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Investigation of mechanical, thermal and rheological properties of fluorocarbon functionalized polystyrene; and, Development of increased flow graphene based polymer nanocomposite membranes

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Investigation of mechanical, thermal and rheological properties of fluorocarbon functionalized polystyrene; and, Development of increased flow graphene based polymer nanocomposite membranes

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Content INVESTIGATION OF MECHANICAL, THERMAL AND RHEOLOGICAL PROPERTIES OF FLUOROCARBON FUNCTIONALIZED POLYSTYRENE AND DEVELOPMENT OF INCREASED FLOW GRAPHENE BASED POLYMER NANOCOMPOSITE MEMBRANES by Hayriye Merve Yurdacan A Dissertation Presented to the FACULTY OF THE USC GRADUATE SCHOOL UNIVERSITY OF SOUTHERN CALIFORNIA In Partial Fulfillment of the Requirements for the Degree DOCTOR OF PHILOSOPHY (CHEMISTRY) DECEMBER 2015 Copyright 2015 Hayriye Merve Yurdacan ! ii! Dedication To my mom, Sevim Yurdacan and my dad, Dr. Fikri Yurdacan ! iii! ACKNOWLEDGEMENTS First and foremost, I want to thank my advisor Dr. Thieo Hogen-Esch who helped and guided me for many years. It has been a great honor for me to be his PhD student. His enthusiasm and guidance in research always inspired me and helped me to understand how to be a good researcher. Thank you very much for always being open to my ideas. Many thanks for encouraging me to implement my ideas and give me endless courage in my research. I want to extend my sincere thanks to Dr. Oliver Franke for his tremendous help in my research. Sadly, we lost him last year. He and his family will always be in our thoughts and prayers. I am also deeply thankful to Dr. Steve Nutt, for his advice and support through my research. Thank you for letting me to use all necessary instruments for my studies and also for your time to analyze and discuss the results throughout mechanical and thermal studies. I want to thank Dr. Massoud Pirbazari and Dr. Varadarajan Ravindran for their advice and help on membrane filtration studies. I want to thank my qualifying exam and dissertation committee members: Dr. Surya Prakash, Dr. Steve Nutt, Dr. Stephen Bradford, Dr. Sri Narayan. Thank you for your time and support through my studies. Thank you past and current members of Dr. Hogen Esch’s lab: Dr. Jingguo Shen, Dr. Victoria Puinova, Dr. Peng Jiang, Dr. Janet Olsen, Dr. Ming Li, Dr. Bing Xu, Sergey Mukhin and Adam Ung. You were always very helpful and it was a big pleasure for me to have the opportunity to work with you. ! iv! I want to thank Dr. Erickson who has supported me through my PhD studies. We have worked together for many years. Thank you for giving me opportunity to be your head Teaching Assistant for 5 years and also for your guidance and endless support. Special thanks to the members of Dr. Nutt’s lab: Thank you Xiaochen Li for your incredible help in DSC and Rheology measurements and for your tremendous contribution to my research. I want to thank Yuzheng Zhang who helped me with Nanoindentation measurements. I also thank Dr. Lessa Grunenfelder and Bo Jin for their help. I also wish to thank the members of Dr. Prakash’s Lab: Dr. Alain Goeppert, Dr. Miklos Czaun for helping me on TGA measurements. I also thank the many friends from USC: Dr. Yuan Ding, Dr. Marcos Sainz, Laxman Gurung, Dr. Somesh Kumar, Courtney Downes, Seyma Ekiz, Gozde Barim, Gokce Ozcelik and Serra Ongun for making the graduate experience more enjoyable one. A special thanks to the Loker Hydrocarbon Institute for supporting my research and LHI and USC Department of Chemistry collegues and staff: Dr. Robert Anizfeld, David Hunter, Jessy May, Carole Phillips, Magnolia Benitez and Michele Dea. Finally, I would like to thank my mom Sevim Yurdacan and my dad Dr. Fikri Yurdacan who always supported me through my career. I love you so much. Without you, it would not have been possible. Thank you my brothers, Fikret Bartu Yurdacan and Sevket Altug Yurdacan for their endless support. ! v! TABLE OF CONTENTS Dedication ....................................................................................................................................... ii Acknowledgements ........................................................................................................................ iii List of Schemes .............................................................................................................................. ix List of Figures ................................................................................................................................. x List of Tables ................................................................................................................................ xv Abstract ........................................................................................................................................ xvi CHAPTER 1 Viscoelastic Properties of Fluorocarbon End Functionalized Polystyrene Films at the microscale by Nanoindentation .......................................................................1 1.1 Introduction ......................................................................................................................... 1 1.1.1 Nanoindentation ......................................................................................................... 3 1.1.1.1 Dynamic Nanoindentation ................................................................................ 4 1.2 Experimental ....................................................................................................................... 7 1.2.1 Materials .................................................................................................................... 7 1.2.2 Initiator Synthesis ...................................................................................................... 7 1.2.3 Synthesis of Perfluoro End-functionalized Polystyrene ............................................ 8 1.2.4 Sample Preparation .................................................................................................... 9 1.2.5 Characterization ....................................................................................................... 10 1.2.5.1 Nanoindentation Measurement Design ........................................................... 10 1.2.6 Molecular Weights of Perfluorocarbon Functionalized Polymers ........................... 12 ! vi! 1.3 Nanoindentation Studies ................................................................................................... 12 1.3.1 Effect of shorter RF groups on the viscoelastic properties of PS by Nanoindentation ........................................................................................................23 1.4 Conclusions ....................................................................................................................... 31 1.5 References ......................................................................................................................... 32 CHAPTER 2 Rheological Properties ............................................................................................ 48 2.1 Introduction ....................................................................................................................... 48 2.2 Rheology Analysis of PFTD functionalized PS 30K ........................................................ 49 2.3 Rheological Properties of of C 7 F 15 -PS .............................................................................. 53 2.4 Conclusions ....................................................................................................................... 54 2.5 References ......................................................................................................................... 55 CHAPTER 3 Thermal and Optical Properties of Perfluoroalkyl End-Functionalized Polystyrenes Using Optical Transmission and Differential Scanning Calorimetry ..............................................................................................................56 3.1 Introduction ....................................................................................................................... 56 3.2 UV-Visible measurements for PS and RF end-functionalized films ................................ 58 3.3 Termogravimetric analysis of perfluoro end-functionalized PS ....................................... 63 3.4 Differential Scanning Calorimetry analysis of Perfluoro End-functionalized PS ............ 66 3.5 References ......................................................................................................................... 73 ! vii! CHAPTER 4 Optical, Thermal and Viscoelastic Properties of isobutyl POSS-end functionalized PS .....................................................................................................82 4.1 Introduction ....................................................................................................................... 82 4.2 Experimental ..................................................................................................................... 83 4.2.1 Materials .................................................................................................................. 83 4.2.2 Synthesis of POSS Macroinitiator (POSS-Br) ......................................................... 84 4.2.3 Preparation of POSS nanocage end-functionalized hybrid Polystyrene via ATRP ....................................................................................................................... 84 4.2.4 Characterization ....................................................................................................... 85 4.3 Molecular Weights of POSS Functionalized Polymers .................................................... 86 4.4 UV-Visible measurements for PS and POSS end-functionalized films ........................... 86 4.5 Differential Scanning Calorimetry analysis of POSS end-functionalized PS .................. 87 4.6 Thermogravimetric analysis of POSS end-functionalized PS .......................................... 88 4.7 Rheological Properties of POSS-PS Hybrid Polymers ..................................................... 89 4.8 References ......................................................................................................................... 91 CHAPTER 5 Development of Increased Flow Graphene Oxide-Polymer Nanocomposite Membranes ...............................................................................................................95 5.1 Introduction ....................................................................................................................... 95 5.1.1 Reverse Osmosis: Basic Principles .......................................................................... 96 5.1.2 Transport through RO membranes ........................................................................... 97 ! viii! 5.2 Current Reverse Osmosis / NanoFiltration Membranes ................................................... 98 5.3 Experimental ................................................................................................................... 100 5.3.1 Materials and Methods ........................................................................................... 100 5.3.2 Synthesis and Preparation of Graphene Oxide Incorporated PolyAmide(PA) 5.3.2 Membranes ............................................................................................................. 100 5.3.3 Preparation of Graphene Oxide incorporated Blends of Poly(tetrabutylammonium styrene sulfonate copolymer and Polyvinylidenefluoride PBASS- PVDF blend and Polystyrenesulfonic acid copolymer-polyvinylidenefluoride (PSSA-PVDF) blend membranes ..................102 5.4 Water Permeability and Particle Rejection Studies of Graphene Oxide Incorporated 5.4 PolyAmide (PA) Membranes .......................................................................................... 103 5.5 Water Permeability and Particle Rejection Studies of GO incorporated PBASS-PVDF 5.5 blend (PSSA-PVDF) blend membranes .......................................................................... 105 5.6 Conclusions ..................................................................................................................... 110 5.7 References ...................................................................................................................... 111 Bibliography ............................................................................................................................... 118 ! ix! LIST OF SCHEMES Scheme 1.1. Representation of a Triboindenter ........................................................................... 4 Scheme 1.2. Force vs. Displacement of dynamic nanoindentation ............................................. 5 Scheme 1.3. Load Profiles for dynamic testing ........................................................................... 6 Scheme 1.4. Proposed association of PFTD end-functionalized PS (Low MW 10K-15K) in bulk ....................................................................................................................... 22 Scheme 4.1. Structure of POSS ................................................................................................. 83 Scheme 5.1. Reverse Osmosis .................................................................................................. 96 Scheme 5.2. Asymmetric Structure of the Current membranes ................................................. 98 Scheme 5.3. GO PA Membrane Fabrication ........................................................................... 101 Scheme 5.4. PA Membrane Synthesis ..................................................................................... 102 ! x! LIST OF FIGURES Figure 1.1. Optical Microscopy (OM) image of Polystyrene film on silicon wafer after heating at 110 °C and then a)rapid cooling b) slow cooling .................................... 9 Figure 1.2. E’ of a) 100 nm b) 5 µm PS film ............................................................................. 10 Figure 1.3. Optical Microscopy image of indents on PS film ................................................... 11 Figure 1.4. X-Bar chart for PS15K showing the average E’ of various indents from different regions of the polymer as a function of displacement at 45 Hz at ambient temperature ..............................................................................................................15 Figure 1.5. X-Bar chart showing the average E’ of various indents from different regions of the polymer at different frequencies for PS15K Polymer at ambient temperature ..16 Figure 1.6. E’ of individual indents from 3 randomly selected regions of the polymer at different frequencies for PS15K Polymer at ambient temperature ..........................17 Figure 1.7. S chart showing Standard deviation, σ, of E’ under conditions of Figure 1.6 ........ 17 Figure 1.8. X-Bar chart showing the average E’ of various indents from different regions of the polymer at different frequencies for PSC1315K Polymer at ambient temperature ...............................................................................................................18 Figure 1.9. S chart showing the average E’ of various indents from different regions of the polymer at different frequencies for PSC1315K Polymer at ambient temperature .18 Figure 1.10. E’ of individual indents from 3 randomly selected regions of the polymer at different frequencies for PSC13-15K Polymer at ambient temperature ...................19 Figure 1.11. Storage Moduli of PS-15K and PSC13-15K Polymers by NI at 25 °C. The thickness of the films were around 5 µm and the NI depth was ≈300 nm ...............20 Figure 1.12. Loss Moduli of PS-15K and PSC13-15K Polymers by NI at 25 °C. The ! xi! thickness of the films were around 5 µm and the NI depth was ≈350 nm ...............21 Figure 1.13. X-Bar chart showing the average E’ of various indents from different regions of the polymer at different frequencies for PSC710K Polymer at ambient temperature ..............................................................................................................23 Figure 1.14. E’ of individual indents from 3 randomly selected regions of the polymer at different frequencies for PSC7-15K Polymer at ambient temperature ................... 24 Figure 1.15. T-test analysis on PSC7-10K and PS ...................................................................... 25 Figure 1.16. S chart showing the average standard deviation of E’ of various indents from different regions of the polymer at different frequencies for PSC710K Polymer at ambient temperature .............................................................................................25 Figure 1.17. X-Bar chart showing the average E’ of various indents from different regions of polymer at different frequencies for PSC1010K Polymer at ambient temperature .26 Figure 1.18. E’ of individual indents from 3 randomly selected regions of the polymer at different frequencies for PSC10-15K Polymer at ambient temperature ................. 27 Figure 1.19. S chart showing the average standard deviation of E’ of various indents from different regions of the polymer at different frequencies for PSC10-10K Polymer at ambient temperature .............................................................................. 28 Figure 1.20. Storage Moduli of PS-10K and PSC7-10K and PSC10-10K Polymers by NI at 25 °C. ...................................................................................................................... 29 Figure 1.21. Loss Moduli of PS-10K and PSC7-10K and PSC10-10K Polymers by NI at 25 °C ....................................................................................................................... 30 Figure 1.22. Box Plot representation of average E’ of PS-10K, PSC7-10K and PSC10-10K .... 30 Figure 1.23. Box Plot representation of average E’’ of PS-10K, PSC7-10K and PSC10-10K ... 31 ! xii! Figure 2.1. Shear Storage (filled) and loss moduli (empty) for PS-30K (diamond) and PSC13-30K (triangle) polymers. The reference frequency and strain are 1Hz and 0.5 %,respectively. ................................................................................. 49 Figure 2.2. Shear Storage Moduli for PS-30K(diamond) and PSC13-30K(triangle) polymers. The reference frequency and strain are 0.1Hz and 0.5 %, respectively. .............................................................................................................50 Figure 2.3. Shear Loss Moduli for 30K PS (diamonds) and PSC13-30K polymers at frequency and strain of 0.1Hz and 0.5 percent, respectively.. ................................ 51 Figure 2.4a. Master curves of G’ G’’for PS-30K at 180 ° C. ...................................................... 51 Figure 2.4b. Master curves of G’,G’’ for PSC13-30K at 180 ° C. ............................................. 51 Figure 2.5. Frequency dependence of G’ for PS-30K (diamond) PSC13-30K (triangles) at 180 °C. Reference strain is 0.5 % ...........................................................................52 Figure 2.6. Frequency dependence of G´´ for PS-30K (diamond) PSC13-30K (triangles) ....... 53 Figure 2.7. Frequency dependence of G´ (filled) and G´´(empty) of PS-10K (square) and PSC7-10K (spheres) at 120 °C. Reference strain is 0.5 %. .....................................54 Figure 3.1. UV-Visible optical transmittance versus wavelength of the PS-10K (blue), PSC7-10K (red), PSC10-10K(purple) and PSC13-10K (green) polymer films ......58 Figure 3.2. UV-Visible Transmittance versus wavelength of 10K and 30K PSC7 polymer films on micro cover glass. ..................................................................................... 60 Figure 3.3. UV-Visible Transmittance versus wavelength of PSC10-10K and PSC10 -30K polymer films on micro cover glass. ....................................................................... 61 Figure 3.4. UV-Visible Transmittance versus wavelength of PSC13-10K (green) ................... 62 ! xiii! Figure 3.5. TGA curves of PS-10K (light blue) , PS-15K (dark blue), PSC13-10K(red) and PSC13-15K(green) ...................................................................................................64 Figure 3.6. TGA curves of PS 10K, PSC7 10K and PSC13 10K. ............................................. 65 Figure 3.7. TGA curves of PS 15K , PSC7 15K and PSC13 15K. ............................................ 65 Figure 3.8. DSC thermograms of PSC13 polymers with varying molecular weights. PSC13-10K(Blue), PSC13-15K (Green), PSC13-30K(Pink) ..................................67 Figure 3.9. DSC thermograms of PSC7-10K(green), PSC10-10K (red), PS1C13-10K (purple) ..............................................................................................68 Figure 3.10. DSC thermograms of PSC7, PSC10 and PSC13 of 15K polymer ...........................69 Figure 3.11. DSC thermograms of PSC7,PSC10, PSC13 and PS30K .........................................70 Figure 4.1. UV-Visible Transmittance versus wavelength of PS POSS 5K and 10K ............... 86 Figure 4.2. Thermogravimetric analysis of POSS-PS 10k (green), PS-10K (red) polymers under air .................................................................................................................. 88 Figure 4.3. Thermogravimetric analysis of POSS PS polymers of various molecular weights under nitrogen . PS-10K (Pink), POSSPS-10K (Blue) ............................................89 Figure 4.4. Shear Storage moduli for POSSPS-10K(blue) , PS-10K (red) polymers.The reference strain is 0.5 %, respectively .....................................................................90 Figure 4.5. Shear Loss moduli for POSSPS-10K(blue) , PS-10K (green) polymers The reference strain is 0.5 %, respectively .....................................................................90 Figure 5.1. Sulfonated P(BASS-S-CMS)- PVDF annealed at 70 °C ....................................... 107 Figure 5.2. Permeate flux of region I and III of Sulfonated P(BASS-S-CMS)-PVDF annealed at 70 °C .................................................................................................. 107 Figure 5.3. Sulfonated P(BASS-S-CMS)- PVDF annealed at 165 °C before membrane ! xiv! filtration ................................................................................................................. 108 Figure 5.4. Sulfonated P(BASS-S-CMS)- PVDF annealed at 165 °C after membrane filtration ................................................................................................................. 108 Figure 5.5. GO coating on PES support before filtration ........................................................ 109 Figure 5.6. GO coating on PES support after filtration ........................................................... 109 Figure 5.7. Permeate flux data observed during membrane testing using wastewater ............ 110 ! xv! LIST OF TABLES Table 1.1. Mechanical properties of RF end-functionalized PS synthesized using ATRP ........ 12 Table 3.1. RF end-functionalized polystyrenes synthesized using ATRP methods .................. 59 Table 3.2. Summary of glass transition temperatures C 7 F 15 , C 7 F 15 and C 7 F 15 end functionalized polystyrenes and the corresponding PS homopolymers ....................68 Table 4.1. Molecular Weights and Polydispersities of PS and POSS-PS hybrid Polymers ...... 86 Table 4.2. Summary of glass transition temperatures C 7 F 15 , C 7 F 15 and C 7 F 15 end functionalized polystyrenes and the corresponding PS homopolymers ....................87 Table 5.1. Types of Membranes ................................................................................................ 95 Table 5.2. Reaction conditions of prepared membranes .......................................................... 103 Table 5.3. Water flux and NaCl rejection properties of membranes ....................................... 104 Table 5.4. Performance comparison of various membranes .................................................... 105 ! xvi! ABSTRACT Semi-fluorinated polymers are of more recent origin and have several applications due to their unique properties that differ from the corresponding analogous hydrogen or fluorocarbon polymers. They have gained attention in many applications. The presence of perfluorocarbon (RF) (C 7 F 15 to C 13 F 27 ) end groups has been shown to enhance the compatibility of polystyrene (PS) blends through fluorophilic interactions. RF functionalized PS with narrow molecular weight (MW) distribution were synthesized by Atom Transfer Radical Polymerization (ATRP). In Chapter 1, RF end functionalized PS and PS thin films (~ 5µm) deposited on silicon wafers were studied at microscopic level by Tribo nano-indenter at ambient temperature. The effect of RF length from C 7 F 15 to C 13 F 27 on the viscoelastic properties of PS was investigated. Two modes of nanoindentation were tested: ramp load and frequency sweep. Frequency sweep experiments were carried out between 10-100 Hz. Nanoindentation measurements showed increases of about 85 and 300 percent in the storage and loss moduli respectively compared with the unfunctionalized PS samples at ambient temperature. Shorter RF groups had much smaller effects. In Chapter 2, the effect of end RF groups on viscoelastic properties at high temperatures was studied. Polymer melts were tested by rheological measurements. Despite relatively modest increases in E’ and E’’ shear storage modulus (G’) of the C 13 F 27 end functionalized polystyrenes (30K) is between one to two orders of magnitude larger than the corresponding PS homopolymers whereas the shear loss modulus (G’’) was 50% higher between 150 and 180 0 C. In Chapter 3, the effect of RF length, RF content and polymer MW on the glass transition temperature, optical transmittance and thermal degradation behavior were investigated. Polymers were studied by differential scanning calorimetry (DSC), thermogravimetric analysis (TGA) and ! xvii! UV-visible spectroscopy. In Chapter 4, mechanical, optical and thermal properties of POSS end-functionalized PS were investigated. The magnitude of effect was compared with RF end-functionalized PS. In Chapter 5, the water flux and rejection properties of Polyamide and P(BASS-S- CMS)/PVDF polymer blend nanocomposite membranes incorporated with GO were studied. It has been shown that incorporation of 0.5-5wt% of GO markedly improved water flux and fouling behavior of membranes. ! 1! CHAPTER 1 Viscoelastic Properties of Fluorocarbon End Functionalized Polystyrene Films at the microscale by Nanoindentation 1.1 Introduction Semi-fluorinated polymers are materials with many potential applications due to their unique properties that are different from analogous hydrogenated polymers. They have gained much attention in antifouling, super-hydrophobic and self-cleaning coatings. 1-8 Other application are microelectronics where they are mainly used as low dielectric constants materials due to low the polarizability of C-F bond. 9-11 Their unique self-assembly and orthogonal (hydrophobic and lipophobic) properties promote their use in organic semiconductors. 12-13 Their tendency towards self-organization and low surface free energies also makes them promising components in photoresponsive materials 14-15 . Low MW perfluorocarbons (RFs) show weak intermolecular interactions compared with hydrocarbons as expressed in much lower cohesive energy densities. 16 As indicated above, it is well known that RFs have both pronounced hydrophobic 17-26 as well as lipophobic properties. 19- 22,26-29 The term fluorophilic is a less well defined- but more general- term to describe aggregation of RFs or RF groups that is driven by their low cohesive energies (δ ≈ 12 ± 0.7 MPa 1/2 ) compared to aliphatic hydrocarbons of comparable chain lengths (δ ≈ 15 ± 1.0 Mpa 1/2 ) and typical glassy polymers such as polystyrene (δ =18 ± 0.5 Mpa 1/2 ). However, compared to hydrocarbons, the melting points (mp’s) of perfluorocarbons increase much faster with chain lengths. 30 For instance, while the melting points (mp’s) of octane and perfluorooctane are nearly the same, the mp’s of tetradecane and perfluorotetradecane are about 275 and 400 K respectively a difference of about 70 and 125 degrees, respectively compared to the corresponding hydrocarbons. This suggests the occurrence of much stronger interactions for ! 2! longer RFs and presumably RF groups as well. Higher MW (10 5 -10 6 D) perfluorocarbons (Teflons) have extremely high mp’s (~ 500-621 K). This polymer has been shown to exist in four different helical crystalline states I-IV, depending on temperature and pressure 31-32 and appears to be susceptible to rapid helix reversals at elevated temperatures. 31 In agreement with the above, numerous polymers containing RF pendent- or end groups have been shown to exhibit mesogenic properties 32-51 and low surface free energies. 52-56 The effects of association of pendent RF groups on the rheological properties of aqueous polymer solutions 24,25 hydrogels 26,57-59 and non-aqueous polymers 60 have been demonstrated. However, these effects, especially that of RF groups on the mechanical and rheological properties of solid state polymers remain to be studied. In addition, the potential of long perfluorocarbon end groups to compatibilize homogeneous blends of low MW (≤20K) polystyrene(PS) and polybutylmethacrylates(PBMA) through fluorophilic interactions has only recently been demonstrated. 60-63 The TEM morphologies of these blends and other data indicate that the C 13 F 27 , compared to C 7 F 15 and C 10 H 21 end groups, are particularly effective in mediating compatibilities as judged by the approximate size of the domains of blends of RF end functionalized PS and PBMA. 61 Furthermore, studies on blends of end functionalized RF- PSRF/RF-PBMA have indicated that the association of RF end groups, at the same RF content and –length, is more pronounced than that of RF pendent groups. 59 This has been verified for several RF groups and is consistent with end groups have greater mobility and are subject to smaller excluded volume effects. 62 Thus, for 1/1 (wt/wt) RF-PS/RF-PBMA blends having an approximate MW of 20,000, the lamellar sizes of the blends of the C 13 F 27 end functionalized polymers were found to be 61-62 small (< 10 nm) compared to blends of the C 7 F 15 and C 10 H 21 end- functionalized polymers (> 50 nm). The TEM data of the C 13 F 27 end functionalized blends with ! 3! MWs of around 13-15k indicate even smaller domain sizes (1-3 nm). Hence the effect(s) of end- functionalization with longer RF groups on the mechanical properties of well-studied polymers, such as PS would seem to be of interest. Here we report the synthesis of a number of a number of C 7 F 15 , C 10 F 21 and C 13 F 27 end- functionalized low MW (≤30 kD) PS using atom transfer radical polymerization (ATRP) and its effects on the viscoelastic properties of polystyrene (PS). As pointed out above, PFTD group is of special interest as association mediated by this group was shown to be stronger than the smaller RF groups. To the best of our knowledge these systems appear not to have been studied. Nanoindentation tests are employed to investigate the hypothesis that the mesogenic RF groups are more likely to prominently show the effects of formation of semicrystalline RF domains that mediate stronger inter-polymer association. In turn, this may result in improved mechanical properties. Polystyrene (PS) has been chosen as a model as it is a well-known and studied thermoplastic polymer with a number of desirable properties including high glass temperature, good processability and therefore use in applications in polymer electrolyte membrane fuel cells (PEMFCs) 64-68 , electronics 69-73 and in block copolymers or their blends. 74-77 1.1.1 Nanoindentation Nanoindentation (NI) is a new but well-established mechanical testing technique that originally was developed based on depth sensing hardness tests. 78-106 While nanoindentation was primarily developed for metals 81,82,86,91,98-100 , ceramics 101,102 and semiconductors 103,104 it has been successfully adapted for soft materials such as various tissues 88,92,95,107-110 and polymers 93,94,111-116 including polystyrene 115,116 . With the resolution of the devices being pushed into nm regime and the establishment of the Oliver-Pharr model to analyze load displacement curves in 1992 84 , NI quickly became a versatile tool to study thin films as well as bulk materials. 86-95 One ! 4! of the main advantages of NI is that it provides a high depth resolution but also allows measuring properties with a high lateral resolution. Hence NI is a valuable tool for generating a library of different polymers as thin films. 1.1.1.1 Dynamic Nanoindentation There has been a switch from quasi-static to dynamic nanoindentation tests to study the mechanical properties of polymers since dynamic nanoindentation enables the determination of viscoelastic properties such as complex modulus and loss factor for non linear viscoelastic solids or storage and loss modulus for linear viscoelastic solids. Scheme 1.1 represents a Triboindenter. Scheme 1.1. Representation of a Triboindenter 117 ! 5! Scheme 1.2. Force vs. Displacement of dynamic nanoindentation 117 A load-displacement DMA transducer was used to determine viscoelastic properties of polymers (Scheme 1.2). A small sinusoidal force having amplitude F 0 and angular frequency ω, the amplitude of displacement X, and phase shift between force and displacement ϕ, are given by 117 : != ! ! !!!! ! ! !!! ! !! (1) != !"# !! !" !!!! ! (2) Where m is the mass of the indenter, k is the combined stiffness and C is the combined damping coefficient. In the model, the stiffness and damping can be described as follows: != ! ! +! ! (3) !=! ! +! ! ! (4) Where C i and C s are the damping coefficients of the sample and displacement sensor and k i and k s are the stiffness of the sample and spring of the indenter. Air calibration is done for … Where Science Gets Down to Business … World Leader in Nanomechanical Test Instruments Hysitron Incorporated Sinusoidal test Time Force/Displacement Sinusoidal test Time Force/Displacement Elastic Time Time Force Displacement C ∝ φ k ∝ F/x c s A k E 2 ' π = c s A C E 2 ' ' π ω = Analysis of sinusoidal force/displacement data. Applications • Biomaterials • Paints • Contact lenses • Low-k dielectrics • Rubbers • Elastomers • MEMS testing • Fatigue testing Dynamic Specifications The nanoDMA technique increases the sensitivity of the instrument to provide augmented testing capabilities. The noise floor specifications in dynamic mode are seen below. • Frequency range: 1.0-300 Hz • Load noise floor: <100 nN • Displacement noise floor: <1 Å Software The nanoDMA ® software has been developed to automate and simplify a test technique and analysis that is potentially very involved and time-consuming. Dynamic test data is actually a convolution of the dynamic mechanical properties of the testing system and the sample. The software automates a calibration process that allows a real- time correction of data for the characteristics of the testing system. The test results are displayed real-time during the testing and are available for post-test processing. The data analysis tools calculate and display stiffness and loss data or storage and loss moduli as a function of load, load amplitude, displacement, frequency or time. With the nanoDMA ® technique and software, scientists and engineers can quantitatively obtain viscoelastic property measurements at the nanoscale in minutes. nanoDMA Data Analysis DMA-01 The nanoDMA ® software displaying data from a typical test on a polymer thin film. 10025 Valley View Road Minneapolis MN 55344 Tele: (1) 952.835.6366 Fax: (1) 952.835.6166 info@hysitron.com www.hysitron.com ! 6! instrument calibration and C i and k i values are subtracted from measured values to calculate stiffness and damping. The storage modulus E’, loss modulus E’’ and loss tangent tanδ were calculated by: ! ! = ! ! ! ! ! ! , (5) ! !! = !! ! ! ! ! , (6) !"#$= !!! !! = ! !! ! ! ! (7) Where A c is contact area. The indenter area function, A c was estimated using: 84 ! ! = 24.5ℎ ! +! ! ℎ ! +! ! ℎ ! !/! +! ! ℎ ! !/! +! ! ℎ ! !/! +! ! ℎ ! !/!" +! ! ℎ ! !/!" +! ! ℎ ! !/!" + ! ! ℎ ! !/!"# (8) Where A c and h c are the indenter contact area and depth respectively. In this study, two main modes were used: Ramp load and Frequency sweep (Scheme 1.3). Scheme 1.3. Load Profiles for dynamic testing 109 JOM • June 2008 50 www.tms.org/jom.html Figure 4. The setup used for in-vitro testing; inset: contact of the liquid tip with the sample and the meniscus of the test media around the shaft. Figure 2. An AFM image of a three-sided Berkovich tip. Figure 3. Materials response under dy- namic loading for an elastic-plastic and a visco-elastic materi- al. Figure 1. Load profi les for dynamic testing by depth- sensing indentation. ! 7! 1.2 Experimental 1.2.1 Materials All chemicals were purchased from Aldrich and used as received unless otherwise noted. Styrene was distilled under high vacuum after being stirred over CaH 2 overnight. Copper (I) Bromide (99.9%) was purified by stirring in acetic acid overnight and then washed with 2- propanol. Perfluoro alcohols were purchased from Synquest Labs, Inc. Silicon wafers for Nanoindentation studies were purchased from University Wafer(P/100 Prime Grade). 1.2.2 Initiator Synthesis ATRP initiator 1,1,-dihydroperfluorotetradecyl-2-bromo-2-methyl-propanoate, 1,1,- dihydroperfluoroundecyl-2-bromo-2-methyl-propanoate, 1,1,-dihydroperfluorooctyl-2-bromo-2- methyl-propanoate, having perfluorotridecyl, (PFTD) perfluorodecyl (PFD) and perfluoroheptyl (PFH) groups were synthesized using previous methods. 60-62 In a typical reaction, 4.3 mmol perfluoro alcohol, 0.2 mmol 4-(dimethylamino)-pyridine and 5.5 mmol triethylamine were dissolved in a mixture of 20% v/v hexafluorobenzene and 80% v/v toluene. The mixture was heated until the perfluoro alcohol dissolved. 5.5 mmol 2-bromoisobutyroyl bromide were added dropwise and the suspension was stirred at room temperature for 8 hours. The suspension was then filtered and toluene was removed by rotary evaporation. The product was dissolved in dichloromethane and washed three times with NaHCO 3 and then three times with deionized water. The product was then dried over MgSO 4 and the dichloromethane was evaporated by rotary evaporation and dried in vacuum oven at room temperature. Its structure was verified by 1 H NMR and 19 F NMR Spectroscopy. Proton NMR (400 MHz) of 1,1,-dihydroperfluorooctyl-2-bromo-2-methyl-propanoate, in CDCl 3 at 298 K, δ (ppm from TMS) : 4.45 (s, 2H), 1.90 (s, 6H); 19 F NMR (CDCl 3 as solvent, 298 K, ! 8! 400 MHz) δ (ppm from 2,2,2-trifluoroethanol) -81.5(3F, CF 3 ), -114.5(2F, CF 2 ), -122.6(2F, CF 2 ), -122.8(2F, CF 2 ), -123.6(2F, CF 2 ), -124.5(2F, CF 2 ), -126.9(2F, CF 2 ). Proton NMR (400 MHz)1,1,-dihydroperfluoroundecyl-2-bromo-2-methyl-propanoate in CDCl 3 298 K, δ (ppm from TMS): 4.45 (s, 2H), 1.90 (s, 6H); 19 F NMR (CDCl 3 as solvent, 298 K, 400 MHz) δ (ppm from 2,2,2-trifluoroethanol) -81.5 (3F, CF 3 ), -120.5 (2F, CF 2 ), -122.6 (6F, 2 CF 2 ), -122.8 (4F, 2 CF 2 ), -123.6 (2F, CF 2 ), -124.2 (2F, CF 2 ), -126.9 (2F, CF 2 ). Proton NMR (400 MHz) 1,1,-dihydroperfluorotetradecyl-2-bromo-2-methyl-propanoate 1 H NMR (CDCl 3 ) 298 K, δ (ppm from TMS): 4.45 (s, 2H), 1.90 (s, 6H); Fluorine NMR in CDCl 3 at 298 K, δ (ppm from 2,2,2-trifluoroethanol: -81.5 (3F, CF 3 ), -120.1 (2F, CF 2 ), -122.4 (12F, 6 CF 2 ), -122.7 (4F, 2 CF 2 ), -123.5 (2F, CF 2 ), -124.0 (2F, CF 2 ), -126.9 (2F, CF 2 ). 1.2.3 Synthesis of Perfluoro End-functionalized Polystyrene Perfluoro End-functionalized polymers were synthesized by ATRP using CuBr/N,N,N ’ ,N ’’ ,N ’’’ -pentamethyldiethylenetriamine (PMDETA) as catalyst and perfluorocarbon initiators with [Initiator]:[CuBr]:[Ligand] ratios of 1:1:2. In a typical reaction, 0.364 mmol CuBr was added into a 25 ml schlenk flask and then degased and refilled with Argon three times. 0.728 mmol PMDETA and 10 ml toluene was then added and the solution was stirrer for 20 min to form Copper-Ligand complex. Styrene was then added to the flask and the flask was evacuated and refilled with Argon via freeze-thaw cycles. After the addition of 0.364 mmol initiator the flask was placed into oil bath at 90 °C. The reaction was stopped at about 40% conversion of styrene via exposure to air and dilution with THF. The solution was then passed through a silica column to remove the copper complex and concentrated via rotary evaporation. Then the polymer was precipitated in methanol and the final product was dried at 60 °C under vacuum for 1 day. ! 9! 1.2.4 Sample Preparation Thin film samples for Nanoindentation (≈ 5 µm) was prepared by drop casting method. Polymer samples were dissolved in toluene (60 mg polymer/ml toluene) and then filtered by 0.5 µm Teflon filter. The solution was then casted onto clean silicon wafers (about 0.4x0.4 in. square). Samples were dried at 110 °C for 24 hours in vacuum oven to evaporate the solvent completely. The samples are then cooled down slowly to the room temperature. The annealing and cooling temperatures were crucial since for instance, fast cooling rates led to micro crack formation. (Figure 1.1) The silicon wafers were then glued to the sample holder. (a) (b) Figure 1.1. Optical Microscopy (OM) image of Polystyrene film on silicon wafer after heating at 110 °C and then a)rapid cooling b) slow cooling It is notable that the sample preparation is the key step for accurate and precise nanoindentation measurements. For instance, if the polymer film is too thin, the modulus of substrate would affect the results. (Figure 1.2) ! 10! Figure 1.2a shows how storage modulus increased with increasing displacement since PS film was too thin and the measurement was affected due to the storage modulus of silicon wafer substrate. According to the rule of thumbs, the measurement depth should not be more than 1/10 of the sample thickness. Figure 1.2b shows the storage modulus of 5 µm PS film is independent from tip displacement after about 50 nm depth. (a) (b) Figure 1.2. E’ of a) 100 nm b) 5 µm PS film 1.2.5 Characterization Molecular Weights were determined by Size Exclusion Chromatography (SEC) performed using Shimadzu LC-20AT pump, an RID-10A refractive index detector. THF is used as the elution solvent at a 1ml/min flow rate using PS standards. The 1 H and 19 F NMR spectra were recorded by Varian Mercury 2-Channel NMR Spectrometer 400 MHz. 1.2.5.1 Nanoindentation Measurement Design Nanoindentation measurements were carried out by Hysitron Triboindenter equipped with Nano DMA with a three-sided Berkovich tip (radius ≈ 100 nm) at 25 °C. During dynamic test, a sinusoidal load was applied to the sample with the nano probe and displacement is measured. Indent locations were carefully selected using optical microscopy 0! 5! 10! 15! 0! 5! 10! 15! 20! E’#(Gpa)# Displacement#(nm)# 0! 2! 4! 6! 8! 10! 12! 20! 220! 420! 620! 820! E’#(Gpa)# Displacement#(nm)# ! 11! to ensure the surface was smooth enough for testing and the indents are made in the sample. The distance between each indent was chosen as 50 µm. It is notable that indents made too close to each other could result in measurement failures due to crack formation. (Figure 1.3) Figure 1.3. Optical Microscopy image of indents on PS film The maximum indentation loads were adjusted to maintain comparable displacements into the surface. For ramp load measurements, DC load was increased from 75 µN to 877 µN at 45 Hz. For frequency sweep measurements, amplitude of AC and DC loads were set to 2.05 µm and 850 µm, respectively and these static and dynamic loads were kept constant throughout the measurements at each frequency. Measurements were recorded with a maximum displacement into the surface of ~350 nm at frequency range of 10-100 Hz. Each data point is an average of 100 cycles. 3X4 test were indented from randomly selected regions. ! 12! 1.2.6 Molecular Weights of Perfluorocarbon Functionalized Polymers Table 1.1 shows the number average, weight average molecular weights and polydispersity index of perfluorocarbon end-functionalized polymers studied by nanoindentation. These values were measured by SEC using PS as a standard. Table 1.1. Mechanical properties of RF end-functionalized PS synthesized using ATRP methods. Polymer M n Mw/Mn a Rfcontent (%wt) b E’ (Gpa) c E’’ (Gpa) c PS-10K 9.3 1.07 - 4.8 0.39 PS-C7-10K 8.4 1.15 4.4 5.3 0.38 PS-C10-10K 8.4 1.13 6.2 6.7 0.61 PS-15K 13.8 1.05 - 4.8 0.42 PS-C13-15K 13.1 1.10 5.1 8.9 1.46 PS-30K 28.3 1.06 - 5.1 0.42 PS-C13-30K 28.3 1.30 2.4 5.2 0.38 a. Determined by SEC using polystyrene standards. b. Calculated using number average MW. c. Grand averages assumed to be frequency independent between10-100Hz The samples are labeled giving both RF carbon numbers preceded by “C” followed by PS in kilo Daltons. Hence, PSC7-10K indicates a PS MW of around 10k having a C 7 F 15 end group connected to the PS chain through a single methylene group. 61 1.3 Nanoindentation Studies The mechanical characteristics of C 7 F 15 , C 10 F 21 and C 13 F 27 end-functionalized polymers at ambient temperatures by NI were determined. This study was predicated upon the known tendency of perfluorocarbon end groups to associate even when attached to different vinyl polymers such as polystyrene and poly(n-butylmethacrylate) (PBMA). 61 Hence it would appear that this association should also take place in a RF functionalized polystyrene or in other RF functionalized polymers. This association should have ramifications as the association of the very long C 13 F 27 groups in the solid state is bound to be strong and is expected to affect its ! 13! properties. Hence a NI study aimed at probing the heterogeneity of the PFTD functionalized PS and corresponding thin film properties was undertaken. Such a study is of scientific and technological value due the wide range of PS applications. 64-77 For instance, for low MW polymers the solid state is inherently heterogeneous due to the greater mobility of polymer (i.e. PS) end groups. Hence domains must be present where the concentration of polymer end groups in a given volume fluctuates depending on the sampled volume. The magnitude of the fluctuations is expected to increase with smaller volumes. For the case when the end groups bear strongly associating RF groups, the effects on the heterogeneity should also be present and may be even more pronounced, especially as the mobility of the end groups should now be subject to restrictions. As the test was performed with a maximum displacement of ≈ 350 nm, and considering the indented area and the volume this resulted in micro scale rather than nano scale properties clearly involving a large number of polymers chains. Due to the high strength and stiffness of PS, time dependent pile-up effects around the indenter affecting the contact area calculation can be neglected. 109 This assumption was verified by using the reference frequency technique that was already successfully applied to soft tissue samples. 108 NI measurements were carried out on PS films with thicknesses of 5 µm deposited on silicon wafers at varying frequencies. Parameters such as DC load, NI depth, sample thickness, region of the indents were determined in a preliminary set of experiments. The use of NI measurements on polymers having glass transition temperatures below or near the ambient temperature is challenging due to “pile up” effects. It is known that NI depth, presence of microcracks due to annealing conditions during film preparation, film thickness, distance between indents affect the measurements. Hence, tuning the instrument with PS film ! 14! was essential. Before probing any RF end group effects the (MWD) PS sample had a MW of about 15K and a narrow molecular weight distribution both being similar to that of the PFTD-functionalized PS. Storage moduli (E’) of PS homopolymer were first measured as a standard and the results are shown in a Minitab X-bar chart (Figure 1.4-1.5). X-bar and S charts are used for subgroup (size of two or more) to monitor measurement stability, repeatability and predictability of process measurements. X-bar shows the change in the mean/average over a defined variable (frequency in our case) whereas S chart shows the variability of the measurements around the mean. In our experiments we used these charts to have better idea of the effect of frequency and spatial separation of NI measurements on the storage moduli. An X-bar chart shows three features: a centerline, an upper critical limit and lower critical limit. Upper critical limits (UCL) and lower critical limit (LCL) show approximately 3 standard deviations (σ) from the center line. Each point on the X-bar chart represents the mean value of all indents made in different regions of the polymer film at each frequency. Here, ! represents the overall (grand) mean of all indents across the 10-100 Hz frequency range or across tip displacement range. As shown in Figure 1.4 the storage moduli at 45 Hz. decreased significantly with increasing depth (from 5.93 Gpa to 4.98 Gpa). After about 196 nm, storage moduli became stable at around 4.98 Gpa. The variation of storage moduli between indents was also much higher at shallower depths. The average modulus for polystyrene with the grand mean of E’ of all indents between 10-100 Hz was calculated as 4.803 Gpa (Figure 1.5). It was found to be about 4.98 Gpa from ramp load experiments using displacements between 250-320 nm. Average moduli we found for PS standard and the decrease in moduli with displacement ! 15! therein are consistent with the ones reported in the literature 116 . It is noted that these NI measurements were dynamic rather than static. Since polymers have time dependent properties, dynamic NI is more useful than static NI to determine storage and loss moduli with respect to frequency. Figure 1.4. X-Bar chart for PS15K showing the average E’ of various indents from different regions of the polymer as a function of displacement at 45 Hz at ambient temperature As shown in Figure 1.5 the average E’ for PS15K as a function of NI frequency from different regions of the polymer are within UCL and LCL with only random variations being observed. Thus, the storage moduli did not change significantly between 10-100 Hz. The typical trend of 3 individual indents from randomly selected regions was shown in Figure 1.6. 327 312 294 273 256 236 214 190 168 144 122 97 61 6.00 5.75 5.50 5.25 5.00 Displacement (nm) Mean E' (GPa) _ _ X=5.204 UC L=5.545 LC L=4.862 327 312 294 273 256 236 214 190 168 144 122 97 61 0.48 0.36 0.24 0.12 0.00 Displacement (nm) StDev E' (GPa) _ S=0.2653 UC L=0.5226 LC L=0.0081 1 1 1 1 Xbar-S Chart ! 16! Figure 1.5. X-Bar chart showing the average E’ of various indents from different regions of the polymer at different frequencies for PS15K Polymer at ambient temperature As shown in Figure 1.7 the standard deviation (σ) of E’ between 10 and 100 Hz is subject to modest fluctuations. For example at 10 Hz, the standard deviation of all indents from different regions of polymer film was 0.4 GPa, which represents some variation at different regions of the polymer at the same frequency. The average standard deviation between 10-100 Hz was 0.310 or about 6% of the average storage modulus. This could be due to irregularities in packing of amorphous PS given that measurements were carried at the µ scale. 100 90 80 70 60 50 40 30 20 10 5.2 5.0 4.8 4.6 4.4 Frequency (Hz) E'(GPa) _ _ X=4.8026 UC L=5.1976 LC L=4.4076 100 90 80 70 60 50 40 30 20 10 1.6 1.2 0.8 0.4 0.0 Frequency (Hz) Sample Range _ R=0.817 UC L=1.638 LC L=0 Xbar-R Chart of PS15K 100 90 80 70 60 50 40 30 20 10 6 5 4 Frequency (Hz) E'(GPa) _ X=4.757 UC L=5.902 LC L=3.611 100 90 80 70 60 50 40 30 20 10 1.6 1.2 0.8 0.4 0.0 Frequency (Hz) Mov ing Range __ MR=0.431 UC L=1.407 LC L=0 I-MR Chart of Indent 1 ! 17! Figure 1.6. E’ of individual indents from 3 randomly selected regions of the polymer at different frequencies for PS15K Polymer at ambient temperature Figure 1.7. S chart showing Standard deviation, σ, of E’ under conditions of Figure 1.5 100 90 80 70 60 50 40 30 20 10 5.6 5.2 4.8 4.4 4.0 Frequency (Hz) E'(GPa) _ X=4.716 UC L=5.432 LC L=3.999 100 90 80 70 60 50 40 30 20 10 0.8 0.6 0.4 0.2 0.0 Frequency (Hz) Mov ing Range __ MR=0.2694 UC L=0.8804 LC L=0 I-MR Chart of Indent 3 100 90 80 70 60 50 40 30 20 10 6 5 4 Frequency (Hz) E'(GPa) _ X=4.724 UC L=5.867 LC L=3.581 100 90 80 70 60 50 40 30 20 10 1.6 1.2 0.8 0.4 0.0 Frequency (Hz) Mov ing Range __ MR=0.430 UC L=1.404 LC L=0 I-MR Chart of Indent 4 100 90 80 70 60 50 40 30 20 10 0.6 0.5 0.4 0.3 0.2 0.1 0.0 Frequency(Hz) Sample StDev _ S=0.3101 UCL=0.6108 LCL=0.0094 S Chart of PS15K ! 18! Figure 1.8. X-Bar chart showing the average E’ of various indents from different regions of the polymer at different frequencies for PSC1315K Polymer at ambient temperature Figure 1.9. S chart showing the average E’ of various indents from different regions of the polymer at different frequencies for PSC1315K Polymer at ambient temperature. Interestingly, the corresponding PFTD end-functionalized isobaric PS having a 5.1 wt percent RF content shows large increases (85%) in storage moduli having the grand mean of E’ of about 8.9 GPa compared with the PS homopolymer (Figure 1.8). Furthermore, the standard deviation PFTD functionalized PS was about 1.6 GPa which is about 5 times of PS homopolymer (0.3 GPa)(Figures 1.9). This is not unexpected assuming the presence of RF domains that should effect the chain segments adjoining the RF domains. The small sampling 100 90 80 70 60 50 40 30 20 10 11 10 9 8 7 Frequency (Hz) E'(GPa) _ _ X=8.862 UC L=10.979 LC L=6.745 100 90 80 70 60 50 40 30 20 10 8 6 4 2 0 Frequency (Hz) Sample Range _ R=4.380 UC L=8.778 LC L=0 Xbar-R Chart of PSC1315K 100 90 80 70 60 50 40 30 20 10 3.5 3.0 2.5 2.0 1.5 1.0 0.5 0.0 Frequency(Hz) Sample StDev _ S=1.604 UCL=3.159 LCL=0.049 S Chart of PSC1315K ! 19! size (µ size) is consistent with these fluctuations. The typical trend of 3 individual indents from randomly selected regions was shown in Figure 1.10. Figure 1.10. E’ of individual indents from 3 randomly selected regions of the polymer at different frequencies for PSC13-15K Polymer at ambient temperature The increase in storage and loss moduli for 15K C 13 F 27 PS polymers is about 85 and 300 percent, respectively (Figure 1.11-1.12). The dramatic increases are due to a modification of 100 90 80 70 60 50 40 30 20 10 15 10 5 Frequency (Hz) E'(GPa) _ X=9.50 UC L=15.93 LC L=3.06 100 90 80 70 60 50 40 30 20 10 8 6 4 2 0 Frequency (Hz) Mov ing Range __ MR=2.420 UC L=7.906 LC L=0 I-MR Chart of Indent 4 100 90 80 70 60 50 40 30 20 10 12.5 10.0 7.5 5.0 Frequency (Hz) E'(GPa) _ X=8.853 UC L=12.365 LC L=5.341 100 90 80 70 60 50 40 30 20 10 4 3 2 1 0 Frequency (Hz) Mov ing Range __ MR=1.321 UC L=4.315 LC L=0 I-MR Chart of Indent 3 100 90 80 70 60 50 40 30 20 10 10 8 6 4 Frequency (Hz) E'(GPa) _ X=7.593 UC L=10.522 LC L=4.665 100 90 80 70 60 50 40 30 20 10 4 3 2 1 0 Frequency (Hz) Mov ing Range __ MR=1.101 UC L=3.597 LC L=0 I-MR Chart of Indent 6 ! 20! polymer structure not an additive and, to our knowledge, appear unprecedented. Stiffness improvement was also achieved by the incorporation of carbon nanotubes into various polymer matrices. 118-120 However, the storage moduli of these increases were larger than that of the loss moduli. 118-119 On the other hand in the present case the end functionalization of PFTD of the PS- 15K showed 300 percent increases in loss modulus that were much higher compared to the increases in storage modulus. This appears to be novel. Figure 1.11. Storage Moduli of PS-15, PSC13-15K, PS-30K, PSC13-30K Polymers by NI at 25 °C. The thickness of the films were around 5 µm and the NI depth was ≈300 nm. For higher MW (30K) PFTD functionalized PS having lower RF content (2.1 wt%) NI data showed more modest (2%) increases (Figures 1.11-1.12). Smaller effects seen for the 30K polymers at ambient temperatures are attributable to the lower RF concentrations. These effects have also been seen using DSC and optical transmission data. 0! 2! 4! 6! 8! 10! 12! 0! 20! 40! 60! 80! 100! E'(GPa)# Frequency(Hz)# PSC13!30K! PS!30K! PSC13!15K! PS!15K! ! 21! Figure 1.12. Loss Moduli of PS-15K and PSC13-15K, PS-30K, PSC13-30K Polymers by NI at 25 °C. The thickness of the films were around 5 µm and the NI depth was ≈350 nm. The PFTDs groups are expected to be stiff and helical with widths on the order of 0.57 nm based on studies of self assembled partially fluorinated monolayers, 121,122 NMR, 123 and computational studies. 124 Hence, the increased inter-polymer interactions may be due to the formation of PFTD aggregates through the formation of RF crystalline or semi-crystalline “micro-clusters” as discussed below. The increased inter-polymer interactions account for the increased storage moduli at least at the frequencies examined. The enhanced loss moduli indicate energy absorption is proposed to occur through the disruption but rapid reformation of the PFTD helices or clusters thereof. 125 Such a model would predict increases in the elastic and loss moduli persisting at the frequencies we studied. 0! 1! 2! 3! 4! 5! 6! 0! 20! 40! 60! 80! 100! E''#(GPa)# Frequency#(Hz)# PSC13!30K! PS!30K! PSC13!15K! PS!15K! ! 22! Scheme 1.4. Proposed association of PFTD end-functionalized PS (Low MW 10K-15K) in bulk Given the well-known tendencies of large perfluorocarbons to associate in different media and their mesogenic properties it would seem reasonable to consider various modes of RF association. A tentative structure represented in two dimensions that accounts for at least some of the data shows large laterally bound aggregates of PFTD and the smaller C 10 F 21 groups (Scheme-1.4). In this structure the PS chains alternate in “opposite” directions in order to optimize polymer packing. Thus, as the diameter of the presumably helical RF groups is on the order of 0.57 nm, alternating arrangements of Scheme 1.4 would separate the PS chains by about 2 PFTD widths or about 1.0—1.2 nm. Given the small sizes of the lower 15K chains (order of about 1.5-2 nm) such arrangements may be feasible in the solid state where PS chain interpenetration prevails. Smaller effects seen for the 30K polymers at ambient temperatures are plausibly due to the lower RF concentrations as well as the smaller effects given the nature of the RF end-groups. These effects have also been seen using DSC and optical transmission data. The above is also consistent with the much higher melting points (mp.s) of the perfluorotridecane (~100 0 C) compared with that of perfluorodecane (~ 40 0 C) and that of !! !! !! !! !! ! 23! perfluoroheptane (-13 0 C). 30 These increases average about 23 0 C /CF 2 unit and are roughly twice that seen for the corresponding hydrocarbons (~14 0 C/CH 2 ) at comparable carbon numbers (C-7 to C14). In addition the increase in mp.s is particularly large for C 10 F 22 vs C 9 F 20 of 55 0 C. This suggests the presence of incipient, conceivably 13/6 or 15/7 –like, helical conformations for the longer RF groups that have been well documented for Teflon. 31 This comparison, however, is not rigorous as the melting point represents an equilibrium unlike the dynamic (kinetic) behavior seen in the large loss moduli of the PS-C13-15K sample. Finally we point out that the low MW RF-PS polymers would be relatively simple to synthesize through conventional polymerization using radical initiators bearing an RF group. This would be especially the case in the presence of a chain transfer agent in order to prevent termination through radical combination a process that should produce a telechelic RF terminated PS. In addition the RF-PS polymers are expected to display a relatively high fraction of RF groups at the surface and this would in turn affect surface properties plausibly giving super-hydrophobic surfaces. This is being investigated. 1.3.1 Effect of shorter RF groups on the viscoelastic properties of PS by Nanoindentation Figure 1.13. X-Bar chart showing the average E’ of various indents from different regions of the polymer at different frequencies for PSC710K Polymer at ambient temperature 100 90 80 70 60 50 40 30 20 10 6 5 4 Frequency (Hz) E'(GPa) _ _ X=5.298 UC L=6.462 LC L=4.134 100 90 80 70 60 50 40 30 20 10 4 3 2 1 0 Frequency (Hz) Sample Range _ R=2.018 UC L=4.268 LC L=0 Xbar-R Chart of C710K ! 24! Figure 1.13 shows the average E’ of PSC710K polymer by various indents across the polymer thin film deposited on silicon wafers at different frequencies. X bar represents the mean of indents at each frequency whereas X bar-bar shows the mean of group of indents across the 10-100 Hz frequency range. The grand average E’ of PSC710K between 10-100 Hz is calculated to be 5.3 GPa. The typical trend of 3 individual indents from randomly selected regions was shown in Figure 1.14. Figure 1.14. E’ of individual indents from 3 randomly selected regions of the polymer film at different frequencies for PSC7-10K Polymer at ambient temperature 100 90 80 70 60 50 40 30 20 10 6.0 5.5 5.0 4.5 Frequency (Hz) E'(GPa) _ X=5.340 UC L=6.217 LC L=4.463 100 90 80 70 60 50 40 30 20 10 1.00 0.75 0.50 0.25 0.00 Frequency (Hz) Mov ing Range __ MR=0.330 UC L=1.077 LC L=0 I-MR Chart of Indent 2 100 90 80 70 60 50 40 30 20 10 6.0 5.5 5.0 4.5 4.0 Frequency (Hz) E'(GPa) _ X=5.024 UC L=6.074 LC L=3.974 10 9 8 7 6 5 4 3 2 1 1.2 0.9 0.6 0.3 0.0 Frequency (Hz) Mov ing Range __ MR=0.395 UC L=1.290 LC L=0 I-MR Chart of Indent 5 100 90 80 70 60 50 40 30 20 10 7.0 6.5 6.0 5.5 5.0 Frequency (Hz) E'(GPa) _ X=6.064 UC L=7.044 LC L=5.083 100 90 80 70 60 50 40 30 20 10 1.2 0.9 0.6 0.3 0.0 Frequency (Hz) Mov ing Range __ MR=0.369 UC L=1.204 LC L=0 I-MR Chart of Indent 4 ! 25! Difference = µ (PS10K) – µ (PSC710K) Estimate for difference: -0.5170 95% CI for difference: (-0.7011, -0.3330) T-Test of difference = 0 (vs ≠): T-Value = -5.93 P-Value = 0.000 Figure 1.15. T-test analysis on PSC7-10K and PS Compared to PS homopolymer, PSC7-10K showed higher average storage modulus at each frequency level (Figure 1.5, Figure 1.13). Since the effect is relatively small and variation exists between different indents, we ran a 2-sample t test to determine if the storage modulus of PSC7 is statistically different than the storage modulus of PS10K homopolymer. The test showed p value was 0.0 <0.05 indicating the average storage modulus of those polymers were statistically different (Figure 1.15). Figure 1.16. S chart showing the average standard deviation of E’ of various indents from different regions of the polymer at different frequencies for PSC7-10K Polymer at ambient temperature 100 90 80 70 60 50 40 30 20 10 1.8 1.6 1.4 1.2 1.0 0.8 0.6 0.4 0.2 0.0 Frequency(Hz) Sample StDev _ S=0.789 UCL=1.648 LCL=0 S Chart of PSC710K ! 26! Figure 1.16 shows the standard deviation of E’ of PSC7-10K between 10-100 Hz. The average standard deviation was 0.7 GPa for the given frequency range which was higher than average standard deviation of storage modulus of PS homopolymer (0.3 GPa). Given that nanoindenter measures storage moduli at submicron level, this is possibly due to the formation of small perfluorocarbon domains and more heterogeneous structure of PSC7 compared with PS homopolymer. Figure 1.17. X-Bar chart showing the average E’ of various indents from different regions of the polymer at different frequencies for PSC10-10K Polymer at ambient temperature Figure 1.17 shows the average E’ of various indents from different regions of the polymer at different frequencies for PSC10-10K Polymer at ambient temperature. The grand average was calculated as 6.7 GPa with only random variation being observed. (Figure 1.17) The typical trend of 3 individual indents from randomly selected regions was shown in Figure 1.18. Figure 1.19 shows the standard deviation of storage moduli between 10-100 Hz. The average standard deviation was similar to the average standard deviation of PSC7. 100 90 80 70 60 50 40 30 20 10 8.0 7.5 7.0 6.5 6.0 Frequency (Hz) E'(GPa) _ _ X=6.713 UC L=7.735 LC L=5.690 100 90 80 70 60 50 40 30 20 10 4 3 2 1 0 Frequency (Hz) Sample Range _ R=1.773 UC L=3.749 LC L=0 Xbar-R Chart of C1010K ! 27! Figure 1.18. E’ of individual indents from 3 randomly selected regions of the polymer at different frequencies for PSC10-10K Polymer at ambient temperature 100 90 80 70 60 50 40 30 20 10 9 8 7 6 5 Frequency (Hz) E'(GPa) _ X=7.041 UC L=8.785 LC L=5.297 100 90 80 70 60 50 40 30 20 10 2.0 1.5 1.0 0.5 0.0 Frequency (Hz) Mov ing Range __ MR=0.656 UC L=2.143 LC L=0 I-MR Chart of Indent 4 100 90 80 70 60 50 40 30 20 10 7 6 5 Frequency (Hz) E'(GPa) _ X=5.847 UC L=7.222 LC L=4.471 100 90 80 70 60 50 40 30 20 10 1.6 1.2 0.8 0.4 0.0 Frequency (Hz) Mov ing Range __ MR=0.517 UC L=1.689 LC L=0 I-MR Chart of Indent 1 100 90 80 70 60 50 40 30 20 10 9 8 7 6 5 Frequency (Hz) E'(GPa) _ X=6.808 UC L=9.187 LC L=4.430 100 90 80 70 60 50 40 30 20 10 3 2 1 0 Frequency (Hz) Mov ing Range __ MR=0.894 UC L=2.921 LC L=0 I-MR Chart of Indent 2 ! 28! Figure 1.19. S chart showing the average standard deviation of E’ of various indents from different regions of the polymer at different frequencies for PSC10-10K Polymer at ambient temperature Storage and loss moduli of PSC7-10K, PSC10-10K and PS-10K were summarized in Figure 1.20 and 1.21. The elastic moduli of the shorter C 7 F 15 groups were smaller than that of the perfluorodecyl (C 10 F 21 ) groups. For instance, E’ of C 7 F 15 functionalized PS was about 5.3 GPa whereas E’ of C 10 F 21 functionalized PS between 10-100 Hz was 6.7 GPa (Figure 1.13-1.17). Therefore, E’ of PS end functionalized with C 7 F 15 and C 10 F 21 were 10% and 42% higher than PS-10K homopolymer, respectively. In addition, the standard deviation of both PSC7-10K and PSC10-10K between 10-100 Hz were about 0.7 GPa, which is two times higher than PS-10K homopolymer. Loss moduli of the polymers show a similar trend. Average E’’ of C 7 F 15 functionalized PS was similar to PS homopolymer at about 0.38 GPa, whereas average E’’ of C 10 F 21 functionalized PS was 0.61 GPa between 10-100 Hz. (Figure 1.21). Box plots of E’ and E’’ of polymers were shown in Figure 1.22 and 1.23. 100 90 80 70 60 50 40 30 20 10 1.6 1.4 1.2 1.0 0.8 0.6 0.4 0.2 0.0 Frequency(Hz) Sample StDev _ S=0.693 UCL=1.448 LCL=0 S Chart of PSC1010K ! 29! Finally the smaller effects noted for the C7 and C10 polymers are consistent with this model with the smaller RF groups being associated in more weakly bound aggregates. Their storage and loss moduli were found to be much smaller than that of the PFTD groups. (Table 1.1) An increase in E’ with RF length compared to the PS homopolymer is noted for PS-C7-10K and PS- C10-10K but is much smaller compared with the PFTD group (Table-1). Similarly, standard deviation of perfluoro polymers increased from 0.3 with increased perfluorocarbon length up to 1.6 with the introduction of PFTD groups indicating increasing heterogeneity in the polymer. The trends in the loss moduli, E’’, followed a similar trend reaching the highest value with PFTD groups. Figure 1.20. Storage Moduli of PS-10K and PSC7-10K and PSC10-10K Polymers by NI at 25 °C. 0! 1! 2! 3! 4! 5! 6! 7! 8! 9! 0! 20! 40! 60! 80! 100! E'#(GPa)# Frequency#(Hz)# PSC710K! PSC1010K! PS!10K! ! 30! Figure 1.21 Loss Moduli of PS-10K and PSC7-10K and PSC10-10K Polymers by NI at 25 °C Figure 1.22. Box Plot representation of average E’ of PS-10K, PSC7-10K and PSC10-10K 0! 0.5! 1! 1.5! 2! 2.5! 0! 20! 40! 60! 80! 100! E''#(GPa)# Frequency#(Hz)# PSC710K!! PSC1010K! PS10K! PSC1010K PSC710K PS10K 7.0 6.5 6.0 5.5 5.0 4.5 E' (GPa) Boxplot of PS10K, PSC710K, PSC1010K ! 31! Figure 1.23. Box Plot representation of average E’’ of PS-10K, PSC7-10K and PSC10-10K 1.4 Conclusions The effects of the long PFTD groups on the NI measurements are quite clear. Both storage and loss moduli of PSC13 15K increased about 85% and 300 %, respectively. Smaller RF groups were also tested, however the increase in both E’ and E’’ was much smaller compared to PFTD groups. Storage moduli of PS did not show a dependent behavior on frequency, which is quite expected due to measurement temperatures much lower than Tg of PS. However, it is noteworthy that storage moduli of PSC13 had a slight increase with increasing frequency at ambient temperature. These effects were much smaller at higher MWs. PS10-10K PSC7-10K PS 10K 1.2 1.0 0.8 0.6 0.4 0.2 E'' (GPa) Boxplot of PS 10K, PSC7-10K, PS10-10K ! 32! 1.5 References 1. Nicolas, M.; Guittard, F.; Géribaldi, S. Stable superhydrophobic and lipophobic conjugated polymers films. Langmuir 2006, 22, 3081-3088. 2. 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C.; McHargue, C. J., The deformation behavior of ceramic crystals subjected to very low load (nano) indentations. Journal of Materials Research 1992, 7(02), 450-473. 102. Woirgard, J.; Tromas, C.; Girard, J. C.; Audurier, V., Study of the mechanical properties of ceramic materials by the nanoindentation technique. Journal of the European Ceramic Society 1998, 18(15), 2297-2305. ! 45! 103. Mintairov, A. M.; Sun, K., Merz; J. L., Li, C.; Vlasov, A. S.; Vinokurov, D. A.; Oktyabrsky, S., Nanoindentation and near-field spectroscopy of single semiconductor quantum dots. Physical Review B 2004, 69(15), 155306. 104. Kucheyev, S. O.; Bradby, J. E.; Williams, J. S.; Jagadish, C.; Toth, M.; Phillips, M. R.; Swain, M. V., Nanoindentation of epitaxial GaN films. Applied Physics Letters 2000, 77(21), 3373-3375. 105. Lund, A. C.; Hodge, A. M.; Schuh, C. A., Incipient plasticity during nanoindentation at elevated temperatures. Appl. Phys. Lett. 2004, 85 (8), 1362-1364. 106. Schuh, C.; Mason, J.; Lund, A., Quantitative insight into dislocation nucleation from high-temperature nanoindentation experiments. Nature Materials 2005, 4 (8), 617-621. 107. Ebenstein, D. M.; Kuo, A.; Rodrigo, J. J.; Reddi, A. H.; Ries, M.; Pruitt, L., A nanoindentation technique for functional evaluation of cartilage repair tissue. J. Mater. Res. 2004, 19 (01), 273-281. 108. Franke, O.; Durst, K.; Maier, V.; Göken, M.; Birkholz, T.; Schneider, H.; Hennig, F.; Gelse, K., Mechanical properties of hyaline and repair cartilage studied by nanoindentation. Acta Biomaterialia 2007, 3 (6), 873-881. 109. Franke, O.; Göken, M.; Hodge, A. M., The nanoindentation of soft tissue: current and developing approaches. JOM 2008, 60 (6), 49-53. 110. Franke, O.; Göken, M.; Meyers, M.; Durst, K.; Hodge, A., Dynamic nanoindentation of articular porcine cartilage. Materials Science and Engineering: C 2011, 31 (4), 789- 795. ! 46! 111. Iqbal, T.; Briscoe, B. J.; Yasin, S.; Luckham, P. F., Nanoindentation response of poly (ether ether ketone) surfaces—A semicrystalline bimodal behavior. J. Appl. Polym. Sci. 2013, 130 (6), 4401-4409. 112. Oyen, M. L., Sensitivity of polymer nanoindentation creep measurements to experimental variables. Acta Mater. 2007, 55 (11), 3633-3639. 113. VanLandingham, M. R. Review of instrumented indentation; DTIC Document: 2003. 114. VanLandingham, M. R.; Chang, N. K.; Drzal, P.; White, C.; Chang, S. H., Viscoelastic characterization of polymers using instrumented indentation. I. Quasi-static testing. J. Polym. Sci., Part B: Polym. Phys. 2005, 43 (14), 1794-1811. 115. Li, Min, C. Barry Carter, and William W. Gerberich. Nanoindentation Measurements of Mechanical Properties of Polystyrene Thin Films. In MRS Proceedings, Cambridge University Press, 2000, 649, 7-21. 116. Chung, P. C., Glynos, E., & Green, P. F. The Elastic Mechanical Response of Supported Thin Polymer Films. Langmuir 2014, 30(50), 15200-15205. 117. NanoDMA I and II User Manual, Hysitron, 2009 118. McNally, T.; Pötschke, P.; Halley, P.; Murphy, M.; Martin, D.; Bell, S. E.; Quinn, J. P., Polyethylene multiwalled carbon nanotube composites. Polymer 2005,46(19), 8222- 8232 119. Song, Y. S.; Youn, J. R., Influence of dispersion states of carbon nanotubes on physical properties of epoxy nanocomposites. Carbon 2005, 43(7), 1378-1385 120. Gojny, F.; Wichmann, M.; Köpke, U.; Fiedler, B.; Schulte, K., Carbon nanotube- reinforced epoxy-composites: enhanced stiffness and fracture toughness at low nanotube content. Compos. Sci. Technol. 2004, 64 (15), 2363-2371. ! 47! 121. Frey, S.; Heister, K.; Zharnikov, M.; Grunze, M.; Tamada, K.; Colorado, R.; Graupe, M.; Shmakova, O.; Lee, T., Structure of self‐assembled monolayers of semifluorinated alkanethiols on gold and silver substrates. Isr. J. Chem. 2000, 40 (2), 81-97. 122. Lu, H.; Zeysing, D.; Kind, M.; Terfort, A.; Zharnikov, M., Structure of Self- Assembled Monolayers of Partially Fluorinated Alkanethiols with a Fluorocarbon Part of Variable Length on Gold Substrate. The Journal of Physical Chemistry C 2013, 117 (37), 18967-18979. 123. Ellis, D. A.; Denkenberger, K. A.; Burrow, T. E.; Mabury, S. A., The use of 19F NMR to interpret the structural properties of perfluorocarboxylate acids: A possible correlation with their environmental disposition. The Journal of Physical Chemistry A 2004, 108 (46), 10099-10106. 124. Liu, Z.; Goddard, J. D., Predictions of the Fluorine NMR Chemical Shifts of Perfluorinated Carboxylic Acids, C n F 2n+1 COOH (n = 6-8). J. Phys. Chem. A 2009, 113 (50), 13921-13931. 125. Xie, X.; Hogen-Esch, T. E., Copolymers of N, N-dimethylacrylamide and 2-(N- ethylperfluorooctanesulfonamido) ethyl acrylate in aqueous media and in bulk. Synthesis and properties. Macromolecules 1996, 29 (5), 1734-1745 ! 48! CHAPTER 2 Rheological Properties 2.1 Introduction The potential of perfluorocarbon end groups to influence the bulk properties of materials is less well studied but has is also of interest. For instance, the formation of homogeneous (1/1 w/w) blends of C 7 F 15 , C 10 F 21 and C 13 F 27 end-functionalized polystyrene (PS) and polybutylmethacrylates (PBMA) has been demonstrated recently. 1-3 The degree of association of RF end groups, at the same RF content and –length, was shown to be much more pronounced than that of RF pendent groups. 3 The TEM, AFM, DSC and other data indicate that the dihydroperfluorotetradecyl (PFTD=CH 2 C 13 F 27 ) compared to the corresponding CH 2 C 7 F 15 and CH 2 C 10 F 21 end groups, is particularly effective in mediating compatibilities as judged by the approximate size of the domain sizes. 2 Thus, at MW of ~15k, the domains sizes of the PFTD end functionalized PS were found to be much smaller (~ 1-3 nm) compared to that of the smaller RF end groups. 1-3 It is plausible that these RF end group effects would also be important in the bulk properties of PS and other homopolymers. For instance we have found that for C 13 F 27 end-functionalized 15k PS the storage moduli were increased nearly two-fold while the loss moduli were increased about four fold. However, C 13 F 27 end-functionalized 30k PS had similar storage modulus with PS 30k homopolymer at ambient temperature shown by nanoindentation measurements. In this chapter, we investigated the rheological behavior of C 13 F 27 end-functionalized 30k PS at high temperatures. ! 49! 2.2 Rheology Analysis of PFTD functionalized PS 30K Two narrow MW distribution RF (C 7 F 15 and C 13 F 27 ) end-functionalized PS samples and their matching isobaric PS homopolymers were synthesized using slightly different reaction solvents especially for the PFTD end functionalized PS (Table-1.1) as described in the Experimental Section. Unlike the lower MW (15K) analog, dynamic nano-indentation (NI) solid state storage (E´) and loss (E´´) moduli of the 30k PFTD –functionalized polymer (PSC13-30K) at ambient temperatures were only somewhat increased compared with the isobaric PS. However, surprisingly, the high temperatures (150-180 0 C) shear storage (G´) and loss (G´´) moduli at 1 Hz were much larger (one order of magnitude at 152 0 C) compared with the isobaric unmodified PS with the differences increasing at higher temperatures (about 30 times at 175 0 C, Figure 2.1). Similar but smaller increases (two to three fold) in the loss moduli were seen were over the same temperature interval (Figure-2.1). Figure 2.1. Shear Storage (filled) and loss moduli (empty) for PS-30K (diamond) and PSC13- 30K (triangle) polymers. The reference frequency and strain are 1Hz and 0.5 %, respectively. The increases are plausibly not due to semi-crystalline interactions but may attributable to the 1.00E+00 1.00E+01 1.00E+02 1.00E+03 1.00E+04 1.00E+05 1.00E+06 1.00E+07 150 155 160 165 170 175 180 G#(Pa)# # Temperature (°C) ! 50! phase separation of the C 13 F 27 groups into fluorous domains due to the large differences in cohesive energy densities of perfluorocarbons (11.3 -12.7 Mpa -1/2 ) and PS (17-18 Mpa -1/2 ). 4 The low frequencies used suggest the occurrence of relatively slow motions consistent with non equilibrium conditions. At a reference frequency of 0.1 Hz, slightly different results were obtained (Figure 2.2). Figure 2.2. Shear Storage Moduli for PS-30K(diamond) and PSC13-30K(triangle) polymers. The reference frequency and strain are 0.1Hz and 0.5 %, respectively. Thus, at a reference frequency of 0.1 Hz between 135- 155 0 C the storage moduli of the PSC13-30K exceed that of the PS-30K by more than an order of magnitude. In addition, the storage modulus of the PSC13-30K decreases more rapidly at 0.1 Hz compared with 1 Hz. These data seem consistent with the enlarged time frame implicit at the lower frequency. However, the loss moduli of the PSC13-30K still exceed that of the PS30K by a factor of 3-4 persisting over the entire 135-155 0 C range indicating a significant and persistent residual effect the nature of which may also be linked to the formation of RF domains (Figure 2.3). 1 10 100 1000 10000 100000 135 140 145 150 155 160 165 G' (Pa) Temperature ( 0 C) ! 51! Figure 2.3. Shear Loss Moduli for 30K PS (diamonds) and PSC13-30K polymers at frequency and strain of 0.1Hz and 0.5 percent, respectively. Figure 2.4.a. Master curves of G´and G´´ for Figure 2.4.b. Master curves of G´and G´´ PS- 30K at 180 ° C. for PSC13- 30K at 180 ° C. Both polymers exhibited pronounced viscous behavior (G’(ω) <G’’(ω )) at 180 ° C, the difference between G’ and G’’ of PS-30K being greater than that of the PSC13-30K polymers (Figure 2.4). This is largely due to the storage moduli being greater for the PSC13-30K samples as documented above. At high frequencies, G´ and G´´ of PSC13-30K tend to converge whereas 1 10 100 1000 10000 100000 135 140 145 150 155 160 165 G'' (Pa) Temperature ( 0 C) 0.1! 1! 10! 100! 1000! 10000! 100000! 0.1! 1! 10! 100! G#(Pa)# Angular(Frequency((ω)(sec 21 ( G'! G''! 1! 10! 100! 1000! 10000! 100000! 0.1! 1! 10! 100! G#(Pa)# Angular#Frequency#(ω)#sec >1 # G'! G''! ! 52! these values are quite different for PS30K indicating largely viscous behavior, consistent with the literature data. 5 Figure 2.5. Frequency dependence of G’ for PS-30K (diamond) PSC13-30K (triangles) at 180 °C. Reference strain is 0.5 %. The G´ and G´´ values of the PS30K and PSC13-30K at 180 0 C as a function of frequency are shown in in Figures 2.4a and 2.4b. Above roughly 3 Hz the storage moduli PSC1330K exceed that of PS30K by factors about 25. As shown in Figure 2.5, below 3 Hz the G’ values of PSC1330K plunge steeply to give similar to that of PS-30K. Taken together the data indicate significant association in the melt of the PFTD end-functionalized polystyrene. Moreover, PSC13-30K was more frequency dependent compared to PS-30K (Figure 2.5). Futhermore, loss moduli of PFTD-PS was higher over the entire frequency range compared with PS-30K (Figure 2.6). The effects of the long C 13 F 27 groups on the rheology and of NI measurements remains to be fully elucidated. 1 10 100 1000 10000 0 2 4 6 8 10 G' (Pa) Frequency (Hz) ! 53! Figure 2.6. Frequency dependence of G´´ for PS-30K (diamond) PSC13-30K (triangles) at 180 °C. Reference strain is 0.5 %. 2.3 Rheological Properties of C 7 F 15 -PS We investigated the effect of shorter RF groups on rheological properties of PS. Lower MW samples were tested since RF content was higher (4.4 wt%). As expected the PSC7-10K sample had much smaller effects about 50% and 80% higher storage moduli and loss moduli, respectively compared with PS-10K consistent with stronger fluorophilic interactions of the longer RF groups (Figure 2.7). 1 10 100 1000 10000 100000 0 2 4 6 8 10 G'' (Pa) Frequency (Hz) ! 54! Figure 2.7. Frequency dependence of G´ (filled) and G´´(empty) of PS-10K (square) and PSC7- 10K (spheres) at 120 °C. Reference strain is 0.5 %. 2.4 Conclusions The remarkable temperature and frequency dependent effects seen for the PFTD functionalized 30K samples at this time are only partially understood. The nature of the association, especially its small but persistent, effects on the loss moduli are plausibly due to this latter type of association. Hence association prevails at all temperatures but its expression in the storage modulus is hidden by the strongly time dependent association seen at very low frequencies. The important role of the PFTD groups is illustrated unambiguously as similar but smaller rheology effects were obtained with C 7 F 15 end functionalized PS samples. The effects documented above are remarkable as the RF content of the PSC13-30K samples is only about two percent. More detailed studies on the MW dependence of the above effects are being planned. Given the above, the above effects are expected to apply to many if not most other fluorinated polymers. 1.00E+00 1.00E+01 1.00E+02 1.00E+03 1.00E+04 1.00E+05 1.00E+06 0 2 4 6 8 10 G (Pa) Frequency (Hz) ! 55! 2.5. References 1. Shen, J.; Hogen-Esch, T. Block copolymer-like self-assembly of fluorocarbon end- functionalized polystyrene and polybutylmethacrylate. J. Am. Chem. Soc. 2008, 130, 10866- 10867. 2. Shen, J.; Piunova, V. A.; Nutt, S.; Hogen-Esch, T. E. Blends of polystyrene and poly (n-butyl methacrylate) mediated by perfluorocarbon end groups. Polymer 2013, 54, 5790-5800. 3. Shen, J. Perfluorocarbon mediated self-assembly of polymers; University of Southern California: 2009. 4. Brandrup, J.; Immergut, E. H.; Grulke, E. A.; Abe, A.; Bloch, D. R. Polymer handbook; Wiley New York: 1999; Vol. 89. 5. Onogi, S.; Masuda, T.; Kitagawa, K. Rheological properties of anionic polystyrenes. I. Dynamic viscoelasticity of narrow-distribution polystyrenes. Macromolecules 1970, 3, 109- 116. ! 56! CHAPTER 3 Thermal and Optical Properties of Perfluoroalkyl End-Functionalized Polystyrenes Using Optical Transmission and Differential Scanning Calorimetry 3.1 Introduction Low molecular weight (MW) perfluorocarbons (RFs) show weak intermolecular interactions compared with hydrocarbons of comparable chain lengths as expressed in much lower cohesive energy densities (δ ≈ 12 vs 15 Mpa 1/2 ). 1 Accordingly, RFs have both pronounced hydrophobic 2- 11 as well as lipophobic properties. 4-14 However, compared to hydrocarbons, the melting points (mp’s) of perfluorocarbons increase much faster with chain lengths. 15 For instance, while the melting points (mp’s) of octane and perfluorooctane are nearly the same, the mp’s of tetradecane and perfluorotetradecane are about 275 and 400 K respectively a difference of about 125 0 C. This indicates the occurrence of much stronger interactions for longer RFs including high MW (10 5 - 10 6 D) polytetrafluoroethylene (PTFE) that have extremely high mp’s (~ 500-621 K). This polymer has been shown to exist in four different helical crystalline states depending on temperature and pressure. 16-17 Semi-fluorinated polymers are of more recent origin and have several applications due to their unique properties that differ from the corresponding analogous hydrogen or fluorocarbon polymers. They have gained attention in antifouling, super-hydrophobic and self-cleaning coatings. 18-25 Other applications include microelectronics due to their lower dielectric constants, their self-assembly, hydrophobic and lipophobic properties and low surface free energies, all of which are useful in organic semiconductors and insulators. 26-29 They are also low surface free energies 30-34 and of interest in surface active and photoresponsive materials. 35-36 In agreement with the above, numerous polymer containing RF pendent groups have been shown to exhibit mesogenic properties. 17,37-55 The effects of association of pendent RF groups on the rheological ! 57! properties of aqueous polymer solutions 9-10 hydrogels 11,56-58 and non-aqueous polymer solutions 59 have been demonstrated. The potential of perfluorocarbon end groups to influence the bulk properties of materials is less well studied but has is also of interest. For instance, the formation of homogeneous (1/1 w/w) blends of C 7 F 15 , C 10 F 21 and C 13 F 27 end-functionalized polystyrene (PS) and polybutylmethacrylates (PBMA) has been demonstrated recently. 59-61 The degree of association of RF end groups, at the same RF content and –length, was shown to be much more pronounced than that of RF pendent groups. 61 The TEM, AFM, DSC and other data indicate that the dihydroperfluorotetradecyl (PFTD=CH 2 C 13 F 27 ) compared to the corresponding CH 2 C 7 F 15 and CH 2 C 10 F 21 end groups, is particularly effective in mediating compatibilities as judged by the approximate size of the domain sizes. 60 Thus, at MW of ~15k, the domains sizes of the PFTD end functionalized PS were found to be much smaller (~ 1-3 nm) compared to that of the smaller RF end groups. 59-61 It is plausible that these RF end group effects would also be important in the bulk properties of PS and other homopolymers. For instance we have found that for C 13 F 27 end-functionalized 15k PS the storage moduli were increased nearly two-fold while the loss moduli were increased about five fold. In addition, given the low refractive indices (Ris) of low MW perfluorocarbons and their tendencies to aggregate in virtually all other types of media optical density measurements should reflect this if the RF domains are comparable or larger than the optical wavelengths. Here we report optical transmission data of films of these PFTD and other RF end functionalized PS with MW’s between 10 and 30k Dalton that show clear evidence of the formation of PFTD groups that aggregate into large (400 nm) domains of unknown shape but ! 58! approximate sizes. In addition, DSC data indicate sizable decreases in Tg (as high as 15 0 C) of the PFTD end-functionalized PS compared to the isobaric PS. This surprising and seemingly unprecedented finding hints at a high PS mobility mediated by the chain end segments being connected to the PFTD aggregates. 3.2 UV-Visible measurements of PS and RF end-functionalized PS films Given their low polarizabilities, low MW perfluorocarbons have very low refractive indices (Ris) (1.278-1.296 for C 7 F 15 –C 11 F 23 ) compared to that of polystyrene (1.59-1.60) at optical wavelengths 62 so that significant light scattering may be expected to occur depending on the number of domains as well as their sizes and shapes. Figure 3.1. UV-Visible optical transmittance versus wavelength of the PS-10K (blue), PSC7- 10K (red), PSC10-10K(purple) and PSC13-10K (green) polymer films. The optical transmission (OT) values should be dependent on (a) the cohesive energy densities differences between PS and the RF domains and their refractive indices. 59-61 (b) the sizes of the perfluorocarbon (RF) domains sizes that depend in turn on (c) the lengths and the mobilities of the RF end groups that are expected to mediate that of PS chain ends. 0! 10! 20! 30! 40! 50! 60! 70! 80! 90! 100! 380! 440! 500! 560! 620! 680! 740! 800! Transmittance##(%)# Wavelength#(nm)# ! 59! A number of RF- end-functionalized PS samples were synthesized using ATRP methods (Table 3.1). The polydispersities (PDIs) were kept low by using high monomer concentrations and limiting monomer conversions. Table 3.1. RF end-functionalized polystyrenes synthesized using ATRP methods. a Samples Mn a Mw/Mn a RF content (wt %) b Samples Mn Mw/Mn RF content (wt %) b PS-10K 9.3 1.07 - PSC10-15K 12.8 1.08 4.1 PS-30K 28.3 1.06 - PSC10-30K 26.5 1.15 2.0 PSC7-10K 8.4 1.15 4.4 PSC13-10K 9.1 1.09 7.4 PSC7-15K 13.1 1.11 2.8 PC13-15K 13.1 1.10 5.1 PSC7-30K 28.4 1.08 1.3 PSC13-30K 28.3 1.30 2.4 PSC10-10K 8.4 1.13 6.2 a Determined by SEC using polystyrene standards. b. Calculated using weight average MW As shown in Figure 3.1, the PS-10K sample is nearly optically clear (OT ~96%) and is essentially identical to that of the 30K PS with both of these showing wavelength independent Ots between 380-800 nm. However, the PSC7-10K and the PSC10-10K films show slightly smaller Ots (90-93 and 86-90% respectively) in the 380-800 nm range. In addition, the OT- wavelength patterns of the PSC10-10K and PSC7-10K show decreases between 500 and 800 nm (Figure 3.1). These decreases are attributable to refractive index differences of the RF and PS domains as shown below. The large decreases in OT seen for the PSC13-10K and PSC13-30K samples will be discussed below. ! 60! Figure 3.2. UV-Visible Transmittance versus wavelength of 10K and 30K PSC7 polymer films on micro cover glass. The effects on OT of MW at constant RF lengths are shown in Figures 3.2 and 3.3. The transmittance of 10K-PSC7 and 30K-PSC7 polymers is nearly identical above 500 nm. (Figure 3.2) As indicated above, the decreases in OT with RF size are consistent with the presence of RF domains large enough to be effective light scattering centers given the relatively large RI difference (Δ n ≅ 0.35 see above) between PS and RF domains. As the data from Figure 3.1 are based on samples in which both MW and length of RF are varied, the formation of the putative RF domains may due to either increases in RF concentration or length. This was checked by comparing the OT data of PSC10-10K and PSC10-30K polymer films where only the RF concentrations vary especially as there are no significant differences in the OT between the 10K and 30K PS homopolymers. 0! 10! 20! 30! 40! 50! 60! 70! 80! 90! 100! 400# 450# 500# 550# 600# 650# 700# 750# 800# Transmittance#(%)# Wavelength#(nm)# PSC7!10K! PSC7!30K! ! 61! Figure 3.3. UV-Visible Transmittance versus wavelength of PSC10-10K and PSC10 -30K polymer films on micro cover glass. As shown in Figure 3.3 the OT values of the PSC10-30K film are uniformly larger than that of PSC10-10K with the differences being about 5 ± 2 percent in the 400-800 nm range. This difference is not caused by the higher MW of the former as the PS-10K and PS-30K homopolymer have essentially the same OT profiles. Furthermore the OT vs. wavelength pattern of PSC10-30K is qualitatively different from the PSC7-10K and PSC10-10K polymers in that the OT vs. wavelength profile of the PSC10-30K shows no decreases between 500 and 800 nm. This is consistent with the absence or sufficiently large aggregates compared to optical wave lengths at this particular concentration (~2 wt %). These results indicate that the differences between PSC10-10K and PSC10-30K are due to roughly three-fold higher fluorocarbon concentration. This indicates that both concentration and RF lengths are important. Hence it is reasonable to infer that complex OT vs wavelength patterns for the case of the PSC10-10K and PSC7-10K may also be due to the presence of aggregates large enough to influence light diffraction, however slightly. 0! 10! 20! 30! 40! 50! 60! 70! 80! 90! 100! 400# 450# 500# 550# 600# 650# 700# 750# 800# Transmittance#(%)# Wavelength#(nm)# PSC10!10K! PSC10!30K! ! 62! Figure 3.4. UV-Visible Transmittance versus wavelength of PSC13-10K (green), PSC13-30K (red) and PS -30K (blue). More convincingly, the PSC13-10K films were strongly opalescent with much lower Ots values at 50 to 58 percent at 380 and 800 nm, respectively (Figure 3.4). This large effect is unlikely due to RF mass concentrations as in PSC13-10K these are calculated to be only about 20 percent larger than in PSC10-10K (Table 3.1). Hence, this remarkably low OT of the PSC13-10K film is consistent with significant association of the long C 13 F 27 groups into aggregates that are quite large compared with the optical wavelengths (380-800 nm). This reinforces the proposition that the small decreases in OT seen for the smaller end groups of PSC7-10K and PSC10-10K is also be linked to RF lengths as well as concentration with the RF length being the more important for the C 13 F 27 end groups. In striking contrast with the PSC13-10K and the other RF polymers, the Ots of the PSC13- 30K are highly dependent on both MW and wavelength with Ots between 70 and 90 percent in the 380-800 nm range (Figure 3.4). Between 380 and about 600 nm the changes in OT are both larger and highly variable indicating the presence of medium sized aggregates of approximate 0! 10! 20! 30! 40! 50! 60! 70! 80! 90! 100! 380! 440! 500! 560! 620! 680! 740! 800! Transmittance#(%)# Wavelength#(nm)# ! 63! sizes in that length scale (≤ 700nm). In this case the calculated concentration of the RF concentration is roughly 2.4 wt percent where aggregates with much smaller (380-700 nm) are present compared with the PSC13-10K sample where the aggregates are plausibly much larger (≥ 700 nm) than the PSC13-30K sample. 3.3 Termogravimetric Analysis of Perfluoro End-Functionalized PS The thermal degradation of PS was studied extensively. 63-67 Likewise radical polymerization, thermal degradation of polystyrene involves radical processes including initiation, propogation, chain transfer and termination steps. 68 There are four main mechanisms of degradation process of polystyrene reported mainly as chain scission, back-biting, disproportionation and C-C scission. A somewhat similar mechanism for RF-end-functionalized PS may occur except that the presence of a perfluorocarbon chain at the end of the polymer that could retard the decomposition. In order to determine the presence of such an effect, we carried out TGA measurements on the shortest (C 7 F 15 ) and longest (C 13 F 27 ) perfluorocarbon end-functionalized polymers. Figure 3.5 shows TGA thermograms of PS homopolymer of 10K and 15K and C 13 F 27 end- functionalized PS of similar MW under air in a temperature range of 250-420 °C. The major loss at onset temperatures ranges from 280-300 °C depending on molecular weight and C 13 F 27 content. The thermal degradation of PS-10K and 15K started at similar temperatures of 280 °C. They showed similar degradation behavior until 320 °C, however after about 320 °C, PS15K showed slower degradation compared to PS-10K. The degradation of C 13 F 27 end-functionalized PS of 10K and 15K started at about 300 °C which is 20 °C higher than PS homopolymer (Figure 3.5). Although PSC13-10K had higher perfluorocarbon content (7.4 % wt) than PSC13-15K, it showed faster degradation indicating MW had a dominant effect on degradation behavior. ! 64! Figure 3.5. TGA curves of PS-10K (light blue), PS-15K (dark blue), PSC13-10K(red) and PSC13-15K(green). Figure 3.6 shows the TGA thermograms of PS-10K, PSC7-10K and PSC13-10K. The degradation of both C7 and C13 PS started about 300 °C, 20 °C higher than PS homopolymer whereas PSC7-10K and PSC13-10K showed similar degradation behavior. The end point of thermal degradation was at 400 °C for PS and 414 °C for both PSC7 and PSC13. The thermal stability of polystyrene increased by either end functionalization with C7 or C13 groups. In the literature, it has been shown that end chain functionalization could have substantial effect on the thermal stability of the polymer. 69-70 The increasing effect in thermal stability could be due to perfluorocarbon functionalization or due to the presence of bulky groups at the chain end. 0! 20! 40! 60! 80! 100! 250! 270! 290! 310! 330! 350! 370! 390! 410! Mass(%)# Temperature#(°C)# ! 65! Figure 3.6. TGA curves of PS-10K, PSC7-10K and PSC13-10K. Figure 3.7 TGA curves of PS-15K , PSC7-15K and PSC13-15K. The thermograms of PS-15K, PSC7-15K and PSC13-15K behaved somewhat similar to 10K polymers (Figure 3.7). Polymer degradation started at 280 °C for PS15K and at 300 °C for PSC7 and PSC13 of 15K. A 20 percent weight loss was observed at 345 °C and 359 °C for PS10K and 0! 10! 20! 30! 40! 50! 60! 70! 80! 90! 100! 250! 270! 290! 310! 330! 350! 370! 390! 410! Mass#(%)# Temperature#(°#C)# PSC7!10K! PS!10K! PSC13!10K! 0! 10! 20! 30! 40! 50! 60! 70! 80! 90! 100! 250! 270! 290! 310! 330! 350! 370! 390! 410! Mass#(%)# Temperature#(°#C)# PSC7!15K! PSC13!15K! PS!15K! ! 66! for both PSC7 and PSC13, respectively. The similar degradation behavior of PSC7 and PSC13 persisted until 385 °C after which PSC13-15K degraded faster. 3.4 Differential Scanning Calorimetry analysis of Perfluoro End-functionalized PS Thermal properties of perfluorocarbon (RF) end-functionalized polymers in the context of the above OT data are especially interesting for several reasons. First, the effect of RF length and content on the glass transitions of PS is expected to be significant in that the chain end RF groups are known to associate to varying degrees in virtually all media other than fluorocarbon solvents consistent with the above OT and other studies. 9-11 DSC measurements were conducted with a Differential Scanning Calorimeter with a Differential Scanning Calorimeter TA Instruments Q2000 type. Samples of about 10 mg in a sealed alumina pan were measured under N 2 flow. The sample was heated to 120 °C at a rate of 10 °C/min then cooled to 30 °C following by reheating to 180 °C. at a rate of 5 °C/min. Measurements were collected for the second run. For the case of RF end functionalized PS end groups, RF association has been shown to be stronger than for pendent groups at the same RF concentration. Hence such association may affect chain end mobilities and hence glass transition temperatures especially for smaller MW polymers where the RF concentrations tend to be high as seen above. In addition the degree of RF association is known to increase with increasing RF lengths. 59-61 Thus RF association may affect the larger scale segmental motions implicated in the glass transition and/or crystallization/melting properties. As the RF content, by definition, depends on MW for the case of RF end functionalized PS and other RF –functionalized polymers, the effects of (a) MW (b) RF length and (c) RF content ! 67! effects will have to be evaluated independently. Hence, it is plausible that this could also be expected to be expressed in the thermal and other i.e. mechanical properties. The above optical transmission data of the above RF-functionalized PS have indicated the presence of large RF domains (> 400 nm) for the C 13 F 27 end functionalized PS at concentrations of roughly 2-6 wt percent RF content. Figure 3.8. DSC thermograms of PSC13 polymers with varying molecular weights. PSC13- 10K(Blue), PSC13-15K (Green), PSC13-30K(Pink). Measurements were taken at second run experiments and the samples were heated to 180 °C at a rate of 5 °C under nitrogen. As shown in Figure 3.8 the glass transition temperature of the PSC13-10K polymer surprisingly has shifted toward a much lower glass transition (76.2 °C, see Table 3.2) compared with the isobaric homopolymer having a Tg of 91.3 °C a difference of around 15 degrees. At a MW of about 30K the effects were somewhat smaller but still robust with a decrease in Tg being almost 11 degrees (Table 3.2). Similar results are shown for the C 7 F 15 and C 10 F 21 end functionalized polymers (Table 3.2 and Figures 3.9). In general the differences increased with increasing RF lengths and decreasing chain lengths (Table 3.2). For instance, a comparison of 0.15 0.05 Heat Flow (W/g) 50 70 90 110 130 150 170 190 Temperature (°C) PS C13 13K ––––––– PS C13 22K EXP274 – – – – ps c13 9k ––––– ∙ Exo Up Universal V4.7A TA Instruments ! 68! the PSC10-10K and the PS-10K gives a decrease in Tg of about 7 °C. The smallest ΔT g values are seen of the highest MW and smallest RF lengths (ΔT g = 2.6 and 3.3 °C for the PSC7-30K and PSC10-30K respectively). Thus the effects here were analogous to changes in OT in that RF length is paramount rather than RF concentration. Table 3.2 Summary of glass transition temperatures C 7 F 15 , C 7 F 15 and C 7 F 15 end functionalized polystyrenes and the corresponding PS homopolymers. Figure 3.9. DSC thermograms of PSC7-10K (green), PSC10-10K (red), PSC13-10K (purple). Measurements were recorded at second run. 0.2 0.1 0.0 Heat Flow (W/g) 50 70 90 110 130 150 Temperature (°C) PSC7 10K ––––––– PSC10 10K ––––––– PSC13 10K ––––––– Exo Up Universal V4.5A TA Instruments Samples RF wt% T g ( 0 C) Samples RF (wt%) T g ( 0 C) PS-10K _ 91.3 PSC13-15K 5.1 90.0 PSC7-10K 4.4 87.4 PS-30K _ 101.6 PSC10-10K 6.2 84.3 PSC7-30K 1.3 99.0 PSC13-10K 7.4 76.2 PSC10-30K 2.0 98.3 PSC7-15K 2.8 98.0 PSC13-30K 2.4 90.6 PSC10-15K 4.1 93.8 ! 69! Figure 3.10. DSC thermograms of PSC7, PSC10 and PSC13 of 15K polymer. Measurements were recorded at second run. The DSC thermograms of PSC7, PSC10 and PSC13 at a PS MW of 15K show Tg values of 98 °C, 93.8 °C and 90 °C, respectively (Figure 3.10). LC formation could not be seen for these polymers. The DSC thermograms of PSC7-30K, PSC10-30K, PSC13-30K and PS-30K showed Tg values of 99, 98.3, 90.6 and 101.6 0 C respectively (Figure 3.11). The effects of the C 7 F 15 and C 10 F 21 end groups were relatively small (2 and 3 0 C) but the Tg of PSC13-30K was about 11 °C lower than the PS homopolymer. Given the above it is clear that the effects on Tg seem to reflect RF length rather than content. 0.2 0.1 0.0 Heat Flow (W/g) 50 70 90 110 130 150 Temperature (°C) PS C7 15K ––––––– PS C10 15K ––––––– PS C13 15K ––––––– Exo Up Universal V4.5A TA Instruments ! 70! Figure 3.11. DSC thermograms of PSC7, PSC10, PSC13 and PS30K. Measurements were recorded at second run. As C 7 F 16 has a low melting point (mp. =-78 0 C) the presence of very small domains is consistent with expectations. 15 Furthermore, the formation of small RF domains has been implicated for the case of blends of PS and polybutylmethacrylate having 1,1- dihydroperfluorooctyl end groups. 59-61 For C 13 F 27 end functionalized PS the decrease in transmittance is consistent with the formation of larger RF crystalline or semi-crystalline C 13 F 27 aggregates. Considering the well-known tendencies of large perfluorocarbons to associate in many media it seems reasonable to expect that this should correlate with higher degrees of association and possibly different RF association modes as RF lengths increase (Scheme 1.4). Given the polymer structure it is clear that at a constant polymer mass, the concentration of perfluoro end groups decreases inversely with the number average MW, i.e. a factor of three in going from a MW of 10k to 30k. 0.2 0.1 0.0 Heat Flow (W/g) 50 70 90 110 130 Temperature (°C) PSC7 30K ––––––– PSC10 30K ––––––– PSC13 30K ––––––– Exo Up Universal V4.5A TA Instruments ! 71! It is surprising that the RF end-functionalized PS show decreases rather than increases in Tg with increasing perfluorocarbon lengths at constant MWs (Figures 3.9-3.11). Like the OT effects discussed in the preceding section the effects on Tg seem to reflect RF length rather than –content (Table 3.1) Decreases from a few degrees of about 15 0 C are significant. The attachment of the relatively rigid perfluorocarbon groups may be expected to increase the Tg’s. For instance, the melting points of the corresponding C 7 F 16 , C 10 F 22 and C 13 F 28 perfluorocarbons are known as about -80, 43 and 104 0 C respectively. 15 However this comparison is at best qualitative as the mp’s reflect solid-liquid equilibria involving the unmodified perfluorocarbons while the Tg values reflect the effects of PS glass transition temperatures mediated by the association of the RF end groups that may or may not be crystalline at the temperatures involved. Assuming that a C 7 F 15 end functionalized PS plausibly may have the perfluorocarbon end-groups associated into a liquid-like domains this is consistent with small decreases in Tg. However the lower MW PSC13 polymers show larger decreases in Tg compared with the unfunctionalized PS. This could be due to the increased disruption in PS chain packing due to the presence of large liquid crystalline PTFD domains and, thus, perhaps, excluded volume effects. On the other hand the presence of large RF groups at the higher temperatures may indicate a liquid like behavior of the RF groups at these higher temperatures that imparts greater mobilities to the longer RF groups. This behavior has been seen in polytetrafluoroethylene where rapid conformational helical chain reversals have been observed. 16 The crystallization temperature seen for PSC13-10K is a clear indication of presence of perfluorocarbon crystalline or semi-crystalline structures (Figure 3.8). At higher molecular weights the crystallization temperature was not seen probably due to the decrease in the perfluorocarbon content in the polystyrene chain, for ! 72! example, compared to 10K polymers, the perfluorocarbon content of 15K decreased by 50 wt% (Figure 3.10-3.11). As indicated above this is also seen in the OT data. 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However, their mechanical and thermal properties can be improved by reinforcing them with carbon nanotubes, nanoparticles, fibers or other fillers. 1-2 Although stiffness of a polymer can be improved by fillers or additives it may decrease the toughness of the polymeric material. 3 Polyhedral Oligomeric Silsesquioxanes (POSS) are nanostructures having the empirical formula RSiO 1.5 , where R may be an alkyl, hydroxyl, hydrogen etc. POSS contains a silica cage and functional groups attached to the corners of it (Scheme 4.1). POSS nanocages have diameters of 1-3 nm and they are considered to be the smallest silicon based structures. 4 POSS nanoparticles have gained much attention due to desirable properties. POSS functionalized materials can potentially be used in many applications such as polymer electrolytes, surface coating and biomedical applications. 5-8 It has been shown that incorporation of POSS nanoparticles in a polymer matrix increase modulus, strength and thermal stability whereas they may reduce the viscosity of the polymer. 4,9-11 In addition, they can be processed easily and they are non-volatile and environmentally friendly. ! 83! Scheme 4.1. Structure of POSS 12 In previous chapters, it has been shown that self-assembly and association of RF groups lead to increased thermal and viscoelastic properties of PS. Similarly, POSS groups are known to form aggregates when they were blended into a polymer matrix or copolymerized. 13-15 Previously, the degree of association of RF end groups was shown to be much more pronounced compared with RF pendent groups 16-18 It is possible same principle also applies for POSS polymers. We synthesized POSS end functionalized PS with different MWs to investigate the effect of isobutyl POSS end functionalization on thermal, optical and rheological properties of PS. 4.2 Experimental 4.2.1 Materials All chemicals were purchased from Aldrich and used as received unless otherwise noted. Styrene was freshly distilled after being stirred over CaH 2 overnight. Copper Bromide (CuBr) was purified by stirring over acetic acid overnight, then washed with 2-propanol and dried in vacuum oven at 40 °C. ! 84! 4.2.2 Synthesis of POSS Macroinitiator (POSS-Br) Macroinitiator was synthesized according to the procedure described in literature. 19 A schlenk flask was evacuated and filled with Argon three times and cooled to -78 °C. 2- PSS-(3- Hydroxypropyl)-heptaisobutyl substituted (POSS-OH) (0.86 mmol), 20 ml dry dichloromethane and triethylamine (1.68 mmol) were added to the reaction flask. Then, 2-Bromoisobutyryl bromide (1.89 mmol) was added slowly and the mixture was stirred at -78 °C for 3 hours and then 2 hours at room temperature. After filtration, the filtrate was diluted with 200 ml of dichloromethane. After standard NaHCO 3 -water work-up, the solution was dried over MgSO 4 , it was filtrated and concentrated by rotary evaporator. It was then precipitated in large excess of methanol, filtered and dried in vacuum oven at 40 °C overnight. 1 H NMR (400 MHz, Chloroform-d) δ 4.14 (dd, J = 7.0, 6.3 Hz, 2H), 1.94 (d, J = 0.7 Hz, 6H), 1.85 (dtd, J = 13.4, 6.7, 1.9 Hz, 8H), 1.80 -1.73 (m, 2H), 0.96 (dd, J = 6.6, 0.7 Hz, 48H), 0.70 -0.64 (m, 2H), 0.61 (m, J = 7.0, 3.2, 0.8 Hz, 16H). 4.2.3 Preparation of POSS End-functionalized hybrid Polystyrene via ATRP Polymers were synthesized by ATRP using [Initiator]/[CuBr]/[Ligand] ratios of 1:1:2.In a typical example, CuBr and a dry magnetic stirrer were added into a dry Schlenk flask and is sealed with a rubber septum before three vacuum/Argon cycles. Styrene, N,N,N’,N’,N’’- pentamethyldiethylenetriamine (PMDETA) and toluene were introduced into the flask under argon atmosphere and stirred for 10 min to allow the formation of copper-ligand complex. POSS macroinitiator was dissolved in 0.5 ml of toluene and added to the flask. The reaction was then degassed by three freeze-thaw cycles and after the equilibration to the ambient temperature, the flask was immersed into the oil bath at 80 °C. The reaction was allowed to proceed overnight and then the reaction was stopped by exposure to the air and was cooled down to the room ! 85! temperature. The solution was diluted with THF and copper catalyst was removed by passing through silica gel column. The filtrate was concentrated by rotary evaporator. The polymer was then precipitated in large excess of methanol and dried in vacuum oven at 70 °C. 4.2.4 Characterization The chemical structures of POSS-OH, POSS-Br and POSS functionalized polymers were characterized by 1 H nuclear magnetic resonance spectroscopy ( 1 H NMR) (Varian Mercury 400 MHz) using CDCl 3 as solvent. The molecular weights and molecular weight distributions were determined by Size Exclusion Chromatography (SEC) performed using Shimadzu LC-20AT pump, an RID-10A refractive index detector. THF was used as the elution solvent with a flow rate of 1 ml/min and the instrument was calibrated using polystyrene standards. The thermal stability of POSS end-functionalized PS was obtained by thermagravimetric analysis (TGA) performed both under air and N 2 with the rate of temperature rise 10 °C/min using Shimadzu TGA-50. DSC measurements were conducted with a Differential Scanning Calorimeter with a Differential Scanning Calorimeter TA Instruments Q2000 type. Samples of about 10 mg in a sealed alumina pan were measured under N 2 flow. All DSC data were collected during second heating run at a scanning rate of 5 °C/min from 30 °C to 180 °C under nitrogen atmosphere after heating samples from 30 °C to 120 °C at 10 °C/min and then cooling to 30 °C. The rheological properties of the polymers were studied with TA instruments AR2000EX parallel plate Rheometer. Viscoelastic regime was determined measuring the dynamic shear storage and loss moduli, G’(ω) and G’’(ω). Rheological measurements were carried at 120 °C. ! 86! 4.3 Molecular Weights of POSS Functionalized Polymers A number of POSS end-funtionalized PS samples were synthesized using ATRP methods (Table 4.1) .The polydispersities (PDIs) were kept low by using high monomer concentrations and limiting monomer conversions. Table 4.1. Molecular Weights and Polydispersities of PS and POSS-PS hybrid Polymers Sample M n a PDI a POSS content b (wt%) PS 10K 8950 1.06 0 PS 25K 22800 1.06 0 POSS-E-PS 5K 5248 1.10 15.1 POSS-E-PS 10K 8465 1.07 9.7 POSS-E-PS 25K 21617 1.30 2.9 a. Determined by SEC using polystyrene standards. b. Calculated using weight average MW 4.4 UV-Visible measurements of PS and POSS end-functionalized PS films Refractive indices of isobutyl functionalized POSS was given as 1.46-1.48 20 which is slightly smaller than that of PS (1.59-1.60). 21 Figure 4.1. UV-Visible Transmittance versus wavelength of PSPOSS-5K and 10K 0! 20! 40! 60! 80! 100! 400# 500# 600# 700# 800# 900# 1000# Transmittance#(%)# Wavelength#(nm)# PS!POSS!5K! PS!POSS!10K! ! 87! The optical transmission (OT) values should be dependent on (a) refractive indices differences between PS (b) the sizes of the POSS domains sizes (c) and the mobilities of the POSS end groups. As discussed above PS-10K sample is almost optically clear (OT ~96%). PS- POSS of 10K showed similar optical transmittance to PS-10K whereas PS-POSS 5k had slightly lower transmittance (OT ~93%) (Figure 4.1). Given the high POSS content (10 %wt- OT 96%) the transmittances were much higher compared with PSC13 polymers (7.4%-OT 58%) at similar MW of 10k. This is attributed to weak association of POSS groups compared with C 13 F 27 groups. It is notable that the refractive indices difference between RF and PS was also much higher compared with the difference between POSS and PS. 4.5 Differential Scanning Calorimetry analysis of POSS end-functionalized PS It has been shown that POSS functionalization on hard thermoplastics 22-24 usually resulted in increased in Tg whereas for resins either an increase or decrease in Tg was observed. 25-27 Table 4.2. Summary of glass transition temperatures C 7 F 15 , C 7 F 15 and C 7 F 15 end functionalized polystyrenes and the corresponding PS homopolymers. Polymer POSS wt % Tg (°C) Delta Tg (°C) PS-5K a - 82.4 - PS-10K - 91.3 - PS-25K - 101.3 - POSS-PS 5K 15.1 76.2 6.2 POSS-PS 10K 9.7 88.5 2.8 POSS-PS 25K 2.9 100.5 0.8 a. Tg value is obtained from the literature 28 ! 88! As shown in Table 4.2 glass transition temperature of PS-POSS shifted towards a lower value compared with corresponding PS. The biggest decrease was seen for POSS-PS 5k with the highest POSS content (15.1 wt%) where its glass transition temperature was about 6 °C lower than PS 5k. Tg of higher MW POSS-PS had smaller decreases, for instance, Tg of POSS-PS 10k and POSS-PS 25k was 2.8°C and 0.8 °C lower compared with PS-10k and PS-25k, respectively. 4.6 Thermogravimetric analysis of POSS end-functionalized PS TGA curves of POSSPS-10k and PS-10k under air were shown in Figure 4.2. The degradation of PS started at about 280 °C. When similar MW of PS and POSSPS-10k was compared, starting point of thermal degradation was 20 °C higher for POSS-PS whereas this difference become more prominent showing the rate of degradation was much faster for PS-10k compared with POSSPS-10k. For instance, the ratio of residual polymer weight was 40 wt% at about 380 °C and 415 °C for PS-10k and POSSPS-10k, respectively. (Figure 4.2) More surprisingly, thermal degradation of POSSPS-10k started at about 60 °C higher compared with PS-10k under nitrogen. (Figure 4.3) Figure 4.2. Thermogravimetric analysis of POSSPS-10k (green), PS-10K (red) polymers under air 0! 10! 20! 30! 40! 50! 60! 70! 80! 90! 100! 250! 270! 290! 310! 330! 350! 370! 390! 410! 430! 450! Weight#%# # Temperature (° C) ! 89! Figure 4.3. Thermogravimetric analysis of POSS PS polymers of various molecular weights under nitrogen . PS-10K (Pink), POSSPS-10K (Blue). 4.7 Rheological Properties of POSS-PS Hybrid Polymers In order to investigate the effect of POSS end functionalization on the viscoelastic properties of PS rheological measurements were carried out. Figure 4.4 shows the shear storage moduli of PS and POSSPS-10k. At low frequency (0.1 Hz) G’ of POSS-PS was 6 times lower than G’ of PS. About after 1 Hz, G’ of PS and POSS-PS became almost identical. Similarly G’’ of POSS- PS and PS were similar between 0.1-10 Hz (Figure 4.5). ! 90! Figure 4.4. Shear Storage moduli for POSSPS-10K(blue) , PS-10K (red) polymers at 120 °C. The reference strain is 0.5 %, respectively. Figure 4.5. Shear Loss moduli for POSSPS-10K(blue) , PS-10K (green) polymers 120 °C. The reference strain is 0.5 %, respectively. 1! 10! 100! 1000! 10000! 100000! 1000000! 0# 1# 2# 3# 4# 5# 6# 7# 8# 9# 10# G'(Pa)# Frequency(Hz)# # 1! 10! 100! 1000! 10000! 100000! 1000000! 0# 2# 4# 6# 8# 10# G''(Pa)# Frequency#(Hz)# ! 91! 4.8 References 1. Coleman, J. N.; Khan, U.; Blau, W. J.; Gun’ko, Y. K. Small but strong: a review of the mechanical properties of carbon nanotube–polymer composites. Carbon 2006, 44, 1624-1652. 2. Nunes, R.; Fonseca, J.; Pereira, M. Polymer–filler interactions and mechanical properties of a polyurethane elastomer. Polym. Test. 2000, 19, 93-103. 3. Morgan, A. B.; Wilkie, C. A. Flame retardant polymer nanocomposites; John Wiley & Sons: 2007; . 4. Li, G.; Wang, L.; Ni, H.; Pittman Jr, C. U. Polyhedral oligomeric silsesquioxane (POSS) polymers and copolymers: a review. Journal of Inorganic and Organometallic Polymers, 2001, 11(3), 123-154. 5. Choi, J.; Lee, K. M.; Wycisk, R.; Pintauro, P. N.; Mather, P. T. Sulfonated polysulfone/POSS nanofiber composite membranes for PEM fuel cells. J. Electrochem. Soc. 2010, 157, B914- B919. 6. Decker, B.; Hartmann-Thompson, C.; Carver, P. I.; Keinath, S. E.; Santurri, P. R. Multilayer Sulfonated Polyhedral Oligosilsesquioxane (S-POSS)-Sulfonated Polyphenylsulfone (S-PPSU) Composite Proton Exchange Membranes. Chemistry of Materials 2009, 22, 942-948. 7. Devaux, E.; Rochery, M.; Bourbigot, S. Polyurethane/clay and polyurethane/POSS nanocomposites as flame retarded coating for polyester and cotton fabrics. Fire Mater. 2002, 26, 149-154. 8. McCusker, C.; Carroll, J. B.; Rotello, V. M. Cationic polyhedral oligomeric silsesquioxane (POSS) units as carriers for drug delivery processes. Chemical communications 2005, 996-998. 9. Pan, M.; Gorzkowski, E.; McAllister, K. In In Dielectric properties of polyhedral oligomeric silsesquioxane (POSS)-based nanocomposites at 77K; IOP Conference Series: Materials Science ! 92! and Engineering; IOP Publishing: 2011; Vol. 18, pp 082006. 10. Zhou, Z.; Cui, L.; Zhang, Y.; Zhang, Y.; Yin, N. Preparation and properties of POSS grafted polypropylene by reactive blending. European Polymer Journal 2008, 44, 3057-3066. 11. Huang, J. C.; He, C. B.; Xiao, Y.; Mya, K. Y.; Dai, J.; Siow, Y. P. Polyimide/POSS nanocomposites: interfacial interaction, thermal properties and mechanical properties. Polymer 2003, 44(16), 4491-4499. 12. Ayandele, E.; Sarkar, B.; Alexandridis, P. Polyhedral oligomeric silsesquioxane (POSS)- containing polymer nanocomposites. Nanomaterials 2012, 2, 445-475. 13. Yei, D.; Kuo, S.; Su, Y.; Chang, F. Enhanced thermal properties of PS nanocomposites formed from inorganic POSS-treated montmorillonite. Polymer 2004, 45, 2633-2640. 14. Misra, R.; Alidedeoglu, A. H.; Jarrett, W. L.; Morgan, S. E. Molecular miscibility and chain dynamics in POSS/polystyrene blends: control of POSS preferential dispersion states. Polymer 2009, 50, 2906-2918. 15. Hirai, T.; Leolukman, M.; Hayakawa, T.; Kakimoto, M.; Gopalan, P. Hierarchical nanostructures of organosilicate nanosheets within self-organized block copolymer films. Macromolecules 2008, 41, 4558-4560 16. Shen, J.; Hogen-Esch, T. Block copolymer-like self-assembly of fluorocarbon end- functionalized polystyrene and polybutylmethacrylate. J. Am. Chem. Soc. 2008, 130, 10866- 10867. 17. Shen, J.; Piunova, V. A.; Nutt, S.; Hogen-Esch, T. E. Blends of polystyrene and poly (n-butyl methacrylate) mediated by perfluorocarbon end groups. Polymer 2013, 54, 5790-5800. 18. Shen, J. Perfluorocarbon mediated self-assembly of polymers; University of Southern California: 2009. ! 93! 19. Ohno, K.; Sugiyama, S.; Koh, K.; Tsujii, Y.; Fukuda, T.; Yamahiro, M.; Oikawa, H.; Yamamoto, Y.; Ootake, N.; Watanabe, K. Living radical polymerization by polyhedral oligomeric silsesquioxane-holding initiators: precision synthesis of tadpole-shaped organic/inorganic hybrid polymers. Macromolecules 2004, 37, 8517-8522. 20. Polyhedral Oligomeric Silsesquioxane Handbook, Phantom Plastics 2010, http://phantomplastics.com/wp-content/uploads/2013/08/POSS-Handbook.pdf 21. Brandrup, J.; Immergut, E. H.; Grulke, E. A.; Abe, A.; Bloch, D. R. Polymer handbook; Wiley New York: 1999; Vol. 89. 22. Haddad, T. S.; Lichtenhan, J. D. Hybrid organic-inorganic thermoplastics: styryl-based polyhedral oligomeric silsesquioxane polymers. Macromolecules 1996, 29, 7302-7304. 23. Miyamoto, K.; Hosaka, N.; Kobayashi, M.; Otsuka, H.; Yamada, N.; Torikai, N.; Takahara, A. Dewetting inhibition and interfacial structures of silsesquioxane-terminated polystyrene thin films. Polym. J. 2007, 39, 1247-1252. 24. Shao, Y.; Aizhao, P.; Ling, H. POSS end-capped diblock copolymers: synthesis, micelle self- assembly and properties. J. Colloid Interface Sci. 2014, 425, 5-11. 25. Pittman Jr, C. U.; Li, G.; Ni, H. In In Hybrid inorganic/organic crosslinked resins containing polyhedral oligomeric silsesquioxanes; Macromolecular Symposia; Wiley Online Library: 2003; Vol. 196, pp 301-325. 26. Villanueva, M.; Martín-Iglesias, J.; Rodriguez-Anon, J.; Proupín-Castiñeiras, J. Thermal study of an epoxy system DGEBA (n= 0)/mXDA modified with POSS. Journal of thermal analysis and calorimetry 2009, 96, 575-582. 27. Zhang, Z.; Liang, G.; Wang, X. The effect of POSS on the thermal properties of epoxy. Polymer Bulletin 2007, 58, 1013-1020. ! 94! 28.DSC Measurements of Polystyrene, Hitachi High-Tech Science Corporation 1995, http://www.hitachihightech.com/file/global/pdf/products/science/appli/ana/thermal/application_T A_068e.pdf ! 95! CHAPTER 5 Development of Increased Flow Graphene Oxide-Polymer Nanocomposite Membranes 5.1 Introduction The water shortage has become a worldwide concern since only small portion of water is fresh in the earth. As population continues to grow, few resources are available to supply clean water. As a result, solutions such as water reuse and water desalination have become leading technologies to supply fresh water for future populations in the earth. 1-7 Water reuse has been implemented to supply water for uses such as industrial processes and irrigation. 5-7 Desalination has currently become an important source of drinking water production. 1-4 With desalination technologies, it is possible to produce fresh water from seas, oceans and brackish water. The salinity of feed for desalination is reported between 1000 mg/L total dissolved solids (TDS) to 60,000 mg/L TDS. 8 Desalination processes are divided into two main categories: thermal process or membrane process. Although thermal processes has been the primary technology in Middle East, as they are very energy intensive, membrane processes have recently become more popular. 9,10 Table 5.1 Types of Membranes 11 Filter Symbol Pore size (µm) Type of materials removed Microfiltration MF 1.0-0.01 Sand, large bacteria and large viruses Ultrafiltration UF 0.01-0.001 Suspended particles, bacteria, Viruses, Nanofiltration NF 0.001-0.0001 Suspended particles, Bacteria, viruses, divalent anions Reverse Osmosis RO <0.0001 Monovalent ions ! 96! Microfiltration (MF), Ultrafiltration(UF), Reverse osmosis (RO), nanofiltration (NF) are membrane processes used for desalination. Table 5.1 shows pores sizes and typical materials that can be removed by using applicable technologies. 11 As shown in Table 5.1, RO membranes are the most effective to remove monovalent ions, such as sodium and chloride. 5.1.1 Reverse Osmosis: Basic Principles Reverse Osmosis has gained much attention in industry in last decades, since it seems to be an efficient method to purify the water. It is well known that when a salt solution is separated from a solution containing a lower salt concentration by a semi permeable membrane (permeable by water but not by salt) a pressure differential is established across the membrane (Figure 5.1- A). The reverse happens in reverse osmosis as shown in Figure 5.1-B. This process will continue until an equilibrium is reached. In reverse osmosis, water is moved through the membrane against the concentration gradient, thus, from lower concentration to higher concentration (Scheme 5.1). A B Scheme 5.1. Reverse Osmosis 12 ! 97! RO processes involve a “feed” water source, high pressure pumps and a semipermeable membrane. It typically involves steps such as pretreatment, membrane separation and post treatment steps. Current RO processes generally uses multi-pass system in which multiple membranes are connected in parallel or more commonly in series, because current RO membranes do not have enough ion rejection efficiency to produce clean water in single-pass. The use of more than one membrane should result in increased ion rejection. 5.1.2 Transport through RO membranes RO membranes are generally asymmetric membranes with pore sizes about less than 0.7 nm. These membranes are able to remove hydrated salt ions (diameter ~7-8 Å) from water, regardless of their charge. Although the exact mechanism of transport is still not well understood two models are used to explain the permeation process; pore flow and solution-diffusion model. 13 The most accepted mechanism for water transport through RO membranes is solution- diffusion model. 13-17 In this model, after water molecules have absorbed onto the membrane surface the concentration gradient across the membrane leads to the diffusion to the permeate side of the RO membrane. Then water molecules then desorbs from the membrane. Mass transport through RO membranes can be described as follows 18-20 : ! ! = ! !"−!П (1) Where N A is liquid (water) flux across the membrane, L is the permeability coefficient, Δp is the transmembrane pressure difference, and ΔП is the osmotic pressure difference. The key concept is the membrane salt rejection which is a measure of overall membrane system performance. Salt rejection through an RO membrane is given by 21 : != 1− ! !"#$"%&" ! !""# !100% (2) Where C is the ion concentration (molar units), R is the rejection. ! 98! 5.2 Current Reverse Osmosis / NanoFiltration Membranes An ideal RO membrane should maximize water permeability, have high salt rejection, and resist fouling and degradation by chemicals. First, cellulose acetate was used as RO membrane in the 1960s. However, they suffered from low flux and biological degradation. 22 Current state-of-art RO membranes are typically polyamide thin-film composite membranes. These membranes are generally consisted of two layers: dense layer and porous sublayer. 23 (Scheme 5.2) The dense layer is usually polyamide that does actual desalination. It is fabricated onto a porous polysulfone that gives mechanical support. 24 However, these membranes are susceptible to degradation from the hypochlorite ion used to prevent biological fouling. 25 Other limitation is the difficulties in making polymeric materials having controllable pore sizes ≤1 nm. Nanostructured materials including zeolites 26-27 , carbon nanotubes (CNTs) 28-29 have also gained attention since they can act as high flux molecular sieving membranes for water desalination. However, they also have some limitations such as permeability of ions around zeolite crystal. Similarly, the diameters CNT pores naturally were too large to filter salt ions. Scheme 5.2. Asymmetric Structure of the Current membranes 23 The other limitation of existing PA membranes is the membrane fouling due to the binding of proteins onto hydrophobic surfaces. Anti-biofouling properties can be introduced to ! 99! the hydrophobic membranes by using hydrophilic or zwitterionic surface coatings. 30-33 Poly(ethylene glycol) (PEG)/ Oligo(ethylene glycol) (OEG) surfaces are the most widely used antifouling materials as hydrophilic coatings. 33 However, zwitterionic-based materials can attach water molecules more strongly than PEG/OEG chains via electrostatically induced hydration. 34-36 Because of the various limitations of existing membranes there is still a need for finding alternative membranes that will have better water flux and salt rejection in addition to good resistance to fouling and degradation. Graphene has recently gained attention due to its unique properties. Graphene is a single layer of graphite with one atom thickness, consisting of a lattice of hexagonally arranged sp 2 -bonded atoms. 37-38 It is an inexpensive and strong material. 39-41 For example, compared to steel, graphene has been reported as six times lighter, two times harder, and giving ten times higher tensile strength. 42 These properties suggest the potential of graphene to create thin membranes with high water permeation. However, it is impermeable to molecules as small as helium in its pristine state 43 , therefore there is a need to produce controllable sub-nm pores to obtain water passage through membrane. Recent simulations and experimental studies suggest that subnanometer pores can be controllably generated by the methods such as oxidation 44 , electron beam irradiation 45-46 , ion bombardment 47-51 , or by doping. 52 Some theoretical studies have been carried out on water transport and ion rejection of graphene sheet. 53-54 Sint et al. studied transport of ions through ~0.5 nm pores in graphene terminated with nitrogen or hydrogen using molecular dynamics (MD) simulations. 55 Then, Suk et al. explored the theory of water transport across 0.75-2.75 nm diameter pores in graphene, and compared water flux in graphene with the flux in 2-10 nm thick carbon nanotubes(CNTs) membranes with similar diameters. 56 These molecular dynamics (MD) simulation studies showed the potential of graphene as membrane material for desalination in the future. Especially ! 100! graphene oxide (GO) has gained much attention since it offers excellent potential for making nanocomposite materials having improved stability, hydrophilicity and superior antibacterial and antifouling properties. 57-61 GO membranes were shown to be impermeable to liquids, vapors and gases including helium whereas they had unimpeded permeation of water. 62-63 However, pure GO membranes are easily leached out of their support under high hydrophilic environment and transmembrane pressure. In order to overcome this problem, we incorporated GO in PA matrix in order to investigate its potential to improve water flux without leaching problems. 5.3 Experimental 5.3.1 Materials and Methods All materials are purchased from Alfa Aesar unless otherwise noted. Polyethersulfone (GE Osmonics YMPTSP3001 molecular weight cut off is about 5000 daltons) was obtained from Sterlitech Corporation. Graphene Oxide solution was purchased from Graphene Supermarket. 5.3.2 Synthesis and Preparation of Graphene Oxide Incorporated Polyamide (PA) Membranes Membranes were hand-cast on polyethersulfone(PES) ultrafiltration membranes (UF) provided by Sterlitech via interfacial polymerization. Polysulfone support membrane was put in deionized water for 12 hours, then removed from the water and taped on a flat glass plate. In a flask, 4 g of m-phenylene diamine (MPD) was dissolved in 120 ml of water. In a separate flask, 0.5 wt% graphene oxide (purchased from graphene supermarket) was prepared in 80 ml of water and ultrasonicated for 45 min. Then m-phenylene diamine solution was mixed with graphene oxide solution making a total of 200 ml of 2% (v/w) m-phenylene diamine and 0.5 wt% graphene oxide aqueous solution. M-phenylene diamine and graphene oxide solution were then poured onto the membrane and allowed to contact to Polysulfone support for at least 3 min. Then ! 101! the excess solution rolled firmly with a rubber roller. Then, 0.1% (w/v) Trimesoyl chloride (TMC) in hexane was poured onto the membrane. After 1 minute reaction, the TMC solution was poured off and the resulting membranes were rinsed using hexane to wash away residuals. After membrane was air dried at room temperature or annealed at 60 °C for 10 min depending on the experiment, it was immersed in deionized water until use. The synthesis and fabrication of Polyamide (PA)-Graphene Oxide nanocomposite membranes are shown in Scheme 5.3-5.4. Scheme 5.3. GO PA Membrane Fabrication 1/1 FIGURE 1 FIGURE 2 FIGURE 3 PDA TMC PA ! 102! Scheme 5.4. PA-GO Membrane Synthesis 5.3.3 Preparation of Graphene Oxide incorporated Blends of Poly(tetrabutylammonium styrene sulfonate copolymer and Polyvinylidenefluoride (PBASS copolymer- PVDF) blend and Polystyrenesulfonic acid copolymer-polyvinylidenefluoride (PSSA copolymer-PVDF) blend membranes Membranes were prepared by solution casting according to the procedure reported. 64 First, GO was sonicated for 45 min. 15% wt P(BASS-S-CMS) and PVDF 80% were dissolved in N,N- Dimethylformamide (DMF) at separate beakers at ambient temperature. After dissolution of P(BASS-S-CMS), sonicated 5% wt GO in DMF was added to the beaker slowly while stirring the solution vigorously in order to prevent GO sheets from “stacking up”. After 30 minutes, PVDF dissolved in DMF was added to the beaker and the solution was stirred for another 15 minutes at ambient temperature. The polymer blend containing 5% wt GO was then poured into a petri dish and transferred to the oven preheated at 70 °C or 165 °C depending on desired experimental conditions. After annealing for 2 hours, the membranes were quickly quenched in water at 25 °C. The corresponding PSSA-PVDF blend membranes were prepared according to the literature. 64 1/1 FIGURE 1 FIGURE 2 FIGURE 3 PDA TMC PA ! 103! 5.4 Water Permeability and Particle Rejection Studies of Graphene Oxide Incorporated Polyamide (PA) Membranes The initial work on the project involved the development of polymer synthesis protocols and synthesis conditions including reaction times and curing procedures. Subsequently, graphene oxide was incorporated as nanomaterial into the polymers at a set proportion. The membranes used in the series of preliminary tests were prepared by interfacial polymerization with monomer on a commercial polyether sulfone (PES) ultrafiltration membrane support. Several membranes were prepared with variations in annealing temperature, or reaction time in the preparation protocol and their effects on the membrane filtration properties were also studied (Table 5.2). A flow rate of 100 mL/min corresponds to a membrane permeate flux of 2500 L/m 2 /hr based on the membrane film dimensions of for the small cell tests. Table 5.2. Reaction conditions of prepared membranes Membrane No Annealing at 68 °C GO in polymer matrix GO as additional layer 3 Yes No No 4 Yes Yes No 5 No Yes No 8 Yes Yes Yes 9 No Yes Yes It was observed that the permeate flux of membrane 4 increased almost 3 fold compared to membrane 3 that is a polyamide control membrane made in the same conditions (Table 5.3). Membrane 5 that was similar to membrane 4 except it was not annealed after the reaction showed relatively higher water flux. Membrane 8 contained graphene oxide both in the polymer ! 104! matrix and on the surface of the membrane as additional layer and it was observed that additional coating of GO decreased flux considerably compared to membranes 4 and 5. Membrane 9 was similar to membrane 8 except it was not annealed after membrane preparation and its water flux was similar to that of membrane 8. It is notable that the permeate flux of pristine polyamide (PA) membranes decreased more rapidly compared to GO incorporated PA membranes. For instance, the initial flux of membrane 3 decreased 4 fold whereas membrane 4 that was prepared at the same conditions with membrane 3 except GO incorporation in the polymer matrix showed only 30% decrease in permeate flux in 3 hours. In addition, GO incorporation did not effect the rejection significantly (Table 5.3). Table 5.3. Water flux and NaCl rejection properties of membranes Membrane #3 #4 #5 #8 #9 TMP (psi) 60 Initial flux (L/m 2 /h) 300 325 450 150 200 Flux after 1hr (L/m 2 /h) 75 250 275 125 125 Flux after 3hr (L/m 2 /h) 75 250 275 125 125 Normalized flux after 3 h 0.25 0.77 0.61 0.83 0.63 Initial conductivity (µS/cm) 2030 2080 2060 2030 2070 Conductivity after 1hr (µS/cm) 1700 1830 1780 1740 1720 Conductivity after 3hr (µS/cm) 1710 1820 1760 1700 1710 Rejection after 3 hr (%) 15.8 12.5 14.6 16.3 17.4 Since NaCl rejection of membranes was weak, we used wastewater having total organic contaminants (TOC) that are bigger as filtration particles. In the present study, the membranes ! 105! designated as membrane 10 was a polyamide membrane without GO whereas membrane 11 had GO in the polymer matrix prepared at the same conditions (Table 5.4). Table 5.4. Performance comparison of various membranes Membrane #10 #11 PES UF Time (h) Permeate flux (L/m 2 /h) 0 265 275 300 0.5 150 140 150 1 50 50 40 2 18 40 40 3 10 40 30 Organic (TOC ) rejection ( %) 44.4 55.9 3.4 The results presented in Table 5.4 provide a performance comparison of various membranes based on permeate flux and TOC rejection. Membrane 10 and 11 had similar initial fluxes about 270 L/m 2 /h, however after 3 hours the permeate flux of membrane 11 was 4 fold higher than membrane 10. TOC rejection also increased with incorporation of GO from 44.4% to 55.9%. UF PES commercial membrane had water flux similar to GO incorporated polyamide membranes except it had much lower TOC rejection, about 3.4%. 5.5 Water Permeability and Particle Rejection Studies of GO incorporated PBASS copolymer- PVDF blend PSSA copolymer-PVDF blend membranes It has been shown that P(BASS-S-CMS)- PVDF blends formed small domains and after acid treatment PSSA copolymer-PVDF membranes showed appreciable potential for water filtration and reclamation applications at high sulfonic acid content (10-40 %). 64 However, it is known ! 106! that high sulfonic acid content membranes could suffer from swelling .This swelling might enlarge the pores leading to decay in rejection properties at long term. For this purpose, we incorporated GO into the blend matrix to enhance the water flux at lower sulfonic acid content in order to compare them with membranes 10 and 11. Desirable properties of GO for water filtration applications were discussed above. Figure 5.1 shows PSSA-copolymer-PVDF blend membranes containing 5 wt% GO annealed at 70 °C. It was observed that membranes were not uniformly coated and some regions had macro defects, as shown by the white cloudy parts. Region I did not have remarkable white spots and looked homogeneous. Region II had GO rich and poor regions, two white circles and region III had one remarkable circle. Therefore Region I and III were used for membrane filtration to see the effect of white cloudy portion on water permeability. The flux pattern for Region I showed similar trends to GO incorporated polyamide membranes in the sense that it came to a stable value at about 400 L/m 2 /hr and did not show a further decrease (Figure 5.2). On the other hand, Region III had the initial flux higher than any other prior membrane filtrations possibly due to blend incompatibilization and graphene oxide sheet stacking. Opaque membranes prepared by annealing at low temperatures were also reported previously. 64 The domain size in region III seemed to vary greatly throughout the membrane possibly diminishing the water permeation with time. Rapid decrease in water permeation in region III can also be explained by fouling. The permeation of GO rich region, region I, was not significantly affected over time. It is also noteworthy that the membranes were brittle and very hard to handle. ! 107! Figure 5.1. PSSA-copolymer-PVDF annealed at 70 °C Figure 5.2. Permeate flux of region I and III of PSSA-copolymer-PVDF annealed at 70 °C Figure 5.3 shows the PSSA copolymer-PVDF blend membranes incorporated with GO annealed at 165 °C. The membranes were much more transparent and stronger compared to the ones annealed at 70 °C. The structure of membranes was visually more homogeneous with no 0 500 1000 1500 2000 2500 3000 3500 0 1 2 3 4 5 6 7 8 9 10 11 12 13 Flux (L/m 2 /hr) Time (hr) Region I Region III ! ! ! I II III I II III ! 108! apparent defects as seen in the previous membranes. The membranes showed similar permeated flux of 400 L/m 2/ /h with region I of membranes annealed at 70 °C (Figure 5.2). That was attributed to the enhanced blend compatibility at high temperatures. The membranes were also tested towards GO leaching. It is clear that there was no GO leaching or membrane damage after high pressure filtration. (Figure 5.4) As a comparison, only GO coating without polymer incorporation on polyethersulfone (PES) support has been shown in Figure 5.5 and 5.6 and clear GO leaching could be seen after filtration. Figure 5.3. PSSA copolymer-GO-PVDF annealed at 165 °C before membrane filtration Figure 5.4. PSSA copolymer-GO-PVDF annealed at 165 °C after membrane filtration ! 109! Figure 5.5. GO coating on PES support before filtration Figure 5.6. GO coating on PES support after filtration More surprising results were seen for P(BASS-S-CMS)-GO-PVDF blends. Although P(BASS-S-CMS)- PVDF blend membranes are expected to be highly hydrophobic and did not show any water permeation even at 60 psi, GO incorporated P(BASS-S-CMS)- PVDF blend membranes showed considerably high water permeation. (Figure 5.7) The water flux was stable after 0.5 hour showing minor fouling indicating the major improvement in water flux and fouling with GO incorporation. ! 110! Figure 5.7. Permeate flux data observed during membrane testing using wastewater 5.6 Conclusions Hybrid nanocomposite membranes have attracted considerable attention as promising candidates for water filtration systems. Compared to pristine polyamide membranes, the water flux of polymeric nanocomposite membranes with 0.5 wt.% GO has increased by 3-fold. Moreover, These nanocomposite membranes showed superior anti-fouling properties and high mechanical strength. PSSA copolymer-PVDF blend membranes incorporated with GO had water flux of 400 L/m 2 /h. 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Abstract (if available)

Abstract Semi-fluorinated polymers are of more recent origin and have several applications due to their unique properties that differ from the corresponding analogous hydrogen or fluorocarbon polymers. They have gained attention in many applications. The presence of perfluorocarbon (RF) (C₇F₁₅ to C₁₃F₂₇) end groups has been shown to enhance the compatibility of polystyrene (PS) blends through fluorophilic interactions. RF functionalized PS with narrow molecular weight (MW) distribution were synthesized by Atom Transfer Radical Polymerization (ATRP). ❧ In Chapter 1, RF end functionalized PS and PS thin films (~ 5μm) deposited on silicon wafers were studied at microscopic level by Tribo nano-indenter at ambient temperature. The effect of RF length from C₇F₁₅ to C₁₃F₂₇ on the viscoelastic properties of PS was investigated. Two modes of nanoindentation were tested: ramp load and frequency sweep. Frequency sweep experiments were carried out between 10-100 Hz. Nanoindentation measurements showed increases of about 85 and 300 percent in the storage and loss moduli respectively compared with the unfunctionalized PS samples at ambient temperature. Shorter RF groups had much smaller effects. ❧ In Chapter 2, the effect of end RF groups on viscoelastic properties at high temperatures was studied. Polymer melts were tested by rheological measurements. Despite relatively modest increases in E′ and E″ shear storage modulus (G′) of the C₁₃F₂₇ end functionalized polystyrenes (30K) is between one to two orders of magnitude larger than the corresponding PS homopolymers whereas the shear loss modulus (G″) was 50% higher between 150 and 180 ℃. ❧ In Chapter 3, the effect of RF length, RF content and polymer MW on the glass transition temperature, optical transmittance and thermal degradation behavior were investigated. Polymers were studied by differential scanning calorimetry (DSC), thermogravimetric analysis (TGA) and UV-visible spectroscopy. ❧ In Chapter 4, mechanical, optical and thermal properties of POSS end-functionalized PS were investigated. The magnitude of effect was compared with RF end-functionalized PS. ❧ In Chapter 5, the water flux and rejection properties of Polyamide and P(BASS-S-CMS)/PVDF polymer blend nanocomposite membranes incorporated with GO were studied. It has been shown that incorporation of 0.5-5wt% of GO markedly improved water flux and fouling behavior of membranes.

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Creator Yurdacan, Hayriye Merve (author)

Core Title Investigation of mechanical, thermal and rheological properties of fluorocarbon functionalized polystyrene; and, Development of increased flow graphene based polymer nanocomposite membranes

School College of Letters, Arts and Sciences

Degree Doctor of Philosophy

Degree Program Chemistry

Publication Date 10/01/2015

Defense Date 07/17/2015

Publisher University of Southern California (original), University of Southern California. Libraries (digital)

Tag desalination,membranes,nano indentation,OAI-PMH Harvest,perfluorocarbon,polymer,water filtration

Format application/pdf (imt)

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Hogen Esch, Thieo (committee chair)

Creator Email hmerveyurdacan@gmail.com,yurdacan@usc.edu

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c40-189961

Unique identifier UC11277226

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Legacy Identifier etd-YurdacanHa-3963.pdf

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Document Type Dissertation

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Rights Yurdacan, Hayriye Merve

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Access Conditions The author retains rights to his/her dissertation, thesis or other graduate work according to U.S. copyright law. Electronic access is being provided by the USC Libraries in agreement with the a...

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