Fabrication of Bioinspired Nanostructured Materials Via Colloidal Self Assembly

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Title:
Fabrication of Bioinspired Nanostructured Materials Via Colloidal Self Assembly
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english
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Huang,Wei-Han
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University of Florida
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Doctorate ( Ph.D.)
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University of Florida
Degree Disciplines:
Chemical Engineering
Committee Chair:
Jiang, Peng
Committee Members:
Ziegler, Kirk
Tseng, Yiider
Chen, Youping

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Subjects / Keywords:
bioinspired -- colloidal -- composite -- gibbsite -- graphene -- nanostructure
Chemical Engineering -- Dissertations, Academic -- UF
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Chemical Engineering thesis, Ph.D.
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theses   ( marcgt )
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Abstract:
Through millions of years of evolution, nature creates unique structures and materials that exhibit remarkable performance on mechanicals, opticals, and physical properties. Recently, a large variety of novel materials have been enabled by natural-inspired designs and nanotechnologies on practical applications. We have utilized bottom-up approaches to fabricate nacre-like nanocomposites with ?brick and mortar structures. Gibbsite-polymer composite displays higher tensile strength and modulus compared to pure polymer. Ultrastrong graphene-oxide-based nanocomposite exhibits high tensile strength and amazing toughness over graphene oxide. We further explored the self-assembly of spherical colloids and the templating nanofabrication of moth-eye-inspired broadband antireflection coatings. Natural optical structures and nanocomposites teach us a great deal on how to create high performance artificial materials. The bottom-up technologies developed in this thesis are scalable and compatible with standard industrial processes, promising for manufacturing high-performance materials for the benefits of human beings.
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In the series University of Florida Digital Collections.
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Includes vita.
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Description based on online resource; title from PDF title page.
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This bibliographic record is available under the Creative Commons CC0 public domain dedication. The University of Florida Libraries, as creator of this bibliographic record, has waived all rights to it worldwide under copyright law, including all related and neighboring rights, to the extent allowed by law.
Statement of Responsibility:
by Wei-Han Huang.
Thesis:
Thesis (Ph.D.)--University of Florida, 2011.
Local:
Adviser: Jiang, Peng.

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UFE0043111:00001


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1 F ABRICATION OF BIOINSPIRED NANOSTRUCTURE D MATERIALS VIA COLLOIDAL SELF ASSEMBL Y By WEI HAN HUANG A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 201 1

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2 2011 Wei han H ua ng

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3 To my family and my love, Chia yi ng, who give me endless support

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4 ACKNOWLEDGMENTS First, I would like to thank my advisor, Dr. Peng Jiang for his directio n He always guided me to discover new ideas in scientific research and give me the freedom and support to do various exploration s. Your invaluable discussions and suggestions not only enrich my knowledge also inspire my research. Also the most important o ne is that He offer ed me a chance to study in University of Florida. I would like to a cknowledge my committee members : Dr. Yiider Tseng Dr. Kirk Zeigler and Dr. Youping Chen. Dr. Tseng is my mentor in scien ce and spirit He encouraged me and showed me ho w to be a scientist and engineer. Dr. Kirk Zeigler and Dr. Youping Chen provided me enormous background knowledge for my study. To all my group members, especially Tzung hua Lin, Inkook Jun and Xuan Dou, I appreciate your greatest help and collaboration T o Dave Jackson, Kerry S iebein, Paul Carpinone Gill Brubaker I am grateful to have you guide me operate instrument s and retrieve data. To all my friends, especially Kai Wei Wang Ming che Yang, Yang yoa Lee and Chien Fong Lo we share lots of good time an d bad time together with their help in all aspects of my research. I express my gratitude to my parents and my sister who supported me endlessly throughout my Ph.D. and always encourage d me to continue my dreams. I especially thank my love Chia ying for accompany ing and support ing me at this hard time I will always love you. Finally, I acknowledge the University of Florida for financial support and friendly environment for carrying out research.

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5 TABLE OF CONTENTS page ACKNOWLEDGMENTS ................................ ................................ ................................ .. 4 LIST OF TABLES ................................ ................................ ................................ ............ 7 LIST OF FIGURES ................................ ................................ ................................ .......... 8 ABSTRACT ................................ ................................ ................................ ................... 11 CHAPTER 1 INTRODUCTION ................................ ................................ ................................ .... 14 1.1 Strong Natural Materials ................................ ................................ ................... 15 1. 2 Synthesize Bioinspired Composite by Assembly Techniques ........................... 21 1.3 Complex Periodic Nanostructures ................................ ................................ ..... 23 2 CONVECTIVE SELF ASSEMBLY GIBBSITE NANOPLATELET ........................... 28 2.1 Introduction and Background ................................ ................................ ............ 28 2.1.1 Gibbsite Nanoplatelets ................................ ................................ ............ 28 2.1.2 Convective Assembly ................................ ................................ .............. 29 2.2 Experimental ................................ ................................ ................................ ..... 30 2.2.1 Synthesis of Gibbsite Platelet ................................ ................................ .. 30 2.2.2 Characterization of Gibbsite Platelets ................................ ...................... 31 2.2.3 Convective Self assembly Gibbsite ................................ ......................... 31 2.3 Result and Discussion ................................ ................................ ...................... 32 3 DIP COATING ASSEMBLE GIBBSITE NANOPLATELET ................................ ..... 40 3.1 Background ................................ ................................ ................................ ....... 40 3.2 Experimental ................................ ................................ ................................ ..... 40 3.3 Result and Discussion ................................ ................................ ...................... 41 4 FABRICATION OF GIBBSITE POLY MER NANOCOMPOSITE VIA ELECTROPHORETIC DEPOSITION ................................ ................................ ..... 45 4.1 Background ................................ ................................ ................................ ....... 45 4.2 Experimental ................................ ................................ ................................ ..... 45 4.2.1 Electrophoretic deposition ................................ ................................ ....... 45 4.2.2 Mechanical Test ................................ ................................ ...................... 46 4.3 Result and Discussion ................................ ................................ ...................... 47 4.3.1 Electrophoretic deposition ................................ ................................ ....... 47 4.3.2 Polymer filled composite ................................ ................................ ......... 48 4.3.3 Mechanical Test ................................ ................................ ...................... 50

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6 5 FABRICATION OF GIBBSITE PVA NANOCOMPOSITE VIA ELECTROPHORETIC CO DEPOSTION ................................ ................................ 58 5.1 Background ................................ ................................ ................................ ....... 58 5.2 Experimental ................................ ................................ ................................ ..... 58 5.2 Results and discussion ................................ ................................ ..................... 59 6 FABRICATION OF GRAPHE NE OXIDE PVA NANO COMPOSITE S VIA FILTRATION ................................ ................................ ................................ ........... 65 6.1 Introduction and Background ................................ ................................ ............ 65 6.1.1 Graphene/Graphene Oxide for Composite App lication ............................ 65 6.1.2 Graphene based material ................................ ................................ ........ 67 6.1.3 Vacuum Infiltration ................................ ................................ ................... 68 6.2 Experimental ................................ ................................ ................................ ..... 68 6.2.1 Synthesize Graphene Oxide ................................ ................................ .... 68 6.2.2 Preparation and Test of GO/PVA Composites ................................ ........ 69 6.2.3 Instruments ................................ ................................ .............................. 70 6.3 Result and Discussion ................................ ................................ ...................... 71 6.3.1 Dispersion of GO She ets with PVA ................................ ......................... 71 6.3.2 Thermal Properties of GO/PVA Nanocomposites ................................ .... 72 6.3.3 Mechanical Properties of Graphene oxide/PVA Nanocompo sites ........... 73 6.3.4 Mechanical Property Simulation ................................ .............................. 76 7 FABRICATION OF PERIODIC BINARY NANOSTRUCTURE ................................ 94 7.1 Background ................................ ................................ ................................ ....... 94 7.2 Experimental ................................ ................................ ................................ ..... 94 7.2.1 Preparation of Double Layer Colloidal Crystals b y Spin Coating ............ 94 7.2.2 Fabrication of Binary Gold Nanohole Arrays by Using Double Layer Colloidal Crystals as Templates ................................ ................................ .... 95 7.2. 3 Replication of Dimple Lithography ................................ ................................ ................................ ... 95 7.2.4 Specular Reflection Measurements ................................ ......................... 96 7.2.5 Ins trumentation ................................ ................................ ........................ 96 7.3 Result and Discussion ................................ ................................ ...................... 97 7.3.1 Fabricate Binary Nanostructures Via Spin Coating ................................ 97 ................................ 100 8 CONCLUSIONS AND FUTURE WORK ................................ ............................... 109 Conclusion ................................ ................................ ................................ ............ 109 Future Work ................................ ................................ ................................ .......... 111 LIST OF REFERENCES ................................ ................................ ............................. 112 BIOGRAPHICAL SKETCH ................................ ................................ .......................... 122

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7 LIST OF TABLES Table page 6 1 Tensile strength (MPa) ................................ ...................... 78

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8 LIST OF FIGURES Figure page 1 1 Performance of engineered and biological materials in the stiffness toughness domai n ................................ ................................ .............................. 18 1 2 Overall view of hierarchical structure of abalone shell and the comparation of inorganic and organic components. ................................ ................................ .... 19 1 3 Models of biocomposites and e xperimental tensile stress strain curve ............. 20 1 4 Demonstration of Lotus effect ................................ ................................ ............. 26 1 5 Nanostructure on the superhydrophobi c wings of cicada ................................ ... 26 1 6 Nano scale nipples on Moth eye ................................ ................................ ........ 27 2 1 TEM image and SAED pattern of gibbsite plat elets ................................ ............ 34 2 2 SEM image s of stacked gibbsite nanoplate ................................ ........................ 35 2 3 Sketch of the particle and water fluxes in the vicinity of monolaye r particle arrays growing on a substrate plate that is being withdrawn from a suspension.. ................................ ................................ ................................ ....... 36 2 4 SEM images of deposited gibbsite layers. ................................ .......................... 37 2 5 Thickness of deposit film versus the volume fraction of gibbsite suspension. .... 38 2 6 XRD patterns of gibbsite layers via convective self assembly ............................ 39 3 1 The s cheme of dip coating gibbsite platelets on substrate ................................ 42 3 2 P hotograph s of dip coating gibbsite ................................ ................................ ... 43 3 3 XRD patterns of multiple time coating of gibbsite ................................ .............. 44 3 4 Thickness of deposit versus dipping time ................................ ........................... 44 4 1 Schematic diagram of electrophoretic deposition of gibbsite nanoplatelets and photograph of dried gibbsite layers and sandwich cell. ............................... 52 4 2 versus the conc entration of colloidal gibbsite suspensions ................................ ................................ .......................... 53 4 3 Electrophoretic assembly of gibbsite nanoplatelets ................................ ............ 54 4 4 Free standing gibbs ite ETPTA nanocomposite ................................ .................. 55

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9 4 5 Thermogravimetric analysis of the gibbsite ETPTA nanocomposite ................... 56 4 6 The SEM image of p orous mem brane with stacked hexagon shaped pores ...... 56 4 7 gibbsite ETPTA nanocomposite, and TPM ETPTA nanocomposite ................. 57 5 1 Electrodeposited gibbsite PVA nanocomposite ................................ .................. 62 5 2 XRD patterns of an electrodeposited gibbsite PVA composite ................... 63 5 3 Deposit weight versus electrophoretic duration ................................ .................. 63 5 4 Thermogravimetric anal ysis of the nanocomposite ................................ ............. 64 6 1 Mother of all graphitic forms and the proposed schematic (Lerf Klinowski model) of graphene oxide structure ................................ ................................ .... 79 6 2 The flow chart of filtration process and s chematic diagram of flow induced pumping system ................................ ................................ ................................ 80 6 3 Images of elastic, free standing GO/PVA compo site and t ensile test ................. 81 6 4 The AFM image of GO sheets ................................ ................................ ............ 81 6 5 Photographes of p ure GO dispersion and GO/PVA suspension ........................ 82 6 6 The SEM image of cross sec tional GO/PVA composite. ................................ .... 82 6 7 The XRD patterns of casted PVA, filtrated GO paper, and GO PVA nanocomposite. ................................ ................................ ................................ .. 83 6 8 TGA results of GO/PVA composites, GO/P VA composites with 2ml of GA, GO paper and casted PVA. ................................ ................................ ................ 85 6 9 The stress strain curve of reference mater ials and GO/PVA nano composites .. 86 6 10 The stress strain curves of nanocomposites: A) GO/PVA without GA, B) GO/PVA with 2ml of GA, C) GO/PVA with 3ml of GA ................................ ......... 87 6 11 The stress strain curves of GO/PVA composites ................................ ................ 88 6 12 The U ltimate tensile strength and ultimate tensile strain of GO/PVA composite ................................ ................................ ................................ .......... 89 6 13 the regions of the stress strain diagram of PVA/GO composite and the schematic models of three regions ................................ ................................ ..... 90

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10 6 14 Model of long polym ers and force extension curves for different kinds of polymers and a sche matic illustration of protein modules ................................ .. 91 6 15 The simulations of ultimate tensile strength of GO/PVA composite .................... 93 6 1 6 The schem atic models of operative failure mode: platelets fracture and platelet pull out. ................................ ................................ ................................ .. 93 7 1 Schematic illustration of the templating procedures for fabricating intercalated hexagonal arrays ................................ ................................ .............................. 103 7 2 SEM images of intercalated hexagonal arrays ................................ ................. 104 7 3 SEM images of binary arrays after a second oxygen plasma etch ................... 105 7 4 SEM image of a binary gold nanohole array ................................ ..................... 106 7 5 SEM images illustrating the progre ssive fabrication procedures for binary ................................ ................................ 107 7 6 Experimental and RCWA simulated normal incidence specular reflection fro m a bare glass substrate ................................ ................................ .............. 108

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11 Abstract of Dissertation Presented to the Graduate School of the University of Florida in Partial Fulfillment of the Requirements for the Degree of Doctor of Philosophy FABRICATION OF BIOINSPIRED NANOSTRU CTURE D MATERIALS VIA COLLOIDAL SELF ASS EMBL Y By Wei Han Huang August 2011 Chair: Peng Jiang Major: Chemical Engineering Through million s of years of evolution, n ature creates unique structure s and materia ls that exhibit remarkable performance on mecha nic al s, optic al s, and physic al p roperties For instan ce nacre ( mother of pearl) bone and tooth show excellent combination of strong mineral s and elastic protein s as reinforced materials S tructured b light or absorb light without dyes. Lotus leaf and cicada are superhydrophobic to prevent water accumulation The principle s of particular biological cap abilities, attributed to the highly sophisticated structures with complex hierarchical designs have been extensively studied Recently, a large variety of novel materials have been enabled by natural inspired designs and nanotech nologies Th e se advanced materials will have h uge impact on p ractical application s We have utilized bottom up approach es to fabricate nacre like nano composite s with brick and mortar structures First we use d s elf assembly processes including convective self assembly, dip coating, and electrophoretic deposition to form well oriented layer structure of synthesized gibbsite ( aluminum hydroxide ) nanoplate let s Low viscous m onomer was permeated into layered nanoplatelets and followed by

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12 photo curing. G ibbsite polymer composite displays 2 times higher tensile strength and 3 times higher modulus when compared with pure polymer. M ore improvement occur red when surface modified gibbsite platelets were cross linked with the polymer matrix. We observe d ~ 4 times higher strength and nearly 1 order of magnitude higher modulus than pure polymer. To further improve the mechanical strength and toughness of inorganic organic nanocomposites we exploited ultrastrong graph ene oxide (GO), a single atom thick hexagonal carbon sheet with pendant oxidation groups GO nanoc omposite is made by co filtrating GO/ polyvinyl alcohol suspension on 0.2 m p ore sized membrane It shows ~2 times higher strength and ~15 times higher ultimate strains than nacre and pure GO paper (also synthesized by filtration ) Specifically, it exhibits ~30 times higher fracture energy than filtrated graphene paper and nacre, ~ 100 times tougher than filtrated GO paper. Besides reinforced nanocomposites, we further explored the self assembly of spherical colloids and the templating nanofabrication of moth eye inspired broadband antireflection coatings. B inary crystalline structu res can be easily accomplish ed by spin coating double layer nonclose packed colloidal crystals as templates followed by colloidal templating T he polymer matrix between self assembled colloidal crystal has been used as a sacri fi cial template to de fi ne the resulting periodic binary nanostructures, including intercalated arrays of silica spheres and polymer posts, gold nanohole arrays with binary sizes, and dimple The binary anti coatings.

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13 N atural optical structure s and nano composite s teach us a great deal on how to create high performance a rtificial material s The bottom up technologies developed in this thesis are scalable and compatible with sta ndard industrial processes, promising for manufacturing high performance materials for the benefits of human beings.

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14 CHAPTER 1 INTRODUCTION Some materials and structures in nature exhibit great properties which are similar or superior to man made ones. Th ey usually apply alternative strategies compared to traditional artificial materials and provide inspiring ideas for the creation of new substance s. Therefore we would like to use selected components and convenient assembly approaches instead of biological materials and growth for fabrication bio memetic materials. H owever, the se biomaterials always have a hierarchical complex structure ranging from micrometer to nanometer scale which is hard to imitate with general process. Recently self assembly of nano s cale colloidal particle provide a way to manufacture well ordered nanostructure s. The se research achievements provide lots of potential applications in e lectronic s, optic s, mechanics among other fields. Several examples will be introduced in this thesis: b inary array nanostructure, made by colloidal templating, exhibits and offers applications for plasmonic devices exhibit improved performance over unitary hemispherical nipples. Na nocomposites, which combined strong and tough characteristics, are made by polymers and gibbsite platelets or graphe n e oxide sheets. Here we use d electrostatically stabilized gibbsite nanoplatelets with well iented assembly of plate like colloids. Convective assembly, dip coating and electrophoretic deposition ha ve been applied to assemble gibbsite nanoplatelets t o large area, orientated layer films. A simple spin coating process ha s also been developed to in ltrate the interstitials between the assemb led nanosheets to form a rtificial nacreous nanocomposites. The resulting self were transparent and exhibit

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15 chemica lly functionalize the surface of the gibbsite nanoplatelets to facilitate the formation of covalent linkage between the ceramic platelets and the polymer matrix. This further reinforces these biomimetic nanocomposites. E ven so this strategy of fabricating nanocomposite s is still restricted by low viscous monomer which is easy to infiltrate. Polyvinyl alcohol ( PVA ) a high molecule weight strong polymer, was introduced to make G ibbsite PVA composite by co electrodeposition. Although t he resulting self stand ing nanocomposite lm was optically transparent and fl exible with high weight percentage of inorganic platelets its mechanical performance is still poor. However, it gave us a new direc tion to blend polymer with strong, high aspect ratio materials by the co deposit ion proc esses. Recently, g raphene and graphe n e oxide (GO) dre w my attention because of their ultrastrong mechanical properties Water suspensions of GO and PVA are able to be well mixed under low concentration without sever agglomeration; and for m a well order ed homogenous nanocomposite GO/PVA nanocomposite, like nacre, is an example of combining the strength of inorganics and the elasticity of organics to a high toughness material. 1.1 Strong Natural M aterial s Materials with s t rong and flaw t olerant character istics are considered as optim um one s to apply on mechanical fields In reality these two properties are usually alternative. The mechanical performances of synthetic materials, illustrated in Figure 1 1 A demonstrate the characteristic c urves in the stiffness toughness diagram with [1] Ceramics are strong but easy to fracture with surface flaws and cracks M ost composites and polymers are flaw tolerant but not a m atch on strength of metal alloys Steel and some me tal alloys, recognize d as tough

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16 substance s have been wide ly used for infrastructure, tools, machines and weapons However m ost engineered materials lie on an inverse relation between increasing toughness as decreasing stiffness. In contrast, many biolog ical materials exhibit remarkable combinations of stiffness strength and fracture toughness. The mechanical performances of biological materials, illustrated in Figure 1 1 B demonstrate the characteristic relationship in the stiffness toughness diagram [1] Protei n and skin on t he upper left corner of map show high toughness with low stiffness which is similar to ela stic polymers. Calcite and calci um phosphate on t he lower right corner of the chart are stiff but brittle minerals Bone, dentin, and mollusk shell se en in the upper right part of the map perform both stiffness and toughness for structural support or armored prot ection Generally, the se materials are buil t with relatively complex s tructures organized over multi scale hierarchical structures and turn b rittle minerals into much tougher material with a few percent additions of biopolymers Red Abalone nacre is an excellent example of such high performance natural materials. I t is composed of at least 95% of brittle minerals such as calcium carbonate (CaCO 3 ) in the form of hexagonal aragonite platelets separated by thin, intermediate biopolymer layers consisting of proteins or polysaccharides (Figure 1 2A ) [2] The structure of nacre is established with to combine the strength of the mineral platelets with the high elasticity of the p rotein layers While nacre is under tensile force in longitudinal direction of platelets as Figure 1 3 A m ost of the load can be carried by the mineral platelets whereas the protein transfers load via the high shear zones. [3] Therefore the loading is scatter ed to the entire material and the whole structure of nacre will elongate until fractured w ith r elatively large failure strains

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17 (Figure 1 3B ) B esides the ordered layer structure there are additional factors that mechanical properties which exceed the simple additive rule s for nano composites such as h ierarch ical structures over several length scales thickness of the aragonite platelets nanoasperities on the platelets and mineral bridges between them. [4 8] As a result, the improvement of mechanical properties is signi ficant: Although identical due to the small protein weight fraction 2 fold increase in strength and 8 fold increase of fracture toughness over monolithic CaCO 3 are observe d (Figure 1 2B ) [2] With experimental and theoretical research in these microstructure and mechanisms a n ovel concept was generate d to design strong tough artifi cial materials by mimicking natural materials

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18 Fig ure 1 1 Performance of engineered and biolog ical materials in the stiffness toughness domain. a ) depict toughness and stiffness values for synthetic materials, such as metals, alloys and ceramics b ) compares the toughness and stiffness properties for a number of biological materials such as antler, dentin, bone and enamel. [1]

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19 A B Figure 1 2. A ) Overall view of hiera rchical structure of abalone shell, showing mesolayers, mineral tiles, and tile pullout in a fracture region. [4] B) the compared to values of th eir inorganic (aragonite, CaCO3) and organic (protein) components. [2]

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20 A B Figure 1 3 A) Models of biocomposites. (a) P erfectly staggered mineral inclusions embedded in protein matrix. (b) A tension shear chain model of biocomposites in which the tensile regions of protein are eliminated to emphasize the load transfer within the composite structure. [3] B ) Experimental tensile stre ss strain curv e for nacre and B ) associated deformation modes. [1]

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21 1.2 Synthesize Bio inspired Composite by Assembly Technique s Nacre a tough bio material which perfectly combine hard and soft constituent provide s a blueprint to create high performance composite Bioinspired c omposite also follows the rule of mixture Theoretically the maximum mechanical properties of composite depend on volume fraction s and mechanical properties of individual components. [9] In order to reach the best performance, t he structural parameters : the platelet aspect ratio platelet dispersion, platelet alignment, and polymer to platelet interfacial stress transfer must be o ptimize d to take advantage s of stiffness of platelets and elasticity of polymer to create strong tough material [10] N ano scale platelet particle and elastic polymer are used as main components to replace mineral and pro tein t o mimic biological material s A g ood filler material for biomemetic composite is consi dered as a strong, high aspect ratio, and single crystal inorganic material. Ceramics such as clays and layered silicates [7 11] or inorganic oxide flake like Titanium dioxide [ 2] and Aluminium oxide [9, 12] have been used in research and industry Polymer s for biomimic composite should offer elasticity, slight stiffness, and good adhesive to filler. Polyvinyl alcohol (PVA) [13] chitosan [9] polyelectroly Poly(methyl methacrylate) (PMMA ) [12, 14] P oly ( diallyldimethylammonium chloride ) (PDDA) [15, 16] and etc. have been used for laboratory research. P olymer to platelet inter force s are able to strengthen with surface modification of platelet or chemical cross linking between two components. [13, 17, 18] However w ell dispersion and alignment of nano scale platelets in polymer matrix are still difficult due to aggregation of nano sized components [10] H igh viscosities of p articles matrix mixture due to t h e relatively strong inter forces between them make the processing of these materials quite difficult Therefore nano composite s with high volume fraction filler

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22 usually results relatively weak mechanical property compared with those predicted by theory [10] Oriented assembly of platelet like particles is particularly interesting as it enables the production of polymer composites with greatly improved mechanical properties. There are several sequential approaches to mimic nacre structure such as layer by layer (LbL) [9, 13, 16] ice templating [12, 1 9] electropho r e tic co deposition [20, 21] doctor blade coating and paper making [17] W ith the aid of assembly approaches, t he mechanical properties are remarkably high and sometimes can exceed those of nacre For instant, Nicholas Kotov and c olleagues demonstrate the outstanding mechanical properties of a multilayer montmorillonite (MTM) clay/polymer nanocomposite, prepared by a layer by layer assembly process and reinforced with glutaraldehyde crosslinker [13] Another route from Lorenz J. Bonderer is based on repeated monolayer transfer of micrometer sized Al 2 O 3 platelets and spin coating of chitosan layers. S cant toughness improvements could be achieved with relative low fraction of platelets. [9] Although the layer by layer repeating process has some drawback s such as being inherently laborious and time consuming i t still offer s a feasible approach to fabricate highly orie nted, ultrastrong and stiff nanocomposites Other assembly techniques also give possibilities of more efficient, high thr oughput fabrication with relative accept able characteristic s.

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23 1. 3 Complex Periodic Nanostructures Some natural materials with p er iodic n anostructures show ultimate performance s that are incomparable to man made substances For example, the structure on a lotus leaf (Figure 1 4 ) and nanopillars on cicada wings (Figure 1 5 ) makes them superhydrophobi c and nonwetted by de w and rain and uncontaminated by the dust [22 26] The hierarchical structure of moth s eye cover ed by sub micro nipples (Figure 1 6 ) eliminates reflections which allows moth to see well in the dark and elimilates r eflections which give its location away to predators [27 29] Facile production of complex periodic nanostructures over a large area is of great technological importance in pushing those unique competence s to practi cal applications. Except imitation of p eriodic n a tural structures complex periodic nanostructures are also valuable in developing novel electronic, optical, and magnetic devices. [30 32] For instance, subwavelength structured metal nanohole arrays exhibit extraordinary optical transmission, which is promising for realizing miniaturized devices for all optical integrated circuits. [33 35] Other than simple hexagonal and square arrays, complex nanohole geometries (e.g., binary arrays) have attracted great scientific and technological interest as they are promising structures for developing more efficient optical devices and biosensors, [36, 37] as well as a fundamental understanding of surface plasmo n enabled optical transmission. [38, 39] Dielectric binary periodic nanostructures have also been explored for creating superhydrophobic and superoleophobic surfaces, [40, 41] photonic crystals, [42] and ant [43] Creation of periodic b inary nanostructures with subwavelength scale resolution is not a trivial task for the current nanofabrication techniques, such as electron beam lithography and focused ion beam. Although arbitrary nanostructures can be fabricated,

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24 attaining high throughpu t and large area fabrication in an inexpensive manner continues to be a major challenge with these top down techniques. [44] Bottom up self assembly and subsequent templating nano fabrication provides a much simpler and inexpensive alternative to nanolithography in fabricating periodic nanostructures. [45 49] Self assembled block copolymers and surfactants enable the creation of sub 50 nm periodic structures. [50, 51] Colloidal lithography provides a ve rsatile approach for making a large variety of periodic microstructures by using self assembled monolayer or multilayer colloidal crystals as deposition or etching masks. [45 48] Binary periodic nanostructures, such as metal nanoparticle arrays, have recently been demonstrated by using angle resolved colloidal lithograp hy. [52, 53] Unfortunately, most of the current co lloidal assembly techniques suffer from low throughput. It usually takes hours to days to assemble a centimeter sized crystal. In addition, current colloidal assembly techniques (e.g., convective self assembly, [54 56 ] g ravitational deposition, [57] and p [58] ) are not compatible with standard microfabrication, limiting the on chip integration of practical devices. Moreover, only close packed colloidal crystals are available thr ough traditional self assembly. [59, 60] But in many cases, nonclose packed crystals are preferred to expand the versat ility of colloidal lithography. [61 63] Furthermore, in conventional colloida l lithography, only self assembled colloidal spheres are used as structural templates. We introduce spin coating [64, 65] and colloidal templating technology [62, 66, 67] for fabricating periodic binar y nanostructures and enabling on chip creation of a wide range of periodic microstructur es, such as metal hole arrays, [62, 67] attolitre microvial arrays, [68] and two d imensional (2 D) magnetic dots, [66] which are not easily available

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25 by traditional top down and bottom up approaches. Double layer, nonclose packed colloidal crystals prepared by a spin coating technology are used as etching masks to lithography, the polymer ma trix between the self assembled colloidal crystal has also template. This bottom up technique is compatible with standard top down microfabrication, enabling scalable production of periodic binary nanostructures that have impo rtant technological applications ranging from plasmonic cells.

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26 Figure 1 4 Demonstration of Lotus effect: structure on lotus leaf cause superhydrophobic ability which make drop s of water without wetting. Figure 1 5 Nanostructure on the superhydrophobic wings of cicada (Cicada orni). [22, 26]

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27 Figure 1 6 Nano scale nipples on Moth eye eliminate the reflection of light. [27]

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28 CHAPTER 2 CONVECTIVE SELF ASSEMBLY GIBBSITE NA NOPLATELET 2.1 Introduction and Background 2.1 1 Gibbsite Nanoplatelets G ibbsite (Al(OH) 3 ) n anoplatelets with uniform shape and dispersibilty provide a model as a nanoplatelet filler of a nanocomposite. S ynthesized by hydrolysis of Al(OH 2 ) 6 3+ at 85 C [69, 70] gibbsite nanoplatelets exhibit a wel l defined hexagon shape and t he aspect ratio ( > 10) which is close to that of natural aragonite platelets in nacre. [71] The diameter and thickness of the gibbsite nanoplatelets can be further controlled by se eded growth. [72] Figure 2 1 shows a typical TEM ima gibbsite nanoplatelets. The as synthesized inorganic platelets have well hexagonal shape and are relatively uniform in size. The average diameter of the nanoplatelets is ca. 200 nm. The thickness of the gibbsite platelets is estimated to be ~10 nm by using cross sectional SEM analysis as shown in Figure 2 2A Those platelets are well dispersed and aggregated particles are rarely seen on TEM grids. Electrostatic repulsion is responsible for the observed colloidal stability and the zeta p otential ( ) of gibbsite particles in deionized water is me asured to be +40.5 2.3 mV by fi tting experimental The gibbsite structure is a stacking of Al OH layers and each Al 3+ is surrounded by six hydroxyl groups. The reaction of s urface hydroxyl groups with water makes the nanoplatelets highly charged and well dispersed in water and alcoholic suspensions. The surface hydroxyl groups also facilitate the chemical modification of the particle surface to render different functionality [73] Selected area electron diffrac tion from individual nanoplatelet (Figure 2 1 ) indicates the as synthesized gibbsite particles are sin gle crystalline. By using

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29 gibbsite nanoplatelet as a model system, Lekkerkerker et al. have extensively exploited the liquid crystal phase transition in suspensions of plate like particles [74 77] Opal like columna r gibbsite colloidal crystals have also been demonstrated by forced sedimentation. [78, 79] In summary, synthesized gibbsite demonstrate th e se properties: Shape: High aspect ratio platelet, low size deviation D ispersion : High e lectrostatic repulsion Flawless: Uniform, single crystalline structure Th e se properties are similar to other common nanoplatelet filler s like montmorillonite clay (MTM) and layered double hydroxides (LDHs) [80] and satisfy the requir ements of high performance organic inorganic nanocomposites. 2.1.2 C onvective A ssembly To grow thin particle layers, the simplest way is to dip a wettable solid plate into a suspension of particles and keep it stationary. Monolayers and successive multilay ers of particle arrays spontaneously start to form on the plate surface from the plate suspension air contact line down to the bulk susp b ension. This technique is so called convective assembly (also known as vertical deposition) [81 84] The profile of growing particle arrays from a bulk suspension onto a flat substrate in the vicinity of the array's leading edge is shown in Figure 2 3 The formation of layered arrays can be split in two main stages: (1) convective transfer of particles from the bulk of the suspension to the thin wetting film due to solvent evaporation from the film surface and (2) interactions between the particles that lead to specific ordering. The schematic can be used to derivate the following e quation (equation 2 1) [81] : (2 1)

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30 Where is the product of the array growth rate. is the value of the coefficient of proportio nality which depends on the particle particle and particle substrate interactions and should vary from 0 to 1. is average evaporation length. is the evaporation flux,. is the particle volume fraction in the suspension, is the thickness of the array, and is array density. Under the proper condition, evaporation of solvent leads to the deposition of the ordered three dimensional packing of mo re or less uniform thickness on the substrate. After the deposit array is dried, the particles adhere to each other and to the substrate so the film of packed particle can be easily handled. E quation 2 1 c an be deriv ed to get equation 2 2 : (2 2) If the evaporation rate of solvent is stable, the factors in brackets are constant. The thickness of the array is proportional to particle volume fraction (only for low volume fraction ). 2.2 Experimental 2.2 .1 Synthesis of Gibbsite Platelet The gibbsite nanoplatelets are synthesized using the following method [69] To 1 L of deionized water, hydrochloric acid (37%, 0.0 9 M), aluminum sec butoxide ( 95%, 0.08 M), and aluminum isopropoxide ( 98%, 0.08 M) are added. The mixture is stirred for 10 days and subsequently heated in a polyethylene bottle in a water bath at 85 C for 72 h. After cooling to room temperature, dispe rsions of gibbsite nanoplatelets are centrifuged at 3500 g for 6 h then the sediments are redispersed in deionized water. To

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31 completely remov e the unreacted reactants and concentrating the nanoplatelets, this process is repeated five times. 2.2 .2 Character ization of Gibbsite Platelets The platelets were further characterized with Scanning Electron Microscopy (SEM), Laser Diffraction for particle size analysis, X Ray Diffraction (XRD), Transmission Electron Microscopy (TEM), Atomic force microscopy (AFM), an d Electrophoretic Mobility (Zeta Potential). The surface morphology of synthesized platelets was investigated using SEM ( Figure 2 2 )and TEM images( Figure 2 1 ) [18] The platelets are hexagonally shaped and are relative ly uniform in size. The diameter of the nanoplatelets is measured to be 188 40 nm by averaging over 100 particles from the TEM micrographs. TEM images also reveal that the nanoplatelets tend to align parallel to the surface of TEM grids. It is very rare to find nanoplatelets oriented perpendicularly to the TEM grid surface as shown by the red arrow pointing to such a particle in Figure 2 1 AFM experiments show the platelet thickness ranges from 10 to 15 nm. The aspect ratio of major dimension to thicknes s ranged from 22:1 to 10:1. The purified gibbsite nanoplatelets are electrostatically stabilized, and the zeta potential ( ) of the colloids in deionized water is measured to be +40.5 aqueous and alcoholic dispersions, and aggregated particles are rarely seen in TEM images. The selected area elec tron diffraction (SAED) pattern from a single platelet ( Figure 2 1 ) indicate s that the as made gibbsite nanoplatelets are single crystal. 2.2.3 Convective S elf assembly Gibbsite By applying convective self assembl y platelet like particle could also form oriented layer structures. A basic apparatus show as Figure 2 3 B : clean glass slide or

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32 silicon wafer as substrate is placed into 15 ml of purified gibbsite suspension in a clean scintillation vial. The gibbsite nan oplatelets are synthesized as previous describe [69] Gibbsite suspension is centrifuged and redispered by ultrasonicator to designed volume fraction with DI water or 200 proof ethanol as solution (For gibbsite ethanol suspension, centrifuge disperse process is repeated for several times to remove most of water) Glass slide and silicon wafer are cut to 3 x 0.5 cm 2 to fit 20m L glass vial then they are immersed in a 80C 3:1 mixture of ammonium hydroxide with hydrogen peroxide (Piranha solution) for 1 hour to remove organic residues and make the surface hydrophilic For gibbsite water suspension, the experiment is heated at 80 C in vacuum oven overnight to deposit gibbsite layer. For gibbsite ethanol suspension the vial of gibbsite ethanol suspension is placed on a vibration free bench in a temperature controlled laboratory (22 1 C) for 5 7 days and c overed by a 1200 m L crys tallizing dish to keep out external airflow and contamination. Scanning electron microscopy (SEM) is carried out on a JEOL 6335F FEG SEM. X ray diffraction (XRD) spectrum is obtained with P h ilips XRD 3720 equipment. A Cu K 1 (k = 1.54049 ) radiation is used from 10 to 70 with a scan rate of 2.4 /min. 2.3 Result and Discussion Figure 2 4 shows SEM images of the cross section view of deposited gibbsite layers made from 0.75 to 2 vol% of gibbsite water suspension Each of the SEM samples is taken in the center section of deposition with the clear cut of substrate. Th ough th e sample preparation cause s some damage s and collapse s the alignment of gibbsite layer structure is obvious in most parts. The t hickness of gibbsite platelet layers is thicker as the v olume fraction of the gibbsite in suspension increase s Figure 2 5

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33 show s the plot of thickness versus volume fraction of gibbsite platelets in suspension. The red lines show approximate linear relationship s between thickness of deposit and volume fraction of gibbsite in suspension that fit the prediction (E quation 2 2) The layer growth rate from gibbsite water suspension is larger than that from gibbsite ethanol suspension because the evaporation rate of water at 80C is higher than that of ethanol at room temperature The evaporation rate can be determined by estimating time for evaporating solvent (hours for water compare to days for ethanol). The oriented assembly of high aspect ratio gibbsite nanoplatelets is further confirmed by X ray diffraction. Figu re 2 6 displays a XRD spectrum of 2 vol% of gibbsite water suspension shows in Figure 2 4 We only observe (002) and (004) diffraction peaks from gibbsite single crystals. As the crystallographic c axis of single crystalline gibbsite is normal to the plate let surfaces, the (002) and (004) reflection are from gibbsite platelets oriented parallel to the substrate surface. [85] This strongly supports the presence of macroscopic alignment of gibbsite nanoplatelets in the method of convective self assembly. With the r esult we find that the technique of convective self assembly could easily make good orientated gibbsite deposit with micrometer scale thickness. However, due to small scale apparatus volume fraction of gibbsite increased dramatically with solvent evaporat ion. It results the thickness diff erence between top and bottom of a sample Large scale sampling and stable experiment condition s could eliminate the edge to preserve the uniform region.

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34 Figure 2 1 TEM image and SAED pattern of gibbsite platelets, the red arrow points nanoplatelets oriented perpendicularly t o the TEM grid surface. [18]

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35 A B Figure 2 2 SEM image s of A) cross section and B) top view of stacked gibbsite nanoplate

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36 A B Figure 2 3 A) Sketch of the particle and water fluxes in the vicinity of monolayer particle arrays growing on a substrate plate that is being withdrawn from a suspension. The inset shows the menisci shape between neighboring particles. Here, is the substrate withdrawal rate, v c is the array growth rate, j w is the water influx, j p is th e respective particle influx, j e is the water evaporation flux, and h is the thickness of the array. [81] B) Apparatus of convective self assemble gibbsite

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37 Figure 2 4 SEM images of deposited gibbsi te layers. The volume fraction s of the gibbsite in DI water range from 0.75% to 2 %.

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38 A B Figure 2 5 Thickness of deposit film versus the volume fraction of gibbsite suspension in A) water and B) ethanol.

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39 Figure 2 6 XRD patterns of gibbsite layers via convective self assembly the (002) and (004) reflection are from gibbsite platelets oriented parallel to the substrate surface.

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40 CHAPTER 3 DIP COATING ASSEMBLE GIB BSITE NANOPLATELET 3.1 Background Dip coating is a simple and popular way of creating thin films for research purposes. Uniform films of liquid for solidification into a coating can be applied onto flat or cylindrical substrates The film is deposi ted by withdrawing a substrate from a liquid bath With steady withdraw velocity; film thickness is set by the competition among viscou s force, capillary (surface tension) force and gravity. Thickness and uniformity can also be sensitive to flow conditions in the liquid bath and atmosphere. Generally the faster substrate is withdrawn, the thicker the film is deposited. This can be achieve d by using volatile solutes and combining rapid enough drying with the basic liquid flow. [86, 87] 3.2 Experimental The cycle of dip coating is followed by the process of dippin g, withdraw and drying ( Figure 3 1). We use modified gibbsite ethanol suspension for dip coating. The surface hydroxyl groups of gibbsite nanoplatelets can be easily modified by reacting with 3 (trimethoxysilyl)propyl methacrylate (TPM) through the well e stablished silane coupling reaction. [88] 10 ml of TPM was mixed for 1 hour in 100 ml of a water methanol solution in order to fully hydrolyze the silane species (water versus methanol in volume ratio = 3:1). Surface modification was then accomplished by adding a 100m L dispersion of gibbsite (~1vol% aqua solution) into the hydrol yzed TPM solution. The suspension was stirred for 30 minutes at 40 C. The modified platelets were washed by repeated centrifugation with pure ethanol and finally concentrate to a stock suspension in 200 proof ethanol. We use a programmable syringe pump to imitate dip and withdraw

PAGE 41

41 process, which can provide precise control on velocity. In a sample preparation, a clean glass slide was immersed in TPM modified Gibbsite suspension of 4.5 and 3.5wt% with 10cm/min of di pping rate and withdrawing rate ( Figure 3 2 A). After withdrawal glass slides are static and dried in air. In order to fabricate thicker film, this dip withdraw dry cycle could be repeated for several times. Scanning electron microscopy (SEM) is carried out on a JEOL 6335F FEG SEM. X ray diffracti on (XRD) spectrum is obtained with Philips XRD 3720 equipment. A Cu K 1 (k = 1.54049 ) radiation is used from 10 to 70 with a scan rate of 2.4 /min. 3.3 Result and Discussion Dip coating is a simple coating process to make film deposit from a liquid ba th of sol gel or suspension. The film thickness depends mostly on withdrawal speed and viscosity of liquid [86] We use constant withdrawal speed 10c m/min to coat a thin film of gibbsite ethanol mixture on glass slide. Figure 3 2B shows the nearly transparent region still contains ethanol before it completel y dries, which means e thanol on the thin film is gradually vaporiz ed from the r im to the center of glass slide This phenomenon, similar to convective assembly [81] indicates that gibbsite arrays growing on a substrate plate are withdrawn from a suspension due to ethanol evaporation. This procedure gi ve s a single deposition cycle and the cycle could then be repeated as necessary to obtain the desired thickness We observe samples that are dip ped in 3.5wt% and 4.45wt% gibbsite ethanol sus pension for 12, 8, 4 times, the se s amples still exhibit similar white smooth surface. The X ray diffraction pattern (Figure 3 3) could show that the ordering of platelet layers remains after multiple dipping The intensity also increasing with

PAGE 42

42 coating times might indicate more gibbsite on substrate This theory is proved by t he thickness of gibbsite deposit increases with dipping times ( Figure 3 4 ) Dip coating could form highly uniform s ediment on entire substrate by simple dip withdraw process es The thickness of the deposit ed layer coul d be controlled by withdraw speed [86] and volume fraction of platelet. Repe tition of dip coating is able to achieve th icker layers with ten of micrometer scale thickness and preserve s packing orientation. Even this repeating process seems low efficiency; it is still a well controlled, uniform, and relatively faster process compared to convective self assembly process The dip coating process can be modified to continuous process to replace repeating dip withdraw steps. This continuous process is possible to fabricate multiple layers of coating of same particle as well as alternative particles for more applications with lar ge scale and high throughput. Figure 3 1. The s cheme of dip coating gibbsite platelets on substrate. A clean glass dips in and withdraws from gibbsite ethanol suspension with liquid suspension film coated. After the film dries, gibbsite platelets ar e deposited on glass slide.

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43 A B C Figure 3 2 A)photograph of dip coating gibbsite ethanol suspension with syringe pump and B)ethanol is evaporating from edge to center. C) Samples dip in 3.5wt% and 4.45wt% gibbsite ethanol suspension for 12, 8, 4 times.

PAGE 44

44 Figure 3 3. XRD patterns of multiple time coating of gibbsite samples Figure 3 4 Thickness of deposit versus dipping time under 4.45wt% and 3.5 wt% of gibbsite ethanol suspension measured by SEM image s

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45 CHAPTER 4 F ABRICATION OF GIBBSITE POLYMER NANO COMPOSITE VIA ELECTROPHORETIC DEPO SITION 4.1 Background Electrophoresis is a well established technology in assembling spherical colloids into highly ordered colloidal crystals [89 91] In this methodolo gy, charged colloids are attracted by electrical force toward the counter electrode and then deposited on the electr ode surface by particle coagula tion [92] Electrodeposition is a simple, inexpensive, and scalable technology that enables rapid production of thick films over large areas. In addition, deposition of metals and conducting polymers in the interstitials of colloids is easily achieved by electrophoresis. This will significantly expand the available materials for the fabrication of layered nano composites. Electrophoretic assembly of nanoclays has previously been tested, but the entrapment of nonplate particles caused by the agglomeratio n of nanoclays deteriorates the layered structure. [93] 4.2 Experimental 4.2.1 Electrophoretic deposition Electrophoretic depositio n of gibbsite platelets is carried o ut using a parallel plate sandwich cell, which consists of an indium tin oxide (ITO) working electrode, a gold counter electrode, and a polydimethylsiloxane (PDMS) spacer, respectively (Figure 4 1 ) The gold electrode is prepared by sputtering deposition of 20 nm of chromium and 200 nm of gold on a (100) silicon wafer. PDMS is used as a spacer to get an active area and a cell gap of 2.2 mm. The gibbsite nanoplatelets are synthesized as previous describe [69] Aqueous suspensions of gibbsite nanoplatelets with different weight percentage are used. The bath solution is gibbsite nanoplatelets dispersed in a water ethanol mixture with volumetric ratio of 1:2. A c onstant voltage of 2.5 V (ITO ve rsus Au)

PAGE 46

46 is applied for 30 min to deposit the positively charged gibbsite nanoplatelets onto the ITO cathode After deposition, the electro proof ethanol and then dried with compressed air. After the oriented assembly, polymer aligned nanoplatelets with photo curable monome rs, followed by photopolymerization. We chose a nonvolatile monomer, ethoxylated trimethylolpropane triacrylate (ETPTA, M.W. 428, viscosity 60 cps, Sartomer USA), to form the nanocomposites. The monomer with 1% photoinitiator (Darocur 1173, Ciba Geigy, Sar tomer USA) is infiltrated into the electroplated gibbsite film, followed by spin coating at 4000 rpm for 1 min to remove exces s monomer, and then polymerized by exposure to ultraviolet radiation. Scanning electron microscopy (SEM) is carried out on a JEOL 6335F FEG SEM. X ray diffraction (XRD) spectrum is obtained with Philips XRD 3720 equipment. A Cu K 1 (k = 1.54049 ) radiation is used from 10 to 70 with a scan rate of 2.4 /min. Thermogravimetric analysis (TGA) is carried out in air with a Perkin Elme r Thermo Gravimetric Analyzer and a platinum crucible between 20 C and 800 C at a heating rate of 5 C/min. 4.2.2 Mechanical Test ETPTA, and TPM ETPTA) are tested using an Instron model 1122 load frame upgraded with an MTS ReNew system and equipped with a 500 g load cell at a crosshead speed of 0.5 mm/min. Testing samples with widths of 1.5 mm and thickness ranging from 30 to 80 m are adhered on homemade sample ho lders with a 20 mm gap using polyurethane monomer as an adhesive and then UV cured. The

PAGE 47

47 thickness of the tested samples is measured by cross sectional SEM to calculate the nal tensile strength. 4.3 Result and Discussion 4.3 .1 Electrophoretic deposition E lectrophoretic deposition method results in the accumulation of ceramic particles at the electrode. Deposit formation is achieved via particle coagulation which come from the application of the classical Derjaguin Landau Verwey Overbeek (DLVO) [94, 95] theory of colloidal stability. The DLVO theory consider s two main forces: double layer repulsion and van Ethanol is add ed to the aqueous dispersions to reduce the dielectric constant of the solvent and thus reduce the electrical double layer thickness of the particles to promote colloidal coagulation on the ITO electrode. [92] Without ethanol, no particle deposits will be adhered on the working electrode after disassembling the electrical cell. The ethanol can also reduce cracking and porosity in the electrophor etically deposited films. The applied electric fi eld strength is ~ 1100 V/m. The electrophoretic velocity of the gibbsite nanoplatelets is estimated to be ~ 7.5 m/s by using the S moluchowski equation (equation 4 1 ): ( 4 1 ) wh 0 is the permittivity of the vacuum, is the solution viscosity, and E is the applied electric field strength. [96] For a 2.2 mm thick sandwich cell, the esti mated time to deposit most particles on the ITO electrode is about 5 min utes agreeing with our experimental observation. The duration of the electrophoretic process is 30 minutes which is long enough ; almost all gibbsite platelets

PAGE 48

48 can be deposited on the ITO electrode. The thickness of the deposits is then linearly proportional to the particle volume fraction of the suspension as shown in Figure 4 2 After electrophoretic deposition, the gibbsite deposits on the ITO cathode are washed with ethanol and the n dried with compressed air. The deposits can be easily peeled off from the ITO surface by using a sharp razor blade, resulting in the formation of self standing fi lms as shown in Figure 4 3 The fi lm is opaque and brittle, and the side facing the ITO cath ode is smoother than the side facing the suspension. The size of the resulting fi lms is solely determined by the dimension s of the ITO electrode. Figure 4 3A depicts a sample with 1.6 0.6 in. 2 size deposited on a 2 1 in. 2 ITO electrode. Figure 4 3B sho ws a top view SEM image of the suspension side of the sample in Figure 4 3A The hexagonal gibbsite nanoplatelets are densely packed and aligned parallel to the electrode surface. The alignment of gibbsite nanoplatelets is further confirmed by the layered structure as shown in the cross sectional SEM image of Figure 4 3C C onvincing evidence of the orientated deposition co mes from the XRD patterns shown in Figure 4 3D Only (002) and (004) peaks are observed in the XRD spectrum. As the crystallographic c ax is of single crystal gibbsite is normal to the plat elet surfaces, the (002) parallel to the electrode surface. [85] Analysis of the half height width of the (002) peak with the Scherrer equation yields an average platelet thickness of 15.1 nm, agreeing with AFM measurement. 4.3 .2 Polymer filled composite P olymer gibbsite nanocomposites are made by fi lling the interstitials between the aligned nanoplatelets with photo curable monomers, spin coating to remov e excess monomer, and polymerized by exposure to ultraviolet radiation. The resulting

PAGE 49

49 nanocomposite fil m beco mes highly transparent (Figure 4 4A ) due to the matching of the refractive index between the gibbsite platelets and the polymer matrix. The normal incidence tr ansmission measurement ( Figure 4 4B ) shows that the free standing nanocompo site film exhibits high transmit tance (>80%) for most of the visible wavelengths. As the reflection (R) from an interface between two materials with refractive indices o f n 1 and n 2 equation (equation 4 2 ) : [97] ( 4 2 ) we can estimate the normal incidence reflection from each air nanocomposite interface to be about 4%. Thus, the optical scattering an d absorption caused by the nanocomposite itself is approxi mately 10%. This suggests that the polymer m atrix has nanoplatelets. The cross sectional SEM image in Figure 4 4C shows that the nano composite retains the layered structure of the o riginal electroplated gibbsite fi lm, and thin wetting layers of ETPTA ( ~ 1 m thick) are observed on the surfaces of the fi lm. The oriented arrangement of the nanoplatelets is also mainta ined throughout the po lymer infil tration process as confi rmed by the distinctive (002) and (004) peaks of the XRD spectrum shown in Figure 4 4D The ceramic weight fraction in the gibbsite ETPTA nanocomposite film is determined by thermogravimetric analysis (TGA) as shown in F igure 4 5 From the TGA curve and the corresponding weight loss rate, it is apparent that two thermal degradation processes occur. One happens at ~ 250C and corresponds to the degradation of the polymer matrix, while another occurs at ~ 350 C and is due to the decomposition reaction of gibbsite : On the basis of the residue

PAGE 50

50 mass percentage (45.65%) and assuming the ash is solely Al 2 O 3 we can estimate the weight fraction of gibbsite nanoplatelets ~ 0.70. Considering the density of gibbsite (~2.4 g/cm 3 ) and ETPTA (~1.0 g/cm 3 ), the volume fraction of gibbsite nanoplatelets in the nanocomposites is approximately 0.50. The complete infiltration of ETPTA between th e electroplated gibbsite platelets is further confirmed by the selective dissolution of gibbsite in a 2% hydrochloric acid aqueous solution. Figure 4 6 shows t his result in the formation of a self standing porous membrane with stacked hexagon shaped pores, which are a negative replica of the assembled gibbsite platelets. 4.3 .3 Mechanical Test The mechanical properties of the biomimetic polymer nanocomposites are evaluated by tensile tests. We compare the tensile strength for three types of thin films, incl uding pure ETPTA, gibbsite ETPTA, and TPM ETPTA. Gibbsite platelets were modified with 3 (Methacryloyloxy)propyl trimethoxysilane(TPM). [88] This results in the formation of surface can be cross linked with the acrylate based ETPTA matrix. The colloidal stabilit y and the surface charge of the resulting nanoplatelets are not affected by this surface potential measurement. Figure 4 7 shows the tensile stress versus strain curve s for the above three types of fi lms. The gibbsite ETPTA nanocomposite displays ~ 2 times higher strength and ~ 3 times higher modulus when compared with pure ETPTA polymer. Even more remarkable improvement occurs when TPM gibbsite platelets are cross linked with the ETPTA matrix. We observe ~ 4 ti m es higher strength and nearly 1 order of magnitude higher modulus than pure polymer. This agrees with early studies that reveal the crucial role

PAGE 51

51 played by the covalent linkage between the ceramic fi llers and the organic matrix in determining the mechanical properties of the artificial nacreous composites. [13] We also conduct a simple calculation to evaluate if the measured mechanical properties of the gibbsite ETPTA nanocomposites are reasonable. For a polymer matrix having a yield shear strength y and strong bonding to the gibbsite nanoplatelet surface (e.g., TPM modified gibbsites), c ) can be calculated using the volume fraction of nanoplatelet (V p ), the nanoplatelet aspect ratio (s), and the tensile stren p m ), as [9, 98] (equation 4 3 ) : ( 4 3 ) For the gibbsite nanoplatelet which has a relatively small aspect ratio (s ~ 12 18), th e factor in equation 3 5 can be estimated as [9] (equation 4 4 ) : ( 4 4 ) From the TGA analysis, the volume fraction of gibbsite nanoplatelets in the po lymer nanocomposite is ~ 0.50. If we take s = 15, eq uation 3 5 can then be simplified as (equ ation 4 5 ): ( 4 5 ) For acrylate based polymer (like ETPTA), the yield shear strength should be close to its tensile strength. Equation 3 7 c ~ This indicates that the strength of the nanocomposite is about fourfold of the strength of the polymer matrix, agreeing with our experi mental results (Figure 4 7 ).

PAGE 52

52 A B Figure 4 1 A) Schematic diagram of elect rophoretic deposition of gibbsite nanoplatelets in a sandwich cell which is composed of Au anode, ITO cathode and PDMS spacer. B) photograph of dried gibbsite layers and sandwich cell.

PAGE 53

53 Figure 4 2 colloidal gibbsite suspensions. The thickness standard deviation for all samples is ca. 10%. [18]

PAGE 54

54 Figure 4 3 Electrophoretic assembly of gibbsite nanoplatelets. (A) Photograph of a free standing gibb s view SEM image of the sample in(A). (C) Cross sectional view of the same sample. (D) XRD patterns of the gibbsite lm in (A). [18]

PAGE 55

55 Figure 4 4 Free standing gibbsite ETPTA nanocomposite. (A) Photograph of a incidence transmissio n spectrum of the sample in (A). (C) Cross sectional SEM image of the same nanocomposite lm. (D) XRD patterns of the nanocomposite sample. [18]

PAGE 56

56 Figure 4 5 Thermogravimetric analysis of the gibbsite ETPTA nanocomposite [18] Figure 4 6 P orous membrane with stacked hexagon shaped pores as a negative replica of the assembled gibbsite platelets.

PAGE 57

57 Figure 4 7. gibbsite ETPTA nanocomposite, and TPM ETPTA nanocomposite. [18]

PAGE 58

58 CHA PTER 5 FABRICATION OF G IBBSITE PVA NANOCOMPOSITE VIA ELECTROPHORETIC CO DEPOSTION 5.1 Background Co deposition of inorganic nanoparticles and polymers (ionic or non ionic) by electric field has also been developed for scalable prod uction of nanocomposites. [21, 99, 100] Here we use a simple electro phoretic deposition technology that enables the creation of inorganic organic nanocomposites with oriented layered nanostructures in a single step. [18, 92] To resolve the colloidal aggregation issue faced by using nanoclays as building blocks, electrostatically stabilized gibbsite nanoplatelets with high aspect ratio are employed as a model system. Though gibbsite nanoplatelets have been utilized in exploring the liquid crystal phase transition in suspensions of plate like particles [75, 76, 101] no attempts have been take n to assemble them into ordered nanocomposites. 5.2 Experimental The gibbsite nanoplatelets are synthesized as previous ly describe d [69] The as made particles are purified in ultrapure w ater (18.2 M cm 1 ) by multiple centrifugation/ redispersion cycles (fi ve times). To prepare the electrophoretic bath solution, 1 ml of 5 wt% polyvinyl alcohol (PVA, Mw 89,000~98,000, Sigma Aldrich) aqueous solution is firstly mixed with 9 ml of 2 wt% gibb site nanoplatelet solution. 20ml of 200 proof ethanol is then added into the abov e suspension. Electrophoretic deposition of PVA gibbsite is performed in a sandwich cell placed horizontally. The bottom and the top of the cell are an ITO working electrode a nd a gold counter electrode, respectively. The ITO coated glass slides (2.5 x 2.5 cm 2 ) with a sheet resistance of 8 are received from Delta Technologies. The gold electrode is prepared by sputtering deposition of 20 nm of

PAGE 59

59 chromium at a deposition rate of 1.6 /s, and then 200 nm o f gold at a deposition rate of ~5 /s on silicon (100) wafer (test grade, n type, Waferne t). Pol ydimethylsiloxane (PDMS, SYLGARD 184, Dow Corning) spacer is used to obtain an active area of 1.5 x 1.5 cm 2 and a cell gap of 2.16 mm. A constant voltage of ~ 2 .5 V (ITO vs. Au electrode) is fi nally applied using an EG&G Model 273A potentiostat/galv a nostat to deposit gibbsite nanoplatelets on the ITO cathode. A fter the electrophoretic deposi tion, the as deposited PVA gibbsite lm is dried in an oven at 80 C. Scanning electron microscopy (SEM) is carried out on a JEOL 6335F FEG SEM. X ray diffraction (XRD) spectrum is obtained with Philips XRD 3720 equipment. A Cu K 1 (k = 1.54049 ) radiation is used from 10 to 70 with a sc an rate of 2.4 /min. Thermogravimetric analysis (TGA) is carried out in air with a Perkin Elmer Thermo Gravimetric Analyzer and a platinum cr ucible between 20 C and 800 C at a heating rate of 5 C/min. 5.2 Results and discussion The electrophoretic deposi tion of gibbsite PVA nanocomposite is carried out using a parallel plate sandwich cell consisting of an ITO cathode, a gold anode, and a PDMS spacer. An ITO anode can also be used to replace the gold e lectrode to conduct the electro phoretic deposition. The positively charged gibbsite nanoplatelets are attracted by the electrical force toward the ITO cathode. As the synthesized gibbsite platelets have positively charged surfaces and almost electrically neutral edges due to their different isoelec tric point ( pH 10 and 7, respectively) [69] they tend to re orient in the electric fi eld with their surfaces facing the ITO electrode. The high molecular weight PVA (Mw 89,000 98,000) is neutrally c harged in the electrophoretic bath. They can be absorbed on the surfaces of gibbsite nanoplatelets and function as water soluble

PAGE 60

60 binders to cement electrodeposited ceramic particles together [92] Ethanol (~ 50% of total volume) is also added to the aqueous colloidal suspensions to reduce the dielectric constant of the sol vent, and thus reduce the electrical double layer thickness of the particles to further promote colloi dal coagulation on the ITO cathode [92] Fig ure 5 1A shows a photograph of a gibbsite PVA nanocomposite formed on an ITO cathode. The fi lm can be easily peeled off from the electrode surface by using a sharp razor blade. The resulting self standing film is fl exible and transparent. Top view SEM image in Fig ure 5 1B illustrates the gibbsite nanoplatelets are preferentiall y oriented with their crystallo graphic c axis perpendicular to the electrode surface. It is very rare to fi nd edge on platelets. The ordered layered structure is clearly evident from the cross section a l SEM images as shown in Figure 5 1 C and D. The orient ed assembly of high aspect ratio gibbsite nanoplatelets is further confi rmed by X ray diffraction. Fig ure 5 2 displays a XRD spectrum of an electrodeposited gibbsite PVA nanocomposite on an ITO electrode. The diffraction peaks from (222), (400), (441), an d (662) planes of the ITO substrate are clearly appeared. Other than ITO diffraction peaks, we only observe (002) and (004) peaks from gibbsite single crystals. As the crystallographic c axis of single crystalline gibbsite is normal to the platelet surface s, the platelets oriented parallel to the electrode surface [85] This strongly supports the macroscopic alignment of gibbsite nanoplatelets in the electrophoretically deposited nanocomposite s. Analysis of the half height wid th of the (002) and (004) peaks with the Scherrer equation [85] yields an average platelet thickness of 10.3 nm, agreeing with cross sectional SEM measurement.

PAGE 61

61 The current electrophoretic deposition technology enables large scale assembly of ordered nanocomposit e lms in a very short time Figure 5 3 exhibits the relationship between the measured weight of deposits on ITO catho de and the electrophoretic dura tion. A weight plateau is reached in ca. 8 min. Experimental obser vation shows almost all gibbsite nanoplate lets have already been deposited in this time interval and the electrophoretic bath changes from turbid to clear. The e lectrophoretic velocity of gibbsite nanoplatelets is estimated to be ~ 8.2 m/s by using the Smoluchowski equation 0 is the permittivity of the vacuum, is the solution viscosity, and E is the applied electric field strength (~1200 V/m in our experiments) [96] For a 2. 16 mm thick sand wich cell, the estimated time to deposit all particles on the ITO electrode is ca. 5 min, reasonably agreeing with the experimental observation. The weight fraction of the inor ganic phase in the electrodeposited nanocomposites can be determi ned by thermogravimetric analysis. Fig ure 5 4 shows the TGA curve and the corresponding weight loss rate for the nanocomposite film as shown in Fig ure 5 1 An apparent thermal degradation process occurs at ~ 250 C that corre sponds to the degradation of t he PVA matrix and the decomposition reaction of gibbsite [69] Based on the residue mass percentage (53.9 6%) and assuming the ash is so lely Al 2 O 3 we can estimate the w eight fraction of gibb site nano platelets i to be 0.825.

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62 Figure 5 1. Electrodeposited gibbsite PVA nanocomposite. (A) Photograph of a composite fi lm on an ITO electrode. (B) Top view SEM image of the sample in (A). (C) Cross sectional SEM image of the sample in (A ). (D) sectional image. [20]

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63 Figure 5 2. XRD patterns of an electrodeposited gibbsite PVA composite on ITO electrode. [20] Figure 5 3. Deposit weight on ITO electrode versus electrophoretic duration. [20]

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64 Figure 5 4. Thermogravimetric analysis of the nanocomposite sample as shown in Figure 5 1. [20]

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65 CHAPTER 6 FABRICATION OF GRAPHENE OXIDE PVA NANO COMPOSITE S VIA FILTRATION 6.1 Introduction and Background We introduce a novel self assembly technology to fabricate graphe n e oxide (GO) / poly(vinyl alcohol) (PVA) nano composite. Stable GO/ PVA dispersion s without agglomeration are filtrated to form well ordered GO/PVA nano composite s The resulting nanocomposite is further reinforced by glutaraldehyde solution (GA) to for m in situ cross li nking between GO and PVA without disrupt ing the aligned structure. As a result, the mechanical toughness and ultimate tensile strain could be significantly increased with small amount of PVA (~10 to 30 wt% based on suspension mixture ) in nanocomposite. The ultimate tensile stress, with the aid of GO/PVA cross lin king sits between that of pure GO paper and reduced graphe n e oxide paper. We demonstrate the mechanical propertie s of GO/PVA nano composites can be further enhanced by using simple and scalable infi ltration concepts to combine GO/PVA with in situ cross linking 6. 1. 1 Graphene /Graphene Oxide for Composite Application Graphene, one atom thick planar sheet of sp 2 bonded carbon atoms has attracted great attention from both the experimental and theoretic al scientific communities in recent years. [102, 103] The two dimensional honeycomb crystal lattice structure of grapheme, as a prototype of carbon allotrope s, exhibit tremendous cap abilities and prospects as well as other s such as buckyball and carbon nanotube.(Figure 6 1 A ) [102] It has a large theoretical specific surface area (2630 m 2 g 1 ), high intrinsic mobility (200,000 cm 2 v 1 s 1 ) [104, 105] ~1. 0 TPa) [106] excellent thermal conductivity ( ~ 5000 Wm 1 K 1 ) [107] and optical transmittance (~ 97.7%) [108]

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66 The se properties provide versatile applications such as field effect transistors [109 115] s ensors [116 120] transparent conductive films or electrodes [108, 121 125] lithium ion batteries [126 128] and graphene polymer nanocomposites [129 133] To synthesize graphene, a variety of methods have been developed such as chemical vapor deposition (CVD) [134 136] ar c discharge [137, 138] epitaxial growth on SiC [139 142] and the mechanical cleavage method [143, 144] Although those methods can produce small amounts of large size, defect free graphene sheets for fundamental studies and electronic applicat ions, they are not appropriate for polymer nanocomposites that require a large amount of graphene sheets preferably with modified surface structure. [145] Graphene oxide (GO), produced by chemical oxidation and exfoliation of graphite, [146] are suitable for mass production required for polymer composite applications and industrial manufacturing Chemical oxidation/exfoliation processes offers the capability of large scale fabrication of low cost raw materials for inorganic organic nanocomposites GO also exhibits possibilities of chemical modification on surface. According to recent studies, GO are oxidized graphene sheets having their basal planes decorated mostly with epoxide and hydr oxyl groups, as well as carbonyl and carboxyl groups l ocated presumably at the edges (Figure 6 1B ) [147 150] These oxygen functionalities render the graphene oxide hydrophilic and provide good dispersity in water and in organic solvents after chemical modification, which facilitates to produce graphene polymer nanocomposites. [151 153] P artial reduction of graphene oxide is capable to restore electrical conductivity and thermal stability. [154]

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67 6.1.2 Graphene based material Graphene oxide and other graphe n e based derivatives have been used in polymer nanocomposites and hold potential for a variety of applications. N anocomposites with GO d erived mate rials as fi ller s have shown dramatic improvements in properties such as elastic modulus, tensile strength, electrical conductivity, and thermal stability. [129, 133, 147, 155] Besides these improvements a re often observed at low loadings of fi ller s due to the large interfacial a rea and high aspect ratio of these materials T o achieve large property enhancements in nanocomposites, GO must be exfoliated and well dispersed in the polymer matrix. [156] T echniques to disperse GO derived filler s into polymers such as solution mixing, melt blending, or in situ polymeriza tion are d epend ed on different kinds of polymer matrix [147] However, the en hancement of mechanical properties for graphe n e based nanocomposite is still limited by high fraction of polymer matrix. High fraction of GO derived filler s might cause high viscosities and aggregation in composite formation process which results relatively weak mechanical properties [157] Recently pure graphe n e oxide and reduced graphe n e oxide (RGO) paper s create by filtration of dispersed sheets, exhibit superior properties than many other paperlike materials in strength and stiffness( Youn modulus of ~ 41.8 GPa, tensile strength of ~ 293.3 MPa ) [158, 159] Th e ultrastrong mechanical proper ties rely on the well ordered structure of nanoscale sheets and interaction between sheets such as H ydrogen bond ing and van der Waals force C ross linking at the molecular level also has been introduced t o enhance the mechanical and physical properties of graphene nanocomposites Polyallylamine and d ivalent Ions are used to form ionic or covalent cross linking with graphe n e oxide [160 162] Unfortunately, the improvements are not

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68 significant and still incomparable to pure RGO paper. Although GO/RGO paper has high strength and elastic modulus, its tensile strain is still too low (~0.8%) and relative inelastic. To achieve a high tough material ( the ability of a material to absorb energy and plastically deform without fracturing ) small amount of polymer must be interfuse d in GO/RGO paper. 6.1. 3 Vacuum Infiltration Vacuum infiltra tion, aka paper making, is one of button up assembling technique s to mimic nature hierarchical architecture The theory is using vacuum pressure as driving force to remove solvent and form aligned structure of building block. For instant, Cellulose nanopap er fi lms has been made by Henriksson and co workers by employing v acuum infiltration assembly [163] This nanopap er sample shows very high tough ness in uniaxial tension and this is associated with a strain to failure as high as 10%. T he modulus ( 13.2GPa) and ultimate tensile strength ( 214MPa) are remarkably high and can sometimes exceed those of wood made by cellulose Walther and co workers obtain biomimetic composites by coating nanoclay platelets with a thin layer of PVA, in combination with p aper making processing [164] This allows ma terials with Young modulus of ~ 45 GPa and tensile stre ngth of ~ 250 MPa Now it is frequently used to form uniform graphe n e film [154, 159, 162, 165, 166] and graphene based composite [154, 160, 161, 167] 6.2 Experimental 6.2.1 S ynthesize Graphene Oxide M onolayer g raphene oxide (GO) sheets w ere synthesized from natural graphite powder (Grade 3243 Asbury Graphite Mills Inc. ) by using modified Hummers method [146, 168] The graphite powder ( 2 g) was put into an 80 C solution of concentrated

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69 H 2 SO 4 ( 3 m L ), K 2 S 2 O 8 (1 g), and P 2 O 5 ( 1 g). The resultant dark blue mixture was thermally isolated and allowed to cool to room temperature over a period of 6 h. The mixture was then carefully diluted with d istilled water, filtered, and washed on the filter until the rinse water pH became neutral. The product was dried in air at ambient temperature overnight. This preoxidized graphite was then subjected to oxidation by oxidized graphite powder was put into cold (0 C) concentrated H 2 SO 4 (46 m L ). K MnO 4 (6 g) was added gradually with stirring and cooling, so that the temperature of the mixture was not above 20 C. The mixture was then stirred at 35 C for 2 h ours and distilled water ( 92 mL ) was added. In 15 min, the reaction was terminated by the addition of a large amount of distilled water (2 80 m L ) and 30% H 2 O 2 solution (5 m L ), after which the color of the mixture changed to bright yellow. The mixture was filtered and washed with 1:10 HCl solution ( 500 m L ) in order to remove metal ions. The resulting ~0.6wt% GO sol was used to prepare exfoliated GO. The slurry of GO is diluted to 2mg/m L by dispers ing in DI water via strong ultrasonication for 1 hour and centrifuged to remove im purities. GO suspension was purified by dialysis for 24 hours to completely remove metal ions acids and other impurities The resulting homogeneous brown dispersion which contained ~2mg/m L GO, was stable for a period of months and was used for film preparation. 6.2 .2 Preparation and Test of GO /PVA C omposites GO/PVA composite was made by filtration of the as prepared colloidal mixture of graphene oxide dispersion and dissolved PVA solution through an Anodisc membrane fi lter ( 25 mm in diameter, 0.2 u m pore size, Wha tman ) followed by addition of aqueous glutaraldehyde solution (1wt%, d iluted with DI water) to the GO/PVA composite film (Figure 6 2A ) Poly (vinyl alcohol) (MW 89,000~98,000) a nd

PAGE 70

70 glutaraldehyde solution(GA 25 wt % in water) were obtained from Aldrich. PVA wa s previously dissolved in DI water at 90 C to obtain 1mg/m L solution. GA was diluted to a de sir ed concentration by adding DI wate r A flow induced pump ing system was made to provide a stable low vacuum pressure to filtrate the suspension s (Figure 6 2B ) After filtration, the GO/PVA composite film was suction dried for 24 hours then blow dried with compressed air. Aft er being removed from filter, the film was air dried and slightly compressed for 12 hours Peeling from the filtering membrane easily acco mplished after dr ying Mechanical properties of GO/PVA composites were tested by INSTRON Tensile Tester model 1122 at a crosshead speed of 0.5 mm/min. Figure 6 3 display s the free standing GO/PVA composite and its sample s for tensile test. Testing sample s were cut by a razor blade to have width of 1.5 mm and t hickness ranging from 10 to 30 m and adhered on home made sample holders using commercial glue as an adhesive. The thickness of the tested samples wa s measured by cross sectional SEM images to calcu late the final tensile strength. 6.2.3 Instruments Atomic force microscopy (AFM) wa s conducted on a Digital Instruments Dimension 3100 unit. Scan ning electron microscopy (SEM) wa s carried out on a JEOL 6335F FEG SEM. X ray diffraction ( XRD) spectrum wa s o btained with P h ilips XRD 3720 equipment. A Cu K 1 (k = 1.54049 ) ra diation wa s used from 4 to 70 with a scan rate of 2.4 /min. U ltrasonication for dispersing GO wa s conducted by using Misonix Sonicator 3000 Ultrasonic Cell Disruptors with Microtip Probes Tensile strength strain test was accomplish ed by INSTRON Tensile Tester model 1122 load

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71 frame with MTS ReNew system Tests were equipped with a 500 g load cell and performed at a rate of 0.5 mm / min Thermogravimetric analysis ( TGA ) wa s characterized by Mettler Toledo TGA/SDTA 851 e All measurements were conducted under dynamic nitrogen flow, over a temperature range of 25 900 o C with a slow ramp rate of 1 o C / min to prevent sample loss. 6.3 Result and Discussion 6. 3. 1 Dispersion of GO Sheet s with PVA GO used in this work was prepared from graphite by the modified Hummers method as previous described and then blend ed with low concentration PVA solution before filtration GO, which contains many oxygen containing functional groups, can be well dispersed in water at the level of individual sheets. Figure 6 4 show a typical tapping mode atomic force microscopy image of GO sheets deposited onto a piece of silicon wafer as substrate from an aqueous dispersion. AFM images revealed that GO sheets ha ve heights of ~0.8 nm, which is characterist ic of a single gra phe n e sheet [157, 169] T h is indicates graphene oxide was fully exfoliated into individual sheets by ultrasonic treatment. PVA, which is also water soluble due to its hydroxyl groups, is expected to form homogenous di spersion s with graphene oxide in molecule scale Figure 6 5 displays pure GO dispersion and well mixed GO/PVA suspension after storage in room temperature for 48 hours Both images show homogen e ous dispersion without significant agglomeration R elative low concentration of GO/PVA suspension and high electric repulsi on between negatively charged GO sheets help to prevent aggregation between GO and PVA. [166] GO/PVA nanocomposites were fabricated through a simple paper making(filtration) method as described in the experimental section. Unlike other drying related assembly

PAGE 72

72 process es which might result in severe agglomerations vacuum filtration is a n appropriate approach to mitigate colloidal agglomerations and en hance t he oriented assembly of GO. The dispersion state of GO sheet s in composite was evaluated by scanning electron microscopy Figure 6 6 shows the cross sectional SEM image of a GO/PVA nano composite(~70wt% GO). The well ordered layer structure in SEM image ver if ies t h at GO sheets are c learly well dispersed in the entire nano composite X ray diffraction is an important tool for determining t he laye r ed structure in the assembled GO PVA nano composite. The peak in the X ray spectrum of graphene oxide sheet correspo nds to the layer to layer distance (d spacing) which can be calculated by Bragg s law Figure 6 7A show s the XRD patterns of casted PVA filtrated GO paper and GO PVA nanocomposite. The peak s of GO are shifted to small er angle s with larger amount of PVA l oading in composite. It indicates the GO layer spacing is increasing and filled by PVA polymers. M ost importantly t he XRD results also demonstrate that the regular and periodic struct ure of graphe n e oxide remains while it blends with PVA. GA cross linked GO PVA composite with similar XRD results as non crosslinked composite also represent that the orientation of composite remains after crosslinking process ( Figure 6 7B and 6 7C) 6. 3.2 Thermal Properties of G O /PVA Nanocomposites Thermogravimetric analysis is commonly employed in research and testing to determine the thermal properties of materials such as degradation temperature absorbed moisture content, the level of inorganic and organic components TGA result for pure PVA GO, and GO/PVA nanocomposite a re shown in Figure 6 8 We can observe two steps in the mass loss for both GO paper and GO/PVA composite. The first significant mass loss below 180 o C can be attributed to the evaporation of adsorbed

PAGE 73

73 water which is ~ 11~ 19 wt% of the specimen s Above 180 o C sharp weight loss (~22 wt % ) indicat es that a large amount of oxygen containing groups have been removed Most carbon are left over 230 o C and slowly dissipated under higher temperature T he TGA curve s of the nanocomposite s are l ifted with more PVA ble nded compared with that of pure GO paper due to the a dsorption of water. W ater is able to be absorbed on the surface of GO paper as well as in the interval of GO sheets. W hile PVA polymers occ upy the space between GO sheets, less water is incorporated in G O/PVA nano composites and this change the TGA curve s. In Figure 6 8 B, the composites blended with various amount s of PVA, display relative unanimous TGA curves due to crosslinked network like structure. Interestingly, GO/PVA composites show indistinctive d ifference to pure GO paper while the de composition temperature of PVA polymer is higher than that of GO paper. ( Figure 6 8 C) PVA has better thermal stability due to its tangled, macro sized polymer structure fastened by H ydrogen bond s. On the other hand, t he H ydrogen bond ing in nano composite is limited by OH group on GO sheet s and the density of OH group on GO is less than that on PVA polymer (t he C:O ratio of GO sheet is between 2.1 and that of PVA is ~2). [146] So the molecul e scale interaction of PVA chains with GO sheets is nearly equivalent to that between GO sheets, which could reasonably explain why GO/PVA composites has simil ar t hermal p roperties to pure GO paper. 6. 3.3 Mechanical Properties of Graphene oxide /PVA Nanocomposites The m echanical performance of the nanocomposite depend s on the large aspect ratio of the graphene oxide sheets the molecular level dispersion of th e PVA polymer, and strong interfacial adhesion between GO and PVA The binding of components is due to hydrogen bonding and van der Waals force between GO sheets and PVA

PAGE 74

74 polymer GA can also form strong chemical crosslinking s with hydroxyl group on GO and PVA to enhance the interaction. The best stress strain curve for nanocomposites within the tested range of GO /PVA weight ratios is give n in Figure 6 9 with other reference materials This GO/PVA composite, which is 5:2 in GO/PVA weight ratio and cross link ed by 2mL of GA shows ex tremely high toughness and relative high strength (275 33 MPa of tensile strength and 14.5 3.4 % of tensile strain) compared with its components Specifically, it exhibits over 2 folds of ultimate strength to filtrated GO paper a nd nacre, ~30 times tougher than filtrated graphene paper and nacre ~100 times tougher than filtrated GO paper. [14, 158, 159] The mechanical performance of the GO /PVA nanocomposite signi fi cantly surmount s that of t he pure GO paper and PVA polymer To evaluate the reason s of the high toughness, specimens of various amounts of PVA and GA are systematically investigated Figure 6 10A displays stress strain curves of nanocomposite consisting of 5mg of GO with various am ount of PVA from 0.5mg to 3mg (from original suspension). Its ultimate tensile strength slightly increases with the amount of PVA but the ultimate tensile strain significantly increases with that of PVA. It shows that t he PVA polymer could provide large el ongation to the resulting comp osite but not the strength To harden and improve the strength of the nano composite s one of the feasible approaches is to use crosslinkers to bind organic and inorganic component s. [13, 17, 162] Figure 6 10 B and C present the stress strain curves of GO/PVA nanocomposite which was crosslink ed with GA Those with adequate PVA (>0.5mg) illustrate clear stress strain relationships with re spect to the larger ultimate strain. Comparing to orig inal GO/PVA nano composite, proper amount of GA crosslinking s offers extra support to composite Therefore, composite is less

PAGE 75

75 sensitive to flaws and is able to achieve higher tensile strength. Figure 6 1 1 shows how the GA influence s the mechanical behaviors of GO/PVA composite. For instance in Figure 6 11A C composite s turn to hardening with the increase of amount of GA Howe ver, composite with high fraction of PVA (~3mg PVA with 5mg GO) does n o t show further mechanical improvement after GA filtration (Fig ure 6 11D) PVA/GO composites with overdosed GA are brittle and difficult to conduct mechanical tests ( Not shown here) Overall, the mechanical properties could be optimized by specific GO/ PVA/GA ratios The mechanical properties of GO /PVA composite, inclu ding ultimate strength and strain, are presented in Figure 6 12 for comparison. On the other hand, the fully developed stress strain curves for GO/PVA nano composite s are coinciden tally similar to each other. T he typical stress strain diagram of PVA/GO com posite is shown in Figure 6 1 3A The ini ti al straightening during the tensile loading is quite small which represents that the whole structure is relatively homogeneous T he stress strain curve displays linear in the elastic region ; t he saw tooth like cur ve and smooth rising curve which are considered as plastic regions, sequentially appear along with the further stretch of the sample The rupture of composite samples is accompanied without necking and pro duces almost straight and flat fracture surfaces Those can be explained from t he schematic models in Figure 6 13B. We speculate that there are three types of arrangements for the GO nanosheets and PVA polymer upon increasing strain, based on the hypotheses of the tension shear chain model and the good d ispersion of GO/PVA composite at the molecule level. In the s traightening and elastic region m ost of the load is carried by the GO sheets whereas the PVA transfers load via the high shear zones between GO sheets Whe n it reaches

PAGE 76

76 the yield point, folded PV A polymer chains between GO sheets start to straighten T herefore, sequential unraveling of individual folds of a PVA chain and regeneration of hydrogen bond s might be the origin of the observed saw tooth pattern s Figure 6 14 demonstrates the similar beha vior from biological materials due to opening of intra chain loops or folded domains within a protein molecule, or release of sacrificial inter chain bonds holding a crosslinked multi chain matrix together. [170 172] Besides, still tangled PVA chains interacted with each other might cause s overall increasing trend of strength in initial plastic region. In the second plastic region, most PVA chains are predicted to be unfolded and less twist ed due to substantial slippi ng of GO sheets T herefore, t he resistant forces of deformation are the relative weak f riction force The composite will keep extending until the PVA strengthen and lose its adhesion in any weak spot, then it breaks without further elongation. Moreover, ad dition of GA will form cross linking between GO and PVA as well as between PVA and PVA. This results in the hardening of the entire c omposite with increase d strength but decrease d elasticity, hence overdose of GA could make composite brittle and easy to fr acture 6. 3.4 Mechanical Propert y Simulation The simulations of the ultimate tensile strength of GO/PVA nano composite shown i n Figure 6 15 are based on a shear lag mechanical model according to which the applied load is fully transferred to the inorgan ic nano platelets through shear stresses developed in the organic matrix. [9, 98] For a polymer matrix having a yield shear strength ( y ) and strong bonding (e.g., hydrogen bond and cross link ing through GA ) to the su rface of graphene sheets (interfacial strength i y ) the tensile strength of the composite can be calculated u sing the volume fraction of graphene sheets ( V s ), the

PAGE 77

77 tensile strength of the graphene sheets ( s ), and of the polymer matrix ( p ), as (equati on 6 1 ): ( 6 1 ) with the factor being a function of the sheet aspect ratio ( s ) t he yield shear strength and the tensile strength of the polymer matrix The factor depends on the operative failure mode, as shown in Figure 6 16, which is determined by the aspect ratio of the graphene sheet ( s ) and critical aspect ratio ( s c ) which is equal to the ratio of s versus y For aspect ratios higher than a critical value ( s > s c ), the composite fails because of the fracture of ino rganic sheets. On the other hand, for aspect ratios lower than the critical condition ( s < s c ) the continuous matrix yields before the sheets break, thus cause sheets pull out and matrix plastic slide before the rupture of the composite Therefore, t he fac tor is given by the following relation: Sheet fracture: ( s > s c ) Sheet pull out: ( s < s c ) We use graphene sheet (GS), graphene oxide paper (GOP), and reduced graphene paper (RGP) as reference material of GO ; their tensile strength s ) is tested as 130GPa, 135MPa, and 293MPa respectively [106, 158, 159] The yield shear strength ( y ) and the tensile strength of the PVA m ) are assumed to be 25 and 40 M Pa, respectively In GS/PVA simulation, most events are fit by sheet pull out model due to high critical aspect ratio (s c = 5200). The tensile strength of composite is rapidly increased with the volume fraction and aspect ratio of graphene sheets. In

PAGE 78

78 GOP/P VA and RGP/PVA simulations, pull out model and fracture model coexist due to relative small critical aspect ratio (s c = 5.4 and 11.7, respectively). In Table 6 1, we have simulation results of the tensile strength of composite from referenced graphenes wit h experimental volume fraction and aspect ratio (s~1000) The results show that the tensile strength of GO/PVA composite is closed to that of RGP/PVA but much lower compared to that of GS/PVA. Table 6 1. Tensile strength of (MPa) Vo lume fraction of graphene (%) (Weight ratio) GS/PVA GOP/PVA RGP/PVA GO/PVA ( Experimental ) 50 (5:3) 48,120 87.3 165.6 214.7 17.0 60 (5:2) 57,740 96.8 190.7 226.5 27.4 75 (5:1) 72,160 111.0 228.5 275.3 33.5 85 (5:0.5) 81,770 120.4 253.6 288.5 14.6 *The simulation results of tensile strength of nanocompsoites are based on the aspect ratio of graphene sheets is 1000. Other parameter s are shown as described The calculated volume fractions are based on weight ratio of graphene and PVA. the density of PVA ma trix and graphene nanosheets are taken as 1.3 g/cm 3 and 2.2 g/cm 3 [157]

PAGE 79

79 A B Figure 6 1. A) Mothe r of all graphitic forms. Graphene is a 2D building material for carbon materials of all other dimensionalities. It can be wrapped up into 0D buckyballs, [102] B) A proposed schematic (Lerf Klinowski model) of graphene oxide structure. [147]

PAGE 80

80 A B Figure 6 2 A) The flow chart of filtration process. B) Schematic diagram of flow induced pumping system

PAGE 81

81 Figure 6 3 A) and B) Images of elastic, free standing GO/PVA composite. C) Tensile test samples, which are cut by double ra zor blade Figure 6 4 The tapping mode AFM image of GO sheets deposited on silica wafer from an aqueous dispersion; the height difference is ~ 0.8 nm

PAGE 82

82 Figure 6 5 A) Pure GO dispersion (2mg/ml in water), and B) W ell mixed GO/PVA suspension (5mg of GO and 2mg of PVA), which had been static for 48 hours Figure 6 6 T he SEM image of cross sectional GO/PVA composite (~70wt% GO).

PAGE 83

83 Figure 6 7. T he XRD patterns of casted PVA filtrated GO paper and GO PVA nanocomposite with A) No GA, B) 2ml GA added, C) 3ml GA added.

PAGE 84

84 Figure 6 8 TGA result s of A) GO/PVA composites, B) G O/PVA composites with 2ml of GA and C) GO paper and casted PVA.

PAGE 85

85 Figure 6 8 Continue d

PAGE 86

86 F igure 6 9 The stress strain curve of reference materials and GO/PVA nano composites which is 5:2 in GO/PVA weight ratio and cross linked by 2ml of GA GO/PVA nanocomposite shows extremely relative high toughness (fracture energy) compared to GO paper, graphene paper and nacre

PAGE 87

87 Figure 6 10. The stress strain curves of nanocomposite s: A) GO/PVA without GA, B) GO/PVA with 2ml of GA, C) GO/PVA with 3ml of GA

PAGE 88

88 Figure 6 11. The stress strain curves of nanocomposite s: A) 5mg GO plus 0.5mg PVA, B) 5mg GO plus 1mg PVA, C) 5mg GO plus 0.5mg PVA, D) 5mg GO plus 0.5mg PVA, (from original suspension).

PAGE 89

89 Figure 6 12 The mechanical properties of GO /PVA composite including A) ultimate tensile strength and B) ultimate tensile strain. Samples include 2.5mg of GO with 0.5, 1, 2 and 3mg of PVA.

PAGE 90

90 Figure 6 13 A) There are three regions in t he stress strain diagram of PVA/GO composite: Elastic, 1 st Plastic, and 2 nd Plastic region B) t he schematic models of three types of arrangements upon increasing strain The black plane and blue curves represent the GO nanosheets and PVA polymer repectively.

PAGE 91

91 A B Figure 6 1 4 A) Model of long polymers and forceextension curves for different kinds of polymers. a, Diagram of long polymers behaving as entropic springs. The low er molecule is compacted, with many domains that are held together w ith intermediate strength sacri fi cial bonds. b, Force extension curves for three different kinds of molecules. [171] B) A schematic illustration of protein modules deforming between mineral platelets in the biological nanostructure and the inherent force extension relation of protein with cross linking mechanism of Ca 2+ formed sacri fi cial bonds.

PAGE 92

92 Figure 6 15 The simulations of ultimate tensile strength of GO/PVA composite are based on a shear lag mechanical model and using A) graphene sheet (GS), B) graphene oxide paper (GOP) and C) reduced graphene paper (RGP) as reference material of GO.

PAGE 93

93 Fig ure 6 15 Continued Figure 6 1 6 T he schematic models of operative failure mode : platelets fracture and platelet pull out.

PAGE 94

94 CHAPTER 7 FABRICATION OF PERIO DIC BINARY NANOSTRUC TURE 7.1 B ackground We developed a simple spin c oating technology that enables large scale production of colloidal crystals with nonclose packed structures. [64, 65] The technology up colloida l self assembly with the scalability and compatibility of standard top down microfabrication. Unitary microstructures have been demonstrated by using spin coated monolayer colloidal crystals as templates. Here we show the scalable fabrication of binary per iodic nanostructures by utilizing double layer nonclose packed colloidal crystals as templates. Contrary to traditional colloidal lithography and our previous templating approaches, the polymer matrix between the spin coated colloidal crystal has also been used as a intercalated arrays of silica spheres and polymer posts, gold nanohole arrays with binary sizes, and dimple ary coatings. 7.2 Experimental 7.2.1 Preparation of Double Layer Colloidal Crystals by Spin Coating Monodispersed silica spheres were synthesized by following the standa rd Stober process. [173, 174] The as made sil proof ethanol and then redispersed in ETPTA monomer, using a Thermdyne Maxi Solution Mixer. The final particle volume fraction was controlled to be ~20%. A 1% (weight) sample of Darocur 1173 was added as the photoinitiator The transparent colloidal suspension

PAGE 95

95 was dispersed on a silicon wafer which had been freshly primed by 3 acryloxypropyl trichlorosilane (APTCS, purchased from Gelest). The recipe for spin coating a double layer colloidal crystal consisting of 320 nm sili ca spheres was 200 rpm for 1 min, 300 rpm for 1 min, 1000 rpm for 30 s, 3000 rpm for 10 s, 6000 rpm for 10 s, and 8000 rpm for 90 s. The monomer was then photopolymerized for 4 s econds using a Pulsed UV Curing System. 7.2.2 Fabrication of Binary Gold Nan ohole Arrays by Using Double Layer Colloidal Crystals as Templates rate, and 100 W for 2 min to partially remove ETPTA polymer matrix for releasing the top layer silica spheres of the spin coated double layer colloidal crystal. The released silica particles were then removed by sweeping, using a cleanroom Q water. The residual ETPTA polymer matrix was etched under the same oxygen plasma etch conditions as shown above f or 1 min. A 2 nm thick layer of chromium followed by a 30 nm thick layer of gold was then deposited by using a Kurt J. Lesker CMS 18 off by Q water. 7.2.3 Replication of Dimple nanohole arrays, both the bottom and top layer silica spheres were removed by dissolution (PDMS, Sylgard 184, Dow Corning) precursors were mixed and degassed and then poured over the resulting polymer post arrays with binary dimple nipple structures. After curing at 80 C

PAGE 96

96 and put on top of ETPTA monomer supported by a glass slide. The ETPTA monomer was photo polymerized for 4 s econds using a Xenon pulsed UV curing system. The to complete the fabrication process. To fabricate sol ~ 900 nm thick) of SOG precurs or (512B, Honeywell Electronic Materials) was spin coated at 800 rpm for 60 s on a glass slide before putting it on th e PDMS mold. After baking at 120 C for 5 min and peeling off PDMS, glass dimple 7.2.4 Specular Refl ection Measurements A HR4000 High Resolution Fiber Optic UV probe illuminate the sample. The beam spot size was about 3 mm on the sample surface. Measurements were performed at normal incidence and the cone angle of collection was l and the reference spectrum. The reference spectrum was the optical density obtained from an aluminum absolute re ectivity was the average of several measurements obtained from different spots on the sample surface. 7.2.5 Instrumentation Scanning electron microscopy was carried out on a JEOL 6335F FEG SEM. A thin layer of gold was sputtered onto the samples prior to imaging. A standard spin coater (WS 400B 6NPP Lite Spin Processor, Laurell) was used to spin coat colloidal suspensions. The polymerization of ETPTA monomer was carried out on a Pulsed UV Curing System (RC 742, Xenon). Oxygen plasma etch was carried out on a Unaxis S huttlelock RIE/ICP reactive ion etcher. A Kurt J. Lesker CMS 18 Multitarget Sputter

PAGE 97

97 Optics HR4000 UV vis spectrometer. 7.3 Result and Discussion 7.3.1 Fabricate B inary N anos tructures Via Spin Coating In the current colloidal templating approach, we use double layer, nonclose packed colloidal crystals prepared by the well established spin coating tec hnique as structural templates. [64, 65] The spin coating methodology is based on shear aligning concentrated colloidal suspensions in a nonvolatile monomer. Contrary to conventional colloidal self assemblies, spin coating enables rapid production of wafer scale (up to 8 in. in diameter) colloidal crystals with remarkably large domain sizes. Another unique property of the spin coated colloidal crystals is their unusual nonclose pa cked crystalline structures. The spheres within each colloidal layer are hexagonally arranged and are separated by ~ 1.4 times the particle diameter The thickness of the shear aligned colloidal crystals can be precisely controlled by adjusting the spin coating conditions (e.g., spin speed and time). Most important, the shear aligned colloidal crys tals are embedded in a polymer matrix, which has a highly uniform thickness over wafer scale areas. In our previous templating approaches [62, 66, 175] the polymer matrix is completely removed and only the released colloidal arrays are used as structural templates. Here we demonstrate a new approach that uses both colloidal spheres and polymer matrix as templates to generate periodic binary nanostructures that are not easily available by traditional top down and bott om up technologies. Figure 7 1 shows the schematic outline of the current colloidal templating procedures for fabricating intercalated arrays of silica spheres and polymer posts and periodic gold nanohole arrays with binary sizes. The spin coating techniqu e is fi rst

PAGE 98

98 applied to make double layer, nonclose packed silica colloidal crystals embedded in a thin layer of poly(ethoxylated trimethylolpropane triacrylate) (PETPTA) matrix. The polymer matrix can then be partially removed by a brief oxygen plasma etch process to release the top layer spheres. During this process, silica particles function as etching masks to protect PETPTA matrix underneath them from being etched. As the top layer silica spheres only ll in the triangularly arranged crevices made by the nontouching bottom layer spheres, [64] intercalated binary arrays of polymer posts and silica spheres are formed after removing the exposed top layer spheres by simply sweeping, using a cleanroom Q tip unde layer spheres is not affected by this sweeping process as they are still partially embedded in the remaining polymer matrix. The residual PETPTA matrix can be further plasma etched to expose the substrate for the subsequent metal deposition. In this process, the size of the polymer posts is reduced due to the isotropic plasma etching. The resulting intercalated silica spheres and polymer posts can fi nally be used as deposition masks during sputtering deposition of a thin layer of chromium and gold. After lifting off the templating particles and polymer posts, periodic gold nanohole arrays with binary sizes resulted. Figure 7 2 show s the top and side view scanning electron microscope (SEM) images of the intercalate d binary arrays of silica spheres and polymer posts prepared by using a double layer colloidal crystal consisting of 320 nm silica spheres as template. The long range hexagonal arrangement of the protruded polymer posts is con fi rmed by the Fourier transfor m of the SEM image as shown in the inset of Figure 7 2A. It is clearly evident that each polymer post is surrounded by three silica particles, and vice versa. The size of the polymer posts is determined by the oxygen plasma etch

PAGE 99

99 durations longer etching r esults in smaller polymer posts. Figure 7 2B shows a sample that has been plasma etched at 40 mTorr pressure, 40 SCCM oxygen fl ow rate, and 100 W for 2 min. The measured diameter of the polymer posts is ~ 250 nm, smaller than the templating silica spheres o f 320 nm. Even smaller polymer posts with diameter close to 100 nm can be fabricated by using longer etch time. The dimple shaped ends of the polymer posts as seen in Figure 7 2B are marks left by the templating top layer spheres. The side view SEM also re veals the bottom layer silica spheres are still partially embedded in a thin layer of polymer matrix. The remaining polymer matrix is caused by the polymer wetting layer ( ~ 100 nm thick) between the spin coated silica colloidal crystals and the silicon sub strate. [67] It forms a continuous fi lm covering the sub strate surface, preven ting the fi nal lift off of deposited metals in the fabrication of binary metal nanohole arrays. We therefore conduct another brief oxygen plasma etch to completely remove the polymer matrix between the silica spheres and the polymer posts. Figure 7 3 shows the SEM images of the same binary array sample as in Figure 7 2 after an additional 1 min of oxygen plasma etch. Both the silica spheres and the polymer posts retain the original hexagonal long range ordering and the intercalated arrangement, though the h eight and size of the polymer posts are reduced due to the isotropic plasma etching. The size of the polymer posts in Figure 7 3A is measured to be ~ 110 nm, much smaller than the original silica templates. After the second plasma etching process, the silic on substrate is completely exposed between the silica spheres and the polymer posts. The intercalated silica spheres and polymer posts can then be used as deposition masks during sputtering deposition of a30 nm thick layer of gold by using a Kurt J. Lesker CMS 18 Multitarget

PAGE 100

100 Sputter. The typical deposition rate is 5 /s and a 2 nm thick layer of chromium is deposited as the adhesion layer. As the adhesion of the silica particles and the polymer posts on the silicon substrate is weak, they both can be easily removed by brief ultrasonication in water or by simply sweeping with a cleanroom Q tip under fl owing water. This results in the formation of periodic binary gold nanohole arrays as shown by the SEM imag e in Figure 7 4. The intercalated hexagonal arrangeme nt of big holes templated from silica spheres and small evident. Similar to the binary silica sphere polymer post template, each big hole is surrounded by three small holes, and vice versa (see the magni fi ed SEM ima ge in the inset of Figure 7 4). The same templating process can be easily extended to many other metals and dielectrics, such as Ag, Pt, Pd, Ni, Cr, Al, and ITO. The binary metal nanohole arrays can be further used as second generation etching masks to tra nsfer the binary periodic patterns to the underneath substrates by dry etch. 7.3.2 B inary N anostructures for C oatings The current colloidal templating technique is taken one step further by demonstrating the production of binary dimple coatings are widely utilized in eliminating panel displays, increasing the transmittance of optical lenses, reducing glaring from automobile dashboards, and enhancing the conversion ef fi ciency of solar cells. [176 178] The binary an corneas of moths, which consist of hexagonal arrays of nonclose packed subwavelength nipples, similar to the protruded polymer posts as shown in Figure 7 2B [29] We have recently developed a soft lithography like templating technology for fabricating polymer moth eye ARCs with p eriodic arrays of unitary hemispherical

PAGE 101

101 nipples. [175, 179] However, the height of the resulting nipples is at most the radius of the templating spheres, limiting performance of the unitary ARCs. By contrast, the optical thickness of the binary ARCs can almost be doubled. We expect this will lead To generate binary ARCs, we first assemble double layer, nonclose packed colloidal crystals ( Figure 7 5A ) by using the spin coating technology. The polymer matrix is then partially removed and the exposed silica spheres are dissolved in a 2% dimple nipple array as shown by the top and side view SEM image in Figure 7 5 panels B and C, respectively. The resulting periodic binary structure can then be tra nsferred to a poly(dimethylsiloxane) (PDMS) mold and then replicated to a thin layer of polymer (e.g., ETPTA) or sol gel glass on a glass substrate by the well establish ed soft lithography technique. [180] Figure 7 5D shows a side view SEM image of a replicated ETPTA binary array. The secular optical re fl ection of the templated polymer binary ARCs is evaluated by using UV vis re fl ectivity measurement at normal incidence. The solid lines in Figure 7 6 show and a glass slide covered with a polymer dimple nipple ARC. For comparison with the new normal with a unitary ET PTA hemispherical ARC as reported in our previous work is also shown (blue curve). [179] The glass control sample exhibits ~ 4% single surface coate d glass slide displays <0.5% re fl ectivity for the whole visible spectrum. It is also evident that the binary ARC shows improved

PAGE 102

102 unitar y hemispherical ARC for most of the visible wavelengths. by theoretical calculation by using a rigorous coupled wave analysis (RCWA) model [181] A coordinate system is set up with the z axis perpendicular to the templat ed dimple nipple surface so that the bottoms of the dimples are at z = 0 and the nipple tops at z = h We divide the dimple nipple arrays into N = 100 horizontal layers with thickness (h)/(N) The dimple and nipple lattices are assumed to be hexagonal and the distance between the centers of the neighboring dimples and nipples is where D is the diameter of templating silica spheres Other structural parameters (e.g., dimple size and depth, nipple depth, size, and shape) are obtained from the side view SEM image as shown in Figure 7 5D. The fraction of ETPTA in each horizontal layer can then be determined by simple geometrical calculation. On the basis of the effective medium theory, the effective refractive index n(z*) of the layer at height z* can be approximated by where = 1.46, = 1.0, and q = 2/3. [29, 182] by solving the Maxwell equation to express the electromagnetic (EM) eld in each layer and then matching EM boundary conditions between neighboring layers for the determination of system. The dotted lines in Figure 7 6 show the RCWA simulated substrate and a binary ETPTA ARC. The simulated spectra agree reasonably well with the experimental results.

PAGE 103

103 Figure 7 1. Schematic illustration of the templating procedures for fabricating intercalated hex agonal arrays of silica spheres and polymer posts and binary arrays of metal nanoholes by using double layer nonclose packed colloidal crystal as template. [117]

PAGE 104

104 Figure 7 2. SEM images of intercalated hexagonal arrays of polymer posts and silica particles after removing the top layer silica spheres of a do uble layer colloidal crystal template: (A) top view and (B) side view. The inset in panel A shows the Fourier transform of the image [117]

PAGE 105

105 Figure 7 3. SEM images of binary arrays of silica sp heres and polymer posts after a second oxygen plasma etch: (A) top view and (B) side view [117]

PAGE 106

106 Figure 7 4. Top view SEM image of a binary gold nanohole array templa ted from a double layer colloidal crystal consisting of 320 nm silica spheres. The inset shows a magnified image. [117]

PAGE 107

107 Figure 7 5. SEM images illustrating the progressive fabrication procedures f or binary view image of a double layer colloidal crystal template consisting of 340 nm silica spheres; (B) top view image after removing silica spheres by wet etch; (C) side view image of the sample in panel B ; and (D) replicated ETPTA binary array from the sample in panel B [117]

PAGE 108

108 Figure 7 6. Experimental (solid) and RCWA simulated (dotted) normal incidence specular reflection from a bare glass substrate (black), a glass slide covered with a templ glass slide coated with an ETPTA unitary hemispherical nipple array (blue ). [117]

PAGE 109

109 CHAPTER 8 CONCLUSIONS AND FUTURE WORK Conclusion Gibbsite nanoplatelets are assembled to form well ordered layered struct ure via different bottom up approaches. W e have applied several technolog ies for assembling gibbsite nanoplatelets into large area, self standing fi lms. These nanosheets with high aspect ratio are preferentially aligned parallel to the substrate Gibbsite polymer nanocomposites are made by infiltrating monomers into layered structure or co deposition. ETPTA monomer can be infiltrated into gibbsite film and polymerized to form nanocomposite. Gibbsite and PVA can be deposited together to form elastic, optical ly transparent nanocomposites. The tensile st rength and the stiffness of these biomimetic nano composites are significantly improved compared to pure polymer films; th o se of crosslinked nanocomposite is even higher than that of no n crosslinked and pure poly mer. The theoretical prediction based on a shear lag model agree s with our experimental results The monomer infiltration process is limited by low viscous monomers. Co deposition with controlled colloidal agglomeration is the way to blend the strong and e lastic polymers with nanoplatelets. But t he mechanical properties of both gibbsite based nanocomposites are not comparable to that of nacre. Ultrastrong GO sheets and elastic PVA polymer are combined to form tough GO/PVA nanocomposite by co infiltration. We have developed a simple paper making (vacuum filtration) approach to assemble graphene oxide/PVA mixture suspension into free standing composite film. The high aspect ratio GO sheets are well aligned parallel to the filter membrane and formed highly ord ered load distributed structure. GA form bindings to inorganic component which can strengthen composite; however, the cross

PAGE 110

110 linking network makes the polymer inelastic and weakens the toughness of composite. Nanocomposite can obtain maximum mechanical prop erties with proper ratio of GO, PVA, and crosslinkers. GO/PVA co mposite exhibits over 2 fold ultimate strength to filtrated GO paper and nacre, ~30 times tougher than filtrated graphene paper and nacre, ~100 times tougher than filtrated GO paper under spec ific ratio of components Simulation results agree with our experimental values and show the potential to obtain graphene nanocomposite with better mechanical properties. Besides reinforced nanocomposites, we explored the self assembly of spherical collo ids and the templating nanofabrication of moth eye inspired broadband antireflection coatings. We combine spin coating assembly and templating of colloidal crystal to generate complex periodic binary structure. Its applications includ e intercalated arrays of silica spheres and polymer posts, gold nanohole arrays with binary sizes, and dimple find that binary coatings. N atural nano com posite s and optical structure s teach us a great deal on how to create high performance a rtificial material s The bottom up technologies developed in this thesis are scalable and compatible with standard industrial processes, promising for manufacturing hig h performance materials for the benefits of human beings.

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111 Future Work I nspired by natural materials, we find our own way to create high performance materials. H owever, they have some drawbacks : GO sheets are unstable over 200 o C. The elastic modulus of GO /PVA composite is lower than that of pure graphe n e paper and other metal alloy. Therefore, t he following are several recommendations for future work: To improve thermal stability of the composite, w e will u se chemically reduced graphene oxide ( RGO ) sheet s to replace GO RGO shows similar mechanical behaviors, better conductivity and thermal stability compared to g raphene oxide. Recent studies also show that RGO sheets can remain dispersed at least 20 hours before aggregation, which make it possible to fab ricate RGO /PVA composite by paper making approach. [166] To strengthen composite and preserve its elasticity, w e will use m odifiers and polymer with functional groups which form chemical bond with graphe n e sheets t o r einforce interaction between organics and inorganics B esides, 3 0 nm scale silica particle can be blended in composite to enhance fraction between graphene sheets. We might apply polymers with special functional groups (such as c hitosan and p olyallylam ine ) g raphene sheets with aspect ratio over 1000 which could enhance tensile strength of the composite, based on theory. PVA polymers with different molecule weight can also be evaluated.

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122 BIOGRAPHICAL SKETCH We i han Huang (Erik) is from Taiwan, received his Bachelor of Science in Chemical Engineering in National Taiwan Uinversity in August 2003. He studied in UF since August 2006 and joined Dr. Peng Jiang's group in summer 2007. He receive s his Ph.D. from the De partment of Chemical Engineering at U niversity of Florida in August 2011