structure, thermodynamics of nanocomposite formation and different phase behavior models will be
discussed,
followed
by
review
of hydrodynamics
of
clay
dispersion.
Micromechanical models for the prediction of physical properties such as tensile modulus and strength, gas permeability will then be discussed. The last section will deal with crystallization behavior and kinetics of polymer used in nanocomposites.
2 LITERATURE REVIEW
5
2.1 LAYERED SILICATE STRUCTURE The layered silicates used in nanocotllPosites belong to the same structural family of mineraIs talc and mica (12). The most commonly used layered silicates are montmorillonite, hectorite and saponite. AlI of these silicates are characterized by large active surface area (700-800 m2/gm) and a moderate negative surface charge. Their crystal lattice consists of 1 nm thick layers which are made up of two tetrahedral sheets of silica fused to an octahedral sheet of alumina or magnesia (Figure 2.1.1, Figure 2.1.2). The lateral dimensions of these layers vary from 30 nm to several microns. Stacking of the layers leads to a regular van der Waals gap between them, called the interlayer or gallery. Because of the relatively weak forces between the layers, intercalation of various molecules, even polymers, between the layers is facile. Pristine mica-type layered silicates usually contain hydrated Na+ or K+ ions. Ion ex change reactions with cation surfactants render the normally hydrophilic silicate surface organophilic, which makes possible intercalation ofmany polymers. Isomorphic substitution within the layer generates negative charges that are normally counterbalanced by hydrated alkali or alkaline earth cations residing in the interlayer. These non-covalent interactions between the sheets are van der Waals forces. The interactions respnsible for these forces became c1ear with the work of Keesom, Debeyand London (8,13,14,15) as, respectively, interactions between two permanent dipoles (orientation forces), a permanent dipole and an induced dipole (induction forces), and a fluctuating dipole and an induced dipole (dispersion forces). While these three kinds of interaction have different origins, the interaction energies for aIl three vary as the inverse of the distance raised to the sixth power (8). Uorientation
= -f.1D
4 /
kTd
6
(2.1) (2.2)
Udispersion
3 2 6 = - 4 a hv /d
(2.3)
2 LITERATURE REVIEW
6
where /iD is the dipole moment, k is the Boltzmann constant, T
is the absolute
temperature, a is the polarizability, d is the interatomic distance, v is the main dispersion frequency and h is Plank' s constant.
~ ~:~:~~~;r,~~,-r··_·· ~ :(rMI ..... _L _, _'co lnterl,aver
Na+
l
...............................................................-.
~~~~
0:0
@:OH o .• :Sî (AI)
•
:AI,Fe,Mg
Figure 2.1.1: Structure of 2: 1 layered silicate (1,12)
DISTANCE 8ETWEEII ATOM CEl/TEllS TETRAHEDAAI.
CttARGE
.. 0
060 A
-12
4si SHEET
+111
1.60 A -10 4 0 20H
OCTAHEDRAL +12
4AI
2.20 A SHEET
2 OH -10 40 TETRAHEDRAL 4si .0
SliEET
+16
-12
--,.. __ ... --- 44
SURFACE
0
0
C§I
Ott
•
51
Â
AI
!UII A FORMULA
OF THE 8.' A 2
+44
UNIT CEI.L:
OF UNIT CELL'
C - SPACIIIG '.2 A
[ AI 210HI2 (Siz'Os'zJ 2 UI/IT cHL WEIGHT: 120 HYOROXYL WAHR; s%
Figure 2.1.2: Clay particles consisting of stacks of identical units (14,15)
2 LITERATURE REVIEW
7
2.1.1 THE HAMAKER APPROXIMATION For two atoms or molecules (i) separated by a short distance in vacuum, the dispersion interaction can be expressed as: (2.4)
where Aii is commonly termed as Hamaker constant, n is the number of atoms per unit
volu~e and Cii =(3/4)a2hv. From Equation 2.3, it follows that: (2.5)
where dis the distance between the atoms i. Extending the interaction to two different atoms i and j :
Ci) =~CiiCjj
(2.6)
It follows then that the Hamaker constant for two different atoms is given by:
Ai) =~AiiAjj
(2.7)
Hamaker calculated the dispersion (van der Waals-London) interaction energy for larger bodies by pair-wise summation of the properties of the individual molecules. Using macroscopic approximation, the total dispersion energy for two semi-infinite fiat paralle1 bodies, separated by a distance d (for d greater than a few atomic diameters), in air or in vacuum, becomes:
Udispersion = - Aii /12trd
2
(2.8)
where A is the Hamaker constant for material i. When two or more materials interact, the total Hamaker constant is determinedby a geometric combining rule. For two atoms of the same material 1, in medium 3 (e.g. two individual clay particles in an aqueous medium or with a surfactant between them) the combining rule gives: (2.9)
or
(2.10)
2 LITERATURE REVIEW
8
For two different objects 1 and 2, in medium 3: A132 =
(fA:: - ~A33 X~A22 - ~A33)
(2.11)
or
(2.12)
For the case where two objects of the same material are embedded in another material, the Hamaker constant, Al3J is always positive. For two different materials, the Hamaker constant, A 132 can be negative when, AlI
> A33 > A22 or when, A Il < A33 < A22.
2.1.2 THE LlFSHITZ APPROACH Lifshitz approached the problem of van der Waals interaction by examining the macroscopic properties of materials, as opposed to the Hamaker treatment of summing individual atomic interactions. The derivation is based on Maxwell's equations, modified to allow rapid temporal fluctuations. The expression for the Hamaker constant Aii is given as:
(2.13)
where k is Boltzmann's constant, T is the absolute temperature,
Wn
is the frequency, and
e (iw n ) is the dielectric susceptibility along the complex frequency axis. 2.1.3 POL.AR FORCES . In aqueous media, and especially for solid surfaces which are rich in oxygen such as silicate mineraIs, the principle polar interaction is hydrogen bonding, involving donors and accepters. These interactions constitute Lewis Acid-Base interactions, ionic double layer, and electrostatic interactions. Since melt processing involves a matrix of nonaqueous organophilic polymer melts, polar forces play an insignificant role in melt exfoliation. Hence, a detailed discussion of polar forces is not presented here.
2 LITERATURE REVIEW
9
2.2 NANOCOMPOSITE FORMATION AND STRUCTURE Scientists have known for about 40 years that sorne polymers interact strongly with montmorillonite and that the clay surface can act as an initiator for polymerization (16,17). Patents for claylNylon-6 composites were issued only in 1980, when clay/polymer nanocomposites were commercialized (18). The improvement that led to commercialization was the appropriate dispersion of the clays at the nanometer scale. The first step in achieving nanoscale dispersion of clays in polymers is to open the galleries and to match the polarity of the polymer or monomer so that it will intercalate between the layers (19). This is done by exchanging an organic cation for an inorganic cation. The larger organic cations swell the layers and increase the hydrophobie (or organophilic) properties of the clay (20), resulting in an organically modified clay. The organically modified clay can then be intercalated with polymer by several routes (Figure 2.2.2). Highly polar polymers such as Nylon and polyimides are more easily intercalated than nonpolar polymers such as polypropylene or polyethylene (9,10,11). Solution processing involves dispersion of organically modified clay and polymer in a common solvent followed by emulsion or suspension polymerization (21). In in-situ polymerization, monomer intercalates directly into the organically modified clay galleries, and the monomer can either adsorb onto the layer surface (22) or be anchored by a free radical reaction (23). Melt intercalation involves mixing the clay and polymer melt, with or without shear (1,19). Other methods such as monomer modification (24,25) and covulcanization (26) are specific for acrylic resins and elastomers respective1y. Nanocomposites can have several structures (Figure 2.2.3). Intercalated nanocomposites incorporate tactoids with expanded interlayer spacing, but the clay galleries have fixed interlayer spacing. Exfoliated nanocomposites are formed when the individual clay layers break off the tactoid and are either randomly dispersed in the polymer (a disordered nanocomposite) or left in an ordered array.
2 LITERATURE REVIEW
10
Lan et al. (27,28,29,30) have found that, as the degree of exfoliation increases by changing the length of the alkylammonium intercalating chain, the modulus and strength of nanocomposite improves. As the polymer intercalates and sweUs the gaUeries, the layers and the area of interaction between the polymer and fiUer increase, and the modulus increases significantly. The types of matrix/fiUer interactions responsible for property improvement in nanocomposites are be1ieved to be between the matrix and clay surface (type A in Figure 2.2.3). In epoxy/clay thermoset nanocomposites, direct interaction of the matrix with the clay basal planes leads to larger increase in the tensile modulus.
Clay b:iJrganic cation i organk
Ctlli()ll
cxclllmgc
'-,
]
oligO~;'i'::~:;l"ioO;:;lYm,
-
Polymer melt or solution intercalation
Polymerizntion
lntcrcalated nanocomposite
Figure 2.2.1: Schematic of the basic steps in nanocomposite formation (19)
2 LITERATURE REVIEW
A. Conventiona!Cornposite
Il
B. Imcrcalllwd Nanocmn!1osüc
with Tactoids
r
_
Nanocornposile
Figure 2.2.2: Schematic of microstructures of layered silicate nanocomposites (19) (A) a conventional micro composite with tactoids, (8) an intercalated nanocomposite, (C) an ordered exfoliated nanocomposite, (0) a disordered exfoliated nanocomposite
Po/ymer Matrix
Clay Layer
Figure 2.2.3: Schematic of types of interfacial interactions occurring in polymerorganoclay nanocomposites (30). (A) Direct binding of polymer to the basal siloxane oxygens, (8) Dissolving of onium ion chains in the polymer matrix, (C) Polymer binding to hydroxylated edge sites.
2 LITERATURE REVIEW
12
Okada et al. (20) found that the excellent properties of nylon nanocomposites result from the enormous interfacial area and ionic bonds between the organic polymer and inorganic silicate. Nylon-6 molecules have -NH2 and -COOH end groups. Titration of nanocomposite revealed that the concentration of -COOH ends was much higher than that of -NH2 ends. The results indicate that the nylon molecules are ion-bonded to silicate at the -NH2 end. Each silicate works as a crosslinker and restricts molecular motion. The filler can constrain the mobility of the polymer chains and modify their relaxation spectra (31), which can change the glass transition temperature (31,32) and tensile modulus of the matrix. In general, the greatest physical property enhancements for polymerie nanocomposites are achieved with less than 4 vol% addition of nanoscale dispersion of 1 nm thick silicate layers with diameter between 20 and 500nm (1). These enhancements appear to be a general phenomenon related to the nanoscale dispersion of layers, but the degree of property enhancement is not universal for aIl polymers. The strength of interactions between the polymer and the silicate, as weIl as the size and rigidity of the inorganic partic1es, have been shown to influence the extent of enhancement and viscoelastic behavior associated with processing (33,34,35). The fabrication of nanocomposite involves transformation of initial heterogeneous, microscale morphology into a homogeneous morphology on the nanoscale.
2.3 THERMODYMAMICS OF NANOCOMPOSITE FORMATION The formation and equilibrium structure of polymer layered silicate nanocomposites with organically modified layered silicates depends on the nature of the polymer (polar or apolar) and the charge carrying capacity of the layered silicate, as well as the chain length and structure of cationic surfactant. In general, interplay of entropie and enthalpie factors determines the outcome of polymer intercalation (36,37). Confinement of the pol ymer inside the inerlayers results in a decrease in the overall entropy of the pol ymer chains. However, the entropie penalty of polymer confinement may be compensated by
2 LITERATURE REVIEW
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the increased conformational freedom of the tethered surfactant chains in a less confined environment, as the layers separate (Figure 2.3.1). Vaia and Giannelis (36,37) developed a semi quantitative mean-field model to describe pol ymer melt intercalation in organically-modified mica-type silicates. The free energy change associated with layer separation and polymer incorporation was separated into two components, the internaI energy change associated with the establishment of new intermolecular interactions as M, and the entropy change associated with configurational changes of various constituents as !l.S. The total change in Helmholtz free energy, M', associated with layer separation from unintercalated interlayer of gallery height ho to a polymer-intercalated interlayer of gallery height h is:
M' = F(h)-F(ho)= M -TM
(2.14)
Negative value for M'indicates layer separation is favorable, and positive value implies the initial unintercalated state is favorable (Figure 2.3.2). Major factors contributing to the free energy change are relative confinement of polymer, configurational changes of tethered chains and establishment of new intermolecular interactions between the polymer, the tethered chains and the silicate surface. By separating and developing expressions for M and M independently, the effect of intermolecular interactions on conformational freedom of the pol ymer and the organically modified silicates can be accounted. For small increase in gallery height, the total entropy change is small, modest changes
in
the
system' s
total
enthalpy
will
determine
if intercalation
is
thermodynamically possible. The enthalpy of mixing can be c1assified into two components - apolar and polar. Apolar will be unfavorable. A favorable enthalpy change is accentuated by maximizing the magnitude and number of polymer-surface interactions and minimizing the apolar interactions.
2 LITERATURE REVIEW
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Â
+
@9Y Figure 2.3.1: Schematic representation of the system components before and after the intercalation (36)
-Total Entropy -
~SurfllCC
_. -Polymer
~--
..c
-----'-------
\ -1
-.- -- -'- -2~~~~~~~~~~~~~~
0.0
0,5
L5 h~ho'
2.5
nm
Figure 2.3.2: The change of entropy per area versus the change in gallery height, for the polymer and the surfactant functionalized surface (36).
This model has ability to determine the effects of various aspects of pol ymer and organically modified silicate on hybrid formation. This mode1 is not applicable to situations where the interlayer is not completely occupied by tethered chain segments, which is the case for many silicates with low charge densities or modified by short aliphatic
chains.
Assumptions
such
as
separation of configurational
terms,
intermolecular interactions, and entropie behavior of the constituents limit the usefulness of the mode!.
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Simha et al. (38) used the lattice-cell theory and put forward a model for nanocomposite system, which comprises rigid solid clay platelets and flexible polymer melt as matrix. The idealized clay particle is 100 nm in diameter and 1 nm in thickness. The computation of binary parameters yielded following values for energetic and volumetrie interaction parameters: For PA-6 [IJ] and clay [zz] e;1 = 32.09 (kJ/mol) e;2
=3063.55 =95.4676xe;1
(kJ/mol)
V;1 = V;2 = 24.89 (mL/mol)
For PA-6/clay nanocomposite [12] e;2 = 313.54 = 9.7707xe;1 (kJ/mol) V;2 = 33.53 = 1.3473xe;1 (mL/mol)
This shows that solid-solid interactions are 95 times stronger than liquid-liquid ones, and PA-6--clay interactions are 10 times stronger than those of between PA-6 polymer chains. The adsorption of PA-6 increases the PA-6-clay cell size by 35 %. The addition of 1.6 wt% of exfoliated montmorillonite nanoparticles reduced the hole fraction (free volume) by 14 %. Hole fraction is a sensitive indicator of structural changes. The compated values of specifie volume at different pressures and temperatures agree with experimental values. Since lattice theory of polymer melt intercalation (36,37) predicts that the polymer-surface interactions are dependent upon the polarity of the polymers. Polar polymers will favor the intercalation where as non-polar polymers will not form the intercalated structure with clay. The formation of nanocomposites from apolar polymer such as polypropylene was not realized even by using organically modified clay. Kawasumi et al. (39) reported an approach to prepare polypropylene nanocomposites by using functional polyolefin oligomer, e.g. maleic anhydride grafted polypropylene oligomer, as compatibilizer. Kuznetsov and Balazs (40) modeled end-functionalized polymers confined between two surfaces and predicted the fabrication of pol ymer-clay nanocomposites
2 LITERATURE REVIEW
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based on Scaling theory. The chains in the gallery can assume one of the following four conformations: loops, bridges, tails, and free chains as shown in Figure 2.3.3 The nonfunctionalized chains are referred to as P-polymers and functionalized chains as Npolymers. The melt contains
f/J volume fraction of N-polymers and (l-f/J) volume fraction
of P-polymers. The free energy of chains between the sheets can be written as (2.15)
F ads -
sticker surface adsorption energy for functionalized chains in loop, bridge and tail conformations,
FinI -
interaction energy between monomers of the N and P chains,
F comp - compression energy for chainsat relatively short H, F el -
e1ongation energy for loops, tails and bridges,
Fenl -
entropic energy for mixing the different types of chain configurations,
F dem - demixing energy associated with extracting N and P chains from the surrounding melt and localizing them in the confined layers. The free energy described ab ove is dependent on the following parameters: the energy per contact between a sheet and a functionalized chain end,
f/J,
eT,
the volume
fraction of functionalized chains in the bulk melt, N and P , the numbers of segments in each functionalized and nonfunctionalized chains respectively, and
%PN ,
the Flory-
Huggins interaction energy parameter for interactions between monomers of different chain types (N and P). The free volume is minimized with respect to the numbers or volume fractions f/JP,f/Jt,f/Jt,f/Jb,f/Jt for the different polymers (f/Jp and
f/JN= f/Jt+f/Jr+f/Jb+f/Jt) and the different
conformations (loops, tails, bridges and free chains) of N-chains. The volume fractions satisfy .the incompressibility constraint, so one of the volume fractions above is completely determined by the others. Thus the minimization has to be fulfilled with respect to the four volume fractions for the melt composed of two types of polymers or
2 LITERATURE REVIEW
17
with respect to three volume fractions for the melt composed of a single type of Npolymer. Using scaling theory it is possible to investigate the equilibrium properties of melt composed of two surfaces and polymer chains. Equilibrium free energy for a pol ymer melt confined between two paraUel sheets at a di.stance H, can be calculated. In this way the free-energy profile is obtained and global and local minima are analyzed for different values of the system parameters (Figure 2.3.4). Balazs et al. used scaling theory to predict the phase behavior in polymer/c1ay systems (41,42,43,44,45,46,47,48). Kim et al. (49) extended self-consistant mean-field theory to map the degree of dispersion as a function of independent variables such as magnitude of the interaction parameters, molecular weights, composition etc. using ID and 2D models. The simulation results for polyolefin nanocomposites showed that intercalation and exfoliation is expected within limited ranges of the independent variables. The optimum ranges of the compatibilizer and intercalant concentrations were identified.
H
Figure 2.3.3: Sketch of the nonfunctionalized P-chains and the different conformations
of the end-functionalized N-chains in the polymer melt confined between two Clay sheets, equilibrated at a distance H (40).
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18
-0.8 ..\F/A 0=0.01
-12
4> '" 0.05 (1,.
-1.6
0.1
1...-_ _- ' -_ _--'-_ _ _"""-_ _....1
o
10
20
30
40
H
Figure 2.3.4: Free energy per unit area, l!..F/A, as a function of surface separation, H. The length of the functionalized polymers is fixed at N
= 100
nonfunctionalized polymers is given by P = 300. The plots are for X shown for three different values of
and the length of
= -75.
Results are
t/J, the volume fraction of functionalized chains (50).
2.4 KINETICS OF POLYMER MELT INTERCALATION It is rather surprising that polymer chains can intercalate layered inorganic compounds unassisted by shear (51,52). It implies that polymer chains can undergo large center of mass displacement in almost two dimensional interstices as the distances between the confining surfaces are substantially smaller than unperturbed radius of gyration of the polymer and are comparable to the monomer size. The reduction in free energy by the intercalate formation and the concentration gradient during the intercalation process give rise to an enthalpic force which drives the pol ymer into the interlayer galleries. On the other hand, the conformational energy cost of stretching the chains, in addition to the topographical constrains and the adsorption on the surfaces are expected to impose severe limitations on diffusion of chains into two dimensional slit. Vaia et al. (53) observed that the kinetics of intercalation under quiescent conditions (absence ofextemal shear) are quite rapid. The evolution of the XRD during annealing can be modeled to determine the apparent diffusivity, D, of the polymer
2 LITERATURE REVIEW
19
within the silicate gallery. The ratio ofthe amount of intercalated pol ymer at time t, Q(t), to that at equilibrium Q( 00) is (54,55): Q(t) = 1-
Q(oo)
i
~exp(-~a 2
m=1 a~
a
tJ
m
(2.16)
where D/t? - effective diffusional coefficient ii - mean size of the silicate surface
am - roots of zeroth order of Bassel function ( Jo(a)=0 ) The apparent diffusivity for the intercalation of polystyrene m organically modified clay was the same order of magnitude (lO-fI cm2/s at 170°C) (53,56) as the self diffusion coefficient of polystyrene determined at comparable temperatures and molecular weights. The activation energy of melt intercalation was 166 kJ/mol, which is comparable to the activation energy for self diffusion of polystyrene (l67 kJ/mol) (57,58). The initial results of their study suggest that the process of intercalation is dictated by the transport of pol ymer to the silicate agglomerates and not by which the pol ymer is moved inside the galleries of the silicates. However Manias et al. (59) proved that this is not the case always. It was found that the effective diffusion coefficient depends on the surfactant used, molecular weight of the polymer and annealing . temperature. Because the surfactant can only affect the polymer motion inside the galleries, one concludes that this motion of the intercalated polymer is the process, which dictates the intercalation kinetics. Further for sorne systems the intercalating polymers were found to possess a mobility that was much faster than the self-diffusion coefficient of the corresponding polymer in the bulk. Since intercalation is a process where polymers are moving down a concentration gradient this phenomenon could be expected. In bulk the polymer motion is entropie in origin. The pol ymer diffusion coefficient increases considerably with the length of the surfactant-typically 10 fold going from 12 to 18 carbons (59). Polymer diffusion coefficient for polystyrene at 170 oC for Mw 35000 to 900000 was found to bear
2 LITERATURE REVIEW
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following relationship with pol ymer length N. N ranges from 2 to 52 entanglement lengths: D
eff
oc
N-I. 09 ±O,05
'
(2.17)
Zang et al. (52) investigated melt intercalation kinetics for Nylon1010 in organically modified montmorillonite, assuming the intercalation process to be firstorder kinetics process bearing the relation:
F(t) = 1- exp(k;t)
(2.18)
where ki is the rate constant of intercalation. The rate constants at various temperatures were estimated to calculate activation energy which was estimated to be 124 kJ/mol for Nylon 1010 intercalation process. Chen et al. (60) investigated the interplay between thermodynamics and kinetics for
polystyrene melt intercalation process. They conc1uded that at low temperatures kinetics plays major role, where as at high temperatures thermodynamics of the process is more important. Poly(styrene-block-isoprene) copolymer intercalation behavior shows that with the increase in size of polystyrene block, the intercalation kinetics was slower (61). Most of the studies on intercalation kinetics use X-ray diffraction method to track the intercalation. Li et al. (62) used rheological approach to study meIt intercalation of pol)'propylene nanocomposites. It was seen that the increasing of intercalated tactoids detached from the primary partic1es resulted in the continuous enhancement of low frequency modulus and viscosity and the formation of the percolation networks. The rheologicalparameters ofviscosity and storage modulus were used to caIculate diffusivities. Apparent diffusivity values were of the order of 10- 12 cm 2/s, and activation energy for maleic anhydride grafted polypropylene intercalation
was 84 kJ/mol. Molecular dynamics simulations of intercalation process by Lee et al. (63,64,65,66,67) are consistent with a diffusive description of the transport, and show qualitative agreement with time-dependent X-ray diffraction measurements of intercalation kinetics in layered nanocomposites. Using computer simulations, Hackett et al. (68) showed that the intercalated polymer chains in PEO-Iayered silicate
2 LITERATURE REVIEW
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nanocomposites are arranged in discrete subnanometer layers parallel to the silicate layers. Structures and thermodynamics of nondilute polymer solutions confined between parallel plates were also studied using Monte Carlo simulations (69,70). There are attempts made to predict binding energy for nylon-6 (71) and nylon-66 (72) organoclay composites by molecular modeling technique.
2.5 HYDRODYNAMICS OF DISPERSION AND DISTRIBUTION OF CLAY IN POLYMER MATRIX One of the key limitations in the commercialization of nanocomposites is processing. Without proper dispersion and distribution of fil 1ers, the high surface area is compromised and aggregates can act as defects, which limit properties. The dispersion of clay in polymer matrix is influenced by two factors, namely the choice of organic treatment for clay and the processing or mixing method used. It is weIl documented (36,37) that the choice of the organic treatment to the clay influences the degree of dispersion. A key to the benefits of clay in nanocomposites, but the challenge in the processing to make nanocomposites, is that there are greater than 3000 platelets in each 6 JlIl1 particle. For most applications, it is generally believed maximum benefits are achieved when the platelets are weIl dispersed (59). To achieve the potential property improvements requires sorne degree of delamination and dispersion. These are dependent upon a combination of the proper chemical treatment and optimized processing. The proper balance of dispersive and distributive mixing results in optimized delamination and dispersion. Distribution of nanofiller describes the homogeneity throughout the sample, and dispersion describes the level of agglomeration (Figure 2.5.1).
2 LITERATURE REVIEW
•
22
• • • • •• •• • • •• •
.--
..
(al
(c) ' - - - - - - -
;. •
·,. -.-.... -..... • •••••
1 ••••••• •
(b) , _ _ _ _ _ _ _- - - l
(d1~. ~
Figure 2.5.1: A schematic illustration of the difference between dispersion and distribution. (a) Good distribution but poor dispersion, (b) Poor distribution poor dispersion,
(C') Poor distribution good dispersion, (d) Good distribution good dispersion (19)
2.5.1 MODELS FOR DISPERSION PROCESS Tadmor (73) analyzed dispersive mixing in polymer processing by modeling agglomerates as dumbbells consisting of two unequal beads connected by rigid connector. The force in the connector was calculated when the dumbbell is placed in a general homogeneous velocity field of a Newtonian fluid. Rupture occurs when the force in the connector exceeds a certain threshold value. In simple shearing flow and steady elongation flow the maximum force in the connector is proportional to the local shear stress and the product ofbeads iadii. In the former the maximum value is obtained when dumbbell is 45° to the direction of flow and in the .latter when dumbbell is aligned to the direction offlow.
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Manas-Zloczower and Feke (74,75) developed dispersive mixing model for rupture of agglomerates. The dispersion process was analyzed by considering the relative motion of fragments subject to the effective van der Waals and hydrodynamic forces through fragment trajectory analysis. The results show that agglomerate size do es not have influence on the kinetics of the separation process and sm aller agglomerates were found to separate earlier, and break to a greater extent, than larger agglomerates. Pure elongation flow field was the most efficient in particle separation followed by simple shear. Uniaxial and biaxial flow fields were less efficient. Coury and Aguiar (76) reviewed two classical theories for rupture of agglomerates. According to Rumpfs theory (77), the limiting strength of an agglomerate is reached when the separation forces imposed by the normal stress equal the adhesion forces. It is therefore assumed that the agglomerate rupture occurs with simultaneous collapse of the interparticle links at the rupture surface. Kendall (78,79) argued that the assumption of rupture of the agglomerate as suggested by Rumpf overestimates its strength and proposed a mechanism similar to the failure of brittle materials where the rupture occurs from the build-up of tensions in defects already present in the brittle solid. This means much smaller energy consumption than that implicit in Rumpfs model. Kendall further explained the phenomenon of adhesion in agglomerates (80) and composites (81) based on this theory. Relation between molecular adhesion at nonometer scale and elastic deformation of solids presented by Kendall (82) could be useful in explaining exfoliation of clay particle. Kinloch et al. (83) modeled the fracture behavior of adhesive joints using Kendall approach. Steven-Fountain et al. (84) used similar approach to explain effect of flexible substrate on pressure-sensitive adhesive performance. Garrivier et al. (85) presented peeling model for cell detachment from cytoplasmic membrane using the analogy of adhesion and fracture. Complex dynamics in peeling adhesive tape was modeled by Ciccotti et al. (86) as a two-dimensional fracture propagation. Niedballa and Husemann (87) modeled deglomeration of fine aggregate particles in an air stream. It is necessary for deglomeration that the van der Waals forces are
2 LITERATURE REVIEW
24
smaller than the dispersion force. The derived dispersion model is able to predict whether the agglomerate can be dispersed due to the stress in the flow. Serville et al. (88) reviewed the interparticle forces in fluidization and agglomeration with respect to van der Waals forces, liquid bridges and sintering. Dispersionmechanism of coagulated particles in liquid flow was studied by Endo and Kousaka (89). They showed that in a shear flow field, both dispersion and shear coagulation occur simultaneously. Schaefer's (90) study on growth mechanism in melt agglomeration in high shear mixer shows that agglomeration is controlled by the balance between the agglomerate strength and the shearing forces. Reddi and Bonala (91) model represents clusters of clay platelets as particles similar to spherical sand grains and express cohesion in terms of hydrodynamic forces associated with critical shear stress. Fedodeyev (92) modeled molecular interactions of di sc-shape flat body and calculated attraction force between the clay platelets. Park and Jana (93) investigated the mechanism of exfoliation of nanoclay particles in epoxy-clay nanocomposites. The elastic force exerted by cross-linked epoxy molecules inside the clay galleries was found responsible for exfoliation of clay layers from the intercalated tactoids. Ginzburg et al. (94) proposed a Kink model to describe the dynamics of pol ymer melt intercalation in the gallery between the adjacent clay sheets. The intercalation process is thought to be driven by the motion of localized excitations (kinks) which open up the tip between the clay sheets. Kinks appear due to the interplay between double-well potential of the clay-clay long range interaction, bending elasticity of the sheets, and external shear force. This model was further used to describe the structural transitions in polymer-clay nanocomposites (95).
2.5.2 DISTRIBUTIVE MIXING Efficient distributive mixing can be achieved by frequent splitting and reorientation, and with high level of shear strain and elongation strain. Agassant and Poitou (96) analyzed distributive mixing by using kinematic approach. They associated mixing efficiency with Lyapunov exponent. Lyapunov exponent is a function of orientation of particle
2 LITERATURE REVIEW
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with respect to its initial position and time. In a good distributive mixing device fluid elements have uniform exposure to the shearing and elongational mixing actions with minimum demixing effects. Demixing occurs when a fluid element is deformed in one direction and later in the opposite direction. Such effects take place in screw extruders. Concept of 'chaotic mixing' is also used to analyze and improve the efficiency of mixing. Mixing is often analyzed in terms of the deformation of an element of the minor component in the major component. In certain types of flow, materiallines can undergo exponential stretching. A convenient way to achieve this is to force the flow to vary with time in a periodic manner. To accomplish effective mixing, such a flow must be able to stretch and fold a region of fluid and return it to its initial location. This process of stretching and folding is referred to as a 'horseshoe map' (Figure 2.5.2), as described by Smale (97). This is one of the identifying features of chaotic flow.
~ ;
• •
ü
III LJ Figure 2.5.2: Representation of Smale horseshoe function
Two-dimensional flows always consist of the same building blocks: hyperbolic points and elliptic points (Figure 2.5.3). In time-dependent cavity flow, a region of outflow associated with a hyperbolic periodic point can cross the region of inflow of the same or another hyperbolic point. A point at which the inflow and outflow of a single hyperbolic point intersect is called a transverse h~moc1inic point. When crossing occurs from flows of two different hyperbolic points, it is called a transverse heteroclinic point. Homoc1inic and heteroc1inic intersections are identifying features of chaos. Positive
2 LITERATURE REVIEW
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Lyapunov exponents in sorne region of the flow, presence ofhomoc1inic or heteroc1inic points and presence of Smale horseshoe function are the criteria used to identify chaotic mixing (98). Cheng and Manas-Zloczower (99) analyzed chaotic features of distributive mixing in extruders. They found that the tangential twin-screw extruder was a better mixing device thanthe single-screw extruder.
, ELLIPTIC POINT
HYPERBOLIC POINT
ELLlPTlC POINT
Figure 2.5.3: Elliptic and hyperbolic points
Many researchers are interested in developing mixing devices that would incorporate chaotic features. Ling (100) designed enhanced mixing simulator (EMS) (Figure 2.5.4). In this device couette flow is perturbed laterally, so that uniform and quick mixing can be achieved. Numerical simulation of the mixing was also performed. Zumbrunnen et al. (101,102,103,104) designed three-dimensional chaotic mixing device for Newtonian fluids (Figure 2.5.5). A mixing chamber was fabricated with a cylindrical glass cavity and rotating upper and lower circular disks. Phase-space trajectories, retum maps and Lyapunov exponents were used to characterize the mixing process and to confirm chaotic behavior. Their results show that fine-scale, filamentary or lamellar structures can also be formed from initially coarse c1usters of added powders or solid fibres where continuous phase is liquid. Lamella thickness of 1 nm was achievable with just 20 periods (rotations) (102). They used this efficient mixing technique to enhance toughness in polymer blends (103) and to produce electrically conducting polymers with reduced carbon black concentration (105,106,107,108).
27
2 LITERATURE REVIEW
Figure 2.5.4: Enhanced Mixing Simulator (100)
,,- Locus of \ Point5 \
Hyp~rbolit~
::'OCUS
\
ot ft111ptic
t?Ol nt iII
Lccua ()t tlhpt.lc. PoinU
Figure 2.5.5: Schematic representation of chaotic mixing apparatus and loci of elliptic and hyperbolic points in chaotic mixing (101, 102, 103)
Solomon et al. (109) studied chaotic mixing of immiscible impurities in a twodimensional flow. Ganeshan et al. (110) studied chaotic heat transfer enhancement in rotating eccentric annular-flow systems. Statistical evolution of chaotic fluid mixing is given by Glimm et al. (111). Visualization study of three-dimensional chaotic mixing was carried out by Fountain et al. (112) by injecting fluorescent dye stream. lnamdar et al. (113) reported a new continuous flow blending process, where morphology develops
2 LITERATURE REVIEW
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progressively and more controllably by chaotic mixing. Patents by General Electric Company (114,115) employ chaotic mixing into single screw extruder by modifying the screw geometry. Kruijt et al (116) used mapping method to analyze flow in multiflux static mixers. They showed that the flow in static mixers has chaotic features. Ling (117) studied mixing in the Kenics static mixer from the viewpoint of chaos theory and suggested optimal element design. Kim and Kwon (118,119) designed a screw called 'chaos screw' for single screw extruder. They introduced spatially periodic barriers in the channel to break closed streamlines in regular flows, which induces the chaotic mixing. Utracki and Luciani (120) designed an extensional flow mixer (EFM), which enhances mixing by using extensional flow field. Paul et al. (121) developed a new dispersive and distributive static mixer called DDSM for compounding highly viscous material which incorporates extensional flow also. Manas-Zloczower and Cheng (122) obtained flow patterns in various types of processing apparatus based on 3D isothermal flow simulations. Dispersive efficiency was studied in terms of shear stresses and elongational flow components' distributions. Distributive mixing efficiency was characterized by length stretch distributions. Features of chaotic flow were also analyzed. 2.5.3 MELT PROCESSING NANOCOMPOSITES
In a thermodynamically compatible polymer/clay system, formation of nanocomposite can be greatly assisted by suitable processing parameters and mixing system. Paul et al. (123,124,125,126) demonstrated the importance of processing in the preparation of nanocomposites by melt compounding. They used two different clay treatments and four different types of extruders with multiple screw designs. The results (Figure 2.5.6) showed that the best delamination and dispersion resulted in medium shear mode of each extruder.
2 LITERATURE REVIEW
29
30 25 c::
.2
:i!O
~ 1'/)
15
~
i5
C"~
C 0 CoR V
CnRNI
0 53
:; W
1-
10 5 o~~~~~~~~~~~~~~~~~~
{l
20
40
60
80
100
120
140
160
180
Mean Residence Time Figure 2.5.6: Degree of dispersion (by Transmission Electron Microscopy) of various
nanocomposites of PA-6/Cloisite15A, plotted against mean residence time in sec (123). CnRI-Counterrotating Intermeshing TSE, CoR-Corotating TSE, CnRNI-Counterrotating non-intermeshing, SS-Single Screw Extruder, LS, MS, HS - Low, Medium, and High Shear
--------
Shear ,---,-_._-~~-->
---
- --
Shcaring 'of platclct stacks lcads to sl1Hlllcr tactoids
--
~~«(-----D1l1uswn
"laH'lets pet'! apart by combined diffusion/shear proccss
Figure 2.5.7: Mechanism of clay platelet exfoliation in the melt compounding of
nanocomposites (124, 126).
2 LITERATURE REVIEW
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Two mechanisms for the exfoliation of clay platelets were proposed (Figure 2.5.7). The first mechanism proposed that, the stacks ofplatelets are decreased in height by sliding platelets apart from each other. This process requires shear intensity. In the second proposed mechanism, the polymer chains enter the clay galleries pushing the ends of the platelets apart. This pathway does not require high shear intensity but involves diffusion of polymer into the clay galleries (driven by either physical or chemical affinity of the polymer for the organoclay surface) and is thus, facilitated by residence time in the extruder. It appears that both pathways are involved in different degrees, depending upon the the nature of the compatibilizer. The effect of matrix molecular weight on the exfoliation of clay was also assessed (126). It was seen that nanocomposites based on the higher molecular weight polyamide yielded superior composite properties with higher degree of clay exfoliation. This was attributed to higher viscosity and thereby higher shear stress with high molecular weight pol ymer matrix. Ko et al. (127,128) studied effects of shear on melt exfoliation of clay in polyamide matrix. Though the thermodynamics, especially the diffusion of polymer chain into the silicate layers played primary role, when shear stress was applied exfoliation took place in a much shorter time. Shen et. al. (129) found that in case of polyethylene oxide(PEO)/organoclay system melt intercalation was faster at higher temperature. Pressure of about 70 MPa was necessary but further increase up to 210 MPa had no· influence on intercalation process. Huang et al. (130) used different arrangements, such as fluted mixing elements, chaos screw and· Kenics static mixer with single screw extruder, to prepare polypropylene nanocomposites. The highest intercalation was obtained with the Kenics static mixer attachment. Dolgovskij et al. (131) prepared polypropylene nanocomposites using five different types of mixers: an internaI mixer, two lab-scale co-rotating vertical twin-screw mixers, co-rotating twin-screw extruder and a multilayer extrusion system. Unexpectedly the mixers with lowest shear, such as vertical twin-screw mixers and the multilayer extruder, showed the highest intercalation.
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Meharabzadeh and Kamal (132,133) reported preparation of nanocomposites based on PA-6/HDPE blends by melt processing. Nanocomposites were characterized using DSC, FTIR, SEM analyis of fracture surface, and mechanical property testing. It was observed that nanoclay reduces the size of crystals. Further they also reported (134,135)
production
of
PA-66/Clay,
HDPE/Clay
and
HDPE/PA-66/Clay
nanocomposites by melt processing. The results showed exfoliation in twin-screw extruder is enhanced by incorporation of mixing and shearing elements and high residence time. Utracki and Kamal (136) gave comprehensive review of the use of clays in polymeric nanocomposites in terms of availability, cost, and aspect ratio. Formation of nanocomposites in a hydrophobic, high molecular weight polymer and compatibilization and diffusion-controlled mixing are also discussed. A recent patent from Mitsubishi Gas Chemical Company (137) describes the melt process of producing polyamide nanocomposites with uniform distribution of clay fil 1ers. In this process, the polyamide and clay are tirst compounded by dispersive
mixing, and then further polyamide is added and mixed by distributive mixing. The process is carried out in a corotating intermeshing twin-screw extmder, so designed as to effect dispersive mixing and distributive mixing. The products have better mechanical and barrier properties, while retaining transparency.
2 LITERATURE REVIEW
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2.6 PHYSICAL PROPERTIES AND CHARACTERIZATION OF NANOCOMPOSITES Polymers are usually reinforced with fillers to improve mechanical properties. The degree of reinforcement depends on the rigidity and aspect ratio of the fiUer and the adhesive strength between the fiUer and the polymer matrix (138,139). Molecular composites with rigid rod-like polymer reinforcing fillers, such as liquid crystalline polymers, were also termed as nanocomposites (140). Polymer layered silicate nanocomposites are known to have excellent properties. They are light weight with high stiffness and strength. They exhibit outstanding barrier properties, without requiring a multiple layer design. They have high heat deflection temperature and high flame retardancy. Furthermore, the optical clarity of the matrix is retained to a significant extent. 2.6.1 MECHANICAL PROPERTIES
Nanocomposites consisting ofhighly anisotropie clay platelets dispersed in a polymerie matnx are of interest for many industrial applications. The platelet-shaped clay nanofillers have thickness of 1 nm. Their aspect ratios (diameter/thickness) range from 10-1000 and elastic (Young' s) moduli are ::::: 100 times those of typical thermoplastic. Furthermore, because the functional groups of sorne polymers such as nylon-6 may interact with the negatively charged silicate sheets, the affinity between filler and the matrix could be fairly good. The superior mechanical properties of polyamide nanocomposites originate from the microstructure and polar interactions, as well as the excellent mechanical properties of silicate sheet, such as high strength and tensile modulus (1). Empirical equations for estimating the effective moduli of unidirectionally aligned short fiber composites were proposed by Halpin, Tsai and Kardos (141). Similar empirical equations were developed by Padawer and Beecher (142), and Riley (143) for flake-like inclusions. Tandon and Weng (144) estimated the, effective moduli of a
2 LITERATURE REVIEW
33
unidirectionally aligned two phase composite, using Mori-Tanaka method (145). Brune and Bicerano (146) of Dow Chemical Company, used the micromechanics approach, to predict Young's modulus of nanocomposites reinforced by platelet-shaped fillers. Micromechanics considers simple geometries and assumes perfect adhesion between components. The model predicts three properties of nanocomposites: buckling of platelets under compressive loading, the reduction in the reinforcement efficiency of clay platelets as a result of incomplete exfoliation and the reduction in reinforcement efficiency as a result of deviation of platelet orientation from perfect biaxial in-plane. It was seen that platelet orientation has significant effect on Young's modulus. A minimum value exists when the particle symmetry axis is about 40° away from the applied normal stress and a maximum at 90°. Varlot et al. (147) showed that the orientation of the montmorillonite sheets and polyamide lamellae play a major role in determining the mechanical properties of nanocomposites. Masenelli-Varlot (148) et al. examined injection-molded PA-6/montmorillonite nanocomposite samples for mechanical properties and showed that montmorillonite sheets preferentially orient themselves around the injection axis. They used the Rule of Mixtures for series coupling and parallel coupling of platelets. The linear dependence of Young's modulus on the filler loading suggests a parallel coupling. Tensile tests show that the series contribution can be neglected. The system reacts as though the montmorillonite sheets were infinitely long. This behavior may be attributed to the microstructure of the matrix. The crystalline lamellae of the matrix (extra 'Y phase) grown from the montmorillonite sheets enhance the transfer of the load from one sheet to another. The largest reinforcement was obtained by compressing the nanocomposites along a direction parallel to the injection axis. Hui and Shia (149,150) derived simple equations for the effective moduli of unidirectional aligned composites and compared the results with established models. For a perfect interface, they proposed the following equations for the Young's modulus of composite containing aligned platelet inclusions:
34
2 LITERATURE REVIEW
Ec _
1
Em - 1-
(2.18)
:[~ + ç!Al
where the empirical constants
çand Ain the equation are defined as, 2
ç=f/J+
A=
Em +3(1_f/J)[(1-g)A -(gI2)] and 2 Er -Em A -1
g=1( A
2
(I-Il{ 3(A' +~;~~-2A'l
Ec, Em, El are the Young's moduli of composite, matrix and filler respectively, A is the aspect ratio defined as thicknessllength and f/J is the volume fraction of filler. Buxton and Balazs (151) used a three-dimensional lattice spring model to evaluate effects of filler geometry on the mechanical behavior of the composite. The increased reinforcement efficiency in nanocomposites, as a consequence of platelet exfoliation, was elucidated. All the theoretical models overestimate the strength of the composite. This difference could be accounted for by considering partial exfoliation of clay platelets, deviation from in-plane orientation and imperfect bonding between matrix and inclusions. 2.6.2 BARRIER PROPERTIES Nanocomposites have excellent barrier properties against oxygen, nitrogen, carbon dioxide, water vapor, gasoline etc. Dispersed platelets of the silicate sheet block the shortest path of gas molecules and force them to take a roundabout way. As a result, the permeation pathway is elongated. The reduced gas and liquid permeability makes them attractive membrane materials. Nielsen (152) calculated the tortuosity factor
'l"
for a plate-like material with
aspect ratio A, defined as the ratio of the length to thickness of filler particle as, (2.19)
2 LITERATURE REVIEW
35
where fjJ is the volume fraction of plate like filler. The relative permeability coefficient
Pc/Pp, where Pc and Pp are the permeability coefficients of the composite and pol ymer matrix, respective1y, is given by, (2.20)
A more rigorous analysis for random in-plane arrangement of plates is given by Cussler et al. (19, 153), leading to the following relation:
(2.21 )
where pc is a geometric factor that depends on the distribution of platelets. In either case, higher aspect ratio leads to a larger decrease in permeability. The barrier is sensitive to the degree of dispersion and the alignment of the plate1ets. Bharadwaj (154) modified the above equation equation by introducing an order parameter SB, defined as, (2.22) where () represents the angle between the direction of preferred orientation and the sheet normal vectors. P jPp, is given as, (2.23)
Messersmith and Gianne1is (155) studied water vapor permeability of polY(Bcaprolactone) layered silicate nanocomposite and showed that, permeability reduces linearly with increase in clay content. Akkapeddi et al. (156) showed that the permeability of nylon-6 improved significantly by incorporation of exfoliated nanoclay. Koch et al. (157), Mueller et al. (158) have patented nanocomposite formulations with improved barrier properties.
2 LITERATURE REVIEW
36
2.6.3 RHEOLOGY In case of nanocomposites, the melt rheological properties are dictated by a combination of the mesoscopic structure and the strength of the interaction between the pol ymer and the layered silicate. The steady shear response of layered silicate nanocomposites has important consequences on the potential processability of the materials. The viscosity of the nanocomposites is enhanced considerably at low shear rates, and increases with increasing silicate loading at a fixed shear rate (159). The nanocomposites display shear thinning behavior at low shear rates, where the pure polymer displays a shear independent viscosity. The same trends are also observed in linear dynamic oscillatory shear measurements. Huh and Balazs (160) investigated the rheological behavior of end-functionalized chains confined between two surfaces, using Monte Carlo simulations. For higher functionality, the chains form a transient network that effectively blocks the stretching of the chains under shear. The melt state viscoelastic properties ofpoly(ethyl-vinyl acetate) co-pol ymer nanocomposites were studied by Bhattacharya et al. (161) to examine the influence of clay in altering the flow properties. Higher vinyl acetate content resulted in higher exfoliation and significant increase in viscosity at low shear rate. Low shear rate viscosity also increased with increase in clay content. Lee and Han (162) assessed the . effects of polymer matrix/organoclay compatibility and the gallery distance of organoclay on the rheological behavior. They used poly(ethyl-vinyl acetate) (EVA), poly( ethyl-vinyl alcohol)(EVOH) and poly( ethyl-vinyl acetate- ethyl-vinyl alcohol) (EVAOH) copolymer nanocomposites with different clays. EVOH and Cloisite 30B, both have hydroxyl· groups and formed hydrogen bonds, which resulted in the increase in dynamic storage modulus when temperature was increased from 120°C to 180°C. Wagener and Reisinger (163) used rheology to compare the degree of exfoliation of poly(butylenes terephthalate) (PBT) nanocomposites. Apparent viscosity versus shear rate curves, taken under low amplitude were fitted· to power law expression to calculate shear thinning exponents. The shear thinning exponent was close to zero for the neat polymer and had higher negative values with increasing degree of exfoliation. Li et al.
37
2 LITERATURE REVIEW
(62) used rheology to study intercalation kinetics. Krishnamoorti and Giannelis (164) gave account of the rheology of poly(E-caprolactone) and nylon-6 nanocomposites. The storage and loss moduli increase at aIl frequencies with increasing silicate loading. At low frequencies, solid-like response was observed. Solomon et al. (165) did similar experiments with polypropylene nanocomposites. Clay loading changed the viscoelastic response significantly at low frequencies. Galgali et al. (166) also found a large increase in zero shear viscosity of polypropylene nanocomposites. Tanoue et al. (167,168) gave the account of compounding and characterization of polystyrene nanocomposites with special emphasis on rheological properties. The low strain frequency sweep showed that the storage and loss modulii increase with organoclay content. The extended low shear viscosity data were used to calculate the intrinsic viscosity and the aspect ratio. The melt compounded nanocomposites gave aspect ratio of clay as 16, whereas solution method resulted in aspect ratio of 269.
2.6.4 CRYSTALLIZATION The crystallization process affects polymer properties through the crystal structure and morphology established during the solidification process. Crystallization kinetics and the morphology of crystallized product are strongly influenced by cooling rate, system pressure and the presence of nucleating agents. In the case of nanocomposites the nature of the matrix polymer, the structure of clay and clay modifier, the degree of intercalation/exfoliation are additional factors that affect the crystallization process. Fomes and Paul (169) observed that very low levels of clay in polyamide-6 nanocomposites result in dramatic increases in crystallization kinetics relative to pure polyamide. The largest enhancement of crystallization rate was observed for high molecular weight polyamide. Mathias et al. (170) reported that clay induces the generation of 1 crystal phase in nylon-6, while maintaining the same percent crystallinity.
Kojima
et al.
(171,172)
studied
crystallization of polyamide-6
nanocomposite by annealing under pressure. After annealing under elevated pressure, the fraction of the 1- form decreased.
2 LITERATURE REVIEW
38
Kamal et al. (173) studied crystallization kinetics of polyamide-6 and its nanocomposites using a high pressure dilatometer. The Avrami exponents and Avrami rate constants were evaluated, for the a-form and the y-form crystals. The rate of crystallization was higher for nanocomposites. Infrared studies were carried out to confirm the formation of different crystal forms. Bureau et al. (174) also reported the presence of the y-form in polyamide-6 nanocomposites. Liu and Wu (175, 176) reported unusual crystallization behavior in polyamide-6 nanocomposites. In contrast to polyamide-6 and other semicrystalline polymers, increased cooling rates resulted in higher crystallinity in polyamide-6 nanocomposites. The y crystalline form was dominant in rapidly cooled nanocomposites. Borse et al. (177) studied isothermal crystallization kinetics of polyamide-6, polyamide-66, and their nanocomposites, using differential scanning calorimetry and a high-pressure dilatometer. Calorimetry data for all samples followed the Avrami equation. High-pressure dilatometry yielded two regions for polyamide-6 and its nanocomposites. The initial region was attributed to the formation of mixed y-form and a-form crystals and the later region to the a-form only. For polyamide-66 and its nanocomposite, only one region was observed. Wu and Wu (178) used X-ray diffraction and differential scanning calorimetry to investigate structural changes in polyamide-6, polyamide-66, and their nanocomposites. Addition of clay increased the crystallization rate. Formation of the y-form in polyamide-6 depended on the rate of cooling from the melt and the presence of clay. Higher cooling rate and the presence of clay resulted in higher y-form. Polyamide-66 and its nanocomposites showed only a-form. Liu et al. (179) also reported increased rate of crystallization in polyamide-66 with the addition of clay. However, they also reported the presence of the y-form in polyamide-66 nanocomposites. Wu et al. (180) synthesized and characterized nylon-12/clay nanocomposites. They reported that clay increased the rate of crystallization, but overall crystallinity was reduced. Crystallization behavior in polypropylene nanocomposites has also been studied. Ma et al. (181) found that the crystallinity of polypropylene/clay nanocomposites
39
2 LITERATURE REVIEW
decreased with the increase of the clay concentration and the spherulite size was smaller. The rate of crystallization was higher in nanocomposites. Maiti et al. (182,183) also reported similar results and found that the inclusion of clay results in higher 1- form in polypropylene/clay nanocomposites. Hambir et al. (184) found increased crystallization rate as well as higher storage modulus for polypropylene nanocomposites. Wu et al. (185) found that clay favored the formation oftrans
~-form
in syndiotactic polystyrene.
2.6.5 CHARACTERIZATION METHODS The distinguishing feature of polymer-Iayered silicate nanocomposites is their morphology, in which the size of the inorganic particles is comparable to the spacing between particles, both of which are nanoscopic. Thus, detailed characterization of the nanoscale morphology ofboth the layered silicate and the polymer is critical to establish structure-property relationships for these materials. Conventionally,
morphology characterization involves
wide-angle
X-ray
diffraction and transmission electron microscopy. Microscopy is useful to provide realspace information on the spatial distribution of phases and structures. In contrast, diffraction and scattering rapidly provide globally averaged information. Due to the periodic arrangement of the silicate layers, both in pristine and the intercalated states, the . choice of X-ray diffraction in determining the interiayer spacing is obvious (Figure 2.6.1). However, in the absence of registry, as in an exfoliated or delaminated nanocomposite, as well as in a disordered nanocomposite, XRD does not provide definite information regarding the structure of the nanocomposite (159). In order to provide quantitative information in XRD silent nanocomposites, TEM is an extremely useful technique. TEM also provides information on the homogeneity of mixing. The wave length  for X-ray machines using Cu-Kal is 0.1540562 nm. The equation for d-spacing is given as, 1/28 =0.11331d -0.00012773
(2.24)
2 LITERATURE REVIEW
40
The interlayer basal spacing is determined from the XRD spectrum as arbitrary intensity versus 2 (J. To provide the most accurate analysis of pol ymer-clay nanocomposites, it is sensible to use both XRD and TEM as complimentary tools (186).
• • • • • • Figure 2.6.1: Bragg's law for constructive interference
To examine the interactions between the clay platelets, the intercalating agent and pol ymer, Fourier Transfer Infrared Spectroscopy (FTIR) is used. Formation of different crystal structures such as a-form or y-form cart also be elucidated by FTIR. Solid-state IH , 13e, 15Nand 29Si Magic Angle Spinning (MAS) Nuclear Magnetic Resonance (NMR) spectroscopy at respective frequencies could be a useful tool for characterization ofnanocomposites. The Young's modulus was found to be proportional to the chemical shift of 15N in ammonium-clay complex (187). DifferentiaI scanning calorimetry _is a useful tool to study crystallization kinetics and melting behavior of different crystalline forms of polymer and nanocomposite. High-pressure dilatometry is used to study the effect of pressure on crystallization kinetics and crystallization at high pressure. Thermogravimetric analysis is used to study the degradation at elevated temperatures. Melt rheology is used to characterize processing behavior of the nanocomposite. It has also been used to study intercalation kinetics (62). Standard test methods that are used for testing mechanical properties of polymeric materials and plastics are suitable for nanocomposites.
3 OBJECTIVES
41
3 OBJECTIVES The incorporation of organoclay into thermoplastics by polymer melt compounding processes is a promising approach for forming nanocomposites that would greatly exp and the commercial opportunities for this technology. If technically possible, melt compounding would be significantly more economical and simple than in situ polymerization processes (125). Incorporation of clay into the polymer matrix results in significant increase of melt viscosity in the low shear region. In situ polymerization process involves reacting under slow agitation at low shear. The increase in viscosity with the addition of clay poses difficulty in processing the composite in the reactor. This has limited the range of clay concentrations used in nanocomposites synthesis by the insitu polymerization process. The dispersion of fiUer agglomerates can be achieved when the cohesive forces of the agglomerates are exceeded by the hydrodynamic separating forces applied by the matrix fluid (76-80). In case of intercalated organoclay, the amount of exfoliation appears to be strongly affected by the mixing conditions (123-126). The degree of dispersion is govemed by matrix viscosity, average shear rate and mean residence time during the mixing process. The general objective of the proposed research is to understand the fundamental issues involved in the melt exfoliation of layered silicate clay in nanocomposites and to design an appropriate processing system for this purpose. In particular, the foUowing . objectives have been identified.
•
To study the exfoliation and intercalation process further and de scribe quantitatively the basic physical and mechanical issues associated with the process. To develop working mathematical model to de scribe important issues required for process/equipment design.
•
To construct an experimental setup based on model predictions, for melt exfoliation of different clays in polyamide and polystyrene matrices.
•
To evaluate and optimize the extent of exfoliation obtained in the above systems.
3 OBJECTIVES
•
42
To characterize and evaluate the property enhancements in the nanocomposites produced.
•
To study the effect of nanoc1ay dispersion on the crystallization behavior of polyamide-6.
4 MODELING AND PROCESS DESIGN
43
4 MODELING AND PROCESS DESIGN In this section, the binding energy or adhesive force between the platelets in a clay particle is estimated and the corresponding shear force required to separate the platelets is compared to the hydrodynamic force available during processing. The results are subsequently used to design an efficient melt dispersion process.
4.1 INTERACTION FORCES BETWEEN CLAY PLATELETS Most of the models for agglomerate dispersion (188) are based on the assumption that the agglomerates are ruptured, when the stress on the agglomerates, caused by the surrounding polymer is larger than the adhesive forces between the agglomerate. The adhesion acts against the stress. The estimation of adhesion between platelets in a clay particle is based on evaluating the van der Waals adhesive forces, as described in the following discussion. 4.1.1 ADHESION BETWEEN PLATELETS IN A CLAY PARTICLE
The layered silicates used in nanocomposites belong to the same structural family of mineraIs as talc and mica (12). Their crystallattice consists of 1 nm thick layers called platelets (Figure 2.1.1, Figure 2.1.2). The lateral dimensions of these platelets vary from 30 nm to several microns. The average size of a clay particle is 6-13 microns (189). Stacking of the layers leads to a regular van der Waals gap between them called the interlayer or gallery. A clay particle may be treated as a bundle of fiat platelets stacked together, having sorne imperfections at the edges as shown in Figure 4.1.1. The gallery spaces in pristine clay have Na+ cations. Organically modified clay commonly has a quaternary ammonium modifier between the platelets, replacing Na+ cations. Adhesion is the interparticle force causing aggregation. It can be defined as the force required to pull two particles apart. The intermolecular interaction energy is considered as the sum of two contributions (190). One is due to the electromagnetic effects of electron clouds and leads to van der Waals interactions. The second is due to surface charge effects and leads to electrostatic interactions. These two contributions are
4 MODELING AND PROCESS DESIGN
44
additive. The van der Waals contribution is univers al and exists in ail systems. The electrostatic contribution depends on the polarity of the liquid media and density of ions. Usually, it is a short-range force originating from molecular forces such as van der Waals attractions of the particle surfaces (191), that is prevalent in dry particles and in polymeric systems. Maugis (13) gives the account of van der Waals forces between solids. Breitmeier and Bailey (192) measured interaction forces between mica surfaces at small separations in polar and non-polar liquids. The interactions are due to dispersion forces and the electrostatic attraction arising from the ions in the cleavage plane. The intermolecular attraction acting in the gap causes the thin sheets to be drawn towards each other. Although ionic forces make the major contribution to the total energy, after separation of 6-8 Â the van der Waals forces dominate and are most effective at large separations.
Figure 4.1.1: Model Clay Particle
Assuming that van der Waals forces are additive, de Boer (193) and Hamaker (194) computed, by simple integration, the energy and the interaction force between two parallel plates, two spheres, or a sphere and a plane.
45
4 MODELING AND PROCESS DESIGN
Figure 4.1.2: Van der Waals forces between solids (13).
Consider a molecule M at a distance d from an infinite haIf-crystal having n atoms per unit volume (Figure 4.1.2). The normal component of the force exerted on M by a molecule P situated at a distance r in a plane of ordinate z is given by London (195) as, 6C cos rplr7 . The contribution from the annulus of volume
dV = 2m 2 sincp dcp dz cos 3cp
(4.1)
IS
6C cos cp .
7
z
cos 7 cpndV=
12mzC . 5 cos 5 cpsmcpdcp z
(4.2)
where, C is a London constant (about 10-79 J m6) given as, C=
2
1 3a hvo (47Z"Eo 4
(4.3)
Y
a is the polarizability of the molecule and absorption band in the ultraviolet region,
E
Vo
is the frequency corresponding to an
is dielectric susceptibility, and h is Plank's
constant. By integration, the interaction force between a molecule and a half-space is
f r/2 coss tpsincp dcp
dz FMP{d) = 12mzC -s z
(4.4)
46
4 MODELING AND PROCESS DESIGN
(4.5)
i.e. Integration of this force over distance d gives interaction energy,
(4.6) Therefore, the interaction force between two half-spaces separated by a distance dis
() = f n:nC 2z
Fppd
2
n:n C
~dz=--3-
(4.7)
6d
The Hamaker constant A for two similar bodies is given as, Au two dissimilar bodies, it is A 12
=rln2C, and for
= n2n1n2C12 , which is about 10- 19 J for interactions in a
vacuum. Hence the force and interactions energy between two planes, per unit area, is given respectively by:
A Fpp (d)=--3 6n:d
(4.8)
Upp (d)-A 12n:d 2
(4.9)
Similar ca1culations can be done for different geometries such as spheres of equal and unequal radii (14). For the platelets in a clay particle (Figure 4.1.1), the geometry is that of two plates of equal thickness. In this case, if we denote the thickness of platelet as b, the adhesive force and interaction energy are, respectively, given as (14,15):
(4.10)
U__ AU (_1 + -
12Jl' d
2
1
(d + 2J)2
-
2
(d +
JY
J
(4.11)
In this case, Au is the Hamaker constant between the platelets of unmodified clay. Medout-Marere (190) measured values of the Hamaker constant for different
47
4 MODELING AND PROCESS DESIGN
materials by immersion calorimetry in apolar liquids. The Hamaker constant for Montmorillonite is given as 7.8 x 10-20 J. In the case of modified clay, the effective Hamaker constant between the platelets with an organic modifier between them can be written as (14,15), (4.12) where A22 is the Hamaker constant of the organic modifier. The Hamaker constant for saturated long chain hydrocarbons like tallow, is around 5xlO-2o J (14). The effective Hamaker constant between the platelets of organically modified clay is then, A 121
;:::
O.31xlO-2o J. It should be noted that the
effective Hamaker constant between two bodies is reduced significantly, due to the presence of a medium between them. In this case, it is reduced by a factor of 16. The interaction energy between the clay platelets was ca1culated by using equation 4.11. The interaction energy is negative for attraction. Figure 4.1.3 shows that the attractive energy between the platelets decreases, with an increase in gallery spacing. Attractive interaction is much higher between the platelets of unmodified clay. For the organically modified clay, the attraction between the platelets does not change significantly, and it is very small beyond a gallery spacing of 3 nm, which is usually the gallery spacing for modified clays. For unmodified clay, usual gallery spacing is 1 nm, and the attractive interaction is very high at this spacing. Figure 4.1.4 shows a plot of van der Waals forces between the platelets, with respect to the gallery spacing, calculated using equation 4.10. It is seen in this plot that the attractive forces between the platelets of unmodified clay are much higher than those of organically modified clay. We shall use these results in the next section to predict the forces required for breaking of clay particle into tactoids.
48
4 MODELING AND PROCESS DESIGN
0.0000
-
'"E
-0.0002
.......
3-
::>
-0.0004
>-
21 Cl) s::::
-0.0006
Cl)
s::::
0
~
-0.0008
~
Cl)
.E -0.0010 - - . - Unmodified Clay --9- Organoclay
-0.0012 -0.0014 0
2
4
10
8
6
Gallery spacing d (nm)
Figure 4.1.3: Attractive van der Waals interaction energy between platelets of
unmodified and organically modified clay platelets versus gallery spacing.
-
'"E ~
10" - - . - Unmodified Clay - -9- Organoclay
107
IJ...
('Il Cl) .... ('Il
~
106
::J
.... Cl)
c.. 105 t/)
Cl)
~
.E t/)
10'
ëii ~ G>
"C
103
s::::
('Il
> 102 0
2
4
6
8
10
Gallery spacing d (nm)
Figure 4.1.4: Attractive van der Waals forces per unit area between platelets of
unmodified and organically modified clay.
49
4 MODELING AND PROCESS DESIGN
4.1.2 BREAKING CLAY PARTICLES INTO TACTOIDS AND PLATELETS The process of dispersion of solids into polymeric melts can be described by three sequential steps (74): •
rupture of agglomerates of solids,
•
separation of fresh fragments away from each other,
•
distribution of the separated solids throughout the melt.
Depending upon the physical characteristics of the solidlliquid system and upon flow patterns within mixing equipment, each of these steps may be the most critical in determining the quality of the mixing operation. According to Rumpf (77), the limiting strength of an agglomerate is reached, when forces imposed by the normal stress equal the adhesion forces. Thus, agglomerate rupture occurs with the simultaneous collapse of interpartic1e links at the rupture surface. The rupture stress may be expressed as: (4.13)
where Fa is a partic1e-partic1e adhesion force and n is the average number of contact points per unit area in the cross section of the agglomerate. For spherical partic1es, Rumpf proposed the following expression for estimating n: n = 1.1(1-c)c-1d;2
(4.14)
where e is the agglomerate porosity and dp is the diameter of the particle. Kendall (196) argues that Rumpf overestimates the strength. He proposed a mechanism similar to the failure of brittle materials, assuming that the rupture occurs from the buildup of tensions in defects aIready present in the brittle solid. In the case of agglomerates, these defects would be small cracks within the structure. Once nuc1eated at these points, the cracks propagate through the agglomerate, consuming the energy necessary to create new surfaces. This requires much smaller energy consumption than implied in Rumpf' s model. He proposed an expression for the rupture stress, in terms of elastic modulus and crack length for spherical partic1es.
4 MODELING AND PROCESS DESIGN
50
Coury and Aguiar (76) used two different kirids of dry agglomerates of the same material, filtration cakes and tumbling drum granules, and evaluated their rupture stresses experimentally. These values were then used for comparing the theories of Rumpf and Kendall. The results indicate that neither theoretical approach could represent the two practical situations. They used the peeling model derived by Kendall (197) to estimate the width of filter cake removed from the clotho The theoretical calculations agree with the experimental results. Tadmor (73) analyzed dispersive mixing by modeling agglomerate as dumbbells consisting of two unequal beads connected by a rigid connector. The force in the connector was calculated, when the dumbbell is placed in a homogeneous velocity field of a Newtonian fluid. Manas-Zloczower and Feke (74,75) extended this model and showed that, even after long times in simple shear flow, ,all agglomerates were not broken. They further investigated the influence of agglomerate structure on the rupture process, using computer simulation (198,199). The results showed that the structure of the agglomerates had a considerable influence on the fracture behavior. The critical shear stresses that must be exceeded, in order to break down the agglomerates, were generally overestimated using the planar model. Following the approach of Powell and Mason (200), Rwei et al. (201,202) developed a model for the description of erosion. It was observed that the erosion process is more graduaI and initiates at lower applied shear stresses than rupture. The erosion process is characterized by the continuous detachment of small fragments from the outer surface of the agglomerate. The strength of the flow field does not affect the kinetics of the dispersion process. These results are similar to those obtained by Paul et al. (123,124,125,126) for the exfoliation of nanoclay in polyamide-6 matrix discussed in section 2.5.3. The process of erosion is very much similar to the peeling mechanism proposed by Paul et al. (Figure 2.5.7). A quantitative model similar to erosion will be presented, in the following section, for the exfoliation of nanoclay in polymer matrix. Exfoliation of nanoclay can be schematically represented, as shown in Figure 4.1.5. When polymer chains have strong affinity (i.e. tendency to form hydrogen bonds)
4 MODELING AND PROCESS DESIGN
51
towards the organic modifier between the platelets, polymer chains will enter the gallery spacing. This initiates peeling of platelets from clay particles at angle
f).
In the absence
of strong affinity, the platelets might be sheared, and the peeling angle is 0°. We calI this as 'lap shearing', which can be considered as a special case of peeling.
--------~~
------------------------.~ F
(b)
(a)
Figure 4.1.5: Schematic representation of exfoliation process (a) Peeling (b) Lap Shearing
4.1.2.1 PEELING MODEL Peeling (Figure 4.1.5(a)) can be modeled as shown in Figure 4.1.6. The width and thickness of the platelet being peeled are b and J, respectively. The plateleUs pulled by force F at an angle
f)
from clay particle. The peeled length of platelet is 1. The adhesive
fracture energy G per unit crack extension, may be derived from the energy balance (83,84), such that:
G
= l(dU b
ext _
dl
dUs _ dUk dl dl
_
dU d ) dl
(4.15)
where Uext is the external work, Us is the stored strain energy, Uk is the kinetic energy, and Ud is the energy dissipated during bending or stretching of the peeling arm.
52
4 MODELING AND PROCESS DESIGN
Width b
F
Figure 4.1.6: Peeling Model
If the peel rate is slow, increments of kinetic energy are assumed to be
negligible. If the peeling angle and thickness of the tape do not vary, the energy stored in bending remains constant, and its contribution to G is negligible. Under the action of force F, the platelet of width b extends by !ll. If Young' s modulus of platelet is E, then:
E=~
(4.16)
!ll = Fl E&
(4.17)
&!ll
and
The stored strain energy is thus,
=~F!ll=
U s
and
dUs dl
2
2
F l 2E&
F2 2E&
--=--
(4.18)
(4.19)
Compared to the original position (before peeling), the load F has moved by the distance (l + !ll - l cos () ), and the external work or its potential energy is gi ven as:
4 MODELING AND PROCESS DESIGN
and
53
Uext = Fz(I-COSB+
~l)
(4.21)
Uext = Fz(I-COSB+
:&)
(4.22)
t dU F ) dt = F ( l-cosB+ E&
(4.23)
Combining equations 4.15, 4.19 and 4.21: dU s )= F(I-COSB+~)- F 1 dl b E& 2E&2 2
G=.!.(dUext b dl F
_
F2
G=-(I-cosB)+ 2 b 2E&
(4.24)
(4.25)
It can be noted that the adhesive fracture energy given by equation 4.25 is
independent of the length of the platelet but depends on its width. This means that the energy required to start peeling is independent of the area of the platelet but depends upon its width and thickness. Equation 4.11 in Section 4.1.1 gives the attractive interaction energy between the platelets and equation 4.25 gives the adhesive fracture energy between the platelets. At equilibrium, both terms will be equal, and peeling or exfoliation will result only if the adhesive fracture energy is greater than the attractive interaction energy. With equations 4.11 and 4.25, it is possible to estimate F, the shear force required for breaking clay particles into tactoids or for the exfoliation of clay particle. This shear force can be compared with the available shear force during processing.
4.1.2.2 LAP SHEARING MODEL When polymer chains have low or no affinity towards the organic modifier between the gallery spaces, polymer chains may not diffuse between the platelets. In this case, it can
54
4 MODELING AND PROCESS DESIGN
be considered as lap shearing. A special case of peeling at 0 0 is termed as 'lap sherujng', as shown in Figure 4.1.7. In this case adhesive fracture energy is given by: . F2 G=--
(4.26)
2Eai
Width b
Crack Initiates
F
Figure 4.1.7: Lap Shearing model
4.1.2.3 MODELING INTERCALATION PHENOMENON Kendall (203) modeled dislocation in lap joints similar to those shown in Figure 4.1.7, for thin films in the presence of small transverse force T on the extremity. The dislocation proceeds as shown in Figure 4.1.8. The transverse load represents the pressure due to the matrix polymer during processing.
l Tt (a)
(b)
-
•
o,,~----------~ ~ - 1...._ _ _ _ _ _ _ _ _ _---'
x
(c)
(d)
Figure 4.1.8: Dislocation in lap joint
... ... ...
4 MODELING AND PROCESS DESIGN
55
When a lap joint is pulled, the free end of platelet lifts up and turns by an angle (J.
Kendall computed it to be: 1
B-
-1+J8
~ 3F
(4.27)
bEô -
Applying now a small transverse force Tf on the extremity of the freed platelet of length c, it takes the shape defined by the following equation: d2 y _ dx 2 -
Mf _ Tf{c-x) El El
-
(4.28)
where Mf is the bending moment and 1 is the moment of inertia, 1 =bf53/12. Integrating twice, with the conditions dy/dx =tan (J and y =0 at x =0, we get: y
The extremity x
=
Tf (x 3 -3cx 2 )+ xtanB. 6EI
(4.29)
=c touches the stretched band for a length c =Cr given by,
CT =
tEl 0 tan
.
(4.30)
Tf
or, with tan (J :::::: (J,
C2 T
Ebô
3
/3F
1
l+J8Vbi8
= 4Tf
(4.31)
The introduction of transverse force Tf modifies the value of G given by equation 4.26, due to the storage of elastic energy in the bent band. We must subtract potential energy of force Tf and the elastic energy in the bend from equation 4.26. The potential energy at x = c, using equation 4.29 can be given as, .
Up
Tf
2 3
C
= Ty = - - - + c tan B 3EI
(4.32)
and the elastic energy as, .
2
('Mf
U = lJ ~= E
T
2 3
c 6EI f
Hence replacing 1 =bf53/12, the total additional energy is,
(4.33)
4 MODELING AND PROCESS DESIGN
2T
UT = -
f
56
2e 3
Ebo
+ eTf tan ()
3
(4.34)
The rate of change of UT with respect to freed length of platelet e is given as, (4.35)
de
By the energy balance similar to equation 4.15, now equation 4.26 can be written as, 2
F _.!.(dU
G_ - 2b 2 E8
b
T )
de
(4.36)
(4.37)
(4.38) For a given value of F, the crack will move initially faster or slower according to the sign of Tf, but, in any case, it will accelerate as observed by Kendall (203). When the peeled band of the platelet reaches the length CT and touches the back to the substrate, a dislocation is formed, and the cracking slows down and then stops. It was shown that the contact is well established only after the length of Cs =2CT. From this point, the platelet will remain stuck to the substrate. Further it was shown that, G=
F 2 C2 S 2b 2 Eo e 2
(4.39)
If pulling continues, the platelet, which is in contact, lifts up once more and then
touches the substrate again, and new dislocation is formed. Thus a succession of crack propagation and resticking phases take place. Experimental results by Kendall are in good qualitative agreement with this theory. This phenomenon might result in increasing the gallery spacing, which we term as intercalation. Polymer/organoclay systems with low affinity may intercalate, but they are incapable of exfoliating.
4 MODELING AND PROCESS DESIGN
57
4.2 MODEl APPLICATION AND RESUlTS A clay particle, under the influence of shear in a polymer melt, may be schematically represented as shown in Figure 4.2.1. The lateraI dimensions L and b are the length and the width, respectively, of the particle or platelet. In the remaining discussion, we assume that L and b are equal. The thickness of an individuaI platelet is 1 nm. The thickness of the tactoid being peeled from the particle surface is
~.
In the following
caIculations, L and b vary from 10 to 10000 nm (0.01 to 10 micron), and ~ varies from 1 to 5000 nm (0.001 to 5 micron). The gallery spacing d varies from 1 to 10 nm. The peeling angle () varies from 0 to 10°. The Young' s modulus of montmorillonite clay platelet is taken as 170 GPa (143). L
)
-+F
Figure 4.2.1: Schematic representation of a clay particle consisting of layers of stacked
platelets.
Using Equations 4.11 and 4.25, the force required to peel the tactoid ofthickness Ô and area bL, from the surface of the clay particle, can be calculated. This force is
directly proportional to the Hamaker constant between the platelets, elastic modulus of platelets, width b and thickness Ô of tactoid, and inversely proportionaI to the gaIlery spacing d. When divided by the area of the tactoid (bL), it gives the shear stress required for exfoliation. Unless mentioned otherwise, the peeling angle is taken as 0°. Clay particles can be broken into tactoids by the mechanisms shown in Figure 4.2.2. In Figure 4.2.2(a), it is shown that the particle breaks into halves consecutively,
58
4 MODELING AND PROCESS DESIGN
and, thereby, the size is reduced. In Figure 4.2.2(b) size reduction is by surface erosion or peeling of tactoids from the surface.
(a) particle breaking consecutively into halves ... "·:··~······S'·····~······5·····§·-·5"":5,,,,,§,,:,,:-E.::,,S;::,,,,.~: .. ··
(b) reduction of particle size by surface erosion or peeling Figure 4.2.2: Mechanisms of breaking clay particle into smaller tactoids
To illustrate the frrst mechanism, consider the breaking of a clay particle of thickness 1000 nm into two halves i.e. 8, the thickness of tactoid, taken as 500 nm. The width and the length of the particle are also 1000 nm each. Figure 4.2.3 shows the shear stress required for this process. The dotted line shows the maximum shear stress that may be available in extrusion processing. Hiemenz and Rajagopalan (204) indicated that typical shear rates in polymer extrusion are 1 to 100
S-l.
The viscosity of most polymers
during extrusion is between 1000 and 2000 Pa.s. This suggests that the maximum available shear stress is 2 x 105 N/m2 • It can be seen that the shear stress required to break the organoclay particle is much smaller than the shear stress required to break
4 MODELING AND PROCESS DESIGN
59
unmodified clay particles. However, in both cases, the required shear stress is much higher than the available shear stress. This means that the clay particle cannot be broken by the mechanism shown in Figure 4.2.2(a).
106
,----------------------------------------,
......... Unmodified Clay --9- Organolay
.......
'"
~ ln ln
107
~
êi5.... CIl
Ol
.r:.
en
"C
~
'5
106
0Ol
a:
Maximum Available Shear Stress during Processing 105+-------~------~------~------~------~
o
2
4
6
8
10
Gallery Spacing d (nm)
Figure 4.2.3: Shear stress required to break 1000 nm thick clay particle into two halves.
We shall now consider the mechanism shown in Figure 4.2.2(b) to calculate the shear stress required to remove tactoids of variable thickness from the particle surface. The dimensions of the clay particle are the same as before. The thickness
~
of the
tactoids varies from 1 nm to 5000 nm. The gallery spacing for unmodified clay and organoclayare 1 nm and 3 nm, respectively. Again, the results from Figure 4.2.4 show that unmodified clay particles cannot be reduced in size by surface erosion. For an organoclay particle, however, tactoids of 15 nm or less in thickness can be removed from the surface of the clay particle during extrusion processing. This shows that the likely mechanism of size reduction in organoclay is that of surface erosion or surface peeling, as shown in Figure 4.2.2(b).
60
4 MODELING AND PROCESS DESIGN
109~~~--=-~~--~--~~----__~~~--~~~
-+- Unmodified clay (b=200) ~
........ 108
'"E
--
Unmodified Clay (b=500 nm) Unmodified Clay (b=1000 nm) -~ Organoclay (b=200) --8- Organoclay (b=500 nm) --6- Organoclay (b=1000 nm)
~ en en ~
-ca
107
en
106
en
a> .s::.
"0
~
'5 ca>
a:
105
10
100
1000
Thickness ôof peeled tactoid (nm)
Figure 4.2.4: Shear stresses required for peeling of tactoids of variable thickness from
the surface of the clay particle with surface area 200x200 nm2 , 500x500 nm2 and 1000x1000 nm 2
Further we shan consider the dependence of the required shear stress on the surface area of the particle or the platelet. It can be seen from Figure 4.2.4 that, when the width b of the particle is varied from 200 nm to 1000 nm, the shear stress required for removing the tactoids from clay particles decreases. For higher surface area, lower shear stresses are required for exfoliation, since the effective shear force available is higher. With the available shear stress, unmodified clay particle cannot be reduced in size by surface peeling. For an organoclay particle of 500 nm width, tactoid of 5 nm or less in thickness cau be removed from the surface during extrusion processing.
61
4 MODELING AND PROCESS DESIGN
1010 - . - Unmodified Clay (15= 50 nm) Unmodified Clay (15= 5 nm) Unmodified Clay (15= 1 nm) -~ Organoclay (15= 50 nm) Organoclay (15 5 nm) Organoclay (15 = 1 nm)
-À-
109 ........ '"E
~
10S
-
107
en en ~
en
=
•
==Et;~~~~FHtii]~~~,~ .
iCI>1
..r:::
en 10S "0
~
"S c-
H)5
'
· ·· . · ... ~th~set~=4tc
Maximum available,shear stress....; .... .:... . "
CI>
a::: 10"
:, &-A'
, .....
103 10
100
1000
10000
Width b of clay particle (nm)
Figure 4.2.5: Shear stresses required for peeling clay tactoids of different .thickness and area from the clay surface.
Figure 4.2.5 shows the shear stress required to peel 1 nm thick clay platelet as weIl as tactoids 5 nm and 50 nm thick from the surface of the particle for variable surface area. In this case aIso, the required shear stress reduces with the increase in width b, and in tum the surface area of the tactoid, since the available effective shear force increases with increase in surface area. In the case of unmodified clay, a platelet can not be peeled off, unless its width is more than 3000 nm (3 micron). On the other hand for an organoclay, the width b of the platelet being peeled could be as smaII as 150 nm. GeneraIly, the clay particles have lateraI dimensions of 500 to 1000 nm. In that case, platelets cannot be peeled from an unmodified clay particle. However, it is possible to peel platelets from organoclay surface. The shear stress required for peeling increases with the increase in tactoid thickness and decreases with the surface area of the particle. For a compatible system of polymer/organoclay, where there is tendency to form hydrogen bonds between the organic modifier and the polymer matrix, polymer chains have strong affinity towards the organic modifier. In this case, the peeling of platelet at
4 MODELING AND PROCESS DESIGN
62
sorne angle may take place as shown in Figure 2.5.7 and Figure 4.1.5. The shear stress required for peeling can be calculated in terms of the peeling angle. Figure 4.2.6 shows the shear stresses required to peel 1 nm thick platelets from the surface of clay particles of widths 100, 500 and 1000 nm, at different peeling angles. The shear stresses required to peel a 1 nm platelet from organoclay paiticle are below the shear stresses available during processing, and, with the increase in peeling angle, the required shear stresses decrease substantially. It can be seen that, to initiate peeling in unmodified clay, the peeling angle needs to be above 6° for a platelet of width 1000 nm. 107
-
106
'".€
~
~-------------=------~ Maximum available shear stress
-
~
en 105 en
CD .... ....
( J)
ct! CD
""
'&.. ~
'6> .......... â
(J)
....
·5
--
- ~----
..... "s~~-~-_
103
--~-
---+-
---- -v--__ _
-----A-. --
- .....(9.._ Ir-----------------~-~
Unmodified Clay (b=100 nm) ----E>- __ ----.t.-- Unmodified Clay (b=500 nm) -+- Unmodified Clay (b=1 000 nm) --9-. Organoclay (b=1000 nm) -~. Organoclay (b=100 nm) -...é.-. Organoclay (b=500 nm)
CT CD
CI:
..........,..... ~
"U.. ~
1Q4
..c:
"0 CD
~~
--0.... """"
102
101 0
2
3
4
5
6
~
==--A----__ - -
7
e- - - - __ _
8
9
10
Peeling Angle (0)
Figure 4.2.6: Shear stress required for peeling of 1 nm thick platelet fram the surface of a clay particle at various peeling angles.
The clay particle consists of stacked platelets. These platelets are not perfectly stacked, and there are defects at the edges. These defects may result in peeling of sorne platelets àt an angle. Although this phenomeIion may result in removal of sorne of the platelets from the unmodified clay particle, this would be generally a very small fraction.
4 MODELING AND PROCESS DESIGN
63
Since polymer chains have affinity towards compatible clay surfaces, the polymer chains entering the organoclay galleries may initiate peeling at any angle above 0°. In conclusion, the mechanism of peeling and lap shearing is similar to the surface
erosion. It has been demonstrated above that the peeling of the platelets from unmodified clay requires high shear stresses that may not be attained in polymer processing equipments. However, peeling of organoclay particle requires much lower shear stress, which is within the reach of the extrusion process.
4.3 PROCESS DESIGN The process of dispersion of solids into polymeric melts depends on the rupture of agglomerates of solids, separation of fresh fragments away from each other and the distribution of the separated solids throughout the melt (74). In order to design an efficient melt exfoliation process, it is necessary to consider the above three requirements. The aim is to design a process, which provides sufficient shear stress and residence time to remove the platelets from clay particles, separates the platelets from each other and distributes them efficiently through out the polymer matrix. It has been indicated in last section that the exfoliation process has similarities to
the surface erosion. Manas-Zloczower et al. (198,201) proposed a kinetic model for agglomerate rupture by surface erosion and derived the following relationship, ln R(t) = Kt y Ro
(4.40)
where R(t) is the radius of the agglomerate at time t, Ro is the original radius, and y is the applied shear rate. They suggested that the constant K depends on the flow geometry, the applied shear stress, and the cohesive strength of the agglomerate. Thus, the process of agglomerate dispersion by erosion was time dependent. Similarly, the degree of exfoliation is expected to depend upon processing time. In fact, Paul et al. (123) have shown, using different extrusion systems, that the degree of exfoliation does depend upon the residence time in processing equipment (Figure 2.5.6).
4 MODE LING AND PROCESS DESIGN
64
Cho and Kamal (205) proposed a model, based on hydrodynamic considerations, for the exfoliation process. They showed that, for a simple shear flow, besides the factors such as applied shear stress and viscosity of the polymer matrix, dispersion depends upon the angle that the plate1et makes with the flow direction. The stretching stress, experienced by the connector between the centers of two adjacent platelets, is highest at an angle of 45°. A similar conclusion was made by Tadmor (73) for the dispersion of agglomerates made of spherical particles, in simple shear flow. Thus, to enhance dispersion, clay particles in flow stream must change its direction frequently. Every time the angle is 45°, the force in the connector between two particles is highest and the possibility of particle rupture or plate1et separation is highest. Chaotic mixing produces efficient distributive mixing (Section 2.5.2). Chaotic mixing has been used by Zumbrunnen et al. (l 02-1 08) as an effective distributive mixing method to produce e1ectrically conducting polymer systems. In summary, to design an effective melt dispersion process and to produce nanocomposites with significant property enhancement, the following factors are very important: •
Compatibility between clay and the pol ymer matrix;
•
Sufficient shear stress;
•
Sufficient residence time for exfoliation;
•
Frequent changes in the flow direction;
•
Efficient spacial distribution by chaotic mixing.
Schrenk et al. (206,207) described the multilayer co-extrusion method and apparatus in their patents. They used the ISG (Interfacial Surface Generator) static mixer, in combination with extruders. The pol ymer melt from the extruders passes through the static mixer and then into the mold. In this work, we propose to use a similar mixing system, consisting of a twin-screw extruder and static mixer, for the purpose of achieving melt dispersion of nanoclay in a polymer matrix. We shall analyze this system in terms of the above mentioned factors.
4 MODELING AND PROCESS DESIGN
65
Static mixers are commonly used to enhance distributive mixing in the compounding of polymers. These mixers consist of a series of motionless inserts installed in pipes, whose purpose is to divide and recombine the fluid streams sequentially, until good mixing is achieved. Static mixers produce plug flow with complete uniformity at a particular cross section, with little axial mixing. The static mixer is distinguished by its ability to produce a uniform controlled degree of shear and flow. The commercial available static mixers are namely: Kenics (Chemineer), Inliner (Lightnin), Cleveland (EMI) , SMX (Koch-Sulzer), LPD and ISG (Ross Engineering Inc.). The pressure drop, the efficiency of mixing, the viscosity of the fluid and the design determine the selection of the type of the static mixer. 4.3.1 SELECTION OF THE STATIC MIXER The frrst criterion for the selection of a static mixer for the melt dispersion of nanoclay in a polymer matrix is that the shear stress available in the mixer should be sufficient to separate the platelets in the clay particle. Figure 4.3.1 shows the static mixers that are generally used in industry (208). The Kenics static mixer from Chemineer (Dayton, Ohio, USA) is constructed of right and left-handed twisted plates. The elements are altemated and oriented so that each leading edge is at 90° of the trailing edge of the next one. The Inliner mixer from Lightnin (Rochester, NY, USA) is formed of a combination of individual segments positioned to divide and rotate the flow altemately 180° right and then left. The principle is similar to that of Kenics. The LPD mixer from Charles Ross & Son (Hauppauge, NY, USA) consists of series of semi-elliptical plates perpendicular to each other and assembled at 45° and 135° from the flow axis. The Cleveland mixer from EMI (Clinton, CT, USA) is similar to the LPD mixer, except that triangular plates connect ellipses.
4 MODELING AND PROCESS DESIGN
66
The SMX mixer from Koch-Sulzer/Sulzer Chemtech USA Inc. (Pasadena, TX, USA) is made of stacked lamella, such as an intricate network of flow channels is formed. Again, each element is rotated at 90°, one to the next. The ISG (Interfacial Surface Generator) static mixer from Charles Ross & Son (Hauppauge, NY, USA), consists of mixing elements shaped such that two adjacent elements form a tetrahedral chamber. Four holes bored through each element provide the flow paths. The performance of a mixer can be appraised with several criteria. Historically, static mixers have been frrst compared on the basis of the pressure drop that they generate for a given flow rate and tube radius. Mixing efficiency parameters have then been introduced, based on the variance of concentration, the residence time distribution, the chaotic nature of the flow and the type of deformation inducing mixing. The pressure in a static mixer is related to its energy consumption. This energy is transformed into shear stress, which is an indication of its capability to break the agglomerates.
4 MODELING AND PROCESS DESIGN
67
(b)
(a)
(d)
(e}
{f)
Figure 4.3.1 : Types of statie mixers (208) (a) Kenies, (b) LPD, (e) Inliner, (d) Cleveland, (e) SMX, (f) ISG
4.3.1.1 PRESSURE DROP AND SHEAR STRESS IN STATIC MIXERS
The energy consumption in a static mixer can be related to the pressure drop it generates for a given flow. The pressure drop correlations are generally presented in three different ways. Firstly, a Z factor has been introduced, which is defined as the ratio of the pressure drop through the static mixer ll.P to the pressure drop in the empty tube, Met. of the same length, namely:
Z= M ~t
(4.41)
4 MODELING AND PROCESS DESIGN
68
Another parameter is the friction faetorf/2 or the Newton number, Ne: (4.42) where Di is the inside diameter and l is the length of statie mixer, v is the velocity and p is the density of the fluid. Sometimes, the produet Ne. Re = Kp is used by analogy with the power constant of meehanieal agitators, where Re is Reynolds number defined as Re=pvDlp. The pressure drop is then obtained by the following formula (Q is the volumetrie flow rate): AP = K f.1vR P
D.2
=K
4Q R pf.1 7lD.4
1
(4.43)
1
Rauline et al. (208) gave a detailed survey of Z and Kp values found in literature for different statie mixers. A summary of their findings is presented in Table 4.3.1. It ean be seen that the values of the Z factor, as well as the power factor Kp, are highest for the ISG statie mixer. The Kenies, Inliner and LPD statie mixers have similar Z and Kp values whieh are lower than those for the SMX and ISG statie mixers. Here, it ean be qualitatively eonc1uded that shear stress generated in the ISG static mixer is the highest, among the above mixers. Rauline et al. further eomputed Kp, AP and mean shear rate values for different statie mixers for a Newtonian fluid, in the laminar flow regime (Reynolds number 5xlO-4) for a flow rate of 15 kg/hr. Their eomputed values are shown in Table 4.3.2. It ean be seen that the eomputed shear rate is also highest for the ISG statie mixer. The shear rate of 105 rate of 100
S-l
S-l
for the ISG statie mixer is comparable to the shear
indieated for extrusion processing (204) of polymer melts. Using these
values of shear rate for different statie mixers and the viseosities at these shear rates, the available shear stresses in different statie mixers ean be estimated. The rheological behavior of PA-6 used in the present study at 240°C is shown in Figure 4.3.2. The shear stress ean be estimated using these data. Figure 4.3.3 shows the comparative estimate of shear stresses available during proeessing using different statie mixers, and the shear stresses required for the delamination of unmodified and organoc1ay partic1es. It ean be seen that the polymer proeessing with extruder or statie
4 MODELING AND PROCESS DESIGN
69
mixer does not generate sufficient shear stress to exfoliate unmodified clay particles. However, the shear stress generated in the extruder and ISG statie mixer is sufficient to delaminate the taetoids of thiekness 5 nm or less from the organoclay particles. The shear stresses generated by other statie mixers, sueh as the SMX, Cleveland, LPD, Inliner and Kenies, are not sufficient for delamination of platelets from organoclay particles. From this figure, it is clear that the ISG statie mixer is the most likely deviee to aehieve melt dispersion of nanoelay in the polymer matrix. Table 4.3.1: Characteristic values for different static mixers· (208)
Type of Statie Mixer
Z
Kp
Emptytube Kenies Inliner LPD SMX ISG
1 5-8 7-9 6 10-60 250-300
32 140-260 240-290 190-280 600-1900 6000-9600
Table 4.3.2: Computed characteristic values for static mixers (208)
Type of Statie Mixer
!li> (MPa)
Kp
Mean Shear Rate (S-I)
Kenies Inliner LPD Cleveland SMX ISG
0.71 0.70 0.75 0.51 1.8 20
255 300 225 190 1120 8460
10 10 8.8 8.4 21 105
70
4 MODELING AND PROCESS DESIGN
Ci) ai
!!::..
~
1500
r---------------------,
1000
+ ......................:....... :...............:..-.... , ..........:.......oc ....... :...........•........ .-.. ,
................. :....... :..................•
900 800
UI
8 UI :>
700 600 500
400 f - - - - - - ' - - - - ' - - - - - i - - - ' - - - ' - - ' - - - + - - - ' - - - ' - - ' - - - j 10 100 1000
Shear rate (S·I)
Figure 4.3.2: Capillary rheometry data for the viscosity of PA-6 at 240°C
107.---~---~~_,-~~--_-_~-_--~
:
1!:
i::
-
N
~ 106
-
--..."...,,"'"
__
--
:
~
__
L~~~~W
____ L __
~
__
;
j
L~:
: SMX $tatie MIxer ----~--~-~-~~-~~~-----~--~-~-~~~t~ pther Staiè Mixers ! • - -G- -
Unmodified Clay Organoclay
10
100
Thickness 80f peeled tactoid (nm)
Figure 4.3.3: Shear stresses encountered during processing polymer melt with different static mixers
4.3.1.2 EXTENSIONAL EFFICIENCY AND STRETCHING IN STATIC MIXERS Manas-Zloczower (209) defined extensional efficiency ae as:
4 MODELING AND PROCESS DESIGN
71
(4.44)
where
I~
and
I@
are the norm of the rate of deformation tensor and the norm of the
vorticity tensor, namely:
and
t= ~ ~V +ŒVY)
(4.45)
~ {yV -(ITY)
(4.46)
lü=
The norm is defined as:
I~=~ 1ft}'
(4.47)
'l
The coefficient ae is equal to 0 for pure rotation, 0.5 for simple shear flows and 1 for pure extension. Rauline et al. (208) conducted a comparative study of extensional efficiencies of different static mixers. Their results are shown in Figure 4.3.4. The abscissa represents the dimensionless length of the mixer. The ordinate is the average extensional efficiency at the cross section plane located at a given abscissa. The average values of extensional efficiencies lie between 0.53 and 0.57 for all the static mixers. The average value is not the only way to analyze this criterion. To break agglomerates, high extensional efficiency is needed. It can be seen that, in the case of ISG mixers, the highest extensional efficiency of 1.0 is attained at the entrance and the exit of the mixer element. This might result in the breaking of clay particles into tactoids. The authors (208) also gave the stretching distribution for different static mixers. The stretching is defined by Ottino (210) as:
=
 s
lim
IIdxll
IIdXll~o IldX11
(4.48)
where dX is an initial segment and dx is the same segment at time t. The stretching values obtained for Kenics, Inliner, LPD and Cleveland static mixers range between 0.18 and 0.57. The stretching values for SMX and ISG static mixers are 4.2 and 3.8,
72
4 MODELING AND PROCESS DESIGN
respectively. The distribution of stretching values shown in Figure 4.3.5 indicates that the ISG static mixer is more effective. 0.65
~~~,""----~_._-,-_.",
rT---;:::==;-';7' {HiS Ci
0.6
r 0.6 1· 055
l
l /'1' 05
,/ 1 ;...--..----.,.........--,---.-,---.....1
tu
0..10.6
O.S
0,2
(),M ..-----------~
0,4
0.6
üJl
r----------~,
Qi
ù.53
0.6 G.55
~~t"l"-~TI\
+----.------------.--]
o
0.2
OA
0.6
DJJ
OA
!16
(Ul
1
{lM
j i"
etb
~
055
i
-Y
V'l°.5
ü5 " G.45
CHtS
o
0:2
OA
0.6
0.8
°
0,2
OA
0.6
0.8
Figure 4.3.4: Extensional efficiencies of static mixers (208)
10%
ISO: 3.8
8%
4%
2% O%+-------~~~~~-~b-~~.".,.~~~.~~~~~-
"5
(1
:5
ln ft"
10
Figure 4.3.5: Stretching distribution for SMX and ISG static mixers (208)
4 MODELING AND PROCESS DESIGN
73
4.3.1.3 DISTRIBUTIVE MIXING IN STATle MIXERS Only a combination of folding and reorienting changes the efficiency of the distributive mixing from a linear dependence on the total shear to an exponential dependence. In a chaotic flow, stretching Âs increases in an exponential way. Stretching and Lyapunov exponent lÎ are related through the following relation:
& = lim 10gÂs L
t~~
t
(4.49)
The Lyapunov exponent is a signature of chaos, and it is also a characteristic of the mixing efficiency. The higher the exponent, the more chaotic the system, and more efficient is the mixer, in principle. The Lyapunov exponent for the Kenics mixer was estimated to be about 0.5 and that of the SMX mixer was 1.9. It is zero for an empty tube (211). Time is the only factor that influences the Lyapunov exponent. Since, for the same diameter and velocity, the time needed for the SMX mixer is 3.8 times that of Kenics mixer,
bsMX
= 3.8hKenics.
The potential of chaotic mixing is well established as a processing technique to create very fine scale structure of the order of 1 nm (102). Schrenk et al. (212, 213, 214) produced coextruded structures with nanolayers of polymer films, by using a ISG static mixer assembly. Chaotic mixing as per definition is based on a repetition of the same motion characterized by stretching, folding and staggering. The static mixer mimics the same phenomena, as shown in Figure 4.3.6. The number of layers or the striation thickness produced at the end of the mixing process gives the degree of mixing. The number of layers produced in a static mixer is ksn, where ks is the number of channels in a mixer element and n is the number of elements. For example, in the present study, an ISG mixer, which has 4 channels, is used. With just 10 elements, 1.05 million layers would be produced at the end of processing. This gives an idea about the extent of distributive mixing efficiency in a static mixer. Figure 4.3.7 shows a schematic of distributive mixing in the ISG static mixer.
74
4 MODELING AND PROCESS DESIGN
Static DÛxer
stretclûng
Chaotic Mixing
cuttingstretching &stacking
stretching;
cutting &stacking
stretciring
n
2 +1
Figure 4.3.6: Sehematie representation of baker's transformation in a statie mixer and ehaotie mixing (116)
Figure 4.3.7: Distributive mixing in the ISG statie mixer
The effectiveness of the ISG static mixer in producing minocomposites can be summarized as following: •
It is capable of imparting sufficient shear stress to delaminate organoclay
particles. •
The ISG static mixer, attached to extruder, will result in increased residence time, which is necessary for polymer diffusion and surface peeling of the clay particles.
•
The operation of the ISG static mixer incorporates the efficient chaotic mixing mechanism.
4 MODELING AND PROCESS DESIGN
•
75
The high level of distributive mixing in the ISG statie mixer is demonstrated by the faet that it produees 1.05 million layers withjust 10 mixing elements.
•
Flow in the ISG statie mixer involves a high level of extensional efficieney and stretehing.
•
The polymer melt, while passing through the ISG statie mixer, changes direction frequently. A frequent change in flow direction is required for breaking and delaminating the clay particles.
•
Installation of static mixer elements is easy. The mixing process is continuous and can be easily scaled up for industrial production.
A schematie of the proposed design for produeing nanocomposites by melt proeessing is shown in Figure 4.3.8. It eonsists of a twin serew extruder, an ISG statie mixer assembly and a die. The statie mixer elements are housed in a metallie eylinder. The temperature of the eylinder is manipulated with external heating bands and temperature eontrollers. The melt pressure is monitored at the end of the extruder serew, with the help of a pressure transdueer. This gives the pressure difference aeross the statie mixer and the die. Feed Hopper
Static Mixer
Extruder
Figure 4.3.8: Schematic representation of the proposed melt processing system for nanocomposite production
5 EXPERIMENTAL
76
5 EXPERIMENTAL 5.1 PROCESSING EQUIPMENT As described in the previous section and shown in Figure 4.3.8, a static mixer was attached to the twin-screw extruder. Bither slit die or strand die, was attached to the other end of the static mixer. The processing arrangement with the twin-screw extruder, ISG static mixer and die to produce nanocomposites is referred to as 'System A'. Nanocomposites were also produced without using the static mixer attachment. The processing arrangement with only the twin-screw extruder and die is referred to as 'SystemC'. Samples leaving the slit die were cooled by air fans located above and below the take-off guide rolls and then pulled by horizontal take-off system. The take-off system has two tractor belts with adjustable gap. The take-off speed was adjusted according to the extrusion flow rate, so as to avoid stretching of the ribbon. The resulting ribbon is 60-70 mm wide and 0.4-0.8 mm thick, which could be used as is for testing mechanical properties. Sorne of the samples were further cut into small pieces and subsequently compression molded, for testing mechanical properties. Samples made with the strand die were cooled in a water trough and palletized using Berlyn pelletizer (Worcester, MA, USA). The pellets were injection molded into plaques.
5.1.1 TWIN-SCREW EXTRUDER A co-rotating twin-screw extruder with intermeshing screws, Model ZE 25, from Berstorff GmBH (Hannover, Germany) was used in the present study. The barrel diameter of the extruder was 25 mm, and the ratio of screw length to screw diameter was 30. The extruder barrel is 'divided into five zones. The temperature of each zone is controlled individually by an electrical heating band and a cooling fan. The polymer and the clay were fed to the extruder hopper by separate feeders. The twin-screw volumetric feeders Model T20 from K-Tron Corporation (Pitman, USA) were used for this purpose. Figure 5.1.1 shows the configuration of the screw elements used for the twinscrew extruder during all the extrusion experiments. Figure 5.1.2 gives the nomenclature
77
5 EXPERIMENTAL
of the different screw elements used for this configuration. The melting section (Zone 1 and 2) contains three kneading elements followed by a mixing element and blister ring. Zone 4 has one more kneading element and zone 5 has five short pitch conveying elements. This is a commonly used configuration for polymer melt blending and processing.
0
MM
50
::J --'
ru
ru
" "r---:
"::::::
ln
ru w
Vl
100
--' ru ~
"
ln
ln
(Y)
(")
w
Vl
r---:
::J
150
ru ~
ru
--' ru
ln
ln
Ln
--'
200
250
::J
z
ln
0
--' ru
300
350
:=;:;
::J
ln ~
::J
400
::J 0
450
ln
"" ::J ru ru
" " "'" ru" " ru" " r---:" ":::::: ":::::: " " :::::: " " " ru "ru r---: ru r---:" ru" c- ru" r---: '- r---: ru ~
ln
Ln
Ln
ln ln
Ln
ln
w
w
w
:Jo
ÇQ
:Jo
(/)
(/)
(/)
(/)
Y'
(/)
ln
Ln
(Y)
(")
(")
ÇQ
Y'
Ln
:Jo Vl
ÇQ
Y'
ln
ln
:Jo
ÇQ
(/)
N
+' lJ1
P5
::J ru
"
"
--' ru
"
"Lr!c--
ln
Ln
(")
(")
(")
w
:Jo
(U
(")
500
w (/)
r---:
Vl
(/)
550
::J ru
650
600
--' ru
z.
ln
0
ru~
"",,--'
700
::J ::J
--'
w
:Jo
750
--'
'"Ln" Lr!ru ru ru ru ru " " "ln c-" ~ " ":::::: " "ln ,-" c-Ln (Uln ln " r---: +,ru " Ln " Ln ln" ru ru ru ru ru (Y)
:Jo Vl
(")
:Jo
Vl
lJ1
ÇQ
:.::::; w
Y'
ÇQ
Vl
Vl
W Vl
(/)
:Jo (/)
Figure 5.1.1: Configuration of screw elements
~ S'vi 37,5/1/2 LI
SW = Self Wiping, SE = Simple Element
.~
-/ -/ -=length / no. of pitch 1 no. of threads
S'vi 25/1/2 LI
LI
~
~
SE 37,5/1/2 LI
SE 37.5/1/2 LI
[ill[]
[]JjjJ
[tt[]
KB 37,5/5/90 NI
KB 37,5/5/45 LI
S'vi 37,5/5/45 RI
Il
~
ZB 37,5/5/10 LI
ZB 15/2/10 LI
~
~
S'vi 7,5/0.5/2 RI
Blister 7,5/24,5
= Left handed, RI = Right handed
KB =Kneading Block -/ -/ -=length / no. of kneading blocks /
shifted angle between successive blocks
=Non conveying, ZB = mixing element NI
-/ -/ -=length / no. of mixing discs / no. of teeth at the periphery Blister 7.5/24.5
-/ -=length of the element / diameter of the disc
Figure 5.1.2: Nomenclature of the screw elements used
5 EXPERIMENTAL
78
5.1.2 STATIC MIXER The static mixer selected for the present study is known as the ISG (Interfacial Surface Generator) motionless or static mixer. The elements of the static mixer were purchased from Ross Engineering Inc (Hauppauge, NY, USA). The housing for mixer elements and other accessories were fabricated at the workshop of the Department of Chemical Engineering, McGill University. The heating bands were purchased from Zesta Engineering limited (Mississaga, ON, Canada). The temperature indicators and controllers were assembled in-house. The static mixer elements with diameters 1 inch or 1.5 inch were suitable as an attachment to the extruder used in the present study. The selection of the static mixer elements depends upon the anticipated pressure drop during processing. The pressure drop created should be within the safe operating limit of the processing equipment under the processing conditions such as flow and viscosity of the polymer melt. 5.1.2.1 ESTIMATION OF PRESSURE DROP The anticipated pressure drop during processing polymer melts through the static mixer was estimated as per manufacturer' s guidelines according to the following procedure. 1. Computation of Reynolds number, Re = 6.31w
J1Di
(5.1)
where w =mass flow rate, lb/hour Di = inside diameter of mixer housing, inches
J.L =absolute viscosity, centipoises
For the flow rate of 2-5 kglhr (4.4-11 lblhr) of polymer melt with viscosity 1000 Pa.S (106 centipoise), the Reynolds number for 1.0 and 1.5 inch diameter ISG static mixer ranges from 1.85xlO-5 to 7.0 xlO-5 • In this laminar flow region, Figure 5.1.3 is used to estimate the pressure drop through static mixer elements. 2. The estimated flow rate of polymer melt during nanocomposite processing is 2-5 kglhr (i.e. == 0.01-0.02 gallons/min). From Figure 5.1.3, the estimated pressure drop
5 EXPERIMENTAL
79
per element for 1.5 inch diameter static mixer is 0.2-0.4 psi, and that for 1.0 inch diameter is 0.6-1.0 psi, for the fluid with 10000 cps viscosity. 3. For the 10 elements used in this study, the total estimated pressure drop in the static mixer is 2-4 psi for the 1.5 inch diameter elements, and 6-10 psi for the 1 inch diameter elements, for the fluid with viscosity 10000 cps. 4. The actual total pressure drop for the polymer melt of viscosity 1000 Pa.S (106 cps) passing through 10 elements of the ISG static mixer is estimated to be 200-400 psi for 1.5 inch diameter and 600-1000 psi for 1.0 inch diameter static mixer. The pressure drop for the 1.5 inch diameter elements is within the reasonable operating limit of the twin-screw extruder. The selected twin-screw extruder is capable of developing 2000 psi pressure under safe operating conditions. The pressure drop for 1.0 inch diameter static mixer is on the higher side and may pose limitation on the flow rate and/or viscosity of the polymer melt being processed.
~
,..1 ~
10.0
~ "" QI(
!t .~ CI
~ ~
HI
flOW - GAllONS PEi MINUTE 1-0.625 IN. PIA, - - - 1 IN. PIA.
1-6 IN. DIA,
- - -
a IN. DIA,
• •••• lSIN. DIA. ••••• 10 IN. OIA,
lIN. DIA. -l21N, DIA.
- - - :1 IN. DIA. - - -16 IN, DIA.
•.•.. 411'
'ai
:; E
0.2 -1'- System A, 2.5 Kg/hr
:J
Ü
___ System A, 1.7 Kg/hr
. ,' 0'
0.0
",0" System C, 1.7 Kg/hr ",0" System C, 1.0 Kg/hr
0.4
0.6
1.0
0.8
1.2
1.4
1.6
1.8
Normalized Time ( VRT)
Figure 6.1.5: Cumulative residence time distribution curves for System A and System C under different feed flow conditions
Table 6.1.1 Mean residence time and variance for System A and System C
System A
System C
Feed Rate
MeanRT
Variance
MeanRT
Variance
(Kg/hr)
(Sec)
cr?
(Sec)
cr?
1.0
449
5031
281
3482
1.7
333
2644
212
2409
2.5
236
1496
157
689
2.85
212
1034
6 RESULTS AND DISCUSIONS
105
6.2 X-RAY DIFFRACTION Mter the clay is organically modified, the most common technique used to analyze the clay is X-ray diffraction (XRD), which allows the interlayer d-spacing to be measured. Because XRD has been successfully used to analyze organically modified clays, it has been employed to observe changes in d-spacing when polymerie layered silicate nanocomposite materials are prepared. XRD is most useful for the measurement of the d-spacing of ordered immiscible and ordered intercalated nanocomposites, but it may be insufficient for the measurement of disordered and exfoliated materials that give no peak (186,221). More specifically, the lack of peak may be misinterpreted in cases where no peak is seen. Many factors, such as concentration and order of the clay, can influence the XRD patterns of layered silicates. TEM and XRD become complimentary techniques, filling the gaps of the information that other technique cannot obtain. In the following sections the results of Wide angle X-ray Diffraction (XRD)
patterns of PA-6 and PS nanocomposites will be discussed, followed by the detailed discussion of interpretation of exfoliation and intercalation. 6.2.1 PA-6 NANOCOMPOSITES
As mentioned earlier in this section, PA-6 nanocomposites were prepared with three different clays, using System A and System C. Figure 6.2.1 shows the XRD plots of PA-6/Cloisite 30B nanocomposites with 2.1 %wt clay, prepared using System A and System C. The XRD plot of in-situ polymerized commercial PA-6 nanocomposite from Ube Industries Ltd. is also shown, for comparison. As described by Kato and Usuki (1), the in-situ polymerized PA-6 nanocomposite has very well-dispersed silicate layers, which exhibit no diffraction peak, but a graduaI increase in diffraction intensity at lower angles is observed. A similar XRD pattern can be seen for the System A sample. However, the System C sample shows a broad diffraction peak around the peak position of Cloisite 30B, and the absence of graduaI increase in intensity towards lower angles. This is clear evidence that System A has formed a higher degree of clay dispersion, compared to System C.
6 RESULTS AND DISCUSIONS
106
20000~---,~--------------------------------------~
15000
>.
ëii c::
(1)
c::
10000
>.
~1
>
- 20000
"iii
60000
t:
(1)
t:
(1)
.!: 15000
~ ....
40000
1
X
~
"iii
10000
E >al
.... 1
X
20000
5000
0 0
2
4
6
8
10
28
Figure 6.2.2: XRD plots of PA-6/Cloisite 308 nanocomposite samples made using
System A with different clay contents
30000
100000 1.8 nm Cloisite 308
25000
~
80000
20000
scale
"iii
-
60000
t:
(1)
E
>-
t:
(1)
15000
E >-
>al ....
X
-
"iii
40000 10000
~1
X
20000
5000
0 0
2
4
6
8
10
28
Figure 6.2.3: XRD Plots of PA-6/Cloisite 308 nanocomposites made using System C
with different clay contents
6 RESULTS AND DISCUSIONS
108
The structure of the polymeric nanocomposites formed during melt processing depends highly on the chemical structure of the clay modifier used. As seen in Section 5.3, organoclay Cloisite 30B contains organic modifier with two hydroxyl functional groups attached to the quaternary ammonium functional group. These hydroxyl groups have tendency to form hydrogen bonds with polyamide chains. This is considered as a critical factor forming the formation of effective intercalatedlexfoliated structure. On the other hand, organoclay Cloisite 15A has higher initial d-spacing but lacks the presence of functional groups capable of forming hydrogen bonds. The phenomenon of formation of exfoliated and intercalated nanocomposites is discussed in more detail at the end of this section (Le. Section 6.2). Figure 6.2.4 shows the XRD plots of PA-6/Cloisite 15A nanocomposites made with System A and System C for 2.5% wt inorganic content of clay. It can be seen that, in both cases, the gallery spacing increases marginally from 3.33 nm to 3.68 nm. The peak at 1.84 nm is a secondary peak of 3.68 nm primary peak. The X-ray diffraction pattern results from the resonance of the waves diffracted from the consecutive layers. It results in the bright and dark intensity pattern. The primary peak is usually the most prominent and is followed by consecutive peaks of diminishing intensity. The virtual dspacing with respect to the secondary peak is exactly half of the primary peak. The increase in gallery spacing by 0.35 nm does not mean that polymer chains are intercalated. Though the peak position is same for both the systems, the intensity of the diffraction peak is lower for the system A sample. If the larger layered silicate particles are broken into smaller tactoids, which are placed away form each other (distance more than 8 nm) in the polymer matrix, the X-ray diffraction intensity could be substantially lowered for the same concentration of clay. In case of PA-6/Cloisite 15A nanocomposites, the clay is in the form of tactoids and sorne exfoliated platelets (which is further evidenced by TEM). System A produces smaller but more tactoids than System C. This results in lowering the X-ray diffraction intensity.
6 RESULTS AND DISCUSIONS
109
60000
50000 3.33nm ~
40000
ii)
c: (])
System C
ë 30000 >-
System A
I!! 1
X
20000
10000
0 0
2
6
4
10
8
29 Figure 6.2.4: XRD plots of PA-6/Cloisite 15A (2.5% wt) nanocomposite.s made using System A and System C 60000 3.68 nm 50000
-
40000
E
30000
>c: (])
i i)
-
>-
I!! 1
X
20000
10000
0 0
2
4
6
8
10
29 Figure 6.2.5: XRD plots of PA-6/Cloisite 15A (4.1% wt) nanocomposites made using System A and System C
6 RESULTS AND DISCUSIONS
110
Figure 6.2.5 glves similar XRD plots for 4.1 % wt. clay content. With the increase in clay content, the X-ray diffraction intensity is higher in this case. The System A sample shows substantially lower X-ray diffraction intensity than that of the System C sample. 10000...-----nr--------------------,
8000
.2:-
'iii c:
-
6000
Q)
.E >-
f:! 1
4000
X
2000
O+----~---_.----~---~---~ 2 4 6 8 10
o
29 Figure 6.2.6: XRD plots of PA-6/Cloisite Na+ (4.2% wt) nanocomposites made using
System A and System C
In Figure 6.2.6, it is seen that even in the case of unmodified Cloisite Na+ clay, melt processing results in the increase in gallery spacing from 1.1 nm to 1.77 nm. The difference between the X-ray diffraction intensities of the System A sample and the System C sample is insignificant. Alhough there is a small increase in d-spacing, compared to the pristine clay, it cannot be concluded that there is useful intercalation by the pol ymer chains.
6.2.2 POL YSTYRENE NANOCOMPOSITES Polystyrene is a non-polar polymer with no functional groups capable of forming hydrogen bonds. Cloisite 10A contains organic surfactant with styrene functional group,
6 RESULTS AND DISCUSIONS
111
which may render it more compatible with the polystyrene matrix. It was observed that the gallery spacing of Cloisite lOA substantially increases when melt processed with polystyrene matrix polymer (51-53). Polystyrene-based nanocomposite systems provide a good oppoctunity to study the effect of different processing parameters on the intercalation behavior. In this work, the melt processing of polystyrene nanocomposites with Cloisite lOA organoclay was evaluated, using System A and System C. The effect of using maleic anhydride grafted polystyrene (Dylark 332) as compatibilizer, on the intercalation behavior, was studied using the same organoclay. The effect of processing temperature was also evaluated, for the polystyrene nanocomposites with and without compatibilizer. Both organoclays Cloisite 15A and Cloisite 30B were tested in the polystyrene matrix. AIso, both low molecular weight polystyrene PS3900 and the higher molecular weight PS1301 were included in the study. Figure 6.2.7 shows .the XRD patterns for PS1301/Cloisite lOA (3.2% wt) nanocomposites prepared at 210°C, using System A and System C. In the case of System A, the XRD peaks are at d-spacing of 3.84 nm and 1.67 nm. The System C sample shows d-spacing of 3.4 nm and 1.67 nm. The increase in d-spacing is higher for the System A samples. The d-spacing of 1.67 nm results from the collapsed clay structure (167). In Figure 6.2.8, it is seen that the peak intensities are sharper. when the Dylark
compatibilizer is incorporated in the system. This plot shows the clear difference between the System A and the System C samples. System A is more effective in increasing the d-spacing of Cloisite lOA organoclay. The peak positions are still the same for both systems, but the intercalation peaks are of higher intensity than those of the collapsed peaks. The increase in intensities can be attributed either to the increase in the number of intercalated clay particles or the increase in orientation of clay tactoids. Galgali et al. (222) showed that, in the case of compatibilized polypropylene nanocomposites, clay tactoids were more oriented and ordered. This may result in sharper X-ray diffraction peaks in compatibilized nanocomposites.
112
6 RESULTS AND DISCUSIONS
20000 2.05 nm 1.67 nm 15000 ~
"iii
1::
Q)
..!:
10000
...>CIl 1
X 5000
o o
2
6
4
8
10
29
Figure 6.2.7: XRD plots of PS1301/Cloisite 10A (3.2%wt) nanocomposites made using System A and System C at 210°C
30000 System A 25000
-
20000
..!:
15000
>-
"iii
-
1.67 nm System C
1:: Q)
1.73 nm System A
...>CIl 1
X
10000
5000
/PS
0 0
2
4
6
8
10
29
Figure 6.2.8: XRD plots of polystyrene nanocomposites with Dylark compatibilizer. PS1301 +2%wt Dylark332/Cloisite 10A(3.2% wt) nanocomposites made at 210°C
6 RESULTS AND DISCUSIONS
113
Figure 6.2.9 shows the XRD patterns for the samples prepared with and without compatibilizer, using System A. The intensity of the intercalated diffraction peak is higher in the compatibilized sample and the intensity of the coIlapsed peak is reduced. Samples were also prepared with only Dylark 332 and Cloisite lOA (3.2% wt) using System A at 210°C, in order to evaluate the extent of the influence of maleic anhydride grafted compatibilizer on the amount of intercalation/exfoliation. It was seen that these samples were very brittle and weak. Probably the reaction between the organoc1ay and the maleic anhydride group resulted in the evolution of gases (e.g. carbon dioxide), resulting in weak and degraded samples. The XRD pattern for one of these samples is shown in Figure 6.2.10. The increase in d-spacing in the pure Dylark nanocomposite was lower than for the polystyrene compatibilized with 2% wt Dylark nanocomposites. Moreover, the coIlapsed peak was at lower d-spacing. This shows that the pure Dylark was not effective in enhancing intercalation, under the conditions employed in this study. The effect of the processing temperature on intercalation behavior was studied for the compatibilized and the uncompatibilized polystyrene nanocomposites, with Cloisite lOA. AlI the samples were prepared using the System A. Two temperatures (210°C and 240°C) were set at the front zones of the extruder, static mixer attachment and at the die. Figure 6.2.11 gives the XRD patterns for uncompatibilized polystyrene nanocomposite samples. The sample prepared at 210°C shows peaks at 3.84 nm and 1.67 nm. The sample prepared at 240°C shows only one peak at 1.5 nm, indicating the coIlapse of the layered platelets. Also the d-spacing of the coIlapsed platelets is lower at 240°C than that of at 210°C. Figure 6.2.12 shows similar plots to those of Figure 6.2.11 for the compatibilized polystyrene nanocomposites. The sample prepared at 210°C shows c1ear and sharp peaks at 3.84 nm and 1.73 nm, indicating the increase in dspacing as weIl as coIlapse of the peaks. The sample prepared at 240°C shows broad peak around 3.84 nm and sharp peak at 1.57 nm. It seems that high processing temperature leads to reduced intercalation and higher coIlapse of the layered organoc1ay platelet, possibly due to degradation of the modifier.
6 RESULTS AND DISCUSIONS
114
30000
25000
PS/2% Dylark/Cloisite 10A
20000 ~
Iii
c:: Q)
E
~ X
15000 Cloisite 10A 10000
5000
0 0
2
6
4
8
10
28
Figure 6.2.9: Effect of compatibilizer on the formation of PS nanocomposites. PS1301/Cloisite 10A (3.2% wt) nanocomposites prepared with and without Dylark 332 compatibilizer at 210°C using System A. 30000
i-.---------;::============il
25000
~
3.84
- - Dylark/Cloisite 10A PS1301/2%Dyl/Cloisite 10A - - - Cloisite 10A
nm
20000
Iii
c::
Q)
ë
~ X
15000
10000
5000
-----------
0 0
2
6
4
8
10
28
Figure 6.2.10: Formation of nanocomposite with Dylark 332 and Cloisite 10A (3.2% wt). Extruded at 210°C using System A
6 RESULTS AND DISCUSIONS
115
20000.---~---------------------------------------.
15000
:È' (J) c
~ c
/
10000
2.05 nm 1 .67 nm /\ 1.5 nm f \ 1 \ 1 \ f f /
~ X 5000 - - Extruded at 210°C - - - Extruded at 240°C - - - Cloisite 10A
-----------
O~=======T======~~------~--------,_------~
o
2
4
6
8
10
28
Figure 6.2.11: Effect of processing temperature; PS1301/Cloisite 10A (3.2% wt) Extruded using System A at 210°C and 240°C 30000
25000
-
- - Extruded at 210°C - - - Extruded at 240°C - - - Cloisite 10A
3.84 nm
20000
>-
i i)
c
Q)
E
15000
>-
~1
>
1
X
3.33 nm 40000
20000
......
0 0
2
4
6
-------8
10
28
Figure 6.2.14: Effect of compatibilizer; PS1301/Cloisite 15A nanocomposites with and without Dylark 332 compatibilizer made using System A at 210°C
118
6 RESULTS AND DISCUSIONS
30000
i--'---?iR'fiÏrn---;::::===========::::;T 100000 - - PS1301/2%DylarklCloisite10A - - PS1301/2%DylarklCloisite15A - - - PS1301/2%DylarklCloisite 308
25000
80000 scale
>.
20000
=55 c:
-
.!: 15000
~
/
\ \ \ \ \
1
X
60000
1.5 nm r\ 1.73 nml \
Q)
10000
40000
\ \
20000
5000
0
+-----.-----.----~----.----_+O
0
2
4
8
6
10
29
Figure 6.2.15: Using different organoclays with PS1301 matrix; Samples prepared using System A at 210°C
30000 3.84 nm 25000
ïn~ c:
--- - .. ... ...
PS3900/2%DylarklCloisite 10NSystem A PS3900/2%DylarklCloislte 10NSystem C PS3900/Cloisite 10NSystem A Cloisite 10A
20000 2.05 nm
Q)
.!: 15000
-,,~--1.73
>.
nm
l!! 1
X
10000
5000
0 0
2
4
6
8
10
29
Figure 6.2.16: PS3900/Cloisite 10A (3.2% wt) nanocomposites with and without compatibilizer using System A and System C at 210°C.
6 RESULTS AND DISCUSIONS
119
6.2.3 FACTORS INFLUENCING EXFOLIATION The above XRD results show that polyamide-6 nanocomposites with Cloisite 30B clay gave well-dispersed potentially exfoliated structure, but the same polymer with Cloisite 15A organoclay showed mainly evidence of intercalation. Similarly, in the case of polystyrene, Cloisite lOA exhibited larger increase in d-spacing, compared to the other organoclays. Polystyrene does not show indications of potential exfoliation, even in the presence of compatibilizingagent. Qualitatively, the modifier employed in Cloisite 30B is capable of forming hydrogen bonds with the polyamide-6 matrix. Thus, it results in a high degree of dispersion and potential exfoliation of clay platelets. In all other cases considered in this study, the constituents are not capable of forming hydrogen bonds. Hansen solubility parameters have been used to explain the dispersion of organically modified clays in organic solvents (223, 224, 225). The Hansen solubility parameter ~ is defined as: (6.1) where subscripts d, p and h refer to dispersive, polar and hydrogen bonding components. To produce exfoliated polymer-clay nanocomposites, the interactions between the clay surfaces and the intercalants need to be understood. The organic molecules can penetrate the spaces between the layers of the clay by interacting with the clay surfaces in the following ways (226): . 1. Cationic bonding, in which the protonated alkylammoniums replace the sodium ions in MMT layers; 2. Ion-ion dipole interactions, in which the polar organic molecules are related to the sodium ions in the MMT layer; 3. Dipole-dipole interactions, which inc1ude the hydrogen bonding that associates polar organic molecules with hydroxyl groups or oxygen in the clay layers.
The total solubility parameter
~
of an MMT layer whose diameter is less than 2
pm is so high that it is not miscible or soluble in organic liquids, monomers or polymers.
6 RESULTS AND DISCUSIONS
120
Liquids interact differently with MMT, because they have different component values in the Hansen solubility parameters. Table 6.2.1 gives the values of the Hansen Solubility Parameters (HSP) for sorne selected polymers, solvents and organic compounds, from the literature. The solubility parameters for Cloisite organoclay modifiers are ca1culated following the group contribution method described by Hansen (227). The values of the dispersion parameters (t:>d) are in close range. The values of the polar parameters (t:>p) and hydrogen bonding parameters (t:>h) depend upon the nature of the functional groups. Non-polar compounds, such as n-Hexane and polyethylene, have low values of polar components. Similarly, compounds with no ability to form hydrogen bonds, such as nHexane and PTFE, show low values of the hydrogen bonding component. On the other hand, polyamide has relatively higher values of t:>d and t:>h. Water has the highest values. The organic modifier in Cloisite 30B has relatively high polar and hydrogen bonding parameters, while the modifier in Cloisite 15A is non-polar and has a very low hydrogen bonding component. Table 6.2.1: Hansen Solubility Parameters (J/cm 3) 1/2 (227)
t:>d
t:>p
t:>h
t:>
Polyamide
16.0
11.0
24.0
30.9
Polystyrene
18.6
10.5
7.5
22.6
Polyethylene
17.1
3.1
5.2
18.1
PTFE (Teflon)
17.1
8.1
1.3
19.0
Water
15.5
16.0
42.4
47.9
n-Hexane
14.9
0.0
0.0
14.9
Cloisite 30B modifier
18.4
1.7
4.6
19.1
Cloisite 15A modifier
17.3
0.3
0.7
17.3
Cloisite 10A modifier
17.6
0.5
1.3
17.7
Compound
Ho and Glinca (224) studied the effect of solvent solubility parameters on organoclay dispersions. They concluded that with the same value of t:>d, increasing the values of t:>p and t:>h changes the clay platelet formation from tactoids to complete
6 RESULTS AND DISCUSIONS
121
exfoliation, indicating that polar and H-bonding forces of the solvent molecules affect the tactoid formation and structure of clay in the suspension. Choe et al. (223) also concluded that liquids with high
~
values exp and the doOl spacing of sodium
montmoriUonite, more than liquids with low C>h values. The hydrogen bonding components (C>h) and polar components (C>p) of liquids were found to be the primary parameters for the dispersion states and for the basal spacing expansion. Since polyamide-6 has high polar and hydrogen bonding components, it is likely to cause exfoliation of Cloisite 30B. Cloisite 15A and Cloisite lOA, however, are nonpolar and cannot form hydrogen bonds with polyamide-6. Thus, the clay platelets are not exfoliated and tactoids are formed. Polystyrene has very low hydrogen bonding component compared to polyamide. This leads to tactoid formation in aU the organoclays. Another consideration relates to the effective Hamaker constant of the system. Neumann et al. (228) showed in particle engulfment experiments that, if the effective Hamaker constant of the system is negative, the particles are rejected by the polymer melts. The effective Hamaker constant A 132 for the system of the pristine clay [1] and the polymer[2] with the organic modifier[3] between them is given by Equation 2.11:
A132 = (JAu - JA XJ~2 - JA 33
33)
where An is the Hamaker constant of pristine montmoriUonite, A 22 is the Hamaker constant of the polymer and A33 is the Hamaker constant of the organic modifier. The organic modifier is between the clay particle and the polymer melt. If the value of A 132 is positive, there is attraction between the clay[1] and the polymer[2], and the clay particles are engulfed by the melt' which results in good dispersion and probably leads to exfoliation. If this value is negative, the particles are rejected by the polymer melt and large tactoids are formed. The Hamaker constant values for the organic modifiers can be estimated by using the group contribution method described by Vial and Carre (229) and the surface tension data from Jasper (230). Table 6.2.2 shows that the Hamaker constant values of non-polar compounds such as polyethylene and PTFE are low. The Hamaker
122
6 RESULTS AND DISCUSIONS
constant values of the organic modifiers also faH in the range of 5.3-6.0 x 10-2oJ, since these are saturated long chain alkane compounds based on Tallow. The Hamaker constant values are higher for amides and alcohols, since they have polar functional groups capable of forming hydrogen bonds. The effective Hamaker constant, in the case of polyamide as the matrix, is always positive and much higher compared to the other systems. This indicates that the organoclay is likely to be better dispersed in a polyamide matrix. Systems involving a polystyrene matrix have a low positive value. Polyethylene shows a value close to zero or negative, .and PTFE gives a negative value. It can be concluded that non-polar polymer matrices will result in po or dispersion of organoclay. Table 6.2.2: Values of Hamaker Constant (228,231,232,233,234)
Substance
Aii (l0-20 J)
A 132 (l0-20 J)
Montmorillonite (A ll )
7.8
Cloisite 30B Modifier (A33)
6.0
Cloisite 15A Modifier (A33)
5.7
Cloisite 10A Modifier (A33)
5.3
Polyamide (A2Û
12.0
0.35 - 0.57
Polystyrene (A22)
6.5
0.034 - 0.12
Polyethylene (A22)
5.1
(-0.021) - (-0.066)
Polytetrafluoroethylene (A 22 )
4.0
(-0.15) - (-0.16)
6.2.4 FACTORS INFLUENCING INTERCALATION The intercalation behavior of polystyrene into the gallery spaces of organoclay has been extensively studied by Vaia et al. (37,51,53). The intercalation kinetics and polymer diffusion studies were based on the observation of the increase in gallery spacing, as indicated by X-ray diffraction. Unfortunately, XRD and even TEM resuIts on the intercalated nanocomposites show only the increase in gallery spacing. They do not confirm that the polymer chains are actually intercalated between the clay galleries.
6 RESULTS AND DISCUSIONS
123
Since none of the techniques available can confirm this fact, the above results have been taken as an indirect evidence of intercalation of polymer chains into the gallery spacing. As mentioned by Giannelis et al. (159), it is rather surprising that polymer melt can intercalate layered inorganic compounds unassisted by shear or solvents. Intercalation implies that polymer chains can undergo large center of mass displacement in almost two dimensional interstices, as the distances between the confined surfaces are substantially smaller than the unperturbed radius of gyration of the polymer and are comparable to the monomer size. Vaia et al. (53), using in-situ XRD, observed that the kinetics of intercalation, even under quiescent conditions (absence of external shear) , are quite rapid. The apparent diffusivity for the intercalation of polystyrene in organoclay was found to be of the same order of magnitude (10. 11 cm2/s at 170°C) as the self diffusion coefficient of polystyrene.
Furthermor~,
the activation energy of melt intercalation was 166 kJ/mol,
which is comparable to the activation energy measured for self diffusion of polystyrene (167 kJ/mol). It was also found that the effective diffusion coefficient depends markedly on the surfactant used, for the same polymer and annealing temperature. This is somewhat strange, because the surfactant can only affect the polymer motion inside the galleries. For sorne systems, the intercalating polymers were found to possess a mobility that was much faster than the self diffusion coefficient of the corresponding polymer in the bulk (235,236), or in a thin film (237). The above studies suggest that the intercalation of polystyrene chains into the gallery spaces of organoclay is almost instantaneous and uniform throughout. However, this does not lead to the exfoliation of the platelets. On the contrary, the exfoliation of the organoclay platelets in polyamide matrix, described by Paul et al. (123-126) and shown in Figure 2.5.6, is dependent on the residence time. The higher residence time gave higher dispersion of the platelets. The increase in the gallery spacing of sodium montmorillonite, when treated with organic modifiers is due to the intercalation of the long chain modifiers between the gallery spaces. The quaternary ammonium organic modifier compounds possess welldefined polar and non-polar regions and are called amphiphilic. They replace the small
6 RESULTS AND DISCUSIONS
124
sodium cations present in the gallery spaces. Amphiphilic molecules form monomolecular interfaciallayers that separate polar and no-polar phases (14). The usual way to prepare a monolayer for study in a Langmuir balance is to dissolve the amphiphilic molecules in a volatile organic solvent and disperse drops of solution onto the air-water interface. Layered silicates, that are susceptible to swelling, show Langmuir-type shape with the adsorption of monoalkyl and dialkyl dimethylammonium surfactants (238), as shown in Figure 6.2.17.
Figure 6.2.17: Orientation of n-alkylammonium ions in the interlayers of layer silicates. 1) Monolayer of short-chain cationic surfactants; 2) Monolayer of close-packed cationic surfactant; 3) Double layer of cationic surfactant parallel to the surface; 4) Pseudo triple layer of cationic surfactants parallel to the surface; 5) Paraffin-like structure (238)
The lowering of surface tension of the interface by adsorbed molecules is given by equation 6.2: Ils =
ro - ra
(6.2)
which represents the difference between Yo, the surface tension with no amphiphile present and Ya, the value with adsorbed amphiphile. Ils represents the force per unit
6 RESULTS AND DISCUSIONS
125
length needed to prevent the film from spreading. It is the two-dimensional equivalent of pressure. Thus, Ils is called the spreading pressure. Ils varies with the surface concentration of the amphiphilic molecules. The amphiphilic concentration may be manipulated by adding or subtracting amphiphilic molecules to and from the system, or by taking those already present and compacting or expanding them.
6.2.4.1 THE PHYSICAL STATES AND COLLAPSE OF MONOMOLECULAR FILMS Two-dimensional monolayers can exist in different physical states, which bear resemblance to the solid, liquid and gaseous states in three-dimensional matter. Surface films are classified according to the lateral adhesion between the film molecules, including end-groups (239). Monolayers can be roughly classified as: 1. Condensed State - Films in which the molecules are closely packed and steeply oriented towards the surface. 2. Liquid-expanded state - Films that are still coherent but occupy a much larger area than condensed films. They have no real three-dimensional equivalent, since they act as highly compressible liquids. 3. Gaseous or Vapor State - Films in which the molecules are separate and move about the surface independently, the surface pressure being exerted on the barriers containing the films by a series of collisions.
Figure 6.2.18 shows these states on a two-dimensional spreading pressure versus area per molecule isotherm (14, 15). For large surface area, the amphiphilic molecule concentration is low and the spreading pressure is very low. The monolayers exhibit gaslike behavior with two-dimensional equivalent of ideal gas law.
nsAI N s = nsa o = kT
(6.3)
where N s is the number of amphiphilic molecules on the surface and ao the effective area per molecule.
126
6 RESULTS AND DISCUSIONS
-
~
g
=
l1v - 0.01
mN/m
--.....
t.,-G /
•
........
l
\
0.20
0.25
0.50
10.0
80(nm2 molecules·') Figure 6.2.18: Schematic of two-dimensional spreading pressure versus area per molecule isotherm. At sufficiently high pressures
ne. the solid monolayer film collapses.
Plots of the compressibility factor Z (defined as IIsao / kT) reveal deviations from ideal conditions, for which Z=1. Plots for series of different n-carboxylic acids are shown in Figure 6.2.19. 2
N
Spreading Pressure
n.
(mM m-1)
Figure 6.2.19: Plots of compressibility factor versus spreading pressure (239)
6 RESULTS AND DIS eus IONS
127
Analogous to three-dimensional pressure-volume plots for gas, negative deviations (Z < 1) occur at low spreading pressures and positive deviations (Z > 1) occur at high spreading pressures. Because the strength of the attraction increases with molecular size, the deviations become more pronounced as the alkyl chain lengthens. As seen in Figure 6.2.18, when pressure TIs is high and ao is low, we observe solid-like behavior. In this case, the hydrocarbon chains are oriented vertically and closely packed together. The film is now relatively incompressible, and attempts to increase TIs still further cause the monolayer film to buckle or collapse at critical spreading pressure TIc. This phenomenon is illustrated in Figure 6.2.20 (240).
Figure 6.2.20: Collapse of monolayer film under pressure (240)
The critical collapse pressure for monolayer films is 20-50 mN/m, according to Figure 6.2.18. For monolayer films with thickness 1-2 nm, this pressure corresponds to a bulk pressure of 1-5x107N/m2 (100-500 atm). Under a static pressure of this magnitude, the monolayer .films between the platelets will collapse thereby reducing the gallery height. Under normal polymer processing conditions of high viscosity polymers, pressure values around 1000 psi (68 atm) are often encountered. The pressure between the nip regions of screws of intermeshing twin-screw extruders or batch mixers is exceedingly high. The higher processing temperatures as weIl as some factors mentioned below reduce the critical collapse pressure of monolayers, resulting in the collapsed structure.
6 RESULTS AND DISCUSIONS
128
6.2.4.2 FACTORS INFLUENCING THE PHYSICAL STATE AND COLLAPSE
OF MONOLA YER FILMS The physical state of a monolayer depends on the lateral cohesive forces between the constituent molecules. Lateral cohesion also depends on the geometry and orientation of the film molecules. The factors that favor the formation of expanded films are also responsible for the reduction in critical collapse pressure. The following factors are favorable for the formation of expanded films (239). 1. Bulky head groups, which prevent efficient packing. 2. More than one polar group. 3. More than one hydrocarbon chain oriented in different directions from the polar part of the molecule. 4. Bent hydrocarbon chains. 5. Branched hydrocarbon chains. Cloisite lOA nanocomposites with polystyrene
alw~ys
show increased gallery
spacing, as weIl as the collapsed structures. During melt processing of nanocomposites, the temperature throughout the matrix increases. This leads to an increase in the pressure in the gallery spaces, which results in the increase of d-spacing and the close packing of the modifier monolayer in the gallery spaces. In this case, the increase in gallery spacing should be relatively fast and uniform throughout. The bulky styrene head group in Cloisite 10A organic modifier resists close packing of the modifier monolayer between the gallery spaces, resulting in reduced critical collapse pressure. The collapsed structure observed in PS/Cloisite 10A nanocomposites could be attributed to the phenomenon of reduced critical collapse pressure. Figure 6.2.11 shows that only collapsed structure is observed when processing temperature is increased to 240°C. The resulting pressure at this processing temperature must have exceeded the critical collapse pressure of the monolayer. The use of Dylark compatibilizer increased the peak heights or the intensity of the peaks, but it did not produce increased gallery spacing. The compatibilizer probably resulted in increased orientation of clay particles. It did not appear to take part in the
6 RESULTS AND DISCUSIONS
129
intercalation. In all the polystyrene nanocomposite samples made at 210°C using System A, the gallery spacing was increased to 3.84 nm in comparison to System C, where gallery spacing was 3.4 nm. System A processing increases the residence time. This probably results in expanding gallery spaces. Fornes et al. (241) studied intercalation of PA-6 into organoclays modified with surfactants having no alkyl tail group, with one alkyl tail group, and with two alkyl tail groups. The surfactant with one tail group resulted in exfoliated structure, but the surfactant with two tail groups produced intercalated and collapsed structures. More than one hydrocarbon chains, oriented in different directions from the polar part of the molecule, also results in lower critical collapse pressure. The authors suggested that increasing the basal spacing of the organoclay does not necessarily lead to more exfoliation. Xie et al. (242) studied the thermal degradation of alkyl quaternary ammonium montmorillonite and observed the changes in gallery spacing with the increase in temperature. They found that the changes in the organization of al~uminosilicate layers due to the increase in temperature reflected the production of volatiles and associated increase of internaI pressure within the interlayer, leading to the expansion of the structure. The decomposition was not restricted to unconfined surfactant, but also included intercalated confined surfactant. Significant increase in gallery height of organoclay was observed when temperature was increased from 200°C to 280°C. The organization of the silicate layers appeared to initially .disorder during thermal decomposition and, depending on the decomposition route and temperature, collapse to varying degrees after completion of volatile release. In conclusion, the increase in gallery spacing of organically modified clay does not necessarily mean that the polymer chains are intercalated. The compacting of the monolayers of the organic modifier between the gallery spaces could result in increasing the distance between the platelets. The compacting of monolayer could be attributed to the high temperatures and pressures encountered during melt processing.
6 RESULTS AND DISCUSIONS
130
6.3 TRANSMISSION ELECTRON MICROSCOPY Characterization of polymer nanocomposites with transmission electron microscopy (TEM) differs from the TEM characterization of other polymer systems in that it requires no staining for sufficient contrast. The organoclay structures in a polymer nanocomposite naturaUy appear as dark features in the TEM micrographs. The XRD results by themselves cannot be used to adequately de scribe nanoscale dispersion of the layered silicates present in polymer nanocomposites. XRD results, when properly combined with TEM results, give a much clearer picture of the aetual nanoscale dispersion and overall global dispersion of the clay in the polymer (221). These two techniques provide information to help derive meaningful relationships between the polymeric layered silicate nanostructure and macroscale properties. The nanocomposite systems are generally classified as immiscible or exfoliated. The immiscible systems are in fact microcomposites rather than nanocoptposites. The exfoliated systems faU into two categories, exfoliated ordered and exfoliated disordered. The greatest clarification is needed for the intercalated definition. The TEM micrographs of intercalated systems show increased gallery spacing but cannot confirm the intercalation of polymer chains into the galleries. Yoon et al. (243) have shown that, in exfoliated systems, platelets have preferential orientation. The platelets are aligned in the flow direction. AU the TEM imaging samples in the present study were microtomed in the direction transverse to the flow direction and viewed parallel to the flow direction. 6.3.1 TEM MICROGARPHS OF PA-6 NANOCOMPOSITES The micrographs shown in Figure 6.3.1 and Figure 6.3.2 are those of PA-6/Closite 30B nanocomposites. They represent the exfoliated structure. The individu al silicate sheets appear as relatively straight to slightly curved dark lines. OccasionaUy sorne tactoids of clay are also seen. Qualitatively, the difference between the System A and the System C micrographs as weU as the difference between the specimens with 3.1 %wt clay and 5.0%wt clay are not obvious. The low magnification micrographs show that, in samples
6 RESULTS AND DISCUSIONS
131
produced with System A, clay platelets are more oriented in the flow direction than those produced with System C. The micrographs of .PA-6/Cloisite 15A nanocomposites are shown in Figure 6.3.3. This is a partially exfoliated system, and clay tactoids of different sizes are observed. System A micrographs show a distribution of individu al platelets, small tactoids and large tactoids. System C micrographs show larger and more tactoids than that of the System A. The nanocomposites of untreated clay were not expected to show exfoliated structure. As seen in TEM micrographs (Figure 6.3.4), only large clay particles are observed. Sorne of the clay tactoids or platelets are seen separating from large clay particles. Figure 6.3.4(d) suggests that the mechanism of breaking clay particle into tactoids and platelets is by erosion or peeling of large particle. Sorne of these specimens were analyzed by Mollet (244), by employing the method of cycloid test lines (245) and semi-automated quantitative image analysis. The cycloid test lines method produced data regarding surface density, specific surface area, the degree of exfoliation of clay particles, and the distribution of the clay particle sizes. The semi-automated quantitative image analysis produced data regarding particle length distribution. The aspect ratios of the nanoclays were estimated by combining the results from both analyses. In cycloid test lines method (245), the cycloid line grid is superimposed on the
micrograph randomly. The number of intersections between the test lines and the polymer/clay interface and the total number of test points hitting the micrographs (Appendix C) were recorded for four quadrants. The procedure was repeated with 10 independent pictures. The surface density of the interface is equal to twice the linear interception density. The surface density is the area generated at the interface between the particles and the matrix per unit volume of composite. It depends upon the clay content in the composite. The specific surface area is the surface area of the particles per unit mass of the particles and is independent of the clay content in the composite. The surface density
6 RESULTS AND DISCUSIONS
132
is the product of the specific surface area, the volume fraction of the silicate particles and the density of the silicate particles. In this analysis, the volume fraction of the silicate particles was estimated from TEM micrographs, as the ratio of the length of test line falling on the particles to the total length of the test line. The density of sodium montmorillonite clay was taken as 2.86 gmlcm3 (189). The specific surface area of the fully exfoliated montmorillonite was taken as 786 m2/gm as determined by the selective molecular absorption in aqueous suspension (246). The ratio of the specific surface area determined from TEM analysis to the maximum surface area of fully exfoliated montmorillonite clay gives an estimate of the extent or degree of exfoliation. Table 6.3.1 summarizes the result of TEM analysis regarding the volume fraction, surface density, specific surface area and the degree of exfoliation for sorne of the PA-6 nanocomposites prepared. The specific surface area and the degree of exfoliation are higher for the PA-6/Cloisite 30B (5%wt) nanocomposite prepared using System A than those prepared using System C. The System A specimen shows 94% exfoliation, whereas the System C specimen shows 70%. At lower clay content (3.1 %wt), even the System C sample shows 97% exfoliation. It can be concluded that System A makes it possible to achieve high dispersion at higher levels of clay content in PA-6 nanocomposites. The distribution of clay particles shown in Figure 6.3.5 for PA6/Cloisite 30B(5%wt) nanocomposites shows higher exfoliation of clay particles in System A samples than in System C samples. System A specimens contained 75% of the particles with single platelets, whereas in System C specimens, only 50% of the particles consisted of a single platelet. The average thickness of the clay particles in System A nanocomposites was 1.9 nm, while it was 4.8 nm in System C nanocomposites. In the case of partially compatible organoclay Cloisite 15A, the X-ray diffraction
patterns and TEM micrographs do not show exfoliated structure. The specific surface area and the degree of exfoliation in these samples are low. The difference between System A and System C samples is not very significant. Moreover, further processing, such as compression molding, did not change the degree of exfoliation to any significant extent for PA-6/Cloisite 15A nanocomposite.
133
6 RESULTS AND DISCUSIONS
(a) System A
(b) System C
(c) System A
(d) System C
Figure 6.3.1: TEM micrographs of PA-6/Cloisite 308 (3.1 %wt) nanocomposites
(a) Low magnification System A, (b) Low magnification System C, (c) High magnification System A, (d) High magnification System C
6 RESULTS AND DISCUSIONS
134
(a) System A
(b) System C
(c) System A
(d) System C
Figure 6.3.2: TEM micrographs of PA-6/Cloisite 30B (5.0%wt) nanocomposites (a) Low magnification System A, (b) Low magnification System C, (c) High magnification System A, (d) High magnification System C
135
6 RESULTS AND DISCUSIONS
(a) System A
(b) System C
(c) System A
(d) System C
Figure 6.3.3: TEM micrographs of PA-6/Cloisite 15A (2.5%wt) nanocomposites (a) Low magnification System A, (b) Low magnification System C, (c) High magnification System A, (d) High magnification System C
6 RESULTS AND DISCUSIONS
136
(a) System A
(b) System C
(c) System A
(d) System C
Figure 6.3.4: TEM micrographs of PA-6/Cloisite Na+ (4.2%wt) nanocomposites (a) Low magnification System A, (b) Low magnification System C, (c) High magnification System A, (d) High magnification System C
6 RESULTS AND DISCUSIONS
137
Table 6.3.1: Surface density, specifie surface area and the degree of exfoliation of PA-6 nanocomposites (244)
Material
Volume fraction
Surface density (gm-l)
Sodium Montmorillonite
Specifie surface area (m2/gm)
Degree of exfoliation (%)
786
100
PA-6NC Ube (2%clay)
0.012
18
402
51
PA-6/Cloisite30B (5%wt) System A
0.021
55
736
94
PA-6/Cloisite30B (5%wt) System C
0.025
48
549
70
PA-6/Cloisite30B (3.1 %wt) System C
0.011
29
760
97
PA-6/Cloisite15A (2.5%wt) System A
0.009
8
245
31
PA-6/Cloisite15A (2.5%wt) System C
0.010
11
292
37
PA-6/Cloisite15A (2.5%wt) System C (compression molded)
0.010
10
265
34
In the semi-automated image analysis, a number of gray scale TEM micrographs were converted into binary black and white images, using image processing software (Adobe photoshope 6.0) and analyzed by using image analysis software (Scion Image Beta 4.02). The particle length distributions were thus obtained. The average particle lengths obtained from this analysis are given in Table 6.3.2. AIso, the aspect ratios of the particles were estimated from the average lengths and average thicknesses obtained. Table 6.3.2 shows that PA-6/Cloisite 30B(5%wt) nanocomposites made using System A and System C contain clay particles with similar lengths but different thicknesses. The aspect ratio of the nanoclay particles in System A nanocomposites is larger than that obtained in the System C nanocomposites. These data clearly demonstrate the greater
6 RESULTS AND DISCUSIONS
effectiveness of System A
In
138
producing exfoliated structures
In
PA-6/clay
nanocomposites.
80
80
~ Avg no of layers per slack: 1.8 Avg Ihickness of slacks: 4.8 nm No of stacks: 85 Interlayer distance: 4.1 nm
~
Avg no of layers per stack: 1.4 Avg Ihickness of stacks: 1.9 nm No of slacks: 94 Interlayer distance: 2.5 nm
70
70 ~
60
ê
"0
.~ 30
a;
§ 20 Z
10
o 2
3
4
5
>5
2
Figure
6.3.5:
Distribution
3
4
~
5
No of layers composing Ihe slacks. N,
No of layers composing Ihe slacks. N,
of clay
particle size
in
PA-6/Cloisite
30B(5%wt)
nanocomposite (244)
Table 6.3.2: Aspect ratios of clay particles in PA-6 nanocomposites (244)
Sample
Length (nm)
Thickness (nm)
Aspect ratio
Silicate sheet
140
0.94
149
PA-6 Nanocomposite (2% clay) Ube Inc.
196
5.6
35
PA-6/Cloisite 30B(5%wt) System A
140
1.9
74
PA-6/Cloisite 30B(5%wt) System C
139
4.8
29
PA-6/Cloisite 30B(3.1 %wt) System C
117
2.5
47
6 RESULTS AND DISCUSIONS
139
6.3.2 TEM MICROGRAPHS OF POL YSTYRENE NANOCOMPOSITES
The low magnification TEM micrographs of PS13DI/Cloisite lOA (2%wt) without compatibilizer and PS13Dl/Cloisite IDA (2%wt) with 2% Dylark compatibilizer aie shown in Figure 6.3.6 and Figure 6.3.7, respectively. The higher magnification TEM micrographs ofthese composites are shown in Figure 6.3.8 and Figure 6.3.9. According to X-ray diffraction (see the previous section) and the TEM micrographs, polystyrene/clay nanocomposites do not show exfoliated structures. The pol ymer matrix contains tactoids of varying dimensions. These are essentially microcomposites, rather than nanocomposites. In the composites without Dylark as compatibilizer (Figure 6.3.6), the clay particles are darker compared to the composites with Dylark compatibilizer (Figure 6.3.7). The compatibilizer probably swells the clay particles more and results in less compact tactoids. The high magnification micrographs of PS13DI/Cloisite IDA nanocomposites show considerable increase in gallery spacing. The TEM as well as X-ray diffraction studies show the increase in gallery spacing in PS/Cloisite 1DA nanocomposites, but these techniques cannot confirm the intercalation of the pol ymer chains into the gallery spacings. In spite of a significant increase in gallery spacing, individual platelets are not exfoliated. The increase in gallery spacing alone do es not necessarily signify the intercalation of pol ymer into the clay galleries. As stated in Section 6.2.4, the increase in gallery spacing of organoclay could be the result of expansion of gallery spacing due to high processing temperature and pressure. The organic mono layer between the galleries may become oriented and re-arranged to form a closely packed condensed state at high temperature and high pressure. This may result in an increase in gallery spacing. The addition of maleic anhydride grafted Dylark compatibilizer results in swelling of clay partic1es, but it does not produce exfoliated structures.
6 RESULTS AND DISCUSIONS
140
(a) System A
(b) System C
Figure 6.3.6: TEM micrographs of PS1301/Cloisite 10A(2%wt) at low magnification
(a) System A
Figure
6.3.7:
(b) System C
TEM
micrographs
of
compatibilizer(2%wt) at low magnification
PS1301/Cloisite
10A(2%wt)
with
Dylark
141
6 RESULTS AND DISCUSIONS
(a) System A
(b) System C
(c) System A
(d) System C
Figure 6.3.8: TEM micrographs of PS1301/Cloisite 10A(2%wt) nanocomposites (a) Low magnification System A, (b) Low magnification System C, (c) High magnification System A, (d) High magnification System C
6 RESULTS AND DISCUSIONS
Figure
6.3.9:
TEM
micrographs
142
of
PS1301/Cloisite
10A(2%wt)
compatibilizer produced with System A at different magnifications
with
Dylark
6 RESULTS AND DISCUSIONS
143
6.4 WATER ABSORPTION It is weIl known that moi sture induces property changes in polymers. At room
temperature and above. water is considered to act as a softening agent (spacer between chains). The absorbed water reduces moduli and the main glass transition temperature, enlarges fracture strain and impact strength (247). The hydrophobicity and hydrophilicity of constituents and crystallinity of the polymer are the three most important factors that determine water absorption. Non-polar polymers, such as polyethylene and polystyrene, are hydrophobic. On the other hand, the presence of amide groups in the polyamide backbone favors water absorption by forming hydrogen bonding with water molecules, owing to its higher polarity (248). Pristine montmorillonite clay is hydrophilic and the organic modifier is amphiphilic. The incorporation of organoclay into polymer matrix also affects the crystallinity depending upon the processing conditions and methods used. This made it interesting to study the water absorption of the nanocomposites.
6.4.1 PA-6 NANOCOMPOSITES The ASTM D570 water absorption test pre scribes the period of test as one day or till equilibrium absorption is achieved. Specimens of PA-6 and PA-6 nanocomposites with different clays were prepared in three different processing ways. Specimens in the form of a ribbon exiting the slit die and cooled by air fan are designated as 'extruded' samples. Ribbon samples which were cut into small pieces and compression molded are designated as 'compression molded' samples. The compression molding of PA-6 samples was carried out at 225°C. The polymer melt was cooled by water flowing in the mold cooling system. The samples that were extruded using capillary die were pelletized and injection molded into plaques. These samples are designated as 'injection molded' samples. The equilibrium water absorption of PA-6 is reported as lO.5%wt (247). The extruded samples reached the equilibrium water absorption in one day. The compression molded samples required 4 days, whereas injection molded samples took over 60 days to
6 RESULTS AND DISCUSIONS
144
reach the equilibrium water absorption. The standard deviation of the percentage water absorption values about the mean was 0.3. Figure 6.4.1 shows the water absorption of extruded nanocomposite samples after keeping the specimens in water for one day. It shows that the incorporation of the unmodified montmorillonite clay into the polymer matrix increased the water absorption. The water absorption values increased with the amount of clay used. The modifier in Cloisite 30B contains hydrophilic hydroxyl functional groups and hydrophobie organic tail. These two opposing effects left the water absorption of the PA-6 nanocomposites with this clay almost unchanged. The organophilic modifier of Cloisite 15A clay reduced the water absorption of PA-6/Cliosite 15A nanocomposites. The difference between System A and System C samples was not very significant, and it fell within the range of experimental standard deviation. Figure 6.4.2 gives the water absorption values of the compression molded specimens, after keeping them in water for one day. It can be seen that the water absorption values are significantly lower than those of slit die extruded specimens (Figure 6.4.1). Water absorption in compression molded samples did not reach equilibrium in one day (Figure 6.4.2). Water absorption values were lower for the organoc1ays than for the pristine clay. The injection molded samples of PA-6 and PA-6/Cloisite 30B (5%wt) nanocomposite absorbed only 2%wt water after immersion in water for one day. Water absorption of these specimens was checked periodically over 120 days. Figure 6.4.3 shows that during the first 3 days, water absorption increased rapidly, then more slowly to reach equilibrium. Water absorption of PA-6 nanocomposite was always lower than that of PA-6. Water absorption by the nanocomposite was 43% lower than that by unfilled PA-6 after one day and 34% lower after 30 days.
6 RESULTS AND DISCUSIONS
145
11.4 11.2 11.0
j' ~
10.8
Cloisite 308 System A -6- Cloisite 308 System Cloisite 15A System A -8- Cloisite 15A System Cloisite Na System A -ê- Cloisite Na System C
/ /
.fir--
c
0
a
10.6
"-
0
(/)
«"-
10.4
ëti
10.2
oC
CI)
P.
/
_-------I!f
~
/'
/'
/'
/'
/
3:
10.0 9.8 9.6
0'
2
6
4
Clay (%wt)
Figure 6.4.1: Water absorption of Slit Die extruded PA-6 nanocomposites after 1 day
4.0 3.8
j' ~ ~
/'
-------
/'
/'
/'
/'
/'
/'
/'
,A
3.6
c 0
~ 0
3.4
oC
«"-
3.2
ëti
3.0
(/)
CI)
3:
2.8 2.6
Cloisite 308 System A -9- Cloisite 308 System Cloisite 15A System A -8- Cloisite 15A System Cloisite Na System A -ê- Cloisite Na System C 0
2
4
6
8
Clay (%wt)
Figure 6.4.2: Water absorption of compression molded PA-6 nanocomposites (1 day)
146
6 RESULTS AND DISCUSIONS
12.------------------------------------------, .....8- - -G- - --G- - -Q- - --0-- - --0-- -
10
-e--
0 .....
/
îct-
-
8 /
P
/
/
/0 flJ/ (2)
/
Q)
1
2
~
-B- PA-6 ........- PA-6/Cloisite 308 (5%wt)
o o
10
20
30
40
50
60
70
80
90
100
110 120
Number of Days
Figure 6.4.3: Water absorption of injection molded PA-6/Cloisite 308
(5o/~wt)
nanocomposite prepared using System A
The extruded samples showed the highest water absorption values and reached equilibrium within 24 hours, whereas compression molded samples required over 4 days to reach equilibrium. This difference is mainly attributed to the differences in the thicknesses of the specimens. The slit die extruded samples were 0.4 mm thick whereas compression molded samples were 0.85 mm thick. The injection molded samples were 3.35 mm thick and required long time to reach the equilibrium moi sture absorption. The difference in the morphology and the crystalline content of the samples prepared using different processing methods may also contribute to the differences in the moi sture absorption values. The diffusion coefficients of water in PA-6 nanocomposites were computed by using Fickian theory (249). The rectangular samples, immersed in distilled water at 20°C, were removed periodically, blotted dry, weighted and re-immersed in water. The process of one-dimensional, unsteady diffusion is govemed by (250): 2
de =Dd e dt dX 2
(6.4)
6 RESULTS AND DISCUSIONS
147
where c is the concentration of the diffusing species,
t
is time, x is the position in the
diffusion direction and D is the diffusion coefficient. D is taken to be constant at a given temperature. The solution of this equation, with constant boundary conditions, yields the moi sture uptake Mt (250) Mt _ Mo -
[1- ~~(2n+1)21i2 8
ex [- D(2n + 1)21i p 41 2
2
t]]
(6.5)
where Mao is the equilibrium increase in sample mass and 21 is the sample thickness. During the initial stages of diffusion, the solution becomes:
-4(
Mt Mo -
1
Dt 1i(21f
)2
(6.6) 1l2
The diffusion coefficient can be computed from the initial slope of Mt vs. t Mo 21 using short time water-uptake data. Typical plots are shown in Figure 6.4.4. 0.4
,
ir========================:;---------:::I o
.Â
•
•
PA-6 PA-6/Cloisite 308 (3.1 %wt) System A Pa-6/Cloisite 308 (5.0%wt) System A
0.3
'E
:::2!:- 0.2 ~
:::2:
0.1
0.0
o
50
100 12 / /2/
t
150
200
(s1/2/mm )
Figure 6.4.4: Water diffusion in slit die extruded PA-6 and PA-6/Cloisite 308
nanocomposites prepared using System A
6 RESULTS AND DISCUSIONS
148
The diffusion coefficients computed from the sI opes of plots similar to those shown in Figure 6.4.4 are given in Table 6.4.1 for different PA-6 nanocomposites. Table 6.4.1: Moisture Diffusion coefficients of PA-6 nanocomposites
Diffusion Coefficient (10- 7 mm2/s) Sample
Slit die extruded
Compression molded
PA-6
5.40
6.29
PA-6/Cloisite 30B (3.1 %wt) System A
5.12
5.56
PA-6/Cloisite 30B (5.0%wt) System A
6.42
6.68
PA-6/Cloisite 30B (3.1 %wt) System C
4.16
5.31
PA-6/Cloisite 30B (5.0%wt) System C
6.04
6.68
PA-6/Cloisite 15A (2.5%wt) System A
5.10
5.55
PA-6/Cloisite 15A (4.1 %wt) System A
5.50
4.47
PA-6/Cloisite 15A (2.5%wt) System C
4.89
4.94
PA-6/Cloisite 15A (4.1%wt) System C
4.88
4.95
PA-6/Cloisite Na+ (4.2%wt) System A
4.30
5.38
PA-6/Cloisite Na+ (6.9%wt) System A
4.48
4.71
PA-6/Cloisite Na+ (4.2%wt) System C
4.65
5.45
PA-6/Cloisite Na+ (6.9%wt) System C
4.79
5.67
The diffusion coefficient of PA-6 changes marginally by the addition of clay. Overall, in case of PA-6 nanocomposites, the diffusion coefficients were higher for Cloisite 30B clay, followed by Cloisite 15A, and they were lowest for the pristine montmorillonite clay. In the case ofCloisite 30B at clay content of3.1 %wt, the diffusion coefficients were lower than those of PA-6. However, at higher clay content (5.0%wt), the diffusion coefficients were higher than pure PA-6. The diffusion coefficients for specimens produced with the System A were very close to those of specimens produced with System C. However, it appears that for composites incorporating organoclays, the
6 RESULTS AND DISCUSIONS
149
with System C. However, it appears that for composites incorporating organoclays, the diffusion coefficients were higher for System A samples. On the other hand, for the pristine clay nanocomposites, they were higher for System C samples.
6.4.2 POLYSTYRENE NANOCOMPOSITES Polystyrene is a non-polar and hydrophobie amorphous polymer. Water absorption by polystyrene is very low. In most cases, even with the incorporation of clay, water absorption was less than O.25%wt. The accuracy and the reproducibility of the water absorption test are poor, at such low level of water absorption. Figure 6.4.5 shows the results of the water absorption of the slit die extruded PSl301/Cloisite 10A nanocomposite, after keeping the specimens in water for 4 days. It can be seen that water absorption increases from O.05%wt to O.25%wt with increasing clay content. Similar results were obtained for the other organoclays. The addition of Dylark compatibilizer to the polystyrene matrix did not have significant effect on the water absorption. Water absorption for compression-molded samples was also low, in the ofO.l-0.2%wtrange. 0.25.,-------------------------, ~
PS/Cloisite 10A System A
- -9- PS/Cloisite 10A System C
)J
0.20 /
--
Ô 0.15
a....o ~
li;
__ 0
el
/
/ /
0.10
'tiî
;:
0.05
0.00 - ' - - - - - . , - - - - - - , - - - - - , . - - - - - - - , - - - - - - - i o 2 3 4
Clay (%wt)
Figure 6.4.5: Water absorption of PS13D1/Cloisite 1DA nanocomposites after 4 days
6 RESULTS AND DISCUSIONS
150
6.5 THERMOGRAVIMETRIC ANALYSIS The clay content in the nanocomposites was usually calculated from the feeder flow rates of the organoclay and the polymer. Sorne selected samples were analyzed by thermogravimetry, in order to confrrm the clay content in them. The results in Figure 6.5.1 show the gravimetric analysis of the PA-6/Cloisite 30B extruded samples, with a calculated clay concentration of 5.07%wt inorganic content based on the feeder flow rates. The results are the average of three different runs. AlI the extruded samples of PA-6 and its nanocomposites show a loss of about 3% of weight during the initial heating between 120°C to 220°C. This loss of weight is attributed to the absorbed water in the polymer. The onset of thermal degradation was at 475°C. AlI the volatile contents were 10st from PA-6 at as it reached 550°C. In the nanocomposites this occurred at 580°C. The average inorganic content in PA-6/Cloisite 30B nanocomposites, prepared using System A as weIl as System C, was 4.95%wt. The difference between the calculated and the gravimetric clay content in the nanocomposite was 0.12%wt. This suggests that the error in the calculated value is around 2.5%.· Figure 6.5.2 shows the results of the gravimetric analysis for injection molded PA-6 and PA-6/Cloisite 30B (5.07%wt) nanocomposite. The gravimetric analysis of the extruded PA-6 nanocomposite is also included for reference. The injection molded samples loose only 0.4-0.6% during the initial heating. This confrrms that the extruded samples tend to absorb more water, as indicated in the water absorption experiments. In injection molded samples, gravimetric analysis gives the inorganic content of clay as 5.3%wt. The estimated error in the calculated value is 4.5%. Figure 6.5.3 shows the gravimetric analysis for polystyrene PS1301 and its nanocomposite with Cloisite 10A. The results show that the incorporation of organoclay in polystyrene matrix increases the onset temperature of thermal degradation. The onset of thermal degradation of polystyrene was at 440°C, and that of the nanocomposite was at 480°C. The clay content in the nanocomposite by thermogravimetry was found to be 4.6%wt. The error in the calculated value was 2.2%.
6 RESULTS AND DISCUSIONS
151
100r-------__________________
''\
.. . \ ... \ \ \ \
80
:s:
40
20
o
\
- - - PA-6 PA-6/Cloisite 308 System A ........ Pa-6/Cloisite 308 System C 100
200
300
,,-----
\
400
500
.
600
Temperature oC
Figure 6.5.1: Thermogravimetric analysis of extruded PA-6 and PA-6/Cloisite 30B (5%wt) nanocomposite
100t-----~~~~--~__- -_______
" ............................... .....
:s~
"\1
"\. \
80
,
\ ..
...
\ .... \ ..• \ .•. 1 .: 1 ..
60
1·· 1 .. 1 .
40
1 .. 1 . 1 \ 1
20
o
,'-----
\
- - - PA-6 (Injection Molded) PA-6/Cloisite 308 (Injection Molded) ........ PA-6/Cloisite 308 (Extruded) 100
200
300
400
500
600
Temperature (oC)
Figure 6.5.2: Thermogravimetric analysis of PA-6 and PA-6/Cloisite 30B nanocomposite; injection molded and extruded
6 RESULTS AND DISCUSIONS
152
.......- __
100
t---------~
.... " \
80
\
\ \
\ \
\
#- 60 .E Cl
\
\ \
li)
~
\ \
40
\
\ \
\ \ \ \
20 - - - PS1301 PS/Cloisite 10A System A ........ PS/Cloisite 10A System C
o 100
200
300
\
400
'----
500
600
Temperature oC
Figure 6.5.3: Thermogravimetric analysis of PS1301 and PS1301/Cloisite 10A (4.5%wt) nanocomposites
The above results show that the calculated values of the clay contents in nanocomposites, based on feeder calibration, are close to the actual values obtained by thermogravimetric analysis. The average standard deviation in the thermogravimetric analysis of three identical samples was O.7%wt of the clay content. The estimated error in the calculated clay content is 2.2-4.5%.
6 RESULTS AND D1SCUS10NS
153
6.6 MECHANICAL PROPERTIES Polymerie materials are often reinforced by stiff flUers to improve mechanical properties. The efficiency of reinforcement depends on the filler aspect ratio, the filler mechanical properties, orientation of filler particles, and the adhesion between the matrix and the filler (251). It has been considered that single clay layers should be ideal reinforcing agents, due to their high aspect ratio and also because the nanometer filler thickness approaches the length scale of polymer chains. This section presents the results regarding mechanical properties, such as tensile strength, tensile môdulus, flexural strength, flexural modulus and impact properties of different nanocomposites produced using System A (with static mixer attachment) and System C (without static mixer attachment). Initially, the results obtained with the polyamide matrix are presented, followed by the results with the polystyrene matrix. Finally the results are fit to the different models available in the literature to calculate different parameter~ related to the physical properties of the nanocomposites and to compare these data to TEM results. 6.6.1
MECHANICAL PROPERTIES OF POLVAMIDE-6 NANOCOMPOSITES
Polyamide-6 nanocomposites were prepared using three different organoclays. Among these, Cloisite 30B is the most compatible organoclay with the polyamide matrix. It yielded a predominantly exfoliated structure, according to XRD and TEM results. Cloisite 15A has the highest gallery spacing, but it contains no polar or hydrogen bonding groups. 1t does not yield exfoliated structure. However, as seen from the TEM micrographs, the large clay particles are broken down into smaller tactoids. Cloisite Na+ is the unmodified clay, and it is the least compatible with the polyamide matrix among the three clays. The structure of the nanocomposites prepared using these clays is reflected in their mechanical behavior.
6 RESULTS AND DISCUSIONS
154
6.6.1.1 TENSILE PROPERTIES OF PA-6 NANOCOMPOSITES EXTRUDED USING THE SLIT DIE
This section reports the results regarding tensile modulus and strength of the extruded samples, prepared using the sHt die followed by cooling of the melt with air fans. The sample ribbons obtained were cut into the desired shape and used for mechanical testing. The reported tensile strength is equal to the measured yield stress. Figure 6.6.1 shows the tensile moduli of the PA-6/Cloisite 30B nanocomposites prepared using System A and System C. The tensile modulus increases with the increase in clay content. System A samples show a higher increase in the tensile modulus than System C samples. The difference becomes more significant at the higher clay content leve1s. Vaia (1) estimated that, for an exfoliated system, the mean distance between clay platelets is 100 nm for 1%vol of the filler. It decreases to 25 nm for 4%vol filler. In a compatible system of organoclay and pol ymer matrix, a fairly good distribution of the exfoliated plate1ets could be obtained by simple extrusion processing at lower concentration of clay. As the concentration of the platelets increases, the exfoliation and the distribution of the exfoHated platelets become difficult. Therefore, in the compatible system, the efficiency of Systerl1 A is more noticeable at the higher clay concentrations. Figure 6.6.2 shows that also the yie1d stress of PA-6/Cloisite 30B nanocomposites, . prepared using the System A, exhibit larger increases above a clay concentration of 4%wt. In a partially compatible system of organoclay and polymer matrix, such as PA6/Cloisite 15A, the breaking of large clay particles into small tactoids and the efficient distribution of the' tactoids in the polymer matrix become more important for the improvement of the mechanical behavior ofthe nanocomposites. Figure 6.6.3 shows that System A produces nanocomposites with larger improvement in tensile modulus. Moreover, Figure 6.6.4 shows that the System C samples exhibit no improvement in tensile strength with the addition of Cloisite 15A organoclay to the PA-6 matrix. Figure 6.6.5 and Figure 6.6.6 show the results for the mechanical properties of the nanocomposites produced with unmodified Cloisite Na+ clay. Cho and Paul (125)
6 RESULTS AND DISCUSIONS
155
reported that the composite of polyamide-6 with untreated sodium montmorillonite shows higher strength and modulus than neat polyamide-6, although there is no evidence of exfoliation of the clay platelets. Our results also show that the untreated clay raises the modulus and the strength of polyamide-6. Figure 6.6.5 shows that System A is more effective in raising the modulus of the composite than System C. System C is ineffective in this case up to a clay loading of 5.5%wt, but System A sampi es show consistent increase in modulus with the increase in clay content. Similarly, as shown in Figure 6.6.6, System A is effective in increasing the tensile strength of the composite with untreated clay. System C sampi es show a modest increase only beyond a loading of 5.5%wt ofuntreated clay. For Figures 6.6.1 to 6.6.10, showing the mechanical properties of the nanocomposites at different clay loadings, the data points are connected by smooth lines for the purpose of showing the trends. drawing the lines.
Sigmaplot polynomial fitting was used in
156
6 RESULTS AND DISCUSIONS
4400 4200
•
0
System A Systeme
4000
Cil
c..
~
3800
!Il ::l
:; 3600 "0 0
::2 3400 J!? "iii 3200 c::
~
3000 2800 2600 0
2
3
4
5
Clay (%wt)
Figure 6.6.1: Tensile Modulus of PA-6/Cloisite 30B nanoeomposites
92 90
• 0
System A System e
88
Cil 86 c..
~
! Il !Il
~
û5
84 82
"0
ID
>=
80 78 76 74 0
2
3
4
5
Clay (%wt)
Figure 6.6.2: Tensile strength of PA-6/Cloisite 30B nanocomposites
6 RESULTS AND DISCUSIONS
157
4000
•
3800
System A System C
0
-5
cu 3600
Q.. rJ)
::J
"5
3400
"0
0 ~ ~
3200
"iii
c: Q)
1-
3000 2800 2600 0
2
3
4
Clay (%wt)
Figure 6.6.3: Tensile modulus of PA-6/Cloisite 15A nanocomposites 86
• 84
System A System C
0
Cil 82 Q.. ~
---enen
-
Q) 80 ....
en
"0
ëi)
>=
--- ~___!_-----------î--f--î î l
78
76
74~--~--------~------~--------~--------r-~
o
2
3
4
Clay (%wt)
Figure 6.6.4: Tensile strength of PA-6/Cloisite 15A nanocomposites
158
6 RESULTS AND DISCUSIONS
3300 • o
3200
as a..
3100
I II :J
3000
"0 0 ~ ~
2900
-
System A System C
~
:;
"iii
c:: Q)
1-
2800 2700 2600
o
2
4
6
8
Clay (% wt)
Figure 6.6.5: Tensile modulus of PA-6/Cloisite Na+ nanocomposites 86~------------------------------------------~
84
• o
System A System C
-6-
cu 82
a..
III III
e-n
Q) 80 ....
"0
ID
>=
78
76
74 0
2
4
6
Clay (% wt)
Figure 6.6.6: Tensile strength of PA-6/Cloisite Na+ nanocomposites
8
6 RESULTS AND DISCUSIONS
159
Figure 6.6.7 compares the tensile moduli of the nanocomposites produced using System A. It shows that Cloisite 15A is the most effective clay at lower concentrations. Cloisite 15A contains higher levels of organic modifier (43%) and has higher gallery spacing (3.15 nm) than Cloisite 30B (organic content 30%, and gallery spacing 1.85 nm). It may be easier to break down the large particles of Cloisite 15A clay into smaller tactoids than those of Cloisite 30B, but due to the absence of polar groups in Cloisite 15A, it may be difficuilt to exfoliate it. Since Cloisite 30B exfoliates in PA-6 matrix, it shows improvement in the modulus values at higher loadings. Cloisite 15A does not show enhancement in the tensile modulus after 3%wt concentration. In the case of untreated Cloisite Na+ clay, modest improvement in the tensile modulus is observed with the increase in clay loading only with the System A. The tensile moduli of the nanocomposites with different clays prepared using the System C (Figure 6.6.8) show similar trends but the overall values are significantly lower than those of the System A. Comparing Figure 6.6.9 and Figure 6.6.10 it can be concluded that the System A is more effective in improving tensile strength of the nanocomposites. Cloisite 15A and Cloisite Na+ nanocomposites produced by using System C show little improvement in tensile strengths. Cloisite 30B seems to be the most effective organoclay in improving tensile strength of the polyamide-6 nanocomposites.
160
6 RESULTS AND DISCUSIONS
4400 4200 4000
as a..
-
•• •
Cloisite 308 Cloisite 15A Cloisite Na
3800
~
CI)
3600
:::J
"3 "0
0
3400
~
.!Q 3200 "iii c:
~
3000
2800 2600 2400 0
2
4
6
8
Clay (% wt)
Figure 6.6.7: Tensile moduli of PA-6 nanocomposites produced with different clays
using System A 4400 4200 4000
as a..
5 CI)
Cloisite 308 Cloisite 15A Cloisite Na
0 0 t::.
3800 /
m /~
3600
:::J
"3 "0
0
~
/
/
3400
~-t~~î
~
.!Q 3200 "iii c: Q) 3000 1-
2800
ll/!/
-t::::::::::/
! -----1 --4---!-----j----lA"
2600 2400 0
2
4
6
8
Clay (%wt)
Figure 6.6.8: Tensile moduli of PA-6 nanocomposites produced with different clays
using System C
6 RESULTS AND DISCUSIONS
161
92 90 88
«1 a..
6 U) U)
...
••
..
Cloisite 308 Cloisite 15A Cloisite Na
86
84
CD
ëi5
82
"0
Qi
>=
80 78 76 74 0
2
4
6
8
Clay (%wt)
Figure 6.6.9: Tensile strength of the PA-6 nanocomposites with different clays using
System A 92 90
0
Cloisite 308 Cloisite 15A Cloisite Na
0
t:. 88
«1 a..
-...
J/f-ï
86
~
U) U)
-
/
84
CD
en
82
"0
Qi
>=
80
/
/
lAi 1---'-~
-~~~!~---~=!~~.-/'/'
78 76
I/
1
-!"
/'
/'
~-------l-Y-
74 0
2
4
6
8
Clay (% wt)
Figure 6.6.10: Tensile strength of the PA-6 nanocomposites with different clays using
System C
6 RESULTS AND DISCUSIONS
162
Table 6.6.1 provides a comparison among the tensile moduli and the tensile strengths of polyamide-6 nanocomposites, prepared using different clays at similar clay contents. The clay contents used for the comparison in this Table are 2%wt and 4% wt. These results show that the System A is more effective than the System C in improving the mechanical behavior of the nanocomposites.
Table 6.6.1: Comparison of tensile moduli and tensile strengths of different nanocomposites produced using System A and System C
Inorganic Clay Content (%wt)
Cloisite 30B System A
1
System C
Cloisite 15A System A
1
System C
Cloisite Na + System A
1
System C
Tensile Modulus (MPa) (average standard deviation 60 MPa) 0
2760
2760
2720
2740
2730
2750
2.0
3080
2990
3330
3140
2860
2750
4.0
3670
3460
3580
3180
3040
2770
Tensile Strength (MPa) (average standard deviation 0.6 MPa) '«
0
75.5
75.8
76.0
76.0
76.0
76.0
2.0
79.7
79.6
81.2
76.8
78.4
76.1
4.0
86.8
86.6
82.8
77.2
80.9
76.6
6 RESULTS AND DISCUSIONS
163
6.6.1.2 EFFECT OF COMPRESSION MOLDING ON TENSILE PROPERTIES OF PA-6 NANOCOMPOSITES In semi-crystalline polymers, the crystallinity of the polymer matrix affects the mechanical behavior. The crystaUization behavior can be directly influenced by the presence of the nanoclay particles. The degree of orientation of polymer crystallites is found to increase linearly with increasing clay
~ontent
(252). AIso, it has been found
that melt processing has an impact on the morphology of polymer/clay nanocomposites (124). Water absorption results in Section 6.4.1 show that water absorption by compression molded sample is slower than that in samples obtained by slit die extrusion. This is probably the result of compact structure and higher crystallinity induced during the compression molding. In this section, we study the effect of compression molding on the mechanical properties of polyamide-6 nanocomposite samples. The sample ribbons extruded with the slit die were cut into small pieces and placed in a rectangular mold. Compression molding was carried out at 225°C followed by cooling of polymer melt by passing tap water through the cooling jacket. The samples were cut into the desired shape and used for testing the mechanical properties. Figure 6.6.11 compares the tensile moduli of slit die extruded samples and the corresponding compression molded samples. AlI the nanocomposite samples show an increase in the tensile modulus values after compression molding. The samples prepared using System A still have higher tensile modulus than those prepared using System C. A similar trend can be observed for the increase in the tensile strength after compression molding (Figure 6.6.12). The highest tensile strength of 95 MPa is obtained for the PA6/Cloisite 30B (5%wt) compression molded samples. System A samples are superior to System C samples in this case also. Sheng et al. (251) also reported an improvement in the mechanical properties of nanocomposites after post-processing.
164
6 RESULTS AND DISCUSIONS
5000~==============~------------------------1
mw#ijjJ Compression Molded lZZI Extruded
4500
tii' ~
......
4000
en
::l
"3
"8
3500
~ ~
.~
3000
~ 2500
PA-6
3.1% Clay
5.0% Clay
System A
3.1% Clay
5.0% Clay
System C
Figure 6.6.11: Tensile moduli of extruded and compression molded PA-6/Cloisite 308 nanocomposites
100 M,W[M Compression Molded
IZZJ Extruded
95
tii'
a...
90
6
en en
.....~
85
Cf)
"0
Qi
>=
80
75
70 PA-6
3.1% Clay
5.0% Clay
System A
3.1% Clay
5.0% Clay
System C
Figure 6.6.12: Tensile strengths of extruded and compression molded PA-6/Cloisite 308 nanocomposites
6 RESULTS AND DISCUSIONS
165
Aiso in the case of PA-6/Cloisite 15A nanocomposites, compression molding produces an increase in the tensile modulus (Figure 6.6.13). However, the results are different for the tensile strength test (Figure 6.6.14). The tensile strength increases after compression molding, only in case of 2.5%wt nanocomposites. For the System A sample with of 4.1 %wt clay content, the tensile strength decreases after compression molding. For the System C sample, while it increases after compression molding, the absolute value of tensile strength is lower than that for the 2.5%wt sample. Cloisite 15A clay contains a high amount of organic modifier (43%). During compression molding, the excess organic modifier as weIl as the modifier between the clay gaIleries might mix with the polymer melt and act as plasticizer. This will obviously be more pronounced at higher clay content. The reduced tensile strength of the Cloisite 15A nanocomposites may be attributed to this plasticizing effect. In this case also, System A samples show higher tensile modulus and higher tensile strength than System C samples.
4000 téfJliMl Comp Molded
3800
[Z2] Extreuded
(ij' 3600
a.
5
en 3400
::::1
:; "C
0
::2:
3200
~
iii 3000 c: CD 1-
2800 2600 PA-6
2.5%
4.1%
2.5%
4.1%
Clay
Clay
Clay
Clay
System A
System C
Figure 6.6.13: Tensile moduli of extruded and compression molded PA-6/Cloisite 15A nanocomposites
6 RESULTS AND DISCUSIONS
166
88 WN&d Comp Molded l2':Z:l Extruded
86 84
ê?
c..
5 CI) CI)
-
82 80
~
en
78
"0
ID
>=
76 74 72 70 PA-6
2.5% Clay
4.1% Clay
System A
2.5% Clay
4.1% Clay
Systeme
Figure 6.6.14: Tensile strength of extruded and compression molded PA-6/Cloisite 15A nanocomposites
In the case of untreated clay Cloisite Na+ the mechanical properties improve
substantially after compression molding (Figure 6.6.15 and Figure 6.6.16). The tensile moduli of System A as weIl as System C samples increase after compression molding. AIso, the difference among the modulus values for the samples obtained using the two processing systems becomes narrow. The tensile strength of the Cloisite Na+ nanocomposite samples with 4.2%wt clay increases significantly after compression molding, and it becomes even higher than that of Cloisite 15A nanocomposites. However, with further addition of clay (6.9%wt), the tensile strength does not improve. It seems that the tensile strength obtained by using 4.2%wt of untreated clay (88 MPa) is
the limiting value for this nanocomposite. In this case, the difference between the tensile strengths of the System A and the System C samples diminishes after compression molding.
167
6 RESULTS AND DISCUSIONS
3800~============~----------------------------1
3600
~Hih!;M Com p Molded [ZLI Extruded
t? [l. 3400 ~ CI)
::J
"3 3200 "0
o
~
.5!! 3000 "iii c:
~
2800
2600 PA-6
4.2%
6.9%
4.2%
6.9%
Clay
Clay
Clay
Clay
System A
System C
Figure 6.6.15: Tensile strength of extruded and compression molded PA-p/Cloisite Na+ nanocomposites 95~------------------------------------------~
fikwwU Comp. Molded
E::ZJ Extruded 90
t? [l. ~ 85 CI) CI)
~
CI)
"0
Qi
80
>= 75
PA-6
4.2%
6.9%
Clay
Clay
System A
4.2%
6.9%
Clay
Clay
System C
Figure 6.6.16: Tensile strength of extruded and compression molded PA-6/Cloisite Na+ nanocomposites
6 RESULTS AND DISCUSIONS
168
Overall the tensile moduli of all the PA-6 nanocomposites increase after compression mo1ding. In the case of Cloisite 30B and C10isite Na+, the tensi1e strength increases after compression mo1ding. The tensile modu1us at higher content of Cloisite 15A is reduced after compression mo1ding, possib1y due to the plasticizing effect of excess modifier. System A samp1es still perform better in most cases. Table 6.6.2 summarizes the tensile properties of the slit die extruded and the compression molded PA-6 nanocomposite samples with the different clay types. The average standard deviation of tensile modulus around the mean was 60 MPa and that of tensile strength was 0.9 MPa.
Table 6.6.2:
Tensile properties of extruded and compression molded PA-6
nanocomposites with different clays
Tensile Modulus (MPa) Clay (%wt)
Extruded with slit die
Tensile Strength (MPa)
Compression Molded
Extruded with slit die
Compression Molded
2960
76.0
80.0
3780 4340 3560 4210
83.9 89.2 83.2 87.0
91.7 94.9 90.4 91.6
3600 3730 3400 3650
81.5 82.6 76.0 77.6
83.9 81.8 82.5 80.2
3450 3620 3360 3600
81.0
88.2
84.5 76.8 80.0
87.7 86.7 88.2
PA-6
o
2720 PA-6/Cloisite 30B Nanocomposite 3.1 3260 System A 5.0 4140 3.1 3140 System C 5.0 3795 P A-6/Cloislte 15A nanocomposites 2.5 3420 System A 4.1 3555 2.5 3150 System C 4.1 3250 PA-6/Cloisite Na+ nanocomposite 4.2 3120 System A 6.9 3225 2760 4.2 System C 2860 6.9
6 RESULTS AND DISCUSIONS
169
6.6.1.3 EFFECT OF CLA y LOADING ON ELONGATION AT BREAK Percent elongation at break for the nanocomposite samples were recorded while testing the specimens for tensile strength. AU the PA-6 and and PA-6 nanocomposite samples were vacuum dried at 80°C for 24 hrs prior to testing. Slit die extruded specimens were 0.4 mm thick, compression molded specimens were 0.85 mm thick and injection molded specimens were 3.35 mm thick. The difference in the cross sections of the samples as weU as the difference in their processing methods results in different elongation at break values. The results are shown in Table 6.6.3. The standard deviation in percent elongation values, between the the tested samples, was 2-4.
Table 6.6.3: Percentage elongation at break for PA-6 nanocomposites
% Elongation at break
Sample
;'
SUt die extruded
Compression molded
PA-6
28
43
PA-6/Cloisite 30B (3.1 %wt) System A
25
34
PA-6/Cloisite 30B (5.0%wt) System A
25
21
PA-6/Cloisite 30B (3.1 %wt) System C
24
34
PA-6/Cloisite 30B (5.0%wt) System C
19
14
PA-6/Cloisite 15A (2.5%wt) System A
23
36
PA-6/Cloisite 15A (4.1 %wt) System A
22
25
PA-6/Cloisite 15A (2.5%wt) System C
28
33
PA-6/Cloisite 15A (4.1 %wt) System C
27
27
PA-6/Cloisite Na+ (4.2%wt) System A
24
30
PA-6/Cloisite Na+ (6.9%wt) System A
21
14
PA-6/Cloisite Na+ (4.2%wt) System C
31
25
PA-6/Cloisite Na+ (6.9%wt) System C
28
18
f
'
6 RESULTS AND DISCUSIONS
170
The injection molded PA-6 specimen gave 73% elongation at break, whereas injection molded PA-6/Cloisite 30B(5%wt) specimen resulted in 23% elongation in flow direction and 32% elongation across the flow direction. In the case of slit die extruded samples; the incorporation of clay does not alter
elongation at break to any significant extent. In case of compression molded samples, however, elongation decreases with the addition of clay. In case of injection molded samples, the elongation at break decreases substantially. 6.6.1.4 EFFECT OF EXTRUDER SCREW SPEED ON THE TENSILE PROPERTIES OF PA-6 NANOCOMPOSITES
Extruder screw speed is an important factor in the melt processing of polymers by twin screwextrusion,. Generally, higher RPM (revolutions per minute) of the screw imparts higher shear to the polymer melt and improves mixing. The addition of the static mixer assembly to the extruder in melt processing of polymer/clay nanocomposites produces high levels of distributive and chaotic mixing. This might reduce the dependence of mixing efficiency on screw RPM, in the case of the System A processing. PA-6/Cloisite 30B nanocomposites were prepared with two different clay concentrations at different screw speeds, using System A and System C. The samples were extruded using the slit die and air cooling. Figure 6.6.17 shows the effect of screw RPM on the tensile strength of 3.1 %wt Cloisite 30B nanocomposites. In case of System A, the maximum strength of the nanocomposites is realized at 100 RPM screw speed. On the other hand, the tensile strength of System C samples increases gradually with the increase in screw speed from 50 RPM to 250 RPM. This shows that System A is effective in both dispersive and distributive mixing of the clay platelets in the polymer matrix, thus reducing the need for higher screw speeds. Figure 6.6.18 shows the results for 5.0%wt Cloisite 30B nanocomposites. System A gives nanocomposites with higher tensile strength than those of System C, for all the screw speeds ranging from 50 to 250 RPM ..
171
6 RESULTS AND DISCUSIONS
90
• 0
System A System C
l
85
l
,V"
as a..
6
en 80 en
-
././/!. . . . - . . .
~
en
./
"C
Ci)
>=
75
r
70
o
/
T
t -- __
.
I.
4
./
/-f
/
50
100
150
200
250
300
Extruder RPM
Figure 6.6.17: Effect of extruder RPM on the tensile strength of PA-6/Cloisite 30B (3.1 %wt) nanocomposites 96
•
94
0
System A System c
92
as a.. 6
90 88
///f/////1-----~-----f
en 86 en CD ë75 84 "C
Ci)
>=
82
r
80 78 76
/
74
0
50
100
150
200
250
300
Extruder RPM
Figure 6.6.18: Effect of extruder RPM on the tensile strength of PA-6/Cloisite 30B (5.0%wt) nanocomposites
/
6 RESULTS AND DISCUSIONS
172
6.6.1.5 IMPACT PROPERTIES OF PA-6 NANOCOMPOSITES
Impact performance of polymerie materials is of great importance in many applications. In the case of polymer blends, impact strength is a good indicator of the quality of adhesion between blend components (253). Kharbas et al. (254) reported that for nanocomposites, the impact strength of microcellular injection molded parts, can be greater than that of solid parts, provided that the microceU sizes and ceU densities are controlled weIL In general they observed that nanocomposites have typically lower impact strength than their neat resin counterparts. Fomes et al. (126) found that the impact strength of PA-6 nanocomposites with a compatible organoclay is relatively independent of clay concentration for high and medium molecular weight composites, but gradually decreases for low molecular weight composites. Wang et al. (255) used partially compatible organoclay, Nanomer 1.30 TC to prepare PA-6 nanocomposites. They reported a significant decrease in Izod impact strength of nanocomposites with the addition of clay. Earlier work on the impact properties of nanocomposites (7, 34, 256) showed that PA-6/clay nanocomposites are superior to the PA-6 matrix, in terms of strength and modulus without sacrificing impact strength. Overall, nanocomposites do not show significant improvement in the impact properties, but it would be desirable that nanocomposites do not suffer a significant decrease in impact strength by the addition of clay to the polymer matrix. Figure 6.6.19 shows the impact strength of PA-6/Cloisite 30B nanocomposites in terms of the ultimate force required to break the specimen. It can be seen that the impact strength of the nanocomposites prepared using System A gradually increases with the addition of clay. On the other hand, the impact strength of the specimens prepared using system C start decreasing above 3%wt clay content. These results also show that system A is effective even at relatively high clay content.
6 RESULTS AND DISCUSIONS
173
550,-------------------------------------------~
-*- System A 500
Z
- -6- System C
450
'-" Q)
--------
~
//t--
.......
.......
0
u. 400 Q) 1ii
. . .!---!
'~
E
E
::> 350
300
250 0
2
3
4
5
6
Clay (%wt)
Figure 6.6.19: Impact strength of PA-6/Cloisite 308 nanocomposites
550
-*- System A --6- System
500
Z
450
'-" Q)
~
0
u. 400 Q) 1ii
E
E
::> 350
300
250 0
2
3
4
5
Clay (%wt)
Figure 6.6.20: Impact strength of PA-6/Cloisite 15A nanocomposites
174
6 RESULTS AND DISCUSIONS
Figure 6.6.20 gives the impact strength ofPA-6/Cloisite 15A nanocomposites. In this case, the impact strength of the samples produced using System A started decreasing above 3%wt clay content. However, System C samples exhibited graduai decrease in impact strength with the increase in clay content. Both TEM and XRD results show that Cloisite 15A is not effective in producing exfoliated structure and that the bulk of the nanocomposite contains clay tactoids and agglomerates. This influences the impact properties. Yet, the effective dispersion and distribution obtained by System A still yields better impact properties than obtained by System C. The impact strength of nanocomposites containing untreated clay (Figure 6.6.21 decreased with the addition of clay. Specimens prepared with both the Systems A and C behaved in a similar manner. Untreated clay undergoes negligible exfoliation or breakdown into smaller tactoids. The presence of large clay particles in the polymer matrix lowers the impact strength of the composites. Overall, the nanocomposites with Cloisite 30B showed higher impact strength than nanocomposites incorporating Cloisite 15A or Cloisite Na+. 550
500
-Z
~ System A --9- System C
450
(J)
~
..........
~ 400 (J)
ro
..................... ~--
E
E ::J
~
350
~ ~
300
250 0 ,-..-
2
3
4
5
6
7
Clay (%wt)
Figure 6.6.21: Impact strength of PA-6/Cloisite Na+ nanocomposites
r
'
6 RESULTS AND DISCUSIONS
175
6.6.1.6 INJECTION MOL DING OF PA-6 NANOCOMPOSITES PA-6/Cloisite 30B (5%wt) nanocomposite was prepared using System A and the capillary die. The sample strand was cooled in a water trough and subsequently pelletized. The pellets were injection molded into plaques of dimensions 100x62x3 mm. Specimens of the required size were cut from the plaques in the flow direction and across the flow direction, in order to evaluate the mechanical properties. The mechanical properties of injection molded PA-6 and PA-6 nanocomposite are shown in Table 6.6.4. The standard deviations of tensile modulus, tensile strength, flexural modulus and flexural strength about the respective means were 125, 1.1,60, 1.5 MPa, respectively. The above mechanical properties of PA-6 in the flow and across flow directions do not differ significantly. In the case of nanocomposite, they are higher in the flow direction. This can be attributed to the alignment of the clay platelets in the flow direction (257,258).
System A processing yielded improvement of 72% in tensile
modulus and 75% in flèxural modulus, by the addition of 5%wt of organoclay to the PA6 matrix. The tensile strength improved from 79.4 MPa to 95.9 MPa and the flexural strength improved from 123.5 MPa to 163.9 MPa. Table 6.6.4: Mechanical properties of injection molded PA-6 nanocomposite
Tensile Modulus (MPa)
Tensile Strength (MPa)
Flexural Modulus (MPa)
Flexural Strength (MPa)
PA-6 (In Flow Dirn.)
2735
79.4
2803
123.5
PA-6 (Across Flow Dim.)
2671
79.3
2790
124.6
PA-6 Nanocomposite
4700
95.9
4890
163.9
(72 % Î)
(21 % Î)
(75 % Î)
(33 % Î)
4120
90.7
4265
147.6
(54 % Î)
(14 % Î)
(53 % Î)
(18 % Î)
5%wt clay (In Flow Dirn.) PA-6 Nanocomposite 5%wt clay (Across Flow)
6 RESULTS AND DISCUSIONS
176
The properties of nanocomposites depend on the method of processing. Figure 6.6.22 and Figure 6.6.23 show the tensile moduli and the tensile strengths of PA-6 and PA-6/Cloisite 30B (5%wt) nanocomposites prepared using System A, followed by compression molding or injection molding. Compression molding samples were based on ribbons prepared using slit die. The ribbons were cut into small pieces, and the pieces were randomly placed on the molel. Thus, compression molded sample contained clay platelets with no specific orientation and the sample may be considered isotropic. Injection molding samples were based on strands prepared using capillary die. The strands were cut into pellets, which were used for molding .. During injection molding, the platelets tend to be oriented in the flow direction. The tensile modulus and strength of PA-6 increase after compression molding. Since the cooling of the polymer melt is slow in compression molding, the resulting specimen has higher crystallinity. Aiso the samples are crystallized at higher pressure, which gives close packing of the chains. This results in higher modulus and strength. During injection molding, cooling is rapid, compared to compression molding. This may result in lower crystallinity. Thus, the above properties are somewhat lower than for compression molded specimens. In the case of PA-6 nanocomposites also, compression molding increases tensile modulus and strength. Besides the increase in crystallinity, this may be attributed to the improved dispersion or sample uniformity, as a result of the additional processing steps (251). The tensile modulus and strength of the injection molded PA-6 nanocomposite specimen in the flow direction were higher than those for the compression molded specimen. These properties were lower across the flow direction, when compared to those of the compression molded specimens. This strongly indicates the orientation of clay platelets in the flow direction and the dependence of mechanical properties on the orientation of the platelets.
6 RESULTS AND DISCUSIONS
177
500°T.=======================~------------------1
c::::J Extruded with Slit Die ~
4500
~ ~
Compression Molded Injection Molded (in flow dirn.) Injection Molded (across flow dirn.)
(ù
~
4000
!/)
::J
:; '0 o
3500
~
.!!l .~ 3000 CI)
12500
PA-6
PA-6/Cloisite 308 (S'Yowt)
Figure 6.6.22: Tensile modulus of PA-6/Cloisite 308 (5%wt) nanocomposites processed using System A followed by compression molding or injection molding
100T.=========================~-------------------
c::::J Extruded with Slit Die ~
95
~
~
(ù
90
~ gj !!!
85
a..
Compression Molded Injection Molded (in flow dirn.) Injection Molded (across flow dirn.)
Ci) '0 Ci)
>=
80
75
PA-6
PA-6/Cloisite 308 (S'Yowt)
Figure 6.6.23: Tensile strength of PA-6/Cloisite 308 (5%wt) nanocomposites processed using System A followed by compression molding or injection molding
6 RESULTS AND DISCUSIONS
178
6.6.2 MECHANICAL PROPERTIES OF POLYSTYRENE NANOCOMPOSITES
Polystyrene nanocomposites were prepared using three organoclays: Cloisite lOA, Cloisite 15A and Cloisite 30B. Cloisite lOA contains organic modifier with a styrene functionaI group. This organoclay is expected to be the most compatible among the above organoclays with the polystyrene matrix. Polystyrene is non-polar and does not contain any functionaI groups that are capable of forming hydrogen bonds. The adhesion between clay particles and the polymer chains is poor, in this case. To increase the compatibility between the polymer matrix and the organoclay, maIeic anhydride grafted polystyrene, Dylark 332, was added to the polymer. XRD results showed that the gaIlery spacing of Cloisite 10A organoclay increased from 2.05 nm to 3.84 nm with System A processing and to 3.4 nm with System C processing. For other organoclays, the change in gallery spacing was not significant. TEM micrographs show that the organoclays are not exfoliated in the polystyrene matrix. However, gallery spacing between the clay platelets in the tactoids is expanded, in the case of organoclay Cloisite 10A. Figure 6.6.24 shows the tensile moduli of PS1301lCloisite lOA nanocomposites with different clay loadings for Systems A and C. The tensile modulus increases slightly (about 7%) up to 2%wt of clay loading and then starts decreasing with further addition of the clay. The difference between Systems A and C is not significant and faIls within the range of the standard deviation. Figure 6.6.25 summarizes the tensile strength data for PS1301lCloisite lOA nanocomposites. The tensile strength decreases with the addition of organoclay to the polymer matrix. The vaIues are higher for System C. The lower values of the tensile strength for System A are probably due to the degradation of polymer chains, which are exposed to higher residence time in System A.
6 RESULTS AND DISCUSIONS
3400
as a..
179
___ System A
- -e-
System C
3200
~ (/J
""
:J
"3
"8
3000
~
..!!l "iij
c: 1-
ID
2800
'r
" "
/
2600
2
0
3
4
Clay (%wt)
Figure 6.6.24: Tensile modulus of PS1301/Cloisite 10A nanocomposites
52~------------------------~--~========~ ___ System A
50
- -e-
System C
48
as a..
-- ------
46
~ 44
----
(/J
~ 42
û5 "C
40
>=
38
ëD
36 34 32~---r---------r---------r---------r--------~
o
2
3
4
Clay (%wt)
Figure 6.6.25: Tensile strength of PS1301/Cloisite 10A nanocomposites /
6 RESULTS AND DISCUSIONS
180
Similar results are obtained for polystyrene nanocomposites prepared using maleic anhydride compatibilizer, Dylark 332. Figure 6.6.26 summarizess the tensile moduli of PS1301/Dylark(2%wt)/Cloisite lOA nanocomposites. The tensile modulus increases (about 7.5%) up to 2%wt clay loading and decreases with further addition of clay. Here also, the difference between System A processing and C samples is insignificant. Figure 6.6.27 shows that the tensile strength decreases with the addition of clay to the polymer matrix. The values of the tensile strength are lower for System A than for System C. Since polystyrene is a glassy polymer at room temperature, the flexural modulus and the flexural strength of the extruded ribbons were evaluated. The flexural modulus of PS1301 nanocomposites shows a similar trend to that of the tensile modulus (Figure 6.6.28). After the initial slight increase (about 8%) up to 2%wt clay loading, the flexural modulus starts decreasing with further addition of clay. The flexural strength (Figure 6.6.29) increases slightly up to 2%wt clay, and then starts decreasing. Table 6.6.5 and Table 6.6.6 give a summary of the tensile properties of different polystyrene nanocomposites with organoclays Cloisite lOA, Cloisite 15A and Cloisite 30B, with and without Dylark compatibilizer, processed using Systems A and C. The average standard deviation of the tensile modulus about the mean values was 200 MPa and that for the tensile strength is 2.0 MPa. There is no significant improvement in the tensile modulus as weIl as the tensile strength of the composites prepared using different clays. In spite of increasing the gallery spacing, Cloisite lOA does not contribute to the mechanical properties of the polystyrene nanocomposites. It seems that the Dylark compatibilizer also does not increase the adhesion between the clay and the polymer matrix to any significant extent. In polystyrene nanocomposites, the clay tactoids are large This results in the failure to increase the modulus. Since the polymer is non-polar and has no hydrogen bonding ability, the tensile strength of the composites is low. The compression molding of polystyrene nanocomposites showed no further improvement in mechanical properties
6 RESULTS AND DISCUSIONS
181
3600 ____ System A -()-
System C
3400
êif
a..
5
(/)
3200
-----
::J
"S
"
"-
"8 ~ ~
3000
"iii
c: Q)
1-
2800
2600~--~--------~--------~---------.--------~
o
2
3
4
Clay (%wt)
Figure 6.6.26: Tensile modulus of PS1301/Dylark (2%wt)/Cloisite 10A nanocomposites
50 ____ System A
48
- ()- System C
46
êif 44
a..
5 (/)
~
Ci)
42
-
_ A, . . . . . . . Y .....
40
,
"0
Qi
>=
38 36 34 32 0
2
3
4
Clay (%wt)
Figure 6.6.27: Tensile strength of PS1301/Dylark (2%wt)/Cloisite 10A nanocomposites
/
6 RESULTS AND DISCUSIONS
182
3800 _ System A --9- System C
3600
Cil
a..
::::?! ( /)
::::1
3400
"3
"8
----
::::?!
li!
3200
j
LI..
3000
2800~---r--------~--------~------~--------~
o
4
3
2
Clay (%wt)
Figure 6.6.28: Flexural modulus of PS1301/Cloisite 10A nanocomposites
60~------------------------~========~1
_
System A
- -9- System C
Cil
55
a..
6
.c:
ë> r::: ~
50
êi5 ëij
~
iL 45
40~---r--------~--------.--------.--------~
o
2
3
4
Clay (%wt)
Figure 6.6.29: Flexural strength of PS1301/Cloisite 10A nanocomposites ,/
183
6 RESULTS AND DISCUSIONS
Table 6.6.5: Tensile properties of PS nanocomposites prepared using System A
Cloisite 10A Clay Tensile content Modulus
PS1301 PS130Il 2% Dylark PS3900 PS3900/ 2% Dylark
Cloisite 15A
Cloisite 30B
Tensile Strength
Tensile Modulus
Tensile Strength
Tensile Modulus
Tensile Strength
{%wt~
{MPa~
{MPa~
{MPa~
{MPa~
{MPa~
{MPa~
0 1.1 2.6 0 1.1 2.6 0 1.1 2.6 0 1.1 2.6
2976 3068 3096 3015 3216 3220 2823 2896 2890 2952 2859 2954
47.0 42.8 40.4 . 41.0 42.9 37.3 29.8 31.2 28.4 33.2 30.9 30.0
2976 3106 3369 3015 3285 3400 2823 2951 3130 2952 2949 2915
47.0 45.4 43.7 41.0 44.4 45.0 29.8 31.7 32.8 33.2 31.5 31.7
2976 3254 3280 3015 3064 3182 2823 3086 3036 2952 2645 2746
47.0 45.6 44.1 41.0 45.5 43.7 29.8 32.8 33.6 33.2 33.0 32.4
Table 6.6.6: Tensile properties of PS nanocomposites prepared using System C
Cloisite 10A
PS1301 PS1301/ 2% Dylark PS3900 PS3900/ 2% Dylark
Clay Tensile content Modulus {%wQ (MPa~ 0 2827 1.1 3010 2.6 2933 2980 0 1.1 3165 2.6 3180 2772 0 3342 1.1 3171 2.6 2961 0 3284 1.1 2.6 3018
Cloisite 15A
Cloisite 30B
Tensile Strength
Tensile Modulus
Tensile Strength
Tensile Modulus
Tensile Strength
{MPa~
{MPa~
{MPa~
{MPa~
{MPa~
50.1 46.6 44.4 43.5 43.4 41.4 31.9 28.3 29.9 28.0 26.8 27.3
2827 3031 2984 2980 3002 2983 2772 3032 2929 2961 2927 2969
50.1 46.4 43.1 43.5 46.6 44.2 31.9 30.1 32.7 28.0 30.1 25.7
2827 2859 2788 2980 2903 3016 2772 2858 2959 2961 3027 3010
50.1 45.6 44.1 43.5 43.3 45.4 31.9 30.3 32.3 28.0 31.8 24.8
6 RESULTS AND DISCUSIONS
184
6.6.3 MODELING TENSILE PROPERTIES OF POLYAMIDE AND POLYSTYRENE NANOCOMPOSITES
The properties of particulate filled polymers are detennined by the composition and characteristics of the individual components, the interactions between the components and the structure of the composite (259). These four factors are equally important and their effects are interconnected. The specific surface area of the filler, for example, detennines the size of the contact surface between the filler and the polymer. Surface energies influence structure, and composition impacts on both properties and the mode of deformation. The reinforcing effect of the filler increases with decreasing matrix stiffness. In a stiff matrix like polystyrene, large stresses develop around the inclusions and the probability of dewetting increases. In such matrices, dewetting is usually the dominating micro-mechanical deformation process. The extent of stress transfer also depends on the strength of adhesion between the components. If this is weak, the separation of the interfaces takes place even under small externalload (260, 261). Many of the particles used to reinforce polymers are anisotropic. The fillers and reinforcements are often differentiated by their degree of anisotropy (aspect ratio). Platelet-like fillers, like clays,· reinforce polymers more than spherical fillers. Anisotropie fillers may become oriented during processing, thus enhancing the reinforcement effect in the direction of orientation.
This produces an anisotropic
distribution of composite properties. Surface free energy (surface tension) of the fillers detennines both matrix/filler and particle/particle interaction. The former has a pronounced effect on the mechanical properties, and the latter detennines aggregation. Both interactions are modified by surface treatment. Particle size, specific surface area and surface energetics influence interfacial interactions. Separation of the matrixlfiller interface is easy, in the case of large particles and weak adhesion. Debonding occurs under the effect of small external /'
load. Small particles form aggregates whieh cause deterioration in mechanical properties of the composites. Specific surface area, which depends on the particle size distribution
6 RESULTS AND DISCUSIONS
185
of the tiller, determines the size of the contact surface between polymer and the tiller. The size of this surface plays a crucial role in interfacial interactions and the formation of the interphase. Exfoliated silicate platelets possess very large surface areas and high aspect ratios. In addition, these platelets have exceptionally high modulus in tension, relative to most polymers and even many other tiller types. Such properties enable the platelets, when dispersed in a polymer matrix, to carry a signiticant part of the applied load. In the following sections, different composite models are used to predict the characteristics of the reinforcement, using the experimentally observed tensile properties of the nanocomposites. 6.6.3.1 MODELING TENSILE MODULUS OF NANOCOMPOSITES
Over several decades, various models have been proposed for predicting the properties of composite materials, based on the properties of the pure components and the morphology of the composite. These models usually assume that each component acts independently of the others. While the general objective of such models is to predict performance of the composite, they provide a tool for evaluation and optimization of the contributions of component properties, such as matrix and tiller modulus, volume fraction, tiller aspect ratio, tiller orientation etc. Since the aspect ratio of partic1es in a nanaocomposite depends on the extent of exfoliation, the estimation or measurement of aspect ratio has been the subject of great interest in the nanocomposite tield (257). Halpin and Tsai (262) developed a composite theory for predicting the stiffness of unidirectional composites as a function of aspect ratio. The longitudinal and transverse moduli are expressed in the following general form: (6.7)
6 RESULTS AND DISCUSIONS
186
where 11 = Er -1 ,A is the aspect ratio (ratio of width or length to thickness), Ec is the Er+ 2A
modulus of the composite, Em is the modulus of the matrix,
~is
volume fraction of tiller,
and Er is the ratio of platelet to matrix modulus. It should be noted that as A
~
0, the Halpin Tsai theory converges to the inverse
rule of mixtures (lower bound). (6.8) where Et is the tiller modulus. Conversely, when A
~
00
the theory reduces to rule of
mixtures (upper bound). (6.9)
Dispersed clay platelets resemble disks (257). Thus, the Halpin Tsai model can be applied to the polymer/clay nanocomposites. In particular, the model may be used to estimate the effective aspect ratio of the clay particles by employing experimental data regarding the modulus of the nanocomposite. Brune and Bicerano (146) used the Halpin-Tsai equation to predict the modulus of exfoliated and partially exfoliated polymer/clay nanocomposites. If exfoliation is incomplete, they consider the system as a composite, which consists of matrix and pseudoparticles, which are incompletely exfoliated stacks of individu al platelets. The Halpin-Tsai equation can then be applied to this model system. Ecomposite _ 1
+ 2A'1]'~'
---',"""",- ,and 1]
Ematrix
1-1] ~
, E ' -1
=
,r
Er
+ 2A
,
(6.10)
where A' is now the aspect ratio of the platelet stack, fjl is the volume fraction of platelet stack to matrix, and E' r is the ratio of the modulus of the platelet stack to that of matrix. The following equations result from the geometry of the stacks in which, N is the number of platelets in the stack, and slt is the ratio of the spacing between platelets in a stack to the thickness of a platelet. The results from this equation show that the ratio of
/
6 RESULTS AND DISCUSIONS
187
composite modulus to the matrix modulus decreases strongly with increase in N but weald y depends on slt.
J
A , =A- [ - - -1- - A S N 1+(1-1/ N)"( A
fi = {b(1+ (1-1/ N)7 ) (1-11 N)~
,1
Er = Er
[
(6.11)
t
+----=-S 1+(1-1IN)"( 1+(1-11 N)"( S
J N ~ N+(l-N)(f)C~; J
Padawer and Beecher (142) and Riley (143) developed empirical equations for flake-like inclusions, based on a simple rule of mixtures. The rule of mixtures predicts composite modulus as given in equation 6.9. The modified rule of mixtures uses the Modulus Reduction Factor (MRF) to compensate for the large discrepancy in the calculated value of the composite modulus using the simple rule of mixtures and the observed value. This discrepancy arises from extremely large difference between the filler and the matrix moduli. The modified rule of mixtures has the form: (6.12)
The modulus reduction factor given by Padawer and Beecher (142) for flakes is: (MRF) =
(1-
ta:hU)
u_l~ AV~
(6.13)
where Gm is the shear modulus of the matrix. A different modulus reduction factor, proposed by Riley (143) for flakes has the form: (MRF)
= 1- ln(u + 1) u
(6.14)
188
6 RESULTS AND DISCUSIONS
The aspect ratios of the clay fillers in nanocomposites can be predicted using the modified rule of mixtures and the modulus reduction factor. For a perfect interface, Hui and Shia (149, 150) derived equation (2.18) for estimating the Young's modulus of composites containing aligned platelet inclusions:
Ec _
1
~ ~ + ç! A ]
E. - 1- [ q=f/J+
Em Er-Em
(2.18)
+3(l_f/J)[(l_g)~2_(g/2)] A -1
and
7t
g=-A 2
The aspect ratio A in the above formulae is defined as the ratio of the thickness to the width. The experimental data on tensile moduli of the nanocomposites were fit to the models namely, Halpin-Tsai, Modified Rule of Mixtures (MROM) and Hui-Shia. The tensile modulus of the clay plate1ets was taken as 170000 MPa (149,257). Figure 6.6.30, Figure 6.6.31 and Figure 6.6.32, for polyamide-6 nanocomposites, show that the relative moduli of the nanocomposites produced using System A are always higher than those produced using System C. The fitting of data to the mode1 equations is fairly good for PA-6 nanocomposites with R2 values ranging from 0.87 to 0.94, except for the PA6/Cloisite Na+ composites produced with the System C, in which case R 2 is 0.5. The fitting lines shown in these figures are the same for aU the composite models used, but the fitting parameter, the aspect ratio was different for different models. The aspect ratio values obtained with the different mode1s are given in Table 6.6.7. The fitting of the data of PS nanocomposites to the model equations was very poor, since the tensile modulus decreases above 2%wt of clay content. The R 2 values in this case were around 0.05. Figure 6.6.33 gives the fitting results for PS/DylarkiCloisite 10A nanocomposite and Table 6.6.8 gives the values of the aspect ratios obtained.
189
6 RESULTS AND DISCUSIONS
1.6~----------------------------------------~
1.5
• Expt. Data PA-6/30B (System A) - - Estimated from model (System A) o Expt. Data PA-6/30B (System C) - - - Estimated from model (System C)
•
1.4
E
o
1.3
w ........., W
1.2
1.1
1.0
0.9 0.000
0.005
0.010
0.015
0.020
0.025
Vol. Fraction
Figure 6.6.30: Experimental data on tensile modulus and model predictions PA-6/Cloisite 308 nanocomposites (Ff: System A=0.9378, System C=0.9228).
1.4
1.3
1.====================::::::;---------, • Expt. Data PA-6/15A (System A) - - Estimated from Model (System A) o Expt. Data PA-6/15A (System C) - - - Estimated from Model (System C)
•
• 1.1
1.0
0.000
0.002
0.004
0.006
0.008
0.Q1 0
0.012
0.014
0.016
0.Q18
Vol. Fraction
Figure 6.6.31: Experimental data on tensile modulus and model predictions PA-6/Cloisite 15A nanocomposites (Ff: System A=0.9168, System C=0.8776).
190
6 RESULTS AND DISCUSIONS
1.25,-;:::===========::::;-----------,
1.20
• Expt. Data PA-6/Na+ (System A) - - Estimated trom Model (System A) 0 Expt. Data PA-6/Na+ (System C) Estimated trom Model (System C)
1.15
•
o
•
1.05
_----
o
1.00 0.000
0.005
o
----
---- --- ---------
0
0.010
0
0 0.015
0.020
0.025
0.030
Vol. Fraction cp
Figure 6.6.32: Experimental data on tensile modulus and model predictions PA-6/Cloisite Na+ nanocomposites (Ff: System A=0.9089, System C=0.4871). 1.10
1.08
o
• o
1.06 E
LU
"'"'-0
1.04
LU
1.02 • Expt. Data PS/Dylark/10A (System A) - - Estimated trom Model (System A) o Expt. Data PS/Dylark/10A (System C) - - - Estimated trom Model (System C)
1.00
0.98 0.000
0.002
0.004
0.006
0.008
0.010
0.012
Vol. Fraction cp
Figure 6.6.33: Experimental data on tensile modulus and model predictions PS1301/Cloisite 10A/Dyalrk (2%wt) nanocomposites (Ff "" 0.2)
191
6 RESULTS AND DISCUSIONS
Table 6.6.7: Aspect ratios of the clay fillers for polyamide-6 nanocomposites calculated using different composite models
P A-6/Cloisite 30B
PA-6/Cloisite Na+
PA-6/Cloisite 15A
Model System A
System C
System A
System C
System A
System C
Halpin-Tsai
16.7
10.4
16.1
6.3
3.1
0.4
MROM
130
80
125
49
27
6.7
Hui-Shia
47
29.2
45
17.7
8.9
1.4
TEM Results
74
29-47
Table 6.6.8: Aspect ratios of the clay fillers for polystyrene nanocomposites calculated using different composite models
PS1301lCloisite IOA Model
PS1301lCloisite IOA! Dl:lark {2%wt}
System A
System C
System A
System C
0.67
3.1
1.0
1.5
MROM
11
27
13.6
17
Hui-Shia
2.5
9
1.3
1.3
Halpin-Tsai
Comparison of the above results with the aspect ratios obtained from TEM analysis shows that the Halpin-Tsai equation underpredicts the aspect ratios, whereas MROM overpredicts them. The Hui-Shia equations give aspect ratio values between Halpin-Tsai and MROM predictions, and close to those obtained by TEM analysis. Table 6.6.7 shows that the aspect ratios are always higher for System A specimens. The ratio of the aspect ratios of the System A and the System C specimens increases as the compatibility between the components (polymer and clay) becomes lower. For example this ratio in PA-6/Cloisite 30B is around 1.6, for PA-6/Cloisite 15A it is 2.5 and for PA-
6 RESULTS AND DISCUSIONS
192
6/C10isite Na+ it is 4-7. This means that System A is more effective in producing nanocomposites, when the compatibility between the components is poor. The aspect ratios are low, in case of PS/Cloisite nanocomposites (Table 6.6.8). Although the values for System C specimens appear to be higher than the values for 'System A specimens, it cannot be concluded that System A was less effective, since the overall fitting of the data to the mode1s was very po or. The clay is not exfoliated in this case. Probably the higher residence time in System A leads to higher degradation of the polystyrene matrix. This could lead to deterioration of the mechanical properties. 6.6.3.2 MODELING TENSILE STRENGTH OF NANOCOMPOSITES It is generally accepted that the strength of micro-particulate filled polymer composite
materials is strongly influenced by the interfacial adhesion between the filler surface and the polymer matrix (263). Unlike fiber-reinforced composites, the interfacial interaction in particulate-filled composites is hard to measure directly, for example, by the fiber pull-out technique. The tensile modulus of the composites is more sensitive to the aspect ratio of the filled particles, whereas the tensile strength is more sensitive to the interfacial adhesion. The work of adhesion Wa can be quantified to determine interfacial bond strength between the silica surface and the pol ymer matrix as follows (264): (6.15) where W/ represents dispersion forces and Wah the forces due to hydrogen bonds. In the case of polymers like polyethylene and polystyrene, which do not have the capacity to form hydrogen bonds, only dispersion forces are responsible for interfacial adhesion. The work of adhesion is equal to the dispersion energy between the two neighboring particles given by equation 2.8:
where Au is the Hamaker constant of the material and d is the separation distance between the two entities. The value of d, the interatomic eut-off distance, is usually taken as 0.165-0.185 nm (191, 228, 229). In order to estimate the dispersion forces
6 RESULTS AND DISCUSIONS
193
between the clay platelets and polymer matrix, the effective Hamaker constant between the two entities needs to be estimated. Figure 6.6.34 shows a schematic representation of the adhesion between the clay platelets and the pol ymer matrix. The effective Hamaker constant, A 12, between the untreated clay platelet and the polymer matrix is given by: (6.16) where
All
and
A22
are the Hamaker constants of the pristine clay and the polymer,
respectively. The effective Hamaker constant, Ai32, between the organically modified clay platelet and the pol ymer, is given by: Am =
(~A33 -KX~A33 -~A22)
(6.17)
where A33 is the Hamaker constant of the organic modifier. The values of the Hamaker constants for different materials are given in Table 6.2.2. These values make it possible to estimate the dispersion component of the work of adhesion.
Organic Modifier (3) Clay (1)
Clay (1)
(a)
(b)
Figure 6.6.34: Schematic representation of adhesion between a clay platelet and the polymer matrix: (a) adhesion between untreated clay, (b) adhesion between organically modified clay and the polymer matrix
6 RESULTS AND DISCUSIONS
194
The improvement in the tensile properties of PA-6 nanocomposites is attributed to the strong interaction between the matrix and the silicate layers via formation of hydrogen bonds (265), as shown in Figure 6.6.35. 16.7 If
.
,0\ p\ 10' ;0\ ;0\ si1 Si1 S'!1 \s·l s", s/1 ,/ "l ,,~$ "~i "1 ' i " O' ""
'".
Figure 6.6.35: Schematic illustration of formation of hydrogen bonds in PA-6/MMT nanocomposites (265)
The bond energy ofN .... H hydrogen bonds is 12-20 kJ/mole (266). The specific surface area of exfoliated MMT clay is 700-760 m 2/gm (225,246), and the cation ex change capacity is 0.92-0.95 mequiv/gm (225). The specific surface area of pristine clay is small, since the clay platelets are not exfoliated. Consider the case of PA-6 nanocomposites made from organically modified clay. The organic modifie'r in Cloisite 15A has no functional groups which will form , hydrogen bonds, and the work of adhesion cornes only from the dispersion forces. On the other hand, Cloisite 30B has two hydroxyl groups per molecule of organic modifier, which are capable of forming hydrogen bonds with the PA-6 matrix. If the specific surface area of the exfoliated clay is 760 m2/gm, the maximum amount of organic modifier associated with it would be (0.95/760) mequivalentlm2• Since there are two hydroxyl groups per molecule of modifier, the PA-6 matrix associated with this clay via hydrogen bonds wou1d be (2 x O.95/760) mequivalentlm2 • Then the maximum total bond energy of hydrogen bonds in the system would be (2 x O.95 x20/760 x 1000) kJ/m 5
2
2
works out to be 5x10- kJ/m or 5xlO- J/m
2
2
•
This
•
The work of adhesion from hydrogen bonding can be estimated for pristine clay by a different method. The chain length ofPA-6 between two hydrogen-bonding sites is
6 RESULTS AND DISCUSIONS
195
1.67 nm (265). In a unit square meter there will be 3.58xlQ17 hydrogen bonds. Since the bond energy is 12-20 kJ/mole, which is for the 6.022x1Q23 (Avogadro's number) bonds, the bond energy per square meter would be 1.19x 10-5 kJ/m 2 or 1.19x 10.2 J/m2• Table 6.6.9 summarizes the calculated values of effective Hamaker constants, work of adhesion from dispersion forces, work of adhesion from hydrogen bonding and the total work of adhesion between the clay platelets and the polymer matrix for different clay polymer systems. It shows that the major contribution to the total work of adhesion in the P A-6lCloisite 30B cornes from hydrogen bonding, which is absent in the case of PA-6lCloisite 15A. The total work of adhesion is highest for the P A-6lCloisite Na+, which is untreated clay. In practice, untreated clay do es not exfoliate in polymer matrices and the specifie surface area is low. Therefore, the total adhesion force between clay particles and the polymer matrix is low. This shows that if the untreated clay can be exfoliated by sorne means, it would result in nanocomposites with the highest mechanical properties. Since polystyrene is a non-polar polymer it has a low Hamaker constant value (Table 6.6.2), and in tum the resultant effective Hamaker constant is low. It do es not have hydrogen bonding functional groups; therefore the total work of adhesion in this case is the lowest among aH the nanocomposite systems reported in Table 6.6.8. Table 6.6.9: Work of adhesion for different clay/polymer systems
Clay/Polymer System
Effective Hamaker Const. A 132 J
J/m2
P A-6lCloisite 30B
0.3484x 10-20
0.279xl0-2
PA-6lCloisite 15A
0.4364x 10-20
0.349xl0-2
PA-6lCloisite Na+
9.675xl0-2O
7.75xlQ-2
PS/Cloisite 10A
0.1214xlQ-20
0.0972xlQ-2
W/
Total Wa J/m 2 5.0xlQ-2 0.349xl0-2 1.19x 10-2
8.94xlQ-2 0.0972xl0-2
6 RESULTS AND DISCUSIONS
196
Pukanszky et al. (259,261,267) estimated the composite yield stresses, using the following relation: ) a c =am 1-~ exp ( B~
(6.18)
1+2.5~
where
O"c
and
O"m
are composite and matrix yield stresses,
~ is
the volume fraction of the
fiIler, and the parameter B can be evaluated from experimental data. The equation consists of two parts. The
(l-~)/(1 +2.5~)
component takes into consideration the
decreas.e in the effective load bearing cross section (268), and the exponent describes aIl other effects, which result in an increase of yield stress. The parameter B characterizes the interfacial interaction, inc1uding the interlayer thickness, the interfacial strength, and the specific surface area of the filler partic1es. They proposed the following relationship between parameter B and the above factors: 0".
B=(1+lpfAf)1n al
(6.19)
m
where 1 is the thickness of the interphase, which is proportional to the interfacial adhesion /12, according to the relation l=k/12. The quantities Ph Ah and Di represent the density of the fiIler, the specific surface area of the fiIler, and the yield stress of the interphase, respectively. It is difficult to quantify the length 1 and the yield stress Di of the interphase. In spite of this, the parameter B is useful in quantifying the effects of these parameters on the yield stress of the composite. Rong et al. (263) used this model to analyze the interfacial interactions in polypropylene nanocomposites. He showed that higher values of the parameter B are obtained, when the interfacial adhesion is high. Sumita et al. (269) used this model to describe dynamic mechanical properties of polypropylene composites filled with ultrathin partic1es. In the present study, the data on yield stress of PA-6 and PS nanocomposites were fit to equation 6.18, and the parameter B was evaluated in aIl cases. Figure 6.6.36Figure 6.6.40 show the results of fitting the model equation to yield stress data. This fitting was good in most cases, with R2 values ranging from 0.75 to 0.97. Table 6.6.10 gives the values for fitting parameter B for different nanocomposite systems.
197 .
6 RESULTS AND DISCUSIONS
1.20 • Expt. Data PA"6/30B (System A) - - Estimated trom Model (System A) o Expt. Data Pa-6/30B (System C) - - - Estimated trom Model (System C)
1.18 1.16
•
1.14 1.12 E
~ \:)
1.10 1.08 1.06 1.04 1.02 1.00 0.98 0.005
0.000
0.010
0.015
Vol. Fraction
0.020
0.025
Figure 6.6.36: Experimental data on tensile strength and model predictions PA-6/Cloisite 30B nanocomposites (Ft: System A=0.9670, System C=0_9541). 1.12
r.:======================:;-----------, •
1.10
Expt. Data PA-6/15A (System A) Estimated trom Model (System A) o Expt. Data PA-6/15A (System C) - - - Estimated trom Model (System C) -
•
1.08
E
1.06
•
~ \:)
•
1.04
1.02
o
o
0
--- ---- --- ---
1.00
-----------Q.-_. o
o
0.98
0.000
0.002
0.004
0.006
0.008
0.010
Vol. Fraction
0.012
0.014
0.016
0.Q18
Figure 6.6.37: Experimental data on tensile strength and model predictions PA-6/Cloisite 15A rianocomposites (Ft: System A=0.7645, System C=0_5190)
6 RESULTS AND DISCUSIONS
198
1.14..,.------------------------. •
Expt Data PA-6/Na+ (System A) Estimated from Model (System A) o Expt. Data PA-6/Na+ (System C) - - - Estimated trom Model (System C)
1.12
-
1.10 1.08 ~E
"'u. 1.06 ~
o
-----
1.04 1.02
-- --
1.00 0.98
------
o
_---- 0
o
o
o
-'------r----,---..,.-----....,.------r----,---..,.-----I 0.000
0.005
0.010
0.015
Vol. Fraction
0.020
0.025
0.030
0.035
Figure 6.6.38: Experimental data on tensile strength and model predictions PA-6/Cloisite 308 nanocomposites (Ff: System A=0.9781 , System C=0.7456) 1.05
1.00
0.95
E
~ ~
0.90
0.85
0.80 •
0.75
Expt. Data PS/1 OA (System A) Estimated from Model (System A) o Expt. Data PS/10A (System C) - - - Estimated trom Model (System C) -
•
0.70 ~===;:===r===:::;===;::=:::::....---r----,---..,.----I 0.000 0.002 0.004 0.006 0.008 0.010 0.012 0.014
Vol. Fraction
Figure 6.6.39: Experimental data on tensile strength and model predictions PS1301/Cloisite 10A nanocomposites (Ff: System A=0.9444, System C=0.9565)
6 RESULTS AND DISCUSIONS
1.10
199
-r----------------------,
•
1.05
o
•
1.00
o
•
0.95
•
Expt. Data PS/10A/Dyl (System A) Estimated from Model (System A) o Expt. Data PS/10A/Dyl (System C) - - - Estimated from Model (System C)
0.90
-
0.85
o
..l.b===r===;=====;r====r=""'---..,------,-----r----/ 0.000
0.002
0.004
0.006
0.008
0.010
0.012
0.014
Vol. Fraction cj>
Figure 6.6.40: Experimental data on tensile strength and model predictions
PS1301/Cloisite 10A/Dylark (2%wt) nanocomposites (Ff: System A=O.5143, System C=O.7386) Table 6.6.10: Values of parameter B for different nanocomposites
PA-6/30B System Parameter B
PA-6/15A
PA-6INa+
PS/10A
PSIlONDyl
A
C
A
C
A
C
A
C
A
C
11.43
10.39
9.45
4.44
7.24
4.87
-15.78
-10.48
-2.79
-4.4
The highest values of parameter B result for the most compatible PA-6/Cloisite 30B nanocomposites system. PA-6/Cloisite 30B nanocomposites also show the highest values of the aspect ratios for the filler particles. Higher aspect ratios of filler particles and' better adhesion between filler and matrix, in this system, results into higher values for the parameter B. Similarly, System A nanocomposites, which contain fillers with higher aspect ratios, result into the higher values of parameter B. The work of adhesion is high for untreated clay, but the aspect ratios are low. This is reflected in parameter B, which is moderately high.
PS/Cloisite 10A nanocomposites have both low work of
6 RESULTS AND DISCUSIONS
200
adhesion and aspect ratios, and the parameter B is negative. The addition of Dylark compatibilizer to this system .enhances the adhesion between the components. This results in somewhat higher but still negative values for the parameter B.
6.7 BARRIER PROPERTIES OF NANOCOMPOSITES The reduced gas and liquid permeability of nanofiUed polymers makes them attractive membrane materials. The large reduction in permeability to gases through composites may be explained by Nielsen's (152) tortuosity model, described in section 2.6.2 and equation 2.20. Cussler's (153) model (equation 2.21) introduces a geometric factor Ji and the aspect ratio to predict the permeability of composite materials. In the present study, thin films of different nanocomposites were prepared by repeated compression mol ding below the melting temperature of the polymer matrix. The oxygen permeability of these thin films were determined and fit to the models to estimate the aspect ratio of the fiUer particles. The oxygen permeability coefficient of polystyrene is very high (- 100 times) compared to that of PA-6.
Both PA-6 and PS nanocomposites show very smaU
difterence between the oxygen permeability values obtained using different clays Figure 6.7.1 shows that the permeability of PA-6 matrix is reduced from 0.9 to 0.2 (cc.mmlm2 .day.atm) by the incorporation of 5%wt of Cloisite 30B organoclay. The permeability values are slightly lower for System A nanocomposites. In the case of PS/Cloisite 10A nanocomposites (Figure 6.7.2), the permeability is reduced from 106 to 75 (cc.mmim2 .day.atm) by the addition of3.5%wt ofthe clay. The permeability is lower for System A nanocomposites.
The incorporation of Dylark compatibilizer in the
polystyrene matrix increased the oxygen permeability of the polymer. Figure 6.7.3 shows that the permeability values of PS nanocomposites incorporating Dylark compatibilizer are higher than those without compatibilizer at aU clay concentrations. They did not cause a reduction in permeability up to a clay content of2.6%wt. The permeability data were fit to the Nielsen model, and the aspect ratios were estimated. There are no specific guidelines available for the estimation of the geometric
6 RESULTS AND DISCUSIONS
201
parameter f.1c of the platelets in Cussler's Model. In the present study, it was assumed that f.1c =1. 1.2~------------------------------------------~
1.0
____- System A
--e-
0.2
Systeme Estimated trom Model System A - - - Estimated trom Model System e
0.0~====T============r====~------~----------~
o
2
4
6
Clay (%wt)
Figure 6.7.1: Oxygen permeability coefficient of PA-6/Cloisite 30B nanocomposites 140 . - - - - - - - - - - - - - - - - - - ------, .......
c:
Q)
'u ;;: Q)
120
E
o .ctl. . . :>.
(.)
~ ctl
._
"0
.0 . ctlC\l
E .... E ~ E c u (.)
---- ----
100
Q)
E -.
80
Q)
~
o
-e-
System A System e Estimated trom Model System A - - - Estimated trom Model System C
--0
60
o
1
2
3
4
Clay (%wt)
Figure 6.7.2: Oxygen permeability coefficient of PS1301/Cloisite 10A nanocomposites
6 RESULTS AND DISCUSIONS
202
160.----------------------------------------,
-
l
140
c: CI>
'5 ;;:::
~ E Co) 1ti•
~~
----------------- ----- -.....
120
:="0 ~",'
.....
CI>..§
E E (fi E 100 o.. c: Co) CI> Co)
~
0
80
"-
"-
\1
-.-- PS/10A --e-- PS/10A/Dylark Estimated trom Model PS/10A - - - Estimated trom Model PS/10AlDylark
60 0
2
3
4
Clay (%wt)
Figure 6.7.3: Comparison of the oxygen permeability coefficients of PS/Cloisite 10A
and PS/Dylark (2%)/Cloisite 10A nanocomposites produced using System A
Table 6.7.1: Aspect ratios of different nanocomposites predicted from oxygen
permeability data
P A-6/Cloisite 30B
PS1301/Cloisite lOA
PS/Dyl/l0A
System A
System C
System A
System C
System A
Nielsen Model
128.0
97.8
48.2
8.0
26.3
Cussler Model
66.3
57.9
47.8
21.4
36.4
Aspect ratios from composite models for tensile modulus Halpi-Tsai
16.7
10.4
0.67
3.1
1.0
MROM
130
80
11
27
13.6
Hui-Shia
47
29.2
2.5
9.0
1.3
TEM
74
29-47
6 RESULTS AND DISCUSIONS
203
The results in Table 6.7.1 show that the aspect ratios of PA-6/Cloisite 30B nanocomposites predicted from oxygen permeability values using the Nielsen model are close to those predicted using MROM from tensile modulus values (section 6.6.3.1), and those of the Cussler model for permeability are close to the Hui-Shia Model for tensile modulus. The aspect ratios of the clay particles in the System A samples are larger than those of the System C. The R2 values for fitting the data to the model predictions were 0.66-0.74 for the Neilsen model and 0.85-0.95 for the Cussler model. The aspect ratios estimated for polystyrene nanocomposites using oxygen permeability data are higher than those predicted using tensile modulus data. The incorporation of clay lowers the permeability ofpolystyrene, but it does not change the tensile modulus significantly.
6.8 RHEOLOGY OF NANOCOMPOSITES Characterization of the rheological behavior and properties of pol ymer and composite melts is important for the design of pol ymer conversion equipment, the estimation of power requirements, the optimization of processing conditions, and understanding of structure development
~uring
processing. AIso, since polymer melt flow behavior is
strongly influenced by the composition, concentration, and morphology of the filler, the study of rheological behavior of pol ymer composite melts is important for the development of formulations with adequate industrial processibility (270). Capillary rheometry was used to estimate the viscosity of the nanocomposite melts at different shear rates. The presence of particles in the melt increases the viscosity and influences the flow field (204). This is illustrated in Figure 6.8.1. In case (a) the velo city profile near wall is shown for pure polymer melt. The variation in velocity in the various layers is indicated by the arrows of different lengths. Case (b) for composite flow at low shear rate, exfoliated platelets, which are randomly oriented, cut across several layers in the flowing polymer melt. This slows down the fluid, and the overall velo city distribution becomes less sharp, reflecting a lower average velo city under the same driving pressure gradient. Thus, an increase in viscosity is observed. In case (c), under high shear, the platelets become oriented in the flow direction and align
6 RESULTS AND DISCUSIONS
204
themselves parallel to the layers. The adhesion between two dissimilar entities such as polymer chains and the clay platelets is lower than the adhesion between the layers of the polymer matrix. This reduces the friction between the adjacent layers and results in lower viscosity.
(a)
(b)
Figure 6.8.1: Sehematie representation of flow pattern of polymer and polymerie nanoeomposite me!t: (a) flow of pure polymer melt near stationary wall, (b) flow of polymer melt eontaining exfoliated platelets randomly oriented under low shear, (e) flow of polymer melt eontaining exfoliated platelets oriented under high shear.
Gu et al. (271) reported that the viscosity of polypropylene nanocomposite was higher than that of the unfilled pol ymer under low shear, but displayed higher shear thinning behavior. The shear thinning began at lower shear rate. Similar observations were reported by other researchers (159-166).
6.8.1 RHEOLOGY OF PA-6 NANOCOMPOSITES Figure 6.8.2 shows the rheological behavior of P A-6 nanocoposites with different clays. the viscosity of nanocomposites with Cloisite 30B and Cloisite 15A organoclays at low shear rates is significantly higher than that of pure pol ymer. At high shear rates, shear
6 RESULTS AND DISCUSIONS
205
thinning behavior is observed. The viscosity behavior of the composite with untreated clay (Cloisite Na+) is similar to that of the unfilled polymer. It do es not show higher viscosity at low shear rates and do es not exhibit a more pronounced shear thinning behavior. 3000 ~----------------------------------------~ .
li 2000
li!
• v•
v
û) rel
ft:.. :i!'
"iii 0
(.)
III
:>
1000 900 800 700
9 ~
~
• • 9
600
·10
500
•
400 • 300
o 6.
V
PA-6 PA-6/Cloisite 308 (5.0%wt) System A PA-6/Cloisite 15A (4.1%wt) System A PA-6/Cloisite Na· (6.9%wt) System A
200 100
10
1000
Shear rate (S·1)
Figure 6.8.2: Steady shear viscosities of PA-6 nanocomposites at 240°C: Effect of using clays with different surface modifiers on the viscosity
Figure 6.8.3 shows the effect of clay content on the rheological behavior of the PA-6/Cloisite 30B nanocomposites made with the System C. Higher clay content in nanocomposite resulted in higher low shear viscosity, as well as more pronounced shear thinning. Similar behavior was observed in PA-6/Cloisite 15A nanocomposites and for both processing Systems A and C (Figure 6.8.4). The effect of processing with System A and with System C on the viscosity of the PA-6/Cloisite 15A nanocomposites is shown in Figure 6.8.4. Overall, the viscosities of System A nanocomposites were lower than those of System C nanocomposites. This difference was more pronounced at higher clay contents. Similar behavior was observed in Cloisite 30B nanocomposites. This could be the result of higher degradation of the pol ymer chains in System A processing, since it involves higher residence time.
6 RESULTS AND DISCUSIONS
206
4000 3000
0 0
2000
A A
êii ctl
e:. ~
ëil 0 u
en
:>
0
A
0
0
:0
1000
~
.Ow
", .. .l.,.
.~
:0
800 700 600
;0
~ (]
500
~
400
0
300
A D
0
PA-6 PA-6/Cloisite 30B (3.1%wt) System C PA-6/Cloisite 30B (5.0%wt) System C
(]
200 1
10
100
1000
5hear rate (5. 1)
Figure 6.8.3: Steady shear viscosities of PA-6/Cloisite 30B nanocomposites at 240°C: Effect of clay content on the viscosity 4000 3000
A A
2000
~
A A
A
i
• • •~
êii ctl
e:.
-
1000 >- 900 '1n 800 0 U en 700 600
A
i
•..
fA
:>
lA
500
•
400
A 0 A
300
•
200 1
PA-6/Cloisite PA-6/Cloisite PA-6/Cloisite PA-6/Cloisite PA-6
15A (2.5%wt) 15A (4.1%wt) 15A (2.5%wt) 15A (4.1%wt)
10
System System System System
A A C C
100
5hear rate (5.
1
•
~
•\ 1000
)
Figure 6.8.4: Steady shear viscosities of PA-6/Cloisite 15A nanocomposites at 240°C: Effect of processing with System A and System C on the viscosity
6 RESULTS AND DISCUSIONS
207
6.8.2 RHEOLOGY OF POLYSTYRENE NANOCOMPOSITES
Unlike PA-6, polystyrene matrix do es not yield exfoliated nanocomposite structure. This is reflected in the rheological behavior of PS composites. Figure 6.8.5 shows that the shear thinning behavior occurs in polystyrene as well as its nanocomposites. The viscosity at low shear rates of Cloisite 30B and Cloisite 15A composites is higher than the pure pol ymer, and it is lower in case of Cloisite 10A composites. X-ray diffraction results showed collapsed peaks for of PS/Cloisite 10A nanocomposites. The bulky styrene group in the modifier molecule tends to escape from the clay galleries resulting in the collapse of layered clay platelets. The modifier that escapes from the clay galleries probably acts as a plasticizer and reduces the viscosity of the polymer. Figure 6.8.6, shows that the viscosity of PS/Cloisite 10A nanocomposite at low shear rates decreases with increasing clay content. Figure 6.8.7 shows that the addition of 2%wt Dylark compatibilizer has no effect on the viscosity of PS nanocomposites and that the viscosity at low shear rates decreased upon the addition of Cloisite 10A clay. 6000 5000 4000
Cl
v ~
3000
--
~ ~
:'j
Cl
l
2000
Cl
enCIl a..
i 1000
i
>-
cn 0
u CI)
:>
700 600 500 400
~
fi
300 200
• ~
Cl
v
PS1301 PS/Cloisite 10A (2.6%wt) System A PS/Cloisite 308 (3.0%wt) System A PS/Cloisite 15A (2.5%wt) System A
~
100 10
100
1000
Shear rate (S-1)
Figure 6.8.5: Steady shear viscosities of PS1301 nanocomposites at 210°C: Effect of using crays with different modifiers on the viscosity.
6 RESULTS AND DISCUSIONS
6000 5000 4000
R
3000
0
~ 0
2000
è
0
Ci)
~
CIl
-a..
>·Cii 0
t.l
(J)
:;
208
1000
~
700 600 500 400
C 0
e
300 200
0
t. 0
PS1301 PS/Cloisite 10A (2.6%wt) System A PS/Cloisite 10A (3.5%wt) System A
It
100 10
100
1000
Shearrate (S·1)
Figure 6.8.6: Steady shear viscosity of PS1301/Cloisite 10A nanocomposites at 210°C: Effect of clay concentration on the viscosity.
5000 4000 3000 2000
Ci)
e:..
CIl
1000
~
700 600 500 400
·Cii 0
t.l
(J)
:;
300 •
200
o ... v
PS1301 PS/Dylark(2%) PS/Cloisite 10A (3.5%) PS/Dylark(2%)/Cloisite10A (3.5%)
100 100
10
1
1000
Shear rate (S·1)
Figure 6.8.7:
Steady
shear
viscosity
of
PS1301/Dylark
(2%wt)/Cloisite
nanocomposites at 210°C: Effect of clay concentration on the viscosity.
10A
6 RESULTS AND DISCUSIONS
209
6.8.3 ESTIMATION OF SHEAR STRESS IN THE STATIC MIXER
The steady shear viscosity data on nanocomposites were used to estimate the shear stresses in the static mixer. The pressure drop in the static mixer was recorded in each of the experiments. The pressure transducer was placed at the exit of the extruder. The pressure indication in System C refers to the pressure drop in the die, while the pressure indication in System A refers to the pressure drop in static mixer and the die. The difference between these two pressures gives the pressure drop in the static mixer. Table 6.8.1 gives the values of pressure drop in the ISO static mixer for different screw speeds of the extruder at 240°C for PA-6/Cloisite 30B nanocomposites. The viscosities for the respective nanocomposites are calculated as described in Sections 4.3.1.1 (equation 4.43) and 5.1.2 (equation 5.3). The feed rate of PA-6 in these experiments was 1.7 Kg/hr. Depending upon the variations in processing conditions, the pressure varied in the range of ± 5%. The viscosities were calculated using differeÎ1t power factors (Table 4.3.1) reported in the static mixer literature (208) and the empirical relation given by the manufacturer. The pressure drop in the static mixer was not influenced by the extruder speed. The corresponding shear rate in the static mixer can now be esCmated using the rheological data obtained by capillary rheometry at the corresponding temperature. Using the viscosity and the shear rate, shear stress in the static mixer can be calculated. The estimated shear stress
In
the ISO static mIxer for PA-6/Cloisite 30B
nanocomposite is shown in Figure 6.8.8, for different clay contents. The available shear stress in ISO static mixer is
~ighest
for the pure pol ymer. With the addition of clay, the
viscosity of the nanocomposite melt increases but the available shear stress decreases. Figure 6.8.8 compares the available shear stress to the minimum shear stress required for the exfoliation of organoclay and untreated clay. The minimum required shear stress is
calculated based on the peeling of single platelets from a clay particle of width and length 1000 nm, using the model described in Section 4.1.2.1.
6 RESULTS AND DISCUSIONS
210
Table 6.8.1: Pressure drop in 1SG statie mixer for PA-6/Cloisite 30B nanoeomposite Extruder speed
Clay
System System A C
RPM
(%wt)
Pressure Pressure (psi) , (psi) (psi)
50
100
150
200
250
Viscosity (Pa.S)
tlP 5.2
Kp 9600
Kp 7200
Kp 8460
Avg.
Eqn.
0.0
560
480
80
499.6
705.3
529.0
621.5
588.8
3.1
740
570
170
1061.7
1498.7
1124.0
1320.7
1251.3
5.0
1100
590
510
3185.2
4496.2
3372.1
3962.2
3753.9
0.0
550
460
90
562.1
793.4
595.1
699.2
'662.5
3.1
740
530
210
1311.5
1851.4
1388.5
1631.5
1545.7
5.0
1100
560
540
3372.5
4760.6
3570.5
4195.3
3974.7
0.0
510
450
60
374.7
529.0
396.7
466.1
441.6
3.1
740
520
220
1374.0
1939.5
1454.6
1709.2
1619.3
5.0
1050
540
510
3185.2
4496.2
3372.1
3962.2
3753.9
0.0
490
430
60
374.7
529.0
396.7
466.1
441.6
3.1
760
520
240
1498.9
2115.8
1586.9
1864.6
1766.5
5.0
1070
600
470
2935.3
4143.5
3107.6
3651.4
3459.5
0.0
500
420
80
499.6
705.3
529.0
621.5
588.8
3.1
700
520
180
1124.2
1586.9
1190.2
1398.4
1324.9
5.0
1020
580
440
2748.0
3879.0
2909.3
3418.4
3238.7
-
It is interesting to note in Figure 6.8.8 that the available shear stress in the ISO
static mixer decreases with the increase in clay content. The exfoliation of organoclay beyond the clay content of 2.5%wt is probably not feasible in the static mixer. The improvement in the properties of nanocomposites prepared using the ISO static mixer above a clay content of 2.5%wt may be attributed to the effective distribution of the clay particles and tactoids, extensional stresses and frequent changes in flow direction during folding action. The minimum shear stress required for the exfoliation of untreated clay is very high and is hardi y attainable in static mixer. The improvements obtained in
211
.6 RESULTS AND DISCUSIONS
mechanical properties, in this case, are also mainly due to the effective distribution of clay particles by static mixer processing.
106~----------------------------------------~
Minimum shear stress required for exfoliation of untreated clay
-
'".ê
~ 105 (J)
!fi.... û5.... as Q)
.s:::.
CI)
104
o
2
4
6
Clay (%wt)
Figure 6.8.8: Relation between shear stress in ISG static mixer and clay content in
PA-6/Cloisite 308 nanocomposite
The pressure drop values across the 10 element ISO static mixer during the processing of the nanocomposites of PA-6 and polystyrene with different clay types are shown in Table 6.8.2, along with the corresponding calculated values of viscosity and the available shear stresses. The shear stress required for the exfoliation of organoclay (Figure 6.8.8) is "" 27x 103 N/m2 (27 KPa). Only in the case of PA-6, the shear stress is sufficient to exfoliate the organoclay. In the case of polystyrene, the pressure drop across the static mixer is high and the corresponding available shear stress is as low as 15.3 KPa. The shear stresses available during the processing of PS nanocomposites with ISO static mixer are lower than the required shear stresses. The addition of 2%wt Dylark compatibilizer has no effect on the pressure drop across the static mixer. It is very unlikely that any of the organoclay types employed in this study could be exfoliated in a polystyrene matrix, using the ISO static mixer.
6 RESULTS AND DISCUSIONS
212
Table 6.8.2: Pressure drop in ISG static mixer for different nanocomposites at extruder speed of 200 RPM and the available shear stresses
A
System C
~p
Viscosity (PaS)
Shear Stress
(%wt)
Pressure (psi)
Pressure (psi)
(psi)
Kp 8460
103 (N/m2)
0
490
430
60
466.1
1071
PA-6/Cloisite 30B (240°C)
5.0
1070
600
470
3651.4
6.158
PA-6/Cloisite 15 A (240°C)
4.1
1160
600
560
4350.6
4.351
PA-6/Cloisite Na+ (240°C)
6.9
730
600
130
1010
13.95
0
1080
600
480
3170
15.3
PS 130 lICloisite 10A (210°C)
3.5
1380
700
680
4490.5
8.583
PS 13 0 lICloisite 15A (210°C)
3.2
1400
700
700
4622.5
15.98
PS1301lCloisite 30B (210°C)
4.0
1380
750
630
4160.3
21.09
0
1230
700
530
3500
12.66
3.5
1400
700
700
4622.5
4.664
Nanocomposite
PA-6 (240°C)
PS1301 (210°C)
PS1301l2%Dylark (210°C) PS 130 1I2%Dylarki Cloisite 1OA(21 O°C)
Clay
System
x
213
6 RESULTS AND DISCUSIONS
6.9 CRYSTALLIZATION KINETICS OF PA-6 AND ITS NANOCOMPOSITES The crystalline structure of PA-6 is rather complex. PA-6 has sheet-like structure due to the hydrogen bonds that are fonned within specifie crystallographic planes (272). In addition, PA-6 exhibits three crystalline fonns that generally coexist in various amounts, depending on processing conditions. The stable monoc1inic a-fonn has a fully extended planar zigzag chain confonnation, with H-bonds lying between anti-parallel chains. The monoc1inic y-fonn has a chain twist in the amide groups with respect to the methylene segment, and the pleated sheets of parallel chains are joined by hydrogen bonds. The metastable pseudo-hexagonal
~-fonn
inc1udes stacking of parallel and antiparallel
chains, para-crystalline disorder, faults in H-bond sheet-like setting, and H-bonded layers nonnal, instead of parallel, to the chain axis. A mesomorphic structure with random distribution ofthe H-bonds has been proposed for this fonn (273). The
~-fonn
is
similar to the y-fonn from a crystallographic standpoint, and it is c10sely related to the amorphous component from the standpoint of chain confonnation (272). Gogolewski et al. (274, 275) studied the crystallization of PA-6 under elevated pressure and conc1uded that the chain-extended crystals (a- fonn) may not necessaril y grow directly in the melt under pressure, but that the fonnation of imperfect crystals of the folded-chain type (y,
~-fonn)
might be an intennediate step of the chain extension.
The a-fonn melting temperature is about 220°C and the y-fonn melting temperature is about 210°C (276). Kojima et al. (171, 172) crystallized PA-6 and its nanocomposite (PNC) by injection molding under elevated pressure, followed by annealing at 200 to 300°C under elevated pressure. The products were analyzed by DSC, wide-angle X-ray diffraction and FTIR. In the case of samples annealed under relatively low pressures (0.12 GPa), two endothennic peaks, due to the ordinary melting temperatures of PA-6 crystal were observed (PNC, 212°C y-fonn and 223°C a-fonn and PA-6, 215°C y-fonn and 225°C a-fonn). With increasing pressure, the lower temperature endothenn due to
6 RESULTS AND DISCUSIONS
214
the me1ting ofy-form crystals disappeared. It was concluded that pressure accelerates the transformation of the y-form to the a-form. The authors reported another high me1ting phase (me1ting point 240°C), in PNC injection molded at elevated pressures. This phase represented 2-3% of the sample, and it was attributed to the ion bonding of smaller molecular weight PA-6 with the negative charge of montmorillonite. The high melting phase is a characteristic of PNC. The majority of high-pressure crystallization studies have considered the crystallization products formed under specific crystallization conditions, rather than the kinetics of the process. In the following, we focus on the study of the crystallization kinetics of PA-6 alone and in the nanocomposite under a range of pressures and temperatures. PA-6 (Ube 1015B) and a PA-6-nanocomposite (Ube 1015C2) containing 2 wt% montmorillonite organoclay were used for the crystallization kinetics study using high pressure dilatometry. Isobaric PVT experiments were carried out to obtain the melting points and the crystallization temperatures of PA-6 alone and in the nanocomposite at different pressures. The melting point and crystallization temperatures were defined as the temperatures at which the rate of change of volume was highest during heating and cooling, respective1y. The Avrami equation (5.12) was used to fit the isothermallisobaric crystallization data. Figure 6.9.1 and Figure 6.9.2 show the pressure dependence of melting point (Tm) and crystallization temperature (Tc). Figure 6.9.3 shows a plot ofmelting points and
crystallization temperatures against pressure. The results show that PA-6 alone and PA-6 in the nanocomposite' have close Tm values, but PA-6 in the nanocomposite has lower Tc values than PA-6 altme, at all pressures. Lowering of the crystallization temperatures of the nanocomposite could be attributed to the stabilization of the y-form. The y-form, which has lower melting point than the a-form, is expected to have a lower
crystallization temperature. During the melting experiments, the y-form is transformed into the a-form before the melting point is reached and the melting temperatures obtained for PA-6 and its nanocomposite are close.
6 RESULTS AND DISCUSIONS
215
16.-----------------------------------,
14
b -: 12
j ÊC> 10
"'~
g ci
6
CI)
'5 Q)
i
4
0
*
a: 190
210
230
250 Temperature (oC)
270
290
190
210
230
270
250
290
Temperature ('C)
Figure 6.9.1: Rate of change of specifie volume of PA-6 and PA-6 nanocomposite during Isobaric heating ~
"
1
-5
C>
"'~-10 ~
>
ci
~o ·15
f
.!:o
~ -15 0
Q)
~
"'.!:
-20
fil a:
-20
0
fil a:
PA-6
PNC
-25 +----.----.-----...__--..-----.----.-----...__--1 160 170 180 190 200 210 220 230 240
-25 160
170
Temperature ('C)
179 10MPa 180
190
200
210
220
230
240
Temperature ("C)
Figure 6.9.2: Rate of change of specifie volume of PA-6 and PA-6 nanocomposite during Isobaric cooling 260
240
~ e 220 2
~
~
~ 200
1-
,. PA-6TM! '.PA-6TC' ,t.PNCTM, loPNCTC
180
1
160 0
50
100
150
200
Pressure (MPa)
Figure 6.9.3: Melting and crystallization temperatures of PA-6 and nanocomposite
6 RESULTS AND DISCUSIONS
216
During the isobaric heating experiment (Figure 6.9.1), the presence of the characteristic high melting phase in PA-6 nanocomposites was observed. The peaks at 260°C for the 10 MPa isobaric run and at 291°C for the 50 MPa isobaric run are observed only in case of PA-6 nanocomposite. The isobaric/isothermal crystallization kinetics data were plotted as ln [-ln (1-X)] against ln t, where X is the relative crystallinity and t is the time in seconds for the range of pressures (50-200 MPa). The results, shown in Figures 6.9.4 and 6.9.5, indicate that there are two distinct regions of crystallization kinetics. The material behavior in the two different regions was fitted using a linear fit. The values of the correlation coefficient (R
2
),
in most cases, for the linear regression range between 0.94 and 0.99. The
correlation was not very good in sorne cases at high supercooling. The si opes of the fitting lines give the values of the Avrami exponent (n) and the y-intercepts represent the Avrami rate constant (K) in the the Equation 5.12.
1- X = exp(-Kt n ) At low crystallization temperatures, the two regions are not distinct and may be represented by a single line. In these cases (205.7°C at 50 MPa and 225.6°C at 150 MPa for PNC), the rates were assumed tentatively for the a-form. However, they probably represent the combined rates for the a-form and the y-form. At higher crystallization temperatures, these regions are very distinct. It is suggested that the initial region represents the formation of folded-chain crystals (y-form), which are grown from the melt during the early stage of crystallization under pressure (274) and the transformation ofthese crystals into chain extended crystals (a-form), as weIl as the growth of the chain extended crystals (a-form). The later region represents the growth of the thermodynamically stable chain-extended crystals (a-form). It is known that y-form crystals are observed on quenching from the melt; while
the a-form results from slow cooling of the melt (276). This suggests that the formation of the y-form crystals is a prerequisite for the a-form crystal formation. The choice of the crystal form in the even polyamides depends on the chain repeat length (277).
6 RESULTS AND DISCUSIONS
217
Polyamides with the longer repeat unit have greater tendency to assume the y-form. This tendency probably relates to the reduction of stress or steric hindrance caused by the rotation of the amide group around the C-CO and C-NH single bonds to form a hydrogen bond between the adjacent paralle1 chains instead of the adjacent antiparalle1 chains. Because every chain has statistically an equa1 chance to meet a paralle1 adjacent chain and an antiparallel one, if the conditions are favorable for the formation of the hydrogen bond between parallel chains, it would be expected that PA-6 crystallize in the y-form. This may lead the crystallization to start with the formation of the sterically favorable y-form and conc1ude with the thermodynamically stable a-form. This has been further confirmed by FTIR analysis. Since the a-form is thermodynamicallY stable, the later region in Figures 6.9.4 and 6.9.5 most likely represents the kinetics of the a-form only. The contribution of the a-form crystallization to the initial region was subtracted to obtain the contribution of yform crystallization in the initial region.
The new slopes and y-intercepts obtained
represent the values of the A vrami kinetic parameters n and K for y-form crystallization.
PA-6
o ~-1
~-2 -3
203'C
205'C
.
207.5'C
-3
210'C
-5
-4
-5
-6.i------------_---5
2 ln
~sec)
200.7'C
211.3'C
~~-----------------~ 4
Inl(sec)
Figure 6.9.4: Crystallization kinetics of PA-6 and PA-6 nanocomposite at 50 MPa
6 RESULTS AND DISCUSIONS
218
PA-6
..
PNC
sr-1
~-1
:s~-2
'Ë"
~-2
..
-3
-3
224.4°C •
-4
225.6 oC -4
• 232.PC
229.3 oC
-5
-5
-6"------------------' 4
4.5
5.5
6
Int(sec)
6.5
7.5
-6 2
5 Int(sec)
Figure 6.9.5: Crystallization kinetics of PA-6 and PA-6 nanocomposite at 150 MPa
Table 6.9.1 shows that the Avrami exponent n was between 1.0 and 3.2 for the "(form ofPA-6, and between 0.9 and 2.6, for the "(-form in PNC. For the a-form ofPA-6, n was between 1.0 and 2.1, and in the PNC, it was between 1.2 and 2.6. The Avrami rate
constants for PA-6 alone and in the PNC are dependent on the crystallization temperature as well as crystallization pressure. The rate constants are substantially higher for the PNC, especially at higher pressures. Hinrichsen and Lux (278) reported a value of3.0 for the Avrami exponent (n) for PA-6 crystallization and between 1.2 and 6.0 for glass fiber reinforced composites. Turska and Gogolewski (279) reported values of n for PA-6 crystallization at atmospheric pressure at temperatures above 210°C to be between 4.0 and 6.0 and below 210°C between 2.8 and 5.0. Yang et al. (280) reported the value of n for PA-6 alone as 4.0 and in PNC with 10% clay as 3.0. The different values of the Avrami exponent for PA-6 reported in the literature may be accounted for by considering the crystallization process to consist of two stages with the formation of a and y-forms in the initial stage and the formation of a-form alone in the later stage. The crystallization half time values for the y-form are lower than those ' for the a-form, indicating that the rate of crystallization is higher for "(-form. Table 6.9.1 shows that, under similar crystallization conditions, the crystallization half time values are lower for the PNC. The nanoclay seems to act as a nucleating agent, and thus it increases the rate of crystallization in PA6 nanocomposite.
6 RESULTS AND DISCUSIONS
219
Table 6.9.1: Avrami parameters and crystallization half-time for y-form and a-form
crystallization of PA-6 and PNC
PA-6 y-form Crystals Supercooling
Press.
Crystn.
MPa
T(°C)
Tm-T
210.0 207.5 205.0 203.0 221.7 220.0 217.0 215.0 228.2 225.8 224.4 222.3 233.2 231.5 227.5
12.5 15.0 17.5 19.5 13.3 15.0 18.0 20.0 16.0 18.4 19.8 21.9 24.8 26.5 30.5
50
100
150
200
Temp.
n
K
2.1 1.1 1.1 0.9 2.5 2.2 0.9 0.6 3.3 2.7 1.6 0.9 1.2 1.3 0.6
3.2 x 10-6 1.5 x 10-3 4.6 x 10-3 1.7 x 10-2 1.7 x 10- 1 2.5 x 10-6 9.0 x 10-3 4.5 x 10-2 7.0 x 10- 10 3.1 x 10-8 1.0 x 10-4 1.1 x 10-2 5.2 x 10"4 4.9 x 10-4 8.8 x 10-2
f1l2
a-form Crystals Sec
n
K
1.8 1.4 1.0 0.8 1.1 1.5 1.1 0.8 2.1 2.1 1.6 1.2 1.7 1.7 1.1
1.1 x 10-5 2.1 x 10-4 4.6 X 10-3 1.9 x 10-2 1.3 x 10-5 6.4 x 10-5 2.3 x 10-3 1.4 x 10-2 6.8 x 10- 1 6.6 x 10-7 8.4 x 10-5 2.0 x 10-3 7.7 x 10-6 1.2 x 10-5 6.9 x 10-3
f1l2
Sec
{oq 330 210 90 70 450 320 110 70 560 510 290 120 420 280 30
530 350 170 110 720 490 210 120 760 720 320 130 810 710 60
PNC y-form Crystals Press. MPa 50
100
150
200
Crystn.
Temp.
Supercooling
T(°C)
Tm-T
211.3 208.7 205.7 222.4 221.0 218.0 232.1 229.3 227.0 225.6 240.5 238.8 237.3
12.2 14.8 17.8 11.8 13.2 16.2 11.2 14.0 16.3 17.7 17.3 19.0 20.5
a-form Crystals
n
K
1.8 0.9
7.9x 10-6 5.3 x 10-3
590 220
2.0 1.9 0.8 2.6 2.0 1.0
1.2 x 1.6 x 2.1 x 8.1 x 2.0 x 7.4 x
10-6 10-5 10-2 10-8 10-5 10-3
840 260 80 450 190 120
1.6 1.7 1.4
4.8x 10-5 6.7 x 10-5 2.3 x 10-3
340 210 60
f1/2
Sec
n
K
{oq 2.1 1.6 0.8 2.5 2.6 1.3 2.1 1.9 1.4 1.3 1.4 1.3 1.2
5.8 x 9.3 x 3.6 x 1.1 x 2.3 x 1.6 x 1.6 x 1.7 x 5.9 x 3.5 x 1.6 x 5.3 x 2.3 x
f1/2
10-7 10-5 10-2 10-8 10-7 10-3 10-6 10-5 10-4 10-3 10-4 10-4 10-3
Sec
810 240 40 1180 340 120 480 240 140 60 360 240 110
6 RESULTS AND DISCUSIONS
220
Figure 6.9.6 shows the infrared spectra of PA-6 and PNC for slow cooled and quenched specimens. The characteristic IR frequencies (cm- 1) for PA-6 are shown in Table 6.9.2. Both, the quenched as weIl as the slow-cooled PA-6 samples, show the presence of the a-form with absorption peaks at 690, 833, 927, 960 and a shoulder at 951 cm- l. The quenched PA-6 sample has a weak peak at 625 cm- 1 and a weak shoulder around 777
cm-l, indicating the presence of the y-form. At 975 cm-l, the quenched
sample has a small peak and the slow-cooled sample has a shoulder indicating the presence of the mesomorphous form. The FTIR spectrum of the slow-cooled PNC sample is very similar to that of PA-6, showing the strong presence of the a-form and weak presence of the y-from. The spectrum of the quenched PNC sample, unlike aIl other samples, shows strong peaks at 626,777,916, 975 and a shoulder at 1000 cm-l, indicating the presence of the y-form and mesomorphous form. The peaks related to the a-form such as 690, 833, 927, 951 and 961 cm- l are either absent or very weak. In the case of PA-6, quenching preserves sorne of the y-form crystals, showing that the formation of y-form is a prerequisite to the formation of a-form crystals. This is further evident from the spectrum of the quenched PNC sample. The nanoclay stabilizes the yform crystal s, which are grown from the melt during the early stage of crystallization. Either it is the high rate of crystallization in the presence of nanoclay that prevents the transformation of the y-form into the a-form, or the nanoclay particles interfere with the transformation process, which involves the shifting ofR-bonds. Table 6.9.2: Characteristic IR frequencies (cm- 1) for PA-6
Crystalline Form
Wave Number (cm- 1)
Reference
a + y + mesomorphous a+y a + amorphous a Y+ mesomorphous
1028-1031, 1079 522, 730 575-580 690,833,927, 951 (weak shoulder), 961 918,975 623-630" 1000(weak shoulder), 777
(281,282,283 ) (284) (282,284,285,286) (282-286) (281-283) (281-286)
Y
6 RESULTS AND DISCUSIONS
221
PM 0Jerd1ed ••• PA-69OJVCooIed
-
.
1 Il
i
\
"U
• 1027.87
1 : 1 t
.~
"
~:
~I \'
\
S29.52
F • •
958.45
1 1
"I l Il
"
522.61
1100
1000
800
700
600
500
WiN9 rurber(/cm)
- Pf\C0Jerd1ed ••• Pf\C 90JV CooIed
l
~
833.1
t F
" " ~
gn~
.
975.8
~ 1
\; 578.54 1100
1000
gx)
800
700
600
Wave N.Irtler (lem)
Figure 6.9.6: FTIR spectra of PA-6 and PA-6 nanocomposites: Slow cooled and quenched
500
7 CONCLUSIONS
222
7 CONCLUSIONS 1. The clay particles were modeled as the stacks of parallel singular clay platelet units. The adhesive energy and the adhesive force between the platelets were estimated using Hamaker approach. It was shown that the breaking of the clay particles into smaller units (tactoids) requires high shear stresses; that may not be accessible in common extrusion equipment. Erosion or surface peeling appears to be the likely mechanism of size reduction of clay particles. The peeling mechanism requires lower shear stresses, which are achievable in pol ymer extrusion. The exfoliation of the clay platelets in pol ymer melts by peeling of platelets from the surface ofthe clay particles necessitates higher residence time for processing. 2. In a compatible polymer/clay system, the polymer chains have strong affinity towards the organic modifier residing between the platelets. The polymer chains intercalate between the galleries and the peeling of platelets sets in. The peeling of . tactoids of untreated clay requires higher shear stresses than those of organoclay. The required shear stress for peeling of tactoids from the particle surface decreases with increase in the tactoid surface area and increases with the tactoid thickness. The required shear stress for peeling significantly decreases with increasing peeling angle. 3. A continuous process for producing nanocomposites was designed based on the modeling results. The process employs an ISO (lnterfacial Surface Oenerator) static mixer, placed between a twin-screw extruder and a die. The new processing system (System A) has the following features. • Imparts sufficient shear stresses for the exfoliation of organoclay in the polymer melt. • Increases the residence time for processing nanocomposites. • Incorporates frequent changes in flow direction for effective peeling. • Imparts effective spacial distribution of clay platelets by chaotic mixing, efficient stretching and extension. • Easy to install, easy to scale up and inexpensive.
223
7 CONCLUSIONS
4. Wide angle X-ray diffraction and transmission e1ectron microscopy of PA-6 nanocomposites showed that the organoclay was better exfoliated in the samples prepared by using System A than in the samples prepared using conventional twinscrew
extrusion
(System
C).
In
the
case
of polystyrene/Cloisite
10A
nanocomposites, System A samples showed higher increase in gallery spacing and less collapsed structures in the organoclay than System C samples. 5. The efficiency of the proposed nanocomposite processing system (System A) was evaluated by comparing the physical properties of the nanocomposites produced by this process with those produced using conventional twin-screw extrusion (System C) under the same processing conditions. In the case of PA-6 nanocomposites with organoclay and untreated clay, System A samples showed significant improvement in the yie1d stress and the tensile modulus compared to System C samples. The oxygen
permeability
values
of
PA-6
nanocomposites
and
polystyrene
nanocomposites produced using System A were lower than those of the System C nanocomposites. The water absorption of the PA-6/organoclay nanocomposite was lower than that ofPA-6. 6. The aspect ratios of the organoclays in nanocomposites were estimated using tensile modulus data, in combination with different composite mode1s available in the literature. The estimated aspect ratios of the filler for PA-6 nanocomposites were higher for the System A specimens. The aspect ratios of the filler clays were also estimated using oxygen permeability data and the tortuosity mode1s. The estimated aspect ratios were higher for the System A specimens. The yield stress data of nanocomposites . were used to calculate the values of parameter B in Pukanszky model which reflects the interfacial interaction, the interfacial strength and the specific surface area of the fiUer particles. Parameter B was positive for PA-6 nanocomposites.
System A specimens showed higher values than System C
specimens, indicating higher specific surface area of the dispersed clay platelets. The parameter B was negative in polystyrene nanocomposites, indicating po or adhesion between the clay particles and the polymer matrix.
224
7 CONCLUSIONS
7. PA-6/organoclay nanocomposites showed an increase in viscosity at low shear rates as well as more pronounced shear thinning behavior, in comparison with the unfilled pol ymer, which is characteristic of nanocomposites containing exfoliated clay. Overall, viscosities of System A samples were lower than those of System C samples at
comparable
shear
rates.
PA-6/untreated
clay
composites
and
polystyrene/organoclay composites did not show the above rheological behavior. 8. The shear stresses in the ISG static mixer were estimated using the rheological data and experimental data regarding the pressure drop across the static mixer elements during processing. The calculated shear stresses encountered in ISG static mixer during the processing of PA-6 nanocomposites are sufficient for the exfoliation of the organoclay. The available shear stresses in the ISG static mixer decrease with the increase in clay content due to the changes in the rheological behavior of the PA-6 nanocomposites. In the case of processing of polystyrene nanocomposites, the available shear stresses in the ISG static mixer are lower than the required shear stresses for organoclay exfoliation. 9. The melting and crystallization behavior, and the crystallization kinetics of PA-6 alone and in nanocomposite at elevated pressures were studied using high pressure dilatometry. The melting temperatures of PA-6 in the nanocomposite were the same as for PA-6 alone. However, the crystallization temperatures were lowered in the presence of clay. The crystallization of PA-6 at elevated pressures occurs in two stages, with the formation of the y-form and the transformation of the y-form into the a-form in the initial stage, and the formation of a-form alone in the later stage. The Avrami parameters and the crystallization half times were determined for the crystallization of PA-6 in nanocomposites and PA-6 alone. The nanoclay seems to act as a nucleating agent, thus raising the rate of crystallization of PA-6 in nanocomposite. The presence of the y-form during the initial stages and its transformation into the a-form during the crystallization pro cess was confirmed by infrared spectroscopy of quenched and slow cooled sampi es of PA-6 and PA-6 nanocomposites.
8 RECOMMENDA TIONS FOR FUTURE WORK
225
8 RECOMMENDATIONS FOR FUTURE WORK 1. The exfoliation of the clay platelets in non-polar polymers such as polyethylene or
polystyrene, and the production of the nanocomposites with these polymers, havirig substantial improvements in their physical properties has been still an elusive task. By using the solubility parameters and the Hamaker constants of the components it is possible to predict the formation and the structure of the nanocomposites. This information could be used to make. tailor made organic modifiers to treat the inorganic clays. As an example, fluorine compounds have low Hamaker constants and the use of these compounds as organic modifiers will result in positive effective Hamaker constant of clay/modifier/pol ymer system. This will result in engulfment of the clay particles by the polymer matrix and better adhesion between the clay and the polymer chains. 2. Clays which are resistant to high temperature, especially phosphonium modified organoclay could be tried for systhesizing nanocomposites. This will avoid degradation of clay as well as polymer at high processing temperatures.
9 ORIGINAL CONTRIBUTIONS TO KNOWLEDGE
226
9· ORIGINAL CONTRIBUTIONS TO KNOWLEDGE 1. Based on the classical theories of interparticle interactions, a simple mathematical model was formulated, for the exfoliation of clay platelets in a polymer matrix during melt processing of the nanocomposites. This mode1 facilitates the estimation of the
~hear
stresses required for the delamination of the untreated clay and the
organoclay. 2. A novel continuous process was deve10ped and implemented for the production of the nanocomposites. The process incorporates chaotic mixing, efficient stretching of pol ymer melt, effective distributive mixing, shear stresses sufficient for the exfoliation of the platelets, frequent changes in flow direction for the effective peeling of the plate1ets, and higher residence time. The efficiency of the pro cess was demonstrated. The process was used to obtain and demonstrate a high degree of exfoliation in P A-6/Montmorillonite clay nanocomposites. 3. The crystallization kinetics of PA-6 and PA-6 nanocomposites were studies at elevated pressures.
The different values of the Avrami exponent for PA-6 reported
in the literature were accounted for by considering the crystallization process to consist of two stages, with the formation of both ex and y-forms in the initial stage and the formation of the ex-form alone in the later stage.
10 REFERENCES
227
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1. Sandermann, A. Keller, J. Polym. Sei., 19,401 (1956).
274.
S. Gogolewski, A. J. Pennings Polymer, 18,654 (1977).
275.
S. Gogolewski, A. J. Pennings Polymer, 16,673 (1975).
276.
N. Hiramatsu, S. Hirakawa, Polym. J., 14, 165 (1982).
277.
H. Arimoto, M. Ishibashi, M. Hirai, Y. Chatani, J. Polym. Sei., Part A, 3,317 (1965).
278.
G. Hinrichsen, F. Lux, Polym. Bull. (Berlin), 24(1), 79 (1990).
279.
E. Turska, S. Gogolewski, Polymer, 12(10),629 (1971).
280.
F. Yang, Y. Ou, Z. Yu, J. Appl. Polym. Sei., 69, 355 (1998).
281.
G. Rotter, H. Ishida, J. Polym. Sei.: PartB: Polym. Phys., 30,489 (1992).
282.
K. H. Illers, H. Haberkom, P. Simak, Makromol. Chem.., 158, 285 (1972).
10 REFERENCES 283.
240
L. Panel Pierron, R. Seguela, J.-M. Lefebvre, V. Miri, C. Depecker, M. Jutigny, J. Pabiot, J. Polym. Sei.: PartE: Polym. Phys., 39, 1224 (2001).
284.
M. I. Kohan, Nylon Plastics Handbook, Hanser Publishers (1995).
285.
I. Matsubara, J. M. Magill, Polymer, 1, 199 (1966).
286.
I. Matsubara, Y. Itoh, M. Shinomiya, Polym. Lett., 4,47 (1966).
i
APPENDICES
APPENDIX-A CALCULATION OF HANSEN SOLUBILITV PARAMETERS
Solubility parameters are used to predict compatibility of polymers, chemical resistance, permeation rates, to characterize the surfaces of pigments, fibers, and fillers. The Hilderbrand solubility parameter is defined as the square root of cohesive energy density: l
o=(~)2 V is the molar volume of the pure solvent, and E is its energy of vaporization. The total energy of vaporization consists of several parts. These arise from (atomic) dispersion forces, (molecular) permanent dipole-permanent dipole forces, and (molecular) hydrogen bonding (electron exchange). The non-polar interactions derive from atomic forces and are called as dispersion interactions. The energy of vaporization due to these forces is dispersion cohesive energy, ED • The permanent dipole-permanent dipole interactions cause polar cohesive energy, Ep • These are molecular interactions. The third cohesive energy source is hydrogen bonding, EH. This is called as an electron exchange parameter. The basic equation which govern the assignment of Hansen parameters is that the total cohesive energy, E, must be the sum of the individual energies which make it up. E=ED+Ep+EH Dividing this by molar volume gives the square of the total Hilderbrand solubility parameter as sum of squares of Hansen D, P, and H componenets. EN = EoN + EpN + EHN
82
=
8D2 + 8P2 + 8H2
Table A.I gives the values of the components of Hansen Solubility Parameters for different functional groups. These values can be used to ca1culate Hansen Solubility Parameters of organic modifiers.
APPENDICES
ii
Table A.1: Group contribution parameters for the components of Hansen Solubility Parameter (227) Group Contributions to Partial So!ubilîty Parilmeters P~ar
Mo"" VOhlfUt?-, ~
funttionat
•••........
~.~---
Aliph.tk Aromalic"
Group
EI~tron
Paramder/ Arom3tk
Aliphatk
ü
o o o
CH J .
J.\5
1.l25
16.1 ·,"IJ)
us"
î)
(j
820
(j
()
ù
o
f~H"2
35ü
-~9.:~
:.:.: orelJu
-CH
:
SOI)"}: WO
H>fJ± 20'
Sfjm~
1.10n 1.650 1,760
t ,310 ± .500
L250:t WU
J .87tJ ± 600
L#O± WU
.VjI'}±
}:iOO.1: J50 ~()O±
j,ii50.t. 140
150
"
l.IXI()± 150
fi ()
;!$f>
(j
L~W
j,250 ±.lOO SOO.t.1'1)" 3.\0 ± 15(~ 1,250* 100 gO\') i :.'!5(j
..,)±),~C'
8{.IO:t ::!OO
H)I)±;O
~Û'\.1;t, 200": 4üuoojw:em C amms.
tl)
,"'akulal~, whh V for
tüW ;.;m:npouud.
CALCULATION FOR ORGANIC MODIFIER OF CLOISITE 308
Cloisite 30B is modified with methyl tallow bis-2-hydroxyethyl quaternary ammonium compound. CH2CH20H 1
CH3 -W-T 1
CH2CH20H
APPENDICES
111
T is tallow which contains =65% C18, :::.30% C16, and =5% C14. 'T' contains average 16.2, -CH2-, and one, -CH3 groups. This makes total 20.2, -CH2-, groups, 2, -CH3, groups, 2, hydroxyl groups and one amine group in Cloisite 30B modifier. From Table A.l we can get values for AV and AV t!1 for D, P and H parameters for each of these functional groups and then calculate
bD,
Ôp and
41.
Table A.2: Partial solubility parameters of functional groups
Functional Groups
~
ti>
&
8
(cal/cm3) 1/2
(cal/cm3)112
(cal/cm3)1/2
(cal/cm3)112
-CH3
5.8
0
0
5.8
-CH2-
8.56
0
0
8.56
-OH
13.3
8.4
21.6
26.7
Phenyl
10.27
0.83
0.83
10.3
Amine
16.8
4.7
12.9
21.1
For Closite 30B modifier partial solubility parameters can be ca1culated as following: ~ =(20.2x8.56 + 2x5.8 + 2x13.3 + 16.8)/25.2 =9.0441 (cal/cm3)112 =18.49 (J/cm3)112
a.. =(20.2xO + 2xO + 2x8.4 + 4.7)/25.2 =0.8532 (cal/cm3)112 =1.745 (J/cm3)112 ~
=(20.2xO + 2xO + 2x21.6 + 12.9)/25.2 =2.2262 (cal/cm3)112 =4.5537 (J/cm3)112
a= (18.492 + 1.7452 + 4.5537 2)1/2 = 19.12 (J/cm3)1/2 Similar calculations were done for Cloisite 15A, and for Cloisite lOA. The ca1culated values of Hansen Solubility Parameters are shown in Table 6.2.1.
APPENDICES
IV
APPENDIX- B CALCULATION OF HAMAKER CONSTANT OF ORGANOCLAY The group contribution method to detennine Hamaker constant of a chemical compound consists of attributing to each chemical group of a molecule or macromolecule a contribution to the Hamaker constant and to the surface free energy (229). The calculated values of their surface free energies are found to be fairly good agreement with those obtained from wettability measurements or by surface tension measurements of molten polymers. The surface free energy of a compound containing n groups can be written as:
1
Yn = -[xYx + yyy + ZYz + ........ ]
n where x, y, z, .... are the number of groups X, Y, Z, .... fonning the molecule. Thus x + y + Z +..... = n; and ]X, 71', 1Z .... represent the contribution of each group to the surface energy Yn. Considering that the surface free energy results from Van der Waals' forces Yn can be also expressed as a function of the Hamaker constant, An, with the following relationship:
where Do is interatomic cut-off distance which approximately equals 0.165 nm. The group contribution values for surface energies at 20°C for different chemical groups are given in Table B.l. Table 8.1: Group contribution values of YG at 20°C (mJ/m 2)
Group
-6.4
-CHz-
Tertiary amine
Phenyl
-CHzOH
30.8
38.92
27.56
47.5
APPENDICES
v
CALCULATION FOR ORGANIC MODIFIER OF CLOISITE 30B
Cloisite 30B organic modifier contains 2, -CH3 groups, 18.2, -CH2- groups, 2, -CH20H groups, and one tertiary amine group. Total number of groups is 23.2. 'YM30B
=[2x(-6.4) + 18.2x30.8 + 2x47.5 + 38.92] /23.2 =29.38 rnJ/m2
Hamaker Constant A M30B is given as: AM30B
=241t (O. 165xlO-9 )2 x 29.38 = 6.03xl0- 17 rnJ = 6.03xl0-20 J
Similar calculations were done for Cloisite 15A and Cloisite lOA organiç modifiers. The calculated values are shown in Table 6.2.2. CALCULATION FOR EFFECTIVE HAMAKER CONSTANT BETWEEN PLATELETS OF CLOISITE 30B CLAY
Effective Hamaker constant between two identical particles with a medium between them is given as:
For Cloisite 30B it will be:
~31 = (.J7.8 - .J6.03 ) X 10-20 = 0.1137 X10-20 J where 7.8xlO-2o J is the Hamaker constant between pristine Montmorillonite clay platelets [1] and 6.03xlO-2o J is the Hamaker constant of organic modifier [3]. CALCULATION FOR EFFECTIVE HAMAKER CONSTANT BETWEEN PA-6 AND PLATELET OF CLOISITE 30B CLAY
Effective Hamaker constant between two dissimilar particles with a medium between them is given as: A132
= (~Au - ~ A33 X~ A 22 - ~ A33 )
For PA-6/Cloisite 30B system it will be given as:
~32 = (.J7.8 -.J6.03
X.Ji2 -.J6.03 )XI0-
where Hamaker constant ofPA-6 is 12.0xlO-2o J.
20
= 0.34xl0-
20
J
......---~
1
APPENDICES
VI
APENDIX - C ANALVSIS OF TRANSMMISSION ELCTRON MICROGRAPHS C.1
ESTIMATION OF SURFACE DENSITY
Estimation of surface density requires a series of linear test probes. Higher the surface area per unit volume of the feature of interest, the larger the number of intersections between the surface and the set of linear test probes. The relationship between the number of intersections and the surface area per unit volume is given by the equation
Sv=2h which means that surface density is equal to twice the number of intersections between the surface and the linear probe, per unit length of test line in the reference space. This equation is an unbiased estimator of surface density only if either the surface is isotropic, the lines are isotropie or both are isotropic. In this context the lines must be isotropic in 3D space. An isotropic line probe is a line that has been generated with isotropic orientation in 3D space. In practice there are two ways of generating isotopic line probes: (i) on vertical uniform random (VUR) sections and (ii) on isotropic uniform random (!UR) sections. In practice VUR sectioning is, in the majority cases, the method of choiee. A VUR section is not an isotropic section in 3D space. What is required is a set of lines oriented so that their length density is proportion al to sine 8. When a cycloid is aligned so that is minor axis is parallel with the vertical direction it has a length density that is proportional to sine 8. A cycloid can be generated by following one point on the circumference of a circle as that circle is rolled along a line. Cycloid grid has a series of associated points, such that there is a known length of line per point. The combination of a plane generated with a VUR protocol and a grid of cycloids oriented in this manner is equivalent to a collection of !UR lines in 3D space.
APPENDICES
vii
1
Min()!'
li2
axh
~---Jl Figure C.1: Cycloid curve (245)
Grid Cl
Figure C.2: Cycloid test grid (245)
APPENDICES
Vlll
Figure C.3: Superimposing cycloid test grid on TEM micrograph
ESTIMATION PROCEDURE
A cycloid test system, with its minor axis parallel to the vertical direction is randomly translated in the x and y direction on the image. The test system has a known length of cycloid per point (lIp). For example in the image shown in Figure C.3, l/p is 0.45 J.UI1. In order to estimate surface density, two counts need to be made on this image, (i)
the number of intersections between the cycloid tines and boundary of interest (l)
(ii)
the number of points that land within the reference space (P).
Surface density is then estimated from,
ix
APPENDICES
n
2I!i
SV (Y reif) = ' l
i=1
n
-I~ p i=1
For ex ample in Figure C.3, there are 15 points hitting the reference space and 82 intersections. The surface density is, Sv
=(2x82) / (0.45x15) =24.3 Jllll-I
In each case ten measurements were made after randomizing the test grid. The volume fraction of the silicate particles was estimated from TEM micrographs as the ratio of the length of test line falling on the particles to the totallength of the test line. C.2
ESTIMATION OF PARTICLE LENGTH
The particles from gray scale image were manually traced on a transparent layer overlaid on micrograph using image processing software (Adobe photoshop 6.0). The binary image is processed by using image analysis software (Scion Image Beta 4.02). The particle length distribution is created from the data obtained.
Figure C.4: Converting gray image into binary black and white image (244)
~..
APPENDICES
x
APPENDIX- 0 PVT APPARATUS
The fully automated GNOMIX high-pressure dilatometer shown in Figure D.1 was used in the present study. Sample weighing 1-2 grams is placed in a steel cell with flexible bellows at one end. The total volume of the cell is 7.0-7.5 cm3 • The cell is filled with mercury as confining fluid. The sealed cell is placed inside a vessel that is under hydraulic pressure of speCial silicon oil. The hydrostatic pressure of silicon oil is produced by a high pressure motorized or manual pump. This pressure is transmitted to the contents of the cell. A linear variable differential transducer (LVDT) located below the pressure vessel, measures the deflection of the bellows as a result of temperature and/or pressure changes. The deflection is measured with a resolution of 0.001 mm, representing a volume change of approximately 0.0001 cm3 • Using the known PVT properties of the confining fluid (mercury), the deflections are converted to volume changes of the sample. The pressure vessel is sealed at two places (at the interface to the base and to the thermocouple block on top) by vertical force of 25 ton hydraulic ram, mounted on a support plate above the PVT apparatus assembly. The thermocouple block on top inc1udes a high pressure thermocouple (type K) which measures sample temperature. A .separate thermocouple is used for temperature controller, which is located in a hole drilled about 1.3 cm into the side of pressure vessel. A temperature calibration is done to relate the sample temperature to the controlled temperature. The PVT experiments can be done in three modes, isothermal, isobaric and isothermal/isobaric. During isothermal experiments the sample is held at predetermined temperature and pressures are varied to get change in volume. In isobaric mode, pressure is held constant and sample is heated or cooled at predetermined rate while change in volume is recorded. The isothermal/isobaric mode is used for the crystallization kinetics experiments where pressure and temperature are held at certain predetermined values and the change in volume is recorded with respect to time.
APPENDICES
xi
Oll CATCH PAN
/ puacp $I'IUT OFF VAt. REU!F V,U.lfE
QAVO!!
Sltl.l.OViS f>lflZOMET!R
1
! D·1-----PUMP
~..
Figure 0.1: Schematic drawing of GNOMIX PVT apparatus
APPENDICES
xü
PVT BEHAVIOR OF PA-6 AND PA-6 NANOCOMPOSITES
Isobaric PVT experiments were carried out to obtain the melting points and the crystallization temperatures of PA-6 and PNC at different pressures. The melting point and crystallization temperatures were defined as the temperatures at which the rate of change of volume was highest during heating and cooling, respectively. Figure D.2 and Figure D.3 show the typical PVT behavior of PA-6 and PA-6 nanocomposite (Ube sample with 2%wt clay) respectively during isobaric heating and cooling. Isobaric experiments were performed at heating and cooling rates of 2.5°C/min and at pressures of 10, 50, 100, 150 and 200 MPa. For each isobaric heating experiment (to determine Tm), the sample used was obtained by cooling from the melt (250°C) at 2.5°C/min under 10 MPa pressure. This ensures that the sample used has the same thermal history. The accuracy of volume measurement is 0.002 cm3/gm at 25-250°C and 0.004 cm3/gm above 250°C, and the sensitivity is 0.0002 cm3/gm. The fluctuations in pressure were within 0.1 MPa. The reproducibility of the specific volume measurement was 0.003 cm3/gm. These data provide information about the melting point Tm and the crystallization temperature Tc at different pressures. To obtain the exact values, the rate of change of specific volume is plotted against temperature. The maxima in the heating runs and the minima in the cooling runs give Tm and Tc, respectively. These plots are shown in Figure 6.9.1. The local curve smoothing technique by usingbisquare weighting and polynomial regression was used to reduce the noise in the data shown in Figure 6.9.1. The Pearson correlation coefficients between the original data and the smoothed data were ca1culated. In case of PA-6, they were in the range 0.97-0.99 for the heating curves, and 0.95-0.99 for the cooling curves. In case of nanocomposite, they were in the range 0.96-0.98 and 0.98-0.99 for the heating curves and the cooling curves, respectively. The reproducibility of Tm and Tc was ±1 oC. Crystallization kinetics experiments were done under isothermallisobaric mode. The typical data obtained for PA-6 nanocomposite crystallization at 150 MPa and different crystallization temperatures are shown in Figure D.4.
,,,..---
APPENDICES
X1ll
0.20
10 MPa 50MPa
0.15
100 MPa
ÊCl
,;;--
150 MPa
E .3- 0.10
200 MPa
CI)
Cl
c::
ca
.r:::.
Ü (5
0.05
> u
~CI) Cl.
0.00
en
-0.05
+----,----,-----,-------,-----,----,-------1
o
50
100
150
200
250
300
350
Temperature (oC)
Figure 0.2: Specifie volume change during isobaric heating and cooling runs of PA-6.
0.20 - , - - - - - - - - - - - - - - - - - - - - - - - - - - - - - , 10MPa 50MPa 0.15
Ê
100 MPa
Cl
M--
E
u
150 MPa
-; 0.10 Cl
c::
200 MPa
ca
.r:::. Ü
ci
>
0.05
o CI)
Cl.
en
0.00
-0.05
+----,----,-----,--------,----,----,------j
o
50
100
150
200
250
300
350
Temperature (oC)
Figure 0.3: Specifie volume change during isobaric heating and cooling runs of PA-6 nanocomposite.
APPENDICES
xiv
0.075
~ -
o v o
E 0.070 E 0
Crstallization at 229°C Crystallization at 232°C Cryatllization at 227°C
0.065
Q)
E ::::J
0
>
0.060
0
'+=
'0 0.055 Q) a.
en c
Q)
0.050
C>
c
CIS
..c
()
0.045
0.040 0
500
1000
1500
2000
Time (Sec) Figure 0.4: Isothermal/isobaric crystallization of PA-6 nanocomposite at 150 MPa
APPENDICES
xv
APPENDIX- E RHEOLOGY OF NANOCOMPOSITES
The rheological propepies of the polymers and their nanocomposites were determined by steady shear viscosity measurements using Instron capillary rheometer (Model TTCM). The capillary die with length to diameter ratio 60 was used to avoid the entrance effect correction. The shear rate was varied by changing the crosshead speed. To determine the true wall shear rate and viscosity of polymer melts, the Rabinowitch correction was applied. The viscosity-shear rate data were fit to power law equation,
17 =
Kr
n 1 - )
where K is power law constant and n is power law index. The values of these parameters obtained for different nanocomposites are given in Table E.l and Table E.2. Table E.1: Power law parameters for PA-6 nanocomposites at 240°C
Processing System
K (Pa sn)
n
PA-6
1618
0.84
PA-6/Cloisite 30B (3.1 % wt)
1609
Material
System A
0.79 ..
PA-6/Cloisite 30B (5.0% wt)
3594
0.62
PA-6/Cloisite 30B (3.1 % wt)
2526
0.75
PA-6/Cloisite 30B (5.0% wt)
4473
0.61
PA-6/Cloisite 15A (2.5% wt)
2169
0.75
PA-6/Cloisite 15A (4.1 % wt)
3516
0.62
PA-6/Cloisite 15A (2.5% wt)
2333
0.74
PA-6/Cloisite 15A (4.1 % wt)
4351
0.62
PA-6/Cloisite Na+ (6.9%wt)
1469
0.86
System C
System A
SystemC ,r·
APPENDICES
xvi
Table E.2: Power law parameters for PS nanocomposites at 21 QOC
K (Pa sn)
n
PS-1301
6845
0.51
PS-1301lCloisite lOA (2.6%wt)
6019
0.55
PS-1301/Cloisite 10A (3.5%wt)
4962
0.57
PS-1301/Cloisite 30B (3.0%wt)
0.47
PS-1301/Cloisite 15A (2.5%wt)
9797 , 8812
PS-130112% Dylark
6284
0.54
PS-130112% DylarklCloisite lOA (3.5%wt)
4640
0.57
Material
0.48
xvii
APPENDICES
APPENDIX- F SLiT DIE Kama! et al. (215) achieved laminar morphology in PPIEVOH blends with good barrier properties using specially designed die and process manipulation. _The die design incorporated converging and diverging sections in order to help in stretching dispersed phase partic1es in two directions (machine and transverse directions). In the present study this die shown in Figure F.l was used with (System A) and without static mixer (System C) assembly to produce sample ribbons of different nanocomposites.
q .'
F
_t..
.:. c
o
x vie.., Z-Z
Figure F.1: Schematic of the die unit
B
A
-
APPENDICES
XVlll
In sections A and C, the flow field is dominated by converging extensional flow;
in section B shear flow is dominant, in section D (in X-Y plane) shear flow is dominant, while in X-Z plane, diverging extensional flow prevails. In section E, converging extensional flow becomes dominant again, while in section F shear flow prevails. Therefore, in sections A, C and E the fluid particles are stretched in the flow direction and in section D in the transverse direction. This biaxial stretching is expected to produce platelets and to develop laminar morphology in polymer blends.
