ABSTRACT Title of dissertation: SYNTHESIS AND CHARACTERIZATION OF LOW FLAMMABILITY POLYMER/LAYERED SILICATE NANOCOMPOSITES Xin Zhang, Doctor of Philosophy, 2009 Dissertation directed by: Professor Robert M. Briber Department of Materials Science and Engineering There has been signi cant interest in the applications of polymer nanocom- posites in a variety of areas. Polymer/layered silicate nanocomposites have been of interest because of relatively low raw material cost and improved materials proper- ties such as higher Young?s modulus, higher thermal deformation temperature, lower small molecule permeability, lower density (compared to metals and traditional glass ber reinforced composites) as well as low ammability. The relationships between the ammability and the dispersion of the layered silicate platelets inside the poly- mer matrix is just being established. The complete set of factors that a ect the ammability of polymer/layered nanocomposites are not fully identi ed. In this thesis polymer/layered silicate nanocomposites with di erent degrees of platelet dispersion were synthesized. The structure of the nanocomposites was characterized by X-ray di raction (XRD), small angle X-ray scattering (SAXS), and transmission electron microscopy (TEM). The ammability of these nanocomposites was characterized by TGA, cone calorimetry and gasi cation. By coupling the structural and ammability data it has been concluded that forming a nanometer scale dispersed structure signi cantly improves the ammability but the details of the degree of dispersion are not critical. The improvement in the ammability arises from the formation of a residue or char layer at the surface of the nanocomposite. This residue layer acts as a radiation shield and as a physical barrier preventing the polymer degradation products from escaping and acting as fuel. It is observed that the stability of the residue layer formed during combustion has major impact on the ammability. This thesis also describes work to improve the ammability of the polymer/layered silicate nanocomposites by enhancing char/residue formation in order to improve the residue layer stability. SYNTHESIS AND CHARACTERIZATION OF LOW FLAMMABILITY POLYMER/LAYERED SILICATE NANOCOMPOSITES by Xin Zhang Dissertation submitted to the Faculty of the Graduate School of the University of Maryland, College Park in partial ful llment of the requirements for the degree of Doctor of Philosophy 2009 Advisory Committee: Professor Robert M. Briber, Chair/Advisor Professor Takashi Kashiwagi Professor Peter Ko nas Professor Lourdes Salamanca-Riba Professor James Quinterie c Copyright by Xin Zhang 2009 Dedication To my wife. ii Acknowledgments I would like to express my sincere thanks to my advisor, Dr. Robert M. Briber, who has given me the opportunity to work in his group, for his advice, encouragement, and support on research and also support during my hard time of my life. A lot of appreciation also goes to Dr. Takashi Kashiwagi, who also contributed a lot of work in this research. I would also thank all committee members, Dr. Lourdes Salamanca-Riba, Dr. James Quintiere, and Dr. Peter Ko nas for all the time and e orts put on my thesis. I would like to thank Dr. Stephen Henderson for all the help on the small angle X-ray scattering and machine shop work. Special thanks go to Dr. Bani Cipriano and Dr. Srinivasa Raghavan?s group for all the works we have done together. A lot of thanks go to all members of Dr. Briber?s group and Dr. Ko nas? group. A lot of thanks go to my friends, Dr. Xiang Wang, Dr. Li Ling, Dr. Zhengkun Ma, Dr. Kehai Zhang, Dr. Lang Chen, Dr. Jason Hattrick-Simpers, Dr. Liyang Dai and Dr.Jianhua Li for all the helps they gave to me. Finally, I would like to thank my family for all the support, especially to my wife, Lei Zheng, whose love supported me through the hardest days in my life, to my grandmother who supported her two sons to college by weaving, later one became a diplomat and the other one became a professor. iii Contents List of Tables vii List of Figures viii List of Abbreviations xiii 1 Introduction 1 1.1 Polymer Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . 1 1.1.1 Nanocomposites with 0-Dimensional Fillers . . . . . . . . . . . 2 1.1.2 Nanocomposites with 1-Dimensional Fillers . . . . . . . . . . . 3 1.1.3 Nanocomposites with 2-Dimensional Filler . . . . . . . . . . . 4 1.2 Clay and Organoclay . . . . . . . . . . . . . . . . . . . . . . . . . . . 8 1.3 Silane Coupling Agent . . . . . . . . . . . . . . . . . . . . . . . . . . 10 1.4 Flammability of Polymers . . . . . . . . . . . . . . . . . . . . . . . . 12 1.4.1 Combustion and Pyrolysis of Polymers . . . . . . . . . . . . . 12 1.4.1.1 So Called "Self-Extinguishing" Polymers . . . . . . . 17 1.4.1.2 Flame Retardant Additives . . . . . . . . . . . . . . 18 1.4.2 Flammability Measurements . . . . . . . . . . . . . . . . . . . 20 1.5 Flammability of Polymer Nanocomposites . . . . . . . . . . . . . . . 21 2 Experimental 23 2.1 Research Outline . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 23 2.2 Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 23 2.2.1 Resin Polymer . . . . . . . . . . . . . . . . . . . . . . . . . . . 23 2.2.2 Additives . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 24 2.2.2.1 Cloisite 30B . . . . . . . . . . . . . . . . . . . . . . . 24 2.2.2.2 Cloisite 15A . . . . . . . . . . . . . . . . . . . . . . . 24 2.2.2.3 Somasif MAE120 . . . . . . . . . . . . . . . . . . . . 25 2.2.2.4 Synthesis of PMMA Grafted Layered Silicate . . . . 25 2.3 Synthesis of Layered Silicate/Polymer Nanocomposite . . . . . . . . . 27 2.3.1 In-situ Polymerization . . . . . . . . . . . . . . . . . . . . . . 27 2.3.2 Solution Mixing . . . . . . . . . . . . . . . . . . . . . . . . . . 28 2.3.3 Synthesis of Nanocomposites with PMMA Grafted Clay . . . 28 2.4 Structural Characterization of Nanocomposites . . . . . . . . . . . . . 28 2.4.1 X-ray Di raction . . . . . . . . . . . . . . . . . . . . . . . . . 30 2.4.2 Small Angle Scattering . . . . . . . . . . . . . . . . . . . . . . 31 2.4.3 Transmission Electronic Microscopy . . . . . . . . . . . . . . . 33 2.5 Flammability Measurement . . . . . . . . . . . . . . . . . . . . . . . 34 2.5.1 Thermogravimetric Analysis . . . . . . . . . . . . . . . . . . . 34 2.5.1.1 Degradation Kinetics . . . . . . . . . . . . . . . . . . 35 2.5.2 Cone Calorimetry . . . . . . . . . . . . . . . . . . . . . . . . . 37 2.5.3 Gasi cation . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39 2.6 Rheology Measurement . . . . . . . . . . . . . . . . . . . . . . . . . . 40 iv 3 Results and Discussion 41 3.1 X-ray Di raction from Clays . . . . . . . . . . . . . . . . . . . . . . . 41 3.2 In-situ Polymerization Samples . . . . . . . . . . . . . . . . . . . . . 41 3.2.1 Dispersion of Clay in MMA Monomer . . . . . . . . . . . . . . 41 3.2.2 Polymerization of the Nanocomposites . . . . . . . . . . . . . 43 3.2.3 Structure of in-situ Polymerized Samples . . . . . . . . . . . . 45 3.2.3.1 X-ray Di raction . . . . . . . . . . . . . . . . . . . . 45 3.2.3.2 Transmission Electron Microscopy . . . . . . . . . . 45 3.2.4 Flammability measurement of in-situ Polymerized Samples . . 47 3.2.5 Summary of in-situ Polymerized Nanocomposites . . . . . . . 52 3.3 Solution Mixing Nanocomposite Structure . . . . . . . . . . . . . . . 52 3.3.1 Transmission Electron Microscopy . . . . . . . . . . . . . . . . 52 3.3.2 X-ray Di raction . . . . . . . . . . . . . . . . . . . . . . . . . 63 3.3.2.1 X-ray Di raction from Nanocomposites . . . . . . . . 63 3.3.3 Small Angle X-ray Scattering . . . . . . . . . . . . . . . . . . 65 3.4 Discussion on the Dispersion of Layered Silicate in Polymer . . . . . . 68 3.5 Flammability Tests . . . . . . . . . . . . . . . . . . . . . . . . . . . . 71 3.5.1 Thermogravimetric Analysis of Nanocomposites Made by the Solution Mixing Method . . . . . . . . . . . . . . . . . . . . . 71 3.5.1.1 Clay Samples . . . . . . . . . . . . . . . . . . . . . . 71 3.5.1.2 Nanocomposite Samples . . . . . . . . . . . . . . . . 73 3.5.2 Gasi cation Experiments . . . . . . . . . . . . . . . . . . . . . 103 3.5.3 Cone Calorimetry . . . . . . . . . . . . . . . . . . . . . . . . . 107 3.5.4 E ect of Clay Dispersion . . . . . . . . . . . . . . . . . . . . . 115 3.5.4.1 Solvent Introduced Dispersion Change . . . . . . . . 115 3.5.4.2 PMMA Grafted Clay Introduced Dispersion Change 119 3.5.5 Residue Layer . . . . . . . . . . . . . . . . . . . . . . . . . . . 120 3.6 Rheology Measurements . . . . . . . . . . . . . . . . . . . . . . . . . 125 3.7 Glass Transition Measurement . . . . . . . . . . . . . . . . . . . . . . 125 3.7.1 Di erential Scanning Calorimetry . . . . . . . . . . . . . . . . 125 3.8 Discussion on the Flammability of Nanocomposites . . . . . . . . . . 131 3.8.1 Mechanism of Low Flammability Nanocomposites . . . . . . . 131 3.8.1.1 Structure Integrity . . . . . . . . . . . . . . . . . . . 133 3.8.2 E ect of Heat Conductivity on the Fire Safety . . . . . . . . . 133 4 Phosphorus Containing Nanocomposite 136 4.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 136 4.2 Char Forming Compound . . . . . . . . . . . . . . . . . . . . . . . . 137 4.3 Synthesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 138 4.3.1 Adding phosphorus to Cloisite 30B Clay . . . . . . . . . . . . 138 4.4 Structure Characterization with TEM . . . . . . . . . . . . . . . . . . 138 4.5 Flammability Test . . . . . . . . . . . . . . . . . . . . . . . . . . . . 139 4.5.1 Gasi cation . . . . . . . . . . . . . . . . . . . . . . . . . . . . 139 4.5.2 Cone Calorimetry . . . . . . . . . . . . . . . . . . . . . . . . . 139 4.6 Residue Characterization . . . . . . . . . . . . . . . . . . . . . . . . . 144 v 4.7 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 148 4.7.1 Char Formation . . . . . . . . . . . . . . . . . . . . . . . . . . 148 4.7.2 Position of the phosphorus Compound During Combustion . . 149 5 Conclusions 150 6 Future Work 154 6.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 154 6.2 Proposed Experiment . . . . . . . . . . . . . . . . . . . . . . . . . . . 156 6.2.1 Material Selection . . . . . . . . . . . . . . . . . . . . . . . . . 156 6.2.2 Organic Modi cation of Nanoparticles and Nanocomposite Synthesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 157 6.2.3 Structure Characterization . . . . . . . . . . . . . . . . . . . . 160 6.2.4 Flammability Characterization . . . . . . . . . . . . . . . . . . 160 Bibliography 161 vi List of Tables 2.1 Solution Precipitated Samples . . . . . . . . . . . . . . . . . . . . . . 29 2.2 Relationship Between the Heat Flux and Fire Scale . . . . . . . . . . 39 3.1 Silicate Stack Thickness Based on Scherrer Equation . . . . . . . . . 65 3.2 Kinetic Parameters Derived from Pure Polymer TGA Data . . . . . . 98 3.3 Kinetic Parameters Derived from Cloisite 15A Low Molecular Weight Nanocomposite TGA Data . . . . . . . . . . . . . . . . . . . . . . . . 99 3.4 Kinetic Parameters Derived from Cloisite 30B Low Molecular Weight Nanocomposite TGA Data . . . . . . . . . . . . . . . . . . . . . . . . 100 3.5 Kinetic Parameters Derived from MAE120 Low Molecular Weight Nanocomposite TGA Data . . . . . . . . . . . . . . . . . . . . . . . . 101 3.6 Kinetic Parameters Derived from PMMA Grafted Clay Low Molecu- lar Weight Nanocomposite TGA Data . . . . . . . . . . . . . . . . . . 102 3.7 Kinetic Parameters Derived from TGA Data of High Molecular Weight Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 103 3.8 Results from Gasi cation Experiments for Samples with 5wt.% Filler 111 3.9 Results from Cone Calorimetry Experiments . . . . . . . . . . . . . . 114 3.10 Glass Transition Temperatures from DSC . . . . . . . . . . . . . . . . 127 4.1 EDX Analysis of the Residue from Cloisite 30B Nanocomposites Con- taining phosphorus Silane . . . . . . . . . . . . . . . . . . . . . . . . 148 6.1 Commercially Available Calcium Phosphates . . . . . . . . . . . . . . 158 6.2 Commercially Available Inorganic Phosphates . . . . . . . . . . . . . 159 vii List of Figures 1.1 Intercalated Structure . . . . . . . . . . . . . . . . . . . . . . . . . . 7 1.2 Exfoliated Structure . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 1.3 Hydrous Silanization Process . . . . . . . . . . . . . . . . . . . . . . . 11 1.4 Anhydrous Silanization Process . . . . . . . . . . . . . . . . . . . . . 12 1.5 Three Types of Flow During Combustion . . . . . . . . . . . . . . . . 13 1.6 PMMA Degradation Mechanism . . . . . . . . . . . . . . . . . . . . . 15 1.7 PS Degradation Mechanism . . . . . . . . . . . . . . . . . . . . . . . 15 2.1 Structure of Surfactant Used in Cloisite 30B Clay . . . . . . . . . . . 24 2.2 Structure of Surfactant Used in Cloisite 15A Clay . . . . . . . . . . . 25 2.3 Silane A174 Structure . . . . . . . . . . . . . . . . . . . . . . . . . . 26 2.4 Short PMMA Chain Grafted on Clay . . . . . . . . . . . . . . . . . . 27 2.5 Shape Factor . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 37 2.6 Ozawa Method for Calculating Degradation Kinetic Parameters . . . 38 2.7 Schematic of Gasi cation Apparatus . . . . . . . . . . . . . . . . . . 40 3.1 XRD Data from Cloisite 15A Clay with a Peak at 2.27 nm, Cloisite 30B Clay with a Peak at 1.67 nm and Peaks from Somasif MAE120 Synthetic Mica (2.76 nm, 1.96 nm, 1.38 nm and 1.18 nm) . . . . . . 42 3.2 XRD Data from in-situ Polymerized Nanocomposites with Peaks La- beled. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 46 3.3 TEM of in-situ Polymerized Cloisite 15A Nanocomposites . . . . . . 48 3.4 TEM of in-situ Polymerized Cloisite 30B Nanocomposites . . . . . . 48 3.5 TEM of in-situ Polymerized MAE120 Nanocomposites . . . . . . . . 49 3.6 TEM of in-situ Polymerized MAE120 Nanocomposites Showing Clay Stacks Being Spread Apart . . . . . . . . . . . . . . . . . . . . . . . 50 viii 3.7 Gasi cation Data from in-situ Polymerized Nanocomposites . . . . . 51 3.8 Clay Stack Count from the Cloisite 15A/PMMA Nanocomposites . . 54 3.9 Clay Stack Count from the Cloisite 30B/PMMA Nanocomposites . . 55 3.10 Clay Stack Count from the MAE120/PMMA Nanocomposites . . . . 56 3.11 Clay Stack Count from the PMMA Grafted Clay/PMMA Nanocom- posites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57 3.12 Intercalated Structure Produced with Cloisite 15A Clay . . . . . . . . 58 3.13 Exfoliated Structure Produced with Cloisite 30B Clay . . . . . . . . . 59 3.14 Intercalated Structure Produced with MAE120 Clay . . . . . . . . . . 60 3.15 Exfoliated Structure from Nanocomposite with PMMA Grafted Clay 61 3.16 Exfoliated Structure from Nanocomposite with PMMA Grafted Clay 62 3.17 XRD Data from Precipitated Nanocomposites . . . . . . . . . . . . . 64 3.18 Small Angle X-ray Scattering from Nanocomposites . . . . . . . . . . 66 3.19 Fitting of PMMA Grafted Sample with Average Stack Model . . . . . 67 3.20 Fitting of Cloisite 30B Sample with Average Stack Model . . . . . . . 68 3.21 Fitting of PMMA Grafted Sample with Multiple Stack Summation Model . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69 3.22 Fitting of Cloisite 30B Sample with Multiple Stack Summation Model 69 3.23 Fitting of Cloisite 15A Sample with Multiple Stack Summation Model 70 3.24 TGA Data from Various Clay Samples . . . . . . . . . . . . . . . . . 72 3.25 TGA Data from Pure PMMA Samples . . . . . . . . . . . . . . . . . 74 3.26 Derivative of TGA Data from Pure PMMA Samples . . . . . . . . . . 75 3.27 TGA Data from Cloisite 30B/PMMA Low Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 76 3.28 TGA Data from Cloisite 30B/PMMA Low Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 77 ix 3.29 TGA Data from Cloisite 15A/PMMA Low Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78 3.30 TGA Data from Cloisite 15A/PMMA Low Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 79 3.31 TGA Data from MAE120/PMMA Low Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 80 3.32 TGA Data from MAE120/PMMA Low Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 81 3.33 TGA Data from Grafted Clay/PMMA Low Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 82 3.34 TGA Data from Grafted Clay/PMMA Low Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 83 3.35 TGA Data from Pure PMMA High Molecular Weight Samples . . . . 84 3.36 TGA Data from Pure PMMA High Molecular Weight Samples . . . . 85 3.37 TGA Data from Cloisite 30B/PMMA High Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 86 3.38 TGA Data from Cloisite 30B/PMMA High Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 87 3.39 TGA Data from Cloisite 15A/PMMA High Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 88 3.40 TGA Data from Cloisite 15A/PMMA High Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 89 3.41 TGA Data from MAE120/PMMA High Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 90 3.42 TGA Data from MAE120/PMMA High Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 91 3.43 TGA Data from Grafted Clay/PMMA High Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 93 3.44 TGA Data from Grafted Clay/PMMA High Molecular Weight Nanocom- posite Samples . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 94 3.45 Kissinger Analysis of Low Molecular Weight PMMA and Nanocom- posite Degradation Kinetics . . . . . . . . . . . . . . . . . . . . . . . 97 x 3.46 Mass Loss Rate Curve from Gasi cation Experiments . . . . . . . . . 104 3.47 Residue After Gasi cation Experiments . . . . . . . . . . . . . . . . . 105 3.48 E ect of Bubbles on Thin Residue Layer . . . . . . . . . . . . . . . . 107 3.49 E ect of Bubbles on Thick Residue Layer . . . . . . . . . . . . . . . . 108 3.50 E ect of Disk Warping on the Flammability . . . . . . . . . . . . . . 108 3.51 Video images captured during the gasi cation of a high molecular weight sample, showing stacks in thin residue (5s, 10s), swelling (28s), and large cracks formed during swelling (28s, 30s). . . . . . . . . . . . 109 3.52 Video images captured during the gasi cation of a low molecular weight sample, showing cracks in thin residue (5s), island like struc- ture formed because of bubbles pushing residue (15s), swelling (25s), and large cracks formed during swelling (25s, 30s). . . . . . . . . . . . 110 3.53 Heat Release Rate Curve from Cone Calorimetry Experiments . . . . 112 3.54 Residue After Cone Calorimetry Experiment . . . . . . . . . . . . . . 113 3.55 Mass Loss Rate Curve from Gasi cation Experiment of Samples Cast from Chloroform Compared to Samples from THF . . . . . . . . . . . 116 3.56 Intercalated/Microcomposite Structure Formed when Cloisite 15A/PMMA Sample Cast from Chloroform . . . . . . . . . . . . . . . . . . . . . . 117 3.57 Intercalated Structure Produced from MAE120/PMMA Cast from Chloroform. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 118 3.58 Mass Loss Rate Data from Nanocomposites with PMMA Grafted Clay120 3.59 Heat Release Rate from Cone Calorimetry Experiment of Nanocom- posite Samples with PMMA Chain Grafted Clay in Comparison with Nanocomposite Samples with Cloisite 30B Clay and Samples of Pure PMMA Polymer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 121 3.60 XRD Data from the Residues . . . . . . . . . . . . . . . . . . . . . . 122 3.61 SEM Observation of Cross Section of a Residue from Gasi cation Experiment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 124 3.62 Viscosity of Di erent Nanocomposites . . . . . . . . . . . . . . . . . . 126 4.1 Structure of Diethylphosphatoethyltriethoxysilane . . . . . . . . . . . 138 xi 4.2 TEM of phosphorus Containing Cloisite 30B Clay/PMMA Nanocom- posites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 140 4.3 TEM of phosphorus Containing Cloisite 30B Clay/PMMA Nanocom- posites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 141 4.4 Gasi cation of phosphorus Containing Cloisite 30B Clay/PMMA Nanocom- posites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 142 4.5 Cone Data of phosphorus Containing Clay/PMMA Nanocomposites . 143 4.6 Mass Loss During Cone Experiments of phosphorus Containing Clay/PMMA Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 145 4.7 Residue from 10wt% of Cloisite 30B/5wt% phosphorus Containing Silane . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 146 4.8 Residue from 10wt% of Cloisite 30B/5wt% phosphorus Containing Silane . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 147 xii List of Abbreviations AIBN Azobisisobutyronitrile AFM Atomic Force Microscope CEC Cation Exchange Capacity DMA Dynamic Mechanical Analysis DMF Dimethylformamide DMSO Dimethylsulfoxide DSC Di erential Scanning Calorimetry EDX Energy Dispersive X-ray Analysis GSD Gaussian Standard Deviation HRR Heat Release Rate IC Integrated Circuit MLR Mass Loss Rate MMA Methylmathacrylate NMR Nuclear Magnetic Resonance PAN Polyacrylonitrile PC Polycarbonate PCB Printed Circuit Board PE Polyethylene PEO Polyethylene Oxide PHRR Peak Heat Release Rate PMA Polymethylacrylate PMLR Peak Mass Loss Rate PMMA Poly (Methyl Methacrylate) POSS Polyhedral Oligomeric Silsesquioxane PP Polypropylene PS Polystyrene PVC Polyvinyl Chloride SAS Small Angle Scattering SANS Small Angle Neutron Scattering SAXS Small Angle X-ray Scattering SEM Scanning Electron Microscopy SLD Scattering Length Density TEM Transmission Electron Microscopy TGA Thermogravimetic Analysis TSCA Toxic Substances Control Act THF Tetrahydrofuran TPO Thermal Plastic Ole n UL Underwriters Laboratories UV Ultra Violet XRD X-ray Di raction xiii Chapter 1 Introduction 1.1 Polymer Nanocomposites Most polymers, especially those used for structural applications, often contain large amounts of ller. Fillers are used to increase the modulus, increase the thermal or electric conduction, increase wear resistance, increase stability against weathering, UV, increase ame retardancy, increase processability, decrease thermal expansion or just reduce the cost of the material. Common inorganic llers include mica, calcite, clay, glass bers, glass beads and carbon black. Organic llers are commonly termed additives. Polymer nanocomposites are polymers containing nanometer size llers. The llers can be inorganic or organic, but small organic molecules like lubricants, ame retardant additives or dyes are generally not classi ed as llers for nanocomposites. Compared to traditional polymer based composite materials, polymer nanocom- posites have a number of advantages. For mechanical properties, nanocomposites have been shown to achieve the modulus of traditional composites at a lower con- centration, which also results in a lower density of the nal product and potentially lower cost. Due to the large surface area of the llers, there is a signi cant amount of interfacial area introduced into the composites, which in turn could be large enough 1 to a ect the conformation of the polymer molecules. Also due to the size of the ller and depending on the dispersion, it can be hard to classify the nanocomposites as a traditional two phase mixtures [1]. The nanoparticles used in polymer nanocomposites can be either 0-dimensional, 1-dimensional or 2-dimensional. 1.1.1 Nanocomposites with 0-Dimensional Fillers 0-dimensional llers are particles where all three dimensions in the nanometer scale. The overall shape of the ller is a generally equi-axed particle and can be termed 0-dimensional. Di erent 0-dimensional nanoparticles have been used to make nanocomposites for industrial applications, for example, nanometer scale silica gel and polyhedral oligomeric silsesquioxane (POSS). Nanometer scale silica gel has long been used as a ller for polymer based IC packaging materials to control the viscosity of the raw material, to increase the modulus and decrease the cure shrinkage [2]. Surface modi ed silica gel has also been used as ame retardant material in the printed circuit board (PCB) materials like the new Hitachi Chemical RO resin. The amount of silica gel can be as much as 80 wt% in this type of resin. Due to the surface modi cation on the silica gel, the processability is still good [1]. POSS molecules are a group of well studied nanoscale particles that can be used in nanocomposites [3]. A POSS molecule is essentially a cage made of silicon 2 and oxygen with organic groups attached to the silicon atoms via covalent bonding [4]. Di erent organic groups have been used with POSS, even di erent monomers have been successfully attached to the POSS molecules. As a result, it is possible to directly polymerize the POSS molecule on polymer chains [5]. CdSe nanoparticle based 0-dimensional nanocomposites have been developed for their light emitting properties [6]. Zinc oxide nanoparticles have also been used as a UV stabilizing component. [7] Metal, ceramic and salt 0-dimensional nanoparticles have also been used as llers in various applications such as electromagnetic wave shielding, magnetic memory, super capacitors and etc. [8, 9] Direct mixing in solution, melt or in-situ polymerization are the typical meth- ods used to make nanocomposites with 0-dimensional llers. Nanocomposites with 0-dimensional llers can also be synthesized by block copolymer templating and sol-gel processes [10, 11, 12]. 1.1.2 Nanocomposites with 1-Dimensional Fillers 1-dimensional llers have two dimensions at the nanometer scale. The overall shape of the ller is a 1-dimensional wire such as a nano- ber, nano-wire or nano- tube [13, 14, 15]. The most studied systems are carbon nanotubes, which include multi-walled carbon nanotube and single walled carbon nanotubes [16]. The carbon nanotube polymer nanocomposites combine the properties of the matrix polymer and the carbon nanotube. Generally properties of high thermal conduction, graphite grade electrical conductivity are found, while the processabil- 3 ity of the polymer matrix is retained. Changes in the rheology upon addition of nanotubes to the polymer has been reported [17]. The nanotubes can also be lled with di erent salts or metals to alter the magnetic, optical or electrical properties of the nal material [18]. The mechanical properties of carbon nanotube/polymer nanocomposites are strongly a ected by the interface between the carbon nanotube and the polymer. The surface of carbon nanotubes is hydrophobic and as a result, non-polar polymers like polypropylene (PP) trend to be more compatible with the nanotubes [19, 20] while polar polymers like PMMA or nylon are less compatible. Consequently the mechanical properties of nanocomposites which are made with polar polymers can be worse than the pure polymer. To prevent this, a range of surface modi cation methods have been developed in an attempt to make the carbon nanotubes more compatible with polar polymers [18, 21]. Methods to prepare carbon nanotube/polymer nanocomposite include in-situ polymerization [22], solution casting and melt compounding. Nanotubes made with materials other than carbon have also been reported [23, 24, 25]. Clay also exists in a tube form naturally [26] and these nanotubes have potential applications in polymer nanocomposites [27]. 1.1.3 Nanocomposites with 2-Dimensional Filler 2-dimensional llers have only one dimension at the nanometer scale, with the overall shape of the ller being a 2-dimensional disk or platelet. Examples include 4 single layers of clay, graphite [28], and layered double hydroxide [29, 30, 31, 32, 33]. Clay/polymer nanocomposites are the most important class of nanocomposites with 2D llers. There is considerable interest on using modi ed clay mixed with polymers dating back to the 1980s and even earlier [34, 35]. Toyota Central Research Lab- oratory rst synthesized clay/Nylon 6 nanocomposites which exhibited a relatively high modulus and heat distortion temperature [36, 37, 38]. It has also been proved that the reenforce e ect is due to the clay particles added and not the change in the crystalline phase of PA6 [39]. This nanocomposite achieved signi cant commer- cial success. This nanocomposite has been used to make timing belt covers, engine covers and other under-the-hood car components by Toyota and Mitsubishi [36]. In 2004, clay/PP nanocomposites were used in the 2004 Acura TL head rest and Chevrolet Impala body side mount. In 2005 the Hummer H2 SUT had clay/TPO (thermoplastic ole n) nanocomposites as trim, center bridge, sail panels and box rail protectors [40, 41, 42]. 2-Dimensional ller nanocomposites, especially layered silicate/polymer nanocom- posites, are the most important class of nanocomposites. The reasons are that the raw materials (clay and mica) are readily available and relatively cheap, the syn- thesis process is simple and the modulus of nanocomposites is readily increased compared to the polymer and many di erent polymer can be used to synthesize nanocomposites [43, 44, 45, 46, 47, 48, 49, 50, 51, 52, 53, 54, 55, 56, 57, 58, 59, 60]. Other advantages include lowered small molecule permeability which leads to the potential of using it as a food packaging materials, fuel tanks [61, 41, 62, 63] and low dielectric materials [64]. 5 Clay minerals have long been mixed with polymeric materials. Kaolinite has been used to replace asbestos in thermoset polymers. Kaolinite can also improve ow, shrinkage and water resistance of thermoset polymers. In thermoplastic ma- terials, the clay can be used to lower the viscosity, and improve the mechanical properties of polymers such as polyesters and nylons. Clay has also been used to replace carbon black in rubber. Montmorillonite clays are preferred as ller in paint- ing applications as they are easily modi ed by ion exchange to be compatible with di erent paints [65]. Based on the dispersion of the silicate in the polymer matrix, layered sili- cate/polymer nanocomposites can be classi ed with two di erent structures, exfoli- ated and intercalated [43, 45, 66]. An intercalated structure is where polymer enters the spaces between silicate layers, termed the gallery space, but the disruption is not large enough to lose the overall stack structure of the silicate (schematic shown in gure 1.1) [67]. When unmodi ed clay with +2 charged cations is dispersed in water, water enters the gallery space of the clay but the stack structure is preserved due to the strong interaction between neighboring clay platelets via the divalent cations [68]. Exfoliated means that the layered silicate has lost the original stack structure and single layers of the silicate are dispersed in the polymer matrix (schematically shown in gure 1.2). When clay with +1 charged cations is dispersed in water, it usually forms an exfoliated structure, where there is essentially no interaction between two neighboring platelets ( there are only weak interactions via hydrogen bonding between water molecule network) [68]. 6 Figure 1.1: Intercalated Structure Figure 1.2: Exfoliated Structure Based on the orientation of the silicate platelets, the exfoliated structure can be classi ed into di erent sub-structures. The silicate platelets can either be randomly dispersed or aligned. When isotropic properties are required, a totally random exfoliated structure is preferred. However, if enhanced properties are required in particular direction, an oriented exfoliated or intercalated structure may work better. The silicate particles can also potentially touch each other via edge-to-edge or edge-to-center interactions, which allows for more complexity in the structure of the dispersed layered silicate. Structures with these interactions usually appear in silicate/small molecules mixture and are not common in polymer nanocomposites. 7 1.2 Clay and Organoclay Clay has long been used as a ller for polymers, generally to enhance the modulus and wear properties. It has also been used to lower the viscosity. Most clay minerals exist in a layered form, but may also exist as needles (hal- loysite, sepiolite) or as spheres (halloysite). The layered silicates contain silicon- oxygen layers which are composed of tetrahedral blocks and aluminum-oxygen- hydroxy layers which are composed of octahedral shaped blocks [69]. The ratio between tetrahedral layers and octahedral layers can be 1:1 (kaoli- nate), 2:1(mica type clay minerals, including montmorillonite ) or 2:1:1 (chlorite) [70]. Some of the aluminum atoms at octahedral positions can be replaced by +2 iron or magnesium atoms, and some of the silicon atoms at tetrahedral positions can be replaced by aluminum atoms. So the clay platelet being negatively charged. To compensate this negative charge, some cations are attracted to the clay platelets. Lithium, sodium, potassium, ammonium, calcium are the cations commonly found in clay platelets. The clay platelets are stacked together to lower the total surface energy. The stacking leads to at least two di erent interesting properties of clay. The rst one is swelling of clay in water, and the second is the large inner surface area (the total surface of the clay platelets including area inside the clay stack). The oxygen layer on the surface of clay is organophilic, but clay is com- monly described as hydrophilic, which comes from two sources. The rst rises from cations that are attracted to the clay platelets and trapped in the gallery space. 8 These cations attract water molecules and shield the oxygen layer from free water molecules. The second reason is that hydroxyl groups are exposed at the edge of the clay platelets, which can form hydrogen bonding with water. When the cations adsorb water, the gallery space expands. In the case of monovalent cations like Na+ or K+, the gallery space can expand in nitely (even- tually exfoliating), while the case of Ca+2 may lead to a 45-145% volume increase [68]. Divalent Ca+2 behaves di erently from monovalent Na+ is because the diva- lent cation is attracted to platelets on both sides of the gallery space, while the Na+ is only attracted to only one side. The average distance between cations in the case of Na+ clay is around 0.7 to 0.8 nm [71, 72]. The oxygen layer on the surface of the clay is considered to be organophilic [73] or have the ability to form very weak hydrogen bonding with species with hydroxyl groups. Small organic molecules like glycerol or alcohols readily adsorb to the inner surface of clay [74, 75, 76]. Another important property of clay is that the cations are exchangeable. By replacing the inorganic cations with organic cations, the clay can be made more hydrophobic. The inorganic - organic ion exchange reaction is commonly done with a sodium - ammonium cationic surfactant exchange reaction in water solution to produce a clay with an organic surfactant surface. When the clay is exposed to surfactant molecules in water, the surfactant will be \sucked" into the clay. The exchange of the organic surfactant is both enthalpically and entropically favored. By leaving the water and entering the gallery space, the surfactant molecules increase the system entropy by releasing adsorbed water molecules from the surfactant tail, 9 and decrease the system enthalpy through charge interaction and the van de Waals interactions between the surfactant molecule and oxygen layer of the clay [77]. The unit of cation exchange capacity (CEC) is meq/100g, which stands for millil-equivalents per 100 gram of clay, which is the same as 10 5 mole of charge per gram of clay. The CEC of montmorillonite is around 100, that of mica ranges from 20 to 40, while the kaolinite has a value of 10 or lower. The exact number depends on the geographical origin of the clay and the pH value of the clay solution [78]. The organically modi ed clay, also called organoclay, has been widely used as an absorbent of oil or organic [79]. Early attempts to make polymer organoclay composites can be traced back to the 1950?s [80], although those composites were not termed nanocomposites. Clays behaves di erently from salts, acids or bases. For example, clay surface shows strong or moderate acidity, but the clay solution doesn?t show acidity [81]. Such acidity has been attributed to the low ammability of polymer/layered silicate nanocomposites [82]. 1.3 Silane Coupling Agent Silane coupling agents are usually alkoxysilanes, which means the silane con- tains Si-O-CH3 groups or Si-O-CH2-CH3 groups. Alkoxysilanes can loose the methoxy (-O-CH3) or ethoxy (-O-CH2-CH3) groups and form silanol (Si-OH) groups or Si-O- X type covalent bonding, where X is Si or other metals. Silane coupling agents have been used to promote wetting between polymers and inorganic materials like mica, 10 Figure 1.3: Hydrous Silanization Process glass, silica, alumina, copper, tin, talc, steel, etc. The silane coupling agent also can bond to the polymer via covalent, hydrogen bonds or van de Waals interaction [83]. Silane coupling agents can bond to inorganic material via covalent bonds or hydrogen bonding [84]. The covalent bonding between the inorganic material and silane coupling agent is formed by condensation. When there is water involved in the reaction, the silane coupling agents rstly form silanols, then the silanols form covalent bonding with hydroxyl groups on the surface of inorganic materials via dehydration. The process of hydrous silanization is shown in gure 1.3. When there is no water involved in the reaction, silane can form covalent bonding with hydroxyl groups on the surface of inorganic materials via condensation without forming silanol. The process of anhydrous silanization is shown in gure 1.4. In this research silane coupling agent was used to attach short polymer chains to the surface of silicate. Attaching short polymer chains to the surface of clay has 11 Figure 1.4: Anhydrous Silanization Process been achieved by making ammonium salt containing short polymer chains [85] and grafting PDMS on to the surface of clay in solution [86]. The advantage of silane over the ammonium salt containing short polymer chain is that silane containing polymer chains are more stable. The advantage of silane over PDMS approach is that silane can be adopted to work with di erent polymers. 1.4 Flammability of Polymers 1.4.1 Combustion and Pyrolysis of Polymers Three factors are required for any kinds of combustion. They are an oxidizing agent, fuel, and an ignition source. When a high temperature ignition source is applied to a polymer, the polymer begins to decompose. During decomposition, molecules of monomer or oligomers are given o . In most cases these small molecules 12 Figure 1.5: Three Types of Flow During Combustion are a ammable fuel. Then fuel is mixed with oxygen in the air. Generally a speci c fuel concentration is required for combustion. If the fuel/oxygen ratio is either too high or too low, combustion will not occur. At last the fuel molecules in air combust. In this step, energy is released and some of this thermal energy is radiated back to the condensed polymer phase and further facilitates the decomposition. There are also two mass ows and one energy ow that must be maintained to sustain the combustion. The rst mass ow is the fuel ow from the condensed phase to the combustion zone in the gas phase, the second mass ow is the ow of air to provide oxygen, and the energy ow is the radiation from the combustion zone in the gaseous phase to the decomposition zone in the condensed phase. These ows are shown schematically in gure 1.5. The polymer decomposition is an important step in the combustion process. There are two di erent major paths for polymer decomposition without char forma- 13 tion. Polymers of the rst type yield monomer. Polymers falling into this decom- position path include PMMA, polyoxymethylene, poly- -methylstyrene and Te on. This type of polymer degradation is also called \unzipping". The second decom- position path yields monomer, dimer, trimer and etc. Polymers falling into this second decomposition type include polystyrene (PS), polymethylacrylate (PMA), polypropylene (PP), polyethylene (PE). This type of polymer degradation is also termed the \unbutton" [87]. The degradation mechanisms are shown in gure 1.6 for \unzipping" and gure 1.7 for \unbutton". The hydrogen plays an important role in the degradation of polymers. If the hydrogens of polystyrene are replaced by deuterium, the number of monomers in the thermal degradation product increases greatly. If the hydrogens of polystyrene are replaced by methyl group, the degradation product will be only monomer. In the unbuttoning degradation mechanism, the free radical abstracts the hydrogen of other polymer chain or hydrogen on the same polymer chain, and then the free radical cleaves the backbone and creates one double bond. The \unbuttoning" reaction mechanism is also called bite back. By bite back, the polymer may be cut into either monomer, dimer, trimer or etc. For polymers without hydrogens, the free radical can not bite back, thus those polymers generally degrade by unzipping. Unzipping and unbuttoning type polymers have di erent properties during thermal degradation. The unzipping type polymer maintains relatively high molec- ular weight at the beginning of the thermal degradation, properties related with molecular weight like viscosity, modulus also remain relative stable. But the degra- dation product is highly volatile and easily enters the gaseous phase. The molecular 14 Figure 1.6: PMMA Degradation Mechanism Figure 1.7: PS Degradation Mechanism 15 weight of an unbutton type polymer drop dramatically at the beginning of the poly- mer degradation, and consequently properties like viscosity and modulus change dra- matically. However, a large amount of the degradation products are large molecules that remain in the condensed phase and burn slowly. Defects in the polymer chains also play an important role in the degradation of polymers. Head-to-head defects are generally unstable sites which degrade at lower temperature than the random scission temperature of the main chain. There are only a small amount of head-to-head defects in PMMA chains made by free radical polymerization. On the other hand, there are a relatively larger amount of head-to-head defects in PS chains made by free radical polymerization. Unsaturated bonds also have an e ect on polymer stability. The carbons connected with double bonds are in SP3 bonded. The carbons connected to C- C single bonds are in SP2 bonded. In bonds adjacent to a C=C double bond, the single bonds are made by one electron in SP2 status and one electron in SP3 status. As a result, the single bonds neighboring a double bond are weakened and relatively easy to break. There are several properties that are readily quanti ed to describe the com- bustion hazard of polymers [88]. They are: (1) Ignition time, which is a measure of the time required to ignite the polymer. (2) Flame spread speed, which is a measure of the velocity at which the ame travel on the surface of the polymer. (3) Fire penetration, which is a measure of how quickly a controlled re (heat source) can penetrate through a material. 16 (4) Extinction, which is a measure of how easily a burning material can be extinguished or if it exhibits self-extinction. (5) Smoke evolution, which is a quantitative measure of the amount of smoke released by a material during combustion. (6) Toxic gas evolution, which is an important factor in determining the com- bustion hazard of a particular material. 1.4.1.1 So Called "Self-Extinguishing" Polymers The \self-extinguishing" polymer is a claim made by the plastic industry. There are no real "self-extinguishing" polymers, there are many polymers with very low ammability, for example polycarbonate, polyvinyl chloride, some nylons and many thermoset polymers. In 1973, the US Federal Trade Commission led a com- plaint against the use of the term "self-extinguishing" [89, 90]. All these polymers have chemical groups which limit combustion through char formation and speci c chemistry. Polyvinyl chloride (PVC) is in the class of self extinguishing polymers. PVC loses hydrogen chloride from its backbone during the rst step of degradation. The hydrogen chloride scavenges free radicals in both the condensed phase and gaseous phases. The PVC back bone without hydrogen chloride groups is essentially polyene. The polyene wraps around itself and forms char at higher temperature. The sta- bility and mechanical properties of pure PVC is very poor so PVC products on the market generally have large amounts of additives. Many PVC products actually 17 have increased ammability because of the additives. Polycarbonate (PC) has phenyl rings in its backbone, and during thermal degradation the phenyl rings form char. 1.4.1.2 Flame Retardant Additives Di erent methods have been developed to reduce the ammability of poly- mers. The most e ective method developed so far is to add ammability retardant additives. Flammability retardants reduce the ammability of the polymer by either physical or chemical processes. Physical processes include cooling, forming a pro- tective layer or diluting the concentration of ammable small molecules and oxygen. Chemical processes include reaction in the gas or condensed phase [91]. Three major classes of additives have been widely used with polymers. The rst class is halogen based additives. The second class is phosphorus based additives. The third class isAl(OH)3 orMg(OH)2 based additives. Other less prominent ame retardants include antimony oxide, nitrogen compounds based ame retardants and intumenscent ame retardants. Halogen based additives quench free radicals formed during the polymer de- composition process and stop the oxidation of fuel molecules in the gas phase. Halogen based ame retardants are often used together with antimony oxide to decrease ammability. Antimony oxide evaporates during combustion and absorbs energy. One problem is that antimony oxide vapor is poisonous. Addition of halogen based additives is an e ective way to reduce the amma- 18 bility of polymers, but many ame retardants release corrosive or harmful products during combustion, some halogen compounds may be carcinogenic. Some ame retardants may release dioxin or other poisonous byproducts during smoldering. Because of these reasons, some of these compounds have been outlawed in many countries (particularly in Europe). In the United States, halogen based additives are currently being phased out. The only major producer in the United States, Great Lakes Chemical Corporation (now Chemtura Corporation), has stopped manufac- turing most halogen based ame retardants [92, 93, 94, 95, 96, 97, 98]. States like Oregon and California have taken very restrict regulation against the halogen based ame retardants. Phosphorus based ame retardant additives include red phosphorus, phosphate and phosphorus compounds. Phosphorus based ame retardant additives have been successfully used in polymeric materials such as ame retardant polyurethane foam for building construction, polyimide and epoxy [99, 100, 101]. After the banning of halogen based ame retardants, phosphorus based ame retardants have been used in electronic applications like printed wire board and integrated circuit packaging materials [102, 1]. In 1786, King Louis XVIII commissioned Gay-Lussac to search for methods that could protect fabrics in the theaters. Around 1820, Gay-Lussac reported that a mixture of ammonium phosphate and ammonium chloride was the most e ective [65]. These products are still on the market under variant trade names like Fire-Trol, FR CROS, ANTIBLAZE ML. Recently it has been reported that phosphorus based ame retardant addi- 19 tives can leach out from the polymer matrix and cause shorts in copper wiring in integrated circuits [103]. Al(OH)3 or Mg(OH)2 based ame retardants can form a protective layer on the surface of the polymer, which can partially stop the transfer of energy from the gas phase to the condensed phase. During the combustion process, the Al(OH)3 or Mg(OH)2 based ame retardants decompose and release water, which absorbs a large amount of energy. The water vapor also dilutes the oxygen in the air. Without enough energy to break the polymer and enough oxygen, the combustion process cannot be sustained. The degradation products of Mg(OH)2 also suppress smoke. One common drawback of the Al(OH)3 or Mg(OH)2 based ame retardants is that high concentrations (around 65%) are required to e ectively reduce the ammability. As a result, the mechanical properties of polymers with Al(OH)3 or Mg(OH)2 based ame retardants are often signi cantly degraded. A successful ame retardant can not only prevent the ignition, but also prevent the ame spreading and penetration. 1.4.2 Flammability Measurements To measure whether a retardant is e ective, a quantitative standard is needed to measure the ammability of the polymer. There are several measurement systems used in the United States in the build- ing and construction, aerospace, automobile and home electronics industries. One of the most important methods is Underwriters Laboratories UL-94 speci cation, 20 which is used to measure the ammability of materials used in home electronics. UL-94 tests the tendency of polymer to extinguish or the ame to spread the ame after the polymer has been successfully ignited. There are three di erent series of UL-94. The rst series is designed to test the materials used in consumer electronic products, building structural parts or enclosures. The second series are designed to test foam materials. The third series are designed to measure the ammability of thin lm materials. The other area of polymer ammability is the ability to resist ignition, espe- cially electrical-related ignition. There are three kinds of electrical-related ignition, hot wire, high current arc and high voltage arc. UL-764A is a test designed to study the ignition resistance of polymers. 1.5 Flammability of Polymer Nanocomposites Most polymer based nanocomposites show reduced ammability compared to their pure matrix polymers, including 0-dimensional, 1-dimensional and 2-dimensional llers [104, 105, 106, 107, 108, 109, 110, 111, 112, 113]. Because there is no signi cant change in the combustion products of nanocom- posites compared with pure matrix polymer, it is believed that the reduced amma- bility of nanocomposites arise from physical reasons, particularly the shielding e ect of the protective char layer formed on the surface of the residue. Compared with other ammability retardants, nanocomposites reduce the re risk while improving the mechanical properties without adding harmful chemical 21 additives. However, the use of a nano ller alone has not yet produced a nanocom- posite which has low enough ammability to eliminate the need for ame retardant additives. 22 Chapter 2 Experimental 2.1 Research Outline Poly(methyl methacrylate) based polymer nanocomposites were synthesized, and their ammability and structure were examined. This work focused on the e ect of molecular weight, clay distribution and clay aspect ratio on the ammability of the PMMA nanocomposites. 2.2 Materials 2.2.1 Resin Polymer PMMA was selected as the matrix polymer in this study for two reasons. The rst reason is that PMMA decomposes via unzipping, and consequently the molec- ular weight of the polymer and the change of melt viscosity during the combustion process is less dramatic than the polymers that degrade to oligomers. The second reason is that PMMA is a widely used plastic with high transparency, good me- chanical properties and good UV resistance and hence studying the ammability of PMMA based nanocomposites has important industrial and commercial value. 23 Figure 2.1: Structure of Surfactant Used in Cloisite 30B Clay 2.2.2 Additives 2.2.2.1 Cloisite 30B Cloisite 30B organically modi ed clay was purchased from Southern Clay Products Inc. It is an organically modi ed montmorillonite with literature lateral dimensions of 100 nm to more than 1000 nm [114]. The organic modi er is methyl, tallow, bis-2-hydroxyethyl, quaternary ammonium. The tallow contains 65% C18, 30% C16 and 5% C14. The structure of the surfactant is shown in gure 2.1 [115]. Cloisite 30B clay forms a transparent dispersion in dimethylformamide (DMF) and dimethylsulfoxide (DMSO). 2.2.2.2 Cloisite 15A Cloisite 15A was purchased from Southern Clay Products Inc. It is an or- ganically modi ed montmorillonite with literature lateral dimensions of 100 nm to more than 1000 nm[114]. The organic modi er is dimethyl, dehydrogenated tallow, 24 Figure 2.2: Structure of Surfactant Used in Cloisite 15A Clay quaternary ammonium. The tallow contains 65% C18, 30% C16 and 5% C14. The structure of the surfactant of Cloisite 15A is shown in gure 2.2 [116]. Cloisite 15A clay forms a transparent dispersion in styrene monomer and chloroform. 2.2.2.3 Somasif MAE120 The Somasif series of synthetic micas were provided by Unicoop Japan. Flu- orine is used to replace the hydroxyl groups on the mica. In the case of MAE120, an organic modi er was used to replace the sodium cations, making the synthetic mica hydrophobic. The MAE120 synthetic mica disperses well in benzene, xylene, chloroform and dichloromethane [117]. The lateral dimension of platelet was measured by AFM as 1-3 um. 2.2.2.4 Synthesis of PMMA Grafted Layered Silicate Short PMMA chains were grafted to the surface of Cloisite 30B clay by silane chemistry. Methyl methacrylate (MMA) monomer, 3-(Trimethoxysilyl)propyl methacry- 25 Figure 2.3: Silane A174 Structure late (silane A174, gure 2.3), Azobisisobutyronitrile (AIBN), and water were dis- solved in DMF with mole ratio of 10:1:1:1. Cloisite 30B clay and Silane A174 at a weight ratio of 2:1 was also dispersed in DMF with help of a sonicator. Then the two solutions were mixed, stirred, bubbled with nitrogen and sealed. The solution was then kept at 60 C for 3 hours. The mixture was then kept at 100 C in vacuum for 1 hour to remove water and complete the reaction. Silane A174 can polymerize with MMA monomer and form a copolymer of silane and MMA. The silane in the copolymer can bond to hydroxyl groups on the surface of clay via covalent bonding as shown in gure 2.4. 26 Figure 2.4: Short PMMA Chain Grafted on Clay 2.3 Synthesis of Layered Silicate/Polymer Nanocomposite 2.3.1 In-situ Polymerization In-situ polymerization is a method where the layered silicate is mixed with monomer and initiator, then the polymerization reaction is initiated by either heat- ing or UV. PMMA/Layered silicate nanocomposites with high transparency can be made by this method. However, due to the batch nature of the synthesis, the molecular weight repeatability is not very good. In these experiments, methyl methacrylate monomer was mixed with organi- cally modi ed clays by sonicating in a ask for 6+ hours. AIBN initiator was added, and the mixture was then heated and sonicated at 60 C until the viscosity started to increase. The viscous mixture was poured into a glass mold and kept in a nitrogen environment at 60 C for 24 hours to cure. The nanocomposite sheets were then 27 milled into disks with diameter of 75 mm and thickness of 8 mm, a size appropriate for gasi cation and cone calorimetry experiment. 2.3.2 Solution Mixing In this method, layered silicate was rst dispersed in THF. Then the polymer was added to this solvent mixture. After mixing, the nanocomposite was precip- itated in water. In general, the solvent must be a common solvent for both the organic clay and the polymer, and the precipitation agent must be a non-solvent for both the organic clay and the polymer. A list of samples is presented in table 2.1. 2.3.3 Synthesis of Nanocomposites with PMMA Grafted Clay Nanocomposites with PMMA grafted clay were also synthesized by solution mixing. 2.4 Structural Characterization of Nanocomposites Three common methods for structural characterization of nanocomposites are: direct imaging, including TEM (transmission electron microscopy), SEM (scanning electron microscopy), optical microscopy and AFM (atomic force microscopy); scat- tering, including XRD (X-ray di raction) and SAS (small angle scattering), and spectroscopy including rheology, dielectric analysis and NMR (nuclear magnetic resonance). 28 Table 2.1: Solution Precipitated Samples Sample Number Matrix Mw (kDa) Clay Type Clay Concentration (wt%) Solvent 1 25 Cloisite 15A 5 THF 2 350 Cloisite 15A 5 THF 3 25 Cloisite 30B 5 THF 4 350 Cloisite 30B 5 THF 5 25 MAE 120 5 THF 6 350 MAE120 5 THF 7 350 Cloisite 30B 10 THF 8 350 Cloisite 15A 5 Chloroform 9 350 MAE 120 5 Chloroform 29 In general, there are three important factors that need to be identi ed in the structural characterization of nanocomposites. These are the degree of dispersion, the degree of orientation and the distance between clay platelets. The particle distribution at the micrometer scale can be measured using trans- mission optical microscopy and TEM at low magni cation. The degree of dispersion can be quanti ed by measuring the percentage of single platelets, double platelets, triple platelets etc. This can be achieved via high magni cation TEM, but di cul- ties include sample preparation and the low contrast of the platelets in the polymer matrix. The percentage of clay particles containing di erent numbers of platelets can be described by an average number of platelets per stack. The overall dispersion of clay particles can be described by a number average and a weight average distribu- tion of clay particles. Distance between clay particles can be measured by TEM. XRD is can also be used when the distance between clay particles is relative small, while SAS is applicable when the distance between clay particles is on the scale of nanometers. 2.4.1 X-ray Di raction X-ray di raction (XRD) is a useful method to distinguish an intercalated struc- ture from a traditional microcomposite structure. If the clay platelets are in stacks, the XRD spectrum exhibits di raction peaks indicating the distance between the clay platelets. 30 The di raction peak can disappear due to either the formation of a fully ex- foliated structure or signi cant deformation of the clay platelet stacks [109]. The XRD experiments were conducted using a four-circle Bruker D8 Discover XRD system equipped with a single Gobel mirror and a scintillation counter at NIST. 2.4.2 Small Angle Scattering Small angle scattering is a powerful tool for measuring the nanometer scale structure. Scattering from disk has been developed by Kratky and Porod in 1948 as Is(q) = 2 2 Z =2 0 fsin[qHcos( )]qHcos( ) g2f2J1[qRsin( )]qRsin( ) g2sin( )d (2.1) Where J1 is a rst order Bessel function. is the scattering length density di erence between the disk and the matrix. is the volume of the disk. is the angle between the clay and ~q. The structure factor of a stack clay platelets can be expressed as S(q;n) = 1 + 2n nX k=1 (N k) cos(kDqcos?) exp[ k(qcos?)2 2D=2] (2.2) where ? is the angle between q and the axis of the tactoid (ie. stack of clay platelets), n corresponds to the total number of platelets in the stack, and D and D represent the next neighbor center-to-center distance and the Gaussian standard deviation (GSD), respectively. D and D can be deduced from the SAXS data plot. 31 It is safe to assume that the surfactant region and pure polymer region have similar electron densities which results in no contrast from the surfactant. Then the scattering intensity from a tactoid of clay with N clay platelets can be expressed as I(q) = ( )2 VP(q)S(q;N) (2.3) where the is the di erence between the SLD of clay and PMMA. However, the X-ray SLD of organically modi ed clay is a ected by the organic modi er and clay, because the nominal composition and density of organoclay are not known, the X-ray SLD of organoclay cannot be calculated easily. In the nanocomposites, the clay exists in stacks of di erent thickness as well as single layers. The interaction between all clay particles can be described by the correlation function of the clay particles. Two simple methods to solve this problem have been proposed. The rst one is to deduce the average clay stack thickness from the linear region of the log-log plot. The second one is to add the intensity from stacks with di erent thicknesses together. I(q) = X N ( )2 NVP(q)S(q;N) (2.4) where = X N N (2.5) Both methods have their drawbacks. For the average thickness method the meaning of average value is unclear, it is not clear whether it represents an arithmetic mean, weighted mean, median or mode of the clay distribution. Also bimodal distributions cannot be described by single average. For the multiple stack thickness summation 32 method, di erent distributions of stack thickness may t one set of scattering data. Thus, it is useful to couple the scattering data with TEM to determine distribution of clay stacks. However, because of the limited observation area in TEM, it is di cult to obtain enough data. Also it is assumed that the distance between the clay stacks is large and the interaction between di erent clay stacks is negligible. A third model has been recently proposed by Hermes [118]. In this model, clay stacks were treated as single particles. However, this model does not appear to provide any improvement in the analysis of the scattering data than the rst two methods, but does increase the computation di culty. 2.4.3 Transmission Electronic Microscopy TEM can provide direct image information of the clay platelet dispersion. TEM is particularly useful in identifying exfoliated structures [119]. Bright eld microscopy is the most widely used method to observe the nanocom- posite structure. There are several di culties in measuring clay distribution by TEM. The rst one is sample preparation. For this work, ultramicrotoming was used to prepare thin sections for TEM. During the microtoming process, the sections can be compressed and the e ect on the original structure is unknown. The second di culty is the observation of single clay platelets in the polymer matrix due to intrinsically low contrast. TEM bene ts from being coupled with XRD to characterize dispersion of the clay. 33 Transmission electron microscopy samples were prepared using a Leica UC-6 ultramicrotome with a Diatome diamond knife at room temperature. The thickness of samples was estimated to be less than 100 nm. Samples were examined with a JEOL JEM-2100 TEM at 200kV coupled with a Gatan CCD camera. 2.5 Flammability Measurement The ammability of polymer products can be measured at di erent mass scales, i.e. milligrams, grams or kilograms. Thermogravimetric analysis (TGA) experiments can estimate the thermal degradation at the milligram level, while cone calorimetry and gasi cation measures the ammability for sample size of tens of grams. Unfortunately it is not always clear that these experiments correlate well with the ammability in real world cases, such as a fully furnished house or building, due to the complex nature of real world combustion situations. 2.5.1 Thermogravimetric Analysis TGA measures the change of mass as a function of time or temperature. TGA uses a sample size on the order of tens of milligrams. Because of the small sample volume, the temperature gradient and degradation product di usion can generally be ignored. TGA can be used to measure the temperatures at which the degradation starts and ends. TGA data can also be used to calculate the degradation kinetics. However, the TGA experiments must be carried out carefully for samples that produce protective residue layers like nanocomposites. The protective layers can 34 in uence in the transport process and a ect the measurements [120]. TGA experiments were conducted using a TA Q-500 equipment at di erent heating speeds from room temperature to 450 C. 2.5.1.1 Degradation Kinetics The reaction kinetics of homogeneous system is described as d[X] dt = Aexp ERT f([X]) (2.6) where E is the activation energy, T is temperature, R is the universal gas constant, A is the pre-exponential factor, [X] is concentration of the reactant, and t is time. However, equation 2.6 cannot be directly used on polymeric materials. A typical polymer sample is considered to be a heterogeneous system because of the di erence in the distribution of the molecular weight and other parameters. Polymer nanocomposites and composites should also be considered a heterogeneous system. The reaction kinetics of heterogeneous system is described below. The amount of polymer converted into monomer (or degree of conversion) is expressed as = ms mm s (2.7) where is the degree of the conversion, ms is the initial mass, and m is the current mass. The reaction kinetics are described by d dt = Aexp E=RT f( )h( ;T) (2.8) 35 h( ;T) is normally considered to be equal to 1 in most cases for ease of discussion. f( ) can take many di erent forms, one commonly presumed form is f( ) = (1 )n (2.9) where n is the reaction order, then the equation can be expressed as d dt = Aexp E=RT(1 )n (2.10) The relationship between temperature and time can be expressed as T = t +T0 (2.11) where is the heating rate, and T0 is the initial temperature. There are several methods to analyze the thermal degradation kinetic parame- ters. The Kissinger equation is one of the most widely used methods. The Kissinger equation is expressed as ln( =T2m) = ln(nRAWn 1m =E) E=RTm (2.12) where Tm is the temperature at the maximum rate of mass loss expressed in K, R is the universal gas constant, Wm is the weight of sample at the maximum rate of weight loss, and n is the apparent order of the reaction. The reaction order can be derived from the shape factor or shape index s. The shape factor, as illustrated in gure 2.5, is expressed as s = a=b (2.13) where b is the \down hill" slope and a is the \up hill" slope. The relationship between reaction order and shape factor is n = 1:26ps (2.14) 36 Figure 2.5: Shape Factor Using the Kissinger method, the activation energy can be derived by plotting ln( =T2m) vs. Tm at di erent heating rates. The reaction order n can be derived from the shape factor, the pre-exponential factor can be derived from the intersection of ln( =T2m) vs. Tm. A second method that can be used to derive the degradation kinetics is the Ozawa method. At the same degree of conversion, the temperature, heating rate and activation energy have a relationship described as d(log i) d(1=Ti) 0:4567(E=R) (2.15) where i is the heating speed, Ti is the temperatures at the selected conversion rate at di erent heating rates, E is the activation energy, and R is the universal gas constant. The activation energy can be calculated by plot log i vs. 1=Ti. The change of Ti in respect to i is illustrated in gure 2.6. 2.5.2 Cone Calorimetry Cone calorimetry measures the mass loss and heat release during the combus- tion process and uses sample sizes on the order of tens of grams. It is the most 37 Figure 2.6: Ozawa Method for Calculating Degradation Kinetic Parameters widely used bench scale experiment of polymer ammability. Data acquired from cone calorimetry reveals the behavior of polymers exposed to a fully developed ame. In cone calorimetry experiments, the polymer is in a closed environment with a con- stant radiant heat source, which supplies energy from the top and heats the sample continuously during the experiment. After ignition, the loss of mass and the release of heat energy are monitored continuously. The combustion products and oxygen consumption can also be monitored by analyzing the exhaust. The energy ux applied in a cone calorimeter can be controlled to simulate di erent re conditions. A heat ux of 50kW=m2 is reasonable to simulate a full scale re in a room. A list of heat uxes at di erent re situations are given at table 2.2. 38 Table 2.2: Relationship Between the Heat Flux and Fire Scale Flux (kW=m2) Scale 25 Small re 35 Fire before ash over 50 Fire after ash over, aircraft re 75 Con agration 100 Petroleum re A model based on DSC and TGA measurements has been developed to predict the behavior of thin polymer samples, where the thermal conductivity is not an important factor [146, 147]. This model was designed to produce an easier method to predict the ammability of a polymer using (relatively) common TGA and DSC equipment without the requirement of running cone calorimetry experiments. 2.5.3 Gasi cation Gasi cation is an experiment similar to cone calorimetry where the samples are exposed to a radiation heat source in a nitrogen environment. This simulated combustion during a gasi cation experiment does not produce exactly the same data as cone calorimetry, but the overall shape of the mass loss curve will be similar. More importantly, there will be no smoke or ame generated at the surface of sample which allows for direct observation of the sample surface. To study the combustion process of nanocomposite materials which may generate a residue layer, gasi cation 39 Figure 2.7: Schematic of Gasi cation Apparatus experiment can provide information about the formation of this layer. A schematic of the gasi cation equipment built by NIST is shown in gure 2.7 [121]. 2.6 Rheology Measurement Rheology measurements were conducted on a TA RAD III rheometer at 200 C. The frequency sweep tests were conducted in the linear region of the stress-strain curves. 40 Chapter 3 Results and Discussion 3.1 X-ray Di raction from Clays The clay XRD data are shown in gure 3.1. The MAE120 sample showed di raction peaks at 2.76 nm, 1.96 nm, 1.38 nm and 1.18 nm. The XRD data shows that the MAE120 mica has platelets which have a well ordered structure. One explanation is that the mica stacks never became fully exfoliated during the organic modi cation. The Cloisite 15A sample showed a peak at 2.27 nm. The Cloisite 30B showed a broad peak at 1.67 nm. The montmorillonite clay from which the Cloisite 15A and Cloisite 30B were made has been known to have random distribution of crystal orientations except for the [001] direction, so the (001) di raction peak has higher intensity than the other planes [122]. 3.2 In-situ Polymerization Samples 3.2.1 Dispersion of Clay in MMA Monomer Cloisite 15A, Cloisite 30B and MAE120 clay showed di erent behavior when dispersed in MMA monomer. The Cloisite 30B organically modi ed clay forms gel like mixture at 5 wt% of clay dispersed in MMA monomer. The gel like Cloisite 30B dispersion does not ow, 41 Figure 3.1: XRD Data from Cloisite 15A Clay with a Peak at 2.27 nm, Cloisite 30B Clay with a Peak at 1.67 nm and Peaks from Somasif MAE120 Synthetic Mica (2.76 nm, 1.96 nm, 1.38 nm and 1.18 nm) 42 when the container is held upside down. But it does ow when a force is applied. A dispersion of 5 wt% Cloisite 15A organically modi ed clay in MMA is a viscous liquid, and the clay settles out of the dispersion after 24 hours. The dispersion of MAE120 organically modi ed mica is a low viscosity liquid and the mica settles out of the dispersion within minutes after the stirring is stopped. During the PMMA polymerization, the monomer is usually partially poly- merized to a viscous mixture before casting. During the curing of the MMA, the viscosity of the partially polymerized monomer can prevent the Cloisite 15A and MAE120 clay particles from precipitating out. 3.2.2 Polymerization of the Nanocomposites During the polymerization of the nanocomposites, there are several di culties associated with the properties and structures of the nanocomposites. First, because of the interaction between the clay platelets and the monomer, the viscosity of the clay dispersion is much higher than that of the pure monomer. As a result, during the polymerization, the heat dissipation by convection is suppressed and autoacceleration of the polymerization reaction is di cult to control. If autoac- celeration occurs, the temperature increases rapidly. If the temperature exceeds the boiling point of the monomer, a large amount of bubbles will be produced. Second, the clay platelets limit the di usion of monomer in the solidi ed poly- mer. This leads to two di erent events at di erent stages of the polymerization. The rst event is at the end of the free radical polymerization, the viscosity is large 43 and the active chain ends have limited motion. Under these conditions, the poly- mer chains grow by monomers di using to the active chain ends. But the monomer di usion is hindered by the clay platelets and a relatively large amount of monomer remains in the nal sample. The second event happens after the polymerization is nished. Bulk free radical polymerization usually leaves some unreacted monomer which is usually removed by heating the sample. However, the clay platelets prevent the monomer from di using out of the nanocomposite. These two events lead to the same result, a relatively large amount of monomer remains in the in-situ polymer- ized nanocomposite samples, and the concentration is related with the dispersion of clay. Monomers are more ammable than polymers and have a shorter ignition time. Thus the monomer inside the in-situ polymerized samples a ects the ammability experiment. Finally, it is di cult to control molecular weight properly. Due to the batch nature of the synthesis, the molecular weight of sample with the same recipe will be di erent for di erent batches. During the polymerization, di erent parts of the reaction mixture may experience di erent temperatures, resulting in di erent molecular weights. It is also di cult to make low molecular weight samples by this polymerization method. Low molecular weight samples can be produced by using either a very high concentration of initiator or chain transfer agent. Both methods have drawbacks. High initiator concentration increases the reaction speed and likelihood of autoacceleration. A high concentration of chain transfer agent can produce low molecular weight polymer repeatedly but the thiol chain transfer agents have terrible smell which limits their use. Halogen chain transfer agents are ame 44 retardants and obscures the e ect of molecular weight on ammability. 3.2.3 Structure of in-situ Polymerized Samples 3.2.3.1 X-ray Di raction The XRD data of samples made by in-situ polymerization are shown in 3.2. All samples showed very weak di raction peaks. Nanocomposite with Cloisite 30B clay showed peaks at 3.4 nm and 1.7 nm. Nanocomposite with Cloisite 15A clay showed peaks at 3.5 nm and 1.77 nm. Nanocom- posites made with MAE120 synthetic mica showed peaks at 2.1 nm and 1.36 nm. Comparing the data to that from the clay samples, the nanocomposites made by this in-situ polymerization method showed that an intercalated structure had formed. 3.2.3.2 Transmission Electron Microscopy A TEM micrograph of an in-situ polymerized Cloisite 30B nanocomposite is shown in gure 3.4, a TEM micrograph of an in-situ polymerized Cloisite 15A nanocomposite is shown in gure 3.3. Two TEM micrographs of an in-situ poly- merized Somasif MAE120 nanocomposite are shown, one is at relatively high mag- ni cation ( gure 3.5) and one is at relatively low magni cation ( gure 3.6). The low magni cation Somasif MAE120 micrograph shows how a clay stack can be spread apart during the in-situ polymerization. All TEM micrographs show thin stacks of clay platelets with Cloisite 30B nanocomposites showing more single layers. It can be concluded that Cloisite 30B, 45 Figure 3.2: XRD Data from in-situ Polymerized Nanocomposites with Peaks La- beled. 46 Cloisite 15A and Somasif MAE120 nanocomposites have an intercalated structure with Cloisite 30B nanocomposite showing the largest degree of exfoliation. The XRD data also shows that an intercalated structure exists for all three types of nanocomposites synthesized by this in-situ polymerization method. The peak position of both Cloisite 30B nanocomposites and Cloisite 15A samples is lower than the peak positions in their corresponding clay, which indicates that the distance between the clay platelets has increased. In-situ polymerization is useful in producing nanocomposites with a relatively high degree of exfoliation. Before the polymerization is initiated, monomer and initiator can enter the gallery space between the silicate platelets. After the poly- merization is initiated, the polymer chains grow in this region between the silicate platelets and expands the gallery space. 3.2.4 Flammability measurement of in-situ Polymerized Samples The gasi cation data from in-situ polymerized samples are shown in gure 3.7. The nanocomposites showed a lower mass loss rate than the pure polymer samples. The high molecular weight pure PMMA sample showed lower ammability than the low molecular weight pure PMMA sample. However, the high molecular weight nanocomposite samples showed similar ammability behavior to the low molecular weight nanocomposite samples, which can be explained by the presence of residual monomer in the nanocomposite. The Somasif MAE120 nanocomposite 47 Figure 3.3: TEM of in-situ Polymerized Cloisite 15A Nanocomposites Figure 3.4: TEM of in-situ Polymerized Cloisite 30B Nanocomposites 48 Figure 3.5: TEM of in-situ Polymerized MAE120 Nanocomposites 49 Figure 3.6: TEM of in-situ Polymerized MAE120 Nanocomposites Showing Clay Stacks Being Spread Apart 50 Figure 3.7: Gasi cation Data from in-situ Polymerized Nanocomposites 51 sample showed lowest mass loss rate in this group of data, despite its relatively low molecular weight. This may due to the e ect of the large platelet size of the Somasif mica. 3.2.5 Summary of in-situ Polymerized Nanocomposites Although the in-situ polymerization method makes nanocomposites with clay particles well dispersed in the polymer matrix, there were several problems which are di cult to solve. The di culties are repeatability, monomer residue and molecular weight control. The in-situ polymerization method was dropped in favor of solution mixing to produce samples. 3.3 Solution Mixing Nanocomposite Structure 3.3.1 Transmission Electron Microscopy Micrographs of nanocomposites with 5 wt% Cloisite 15A, Cloisite 30B and MAE 120 clay produced by solution mixing were taken using a JEOL JEM-2100 TEM. Two di erent molecular weights were used, a high molecular weight with a nominal Mw of 350kDa and a low molecular weight with a nominal Mw of 25kDa. Charging and electron beam degradation were a problem for the low molecular weight nanocomposites and the TEM samples were unstable and moved during observation, resulting micrographs of poor quality. Only micrographs from the high molecular weight samples (sample numbers 2, 4, and 6) are shown here. Nonetheless, the low molecular weight nanocomposites showed a similar clay dispersion as the 52 high molecular weight nanocomposites. An example TEM micrograph of a 5 wt% Cloisite 15A/PMMA nanocomposite (350k g/mol molecular weight) is shown in gure 3.12. The Cloisite 15A primarily exists as 10-20 nm thick stacks. The distance between layers in the stacks was measured to be 2.3-2.5 nm. In contrast to the 30B samples ( gure 3.13), only few double layer stacks were observed. Also in contrast to the 30B samples, the silicate stacks in the Cloisite 15A samples are generally thicker and less warped. A clay stack count is given in g 3.8 An example TEM micrograph of nanocomposite of a 5 wt% Cloisite 30B/PMMA (350k g/mol molecular weight) is shown in micrograph 3.13. The Cloisite 30B clay shows a mixed morphology of thin stacks containing less than 5 silicate layers along with thick stacks with more than 5 layers. Double layers were widely present in the image(s). The silicate stacks are also highly warped. Compared with Cloisite 15A and MAE120, Cloisite 30B has a relatively high a nity to the PMMA matrix because of its more polar surfactant, which leads to a higher level of clay exfoliation. The higher level of intercalation (and some exfoliation) in the 30B samples results in more deformation of the clay platelets during molding of the samples. The distance between clay layers measured from TEM ranges from 2.3-2.5 nm for the Cloisite 30B while the stack thickness varied from a single layer to 15 nm (8 layers). A clay stack count is given in gure 3.9. An example TEM micrograph of a nanocomposite of 5 wt% MAE120/PMMA (350k g/mol molecular weight) is shown in micrograph 3.14. The stack thickness is 53 Figure 3.8: Clay Stack Count from the Cloisite 15A/PMMA Nanocomposites 54 Figure 3.9: Clay Stack Count from the Cloisite 30B/PMMA Nanocomposites 55 Figure 3.10: Clay Stack Count from the MAE120/PMMA Nanocomposites several hundred nanometers and the PMMA/Somasif MAE120 samples should be classi ed as microcomposites rather than nanocomposites. The micrograph shows two di erent morphologies. One morphology is consistent of stacks with a platelet layer spacing of 2.6 nm and a second one with a spacing of 1.5 nm. The second morphology is consistent with MAE120 platelets without organic modi cation. A clay stack count is given in gure 3.10 An example TEM micrograph of nanocomposite of a 5 wt% PMMA grafted 56 Figure 3.11: Clay Stack Count from the PMMA Grafted Clay/PMMA Nanocom- posites Cloisite 30B/PMMA (350k g/mol molecular weight) is shown in micrographs 3.15 and 3.16 at low and high magni cation respectively. The PMMA grafted Cloisite 30B sample showed well distributed platelets in the PMMA matrix. A large number of single and double layers can be seen in the image. A clay stack count is given in table 3.11 57 Figure 3.12: Intercalated Structure Produced with Cloisite 15A Clay 58 Figure 3.13: Exfoliated Structure Produced with Cloisite 30B Clay 59 Figure 3.14: Intercalated Structure Produced with MAE120 Clay 60 Figure 3.15: Exfoliated Structure from Nanocomposite with PMMA Grafted Clay 61 Figure 3.16: Exfoliated Structure from Nanocomposite with PMMA Grafted Clay 62 3.3.2 X-ray Di raction 3.3.2.1 X-ray Di raction from Nanocomposites The XRD data from solution mixing 5 wt% Cloisite 15A, Cloisite 30B and MAE120 nanocomposites are shown in gure 3.17. The PMMA/MAE120 sample showed peaks at 3.60 nm, and 1.92 nm. These peaks can be indexed as (001), and (002) from the stacks of clay platelets. The primary gallery spacing of PMMA/MAE120 sample is approximately 3.6 nm. The Cloisite 15A and Cloisite 30B samples showed broad peaks at 3.43 nm, and 1.77 nm, which correspond to (001), and (002). PMMA grafted sample showed a shoulder at 3.43 nm. Cloisite 15A, Cloisite 30B and PMMA grafted samples had broader and weaker di raction peaks than MAE120 (with the PMMA grafted sample being being broadest). The XRD patterns indicate that there are signi cant stacks of platelets present in the samples, and the increase in the gallery spacing indicates that the nanocomposite has an intercalated structure. The degree of order of the layered structure is MAE120, Cloisite 15A, Cloisite 30B and short PMMA chain grafted samples from high to low. Comparing the XRD data from the pure clay to that from the nanocomposites, it can be concluded that some intercalation has occurred in all samples. An analysis of crystal size in the direction of X-ray beam was conducted based on Scherrer equation t = 0:89 =Bcos (3.1) Where t is the thickness of crystal, is the wave length of X-ray, B is the half-peak 63 Figure 3.17: XRD Data from Precipitated Nanocomposites 64 Table 3.1: Silicate Stack Thickness Based on Scherrer Equation Sample 2 ( ) B Thickness (nm) # of layers MAE120 2.4 0.1134 70.1 More than 20 Cloisite 15A 2.6 0.2456 32.4 10 Cloisite 30B 5 0.75 10.6 3 width and is the center of the peak. The result of analysis is listed in table 3.1. It should be noted that the peak in the XRD data of PMMA grafted sample is too small and half peak width is hard to measure. So no analysis was done on PMMA grafted sample. 3.3.3 Small Angle X-ray Scattering SAXS data for the nanocomposites are shown in gure 3.18 as a log-log plot. Over the q range from 0.02 to 0.05, the PMMA grafted nanocomposite sample showed a slope of -2.25, the Cloisite 30B nanocomposite sample showed a slope of -2.42, the Cloisite 15A sample showed a slope of -2.73 and the MAE120 nanocom- posite showed a slope of -3.21. The slope of scattering data of the PMMA grafted sample is similar to the slope of exfoliated clay in water [123, 124, 125], where 2 dimensional disk shaped single sheets are expected. The slope of the SAXS data from the MAE120 nanocomposite is consistent with a 3 dimensional object. The slope of the SAXS data from the Cloisite 30B nanocomposite was consistent with a 2 dimensional structure and the 65 Figure 3.18: Small Angle X-ray Scattering from Nanocomposites Cloisite 15A nanocomposite showed a structure closer to a 3 dimensional object. Using the average stack method (equation 2.3), the average number of platelets in a stack is calculated as 1.1, 1.5, 3 and 6 for the PMMA grafted, Cloisite 30B, Cloisite 15A, Somasif MAE120 nanocomposite samples respectively. It should be noted that the model did not t the 15A or Somasif MAE120 samples well as those nanocomposites have a structure that is not captured by the model. The ts using the average stack method is shown in gure 3.19 (the PMMA grafted sample) and gure 3.20 (the Cloisite 30B sample). 66 Figure 3.19: Fitting of PMMA Grafted Sample with Average Stack Model Fits based on the multiple stack summation model (equation 2.4) are shown in gure 3.21 (PMMA grafted sample), 3.22 (Cloisite 30B sample) and 3.23 (Cloisite 15A). The multiple stack summation model can t the nanocomposites which have a high degree of exfoliation, such as the PMMA grafted samples and the Cloisite 30B samples. For the nanocomposites with a mostly intercalated structure such as the Cloisite 15A, the model can show the trends but the ts are not very good. The multiple stack summation model fails to t the MAE120 sample as the model can not capture the microcomposite structure. Based on the multiple stack summation model, the PMMA grafted sample had about half of the clay platelets as single layers and the average stack thickness was 2.2 layers. The Cloisite 30B sample showed 40% of clay platelets in single layers and the average stack thickness is close to 2.6 layers. The Cloisite 15A had essentially 67 Figure 3.20: Fitting of Cloisite 30B Sample with Average Stack Model no single layers and the MAE120 sample had multiple thick stacks. 3.4 Discussion on the Dispersion of Layered Silicate in Polymer By comparing the XRD, TEM and SAXS measurements, it can be found that di erent techniques have their advantages and disadvantages. TEM observation provides a direct image where it is easy to observe the exfoliated structure, but due to possible compression during the sample preparation process (ultramicrotoming), the measurement of the layer distance in the intercalated structure may be less accurate than XRD. The XRD gives an accurate distance between the clay layers and clay stack thickness, but the amount of isolated clay platelets can not be easily deduced from XRD. SAXS measurements can be used to study both the shape and 68 Figure 3.21: Fitting of PMMA Grafted Sample with Multiple Stack Summation Model Figure 3.22: Fitting of Cloisite 30B Sample with Multiple Stack Summation Model 69 Figure 3.23: Fitting of Cloisite 15A Sample with Multiple Stack Summation Model thickness of clay stacks. This can be used to deduce the degree of exfoliation of clay particles in the nanocomposite. In summary, SAXS probably has the largest advantage, however, SAXS is not widely used in industry. In reference to ammability applications, the degree of clay particle dispersion in the polymer matrix should be measured at di erent length scales: the microscale and the nanoscale. The dispersion of clay particles at the microscale can be measured by optical microscopy and can be described by the average dimension (Daverage) of the particle. This factor can be quantitated by the following equation Fmicro = 1 DaverageD allowed (3.2) In the experiments reported in this thesis, all samples with visible particles by optical microscopy have been rejected, which means Fmicro = 1 (3.3) 70 The nanoscale dispersion can be expressed in similar manner, by using the clay stack thickness (Taverage) Anano = 1 Taverage 1T allowed (3.4) Nanocomposites with a stack thicker than 51 nm will give a Anano less than 0. The fully exfoliated nanocomposite will give a Anano of 1. The over all dispersion can be expressed as Atotal = Anano +Amicro 1 (3.5) If the sample has Atotal less than 0 then it will be rejected as not suitable for low ammability applications. In our experiments, the PMMA grafted, Cloisite 30B and Cloisite 15A samples gives numbers Atotal 1 to 0.5 and Somasif MAE120 samples gives Atotal less than 0. 3.5 Flammability Tests 3.5.1 Thermogravimetric Analysis of Nanocomposites Made by the Solution Mixing Method 3.5.1.1 Clay Samples The TGA data acquired from the clays are shown in gure 3.24. The Cloisite 30B, Cloisite 15A and MAE120 samples showed signi cantly dif- ferent degradation behavior. 71 Figure 3.24: TGA Data from Various Clay Samples 72 Cloisite 30B lost about 20% of its mass at 450 C. Cloisite 15A lost about 25% of its mass at 450 C. The degradation of the clay contributes about 1 wt% mass loss for a nanocomposite with 5 wt% clay. MAE120 lost 40% of its mass at 450 C. The clay degradation will contribute about 2 wt% of the mass loss for a nanocomposite with 5 wt% of MAE120 clay. The thermal degradation process of Cloisite 30B organically modi ed clay shows two steps, the rst step is at 257 C, the second step is at 342 C. The degradation of Cloisite 15A can be separated into 3 peaks, the rst one is at 228 C, the second one is at 290 C and the third one is at 341 C. MAE120 shows 4 steps in the degradation process. The rst MAE120 degradation step is at 253 C, the second step is at 298 C, the third step is at 334 C and the fourth step is at 410 C. 3.5.1.2 Nanocomposite Samples Thermal degradation of PMMA has been well studied [126, 127, 128, 129, 130, 131, 132, 133]. The degradation of PMMA has three steps. The rst step is at 150 C to 200 C. The degradation at this step is due to head-to-head defects in the polymer backbone. A second step is at 250 C to 320 C. This step is attributed to unsaturated bonds. The third step is at 300 C to 400 C. This step is the main PMMA chain undergoing random scission. 73 Figure 3.25: TGA Data from Pure PMMA Samples 74 Figure 3.26: Derivative of TGA Data from Pure PMMA Samples 75 Figure 3.27: TGA Data from Cloisite 30B/PMMA Low Molecular Weight Nanocom- posite Samples 76 Figure 3.28: TGA Data from Cloisite 30B/PMMA Low Molecular Weight Nanocom- posite Samples 77 Figure 3.29: TGA Data from Cloisite 15A/PMMA Low Molecular Weight Nanocom- posite Samples 78 Figure 3.30: TGA Data from Cloisite 15A/PMMA Low Molecular Weight Nanocom- posite Samples 79 Figure 3.31: TGA Data from MAE120/PMMA Low Molecular Weight Nanocom- posite Samples 80 Figure 3.32: TGA Data from MAE120/PMMA Low Molecular Weight Nanocom- posite Samples 81 Figure 3.33: TGA Data from Grafted Clay/PMMA Low Molecular Weight Nanocomposite Samples 82 Figure 3.34: TGA Data from Grafted Clay/PMMA Low Molecular Weight Nanocomposite Samples 83 Figure 3.35: TGA Data from Pure PMMA High Molecular Weight Samples 84 Figure 3.36: TGA Data from Pure PMMA High Molecular Weight Samples 85 Figure 3.37: TGA Data from Cloisite 30B/PMMA High Molecular Weight Nanocomposite Samples 86 Figure 3.38: TGA Data from Cloisite 30B/PMMA High Molecular Weight Nanocomposite Samples 87 Figure 3.39: TGA Data from Cloisite 15A/PMMA High Molecular Weight Nanocomposite Samples 88 Figure 3.40: TGA Data from Cloisite 15A/PMMA High Molecular Weight Nanocomposite Samples 89 Figure 3.41: TGA Data from MAE120/PMMA High Molecular Weight Nanocom- posite Samples 90 Figure 3.42: TGA Data from MAE120/PMMA High Molecular Weight Nanocom- posite Samples 91 By comparing the TGA data ( gure 3.41 and 3.42) from the MAE 120 nanocom- posite and TGA data from other nanocomposites with that of the pure polymer, it can be noticed that the MAE 120 samples have distinct thermal degradation charac- teristics. A large mass loss event can be observed in the 250 C to 300 C range . The ratio between the degradation event at 250 C to 300 C range and degradation event at 300 C to 400 C is related to the heating rate. At low heating rates, the degra- dation from 250 C to 300 C is more prominent. The di erent degradation behavior can be explained by the interaction between the clay degradation and the polymer degradation. The TGA data from the MAE 120 synthetic mica ( gure 3.24) showed strong degradation at 250 C, the free radicals produced during the degradation of organically modi ed silicate can attack the polymer chains and lead to a mass loss in the nanocomposites in this temperature range. Polymer degradation due to the surfactant of organically modi ed clay degradation has also been reported in poly- carbonate nanocomposites [134], LLDPE [135] and PA6 nanocomposites [136], and change of thermal degradation behavior of matrix polymer has been reported in PS nanocomposites [137, 138]. The TGA data from nanocomposites with the di erent clays and the low molecular weight matrix polymer are shown in gure 3.25 (Pure PMMA), 3.27 (Cloisite 30B), 3.29 (Cloisite 15A), 3.31 (MAE120) and 3.33 (Grafted Clay). The derivative TGA curves are shown in 3.26 (Pure PMMA), 3.28 (Cloisite 30B), 3.30 (Cloisite 15A), 3.32 (MAE120) and 3.34 (Grafted Clay). 92 Figure 3.43: TGA Data from Grafted Clay/PMMA High Molecular Weight Nanocomposite Samples 93 Figure 3.44: TGA Data from Grafted Clay/PMMA High Molecular Weight Nanocomposite Samples 94 The TGA data from nanocomposites with di erent clays and high molecular weight matrix polymer are shown in gure 3.35 (Pure PMMA), 3.37 (Cloisite 30B), 3.39 (Cloisite 15A), 3.41 (MAE120) and 3.43 (Grafted Clay). The derivative of the TGA curves are shown in 3.36 (Pure PMMA), 3.38 (Cloisite 30B), 3.40 (Cloisite 15A), 3.42 (MAE120) and 3.44 (Grafted Clay). From gures 3.25 and 3.35, it can be observed that the 25k molecular weight PMMA begins to degrade at higher temperatures than the 350k molecular weight PMMA. It is possible that the low molecular weight PMMA is synthesized using a chain transfer agent, thus the chain ends of low molecular weight samples were saturated bonds (due to the chain transfer agents). The high molecular weight PMMA was probably synthesized using simple free radical polymerization which would result in the chain ends being mostly unsaturated bonds. Unsaturated bonds begin to degrade at a lower temperature than saturated bonds. The reaction order and other degradation kinetic parameters calculated by Kissinger method and shape factor methods from low molecular weight samples are listed in table 3.2 (pure low molecular weight PMMA), 3.3 (low molecular weight Cloisite 15A 5 wt% nanocomposite), 3.4 (low molecular weight Cloisite 30B 5 wt% nanocomposite), 3.5 (low molecular weight MAE 120 5 wt% nanocomposite), and 3.6 (low molecular weight PMMA grafted 5 wt% nanocomposite). The activation energies of degradation reaction for di erent samples were also calculated by Ozawa method. Because for of high molecular weight sample, the degradation process in the 250 C - 300 C range and degradation process in the 300 C - 400 C range overlap 95 with each other, the degradation activation energy of high molecular weigh samples were only calculated by the Ozawa method. The calculation results are listed in table 3.7. The data showed that the in the low molecular weight samples, all low nanocom- posites showed higher activation energies than the pure polymer. The high molecular weight samples also showed higher activation energies than the pure polymer, with the exception of the MAE120 sample, which has an activation energy lower than the pure polymer. The reason high molecular weight MAE 120 sample showed a lower activation energy than the pure polymer is that the nanocomposite degradation is controlled by the degradation of the surfactant rather than the matrix polymer. It can also be observed that the reaction order for the nanocomposites de- crease as the heating speed and the mass of the samples increase. The pure polymer samples do not show this behavior over the tested range of heating rates and sample mass. The MAE120 sample, which has a structure between a microcomposite and a nanocomposite also were not a ected by the change in the heating rate and sample mass. This phenomenon can be explained as the nanocomposites forms a residue layer that acted as a small molecule barrier and a heat insulating layer. However TGA is not suitable when such mass and heat transportation barrier exist. Interest- ingly, the formation of such residue/char barrier can be used to identify the quality of the dispersion for the nanocomposites as materials with well dispersed layered silicate show di erent reaction orders compared to poorly dispersed nanocomposites. 96 Figure 3.45: Kissinger Analysis of Low Molecular Weight PMMA and Nanocompos- ite Degradation Kinetics 97 Table 3.2: Kinetic Parameters Derived from Pure Polymer TGA Data Sample Heating Reaction Kissinger A Ozawa Weight Speed Order Activation Pre-exponential Activation (mg) (K/s) Energy (J/mol) Factor (1/s) Energy (J/mol) 2.7953 0.5 0.978 148205 2.43e+11 141555 2.5706 1 0.965 1.7672 2 1.051 2.3850 5 1.098 31.034 2.5 0.972 16.646 5 0.967 16.657 10 1.151 98 Table 3.3: Kinetic Parameters Derived from Cloisite 15A Low Molecular Weight Nanocomposite TGA Data Sample Heating Reaction Kissinger A Ozawa Weight Speed Order Activation Pre-exponential Activation (mg) (K/s) Energy (J/mol) Factor (1/s) Energy (J/mol) 5.4852 0.5 0.957 190823 8.42E+14 191249 2.8123 1 1.000 5.9433 2 0.979 7.4183 5 0.903 17.874 2.5 0.767 18.63 5 0.755 28.807 10 0.652 99 Table 3.4: Kinetic Parameters Derived from Cloisite 30B Low Molecular Weight Nanocomposite TGA Data Sample Heating Reaction Kissinger A Ozawa Weight Speed Order Activation Pre-exponential Activation (mg) (K/s) Energy (J/mol) Factor (1/s) Energy (J/mol) 8.4561 0.5 1.098 210477 4.17E+16 206056 6.9799 1 1.028 8.4127 2 1.051 8.4719 5 0.981 58.94 2.5 0.829 41.018 5 0.786 57.238 10 0.642 100 Table 3.5: Kinetic Parameters Derived from MAE120 Low Molecular Weight Nanocomposite TGA Data Sample Heating Reaction Kissinger A Ozawa Weight Speed Order Activation Pre-exponential Activation (mg) (K/s) Energy (J/mol) Factor (1/s) Energy (J/mol) 4.7121 0.5 0.947 234613 6.674e+18 214270 2.4946 1 0.946 2.1487 2 0.984 1.9742 5 1.025 9.1827 2.5 0.983 24.564 5 1.060 9.9289 10 0.999 101 Table 3.6: Kinetic Parameters Derived from PMMA Grafted Clay Low Molecular Weight Nanocomposite TGA Data Sample Heating Reaction Kissinger A Ozawa Weight Speed Order Activation Pre-exponential Activation (mg) (K/s) Energy (J/mol) Factor (1/s) Energy (J/mol) 4.3296 0.5 0.992 164875 4.61E+12 171669 6.6164 1 0.840 5.6903 2 0.980 8.3840 5 0.862 27.203 2.5 0.726 29.097 5 0.798 26.083 10 0.813 102 Table 3.7: Kinetic Parameters Derived from TGA Data of High Molecular Weight Samples Sample Ozawa Activation Energy (J/mol) Pure PMMA 157159 PMMA C15A 5% 207094 PMMA C30B 5% 196913 PMMA MAE120 5% 129447 Grafted Clay 5% 192725 3.5.2 Gasi cation Experiments Gasi cation results are shown in gure 3.46, and table 3.8 lists the values of ignition time, peak mass loss rate, time to peak mass loss rate and average mass loss rate. Photographs of the sample residue after gasi cation are shown in gure 3.47. The photographs on the left are residues from the low molecular weight samples, and the photographs on the right are from the high molecular weight samples. The pure PMMA samples left no residue after gasi cation or cone calorimetry and are not shown here. The gasi cation data is separated into three groups, pure polymer, low molec- ular weight samples, and high molecular weight samples. The pure polymer showed a continuously increasing mass loss rate until reaching a maxima and then dropped 103 Figure 3.46: Mass Loss Rate Curve from Gasi cation Experiments 104 Figure 3.47: Residue After Gasi cation Experiments 105 quickly to zero. The nanocomposites all showed signi cantly lower mass loss rates compared with the pure polymers. All low molecular weight samples showed higher mass loss rates compared to their corresponding high molecular weight samples, while the mass loss rate of the MAE 120 samples increased after a plateau but the Cloisite 30B and 15A samples dropped slowly after a plateau as shown in gure 3.46. During the gasi cation experiments, decomposition products formed bubbles in the polymer melt and these bubbles rose to the surface of the polymer melt and erupted. The motion of bubbles directly a ects the formation of any protective residue layer at the surface of the degrading sample. At beginning of the experiment, the residue layer is very thin, the eruption of bubbles can break the residue layer as shown in gure 3.48. In the low molecular weight samples, bubbles can keep the residue from forming a continuous lm on the residue even after the residue layer has become thick. Such a case is illustrated in gure 3.49. In the case of high molecular weight samples, because of the relatively high melt viscosity, the disturbance of bubbles was suppressed and generally resulted in the formation of a protective residue layer with better integrity. It is also interesting to note that decomposition gases built up under all sample disks causing the samples to swell up due to the increasing gas pressure. During swelling, the surface area of the sample increased, the protective residue layers often cracked and broke, exposing new polymer melt to the incident radiation, which in turn increased the decomposition rate. The swelling-cracking situation is illustrated in gure 3.50. In general, the high molecular weight samples have a higher storage modulus and as a result are more resistant to swelling. The low molecular weight samples were more prone to swelling, reaching the dramatic 106 Figure 3.48: E ect of Bubbles on Thin Residue Layer situation for the low molecular weight MAE120 sample when the melt erupted catas- trophically due to melt being stretched beyond the elastic limit by the build up of decomposition products. 3.5.3 Cone Calorimetry Cone calorimetry data are shown in gure 3.53. Table 3.9 lists ignition time, peak heat release rate (PHRR), time to peak heat release rate, peak mass loss rate, time to peak mass loss rate. The cone calorimetry behavior showed similar trends as the gasi cation data, but due to the heat from the burning polymer, the degradation 107 Figure 3.49: E ect of Bubbles on Thick Residue Layer Figure 3.50: E ect of Disk Warping on the Flammability 108 Figure 3.51: Video images captured during the gasi cation of a high molecular weight sample, showing stacks in thin residue (5s, 10s), swelling (28s), and large cracks formed during swelling (28s, 30s). 109 Figure 3.52: Video images captured during the gasi cation of a low molecular weight sample, showing cracks in thin residue (5s), island like structure formed because of bubbles pushing residue (15s), swelling (25s), and large cracks formed during swelling (25s, 30s). 110 Table 3.8: Results from Gasi cation Exp erimen ts for Samples with 5wt.% Filler Sample PMMA PMMA Cloisite Cloisite Cloisite Cloisite MAE MAE 30B 30B 15A 15A 120 120 Molecular W eigh t(g/mol) 25K 350k 25k 350k 25k 350k 25k 350k Ignition time (s) 22.5 27.5 21 21.5 30 28 17 22.5 Peak MLR (g=m 2 s ) 40.045 35.747 23.982 17.195 21.267 17.647 26.244 20.326 Time to PMLR (s) 230.5 201.5 206.5 107.5 104 128.5 236 257.5 Av erage MLR (g=m 2 s ) 21.13 19.86 14.26 10.57 11.05 10.95 16.15 13.37 111 Figure 3.53: Heat Release Rate Curve from Cone Calorimetry Experiments 112 Figure 3.54: Residue After Cone Calorimetry Experiment 113 Table 3.9: Results from Cone Calorimetry Exp erimen ts Sample PMMA PMMA Cloisite Cloisite Cloisite Cloisite MAE MAE 30B 30B 15A 15A 120 120 Molecular W eigh t(g/mol) 25K 350k 25k 350k 25k 350k 25k 350k Ignition time (s)* 21 27 24 18 22 18 15 12 PHRR (kW/m2) 1508.8 1111.3 726 559.7 621.2 632.7 1041.7 696 Time to PHRR (s)* 141 144 99 81 78 87 102 102 Peak MLR (g/m2s) 59.06 44.815 29.522 22.916 25.535 26.227 38.993 27.089 Time to PMLR (s)* 141 141 96 96 84 78 99 96 *data were acquired ev ery 3s. 114 process is faster. Image of the residue are shown in gure 3.54. 3.5.4 E ect of Clay Dispersion 3.5.4.1 Solvent Introduced Dispersion Change One possible method to improve the stability of the residue layer is to alter the clay dispersion, although it has been reported that the clay dispersion is a secondary issue a ecting the ammability of nanocomposites [109, 139], this has been explored in further detail. An experiment was designed to examine the e ect of clay dispersion on amma- bility. MAE120 synthetic mica disperses better in chloroform than in tetrahydrofu- ran (THF), so a sample was cast from chloroform and the gasi cation measured to compare with the previous gasi cation data on materials precipitated from THF. The results are shown in 3.55. The TEM of the samples cast from chloroform revealed that the Cloisite 15A sample has a microcomposite structure with relatively poor clay dispersion ( gure 3.56), while the MAE120 sample has a higher degree of intercalation ( gure 3.57). TEM observation of the samples formed by precipitation from THF revealed that the MAE120 sample was more like a microcomposite and Cloisite 15A sample is more like a nanocomposite. The maximum mass loss rate of the MAE120 sample cast from chloroform is lower than the sample from THF, which is consistent with an improvement of the dispersion in the chloroform cast sample. The mass loss rate data for the the 115 Figure 3.55: Mass Loss Rate Curve from Gasi cation Experiment of Samples Cast from Chloroform Compared to Samples from THF 116 Figure 3.56: Intercalated/Microcomposite Structure Formed when Cloisite 15A/PMMA Sample Cast from Chloroform Cloisite 15A sample cast from chloroform is higher than the sample cast from THF, which is also consistent with a decrease in the dispersion. These experiments indicate changing the dispersion of the clay from a micro- composite to an intercalated nanocomposite does a ect the ammability, with the intercalated systems showing improved ammability. 117 Figure 3.57: Intercalated Structure Produced from MAE120/PMMA Cast from Chloroform. 118 3.5.4.2 PMMA Grafted Clay Introduced Dispersion Change The PMMA grafted clay disperses in the PMMA matrix in a more exfoliated manner than the Cloisite 30B. However, the PMMA grafted clay did not produce signi cant improvement in the ammability. The TGA data are shown in gure 3.33 (low molecular weight sample), 3.43 (high molecular weight sample), 3.34 (derivative TGA data for the low molecular weigh sample), 3.44 (derivative TGA data for the high molecular weigh sample). The data shows three peaks during the thermal degradation. The rst peak is in the range from 150 C to 200 C, which is attributed to head to head defects, the second peak is in the range of 250 C to 330 C, which is attributed to unsaturated bonds. The third peak is in the range of 300 C to 400 C, which is attributed to random chain scission. Gasi cation data from the nanocomposite with the PMMA grafted clay is shown in gure 3.58. The mean mass loss rate of the grafted clay nanocomposite is lower than the Cloisite 30B nanocomposite (10.00 g/m2s vs. 10.57 g/m2s in the high molecular weight samples and 12.15 g/sm2 vs. 14.26 g/sm2 in low molecular weight samples, g/m2s is gram per square meter per second), but the peak heat release rate is not reduced. The cone calorimeter data are shown in gure 3.59. The short PMMA chains grafted on the surface of clay are synthesized by free radical polymerization and are expected to have chain ends with double bonds. Although the dispersion showed improved exfoliation, the thermal degradation and 119 Figure 3.58: Mass Loss Rate Data from Nanocomposites with PMMA Grafted Clay ammability was not signi cantly improved, probably due to the presence of the unsaturated chain ends. 3.5.5 Residue Layer XRD data from residue are shown in gure 3.60. The di raction peak from the Cloisite 30B nanocomposite residue and the Cloisite 15A nanocomposite residue 120 Figure 3.59: Heat Release Rate from Cone Calorimetry Experiment of Nanocompos- ite Samples with PMMA Chain Grafted Clay in Comparison with Nanocomposite Samples with Cloisite 30B Clay and Samples of Pure PMMA Polymer 121 Figure 3.60: XRD Data from the Residues 122 is at 1.26 nm and the peak from the MAE120 nanocomposite residue is at 0.98 nm. The unmodi ed Cloisite Na+ clay has a spacing of 0.98 nm and the unmodi ed ME100 mica has a spacing of 0.93 nm. The XRD data shows that there is some char residue trapped inside the gallery space in the layered silicates. A cross section SEM micrograph of the residue after the gasi cation experiment is shown in gure 3.61. The residue shows a porous structure, which is a good structure for heat insulation. The characteristics of the char layer including the structure and uniformity are key factors related to the measured ammability. The e ectiveness of the residue layer as a shield (radiation and/or mass loss) is maximized when the polymer melt surface is completely covered by the residue layer. For this to occur, the residue layer needs su cient silicate to form and remain stable on top of the polymer melt to keep from cracking and disintegrating. If the protective residue layer forms cracks, radiation can penetrate into the polymer matrix and the decomposition products can escape. There are several factors that a ect the stability of the residue layer. First is the viscosity of the polymer resin. The lower the polymer resin viscosity, the more turbulent is the surface of the polymer melt during combustion, which makes it di cult for a uniform residue layer to form. A second factor is the toughness of the residue layer. In most cases, the residue layer is actually composed of porous silicates. By adding a coupling agent, bridges could form between the silicate layers and it may be possible to increase the toughness of the residue layer. The coupling 123 Figure 3.61: SEM Observation of Cross Section of a Residue from Gasi cation Experiment 124 agent could also be a low melting temperature glass or char. 3.6 Rheology Measurements Measurement of the rheology showed that the viscosity of the nanocomposite is higher than the pure polymer resin ( gure 3.62). The increased viscosity can help prevent bubble formation and slow the rising of the bubbles to the surface, which has two e ects, rst, the mass transfer from the polymer melt to the gaseous phase by the bubbles is slowed and second, the polymer melt surface is less turbulent. 3.7 Glass Transition Measurement Increasing the surface or interface area can cause a change in the glass tran- sition temperature. The glass transition temperature can increase or decrease de- pending on speci c interactions at the polymer/clay interface. Layered silicates introduced into the polymer matrix result in a large interface between the silicate and the polymer and can shift the measured glass transition temperature. The shift can provide insight into the nature of the interface and the degree of dispersion. The glass transition temperature is often measured by Di erential Scanning Calorimetry (DSC). 3.7.1 Di erential Scanning Calorimetry The glass transition temperatures of di erent samples measured by DSC are listed in table 3.10. 125 Figure 3.62: Viscosity of Di erent Nanocomposites 126 Table 3.10: Glass Transition Temperatures from DSC Resin Filler Filler Glass Transition Molecular Weight Type Weight Percent Temperature ( C) 25k None 0 113 25k C15A 5 115 25k C30B 5 115 25k MAE120 5 111 350k None 0 118 350k C15A 5 122 350k C30B 5 101 350k MAE120 5 106 350k Grafted 5 113 127 The organically modi ed silicate has little impact on the glass transition of the low molecular weight nanocomposite. The high molecular weight Cloisite 30B nanocomposite sample showed a 17 C lower glass transition temperature than that of the pure polymer. Polymer/layered silicate nanocomposites can be idealized as a polymer lm between two impenetrable walls, and the glass transition temperature can be af- fected by several factors. Those factors include the thickness of the polymer lm and wetability between the polymer lm and wall [140, 141]. An additional factor is hydrogen bonding between the exposed hydroxyl groups on clay and the oxygen atoms on the PMMA [142]. The e ect of lm thickness is related to the wetability between the polymer lm and the walls. If the the polymer lm does not wet the walls, the thinner the lm thickness the lower the glass transition temperature [140]. If the polymer lm has a strong interaction with the wall, the glass transition tem- perature should increase with decreasing lm thickness. If the lm wets the wall but with no strong interactions, the glass transition temperature should not change with lm thickness. Assuming the surface of the layered silicate is covered by the hy- drocarbon tails of surfactant molecules, no hydrogen bonding should form between the silicate surface and the PMMA polymer and the silicate can be assumed to be non-wetting or poor wetting for the PMMA. There are two di erent lm thicknesses in the nanocomposites, one is the polymer lm in the gallery space and a second one is the polymer between the silicate stacks. The thickness of these two lms contributes to the change of the glass transition temperature. The hydroxyl groups on the edge of silicates are exposed and can form hydrogen bonds, which should 128 increase the glass transition temperature. Cloisite 15A and Cloisite 30B clays were both made from Cloisite Na+ clay, so they have the same shape and size distribu- tion and thus should have same amount of hydroxyl groups exposed. The Somasif MAE120 is a modi ed mica with all the hydroxyl groups replaced by uorine, so no hydrogen bonding should form. The glass transition of polymer/layered silicate nanocomposites can be described by: Tg = V1Tg1 +V2Tg2 +VhTgh (3.6) Where V1 is the volume fraction of polymer in the gallery space, Tg1 is the glass transition temperature of polymer in the gallery space, V2 is the volume fraction of polymer between the clay stacks, Tg1 is the glass transition temperature of polymer between the clay stacks, Vh is the volume fraction of polymer forming hydrogen bonding, and Tgh is the glass transition temperature of polymer forming hydrogen bonding. Tg2 can be expressed as [141] Tg(h) = Tbulkg + 2 h 2 (Tinterfaceg Tbulkg ) (3.7) Where h2 is the thickness of the polymer lm between the clay stacks, is the thickness of the interface layer and Tinterfaceg is the glass transition of the interface layer. Tinterfaceg is a ected by the wetability or compatibility between the polymer and the wall. If the polymer is fully compatible with the wall, Tinterfaceg will equal Tbulkg , while if the polymer forms an enthalpic favorable interaction with the wall, Tinterfaceg may be higher than Tbulkg , if the polymer is not compatible with the wall, Tinterfaceg may be lower than Tbulkg . The glass transition temperature of polymer in 129 the gallery space can be calculated in a similar manner, but due to the limited gallery space, the whole lm may be e ecting an interface layer. Also if there polymer and silicates forms an enthalpic favorable interaction, the silicates may become fully exfoliated eliminating the gallery space. In the case of the Cloisite 15A nanocomposite, the Tg1 and Tg2 terms lower the glass transition temperature, while the Tgh item increases the glass transition temperature and dominates the rst two terms. As a result the measured glass transition temperature increases. Because of the better dispersion of the silicates, the average distance between the silicates platelets in the Cloisite 30B nanocom- posites are smaller than in the Cloisite 15A nanocomposites, and the Cloisite 30B nanocomposite has a smaller h2 value and the V2Tg2 item has a larger absolute value and dominates the hydrogen bonding e ect. As a result the glass transition temperature decreased. The PMMA grafted sample has improved polymer-wall compatibility than in the Cloisite 30B nanocomposite, and as a result the observed glass transition temperature increased. For the Somasif MAE120 sample, there is no hydrogen bonding, so the glass transition temperature of the MAE120 sample is lower than the glass transition of the pure polymer. should be a ected by the molecular weight or Rg of polymer and the lower the molecular weight, the smaller [143]. As a result the glass transition temperature for the low molecular weight samples is less a ected by the di erent types of silicates. 130 3.8 Discussion on the Flammability of Nanocomposites 3.8.1 Mechanism of Low Flammability Nanocomposites The experiments presented here show that nanocomposites can exhibit signif- icantly reduced ammability compared to the pure matrix polymer. Several expla- nations have been proposed to explain this observation. The rst explanation is the formation of a residue layer at the upper surface of the nanocomposite during gasi cation, which can shield against radiation entering the polymer matrix and slow the degradation products from entering the gaseous phase and acting as fuel [110, 144]. The second explanation is the clay changes the polymer degradation kinetics by scavenging free radicals due to the iron contained in the clay [145]. A third possibility is that acidic sites created by the surfactant decomposition changes the polymer degradation kinetics [82]. The rst explanation is relative easy to understand and is observed in this work. The second explanation is less likely. The iron ions are buried inside the alumina layer of the clay which is sandwiched between two layers of silicate and there is little chance for the iron ions to leave the silicate. For the iron atoms to behave as a free radical scavengers, the radicals need to penetrate the silicate layer to react with the iron ions. In addition, montmorillonites from di erent geological origins can have Mg+2 rather than iron replacing the aluminum in the octahedral site. Reports on the ammability of montmorillonite based nanocomposites from di erent countries do not vary much [109, 139]. The the nanocomposites with well dispersed Somasif MAE120 synthetic mica (in situ polymerization and chloroform 131 casting) also showed lowered ammability while mica contains little or no iron. The third explanation is that when the surfactant decomposes, a proton is left at the position of the surfactant as an acidic site. The proton can act as a free radical scavenger. If the scenario described in the third explanation occurs, the ammability of the nanocomposites will be correlated with the amount of acidic sites exposed to the polymer matrix and hence correlated with the amount of clay in the nanocomposite and the degree of exfoliation. If this is true then if the amount of clay increases or the degree of exfoliation increases, the number of acidic sites will increase and the ammability would be expected to drop signi cantly. How- ever, nanocomposites with Cloisite 30B and Cloisite 15A showed little di erence in ammability, although there was a di erence in degree of exfoliation. Doubling the amount of clay also showed only a small e ect on the ammability as shown in gure 4.5. Generally the surfactant degrades at lower temperature than the polymer matrix and when the polymer matrix begins to degrade the clay platelets may have already collapsed and aggregated, leaving only the exterior surface of the residue exposed to the underlying polymer. Also, if the clay has an e ect on the polymer degradation reaction kinetics, the ignition time should be a ected by the clay, but no such e ect has been observed in cone calorimetry or gasi cation experiments. The clay can scavenge free radicals and the results are di erent for polymer- ization and thermal degradation. In polymerization, it is assumed that most free radicals are generated at the beginning of the reaction, if a free radical is removed later, there is no new free radical formed to resume the reaction. In thermal degra- dation, free radicals are generated during the full time span of thermal degradation 132 process and even if a free radical is eliminated during thermal degradation, the for- mer site of the free radical is probably relatively unstable and could form a new free radical. Based on this discussion, it is clear that the low ammability of polymer/clay nanocomposites largely arises from physical reasons rather than chemical ones. 3.8.1.1 Structure Integrity Having a polymer melt with a high viscosity and large complex modulus helps to preserve the structural integrity of the polymer components during a re. Structural integrity is important in the real re situations. If the polymer structure collapses during a re, the re can spread more rapidly. If the polymer structure does not collapse, it can help to con ne the re leading to increased safety. The integrity of a polymeric structure is related to properties such as viscosity, modulus, and char forming ability. If the viscosity of the polymeric material is high, the polymer melt does not readily ow or collapse due to gravity. 3.8.2 E ect of Heat Conductivity on the Fire Safety The heat conduction process can be expressed by the empirical equation Q t = A @T @x (3.8) where Q is the amount of heat, t is time, is the heat conductivity, A is the cross section, T is the temperature, and x is the distance in the conducting direction. 133 The heat conductivity was expressed by Debye as = cp sl (3.9) where cp is the speci c heat capacity, is the density, s is the speed of sound in the polymer, and l is the distance between molecules. In the case of polymer/clay nanocomposites the heat conductivity is also a ected by the clay particles, which have a higher heat conductivity than the polymer. The heat conductivity can have a large impact on re safety on both burning through and re spreading. On the burning through aspect the heat conductivity has little impact on the ammability of thin polymer walls, but polymer with high thermal conductivity burns faster[146, 147]. For example, ignition requires the ma- terial reach a critical temperature, and if the heat conductivity is high, the heat can be more readily dissipated. If a re is con ned in an enclosed space, such as inside a TV set, a cabinet or a room, the heat conductivity of the enclosure or wall has an e ect on the ability to contain the re. If the heat conductivity of the enclosure material is high, the material outside of the enclosure may be ignited by the con- ducted heat. If the heat conductivity of enclosure material is low, the environment outside is at less risk. Ideally a structure can be designed that it has a high heat conduction rate at low heat ux and a low heat conduction rate when exposed to a high heat ux. Such structure can consist of a higher heat conduction layer on the surface and a layer beneath the surface layer that can produce bubbles (termed an intumescent material [148]). When the heat ux is low, the surface layer can distribute heat to a 134 larger surface area to prevent ignition. When the heat ux is large, the intumescent layer can produce bubbles and break the continuity of the surface layer. Polymer/layered silicate nanocomposites have a higher thermal conductivity before the ignition. The residue has low heat conductivity because it is mostly loosely packed clay. This type of behavior is consistent with a material with good ame retardant properties, however due to the low silicate concentration in nanocomposites the di erence in thermal conductivity between the nanocomposite and the pure polymer is relatively small. 135 Chapter 4 Phosphorus Containing Nanocomposite 4.1 Introduction To improve ammability it would be bene cial if nanocomposites enhanced char formation during combustion. Char is less ammable compared to most poly- mers. During combustion polymers release ammable volatile small molecules from the thermal degradation process which fuels the combustion. In contrast char com- busts by oxidizing. The oxidization process is a much slower process than the thermal degradation process. Char is usually a porous solid which has a low heat conduction rate and rela- tively high modulus compared to polymer melts. There are two types of char forming polymers, the rst type has side groups that are readily removed during degradation. During the thermal degradation, the side groups are removed from the main chain and leave free radicals. The free radicals then form double bonds on the main chain. The main polymer chain then wraps and forms carbon rings. PVC and polyacrylonitrile (PAN) belong to this group of polymers. The char forming properties of PAN has been used to make carbon bers. The second type of char forming polymer has carbon rings as part of the backbone and during thermal degradation forms char. Polycarbonate belongs to this group of polymers. 136 By introducing a char forming agent, polymers that are not usually char form- ing can be modi ed to form char during thermal degradation. Common char forming agents include halogen, phosphorus and nitrogen compounds. However, if the char forming agent is a small molecule, it can be leached out during the life cycle of the product and lead to decreased ammability resistance and possible environmental issues. One idea would be to attach the char forming agent to the surface of the clay. Since the clay has a large surface area, a relatively large amount of char forming agent can be introduced into the polymer matrix. It is also possible that the char forming agent attached to the platelets would cause the platelets to fuse together by char. This would give rise to residue composed of char and silicate which could be considered as a layer of silicate reinforced char, and should be stronger than a loosely packed silicate residue layer and increase the potential shielding e ect during combustion. Synergetic behavior of a polymer/layered silicate nanocomposite combined with a char forming agent has been reported, but the char forming agent is not chemically attached to the surface of silicates [149]. 4.2 Char Forming Compound Diethylphosphatoethyltriethoxysilane, which has an phosphate containing silane, was selected as the char forming agent in this research. The silane can be grafted onto the surface of clay by silane chemistry. 137 Figure 4.1: Structure of Diethylphosphatoethyltriethoxysilane Diethylphosphatoethyltriethoxysilane has low toxicity and low environmen- tal impact, it has been used as an anti-piling agent for textile. The structure of diethylphosphatoethyltriethoxysilane is shown in gure 4.1. 4.3 Synthesis 4.3.1 Adding phosphorus to Cloisite 30B Clay Cloisite 30B clay, diethylphosphatoethyltriethoxysilane, and DMF were mixed at ratio of 2:1:10 with help of sonicating. The mixture was then heated in an oven at 100 C for 3 hours. Then the solution was kept in vacuum at 100 C for 1 hour to remove volatile molecules produced during silanization and to nish the reaction. The clay dispersion in DMF was then mixed with PMMA solution in THF, the nanocomposite is precipitated in water and dried in vacuum at 60 C. 4.4 Structure Characterization with TEM TEM observation of the nanocomposites with phosphorus modi ed clay sam- ples are shown in gure 4.2 and 4.3 at low and high magni cations to examine the 138 dispersion of clay in the PMMA matrix and details of the clay platelets distribution. It can be observed that there are areas with higher mass density around the clay particles, labeled in gure 4.3, which is assumed as to be cross linked diethylphos- phatoethyltriethoxysilane. 4.5 Flammability Test 4.5.1 Gasi cation The gasi cation data of nanocomposite samples with phosphorus added clay are shown in gure 4.4. Compared with nanocomposites with 10wt% Cloisite 30B nanocomposite, the nanocomposite containing 10wt% of phosphorus modi ed clay showed similar ignition time, peak mass loss rate and time to peak mass loss rate. However the total gasi cation time of phosphorus sample was close to 600 seconds while the total gasi cation time of sample with Cloisite 30B clay was around 500 seconds. The mean mass loss rate of phosphorus added sample is 7.2 g/sm2 while the mean mass loss rat of the Cloisite 30B sample is 9.0 g/sm2 compared to pure PMMA polymer at 19.9 g/sm2. Compared with the sample without phosphorus, the phosphorus added nanocomposite reduced the mean mass loss rate by about 20%. 4.5.2 Cone Calorimetry Cone calorimetry data are shown in gure 4.5. Similar to the gasi cation experiment, the phosphorus added nanocomposites 139 Figure 4.2: TEM of phosphorus Containing Cloisite 30B Clay/PMMA Nanocom- posites 140 Figure 4.3: TEM of phosphorus Containing Cloisite 30B Clay/PMMA Nanocom- posites 141 Figure 4.4: Gasi cation of phosphorus Containing Cloisite 30B Clay/PMMA Nanocomposites 142 Figure 4.5: Cone Data of phosphorus Containing Clay/PMMA Nanocomposites 143 showed similar peak heat release rate, but di erent mean heat release rate. The mean heat release rate of phosphorus added nanocomposite is 224.5 kW/m2, the mean heat release rate of sample without phosphorus is 298.2 kW/m2 while the mean heat release rate of the pure polymer is 621.4 kW/m2. A 25% of reduction in the mean heat release rate was achieved by adding phosphorus to the clay. The mass evolvement data during the cone calorimetry experiment are shown in gure 4.6. The phosphorus added nanocomposite left 11wt% of residue after cone experiment and the sample with same amount of clay but without phosphorus left 7wt% of residue. The increased 4wt% of residue is assumed to be char formed due to the phosphorus compound. 4.6 Residue Characterization Residue from the TGA was examined by SEM. The residue from the sample with 10%wt Cloisite 30B and 5%wt phosphorus silane is shown in gures 4.7 and 4.8 at low and high magni cation respectively. The original sample was a 0.25 mm thick nanocomposite disk. Figure 4.7 showed bubbles in the residue with diameter as large as 0.8 mm. 4.8 shows the detailed structure of a bubble wall. The wall of the bubble consists of loosely compacted clay stacks. The space between the clay stacks is empty. The additional char formed due to phosphorus is probably attached to the surface or trapped inside the clay stacks. EDX analysis was also conducted on the residue with results shown in table 4.1. Only a small amount of phosphorus is observed in the residue. 144 Figure 4.6: Mass Loss During Cone Experiments of phosphorus Containing Clay/PMMA Nanocomposites 145 Figure 4.7: Residue from 10wt% of Cloisite 30B/5wt% phosphorus Containing Silane 146 Figure 4.8: Residue from 10wt% of Cloisite 30B/5wt% phosphorus Containing Silane 147 Table 4.1: EDX Analysis of the Residue from Cloisite 30B Nanocomposites Con- taining phosphorus Silane Element Weight Percentage Atom Percentage O 44.48 47.24 C 26.87 38.01 Si 9.01 2.74 Al 3.69 2.32 Mg 0.69 0.48 P 0.61 0.33 4.7 Discussion 4.7.1 Char Formation The goal of this work was to use phosphorus containing silane to enhance char formation during combustion. The experiments showed some enhanced char formation but the improvement in the char formation was limited. There are three possible reasons for the relatively small e ect of the phosphorus modi cation on the ammability. First, the char forming ability of the phosphorus (in the form of phosphate) on the selected silane is weak. Second, the ratio of phos- phorus to silane is one to one, which may be too small. The sites on the clay for the silane to attach were limited, if there was more phosphorus per silane, more phos- phorus could have been attached on the surface of clay. Third, the absolute amount 148 of phosphorus is relatively small. 4wt% of phosphorus may be added to resins to achieve desired ame retardant ability [150], where there is no more than 0.5wt% of phosphorus added to the PMMA in the phosphorus modi cation experiment. 4.7.2 Position of the phosphorus Compound During Combustion The EDX observation examined sample, where the atom ratio of phospho- rus is 0.33% and weight percentage is 0.61%. The initial phosphorus content of nanocomposite is about 0.5wt%, and if the phosphorous compound did not leave the condensed phase, the weight percentage of phosphorus in residue should be around 5wt%. It is possible that most of the phosphorus has evaporated during the combustion. This could explain the di erence between the gasi cation and cone calorimetry experiments. The gasi cation experiments showed a 20% improvement in the mean mass loss rate and the cone calorimetry experiment showed a 25% improvement in the mean heat release rate. This di erence may because the phosphorous atoms in the gaseous phase acted as free radical scavengers and reduced the heat released from the ame. In the gasi cation experiment, there is no gaseous phase reaction, so such e ect was not observed. 149 Chapter 5 Conclusions Di erent layered silicate/PMMA nanocomposites were synthesized by chang- ing the type of the silicate, modi cation of silicate, polymer molecular weight, and synthesis method to observe the e ect of the dispersion of the silicate, the e ect of the aspect ratio of silicate, and the molecular weight of the matrix polymer on the ammability of layered silicate/PMMA nanocomposites. The silicates used included two types of montmorillonite with di erent organic modi cations (Cloisite 15A and Cloisite 30B) and one organically modi ed synthetic mica (Somasif MAE120). The structure of the nanocomposites was studied by XRD, SAXS, and TEM. By coupling XRD, SAXS, and TEM together, the dispersion of silicate was deter- mined at di erent length scales. Nanocomposites with di erent degrees of dispersion were produced and con- rmed by structural characterization. The ammability of the nanocomposites was characterized by TGA, gasi cation and cone calorimetry. An in-situ polymerization method was used to produce nanocomposites with a high degree of exfoliation, but the polymer molecular weight reproducibility was relatively poor. A solution mixing method produced nanocomposites with controlled polymer molecular weight but the degree of exfoliation was not as high as the samples 150 made using the in-situ polymerization method. To improve the dispersion for the nanocomposites made by solution mixing, a method to graft short polymer chains onto the silicate surface using a silane coupling agent was developed. The XRD data showed that the degree of order of the layered silicates in the nanocomposites from high to low is Somasif MAE120, Cloisite 15A, Cloisite 30B and PMMA grafted respectively. Small angle X-ray scattering (SAXS) data con rmed the degree of exfoliation in the nanocomposites from high to low is PMMA grafted, Cloisite 30B, Cloisite 15A and Somasif MAE120. Low magni cation TEM con rmed layered silicates are well distributed in the polymer matrix. High magni cation TEM con rmed the degree of exfoliation from high to low is consistent with the SAXS data. The structural characterization at di erent length scales con rmed successful synthesis of nanocomposites with di erent degrees of exfoliation by changing silicate and that grafting short polymer chains onto the silicates improves the degree of exfoliation. The dispersion in the polymer matrix is further con rmed by glass transition temperature measurements by DSC. This work con rmed that by forming a well dispersed layered silicate/polymer nanocomposite, the ammability of the nanocomposite can be reduced compared to pure polymer and traditional composite. The peak mass loss rate of nanocomposites can be as low as 53% (low molecular weight) and 48% (high molecular weight) compared to pure PMMA in a gasi cation experiment. The peak heat release rate of nanocomposites can be as low as 41% (low molecular weight) and 57% (high molecular weight) of pure PMMA in a cone calorimetry experiment. As long as a true nanocomposite is formed, the precise degree of exfoliation has little impact on 151 the ammability. Nanocomposites with large aspect ratio silicate showed lowered ammability. Nanocomposites with high molecular weight matrix PMMA polymer also showed lowered ammability. The degradation kinetics of the nanocomposites were also studied, and it was found that nanocomposites showed higher activation energies than the corresponding pure polymers. However the increase in the activation energy is not as important an e ect as the stability of the residue layer. Nanocomposites with low molecular weight have an higher activation energy than nanocomposites with high molecular weight, but the nanocomposites with high molecular weight showed lowered ammability. It is also interesting to note that the sample form factor has a signi cant impact on the nominal reaction order of the nanocomposites as compared to microcomposites or pure polymer. This properties can be used as a quick method to identify the formation of a nanocomposite. The residues after ammability measurements were examined by SEM and XRD. The residues were loosely packed clay stacks with porous structure. The clay platelets collapsed into stacking structures close to the original unmodi ed clay structure. Based on the structural and ammability measurements, it was concluded that the mechanism for lowering the ammability for nanocomposites was from the formation of a shielding layer at the surface of the nanocomposite composed of the collapsed silicate layers. This shielding layer acts to both block heat radiation and to prevent the polymer degradation products from escaping and combusting. As the improvement in nanocomposite ammability is derived largely from the 152 physical presence of a residue layer, one method to further lower the ammability is to introduce chemical reagents that can enhance char and residue layer formation during combustion. A phosphorous compound was selected and a method was de- veloped using silane to couple it to the surface of the silicate. The nanocomposites made with phosphorus modi ed silicate showed enhanced char formation ability and lowered ammability. The nanocomposite with phosphorus modi ed silicate showed a 20% reduction in mean mass loss rate in a gasi cation experiment and a 25% reduc- tion in mean heat release rate in a cone calorimetry experiment. This demonstrated a synergy of chemical ame retardants with layered silicate nanocomposites that can further decrease the ammability. 153 Chapter 6 Future Work 6.1 Introduction Based on the work presented here and in the literature, it can be concluded that layered silicate nanocomposites do not have low enough ammability to pass industry standards. To increase the ame retardancy, one potential method is to synthesize inorganic phosphate nanocomposites. Phosphate can potentially enhance the char forming ability of polymer nanocom- posites. By enhancing char forming during combustion, the heat release rate and mass loss rate could be lowered and the structural integrity could be enhanced. As a result, the overall ame retardant properties of polymer could be improved. There are several requirements for the phosphate if it is to be used as a ame retardant. First, the phosphate must have low or no toxicity, which ensures safety in manufacturing, deployment and recycling of the product. Second, the phosphate should be stable at the processing temperature of the polymer, so the phosphate won?t prematurely degrade and loose its ame retardancy. Third, the phosphate should have low solubility in water, especially for applications as PCB and IC pack- aging materials, otherwise it may leach out and cause problems when exposed to a high humidity environment. Phosphates can also lead to environment issues if they leach out into the environment. 154 It will be an advantage if the phosphate also has a phase transformation (melt, solid state phase transformation, dehydration etc.) at a temperature that is lower than the thermal degradation temperature of the polymer matrix. The phase trans- formation would absorb energy, lowering the temperature of the polymer matrix. In the case of a small and temporary ignition source (like a candle, match or cigarette), the polymer may not ignite because of the phase transformation. In the case of a large ignition source, the ignition could be postponed. One of the problems of traditional phosphorous based ame retardants is that the char formation is not uniform, because the phosphorous based ame retardant phase separates into micron sized domains. Only the surface layer of the ame retardant domains can contribute to char formation. If the phosphorous retardant was dispersed as a nanocomposite, it would contribute more e ectively as char forming agent during combustion. If the phosphate nanoparticles are well dispersed in the polymer matrix, the char formation should also be more uniform. Similar to layered silicates, many inorganic phosphates also have a layered structure and many inorganic phosphates also have cation exchange capacity. So it should be possible to organically modify phosphates by cation exchange and produce hydrophobic organically modi ed phosphates. The organically modi ed inorganic phosphates should have low water solubility and low chance of leaching out in high humidity environments. 155 6.2 Proposed Experiment The proposed experiments can be separated into several steps. The rst step is material selection, the second step is synthesis and organic modi cation of the phosphate nanoparticles, the third step is nanocomposites synthesis, the fourth step is structural characterization and the fth step is ammability characterization. 6.2.1 Material Selection There are many inorganic phosphates which are commercially available, some calcium phosphates with their properties are listed in table 6.1 and some others are listed in table 6.2. It can be concluded that the calcium phosphates have the least toxicity [151, 152], the least water solubility and they are potentially the best choice for consideration in synthesizing a nanocomposite. Octacalcium phosphate can be synthesized by adding phosphate solution to a calcium solution or vice versa [153]. Octacalcium phosphate has a layered struc- ture, and organic modi cation has been reported [154]. The octacalcium phosphate has been mixed with PEO and some other polyelectrolytes to make composites for bone cement applications [155]. It is possible to make polymer/layered octacalcium phosphate nanocomposites by manner similar to making polymer/layered silicate nanocomposites. Barium hydrogen phosphate is also a very interesting material with a melting point of 410 C, which is near the nal degradation temperature of PMMA and lower than the degradation temperature of many other polymers. When exposed 156 to a heat source, barium hydrogen phosphate can absorb energy by melting and shield the surface of polymer by forming a glassy inorganic melt. Barium hydrogen phosphate also has low toxicity. Another potentially interesting inorganic phosphate is zinc phosphate, which has been used as an anticorrosion coating for metals. Both PMMA and PS would be the logical polymer matrix for this study to examine the e ect of ammability on both an unzipping and an unbuttoning type polymer. 6.2.2 Organic Modi cation of Nanoparticles and Nanocomposite Syn- thesis Depending on the phosphate selected, two di erent organic modi cation meth- ods could be used. If the phosphate has a layered structure, it could by organically modi ed by cation exchange. The phosphate would be dispersed in water (or wa- ter/ice mixture if the phosphate dissolve in water), then cation surfactant is added to the dispersion. The modi ed phosphate will become hydrophobic and should phase separate from the water. If the phosphate does not have a layered structure, then the phosphate can be milled into nanoparticles by ball milling or synthesized by a sol-gel method and organically modi ed by silane or Tyzor R . Solution mixing would be the preferred method for making the nanocomposite and the nanocomposite can be retrieved from the solution by either precipitation or evaporating the solvent. 157 Table 6.1: Commercially Available Calcium Phosphates Name CAS No. Structure Water Soluble Toxicity Tm ( C) Tricalcium phosphate 7758-87-4 Ca3(PO4)2 no no, bone and teeth composition [156] 1670 [157] Hydroxyl- apatite 1306-06-5 Ca5(PO4)3(OH) no no, same as above [158] 800 Octacalcium phosphate 13767-12-9 Ca8H2(PO4)6 5H2O low no, bone cement [155, 159] Dicalcium phosphate 7757-93-9 CaHPO4 no no, dietary supplement [160, 161] Monocalcium phosphate 7758-23-8 Ca(H2PO4)2 no no, tooth- pastes [162] 1230 158 Table 6.2: Commercially Available Inorganic Phosphates Name CAS No. Structure Water Soluble Toxicity Tm ( C) Aluminum dihydrogen phosphate 13530-50-2 Al(H2PO4)3 yes TSCA listed ? Aluminum phosphate 7784-30-7 AlPO4 no TSCA listed 1500 Ammonium hydrogen phosphate 7783-28-0 (NH4)2HPO4 yes TSCA listed 155 Barium hydrogen phosphate 10048-98-3 BaHPO4 yes Harmful by inhalation and swallow 410 Barium metaphos- phate 13762-83-9 Ba(PO3)2 no 1560 159 6.2.3 Structure Characterization Optical microscopy would be used to examine the existence of large particles. Transmission electron microscopy can be used to study the dispersion of nanoparti- cles. SAXS would be used to study the dispersion of nanoparticles. If the selected phosphate has a layered structure, XRD should be used to examine the intercalation and exfoliation of the layered phosphate. 6.2.4 Flammability Characterization Due to the expected char forming properties, TGA and cone calorimetry would be the major methods used to study the ammability of the polymer/phosphate nanocomposites. Oxygen index ammability can also be used to examine the e ect of the phosphate on the ammability. 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