A REVIEW and STUDY of VERY HIGH NANOFILLER-CONTENT NANOCOMPOSITES: THEIR PREPARATION METHODS, CHARACTERIZATION and PROPERTIES By

A REVIEW AND STUDY OF VERY HIGH NANOFILLER-CONTENT

NANOCOMPOSITES: THEIR PREPARATION METHODS, CHARACTERIZATION

AND PROPERTIES

JEFFREY PHILIP GEORGE

Submitted for partial fulfillment of the requirements

For the degree of Master of Science

Thesis adviser: Dr. Hatsuo Ishida

Department of Macromolecular Science and Engineering

CASE WESTERN RESERVE UNIVERSITY

August, 2019

CASE WESTERN RESERVE UNIVERSITY

SCHOOL OF GRADUATE STUDIES

We hereby approve the thesis/dissertation of

Jeffrey Philip George

candidate for the degree of Master of Science*.

Committee Chair

Dr. Hatsuo Ishida

Committee Member

Dr. David Schiraldi

Committee Member

Dr. Gary Wnek

Date of Defense

June 6th, 2019

*We also certify that written approval has been obtained for any proprietary material

contained therein.

Table of Contents

Table of Contents...... i List of Figures...... iv List of Tables...... x Acknowledgement...... xi Abstract...... xii Chapter 1: A Review on the Very High Nanofiller-Content Nanocomposites: Their Preparation Methods and Properties with High Aspect Ratio Filler...... 1 1.1 Introduction...... 2 1.2 Definition of Nanofillers...... 4 1.3 Types of Nanofillers...... 6 1.3.1 Clay Nanofiller...... 6 1.3.2 Graphene Oxide Nanofiller...... 8 1.3.3 Other Nanofillers...... 9 1.4 Different Fabrication Techniques...... 9 1.4.1 Layer-by-Layer Deposition Techniques and Its Variants...... 10 1.4.2 Electrophoretic Deposition (EPD) ...... 14 1.4.3 Mechanical Self-assembly Methods...... 15 1.4.3.1 Earlier Mechanical Methods...... 16 1.4.3.2 Doctor Blading (DB) and Vacuum Assisted Filtration (VAF)... 17 1.4.3.3 Solution Casting (SC)…...... 21 1.4.3.4 Gel Casting and Hot-pressing Fabrication Method...... 25 1.4.4 Other Novel Techniques...... 26 1.4.4.1 Chemical Assembly...... 26 1.4.4.2 Gel-film Transformation Method...... 29 1.4.4.3 Ice Templating...... 30 1.4.4.4 Freeze Drying...... 31 1.4.4.5 Wet Spinning...... 33 1.4.4.6 Other Methods for Viscous Systems...... 34 1.4.5 Fabrication of Single Component Nanocomposite Systems...... 36 1.5 Surface Modification...... 37

1.5.1 Surface Treatment of Polymer...... 37 1.5.2 Chemical Treatments of Inorganic Nanofillers...... 39 1.5.3 Grafting Modification of Nanoparticles...... 42 1.6 Mechanical Properties of High Filler-content Nanocomposites...... 46 1.6.1 Effects of Interfacial Interactions...... 46 1.6.1.1 Hydrogen Bonding...... 47 1.6.1.2 Covalent Bonding...... 50 1.6.1.3 Electrostatic Bonding...... 51 1.6.1.4 Synergistic Combinations of Bonding...... 54 1.6.2 Effects of Crosslinking Agents...... 59 1.6.3 Effects of Fabrication Methods...... 65 1.6.4 Effects of Electrochemical Reduction...... 67 1.7 Electrical Properties of High Filler-content Nanocomposites...... 71 1.8 Gas Barrier Properties of High Filler-content Nanocomposites...... 75 1.9 Flame Retardancy Properties of High Filler-content Nanocomposites...... 78 1.10 Conclusions...... 81 Chapter 2: Application of Very High Nanofiller-content Laponite/DNA Nanocomposite coating on Polyurethane Foam through Single-Dip Fabrication 83 2.1 Introduction...... 84 2.2 Experimental Section...... 87 2.2.1 Materials...... 87 2.2.2 Preparation of LAP/DNA suspensions...... 87 2.2.3 Preparation of coated PU foams...... 88 2.2.4 Preparation of LAP/DNA nanocomposite films...... 88 2.2.5 Characterization...... 88 2.3 Results and Discussion...... 89 2.3.6 FT-IR spectroscopy...... 89 2.3.7 SEM and EDX analysis...... 91 2.3.8 Thermal Stability...... 92 2.3.9 Compression Testing...... 95 2.3.10 Flammability...... 97

2.4 Conclusion...... 99 2.5 Acknowledgement...... 99 References...... 100

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List of Figures

1. Figure 1.1: Geometric definition of nanofillers. (a) nanoparticle, (b) nanorod, (c) nanoplatelet, (d-f) secondary particle or aggregate form of each nano-fillers...... 4 2. Figure 1.2: Proposed structural models of MTM/PVA nanocomposites prepared by evaporation-induced assembly. (a) The composites with MTM weight fraction, w, between 0 wt% and 30 wt% have a random structure. The crystallinity of PVA decreases, accompanied by the increase of the constrained region. (b) The composites with w between 30 wt% and 70 wt% have a nacre-like layered structure. PVA is completely constrained by MTM platelets. (c) The structure of composites with w between 70 wt% and 100 wt% is transformed to tactoids...... 5 3. Figure 1.3: The structure of a 2:1 layered silica...... 8 4. Figure 1.4: The structural model of Graphene oxide...... 8 5. Figure 1.5: (A) Schematic diagram of a LbL assembly method. (B) Simplified picture of alternating layers of oppositely charged polymers being adsorbed on substrate. (C) Chemical structures of sodium salt of poly(styrene sulfonate) and poly(allylamine hydrochloride)...... 11 6. Figure 1.6: Free-standing, 300-bilayer PVA/MTM composite film showing high flexibility and transparency...... 12 7. Figure 1.7: The SA-LbL fabrication of CH/AL nanocomposite films... 13 8. Figure 1.8: The scheme showing the electrophoretic deposition assembly to construct nano-laminated film structure...... 15 9. Figure 1.9: Schematic illustrations of the different physical methods used to increase the orientation of the talc tablets...... 17 10. Figure 1.10: A schematic representation of the VAF fabrication of nanocomposite films from a solution of dispersed inorganic nanofiller and polymer...... 18 11. Figure 1.11: SEM and TEM images of various layered PVA/MTM nacre- mimetic paper composites obtained via paper-making process. Panels a-d show films of thicknesses of 0.17, 0.063, 0.044, and 0.026mm,

demonstrating that even thick films can be prepared and the film thickness can be tuned. The digital photograph (e) demonstrates the light translucence of nacre-mimetic papers of various thicknesses (from left to right: 0.026, 0.044, and 0.08 mm) as induced by the highly aligned orientation of the nanoclay platelets. The higher resolution SEM images clearly reveal the strong orientation of the nanoclay platelets parallel to the substrate (f and g)...... 18 12. Figure 1.12: The fracture path of the reduced PCDO/GO composites upon loading. The curled long chain of PCDO is gradually stretched and broken with loading, simultaneously resulting in curving of the edges of the rGO sheets...... 21 13. Figure 1.13: Schematic of the bioinspired assembly process for the fabrication of a CMC/LDH nanocomposite film trough hydrothermal method coupled with solution casting method...... 22 14. Figure 1.14: The fabrication of the aligned bulk nanocomposites by evaporation-induced self-assembly followed by lamination...... 23 15. Figure 1.15: Schematic illustration of Hot-press Assisted Slip Casting method...... 25 16. Figure 1. 16: Schematic representation of Gel-casting and Hot-pressing method...... 26 17. Figure 1.17: Schematic illustration of in-situ polymerization of hexamethylene adipamide on VMT...... 27 18. Figure 1.18: Schematic illustration of in-situ inverse microemulsion polymerization...... 29 19. Figure 1.19: Synthetic route of PAPB/GO hydrogel nanocomposites via Gel-film transformation method...... 30 20. Figure 1.20: Photographs of GO films containing 2 wt % CH (top) and the films with different shapes processed by origami and paper-cutting (bottom)...... 30 21. Figure 1.21: Schematic representation of Ice-templating where porous ceramic scaffolds are fabricated...... 31

22. Figure 1.22: (a) before freeze drying, and (b) after freeze drying. Samples from left to right contain 0wt% content polymer to 100wt% polymer.. 33 23. Figure 1.23: Preparation of PVA-coated rG nanocomposites through wet spinning. (A) aqueous solution containing GO and PVA. (B) PVA-coated rG solution after removing the free PVA. (C) Pre-aligned liquid crystalline PVA-coated rG. (D) Wet-spinning assembly for preparation of continuous nacre-mimic fibers...... 34 24. Figure 1.24: Schematic of the coextrusion assembly...... 35 25. Figure 1.25: (a) Structure of polymer-grafted-nanoparticles at different graft densities. (b) The different structures represented as matrix is absent for the HNP assemblies and added for the swollen and dispersed phase 36 26. Figure 1.26: Illustration for the Preparation of DA-Conjugated NFC from Pulp Fibers and Subsequent Assembly of DA- NFC with MTM to Prepare Transparent Nanocomposite...... 39 27. Figure 1.27: Synthesis of homo-telechelic functionalized PEG precursors...... 39 28. Figure 1.28: SEM images of self-assemblies with two levels of MVMT. (a) Shows the fiber-like assembly and (b) the presence of multiple layers when added with MVMT containing 11% coupling agent. Adding MVMT with 17% coupling agent resulted in nacre-like structures (c and d)...... 40 29. Figure 1.29: Tensile stress vs. strain curves for plain ethoxylated trimethylolpropane triacrylate (ETPTA) film, ETPTA–gibbsite nanocomposite, and ETPTA–PEI–SCG nanocomposite...... 41 30. Figure 1.30: Schematic diagram of the mechanism of formation of PPY/CNT nanocomposites...... 42 31. Figure 1.31: Schematic representations of 'graft-to' and 'graft-from' methods...... 43

32. Figure 1.32: Experimental dispersion diagram of PGMA-grafted TiO2 NPs within an epoxy matrix. The solid and dash lines represent the maximum achievable graft densities and the “brush-to-mushroom” transition, respectively...... 46

33. Figure 1.33: Schematic diagram depicting the ability of different intercalating species to contribute to gallery-spanning hydrogen bond bridges. A) Anhydrous stacked GO sheets can form a small number of hydrogen bonds between surface-bound groups when the sheets are close enough to each other. B) Water molecules, which are both hydrogen bond donors and acceptors, between stacked GO sheets create a network of many hydrogen bonds that can readily adapt to mechanical stresses. C) DMF molecules between stacked GO sheets reduce the intersheet interactions due to the limited ability of DMF to hydrogen bond (it can only accept hydrogen bonds but not donate). D) PVA chains between stacked GO sheets increase the intersheet interactions, not only due to the ability of PVA to hydrogen bond in a similar fashion to water, but also strengthen the bond network with the covalent C–C bonds in the polymer (i.e., between the hydrogen-bonding capable monomer units), creating a very stiff structure. E) PMMA chains between stacked GO sheets are similar to the DMF molecules in case C, because PMMA can only accept hydrogen bonds...... 49 34. Figure 1.34: Stress–strain curves for M-NCs with different clay contents (5.5–23) and chemically crosslinked polymer (M-OR3). Schematic illustrations of typical yielding behaviors with a well-defined necking point are depicted (I, IIa, II, III)...... 50 35. Figure 1.35: Schematic illustration of fabrication procedure of ternary artificial nacre nanocomposites with dual covalent and hydrogen interactions...... 54 36. Figure 1.36: Formation process of covalent bonding between borate and GO nanosheets...... 60 37. Figure 1.37: Fabrication process of the reduced PCDO/GO nanocomposites with 푍푛2+...... 62 38. Figure 1.38: Stress–strain curves obtained by tensile testing of various samples with different counterions...... 64

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39. Figure 1.39: The preparation of self-assembled nacre-mimetic structure involving hydrogen bonds: First the anionic nanoclay (MTM) platelets are coated with cationic polymer (PDDA) via adsorption. Excess polymer is removed by washing. Hydrogen bonding molecules, anionic dGMP, were introduced into structure resulting in polymer-coated clays, which were linked by hydrogen bonding due to recognition sites (D and A) ...... 65 40. Figure 1.40: (a) Schematic illustration of fabrication of PVA/GO nanocomposite films followed by electrochemical reduction. (b) Color change of reduced nanocomposite films...... 69 41. Figure 1.41: Schematic illustration of Microstamping...... 70 42. Figure 1.42: Preparation of conductive nanocomposites with conductive polymer PEDOT:PSS or PIL. The bottom photograph shows a LED connected via conducting PEDOT:PSS/SA nacre-mimetics and powered at 3.5 V...... 72 43. Figure 1.43: Residues collected after vertical flammability test of NFC/ Clay nanocomposites...... 79 44. Figure 1.44: Heat release rate (HRR) and smoke production rate (SPR) plots of uncoated and clay nanopaper (NFC/30% MTM) coated GF/EP composites...... 80 45. Figure 1.45: Temperature profile during cone calorimetry, measured 2 mm under the surface of Coated (clay nanopaper coated wood) and Wood (uncoated wood) samples...... 80 46. Figure 2.1: FT-IR spectrum of (Top to Bottom) Laponite, DNA, Subtracted PU-LAP-90 (No PU), Control and PU-LAP-90...... 90 47. Figure 2.2: SEM images of fractured film surfaces: a) laponite, b) LAP-60, and c) LAP-90 films...... 91 48. Figure 2.3: EDX image mapping the distribution of Phosphorus across the LAP-90 film surface...... 91 49. Figure 2.4: TGA thermograms of control and coated PU foams in nitrogen atmosphere...... 93 50. Figure 2.5: Residue percentage and char yield percentage of coated PU

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foams...... 94 51. Figure 2.6: TGA thermograms of control and coated PU foams in air atmosphere...... 95 52. Figure 2.7: Compression strength of control and coated PU foams...... 97 53. Figure 2.8: Modulus of control and coated PU foams...... 97 54. Figure 2.9: Figure showing the pre and post burning images of (a, e) control PU, (b, f) PU-DNA, (c, g) PU-LAP-60, and (d, h) PU-LAP-90 98

List of Tables

1. Table 1.1: Effect of Hydrogen bonding on Mechanical properties...... 48 2. Table 1.2: Effect of Covalent bonding on Mechanical properties...... 51 3. Table 1.3: Effect of electrostatic interaction on Mechanical properties 52 4. Table 1.4: Effect of electrostatic interaction on positive rims of nanoplatelets on the mechanical properties...... 53 5. Table 1.5: Effect of synergistic interactions on mechanical properties 55 6. Table 1.6: Effect of glutaraldehyde and borate linkages on mechanical properties...... 60 7. Table 1.7: Effect of cationic metal ions on mechanical properties...... 62 8. Table 1.8: Effect of anionic counterions on mechanical properties...... 64 9. Table 1.9: Effect of electrochemical reduction on the mechanical properties of high filler-content nanocomposites...... 70 10. Table 1.10: Electrical properties of high-filler content nanocomposites 74 11. Table 1.11: Gas barrier properties of high filler-content nanocomposites for different humidity levels...... 76 12. Table 2.1: Residue percentage and char yield percentage data of coated PU foams...... 94 13. Table 2.2: Compression strength and modulus data of control and coated PU foams...... 96

Acknowledgement

I first would like to thank Jesus for giving me the strength, peace of mind and health in order to finish my research. I would like to extend my appreciation to my family for their support and encouragement which helped me complete my Masters. I would like to express my deepest gratitude to my professor, Dr. Hatsuo Ishida, for his patient and timely counsel through the course of my Masters. I certainly would not have been able to do well without his guidance and encouragement. I would also like to thank my past and present group members and mentees who have helped me along the way. I would also like to thank UL for their provision of characterized polyurethane foam as well as Dr. David Schiraldi and SCSAM for use of their equipment.

A Review and Study of Very High Nanofiller-Content Nanocomposites: Their Preparation Methods, Characterization and Properties

Abstract by JEFFREY PHILIP GEORGE

In the last twenty years, high nanofiller-content nanocomposites have seen a surge in interest within the research community owing to the interesting properties it offers. However, development of such nanocomposites has always been difficult due to the affinity of nanofillers to aggregate. But recently, novel and innovative ideas have been proposed to achieve fast and scalable processing techniques which yield homogenous and well oriented nanocomposite structures with excellent mechanical, electrical and flame- retardant properties. In the first chapter of this thesis, a detailed review of recent progress in high filler content nanocomposites will be presented along with their novel methodologies and various kinds of properties. This review will then serve as foundation for chapter 2, where a high filler content nanocomposite coating on polyurethane foams is developed through a new method called single-dip coating. The foams are then characterized and tested to evaluate their thermal stability, mechanical behavior and flame- retardant properties.

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Chapter 1. A Review on the Very High Nanofiller-Content Nanocomposites: Their Preparation Methods and Properties with High Aspect Ratio Fillers

Abstract For many years, the inorganic content in nanocomposites has never exceeded 10 vol%. Novel fabrication methods have been employed to incorporate high inorganic loading within the polymer matrix. This review reports recent progress in high filler- content nanocomposites and the various unique techniques used for their fabrication. This review will also highlight the mechanical, electrical, gas barrier and flame-retardant properties of the high filler-content nanocomposites.

1.1 Introduction Polymer composites have been of extensive interest in the past half a century, due to the rising need for materials with a specific blend of properties. Polymer composites are materials that are combinations of two or more organic and inorganic materials, to create a new material with enhanced physical cumulative properties of the parent materials. The polymer acts as the matrix, while the inorganic filler is dispersed homogenously in order to improve the physical properties of the composite. A unique set of properties is developed in the composites, such as enhanced mechanical properties, electrical conductivity, thermal conductivity, gas barrier properties and so on, lacking in naturally-occurring materials, is accomplished in the category of composites. Polymer composites may be further classified on the basis of the dimensions of the materials incorporated. Micro composites involve the incorporation of precursors with dimensions in the micrometer-scale while nanocomposites possess inorganic fillers with at least one of the three dimensions in the nanoscale range. In this current review, polymer nanocomposites will be the focus. Many types of inorganic nanofillers such as clay[1], graphene (G)[2], graphene oxide (GO)[3], alumina flakes (Al2O3)[4], boron nitride (BN)[5], molybdenum disulphide

(MoS2)[6], layered double hydroxides (LDH)[7], carbon[8], nanocellulose[9] and others have been used to enhance polymer properties . These inorganic nanofillers impart new and unique properties due to the nanoscale interfacial interactions between the polymer and the inorganic filler. Over the years, the inorganic nanofiller weight concentration has always been maintained below 10 vol%.[10–14] This was due to the affinity of inorganic fillers to form clusters which act as stress concentrators in the material causing it to be brittle and fragile.[15–18] The polymer chains on a substrate that interact with the surface have been reported to possess a unique structure from the bulk due to their nanoconfinement. The properties of the polymer, e.g., mobility, crystallinity, orientation, etc., can be significantly altered through contact with an inorganic filler.[19,20] This interfacial region with finite thickness is termed interphase[19–21] whose thickness for the nanoconfined layer is roughly on the order of two radii of gyration of the polymer chain [22]. The fraction of the interphase to the total volume for a microcomposite is negligibly low and the contribution of the properties of the interphase to overall composite properties is not obvious other than its

2 role in adhesion. However, nanofillers, having very large surface areas, can increase the fraction of the interphase dramatically if each nanofiller is properly dispersed in the matrix. Larger surface areas of nanofillers correlate to the larger interfacial interaction regions between the inorganic filler and polymer. [21] The difference between composites incorporating nanofillers and larger sized fillers has been well documented in various studies. [5,21,22] However, the majority of papers on nanocomposites with low nanofiller content report significant reduction of the unique properties at higher filling fractions of nanofillers. At these high concentrations, the nanofillers begin to aggregate, losing the high fraction of unique interphase properties. However, more recently, new and novel techniques have been introduced to fabricate well-dispersed high filler-content nanocomposites. The emergence of these new techniques was largely due to inspirations from naturally occurring biomaterials like nacre and bone which are renowned for their high strength, Young’s modulus and toughness.[23– 26] Their distinctive mechanical properties are due to the unique brick-mortar arrangement of organic and inorganic elements, which combines the elasticity of 10–50 nm organic protein layers and the strength of aragonite tablets that are 200–900 nm thick. Nacre is a high filler-content bio-composite with over 95 wt% aragonite platelets and just 5 wt% organic material.[27] The harmony between the inorganic and organic components has inspired many studies on high filler-content nanocomposites. A high filler-content nanocomposite is defined as a nanocomposite containing a polymer matrix loaded with inorganic filler of nanoscale dimensions with a volume content over 10 vol%. The first high filler-content nanocomposite was reported by Wu et al.[28] back in 1993. A high clay content of 30 vol% was incorporated with the poly (ethylene oxide) (PEO) polymer to obtain a good-structured nanocomposite. In the years that followed, various polymers and inorganic nanofillers have been combined in order to synthesize high filler-content nanocomposites with unique set of properties. Over the years, mechanical behavior was observed to improve as homogeneity and interfacial interaction improved. Electrical properties were also observed to improve owing to the creation of conductive pathways by the high content of nanofillers within the polymer matrix. Additionally, unique properties such as gas barrier and flame-retardant properties were also exhibited due to the tortuous diffusion path created by high filling fraction of nanofillers.

In this review paper, we wish to review the recent advances in the field of high nanofiller-content nanocomposites, including the various fabrication methods used, surface modification techniques and their unique properties. 1.2 Definition of Nanofillers The structure of a high filler-content nanocomposite is heavily dependent on the type of filler used. Fillers come in various shapes and sizes. As we are only concerned with nanocomposites in this paper, we will only examine fillers with geometric sizes on the order of magnitude of nanometers. A ‘nanofiller’ does not have any strict definition, but typically has geometric dimensions below 100 nm at least in one dimension.[29]

Figure 1.1: Geometric definition of nanofillers. (a) nanoparticle, (b) nanorod, (c) nanoplatelet, (d-f) secondary particle or aggregate form of each nano-fillers.[29] Nanofillers come in various shapes; mainly (a) nanoparticle, (b) nanorod, and (c) nanoplatelet (as seen in Fig 1.1). In the case of nanorod and nanoplatelet, aspect ratio (L/D), which is the ratio of the longest size L and the shortest size D of the particle, is generally used to define the geometric anisotropy of the nanofiller. These primary particles can aggregate into clusters of particles, and generally must be exfoliated to obtain a homogenous dispersion in a nanocomposite.[29] Depending on the geometric shape of the nanofiller, the volume fraction of the inorganic filler with respect to the polymer is an essential factor for the nanocomposite’s microstructure. For nanoplatelets, it was observed that a high-volume fraction of nanofillers creates a hierarchical layered structure of stacked inorganic platelets, with the polymer matrix present between each layer. The hierarchical ordered structure is, in many cases, mimetic to the brick-mortar structure of biological materials like nacre.[30–35] At a volume fraction lower than the critical value, the polymer matrix dictates the structure’s

4 formation and therefore gives rise to a disoriented, random dispersion of inorganic platelets in the polymer matrix.[36] These phenomena were best shown experimentally by Wang et al.[37] where the relationship between the nanofiller loading, and corresponding structure and performance of the nanocomposite was studied. It was observed that between the range of loading 0-30 wt% of clay nanofiller, the microstructure of the nanocomposite was observed to be random in nature. As the clay nanofiller loading increased from 30 wt% to 70 wt%, an alternating layered structure of clay nanofiller and polymer phase was observed (Fig 1.2).[37] Another type of structure that was observed in nanocomposites with low volume fraction of inorganic filler, is that the polymeric component forms a distinct phase separate from the adsorbed phase on the surface of the 2D nanoplatelets.[28,38–40] Controlling the volume fraction of inorganic nanoplatelets allows for tunability of the microstructure, and consequently its physical properties.

Figure 1.2: Proposed structural models of MTM/PVA nanocomposites prepared by evaporation-induced assembly. (a) The composites with MTM weight fraction, w, between 0 wt% and 30 wt% have a random structure. The crystallinity of PVA decreases, accompanied by the increase of the constrained region. (b) The composites with w between 30 wt% and 70 wt% have a nacre-like layered structure. PVA is completely constrained by MTM platelets. (c) The structure of composites with w between 70 wt% and 100 wt% is transformed to tactoids. [37] Likewise, this effect also holds true for nanorods due to its high aspect ratio and affinity to align in an energetically favorable direction. Above a critical volume fraction of nanorods, the 1D fiber-like fillers align in a hierarchical ordered structure, which is parallel to the surface. [8,41,42] For high volume fraction of nanoparticles, on the other hand, these spherical nanoparticles typically enter the interstitial sites within the polymer matrix, resulting in high interfacial interactions between the polymer and the nanoparticles. [5,43]

The polymer that enters the interstices of the aggregates behaves as though it is part of the fillers, leading to the increase in the apparent volume fraction and, thus, increased viscosity of the filled melt.[44] The Einstein coefficient, which is a constant that reflects the degree of viscosity increase as a function of the volume fraction of the filler, as shown in Equation (1) [45], increases in the order from spherical, rod-like to platelet.

휂 = 휂표 (1 + 푘퐸휙) ………………………………………………………………………. (1) where  is the viscosity of the filled system, o is the viscosity of the matrix, kE is the Einstein coefficient and  is the volume fraction of the filler. This is due to the particle- particle interaction being the point contact for the spherical particles, whereas for rod-like fillers, it can be the point contact or line contact. Complexity of the contact further increases for platelet where the point contact, line contact and area contact are all possible, depending on the volume fraction of the fillers. Naturally, the amount of the polymer confined between the filler particles increases in the order from point, line to area, suggesting that the nacre-like structure has the highest degree of confinement. In nanocomposites, maximum achievable inorganic reinforcement content corresponds to the maximum packing density that can be reached through the ordered structure of the reinforcement phase. For example, in the case of identical spherical particles, maximum packing density is 74 vol% for ‘face-centered cubic’ ordering of particles[46], although this value decreases to around 63.4 vol% for random packing of the spheres[47]. For perfectly aligned continuous fibers in a matrix, maximum reinforcement packing can be achieved with hexagonal ordering of fibers at the cross-section of the composite that is approximately 90.7 vol%, assuming a pore free system.[46] Nanoplatelets, on the other hand, can achieve a maximum packing density as high as 95 % for the brick-mortar structure as seen in nacre.[46,48] Depending on the type, geometric shape and dimension of the nanofiller, the volume content can be tuned in order to obtain the best physical properties. 1.3 Types of Nanofillers 1.3.1 Clay Nanofiller Due to its low cost, versatility and ready availability, clay nanosheets have been identified as a viable inorganic component for high filler-content nanocomposites. These

6 inorganic nanometer sized sheets are typically negatively charged with a number of hydroxyl groups available for crosslinking. Montmorillonite (MTM), vermiculite (VMT), laponite (LAP) (also called hectorite) and saponite (SA) are the most commonly used clays for fabrication of polymer-clay nanocomposites. In Fig 1.3, the structure of a 2:1 layered silicate is shown.[49] The reason why these materials have received a great deal of focus is due to their high aspect ratio and the unique intercalation/exfoliation characteristics.[15] MTM and VMT clays were observed to have higher cation exchange capacities and specific surface areas in comparison to other clay nanofillers. MTM clay nanofillers, especially, have smaller crystallite size and rougher platelet surfaces than other clays.[50] Incorporation of clay nanofillers in polymer-clay nanocomposites also leads to good mechanical, gas barrier and flame retardant properties.[1,49] For these reasons, MTM clays have been the most widely used clay nanofiller in literature. The first clay nanopaper was fabricated from bentonite clay by Hauser in 1938. [51,52] However, the material was somewhat opaque and possessed unfavorable gas barrier properties. Since the 1970s, polymer-clay nanocomposites have been actively studied to obtain nanocomposites with improved physical properties.[53] Most of these composites were mainly composed of the organic phase, with the clays being a minor component (<10 vol%). Increasing clay content above 10 vol%, was observed to yield mechanically weak nanocomposites. This was attributed to the inhomogeneous dispersion of clay sheets resulting in formation of clay agglomerates as defects in a brittle material. [53] Furthermore, by reforming multiple silicate layers and increasing the total thickness of the aggregates, the aspect ratio decreases, negating the advantage of nanocomposite formation. For this reason, many studies have been done to investigate the polymer-clay interface and understand how to achieve homogenous dispersion to obtain improved properties.[28,38,60–63,39,40,54–59]

Figure 1.3: The structure of a 2:1 layered silicate [49]

1.3.2 Graphene Oxide Nanofiller Graphene (G), a 2D nanomaterial, possesses exceptional physical properties due to its pure sp2 hybridization network.[64] It is known for its excellent mechanical properties, exceptional electrical and thermal conductivity . This makes it a promising material for various applications like the field of aerospace, manufacture of flexible supercapacitor electrodes, and in artificial muscle and tissue engineering.[64] Pristine G is typically obtained by the top-down method of mechanical exfoliation and the bottom-up method of chemical vapor deposition (CVD).[64] However, due to its limited ability to scale-up and time-consuming nature, focus has been shifted to its precursor, graphene oxide (GO). GO is usually obtained by Hummers’ method[65] or modified Hummers’ method as described below through chemically functionalizing mineral graphite flakes. However, this method does have some shortcomings, such as low yield and toxic gas production. Hence, modified and improved methods have been reported by various researchers. Those include elimination of NaNO3[66], preoxidation prior to the Hummers procedure[67], modification of KMnO4/NaNO3 ratio[68], and replacement of KMnO4 with K2FeO4 without the use of

NaNO3[69].

Figure 1.4: The structural model of Graphene oxide [70]

The resulting GO sheets, a water-soluble derivative of G, are an ideal candidate for fabricating artificial nacre, due to the presence of many functional groups (hydroxyl and epoxide groups) that are available for crosslinking.[3,70,71] In fact, GO is one of the few platelet nanofillers that has reactive functional groups on the basal plane available for covalent bond formation. (in Fig 1.4, the structural model of graphene oxide is shown[70]). Silicate layers derived from clay do have hydroxyl groups but these groups are buried inside the silicate layer, and the distance from the basal surface makes it unavailable for covalent bond formation. Only the silanol groups exposed at the edge plane are capable of forming covalent bonds. The planar feature of the GO sheets makes it easy to assemble into paper-like structures through self-assembly and filtration techniques. [72,73] Remarkable mechanical properties of the GO monolayered-sheets (Tensile strength: 63 GPa, Young’s modulus: 200–500 GPa) combined with high surface areas, contribute to the fabrication of mechanically exceptional nanocomposites. [70,74,75] 1.3.3 Other Nanofillers Layered double hydroxides (LDH) are a category of ionic lamellar compounds made up of positively charged brucite-like layers with an interlayer zone containing anions and solvation molecules.[76] One popular and effective method used to fabricate LDH-based high filler-content nanocomposites is the electrophoretic deposition (EPD) method due to the ionic nature of the inorganic filler.[7,77,78] Another common inorganic filler is alumina (Al2O3) (also known as aluminum oxide) which is a thermodynamically stable material with a trigonal symmetry.[79] Due to its favorable physical properties, Al2O3 has been commonly used in the fabrication of high performance nanocomposites.[80] The other nanofillers commonly used in the literature are boron nitride (BN) [5,43], carbon nanotubes (CNT) [8,81] and silica (SI) nanoparticles [82,83]. Other more recent nanofillers, such as MoS2 [84] and black phosphorus (phosphorene), might offer unique potential due to their interesting properties; however, no attempts to manufacture high nanofiller-content nanocomposites have yet been reported. Black phosphorus is a graphite analogue and thus there are a vast number of potential applications that can be considered if environmental issues can be overcome. A black phosphorus based nanocomposite partially overcoming this problem has recently been reported.[85] 1.4 Different Fabrication Techniques

In order to synthesize high performance high filler-content nanocomposites, good control must be established over the content of inorganic phase, homogeneity of hard and soft component, development of microstructure, optimization of interface, and growth of fabrication techniques. Various techniques have been developed over time to increase the control of the structure’s nanoscale formation. This started with the layer-by-layer (LbL) technique to have finer control over the microstructure. Soon afterwards, electrophoretic deposition (EPD) and freeze drying became viable techniques to fabricate nanocomposites. Mechanical self-assembly methods like vacuum-assisted filtration (VAF) (or paper making), doctor blading (DB), etc. also found wide applications due to its low cost and large scalability. Other novel techniques like gel-film transformation (GFT), sol-gel technique, and several others will all be summarized below. The high filler-content nanocomposites fabricated by the above-mentioned techniques will also be highlighted. 1.4.1 Layer-by-Layer Deposition Techniques and Its Variants Layer-by-layer (LbL) technique allows the fabrication of films with good nanofiller alignment by alternately immersing a clean substrate into two solutions, which interacts and assembles layer-by-layer upon each other. This was first developed by Decher et al.[86] who alternately assembled oppositely charged polyelectrolytes on top of each other (in Fig 1.5, the schematic diagram of LbL process is represented[86]). This method was taken and employed for the fabrication of nanocomposites with alternating layer of inorganic nanofillers and polymer. The interaction and assembly are driven forward by a driving force which can change depending on the different nanofillers and polymers. The driving forces typically include electrostatic interactions, hydrogen bonding, π − π interactions, covalent bonding, and hydrophobic-hydrophobic interactions. This technique allows for precise control over the different layers stacked, consequently enabling homogenous dispersion of high nanofillers (generally nanoplatelets) within the nanocomposite system. Tang et al.[30] implemented the LbL assembly technique to prepare an electrostatically bound nacre-mimetic structure between cationic poly(diallydimethylammonium) chloride (PDDA) and anionic MTM clay. Cationic PDDA and anionic MTM clay were selected for this purpose owing to their high affinity and compatibility with each other as a result of the strong attractive electrostatic and van der

Waals interactions at the PDDA/MTM interface. The tensile strength and Young’s modulus of the nacre-like systems were measured to be 100±10 MPa and 11±2 GPa, respectively. The improved mechanical behavior of the system was credited to the homogeneity of the nanocomposites brought about by the fabrication technique, as well as the strong interlayer adhesion of ionic bonds and van der Waals interaction which promotes strong matrix connectivity.

Figure 1.5: (A) Schematic diagram of a LbL assembly method. (B) Simplified picture of alternating layers of oppositely charged polymers being adsorbed on substrate. (C) Chemical structures of sodium salt of poly(styrene sulfonate) and poly(allylamine hydrochloride). [86] Subsequently, Podsiadlo et al.[87] replaced PDDA with chitosan (CH), expecting that the stronger cationic nature and higher mechanical properties of the polymer would translate into much superior nanocomposites. Counterintuitively, the nanocomposite had poorer mechanical strength (81 MPa) than the CH polymer itself (108 MPa). This was attributed to the rigid molecular structure of the polymer which limited an efficient load transfer. Alternatively, L-3,4-dihydroxyphenylalanine (DOPA) polymer adhesive, was used to induce covalent crosslinks within the sequentially deposited nanocomposite system.[33] This strong covalent crosslinking in the presence of Fe3+ led to a strength of 200 MPa, which is twice the value of the PDDA/MTM layered nanocomposite, while the toughness value was also increased to 4.2 MJm−3. Hydrogen bonding of poly(vinyl alcohol) (PVA) was similarly used as the driving force in LbL assembly to fabricate

11 homogenous nanocomposites with high inorganic content, and consequently an improved tensile strength and Young’s modulus of 150 MPa and 13 GPa, respectively, were achieved (Fig 1.6).[88,89] The free standing film with 300 bilayers with good flexibility and transparency is shown in Fig.4.[89] Gas barrier properties and flame retardant properties were also achieved owing to the compact hierarchical layered structure and high inorganic loading achieved by the LbL technique.[90,91] Xu et al.[92] developed robust PDDA/MTM films with high modulus of 9.4GPa that also possessed stable underwater superoleophobicity. In order to mimic the structure of nacre, Wei et al.[93] utilized LbL to fabricate calcium carbonate (CaCO3) based multi-layered films with polymers diazo-resins (DAR) and poly (acrylic acid) (PAA). An organic content as little as 6.9 wt% was achieved. Likewise, Finnemore et al.[94] applied this method to assemble CaCO3 platelets with polymers PAA and poly (4-vinyl pyridine). Later, layered double hydroxide (LDH) nanofillers were assembled using the LbL method to obtain mechanically improved nacre- mimetic structures.[95–97] SI particles were used as nanofillers and incorporated along with PDDA and poly(sodium 4-styrenesulfonate) to assemble polyelectrolyte/silicate hybrid nanolaminates through the LbL method.[98] Single walled carbon nanotubes (SWNT) were also successfully assembled by the LbL method through the synergistic interaction of covalent and electrostatic bonding.[8]

Figure 1.6: Free-standing, 300-bilayer PVA/MTM composite film showing high flexibility and transparency. [89] Although the LbL assembly method proved to be a powerful fabrication technique to obtain highly controlled structures, it was extremely laborious and time consuming to execute. Podsiadlo et al.[99,100] introduced a faster LbL method, known as the exponential LbL method (e-LbL) which had a film growth rate over 100 times faster than

12 the conventional LbL method. However, the filler content was low (~10wt%). Cho et al.[101] first utilized the spin-assisted layer-by-layer (SA-LbL) assembly method, which was 10 times faster than the conventional LbL method. The nanocomposite structures obtained were observed to be much more highly ordered than those prepared by the conventional LbL method. Vertlib et al.[102] achieved a high nanofiller content of 60 wt% LAP through the SA-LbL technique, with the Young’s modulus improving 17-fold due to controlled layered structure.

Figure 1.7: The SA-LbL fabrication of CH/AL nanocomposite films [103]

Bonderer et al.[103,104] later utilized the SA-LbL method to fabricate nanocomposites with Al2O3 nanoplatelets and polymers like CH and polyimide (PI) (Fig 1.7). Highly ordered structures with high tensile strength (300 MPa) and strain to failure (25 %), were fabricated at a volume fraction of nanofillers as high as 0.15. Hu et al.[105] fabricated high performance graphene oxide (GO)-based nanocomposites by the SA-LbL method. The driving force was a combination of synergistic bonding interactions (hydrogen bonding, and polar–polar and hydrophobic–hydrophobic interactions) between the silk fibroin (SL) polymer and the GO nanofiller. High mechanical properties with elastic modulus as high as 145 GPa was obtained owing to these synergistic interactions. To this date, the highest performance nanocomposite synthesized by the SA-LbL method, was by Xiong et al.[106], with GO nanofiller and cellulose nanocrystals (CNC). Post electrochemical reduction, exceptional mechanical properties were obtained with tensile strength and Young’s modulus reaching values of 655 MPa and 169 GPa respectively. Shu et al.[107] synthesized a LDH based nanocomposite with PVA polymer through the SA- LbL technique. A filler loading of almost 98 % was achieved, with high mechanical

strength of 169.4 MPa being reported. Likewise, Dou et al.[108] fabricated highly ordered, multi-layered cellulose acetate (CA)/LDH nanocomposite films that possessed high gas barrier properties (oxygen transmission rate (OTR) <0.005 cm 3 m−2day−1atm−1 for dry conditions). Burghard and his colleagues[109–112] developed an interesting combination of techniques, consisting of LbL and chemical bath deposition (CBD) to obtain periodic microstructures of organic and inorganic layers. The LbL technique controlled the layered deposition of the oppositely charged organic polyelectrolyte (PE) layer while the CBD technique controlled the inorganic layer growth which involved particle growth at surface nucleation sites of polyelectrolytes and simultaneous particle deposition from solution.

[111] Various nanocomposite structures like the PE/TiO2 nanocomposite[110,112],

PE/ZnO nanocomposite[109], and PE/ZrO 2 nanocomposite[111] were reported to be fabricated from this technique. The advantage of this combination of techniques is that the density of their inorganic bridges could be regulated by the thickness of their organic layers. The control over molecular structure allowed the nanocomposite films to reproduce similar mechanical properties of nacre, such as improved Young’s modulus and strength combined with improved toughness.[110] 1.4.2 Electrophoretic Deposition (EPD) EPD is a facile, inexpensive and large-area scalable technique which can rapidly produce nanocomposite films. The technique involves two steps: (i) particles suspended in a liquid medium are forced to move towards an electrode by applying an electric field to the suspension (electrophoresis), and (ii) the particles are collected at one of the electrodes and forms a coherent deposit on it (deposition).[113] Fig 1.8 showcases the assembly of alternating layers of nanofillers and polymer through EPD. [114] EPD has been extensively used in many studies to produce highly-ordered large-area nanocomposite films.

Figure 1.8: The scheme showing the electrophoretic deposition assembly to construct nano-laminated film structure.[114] Long et al.[115] implemented this technique to fabricate highly ordered poly(acrylamide) (PACA) intercalated MTM clay structures. A high nanofiller content of 95.3 wt% was incorporated with hardness and Young’s modulus values of 0.95 GPa and 16.92 GPa, respectively, which exceeded those of pure MTM prepared through EPD. Later, Lin et al.[114] intercalated the MTM interlayers with an anodic electrophoretic resin and fabricated an ordered structure through EPD, which saw an increase in reduced modulus from 2.9 GPa for pure MTM clay film to 5 GPa for the nanocomposite film. In a series of studies, Lin et al.[7,77,78,116] assembled LDH nanoplatelets into nacre-like structure by electrodeposition technology. High-filler content was always achieved, but mechanical properties of the nanocomposites were poor (tensile strength of 55 MPa and strain to failure(%) of 0.01 %).[7] Even after modifying the LDH nanoplatelets to induce surface roughness, the strain to failure (%) was the only mechanical property observed to slightly improve to 0.08 %.[116] Similar results were observed by Furman et al., who achieved high filler loading, yet only slightly improved mechanical properties.[117] The reason for poor mechanical properties of electrophoretically deposited nanocomposites as compared to those fabricated by LbL assembly, may be due to poor optimization of the interface and weak alignment of silicate layers. Hence, although EPD possesses the unique desired properties like rapid processing, scalability and low cost, the lack of good mechanical properties limits its feasibility for industrial applications. 1.4.3 Mechanical Self-assembly Methods Although they possess certain powerful advantages, the LbL assembly and the EPD technique still have serious limitations in their large-scale applications. In order to

15 overcome these limitations and provide a new facile, rapid and low-cost method to produce highly ordered structures with scale-up abilities, many authors over the years have come up with new and novel techniques which allow for self-assembly of the nanocomposite. Several macroscopic processes like vacuum-assisted filtration (VAF), doctor blading (DB), painting, simple solution casting (SC), etc. have been implemented to fabricate oriented structures. The general advantage of these methods is that thick films, even bulk materials, are produced rapidly and economically. This allows for extensive application with appreciable results. 1.4.3.1 Earlier Mechanical Methods Almqvist et al.[118] attempted to fabricate well-ordered nacre-like nanocomposites with talc tablets by utilizing various techniques such as centrifugation, shearing cylinders, sedimentation, dipping, spinning cylinder, spinning plate and shearing plate. In Fig 1.9, the effectiveness of the seven methods to achieve high alignment, was observed to increase from centrifugation to shearing plates. Later, Liu et al.[119] implemented five methods of sedimentation, centrifugation, slip casting, filtration and electrophoresis and found that slip casting method gave the best alignment. Chen et al.[120] developed a centrifugal deposition process to rapidly deposit clay and polymer to form a layered polymer-clay nanocomposite. The time to complete this procedure was only about 15 minutes and achieved decent mechanical properties (tensile strength and Young’s modulus reached 80 MPa and 8 GPa, respectively) due to the highly ordered clay nanoplatelets. Jin et al.[121] employed simple centrifugation technique to fabricate electrostatically bound clay- nanofibrillated cellulose (NFC) nanocomposites with an excellent toughness value of 23.1 MJm−3. Park et al.[122] implemented a simple filtration technique to fabricate nanocomposites consisting of GO nanofiller and poly (allylamine) (PALA) which brought about a modest improvement of mechanical properties.

Figure 1.9: Schematic illustrations of the different physical methods used to increase the orientation of the talc tablets [118] 1.4.3.2 Doctor Blading (DB) and Vacuum-assisted Filtration (VAF) Method More focus was subsequently given to the introduction of other facile methods which preserve the excellent alignment of inorganic nanoplatelets. Walther et al.[123] achieved highly aligned materials of PVA/MTM having mechanical properties comparable to that of nacre, with simple and cost-effective techniques like vacuum-assisted filtration (VAF) (paper making) (Fig 1.10), doctor blading (DB) and painting. VAF technique was the most successful of the three, with DB technique and painting achieving less ideal mechanical properties due to imperfect microstructure. As observed in Fig 1.11, thick films could be prepared utilizing the VAF method with the ability to tune film thickness. Strong orientation of nanoclay platelets was also observed in the VAF-prepared nanocomposite films. For these reasons, VAF method was popularly used in the literature to assemble nanocomposites with good alignment. However, the DB technique was also used in some studies to fabricate nanocomposites with high nanofiller-content. Kunz et al.[124] utilized the DB technique to fabricate nanocomposite from hectorite clay and thermoplastic polyurethane (TPU) with a high filler content of 50 wt%. The resultant films possessed excellent mechanical properties with a modulus of 40 GPa. The DB technique was also eventually used to synthesize a bioinspired nanocomposite combining a liquid crystal polymer and LAP clay.[125] The nanocomposites possessed appreciable mechanical properties with a Young’s modulus of 14.5 GPa and hardness of 0.5 GPa.

Figure 1.10: A schematic representation of the VAF fabrication of nano-composite films from a solution of dispersed inorganic nanofiller and polymer[126]

Figure 1.11: SEM and TEM images of various layered PVA/MTM nacre-mimetic paper composites obtained via paper-making process. Panels a-d show films of thicknesses of 0.17, 0.063, 0.044, and 0.026mm, demonstrating that even thick films can be prepared and the film thickness can be tuned. The digital photograph (e) demonstrates the light translucence of nacre-mimetic papers of various thicknesses (from left to right: 0.026, 0.044, and 0.08 mm) as induced by the highly aligned orientation of the nanoclay platelets. The higher resolution SEM images clearly reveal the strong orientation of the nanoclay platelets parallel to the substrate (f and g).[123] Following its success in achieving mechanical properties comparable to those of nacre, the VAF technique was used by Walther et al.[127] to fabricate ordered structures of PDDA/MTM. The material produced had a tensile strength and Young’s modulus that was comparable to those of PDDA/MTM[30] prepared by the LbL method. In the years that followed, Yao et al.[128], Wang et al.[129] and Yao et al.[130] utilized the VAF technique but did not achieve significant improvement in the mechanical properties of the nanocomposites. Kochumalayil et al.[131] finally fabricated water stable polymer-clay nanocomposites via the VAF method. An improved Young’s modulus at 30 GPa and a tensile strength of 147.5 MPa were recorded at 50 % (relative humidity) RH with Young’s modulus only dropping down to 20 GPa at 90 %RH. The high stability was attributed to

18 the strong covalent bonds formed between the MTM clay and xyloglucan (XG). Yao et al.[132] and Liu et al.[133] implemented the VAF technique to form clay nanocomposites with nanofibrillated cellulose (NFC) as the polymer. Yao et al. modified their NFC with dopamine (DA) while Liu et al. used unmodified NFC in their nanocomposites. The end results at constant MTM nanofiller loading of 50 wt% showed significantly different results with the former having higher mechanical strength at 50 %RH as compared to unmodified NFC nanocomposites at standard conditions. This was likely due to the higher interfacial adhesion achieved by modification with DA between the polymer and clay. Owing to this effect, the DA modified NFC/MTM nanocomposites preserved its high mechanical strength of 280 MPa even at 100 %RH. Sehaqui et al.[134] used the VAF technique to create stiff, flexible and flame retardant films with excellent gas barrier properties (oxygen permeability (OP) of 0.20 cm 3mm m−2day−1atm−1 at 80 %RH). The facile and scalable production combined with the robust versatility of the nanocomposite make it potentially usable for packaging, transportation, construction, and insulation applications. Later, Liang et al.[135]utilized the VAF technique to fabricate alginate (ALG)/MTM nanocomposites with high strength of 280MPa as well as high toughness of 7.2 MJm−3. VAF technique was also successfully used to fabricate ternary high filler nanocomposites with mechanical strength comparable to that of natural biomaterial.[136] GO based nanocomposites were also extensively fabricated by the VAF method owing to the high fracture stress of the 2D nanoplatelets.[74] Putz et al.[126] synthesized high GO nanofiller composites using poly (methyl methacrylate) (PMMA) and PVA. High tensile strength of 148.3 MPa, which is comparable to that of nacre, was achieved for PMMA/GO films while high Young’s modulus of 36.4 GPa was achieved for the PVA/GO. In order to increase toughness of the final nanocomposite, Cheng et al.[137] used VAF method to introduce long chain polymer 10,12-pentacosadiyn-1-ol (PCDO) as the organic phase in the nanocomposite. As observed in Fig 1.12, the presence of the long chains of PCDO between layered reduced GO (rGO) nanosheets contributes significantly to the toughness of the nanocomposite. Upon loading, the curled long chain of PCDO gradually stretches and breaks, simultaneously resulting in the curving of the edges of the rGO sheets, and thereby imparting excellent toughness (3.91 MJm−3). Crosslinking was observed to be a major mechanical property improving factor by Tian et al.[138] Utilizing only the

VAF technique gave rise to uncrosslinked nanocomposites with weak hydrogen bonding, whereas supplementing with further crosslinking gave rise to formation of strong covalent bonds, significantly improving Young’s modulus (from 32.2 GPa to 84.8 GPa) and tensile strength (from119MPa to 178.9MPa). Considering the success of SA-LbL prepared SL/GO nanocomposites, Hu et al.[139] utilized the VAF method to fabricate high GO nanofiller (90 wt%) nanocomposites with SL as the soft component. Mechanical properties comparable to those of nacre were achieved which can be further significantly improved with electrochemical reduction. Biocompatible nanocomposite films with appreciable mechanical properties were fabricated using the VAF technique from gellan-gum (GG) and GO filler.[140] Exceptional mechanical strength and toughness values of 272 MPa and 18.1 MJm−3, respectively, were achieved for the VAF-fabricated GO nanocomposites with ALG as the organic component.[141] Wan et al.[142] achieved low mechanical properties for nanocomposites, prepared by the VAF method, with CH and GO nanofiller. But upon electrochemical reduction, the mechanical properties were remarkably improved with strength and toughness improved to 527 MPa and 17.7 MJm−3. A low Tg elastomeric polymer, poly (n-butyl acrylate) (PBA), was incorporated by Wu et al.[143] using VAF with GO nanosheets. The highest mechanical properties for binary nanocomposites, without electrochemical reduction, were achieved with a tensile strength of 187 MPa and toughness of 4.3 MJm−3. Ternary nanocomposites were also successfully fabricated by the

VAF technique. Synergistic combination of 2D MoS2 and GO sheets within a nanocomposite was accomplished by Wan et al.[6] with TPU as the polymer. Amorphous

Al2O3 and GO sheets were also successfully synergistically combined with (carboxy- methylcellulose) CMC as the polymer by Zhao et al.[144] The VAF prepared ternary nanocomposite exhibited dual enhancement of tensile strength and toughness, with an exceptional improvement over the binary nanocomposite of TPU/GO nanocomposite.

Figure 1.12: The fracture path of the reduced PCDO/GO composites upon loading. The curled long chain of PCDO is gradually stretched and broken with loading, simultaneously resulting in curving of the edges of the rGO sheets.[137] A study focused on evaluating the difference between mechanical, electrical and structural properties between the LbL-assembled structures and VAF-developed structures was done by Zhu et al.[2] Nanocomposites were synthesized employing both methods using PVA as polymer and reduced graphene (rG) sheets as inorganic nanofillers. The prepared nanocomposites were observed to have structural differences at the atomic and nanoscale level, but revealed similarities at a micrometer level. The mechanical properties, however, were observed to be similar in both cases of nanocomposites. This was concluded to be due to the thermodynamic state at the polymer/clay interface rather than the manner of packing of the inorganic nanofiller. 1.4.3.3 Solution Casting (SC) Solution casting (SC) (also known as solvent evaporation) was another technique which could induce good alignment of nanoplatelets. The solution casting method is an evaporation induced assembly technique which typically consists of a thoroughly dispersed solution of inorganic filler and polymer which is co-casted onto a petri dish and dried to finally obtain highly ordered nanocomposite systems. Darder et al.[145] first utilized this technique to intercalate CH between the silicate layers of MTM clay. A compact, robust and three-dimensional nanocomposite with high clay-content of 80 wt% was reported. Later, Yao et al.[128] applied this technique to order MTM platelets in a CH/MTM nanocomposite system. The resultant mechanical strength of 99 MPa was higher than that obtained by VAF (76MPa) in the same study. Ebina et al.[146] used this method to produce a transparent nanocomposite film with high nanofiller content. Electrostatically-bound clay-polymer nanocomposite systems with excellent gas barrier properties were achieved (OP 7.40 x 10-4 cm 3mm m−2day−1atm−1). In a similar manner, Guo et al.[147] utilized solution casting to fabricate high clay filler-content nanocomposites with good underwater

21 superoleophobicity. 2D nanosheets of clay, GO and LDH were also successfully incorporated into nanocomposite system by solution casting method, with many nanocomposite systems showing impressive mechanical properties.[35,148–154] Aulin et al.[155] utilized this technique to fabricate VMT-NFC nanocomposites with excellent modulus of 17.3 GPa and tunable oxygen and water vapor shielding properties. Water soluble spider silk was incorporated with hectorite clay by solution casting method to obtain strong nanocomposite films with excellent gas barrier properties, thermal and chemical stability, making it a viable packing material.[156] Following suite, Cui et al.[157] were able to achieve nacre-like microstructures for GO based nanocomposites with high mechanical strength of 205 MPa and toughness of 4 MJm−3. The hydrothermal method was coupled with the solution casting technique (Fig 1.13) by Wang et al.[158] to obtain a densely oriented microstructure of CMC/LDH films with good transparency and an appreciable strain to failure value of 8.4 %. Morits et al.[31] combined evaporation induced-assembly of PVA coated MTM nanoplatelets, with lamination of multiple of these PVA coated MTM films on top of one another (as seen in Fig 1.14). The finished product possessed a flexural strength of 220 MPa and flexural modulus of 25 GPa, owing to the ordered laminated structure formed. Dong et al.[159] fabricated a flame retardant GO nanocomposite paper by incorporating hexachlorocyclotriphosphazene (HCCP) with 66.6 wt% GO nanofiller. Finally, the evaporation induced assembly was also successfully used to fabricate ternary nanocomposites with high mechanical properties, with either clay or GO as the nanofiller.[41,160]

Figure 1.13: Schematic of the bioinspired assembly process for the fabrication of a CMC/LDH nanocomposite film trough hydrothermal method coupled with solution casting method. [158]

Figure 1.14: The fabrication of the aligned bulk nanocomposites by evaporation-induced self-assembly followed by lamination. [31] Polymer-polymer nanocomposites were also fabricated by the solution casting method. Dufresne et al.[161] fabricated nanocomposites with starch polymer amylopectin (AMP) with NFC, AMP being the matrix. The mechanical properties of the nanocomposite were tuned by varying the amount of plasticizer (glycerol)added. A high loading of 50 wt% nanofiber was achieved. Bruce et al.[162] used fibrilised cell wall material from the cell wall fragments of Swede root and incorporated it into a 40 wt% PVA matrix via the solution casting method. A tensile strength of 145 MPa and Young’s modulus of 8.9 GPa were achieved. Shi et al.[163] fabricated polyaniline (PANI)/cellulose nanocomposites with the cellulose being the reinforcement. Four wt% PANI was incorporated with cellulose to give nanocomposites with appreciable mechanical strength of 108 MPa and strain to failure of 11.5 %. Leitner et al [164] achieved appreciable mechanical properties for the NFC based nanocomposite with 10 wt% phenol-formaldehyde (PF) resin matrix, which showed a tensile strength of 127 MPa, a Young’s modulus of 9.5 GPa, and an elongation at break of 2.9 %. Carbon nanotubes-based nanocomposites were fabricated utilizing the solution casting method and variants of this method to obtain films with improved electrical properties. High nanofiller loading was required in order to create conductive pathways for good electrical conduction. Various polymers like poly(m-phenylenevinylene-co-2,5- dioctoxyp-phenylenevinylene) (PmPV), polypyrrole (PPY), poly(3-octylthiphene) (P3OT), PVA, PMMA and so on, were used as matrices for different carbon nanotubes (single-walled carbon nanotubes (SWNT) and multi-walled carbon nanotubes (MWCNT)) to obtain strong and durable films with good electrical properties. Shaffer et al.[165] incorporated a carbon nanotube content of 60wt% into a polymer PVA matrix by

23 employing the SC method. The resultant nanocomposite had an appreciable electrical conductivity of 100 S/m. Later, Kim et al.[166] used PMMA polymer with 40 wt% MWCNT fillers to fabricate nanocomposites that possessed electromagnetic interference shielding properties (~27 dB). TPU was incorporated with 10 vol% of MWCNT fillers to fabricate nanocomposites with a high strain to failure of 1200 % and a good electrical conductivity value of 1000 S/m.[167] Kymakis et al.[168] adopted a variant of the casting method, namely drop casting, where a drop of the solution mixture is dropped on the glass substrate and evaporated to obtain the nanocomposite. The P3OT/SWNT nanocomposite obtained was found to have an electrical conductivity of 1.00 x 10-4 S/m. A variant of the solution casting method termed spin casting was employed where a solution of well dispersed mixture of polymer and nanofiller was dropped onto a spinning glass substrate to obtain a thin electrically conductive film. Using this method, Kymakis et al.[169] fabricated nanocomposites with 30 wt% purified SWNT nanofillers and polymer P3OT. The resultant nanocomposites achieved a higher electrical conductivity of 5.00 x 10-2 S/cm. Curran et al.[170] employed this method to fabricate a nanocomposite film of poorly conductive polymer PmPV and CNT. The electrical conductivity improved by a factor of 108 over the initial pure polymer value to achieve the value of 0.01 S/m which was later improved by Coleman et al.[171,172] to achieve a value of 3.00 S/cm. Poly(3-hexylthiophene) (P3HT) was incorporated with 30 wt% SWNT nanofiller by employing the spin casting method to achieve an electrical conductivity of 1.00 x 10-4 S/cm.[173] Spin casting was again employed by Musumeci et al.[174] to fabricate P3HT/MWCNT nanocomposites which possessed a conductivity value of 0.5 S/m. This method was also used by Livanov et al.[175] to fabricate high content Al2O3 based nanocomposites with polymer PVA and PMMA in its interlayers. A technique similar to the solution casting method was slip-casting method in which a slurry of the nanocomposite dispersion is poured into a mold. Once the slurry filled the mold, the bulk portion of the slurry is drained, leaving only the materials that are near the mold surface remaining. The slurry that remained on the wall of the mold surface layer dries to form a solid resembling the mold shape. Abba et al.[34] utilized this method to fabricate parallel aligned microstructures of CH/Al2O3 nanocomposite films which

24 possessed appreciable mechanical properties in high humidity. Ekiz et al.[46] developed an efficient and relatively simple hybrid conventional method called “Hot-press Assisted Slip Casting” (HASC) through which bulk nacre-like nano-laminar composites could be achieved. HASC involves a liquid matrix pressured to flow through a porous mold, thereby decreasing the volume fraction of the matrix phase and forcing the Al2O3 sheets to align perpendicular to the applied pressure direction (Fig 1.15). The HASC-processed nanocomposite was reported to have a work of fracture (WOF) of 254 J/m2 owing to the excellent alignment brought about by the method. Sellinger et al.[25] used simple sol-gel dip coating technique to mimic nacre-like structure.

Figure 1.15: Schematic illustration of Hot-press Assisted Slip Casting method [46] 1.4.3.4 Gel Casting and Hot-pressing Fabrication Method In order to replace the laborious LbL technique, Bonderer et al.[176,177] reported the combined gel casting and hot-pressing fabrication method. As seen in Fig 1.16, the method involves a two-step process. (i) The polymer is dissolved in an organic solvent

Al2O3 platelet mixture at high temperature which is then cooled to form a gel network trapping the platelets in position. The cast gel films are then dried to improve orientation of the platelets. (ii) The cast and dried nanocomposites are cut into pieces, stacked and hot- pressed to further improve the platelet orientation and increase the density of entanglements in the polymer. Bonderer et al. used polypropylene (PP)[176] and TPU[177] as the polymer matrix to be used with the Al2O3 platelets. Remarkable strengthening was observed within each structure owing to the good alignment introduced by this combined technique.

Figure 1.16: Schematic representation of Gel-casting and Hot-pressing method [177] All the above mechanical assembly methods have the advantages of being fast, economical with easy scale-up ability. However, the one major disadvantage is the lack of fine control over the nanoscale formation of layered structure. 1.4.4 Other Novel Techniques Over time, many research studies have implemented new, novel or innovatively combined techniques to achieve a homogenously dispersed nanocomposite. Such techniques and the various literature implementing them will be highlighted below. 1.4.4.1 Chemical Assembly Chemical assembly, combining methods like evaporation-induced assembly and in- situ polymerization of monomers on the surface of exfoliated nanofillers, has also risen as a viable route for preparation of nacre-mimetic structures. This technique provides control of microstructure formation from the molecular level and thereby allows tunability of structure by changing certain molecular-level parameters. Modified VMT clay (MVMT) were utilized by Ma et al.[178] as reaction sites for in-situ polymerization of nylon (PA66) precursors (Fig 1.17). Highly thermally stable nanocomposites were built because of the strong covalent crosslinking achieved between inorganic MVMT and polymer matrix. Haraguchi et al.[179] utilized this method to fabricate nanocomposites of poly(2- methoxyethyl acrylate) (PMEA)/LAP nanocomposites with high filler loading. Extremely high strain to failure values of ~1000 % were obtained for these nanocomposites due to

26 the stable 3D network formed between the polymer and clay. A ternary GO/clay based nanocomposite was fabricated by Teng et al.[180] by a combination of evaporation and in- situ polymerization to achieve nanocomposites with high strain to failure values of ~700 %. In-situ polymerization was also utilized by Zhang et al.[9] to incorporate PANI polymer with 60 wt% rod-like nanocellulose to fabricate polymer-polymer nanocomposites.

Figure 1.17: Schematic illustration of in-situ polymerization of hexamethylene adipamide on VMT. [178] Carbon-based nanocomposites were also typically fabricated through in-situ polymerization of monomer on the surface of the carbon nano-filler. Pyrrole monomer was in-situ polymerized on the surface of CNT by Fan et al.[181] to yield nanocomposites with 50 wt% loading of nanofillers. The developed nanocomposites possessed high electrical conductivity with a value of 1600 S/m. Grimes et al.[182] fabricated carbon based nanocomposites by this method and observed to have an effect on its complex permittivity spectra as nanofiller loading increased. At 500 MHz, as the nanofiller loading of the CNT was increased from 0 to 23 %, the real permittivity was found to increase by a factor of ~35, and the imaginary permittivity by a factor of 1200. Sainz et al.[183] utilized in-situ polymerization to synthesize PANI/CNT nanocomposites. The resultant nanocomposites possessed luminescent properties with electrical conductivity reaching 100 S/m at a nanofiller loading of 50 wt%. Thiophene molecules were in-situ polymerized on the surface

27 of CNT by Karim et al.[184] to achieve nanocomposites with an electrical conductivity of 41 S/m. In-situ chemical oxidative polymerization directed by surfactants were also used to fabricate nanocomposites. The main advantage of this in-situ technique directed by surfactant is to increase the solubility of CNT in water and attain a more homogenous dispersion in the polymer matrix. The more homogenous dispersion yields nanocomposites with improved conductive pathways thereby increasing electrical conductivity. Zhang et al.[185] demonstrated this by fabricating carbon based nanocomposites by in-situ chemical oxidative polymerization of pyrrole on the surface of the CNT. The surfactant used was cetyltrimethylammonium bromide (CTAB) which improved the dispersion of carbon nanofiller in the polymer matrix. An electrical conductivity of 2300 S/m was achieved with 25 wt% loading as opposed to the 1600 S/m achieved by Fan et al.[181] at a loading of 50 wt%. PANI/CNT nanocomposites were fabricated through this method by Long et al.[186] and also observed to have an improved electrical conductivity of 127 S/m at 24.8 wt% as compared to 100 S/m achieved by Sainz et al.[183] for a loading of 50 wt%. In-situ inverse microemulsion polymerization (Fig 1.18) was another method that was developed in order to fabricate homogenous carbon-based nanocomposites.[187] This method is a simple and reproducible procedure in which the thickness and adherence of the coating are easily regulated by the monomer concentration and the reaction conditions. Yu et al.[188] utilized this in-situ polymerization technique to fabricate carbon based nanocomposites with PPY as the polymer and surfactant sodium dodecyl benzenesulfonate (SDBS). The resultant nanocomposites possessed an appreciable improvement in its electrical conductivity to achieve a value of 40 S/m.

Figure 1.18: Schematic illustration of in-situ inverse microemulsion polymerization. [188] Albeit that this form of chemical assembly allows for fine control at molecular level, expanding this to large scale application still remains a problem. 1.4.4.2 Gel-film Transformation Method Zhang et al.[189] reported a new, facile and cost-effective gel-film transformation (GFT) technique to fabricate high performance nanocomposite films through synergistic effects based on hydrogen bonding and π − π interaction. As seen in Fig 1.19, addition of trace amounts (1-8 wt%) of designed gel-promoters poly(acrylic acid-co-(4- acrylamidophenyl) boronic acid) (PAPB) to aqueous dispersions of GO, reinforces the GO network by enhancing the interaction between GO sheets, resulting in the formation of a hydrogel. These hydrogels were then cast dried at room temperature to obtain the gel-film transformed GO films (g-GO). The g-GO films with 2 wt% PAPB exhibited a tensile strength and strain to failure (%) of 207 ± 20 MPa and 3.48 ± 0.55 %, respectively, which is 70 % and 203 % higher than the values of pure GO films. The outstanding mechanical properties were achieved as the result of the unique intimate wrinkle-interlocking of the GO sheets formed by GFT technique, which is held together by synergistic interactions of hydrogen bonding and π-π interaction. In light of the success of the GFT process reported by Zhang et al.[189] in the formation of PAPB/GO, Tan et al.[190] researched the various polymers with low critical gel concentrations (CGCs) which promotes the formation of a GO hydrogel. To study this gelation phenomena, different polymer gel promoters such as CH, poly(ethylenimine) (PEI), fourth generation amine-terminated poly (amido amine)

29 dendrimer (G4NH2), poly(vinylpyrrolidone) (PVP), PVA, poly (ethylene oxide) (PEO), and sodium carboxy-methylcellulose (CMC) were used. It was observed that the polycations CH, PEI, and G4NH2, with positively charged amino groups, are more effective gel promotors than the nonionic polymers and polyanions with CGCs as low as 1 wt % relative to GO. These CGCs values are much lower than those of nonionic (PVA, PVP, PEO) and polyanionic (CMC) polymers (10−15 wt %), signifying that electrostatic attraction is more effective than hydrogen bonding for promoting GO gelation. Remarkable mechanical properties were achieved owing to their unique wrinkle-interlocked morphology. As seen in Fig 1.20, the prepared nanocomposite films were smooth and uniform and could be readily attained in large-area format and processed into desired shapes through origami and paper-cutting due to their high toughness and flexibility.

Figure 1.19: Synthetic route of PAPB/GO hydrogel nanocomposites via Gel-film transformation method.[189]

Figure 1.20: Photographs of GO films containing 2 wt % CH (top) and the films with different shapes processed by origami and paper-cutting (bottom).[190] 1.4.4.3 Ice Templating

Inspired by natural materials and their nanostructure, ice templating or freeze casting has been found to be very successful in fabricating nacre-like layered structures.[191–193] Ice templating is a technique developed by Deville et al.[194] to build homogenous, layered and porous scaffold by taking advantage of the physics of ice formation (Fig 1.21). The porous ceramic scaffolds fabricated by freeze casting is then filled with polymer and typically pressed down to obtain a brick-mortar structure. PMMA/Al2O3 nanocomposites were fabricated by Ritchie and his coworkers[195,196] by this method to develop a nacre- like structure with over 80 vol% Al2O3 content. A fracture toughness and tensile strength of 30 MPa m1/2 and 200 MPa, respectively, were achieved. The high mechanical properties were attributed to the presence of the thin polymer layer as a lubricant between the ceramic layers, thereby controlling the sliding of the load bearing ceramic phase.

Figure 1.21: Schematic representation of Ice-templating where porous ceramic scaffolds are fabricated.[191] 1.4.4.4 Freeze Drying (Lyophilization) Freeze drying is a technique where a solution is frozen and the surrounding temperature is slowly dropped in order to allow the frozen water to sublime from solid state to gas state. Aerogels have popularly been fabricated by this method in order to achieve low density lightweight material by swapping out the liquid phase with gas phase (in Fig 1.22, the images of before and after freeze drying of polymer-clay aerogels are shown). Alhassan et al.[197] utilized this method to fabricate PVA/MTM aerogel nanocomposites with a clay content of 50 wt%. Alhassan et al.[198] then incorporated PVA with 30 wt% LAP clay to fabricate an aerogel with a compressive Young’s modulus value of 12 MPa.

Pojanavaraphan et al.[199] incorporated natural rubber (NR) with 66.6 wt% clay with this method and crosslinked it further to obtain NR aerogel nanocomposite. The aerogel system possessed a compressive Young’s modulus of 1.8 MPa which was 26 times higher than that of neat control. Chen et al.[200] used the freeze casting method to synthesize PVA/MTM aerogel nanocomposites with divinylsulfone (DVS) as the crosslinker. A 60 wt% clay loading was achieved with a compressive Young’s modulus of 22.9 MPa being achieved. Later Chen et al.[201] fabricated PVA/MTM aerogel nanocomposites with a clay loading of 80wt%. Gamma irradiation was applied in order to crosslink the aerogel nanocomposites and a compressive Young’s modulus of 13.18 MPa was achieved. Freeze drying was used by Donius et al.[202] to incorporate NFC with 71 wt% MTM clay to obtain an aerogel with impressive compressive Young’s modulus of 38.5 MPa. Wang et al.[203] incorporated poly(furfuryl alcohol) (PFA) with 20 wt% MTM clay to synthesize aerogel nanocomposite with flame retardant properties. Shang et al.[204] synthesized aerogels with different nanofillers like clay and metal hydroxides using the freeze-drying method. 50wt% nanofiller loading was achieved in each case with good flame-retardant properties being achieved for the ultra-low dense materials. The peak heat release rate (pHRR) of the various aerogel nanocomposites was observed to decrease by 70 % when combined with the different nanofillers, indicating a significant decrease in flammability. Polymer-polymer nanocomposites were also fabricated using the freeze-drying method. Granular amylopectin potato starch (AMP) was synthesized with a range of loading of NFC to fabricate AMP/NFC nanocomposites.[205–207] A well distributed NFC network within the AMP/NFC nanocomposite was observed to reduce the moisture uptake of the system.

Figure 1.22: (a) before freeze drying, and (b) after freeze drying. Samples from left to right contain 0wt% content polymer to 100wt% polymer.[198] 1.4.4.5 Wet Spinning Kou et al.[208] employed the method of wet spinning, where the polymer-inorganic dispersion is drawn out with a spinneret in its wet state, to form nanocomposite fibers (Fig 1.23). The mechanical strength of the synthesized PVA-coated rG fibers increased 95 % from 86 MPa to 161 MPa, when compared to pure PVA fibers. PEI and other amines are known to effectively interact with CNT via physisorption.[8,187,209–212] Due to this affinity, a solution of PEI with CNT were subjected to fiber spinning by Munoz et al.[213] to obtain fibers with impressive mechanical and electrical properties. A tensile strength of 100 MPa/g/cm3 and electrical conductivity of 2x103 S/m was reported for the nanocomposites. Shin et al.[42] introduced twisting while simultaneously spinning the wet polymer/inorganic filler solution to obtain fibers with spring-like structures. Synergistically toughened ternary nanocomposite fibers, with PVA and inorganic components, rGO, and single wall carbon nanotubes (SWNT), were developed with an exceptional tensile strength of 575 MPa and toughness of 1380 MJm−3.

Figure 1.23: Preparation of PVA-coated rG nanocomposites through wet spinning. (A) aqueous solution containing GO and PVA. (B) PVA-coated rG solution after removing the free PVA. (C) Pre-aligned liquid crystalline PVA-coated rG. (D) Wet-spinning assembly for preparation of continuous nacre-mimic fibers. [208] 1.4.4.6 Other Methods for Viscous Systems For resins, like epoxies (EP) and benzoxazines, techniques different from the ones mentioned above must be implemented due to their viscosity and thermoset nature. Zhang et al.[83,214] reported the use of a special sol-gel technique to fabricate nanocomposites with high nano-silica loading of 40 wt%. Strong particle-matrix adhesion was observed which translated into improved mechanical properties. Multilayer co-extrusion was also used as an effective method to fabricate high filler-content nanocomposite films from highly viscous materials (Fig 1.24). Wilkerson et al.[215] co-extruded Al2O3 nanoplatelets with metal compliant phase Nickel (Ni) while Gupta et al.[216] coextruded phosphate glass (Pglass) with poly(propylene-graft-maleic anhydride) (PPgMA). Lamination curing method was used by Zheng et al.[4] where Kevlar (KEV) fibers impregnated with EP resin and Al2O3 nanoplatelets were alternatively stacked and cured under pressure to obtain multi-laminated nanocomposites. For polybenzoxazine nanocomposites, melt mixing was used to fill the matrix with high SI content (30 wt%).[217] As a result of homogenous dispersion of the spherical SI nanofillers, Young’s modulus and microhardness were significantly enhanced. Another technique used to disperse high content nanoparticles in matrix was compounding, which was achieved for BN nanoparticles in silicon rubber (SR).[5]

Figure 1.24:Schematic of the coextrusion assembly. [215] A hot compression molding method, where the system is compressed and polymerized under pressure, was used to fabricate polymer-polymer nanocomposite with NFC nanofiber from kraft pulp dispersed in resin matrix.[218] A high strength of over 300 MPa was achieved for nanocomposites with over 70 wt% NFC nanofibers. Nakagaito et al.[219] incorporated 80 wt% alkali-treated NFC with resin matrix and achieved a strain to failure 2 times higher than that of untreated nanocomposites. The improvement in mechanical properties was attributed to the high ductility brought about by treatment of the NFC nanofibers, which transforms the nanofibers to a more amorphous structure. The high fiber content was also found to have an effect on the thermal properties of the nanocomposites where the thermal expansion values decreased rapidly when fiber content was higher than 60 wt%.[220] Nakagaito et al.[221] also achieved the incorporation of 90 wt% hydrophilic NFC nanofibers in hydrophobic poly(lactic acid) (PLA) which were compression molded to achieve fabrication of nanocomposite films. The mechanical properties were found to increase linearly with NFC content with a maximum tensile strength of 180 MPa which was achieved for 90 wt% NFC content. Similar process was followed by Henriksson et al.[222] where NFC fibers were immersed in 13 wt% water soluble melamine formaldehyde (MF) resin and then polymerized under pressure to give semitransparent films. Later, however, it was observed that fiber-resin nanocomposites can be fabricated with no pressure.[223] The NFC nanofibers were simply immersed in the acrylate resin and polymerized to obtain transparent nanocomposites. Okahisa et al.[224] used the same method and fabricated NFC/acrylic resin nanocomposites with a fiber content of 40 wt%. Chitin nanofibers (CNF) was immersed in acrylate resin and polymerized to give

35 nanocomposites with 25 wt% fiber content.[225] The Young’s modulus of CNF reinforced resin films was increased 115-fold and the tensile strength 3-fold when compared with neat acrylic resins, to obtain values of 3.46 GPa and 29.9 MPa, respectively. 1.4.5 Fabrication of Single Component Nanocomposite Systems Single component nanocomposites are systems consisting of polymer-grafted nanoparticles (PGN) (Fig 1.25(a)) which settle into a self-suspended dispersed structure without the need of an external polymer matrix. It has the advantage of being able to maintain inter-particle distance constant. The definition of single component nanocomposite is limited by the following criteria.[226] (i) The nanoparticle core has nanoscale dimensions; (ii) the nanoparticle core is rigid (inorganic or organic) and that there is a well-defined interface between the core and corona; (iii) the corona consists of un-crosslinked organic molecules attached strongly (covalent or ionic) at a single point of the nanoparticle core; (iv) the whole attached organic fraction within the system is strongly associated to the core (i.e. no separate solvent or matrix). These criteria strongly distinguish single component systems from traditional polymer-inorganic nanofiller nanocomposite systems (Fig 1.25(b)).

Figure 1.25: (a) Structure of polymer-grafted-nanoparticles at different graft densities. (b) The different structures represented as matrix is absent for the HNP assemblies and added for the swollen and dispersed phase. [226] Solution casting was the method most popularly used to prepare the single component films with high nanofiller content. Williams et al.[227] fabricated self-healable superlattice nanocomposites through self-assembly of PACA-grafted-SI nanoparticles. An inorganic loading of 49.9 wt% was achieved with appreciable mechanical strength (82.1 MPa) and toughness (5.3 MJm−3) being reported. The high filler content of 49.9 wt% was observed to be influenced by the molecular weight of the polymer. The filling fraction of

36 the nanofiller was observed to increase with low molecular weight of polymer. Grafting density also had a similar effect with lower grafting density yielding high filler-content single component nanocomposites. Tchoul et al.[228] grafted PS on the surface of TiO2 nanoparticles and utilized solution casting method to achieve fabrication of the films. The inorganic content was varied from 60 to 80 wt% (27 to 50 vol %) by varying the molecular weight of polystyrene (PS) grafted on the TiO2 surface. Choi et al.[229] grafted PS and PMMA polymer on the surface of SI nanoparticles via atom-transfer radical polymerization (ATRP) and fabricated single component films by solution casting. Choi et al.[230] also fabricated particle brush systems through surface-initiated atom-transfer radical polymerization (SI-ATRP) method that was able to form plastic array structures. Polymer- nanoparticle superlattice structures were fabricated by Labastide et al.[231] for photovoltaic applications. Zhang et al.[232] utilized reversible addition-fragmentation chain transfer polymerization (RAFT) along with in-situ polymerization to fabricate single components PS/MTM nanocomposites. Single component nanocomposites hold significant potential owing to its good dispersion and other properties and is a promising area for future research.[233] 1.5 Surface Modification One of the problems that still affect the incorporation of high nanoparticles content into polymer matrices is the tendency for them to agglomerate.[234,235] This leads to poorly dispersed nanocomposite structure with weak mechanical properties. In order to homogenize the system more and to achieve a highly well-dispersed nanocomposite, a possible route is modification of the components before development of nanocomposite. Surface modification has attracted strong attention owing to its ability to produce remarkable integration and an improved interface between the inorganic nanoparticles and the polymer matrix.[236–239] Due to higher nanofiller content in the nanocomposites , the nanofillers are much more likely to agglomerate to create nanocomposites with poor polymer-nanoparticle interface. In conjunction with the unique fabrication methods mentioned above, the individual components (organic or inorganic) were also surface modified to achieve homogenous dispersion of nanofillers in polymer-nanofiller nanocomposite. 1.5.1 Surface Treatment of Polymer

In order to maximize the interaction with ionically charged inorganic nanofillers, the organic components were modified before-hand in order to enhance the interfacial interaction between the polymer and nanofiller. Ho et al.[50] cationized the surface of the NFC nanofiber with trimethylammonium (TMA) so as to bring about an electrostatic interaction with the negatively-charged MTM clay, thereby creating a homogenously dispersed nanocomposite system. Yao et al.[132] modified the surface of NFC nanofiber with dopamine (DA) to obtain DA conjugated NFC which adhere to the MTM clay through the formation of catechol/metal ion chelation and hydrogen bonding. (Fig 1.26) As a result, the DA-NFC/MTM nanocomposite possessed high mechanical strength at high humidity levels owing to its enhanced interfacial interaction. Wu et al.[240] took another approach by modifying the surface of NFC nanofiber with 2,2,6,6-tetramethylpiperidine-1-oxyl radical (TEMPO) which gives rise to abundant anionically charged sodium C6-carboxyl groups on the NFC surfaces. The idea was to cause electrostatic repulsions between the anionically charged MTM nanofillers and polymers in order to achieve high level of dispersion. An excellent tensile strength of 410 MPa was achieved for a clay loading of 50 wt%. Zhang et al.[241] synthesized homo-telechelic functionalized PEO through modification with phenyl, pyrene and di-pyrene groups via esterification reactions. (Fig 1.27) This enhanced the interaction between GO nanofillers and the modified PEO and yielded a well dispersed nanocomposite system.

Figure 1.26:Illustration for the Preparation of DA-Conjugated NFC from Pulp Fibers and Subsequent Assembly of DA- NFC with MTM to Prepare Transparent Nanocomposite.[132]

Figure 1.27:Synthesis of homo-telechelic functionalized PEG precursors.[241] 1.5.2 Chemical Treatment of Inorganic Nanofillers The inorganic nanofillers were also modified to enhance dispersability and to improve the polymer-nanofiller interaction. Möller et al.[242] subjected mica nanoplatelets to a process of cation exchange with cations with low enthalpy of hydration which induces a collapse of the inter-lamellar water and produces non-hydrated, mica-like nanoplatelets. This allows for tunability of the mechanical properties of the filler leading to a stronger reinforcement within the matrix. Chen et al.[120] treated MTM clay with organic ammonium salts in order to achieve the combination of MTM clay with the polymer at the molecular level. Ma et al.[178] attached –NH2 groups onto the exfoliated VMT surface by

39 the use of a coupling agent, (-aminopropyltriethoxysilane), to improve the reactivity of VMT clays. Depending on the level of concentrations of the coupling agent, the nanocomposites self-assembled into fiber like structure (11 % coupling agent) and nacre like structure (17 % coupling agent), thereby enabling control over the size and shape of the final structure (Fig 1.28). In order to enhance the interfacial interaction between nanofiller and polymer, Tian et al.[138] self-polymerized DA on the surface of GO nanoplatelets due to its ability to react with amine groups. The PDA capped GO sheets were then incorporated with PETI polymer to yield nanocomposites with good mechanical properties. LDH nanofillers were also modified to achieve an improved interfacial interaction between polymer and nanofillers. Lin et al.[7] modified the LDH nanoplatelets with 3-(trimethoxysilyl)propyl methacrylate (TPM) by the silane coupling reaction[243]. The TPM modified nanocomposites possessed higher mechanical strength than the unmodified nanocomposite owing to the higher interfacial interaction between polymer and modified nanofiller. Lin et al.[116] again modified the surface of LDH nanoplatelets with sol-gel silica in order to induce surface roughness and further reversed the charge of the silica coated nanofiller through attachment with PEI. The modified nanocomposites were again observed to have much better mechanical properties than the unmodified nanocomposite (Fig 1.29).

Figure 1.28:SEM images of self-assemblies with two levels of MVMT. (a) Shows the fiber- like assembly and (b) the presence of multiple layers when added with MVMT containing 11% coupling agent. Adding MVMT with 17% coupling agent resulted in nacre-like structures (c and d).[178]

Figure 1.29:Tensile stress vs. strain curves for plain ethoxylated trimethylolpropane triacrylate (ETPTA) film, ETPTA–gibbsite nanocomposite, and ETPTA–PEI–SCG nanocomposite.[116] For carbon based nanocomposites, the in-situ polymerization was directed by surfactant, such as cationic surfactant, CTAB, nonionic surfactant, polyethylene glycol mono-p-nonylphenyl ether (Oπ-10), or anionic, SDBS.[185,186,188] In the case of CTAB and Oπ-10 surfactants, CNT was dispersed in a solution containing CTAB or Oπ- 10.[185,186] The surfactant molecules were adsorbed and arranged regularly on the CNT surfaces. Pyrrole was added and absorbed at the surface of CNT and wedged between the arranged surfactant molecules. This is the technique termed as adsolubilization. Once the reagent, ammonium persulfate (APS), was added, pyrrole is polymerized in situ at the surfaces of the CNT (Fig 1.30). In the case of SDBS, the surfactant molecules adhered to the surface of CNT.[188] The reagent, APS, was added to the solution and concentrated in the SDBS molecules. Once the pyrrole monomer was added, they entered the micelle and polymerized, resulting in the adsorption of the PPY polymer at the surface of MCNT. In this way, several types of surfactants were utilized to achieve good dispersion of CNT into an insoluble and infusible polymer matrix.

Figure 1.30: Schematic diagram of the mechanism of formation of PPY/CNT nanocomposites.[185] 1.5.3 Grafting Modification of Nanoparticles Another surface modification technique that is becoming an attractive option, is the grafting of organic components on the surface of nanofillers. This type of modification involves low molecular weight monomers penetrating the aggregated nanoparticles and reacting with the activated sites on the nanofiller surface. The interstitial volume is filled by these organic components and therefore the aggregated inorganic nanofillers becomes more separated.[244] Moreover, the surfaces of the nanofillers become hydrophobic thus further contributing to further separation. There are two prominent methods that have been reported in the literature to graft polymers on the surface of inorganic nanofillers. The first method is the ‘graft-to’ method which involved the grafting of already-prepared end-functionalized polymer on the active sites of the inorganic nanofiller.[245] The other method is called ’graft-from’ which involves the growth of a polymer chain from an initiator on the surface of the inorganic nanofiller.[245–248] In Fig 1.31, the schematic representations of the ‘graft to’ and ‘graft- from’ methods are shown.[249]

Figure 1.31: Schematic representations of 'graft-to' and 'graft-from' methods.[249] Comparing the ‘graft-to’ and ’graft-from’ techniques, it is observed that the ‘graft- from’ approach achieves a higher grafting density than the other.[250] The ‘graft-to’ method involves end-functionalized polymer reacting with the functional groups on the surface of the inorganic nanofillers. However, only a limited amount of polymer chains can be attached to the surface of the nanofiller. This is because the incoming polymer chains must overcome the activation barrier, which appears as soon as the earlier attached chains begin to overlap. In the case of the ‘graft-from’ method, polymerization is initiated from the surface of the nanofiller. The molecules of the monomer can easily penetrate through the already polymerized polymer chains, thereby achieving a higher grafting density. Different types of polymerization methods such as cationic[251,252], anionic[253], and ‘living’ free radical polymerizations[254–257] have been successfully used to graft polymers on the surfaces of substrates.

Various inorganic nanofillers like SI, Al2O3, TiO2, ZnO, carbon black (CB) and magnetic nanoparticles like Fe, Co, Ni and Fe3O4 were modified through grafting. Modification of surfaces of SI nanoparticles through grafting method using ATRP was first proposed by Werne et al.[258] Hayashi et al.[259] modified SI nanoparticles by grafting hyperbranched polymers possessing pendant azo-groups onto their surfaces. The modified SI nanoparticles were subsequently subjected to a radical post-graft polymerization of vinyl monomers from the azo groups of the hyperbranched polymers. Yokoyama et al.[260] grafted a biocompatible material 2-methacryloyloxyethyl phosphorylcholine (MPC) on the SI nanoparticle surface initiated by the presence of azo groups introduced beforehand on

43 the SI surface. Ranjan et al.[261] modified SI nanoparticles via RAFT polymerization of multiple monomers (styrene and methyl acrylate) on the surface of the nanoparticles.

Sidorenko et al.[262] researched the modification of TiO2 nanoparticles. In their study, they achieved the radical polymerization of styrene and methyl methacrylate onto the TiO2 nanoparticles’ surface by using adsorbed hydroperoxide initiators. Wang et al.[263] grafted PMMA polymer chains directly from the TiO2 nanoparticle surfaces in water under sunlight illumination. Shirai et al.[264] went on to investigate the grafting of vinyl polymers from the surface of polymethylsiloxane-coated titanium dioxide modified with alcoholic hydroxyl groups. ZnO nanoparticle surfaces were modified by Tang et al.[265] by grafting poly(methacrylic acid) (PMAA) chains onto the nanoparticle surface in order to create a better dispersion in an aqueous medium. PMMA chains were grafted on the surface of ZnO nanoparticles through a free radical polymerization by Hong et al.[266] in order to lessen aggregation. Liu et al.[267] achieved grafting of poly(hydroethyl acrylate) (PHEA) from the surface of ZnO nanoparticles via copper-mediated SI-ATRP technique. Shirai et al.[268] further studied the graft polymerization of vinyl polymers on the nanoparticle surfaces of SI, TiO2 and CB by using an initiation system consisting of trichloroacetyl groups and molybdenum hexacarbonyl on the nanoparticle surfaces. Bach et al.[269] grafted PMMA from magnetic material, Fe3O4, nanoparticle surfaces via thiol- lactam initiated radical polymerization (TLIRP). In the above-cited literature, the graft polymerized nanoparticles were observed to have less aggregation, thus having better dispersion in complementary solvents. Therefore, surface modification through grafting is a promising route to obtain well dispersed nanofillers in polymer matrices without aggregation. For this reason, many authors have focused their efforts to investigate surface modification of nanofillers through grafting for the sole purpose of improving interfacial interaction within a polymer-inorganic nanofiller nanocomposite. Rong et al.[270] modified Al2O3 nanoplatelets by grafting PS and PACA on its surfaces to achieve a better interfacial interaction between polymer matrix and nanofiller. Magnetic materials like Fe, Co and Ni nanoparticles were all modified by Rong et al.[271], via irradiation graft polymerization in order to tailor the interfacial interaction in magnetic metal nanoparticles filled polymer nanocomposites. Shin et al.[272] achieved modification of SI nanoparticles with poly(ethylene glycol) methacrylate (PEGMA) or

44 poly(propylene glycol) methacrylate. In this study, grafting method was found to separate the aggregated nanoparticles and create good dispersion. Through grafting, there are two systems by which homogenously dispersed nanocomposites can be achieved. The first way is the single component nanocomposite system where polymer-grafted nanoparticles are bound together to form a polymer- inorganic single component nanocomposite system without extra polymer matrix.[226] Williams et al.[227] developed single component systems through self-assembly of polyamide-grafted-SI nanoparticles. Tchoul et al.[228] grafted PS on the surface of TiO2 nanoparticles and through solution casting, achieved self-assembly of the polymer-grafted nanoparticles to obtain a single component system. Choi et al.[230] utilized SI-ATRP to synthesize particle brush systems that were able to form plastic array structures. Ohno et al.[273] synthesized a single component system through formation of a colloidal crystal from PMMA-grafted-SI nanoparticles. Zhang et al.[232] utilized RAFT polymerization along with in-situ polymerization to fabricate single components PS/MTM nanocomposites. Another option is to disperse the polymer-grafted nanoparticles into an external polymer matrix. While low content (<10 vol%) polymer-grafted nanoparticles have been successfully incorporated into external polymer[274–280], literature on successful incorporation of high content polymer-grafted nanoparticles in an external polymer matrix is limited. A successful study was reported by Tao et al.[281] who incorporated 30 wt% of poly(glycidyl methacrylate) (PGMA) brush-grafted TiO2 nanoparticles into an epoxy (EP) matrix. It was observed that at high grafting density, good dispersion within matrix was observed, while below a certain grafting density value, the nanoparticles tend to aggregate (Fig 1.32). Overall, bulk polymer matrices with high-content of polymer- grafted-nanoparticles is an area of interest that holds promise for future research.

Figure 1.32: Experimental dispersion diagram of PGMA-grafted TiO2 NPs within an epoxy matrix. The solid and dash lines represent the maximum achievable graft densities and the “brush-to-mushroom” transition, respectively.[281] 1.6 Mechanical Properties of High Nanofiller Nanocomposites High-performance nanocomposites typically contain a low volume fraction of inorganic content due to the affinity of high-volume fractions of nanofillers to agglomerate and form sites of stress concentrations, thereby making the material brittle. However, due to the rise of the above-mentioned techniques to fabricate well-dispersed and distributed nanocomposites, the mechanical properties of the fabricated high filler-content systems have seen a substantial improvement. In several studies, the mechanical behavior of the nanocomposites has been found to be influenced by several factors like different interfacial interactions, effects of crosslinking agents, inorganic filler modifications and so on; hence the large range of mechanical properties obtained. Below, we will highlight the various factors and the noteworthy mechanical properties achieved as a result. 1.6.1 Effects of Interfacial Interactions The use of nanofillers in the nanocomposites allows for an enhanced interfacial surface area available for interaction with the polymer.[282] The interfacial interaction between the polymer matrix and inorganic nanofiller is an important criterion to be controlled to determine the final mechanical properties. Various types of interfacial bonding like hydrogen bonding, covalent bonding, electrostatic interactions and so on, play an essential role in determining the mechanical behavior of nanocomposites.

1.6.1.1 Hydrogen Bonding Hydrogen bonding, the weak bonding between a proton and an electronegative atom, has been extensively used to obtain high performance nanocomposites. As observed in Table 1.1, most literature favored PVA as the preferred polymer to impart hydrogen bonding between the polymer and nanofiller due to its ability to be a hydrogen bond acceptor and donor with the oxygen groups found on the surface of most fillers. [31,78,89,123,126,149,151,179,190] This was observed by Putz et al.[126] who fabricated polymer-GO nanocomposites from polymers, such as PVA and PMMA. PVA based nanocomposites had an improved Young’s modulus value (36.4 GPa) whereas the PMMA based nanocomposites only had a minimal improvement (7.5 GPa), because PVA can act as both the hydrogen bond acceptor and donor resulting in higher abundance of hydrogen bonds while PMMA could only act as a hydrogen bond acceptor. As observed in Fig.26, the interfacial strength varies depending on whether the organic molecules are hydrogen acceptors or both acceptors and donors. Dimethyl formamide (DMF) and PMMA are observed to be only hydrogen acceptors and therefore are limited in their interaction with the GO sheets (Fig 1.33 C and E) whereas, water and PVA are both hydrogen acceptor and donors enabling to create more interaction (Fig 1.33 B and D). In another study, high mechanical properties were reported at high humidity level (75 %RH) in PVA based nanocomposites with Morits et al. reporting tensile strength and flexural modulus at 220 MPa and 25 GPa, respectively.[31] Wang et al.[37] also fabricated PVA based nanocomposites by the evaporation method to yield nacre-like structures with high mechanical strength of 219 MPa. Strong hydrogen bonding also played an important role in other nanocomposites with other polymeric components. Haraguchi et al.[179], reported a nanocomposite with poly(2-methoxyethyl acrylate) which achieved remarkable strain to failure of ~1000 % at high filler loading (23 wt%) (Fig 1.34). Hydrogen bonding also played a significant role in the structure of ternary nanocomposite of poly(N- isopropylacrylamide) (PNIPAM) and inorganic nanofillers, GO and Clay.[180] As a result, the mechanical properties of the GO/Clay/PNIPAM hydrogel was massively improved −3 with strain to failure reaching as high as 800 % and toughness 5.6 MJm . For Al2O3 based nanocomposites, Bonderer et al.[103] achieved high mechanical strength of 315 MPa owing to the strong hydrogen bonding. LDH based nanocomposites were also observed to

47 achieve high mechanical strength values ~160 MPa for high filler content of 97 wt%.[95,107] Table 1.1: Effect of Hydrogen bonding on Mechanical properties Inorganic Inorg. Organic Tensile Young’s Shear Toug Strain Ref. wt% strength modulus Modul hness to No. (vol% if (MPa) (GPa us (MJ/ Failur specified unless (GPa) m3) e (%) ) otherwis e specifie d)

LAP 30 PMEA 6.2 192MPa 1000 [179]

MTM 70 PVA 150 13 [88]

MTM 70 PVA 165 27.1 1.7 [123]

MTM 70 PVA 105 21.3 0.6 [123]

MTM 23.5 PNIPAM 1.6 43.2 7.38 [129] MTM 50 NFC 410 19 [240]

MTM 70 PVA 219 19 [37]

MTM 40 PVA 220 25 [31] (75%RH (75%R ) H)

GO 72 PMMA 148.3 7.5 3.17 [126]

GO 72 PVA 80.2 36.4 [126]

GO 80 PVA 118 11.4 [151] GO 18.6vol ALG 272.3 18.1 [141] %

GO 90 PVA 236 3.95 3.13 [190]

GO 90 PEO 236 5.06 4.18 [190]

GO 85 CMC 234 4.45 3.7 [190]

GO 85 PVP 209 2.84 2.63 [190]

48 rGO/Clay 21 PNIPAM 0.69 2.52 3.5 750 [180] GO/Clay 26 PNIPAM 0.97 9.69MP 5.6 625 [180] a

Al2O3 15vol% CH 315 10 21 [103]

LDH 97 CH 160 12.7 [95]

LDH 96.9 PVA 169.36 2.1 [107]

Figure 1.33: Schematic diagram depicting the ability of different intercalating species to contribute to gallery-spanning hydrogen bond bridges. A) Anhydrous stacked GO sheets can form a small number of hydrogen bonds between surface-bound groups when the sheets are close enough to each other. B) Water molecules, which are both hydrogen bond donors and acceptors, between stacked GO sheets create a network of many hydrogen bonds that can readily adapt to mechanical stresses. C) DMF molecules between stacked GO sheets reduce the intersheet interactions due to the limited ability of DMF to hydrogen bond (it can only accept hydrogen bonds but not donate). D) PVA chains between stacked GO sheets increase the intersheet interactions, not only due to the ability of PVA to hydrogen bond in a similar fashion to water, but also strengthen the bond network with the covalent C–C bonds in the polymer (i.e., between the hydrogen-bonding capable monomer units), creating a very stiff structure. E) PMMA chains between stacked GO sheets are similar to the DMF molecules in case C, because PMMA can only accept hydrogen bonds. [126]

Figure 1.34: Stress–strain curves for M-NCs with different clay contents (5.5–23) and chemically crosslinked polymer (M-OR3). Schematic illustrations of typical yielding behaviors with a well-defined necking point are depicted (I, IIa, II, III). [179]

1.6.1.2 Covalent Bonding In order to obtain stronger interfacial interactions between the polymer and nanofiller, hydrogen bonding was replaced by covalent crosslinks to enhance mechanical strength. As observed in Table 1.2, many nanocomposites have higher mechanical properties owing to the higher interfacial strength brought about by covalent bonding. Podsiadlo et al.[33] fabricated high clay nanofiller-content nanocomposites with 35 wt% polymer DOPA. The covalent crosslinking at the interface brought about dual mechanical properties of strength (200 MPa) and toughness (4.2 MJm−3). Kochumalayil et al.[131] achieved good mechanical strength values at high moisture levels with a tensile strength and Young’s modulus of 147.5 MPa and 30 GPa, respectively. For GO nanocomposites, polymers, such as PETI[138] and PBA,[143] were incorporated to obtain good mechanical strength in both cases. Nanocomposites with strong covalent crosslinks between polymer and high Al2O3 filler content (80 vol%), were observed to have an improved strength of 200 MPa and a fracture toughness higher than 30 MPa.m1/2.[195,196]

Table 1.2: Effect of Covalent bonding on Mechanical properties Inorganic Inorganic Organi Tensil Young’s Tough Strai Fracture Ref. wt% c e modulus ness n to Toughness No. streng (GPa) (MJ/m Fail (MPa. th 3) ure m1/2) (MPa) (%)

MTM 65 DOPA 200 4.2 [33]

MVMT 60 PA66 19 [178]

MTM 60 HEC 125 3.7 6.8 [134]

MTM 60 PEO 59 5.1 3.6 [134]

MTM 60 XG 147.5 30 (50 [131] (50 %RH) %RH)

MTM 50 HEC 129.3 [147]

GO 96.25 PALA 91.9 33.3 [122]

GO 93.5 PCDO 106.6 2.52 [137]

GO 70 PETI 179 84.8 [138]

GO 96.5 PBA 187 4.3 [143]

Al2O3 80 vol% PMM 200 30 [195] A

Al2O3 80 vol% PMM 200 30 [196] A

LDH 70 ETPT 55 0.01 [7] (GB) A

1.6.1.3 Electrostatic Bonding Ionic interactions also proved to be another strong bonding option which yielded high performance nanocomposites as seen in Table 1.3. Clay and layered double hydroxide nanofillers most often tend to be negative on the surface while positive charge exists on the rims. The anionic property of these filler surfaces have been exploited in the literature to induce an electrostatic interaction by bringing in a positively charged polymer.

[30,50,77,87,116,127] Tang et al.[30] successfully used LbL assembly to construct layers of cationic polymer and anionic clay and developed electrostatically bound nacre-like structures with appreciable modulus of 11 GPa. Podsiadlo et al.[87] replaced PDDA with a stronger polycation CH, expecting development of nanocomposites films with higher mechanical properties owing to stronger electrostatic interactions. However, poorer mechanical strength (81 MPa) than that of CH (108 MPa) was obtained, which was attributed to the rigid structure of the polymer. Walther et al.[127] replaced the monovalent − 2- 3- - anionic counterion Cl with multivalent ions SO4 and PO4 and organoanion StSO3 to explore their effect on mechanical properties. It was observed that a higher valency created a greater number of crosslinking points at the interface, thereby making nanocomposite mechanically stronger. Table 1.3: Effect of electrostatic interaction on Mechanical properties Inorganic Inorganic Organic Tensile Young’s Toughness Strain Ref. wt% strength modulus (MJ/m3) to No. (MPa) (GPa) Failur e (%)

MTM PDDA 100 11 [30]

MTM 80 CH 81 6.1 [87]

MTM 70 PDDA / 106 12.9 2.1 [127] Cl−

MTM 70 PDDA / 110 24.2 0.7 [127] SO42−

MTM 70 PDDA / 151 32.9 0.8 [127] PO43−

MTM 70 PDDA / 119 29.3 0.6 [127] StSO3−

MTM 63 NFC 76 23.1 37 [121]

SA 47 PIL 56 (20 17.2 (20 [35] (LMW) %RH) %RH)

SA 47 PIL 77 (20 21.9 (20 [35] (HMW) %RH) %RH)

LDH ETPTA/ 37 0.08 [116] (GB)/SI PEI

The positive charge on the rims of the inorganic nanoplatelets have also been exploited in the literature to obtain nanocomposite structures with strong electrostatic interactions (Table 1.4).[146,283] The microstructure of those nanocomposite films with electrostatic interaction at the positive edge rims of the clay were observed to have no intercalation as opposed to those with electrostatic interactions at the clay surfaces. High mechanical strength was achieved for these nanocomposites with Das et al.[150] reporting a nacre-like structure with mechanical strength of 251 MPa and Young’s modulus of 24.6 GPa at nanofiller loading of 40 wt%. Mäkiniemi et al.[35] further reported a nanocomposite, with poly(3,4-ethylenedioxythiophene): polystyrenesulfonate (PEDOT:PSS) attached to the rims of clay, possessing stable mechanical strength of 150 MPa at low humid conditions. Table 1.4: Effect of electrostatic interaction on positive rims of nanoplatelets on the mechanical properties Inorga Inorganic Organic Tensile Young’s Toughn Strai Ref. nic wt% strength modulus ess n to No. (MPa) (GPa) (MJ/m3 Failu ) re (%)

MTM 20 CMC 320 21.5 2.6 [15 0]

MTM 40 CMC 251 24.6 1.5 [15 0]

SA 50 PEDOT:P 150 14.5 [35] SS (20%RH) (20%RH)

1.6.1.4 Synergistic Combinations of Bonding Combination of the various individual interactions mentioned above to give synergistic interactions have been extensively studied. In fact, natural biomaterials possesses a synergy of various interactions.[284] A synergy of interactions increases cohesion between the organic and inorganic component yielding a higher performance nanocomposite (Table 1.5). Hydrogen and covalent bonding has been widely used as a synergistic interaction to enhance mechanical properties.[89,96,142,157,160,285] Duan et al.[285] fabricated a ternary nanocomposite GO/NFC/PCDO which possessed dual hydrogen and covalent interactions between the layers of the nanocomposite as observed in Fig 1.35. The synergistic interactions between the organic and inorganic components and the presence of long chain PCDO, gave rise to exceptional dual mechanical properties (tensile strength 315 MPa and toughness 9.8 MJm−3). Crosslinking agents and precursor modification have also been implemented in order to bring about enhancement of interfacial cohesion.[89,132] Remarkable mechanical strength enhancement was observed owing to the high interfacial strength induced by synergy of bonding.

Figure 1.35: Schematic illustration of fabrication procedure of ternary artificial nacre nanocomposites with dual covalent and hydrogen interactions.[285] Alternatively , electrostatic interactions combined with hydrogen bonding has been extensively utilized to improve interfacial interaction.[97,128,133,136,190] Owing to the synergistic electrostatic interaction and hydrogen bonding, Tan et al.[190] obtained high mechanical properties for various cationic polymers with CH/MTM nanocomposite achieving a strength and toughness value of 254 MPa and 6.42 MJm−3, respectively. Xiong et al.[106] modified anionic CNC polymer with positive PEI layer in order to bring about electrostatic interactions in addition to the already-present hydrogen bonding with

54 inorganic filler GO. Exceptional strength and Young’s modulus were achieved with values of 490 MPa and 59 GPa, respectively. Other synergistic interactions were also experimented to observe enhancement in mechanical properties. Hu et al.[105,139] introduced silk fibroin (SL) into GO based nanocomposites and reported high mechanical properties (tensile strength 300 MPa) owing to the synergistic weak interactions (hydrogen bonding, and polar–polar and hydrophobic– hydrophobic interactions) between the SL domains and GO nanosheets. Electrostatic interaction combined with covalent interaction has not been widely utilized except by Mamedov et al.[8] who achieved a tensile strength of 325 MPa for a LbL assembled carbon based nanocomposite. Zhang et al.[189] and Xiong et al.[106] were able to add 휋 − 휋 interaction to the already existing interfacial interactions to build exceptionally strong GO based nanocomposite films. Table 1.5: Effect of synergistic interactions on mechanical properties Inorga Inorga Orga Interfacial Tensile Young’s Toughness Strain nic nic nic Interaction strength modulus (MJ/m3) to Ref. wt% (MPa) (GPa Failure No. unless (%) otherwise specified)

MTM 70 PVA Hydrogen 150 13 [89] + Covalent

MTM 70 PVA Hydrogen 480 125 [89] + Covalent

MTM 70 PVA Hydrogen 320 60 [88] + VDW + Intramole cular ionic

MTM 70 PVA Hydrogen 250 41 [88] + VDW + Intramole cular ionic

MTM 70 PVA Hydrogen 275 50 0.9 [123] + Covalent

MTM 70 PVA Hydrogen 141 34.2 0.5 [123] + Covalent

MTM 76 CH Ionic + 76 10.7 0.97 [128] Hydrogen

MTM 76 CH Ionic + 99 10.7 2.32 [128] Hydrogen

MTM 50 NFC Ionic + 124 8.7 [133] Hydrogen

MTM 89 NFC Ionic + 30 2 [133] Hydrogen MTM 77 PDD Ionic + 67 13.5 (50% 0.58 1.24 [286] A Hydrogen (50%R RH) (50%RH) (50%R H) H) MTM 40 CM Ionic + 125 13.5 (95% 1.3 1.5 [287] C Hydrogen (95%R RH) (95%RH) (95%R H) H)

MTM ALG Ionic + 10.28 [204] Hydrogen MPa

MTM 40 ALG Ionic + 280 7.2 [135] Hydrogen

MTM 40 ALG Ionic + 170 [135] Hydrogen (100C)

MTM 50 NFC Hydrogen 357.8 12.1 (50% [132] + (50%R RH) Covalent H)

MTM 50 NFC Hydrogen 280.1 [132] + (100% Covalent RH)

MTM 45.5 CH/ Ionic + 132 7.8 2.2 [136] NFC Hydrogen

MTM PVA Hydrogen 302 22.8 5.2 [160] / + NFC Covalent GO 23.5v SL Synergist 300 145 2.2 1 [105] ol ic weak interactio ns GO 90 SL Synergist 153 13 2.6 2.8 [139] ic weak interactio ns

rGO 95.4 PDA Hydrogen 204.9 4 [157] + Covalent

GO 98 PAP Hydrogen 207 3.56 3.48 [189] B + pi-pi

rGO 96 PAP Hydrogen 382 7.5 4.31 [189] B + pi-pi

GO GG Ionic + 88.7 25.4 0.84 [140] Hydrogen +

Coordinat ion

GO 96 CH Ionic + 254 6.42 4.85 [190] Hydrogen rGO 96 Ionic + 424 8.98 5.52 [190] Hydrogen

GO 99 PEI Ionic + 253 6.42 3.91 [190] Hydrogen

GO 99 G4N Ionic + 241 5.13 4.03 [190] H2 Hydrogen rGO 94.4 CH Hydrogen 526.7 17.7 [142] + Covalent GO 59.1 CNC Ionic + 490 59 3.9 1.5 [106] Hydrogen rGO 59.1 CNC Ionic + 655 169 [106] Hydrogen +pi-pi rGO 94.7 NFC Hydrogen 314.6 9.8 [285] / + PCD Covalent O

LDH PVA Hydrogen 195 9.5 12 [96] + Covalent

LDH HEP Ionic + 23 [97] Hydrogen

ATH ALG Ionic + 8.63 MPa [204] Hydrogen

SWN 50 PEI/ Ionic + 325 [8] T PAA Covalent

1.6.2 Effect of Crosslinking Agents In the literature, various crosslinking agents have been employed to induce synergistic interactions to further strengthen the nanocomposite material (Table 1.6). These crosslinking agents cause new linkages to form between the components of the nanocomposites, reinforcing the material and enhancing mechanical properties. Glutaraldehyde, boric acid, metal ions and anionic counterions have been utilized in various literature for this purpose. Glutaraldehyde (GA) molecules are small crosslinking agents which are popularly incorporated in high-filler content nanocomposites to enhance mechanical properties. Gao et al.[288] introduced small GA molecules into the gallery regions of GO paper and successfully tailored the interlayer adhesions within the GO paper. The hydroxyl groups on the surface of the GO sheets reacted with the aldehyde groups of the GA molecules through intermolecular acetalization to induce covalent bonding. The mechanical properties of the GO paper were effectively improved and a Young’s modulus of 30 GPa was achieved. Podsiadlo et al.[89] incorporated GA molecules within PVA/MTM nanocomposite prepared by LbL technique. Along with the pre-existing hydrogen bonding in the nanocomposite due to the polymer PVA, covalent bonding was introduced which significantly improved mechanical strength with a tensile strength and Young’s modulus of 480 MPa and 125 GPa, respectively, being achieved. Later, Walther et al.[123] fabricated PVA/MTM nanocomposites with other techniques such as VAF and the doctor blading technique and introduced GA crosslinking to obtain stronger materials. There was more-or-less no change observed in GA-crosslinked nanocomposites prepared by the VAF method. However, doctor blading prepared nanocomposites with GA covalent linkages were observed to have improved mechanical strength from 105 MPa to 141 MPa and Young’s modulus from 21.3 GPa to 34.2 GPa. Graphene based nanocomposites were also reinforced through crosslinking by GA with the GA-crosslinked nanocomposites possessing a new mechanical strength of 222 MPa which is a 39 % increase from its old value.[2] Han et al. incorporated the GA molecules into PVA/LDH nanocomposite layers

59 to introduce covalent bonding. The synergistic interactions of covalent and hydrogen bonding yielded a material with high mechanical strength of 195 MPa.[96]

Figure 1.36: Formation process of covalent bonding between borate and GO nanosheets [289]

Table 1.6: Effect of glutaraldehyde and borate linkages on mechanical properties Inorga Inorganic Orga Crosslinking Tensi Youn Toughn Strai Ref. nic wt% nic agent le g’s ess n to No. stren modu (MJ/m3 Failu gth lus ) re (MPa (GPa) (%) )

GO GA 101 30.4 0.3 [288]

MTM 70 PVA GA 480 125 [89]

MTM 70 PVA GA 141 34.2 0.5 [123]

R G 50 PVA GA 222 13 [2]

LDH PVA GA 195 9.5 12 [96]

GO Borate 160 [289]

MTM 70 PVA Borate 275 50 0.9 [123]

Borate ion was another crosslinking agent used to increase mechanical properties by enhancing interfacial adhesion through introduction of covalent linkages. An et al.[289] enhanced the mechanical properties of GO paper by introducing borate linkages (Fig 1.36) which increased the mechanical strength from 130 MPa to 160 MPa. Walther et al.[123] introduced borate covalent linkages in PVA/MTM nanocomposite films prepared by the VAF method to achieve a high mechanical strength of 275 MPa and Young’s modulus of 50 MPa. Literature confirms that metal ions have been extensively used to reinforce nanocomposites to yield high mechanical properties (Table 1.7). Park et al.[290] incorporated divalent ions Mg2+ and Ca2+ while Lam et al.[291] incorporated Zn2+, to introduce covalent crosslinking between GO sheets and yielded mechanically superior GO paper. Liang et al.[135] utilized Ca2+ metal ions to enhance the mechanical properties of clay based nanocomposites with ALG polymer. The resultant nanocomposites possessed both high strength and toughness of 280 MPa and 7.2 MJm−3, respectively. Ca2+ was also incorporated by Shang et al.[204] in different aerogel nanocomposites with different nanofillers like MTM clay and aluminum hydroxide. Impressive improvements in modulus values were observed for the aerogels; for example, a Young’s modulus value of 10.28 MPa being achieved for the MTM based aerogel. Likewise, Zn2+ was used as a crosslinking agent for PCDO/reduced-GO nanocomposite (Fig 1.37) which yielded excellent dual mechanical properties with a strength of 439 MPa and toughness of 7.6 MJm−3.[292] Podsiadlo et al.[33] fabricated a clay based nanocomposite with L-3,4- dihydroxyphenhylalanine (DOPA) polymer and crosslinked it further by using trivalent Fe3+ions. Simultaneous improvement of strength (200 MPa) and toughness (4.2 MJm−3) were achieved due to increase in interfacial adhesion. Again, Podisadlo et al.[88] introduced intramolecular ionic bonding through incorporation of Cu2+ and Al3+ into the nanocomposite structure. Mechanical strength and Young’s modulus were improved reaching 320 MPa and 60 GPa, respectively, for Cu2+ modified nanocomposites and, 250 MPa and 41 GPa, respectively, for Al3+ modified nanocomposites. Cu2+ was also incorporated into clay-polymer nanocomposites to achieve appreciable mechanical strength (125 MPa) in high humidity conditions.[287] Chen et al.[293], introduced several different types of Mn+ ions to observe the varying effects, of using different crosslinking

61 cations, on the mechanical properties of the nanocomposites. (TiO)2+ and Al3+ were observed to particularly yield high mechanical properties owing to the zigzag nature of (TiO)2+ and the trivalent nature of Al3+. The differences in the bonding styles and different bonding energies of the Mn+ ions were also significant factors for the varying mechanical properties. Those reinforced with (TiO)2+ yielded crosslinked nanocomposites with high strength of 228 MPa and high toughness of 15.7 MJm−3. The Al3+ crosslinked nanocomposites, on the other hand, had high mechanical strength (286 MPa) but relatively lower toughness (5.4 MJm−3). Liang et al.[135] utilized Ca2+ metal ions to enhance the mechanical properties of clay based nanocomposites with ALG polymer. The resultant nanocomposites possessed both high strength and toughness of 280 MPa and 7.2 MJm−3, respectively. Ca2+ was also incorporated by Shang et al.[204] in different aerogel nanocomposites with different nanofillers like MTM clay and aluminum hydroxide (ATH). Impressive improvements in modulus values were observed for the aerogels; for example, a Young’s modulus value of 10.28 MPa being achieved for the MTM based aerogel.

2+ Figure 1.37: Fabrication process of the reduced PCDO/GO nanocomposites with 푍푛 . [292] Table 1.7: Effect of cationic metal ions on mechanical properties Inorg Inorganic Orga Metal Tensile Young’s Toughne Strain Ref. anic wt% nic ions strength modulus ss to No. (MPa) (GPa) (MJ/m3) Failure (%) GO 99 Mg2+ 80.6 21.8 0.33 [290] GO 99 Ca2+ 125 23.3 0.5 [290] GO 99.21 Zn2+ 142.24 35.4 [291]

MTM 65 DOP Fe3+ 200 4.2 [33] A

MTM 70 PVA Cu2+ 320 60 [88]

MTM 70 PVA Al 250 41 [88] MTM 40 CM Cu2+ 125 13.5 1.3 1.5 [287] C (95%RH (95%RH) (95%RH (95%R ) ) H)

MTM ALG Ca2+ 10.28MP [204] a

MTM 40 ALG Ca2+ 280 7.2 [135]

MTM 40 ALG Ca2+ 170 [135] (100C)

GO CM TiO 228.2 15.7 [293] C 2+

GO CM Al3+ 286.4 5.4 [293] C

rGO 97.2 PCD Zn2+ 439.1 7.6 [41] O

ATH ALG Ca2+ 8.63MPa [204]

Anionic counterions have also been used in the literature to bring about supramolecular ionic bondings to further strengthen the interfacial adhesion between the layers and enhance the mechanical properties (Table 1.8). Walther et al.[127] fabricated clay nanocomposites with PDDA and had paid attention to the influence of the anionic counterions present between the polymer and inorganic platelet layers. Interestingly, it was observed that the valency of the polymer ions was directly correlated to the mechanical properties of the nanocomposite. Replacing monovalent chloride ions with multivalent ions, like SO42−and PO43−, was observed to introduce higher number of ions available for ionic crosslinking, thereby improving cohesion and increasing mechanical properties (Fig 1.38). Control over supramolecular bonding within the nanocomposite system allows for tunability of mechanical properties on-demand by simply replacing the existing

63 counterions. Following suite, Martikainen et al.[286] replaced Cl− with anionic 2′- deoxyguanosine 5′-monophosphates (dGMP) to induce supramolecular hydrogen bonding with other dGMP moieties as observed in Fig 1.39. Addition of dGMP increased the modulus, tensile strength, and strain to failure by 33.0 %, 40.9 %, and 5.6 %, respectively, to 13.5 GPa, 67 MPa, and 1.24 %, at 50 %RH. Table 1.8: Effect of anionic counterions on mechanical properties Inorg Inorgani Orga Addi Tensile Young’s Toughne Strain to Refere anic c wt% nic tive strength modulus ss Failure nces (MPa) (GPa) (MJ/m3) (%)

MTM 70 PDD Cl- 106 12.9 2.1 [127] A

MTM 70 PDD SO4 110 24.2 0.7 [127] A 2-

MTM 70 PDD PO4 151 32.9 0.8 [127] A 3-

MTM 70 PDD StSO 119 29.3 0.6 [127] A 3- MTM 77 PDD dGM 67 (50 13.5 (50 0.58 (50 1.24 (50 [286] A P %RH) %RH) %RH) %RH)

Figure 1.38: Stress–strain curves obtained by tensile testing of various samples with different counterions[127]

Figure 1.39: The preparation of self-assembled nacre-mimetic structure involving hydrogen bonds: First the anionic nanoclay (MTM) platelets are coated with cationic polymer (PDDA) via adsorption. Excess polymer is removed by washing. Hydrogen bonding molecules, anionic dGMP, were introduced into structure resulting in polymer- coated clays, which were linked by hydrogen bonding due to recognition sites (D and A). [286] 1.6.3 Effect of Fabrication Methods Fabrication techniques also play a crucial role in determining the mechanical behavior of the prepared nanocomposites. The mechanical properties of the nanocomposites are largely dictated by the microstructure and the even dispersion of polymer and nanofiller in the nanocomposites. The microstructure and dispersion of individual components, in-turn, are dependent on the type of assembly technique used. As is already discussed in the fabrication methods section, cost, scalability, ease of use and quality of end product also dictate the feasibility and practicality of these techniques. The polymer, PVA, has been extensively used in the literature with different techniques to fabricate high filler nanocomposites with various types of nanofiller. For this reason, we have chosen PVA based nanocomposites as an example to explain the effect of fabrication methods on the mechanical behavior of the nanocomposites. Various fabrication techniques like layer-by-layer (LbL) assembly[88,89], vacuum-assisted filtration (VAF)[123], doctor-blading (DB)[123], freeze casting (FC)[197], evaporation[37] and evaporation combined with lamination (E+L)[31] have all successfully been employed to develop nanocomposite films of PVA and MTM clay. A high clay content of 70 wt% has been achieved in each case (except for ice templating and

SC coupled with lamination methods where 50 wt% and 40 wt% clay loading, respectively, were achieved). The PVA and MTM clay used were not modified in any way which is often used to achieve better mechanical properties. PVA/MTM nanocomposites prepared by the LbL assembly technique were observed to have comparable mechanical properties (휎=150 MPa and E=13 GPa) to that of nacre.[88,89] Although this method allows for tunability of structure and thickness of each layer of the structure, the laboriousness and the time-consuming nature of the method is not an attractive trade-off for the mechanical behavior achieved. The faster and cheaper VAF method achieved nanocomposites with similar strength to that of LbL-nanocomposites, but with higher modulus (E=27.1 GPa).[123] Although this method does not allow regulation over the platelet content as LbL does, the ability to scale-up, ease of fabrication and high mechanical properties achieved make it attractive over the LbL technique. This was also proven in the case of PDDA/MTM based nanocomposites which was observed to have comparable mechanical properties when assembled by both LbL[30] and VAF[127] techniques. The DB technique yielded PVA/MTM nanocomposites with lower mechanical strength (휎=105 MPa) but appreciable modulus (E=21.3 GPa).[123] The lower mechanical properties of the DB- prepared nanocomposites could be attributed to the softer drying conditions and weak control over structure assembly. The evaporation technique which involves slow evaporation of solvent from the mixture during self-assembly, achieved the highest mechanical strength of 219 MPa with a good Young’s modulus value of 19 GPa.[37] The higher strength of the evaporation fabricated nanocomposites can be attributed to the slow evaporation which allows ample time for the nanoplatelets to interact with the polymer, resulting in good-structured nanocomposites with alternating layers of polymer and nanoplatelets. For PVA nanocomposites with G nanoplatelets, different fabrication techniques have successfully been employed to incorporate ~35 wt% nanoplatelets of rG. The LbL technique yielded a PVA/rG nanocomposite with good mechanical strength of 160 MPa. Wet spinning was utilized by Kou et al.[208] to fabricate fibers with bioinspired layered structure of polymer and nanoplatelets. The resultant structure possessed a similar mechanical strength of 161 MPa. The performance of the nanocomposites can be tweaked and improved by further altering the nanocomposite structure. Zhang et al.[294] achieved

66 remarkable mechanical properties for PVA/rG based nanocomposite through coiling of the PVA/G fiber. Coiled spring-like PVA fibers were obtained by a combination of drawing and twisting the fiber in its wet state. A filler loading of 66 wt% was achieved for the coiled PVA/rG with a tensile strength 270 MPa and a maximum strain to failure of 330 %. The area under the stress-strain curve was calculated to obtain an exceptional toughness of 460 MJm−3. 1.6.4 Effects of Electrochemical Reduction Electrochemical reduction is a post-fabrication treatment, whereby 휋 − 휋 interaction was introduced into the interlayers by getting rid of excess unreacted moieties present on the surface of graphene (G) and graphene oxide (GO) sheets, thereby reducing the d- spacing between layers making the system more compact. As observed in Table 1.9, this treatment step brought about remarkable improvements in mechanical and electrical properties of the material owing to the 휋 − 휋 interaction and compact nature of the nanostructure. Electrochemical reduction was accomplished by either treatment with hydroiodic acid (HI) or by microstamping which reduced the unreacted moieties and consequently reduced interlayer spacing. In many cases, remarkable enhancements of mechanical behavior were observed when compared to the parent unreduced nanocomposites. Electrochemical reduction by HI is the typical method employed to reduce the interlayer spacing of nanostructure. Li et al.[151] fabricated PVA/GO nanocomposites and further reduced them with HI treatment. The treatment resulted in a change in color of the nanocomposite from brown to black with mechanical strength of the material significantly improved from 118 MPa to 188 MPa (Fig 1.40). Cheng et al.[137] used a long chain polymer PCDO in order to bring about toughness in the nanocomposite system. The PCDO/GO system was put through reduction by HI treatment to yield reduced systems with a more compact structure. The reduced PCDO/GO system was observed to have improvements with its strength increasing from 106.6 MPa to 129.6 MPa, and toughness from 2.52 MJm−3 to 3.91 MJm−3. Later, Gong et al.[292] doped the same reduced PCDO/GO system (with slightly higher inorganic content) with Zn2+ and obtained a tensile strength of 439 MPa and a toughness of 7.6 MJm−3. Cui et al.[157] followed suite by obtaining reduced GO nanocomposites with polymer polydopamine (PDA) with good

67 mechanical strength of 205 MPa. The gel fabrication technique was utilized by Zhang et al.[189] and Tan et al.[190] to fabricate hydrogel based nanocomposites with polymers PAPB and CH. The resultant as-prepared nanocomposites reported excellent mechanical properties (as seen in the table) owing to the roughness of the microstructure. Further reduction treatment with HI acid resulted in exceptional improvements of mechanical properties. The reduced PAPB/GO nanocomposite possessed a tensile strength of 382 MPa (before treatment: 207 MPa) and toughness 7.5 MJm−3 (before treatment: 3.56 MJm−3), while the reduced CH/GO nanocomposite was reported to have a tensile strength of 424 MPa (before treatment: 254 MPa) and toughness of 8.98 MJm−3 (before treatment: 6.42 MJm−3). The VAF method and HI reduction treatment was utilized by Wan et al.[142] to fabricate reduced CH/GO nanocomposite with high mechanical properties. An exceptional tensile strength value of 527 MPa and toughness value of 17.7 MJm−3 were reported. Xiong et al.[295] fabricated GO nanocomposites with polymer CNC with excellent mechanical strength. The strength was further improved after electrochemical reduction with all-time high values of tensile strength and modulus of 655 MPa (before treatment: 490 MPa) and 169 MJm−3 (before treatment: 59 MJm−3), respectively, being reported. Ternary nanocomposite systems were also subjected to electrochemical reduction and achieved substantial mechanical properties. Shin et al. [296] synthesized a reduced ternary nanocomposite system with inorganic nanofillers as GO and SWNT and PVA as the organic component. An incredible toughness value of 1380 MJm−3 was reported. Later, Gong et al.[41] synthesized a reduced ternary nanocomposite system with inorganic nanofillers as GO and double walled carbon nanotubes (DWNT) and PCDO s the organic component. Post-reduction treatment, the mechanical properties were improved with a tensile strength of 374 MPa (before treatment: 238 MPa) and a toughness of 9.2 MJm−3 −3 (before treatment 4.1 MJm ) being reported. Molybdenum disulphide (MoS2) and Al2O3 were also used in conjunction with GO nanofillers to fabricate nanocomposites with high mechanical properties. [6,152] The resultant nanocomposites were electrochemically reduced with HI acid, with the reduced GO/MoS2 nanocomposite possessing a tensile strength and toughness of 235 MPa and 6.9 MJm−3, respectively, while the reduced

GO/Al2O3 nanocomposite recorded a much higher tensile strength and toughness of 587 MPa and 12.1 MJm−3, respectively. Duan et al.[285] synthesized a reduced GO

68 nanocomposite with 2 polymers, NFC and PCDO, which possessed impressive mechanical properties (strength 315 MPa and toughness 9.8 MJm−3) owing to the reduced compact microstructure and synergistic interaction between the inorganic and organic components. Electrochemical reduction by “microstamping” was another method introduced by Hu et al.[139] to reduced selected regions without compromising the mechanical and environmental stability of the GO based nanocomposites. This method involved “microstamping” with a pre-shaped aluminum foil in direct physical contact with the surface of the GO nanocomposite papers and thereby, locally reducing selected regions of the nanocomposite paper (Fig 1.41). Silk fibroin (SL) was the polymer incorporated in the GO nanocomposite and post-reduction treatment by “microstamping” resulted in much improved mechanical strength from 153 MPa to 300 MPa.

Figure 1.40: (a) Schematic illustration of fabrication of PVA/GO nanocomposite films followed by electrochemical reduction. (b) Color change of reduced nanocomposite films. [151]

Figure 1.41: Schematic illustration of Microstamping.[139]

Table 1.9: Effect of electrochemical reduction on the mechanical properties of high filler- content nanocomposites Inorganic Inorganic Organic Tensile Young’s Toughness Strain to Ref.No. wt% strength modulus (MJ/m3) Failure (MPa) (GPa) (%)

GO 80 PVA 118 11.4 [151]

rGO 80 PVA 188.9 10.4 2.67 [151] GO 93.5 PCDO 106.6 2.52 [137]

rGO 93.5 PCDO 129.6 3.91 [137] rGO 97.8 PCDO 439.1 7.6 [292] GO 90 SL 153 13 2.6 2.8 [139]

rGO 90 SL 300 26 2.8 [139]

rGO 95.4 PDA 204.9 4 [157]

GO 98 PAPB 207 3.56 3.48 [189]

rGO 96 PAPB 382 7.5 4.31 [189]

GO 96 CH 254 6.42 4.85 [190]

rGO 96 424 8.98 5.52 [190]

rGO 94.4 CH 526.7 17.7 [142] GO 59.1 CNC 490 59 3.9 1.5 [295]

rGO 59.1 CNC 655 169 [295]

rGO/ 30 (1:1) PVA 575 1380 350 [296] SWNT

rGO/ 90 TPU 235.3 6.9 [6]

MoS2

GO/ 95 PCDO 238.2 4.1 [41] DWNT

rGO/ PCDO 374.1 9.2 [41] DWNT

rGO 94.7 NFC/ 314.6 9.8 [285] PCDO

rGO/ CMC 586.6 12.1 [152]

Al2O3

1.7 Electrical Properties of High Filler-content Nanocomposites Enhancement of electrical properties along with mechanical properties, has been the main objective in several literature. Obtaining mechanically durable nanocomposite films with high and industry-viable electrical properties is a challenge due to the insulating nature of many polymers and even some inorganic fillers. However, inorganic nanofillers like G, GO and CNT have successfully been incorporated together with polymers to yield mechanically strong nanocomposite with appreciable electrical properties (Table 1.10). Nanocomposites with insulative inorganic nanofillers like clay can only be conductive if a conjugated polymer provides conductive pathways within the nanostructure of the nanocomposite film. Mäkiniemi et al.[35] utilized this concept by using conductive polymer, poly(3,4-ethylenedioxythiophene), polystyrenesulfonate (PEDOT:PSS), or poly(ionic liquid) (PIL) with 50 wt% SA clay to achieve a decent electrical conductivity of 580 S/m (Fig 1.42).

Figure 1.42: Preparation of conductive nanocomposites with conductive polymer PEDOT:PSS or PIL. The bottom photograph shows a LED connected via conducting PEDOT:PSS/SA nacre-mimetics and powered at 3.5 V [35] Over the years, G has been renowned for its exceptional electrical properties.[64,297,298] However, GO nanosheets, the water soluble derivative of G , is typically insulating due to the lack of conjugated areas which are usually destroyed in the process of exfoliation from graphite by chemical treatment.[299] GO sheets possess electrical conductivity on the order of ~10−5 S/m.[300] Electrochemical reduction, mentioned above, is a method that not only eliminates the electrically insulating oxygen- containing groups and atomic-scale lattice defects, but also retrieves the lost conjugated network of the graphitic lattice.[301] This reduction step restored the lost conjugated network and thereby achieves an electrical conductivity of ~5.00 x 103 S/m for rGO.[75] Incorporating GO nanofillers with insulative polymers would consequently decrease electrical properties of the resultant nanocomposites. Electrochemical reduction method was found to be incredibly effective in improving the electrical properties of various polymer-GO based nanocomposites. Li et al.[151] incorporated PVA with 80 wt% GO nanofillers and subjected the nanocomposite to electrochemical reduction and achieved a conductivity value of 5.27 x 103 S/m, comparable to that of rGO. Xiong et al.[106] also achieved a similar electrical conductivity value of 5.00 x 103 S/m with polymer CNC and 59.1 wt% GO sheets post-reduction treatment. In cases of rGO based nanocomposite films with polymers SL[139] and PDA[157], the electrical conductivity was improved post reduction but still less than that of rGO films. PNIPAM polymer was incorporated with clay and reduced GO to fabricate nanocomposite hydrogels with an electrical conductivity

72 of 0.37 S/m.[180] Although this value is 5 orders less than that of rGO, it is comparable to the electrical conductivity of G hydrogel (0.49 S/m) synthesized by the hydrothermal method.[302] Cheng et al.[137] fabricated PCDO/GO nanocomposites and achieved an electrical conductivity value of 2.32 x 104 S/m. The significantly high electrical conductivity value could be attributed to the thoroughly reduced GO sheets and the π − π conjugated backbones of the crosslinked polymer PCDO. Wan et al.[142] and Wu et al.[143] followed suite and fabricated GO based nanocomposites with polymers CH and PBA, respectively. In both cases, high electrical properties were obtained by post electrochemical reduction with a value of 1.55 x 104 S/m reported for reduced CH/GO and 1.19 x 104 S/m for reduced PBA/GO. The highest electrical conductivity of 3.37 x 104 S/m was reported by Zhang et al.[189] for reduced PAPB/GO nanocomposites with nanofiller content of 96 wt%. The low content of PAPB, although insulative, had a negligible effect on the reduced nanocomposite film. Carbon nanotubes, single-walled carbon nanotubes (SWNT) and multi-walled carbon nanotubes (MWCNT), are known for their electrical properties and the host of applications it can be used for.[303] These nano-sized fillers have also been extensively used in polymer-carbon nanocomposites to enhance mechanical and electrical properties.[304] Depending on the processing technique, nanofiller content and polymer used, the electrical properties varied. Curran et al.[170] used the polymer, PmPV, with 35 wt% MWCNT to achieve eight orders of magnitude improvement in its electrical conductivity (0.01 S/m). This was further improved by Coleman et al.[171,172] to achieve an electrical conductivity of 3 S/m. Shaffer et al.[165] incorporated PVA polymer with 60 wt% CNT. However, an electrical conductivity of only 100 S/m was obtained for the high filler loading, which is attributed to the adsorbed layer of polymer on the surface of nanotubes, which reduces the quality of electrical contacts between the nanotubes. PANI/CNT nanocomposites were fabricated by Sainz et al.[183] and Long et al.[186] and achieved an electrical conductivity of ~100 S/m owing to the relatively high conductivity of PANI. Karim et al.[184] fabricated carbon based nanocomposites with polymer, PTH, and achieved seven orders of magnitude improvement in electrical conductivity (41 S/m). Poly(ethyl methacrylate) (PEMA), P3HT and P3OT are other polymers which have been incorporated with CNT with all attaining an electrical conductivity of < 1 S/m.

[169,173,174,182] Munoz et al.[213] fabricated PEI/SWNT nanocomposites through fiber spinning method and achieved a high electrical conductivity of 2.00 x 103 S/m. Ten vol% of MWCNT was incorporated with TPU by Koerner et al.[167] and reported a good electrical conductivity of 1.00 x 103 S/m. The conjugated polymer PPY was used in the literature in the fabrication of carbon based nanocomposites and achieved high electrical conductivity.[181,185] Fan et al.[181] fabricated PPY/CNT through in-situ polymerization and achieved an electrical conductivity of 1.60 x 103 S/m. Later, Zhang et al.[185] achieved better interfacial interaction between PPY polymer and CNT nanofillers by utilizing cationic or ionic surfactants in an in-situ chemical oxidative polymerization and obtained a remarkable electrical conductivity of 2.30 x 103 S/m. Yu et al.[188] utilized an anionic surfactant in an in-situ inverse microemulsion polymerization, but achieved a much a lower value of 40 S/m. Table 1.10: Electrical properties of high-filler content nanocomposites Electrical Inorganic Inorganic Organic Conductivity References wt% (S/m) CLAY SA 50 PEDOT:PSS 5.80 x 102 [35] GRAPHENE

OXIDE rGO 80 PVA 5.27 x 103 [151] rGO 93.5 PCDO 2.32 x 104 [137] rGO 97.5 SL 1.35 x 103 [139] rGO 95.4 PDA 1.85 x 103 [157] rGO 96 PAPB 3.37 x 104 [189] rGO 94.4 CH 1.55 x 104 [142] rGO/Clay 21 PNIPAM 0.37 [180] rGO 59.1 CNC 5.0 x 103 [106] rGO 96.5 PBA 1.19 x 104 [143] 3 rGO/ MoS2 90 TPU 4.64 x 10 [6]

rGO/ 95 PCDO DWNT 3.94 x 104 [41] rGO 94.7 NFC/PCDO 1.63 x 104 [285] CARBON

NANOTUBES MWCNT 35 PmPV 0.01 [170] CNT 36 PmPV 3 [172] CNT 50 PPY 1.6 x 103 [181] CNT 36 PmPV 3 [171] CNT 60 PVA 1.0 x 102 [165] SWNT 23 PEMA 0.001 [182] CNT 25 PPY 2.3 x 103 [185] MWCNT 24.8 PANI 1.27 x 102 [186] MWCNT 50 PANI 1.00 x 102 [183] SWNT 75 PEI 2.0 x 103 [213] MWCNT 10vol% TPU 1.0 x 103 [167] Purified 30 PPY MWCNT 40 [188] Purified 30 P3OT SWNT 0.05 [169] SWNT 50 PTH 41 [184] MWCNT 30 P3HT 0.5 [174] SWNT 30 P3HT 1 x 10−4 [173]

1.8 Gas barrier Properties of High Filler-content Nanocomposites Among the various high filler-content nanocomposites, only clay-based nanocomposites were reported to have high gas-barrier properties. The oxygen permeability (OP) was measured at different humidity levels and remarkably low values were found even at high humidity conditions. Inorganic clay papers and most polymeric films tend to be hydrophilic and lose their structural integrity at humidity levels due to swelling, deteriorating their gas barrier properties. Combining the polymer and clay to form

75 polymer-clay nanocomposites leads to tightly crosslinked layered structures which creates a tortuous diffusion path for the gas molecules to move through.[1] Through fine control of number of layers, it was observed that the more the number of layers, the more tortuous the path for gas molecules to travel, thereby further reducing OP .[90] Furthermore, higher the concentration, lower the OP even at humidity value.[125] This was attributed to nacre- like layered structure formed at high concentration of inorganic nanoplatelets, thereby forming a tortuous path for the flow of gas molecules.

Table 1.11: Gas barrier properties of high filler-content nanocomposites for different humidity levels

Oxygen Inorganic Relative Inorganic Organic permeability Ref. No. wt% Humidit (cm 3mm m−2day−1atm−1) y (%) SA 80 PAA 0 7.40 x 10-4 [146] 50 - 80 - 95 - MTM 70 PVA 0 - [123] 50 - 80 3.25 x 10-1 95 - PEI/PA 0 MTM <5.00 x 10-3 A [90] 50 - 80 - 95 9.30 x 10-2 MTM 50 NFC 0 1.00 x 10-3 [133] 50 - 80 - 95 3.5 MTM 50 NFC 0 <5.00 x 10-3 [240] 50 2.00 x 10-2 80 - 95 - VMT 20 NFC 0 - [155] 50 7.00 x 10-3 80 1.50 x 10-1 95 -

MTM 20 XG 0 0 [305] 50 - 80 1.44 x 10-1 95 - HE 50 TPU 0 <5.00 x 10-3 [124] 50 <5.00 x 10-3 80 - 95 - MTM 60 HEC 0 - [134] 50 1.40 x 10-1 80 2.00 x 10-1 95 - MTM 60 PEO 0 - [134] 50 6.00 x 10-2 80 2.90 x 10-1 95 - MTM 60 XG 0 - [131] 50 - 80 6.50 x 10-2 95 - HE 60 SPD 0 - [156] 50 3.00 x 10-3 80 - 95 - MTM 50 NFC 0 3.90 x 10-2 [132] 50 - 80 - 95 5.33 x 10-1 MTM 45.5 CH/NFC 0 0 [136] 50 - 80 - 95 2.5

As observed in Table 1.11, there are several clay-based nanocomposites with very low OP values owing to the high filler-content and consequent layered structure of the nanocomposites. Ebina et al.[146] fabricated PAA/SA nanocomposites with a high filler content of 80wt% and achieved an OP value at 7.40 x 10-4 cm 3mm m−2day−1atm−1 in dry conditions. Walther et al.[123] achieved a low OP value of 0.325 cm 3mm m−2day−1atm−1 at 80 %RH for a PVA/MTM nanocomposite. This was impressive, as MTM clay and PVA are known to be hydrophilic and swell at high humidity.

However, tight bonds formed between polymer and clay during fabrication allows for high gas barrier properties for the nanocomposites. Yao et al.[132] achieved low OP values at high humidity owing to the neatly adhered layered nacre-like structure of high-content nanofillers which was brought about by the additive DA. Kunz et al.[124] reported non- detectable OP values in dry conditions as well as at 50 %RH. Priolo et al.[90] fabricated a ternary nanocomposite with polymers PEI and PAA and achieved a low OP value of 9.30 x 10-2 cm 3mm m−2day−1atm−1 at 95 %RH. Liu et al.[136] also fabricated a ternary nanocomposite and achieved appreciable low OP values at high humidity. 1.9 Flame Retardancy Properties of High Filler-content Nanocomposites In addition to the gas barrier properties, high filler-content nanocomposites were also reported to show flame retardancy properties. Despite the absence of halogens, nitrogen or phosphorus groups, the nanocomposites possessed significant flame shielding properties owing to the presence of high inorganic content which allowed the ability to extinguish once the organic content was burned off. The high loading of interlocked inorganic nanofillers, as understood from previous section, results in a nacre-like brick-mortar structure.[37] Carosio et al.[91] also observed that, higher the clay loading, the lower the burning time. This was attributed to the layered and oriented structure of inorganic nanoplatelets which provides a tortuous path for oxygen and volatiles to pass through, thus making it extremely difficult for thorough burning to take place as seen in Fig 1.43. This was observed by Sehaqui et al.[134] for clay based nanocomposites where a delayed burning of matrix was observed owing to the high clay content of 75 wt%. Much of the literature also reported a self-extinguishing effect once flame source was removed.[128,133,135,150] The high clay content within the nanocomposite creates a protective layer over the polymer matrix thus impeding the burning effect. Integrity of the nanocomposite structure was also preserved which can be attributed to the high content of interlocked nanoclays and their condensation via silanol groups which in turn strengthened the inorganic framework. Self-extinguishing, delayed burning time, and preservation of structural integrity of the nanocomposite films are attractive attributes making them a potential option for coating applications such as painting.[123] A thin layer of coating would go a long way in protecting the underlying flammable surface from burning and thereby give ample time for

78 escape in case of an emergency. Carosio et al.[91] demonstrated this effect by fabricating a nanocomposite coating through the LbL assembly of NFC and 30 % MTM clay. This coating was applied to the surface of glass fiber/epoxy (GF/EP) composites. A significant improvement was observed for total time to ignition (TTI) which was increased from 1 minute 8 seconds to ~5 minutes. As observed in Fig 1.44, the total heat release rate (THR) and smoke production rate (SPR) was significantly decreased, thus making this nanocomposite film a valuable coating for the GF/EP composite. Carosio et al.[306] again demonstrated the value of a transparent fire retardant nanocomposite coating for protection of easily flammable materials like wood. 50 wt% of MTM clay was incorporated with NFC through VAF method to yield a transparent, flame retardant and mechanically durable nanocomposite film. The film was then coated on the surface of the wood surface and subjected to cone calorimetry testing which yielded favorable results. A TTI value of ~6 minutes was obtained for the coated material which was a significant improvement from the 1 minute 20 seconds TTI value for uncoated material. A 33 % reduction of THR was also observed for the coated material making this nanocomposite film an attractive flame- retardant coating (Fig 1.45).

Figure 1.43: Residues collected after vertical flammability test of NFC/ Clay nanocomposites. [91]

Figure 1.44: Heat release rate (HRR) and smoke production rate (SPR) plots of uncoated and clay nanopaper (NFC/30% MTM) coated GF/EP composites.[91]

Figure 1.45: Temperature profile during cone calorimetry, measured 2 mm under the surface of Coated (clay nanopaper coated wood) and Wood (uncoated wood) samples[306] High nanofiller-content aerogel composites also proved to be good flame-retardant materials. Chen et al.[201] fabricated aerogel nanocomposites with 80 wt% clay nanofillers and PVA polymer through freeze drying. The low flammability nature of clay nanofillers and the tortuous path of the layered structure of the aerogel coupled with the high mechanical performance of these nanocomposites make them good flame-retardant

80 materials. Shang et al.[204] coupled low flammability polymer, ALG, with 50 wt% nanofillers such as MTM clay, layered silicates and metal hydroxides. These inorganic nanofillers provided impressive flame suppressive capability with the metal hydroxide nanofillers further being able to dehydrate and reduce the temperature and density of the combustible gas. Wang et al.[203] fabricated an aerogel nanocomposite wit polymer PFA and clay nanofillers through in-situ polymerization and freeze drying. Uncured and vacuum-cured aerogels were observed to possess self-extinguishing property while the air- cured aerogel did not combust with a noticeable flame. The fire resistance can be enhanced even more by introducing fire resistant nitrogen, phosphorus and chlorine groups into the nanocomposite. Walther et al.[127] investigated this effect by coupling the high nano-filler nanocomposite films with fire- 2 - - resistant counterions such as SO4 -, StSO3 , PO43- and Cl . The nanocomposite films were observed to burn with less flames in the order of doped counterions: SO42− ≈ StSO3−>PO43−>Cl−. The nanocomposite films were observed to have self-extinguishing property while maintaining the integrity of its structure after burning. 1.10 Conclusions Papers on high nanofiller content nanocomposites have been reviewed. Their various preparation methods have been summarized along with a quantitative compilation of the different properties. Due to the anisotropic nature of the packed fillers, increased tortuous path, high content of nanoconfined polymer matrix, many unique properties have been observed. It has been observed that the strong interfacial interaction between the nanofiller and polymer matrix coupled with effective processing methods are essential to successfully manufacture high nanofiller content nanocomposites. Although significant strides have been made in achieving improved properties, there is still a lot of scope for further study of the interfacial interaction between nanofiller and matrix. The mechanical behavior of nacre-like nanocomposites is one such area that requires special focus. So far, there have been very few successful fabrication techniques which yield dual improvements in tensile strength and toughness. Therefore, much investigation is required in broadening the range of novel fabrication methods available to synthesize nanocomposites with simultaneous improvement of tensile strength and toughness. There is also a great need for non- halogenated flame-retardant material owing to their toxic nature. High filler-content

81 nanocomposites are a promising development in this area, and its potential can be tapped into with additional research. Gas barrier properties and electrical properties are other areas that can also be further optimized with considerable research. Surface modification of precursors is another young field of research interest which can possibly achieve homogenous dispersion of high-content of nanofillers in polymer matrix. As mentioned earlier, bulk polymer matrices with high-content of polymer-grafted nanoparticles is an area that holds promise for future research. All in all, this emerging field of nanocomposites is now being gradually recognized. The excellent properties and effectiveness of this approach over the traditional low nanofiller content nanocomposites will open unique fields of applications.

Chapter 2.

Application of Very High Nanofiller-content Laponite/DNA Nanocomposite coating on Polyurethane Foam through Single-Dip Fabrication

Abstract

Very high nanofiller content nanocomposite coating, comprising of low content of deoxy-ribonucleic acid (DNA) as the polymer and high content of laponite (LAP) nanofillers as the inorganic filler, has been applied on polyurethane (PU) foam surface to enhance its flame-retardant behavior. The ultra-thin coating is applied via single-dip method. The successful fabrication of nanocomposite-film-coated PU foams is confirmed by different characterization techniques like Fourier transform infrared spectroscopy (FT- IR), scanning electron microscopy (SEM) and Energy-dispersive X-ray spectroscopy (EDX). Thermogravimetric analysis (TGA) showed high char yield formation of about 6.20% for PU foams coated with 90wt% laponite and 10wt% DNA. Flammability tests showed char formation at high laponite content while also maintaining structural integrity. Compression testing also showed change in mechanical behavior for the coated PU foams.

2.1 Introduction

In the year 2017, National Fire Protection Association (NFPA) reported over 3400 fire-related civilian deaths among which home fires caused over 2630 deaths comprising 77% in the United States of America.[307] Polyurethane foams (PU) is a soft, flexible material commonly used in the manufacture of mattresses, upholstered furniture and other indoor comfort-purposed objects. However, due to the high aliphatic content, large surface area, and good air permeability, PU is extremely flammable characterized by susceptibility to ignition, rapid fire spread, and large amount of heat release during combustion.[308] Toxic smoke and gases like CO and HCN are also released during burning making it more hazardous for product users.[309] Therefore, there is huge research interest to impart flame retardance to PU foams. Most conventional techniques involve addition of halogenated additives which are known for their flame-retardant benefits. However, the long-term toxic effects of these additives on human health and environment has outweighed their fire safety benefits.[310] Application of other types of flame retardant additives, such as phosphorous and nitrogen containing compounds, are often confronted by sacrificing physical properties.[311,312] Therefore, there is strong interest in the industry for a non-toxic and environment-friendly approach that deposits a thin layer of coating on the surface of PU foam, thereby preserving its flame retardant characteristics without compromising its physical properties.

Recently, high nanofiller-content nanocomposites have come to be known for their ability to impart flame retardant properties.[313] The combination of high content of inorganic fillers with low content of organic component creates a nacre-like brick-mortar structure. Clay nanofillers seemed to be the popular choice as the inorganic filler due to their thermal stability and ability to stack owing to its nanoplatelet-like structure. Carosio et al. observed that with higher loading of nanoclay fillers, the burning time decreased.[91] This was attributed to the layered and oriented structure of inorganic nanoplatelets which provides a tortuous path for oxygen and volatiles to pass through, thus making it extremely difficult for thorough burning to take place. This was also observed by Sehaqui et al.[134] for clay-based nanocomposites where a delayed burning of matrix was observed owing to the high clay content of 75 wt%. Many literature also reported a self-extinguishing effect

84 once flame source was removed.[128,133,135,150] The high clay content within the nanocomposite creates a protective layer over the polymer matrix thus impeding the burning effect. Integrity of the nanocomposite structure was also observed to be preserved which can be attributed to the high content of interlocked nanoclays and their condensation via silanol groups which in turn strengthens the inorganic framework.

A technique that has been popularly used to achieve fabrication of highly ordered, nacre-like nanocomposite films is layer by layer deposition (LbL).[313] This technique was first developed by Decher et al.[86] to alternately assemble oppositely charged polyelectrolytes on top of each other This method was taken and employed for the fabrication of nanocomposites with alternating layer of inorganic nanofillers and polymer. It was used by Kim et al. to assemble layers of carbon nanofibers and polyacrylic acid to make a coating that had a nanofiller content of over 50 wt%.[314] This coating was applied to PU foam which resulted in the reduction of the peak heat release rate (PHHR) by 40% and simultaneously also eliminated melt dripping. The same group later utilized LbL method to explore clay nanoplatelets as a potential nanofiller and found that higher concentration of clay nanofillers caused a dramatic reduction in flammability. [315,316] Laufer et al. also used this method to obtain coated PU foams with PHHR reduced by 52%.[317] The LbL method has continually drawn scientific and industrial attention due to its environment-friendly processing that leads to well-oriented nacre-like structures that promotes flame retardant behavior. However, the procedure is extremely laborious and time consuming and therefore, unfit for large scale application. A simpler procedure is required that is faster and yet preserves the strong layered morphology that promotes flame retardant behavior.

Few efforts have been devoted to discover rapid, scalable and environment-friendly processing that also preserves the coating’s strong orientational nature like nacre. Carosio et al.[306] utilized vacuum assisted filtration to prepare nanocomposite films with 50wt% clay nanofillers and nano-fibrillated cellulose (NFC). This film was then coated through hot-pressing onto the surface of a wooden substrate. A 33% reduction in the THR was observed for the coated wooden surface. Freeze drying was employed to prepare an alginate/clay aerogel coating which imparted flame characteristics of the coated PU.[318]

The PHRR was reduced from 323 KW/m2 to 20 KW/m2, but the specific modulus was simultaneously increased by 2.24 times. Recently a novel process called ‘single-dip coating’ has been reported that is rapid, scalable and environment-friendly. This is a method where all the components of the coating are mixed together to form a suspension, in which the PU foam is then dip coated. Davis et al. successfully utilized this ‘one-pot’ coating to coat PU foam which reduced the PHHR by 75% with a coating uptake of 155%.[319] Most recently, Liu et al. prepared a very high nanofiller content Laponite/Casein coating on PU foam using the single dip coating process and achieved reduction of average heat release rate (aHRR) by 30.9%.[320] In this paper, we will be utilizing the single-dip coating approach owing to facile and scalable processing capability.

Double helix deoxy-ribonucleic acid (DNA) is a material that has come to be known for its non-toxic, intumescent flame retardant ability.[321] It is composed of two long chain polymers of nitrogen-containing bases: adenine (A), guanine (G), cytosine (C) and thymine (T) with backbones made of sugars and phosphate groups connected through ester bonds. An intumescent material consists of three components: (i) an acid source (i.e. ammonium phosphates or polyphosphates, which release phosphoric acid), (ii) a carbon source (pentaerythritol, arabitol, sorbitol, inositol, cyclodextrins, saccharides, polysaccharides, etc.) and (iii) a blowing agent (guanidine, melamine, etc.) which, upon heating, releases great amounts of expandable or non-combustible gases such as water vapor, ammonia or carbon dioxide. By direct comparison, DNA contains all the three components required of an intumescent material in a single molecule: the phosphate groups able to produce phosphoric acid, the deoxyribose units acting as a carbon source and as blowing agents and the nitrogen-containing bases (A, G, C and T) that may release ammonia. In addition, the availability of DNA has become competitive with those of other chemicals, thanks to the large-scale method developed by Wang et al.[322] who proposed the extraction and purification of DNA from salmon milt and roe sacs. Owing to its intumescent and char forming ability, coupled with easy availability and non-toxic nature, DNA is a favorable candidate for flame retardant coating fabrication.

Laponite (LAP) is a synthetic clay that are disk shaped and has an empirical formula + −0.7 of Na 0.7[(Si8Mg5.5Li0.3)O20(OH)4] .[323,324] The single nanoplatelet of laponite has an

86 approximate diameter of 25 nm and a thickness of approximately 1 nm. It undergoes face- to-face packing of each individual platelet in its aggregated state, but swells or exfoliates in water dispersion affording negative charge on its faces. At the edge of individual platelet, the oxygen atoms and hydroxyl groups are able to accept or release protons depending on the environment property. Owing to its nano-dimensions and ability to exfoliate, LAP has been popularly used in the development of high nanofiller nanocomposites. [125,148,179,198,320]

In this paper, we utilize the single-dip coating approach to create a thin layer of LAP/DNA coating on surface of PU foam. The main focus will be directed to the overall preparation of coating suspensions, the coating of PU foam samples through single-dip coating method and study of properties of the resulting system. Scanning electron microscopy (SEM) and Energy-dispersive X-ray spectroscopy (EDX) was used to study the morphology of the nanocomposite films. Thermogravimetric analysis (TGA) in both nitrogen and air conditions were used to study the thermal stability and char yield of the coated PU foams. Compression testing is employed to analyze the physical properties and finally, flammability of the foams was also studied. This study might provide constructive insights to flame retardant fabrication of PU foams using a non-toxic and environment- friendly approach.

2.2 Experimental Section

2.2.1 Materials

Deoxyribonucleic acid (DNA) (crude oligonucleotides, <50bp) from herring sperm was purchased from Sigma Aldrich. Laponite RD (LAP) was provided by BYK USA INC. Untreated and characterized polyurethane (PU) foam was supplied by Underwriters Laboratory (UL). 1wt% DNA solution was prepared by magnetic stirring and sonication at 30% amplitude for 5hrs. 1wt% laponite suspensions were prepared by magnetic stirring and sonication at 25% amplitude for 1 hr. Sonication was performed with Brandson Ultrasound. Polyurethane foam was cut into 4cm cubes to prepare samples for single-dip coating.

2.2.2 Preparation of LAP/DNA suspensions

Laponite suspensions were prepared by stirring 1 wt% laponite in water with sonication at 25% amplitude for 1 hour. DNA was dissolved in water and stirred with sonication at 30% amplitude for 5 hours to prepare 1wt% DNA solutions. Appropriate amounts of 1wt% laponite suspension and 1wt% DNA solution was stirred and sonicated at 25% amplitude for 1 hour to obtain homogenously mixed LAP/DNA suspensions with different laponite content of 60wt%, 70wt%, 80wt%, 90wt% and 95wt%.

2.2.3 Preparation of coated PU foams

PU foams were cut into the aforementioned dimensions and immersed into LAP/DNA suspension. Once the foams were soaked, the excess suspension was squeezed out using a constant load (25lb weights were used). The foams were then dried in the oven at 50℃ for 24 hours and then kept in the desiccator for 48 hours. LAP/DNA coated PU foams with varying laponite content of 60wt%, 70wt%, 80wt%, 90wt% and 95wt% was prepared. The coated foam was designated as PU-LAP-60, which is a PU foam with LAP/DNA coating at 60wt% laponite content. DNA coated PU foams (PU-DNA) and laponite coated PU foams (PU-LAP) were also prepared using same procedure. Control foams (uncoated PU foams) were left in the desiccator for 48 hours.

2.2.4 Preparation of LAP/DNA nanocomposite films

The prepared LAP/DNA suspensions with different laponite content of 60wt%, 70wt%, 80wt%, 90wt% and 95wt% were each cast on a petri dish and dried in oven at 50℃ for 24 hours. Semi-transparent and brittle films were formed.

2.2.5 Characterization

FT-IR analysis was conducted using an ABB MB3000 FT-IR Laboratory Analyzer. 32 scans were co-added per spectrum at a resolution of 4 cm-1 after purging the spectrometer with dry air. KBr powder and solid samples were ground and pressed to form a 13 mm pellet to collect IR spectra in the absorbance mode Thermogravimetric analysis (TGA) was conducted on a TA instruments 2950 TGA with a heating rate of 10℃/min under nitrogen and air at a flow rate of 60 mL/min. Scanning electron microscopy (SEM) and Energy-dispersive X-ray spectroscopy (EDX) was conducted with a FEI Helios 650 Field Emission Scanning Electron Microscope with Focused Ion Beam with XEDS. SEM

88 samples were sputter-coated with a thin layer of conducting material before testing. Compression testing was performed by an MTS compression machine (Model 5565) at a strain rate of 10mm/min using 1 kN load cell to measure the applied force.

2.3 Results and Discussion

2.3.1 FT-IR spectroscopy

In order to characterize the presence of LAP/DNA nanocomposite coating on the surface of PU foam, model films without the PU substrate were casted by employing the water evaporation procedure which is identical to the flame-retardant coating fabrication. Fig 2.1 shows the FT-IR spectra comparing the spectra of Untreated DNA, Untreated LAP, subtracted PU-LAP-90 (PU spectra subtracted) and LAP-90 film. The subtracted spectra are compared with the laponite and DNA spectra in order to clearly see the characteristic peaks of laponite and DNA. The characteristic laponite bands are observed at 1009 cm-1, 704 cm-1 and 654 cm-1 which correspond to Si-O stretching, Si-O bending and OH bending, respectively.[325] The characteristic DNA bands are observed at 3374 cm-1, 1687 cm-1 and 1063 cm-1 which corresponds to the stretching vibration of hydroxyl groups, the asymmetric stretching vibrations of carboxyl groups and the symmetric stretching vibration of the phosphate groups, respectively.[326,327] Changes are apparent in the shifts of the characteristic absorption bands for both laponite and DNA in the spectra for PU-LAP-90. Characteristic peaks of laponite are observed to shift to lower wavenumbers 987 cm-1 and 648 cm-1 for PU-LAP-90 foam. Characteristic peaks of DNA were more difficult to observe owing to its low content, but the peak corresponding asymmetric stretching vibrations of carboxyl groups is observed to shift slightly higher to 1694 cm-1. In summary, FT-IR spectra confirm the presence of both laponite and DNA in PU-LAP-90 foams. This was also observed to be true for other coated PU foams with different laponite-DNA content.

4000 3500 3000 2500 2000 1500 1000 500

1009 cm-1

654 cm-1 Laponite 704 cm-1

1687 cm-1

3374 cm-1

DNA

987 cm-1

-1

1694 cm-1 648 cm

(No PU) (No

Subtracted PU-LAP-90 Subtracted

Control

PU-LAP-90

4000 3500 3000 2500 2000 1500 1000 500 Wavenumber (cm-1)

Figure 2.1: FT-IR spectrum of (Top to Bottom) Laponite, DNA, Subtracted PU-LAP-90 (No PU), Control and PU-LAP-90.

Figure 2.2: SEM images of fractured film surfaces: a) laponite, b) LAP-60, and c) LAP- 90 films.

Figure 2.3: EDX image mapping the distribution of Phosphorus across the LAP-90 film surface.

2.3.2 SEM and EDX analysis

SEM imaging was employed to characterize the cross-sectional morphology of fractured LAP, LAP-60 and LAP-90 films. As is shown in Fig 2.2 the cross section of

91 laponite film is extremely textured with oriented sheet-like structure. LAP-60 film with 40wt% DNA incorporated shows a poorly oriented sheet-like structure and seems to have a reduced surface roughness. LAP-90 film with 10wt% DNA, on the other hand, seem to have highly oriented nacre-like structure. For most papers reporting nanocomposites, clay nanofillers are dispersed in random orientation whether they are intercalated or exfoliated. Such random orientation does not efficiently block the diffusion of fragmented molecules. In this paper, we observe that the nanoplatelets neatly stack in x-y plane with little evidence for z-direction orientation. This will aid in efficiently blocking the diffusion of fragmented molecules. EDX analysis was utilized to examine the presence of DNA in the LAP-90 film. As can be seen in Fig 2.3, DNA (marked by the presence of Phosphorus) is homogenously dispersed across the film. The layered sheet-like structure in addition to the homogenous dispersion of DNA across the film, suggests the successful development of a nacre-mimetic high nanofiller content nanocomposite. Although further detailed study is needed, there is some evidence from our previous study that the coin-like laponite stack is intercalated. These intercalated laponite stack then align to form layered structure. Therefore, it is proposed that a hierarchical structure of orientation exists where intercalated coin-like laponite fillers with individual silicate layers of approximately 1 nm, also forms a secondary layered structure of the thickness slightly more than 10 nm. In the end, it is safe to conclude that well oriented, compact, homogenous structure of the LAP-90 will effectively block the diffusion of the molecular fragments because of the effective tortuous path effect, thereby reducing the flammability the material.

2.3.3 Thermal Stability

TGA was used to characterize the thermal degradation of PU foam and coated PU foams under nitrogen and air atmospheres. The nitrogen atmosphere pyrolysis allows to simulate the degradation mechanism in condensed phase which is also under an oxygen- deficient condition. As shown by the TGA curve showed in Fig 2.4, the pure PU foam undergoes two main thermal degradation steps. The first degradation step occurs between 190℃ and 390℃ and is believed to be the bond breakage of urethane, thereby generating polyols and isocyanates. Above 390℃, volatilization of regenerated polyols comprises the second degradation step.[328] At around 700℃. The onset temperature was not

92 significantly altered for coated samples while the two-step degradation path was still preserved. This suggest that LAP DNA coating does not chemically change the path way. A trend of delaying weight loss over the two-step degradation process was however observed. This would suggest that the LAP/DNA coating helped improve the char forming process while simultaneously mildly delaying the two-step degradation process. This could be attributed to the alteration of the rheological and physical properties of the liquefied material brought on by the high filler-content nanocomposite coating.

120 ––––––– Control ––––––– PU-DNA ––––––– PU-LAP-60 ––––––– PU-LAP-70 100 ––––––– PU-LAP-80 ––––––– PU-LAP-90 ––––––– PU-LAP-95 ––––––– PU-LAP

Weight (%) Weight 40

-20 0 200 400 600 800 Temperature (°C) Universal V4.5A TA Instruments

Figure 2.4: TGA thermograms of control and coated PU foams in nitrogen atmosphere.

The char yield within the residue of the coated foams were also evaluated as observed in Table 2.1. TGA analysis of LAP/DNA films of different content ratios (not shown) were conducted, and percentage of residue formed is used against the percentage of residue formed for the coated foams to calculate the char yield. This data is only suggestive of the char formation potential of LAP/DNA coated PU foams. As shown in

Fig 2.5, char yield is observed to significantly improve with coated PU foams. PU-LAP- 90 with 90wt% laponite in the coating mixture is observed to have the highest char yield of 6.2%. This might be attributed to the highly oriented, compact, nacre-like structure coating formed on the surface of the coated foams, which helps in the formation of char. PU-LAP-70 and PU-LAP-95 were also found to have appreciable improvements in their char yield with values of 3.94% and 3.75%, respectively.

Table 2.1: Residue percentage and char yield percentage data of coated PU foams.

Wt% of Char yield Material Residue% coating Total LAP/DNA residue wt% % PU-LAP-60 5.98 5.46 4.60 1.38 PU-LAP-70 8.29 5.18 4.35 3.94 PU-LAP-80 6.67 5.79 4.98 1.69 PU-LAP-90 10.77 5.26 4.57 6.20 PU-LAP-95 8.98 5.84 5.23 3.75

Char Yield

PU-LAP-60

PU-LAP-70

PU-LAP-80

PU-LAP-90

PU-LAP-95

0.00 2.00 4.00 6.00 8.00 10.00 12.00 14.00 16.00 18.00

Residue% Char yield %

Figure 2.5: Residue percentage and char yield percentage of coated PU foams.

TGA in air atmosphere was also done in order to understand the thermal stability of the foams in normal environmental conditions (Fig 2.6). Uncoated PU foams were observed to degrade in a three-step degradation process. However, the uncoated samples were observed to leave almost no residue in the end. Upon coating with pure DNA, a delay

94 of weight loss was observed in the three-step degradation process. The same was observed with PU-LAP-60 and PU-LAP-70 where the three-step degradation process showed a delay in weight loss. Char yield was also observed to increase with PU-LAP-70 having an increase to around 4%. Above 70wt% of laponite, the trend of delay in weight loss was not observed. However, the char yield significantly increased to around 8% for PU-LAP-90 with PU foams coated with purely laponite coating had a char yield of 11%. Structural integrity of the coated samples with higher content of nanofillers was also observed to be preserved, which was absent in coated samples with low content of nanofillers. Overall, the TGA profiles seem to suggest that LAP/DNA coating made the thermal degradation of the PU foams less severe and laponite concentration seemed to play an important role in this process.

Weight (%) Weight 40

-20 0 200 400 600 800 Temperature (°C) Universal V4.5A TA Instruments

Figure 2.6: TGA thermograms of control and coated PU foams in air atmosphere.

2.3.4 Compression Testing

The modulus and compression strength are understood to be dependent on the

95 thickness of the coated PU foams, owing to coating uptake variance for PU foams of different thicknesses. However, cubic samples of edge size 4cm were prepared as this was representative of the bulk behavior of PU foams in compression. For each coating of different laponite-DNA content, five samples were prepared in order to test reproducibility. As shown in Table 2.2, compression strength and modulus were calculated in order to evaluate the mechanical behavior of the PU foams upon coating with different films of different laponite-DNA content. Compression strength (Fig 2.7) was observed to be the highest for PU-LAP-90 with a value of 9.68 kPa. The improvement in compression strength can be attributed to the homogenous coating of the foams’ surface with high content of nanofillers. However, PU foams coated with only laponite (without DNA) bore foams with lesser compression strength than PU-LAP-90, while those foams coated with DNA suspension (without laponite) had mechanical behavior closer to that of control foams. Therefore, not only is the compression strength significantly improved by the presence of high content of nanofillers, it also seems that synergistic interactions between laponite and DNA also helps in achieving improved strength of the foam. Unfortunately, incorporation of high content of clay nanofillers had a negative impact on modulus (Fig 2.8). The significant rise in stiffness causes increase in deformation recovery time, i.e., the laponite incorporated PU foams take longer time to recover their original shape after compression.

Table 2.2: Compression strength and modulus data of control and coated PU foams.

Material Compression Strength Modulus (MPa) (kPa) Control 2.94 0.035 PU-DNA 3.45 0.042 PU-LAP-60 7.54 0.131 PU-LAP-70 7.78 0.142 PU-LAP-80 8.39 0.182 PU-LAP-90 9.68 0.199 PU-LAP-95 9.30 0.255 PU-LAP 8.37 0.192

Compression Strength 12.00 10.00 8.00 6.00 4.00 2.00

0.00

Compression Strength (kPa) Strength Compression

PU-LAP

Control

PU-DNA

PU-LAP-70 PU-LAP-80 PU-LAP-90 PU-LAP-95 PU-LAP-60

Figure 2.7: Compression strength of control and coated PU foams.

Modulus 0.300 0.250 0.200 0.150 0.100

Modulus (MPa) Modulus 0.050

0.000

PU-LAP

Control

PU-DNA

PU-LAP-70 PU-LAP-80 PU-LAP-90 PU-LAP-95 PU-LAP-60

Figure 2.8: Modulus of control and coated PU foams.

2.3.5 Flammability

The flammability behavior is also understood to be dependent on the thickness of the coated PU foams, owing to coating uptake variance for PU foams of different thicknesses. Coated cubic PU foams with edge dimensions of 4cm were prepared to qualitatively observe the flame retardant effects of coatings with different laponite-DNA content. A bunsen burner was used to ignite the samples.

The control foam with no flame retardant coating was ignited and expectantly had significant dripping (Fig 2.9 a, e). The foam melted away as it burned leaving no solid residue in the end. Incorporation of DNA without laponite onto surface of PU foams seemed to reduce dripping which can be attributed to the intumescent nature of DNA (Fig 2.9 b, f). However, negligible improvement was observed in char yield and structural integrity. Upon incorporation of laponite with DNA in the coating mixture, char yield seemed to be improve with higher content of laponite. A laponite content of 60wt% was observed to have increased char yield formation for foam(Fig 2.9 c, g). Increasing laponite content to 90wt% led to a much higher char yield with no dripping being observed (Fig 2.9 d, h). Higher content of inorganic nanofillers, therefore, seemed to significantly promote char formation, thereby eliminating the melting of PU foam and delaying the fire-spreading rate. The structural integrity was also observed to significantly improve for the PU-LAP- 90 foams. The improvement in flammability properties can be attributed to the highly tortuous path created by the laponite-dna nacre like structure which block diffusion of molecules, thereby reducing the flammability of the material. Therefore, the single-dip coating approach using intumescent DNA coupled with high content of inorganic nanofillers was found to appreciably improve the flammability behavior of PU foams.

Figure 2.9: Figure showing the pre and post burning images of (a, e) control PU, (b, f) PU-DNA, (c, g) PU-LAP-60, and (d, h) PU-LAP-90.

2.4 Conclusion

1wt% LAP/DNA nanocomposite suspensions with high content of inorganic nanofillers was successfully applied on PU foam surface using the single-dip coating approach. SEM analysis showed highly oriented, nacre-like cross sectional morphology for LAP-90, while EDX mapping of phosphorus showed homogenous dispersion of DNA in the nanocomposite film. TGA study of PU-LAP-90 foam was observed to give the highest char yield of 6.20%. Flammability tests showed high char-formation and preservation of structural integrity post burning of the samples. Compression testing exhibited significant rise in modulus while compression strength of PU-LAP-90 foam was enhanced by 229%. Therefore, application of LAP/DNA nanocomposite film on PU foams through single-dip coating method is a potentially suitable approach owing to its merits of green raw materials, abundant materials resource, simple and scalable fabrication process, high char formation, and enhanced mechanical behavior.

2.5 Acknowledgment

The authors are indebted to the partial financial support of Underwriters Laboratories (UL).

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